VSEBINA – CONTENTS
PREGLEDNI ^LANEK – REVIEW ARTICLE
Wear mechanisms and surface engineering of forming tools
Obrabni mehanizmi in in`eniring povr{ine preoblikovalnih orodij
B. Podgornik, V. Leskov{ek. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 313
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Predicting the physical properties of drawn Nylon-6 fibers using an artificial-neural-network model
Napovedovanje fizikalnih lastnosti vle~enih vlaken iz najlona 6 z uporabo modela umetne nevronske mre`e
R. Semnani Rahbar, M. Vadood. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 325
Influence of the impact angle and pressure on the spray cooling of vertically moving hot steel surfaces
Vpliv vpadnega kota in tlaka na ohlajanje z brizganjem na vertikalno premikajo~e se vro~e povr{ine jekla
M. Hnízdil, M. Chabi~ovský, M. Raudenský . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 333
Techniques of measuring spray-cooling homogeneity
Tehnike merjenja homogenosti hlajenja z brizganjem
M. Chabi~ovský, M. Raudenský . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 337
Mineralogical and geochemical characterization of roman slag from the archaeological site near Mo{nje (Slovenia)
Mineralo{ka in geokemi~na karakterizacija rimske `lindre z arheolo{kega najdi{~a pri Mo{njah (Slovenija)
S. Kramar, J. Lux, H. Pristacz, B. Mirti~, N. Rogan - [muc . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 343
Reynolds differential equation singularity using processes of small straining with lubrication
Reynoldsova diferencialna ena~ba pri procesih majhne deformacije z mazanjem
D. ]ur~ija, F. Vodopivec, I. Mamuzi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 349
Prediction of the catastrophic tool failure in hard turning through acoustic emission
Napovedovanje katastrofi~ne po{kodbe kerami~nih vlo`kov pri stru`enju z akusti~no emisijo
M. ^illiková, B. Mi~ieta, M. Neslu{an, R. ^ep, I. Mrkvica, J. Petrù, T. Zlámal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 355
Erosive wear resistance of silicon carbide-cordierite ceramics: influence of the cordierite content
Odpornost keramike silicijev karbid-kordierit proti obrabi pri eroziji: vpliv vsebnosti kordierita
M. Po{arac-Markovi}, Dj. Veljovi}, A. Deve~erski, B. Matovi}, T. Volkov-Husovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 365
Influence of the substrate temperature on the structural, optical and thermoelectric properties of sprayed V2O5 thin films
Vpliv temperature podlage na strukturne, opti~ne in termoelektri~ne lastnosti napr{ene tanke plasti V2O5
Y. Vijayakumar, K. N. Reddy, A. V. Moholkar, M. V. R. Reddy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 371
Deep micro-hole drilling for hadfield steel by electro-discharge machining (EDM)
Vrtanje globokih mikrolukenj v jekla hadfield z metodo elektrorazreza (EDM)
V. Yilmaz, M. Sarýkaya, H. Dilipak . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 377
Surface analysis of electrochromic CuxO films in their colored and bleached states
Povr{inska analiza elektrokromiznih plasti CuxO v njihovih obarvanih in obeljenih stanjih
M. M. Ristova, M. Milun, B. Pejova . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 387
Thermodynamic analysis of the precipitation of carbonitrides in microalloyed steels
Termodinamska analiza izlo~anja karbonitridov v mikrolegiranih jeklih
M. Opiela . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 395
Experimental investigation of the crack-initiation moment of Charpy specimens under impact loading
Eksperimentalna preiskava trenutka iniciacije razpoke pri udarni obremenitvi Charpyjevih vzorcev
V. Kharchenko, E. Kondryakov, A. Panasenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 403
Structural, thermal and magnetic properties of Fe-Co-Ni-B-Si-Nb bulk amorphous alloy
Strukturne, termi~ne in magnetne lastnosti masivne amorfne zlitine Fe-Co-Ni-B-Si-Nb
S. Lesz, M. Nabia³ek, R. Nowosielski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 409
The nano-wetting aspect at the liquid-metal/SiC interface
Vidik nanoomakanja na stiku staljena kovina-SiC
M. Mihailovi}, K. Rai}, A. Patari}, T. Volkov - Husovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 413
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 49(3)311–477(2015)
MATER.
TEHNOL.
LETNIK
VOLUME 49
[TEV.
NO. 3
STR.
P. 311–477
LJUBLJANA
SLOVENIJA
MAY–JUNE
2015
The effect of a superplasticizer admixture on the mechanical fracture parameters of concrete
Vpliv dodatka superplastifikatorja na parametre mehanskega zloma betona
H. [imonová, I. Havlíková, P. Danìk, Z. Ker{ner, T. Vymazal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 417
Properties and structure of Cu-Ti-Zr-Ni amorphous powders prepared by mechanical alloying
Lastnosti in struktura amorfnih prahov Cu-Ti-Zr-Ni, pripravljenih z mehanskim legiranjem
A. Guwer, R. Nowosielski, A. Lebuda . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 423
Interaction of Cr2N and Cr2N/Ag thin films with CuZn-brass counterpart during ball-on-disc testing
Interakcija Cr2N in Cr2N/Ag tankih plasti v paru s CuZn-medenino med preizkusom krogla na disk
P. Bílek, P. Jur~i, P. Dulová, M. Hudáková, J. Pta~inová, M. Pa{ák. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 429
Using simulated spectra to test the efficiency of spectral processing software in reducing the noise in Auger electron spectra
Uporaba simuliranega spektra za preizkus u~inkovitosti programske opreme predelave spektra pri zmanj{anju {uma spektra Augerjevih
elektronov
B. Poniku, I. Beli~, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 435
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Surface behavior of AISI 4140 modified with the pulsed-plasma technique
Lastnosti povr{ine AISI 4140, spremenjene s tehniko pulzirajo~e plazme
Y. Y. Özbek, M. Durman . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 441
Preparation and dielectric properties of thermo-vaporous BaTiO3 ceramics
Priprava in dielektri~ne lastnosti termo-parno porozne keramike BaTiO3
A. Kholodkova, M. Danchevskaya, N. Popova, L. Pavlyukova, A. Fionov . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 447
Hybrid sol-gel coatings doped with cerium to protect magnesium alloys from corrosion
Hibridni sol-gel-nanosi, dopirani s cerijem, za korozijsko za{~ito magnezijevih zlitin
N. V. Murillo-Gutiérrez, F. Ansart, J.-P. Bonino, M.-J. Menu, M. Gressier . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 453
Influence of the tool geometry and process parameters on the static strength and hardness of friction-stir spot-welded
aluminium-alloy sheets
Vpliv geometrije orodja in parametrov procesa na stati~no trdnost in trdoto pri vrtilno-tornem to~kastem varjenju plo~evin iz Al-zlitine
H. Güler . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 457
The stabilization of nano silver on polyester filament for a machine-made carpet
Stabilizacija nanodelcev srebra na poliestrskem vlaknu za strojno izdelavo preprog
K. M. Shojaei, A. Farrahi, H. Farrahi, A. Farrahi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 461
Evaluation of the thermal resistance of selected bentonite binders
Ocena toplotne upornosti izbranih bentonitnih veziv
J. Beòo, J. Vontorová, V. Matìjka, K. Gál . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 465
Development of numerical models for the heat-treatment-process optimisation in a closed-die forging production
Razvoj numeri~nih modelov za optimizacijo postopka toplotne obdelave pri proizvodnji odkovkov v zaprtih utopnih orodjih
L. Male~ek, M. Fedorko, F. Van~ura, H. Jirková, B. Ma{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 471
IN MEMORIAM
Alojz Pre{ern, dipl. in`. metalurgije (1920–2015) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 477
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
WEAR MECHANISMS AND SURFACE ENGINEERING OF
FORMING TOOLS
OBRABNI MEHANIZMI IN IN@ENIRING POVR[INE
PREOBLIKOVALNIH ORODIJ
Bojan Podgornik, Vojteh Leskov{ek
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
bojan.podgornik@imt.si
Prejem rokopisa – received: 2015-01-05; sprejem za objavo – accepted for publication: 2015-03-02
doi:10.17222/mit.2015.005
The manufacturing of parts is faced with ever-increasing demands for higher strength (hardness) and toughness of the work
material, as well as higher productivity and environmental concerns. At the same time, the quality requirements are high and
will continue to grow in the future. This leads to restrictions in terms of the type and use of lubricants and especially to
increased requirements relating to the wear and fatigue resistance of tools.
The main focus on improving the wear resistance and tribological properties of forming tools has mainly been on developing
tool steels with an improved fracture toughness and then modifying the lubricants for better retention and permeability at the
tool/work-piece contact area. Nevertheless, the wear resistance of forming tools can also be successfully improved by
surface-engineering techniques. In recent years, hard coatings in particular have shown enormous potential for improving the
tribological properties of contact surfaces.
This paper reviews the wear mechanisms encountered in forming processes, as well as various surface-engineering techniques
designed to improve the wear resistance and the anti-sticking properties of forming tools. The possible benefits and restrictions
of different surface-engineering techniques are presented for the example of sheet-metal forming, fine blanking and forging.
Keywords: surface engineering, forming, hard coatings, topography, friction, wear
Preoblikovalna industrija se spopada z nenehno nara{~ajo~imi zahtevami po vi{ji trdnosti (trdoti) in `ilavosti preoblikovanih
materialov, kakor tudi ve~ji produktivnosti in skrbi za okolje. Isto~asno se pove~ujejo zahteve po kvaliteti izdelkov, kjer
predvsem kvaliteta povr{ine postaja vedno bolj pomembna. To pomeni omejitve pri uporabi maziv ter seveda ve~je zahteve po
obrabni odpornosti in vzdr`ljivosti preoblikovalnih orodij.
Do nedavnega je izbolj{evanje obrabne obstojnosti in tribolo{kih lastnosti preoblikovalnih orodij temeljilo predvsem na razvoju
bolj `ilavih in ~istej{ih orodnih jekel ter modifikaciji maziv z bolj{imi mazalnimi lastnostmi. V zadnjih letih pa tudi na podro~ju
preoblikovalnih orodij tehnike in`eniringa povr{in, {e posebej nanos trdih za{~itnih prevlek, omogo~ajo nadaljnje izbolj{anje
tribolo{kih lastnosti kontaktnih povr{in.
V prispevku so tako predstavljeni obrabni mehanizmi, ki so jim izpostavljena preoblikovalna orodja, kakor tudi razli~ne tehnike
in`eniringa povr{ine, namenjene izbolj{anju obrabne obstojnosti in tornih lastnosti preoblikovalnih orodij. Problemi in prednosti
uporabe in`eniringa povr{ine so poudarjene na primeru hladnega preoblikovanja plo~evine, {tancanja in kovanja.
Klju~ne besede: in`eniring povr{in, preoblikovanje, trde prevleke, topografija, trenje, obraba
1 INTRODUCTION
The forming industry is confronted with ever-increas-
ing demands to form low-weight, high- and ultra-high-
strength materials, as well as with higher productivity
and environmental concerns.1 Simultaneously, we are
increasing the requirements on the quality of the formed
parts, especially in terms of surface quality.2–4 All these
lead to restrictions regarding the use of the type and
quality of the lubricants and to increased demands on the
wear and fatigue resistance of the tool.
The wear resistance and the efficiency of forming
tools are limited for different reasons, including thermal
and mechanical fatigue and cracking, erosion, corrosion,
abrasive and adhesive wear, and galling.5,6 In addition to
that, tool replacement and reconditioning often depends
on the surface quality of the formed parts.3,7,8 A smooth,
defect-free surface represents a competitive market
advantage, but mostly it is more favourable contact con-
ditions and a better wear resistance. The main obstacles
to obtaining smooth surfaces are the increased surface
roughness and the wear of the tool, especially the
adhesive wear and the galling.9,10 On the other hand, tool
wear and galling also lead to increased contact pressures
and unstable friction in the forming process.
Traditionally, improving the tribological properties
and the wear resistance of forming tools was mainly
based on the increased cleanliness, hardness and fracture
toughness of the tool steel, as well as on the better lubri-
city and stability of the forming lubricants. However,
wear and galling resistance can also be very successfully
reduced with a proper modification or surface engineer-
ing of the tool surface.2,11 Commonly applied surface-
engineering processes, found in many tool applications,
are the thermo-chemical treatments, ranging from basic
surface hardening and plasma nitriding to more recent
processes of deep cryogenic treatment and laser-surface
remelting.12–16 However, in the past two decades, surface
coating processes are gaining an advantage, including
PVD, CVD and PACVD techniques. Thin, hard coatings
Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324 313
UDK 531.43:539.92:621.793/.795 ISSN 1580-2949
Review article/Pregledni ~lanek MTAEC9, 49(3)313(2015)
with excellent tribological properties, like TiN, TiAlN,
CrN, etc., have already outperformed and successfully
replaced traditional high-speed steel tools in the majority
of cutting operations.17,18 Despite this, the majority of
forming tools are still uncoated. Besides the complex
geometry, which makes it difficult to obtain uniform
coating deposition, commercial hard wear-resistant coat-
ings show a relatively high coefficient of friction and a
high tendency for galling against typical formed mate-
rials.19–21 Typical tool steels also have a lower load-carry-
ing capacity than high-speed steel, ceramics and hard
metals. As such, tool steels do not provide sufficient
support for very thin, hard and brittle coatings. However,
in recent years a lot of new deposition methods and
coating types have been developed, which show excel-
lent adhesion, as well as mechanical and tribological
properties.22 In order to exploit the full potential of hard
coatings in terms of improving the tribological properties
of forming tools, coated surfaces must primarily be able
to sustain high-impact dynamic loading without crack-
ing, debonding or spallation. The improved load-carry-
ing capacity of coated systems can be obtained by
increasing the thickness of the coating, which is not easy
to achieve due to the high residual stresses, by employ-
ing support layers, which increases costs, or through the
most efficient duplex technique, combining a classic
thermo-chemical treatment of the steel substrate and
PVD or PACVD deposition of the protective coat
ing.19,23,24
Another option for improving the tribological proper-
ties, mainly the friction behaviour of the forming tools,
either coated or un-coated, is through a surface-rough-
ness and topography optimization. Proper selection and
preparation of the surface topography, using techniques
like polishing, shot-peening and laser surface texturing,
can greatly improve the performance and galling
resistance of the forming tools, and thus reduce or even
eliminate the need for lubrication.7,25,26
2 WEAR MECHANISMS IN FORMING
APPLICATIONS
Different materials can be formed in a desired shape
or product using different processes, like casting, metal
forming and machining. Metal forming is further divided
into rolling, drawing, extrusion, sheet-metal forming, die
casting and forging.27 During forming the tool surface is
subjected to a sliding contact with the formed material,
to high contact stresses and often to elevated tempera-
tures, which all lead to tool wear. The wear mechanisms
that can be found in forming applications include abra-
sive and adhesive wear, mechanical and thermal fatigue,
plastic deformation and corrosion.27,28 Certainly, different
wear mechanisms require different properties of the tool
material and especially of the tool surface. However, in
general, the surface of the tool should be hard and should
maintain this hardness at high temperatures to reduce the
abrasive wear. Furthermore, the tool material should be
tough in order to limit or reduce the fatigue and be heat
resistant with a high thermal resistivity. Typical wear
mechanisms found in different forming operations and
being responsible for tool wear are summarized in Table
1.
2.1 Cold forming
In the case of cold forming the tool failure is caused
by five main wear mechanisms, which are a consequence
of the high contact pressure and the relative motion bet-
ween the tool surface and the formed material (Figure
1).29 The main wear and damage mechanisms in cold
forming are:
• abrasive wear
• adhesive wear
• low-cycle fatigue
• crack propagation
• plastic deformation
The most dominant wear mechanism, found in all
cold-forming operations, is adhesive wear. However,
normally more than just one wear mechanism takes
place, often even all of them can be observed, which
depends on the forming process and the work material.
In the case of punching and fine blanking, characterized
by sharp edges and high impact loads, adhesion is
accompanied by fatigue and chipping, and for drawing
and extrusion, with abrasive wear. On the other hand, in
the cold forming of harder and thicker materials, abra-
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
314 Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324
Table 1: Wear mechanisms in forming processes
Tabela 1: Obrabni mehanizmi v preoblikovalnih procesih
abrasion adhesion low cyclefatigue
high cycle
fatigue
thermal*
fatigue corrosion
plastic
deformation
Punching, Fine blanking x x x x x
Deep drawing x x
Sintering x x x
Rolling x x x x x x x
Extrusion Drawing x x x x x x
Forging x x x x x x x
Injection moulding x x x x
Die casting x x x x
* valid for processes at elevated temperatures
sive wear and plastic deformation will dominate, while
for softer materials it is adhesive wear, known also as
galling.
Adhesive wear as the prevailing wear mechanism can
be found in the cold forming of softer metals, like
aluminium, titanium, copper, austenitic stainless steel,
etc., and is especially critical in sheet-metal forming.
The adhesive wear of forming tools is a result of local
micro welds between the surface of the tool and the
work-piece followed by the gradual transfer and accumu-
lation of the work material on the tool surface, known
also as galling.30 The transferred material causes unstable
friction as well as scratching and a poor, uneven surface
of the work-piece. Removal of the transferred work
material can even cause abrasive wear of the tool surface.
For good adhesive wear resistance the tool surface
should be hard but with sufficient ductility, as smooth as
possible and have low friction against the work material.
In practice, adhesion or galling problems are mostly
tackled by the application of different lubricants. How-
ever, the use of lubricants requires additional cleaning
procedures and presents environmental concerns.
Abrasive wear dominates when forming hard mate-
rials or multiphase materials with hard particles (oxides,
carbides, etc.) which scratch the tool surface and lead to
tool wear. However, as already mentioned, abrasive wear
can also be a result of adhesive wear. During forming the
transferred and adhered work material becomes work
hardened and when detached from the surface it repre-
sents a hard abrasive particle, which can cause three-
body abrasive wear of the tool surface. Abrasive wear
can be found in punching, fine blanking, drawing, extru-
sion and forging. The improved abrasive wear resistance
of forming tools is mainly achieved by increasing the
tool-surface hardness, either through thermo-chemical
treatments or the deposition of hard, wear-resistant
coatings.
Low-cycle fatigue is typical for punching, stamping
and fine blanking, where the cutting edges are subjected
to repeated high contact stresses. The repeated impact
loading, sliding motion and local plastic deformation
result in crack initiation and propagation, which finally
leads to chipping. Due to this chipping the cutting edges
become blunt and, consequently, this leads to increased
stresses, unfavourable sliding conditions between the
tool surface and the work-piece, and to adhesive wear.
The resistance to low-cycle fatigue can be improved by
increasing the ductility but without sacrificing the tool
hardness.
In contrast to other wear mechanisms found in cold
forming, crack propagation leads to instantaneous failure
or even to the fracture of the tool, which is very hard to
predict. Crack propagation is a consequence of stress
concentrations, cracks initiation and the tensile stress
field in the tool material. Since the majority of cracks
start and propagate below the surface, the most important
properties in terms of crack-propagation resistance are
the properties of the base or substrate tool material. It
has to possess high toughness, which on the other hand
also means reduced hardness. Therefore, a suitable com-
promise between the increased toughness and the
reduced tool hardness needs to be obtained.
Plastic deformation, caused by high contact pres-
sures, is a common problem in many cold-forming
operations. When the compressive stress exceeds the
yield strength of the tool material, the tool geometry
becomes distorted and this is often also coupled with
imprints on the tool surface. The tool-geometry
distortion means an incorrect shape of the formed part,
while the presence of imprints leads to a poor surface
quality. In terms of the resistance to plastic deformation,
the most critical tool property is the hardness of the tool
surface.
2.2 Hot forming
In hot-forming processes, including die casting,
injection moulding and hot forging, the wear of the tool
is caused by four types of loading:31
• thermal,
• mechanical,
• tribological,
• chemical.
Thermal loads are caused by repeated heat transfer
from the work-piece and the repeated cyclic heating and
cooling of the tool surface. High-temperature exposure
leads to a reduced tool hardness as well as cyclic heating
and cooling, leading to thermal fatigue. In the same way
as in cold forming, also for hot forming, especially in the
case of hot forging, mechanical loads may result in fati-
gue and plastic deformation. However, the tribological
conditions in the case of hot forming are much more
complex. Oxide layers formed on the work-piece surface
are hard and brittle and as such result in the generation
of hard, abrasive particles, which often cause abrasive
wear of the tool surface. On the other hand, contact
temperatures during hot forming often exceed 700 °C,
thus affecting the hardness and adhesive properties of the
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324 315
Figure 1: Wear mechanisms found in cold-forming processes29
Slika 1: Obrabni mehanizmi, zna~ilni za procese hladnega preobli-
kovanja29
contact surfaces. The higher is the temperature, the
larger will be the drop in the surface hardness and more
reactive is the surface, resulting in more intense abrasive
wear and galling. The situation becomes even worse due
to the constant flow of hot-work material into the contact
with the tool surface. On the other hand, high tempera-
tures also result in chemical loads on the tool, which
may cause different chemical reactions on the tool sur-
face, i.e., oxidation.
Different parts of hot-forming tools are subjected to
different and quite specific combinations of loads and
consequently display different wear mechanisms, as
shown in Figure 2.32 The surfaces exposed to the highest
temperature changes are subjected to thermal fatigue,
and the surfaces with the highest stress concentrations, to
mechanical fatigue. Furthermore, in processes with high
thermal loads (die casting and injection moulding) an
increased surface reactivity leads to adhesive wear and
galling. However, in highly loaded tools for hot forging
and extrusion, more than 70 % of all tool failures are
caused by abrasive wear of the tool surface.
3 SURFACE ENGINEERING FOR IMPROVED
TRIBOLOGICAL PROPERTIES OF FORMING
TOOLS
3.1 Cold sheet-metal forming
Failure and the replacement of tools for cold sheet-
metal forming is mainly caused by adhesive wear and
galling, which is followed by abrasive wear. An impro-
vement in the abrasive wear is directly related to an
increased hardness of the surface, achieved through
different thermo-chemical treatments and hard-coating
deposition techniques.12 Galling problems and unstable
friction, on the other hand, are mainly addressed by the
use of highly additivated and environmentally hazardous
lubricants, which should provide proper lubrication bet-
ween the tool and the work-piece surface. The improved
tribological properties in sheet-metal forming processes
can also be achieved through the proper surface engi-
neering of the tool.2,33 The first group are the thermo-
chemical treatments (nitriding, boriding, vanadizing), the
second is surface texturing and the third is the deposition
of low-friction coatings.
Investigations in the field of thermo-chemical
treatments have shown that increased hardness of the
surface and microstructure refinement obtained through
a deep-cryogenic treatment can effectively improve the
galling resistance of tool steel.16 On the other hand, the
best results are obtained using plasma nitriding.25,34 The
plasma nitriding of tool steel gives up to 40 % lower
friction against austenitic stainless steel and up to 50 %
better galling resistance, as compared to hardened tool
steel, as shown in Figure 3. However, the selection of
the proper plasma-nitriding parameters and post-treat-
ment conditions is crucial.25,35 The presence of a compact
’ (Fe4N) compound layer may enhance the galling resis-
tance against stainless steel, while a combination of
porous  (Fe2-3N) and ’ compound layer and the use of a
non-polished, nitrided surface have a detrimental effect
(Figure 3).
More important than the selection of the thermo-che-
mical treatment is the proper preparation of the tool
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
316 Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324
Figure 3: a) Coefficient of friction and b) critical loads for galling
initiation (Lc1) and transfer layer build up (Lc2) for different surface
treatments of tool steel35
Slika 3: a) Potek koeficienta trenja in b) kriti~na obremenitev za~etka
prenosa (Lc1) ter tvorjenja plasti nerjavnega jekla (Lc2) v odvisnosti od
kemo-termi~ne priprave povr{ine orodnega jekla35
Figure 2: Typical wear mechanisms in hot forming32
Slika 2: Obrabni mehanizmi, prisotni pri vro~em preoblikovanju32
surface.7,26 As shown in Figure 4, even with just simple
polishing of the tool surface, a low and stable coefficient
of friction can be achieved, and the galling resistance can
be greatly improved. Rough surfaces (Ra > 0.25 μm)
result in high and unstable friction, but above all in the
almost immediate transfer of work material to the tool
surface. By reducing the surface roughness below 0.05
μm, low and stable friction and excellent galling resis-
tance across a broad load range can be achieved, even
with reduced lubrication, the use of less additivated
lubricants or without any lubrication (Figure 4).
A microscopic analysis of the contact surfaces of the
tools for cold forming has revealed that galling and the
transfer of work material is initiated at scratches,
irregularities and asperities on the tool surface,36 as
shown in Figure 5a. During forming and load increases
the transferred material becomes accumulated around the
initial galling points, forming a thick, transferred layer
on the tool surface (Figure 5b). Therefore, every imper-
fection on the tool surface represents a potential initial
point for the beginning of the galling and the work-mate-
rial transfer. Polishing the surface, on the other hand,
removes and smoothens the scratches and imperfections
caused during tool manufacturing, thus greatly reducing
the risk of adhesive wear. In forming processes, where
lubricants and lubrication cannot be eliminated, more
favourable tribological properties and a reduced amount
of lubricants can be achieved through the surface
texturing.37–39 Surface structures in the form of channels
or dimples with the proper size and density (Figure 6)
may act as mini-reservoirs, which can effectively feed
the lubricant into the contact and prevent, or at least
reduce, galling.26
The third option for improving the tribological
properties of tools for cold sheet-metal forming is the
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324 317
Figure 5: a) Beginning of galling and b) accumulation of thick trans-
ferred layer of work material on the tool surface
Slika 5: a) Za~etek adhezije preoblikovanega materiala na povr{ino
orodja in b) akumulacija in tvorjenje plasti prenesenega materiala
Figure 4: Effect of surface roughness on: a) the coefficient of friction
and b) the critical loads for stainless-steel transfer7
Slika 4: Vpliv hrapavosti povr{ine na: a) potek koeficienta trenja in b)
kriti~no obremenitev prenosa nerjavnega jekla7
Figure 6: Surface-textured tool surface
Slika 6: Obli~ena kontaktna povr{ina
deposition of a hard protective coating. In the case of
cutting tools, hard ceramic coatings, i.e., TiN, TiAlN,
AlCrN, etc., successfully replace classic materials,
giving a greatly improved abrasive wear resistance and
productivity and, in some cases, even allowing for dry
cutting.40 However, the requirements for the successful
implementation of hard coatings on forming tools are
much tougher and more complex. Besides the more
complex geometry of forming tools and the limited
load-carrying capacity of the tool-steel substrate, typical
commercial hard coatings for cutting applications show a
high friction and a strong tendency to pick up work
material. Therefore, coating selection in the forming
operations mainly depends on the material to be formed
and the coating galling tendency against it.2,32,41
In the case of stainless steel a TiN coating gives a
similar friction to un-coated cold-work tool steel, but
almost instantaneous adhesion and galling. As shown in
Figure 7, similar properties with a high galling tendency
against stainless steel are also displayed by other hard
ceramic coatings. On the other hand, carbon-based coat-
ings, either amorphous diamond-like-carbon coatings
(DLC) or metal-doped carbon-based coatings (W-C,
Ta-C, etc.) provide low and stable friction, even under
dry-sliding conditions.42 At the same time, DLC coatings
considerably increase the critical loads for galling
initiation and transfer-layer build up when it comes to
the forming of stainless steel (Figure 7b).
Although polishing of the contact surface greatly
improves the tribological properties of tool steel, highly
additivated lubricants are still required to prevent galling
and transfer-layer build up when forming stainless steel.
Switching to pure base oil does not allow more than 10
forming cycles, which are more or less limited to low
loads and low deformation rates before adhesion and
galling take place (Figure 8a). However, with the
application of DLC coatings we can effectively prevent
adhesive wear and provide a stable tribological contact
with a low coefficient of friction ( 0.1) with a more
environmentally friendly base lubricant without any
additives,7 as shown in Figure 8b.
When forming aluminium and aluminium alloys,
galling and adhesion to hardened tool steel take place
immediately and at very low loads if the contact is not
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
318 Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324
Figure 8: Friction maps for: a) polished tool steel and b) DLC-coated
surface-lubricated with pure PAO8 base oil7
Slika 8: Karte koeficienta trenja za: a) polirano orodno jeklo in b)
povr{ino, prekrito z DLC-prevleko; bazno olje PAO87
Figure 7: a) Coefficient of friction and b) critical loads for galling
initiation and transfer-layer build up for different hard coatings tested
against stainless steel35
Slika 7: a) Potek koeficienta trenja in b) kriti~ne obremenitve prenosa
nerjavnega jekla pri suhem drsnem kontaktu razli~nih prevlek35
lubricated. Even through surface engineering (polishing,
nitriding, hard coating) galling resistance cannot be im-
proved to any great extent. However, the proper selection
of surface-engineering technique affects the coefficient
of friction and the critical loads for aluminium trans-
fer-layer build up. In contrast to stainless steel, DLC
coatings show no beneficial effect when it comes to
forming aluminium alloys. However, the galling perfor-
mance of DLC coatings also depends on the coating and
the aluminium-alloy type.43–45 In the case of aluminium
alloys a certain reduction in the coefficient of friction
and increased critical loads for the transfer-layer build up
can be expected with the application of nitride-type
coatings (TiN, VN, CrN, etc.) and by nitriding the
tool-steel surface, as shown in Figure 9a. Even smaller
differences between different surface-engineering tech-
niques in terms of galling resistance are found for tita-
nium and titanium alloys, where plasma nitriding shows
the greatest potential (Figure 9b).
3.2 Punching and fine blanking
In punching and fine blanking the cutting elements of
the tool are subjected to high impact loads as well as to a
sliding contact. Impact loads lead to low-cycle fatigue
and the chipping of the cutting edges, while sliding
against steel sheet material causes galling and abrasive
wear. An improvement in the abrasive wear resistance as
well as in the galling resistance can be achieved through
the deposition of hard coatings. At the same time, a sub-
strate material with a high toughness is required in order
to postpone the low-cycle fatigue and the chipping of the
cutting edge. This, on the other hand, means insufficient
load-carrying capacity, which often results in coating
cracking and delamination46. And the harder is the
coating, the higher is the risk of coating failure. For the
coating to show its full potential in reducing friction and
increasing surface wear resistance, we primarily have to
provide sufficient load-carrying capacity for the sub-
strate. This can be achieved by combining a thermo-
chemical treatment of the substrate and hard-coating
deposition, called a duplex treatment.19,47 As shown in
Figure 10a, by plasma nitriding the load-carrying capa-
city of the tool steel can be increased by up to three
times. However, improper preparation of the substrate
(the presence of a ’ compound layer, a high roughness)
or insufficient coating adhesion (TaC) will lead to
immediate flaking of the coating (Figure 10b), regard-
less of the substrate’s properties.
Besides the load-carrying capacity, the tribological
properties and the resistance of the coated surfaces to
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324 319
Figure 10: a) Effect of substrate treatment on the load-carrying capa-
city of coated tool steel and b) flaking of the TaC coating19
Slika 10: a) Vpliv kemo-termi~ne priprave podlage na nosilnost
opla{~enih povr{in in b) lu{~enje prevleke TaC19
Figure 9: Critical loads for galling against: a) aluminium alloy and b)
titanium alloy35
Slika 9: Kriti~ne obremenitve prenosa materiala za: a) aluminijevo in
b) titanovo zlitino35
dynamic loading also depend on the coating type and the
properties of the substrate material,48 as shown in Figure
11. The best resistance is achieved by multilayer coat-
ings composed of a base, load-carrying ceramic layer
(TiAlN, CrN, etc.) and top, low-friction layer (DLC,
Me-C:H, MoS2, etc.). Multilayer coatings are followed
by gradient coatings (i.e., TiCN), while the lowest
resistance to dynamic loading is displayed by monolayer
coatings of TiN, AlCrN, etc. (Figure 11). In terms of
substrate material, due to the poor adhesion and a high
flaking tendency, tungsten carbide is not the most
suitable for coated tools subjected to high dynamic
impact loading. The best substrate material, regardless of
the coating used, is fine-grained, micro-clean tool steels
produced by powder metallurgy (P/M) processes.49,50 In
order to provide a high load-carrying capacity combined
with a superior impact, a wear-resistant tool-steel sub-
strate should have a hardness of at least 64–65 HRc and
a fracture toughness above 12 MPa m1/2.51 Although plas-
ma nitriding provides the highest static load support, its
negative effect on the surface ductility greatly reduces
the resistance of the coated surface on the crack initi-
ation and propagation,52,53 as shown in Figure 12.
The purpose of hard coatings in punching and fine
blanking is not only to increase the wear resistance of the
tool but also to reduce or even eliminate the lubrication.
As shown in Figure 13, regardless of the substrate ma-
terial used, the lowest coefficient of friction ( 0.15) and
the highest potential are shown by carbon-based or
diamond-like-carbon coatings (DLC). On the other hand,
harder ceramic coatings, especially nitride-type multi-
layer coatings, display a much higher friction (> 0.3), but
when deposited on a suitable substrate up to 10 times
less wear (Figure 13).
Experiments with coated punching and fine-blanking
tools have shown that although coatings can greatly
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
320 Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324
Figure 13: a) Effect of substrate material and coating type on: a)
coefficient of friction and b) wear rate in a dry-sliding contact48
Slika 13: a) Vpliv materiala podlage in tipa prevleke na: a) koeficient
trenja in b) stopnjo obrabe pri suhem drsnem kontaktu48
Figure 11: a) Effect of substrate material and coating type on the
load-carrying capacity and b) example of poor coating adhesion in the
case of tungsten carbide substrate48
Slika 11: a) Vpliv materiala podlage in tipa prevleke na nosilnost
opla{~ene povr{ine ter b) primer neustrezne oprijemljivosti na podlago
karbidne trdine48
Figure 12: Effect of substrate treatment type on the crack length and
the density in a gradient TiCN coating deposited on fine-grained
micro-clean P/M tool steel and subjected to a sliding contact at 3500
N52
Slika 12: Vpliv postopka priprave podlage na dol`ino in gostoto
razpok v gradientni TiCN-prevleki, naneseni na finozrnato orodno
jeklo P/M in izpostavljeno drsnemu kontaktu pri 3500 N52
increase the tools’ performance and wear resistance,
neither coating group and type allow the complete elimi-
nation of lubrication.48,54 In the case of hard ceramic
coatings, high friction leads to the adhesion of the work
material to the tool surface and consequently for up to
three times higher extraction forces and even up to ten
times higher punching forces, which then lead to coating
cracking and flaking (Figure 14). On the other hand, for
DLC coatings, although initially reducing the punching
forces, a high coating-wear rate in just a few strokes
leads to failure of the cutting elements. However, the
combination of minimum lubrication, a suitable
fine-grained-steel substrate and its heat treatment, and a
hard, wear-resistant coating can result in the greatly
improved resistance of both sliding surfaces and cutting
edges of the punching and fine blanking tools.
3.3 Hot forging
Tools for hot forging are exposed to the widest and
most demanding loading and contact conditions, which
result in the most severe wear of the tool surface.
Besides the adhesive and abrasive wear, the contact
surfaces are exposed to mechanical and thermal fatigue,
as well as to plastic deformation.31,46 Traditionally,
forging dies are subjected to different thermo-chemical
treatments, mostly to nitriding in order to improve their
wear resistance and lifetime.55 However, due to the very
high contact temperatures, all these treatments have only
a limited effect.46 On the other hand, investigations with
hard coatings, including TiN, TiCN, CrN, TiB2, etc.,
show very promising results.56,57 Coatings prevent ther-
mal fatigue and the formation of intermetallic alloys on
the tool surface, while at the same time they protect the
substrate material from thermal shocks and softening.
Furthermore, with a high hardness, which can be main-
tained even at very high temperatures, coatings also
improve the tool’s abrasive and erosive wear resistance.
As shown in Figure 15, the deposition of gradient or
multilayer coatings can significantly improve the wear
resistance of hot-forging die inserts. When they are only
nitride, the surface hardness of the inserts will drop by
more than 15 % in less than 15000 strokes, eventually
resulting in severe plastic deformation and wear of the
inserts. On the other hand, a combination of plasma
nitriding and a hard PACVD coating was found to pre-
vent a drop in the surface hardness and almost eliminated
the wear of the contact surfaces.56 As shown in Figure
15, even after about 15000 strokes the coated inserts
were able to maintain their initial geometry.
Although the coatings have exceptional thermal and
anti-wear properties, coating deposition on an impro-
perly prepared substrate will lead to a deterioration
rather than an improvement in the performance of forg-
ing dies. Insufficient hardness and a too rough substrate
will cause coating cracking and delamination (Figure
16a), which even accelerates the abrasive wear of the
die. Furthermore, the use of hardened hot-work tool steel
with a coarse microstructure greatly increases the likeli-
hood of cracks initiation and propagation, thus reducing
the tool’s fatigue resistance.5 As in the case of punching
tools, the best results are obtained through duplex
technology, combining plasma nitriding of the tool-steel
substrate and PVD or PACVD coating deposition.57,58
However, besides the substrate hardness, a very import-
ant substrate property is its fracture toughness.59 A dense
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324 321
Figure 15: Forging die insert after 13500 strokes: a) Duplex – TiCN +
nitriding, b) only nitrided56
Slika 15: Vlo`ek kova{kega utopa po 13500 udarcih: a) Duplex –
TiCN + nitriranje, b) nitriran56
Figure 14: Contact surface of punching tool coated with monolayer
AlCrN coating after 4000 strokes: a) fine-grained micro-clean P/M
tool steel substrate and b) WC substrate48
Slika 14: Povr{ina {tancnega no`a, prekritega z AlCrN-prevleko, po
4000 udarcih; podlaga iz: a) visokokvalitetnega jekla P/M in b) kar-
bidne trdine48
’ compound layer may act as an additional support
interlayer, but its brittleness leads to a reduced coating-
cracking resistance as soon as the substrate core hardness
is reduced below 50 HRc. Even for a compound-
layer-free nitrided substrate a high core-hardness level is
the main requirement in providing good load support, but
it needs to be supported by a sufficient fracture-tough-
ness level53 (Figure 17).
4 CONCLUSIONS
The first step in improving the tribological properties
and galling resistance of forming tolls is reducing the
surface roughness. Surface polishing smoothens and
removes the surface irregularities and thus eliminates
potential spots for galling initiation. Increasing surface
hardness or the deposition of a hard coating, although
improving the abrasive wear resistance, exaggerates the
effect of the surface roughness.
In cold sheet-metal forming the main wear mecha-
nism responsible for tool failure is adhesive wear or
galling. The selection of a surface-engineering technique
aimed at improving the anti-galling properties mainly
depends on the material to be formed. In the case of
stainless steel, low and stable friction as well as com-
plete protection of the tool surface against galling under
high loads and boundary lubrication is provided by
carbon-based or diamond-like-carbon coatings. Nitride-
type coatings, on the other hand, are more suitable for
the cold forming of aluminium and aluminium alloys,
and plasma-nitrided tool steel for titanium alloys.
For the successful use of hard coatings and improved
tool performance in punching and fine blanking opera-
tions, first of all the substrate needs to have a sufficient
load-carrying capacity. Besides that, a high hardness
should be coupled with a sufficient fracture toughness.
The best results are obtained when using multilayer coat-
ings on fine-grained micro-clean P/M tool steels. P/M
tool steels give good impact and fatigue resistance, while
an optimal thermo-chemical treatment provides a supe-
rior load support. Although the complete elimination of
lubrication in punching and fine blanking is not yet
possible, the superior wear resistance and friction perfor-
mance of hard coatings allow a reduction in the lubri-
cation quantity and the amount of additives.
Due to the combined mechanical and thermal loads
of hot-forging dies the proper selection and optimal
parameters of substrate thermo-chemical treatment are
crucial for the proper performance of the coated tools.
Instead of the expected improvement an improper
combination of substrate hardness and toughness can
lead to increased wear and accelerated crack initiation
and propagation. The highest potential for improving the
performance of hot-forging dies is shown by the combi-
nation of a plasma-nitrided tool-steel substrate coated
with a multilayer nano-composite coating. However,
before the coating the substrate needs to be polished and
increased substrate hardness combined with sufficient
fracture toughness.
5 REFERENCES
1 P. Groche, M. Christiany, Evaluation of the potential of tool materials
for the cold forming of advanced high strength steels, Wear, 302
(2013), 1279–1285, doi:10.1016/j.wear.2013.01.001
2 B. Podgornik, S. Hogmark, Surface modification to improve friction
and galling properties of forming tools, Journal of Materials Process-
ing Technology, 174 (2006), 334–341, doi:10.1016/j.jmatprotec.
2006.01.016
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
322 Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324
Figure 17: Effect of substrate hardness and fracture toughness on the
load-carrying capacity of TiN/TiB2-coated hot-work tool steel53
Slika 17: Vpliv trdote in lomne `ilavosti orodnega jekla za delo v vro-
~em na kriti~no obremenitev nastanka razpok v TiN/TiB2-prevleki53
Figure 16: Failure of the coated die insert: a) flaking and b) crack
initiation and propagation
Slika 16: Po{kodba opla{~enega kova{kega orodja: a) lu{~enje in b)
nastanek ter {irjenje razpok
3 C. Mitterer, F. Holler, C. Lugmair, R. Nobauer, R. Kullmer, C. Tei-
chert, Optimization of plasma-assisted chemical vapour deposition
hard coatings for their application in aluminium die-casting, Surface
and Coatings Technology, 142-144 (2001), 1005–1011, doi:10.1016/
S0257-8972(01)01065-9
4 A. Attanasio, E. Ceretti, C. Giardini, L. Mazzoni, Asymmetric two
points incremental forming: Improving surface quality and geometric
accuracy by tool path optimization, Journal of Materials Processing
Technology, 197 (2008), 59–67, doi:10.1016/j.jmatprotec.2007.
05.053
5 A. Persson, J. Bergstrom, S. Hogmark, Influence of surface engineer-
ing on the performance of tool steels for die casting, 6th International
Tooling Conference, Karlstad, Sweden, 2002, 841–854
6 R. Jervis, B. Johnson, L. A. Norstrom, O. Sandberg, Properties
profile of a new presswork tool steel, International Conference on
Tooling Materials, Interlaken, Switzerland, 1992, 49–64
7 B. Podgornik, S. Hogmark, O. Sandberg, Influence of surface rough-
ness and coating type on the galling properties of coated forming tool
steel, Surface and Coatings Technology, 184 (2004), 338–348,
doi:10.1016/j.surfcoat.2003.11.002
8 D. Wang, H. Li, H. Yang, J. Ma, G. Li, Tribological evaluation of
surface modified H13 tool steel in warm forming of Ti–6Al–4V
titanium alloy sheet, Chinese Journal of Aeronautics, 27 (2014),
1002–1009, doi:10.1016/j.cja.2014.03.030
9 S. R. Schmid, W. R. D. Wilson, Tribology in manufacturing, In: B.
Bhushan (Ed.), Modern Tribology Handbook, CRC Press, NY 2000
10 S. Gulizia, M. Z. Jahedi, Evaluation of PVD coatings for industrial
applications, 6th International Tooling Conference, Karlstad,
Sweden, 2002, 739–748
11 Y. Hou, W. Zhang, Z. Yu, S. Li, Selection of tool materials and
surface treatments for improved galling performance in sheet metal
forming, The International Journal of Advanced Manufacturing
Technology, 43 (2009), 1010–1017, doi:10.1007/s00170-008-1780-2
12 B. Bhushan, B. K. Gupta, Handbook of Tribology: Materials,
Coatings and Surface Treatments, McGraw-Hill, NY 1991
13 H. Paschke, M. Weber, G. Braeuer, T. Yilkiran, B. A. Behrens, H.
Brand, Optimized plasma nitriding processes for efficient wear
reduction of forging dies, Archives of Civil and Mechanical
Engineering, 12 (2012), 407–412, doi:10.1016/j.acme.2012.06.001
14 S. J. Bull, R. I. Davidson, E. H. Fisher, A. R. McCabe, A. M. Jones,
A simulation test for the selection of coatings and surface treatments
for plastics injection moulding machines, Surface and Coatings
Technology, 130 (2000), 257–265, doi:10.1016/S0257-8972(00)
00697-6
15 B. Podgornik, V. Leskov{ek, J. Vizintin, Influence of deep-cryogenic
treatment on tribological properties of P/M high-speed steel, Mate-
rials and Manufacturing Processes, 24 (2009), 734–738, doi:10.1080/
10426910902809339
16 B. Podgornik, F. Majdic, V. Leskovsek, J. Vizintin, Improving tribo-
logical properties of tool steels through combination of deep-cryoge-
nic treatment and plasma nitriding, Wear, 288 (2012), 88–93,
doi:10.1016/j.wear.2011.04.001
17 K. L. Rutherford, S. J. Bull, E. D. Doyle, I. M. Hutchings, Labo-
ratory characterisation of the wear behaviour of PVD-coated tool
steels and correlation with cutting tool performance, Surface and
Coatings Technology, 80 (1996), 176–180, doi:10.1016/0257-8972
(95)02706-8
18 M. Stoiber, M. Panzenbock, C. Mitterer, C. Lugmair, Fatigue pro-
perties of Ti-based hard coatings deposited onto tool steels, Surface
and Coatings Technology, 142–144 (2001), 117–124, doi:10.1016/
S0257-8972(01)01220-8
19 B. Podgornik, S. Hogmark, O. Sandberg, V. Leskovsek, Wear resis-
tance and anti-sticking properties of duplex treated forming tool
steel, Wear, 254 (2003), 1113–1121, doi:10.1016/S0043-1648(03)
00322-3
20 V. Imbeni, C. Martini, E. Lanzoni, G. Poli, I. M. Hutchings, Tribo-
logical behaviour of multi-layered PVD nitride coatings, Wear, 251
(2001), 997–1002, doi:10.1016/S0043-1648(01)00706-2
21 J. Heinrichs, S. Jacobson, Mechanisms of Transfer of Aluminium to
PVD-Coated Forming Tools, Tribology Letters, 46 (2012), 299–312,
doi:10.1007/s11249-012-9952-5
22 C. Donnet, A. Erdemir, Tribology of diamond-like carbon films:
Fundamentals and Applications, Springer, 2008
23 B. Podgornik, J. Vi`intin, O. Wänstrand, M. Larsson, S. Hogmark,
H. Ronkainen, K. Holmberg, Tribological properties of plasma
nitrided and hard coated AISI 4140 steel, Wear, 249 (2001),
254–259, doi:10.1016/S0043-1648(01)00564-6
24 I. Ebrahimzadeh, F. Ashrafizadeh, High temperature wear and fric-
tional properties of duplex-treated tool steel sliding against a two
phase brass, Ceramics International, 40 (2014), 16429–16439,
doi:10.1016/j.ceramint.2014.07.151
25 B. Podgornik, J. Vizintin, S. Hogmark, Improvement in galling
performance through surface engineering, Surface Engineering, 22
(2006), 235–238, doi:10.1179/174329406X122946
26 B. Podgornik, J. Jerina, Surface topography effect on galling resis-
tance of coated and uncoated tool steel, Surface and Coatings Tech-
nology, 206 (2012), 2792–2800, doi:10.1016/j.surfcoat.2011.11.041
27 J. A. Schey, Tribology in Metalworking: Friction, Lubrication and
Wear, ASM, Ohio 1984
28 C. Subramanian, K. N. Strafford, T. P. Wilks, L. P. Ward, On the
design of coated systems: metallurgical and other considerations,
Journal of Materials Processing Technology, 56 (1996), 385–397,
doi:10.1016/0924-0136(95)01852-2
29 L. Norstrom, B. Johansson, Selection of tool steels for cold-work
applications, International Conference on Tooling, Bochum, Ger-
many, 1989, 193–208
30 A. Gåård, Influence of tool microstructure on galling resistance, Tri-
bology International, 57 (2013), 251–256, doi:10.1016/j.triboint.
2012.08.022
31 R. Seidel, H. Luig, Friction and wear processes in hot die forging,
International Conference on Tooling Materials, Interlaken,
Switzerland, 1992, 467-480
32 O. Barrau, C. Boher, C. Vergne, F. Rezai-Aria, Investigations on
friction and wear mechanisms of hot forging tool steels, 6th
International Tooling Conference, Karlstad, Sweden, 2002, 81–94
33 P. Carlsson, M. Olsson, PVD coatings for sheet metal forming pro-
cesses – a tribological evaluation, Surface and Coatings Technology,
200 (2006), 4654–4663, doi:10.1016/j.surfcoat.2004.10.127
34 H. Paschke, M. Weber, P. Kaestner, G. Braeuer, Influence of different
plasma nitriding treatments on the wear and crack behavior of
forging tools evaluated by Rockwell indentation and scratch tests,
Surface and Coatings Technology, 205 (2010), 1465–1469, doi:
10.1016/j.surfcoat.2010.07.053
35 B. Podgornik, J. Vi`intin, Galling properties of forming tool steels -
How to determine and improvements through surface engineering,
International Conference on Tribology in Manufacturing Processes,
Yokohama, Japan, 2007, 195–202
36 A. Gåård, R. M. Sarih, Influence of Tool Material and Surface
Roughness on Galling Resistance in Sliding Against Austenitic
Stainless Steel, Tribology Letters, 46 (2012), 179–185, doi:10.1007/
s11249-012-9934-7
37 M. Geiger, U. Popp, U. Engel, Excimer Laser Micro Texturing of
Cold Forging Tool Surfaces - Influence on Tool Life, Manufacturing
Technology, 51 (2002), 231–234, doi:10.1016/S0007-8506(07)
61506-6
38 T. D. Ling, P. Liu, S. Xiong, D. Grzina, J. Cao, Q. J. Wang, Z. C.
Xia, R. Talwar, Surface Texturing of Drill Bits for Adhesion
Reduction and Tool Life Enhancement, Tribology Letters, 52 (2013),
113–122, doi:10.1007/s11249-013-0198-7
39 B. Podgornik, M. Sedla~ek, Performance, characterization and
design of textured surfaces, Journal of Tribology, 134 (2012),
041701/1–041701/7, doi:10.1115/1.4007108
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324 323
40 L. Lu, Q. M. Wang, B. Z. Chen, Y. C. Ao, D. H. Yu, C. Y. Wang, S.
H. Wu, K. H. Kim, Microstructure and cutting performance of
CrTiAlN coating for high-speed dry milling, Transactions of Non-
ferrous Metals Society of China, 24 (2014), 1800–1806, doi:10.1016/
S1003-6326(14)63256-8
41 B. Podgornik, S. Hogmark, O. Sandberg, Proper coating selection for
improved galling performance of forming tool steel, Wear, 261
(2006), 15–21, doi:10.1016/j.wear.2005.09.005
42 A. Ghiotti, S. Bruschi, Tribological behaviour of DLC coatings for
sheet metal forming tools, Wear, 271 (2011), 2454–2458,
doi:10.1016/j.wear.2010.12.043
43 T. Horiuchi, S. Yoshihara, Y. Iriyama, Dry deep drawability of
A5052 aluminum alloy sheet with DLC-coating, Wear, 286–287
(2012), 79–83, doi:10.1016/j.wear.2011.07.005
44 J. Heinrichs, M. Olsson, S. Jacobson, Initiation of Galling in Metal
Forming: Differences Between Aluminium and Austenitic Stainless
Steel Studied In Situ in the SEM, Tribology Letters, 50 (2013),
431–438, doi:10.1007/s11249-013-0139-5
45 J. Heinrichs, S. Jacobson, Mechanisms of Transfer of Aluminium to
PVD-Coated Forming Tools, Tribology Letters, 46 (2012), 299–312,
doi:10.1007/s11249-012-9952-5
46 G. Hakansson, Experiences of surface coated tools, International
Conference on Recent Advances in Manufacture & Use of Tools &
Dies and Stamping of Steel Sheets, Olofstrom, Sweden, 2004,
245–254
47 B. Podgornik, J. Vi`intin, V. Leskov{ek, O. Wänstrand, M. Larsson,
S. Hogmark, Wear properties of plasma nitrided and hard coated
42CrMo4 steel, Mater. Tehnol., 34 (2000) 1–2, 17–21
48 B. Podgornik, B. Zajec, N. Bay, J. Vi`intin, Application of hard
coatings for blanking and piercing tools, Wear, 270 (2011), 850–856,
doi:10.1016/j.wear.2011.02.013
49 P. Jur~i, B. [u{tar{i~, V. Leskov{ek, Fracture characteristics of the
Cr-V ledeburitic steel Vanadis 6, Mater. Tehnol., 44 (2010) 2, 77–84
50 P. Karlsson, A. Gåård, P. Krakhmalev, Influence of tool steel
microstructure on friction and initial material transfer, Wear, 319
(2014), 12–18, doi:10.1016/j.wear.2014.07.002
51 B. Podgornik, M. Sedla~ek, V. Leskov{ek, Surface engineering of
fine blanking tool inserts, Project report, Institute of Metals and
Technology, Ljubljana, 2014
52 B. Podgornik, V. Leskov{ek, B. Arh, The effect of heat treatment on
the mechanical, tribological and load-carrying properties of
PACVD-coated tool steel, Surface and Coatings Technology, 232
(2013), 528–534, doi:10.1016/j.surfcoat.2013.06.019
53 B. Podgornik, V. Leskov{ek, F. Tehovnik, J. Burja, Vacuum heat
treatment optimization for improved load carrying capacity and wear
properties of surface engineered hot work tool steel, Surface and
Coatings Technology, 261 (2015), 253–261, doi:10.1016/j.surfcoat.
2014.11.021
54 M. Çöl, D. Kir, E. Eriºir, Wear and blanking performance of AlCrN
PVD-coated punches, Materials Science, 48 (2013), 514–520
55 M. Bayramoglu, H. Polat, N. Geren, Cost and performance
evaluation of different surface treated dies for hot forging process,
Journal of Materials Processing Technology, 205 (2008), 394–403,
doi:10.1016/j.jmatprotec.2007.11.256
56 V. Leskov{ek, B. Podgornik, M. Jenko, A PACVD duplex coating for
hot-forging applications, Wear, 266 (2009), 453–460, doi:10.1016/
j.wear.2008.04.016
57 E. Doege, E. Westkamper, R. Seidel, C. Romanowski, G. Andreis, O.
Morlok, Combined surface treatment of forging tools by plasma
nitriding and PACVD, Progress in tool steels, Bochum, Germany,
1996, 373–384
58 P. Panjan, I. Urankar, B. Navin{ek, M. Ter~elj, R. Turk, M. ^ekada,
V. Leskov{ek, Improvement of hot forging tools with duplex treat-
ment, Surface and Coatings Technology, 151–152 (2002), 505–509,
doi:10.1016/S0257-8972(01)01634-6
59 V. Leskov{ek, B. Podgornik, Multi-functional KIc-test specimen for
the assessment of different tool- and high-speed-steel properties,
Mater. Tehnol., 47 (2013) 3, 273–283
B. PODGORNIK, V. LESKOV[EK: WEAR MECHANISMS AND SURFACE ENGINEERING OF FORMING TOOLS
324 Materiali in tehnologije / Materials and technology 49 (2015) 3, 313–324
R. S. RAHBAR, M. VADOOD: PREDICTING THE PHYSICAL PROPERTIES OF DRAWN NYLON-6 FIBERS ...
PREDICTING THE PHYSICAL PROPERTIES OF DRAWN
NYLON-6 FIBERS USING AN ARTIFICIAL-NEURAL-NETWORK
MODEL
NAPOVEDOVANJE FIZIKALNIH LASTNOSTI VLE^ENIH
VLAKEN IZ NAJLONA 6 Z UPORABO MODELA UMETNE
NEVRONSKE MRE@E
Ruhollah Semnani Rahbar1, Morteza Vadood2
1Department of Textile and Leather, Faculty of Chemistry and Petrochemical Engineering, Standard Research Institute (SRI), Karaj,
P. O. Box 31745-139, Iran
2Department of Textile Engineering, Amirkabir University of Technology, Tehran, Iran
semnani@standard.ac.ir
Prejem rokopisa – received: 2013-08-13; sprejem za objavo – accepted for publication: 2014-07-11
doi:10.17222/mit.2013.128
Low-oriented nylon-6 fibers were drawn in a multistage drawing process, during which the number of drawing steps and the
temperature of each step were changed. The physical properties of these fibers were measured and compared with the values
predicted by a multiple-linear-regression model. Moreover, six input variables and four output variables were used in an
artificial neural network (ANN) to establish the logical relationships between the inputs and outputs. Attempts were also made
to determine the effective parameters for each physical property and explain the observed trends. The results showed that the
models based on the ANN performed well and provided stable responses in predicting combined interactions between
independent variables.
Keywords: drawing process, artificial neural network, modeling, physical properties
Malo orientirana vlakna najlon 6 so bila vle~ena z ve~stopenjskim postopkom, pri ~emer se je pri vsakem vleku spreminjala
stopnja vle~enja in temperatura. Izmerjene vrednosti teh vlaken so bile primerjane z vrednostmi, napovedanimi z modelom
multivariantne linearne regresije. Poleg tega je bilo v umetni nevronski mre`i (ANN) uporabljenih {est vhodnih spremenljivk in
{tiri izhodne, da bi ugotovili logi~ne odvisnosti med vhodnimi in izhodnimi spremenljivkami. Posku{alo se je ugotoviti
u~inkovite parametere za vsako fizikalno lastnost in razlo`iti opa`ene tendence. Rezultati so pokazali, da so modeli na osnovi
ANN dobri in ponujajo stabilne odgovore pri predvidevanju kombiniranih interakcij neodvisnih spremenljivk.
Klju~ne besede: postopek vle~enja, umetna nevronska mre`a, modeliranje, fizikalne lastnosti
1 INTRODUCTION
Synthetic-fiber drawing is a critical process to obtain
fibers with desired properties for final applications.
There are various parameters in this process that should
be controlled to yield a fiber with acceptable technical
specifications. These variables include the draw ratio, the
drawing temperature, the number of drawing steps, the
drawing speed and the distribution of draw ratio in the
multistage drawing1. Because of the complex relation-
ships between the fiber properties and the drawing-pro-
cess variables, there is a need for a sound experimental
design and a careful analysis of the experimental results.
In this way, the relationships between a measured
characteristic of a drawn fiber and the influencing factors
can be identified and optimized. The understanding of
these relationships reduces the processing cost and
provides for reproducibility in a day-to-day production.
Moreover, the risk of excessive downtime for trials is
reduced2–5.
Due to their simplicity, regression-based models and
statistical analyses were extensively used to solve textile
technological problems2–4,6–9. However, they have certain
limitations as they require a specialized knowledge of
both the statistical methods and techniques of experi-
mental design. Moreover, the prediction ability of the
regression analysis may be limited in the case of an
analysis of multidimensional technical problems10,11.
In recent years, artificial neural network (ANN) has
been used as an alternative modeling method in many
different engineering fields to predict the properties of
materials. ANN can be considered as a black box con-
sisting of a series of complex equations for estimating
the outputs on the basis of a given series of input values.
The advantage of ANNs is the ability of representing
complex relationships directly from the data being
modeled, while their representation (modeling) is always
nonlinear12–14.
Many researches were done in the textile industry to
predict the properties of yarns, woven and nonwoven
fabrics and many other characteristics of textile mate-
rials10,13–25. Among these, just few studies are devoted to
melt spinning and drawing of synthetic fibers. It seems
that there is a lack of information concerning the appli-
Materiali in tehnologije / Materials and technology 49 (2015) 3, 325–332 325
UDK 677.017:519.233.4 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)325(2015)
cation of ANN for predicting fiber properties during
multistage drawing. Therefore, in this paper, an ANN
model was used to predict some physical properties of
drawn nylon-6 fibers upon multistage drawing.
2 EXPERIMENTAL WORK
2.1 Materials and methods
A low-oriented nylon-6 multifilament yarn, 340 dtex,
24 filaments, was kindly supplied by Alyaf Co. (Iran).
Nylon 6 (polyamide 6) is made up of linear macromole-
cules whose structural units are linked with an amide
linkage (-NH-CO-group). This synthetic fiber is used for
a variety of different applications ranging from garments
to industrial usages.
Fiber samples were prepared by melt spinning and
the take-up speed of 800 m/min. The undrawn multifila-
ment yarn had a polymerization degree of 148. The
drawing process was performed on an industrial Zinser
draw-twisting machine (Germany), type 520-2. The de-
tails of the drawing process were explained in the
previous work26. The fixed drawing conditions are listed
in Table 1.
Table 1: Fixed operating conditions of the drawing experiments
Tabela 1: Pogoji obratovanja pri preizkusih vle~enja
Temperature
of the feeding
roller (°C)
Temperature
of the third
godet roller
(°C)
Drawing
speed
(m/min)
Intermingling
jet pressure
(bar)
Spindle
speed
(r/min)
Room
temperature
Room
temperature 400 2 4000
Six process parameters of drawing trials included the
first-stage draw ratio (FSDR), the second-stage draw ra-
tio (SSDR), the third-stage draw ratio (TSDR), the first-
godet temperature (FGT), the block-heater temperature
(BHT), and the second-godet temperature (SGT). In the
trials, one-, two- and three-stage drawn fibers and also
the fiber heat-treated without drawing were produced to
have a broad range of fibers with different structures. A
total of 61 different fibers were prepared and the drawing
trials are reported in Table 2.
Table 2: Experimental array for the drawing process
Tabela 2: Pogoji eksperimenta pri postopku vle~enja
Run
F
ir
st
-s
ta
ge
dr
aw
ra
ti
o,
F
SD
R
S
ec
on
d-
st
ag
e
dr
aw
ra
ti
o,
SS
D
R
T
hi
rd
-s
ta
ge
dr
aw
ra
ti
o,
T
SD
R
F
ir
st
-g
od
et
te
m
pe
ra
tu
re
(°
C
),
F
G
T
B
lo
ck
-h
ea
te
r
te
m
pe
ra
tu
re
(°
C
),
B
H
T
S
ec
on
d-
go
de
t
te
m
pe
ra
tu
re
(°
C
),
SG
T
1 1.1 4.126 1.3 100 170 170
2 2.098 1.611 1.626 100 170 170
3 2.8 1.621 1.3 100 170 170
4 1.755 1.5 1.3 100 170 170
5 1.755 2.586 1.3 100 170 170
6 1.615 1.884 1.939 100 170 170
7 1.1 2.488 1.3 100 170 170
8 1.1 2.551 2.102 100 170 170
9 1.1 1.5 1.3 100 170 170
10 1.315 2.074 1.564 100 170 170
11 1.755 1.97 1.3 100 170 170
12 1.348 1.611 2.53 100 170 170
13 1.8 1.5 2.185 100 170 170
14 1.1 1.5 2.102 100 170 170
15 1.1 1.5 3.4 100 170 170
16 1.637 1.5 1.792 150 170 170
17 1.1 1.739 2.3 150 170 170
18 1.1 1.5 1.729 150 150 150
19 1.1 1.615 2.3 30 30 30
20 1.1 3.077 1.3 150 150 150
21 1.1 2.313 1.729 30 170 170
22 1.253 1.975 1.676 60 80 80
23 1.1 3.077 1.3 90 170 170
24 2.1 1.612 1.3 150 170 170
25 1.1 1.5 1.3 30 30 30
26 1.1 1.5 1.3 30 170 170
27 1.1 2.148 1.3 150 150 150
28 1.1 1.5 1.3 90 90 90
29 2.1 1.5 1.397 30 30 30
30 1.1 3.077 1.3 30 170 170
31 1.637 1.5 1.792 30 170 170
32 1.275 1.5 2.3 30 100 100
33 1.52 1.5 1.3 30 170 170
34 1.732 1.591 1.506 90 110 110
35 1.52 2.227 1.3 30 30 30
36 1.1 1.5 1.3 150 170 170
37 2.1 1.612 1.3 30 170 170
38 1.1 1.5 2.3 30 170 170
39 1.1 3.077 1.3 30 30 30
40 1.637 1.5 1.792 30 30 30
41 1.275 1.5 2.3 90 90 90
42 1.1 1.5 1 100 170 170
43 1.1 4.2 1 100 170 170
44 1.1 4.2 1.3 100 170 170
45 2.8 1 1 30 30 30
46 2.8 1 1 100 30 30
47 1.1 4.126 1.3 100 30 170
48 1.1 4.126 1.3 100 60 170
49 1.1 4.126 1.3 100 100 170
50 1.1 4.126 1.3 100 150 170
51 1.008 4.6 1 100 170 170
52 1.1 4.215 1 100 170 170
53 1.1 3.242 1.3 100 170 170
54 2.8 1.648 1.3 100 170 170
55 2.8 1.506 1.3 100 170 170
56 1.187 1.5 2.5 30 30 30
57 2.1 1.564 1.3 30 30 30
58 1.579 1.5 1.879 30 30 30
59 1.1 1 1 30 30 30
60 1.1 1 1 100 30 30
61 1 1 1 100 30 30
The yarn linear density (expressed in dtex) was deter-
mined in accordance with ASTM D 1577-96. The mean
value is the average of five measurements.
R. S. RAHBAR, M. VADOOD: PREDICTING THE PHYSICAL PROPERTIES OF DRAWN NYLON-6 FIBERS ...
326 Materiali in tehnologije / Materials and technology 49 (2015) 3, 325–332
Stress-strain curves were obtained with an EMT-3050
tensile testing machine (Elima Co., Iran). The initial
lengths of the fibers and cross-head speeds were 300 mm
and 500 mm/min, respectively. From the stress-strain
plots, the initial modulus, the tenacity and the specific
work of rupture were evaluated. The reported values for
all the mechanical properties were averaged over at least
ten independent measurements.
Yarn-shrinkage measurements were made according
to DIN 53840 at 130 °C for 10 min. The initial and final
lengths were measured at room temperature and the total
shrinkage was defined as the fraction of the initial sam-
ple length remaining after the exposure to the elevated
temperature. The reported values are the average of eight
tests.
2.2 Artificial neural network
Neural networks are mainly composed of the pro-
cessing elements called neurons with interconnections.
The exclusive structure of an ANN makes it very appro-
priate for modeling a complex system with nonlinear
relations between the parameters. Generally, an ANN
can be made of many layers, namely the input, output
and several hidden layers. The neurons in each layer are
connected with the associated weights to the other
neurons in the next layer. The input layer receives the
input parameters and, through the hidden layers based on
Equation (1), the output can be calculated in the output
layer19. Figure 1 indicates the topology of an ANN with
one hidden layer:
y f w f w x b bk jk
j
m
ij
i
n
i j k= +
⎛
⎝
⎜ ⎞
⎠
⎟ +
⎛
⎝
⎜
⎞
⎠
⎟
= =
∑ ∑
1 1
1 2 (1)
Here wij, wjk, b1j and b2k are the weight between the ith
input neuron and the jth hidden neuron, the weight bet-
ween the jth hidden neuron and the kth output neuron, the
bias for the jth hidden neuron and the bias for the kth out-
put neuron, respectively. The f(x) is the activation func-
tion27,28. All the data is divided into three groups known
as the training, validation and testing sets. The first
group is used to train the ANN. When the network
begins to overfit the data, the training is stopped for a
specified number of iterations (maximum number of
fails) and the weights and biases are returned with the
minimum error on the validation set. The testing group is
used to control the error after the training process21. In
this study, to apply ANNs, the Matlab R2008 software
was used. The ANNs were trained on the basis of the
error back-propagation algorithm using the "Trainlm" func-
tion to avoid an over-fitting error. The activation func-
tions for all the hidden and output layers were con-
sidered as the hyperbolic tangent and linear function,
respectively (this type of ANN is called the perceptron).
3 RESULTS AND DISCUSSION
3.1 Linear regression model
In this research, multiple linear regression (MLR)
was employed to evaluate the performance of different
models. In order to determine the variables that can
directly affect the physical properties, as well as
decreasing the number of variables, a traditional linear-
regression model at the 90 % confidence level was
accomplished. This procedure was only used for three
variables including the FSDR, SSDR and TSDR because
they were varied at multiple levels. The FSDR, SSDR
and TSDR were changed at the 18th, 21st and 17th levels,
respectively. Moreover, theses parameters were indivi-
dually considered with respect to the responses of the
linear-regression model. Table 3 shows the Pearson
correlation coefficient (PCC) and the corresponding
P-value between the measured parameters (FSDR, SSDR
and TSDR) and the physical properties.
Table 3: PCC and P-value between the measured parameters and
physical properties
Tabela 3: PCC in P-vrednost med izmerjenimi parametri in fizikal-
nimi lastnostmi
Shrinkage Specific workof rupture Tenacity
Initial
modulus
PCC P-value PCC
P-
value PCC
P-
value PCC
P-
value
FSDR 0.23 0.08 –0.25 0.05 0.23 0.08 0.25 0.06
SSDR –0.19 0.14 –0.49 0.00 0.59 0.00 0.31 0.01
TSDR 0.38 0.00 –0.32 0.01 0.14 0.29 0.36 0.00
The PCC varies between –1 and 1; the more absolute
value of the PCC indicates the existence of a strong rela-
tion between two parameters. P-values are used for test-
ing the hypothesis of no correlation. Each P-value is the
probability of getting a correlation as large as the ob-
served value by random chance when the true correlation
is zero. If a P-value is less than 0.1, the corresponding
R. S. RAHBAR, M. VADOOD: PREDICTING THE PHYSICAL PROPERTIES OF DRAWN NYLON-6 FIBERS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 325–332 327
Figure 1: ANN with one hidden layer
Slika 1: ANN z eno skrito plastjo
correlation is significant. But the data for the FGT, BHT
and SGT were different and they were varied only at the
5th, 9th and 7th levels, respectively. Hence, a one-way
ANOVA was used to evaluate their effects on the physi-
cal properties and the results are listed in Table 4. The
confidence level was 90 %. As mentioned before, the
parameters with the P-value below 0.1 have a significant
effect on the physical parameters. The greater the
F-value, the greater are the influence and the relevance
of the source factor.
Table 4: ANOVA results identifying the statistical significance of
FGT, BHT, and SGT for the physical properties
Tabela 4: Rezultati ANOVA prikazujejo statisti~no pomembnost FGT,
BHT in SGT na fizikalne lastnosti
Shrinkage
Specific
work of
rupture
Tenacity Initialmodulus
F
G
T
SSE 103.202 62.662 4003.172 386497.497
DF 4.000 4.000 4.000 4.000
MSE 25.801 15.665 1000.793 96624.374
F 3.480 0.703 5.435 2.339
P 0.013 0.593 0.001 0.066
B
H
T
SSE 185.671 240.453 3159.013 710707.421
DF 8.000 8.000 8.000 8.000
MSE 23.209 30.057 394.877 88838.428
F 3.627 1.461 1.841 2.323
P 0.002 0.194 0.090 0.033
SG
T
SSE 176.808 235.435 3412.317 694829.932
DF 6.000 6.000 6.000 6.000
MSE 29.468 39.239 568.720 115804.989
F 4.659 1.971 2.817 3.119
P 0.001 0.086 0.019 0.011
SSE: sum of squares, DF: degree of freedom, MSE: mean squared
error, F: F-statistic (the ratio of the mean squares), P: P-value
According to Tables 3 and 4, the effective factors for
each physical property were determined and shown in
Table 5.
Table 5: Parameters affecting the physical properties
Tabela 5: Parametri, ki u~inkujejo na fizikalne lastnosti
Shrinkage Specific workof rupture Tenacity
Initial
modulus
FSDR + + + +
SSDR – + + +
TSDR + + – +
FGT + – + +
BHT + – + +
SGT + + + +
+ : effective, – : non-effective
After determining the effective parameters, multiple-
linear-regression (MLR) analyses based on Table 5 were
performed. The results (Figure 2) indicated that the
developed MLR models provided different predictions
for the physical properties.
According to Figure 2, MLR can predict the specific
work of rupture and the shrinkage with the highest and
the lowest accuracies, respectively. Table 6 shows the
coefficients of MLR for predicting physical properties.
Actually, each coefficient in MLR is a partial derivative
of the model response with respect to the variable of that
coefficient. So, the contribution of each variable in pre-
dicting the model response can be assessed by checking
the coefficient values.
Referring to Table 6, the TSDR and the SGT have the
highest and the lowest effects on the shrinkage, respec-
tively. A similar trend was observed for the specific work
of rupture. Moreover, for the tenacity and the initial
modulus, the FSDR and the TSDR play the main roles,
respectively. Also, the BHT and the FGT have the lowest
effects on the tenacity and the initial modulus, respec-
tively. Although increasing the order of the regression
equation gives better results, it makes the regression
equation more complicated. For example, using a quad-
ratic regression including linear, interaction and square
terms to predict the initial modulus increases the number
of coefficients up to 28. Therefore, using an ANN model
in this situation can be very beneficial. The ANN con-
tains various parameters such as the number of hidden
layers and the number of neurons in each hidden layer
that directly affect the output of the ANN. Hence, to
determine the best set of the ANN parameters, the trial-
and-error method was applied. The numbers of hidden
layers and neurons in each hidden layer were considered
to be between 1 to 5 and 1 to 6, respectively. In this
study the stopping criteria for the ANN training were the
following three options: the training tolerance level (0),
the maximum number of fails (6 iterations) or the total
number of iterations for training (1000 iterations). The
experimental data were used to train the network; 37, 12
and 12 data sets were randomly chosen for the training,
R. S. RAHBAR, M. VADOOD: PREDICTING THE PHYSICAL PROPERTIES OF DRAWN NYLON-6 FIBERS ...
328 Materiali in tehnologije / Materials and technology 49 (2015) 3, 325–332
Figure 2: Results of the MLR prediction for various physical proper-
ties (R2 = PCC)
Slika 2: Rezultati MLR-napovedovanja razli~nih fizikalnih lastnosti
(R2 = PCC)
validation and testing groups, respectively. To evaluate
the accuracy of the developed ANN, the absolute value
of the PCC between the ANN outputs and the actual
values for the testing set was calculated. A higher PCC
indicates a higher accuracy of the ANN. To remove the
effects of the initial weights and biases on the ANN out-
put, each ANN structure was created five times with the
weights and biases chosen randomly and the ANN with
the highest absolute value of the PCC was considered for
that structure.
Table 6: Coefficients of MLR models for predicting the physical pro-
perties
Tabela 6: Koeficienti MLR-modela za napovedovanje fizikalnih last-
nosti
Shrinkage
Specific
work of
rupture
Tenacity Initialmodulus
Constant 1.777 36.461 –14.612 –697.759
FSDR 1.688 –6.795 15.768 251.684
SSDR – –4.518 10.865 115.788
TSDR 2.899 –6.892 – 270.758
FGT 0.008 – –0.014 0.347
BHT –0.023 – –0.010 –1.231
SGT 0.003 –0.002 0.065 2.084
Unlike the regression method, an ANN can predict
physical properties simultaneously with four neurons in
the output layer, but such results showed an insufficient
accuracy. In this way the prediction ability of an ANN
for each output neuron only involves the weights bet-
ween the last hidden layer and the output layer; so, con-
sidering separate ANNs for each physical property (one
neuron in the output layer) would give a higher accuracy
in the prediction. As the ANN calculations using com-
puters can be performed quickly, to consider all the
terms of the effects, such as linear terms or interactions,
all the input parameters were considered for the ANN
models. As expected, it was found that the structure of
the neural network for each physical property was diffe-
rent and exhibited a high prediction performance of the
models in terms of the PCC, as given in Table 7. The
individual values of the predictions of the models, along
with the experimental values, are shown in Figure 3. By
comparing Table 7 and Figure 2, it can be found that the
prediction potential of the ANN models is superior to
that of the MLR. This means that removing some input
variables could not enhance the quality of the MLR
predictions as compared to the ANN models.
Table 7: Best ANN structure for predicting various mechanical
properties
Tabela 7: Najbolj{a struktura ANN za napovedovanje razli~nih
mehanskih lastnosti
Physical property Maximum PCC Hidden layer
Shrinkage 0.9182 [1-1-2]
Specific work of rupture 0.9907 [3]
Tenacity 0.981 [3-4-2]
Initial modulus 0.9917 [2-1-3]
In Table 7, [1-1-2] in the hidden-layer column means
that the ANN contains three hidden layers with 1, 1 and
2 neurons in the first, second and third hidden layers,
respectively. As can be seen in Figure 2, the PCC value
for the specific work of rupture modeled with MLR is
0.96, indicating that the relation between this parameter
and its effective variables is not complicated. Also,
according to Table 7, a simple ANN can model the
specific work of rupture with the PCC value of almost 1
which indicates a perfect prediction. As observed from
Table 7, there is an increase in the PCC values for the
shrinkage, tenacity and initial modulus compared to the
values of the MLR predictions as shown in Figure 2.
This means that the relations between the parameters are
no longer linear and the MLR prediction is not reliable
any more, whereas the ANN using multilayers could
predict these responses very well.
3.2 Sensitivity analysis of the ANN model
Because of the complex and nonlinear form of the
ANN analysis, a sensitivity analysis was conducted to
study the influence of the input variables on the output.
As a matter of fact, the aim of this analysis was to
evaluate the effects of the changes in each input on the
ANN output. In this process, the value of one input
variable from the initial condition is slightly changed
(10–50 %) and then the output value is predicted, while
all the other variables are set to the selected constant
values. This process is repeated for all the input
variables. The robustness of the model is determined by
examining how well the predictions compare with the
available structural knowledge. Figure 4 shows the
effects of the changes in the input variables from –50 %
to +50 % in the scale of 0.1 for the physical properties. It
is evident from Figure 4a that the shrinkage initially
increases with the TSDR and then reaches a constant
value as the TSDR moves toward higher values. The
changes in the other input variables have no significant
R. S. RAHBAR, M. VADOOD: PREDICTING THE PHYSICAL PROPERTIES OF DRAWN NYLON-6 FIBERS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 325–332 329
Figure 3: Predictions of ANN models for the testing group
Slika 3: Napovedi ANN-modelov za preizku{eno skupino
effect on the shrinkage. The internal stress achieved with
the third-stage draw ratio increases the fiber shrinkage.
This is related to the absence of heat in the third stage of
drawing. A further crystallization after a critical TSDR,
along with an increase in the orientation of the amor-
phous region, led to a negligible change in the shrinkage
value. According to Figure 4b, the effects of the FSDR,
SSDR and TSDR on the specific work of rupture are the
same and the effect of the SGT is also very similar to
them. By increasing these parameters, a decrease in the
specific work of rupture can be achieved. But the BHT
has an opposite effect on this parameter and an increase
in the BHT gives rise to an increase in the specific work
of rupture. After a critical BHT, a reorientation begins
which is accompanied by a decrease in the orientation
and confirmed by a significant increase in the specific
work of rupture. Moreover, the FGT has no appreciable
effect on the specific work of rupture.
Based on Figure 4c and considering the tenacity
variation, the variables can be divided into effective/
non-effective categories; the FGT, BHT and SGT are
categorized as the non-effective group, while the FSDR,
SSDR, and TSDR are considered as the effective group.
By increasing the values of the effective group, the
tenacity increases, and with even higher values the tena-
city almost reaches the plateau. As shown in Figure 4d,
the initial modulus increases due to the increased FSDR,
SSDR, TSDR and SGT, while the FGT and BHT have no
visible effect on the initial modulus. It seems that the
temperature of the block heater mainly affects the crys-
talline structure of the fibers that restricts the chain
orientation in the third stage of drawing. This explains
why the initial modulus did not change significantly as
the BHT increased.
3.3 Important index analysis of the ANN model
To study the contributions of different input variables
to the ANN output model, various methods were intro-
duced, like the weights method29,30. A modified weights
method was introduced by Gevrey et al.31 and their
results are very close to the weights method. In this
study, the modified weights method was implemented to
determine the percentage contribution of each input vari-
able to the physical properties. Referring to the modified
weight method and Figure 1, the relative importance
(RI) is calculated through Equations (2) and (3):
Q
w
w
j
ij
ij
i
n=
=
∑
1
, j = 1, 2, ..., m (2)
RI
Q
Q
i
ij
j
m
ij
i
n
j
m= ×
=
==
∑
∑∑
1
11
100 (3)
It must be noted that only the first hidden layer is
considered in the modified weights method. The RI
values for all the input variables were obtained and
shown in Figure 5.
330 Materiali in tehnologije / Materials and technology 49 (2015) 3, 325–332
R. S. RAHBAR, M. VADOOD: PREDICTING THE PHYSICAL PROPERTIES OF DRAWN NYLON-6 FIBERS ...
Figure 4: Effects of the changing input variables on the mechanical properties: a) shrinkage, b) specific work of rupture, c) tenacity, d) initial
modulus. Since the variation of TSDR is higher than the ones of the other parameters, it is shown on the right vertical axis in Figure 4a.
Slika 4: Vpliv spremenjenih vhodnih spremenljivk na mehanske lastnosti: a) kr~enje, b) specifi~no delo pri pretrgu, c) trdnost, d) za~etni modul.
Ker je spreminjanje TSDR ve~je od drugih parametrov, je to prikazano na desni vertikalni osi na sliki 4a.
Figure 5 indicates that the main effect of the TSDR is
more pronounced on the shrinkage than the other factors
and that among the temperature variables, the BHT plays
the main role. In the case of the tenacity, although all the
draw ratios and temperatures influence this mechanical
property, the tenacity of drawn fibers is primarily
governed by the draw ratios. The initial modulus and the
specific work of rupture correlate with all the variables
except for the FGT. Among them, the second-stage
drawing variables have significant effects on the initial
modulus.
4 CONCLUSION
The results of this investigation suggest that an ANN
can be used quite effectively for a prediction of the
physical properties of drawn fibers. The ANN approach
was better at fitting the measured response in compa-
rison to the regression model. According to the obtained
results, it can be seen that, though only a limited number
of samples were available for the training and testing, the
overall prediction capabilities of the models were good
and would be very useful for optimizing the hot multi-
stage drawing process with regard to the time, quality
and price.
Among the variables, the first-godet temperature had
no significant effect on the physical properties. The
shrinkage was governed by the third-stage draw ratio and
the draw ratios were the variables affecting the tenacity.
The second-stage drawing variables played the main
roles in determining the initial modulus, while the spe-
cific work of rupture was governed by the draw ratios
and the second-godet temperature.
5 REFERENCES
1 D. R. Salem (Ed.), Structure formation in polymeric fibers, Hanser
Publishers, Munich 2000, 118–184
2 R. Yang, R. R. Mather, A. F. Fotheringham, Processing, structure,
and mechanical properties of as-spun polypropylene filaments – A
systematic approach using factorial design and statistical analysis,
Journal of Applied Polymer Science, 96 (2005) 1, 144–154, doi:10.
1002/app.21416
3 R. Yang, R. R. Mather, A. F. Fotheringham, The influence of pro-
cessing parameters on the structural and mechanical properties of
drawn polypropylene fibres: A factorial design approach, Journal of
Applied Polymer Science, 124 (2012) 5, 3606-3616, doi:10.1002/
app.34738
4 R. D. Yang, R. R. Mather, A. F. Fotheringham, The application of
factorial experimental design to the processing of polypropylene
fibres, Journal of Materials Science, 36 (2001) 13, 3097–3101,
doi:10.1023/A:1017909630132
5 B. Younes, A. Fotheringham, R. Mather, Statistical analysis of the
effect of multi-stage hot drawing on the thermal shrinkage and
crystallographic order of biodegradable aliphatic-aromatic co-poly-
ester fibres, Fibers and Polymers, 12 (2011) 6, 778–788,
doi:10.1007/ s12221-011-0778-9
6 F. Arain, A. Tanwari, T. Hussain, Z. Malik, Multiple response opti-
mization of rotor yarn for strength, unevenness, hairiness and imper-
fections, Fibers and Polymers, 13 (2012) 1, 118–122, doi:10.1007/
s12221- 012-0118-8
7 K. Yildirim, Y. Ulcay, O. Kopmaz, Forming Regression-based
Mathematical Models to Predict PET POY Yarn Properties in the
Case of Changing Production Parameters, Textile Research Journal,
80 (2010) 5, 411–421, doi:10.1177/0040517509346437
8 B. Younes, A. Fotheringham, H. M. El-Dessouky, G. Hadda, A Stati-
stical Analysis of the Influence of Multi-Stage Hot-Drawing on the
Overall Orientation of Biodegradable Aliphatic-Aromatic Co-Poly-
ester Fibers, Journal of Engineered Fibers and Fabrics, 8 (2013) 1,
6–16
9 M. D. Teli, R. Chakrabarti, Use of Statistical Methods to Understand
the Effect of Yarn and Fabric Parameters on Desizing Efficiency,
Fibres & Textiles in Eastern Europe, 16 (2008) 2, 95–100
10 L. J. Strumillo, D. Cyniak, J. Czekalski, T. Jackowski, Neural Model
of the Spinning Process Dedicated to Predicting Properties of
Cotton-Polyester Blended Yarns on the Basis of the Characteristics
of Feeding Streams, Fibres & Textiles in Eastern Europe, 16 (2008)
1, 28–36
11 A. Moghassem, A. Fallahpour, M. Shanbeh, An intelligent model to
predict breaking strength of rotor spun yarns using gene expression
programming, Journal of Engineered Fibers and Fabrics, 7 (2012) 2,
1–10
12 A. A. Gharehaghaji, M. Shanbeh, M. Palhang, Analysis of two
modeling methodologies for predicting the tensile properties of
cotton-covered nylon core yarns, Textile Research Journal, 77 (2007)
8, 565–571, doi:10.1177/0040517507078061
13 C. F. J. Kuo, K. I. Hsiao, Y. S. Wu, Using neural network theory to
predict the properties of melt spun fibers, Textile Research Journal,
74 (2004) 9, 840–843, doi:10.1177/004051750407400914
14 G. Allan, R. Yang, A. Fotheringham, R. Mather, Neural modelling of
polypropylene fibre processing: Predicting the structure and
properties and identifying the control parameters for specified fibres,
Journal of Materials Science, 36 (2001) 13, 3113–3118, doi:10.1023/
a:1017913731041
15 G. Yao, J. Guo, Y. Zhou, Predicting the warp breakage rate in
weaving by neural network techniques, Textile Research Journal, 75
(2005) 3, 274–278, doi:10.1177/004051750507500314
16 O. Balcý, R. Tuðrul Oðulata, Prediction of the changes on the
CIELab values of fabric after chemical finishing using artificial
neural network and linear regression models, Fibers Polym, 10
(2009) 3, 384–393, doi:10.1007/s12221-009-0384-2
17 A. Rawal, A. Majumdar, S. Anand, T. Shah, Predicting the properties
of needlepunched nonwovens using artificial neural network, Journal
of Applied Polymer Science, 112 (2009) 6, 3575–3581, doi:10.1002/
app.29687
18 M. M. Yazdi, D. Semnani, M. Sheikhzadeh, Moisture and heat trans-
fer in hybrid weft knitted fabric with artificial intelligence, Journal of
Applied Polymer Science, 114 (2009) 3, 1731–1737, doi:10.1002/
app.30685
Materiali in tehnologije / Materials and technology 49 (2015) 3, 325–332 331
R. S. RAHBAR, M. VADOOD: PREDICTING THE PHYSICAL PROPERTIES OF DRAWN NYLON-6 FIBERS ...
Figure 5: RI for all the input variables versus the physical properties
Slika 5: RI vseh vhodnih spremenljivk proti fizikalnim lastnostim
19 M. Vadood, D. Semnani, M. Morshed, Optimization of acrylic dry
spinning production line by using artificial neural network and
genetic algorithm, Journal of Applied Polymer Science, 120 (2011)
2, 735–744, doi:10.1002/app.33252
20 O. N. Hung, L. J. Song, C. K. Chan, C. W. Kan, C. W. M. Yuen,
Using artificial neural network to predict colour properties of
laser-treated 100% cotton fabric, Fibers Polym, 12 (2011) 8,
1069–1076, doi:10.1007/s12221-011-1069-1
21 D. Semnani, M. Vadood, Improvement of intelligent methods for
evaluating the apparent quality of knitted fabrics, Engineering
Applications of Artificial Intelligence, 23 (2010) 2, 217–221,
doi:10.1016/j.engappai.2009.08.003
22 P. Unal, M. Üreyen, D. Mecit, Predicting properties of single jersey
fabrics using regression and artificial neural network models, Fibers
Polym, 13 (2012) 1, 87–95, doi:10.1007/s12221-012-0087-y
23 K. Nasouri, H. Bahrambeygi, A. Rabbi, A. M. Shoushtari, A. Kaflou,
Modeling and optimization of electrospun PAN nanofiber diameter
using response surface methodology and artificial neural networks,
Journal of Applied Polymer Science, 126 (2012) 1, 127–135,
doi:10.1002/app.36726
24 T. Chen, C. Zhang, L. Li, X. Chen, Simulating the drawing of
spunbonding nonwoven process using an artificial neural network
technique, Journal of The Textile Institute, 99 (2008) 5, 479–488,
doi:10.1080/00405000701608631
25 P. Soltani, M. Vadood, M. Johari, Modeling spun yarns migratory
properties using artificial neural network, Fibers Polym, 13 (2012) 9,
1190–1195, doi:10.1007/s12221-012-1190-9
26 R. Semnani Rahbar, M. R. M. Mojtahedi, Influence of hot multistage
drawing on structure and mechanical properties of nylon 6 multi-
filament yarn, Journal of Engineered Fibers and Fabrics, 6 (2011) 2,
7–15
27 L. Fausett, Fundamentals of neural networks, Prentice Hall, New
Jersey 1994
28 I. Beli~, Neural Networks and Static Modelling, In: M. ElHefnawi,
M. Mysara (Eds.), Recurrent Neural Networks and Soft Computing,
InTech [online]. 2012, 1, 1–21. Available from:
http://www.intechopen.com/books/recurrent-neural-networks-and-so
ft-computing/neural-networks-and-static-modelling
29 A. T. C. Goh, Back-propagation neural networks for modeling com-
plex systems, Artificial Intelligence in Engineering, 9 (1995) 3,
143–151, doi:10.1016/0954-1810(94)00011-S
30 G. D. Garson, Interpreting neural network connection weights, Arti-
ficial Intelligence Expert, 6 (1991) 4, 46–51
31 M. Gevrey, I. Dimopoulos, S. Lek, Review and comparison of me-
thods to study the contribution of variables in artificial neural net-
work models, Ecological Modelling, 160 (2003) 3, 249–264,
doi:10.1016/S0304-3800(02)00257-0
332 Materiali in tehnologije / Materials and technology 49 (2015) 3, 325–332
R. S. RAHBAR, M. VADOOD: PREDICTING THE PHYSICAL PROPERTIES OF DRAWN NYLON-6 FIBERS ...
M. HNÍZDIL et al.: INFLUENCE OF THE IMPACT ANGLE AND PRESSURE ...
INFLUENCE OF THE IMPACT ANGLE AND PRESSURE ON THE
SPRAY COOLING OF VERTICALLY MOVING HOT STEEL
SURFACES
VPLIV VPADNEGA KOTA IN TLAKA NA OHLAJANJE Z
BRIZGANJEM NA VERTIKALNO PREMIKAJO^E SE VRO^E
POVR[INE JEKLA
Milan Hnízdil, Martin Chabi~ovský, Miroslav Raudenský
Heat Transfer and Fluid Flow Laboratory, Brno University of Technology, Technicka 2896/2, 616 69 Brno, Czech Republic, European Union
hnizdil@fme.vutbr.cz
Prejem rokopisa – received: 2013-10-09; sprejem za objavo – accepted for publication: 2014-09-04
doi:10.17222/mit.2013.239
The cooling of vertically moving strips is used very often to obtain the required material properties. Water spray cooling has to
be used when a high cooling intensity is needed.
Our Heat Transfer and Fluid Flow Laboratory is equipped with a testing device which allows vertical movement of a heated
experimental plate (sheet). Two different sizes of flat-jet nozzles were tested with different water pressures and angles of the
water impact (inclination angles of the spraying bar). The water-pressure range was between 2 bar and 9.3 bar and the angle of
the water impact changed from 20 ° to 40 °.
The dependence of the heat-transfer coefficient on the surface temperature was evaluated for each experiment. Interesting
results were obtained from the comparison of these experimental results, showing that the heat-transfer coefficient and the
Leidenfrost temperature increase with the increasing water pressure. Very interesting results were obtained during the tests with
different inclination angles. The highest heat-transfer coefficient was obtained for the angle of 20 ° and the lowest value of the
heat-transfer coefficient was obtained for the angle of 40 ° at the surface temperatures of around 200 °C.
Keywords: spray cooling, flat-jet nozzles, impact angle, water impingement density, Leidenfrost temperature
Ohlajanje vertikalno premikajo~ih se trakov se pogosto uporablja za zagotovitev zahtevanih lastnosti materiala. Kadar je
potrebna velika intenzivnost hlajenja, se uporablja ohlajanje z brizganjem vode.
Laboratorij za prehajanje toplote in tok fluidov je opremljen s preizkusno napravo, ki omogo~a vertikalno premikanje eksperi-
mentalne plo{~e (jeklo). Preizku{eni sta bili dve razli~ni dimenziji {ob pri razli~nih tlakih vode in razli~nih kotih pr{enja vode
(nagibni kot palice za brizganje). Obmo~je tlaka vode je bilo med 2 bar in 9,3 bar, kot vodnega curka pa med 20 ° in 40 °.
Za vsak poskus je bila ocenjena odvisnost koeficienta prehajanja toplote od temperature povr{ine. Dobljeni so bili zanimivi
rezultati iz primerjave eksperimentalnih podatkov, ki ka`ejo, da koeficient prehajanja toplote in Leidenfrostova temperatura
nara{~ata z ve~anjem tlaka vode. Zanimive rezultate smo dobili tudi pri poskusih z razli~nimi vpadnimi koti. Najvi{ji koeficient
prehajanja toplote je bil dose`en pri kotu 20 °, najni`ja vrednost koeficienta prehajanja toplote pa je bila dose`ena pri kotu 40 °
pri temperaturi povr{ine okrog 200 °C.
Klju~ne besede: hlajenje z brizganjem, {obe s plo{~atim curkom, vpadni kot curka, gostota udarca vode, Leidenfrostova tem-
peratura
1 INTRODUCTION
Vertically moving strips are cooled by spraying water
when a high cooling intensity (the heat-transfer coeffi-
cient) is needed. The cooling of these strips has to be
homogeneous to avoid thermal deformation. Falling
reflected water has a significant influence on the cooling
homogeneity1. It can cause local thermal stresses in a
strip and its deformation. The heat-transfer coefficient is
influenced by many parameters such as water pressure,
nozzle distance, water impingement density, water tem-
perature, etc.2–4
A special cooling system which sprays water at a pre-
scribed inclination angle to avoid the falling of the
reflected water (Figure 1) was tested and the influence
of the impact angle and the pressure on the heat-transfer
coefficient was investigated during the cooling of a stain-
less-steel sheet at the initial temperature of 900 °C.
Materiali in tehnologije / Materials and technology 49 (2015) 3, 333–336 333
UDK 536.2:536.28 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)333(2015)
Figure 1: Cooling system (inclination angle of 30 °, distance to the
surface is 100 mm)
Slika 1: Sistem ohlajanja (vpadni kot 30 °, razdalja do povr{ine 100
mm)
2 LABORATORY STAND, EXPERIMENTAL
PROCESS AND CONFIGURATION
An experimental apparatus developed for the vertical
moving of a hot stainless-steel test sheet was used in the
experiments (Figure 2). An electrical furnace which
heated the experimental sheet to its initial temperature
was positioned in the upper part of the testing bench.
This sheet was made of thick austenitic stainless steel 1.5
mm. Five thermocouples were welded to the rear side of
the sheet and they were positioned in the center line, per-
pendicular to the direction of the movement of the sheet.
The test sheet was fixed on the trolley which moved
repeatedly through the water spray. The temperatures
were recorded by a data logger with a frequency of 320
Hz for each thermocouple.
The experiments started by heating the sheet in an
inert atmosphere to the initial temperature (900 °C).
Then the required water pressure was set. The data
logger was switched on and it started to record the tem-
peratures. Then the test sheet repeatedly moved through
the spray till the surface temperature was less than 200
°C.
The recorded data were transferred to a computer and
the inverse task5 was used to compute the surface tem-
perature, the heat flux and the heat-transfer coefficient,
which is usually used as a boundary condition6.
The influence of the water pressure was investigated.
Five tests, each with a different water pressure (2, 4, 6, 8
and 9.3) bar, were conducted. Commercially available
flat-jet nozzles were fixed in a single row with the same
offset angle. The distance between the nozzles was 55
mm. The distance of the nozzle orifices to the test sheet
was 100 mm and the inclination angle was 20 ° during
these tests.
Further, three tests, each with a different inclination
angle (20 °, 30 ° and 40 ° from the vertical direction,
Figure 1) were conducted to investigate the influence of
the inclination angle. These nozzles had a larger size of
the orifice than the nozzles used for the tests with diffe-
rent water pressures. The distance between the nozzles
was 100 mm. The distance of the nozzle orifices to the
test sheet was 100 mm and the water pressure was 4 bar
during these experiments. All the experiments were con-
ducted with the velocity of the movement of 3 m/s.
3 RESULTS
The results shown in this part are the average values
of the heat-transfer coefficient in the impact area com-
puted for the positions of 0 mm to 600 mm (Figure 3).
The measured dependences of the heat-transfer coeffi-
cient on the surface temperature are shown for different
water pressures (Figure 4a) and for different inclination
angles (Figure 4b).
It is evident that the heat-transfer coefficient in-
creases with the increasing water pressure (an increase in
the water pressure also causes an increase in the water
impingement density m L ). Also, the Leidenfrost tempe-
rature (TLeid (°C)), which is defined as the temperature at
which the heat flux reaches its minimum (or the
temperature, at which the film-boiling regime changes to
the transition-boiling regime), increases with the in-
crease in the pressure (Figure 5). The measured Leiden-
frost temperatures are higher than the ones predicted
M. HNÍZDIL et al.: INFLUENCE OF THE IMPACT ANGLE AND PRESSURE ...
334 Materiali in tehnologije / Materials and technology 49 (2015) 3, 333–336
Figure 2: Testing bench for experimental tests of cooling vertically
moving strips (1 – collector with nozzles; 2 – pressure gauge; 3 – test
plate; 4 – motor with rope; 5 – girder; 6 – trolley with position sensor
and data logger for recording temperatures and position of the test
plate; 7 – heater; 8 – water tank; 9 – pump; 10 – control valve)
Slika 2: Preizkusna klop za preizku{anje ohlajanja vertikalno
premikajo~ih se trakov (1 – nosilec s {obami; 2 – merilnik tlaka; 3 –
preizkusna plo{~a; 4 – motor z vrvjo, 5 – nosilec; 6 – vozi~ek s
senzorjem pozicije in "data logger" za registracijo temperature ter
pozicije preizkusne plo{~e; 7 – grelnik; 8 – vodni rezervoar; 9 –
~rpalka; 10 – kontrolni ventil)
Figure 3: Dependence of the HTC on the position, for the surface
temperature of 800 °C – evaluating area of 0–600 mm
Slika 3: Odvisnost HTC od polo`aja pri temperaturi povr{ine 800 °C
– podro~je ocenjevanja 0–600 mm
with the common prediction equations that use the water
impingement density7,8:
T mLeid L= 339604
0 19.  . (1)
T
m d
Leid
L=1400
2
32
0 13

.

(2)
where m L is the water impingement density (kg m
–2 s–1),
d32 is the Sauter mean diameter (m) of water droplets,
 is the density (kg m–3) and  is the surface tension
(N m–1).
The comparison of the measured Leidenfrost tempe-
ratures and the ones predicted with Eqs. (2) and (3) is
shown in Figure 5. The error of prediction with Eqs. (2)
and (3) is growing with the increase in the pressure. This
shows that the water pressure must be taken into account
when an accurate prediction of the Leidenfrost tempe-
rature is necessary. The increase in the pressure increases
the exit velocity of the water from a nozzle and this leads
to destroying the vapour layer at higher surface tempe-
ratures. The new prediction of the Leidenfrost tempera-
ture based on Eq. (1) that takes the pressure into account
is:
T m pLeid = 339604
0 19. (  ) .L (3)
where m L is the water impingement density (kg m
–2 s–1)
and p is the water pressure (bar).
The comparison of the new prediction with the
measured data and the other predictions is shown in
Figure 5.
The second part of the experiments was focused on
the influence of the water impact angle on the cooling
intensity. As it is shown in Figure 4b, the heat-transfer
coefficient and the Leidenfrost temperature are influ-
enced by the change in the impact angle. The Leidenfrost
temperature is the highest for the impact angle of 30 °
(800 °C) and the lowest for the impact angle of 40 °
(540 °C). The Leidenfrost temperature for the impact
angle of 20 ° is 610 °C. The heat-transfer coefficient in
the area of stable film boiling (the surface temperatures
are higher than the Leidenfrost temperature – higher than
800 °C) is almost the same for all the impact angles. The
main effect of the inclination angle on the heat-transfer
coefficient occurs at low surface temperatures (approxi-
mately 200 °C), where the heat-transfer coefficient is the
highest at the angle of 20 ° and the lowest at 40 °.
4 CONCLUSION
Two kinds of tests were done to find the influence of
the water pressure (the flow rate) and the impact angle
on the heat-transfer coefficient. It is evident (from the
results) that the cooling intensity increases with the
increasing water-flow rate (the water pressure). Also, the
Leidenfrost temperature is dependent on the water pres-
sure – a higher water pressure means a higher leidenfrost
point. Different spray impact angles cause changes in the
heat-transfer coefficient at low surface temperatures and
the Leidenfrost temperature also depends on the impact
angle. A new correlation for the prediction of the leiden-
frost temperature was found.
M. HNÍZDIL et al.: INFLUENCE OF THE IMPACT ANGLE AND PRESSURE ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 333–336 335
Figure 5: Comparison of the measured Leidenfrost temperature with
the existing predictions and the new modified prediction of the
Leidenfrost temperature (corresponding water impingement densities
for water pressures of (2, 4, 6, 8 and 9.3) bar are (3.18, 4.5, 5.51, 6.36
and 6.86) kg m–2 s–1)
Slika 5: Primerjava izmerjene Leidenfrostove temperature z obstoje-
~im predvidevanjem in novo modificirano predvidevanje Leiden-
frostove temperature (ustrezne gostote vodnega curka pri tlakih vode
(2, 4, 6, 8 in 9,3) bar so (3,18, 4,5, 5,51, 6,36 in 6,86) kg m–2 s–1)
Figure 4: Dependence of the heat-transfer coefficient on the surface
temperature for the: a) increasing water pressure and b) increasing
inclination angle
Slika 4: Odvisnost koeficienta prehoda toplote od temperature
povr{ine pri: a) nara{~ajo~em tlaku vode in b) nara{~ajo~im vpadnim
kotom
Acknowledgement
This work is an output of research and scientific
activities of the project LO1202 with financial support of
the MEYS under the programme NPU I and the internal
grant of the Brno University of Technology focused on
specific research and development No. FSI-S-14-2437
and by the project No. CZ.1.07/2.3.00/20.0188,
HEATEAM - Multidisciplinary Team for Research and
Development of Heat Processes.
5 REFERENCES
1 M. Chabicovsky, M. Raudenský, Experimental investigation of spray
cooling of horizontally and vertically oriented surfaces, Conference
proceedings of 22nd Conference on metallurgy and materials 1,
Ostrava, 2013, 102–107
2 M. Chabicovsky, M. Raudenský, Experimental Investigation of the
Heat Transfer Coeficient, Mater. Tehnol., 47 (2013) 3, 395–398
3 M. Raudenský, M. Hnízdil, P. Kotrbá~ek, Why oxides intensify spray
cooling, 30th International Steel Industry Conference, Paris, 2012,
92–337
4 M. Hnízdil, M. Raudenský, Influence of Water Temperature on the
Cooling Intensity during Continuous Casting and Hot Rolling,
Conference METAL 2012, Brno, 2012, 1–6
5 M. Raudenský, Heat Transfer Coefficient Estimation by Inverse
Conduction Algorithm, International Journal of Numerical Methods
for Heat and Fluid Flow, 3 (1993) 3, 257–266, doi:10.1108/eb017530
6 J. Stetina, F. Kavicka, T. Mauder, Heat Transfer Coefficients Beneath
the Water Cooling Nozzles of a Billet Caster, 18th International Con-
ference on Metallurgy and Materials, Hradec Nad Moravici, 2009
7 H. Muller, R. Jeschars, Wärmeübergang bei der Spritzwasserkühlung
von Nichteisenmetallen, Metallkunde, 1983, 257–264
8 S. Yao, T. L. Cox, A general heat transfer correlation for impacting
water sprays on high-temperature surfaces, Experimental Heat Trans-
fer, 15 (2002) 4, 207–219, doi:10.1080/08916150290082649
M. HNÍZDIL et al.: INFLUENCE OF THE IMPACT ANGLE AND PRESSURE ...
336 Materiali in tehnologije / Materials and technology 49 (2015) 3, 333–336
M. CHABI^OVSKÝ, M. RAUDENSKÝ: TECHNIQUES OF MEASURING SPRAY-COOLING HOMOGENEITY
TECHNIQUES OF MEASURING SPRAY-COOLING
HOMOGENEITY
TEHNIKE MERJENJA HOMOGENOSTI HLAJENJA
Z BRIZGANJEM
Martin Chabi~ovský, Miroslav Raudenský
Heat Transfer and Fluid Flow Laboratory, Faculty of Mechanical Engineering, Brno University of Technology, Technická 2, 616 69 Brno,
Czech Republic
chabicovsky@LPTaP.fme.vutbr.cz
Prejem rokopisa – received: 2013-10-18; sprejem za objavo – accepted for publication: 2014-07-14
doi:10.17222/mit.2013.253
The cooling homogeneity is one of the most important factors that must be considered in the design of cooling sections for the
hot rolling of thin sheets. Inhomogeneous cooling can lead to undesirable thermal distortion. The cooling homogeneity is
mainly influenced by the water distribution of the cooling section. And so, one way to measure the cooling homogeneity is to
measure the impact-pressure distribution of the cooling section. Another way is to measure the surface-temperature distribution
of a steel sample during the cooling process. There are two ways to measure the surface temperature and the temperature field:
the contact and non-contact measurements. The contact measurement can be performed with thermocouples and the non-contact
method is an optical measurement like the one using an infrared scanner. Each of these methods has their advantages and
disadvantages. Their comparison was made during an experimental measurement of the cooling of a stainless steel sheet using
full-cone water nozzles and a special linear pneumatic sprayer.
Keywords: cooling homogeneity, full-cone nozzles, linear pneumatic sprayer, heat-transfer coefficient
Homogenost ohlajanja je med najpomembnej{imi dejavniki, ki ga je treba upo{tevati pri na~rtovanju ohlajevalnega podro~ja pri
vro~em valjanju tankih plo~evin. Nehomogeno ohlajanje lahko povzro~i ne`eleno toplotno izkrivljanje. Homogenost ohlajanja
je predvsem odvisna od razporeditve vode v obmo~ju ohlajanja. Ena od mo`nosti merjenja homogenosti ohlajanja je merjenje
razporeditve tlaka udarca v obmo~ju ohlajanja. Druga mo`nost je merjenje razporeditve temperature na povr{ini vzorca iz jekla
med postopkom ohlajanja. Obstajata dva na~ina za merjenje temperature povr{ine in temperaturnega polja: kontaktno in
nekontaktno. Kontaktno merjenje se lahko izvr{i s termoelementi, nekontaktno pa z opti~no meritvijo, kot je na primer
infrarde~e vrsti~no tipalo. Vsaka od teh metod ima svoje prednosti in slabosti. Izvr{ena je bila eksperimentalna primerjava
merjenja ohlajanja plo~evine iz nerjavnega jekla v obmo~ju vodnih {ob in s posebnim linearnim pnevmatskim brizgalnikom.
Klju~ne besede: homogenost ohlajanja, obmo~je s {obami, linearni pnevmatski brizgalnik, koeficient prehoda toplote
1 INTRODUCTION
The intensive cooling of thin sheets during hot rolling
or heat treatment is mainly conducted with water. Cool-
ing with water provides high cooling rates compared to
gas cooling, but these can be associated with undesirable
problems of the homogeneity of cooling. Inhomoge-
neous cooling can lead to non-homogeneous material
properties and also to thermal distortion. The cooling
homogeneity is one of the most important factors that
must be considered in the design of water-cooling
sections for the hot rolling or the heat treatment of thin
sheets.
Cooling hot surfaces with water is associated with
different regimes of boiling. The film-boiling regime
occurs when the surface temperature is higher than the
Leidenfrost temperature1 (Figure 1). The surface is
covered with a vapor layer which protects the surface
and lowers the cooling intensity during the film-boiling
regime. When the surface temperature drops below the
Leidenfrost temperature, the vapor layer is broken and
the transient and nucleate boiling regime occurs, which
is characterized by high cooling intensities. A large inho-
mogeneity of the surface-temperature distribution can
mostly be observed when the temperature of one part of
the cooled surface is below the Leidenfrost temperature
and another is above the Leidenfrost temperature. This is
well demonstrated in Figure 1. The temperature record
at time T1 corresponds to the heat-transfer coefficient
HTC T1. This heat-transfer coefficient has a Leidenfrost
Materiali in tehnologije / Materials and technology 49 (2015) 3, 337–341 337
UDK 621.77:536.28 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)337(2015)
Figure 1: Simulation of the cooling process
Slika 1: Simulacija postopka ohlajanja
temperature of approximately 500 °C. And so the
intensive cooling caused by the breaking up of the vapor
layer starts at a time of 44 s. The distribution of the
heat-transfer coefficient HTC T2 has a Leidenfrost
temperature of 450 °C. This temperature relates to the
start of a rapid decrease in temperature T2 at a time of 50
s. This is why the difference in the temperature is almost
300 °C between 50 s and 60 s although the difference
between the heat-transfer coefficients is small, especially
for the surface temperatures higher than 500 °C.
The best way of measuring the cooling inhomoge-
neity during the cooling of thin sheets is to do it in the
plant. However, because this method is very expensive,
laboratory experimental methods simulating the plant
conditions should be taken into account. With the
reported technique we simulated a cooling process
involving a small sheet, and measured the distribution of
surface temperatures during the cooling process either
with the contact method using thermocouples or the
non-contact method using an infrared scanner. As the
cooling intensity is mainly influenced by the water distri-
bution and impact velocity2,3 the cooling homogeneity
can be evaluated by measuring the impact-pressure
distribution of the cooling section. The experimental
methods for the measurement of the cooling inhomo-
geneity are described below.
2 EXPERIMENT
2.1 Measurement of spray-cooling homogeneity
An experimental study was performed to compare
three different methods of measuring the cooling homo-
geneity. The homogeneity of cooling was measured
during the cooling process with thermocouples and an
infrared scanner. The cooling homogeneity was also
investigated during cold tests when the impact pressure
was measured.
The experiments for this comparison were conducted
with three different nozzle configurations, which pro-
vided three different levels of cooling homogeneity.
The configuration with a very good cooling homoge-
neity (Configuration 1) was composed of a row of full-
cone nozzles. The distance between the full-cone nozzles
was 80 mm and the distance from the nozzle orifices to
the test sheet was 250 mm. The water pressure was 130
kPa. The spray angle was 60 ° and the water impinge-
ment density in the impact area was 12.07 kg m–2 s–1.
Configuration 2 was a configuration with a medium
cooling homogeneity and Configuration 3 provided very
inhomogeneous cooling. Configurations 2 and 3 were
composed of a special linear pneumatic spray box which
is well described in4.
The distance between the nozzle orifice and the test
sheet was 250 mm and the spray angle was 60 °. The
water pressure was 130 kPa for Configuration 2 and 600
kPa for Configuration 3. The air pressure was 8.5 kPa for
Configuration 2 and 10.5 kPa for Configuration 3.
The water impingement density was 9.76 kg m–2 s–1
for Configuration 2 and 20.64 kg m–2 s–1 for Configura-
tion 3.
One way to measure the temperature is using the
thermocouples welded onto the rear side of the test sheet
which, with the help of inverse methods, determines the
surface temperatures. The other method uses an infrared
scanner for a direct measurement of the surface tempera-
tures during the cooling process. The apparatus for
measuring the temperatures using an infrared scanner
and thermocouples is shown in Figure 2. Details of the
experimental apparatus and the experimental set-up are
described in5,6. Hot tests were conducted with the
austenitic stainless-steel sheets with a thickness of 2 mm,
a height of 300 mm and a width of 320 mm. A test sheet
was heated in a furnace with a protective atmosphere at a
temperature of 950 °C. After the heating, the test sheet
moved down with a velocity of 3 m s–1 and passed
through the cooling section. Then it moved up through
the cooling section and stopped for 1.5 s above the
cooling section. During this time the infrared scanner
measured the surface temperature. Then the test sheet
moved down and repeated the previous steps until it was
cooled to a temperature of 200 °C. During the experi-
ment, the temperature of the sheet was measured with
five type-K thermocouples with a frequency of 320 Hz.
The thermocouples were welded onto the rear side of the
test sheet. One thermocouple was welded in the center of
the test sheet and the distance between the thermo-
couples was 20 mm (for more details see reference5).
The distance between the infrared scanner and test sheet
was 1500 mm. The frequency of the measurement was
36 Hz and the resolution was 1024 points per line.
M. CHABI^OVSKÝ, M. RAUDENSKÝ: TECHNIQUES OF MEASURING SPRAY-COOLING HOMOGENEITY
338 Materiali in tehnologije / Materials and technology 49 (2015) 3, 337–341
Figure 2: Experimental apparatus for measuring cooling homogeneity
with thermocouples and an infrared scanner
Slika 2: Eksperimentalna naprava za merjenje homogenosti ohlajanja
s termoelementi in infrarde~im vrsti~nim tipalom
Impact-pressure tests were conducted with a pressure
sensor with a diameter of 12 mm and during the test the
impact pressure was measured on the whole area that
was influenced by the spraying nozzles, with a step of 10
mm in the direction of the X-axis and 20 mm in the
direction of the Y-axis.
3 RESULTS AND DISCUSSION
3.1 Thermocouples
The use of thermocouples for measuring the cooling
homogeneity has some restrictions. Thermocouples can
be placed only in a finite number of positions. So, it is
necessary to know the positions, in which inhomogeneity
can be expected. The water distribution can serve as the
first source of information about the expected cooling
homogeneity. An example of the computed water distri-
bution is shown in Figure 3. The information about the
water distribution is necessary when deciding where to
place the thermocouples. The thermocouples were
installed in positions of 0 mm and 40 mm (Figure 3) for
Configuration 1. The position of 40 mm corresponds to
the position with the highest water impingement density
and the position of 0 mm is the position with the lowest
water impingement density. The measured temperatures
are shown in Figure 4. The temperatures are almost the
same for the positions of 0 mm and 40 mm. The surface
temperatures and the heat-transfer coefficients can be
computed at these positions using inverse methods.7–9
The computed dependence of the heat-transfer coeffi-
cient on the surface temperature is shown in Figure 5.
The dependence of the heat-transfer coefficient on the
surface temperature also shows a very good homogeneity
of the cooling, obviously for the surface temperatures
higher than 600 °C.
3.2 Infrared scanner
Although the measurement with an infrared scanner
provides information about the surface temperature of
the overall surface area, and these measured tempera-
tures can be directly used for the evaluation of the cool-
ing homogeneity, this method of measuring the cooling
homogeneity needs a more complicated experimental
set-up. The precision of the non-contact measurement is
not as good as in the case of thermocouples. There are
several factors that have a considerable influence on the
quality of measurement. One of them is emissivity.
Information about emissivity is necessary for obtaining
surface temperatures. Emissivity depends on the proper-
ties of the surface (material, temperature, presence of
oxides and others) and the surface properties change dur-
ing the cooling process. If a sheet is heated in a protec-
tive atmosphere and removed from the furnace, its sur-
face is like a mirror and the emissivity is lower than
during the cooling process due to different surface tem-
peratures and the presence of scales on the surface. The
scales on the surface enhance the emissivity. Uneven
M. CHABI^OVSKÝ, M. RAUDENSKÝ: TECHNIQUES OF MEASURING SPRAY-COOLING HOMOGENEITY
Materiali in tehnologije / Materials and technology 49 (2015) 3, 337–341 339
Figure 4: Measured temperatures for Configuration 1
Slika 4: Izmerjene temperature pri postavitvi 1
Figure 3: Water distribution for Configuration 1
Slika 3: Razporeditev vode pri postavitvi 1
Figure 6: Surface temperatures at given times for Configuration 2
Slika 6: Temperature povr{ine pri danih ~asih pri postavitvi 2
Figure 5: Heat-transfer coefficient for Configuration 1
Slika 5: Koeficient prehoda toplote pri postavitvi 1
scale coverage of the surface is a problem and can lead
to errors in the measurement of the surface temperatures
and an incorrect conclusion regarding the cooling
homogeneity.
The temperature that is measured can also be influ-
enced by the presence of water droplets and vapor in the
air between the sheet and the infrared scanner. The water
remaining on the sheet and the radiation from the
surroundings also prevent the measured temperatures
from reflecting the reality. The presence of vapor and
water droplets in the air or a water layer on the surface
can smooth a low temperature inhomogeneity or cause
the measured temperatures to be lower or higher than the
real temperatures. The radiation from the surroundings
can cause the temperatures measured in some positions
to be higher than in reality. All the above factors should
be considered during the measurement preparation to
obtain realistic results. An example of the temperatures
measured with the infrared scanner is shown in Figure 6.
The position of 0 mm was in the center of the test sheet
with a width of 320 mm and the temperatures were
measured in the centerline of the test sheet. The
temperature was also measured with five thermocouples
during the cooling process and the comparison is shown
in Figure 7. The thermocouples were welded on at the
positions of (–40, –20, 0, 20 and 40) mm. The surface
temperatures that were measured with the thermocouples
and computed with an inverse task were slightly lower
than the temperatures measured with the infrared
scanner. There is only one big difference between the
temperatures measured with the infrared scanner and
those measured with the thermocouples. It is found in the
position of 40 mm at a time of 52 s. The difference is
approximately 120 °C. It can be explained with the water
layer on the surface in this position. The homogeneity of
cooling is almost the same with both methods.
3.3 Impact pressure
The measurement of the impact pressure is an
indirect method for measuring the cooling homogeneity
because the real surface temperature is not measured
during the cooling process. An example of the measured
impact-pressure distribution is shown in Figure 8. Its
average value over the width of the spray configuration
(Y-axis) of Configuration 2 is compared with the surface
temperatures measured for this configuration with the
infrared scanner in Figure 9.
Although the impact-pressure distribution looks
inhomogeneous, the surface temperatures at the times of
12 s and 25 s look homogeneous except at the position of
80 mm. This is why the surface temperatures are higher
than the Leidenfrost temperature and the vapor layer
protects the surface and only a significant inhomogeneity
of the impact pressure is observed during the real
cooling. This was confirmed during the experiment with
Configuration 3 (Figure 10). The surface temperatures
M. CHABI^OVSKÝ, M. RAUDENSKÝ: TECHNIQUES OF MEASURING SPRAY-COOLING HOMOGENEITY
340 Materiali in tehnologije / Materials and technology 49 (2015) 3, 337–341
Figure 9: Comparison of the impact pressure with the measured sur-
face temperatures for Configuration 2
Slika 9: Primerjava dinami~nega tlaka pri izmerjeni temperaturi
povr{ine pri postavitvi 2
Figure 8: Impact-pressure distribution for Configuration 2
Slika 8: Razporeditev dinami~nega tlaka pri postavitvi 2
Figure 7: Comparison of the surface temperatures measured with the
infrared scanner (lines) and with the thermocouples (marks: , , 
and )
Slika 7: Primerjava temperature na povr{ini, izmerjene z infrarde~im
vrsti~nim tipalom (polna ~rta) in izmerjene s termoelementi (oznake:
, ,  in )
Figure 10: Comparison of the impact pressure with the surface-tem-
perature distribution for Configuration 3
Slika 10: Primerjava dinami~nega tlaka z razporeditvijo temperature
pri postavitvi 3
after the second pass through the cooling section (pass 2)
fit well with the impact-pressure distribution (Imp. 2)
only at the position of 80 mm. When the surface tempe-
rature is lower than the Leidenfrost temperature (pass 9)
almost every increase or decrease in the impact pressure
is reflected in the surface-temperature distribution.
4 CONCLUSIONS
The thermocouples provided accurate and reliable
information about the cooling homogeneity, though only
in a finite number of positions. The benefit of the
infrared scanner was that the cooling homogeneity was
measured on all the surface area; but to obtain accurate
values of the measurement, a demanding measurement
preparation was necessary. The homogeneity of the
impact-pressure distribution corresponded with the
homogeneity of the measured surface temperatures only
for the surface temperatures below the Leidenfrost
temperature. During the test with a hot test plate, a large
inhomogeneity of the impact pressure was observed only
for the surface temperatures higher than the Leidenfrost
temperature.
Acknowledgement
This work is an output of research and scientific
activities of this project LO1202 with financial support
of the MEYS under the programme NPU I and the
internal grant of the Brno University of Technology
focused on specific research and development No.
FSI-S-14-2437.
5 REFERENCES
1 S. Nukiyama, The Maximum and Minimum Values of the Heat Q
Transmitted from Metal to Boiling Water Under Atmospheric
Pressure, International Journal of Heat and Mass Transfer, 9 (1966)
12, 1419–1433, doi:10.1016/0017-9310(66)90138-4
2 J. Wendelstorf, K. H. Spitzer, R. Wendelstorf, Spray water cooling
heat transfer at high temperatures and liquid mass fluxes, Interna-
tional Journal of Heat and Mass Transfer, 51 (2008) 19–20,
4902–4910, doi:10.1016/j.ijheatmasstransfer.2008.01.032
3 S. C. Yao, T. L. Cox, A general heat transfer correlation for impac-
ting water sprays on high-temperature surfaces, Experimental Heat
Transfer, 15 (2002) 4, 207–219, doi:10.1080/08916150290082649
4 J. Rivallin, S. Viannay, General principles of controlled water cool-
ing for metallurgical on-line hot rolling processes: forced flow and
sprayed surfaces with film boiling regime and rewetting phenomena,
International Journal of Thermal Sciences, 40 (2001) 3, 263–272,
doi:10.1016/S1290-0729(00)01216-3
5 M. Chabicovsky, M. Raudensky, Experimental Investigation of a
Heat Transfer Coefficient, Mater. Tehnol., 47 (2013) 3, 395–398
6 M. Chabi~ovský, M. Raudenský, Experimental investigation of spray
cooling of horizontally and vertically oriented surfaces, Proc. of the
22nd International Conference on Metallurgy and Materials METAL
2013, Brno, 2013, 198–202
7 M. Raudensky, Heat Transfer Coefficient Estimation by Inverse
Conduction Algorithm, International Journal of Numerical Methods
for Heat and Fluid Flow, 3 (1993) 3, 257–266, doi:10.1108/eb017530
8 M. Pohanka, K. A. Woodbury, A Downhill Simplex method for
computation of interfacial heat transfer coefficients in alloy casting,
Inverse Problems in Engineering, 11 (2003) 5, 409–424,
doi:10.1080/1068276031000109899
9 M. Pohanka, P. Kotrbá~ek, Design of Cooling Units for Heat Treat-
ment, In: F. Czerwinski (ed.), Heat Treatment – Conventional and
Novel Applications, chapter 1, InTech, 2012, 1–20, doi:10.5772/
50492
M. CHABI^OVSKÝ, M. RAUDENSKÝ: TECHNIQUES OF MEASURING SPRAY-COOLING HOMOGENEITY
Materiali in tehnologije / Materials and technology 49 (2015) 3, 337–341 341

S. KRAMAR et al.: MINERALOGICAL AND GEOCHEMICAL CHARACTERIZATION OF ROMAN SLAG ...
MINERALOGICAL AND GEOCHEMICAL CHARACTERIZATION
OF ROMAN SLAG FROM THE ARCHAEOLOGICAL SITE NEAR
MO[NJE (SLOVENIA)
MINERALO[KA IN GEOKEMI^NA KARAKTERIZACIJA RIMSKE
@LINDRE Z ARHEOLO[KEGA NAJDI[^A PRI MO[NJAH
(SLOVENIJA)
Sabina Kramar1, Judita Lux2, Helmut Pristacz3, Breda Mirti~4,
Nastja Rogan - [muc4
1Slovenian National Building and Civil Engineering Institute, Dimi~eva 12, 1000 Ljubljana, Slovenia
2Institute for the Protection of Cultural Heritage of Slovenia, Preventive Archaeology Department, Tom{i~eva 7, 4000 Kranj, Slovenia
3University of Vienna, Institute of Mineralogy and Crystallography, Althanstrasse 14, 1090 Vienna, Austria
4University of Ljubljana, Faculty of Natural Sciences and Engineering, Department of Geology, A{ker~eva 12, 1000 Ljubljana, Slovenia
sabina.kramar@zag.si
Prejem rokopisa – received: 2013-12-02; sprejem za objavo – accepted for publication: 2014-07-25
doi:10.17222/mit.2013.299
Roman slag from the archaeological site near Mo{nje (NW Slovenia) was studied with respect to its mineralogical and
geochemical characteristics. Samples were analysed with light microscopy, X-ray powder diffraction (XRD), scanning electron
microscopy with an energy dispersive spectrometer (SEM-EDS) and Raman microspectroscopy. A chemical investigation was
carried out using inductively coupled plasma-atomic emission spectroscopy (ICP-ES) to determine the major elements. The
investigated slags are recognized as iron slags, grouped into two main classes according to their chemical and mineralogical
compositions. The first group is characterized by high Fe2O3 and low SiO2 contents and the second by high CaO and SiO2 and
low Fe2O3 contents. The phase occurrence obviously depends on the chemical composition as the first group mainly consists of
fayalite, wüstite and magnetite, and the second one of augite, leucite, hedenbergite, monticellite and kirschsteinite.
Keywords: slag, archaeometallurgy, iron slag, Roman slag, archaeological site near Mo{nje (Slovenia)
V prispevku je obravnavana rimska `lindra z arheolo{kega najdi{~a Mo{nje (SZ Slovenija), pri kateri smo dolo~ili mineralno in
kemijsko sestavo. Vzorci so bili analizirani s svetlobno mikroskopijo, rentgensko pra{kovno difrakcijo, SEM-EDS in ramansko
mikrospektroskopijo. Kemijska sestava glavnih elementov je bila dolo~ena z ICP-ES. Preiskana `lindra dokazuje, da izvira iz
pridelave `eleza na tem podro~ju. Glede na mineralo{ko in kemijsko sestavo lahko delimo `lindro na dve skupini. Za prvo
skupino je zna~ilna visoka vsebnost Fe2O3 in nizka vsebnost SiO2, za drugo pa visoki vsebnosti SiO2 in CaO ter nizka vsebnost
Fe2O3. Kemijska sestava vpliva na mineralne faze; tako je za prvo skupino zna~ilen fajalit, wustit in magnetit, za drugo pa avgit,
levcit, hedenbergit, monticelit in kirschsteinit.
Klju~ne besede: `lindra, arheometalurgija, `elezova `lindra, rimska `lindra, arheolo{ko najdi{~e pri Mo{njah (Slovenija)
1 INTRODUCTION
The archaeological site near Mo{nje with a Roman
villa rustica is located in the northwestern region of
Slovenia. Five masonry structures, incorporated in an
embankment, were recognized. The main structure repre-
sents a residential building with seven rooms, including
baths, decorated with a floor mosaic and wall paintings.
The discovery of small finds such as coins, jewellery,
elements of costume and fragments of pottery within
closed stratigraphic layers suggests that the Roman villa
rustica was inhabited during the first and second
centuries AD and was already in ruins by the third
century, whereas the archaeological finds from the mixed
layers indicate that this area has been inhabited since the
early Iron Age1. Excavation also provides evidence of
prehistoric metallurgical activities as numerous slag
occurrences – almost 400 pieces of slag – were evi-
denced at the mentioned location.
Slags form as by-products of metallurgical processes
and they often represent the only relics of prehistoric
metallurgical activities at certain archaeological sites2–4.
Slags vary in terms of size, shape, and chemical and
mineralogical compositions as a consequence of the
technological process. In particular, the compositions
and properties of metallurgical slags are influenced by
the types of ores, the fluxes added, the furnace construc-
tion material (the lining), the use of charcoal and the
cooling conditions5. Iron slags are routinely discovered
at almost all archaeological sites throughout Slovenia6–9
and occasionally blacksmiths’ hearths or smelting fur-
naces are also found10.
Preliminary analyses revealed that the archaeometal-
lurgical samples from the archaeological site near
Mo{nje (Slovenia) differentiated in their mineralogical
and chemical compositions11. Consequently, in this
study, detailed mineralogical and geochemical characte-
ristics of the slag from the Roman villa rustica were
Materiali in tehnologije / Materials and technology 49 (2015) 3, 343–348 343
UDK 66.046.58 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)343(2015)
investigated in order to obtain information on the metal-
lurgical operation and, thus, contribute to better
knowledge and evidence of metallurgical activities.
2 EXPERIMENTAL WORK
2.1 Materials
A set of 12 archaeometallurgical samples were se-
lected for the study. Information regarding the samples is
provided in Table 1. Samples of the slag were selected
from the closed stratigraphic units of the villa rustica
(Figure 1). Thus, the slags from the stratigraphic units
that were determined as antique (i.e., SU 280, SU 337
inside the object or SU 281) or presumably antique (SU
367) and the layers that periodically contained different
finds (from the Roman and recent periods, i.e., SU 126,
SU 38) were sampled and compared. Additionally, vari-
ous layers from the inside objects, such as rubble layers
(SU 337) or different coat layers (SU 280, SU 281, SU
126, 38, 367), fill layers (SU 296) and hearth remnants
(SU 285) from the outside objects were selected. This
selection enabled us to obtain a variety of samples that
provided information regarding the diversity of the finds.
2.2 Methods
Polished cross-sections of the samples were studied
with light microscopy using a Zeiss AX 10 equipped
with an AxioCam MRc5 digital camera.
The mineral composition of the slag samples was
determined with X-ray powder diffraction (XRD), using
a Philips PW3710 X-ray diffractometer equipped with
Cu-K radiation and a secondary graphite monochro-
mator. Data were collected at 40 kV and a current of 30
mA in the range from 2 = 2 ° to 70 ° with a speed of 3.4
r/min. The results were analyzed with the EVA diffrac-
tion software v12 using ICSD powder diffraction files.
For the purpose of the XRD analysis, the samples were
milled in an agate mortar to a particle size of less than 50
μm.
Polished cross-sections of the samples were exa-
mined in the back-scattered-electron (BSE) imaging
mode of low-vacuum scanning electron microscopy
(SEM) and the energy dispersive X-ray (EDS) technique
using a JEOL 5500 LV instrument.
S. KRAMAR et al.: MINERALOGICAL AND GEOCHEMICAL CHARACTERIZATION OF ROMAN SLAG ...
344 Materiali in tehnologije / Materials and technology 49 (2015) 3, 343–348
Table 1: Summary of the investigated samples, their locations and mineral compositions
Tabela 1: Sinteza preiskanih vzorcev z njihovo lokacijo in mineralno sestavo
Sample Location Stratigraphy Dating Mineral composition
A (6597) J/6,7 280 Layer, Roman period;1st–3rd c. AD
Fayalite, wüstite
B (1499) L6 126 Layer, mixed,prehistoric–recent
Fayalite, wüstite, magnetite, quartz, goethite, -Fe?
C (2630) D, E/25, 26 38 Layer, mixed, Romanperiod–recent
Wustite, magnetite, quartz, goethite, lepidocrocite
D (1580) K6 126 Layer, mixed,prehistoric–recent
Fayalite, wüstite, quartz, dolomite, magnetite,
lepidocrocite, graphite?
E (4193) J7 281 Layer, Roman period,1st, 2nd c. AD
Magnetite, goethite, quartz, fayalite, wüstite
F (5006) J7 281 Layer, Roman period Magnetite, quartz, augite, leucite, fayalite, wüstite,kirschsteinite, akermanite?
G (5129) I/28, 29 367 Villa’s entrance,layer, Roman period?
Quartz, augite, kirschsteinite?, fayalite, leucite, cristo-
balite, hercynite, monticellite, magnetite, maghaemite?
H (3531) G6 296 Fill layer, Romanperiod
Quartz, cristobalite, augite, leucite, hedenbergite,
monticellite
I (3823) H6 285 Hearth, Romanperiod
Wuestite, kirschsteinite, pyrolusite, akermanite?, leucite
J (1999) K7 194 Layer, Roman period,1st–3rd c. AD
Quartz, augite, anorthite, hedenbergite
K (4100) K, L/12 0337
Object 2, room 1,
layer, Roman period,
1st–2nd c. AD
Quartz - mullite 3 : 2, cordierite, -Fe?, magnetite, wüstite
L (2641) D24 38 Layer, mixed, Romanperiod–recent
Quartz, diopside, anorthite, leucite, akermanite?, olivine
Figure 1: Plan of the villa rustica near Mo{nje with marked sampling
locations
Slika 1: Na~rt vile rustike pri Mo{njah z ozna~enimi vzor~nimi mesti
Raman spectra of the phases were obtained on the
polished cross-sections with a Horiba Jobin Yvon
LabRAM HR800 Raman spectrometer equipped with an
Olympus BXFM light microscope. The measurements
were made using a 633 nm laser excitation line (6.6 mW
measured behind the objective) and a Leica 100 ° objec-
tive was used. The wavenumber accuracy was better than
1 cm–1 and the spectral resolution was about 2.5 cm–1.
The samples were analysed for their major chemical
components (SiO2, Al2O3, Fe2O3, MgO, CaO, Na2O,
K2O, TiO2 and P2O5) in an accredited commercial Cana-
dian laboratory (Acme Analytical Laboratories, Van-
couver, B.C., Canada) after fusion with a mixture of
lithium metaborate/tetraborate and dissolution in nitric
acid using inductively coupled plasma emission spectro-
scopy (ICP-ES). The accuracy and precision of the slag
analysis were assessed using the reference material
CCRMR SO-18 CSC. The analytical precision and accu-
racy were better than ± 5 % for the investigated ele-
ments. This was indicated by the results of duplicate
measurements in 12 slag samples as well as duplicate
measurements of the standards.
To identify distinct groupings of the investigated
samples, a hierarchical cluster analysis was applied using
Ward’s method and squared Euclidean distance. The
basic statistical parameters for each element and the
calculations mentioned above were computed using the
statistical software program Statistica VII.
3 RESULTS AND DISCUSSION
3.1 Mineral composition
The mineral compositions of the samples were
determined by combining various methods and they are
given in Table 1. For instance, using Raman microspec-
troscopy, tiny crystals that could not be analysed with
XRD due to their small quantity were also detected. The
X-ray powder-diffraction patterns of the selected slag
samples are shown in Figure 2.
S. KRAMAR et al.: MINERALOGICAL AND GEOCHEMICAL CHARACTERIZATION OF ROMAN SLAG ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 343–348 345
Figure 3: Microphotographs of the investigated samples: a) globular wüstite, sample A, b) dendritic wüstite and lamellar fayalite, sample D, c)
abundant leucite in sample G, d) preserved carbonate structure indicating siderite ore, sample I
Slika 3: Mikroposnetki preiskanih vzorcev: a) globularni wüstit, vzorec A, b) dendritski wüstit in lamelarni fajalit, vzorec D, c) levcit, ki se
pojavlja v ve~jih koli~inah, v vzorcu G, d) ohranjena karbonatna struktura, ki nakazuje sideritno rudo, vzorec I
Figure 2: X-ray powder diffraction patterns of the selected samples
Slika 2: Rentgenogrami izbranih vzorcev
The investigated samples were distinguished by their
mineral compositions and different distributions of the
phases, most probably reflecting different stages of the
iron production. The samples were recognized as the
iron slags that could be grouped, in general, into two
main classes according to the mineral composition.
Wüstite and fayalite are the most abundant phases in
the samples of the first group (A–E). Wüstite is particu-
larly abundant in sample A, where it occurs as globular
grains (Figure 3a). Dendritic wüstite is often observed
(Figure 3b). A common feature of fayalite (Figure 4a)
is idiomorphic lath crystals. Besides, goethite and lepi-
docrocite as the weathering products or magnetite could
also be detected in those samples. Quartz is present in
some samples, probably from the smith’s hearth, and
dolomite is identified in sample D. The glassy matrix is
mostly formed of Si, Al, and Fe; small quantities of K,
Na, Ca, and P are also present, as determined with
SEM-EDS. Samples B, C, D and E also contain the
remnants of charcoal particles, which were incorporated
into the slag at the smithing hearth. Larger areas of
metallic iron, somewhat limonitized, are observed in
samples B, C, and E, which could be part of refined
bloom – gromp12.
Concerning the second group of samples (F–L), the
more abundant slag phases are represented by the
pyroxenes augite and hedenbergite, and Ca-rich olivine
monticellite, followed by leucite (Figures 3c and 4b).
Leucite, which is especially abundant in sample G,
shows that the melt contains enhanced quantities of Al
and K with respect to the melt of the samples within the
first group. The samples are depleted in iron oxide,
which was also noted in their microstructure, due to the
absence of wüstite. While wüstite and olivine minerals
predominate within the first group, in the samples of the
second group they are present in small quantities or
absent. Often two generations of fayalite are observed.
Cordierite, mullite, cristobalite, pyrolusite and akerma-
nite are also found, indicating furnace or smith’s hearth
linings. These are the minerals that were formed by
firing at over 900 °C in the solid state and/or partly
liquid state from the melt. However, as observed with the
microscope, in those samples, slag is adhered to the
ceramic (samples H, J and K). The melt of the slag is
surrounded by partially reacted quartz grains; in some
samples (sample H) the grains of quartz sand trans-
formed into the high-temperature-modified cristobalite
during the firing, which suggests a temperature of 1200
°C. According to the SEM-EDS analyses the glassy
phase within this group contains rather smaller amounts
of Fe (up to amount fractions 2 %) and larger amounts of
Mg (about amount fractions 3 %) with respect to the first
group. Furthermore, compared with the olivine of the
first group, olivine laths contain a higher amount frac-
tions of Mg (8 %) and less Fe (16 %). Hedenbergite and
monticellite also crystallized in some samples that
incorporated enhanced values of Ca (around 6.5 %).
Ca-rich minerals, such as augite, hedenbergite and
monticellite, formed due to the enhanced Ca content in
the melt, which could have originated from several
sources such as ore, ash, furnace lining or an addition of
lime fluxes13. The use of sideritic ore for the iron produc-
tion was proved by the preserved carbonate structures in
sample I (Figure 3d). The siderite ore known from the
vicinity of the archaeological site, i.e., Savske jame,
Jesenice, occurring as veins between dolomite and lime-
stone14 could have eventually contributed to the higher
lime levels. Concerning the furnace or smith’s hearth
linings, a lime-rich furnace lining might have resulted in
the formation of lime-rich slag even if the smelted ore
contained little lime13. A high quantity of pores is cha-
racteristic for the samples within the group.
3.2 Chemical analysis
The major elements of the 12 investigated samples
along with their mean and standard deviations are pre-
sented in Table 2. The samples have heterogenic chemi-
cal compositions.
The amount fractions of SiO2 ranges from 5.79 % to
65.24 % in the samples. The highest SiO2 concentrations
are detected in the samples of the second group (F–L).
S. KRAMAR et al.: MINERALOGICAL AND GEOCHEMICAL CHARACTERIZATION OF ROMAN SLAG ...
346 Materiali in tehnologije / Materials and technology 49 (2015) 3, 343–348
Figure 4: Raman spectra of the slag minerals of: a) fayalite (sample
D) and b) leucite (sample G)
Slika 4: Ramanski spekter mineralov `lindre: a) fajalit (vzorec D) in
b) levcit (vzorec G)
Iron (Fe2O3) and alumina (Al2O3) concentrations range
from 6.18 % to 87.30 % and from 2.90 % to 17.36 %,
respectively. The highest Fe2O3 concentrations are deter-
mined in samples A–E (the first group), which are
mostly associated with different types of iron oxides and
hydroxides (weathering) and the presence of wüstite and
fayalite in these samples. According to the increased
abundance of alumosilicates (augite, leucite, cordierite,
mullite, and anorthite) in the studied samples, the sam-
ples of the second group contain the highest amount of
Al2O3. Alkalis have moderate concentrations (0.05–2.48
% for Na2O and 0.16–4.016 % for K2O). The highest
potassium amounts are observed in the samples of the
second group, which are associated with the higher
leucite amounts in the mentioned samples. In addition,
the glassy matrix also hosts the elements that do not
enter crystalline silicates, such as potassium originating
from the charcoal15. MgO ranges between 0.27 % and
6.21 % (the highest values in the samples of the second
group might be related to the occurrence of the Mg-rich
olivine and Mg-enriched glassy matrix), while the CaO
content varies between 0.28 % and 14.42 % and is
enhanced in the samples of the second group where it is
sometimes incorporated in the Ca-rich olivine and
hedenbergite. Titanium (TiO2) and phosphorous (P205)
concentrations range from 0.07 % to 1.10 % and from
0.20 % to 2.53 %, without significant variations amongst
the analysed samples.
The major elemental compositions generally confirm
the mineral compositions of the slag samples. Addition-
ally, the chemical compositions of the major elements
were subjected to a multivariate analysis in order to
distinguish individual groups among the samples studied.
Figure 5 shows the results of the multivariate analysis.
Two features became evident from the evaluation of
Figure 5. First, two samples, F and I, are placed as out-
liers within the first group. This could be because of the
absence or a small quantity of the slag minerals. Second,
the remaining 10 samples are clustered into two different
groups according to the prevailing component: Fe-mine-
rals or alumosilicates. The first group includes five
samples (A–E) characterized by high Fe2O3 and low SiO2
contents. The second group comprises five samples (G,
L, K, H, and J) characterized by high SiO2 contents,
most probably derived from quartz.
The enhanced Al2O3 values result in the crystalliza-
tion of plagioclases and other alumosilicates. In the
samples of the second group Mg is enhanced and it is
incorporated in the glassy phase, especially in samples H
and J. In addition, olivine minerals are also enriched in
Mg in these samples, as identified with SEM-EDS. The
samples of the second group have enhanced values of Ca
in the glassy phase. The enhanced Ca values in the sam-
ples of the second group allow the occurrence of Ca-rich
olivines – kirschsteinite and monticellite – and Ca-rich
pyroxene hedenbergite, and Ca is incorporated in the
glassy phase. The Na values are enhanced in the samples
of the second group and could be incorporated in the
augite minerals, which was proved with the XRD analy-
S. KRAMAR et al.: MINERALOGICAL AND GEOCHEMICAL CHARACTERIZATION OF ROMAN SLAG ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 343–348 347
Table 2: Chemical compositions (mass fractions, w/%) of the samples
Tabela 2: Kemijska sestava (masni dele`i, w/%) vzorcev
Sample SiO2 Al2O3 Fe2O3 MgO CaO Na2O K2O TiO2 P2O5 MnO
A 11.20 3.16 87.30 1.00 2.56 0.07 0.62 0.16 0.58 0.16
B 15.43 3.92 72.77 0.66 1.44 0.16 0.68 0.22 0.42 0.15
C 11.16 2.90 68.98 0.48 2.17 0.17 0.75 0.15 0.59 0.15
D 17.36 4.40 70.23 0.83 1.90 0.18 0.82 0.25 0.25 0.18
E 5.79 1.23 77.93 0.27 0.98 0.05 0.16 0.07 0.58 0.10
F 33.11 7.03 46.68 2.70 8.85 0.64 2.60 0.38 0.47 0.16
G 57.62 14.32 14.27 1.93 4.61 0.95 4.16 0.73 0.20 0.08
H 52.76 11.91 10.03 5.13 11.81 0.61 2.74 0.56 0.69 0.15
I 25.49 7.06 39.64 2.78 14.42 0.48 3.17 0.36 2.53 0.38
J 54.77 14.33 6.18 6.21 10.76 0.61 1.82 0.77 0.54 0.13
K 65.24 17.36 7.73 2.06 1.47 0.74 3.80 1.10 0.32 0.16
L 56.49 10.87 17.50 1.54 6.44 2.48 1.81 0.46 1.04 0.16
Min 5.79 2.90 6.18 0.27 1.44 0.07 0.16 0.07 0.20 0.08
Max 65.24 17.36 87.30 6.21 11.81 2.48 4.16 1.10 2.53 0.38
Mean 33.87 8.21 43.27 2.13 5.62 0.60 1.93 0.43 0.68 0.16
SD 22.06 5.37 31.15 1.86 4.73 0.66 1.35 0.31 0.62 0.07
Figure 5: Investigated-sample dendrogram
Slika 5: Dendrogram preiskanih vzorcev
sis of the samples in the second group. Potassium is
sourced from the charcoal and is enhanced in the sam-
ples of the second group, where it could be incorporated
in the leucite mineral and the glassy matrix. K and Na
are correlated, which could indicate that Na and K were
derived together from wood or charcoal.
4 CONCLUSIONS
The obtained results showed that the investigated
samples from the archaeological site near Mo{nje (Slo-
venia) indicate an iron production.
The samples are recognized as iron slags that can be
grouped into two main slag classes according to their
mineral compositions, representing different stages of
the iron processes. The first group is characterized by
high Fe2O3 and CaO and low SiO2 contents. The second
group is characterized by high SiO2 contents. The phase
occurrence obviously depends on the chemical compo-
sition as the first group consists of fayalite, wüstite and
magnetite and the second of augite, hedenbergite, mon-
ticellite and kirschsteinite. The slag in some samples of
the second group is adhered to the ceramic, indicating
smelting or smith’s hearth lining, which could have been
the source of the enhanced CaO values of those samples.
The preserved carbonate structures indicating roasted
siderite ore prove the use of the local siderite ore for the
iron production.
The results showed that samples A (SU 280), G (SU
367) and L (SU 38) are very particular due to their sur-
face textures or mineral compositions. Sample C from
SU 38 belongs to a similar group of samples (B, D, and
E), which were found in SU 126 and 281. The second
group of similar samples (F, H, I, J and K) was found in
Roman layers SU 194, 281, 296, 285 and 337. Thus, the
groups established on the basis of the chemical-mine-
ralogical analyses are not in complete accordance with
the grouping of the samples based on their stratigraphy.
For instance, among slag samples C, B, D and E, only
one example (E) was found in the layer that is dated to
the Roman period, while the rest are dated to a longer
period – from prehistory to the modern era. However, the
results indicate general varieties in the types of the slag
from the archaeological site of Mo{nje that can be
defined in more detail only with a further study.
Acknowledgements
The above analyses were conducted as part of the
post-excavation process and the study of the archaeolo-
gical site of Mo{nje – Pod cesto (E[D 10036). This work
was financially supported by ARRS Programme Group
P2-0273 and Programme Group P1-0195.
5 REFERENCES
1 J. Lux, M. Sagadin, Emona between Aquileia and Pannonia, Sympo-
sium in memory of Ljudmila Plesni~ar Gec, City Museum of Ljub-
ljana, Slovenia, 2010, 27–28
2 M. L. Wayman, Mater. Charact., 45 (2000), 259–267, doi:10.1016/
S1044-5803(00)00108-X
3 R. Pleiner, Iron in Archaeology: The European Bloomery Smelters,
Archaeological Institute of the Academy of Sciences, Praha 2000,
418
4 R. Pleiner, Iron in Archaeology: Early European Blacksmiths,
Archaeological Institute of the Academy of Sciences, Praha 2006,
379
5 V. F. Buchwald, Iron and Steel in Ancient Times, Royal Danish
Academy of Sciences and Letters, Copenhagen, 29 (2005), 372
6 N. Zupan~i~, M. Mi{i~, RMZ – Mater. Geoenviron., 48 (2001) 3,
447–457 (in Slovene)
7 L. Orengo, P. Fluzin, Iron metallurgical remnants from Trnava,
Trnava, ZVKDS, Ljubljana 2006, 62–66 (in Slovene)
8 J. Lamut, J. Medved, Analyses of slag, Sela pri Dobu, ZVKDS,
Ljubljana 2007, 63–66 (in Slovene)
9 S. Kramar, V. Tratnik, I. M. Hrovatin, A. Mladenovi}, H. Pristacz, N.
Rogan [muc, Archaeometry, (2014), doi:10.1111/arcm.12116
10 M. Horvat, Sela pri Dobu, ZVKDS, Ljubljana 2007, 68 (in Slovene)
11 U. Umek, Mineralogical characterization of Roman slag from
archaeological site Mo{nje, Diplomsko delo, Ljubljana, 2011, 162
(in Slovene)
12 E. M. Nosek, The metallography of gromps, La sidéurgie ancienne
de l´Est de la France dans son cintexte Européen, Ann. Litt. Univ.
Besançon, 536 (1994), 65–73
13 S. Paynter, Archaeometry, 48 (2006), 271–292, doi:10.1111/j.1475-
4754.2006.00256.x
14 B. Berce, Pregled `elezovih nahajali{~ LR Slovenije (Review of iron
ore deposits in P. R. Slovenia), Prvi jug. Geol. Kongr., Ljubljana,
1956, 235–259 (in Slovene)
15 A. Manasse, M. Mellini, J. Cult. Herit., 3 (2002), 187–198,
doi:10.1016/S1296-2074(02)01176-7
S. KRAMAR et al.: MINERALOGICAL AND GEOCHEMICAL CHARACTERIZATION OF ROMAN SLAG ...
348 Materiali in tehnologije / Materials and technology 49 (2015) 3, 343–348
D. ]UR^IJA et al.: REYNOLDS DIFFERENTIAL EQUATION SINGULARITY USING PROCESSES ...
REYNOLDS DIFFERENTIAL EQUATION SINGULARITY USING
PROCESSES OF SMALL STRAINING WITH LUBRICATION
REYNOLDSOVA DIFERENCIALNA ENA^BA PRI PROCESIH
MAJHNE DEFORMACIJE Z MAZANJEM
Du{an ]ur~ija1, Franc Vodopivec2, Ilija Mamuzi}1
1Croatian Metallurgical Society, Berislavi}eva 6, Zagreb, Croatia
2Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
ilija.mamuzi}@public.carnet.hr
Prejem rokopisa – received: 2014-02-01; sprejem za objavo – accepted for publication: 2014-09-03
doi:10.17222/mit.2014.025
Frequently, simplified partial differential equations include transcendental functions with analytical solutions based on a
singularity. Such solutions are characteristic for numerical analyses with strong Solvers and several programs singularities in the
processes of dressing rolling mills with lubrication. The devised dynamical model includes the variability of the gripping angle
and of the rolls radius in a section of continuous rolling. Below the real lubricant layer, in the analysis two apparent lubricant
layers are presumed. The solution of the differential equation of the lubricant layer with the singularity is obtained using
standard mathematical solutions for apparent lubricant layers. On the ring diagram, the transfer over the singularity shows a
stronger disorder, i.e., disharmony, than the transfer over the transcendental point.
Keywords: Reynolds differential equation, singularity, dressing rolling mill, lubrication, geometrical centre
Pogosto poenostavljene diferencialne ena~be vklju~ujejo transcendentne funkcije z analiti~nimi re{itvami na podlagi singularne
to~ke. Take re{itve so zna~ilne za numeri~ne analize z zmogljivimi re{evalci in ve~ singularnimi to~kami v programih pri
procesih dresirnih valjarn z mazanjem. Predlagan dinami~en model vklju~uje razli~nost prijemnega kota in premera valjev na
delu kontinuirne valjalne proge. Pod realno plastjo maziva sta v analizi predpostavljeni dve navidezni plasti maziva. Re{itev
diferencialne ena~be plasti maziva s to~ko singularnosti je dose`ena z uporabo standardnih matemati~nih metod za navidezni
plasti maziva. Na kro`nem diagramu prenos preko singularnosti poka`e ve~ji nered oziroma disharmonijo, kot je prenos preko
transcendentne to~ke.
Klju~ne besede: Reynoldsova diferencialna ena~ba, to~ka singularnosti, dresirna valjarna, mazanje, geometri~na sredina
1 INTRODUCTION
The Reynolds1 differential equation2,3 is used for the
analysis of the processes of the lubricated low reduction
of metals (dressing, cold rolling and drawing) and a
simplified equation is used4,5:
d
d
Rp
x
v v
x
Q
x
=
⋅ +
−
⋅ ⋅6 120	

	

 

( )
( ) ( )
(1)
The approximate solution using the transcendent
equation is:
A
R R R
=
−
⋅
+
⋅ ⋅
⎛
⎝
⎜
⎞
⎠
⎟ +
⋅
−
⋅
⋅ ⋅


 

 



 
 
 

 



3
2 2 2
⎛
⎝
⎜
⎞
⎠
⎟
−
−
⎛
⎝
⎜
⎞
⎠
⎟ = = −





    

 ln
2
R
2
(2)
A
p
v v
=
− − ⋅
⋅ +
1
6 0
exp( )
)

	 

R
(3)
where A is a technological parameter, R is the rolls
radius, μ is the lubricant dynamical viscosity, v0 and vR
are the rolling and circumferential rolls velocity,  is the
rolling angle, (x) is the geometry of the lubricant layer
in the deformation zone, Q is the lubricant consumption,
dp/dx is the axial stressing gradient x, 0 is the thickness
of the lubricant layer of the entry section of the defor-
mation zone,  is the piezo-coefficient of the lubricant
viscosity and p0 is the rolling pressure.
The entry roll in Section I has the radius R1, the
gripping angle 1 and forms a lubricant layer with the
thickness 1 (Figure 1). The exchange parameter  in
Materiali in tehnologije / Materials and technology 49 (2015) 3, 349–354 349
UDK 519.61/.64:621.77 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)349(2015)
Figure 1: Scheme of the calculation of lubricant layer between two
rolling stands. The entry roll I joins in the singularity, i.e., the trans-
cedency point.
Slika 1: Shema izra~una plasti maziva med dvema valjalnima ogrod-
jema. Vstopni valj I dose`e singularno to~ko oz. transcendentno to~ko.
Section II changes these parameters, either as: R2 = 3 ·
R1, 2 = 1/ or R2 = R1/3, 2 = 1 · 
The exit lubricant layer thickness 2 is calculated
using the Solver solution of the transcendent Equation
(2) and it is confirmed with a Monte Carlo numerical
integration of Equation (1).
2 NUMERICAL ANALYSIS OF THE SINGULA-
RITY D AND THE TRANSCENDENT POINT T
As a special case of the solution of the transcendent
Equation (2) is the singularity solution acceptable for
dressing processes. This solution2,6 is:
  


 
 =
⋅
1
2 15
R
R A
  (4)
Figure 2 shows the occurrence of the singularity in
the dressing process with respect to the lubricant-layer
thickness and the gripping angle. With respect to the
singularity, the transcendent point T is situated on the
right-hand side.
The singularity creates a vicinity of unpredictable
behaviour and the numerical analysis is spread around
the singularity. The determinant (5) was used for the
analysis:
 
 



 
T
MS MT
⎡
⎣⎢
⎤
⎦⎥
(5)
where 
 is the thickness of the lubricant layer for
 
MS is the thickness of lubricant layer according to
the Mizuno-Grudev equation for *, 
T is the thickness
of the lubricant layer according to the transcendent Eq.
(2) and 
MT is the thickness of the lubricant layer
according to the Mizuno-Grudev6 equation  = */ for
 = */.
If the singularity is at the interaction of two points,
the value of the determinant (5) tends to zero. In the
absence of a singularity between the initial and the
aimed for point, the value of the determinant (5) is zero.
This allows us to describe all the aimed for points’ values
without a singularity using:




Target =
⋅
⋅
⎛
⎝
⎜ ⎞
⎠
⎟R
R A15
2
2
(6)
where  is a proportionality constant in the technical
signification equal to the deformation degree. In the
analysis, the apparent lubricant layer 0 (column J) is
divided by the aimed analysis into two pseudo layers
(columns H and I in Tables 1 and 2).
3 SOLVER CALCULATION USING EQUATIONS
(2) AND (6) WITH THE INITIAL ON THE
SINGULARITY
The results of calculations using the Solver (Math-
CAD, EXCEL) after the commutation law of multipli-
cation for the condition of the technological process are
listed in Table 1. With respect to the commutation law of
multiplication, the external ring J is the product of the
inner two rings H and I. The imagined is the rolling line
with rolling 10 cages with deformation degrees varying
from cage to cage. The results of the calculations are
listed in Table 2 for increasing values of .
The constructive coefficient of transfer between two
rolling stands is  · R = 0.22145 · 2.
D. ]UR^IJA et al.: REYNOLDS DIFFERENTIAL EQUATION SINGULARITY USING PROCESSES ...
350 Materiali in tehnologije / Materials and technology 49 (2015) 3, 349–354
Table 1: Solver calculations for Eq. (2) for three singularity rings. The variability of  influences the variability of the clutch angle  and the
rolls radius R.
Tabela 1: Izra~uni ena~be (2) z uporabo re{evalca za tri singularne obro~e. Variabilnost  vpliva na variabilnost prijemnega kota  in polmera
valjev R.
 = 2.2 H I J
2 5.6881486648E-05 4.3089901167E-01 2.4510176379E-05
2.2 5.6754502094E-05 4.7504943260E-01 2.6961194017E-05
2.22 5.2837299927E-05 5.1490700355E-01 2.7206295781E-05
2.222 6.2070776615E-05 4.3870573952E-01 2.7230805957E-05
2.2222 5.4395165663E-05 5.4219876699E-01 2.9492991753E-05
2.22222 5.9160153196E-05 4.6033521899E-01 2.7233502077E-05
2.222222 5.2737233458E-05 5.1640036463E-01 2.7233526587E-05
2.2222222 5.4753146188E-05 4.9738747779E-01 2.7233529284E-05
2.22222222 4.7791846331E-05 5.6983630888E-01 2.7233529308E-05
2.222222222 5.6585859142E-05 4.8127800343E-01 2.7233529310E-05
Figure 2: Vicinity of the singularity (singular point) D( 
) and the
transcendent point T
Slika 2: Bli`ina singularne to~ke D( 
) in transcendentne to~ke T
Figure 3 has a mark for the turn of the apparent
lubricant rings H and I with the singularity as the initial
point with respect to the fictive outer lubricant layer for
the cage rolling for ten cages in the rolling line.
The transfer in Figure 4 also indicates the discord-
ance of the inner two rings as in Figure 3, where the
transfer was achieved using the Solver.
4 APPARENT LUBRICANT LAYERS H AND I
The Solver solution of the transcendent equation, the
apparent layers H and I may have different results for an
equal degree of deformation, as listed in Table 3. The
value of the lubricant layer J = H * I is equal, as also
shown by the solution of Eq. (1).
The constructive transfer coefficient between two
rolling cages is  · R = 0.22145/2.
The data in Table 3 are depicted in Figure 5.
The results of the investigation of the columns H and
I are listed in Table 4. The geometrical average of the
column J is obtained using the apparent lubricant layers
K and I. The relation of the column J and the apparent
lubricant columns K and I are also supported by the
standard mathematical averages: arithmetic, harmonic,
geometric, quadratic, etc.
Although having an equal value to the external, i.e.,
third ring J, the lower apparent rings H and I are in
disharmony with the external ring, and only with values
for i = 1 is the harmony achieved. The second series i = 2
creates the inversion of the two internal rings, while the
series i = 3 supports the turn, i.e., with respect to the
external ring J. As explained already, this is supported by
D. ]UR^IJA et al.: REYNOLDS DIFFERENTIAL EQUATION SINGULARITY USING PROCESSES ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 349–354 351
Table 2: Solver calculations for three singularity rings using Eq. (2)
Tabela 2: Izra~uni z uporabo re{evalca za tri obro~e singularnosti z ena~bo (2)
 =  H I J
3 9.0652411487E-01 4.0556839290E-05 3.6765752839E-05
3.1 7.2525287463E-01 5.2383491694E-05 3.7991277934E-05
3.14 9.9122718480E-01 3.8822066788E-05 3.8481487970E-05
3.141 6.2041700114E-01 6.2044951242E-05 3.8493742586E-05
3.1415 9.7002625173E-01 3.9689514361E-05 3.8499870849E-05
3.14159 8.4111868525E-01 4.5773532910E-05 3.8500973821E-05
3.141592 8.2952986160E-01 4.6413034797E-05 3.8500998332E-05
3.1415926 1.0885265884E+00 3.5369834871E-05 3.8501005684E-05
3.14159265 8.8297139988E-01 4.3603910956E-05 3.8501006297E-05
3.141592654 8.0201286093E-01 4.8005472508E-05 3.8501006346E-05
Figure 4: Transfer of similarity for the singularity to the transcendent
point using Eq. (6)
Slika 4: Prenos podobnosti s to~ke singularnosti na to~ko transcen-
dentnosti z ena~bo (6)
Figure 3: Aimed for transcendent point from the singularity after the
Solver and Eq. (2) with:  = 2.222222222,  = 1.107262798/, R =
0.2 · 3, A = 1965512 m–1, R = 0.2 m,  = 1.107402627 rad, μ ·  =
5.232 · 10–9 and p0 ·  = 4.36 s
Slika 3: Ciljana to~ka transcendentnosti na podlagi singularnosti,
izra~unane z uporabo re{evalca in ena~be (2) z :  = 2,222222222,  =
1,107262798/, R = 0,2 · 3, A = 1965512 m–1, R = 0,2 m,  =
1,107402627 rad, μ ·  = 5,232 · 10–9 in p0 ·  = 4,36 s
the mathematical average and the law of commutation of
the hyperbolic multiplication.
Further, the numbers in Table 4 show that for the va-
lues of columns J (J = H * I):
G = A · H (7)
G is the geometrical average for column J
A is the arithmetic average for column I
H is the harmonic average for column H.
The numbers in column J, Table 3 could also be
obtained with the opposite values of the apparent
lubricant layers K and I, e.g.:
G
H I
H I
E=
+
⋅
+
=6 4
4 62
2
1 1
4 0850293966 06
( / ) ( / )
. – (8)
The connection between the apparent layers is
supported by the algebraic opposite identity of several
possible and with respect to Table 3, on the hyperbole
for the first step it is:
( ) ( ) ( ) ( )
( ) ( )
H I H I
H I
i i i i
i i
+ + +
+ +
+
+
+
=
=
+
1
3 3
3 1
3
1
3
3
1
3
2
2 2
2
3
3
2
(9)
D. ]UR^IJA et al.: REYNOLDS DIFFERENTIAL EQUATION SINGULARITY USING PROCESSES ...
352 Materiali in tehnologije / Materials and technology 49 (2015) 3, 349–354
Table 3: Different values in the columns H and I for the equal degree of deformation  = 3
Tabela 3: Razli~ne vrednosti v kolonah H in I za enako stopnjo deformacije  = 3
i  = 3.0 H I J
1 3 3.6080500075E-05 1.1321986636E-01 4.0850293967E-06
2 3 3.6684996110E-05 1.1135422733E-01 4.0850293964E-06
3 3 4.4999094308E-05 9.0780258123E-02 4.0850293966E-06
4 3 4.5445189361E-05 8.9889148972E-02 4.0850293965E-06
5 3 5.1287540856E-05 7.9649547013E-02 4.0850293966E-06
6 3 6.1373811833E-05 6.6559812314E-02 4.0850293966E-06
Table 4: Geometrival averages of columns H and I in interval form
Tabela 4: Geometri~na povpre~ja kolon H in I v intervalni obliki
i H I J
i = 1 to 6 4.5193358310E-05 9.0390038477E-02 4.0850293966E-06
i = 2 to 6 4.5232311633E-05 9.0312196062E-02 4.0850293966E-06
i = 3 to 6 4.3833687801E-05 9.3193833361E-02 4.0850293966E-06
i = 4 to 6 3.9181673469E-05 1.0425867593E-01 4.0850293966E-06
Figure 5: Ring diagram of Table 3 calculated using the Solver from
the initial singularity for two rolling stands
Slika 5: Kro`ni diagram tabele 3, izra~unan z re{evalcem od za~etne
to~ke singularnosti za dve valjalni ogrodji
Figure 6: Disorder on the transcendent point by a large degree of
metal deformation using the initial singularity calculated with the
Solver
Slika 6: Nered na to~ki singularnosti pri veliki stopnji deformacije za
za~etno to~ko singularnosti, izra~unano z re{evalcem
The singularity coexists with two apparent lubricant
layers, which are in the ring diagram incongruous with
respect to the external apparent lubricant layer J and the
disharmony i.e., the disorder may be obtained. The
internal apparent lubricant rings rotate with respect to the
external fixed ring and neither are congruent.
5 LARGE METAL DEFORMATION
The calculation results are listed in Table 5. By trans-
ferring the similarity from cage to cage, the deformation
coefficient varies strongly.
In Figure 6 the data from Table 5 are depicted as a
ring diagram. The great rings disorder by the transfer of
disharmony on the outer ring J was calculated using the
transcendent Eq. (2). The results of this calculation7 were
confirmed by a numerical integration7 Monte-Carlo of
Eq. (1). The apparent column H is constant by rolling on
a line with 11 rolling cages.
6 TRANSFER TRANSCENDENT TO
TRANSCENDENT POINT
The initial roll in Figure 1 is on the transcendent
point T in Figure 2. The Solver calculation of the lubri-
cant layer between two connected cages shows great
stability and rhythmics (Figure 7).
The apparent lubricant layers K and I turn within and
unrelated to the external lubricant layer J. The con-
gruency of all three rings is achieved without discordant
inversions in rings H and I.
The transfer is:
R Ri i+ =1
3   i i /+ =1 i = 3.141592654
Using  = 1 no transfer of similarity coefficient
should occur as both rolls in Figure 1 are in the same
position and no metal deformation occurs on the dressing
line.
7 CONCLUSION
The mixing and penetration of the layers in the H and
I rings by calculation using Eq. (6) are the evaluation of
the stability of technological procedure of rolling on
continuous lines with different degrees of deformation .
The mixing is diminished essentially by Solver calcu-
lations, although the layer inversion could be obtained,
also. Besides inversion, according to law commutation of
multiplication, the layers are inclined to rotation at the
singularity with respect to the outer layers. For proper
use of the Solver program, experience is necessary. The
apparent layers H and I are connected to the column J
and the related geometrical average after Eq. (7) verified
by numerical integration. It was shown that the data in
column J in Table 4 could also be calculated using Eq.
(8), which was not generalised. Eq. (9) is a possible con-
nection of apparent lubricant layers K and I. The trans-
cendent point shows the marked order for the lubricant
layer better than the singularity.
The results of the calculation of lubricant layer using
the Reynolds equation agree well with experimental
results for the processes of dressing of bands and cold
D. ]UR^IJA et al.: REYNOLDS DIFFERENTIAL EQUATION SINGULARITY USING PROCESSES ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 349–354 353
Table 5: Solver calculation for large metal deformation from the singularity to the transcendent point
Tabela 5: Izra~uni z uporabo re{evalca za velike deformacije od to~ke singularnosti do to~ke transcendentnosti
H I J
2 4.9140375541E-01 1.4369767894E-05 7.0613579075E-06
3 4.9140375543E-01 2.8739535788E-05 1.4122715816E-05
4 4.9140375554E-01 4.3109303672E-05 2.1184073723E-05
5 4.9140375543E-01 5.7479071576E-05 2.8245431631E-05
6 4.9140375544E-01 7.1848839468E-05 3.5306789539E-05
7 4.9140375585E-01 8.6218607289E-05 4.2368147446E-05
8 4.9140375543E-01 1.0058837526E-04 4.9429505355E-05
9 4.9140375543E-01 1.1495814315E-04 5.6490863261E-05
10 4.9140375543E-01 1.2932791105E-04 6.3552221172E-05
11 4.9140375543E-01 1.4369767894E-04 7.0613579078E-05
12 4.9140375543E-01 1.5806744683E-04 7.7674936983E-05
Figure 7: Example of rhythmics by ten cages
Slika 7: Primer ritmi~nosti za deset ogrodij
tube drawing8,9 and by investigations of the contact fric-
tion10.
Acknowledgements
The authors are indebted to prof. F. Vodopivec for
useful comments and the Croatian-to-English translation
of the manuscript.
8 REFERENCES
1 D. ]ur~ija, F. Vodopivec, I. Mamuzi}, Mater. Tehnol., 47 (2013) 1,
53–57
2 D. ]ur~ija, I. Mamuzi}, Mater. Tehnol., 41 (2007) 1, 21–27
3 D. ]ur~ija, I. Mamuzi}, Mater. Tehnol., 42 (2008) 2, 59–63
4 D. ]ur~ija, I. Mamuzi}, Metalurgija, 44 (2005) 3, 221–226
5 D. ]ur~ija, I. Mamuzi}, Metalurgija, 44 (2005) 4, 295–300
6 D. ]ur~ija, Mater. Tehnol., 37 (2003) 5, 237–250
7 M. I. Sobol, Die Monte Carlo Methode, Ver. H. Deutsch, Frankfurt a.
Main 1986
8 O. P. Maksimenko, A. A. Semen~a, Issledovanie kontaktno-gidro-
dinami~eskoj smazki pri prokatke, Su~asni problemi metallurgii, 8
(2005), 99–103
9 M. R. Jensen, L. Olovsson, J. Danckert, Numerical model for the oil
pressure distribution in the hydromechanical deep drawing process,
J. Mater. Process. Techn., 103 (2000) 1, 74–79, doi:10.1016/S0924-
0136(00)00421-0
10 P. Heyer, J. Läuger, Correlation between friction and flow of lubri-
cating greases in new tribometer device, Lubrication Science, 21
(2009) 7, 253–268, doi:10.1002/ls.88
D. ]UR^IJA et al.: REYNOLDS DIFFERENTIAL EQUATION SINGULARITY USING PROCESSES ...
354 Materiali in tehnologije / Materials and technology 49 (2015) 3, 349–354
M. ^ILLIKOVÁ et al.: PREDICTION OF THE CATASTROPHIC TOOL FAILURE OF CERAMIC INSERTS ...
PREDICTION OF THE CATASTROPHIC TOOL FAILURE IN HARD
TURNING THROUGH ACOUSTIC EMISSION
NAPOVEDOVANJE KATASTROFI^NE PO[KODBE KERAMI^NIH
VLO@KOV PRI STRU@ENJU Z AKUSTI^NO EMISIJO
Mária ^illiková1, Branislav Mi~ieta1, Miroslav Neslu{an1, Robert ^ep2,
Ivan Mrkvica2, Jana Petrù2, Tomá{ Zlámal2
1University of @ilina, Faculty of Mechanical Engineering, Univerzitna 8215/1, 010 26 @ilina, Slovakia
2V[B-Technical University of Ostrava, Faculty of Mechanical Engineering, 17. listopadu 15, 708 33 Ostrava, Czech Republic
maria.cilikova@fstroj.uniza.sk
Prejem rokopisa – received: 2014-02-07; sprejem za objavo – accepted for publication: 2014-09-03
doi:10.17222/mit.2014.029
The paper deals with a new concept for the detection and prediction of the catastrophic tool failure (CTF) of ceramic inserts
using an acoustic emission (AE) technique and an associated analysis of the chip formation during hard turning of bearing steel
100Cr6. The suggested method is based on the application of two sensors and the ratios of parameters of the acoustic emission
such as AE RMS, AE absolute energy, and AE strength. The specific character of the segmented chips during hard turning is
associated with the raw acoustic signals as well as the extracted AE features. The paper indicates that the conventional data
processing of acoustic emission signals enables the detection of CTF. Tool wear connected with the cutting edge micro-chipping
is related to the slow increase of tool wear (mainly flank wear VB) and stable values of AE features in the normal phase of tool
wear. The CTF alters the AE waveforms as well as the course of the AE features. Conventional AE signal processing enables the
detection of tool breakage. However, approaching CTF itself cannot be reliably predicted. Hence, a new concept of AE
processing based on the ratios of the extracted AE features obtained in the different frequency ranges is suggested.
Keywords: hard turning, acoustic emission, wear, prediction
^lanek obravnava nov na~in odkrivanja in napovedovanja poru{itve (CTF) kerami~nih vlo`kov z uporabo tehnike akusti~ne
emisije (AE), povezano z analizo nastanka ostru`ka med stru`enjem jekla 100Cr6 za le`aje. Predlagana metoda temelji na
uporabi dveh senzorjev in razmerij pri akusti~ni emisiji, kot so AE RMS, absolutna energija AE in AE-mo~. Specifi~na oblika
delca ostru`ka med stru`enjem je povezana s surovim akusti~nim signalom, kot tudi z lo~enimi AE-lastnostmi. ^lanek
obravnava, kako konvencionalna obdelava podatkov signalov akusti~ne emisije omogo~a dolo~itev CTF. Obraba orodja je
povezana s kru{enjem rezilnega roba ter s po~asnim pove~anjem obrabe orodja (ve~inoma obrabe boka VB) in s stabilnimi
vrednostmi AE-lastnosti pri normalni obrabi orodja. CTF spreminja obliko AE-valov, kot tudi AE-lastnosti. Konvencionalna
obdelava AE-signalov omogo~a odkritje loma orodja. Vendar pa napovedovanje nastanka CTF ni mogo~e zanesljivo napovedati.
Predlagan je torej nov na~in obdelave AE, ki temelji na izbranih AE-lastnostih, dobljenih pri razli~nih obmo~jih frekvenc.
Klju~ne besede: stru`enje, akusti~na emisija, obraba, napovedovanje
1 INTRODUCTION
The hard tuning process has found great industrial
relevance as a result of recent developments in machine
tools (especially cubic boron nitride and ceramics). High
flexibility, high removal rates, and the ability to manu-
facture a complex workpiece geometry in one set are the
main advantages of hard turning operations compared to
grinding1. Furthermore, hard turning makes it possible to
avoid coolants and therefore can actually be regarded as
an interesting alternative, even from the ecological point
of view1,2.
On the other hand, the main disadvantages of hard
turning can be found in the formation of white layers
induced in the early stages of tool wear and the high risk
of unexpected catastrophic tool failures (CTFs) at the
end of the normal phase of tool wear (especially in the
case of ceramic inserts3). Macro-chipping of the cutting
edge should be regarded as a serious problem. CTF can
damage the machined surface or tool holder and drama-
tically alter the tool geometry and corresponding work-
piece dimensions. Thus, hard turning operations need the
incorporation of a properly proposed monitoring system
for the detection or, much better, the prediction of CTF.
Lee, Dornfeld, and Wright4 reported on a variety of
sensors, each having a certain degree of applicability, for
the monitoring of the cutting process. Monitoring tech-
niques based on the implementation of dynamometers,
accelerometers, laser interferometers, acoustic emission,
and so on were proposed and properly integrated into
machine control systems5. However, the specific mecha-
nism of chip formation during hard turning means that
specific requirements should be met to suggest the pro-
per concept for monitoring hard turning operations.
It is well known that segmented chips are produced
during hard turning6,7 as a result of the poor plasticity
and high hardness and strength of the work material.
Thus, thermal softening dominates over strain hardening
when the hardness of the work material attains a certain
value. It should also be mentioned that the shear insta-
bility during chip segmentation is a function, not only of
Materiali in tehnologije / Materials and technology 49 (2015) 3, 355–363 355
UDK 621.941:539.538:539.42 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)355(2015)
the hardness of the work material, but also such aspects
as cutting conditions or tool geometry6,7. Figure 1 illu-
strates a segmented chip during the turning of hardened
steel 100Cr6. Poulachon and Moisan8 published an out-
standing study of the different phases of chip segmen-
tation as a series of micrographs. They clearly proved
that segment formation starts with a crack initialization
near the free surface and its propagation towards the tool
tip of the cutting edge (zone Ib; Figure 2). However,
crack propagation transforms into a plastically deformed
zone where microcracks still occur (zone Ia, Figure 2).
The variable chip cross-section is due to the sliding of a
segment along the fully cracked surfaces as well as the
thinning of the microcracked region during sliding.
Severe plastic deformation can also be found at the
tool–chip interface. Thus, the segmented chip is a mix-
ture of zones of extremely high intensity structural
transformation located on segment boundaries (narrow
white bands) and the nearly untouched structure inside
the segment.
In the case of a continuous chip, the rapid shift of the
plastic deformation from the actual shear plane to the
adjacent plane of the incoming work material is driven
by the strain hardening. Therefore, a more pronounced
homogeneity of the continuous chip can be found. As a
result of the different mechanism of chip formation, con-
ventional soft turning is a process during which the ener-
gy needed for chip separation is continuously released,
while the chip segmentation represents the cyclic process
when the energy accumulated in front of the cutting edge
is suddenly emitted when the shear strain near the free
surface attains the value which the structure cannot
afford to exceed7.
This cyclic process is associated with the fluctuation
of the cutting force and the corresponding segmentation
frequency. This segmentation frequency is a function of
the work material properties and the cutting conditions
as well as the tool geometry. Shaw and Vyas6 reported
that the segmentation frequency is about 18 kHz for a
cutting speed of 100 m min–1 and a feed of 0.28 mm, and
the case of carburized steel with a hardness of 62 HRc.
Neslu{an3 reported that the segmentation frequency is
strongly dependent on the feed and cutting speed and can
be found in the range of 14 kHz to 90 kHz. A dynamic
analysis of the processes that fluctuate at a frequency
over 25 kHz with the application of conventional accele-
rometers or dynamometers is difficult to carry out. On
the other hand, acoustic emission (AE) techniques enable
detection of processes that fluctuate at a frequency over
several megahertz.
The AE technique can be successfully adapted for
machining operations9. This technique is sensitive to
events such as dislocation movements, deformation,
inclusion fracture, crack propagation, and so on. The
major AE sources10 in a metal cutting process are the
deformation and fracture of work materials in the cutting
zone, deformation, and mainly the fracture of cutting
tools, collisions, entanglement, and the breakage of
chips.
AE signals can be classified11,12 as being either conti-
nuous-type or burst-type AE signals. Continuous signals
are usually associated with shearing in the primary zone
and at the tool–chip and tool–workpiece interfaces, while
burst-type AE signals are observed during crack growth
in the material, tool fracture, or chip breakage. The
major advantage of the AE technique is its high fre-
quency range (much higher than that of the machine
vibrations and environmental noises)10.
AE techniques were adapted for monitoring turning13,
milling14, drilling15, grinding16,17, and precision machin-
ing18. Suggested concepts based on the AE technique are
based on the sensitivity of the AE signal to various
contact areas and deformation regions. Jemielniak and
Otman19 reported that the skew and kurtosis extracted
from the raw AE signal are more sensitive to CTF than
conventional AE signals and AE features. Li and Yuan20
gives a brief review of the acoustic emission methods for
tool-wear monitoring during turning. Dolin{ek and Ko-
pa~21 investigated the AE signals and associated process
M. ^ILLIKOVÁ et al.: PREDICTION OF THE CATASTROPHIC TOOL FAILURE OF CERAMIC INSERTS ...
356 Materiali in tehnologije / Materials and technology 49 (2015) 3, 355–363
Figure 2: Cutting zone during hard cutting – brief sketch: Ia-
microcracked and plastically deformed shear zone, Ib-cracked region,
II-tool–chip contact, III-tool–workpiece contact
Slika 2: Obmo~je rezanja med stru`enjem – shematsko: Ia-plasti~no
deformirano podro~je stri`enja z mikrorazpokami, Ib-razpokano
podro~je, II-stik orodje-ostru`ek, III-stik orodje-obdelovanec
Figure 1: Segmented chip 100Cr6 (hardened 62 HRc), vc = 100 m
min–1, f = 0.09 mm
Slika 1: Segmentirani ostru`ki 100Cr6 (trdota 62 HRc, vc = 100 m
min–1, f = 0,09 mm
of tool-wear monitoring. They noticed that tool wear is
one of the most influential factors contributing to an
increase in the energy of the AE signal. Govekar,
Gradi{ek, and Grabec22 suggested a system and methods
for the extraction of useful information from signals of
multiple sensors (including the AE technique). Inasaki12
reported that continuous-type AE signals are associated
with plastic deformations in ductile materials, while
burst-type signals are observed during crack growth in
the material. Also, many attempts to predict tool wear
have been investigated, such as the pattern-classification
methodology, fuzzy classifiers, neural networks, and
sensor and data fusion methodology, overviewed in.20
It is well known that the structure of the workpiece
strongly affects the mechanism of chip separation and
the machinability of a job.3,6,23,24 AE techniques adopted
for discontinuous chips are less frequently reported. The
chip formation during machining of hardened steel deter-
mines the criteria for crack initiation and propagation.
Strong elastic waves related to the crack initiation and its
propagation during the segment formation can be
detected through the AE systems as well as the severe
plastic deformation at the tool–chip and tool–workpiece
interfaces. Furthermore, regarding segmentation fre-
quencies, the AE technique is able to detect high-fre-
quency processes emitted by the cutting zone. Therefore,
the suggestion of using the AE technique to monitor hard
turning operations seems to be reasonable.
Uehara25 reported on remarkable patterns in the AE
waveforms due to chip segmentation. The amplitude of
the acoustic emission varies in accordance with the
periodic change of the cutting force. Guo and Ammula26
applied AE for real-time monitoring of the surface
damage in hard machining. The results show that AE
signals can be used to monitor surface integrity; speci-
fically AErms, frequency, and count rate are sensitive to
the existence of a white layer and the corresponding tool
wear and surface roughness. Barry and Byrne27 reported
that in comparison to the AErms during continuous chip
formation, which is between 0.05 V and 0.1 V, the AErms
produced during segmented chip formation is at least one
order of magnitude greater. In addition to the single AE
sensor techniques, multi-sensing techniques were also
developed. Axinte et al.28 reported the application of the
triangulation technique to arrays of acoustic emission
sensors for the location of uneven events occurring dur-
ing machining. The specific character of chip formation
in hard turning connected with shear instability, its cyclic
character, and mixed processes in the cutting zone
(plastic deformation mixed with brittle cracking) suggest
that the AE technique should be adapted. Two sensor
techniques were suggested and verified.29
The application of two different AE sensors is con-
nected with the frequency range of the different pro-
cesses during the formation of a segmented chip. As was
observed, chip formation during hard turning exhibits
two different processes. The first is associated with crack
initiation on the free surface and its prolongation towards
the tool tip (zone Ib, Figure 2). The second type is repre-
sented by severe plastic deformation in the shear zone
(zone Ia) and at the tool–chip (zone II) and tool–work-
piece (III) interfaces. Conventional AE techniques (when
turning soft steels) are usually carried out in the fre-
quency ranges corresponding to the frequency range of
the WD sensor. However, the process of plastic deforma-
tion in the cutting zone during the machining of
hardened steel differs from a continuous chip. The
cutting process is less stable with the characteristic
segmentation frequency (accumulation of energy ahead
of the cutting edge and its sudden relaxation in the form
of a crack initiated on the free surface). As previously
reported,3,29,30 this segmentation frequency does not
usually exceed 100 kHz and so it would be difficult to
monitor the true dynamics of the cutting process via a
WD sensor. Hence, the next D9241A sensor was inte-
grated into the measuring system.
The preliminary studies3,30 proved that this suggestion
was reasonable. It was found that the peak frequency
extracted from the raw AE signal recorded by the
D9241A sensor is strongly correlated with the calculated
segmentation frequencies and allows the true dynamic,
reflecting the cracking process in the shear zone to be
detected. The application of the WD sensor does not
allow the segmentation frequencies to be detected
because the detectable frequency spectrum of the WD
sensor lies above all the segmentation frequencies. The
instability of the cutting process connected with the chip
segmentation (crack initiation and its propagation in the
shear zone) does not interfere with AE signals detected
by the WD sensor. However, while the character of the
raw AE signal and also the values of the extracted AE
features (such as AErms, AEabsolute energy and AEstrength)
obtained from the D9241A sensor stay nearly constant
with varying cutting conditions, the magnitude of the AE
signals for the WD sensor and the extracted AE features
stated above are correlated with the release of energy
accumulated ahead of the cutting edge and the
corresponding intensity of the transformation processes
in the white narrow bands and the more pronounced
thinning region of a formed chip. These preliminary
studies discussed earlier and carried out at varying
cutting speeds (from 25 m min–1 to 200 m min–1) and
feeds (from 0.051 mm to 0.27 mm) indicate the corre-
lation among the real intensity of the transformation
processes, its dynamics (based on metallographic ana-
lyses of the formed chips), and the AE signals recorded
by both AE sensors. These analyses also indicated
parameters suitable for the evaluation of the deformation
processes such as AErms, AEabsolute energy and AEstrength.
As stated above, a specific mechanism of chip
formation takes place during the hard turning. Thus, a
specific character of the emitted acoustic waves should
be expected. Furthermore, a suitable concept for the
detection of the CTF or, much better, its prediction
M. ^ILLIKOVÁ et al.: PREDICTION OF THE CATASTROPHIC TOOL FAILURE OF CERAMIC INSERTS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 355–363 357
should differ from the concept proposed for turning soft
steel. Therefore, this paper discusses a specific aspect of
AE signals during hard turning and suggests a suitable
concept for the CTF prediction.
2 EXPERIMENTAL SETUP
The experimental setup is shown in Figure 3. Two
commercial piezoelectric AE sensors (D9241A – fre-
quency range from 15 kHz to 180 kHz; WD – frequency
range from 100 kHz to 1000 kHz) from Physical
Acoustics Corporation were mounted on the top of the
tool holder (Figures 3 and 4). Semi-solid high-vacuum
grease was used to maintain a good propagation of
signals from the tool holder to the sensor. During the
experiment, the AE signals were amplified, high passed
at 15 kHz, low passed at 1000 kHz, and then sent
through a preamplifier at a gain of 40 dB to the signal-
processing software package. The signals were sampled
in real-time, amplified, digitized, and then fed to the
signal processing unit. The AE signals were post-pro-
cessed using AEwin.
The experimental study was conducted on bearing
steel 100Cr6 of hardness 62 HRc and external diameter
56 mm. The cutting and other conditions were as fol-
lows: vc = 170 m min–1, f = 0.09 mm, ap = 0.25 mm, dry
cutting, CNC Lathe Hurco TM8; the cutting tool was
TiC reinforced Al2O3 ceramic inserts DNGA150408
(TiN coating).
3 RESULTS
The specific information connected with the applica-
tion of two different sensors and their sensitivity to the
different processes in the cutting zone can be applied for
the monitoring of tool wear. Monitoring of the tool wear
for ceramic inserts, especially the prediction of cutting-
edge breakage, is a very sophisticated problem. Labora-
tory and also practical applications of these inserts
indicated relatively stable and low-intensity tool wear in
the normal phase of tool wear with the following
unexpected CTF. The intensity of the tool wear is high in
the initial and catastrophic phases of tool wear, as
illustrated in Figure 5. While micro-chipping of the
cutting edge is a dominant mechanism of tool wear in the
initial and normal phases of tool wear, CTF represents
tool wear in the form of massive breakage when the
shape and tool geometry change dramatically in a very
short time period. It is well known that the tool geometry
strongly affects the stress and the temperature distribu-
tion in the cutting zone and therefore the corresponding
chip form, as illustrated in Figures 6 and 7.
A segmented chip forms during the initial and normal
phases of tool wear, as shown in Figure 6. It forms on
M. ^ILLIKOVÁ et al.: PREDICTION OF THE CATASTROPHIC TOOL FAILURE OF CERAMIC INSERTS ...
358 Materiali in tehnologije / Materials and technology 49 (2015) 3, 355–363
Figure 5: Different phases of tool wear
Slika 5: Razli~ne faze obrabe orodja
Figure 4: Detail of the sensor placement
Slika 4: Detajl namestitve senzorjev
Figure 3: Schematic of the experimental setup
Slika 3: Shematski prikaz eksperimentalnega sestava
Figure 6: Segmented chip in the normal phase of tool wear
Slika 6: Segmentirani ostru`ki pri normalni fazi obrabe orodja
the chamfer in the form of a radius with a highly nega-
tive geometry under the high normal stresses induced
near the tool tip. In the early stages of tool wear a crater
is formed on the tool radius with a rake angle between
–7 ° and –10 °, as indicated in Figure 8. However, the
CTF usually forms a mainly positive tool geometry and
significantly alters the temperature and stress distribution
in the cutting zone and the corresponding chip formation
(Figures 7 and 9). High compressive stresses near the
tool tip can be found when the chip is produced at the
cutting-edge radius. The high compressive stresses ahead
of the cutting edge obstruct the early crack initiation. A
high portion of energy is accumulated ahead of the
cutting edge. A crack on the free surface is initiated
when the energy in front of the cutting edge attains its
ultimate value. The high energy emitted during brittle
cracking corresponds with a more pronounced region of
thinning of the produced chip (a longer cracking zone, as
illustrated in Figure 6) in the initial and also the normal
phase of tool wear. However, the chip becomes more
continuous when the tool geometry is strongly altered
due to the CTF (Figure 8). The zone of crack
propagation Ib is shorter and the thinning region is
reduced.
The character of the chip formed is strongly corre-
lated with the AE signals and the extracted features
(parameters). A conventional segmented chip is formed
in the initial and normal phases of the tool wear.
According to the theory of crack propagation and seg-
ment formation, the change in the amplitude of the
acoustic emission indicates the change in the sliding
velocity at the tool–chip interface and also different pha-
ses of the segment formation (the accumulation of ener-
gy ahead of the cutting edge and its sudden emission;
Figure 10). The observed pulses of the AE signal corres-
pond to the periodic fluctuation (relaxation character) of
the cutting process. The signal level of the AE between
these pulses is relatively small. During the segmented
M. ^ILLIKOVÁ et al.: PREDICTION OF THE CATASTROPHIC TOOL FAILURE OF CERAMIC INSERTS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 355–363 359
Figure 9: CTF of cutting insert
Slika 9: CTF rezalnega vlo`ka
Figure 7: Segmented chip in the catastrophic phase of tool wear
Slika 7: Segmentirani ostru`ki pri katastrofi~ni fazi obrabe orodja
Figure 8: Crater on the cutting-insert radius formed in the initial
phase of tool wear
Slika 8: Krater na rezalnem vlo`ku, nastal v za~etni fazi obrabe orodja
Figure 12: AE-signal in the catastrophic phase of tool wear, D9241A-
sensor
Slika 12: AE-signal pri katastrofi~ni fazi obrabe orodja, D9241A-
senzor
Figure 11: FFT-spectrum for D9241A-sensor in the normal phase of
tool wear
Slika 11: FFT-spekter za D9241A-senzor pri normalni fazi obrabe
orodja
Figure 10: AE-signal in the normal phase of tool wear, D9241A-sen-
sor
Slika 10: AE-signal pri normalni fazi obrabe orodja, D9241A-senzor
chip formation, the chip slides over the rake face with
varying speed.
Figure 10 shows this relaxing and periodic character
of the signal in the normal phase of the tool wear and the
related character of the FFT (Fast Fourier Transforma-
tion) spectrum (Figure 11) with the periodic peaks in
this spectrum (overtones frequencies of the sensor reso-
nance frequency). This character of the FFT spectrum
confirms the dominant periodic character of the recorded
signal and the ability of the D9241A AE sensor to detect
the periodic process typical for the segmented chip
formation proved by preliminary experiments.31 On the
other hand, Figure 12 shows that the character of the AE
signal in the catastrophic phase of tool wear is altered.
This signal is partially deformed and the periodic cha-
racter of the AE signal is violated. The irregular
character of the AE signal corresponds with the irregular
form of the segments produced after the CTF (Figure 7).
Moreover, the FFT spectrum of the AE signal is peak-
free (Figure 13).
Considering the WD AE sensor, all the segmentation
frequencies lie outside the frequency range of this sensor
and the relaxing character of the AE signal is missing.
The FFT spectrum of the AE signal for the WD sensor is
peak-free for all the formed chips (Figure 14). Further-
more, the transformation in the shear zone, where the
cracking zone is suppressed and the shear zone is en-
larged after the CTF, corresponds with the semi-conti-
nuous chip (Figure 15) and also the alteration in the
appearance of the AE signal (Figure 16).
The AE features extracted from the raw AE signal
plotted in the time scale are illustrated in Figures 17 to
19. The values of all the parameters for both sensors stay
nearly constant in the initial and normal phases of tool
wear. Wang and Liu32 carried out the decomposition of
the cutting force during hard turning with respect to the
chip formation and flank wear. They reported a gradual
increase of the shear and the normal force with
increasing flank wear and a gentle drop in the forces
associated with the chip formation. As the flank wear
progresses, the change in the force becomes appreciable.
Figures 17 to 19 illustrate that AErms, AEabsolute energy
and AEstrength vary only a little, despite a gradual increase
in the flank wear VB. This indicates that the processes at
the tool–workpiece interface play only a minor role
M. ^ILLIKOVÁ et al.: PREDICTION OF THE CATASTROPHIC TOOL FAILURE OF CERAMIC INSERTS ...
360 Materiali in tehnologije / Materials and technology 49 (2015) 3, 355–363
Figure 17: Influence of tool wear on RMS values of the AE
Slika 17: Vpliv obrabe orodja na vrednost RMS pri AE
Figure 18: Influence of tool wear on absolute energy of the AE
Slika 18: Vpliv obrabe orodja na absolutno energijo AE
Figure 16: AE-signal in the catastrophic phase of tool wear, WD-sen-
sor
Slika 16: AE-signal pri katastrofi~ni fazi obrabe orodja, WD-senzor
Figure 15: AE-signal in the normal phase of tool wear, WD-sensor
Slika 15: AE-signal pri normalni fazi obrabe orodja, WD-senzor
Figure 14: FFT-spectrum for WD-sensor in the normal phase of tool
wear
Slika 14: FFT-spekter za WD-senzor pri normalni fazi obrabe orodja
Figure 13: FFT-spectrum for D9241A-sensor in the catastrophic
phase of tool wear
Slika 13: FFT-spekter za D9241A-senzor pri katastrofi~ni fazi obrabe
orodja
concerning the AE features and their contribution to the
overall AE signals is low.
Since the CTF can be easily recognized (as shown in
Figures 17 to 19) the gradual increase of flank wear and
the associated process at the tool–workpiece interface are
difficult to detect using the AE technique. Furthermore,
the AE burst-type signal reflecting micro-chipping as a
dominant mechanism of tool wear in the initial and
normal phases of tool wear interferes with AE bursting
due to cracking in the shear zone. Thus, an alternative
approach to the prediction of the CTF has to be
suggested. This suggestion can be based on the corre-
lation between the chip form and the corresponding AE
signal (as well as the extracted AE features) since the
tool geometry strongly affects the stress and the tempe-
rature distribution in the cutting zone and, therefore, the
chip appearance.
As illustrated in Figures 17 and 18, a certain fall of
AErms and AEabsolute energy can be viewed at the end of the
normal phase with a subsequent abrupt increase after the
tool breakage. On the other hand, the intensity of these
changes differs between the two sensors. While a gentle
decrease in the AErms value before the CTF for the 9241A
sensor can be observed (in the time interval from 8 to 10
minutes), a more pronounced decrease is found for the
WD sensor. Furthermore, while the increase in the AErms
value after the CTF is only 21 % for the 9241A sensor, in
the case of the WD sensor the CTF leads to AErms values
that are twice as high. Thus, the idea of the signal ratio
between the features extracted from the different sensors
seems to be a suitable quantity that properly expresses
the transformation processes in the cutting zone with
respect to the tool wear.
Figure 20 indicates the transformation of the AE
features’ ratios between the sensors in connection with
the transformation of processes in the cutting zone at the
end of the normal phase of tool wear. The abrupt in-
crease of the AE features allows the tool breakage itself
to be detected. However, poor sensitivity with regard to
the CTF prediction is also visible. A reliable CTF pre-
diction can be suggested through a rationing of the
extracted features, as indicated by Equations (1) and (2)
and illustrated in Figure 20. The significant peaks occur
before the tool breakage and these peaks warn of an
imminent, approaching CTF. These peaks occur due to
the unbalanced sensitivity to the different processes in
the cutting zone between WD and D9241A, as discussed
above. Moreover, the sensitivity of this approach is
significantly higher than that based on the simple fall of
AE parameters derived from the WD sensor. The
increase in the R1 parameter in the time interval from
8 min to 10 min is about 77 % and the R2 parameter
increases two times in the same time interval. The abrupt
fall of the R1 and R2 ratios after the tool breakage is
connected with considerable changes in the tool geome-
try, the corresponding stress and temperature distribution
ahead of the cutting edge and the associated chip appear-
ance. This new suggested data processing increases the
sensitivity, not only considering the tool-failure predic-
tion, but also considering the CTF detection itself.
The different information can be obtained from
AEstrength. Figure 19 shows that the abrupt increase of
AEstrength can be found only for the low-frequency sensor,
whereas the AEstrength extracted from the high-frequency
WD sensor exhibits a remarkable drop. As was reported,
the information from WD sensors is mainly correlated
with the energy released during the segment formation
and corresponds to the length of the cracked region in
M. ^ILLIKOVÁ et al.: PREDICTION OF THE CATASTROPHIC TOOL FAILURE OF CERAMIC INSERTS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 355–363 361
Figure 19: Influence of tool wear on signal strength of the AE
Slika 19: Vpliv obrabe orodja na mo~ signala AE
Figure 21: Influence of tool wear on the ratio R3
Slika 21: Vpliv obrabe orodja na razmerje R3
Figure 20: Influence of tool wear on the ratios R1 and R2
Slika 20: Vpliv obrabe orodja na razmerje R1 in R2
the shearing zone. Being so, the feature AEstrength is a
quantity more properly corresponding with the appear-
ance of the obtained chips in the different phases of the
tool wear (Figures 6 and 7) than those expressed in
AErms and AEabsolute energy. Figure 7 shows that the
cracked region is reduced after the CTF due to an
alteration of the tool geometry (compared to the initial or
normal phase of tool wear). The AEstrength of the WD
sensor also drops down after the CTF, whereas the AErms
and AEabsolute energy abruptly increases. A different
appearance is also exhibited by the R3 ratio (indicated by
Equation (3) and illustrated in Figure 21. While nearly
constant values of this ratio can be found in the initial
and normal phases of tool wear, an abrupt increase due
to the CTF failure corresponds with transformations in
the processes ahead of the cutting edge (mainly a
reduced cracking region and the corresponding decrease
of the energy released during the brittle cracking):
R1 = AErms (D9241)/AErms (WD) (1)
R2 = AEabsolute energy (D9241)/AEabsolute energy (WD) (2)
R3 = AEstrength (D9241)/AEstrength (WD) (3)
4 CONCLUSIONS
AE techniques are very sensitive for monitoring the
specific processes in the cutting zone. The formation of a
segmented chip during hard turning is a very specific
example of different processes in the cutting zone. This
aspect indicates some conclusions connected with this
experimental study.
• The application of two AE sensors with different
frequency ranges makes it possible to detect the
different processes in the cutting zone.
• While the low-frequency AE sensor 9241A allows
the real dynamic character of chip segmentation to be
detected, the high-frequency AE sensor WD is
sensitive to the real intensity of the deformation
processes in the different zones.
• Both sensors make it possible to detect tool breakage
through the conventional parameters derived from the
AE signals, and the AE signals reflect the transfor-
mation in the formation of the chip produced.
• The sensitivity of the tool-breakage prediction
through the conventional parameters derived from the
AE signals is poor.
• The ratios of AE features such as AErms and AEabsolute
energy between sensors change at the end of the
normal phase of the tool wear. This transformation
forms peaks that occur before the tool breakage,
providing a warning of imminent tool failure.
• The ratio of AEstrength between the sensors changes
after the CTF. The AEstrength itself more properly ex-
presses the true intensity of the processes during chip
separation than the AErms and AEabsolute energy.
• The results of the experimental work presented in this
paper were verified by many similar experiments
carried out under different cutting conditions. A cer-
tain difference was found at low cutting depths, but
the signals ratio also significantly increases the sensi-
tivity for the detection of tool breakage and the pre-
diction of tool failure.
Acknowledgement
This article was edited under the financial support of
KEGA projects n. 005@U – 4/2014, 009@U – 4/2014.
5 REFERENCES
1 H. K. Tönshoff, C. Arendt, R. B. Mor, Cutting of Hardened Steel,
CIRP, 49 (2000), 547–564, doi:10.1016/S0007-8506(07)63455-6
2 S. J. Heo, Environmentally conscious hard turning of cemented car-
bide materials on the basis of micro-cutting, J. of Mech. Sci. and
Technol., 22 (2008), 1383–1390, doi:10.1007/s12206-008-0411-z
3 M. Neslu{an, Hard turning, Edis, @ilina 2010
4 Y. Lee, D. A. Dornfeld, P. K. Wright, Open Architecture Based
Framework in Integrated Precision Machining System, IMECE
Anaheim, 1998
5 M. Teti, K. Jemielniak, G. O. Donnell, D. Dornfeld, Advanced moni-
toring of machining operations, CIRP, 59 (2010), 717–739,
doi:10.1016/j.cirp.2010.05.010
6 M. C. Shaw, A. Vyas, The mechanism of chip formation with hard
turning steel, CIRP, 47 (1998), 77–82, doi:10.1016/S0007-8506(07)
62789-9
7 K. Nakayama, M. Arai, T. Kanda, Machining characteristics of
hardened steels, CIRP, 37 (1988), 89–92, doi:10.1016/S0007-
8506(07)62789-9
8 G. Poulachon, A. Moisan, Contribution to the study of the cutting
mechanism during high speed machining of hardened steel, CIRP, 47
(1998), 73–76, doi:10.1016/S0007-8506(07)62788-7
9 H. K. Tönshoff, M. Jung, S. Mannel, W. Rietz, Using acoustic
emission signals for monitoring of production processes, Ultrasonics,
37 (2000), 681–686, doi:10.1016/S0041-624X(00)00026-3
10 D. A. Dornfeld, Acoustic emission in monitoring and analysis in
manufacturing, Proceedings of AE Monitoring, 14 (1984), 124–131
11 D. A. Dornfeld, Manufacturing process monitoring and analysis
using acoustic emission, Journal of Acoustic Emission, 4 (1985),
123–126
12 I. Inasaki, Application of acoustic emission sensor for monitoring
machining processes, Ultrasonics, 36 (1998), 273–281, doi:10.1016/
S0041-24X(97)00052-8
13 J. Barry, G. Byrne, D. Lennon, Observations on chip formation and
acoustic emission in machining Ti-6Al-4V alloy, International
Journal of Machine Tools and Manufacture, 41 (2001) 7, 1055–1070,
doi:10.1016/S0890-6955(00)00096-1
14 I. Marinescu, D. A. Axinte, A critical analysis of effectiveness of
acoustic emission signals to detect tool and workpiece malfunctions
in milling operations, International Journal of Machine Tool and
Manufacture, 48 (2008), 1148–1160, doi:10.1016/j.ijmachtools.
2008.01.011
15 M. P. Gómez, A. M. Hey, J. E. Ruzzante, C. E. D´Attellis, Tool wear
evaluation in drilling by acoustic emission, Physica Procedia, 3
(2010), 819–825, doi:10.1016/j.phpro.2010.01.105
16 D. E. Lee, I. Hwang, C. M. O. Valente, J. F. G. Oliveira, D. A.
Dornfeld, Precision manufacturing process monitoring with acoustic
emission, International Journal of Machine Tool and Manufacture, 46
(2006), 176–188, doi:10.1016/j.ijmachtools.2005.04.001
17 R. Babel, P. Koshy, M. Weiss, Acoustic emission spikes at workpiece
edges in grinding: Origin and applications, International Journal of
Machine Tool and Manufacture, 63 (2013), 96–101, doi:10.1016/
j.ijmachtools.2012.08.004
M. ^ILLIKOVÁ et al.: PREDICTION OF THE CATASTROPHIC TOOL FAILURE OF CERAMIC INSERTS ...
362 Materiali in tehnologije / Materials and technology 49 (2015) 3, 355–363
18 S. Min, J. Lidde, N. Raue, D. Dornfeld, Acoustic emission based tool
contact detection for ultra-precision machining, CIRP, 60 (2011),
141–144, doi:10.1016/j.cirp.2011.03.079
19 K. Jemielniak, O. Otman, Tool failure detection based on analysis of
acoustic emission signals, Journal of Material Processing Tech-
nology, 76 (1998), 192–197, doi:10.1016/S0924-0136(97)00379-8
20 X. Li, Z. Yuan, Tool wear monitoring with wavelet packet trans-
form-fuzzy clustering method, Wear, 219 (1998), 145–154,
doi:10.1016/S0043-1648(98)00165-3
21 S. Dolin{ek, J. Kopa~, Acoustic emission signals for tool wear iden-
tification, Wear, 225–229 (1999), 295–303, doi:10.1016/S0043-
1648(98)00363-9
22 E. Govekar, J. Gradi{ek, I. Grabec, Analysis of acoustic emission
signals and monitoring of machining processes, Ultrasonics, 38
(2000), 598–603, doi:10.1016/S0041-624X(99)00126-2
23 M. Sekuli~ et al., Influence of material properties on the machinabi-
lity in face milling, Mater. Tehnol., 46 (2012) 6, 601–606
24 O. Topku, M. Übeyli, A. Acir, Effect of the martensite volume frac-
tion on the machining of a dual – phase steel using a milling
operation, Mater. Tehnol., 45 (2011) 2, 145–150
25 K. Uehara, Identification of Chip Formation Mechanism through
Acoustic Emission Measurements, CIRP, 33 (1974), 71–74,
doi:10.1016/S0007-8506(07)61382-1
26 Y. B. Guo, S. C. Ammula, Real-time acoustic emission monitoring
for surface damage in hard machining, International Journal of
Machine Tool and Manufacture, 45 (2005), 1622–1627, doi:10.1016/
j.ijmachtools.2005.02.007
27 J. Barry, G. Byrne, Chip Formation, Acoustic Emission and Surface
White Layers in Hard Machining, CIRP, 51 (2002), 65–70,
doi:10.1016/S0007-8506(07)61467-X
28 D. A. Axinte, D. R. Natarajan, N. Z. Gindy, An approach to use an
array of three acoustic emission sensors to locate uneven events in
machining – Part 1: method and validation, International Journal of
Machine Tool and Manufacture, 45 (2005), 1605–1613, doi:10.1016/
j.ijmachtools.2005.02.005
29 M. [ípek, M. Neslu{an, M. ^illiková, Analysis of segmented chip
when hard turning through acoustic emission, Manufacturing
Technology, 14 (2010), 254–257
30 M. Neslu{an, I. Mrkvica, R. ^ep, M. ^illiková, Influence of cutting
speed on intensity of the plastic deformation during hard turning,
Mater. Tehnol., 47 (2013) 6, 745–755
31 B. Karpuschewski, K. Schmidt, J. Prilukova, J. Beòo, I. Maòková, N.
T. Hieu, Influence of tool edge preparation on performance of
ceramic tool inserts when hard turning, Journal of Materials
Processing Technology, 213 (2013), 1978–1988, doi:10.1016/
j.jmatprotec.2013.05.016
32 J. Y. Wang, C. R. Liu, The effect of tool flank wear on the heat
transfer, thermal damage and cutting mechanics in finishing hard
turning, CIRP, 48 (1999), 53–56, doi:10.1016/S0007-8506(07)
63130-8
M. ^ILLIKOVÁ et al.: PREDICTION OF THE CATASTROPHIC TOOL FAILURE OF CERAMIC INSERTS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 355–363 363

M. PO[ARAC-MARKOVI] et al.: EROSIVE WEAR RESISTANCE OF SILICON CARBIDE-CORDIERITE CERAMICS ...
EROSIVE WEAR RESISTANCE OF SILICON
CARBIDE-CORDIERITE CERAMICS: INFLUENCE OF THE
CORDIERITE CONTENT
ODPORNOST KERAMIKE SILICIJEV KARBID-KORDIERIT PROTI
OBRABI PRI EROZIJI: VPLIV VSEBNOSTI KORDIERITA
Milica Po{arac-Markovi}1, Djordje Veljovi}2, Aleksandar Deve~erski1,
Branko Matovi}1, Tatjana Volkov-Husovi}2
1University of Belgrade, Vinca Institute of Nuclear Sciences, P.O. Box 522, 11001 Belgrade, Serbia
2University of Belgrade, Faculty of Technology and Metallurgy, Karnegijeva 4, POB 3503, Belgrade, Serbia
tatjana@tmf.bg.ac.rs
Prejem rokopisa – received: 2014-04-28; sprejem za objavo – accepted for publication: 2014-06-18
doi:10.17222/mit.2014.071
A cordierite/SiC composite was created in situ with reactive sintering at 1250 °C and 1300 °C. The cordierite precursor was
made from commercially available spinel, alumina and quartz and was mixed with the comercial SiC powder to obtain
composite materials during the sintering. It was found that cordierite particles bind efficiently with the SiC powder during
sintering and that reactive sintering is an effective way to produce ceramics at a relativly low temperature.
The goal of this investigation was to check the possibilities of using the silicon carbide-cordierite composite as a material
resistant to the erosive wear. The fluid dynamic system of the experimental methodology was used here to produce ultrasonic
erosive wear. Two kinds of SiC/cordierite samples were investigated, KS 50 and KS 30, with different mass contents of
cordierite (w = 50 % and w = 30 % of cordierite). The mass loss and the level of surface degradation were measured before and
during the experiment. The level of surface degradation of the samples was monitored using the Image-Pro Plus program for the
image analysis. It was found that after 150 min the mass loss was below 1.3 mg and the surface degradation was below 7 %. The
obtained results indicated that both samples exhibited an excellent erosion resistance during the cavitation experiment.
Keywords: ceramic-matrix composites (CMCs), damage tolerance, non-destructive testing, cavitation-erosion diameter and area
Kompozit kordierit-SiC je bil izdelan in-situ z reakcijskim sintranjem pri 1250 °C in 1300 °C. Kordieritna predoblika, izdelana
iz komercialno dostopnega {pinela, aluminijevega oksida in kremena, je bila zme{ana s prahom komercialnega SiC, da bi dobili
s sintranjem kompozitni material. Ugotovljeno je, da delci kordierita u~inkovito ve`ejo SiC-prah med sintranjem in da je
reakcijsko sintranje u~inkovita metoda za izdelavo keramike pri relativno nizki temperaturi.
Cilj te raziskave je preveriti mo`nost uporabe SiC-kordieritnega materiala, odpornega proti obrabi z erozijo. Za ultrazvo~no
erozijsko obrabo je bila uporabljena eksperimentalna metoda fluidnega dinami~nega sistema. Preiskovani sta bili dve vrsti
vzorcev SiC-kordierit, KS 50 in KS 30, z razli~nima masnima vsebnostima kordierita (w = 50 % in w = 30 % kordierita). Masni
dele` obrabe povr{ine je bil izmerjen pred preizkusom in po njem. Stopnja degradacije povr{ine vzorcev je bila ugotovljena s
programom za analizo slik Image Pro Plus. Ugotovljeno je bilo, da je bila izgube mase po 150 min manj{a od 1,3 mg in
degradacija povr{ine manj{a od 7 %. Dobljeni rezultati ka`ejo, da imata med preizkusom kavitacije oba vzorca odli~no
odpornost proti eroziji.
Klju~ne besede: kompozit s kerami~no osnovo (CMCs), toleranca po{kodb, neporu{ne preiskave, premer in povr{ina erozije pri
kavitaciji
1 INTRODUCTION
Ceramic materials have been used for centuries as the
materials with a very wide range of mechanical and ther-
mal properties suitable for various applications. Today,
their common industrial use is more related to the elec-
tronics for mechanical parts and biomedicine applica-
tions (hip prosthesis, dental implants). One of the main
reasons for this is the improvement of the fracture tough-
ness that allows a ceramic material to perform better
when it is subjected to the operating conditions. Some-
times these operating conditions may involve rapid
changes in the temperature, or a high-temperature appli-
cation as well as the conditions that include cavitation
erosion. This could be related to the technical ceramics,
for example, the operations of bearings, injectors or
valves. Cavitation, i.e., the appearance of vapor cavities
inside an initially homogeneous liquid medium, occurs
in very different situations. Hence, the study of the cavi-
tation and cavitation-erosion mechanisms of technical
ceramics is of importance to improve their performance
in real applications.1–4
Metallic materials are the most common choice when
cavitation-erosion resistance is required. Research results
concerning the applications of different classes of mate-
rials, including ceramics and composite materials for a
similar use, were published in the last decade.1–10 Diffe-
rent types of ceramics based on silicon nitride and zir-
conia, as well as alumina-based ceramic materials were
investigated in the conditions of cavitation erosion.2,4–9 In
numerous papers1–12 related to the investigations of the
erosion rate of these materials, attempts were made to
Materiali in tehnologije / Materials and technology 49 (2015) 3, 365–370 365
UDK 666.3/.7:620.179.1:620.193.16 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)365(2015)
investigate the type of cracking as well as the influence
of the grain sizes and the contents of different com-
pounds and phases on the resistance of the materials to
the erosive wear.
The goal of our investigation was to study the possi-
bilities of using a silicon carbide/cordierite based cera-
mic material as a potential cavitation-resistant material.
2 MATERIALS
The properties of cordierite/silicon carbide based
ceramics that enable a widespread use in high-tempe-
rature conditions are a particularly low coefficient of
thermal expansion, a high thermal conductivity, an
excellent thermal fire resistance and an excellent
resistance to a thermal shock. Their target application is
in the furnaces for the use at the temperatures of over
1000 °C. An addition of silicon carbide improves the
physical and mechanical properties of refractory cor-
dierite. The high thermal conductivity of silicon carbide
reduces the thermal stresses within a cordierite ceramic
body. A mixture of commercially available spinel
(MgAl2O4), quartz (SiO2) and alumina (Al2O3) corres-
ponding to a cordierite stoichiometric composition was
attrition milled using Al2O3 balls and ethyl alcohol as the
media for four hours; henceforth, this mixture was
labeled as KS. Mixture KS was used for preparing the
cordierite/SiC composite ceramics with the mass ratio of
30 : 70 and 50 : 50, respectively. After the milling with
the Al2O3 balls in DI water in a polyethylene bottle for 24
h, the samples were uniaxially pressed and sintered at
1300 °C and 1250 °C, for 3 h, and the ceramic composite
samples were marked as KS 30 and KS 50, respectively.
The typical values of the selected properties of the dense
constituents used in the refractory materials investigated
are listed in the previous papers.13–15
An XRD analysis of the KS 30 sample (Figure 1)
shows a presence of three crystalline phases: SiC, cor-
dierite and sapphirine and a presence of a certain amount
of a glassy phase. In the case of sample KS 50, the phase
analysis shows a presence of four crystalline phases:
SiC, cordierite, sapphirine and enstatite. There is also a
small amount of a glassy phase, which affects the for-
mation of cordierite and sapphirine. One can conclude
that an increase in the added amount of the mixture with
a composition corresponding to stoichiometric cordierite
improves the quantity and crystallization of cordierite.
The microstructures of samples KS 30 and KS 50
(Figure 2) were examined with scanning electron micro-
M. PO[ARAC-MARKOVI] et al.: EROSIVE WEAR RESISTANCE OF SILICON CARBIDE-CORDIERITE CERAMICS ...
366 Materiali in tehnologije / Materials and technology 49 (2015) 3, 365–370
Figure 2: SEM of the: a) KS 30 and b) KS 50 samples
Slika 2: SEM-posnetka vzorcev: a) KS 30 in b) KS 50
Figure 1: XRD of the: a) KS 30 and b) KS 50 samples
Slika 1: XRD-posnetka vzorcev: a) KS 30 in b) KS 50
scopy using a VEGA TS 5130 mm (TESCAN) device;
the samples were coated with a layer of the Au-Pd mix-
ture. A porous structure was observed, with the particles
of different sizes and pronounced necks formed during
the densification. The pores that are not spherical indi-
cate the initial stages of sintering. The pores present have
different sizes and shapes.
3 CAVITATION-EROSION TESTING
The experimental methodology used for the cavi-
tation-erosion testing is explained in7,10,16,17. With respect
to the used equipment, the diameter of the horn was 10
mm and the distance between the horn and a sample was
1 mm. The samples were discs with a diameter of 3 cm
and a height of 1 cm.
4 RESULTS AND DISCUSSION
4.1 Mass loss and level of destruction during the
testing
The mass losses of the test specimens were deter-
mined using the analytical balance with an accuracy of
± 0.1 mg. The measurements were performed after
subjecting each test specimen to cavitation for 30 min.
The duration of the tests was 150 min. The light micro-
scopy technique was applied to analyze the effect of the
erosion and to interpret the results of the cavitation tests
(Figure 3).
The mass loss was below 1.3 mg, which is similar to
metallic15–25 and ceramic5,6,24,25 materials. From Figure 2
it can be seen that sample KS 30 exhibited a better resis-
tance to erosive wear, as after 150 min the mass loss for
KS 30 was 0.12 mg and for KS 50 it was 1.23 mg.
The next experiment consisted of the measurements
of the degradation levels before and during the testing. In
order to carry out an image analysis determining the
level of destruction before and during the testing, the
samples were photographed as shown in Figure 4.
The results of the image analysis of samples KS 50
and KS 30 are given in Figure 5.
Based on the obtained results, similar conclusions
regarding the level of degradation and the mass loss of
the samples can be given. Samples KS 50 exhibited a
better erosion resistance as their level of damage was
4.4 % compared to 6.9 % for the KS 30 samples. The
differences in the sample behavior can be related to the
crack nucleation and propagation during the cavitation
testing.
As can be seen from Figure 5, samples KS 50 exhi-
bited a stable increase in the level of degradation and
similar rates of degradation for the entire time range. At
M. PO[ARAC-MARKOVI] et al.: EROSIVE WEAR RESISTANCE OF SILICON CARBIDE-CORDIERITE CERAMICS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 365–370 367
Figure 4: Photographs of the samples during the testing
Slika 4: Posnetki vzorcev med preizkusi
Figure 3: Mass loss during the experiment
Slika 3: Izguba mase med preizkusom
Figure 5: Level of destruction of the samples (P/Po) versus the time
of experiment (P is the damaged area after the testing and Po is the
undamaged area)
Slika 5: Nivo propada vzorcev (P/Po) v odvisnosti od ~asa preizkusa,
(P je po{kodovano podro~je po preizkusu in Po je nepo{kodovano
podro~je)
the end of the experiment, after 150 min, the level of
degradation was below 5 %.
For the KS 30 samples, two regimes can be observed
on Figure 5. The first regime is characterized by a slow
increase in the level of destruction over the period from 0
min to 60 min. The second regime, after 60 min of the
testing, is characterized by a more rapid change in the
damage level. However, it has to be mentioned that the
total level of degradation, after 150 min of the testing,
did not exceed 8 %, which is an excellent result for the
cavitation resistance.
From these results it can be concluded that samples
KS 30 have a better resistance to crack nucleation and
propagation up to the 60th minute of the experiment.
After this critical time, the crack propagation is very
rapid. The increased cordierite amount and decreased
SiC amount influenced the crack nucleation and propa-
gation, as the higher values of cordierite in samples KS
50 caused lower rates of crack propagation during the
entire time (0–150 min) and the total degradation of the
samples after 150 min of the testing was lower.
4.2 Erosion-ring diameter and erosion-ring area
The samples exposed to cavitation erosion (Figure 3)
were also monitored in order to measure the diameters
and the areas of the erosion-ring regions. These measure-
ments were performed using the following steps:
Step 1: Making appropriate micrographs that include the
erosion ring areas (Figure 3)
Step 2: Measuring the ring diameters: at this point two
types of measurements were taken: diameters (d1)
and (d2) of the cavitation-ring erosion measured with
the Image-Pro Plus program. The effective diameter
of each ring is calculated according to the following
equation:
d
d d
=
+1 2
2
(1)
Step 3: Calculating the effective area of the erosion ring
using the values of the diameters measured in Step 2:
P
d d
av =
π 1 2
4
(2)
Step 4: Using the Image-Pro Plus program to determine
the average erosion-surface area of the ring, (Pmeas).
4.3 Erosion-ring measuremets
The measurements of the average diameter of the
erosion ring were made after (30, 45, 60, 90, 120 and
150) min.
Samples KS 50: The results presented in Figure 6
show that the formation of an erosion ring can be ob-
served after 45 min of the testing. The average erosion-
ring diameter increased from 5.68 mm after 45 min to
6.03 mm after 150 min of the testing.
Samples KS 30: The errosion-ring formation of this
sample was observed after 30 min. A rapid increase in
the ring diameter was observed after 90 min, when the
ring diameter increased from 3.29 mm after 60 min to
7.04 mm after 90 min of the testing. After 90 min the
increase in the ring diameter was very slow so that after
150 min the ring diameter reached only 7.41 mm.
The results for the erosion-ring measuremet are in a
strong correlation with the results for the level of des-
truction (Figure 5) supporting the conclusions about the
influence of cordierite and SiC content on the level of
degraration. The same conlusions are valid if the ero-
sion-ring measurements are taken into account. With
respect to the energy for crack nucleation, lower values
are expected for the KS 30 samples where the formation
of the erosion ring was observed after 30 min, while for
the KS 50 samples it was formed after 45 min. This
difference could be important for some specific applica-
M. PO[ARAC-MARKOVI] et al.: EROSIVE WEAR RESISTANCE OF SILICON CARBIDE-CORDIERITE CERAMICS ...
368 Materiali in tehnologije / Materials and technology 49 (2015) 3, 365–370
Figure 7: Average area of the erosion ring based on diameter measu-
rements and area measurements: a) KS 50 and b) KS 30
Slika 7: Povpre~na povr{ina erozijskega kroga na podlagi meritev pre-
mera in povr{ine: a) KS 50 in b) KS 30
Figure 6: Average diameter of the erosion ring during the testing
Slika 6: Povpre~ni premer erozijskega kroga med preizku{anjem
tions where the specific resistance to either crack
nucleation or crack propagation is requested.
4.4 Average erosion-surface area
According to the procedure described above, Figure
7 shows the values obtained for the average erosion-sur-
face areas based on monitoring the erosion-ring dia-
meters.
Samples KS 50: The formation of the erosion ring
was detected after 45 minutes and the average erosion-
surface area was Pmeas = 20.73 mm2 and Pav = 25.23 mm2.
Both values for the erosion ring area were slowly
increased, reaching the values of Pav = 28.35 mm2 and
Pmeas = 33.42 mm2 after 150 min of the testing. The
results showed the differences caused by implementing
the method for the erosion-ring-area determination
(during the above steps) and these differences apply to
the whole experiment.
Samples KS 30: As for the ring-diameter measure-
ments, after 60 min of the testing the erosion ring was
detected, with the erosion-surface area of Pmeas = 10.56
mm2 and Pav = 8.49 mm2. After 90 min a rapid change in
the erosion area was observed so that after 150 min of
the testing the erosion-surface values raised to Pav =
43.76 mm2 and Pmeas = 52.81 mm2. The differences in the
values obtained with step 4 (Pav) and step 5 (Pmeas) were
increasing during the experiment.
Based on the obtained results it can be concluded that
samples KS 30 are more sensitive to the formation of the
erosion ring. Also, after 90 min of the testing there are
rapid changes in the diameter and the surface of the
erosion ring, but by the end of the testing (150 min) this
increase is slowed down.
Samples KS 50 exhibited a better erosion resistance
as the ring diameter and the average area of erosion were
lower, and their increase was slower over the testing
time.
The formation of the erosion ring for the KS 50 sam-
ples is visible after 30 min. This can be related to the
influence of the crack nucleation caused by the erosion
experiment.
5 CONCLUSION
Ceramic composite samples based on the SiC/cordi-
erite ceramic material were synthesized in order to inve-
stigate the sample resistance to the erosive wear. Two
samples with different mass contents of cordierite and
SiC were used in this investigation: KS 50 (w = 50 % of
cordierite) and KS 30 (w = 30 % of cordierite).
The results for the mass loss as well as for the level
of degradation indicate a similar (or even better) erosion
resistance compared to the metallic and low-alumina
samples.
Both samples exhibited an excellent erosive resis-
tance, but sample KS 50 exhibited the better resistance
of the two as its mass loss was lower, as were the level of
degradation, the erosion-ring diameter and the erosion
area.
These experiments showed that composite ceramic
materials based on SiC/cordierite can be used as
erosion-resistant materials, used for new applications.
Also, an implementation of non-destructive testing
such as an image analysis to determine the erosion-ring
diameter and the erosion-ring area improved the
reliability of predicting the sample behavior in the
conditions of erosive wear.
Acknowledgement
This research was financed by the Ministry of
Education, Science and Technological Development of
Serbia as part of project III 45012.
6 REFERENCES
1 W. J. Tomlinson, S. J. Matthews, Cavitation erosion of structural
ceramics, Ceramics International, 20 (1994), 201–209, doi:10.1016/
0272-8842(94)90040-X
2 U. Litzow, K. H. Z. Gahr, J. Schneider, Cavitation erosion of ad-
vanced ceramics in water, International Journal of Materials
Research: Zeitschrift für Metallkunde, 97 (2006), 1372–1377,
doi:10.3139/146.101380
3 D. Niebuhr, Cavitation erosion behavior of ceramics in aqueous solu-
tions, Wear, 263 (2007), 295–300, doi:10.1016/j.wear.2006.12.040
4 G. Garcia-Atance Fatjó, M. Hadfield, K. Tabeshfar, Pseudoplastic
deformation pits on polished ceramics due to cavitation erosion,
Ceramics International, 37 (2011), 1919–1927, doi:10.1016/
j.ceramint.2011.03.043
5 S. Martinovic, M. Dojcinovic, J. Majstorovic, A. Devecerski, B.
Matovic, T. Volkov Husovic, Implementation of image analysis on
thermal shock and cavitation resistance testing of refractory concrete,
Journal of European Ceramic Society, 30 (2010), 3303–3309,
doi:10.1016/j.jeurceramsoc.2010.07.041
6 S. Martinovic, M. Vlahovic, M. Dojcinovic, T. Volkov-Husovic, J.
Majstorovic, Thermomechanical properties and cavitation resistance
of a high-alumina low-cement castable, International Journal of
Applied Ceramic Technology, 8 (2011) 5, 1115–1124, doi:10.1111/
j.1744-7402.2010.02545.x
7 G. Garcia-Atance Fatjó, M. Hadfield, C. Vieillard, J. Sekulic, Early
stage cavitation erosion within ceramics—an experimental investi-
gation, Ceramics International, 35 (2009), 3301–3312, doi:10.1016/
j.ceramint.2009.05.020
8 M. H. Korkut, Y. Küçük, A. C. Karaoðlanli, A. Erdoðan,Y. Er, M. S.
Gök, Effect of the abrasive grit size on the wear behavior of ceramic
coatings during micro-abrasion test, Mater. Tehnol., 47 (2013) 6,
695–699
9 J. Ba`an, L. Socha, L. Martínek, P. Fila, M. Balcar, J. Chmelaø, Wear
of refractory materials for ceramic filters of different porosity in
contact with hot metal, Mater. Tehnol., 45 (2011) 6, 603–608
10 G. Garcýá-Atance Fatjó, A. Torres Pérez, M. Hadfield, Experimental
study and analytical model of the cavitation ring region with small
diameter ultrasonic horn, Ultrasonics Sonochemistry, 18 (2011),
73–79, doi:10.1016/j.ultsonch.2009.12.006
11 F. G. Hammit, Cavitation and Multiphase Flow Phenomena,
McGraw-Hill, New York 1980
12 R. T. Knapp, J. W. Daily, F. G. Hammit, Cavitation, McGraw-Hill,
New York 1970
13 M. Posarac, M. Dimitrijevic, T. Volkov-Husovic, A. Devecerski, B.
Matovic, Determination of thermal shock resistance of silicon car-
bide/cordierite composite material using nondestructive test methods,
Journal of the European Ceramic Society, 28 (2008) 6, 1275–1278,
doi:10.1016/j.jeurceramsoc.2007.09.038
M. PO[ARAC-MARKOVI] et al.: EROSIVE WEAR RESISTANCE OF SILICON CARBIDE-CORDIERITE CERAMICS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 365–370 369
14 M. Dimitrijevic, M. Posarac, J. Majstorovic, T. Volkov-Husovic, B.
Matovic, Behavior of silicon carbide/cordierite composite material
after cyclic thermal shock, Ceramics International, 35 (2009) 3,
1077–1081, doi:10.1016/j.ceramint.2008.04.029
15 M. Posarac-Markovic, Ph.D. Thesis, University of Belgrade, Faculty
of Technology and Metallurgy (in Serbian)
16 T. Okada, S. Hattori, M. Shimizu, A fundamental study of cavitation
erosion using a magnesium oxide single crystal (intensity and
distribution of bubble collapse impact loads), Wear, 186–187 (1995),
437–443, doi:10.1016/0043-1648(95)07162-8
17 S. Hattori, H. Mori, T. Okada, J. Fluid Eng, Trans. ASME, 120
(1998), 179–185, doi:10.1115/1.2819644
18 T. Okada, S. Hattori, Proc. of the Int. Symposium on Aerospace and
Fluid Science, Sendai, Japan, 1993, 347
19 K. Steller, Proc. of the 6th Int. Conf. on Erosion by Liquid and Solid
Impact, Cambridge, UK, 1983, 121
20 G. Bregliozzia, A. Di Schinob, S. I. U. Ahmeda, J. M. Kennyb, H.
Haefkea, Wear, 258 (2005), 503–510, doi:10.1016/j.wear.2004.
03.024
21 W. J. Tomlinson, A. S. Bransden, Wear, 185 (1995), 59–65,
doi:10.1016/0043-1648(94)06584-5
22 C. J. Lin, J. L. He, Wear, 259 (2005), 154–159, doi:10.1016/j.wear.
2005.02.099
23 M. Dojcinovic, T. Volkov Husovic, Cavitation damage of the me-
dium carbon steel: impelmentation of image anaysis, Materials
Letters, 62 (2008), 953–956, doi:10.1016/j.matlet.2007.07.019
24 S. Martinovic, M. Dojcinovic, J. Majstorovic, A. Devecerski, B.
Matovic, T. Volkov Husovic, Implementation of image analysis on
thermal shock and cavitation resistance testing of refractory concrete,
Journal of European Ceramic Society, 30 (2010), 3303–3309,
doi:10.1016/j.jeurceramsoc.2010.07.041
25 S. Martinovic, M. Vlahovic, M. Dojcinovic, T. Volkov-Husovic, J.
Majstorovic, Thermomechanical properties and cavitation resistance
of a high-alumina low-cement castable, International Journal of
Applied Ceramic Technology, 8 (2011) 5, 1115–1124, doi:10.1111/
j.1744-7402.2010.02545.x
M. PO[ARAC-MARKOVI] et al.: EROSIVE WEAR RESISTANCE OF SILICON CARBIDE-CORDIERITE CERAMICS ...
370 Materiali in tehnologije / Materials and technology 49 (2015) 3, 365–370
Y. VIJAYAKUMAR et al.: INFLUENCE OF THE SUBSTRATE TEMPERATURE ON THE STRUCTURAL, OPTICAL ...
INFLUENCE OF THE SUBSTRATE TEMPERATURE ON THE
STRUCTURAL, OPTICAL AND THERMOELECTRIC PROPERTIES
OF SPRAYED V2O5 THIN FILMS
VPLIV TEMPERATURE PODLAGE NA STRUKTURNE, OPTI^NE
IN TERMOELEKTRI^NE LASTNOSTI NAPR[ENE TANKE PLASTI
V2O5
Yelsani Vijayakumar1, Katta Narasimha Reddy1, Annasaheb Vitthal Moholkar2,
Musugu Venkata Ramana Reddy1
1Thin Films and Nanomaterials Research Laboratory, Department of Physics, Osmania University, 500007 Hyderabad, India
2Thin Film Nanomaterials Laboratory, Department of Physics, Shivaji University, Kolhapur, India
vijay.yelsani@gmail.com
Prejem rokopisa – received: 2014-05-13; sprejem za objavo – accepted for publication: 2014-07-11
doi:10.17222/mit.2014.079
Vanadium pentoxide (V2O5) thin films were deposited using the spray pyrolysis technique. An aqueous solution of ammonium
vanadate with a 0.05 M concentration was used for depositing V2O5 thin films at three different substrate temperatures on glass
substrates. The structural and optical characteristics of the V2O5 thin films were examined with X-ray diffraction (XRD) and
double-beam UV-visible spectrophotometry. The X-ray diffraction study of the V2O5 thin films revealed a polycrystalline nature
of the orthorhombic structure with the preferred orientation of (001). The crystallite size (d) was calculated from the (001)
diffraction peak using the Debye-Scherrer formula. From the optical absorbance measurements, the optical band gap (Eg) was
determined. A scanning electron microscope (SEM) was used to characterize the morphology of the films. Electrical measure-
ments of the films indicated that the resistance decreases with an increase in the substrate temperature. From the thermoelectric
measurements, the Seebeck coefficient was determined.
Keywords: V2O5 thin film, spray pyrolysis, optical band gap, activation energy, temperature coefficient of resistance, Seebeck
coefficient
Tanka plast vanadijevega pentoksida (V2O5) je bila nanesena s tehniko piroliznega brizganja. Za nanos tanke plasti V2O5 na
podlago iz stekla pri treh razli~nih temperaturah podlage je bila uporabljena koncentracija vodne raztopine amonijevega
vanadata 0,05 M. Zna~ilnosti strukture in opti~ne zna~ilnosti tanke plastiV2O5 so bile preiskovane z rentgensko difrakcijo in z
dvo`arkovno UV-vidno spektrofotometrijo. Rentgenska difrakcija tanke plasti V2O5 je odkrila polikristalno naravo ortorombi~ne
strukture s prednostno orientacijo (001). Velikost kristalitov (d) je bila izra~unana iz difrakcijskega vrha (001) z
Debye-Scherrerjevo formulo. Iz meritev opti~ne absorbance je bila dolo~ena pasovna vrzel (Eg). Za karakterizacijo morfologije
plasti je bil uporabljen vrsti~ni elektronski mikroskop (SEM). Elektri~ne meritve tankih plasti so pokazale, da se upornost
zmanj{uje z nara{~anjem temperature podlage. Iz termoelektri~nih meritev je bil dolo~en Seebeckov koeficient.
Klju~ne besede: tanka plast V2O5, pirolizno brizganje, opti~na pasovna vrzel, aktivacijska energija, temperaturni koeficient
upornosti, Seebeckov koeficient
1 INTRODUCTION
Vanadium oxide is of enormous research interest
because of its multivalent nature. The VO2, V2O3 and
V2O5 multivalent oxides exhibit a lot of fascinating and
novel properties. Among these vanadium pentoxide
(V2O5) has been extensively studied and because of its
highest oxidation state in the V – O system, a wide band
gap, a better stability and its electrothermal effects it is
useful for device applications. V2O5 is used in various
devices, such as color filters, smart windows1 and infra-
red detectors,2 as well as gas sensing3 and catalysis.4
Vanadium pentoxide thin films are prepared with
different physical and chemical techniques, namely, ther-
mal evaporation,5 pulsed-laser deposition,6 sputtering,7
inorganic sol-gel method8 and spray pyrolysis.9 Being
simple and less expensive, the spray-pyrolysis technique
(SPT) is a better chemical technique, carried out at a
lower cost, for the preparation of thin films with a larger
area. In addition, it provides an easy way to dope any
element in the ratio of a required proportion through the
solution medium. This method is convenient for pre-
paring pinhole-free, uniform thin films with the required
thickness.10 In the spray-pyrolysis technique, various
deposition parameters like the compressed-air pressure,
the spray rate, the substrate temperature, the distance
between the nozzles and the substrate and the cooling
rate after deposition also affect the physical, electrical
and optical properties of thin films.11 However, few
efforts have been made to systematically investigate the
effects of deposition parameters on the structural,
electrical and optical properties of the vanadium oxide
thin films deposited with SPT.12
In the present investigation, a synthesis of V2O5 made
with the spray-pyrolysis technique was investigated at
Materiali in tehnologije / Materials and technology 49 (2015) 3, 371–376 371
UDK 532.6 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)371(2015)
low substrate temperatures, and the structural, optical
and thermoelectric properties of the films are reported.
2 MATERIAL AND METHODS
Before depositing the V2O5 thin films, the glass
substrates were cut into 2.25 cm × 2.25 cm pieces and
subjected to cleaning and degreasing protocols. 0.05 M
concentrated ammonium vanadate with deionized water
was used as the starting material. The spray solution was
introduced into the air stream by means of a syringe
pump. In the spray system, compressed and purified air
was used as the carrier gas with a 3 kg/cm2 pressure and
the solution spray rate was maintained at 3 mL/min. The
distance between the spray nozzle and the substrate was
fixed at 25 cm. The spray head moved in the horizontal
plane due to a stepper motor to achieve a uniform depo-
sition of the films on the heated substrates maintained at
different temperatures, i.e., (250, 300 and 350) °C. The
substrate temperature was controlled through a digital
temperature controller with an accuracy of ±5 °C.
The crystal structures of the films were studied with
XRD using a Philips Xpert diffractometer with Cu-K
radiation (the X-ray wavelength  = 0.154 nm). Micro-
photography of the films was carried out using a
scanning electron microscope. Optical parameters were
calculated from the absorption spectra recorded against
the wavelength using a Lab India UV-Visible 3000 spec-
trophotometer. The resistance of the films was measured
with the two-point probe method using a Keithley
electrometer (model no. 196) in the temperature range of
27–100 °C. The metallic contacts on the films were
made of silver paint. Thermoelectric power [TEP] was
measured using a home-built system with two copper
blocks, one for the heat source and the other one for the
heat sink to create a temperature gradient and produce
the Seebeck voltage. The whole apparatus was kept in an
enclosure to minimize the air-current disturbances. The
temperature of the hot junction was raised slowly and the
thermo e.m.f. was noted at regular intervals of 5 °C. The
thermo e.m.f. was measured with a Keithley nanovolt-
meter (model no.181).
3 RESULTS AND DISCUSSION
3.1 Structural properties
The X-ray diffraction (XRD) patterns of V2O5 thin
films at different substrate temperatures are shown in
Figure 1. The peaks obtained in the XRD pattern match
the peaks in JCPDS # 89-2482, corresponding to the
orthorhombic V2O5 phase with the lattice-parameter va-
lues of a = 1.154 nm, b = 0.3571 nm and c = 0.4383 nm.
The V2O5 phase formation starts on the films deposited
at the substrate temperature of 250 °C with the (001)
reflection and this reflection was more dominant with the
film deposited at 300 °C. The XRD patterns suggest that
the texture of a V2O5 thin film is oriented along the
c-axis and, on a further increase in the substrate tempe-
rature, up to 350 °C, other reflections – (200), (301) –
also appear. The orthorhombic V2O5 phase is in agree-
ment with the earlier reports on the V2O5 thin films
deposited with the spray pyrolysis and also with other
methods.13,14 The crystallite size of the films was
estimated with the Debye-Scherrer formula for the (001)
reflection:
d =
0 94.
cos

 
(1)
where d is the crystallite size,  is the X-ray wavelength
(0.154 nm),  is the full-width half maximum and  is
the Bragg diffraction angle in degrees.
The variation in the crystallite size with the substrate
temperature is summarized in Table 1. The results show
that the crystallite size varies from 67 nm for the film
deposited at 250 °C to 84 nm for the film deposited at
300 °C and it further changes to 74 nm for the film
deposited at 350 °C. It can also be observed that the
crystallite size increases with the substrate temperature
varying from 250 °C to 300 °C. This could be associated
with the coalescence process being favored in this
temperature range, leading to an increase in the crystal-
Y. VIJAYAKUMAR et al.: INFLUENCE OF THE SUBSTRATE TEMPERATURE ON THE STRUCTURAL, OPTICAL ...
372 Materiali in tehnologije / Materials and technology 49 (2015) 3, 371–376
Figure 1: X-ray diffraction patterns of V2O5 films at different sub-
strate temperatures
Slika 1: Posnetek rentgenske difrakcije tanke plasti V2O5 pri razli~nih
temperaturah podlage
Table 1: Variation in the crystallite size, optical band gap, activation
energy and Seebeck coefficient with the substrate temperature
Tabela 1: Spreminjanje velikosti kristalitov, opti~ne pasovne vrzeli in
Seebeckovega koeficienta s temperaturo podlage
Substrate
temperature
Crystallite
size d/nm
Optical
band gap
Eg/eV
Activation
energy
Ea/eV
Seebeck
coefficient
S/(μV/K)
250 °C 67 2.34 0.15 –70
300 °C 84 2.29 0.13 –66
350 °C 74 2.21 0.12 –65
lite size. A further increase in the substrate temperature
leads to a decrease in the crystallite size which may be
due to the re-crystallization of the material.
The dislocation density ( ) is described as the length
of dislocation lines per unit volume of the crystal. The
dislocation density ( ) of the crystal gives information
about the crystal structure. The dislocation density for
the preferential orientation can be calculated using the
formula below:15
 =
1
2d
(2)
where d is the crystallite size. The dislocation density
obtained from Equation (2) for various crystallite sizes
is found to be 15 · 10–3 nm–2, 12 · 10–3 nm–2 and 13 ·
10–3 nm–2. It can be concluded from the above results
that the smaller the dislocation density the better is the
crystallization of the film.
Figure 2 presents the SEM images of the films depo-
sited at different substrate temperatures. Similar results
were observed for MoO3 thin films.16
It can be seen from the SEM images and analyses of
the topographical profiles that the surfaces of the films
grown at a substrate temperature of 300 °C clearly grew
like a sponge-type structure with macropores.
3.2 Optical properties
The impact of the substrate temperature on the opti-
cal energy-gap (Eg) values was investigated with the
optical-absorbance measurements. The absorbance spec-
tra of the films deposited at different substrate tempera-
tures are presented in Figure 3. By increasing the sub-
strate temperature, the absorbance of the films was
increased. The optical absorption coefficient  was
estimated with the following relation:
=
A
t
(3)
where t is the film thickness and A is the absorbance.
According to the interband absorption theory, the
optical band gap (Eg) of the films was calculated using
the following relation:
hv B hv E= −( )g
m (4)
where B is the probability parameter for the transition,
Eg is the optical band gap of the material, hv is the inci-
dent photon energy, and m is the transition coefficient.
Y. VIJAYAKUMAR et al.: INFLUENCE OF THE SUBSTRATE TEMPERATURE ON THE STRUCTURAL, OPTICAL ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 371–376 373
Figure 2: SEM images of V2O5 films at different substrate temperatures: a) 250 °C, b) 300 °C, c) 350 °C
Slika 2: SEM-posnetki tanke plasti V2O5 pri razli~nih temperaturah podlage: a) 250 °C, b) 300 °C, c) 350 °C
Figure 4: (hv)2 versus hv at different substrate temperatures
Slika 4: Odvisnost (hv)2 od hv pri razli~nih temperaturah podlage
Figure 3: Absorbance versus wavelength at different substrate tempe-
ratures
Slika 3: Odvisnost absorbance od valovne dol`ine pri razli~nih tempe-
raturah podlage
The value of m was taken as 1/2 for direct transitions,
3/2 for direct forbidden transitions, 2 for indirect tran-
sitions and 3 for indirect forbidden transitions.17
The plotting of (hv)1/m versus the photon energy (hv)
and extrapolating it to (hv)1/m = 0 gives the value of Eg.
Figure 4 shows the plots of (hv)2 versus hv for the V2O5
films deposited at different substrate temperatures. The
results obey the above equation with m = 1/2 indicating a
direct transition. The calculated values of the optical
band gap Eg were found to be (2.34, 2.29 and 2.21) eV
for the films deposited at (250, 300 and 350) °C, respec-
tively. These values for the V2O5 thin films investigated
in the present study are consistent with the values
reported in14. The decrease in the optical band gap is
attributed to the microstructural changes caused by a
high substrate temperature. At high temperatures the
interatomic distance decreases, leading to a decrease in
the localized states in the conduction and valance bands.
3.3 Thermoelectric properties
Figure 5 presents the room-temperature resistances
of the films deposited at different substrate temperatures
(Ts). It was found that the room-temperature surface resi-
stance decreased from 180 k to 50 k as the substrate
temperature increased from 250 °C to 350 °C. We main-
tain that the resistances are related to the microstructures
of the films, which strongly depend on the substrate tem-
perature. The film growth is directly related to the diffu-
sion of atoms into the substrates:18
D D
E
kT
= −
⎛
⎝
⎜ ⎞
⎠
⎟
0 exp
d
(5)
Here D is the surface-atomic-diffusion coefficient, Ed
is the activation energy (atom), k is the Boltzmann con-
stant and T is the absolute temperature. In the film
growth mechanism, it is evident that the resistance of the
films is dependent on the substrate temperature which
can be expressed with Equation (5) supported with Fig-
ure 5. At lower temperatures, the atoms may not have
the sufficient energy for the atomic-jump process to
overcome the potential energy of the nucleation sites of
the substrate. At higher substrate temperatures, the mobi-
lity of atoms on the substrate surface is generally higher.
As a result, the diffusion distance of atoms on the surface
increases and the collision process initiates the
nucleation for more atoms joined together, resulting in a
decrease in the room-temperature resistance. To further
investigate the electrical properties of the films, we
measured the resistance and the temperature coefficient
of resistance (TCR) of the films. The variation in the
resistance with different substrate temperatures is plotted
in Figure 6. These plots are in good agreement with the
thermal-activation mechanism evaluated with the
following relation:19
R R
E
kT
=
⎛
⎝
⎜ ⎞
⎠
⎟
0 exp
a
(6)
where R is the resistance, R0 is the constant, Ea is the
activation energy, k is the Boltzmann constant and T is
the absolute temperature. The values of Ea derived
according to the plots shown in Figure 7, using
Equation (6) are (0.15, 0.13 and 0.12) eV. These results
indicate that electrons need less activation energy to
jump from a vanadium site to another one with an
increased temperature as the substrate temperature
increased. They also explain that in the case of a film
grown at a higher substrate temperature, atoms are clo-
sely packed, there are fewer defects and, hence, the
hopping energy related to the thermally assisted tunnel-
ing process decreases.18 The temperature coefficient of
resistance (TCR) can be calculated as:20
Y. VIJAYAKUMAR et al.: INFLUENCE OF THE SUBSTRATE TEMPERATURE ON THE STRUCTURAL, OPTICAL ...
374 Materiali in tehnologije / Materials and technology 49 (2015) 3, 371–376
Figure 6: Resistance of V2O5 thin films as a function of the heating
temperature
Slika 6: Upornost tanke plasti V2O5 kot funkcija temperature ogre-
vanja
Figure 5: Resistance of V2O5 film versus the substrate temperature
Slika 5: Odvisnost upornosti tanke plasti V2O5 od temperature pod-
lage
TCR
R
T
=
ln
d
(7)
where R is the resistance and T is the absolute tempera-
ture.
The TCR values are (1.53, 1.59 and 1.50) · 10–2 K–1
for the substrate temperatures of (250, 300 and 350) °C,
respectively. The increase in the TCR value up to 1.59 ·
10–2 K–1 at the substrate temperature of 300 °C may be
due to a more intense crystallization of the film at this
temperature.
The Seebeck coefficient was determined by measur-
ing the thermo e.m.f (V) as a function of the tempe-
rature difference (T). Typical data from the Seebeck
measurements are shown in Figure 8, the thermo e.m.f.
(V) shows a linear dependence of T. The negative
value of the thermo e.m.f. was consistent with the n-type
semiconductor behavior. The Seebeck coefficient (S)
could be found from the slope of the graphs (Figure 8):21
S
V
T
=
Δ
Δ
(8)
The Seebeck-coefficient value increases as the
substrate temperature Ts is increased to the maximum
value of 65 μV/K at 350 °C. The enhancement of the
Seebeck-coefficient value due to the increase in the
substrate temperature leads to a decrease in the activation
energy, which can be attributed to the improvement in
the crystallinity compared to the low-substrate-tempe-
rature films. The thermoelectric properties of the V2O5
thin films produced are relatively suitable for the IR
sensor applications.
4 CONCLUSIONS
V2O5 thin films were prepared, with spray pyrolysis,
on glass substrates at different substrate temperatures.
XRD patterns of the V2O5 thin films showed a crystalline
orthorhombic structure with the preferential orientation
of (001). The films deposited at a temperature of 300 °C
are well textured and c-axis oriented with good
crystalline properties. The optical band gaps of the films
prepared at different temperatures are found to be (2.27,
2.25 and 2.16) eV. The electrical resistance decreased
with an increase in the substrate temperature. The elec-
trical resistance, TCR and the Seebeck coefficient of the
V2O5 films were also strongly influenced by the substrate
temperature.
Acknowledgment
One of the authors, Y. V. K., thanks the UGC, New
Delhi, for providing the financial assistance in the form
of RFSMS. M. V. R. R. thanks UGC-MRP F.41-907/
2012(SR), New Delhi, and DST-PURSE, Osmania Uni-
versity, Hyderabad, for providing the financial assistance
in the form of a project.
5 REFERENCES
1 K. Nagase, Y. Shimizu, N. Miura, N. Yamazoe, Electrochromic
properties of vanadium pentoxide thin films prepared by new wet
process, Appl. Phys. Lett., 60 (1992) 7, 802, doi:10.1063/1.106523
2 R. T. Rajendra Kumar, B. Karunagaran, D. Mangalaraj, S. K. Na-
rayandass, P. Manoravi, M. Joseph et al., Room temperature
deposited vanadium oxide thin films for uncooled infrared detectors,
Mater. Res. Bull., 38 (2003) 7, 1235–1240, doi:10.1016/S0025-5408
(03)00118-1
3 D. Manno, A. Serra, M. DiGiulio, G. Micocci, A. Taurino, A. Tepore
et al., Structural and electrical properties of sputtered vanadium
oxide thin films for applications as gas sensing material, J. Appl.
Phys., 81 (1997) 6, 2709–2714, doi:10.1063/1.363973
4 M. Ponzi, C. Duschatzky, A. Carrascull, E. Ponzi, Obtaining benzal-
dehyde via promoted V2O5 catalysts, Appl. Catal. A, 169 (1998) 2,
373–379, doi:10.1016/S0926-860X(98)00026-X
5 M. F. Al-Kuhaili, E. E. Khawaja, D. C. Ingram, S. M. A. Durrani, A
study of thin films of V2O5 containing molybdenum from an evapo-
ration boat, Thin Solid Films, 460 (2004) 1–2, 30–35, doi:10.1016/
j.tsf.2004.01.076
Y. VIJAYAKUMAR et al.: INFLUENCE OF THE SUBSTRATE TEMPERATURE ON THE STRUCTURAL, OPTICAL ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 371–376 375
Figure 8: V versus T for different substrate temperatures
Slika 8: Odvisnost V od T pri razli~nih temperaturah podlage
Figure 7: Ln (resistance) versus 1000/T
Slika 7: Odvisnost Ln (upornost) od 1000/T
6 C. V. Ramana, B. S. Naidu, O. M. Hussain, R. Pinto, Low-tempera-
ture growth of vanadium pentoxide thin films produced by pulsed
laser ablation, J. Phys. D: Appl. Phys., 34 (2001) 7, L35,
doi:10.1088/0022-3727/34/7/101
7 C. W. Zou, X. D. Yan, J. Han, R. Q. Chen, W. Gao, Microstructures
and optical properties of -V2O5 nanorods prepared by magnetron
sputtering, J. Phys. D: Appl. Phys., 42 (2009) 14, 145402,
doi:10.1088/0022-3727/42/14/145402
8 Y. V. Kumar, M. V. R. Reddy, K. N. Reddy, Preparation and optical
characterization of V2O5 thin films, International Conference on
Nanoscience, Engineering and Technology (ICONSET), Chennai,
India, 2011, 244–246, doi:10.1109/ICONSET.2011.6167931
9 M. A. Kaid, Characterization of Electrochromic Vanadium Pentoxide
Thin Films Prepared By Spray Pyrolysis, Egypt. J. Solids, 29 (2006)
2, 273–291
10 P. S. Patil, Versatility of chemical spray pyrolysis technique, Mate-
rials Chemistry and Physics, 59 (1999) 3, 185–198, doi:10.1016/
S0254-0584(99)00049-8
11 P. S. Patil, Preparation and characterization of spray pyrolyzed nickel
oxide (NiO) thin films, Appl. Surf. Sci., 199 (2002) 1–4, 211–221,
doi:10.1016/S0169-4332(02)00839-5
12 M. Mousavi, A. Kompany, N. Shahtahmasebi, M. M. Bagheri-Moha-
gheghi, Study of structural, electrical and optical properties of
vanadium oxide condensed films deposited by spray pyrolysis tech-
nique, Adv. Manuf., 1 (2013) 4, 320–328, doi:10.1007/s40436-
013-0045-y
13 A. Kumar, P. Singh, N. Kulkarni, D. Kaur, Structural and optical
studies of nanocrystalline V2O5 thin films, Thin Solid Films, 516
(2008) 6, 912–918, doi:10.1016/j.tsf.2007.04.165
14 R. Irani, S. M. Rozati, S. Beke, Structural and optical properties of
nanostructural V2O5 thin films deposited by spray pyrolysis techni-
que: Effect of the substrate temperature, Materials Chemistry and
Physics, 139 (2013) 2–3, 489–493, doi:10.1016/j.matchemphys.
2013.01.046
15 S. Aydogu, O. Sendil, M. B. Coban, The Optical and Structural
Properties of ZnO Thin Films Deposited by the Spray Pyrolysis
Technique, Chinese Journal of Physics, 50 (2012) 1, 89–100
16 H. M. Martínez, J. Torres, L. D. López-Carreño, M. E. Rodríguez-
García, The Effect of Substrate Temperature on the Optical Pro-
perties of MoO3 Nano-crystals Prepared Using Spray Pyrolysis, J.
Supercond. Nov. Magn., 26 (2013) 7, 2485–2488, doi:10.1007/
s10948-012-1692-0
17 G. A. Khan, C. A. Hogarth, Optical absorption spectra of evaporated
V2O5 and co-evaporated V2O5/B2O3 thin films, Journal of Mate-
rials Science, 26 (1991) 2, 412–416, doi:10.1007/BF00576535
18 Z. Luo, Z. Wu, X. Xu, M. Du, T. Wang, Y. Jiang, Impact of substrate
temperature on the microstructure, electrical and optical properties of
sputtered nanoparticle V2O5 thin films,Vacuum, 85 (2010) 2,
145–150, doi:1016/j.vacuum.2010.05.001
19 R. M. Öksüzoðlu, P. Bilgiç, M. Yýldýrým, O. Deniz, Influence of
post-annealing on electrical, structural and optical properties of
vanadium oxide thin films, Optics and Laser Technology, 48 (2013),
102–109, doi:10.1016/j.optlastec.2012.10.001
20 H. Wang, X. Yi, S. Chen, Low temperature fabrication of vanadium
oxide films for uncooled bolometric detectors, Infrared Physics &
Technology, 47 (2006) 3, 273–277, doi:10.1016/j.infrared.2005.04.
001
21 S. Iwanaga, M. Marciniak, R. B. Darling, F. S. Ohuchi, Thermo-
power and electrical conductivity of sodium-doped V2O5 thin films,
J. Appl. Phys., 101 (2007), 123709, doi:10.1063/1.2739311
Y. VIJAYAKUMAR et al.: INFLUENCE OF THE SUBSTRATE TEMPERATURE ON THE STRUCTURAL, OPTICAL ...
376 Materiali in tehnologije / Materials and technology 49 (2015) 3, 371–376
V. YILMAZ et al.: DEEP MICRO-HOLE DRILLING FOR HADFIELD STEEL ...
DEEP MICRO-HOLE DRILLING FOR HADFIELD STEEL BY
ELECTRO-DISCHARGE MACHINING (EDM)
VRTANJE GLOBOKIH MIKROLUKENJ V JEKLA HADFIELD
Z METODO ELEKTRORAZREZA (EDM)
Volkan Yilmaz1, Murat Sarýkaya2, Hakan Dilipak1
1Manufacturing Department, Technology Faculty, Gazi University, 06500 Ankara, Turkey
2Department of Mechanical Engineering, Sinop University, 57030 Sinop, Turkey
msarikaya@sinop.edu.tr
Prejem rokopisa – received: 2014-06-23; sprejem za objavo – accepted for publication: 2014-09-05
doi:10.17222/mit.2014.091
In this study, a new system for drilling deep micro-holes was designed for Hadfield steel (which is difficult to process with
classical methods) with the electro-discharge-machining method (EDM) and the system was experimentally examined. The tests
were carried out at three different discharge currents (6, 12 and 24) A, three different electrode-tool rotational speeds (200, 400
and 600) r/min, three different dielectric spray pressures (40, 80 and 120) bar, a constant pulse duration (12 μs) and a constant
pulse interval (3 μs). After the tests the effects of the processing parameters on the basic performance outputs (the material
removal rate – MRR), the electrode wear rate (EWR) and the relative wear (RW)) were investigated. Additionally, an analysis of
variance (ANOVA) was also applied to identify the most significant factor. Optimum operating parameters were determined
using the desirability-function analysis through the response surface methodology (RSM). It was found that the most effective
variable affecting the MRR, EWR and RW was the discharge current. The discharge current was found to be the most significant
control factor influencing the performance of the machining process.
Keywords: deep micro-EDM, hole drilling, Hadfield steel, multi-response optimization, ANOVA
V tej {tudiji je bil postavljen nov sistem za vrtanje globokih mikrolukenj v jeklo Hadfield (ki se te`ko obdeluje z navadnimi
metodami) z metodo elektroerozije (EDM) in bil tudi eksperimentalno preizku{en. Poskusi so bili izvedeni s tremi razli~nimi
tokovi (6, 12 in 24) A, pri treh razli~nih hitrostih vrtenja orodja (200, 400 in 600) r/min, pri treh razli~nih dielektri~nih tlakih
razpr{evanja (40, 80 in 120) bar, pri konstantnem trajanju impulza (12 μs) in konstantnem intervalu impulza (3 μs). Preiskovani
so bili u~inki procesnih parametrov na zmogljivost (hitrost odvzema materiala (MRR), stopnja obrabe elektrode (EWR) in
relativna obraba (RW)). Dodatno je bila uporabljena analiza variance (ANOVA) za ugotovitev najpomembnej{ega faktorja.
Optimalni obratovalni parametri so bili dolo~eni z analizo funkcije odziva na metodologijo odziva povr{ine (RSM). Ugotovljeno
je bilo, da je razelektritveni tok najpomembnej{i kontrolni faktor, ki vpliva na zmogljivost procesa obdelave.
Klju~ne besede: globoka mikro-EDM, vrtanje luknje, jeklo Hadfield, optimiranje multiodziva, ANOVA
1 INTRODUCTION
The diameters of the holes are becoming smaller with
the developing micro-mechanical systems and the classi-
cal chip-removal methods are insufficient for obtaining
micro-holes, so the researchers are focusing on new
manufacturing methods. Among these methods, the most
applicable and commercially used one is the processing
using EDM. Easy processing of hard materials and
complex geometries made this method one of the most
preferred uncommon manufacturing methods.1 The most
important characteristic that must be exhibited by the
workpieces to be processed by EDM is electrical con-
ductivity. The characteristics such as the workpiece hard-
ness and toughness that are effective in the processing
with conventional manufacturing methods are not im-
portant with respect to EDM. On the other hand, a good
processing performance depends on the thermal and
electrical conductivities of a material.2–4 Studies of EDM
are generally concentrated on the performance outputs.
In the studies based on the lower-duration, low-cost and
high-quality expectations of the manufacturing industry,
the improvement of MRR, EWR, RW and surface-rough-
ness outputs is emphasized.5–8 Besides these studies on
the EDM system, rapid hole-drilling electro-discharge-
machining machines were also developed to meet the
new development expectations. This new EDM techni-
que became a production technique that is often pre-
ferred in the aviation (cooling holes in plane turbine
blades), automotive (fuel injection) and medical (dental
and surgical implants) areas, used for medical materials,
cutting-tool cooling channels and micro-hole drilling of
hard, brittle and difficult-to-process materials.9–12 In this
method, small-sized processing residuals are removed
from the processing area by means of a dielectric liquid
sprayed at a desired pressure through a tube-type elec-
trode of a small diameter, rotating at a specified speed.
Holes with larger diameters than the one of the used
electrode tool are easily produced. The most important
advantage of the method is that holes can be drilled into
any electrically conductive metal. Micro-EDM is an im-
portant method, especially for small pieces, micro-con-
stituents and the production of micro-tools providing a
good surface quality and high integrity. Besides, it is
Materiali in tehnologije / Materials and technology 49 (2015) 3, 377–386 377
UDK 621.95:519.233.4:621.9.048 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)377(2015)
maintained that micro-production technology can be
developed with the micro-EDM.13 In addition to the
tube-like tools, the use of the tools that are not tube-like
is also increasing and it is maintained that the holes
obtained with the cylindrical tool with an orbital motion
are more uniform than the ones made with the tool not
having a rotational motion.14 It was observed that MRR
increased with the inverse polarity (tool š+’,workpiece
š–’) causing the electrode wear as well.15 It is known that
vibration applications give favorable results with respect
to the increased MRR and MRR increases during
vibrational processes with the shortening of the process
duration.16,17
When the studies in the literature are evaluated, it is
seen that hole-drilling applications using the developing
EDM method are used for the micro-hole drilling pro-
cesses as an alternative method. The current studies
relating to all the metals are developing in the direction
of drilling deep holes with desired sizes and geometries
using the above types of drilling units. In deep micro-
hole drilling applications, the main expectations are to
obtain a higher MRR and lower EWR and RW for a small
diameter and a higher hole length. The aim of this study
is to meet these expectations during the drilling of mi-
cro-holes into Hadfield steel, which is hard to process
with the classical chip-removal methods due to deforma-
tion hardening; the effects of the discharge current, the
electrode-tool rotational speed and the dielectric spray
pressure on the basic performance outputs were investi-
gated. Besides, in the literature, no study on deep micro-
hole drilling of Hadfield steel with EDM was en-
countered.
2 MATERIAL AND METHOD
2.1 Experimental material and equipment
As the test sample, Hadfield steel, which is hard to
process with the classical chip-removal methods (be-
cause of deformation hardening) was used. Because of
its unique properties such as hardness, wear resistance,
strength and low thermal conductivity, this material has
recently been commonly utilized in many engineering
operations involving mining equipment, excavators, rail-
ways, pumping equipment, rolling-mill parts for steel
factories and wear-resistant components of machining
elements18,19. Test samples in the size of 10 mm × 20 mm
× 200 mm were prepared. The chemical composition of
Hadfield material is given in Table 1. In the tests, the
FURKAN brand, "EEÝ M50A" type EDM machine was
used. A head was fixed to the moving head of the elec-
tro-erosion machine to rotate the electrode tool at diffe-
rent rotations and spray the process liquid (dielectric
liquid) to the processing area at the desired pressure. The
fixed head is seen in Figure 1. In the tests, brass elec-
trodes with a length 400 mm, inside diameter 0.18 mm
and outside diameter 0.8 mm were used as the electrode
tool. A fixed electrode and the test samples are presented
in Figure 2.
V. YILMAZ et al.: DEEP MICRO-HOLE DRILLING FOR HADFIELD STEEL ...
378 Materiali in tehnologije / Materials and technology 49 (2015) 3, 377–386
Figure 2: Fixing the electrode
Slika 2: Pritrditev elektrode
Figure 1: Pressure head
Slika 1: Tla~na glava
Table 1: Chemical composition of the workpiece material in mass fractions, w/%
Tabela1: Kemijska sestava materiala obdelovanca v masnih dele`ih, w/%
C Si Mn P S Cr Mo Ni Al Co Cu V Fe
1.08 0.621 13.6 0.0152 0.0004 0.721 0.263 0.286 0.004 0.023 0.183 0.0003 Balance
2.2 Experiments
In this study, tests were made using the EDM method
at three different discharge currents (6, 12 and 24) A,
three different electrode rotational speeds (200, 400 and
600) r/min, three different dielectric spray pressures (40,
80 and 120) bar, a constant pulse duration (12 μs) and a
constant pulse interval (3 μs). A schematic view of the
EEP unit designed for the tests is given in Figure 3. As
seen in this figure, the pressure head is mounted on the
moving part of the EDM machine at the Z-axis. Owing to
the pressure head, the electrode had a rotational motion
and the pressurized dielectric liquid had access to the
processing area through the electrode. The rotational
motion was given to the electrode by a D. A motor on
the pressure head. Due to the power source added to the
system, the rotational speed of the electrode was con-
trolled and it reached the desired r/min values. In this
system, the motion of the electrode at the vertical axis
was provided by another D. A motor included in the
EDM machine. Dielectric liquid reached the pressure
head by means of a pressure pump. The dielectric liquid
pressure was continuously controlled by a manometer
mounted on the by-pass mechanism. The desired dielec-
tric liquid pressure was achieved with a set screw. With a
high-pressure resistant hose, dielectric liquid was
pumped to the pressure head at the pressure of up to 200
bar. Dielectric liquid was pumped to the processing area,
passing through the interior part (0.18 mm) of the elec-
trode fixed to the pressure head by means of a mandrel.
The processes were carried out in a separate tank
mounted in the process chamber. The workpieces in the
process tank were fixed with a clamp. The parallelism of
the clamp and the process tank was controlled with an
assay balance and it was verified that they were also
parallel to the EDM machine base and perpendicular to
the pressure head.
2.3 Determination of the EDM basic performance
outputs (MRR, EWR and RW)
At the end of the tests, the MRR, EWR and RW values
were calculated with the following formulas:
MRR =
Total workpiece wear volume (mm )
Total working
3
time (min)
(1)
EWR =
Total electrode wear volume (mm )
Total working
3
time (min)
(2)
RW
EWR
MRR
= ⋅100 (3)
MRR is material removal rate (mm3/min), EWR is
electrode wear rate (mm3/min), RW is relative wear (%).
3 RESULTS AND DISCUSSION
The hole pictures obtained after the tests are given in
Figure 4 and the test results are in Table 2. The perform-
ance characteristics were specified as MRR/(mm3/min),
EWR/(mm3/min) and RW/%. The results were expressed
graphically in order to be easily discussed and compared.
Table 2: Experimental results
Tabela 2: Eksperimentalni rezultati
Te
st
nu
m
be
r
D
is
ch
ar
ge
cu
rr
en
t
(A
)
D
ie
le
ct
ri
c
sp
ra
y
pr
es
su
re
(b
ar
)
E
le
ct
ro
de
ro
ta
ti
o-
na
l
sp
ee
d
(r
/m
in
)
M
at
er
ia
l
re
m
ov
al
ra
te
M
R
R
/(
m
m
³/
m
in
)
E
le
ct
ro
de
w
ea
r
ra
te
E
W
R
/(
m
m
³/
m
in
)
R
el
at
iv
e
w
ea
r
R
W
/%
1
6
40
200 1.452 0.473 32.595
2 400 1.654 0.457 27.622
3 600 1.801 0.419 23.282
4
80
200 1.433 0.519 36.217
5 400 1.707 0.540 31.664
6 600 1.808 0.561 31.043
7
120
200 2.017 0.613 30.41
8 400 2.291 0.622 27.163
9 600 2.173 0.616 28.344
10
12
40
200 5.228 2.678 51.221
11 400 4.932 2.417 49.015
12 600 5.301 2.715 51.221
13
80
200 5.614 2.875 51.221
14 400 6.524 3.182 48.782
15 600 6.769 4.098 60.534
16
120
200 5.875 3.283 55.878
17 400 7.218 3.309 45.845
18 600 7.126 3.921 55.031
19
24
40
200 8.815 10.567 119.88
20 400 9.950 11.501 115.59
21 600 9.474 9.722 102.62
22
80
200 10.527 10 94.989
23 400 14.385 10.960 76.197
24 600 12.632 9.625 76.197
25
120
200 11.475 7.898 68.835
26 400 14.982 9.842 65.694
27 600 16.452 8.325 50.6
V. YILMAZ et al.: DEEP MICRO-HOLE DRILLING FOR HADFIELD STEEL ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 377–386 379
Figure 3: Schematic view of designed EEP test set-up
Slika 3: Shematski prikaz zasnovanega EEP-preizku{evali{~a
3.1 Variation of the metal removal rate (MRR) with the
processing parameters
MRR is expressed as the amount of the removed
material (chip) per unit time (mm3/min) and it is one of
the most important output parameters in the EDM opera-
tions. In the EDM operations, obtaining high MRR-va-
lues is the main requirement and the studies are focusing
on this issue. The variations in the MRR-values obtained
in the tests with the processing parameters are graphi-
cally expressed in Figures 5 and 6.
When the variation of MRR with the discharge cur-
rent (I) is examined in Figure 5, it is seen that the MRR-
values increase with an increase in the I-values.
Increasing the crater dimensions (occurring on the
surface of the workpiece due to electrical discharges)
with the direct proportional increase in the discharge
energy is the general principle of the electro-erosion
processing method. The reason for this is the evaporation
of a large amount of the material (through fusion) per
unit time from the workpiece surface due to the increas-
ing discharge energy with the increase in the discharge
current. Namely, with the increase in the discharge
current each spark becomes more severe and each time
these sparks pull off a greater area from the workpiece
material. So, with the increasing discharge current (I)
more material is fused and evaporated in a shorter time,
causing an increase in the MRR-values. In the tests, the
MRR-values obtained in the interval of 1.4–2.3 mm3/min
with the 6 A discharge current, increased by approxima-
tely 200 %, to the 4.9–7.2 mm3/min interval, with the
12 A discharge current, and by approximately 90 %, to
the 8.8–16.4 mm3/min interval, with the 24 A discharge
current. When the general principle of the EDM system
is taken to be the chip removing of high-energy sparks
(occurring between the electrode and the workpiece)
from the workpiece surface due to fusion and evapora-
tion, the increase in the MRR-values with the increase in
the discharge-current values is comparable with the
results in3–8. When Figures 5 and 6 are evaluated, it is
seen that as the I-value is increased, keeping the elec-
trode-tool rotational speed (n) constant, there is a gradual
increase in the MRR-values. This shows that in all of the
tests with the increasing discharge-current values the
MRR-values increase without an exception. With the
increasing electrode-tool rotational speed, the MRR per-
formance values also increased. In the 6 A discharge-
current tests at the 40 bar dielectric liquid spray pressure,
with an increase in the electrode-tool rotational speed
from 200 r/min to 400 r/min, the MRR-values increased
by 14 %, and with an increase from 400 r/min to 600
r/min, MRR exhibited an increase of 9 %. This increase
was 19 % and 6 % at the 80 bar dielectric liquid
pressure. It is also valid for the other tests. The increase
in the tool rotational speed provided a continuous and
rapid flow of dielectric liquid to the processing area and,
consequently, the formation of continuous sparks in the
processing area made the processing more efficient and
uninterrupted. The continuous spark discharge became
the most important reason for the increase in the MRR-
values. When the experimental values are considered the
increase in the dielectric spray pressure, together with
the rotational speed, makes a significant contribution to
the effective washing of the processing area. In the 12 A
discharge-current tests, at the 200 r/min electrode-tool
rotational speed, with the increase in the dielectric liquid
spray pressure from 40 bar to 80 bar, the MRR-values
V. YILMAZ et al.: DEEP MICRO-HOLE DRILLING FOR HADFIELD STEEL ...
380 Materiali in tehnologije / Materials and technology 49 (2015) 3, 377–386
Figure 5: MRR – I variation
Slika 5: Diagram MRR – I
Figure 4: Hole pictures after the tests
Slika 4: Videz lukenj po preizkusih
Figure 6: MRR – P variation
Slika 6: Diagram MRR – P
increased by 8 %, and with an increase from 80 bar to
120 bar, MRR exhibited an increase of 5 %. The main
reason for this is a faster motion of the fluid dielectric
liquid in the area of the rotation of the tool. Due to this
rapid motion, the dielectric liquid moves away from the
side spaces between the electrode tool and the workpiece
faster and the hole-drilling operation is carried out more
effectively. When an evaluation of the MRR-values is
made it can be found that the spray-type dielectric-appli-
cation method is very effective. These results show
parallelism with the studies in the literature. With respect
to the spray-type dielectric-liquid applications, the
reports in the literature state that higher MRR-values are
obtained with the lateral-spray type, the reasons being a
low temperature, a smaller amount of the processing-
area contamination and a lower volume of gas in the pro-
cessing space compared to the other methods.15,16 With
the spray-type washing used in the tests, an effective
washing of the processed products can be carried out.
This way a short-circuit formation of a continuous, clean
dielectric liquid decreases and the sparks in the clean
processing space affect the workpiece more effectively,
increasing the MRR. Thus, for the higher MRR-values in
the EDM operations, it was determined that the dis-
charge current, the electrode-tool rotational speed and
the dielectric spray pressure must be selected at high
intervals.
3.2 Variation in the electrode-wear rate (EWR) with
the processing parameters
Another important performance characteristic in the
EDM applications is the EWR-value. The sparks created
between the workpiece and the electrode in the EDM
operations not only fuse the area and cause evaporation
of the workpiece but they also cause evaporation of a
certain area on the electrode tool. This loss of the elec-
trode tool is expressed as EWR and in this study EWR
was calculated as the decrease in the electrode volume
per unit time (mm3/min). EWR primarily depends on the
thermal and electrical properties of the electrode material
and then also on the processing parameters7. The rela-
tionships between the EWR-values and the discharge-
current (I), electrode-tool-rotational-speed and dielec-
tric-spray-pressure values are graphically shown in
Figures 7 and 8.
When the variation in the EWR of the electrode tool
(brass) (Figure 7) with the discharge current (I) is exa-
mined it is observed that the electrode-tool EWR
increases with the increase in the I-value, keeping the
tool rotational speed constant. It was also found that due
to the increase in the EWR of the electrode tool more
material is fused and evaporated, proportionally with the
applied discharge energy. As seen from Figure 4, when
the discharge current is 6 A, the EWR-values are in the
interval of 0.4–0.6 mm3/min and in the case of the 12 A
discharge current the EWR-values increase by approxi-
mately 500 % and take place in the 2.4–4 mm3/min
interval. With the discharge current of 24 A, the
EWR-values increase by approximately 300 % and rise
to the 9.6–11.5 mm3/min interval. According to these
results, the increases in the discharge-current values
caused very high increases in the brass-tool EWR-values.
The reason for this is that the electrode has a fine
structure and its inside space is empty. This fine structure
heats up very rapidly by losing its electrical resistance
with the increase in the tool discharge-current values. In
the tests using 12 A and 24 A this was explicitly evident.
During the processing these currents heated up the fine
electrode material in a very short time and during the
flowing of the sparks from the tool to the workpiece, big
pieces were pulled off due to fusion, causing an increase
in the EWR-values. From Figures 7 and 8 it is seen that
in the tests with the constant dielectric spray pressure
and I-value, with an increase in the rotational speed the
EWR-values had a tendency to increase. Since the
increasing rotational speed provided a continuous, clean
processing liquid in the processing area, more sparks
occurred and an increased spark discharge increased the
V. YILMAZ et al.: DEEP MICRO-HOLE DRILLING FOR HADFIELD STEEL ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 377–386 381
Figure 8: EWR – P variation
Slika 8: Diagram EWR – P
Figure 7: EWR – I variation
Slika 7: Diagram EWR – I
EWR.3,7,10 Consequently, in the continuously cleaned
processing area an effective and high amount of spark
discharge caused more material to fuse and evaporate
from the finely structured tool and, naturally, also caused
the EWR to increase. This increase changes depending
on the physical property of the electrode material, the
discharge energy, the pulse duration, the dielectric spray
characteristic and the type. When the relationship bet-
ween the EWR and the dielectric spray pressure is exa-
mined in Figure 8, it is seen that the EWR-values
increase with the increase in the dielectric spray pres-
sure. With the increase in the dielectric spray pressure
from 40 bar to 80 bar at the 6 A discharge current and
200 r/min electrode-tool rotational speed, the EWR-va-
lues increased by 9 %, and with the increase from 80 bar
to 120 bar, the EWR-values increased by 20 %. The in-
creases were 20 % and 15 % at the 400 r/min elec-
trode-tool rotational speed and 37 % and 9 % at the 600
r/min electrode-tool rotational speed. The reason for the
increase was found to be the rapid cleaning of the
processing area with the increase in the dielectric spray
pressure and the effective spark discharges. Besides, it is
also known from the literature that the processed pro-
ducts are moved away from the processing area more
effectively with the spraying of dielectric from the
tool.20–22 So, the decreases in EWR depending on the
dielectric spray pressure relieve the processing area of
effective spraying in electro-erosion drilling, clean the
processing area more rapidly and, with a higher energy
spark discharge, pull off a smaller amount of electrode
material from the tube-like tool (from the point of
separating the sparks).
3.3 Variation in the relative wear (RW) with the pro-
cessing parameters
RW is a significant output parameter expressing the
relationship between EWR and MRR during each process
in the EDM operations. A graphical expression of the
RW-values calculated with the data from the tests is
given in Figures 9 and 10.
From Figure 9, it is seen that the increases in the
RW-values also occurred, being parallel to the increase in
the discharge current. The reason for this was a higher
increase in the EWR-values with respect to the
MRR-values at high discharge-current values. In the tests
with the 6 A discharge current, the RW-values were in
the interval of 23–33 % and in the tests with the 12 A
discharge current the interval increased to 49–52 %. This
shows that both EWR- and MRR-values increased by the
same, small amount during the increase in the dis-
charge-current value from 6 A to 12 A. The RW-values
exhibited a significant increase with the 24 A discharge
current compared to the results obtained with the 6 A
and 12 A values and the RW-values increased to the
102–120 % interval. This was explained with the fact
that the finely structured and empty electrode tool had a
higher EWR with respect to the MRR at high dis-
charge-current values. As an empty and finely structured
electrode tool rapidly loses its electrical resistance at a
high discharge current, it causes the EWR-values to
increase. So, when RW is calculated with the (EWR/
MRR) × 100 formula, these increases in the EWR-values
that are higher than the MRR-values also cause an
increase in the RW-values. In Figure 9 it is seen that the
RW-values have a tendency to increase with the
rotational speed applied to the electrode tool. This was
explained with the increase in the EWR-values caused by
the increasing electrode rotational speed. In Figure 10 it
is observed that the RW-values decrease with the increase
in the dielectric spray pressure. Processing residuals can
be easily removed from the processing area owing to the
spray-type washing applied in the tests (in line with the
studies in the literature)21–24 and for this reason the
MRR-values increase in the clean processing area. So,
increasing the MRR-values more than the EWR-values
caused a decrease in the RW-values.
V. YILMAZ et al.: DEEP MICRO-HOLE DRILLING FOR HADFIELD STEEL ...
382 Materiali in tehnologije / Materials and technology 49 (2015) 3, 377–386
Figure 10: RW – P variation
Slika 10: Diagram RW – P
Figure 9: RW – I variation
Slika 9: Diagram RW – I
3.4 Analysis of variance (ANOVA)
Analysis of variance (also known as ANOVA) is a
statistical method used to identify individual interactions
of all the control factors in the experimental results. In
the present work, ANOVA was used to determine the
effects of the discharge current, the dielectric spray
pressure and the electrode rotational speed on MRR,
EWR and RW. ANOVA results for the responses are
given in Table 3. The ANOVA analysis was performed at
the 95 % confidence level and 5 % significance level.
The F-values of the control factors indicated the signi-
ficance of the control factors with the ANOVA anal-
ysis.19,25–27 The percentage contribution of each parame-
ter is shown in the last column of the ANOVA table. This
column shows the influence rates of the control factors
for the experimental results. In addition, ANOVA results
are also summarized as a column chart in Figure 11. In
Table 3, the percent contributions of the factors such as
discharge current, dielectric spray pressure and electrode
rotational speed to the MRR were determined as 88.1 %,
4.6 % and 1.7 %, respectively. Therefore, the most
effective variable affecting the MRR was the discharge
current (88.1 %). It was seen that the discharge current
and dielectric spray pressure significantly affect the MRR
at the reliability level of 95 % or the significance level of
5 %, because the P-values of these variables are lower
than 0.05.26 According to Table 3, the percent contribu-
tions of the input parameters to the EWR were found to
be 96.9 %, 0.2 % and 0.2%, respectively, and the error
was 2.7 %. It was determined that the most effective
parameter with respect to the EWR is the discharge
current. Moreover, among the input parameters, only the
discharge current significantly affects the EWR at the
reliability level of 95 % or the significance level of 5 %.
From Table 3, the effects of the control factors on the
RW were obtained as 74.8 %, 6.2 % and 1.3 %, and the
error amounted to 17.7 % of the contribution rate. The
ANOVA table indicated that with respect to the RW, the
most effective parameter is the discharge current.
3.5 Multi-response optimization of the EDM parame-
ters based on RSM
Mono-response optimization is a common and
popular method to solve the problems of optimization
approaches. But the method cannot be used to determine
the optimum combination of the machining parameters
that simultaneously optimize the output parameters.28 To
overcome this problem in the present work, a multi-
V. YILMAZ et al.: DEEP MICRO-HOLE DRILLING FOR HADFIELD STEEL ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 377–386 383
Figure 11: Histogram of the ANOVA results
Slika 11: Histogram rezultatov ANOVA
Table 3: ANOVA results for means
Tabela 3: ANOVA povpre~ni rezultati
Variation of source
Degree of
freedom
(DF)
Sum of squares
(SS)
Mean of
squares
(MS)
F-ratio P-value
Contribution
(%)
MRR
Discharge current 2 478.516 239.258 156.96 0.000 88.1
Dielectric spray pressure 2 24.893 12.447 8.17 0.003 4.6
Electrode rotational speed 2 9.216 4.608 3.02 0.071 1.7
Error 20 30.487 1.524 5.6
Total 26 543.112 100
EWR
Discharge current 2 412.892 206.446 358.86 0.000 96.9
Dielectric spray pressure 2 0.882 0.441 0.77 0.478 0.2
Electrode rotational speed 2 0.912 0.456 0.79 0.466 0.2
Error 20 11.506 0.575 2.7
Total 26 426.191 100
RW
Discharge current 2 14205.4 7102.7 42.16 0.000 74.8
Dielectric spray pressure 2 1175.1 587.5 3.49 0.050 6.2
Electrode rotational speed 2 253.6 126.8 0.75 0.484 1.3
Error 20 3369.5 168.5 17.7
Total 26 19003.6 100
response optimization was performed to determine the
objective values of three responses, i.e., MRR, EWR, and
RW in deep micro-hole drilling of Hadfield steel with
EDM. The multi-response optimization was employed
through the Minitab 16.0 software based on the response
surface methodology (RSM).26 The graph of the multi-
response optimization plotted in Minitab is shown in
Figure 12. In this figure, each column of the plots shows
the machining parameters, and the responses are shown
by each row of the plots.
The objective value of each response, namely y, is
displayed along with the desirability value range, namely
d, which is between 0 and 1 as shown in the figure. If d =
0 or approaches to 0, then the output is clearly undesir-
able. If d = 1 or approaches to 1, then the output perfect-
ly meets the target value. A higher value of desirability
indicates a better optimization.26 The highest desirability
value is favored for the best solution of deep micro-EDM
drilling. The goal, lower value, target value, upper value,
mass, the importance of the factors and the best global
solution were determined for the multi-response optimi-
zation as shown in Table 4. According to Table 4 and
Figure 12, the optimization values are found to be
10.512 mm3/min, 4.737 mm3/min and 50.1315 % for
MRR, EWR and RW, respectively. Individual desirability
values are 0.60448, 0.61041 and 0.72205. Moreover, the
composite desirability value is 0.643461 for all the
responses. The levels of the control parameters are found
to be 15.6831 A for the discharge current, 120 bar for the
dielectric spray pressure and 600 r/min for the electrode
rotational speed for the multi-response optimization in
deep micro-hole drilling of Hadfield steel by EDM.
4 CONCLUSION
In this study, using the EDM method, deep micro-
hole drilling was carried out on Hadfield steel which is
hard to process with classical chip-removing methods
due to deformation hardening. The effects of the process-
ing parameters of deep micro-hole drilling applications
(discharge current, electrode tool rotational speed and
dielectric spray pressure) on the processing performance
outputs (MRR, EWR, RW) were examined experimen-
tally. The results are given below.
• By means of the installed system, the holes with a
diameter 0.8 mm and length 20 mm were drilled into
V. YILMAZ et al.: DEEP MICRO-HOLE DRILLING FOR HADFIELD STEEL ...
384 Materiali in tehnologije / Materials and technology 49 (2015) 3, 377–386
Figure 12: Plot of multi-response optimization
Slika 12: Prikaz multiodzivne optimizacije
Table 4: Best global solutions for multi-optimization
Tabela 4: Najbolj{e celotne re{itve multioptimizacije
Response Goal
Global solution
(Multi-optimization) Lower Target Upper Weight Imp. Predicted Desira-bility
I/A P/bar r/min
MRR Max. 15.6831 120 600 1.4332 16.4525 16.453 1 1 10.512 0.604479
EWR Min. 15.6831 120 600 0.4194 0.4194 11.502 1 1 4.7370 0.610408
RW Min. 15.6831 120 600 23.282 23.282 119.880 1 1 50.1315 0.722049
Composite desirability = 0.643461
Hadfield steel. The most important parameter affect-
ing the drilling durations was the discharge current.
Increasing the electrode-tool rotational speed and
dielectric spray pressure provided significant contri-
butions to an easy drilling of the holes.
• When the MRR results were considered, a significant
increase in the MRR-values was observed, related to
the increase in the discharge-current values. The
MRR-values had a tendency to increase with the
increase in the electrode-tool rotational speed. Owing
to the increasing dielectric spray pressure, the pro-
cessing area always remained clean and the MRR-
values increased accordingly.
• According to the EWR results, the EWR-values
increased with the increase in the discharge current.
The increase in the EWR-values was explained with a
finely structured and empty electrode tool. The
EWR-values increased with the increasing electrode-
tool rotational speeds and decreased with the in-
creasing dielectric spray pressure.
• When the RW-values were taken into consideration,
the RW-values also increased depending on the
increase in the discharge current. This was explained
with the fact that the EWR-values increased more
than the MRR-values because of the rise in the
discharge current. The RW-values increased with the
increase in the electrode-tool rotational speeds but
decreased with the increasing dielectric spray pres-
sure.
• From the results of ANOVA for all the responses, it
was found that the most effective variable affecting
the MRR, EWR and RW was the discharge current.
Moreover, it was seen that the discharge current
significantly affected the MRR, EWR and RW at the
reliability level of 95 % or the significance level of
5 %.
• From the multi-response optimization results based
on RSM, the optimum values were found to be
10.512 mm3/min, 4.737 mm3/min and 50.1315 % for
MRR, EWR and RW. The levels of the control
parameters for the optimum results were found to be
15.6831 A for the discharge current, 120 bar for the
dielectric spray pressure and 600 r/min for the
electrode rotational speed in deep micro-hole drilling
of Hadfield steel with EDM.
5 REFERENCES
1 V. Yilmaz, Experimental investigation of drillability of micro holes
using electro discharge machining, Ph. D. Thesis, Gazi University
Graduate School of Natural and Applied Sciences, Ankara, 2013
2 P. C. Pandey, H. S. Shan, Modern machining processes, Tata
McGraw-Hill Publishing Company Limited, New Delhi 1980,
84–114
3 R. K. Springborn, Non-traditional machining processes, American
Society of Tool and Manufacturing Engineers, USA 1967
4 H. M. Chow, B. H. Yan, F. Y. Huang, J. C. Hung, Study of added
powder in kerosene for the micro-slit machining of titanium alloy
using electro-discharge machining, Journal of Materials Processing
Technology, 101 (2000), 95–103, doi:10.1016/S0924-0136(99)
00458-6
5 Z. E. Ergün, C. Çoðun, Experimental investigation on workpiece sur-
face characteristics in electric discharge machining (EDM), J. Fac.
Eng. Arch. Gazi Univ., 21 (2006) 3, 427–441
6 A. Özgedik, C. Çoðun, An Investigation on the effect of the applied
machining tank vibrations on the machining performance in
electrical discharge machining, Electronic Journal of Machine
Technologies, 8 (2011) 3, 13–25
7 A. Özgedik, C. Çoðun, Experimental investigation of the electrode
front surface wear in electrical discharge machining, Engineer and
Machinery, 521 (2003), 21–28
8 C. Çoðun, B. Kocabaº, A. Özgedik, Experimental and theoretical
investigation of workpiece surface profiles in electrical discharge
machining (EDM), J. Fac. Eng. Arch. Gazi Univ., 19 (2004) 1,
97–106
9 V. Yilmaz, H. Dilipak, Method of electro discharge machining
(EDM) micro hole drilling system design, IV. UTIS 2013, Kuþadasý,
Turkey, 2013, 151–159
10 F. N. Leao, Optimisation of EDM fast holedrilling through evaluation
of dielectric and electrode materials, Proceedings of COBEM 2005,
18th International Congress of Mechanical Engineering, Ouro Preto,
MG, 2005
11 P. Kuppan, A. Rajadurai, S. Narayanan, Influence of EDM process
parameters in deep hole drilling of Inconel 718, International Journal
of Advanced Manufacturing Technology, 38 (2008) 1–2, 74–84,
doi:10.1007/s00170-007-1084-y
12 T. Asokan, S. S. Reddy, P. D. E. Costa, Electrical discharge drilling
of titanium alloys for aerospace applications, Proceedings of 19th
AIMTDR conference, IIT Madras, Chennai, 2000, 161–165
13 D. T. Pham, S. S. Dimov, S. Bigot, A. Ivanov, K. Popov, Micro-
EDM-recent developments and research issues, Journal of Materials
Processing Technology, 149 (2004), 50–57, doi:10.1016/j.jmatprotec.
2004.02.008
14 E. Bamberg, S. Heamawatanachai, Orbital electrode actuation to
improve efficiency of drilling micro-holes by micro-EDM, Journal of
Materials Processing Technology, 209 (2009), 1826–1834,
doi:10.1016/j.jmatprotec.2008.04.044
15 B. H. Yan, C. C. Wang, The machining characteristics of Al2O3/
6061Al composite using rotary electro-discharge machining with a
tube electrode, Journal of Materials Processing Technology, 95
(1999) 1–3, 222–231, doi:10.1016/S0924-0136(99)00322-2
16 W. Koenig, R. Weill, R. Wertheim, W. I. Jutzler, The flow fields in
the working gap with electro-discharge-machining, Annals of the
CIRP, 25 (1977) 1, 71–76
17 O. Yýlmaz, A. T. Bozdana, M. A. Okka, Ý. H. Filiz, An intelligent and
automated system for EDM hole drilling of super alloys, 5th Inter-
national Conference on Responsive Manufacturing – Green Manu-
facturing, Ningbo, China, 2010
18 D. Canadinc, H. Sehitoglu, H. J. Maier, Y. I. Chumlyakov, Strain
hardening behavior of aluminum alloyed Hadfield steel single
crystals, Acta Materialia, 53 (2005), 1831–1842, doi:10.1016/
j.actamat.2004.12.033
19 T. Kivak, Optimization of surface roughness and flank wear using the
Taguchi method in milling of Hadfield steel with PVD and CVD
coated inserts, Measurement, 50 (2014), 19–28, doi:10.1016/
j.measurement.2013.12.017
20 K. H. Ho, S. T. Newman, State of the art electrical discharge machin-
ing (EDM), International Journal of Machine Tools & Manufacture,
43 (2003), 1287–1300, doi:10.1016/S0890-6955(03)00162-7
21 S. L. Chen, B. H. Yan, F. Y. Huang, Influence of kerosene and di-
stilled water as dielectrics on the electric discharge machining cha-
racteristics of Ti-6Al-4V, Journal of Materials Processing Tech-
nology, 87 (1999) 1–3, 107–111, doi:10.1016/S0924-0136(98)
00340-9
22 N. Mohri, M. Suzuki, M. Furuya, N. Saito, Electrode wear process in
electrical discharge machining, Annals of the CIRP, 44 (1995) 1,
165–168, doi:10.1016/S0007-8506(07)62298-7
V. YILMAZ et al.: DEEP MICRO-HOLE DRILLING FOR HADFIELD STEEL ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 377–386 385
23 M. Znidarsic, M. Junkar, Deep Small Hole Drilling with EDM, Pro-
ceedings of 4th International Conference on Advanced Manufactu-
ring Systems and Technology, Udine, 1996, 527–533
24 H. S. Liu, B. H. Yan, C. L. Chen, F. Y. Huang, Application of micro-
EDM combined with high-frequency dither grinding to improve
micro-hole machining, International Journal of Machine Tools &
Manufacture, 46 (2006), 80–87, doi:10.1016/j.ijmachtools.2005.03.
017
25 A. Kadirvel, P. Hariharan, Optimization of the die-sinking micro-
EDM process for multiple performance characteristics using the
Taguchi-based grey relational analysis, Mater. Tehnol., 48 (2014) 1,
27–32
26 M. Sarikaya, A. Gullu, Taguchi design and response surface me-
thodology based analysis of machining parameters in CNC turning
under MQL, Journal of Cleaner Production, 65 (2014), 604–616,
doi:10.1016/j.jclepro.2013.08.040
27 M. Sarikaya, H. Dilipak, A. Gezgin, Optimization of the process
parameters for surface roughness and tool life in face milling using
the Taguchi analysis, Mater. Tehnol., 49 (2015) 1, 139–147
28 M. Sarikaya, V. Yilmaz, H. Dilipak, Modeling and multi-response
optimization of milling characteristics based on Taguchi and gray
relational analysis, Proc IMechE Part B: J Engineering Manufacture,
(2015), doi:10.1177/0954405414565136
V. YILMAZ et al.: DEEP MICRO-HOLE DRILLING FOR HADFIELD STEEL ...
386 Materiali in tehnologije / Materials and technology 49 (2015) 3, 377–386
M. M. RISTOVA et al.: SURFACE ANALYSIS OF ELECTROCHROMIC CuxO FILMS ...
SURFACE ANALYSIS OF ELECTROCHROMIC CuxO FILMS IN
THEIR COLORED AND BLEACHED STATES
POVR[INSKA ANALIZA ELEKTROKROMIZNIH PLASTI CuxO V
NJIHOVIH OBARVANIH IN OBELJENIH STANJIH
Mimoza M. Ristova1,2, Milorad Milun3, Biljana Pejova4
1Faculty of Natural Sciences and Mathematics, Institute of Physics, P.O. Box 162, Skopje, R. Macedonia
2Materials Science Division, Lawrence Berkeley National Laboratory, Berkeley, CA, USA
3Institute of Physics, Bijenicka Cesta 46, Zagreb, Croatia
4Faculty of Natural Sciences and Mathematics, Institute of Chemistry, P.O. Box 162, Skopje, R. Macedonia
mima.ristova@gmail.com
Prejem rokopisa – received: 2014-06-25; sprejem za objavo – accepted for publication: 2014-07-04
doi:10.17222/mit.2014.092
CuxO is known as an electrochromic material with a possible applicability for solar-light modulation. The reversible transition
between the two different oxidation states, CuO and Cu2O, is responsible for the visible-light switching ability. CuxO films in
their as-prepared, colored and bleached states were subjected to a surface analysis in order to relate the bleaching/coloring
effects to the quantified Cu-oxide transition. An XPS analysis on the Cu2p electrons of the as-prepared, bleached and colored
films showed that the Cu-ion quantity reversibly turning from CuO to Cu2O during the electrochromic cycling was about 3.4 %.
An analysis of the XRD patterns of the film's three states confirmed that a small portion of the surface Cu-atoms probably
participate in the coloration/bleaching process. Scanning electron microscopy (SEM) images revealed obvious changes in the
surface morphology due to bleaching and coloration transitions, particularly in the grain size and porosity of the CuxO films.
The surface morphology of the films was also studied with the atomic force microscopy (AFM). This technique allowed
significant conclusions to be derived relating to the surface roughness as well as the compositional homogeneity of the films
before and after electrochemical treatments. These results appeared to be complementary to those derived from the X-ray
diffraction patterns. One may assume that the coloration centers are located at very few film's monolayers of the interface with
the electrolyte.
Keywords: electrochromism, Cu2O, CuO, XPS, XRD, SEM, AFM
CuxO je poznan kot elektrokromizni material z mo`nostjo uporabe za modulacijo son~ne svetlobe. Reverzibilni prehod med
dvema oksidacijskima stanjema, CuO in Cu2O, vpliva na mo`nost preklapljanja vidne svetlobe. Tanke plasti CuxO v izhodnem,
obarvanem ali obeljenem stanju so bile analizirane na povr{ini, da bi ugotovili u~inke barvanja/beljenja pri kvantificiranem
prehodu Cu-oksidov. XPS-analiza z elektroni Cu2p pripravljene, obeljene in obarvane plasti je pokazala, da se okrog 3,4 %
koli~ine Cu-ionov med elektrokromizno obdelavo reverzibilno odmika od CuO proti Cu2O. Analiza XRD-sledov treh stanj traku
je potrdila, da verjetno majhen dele` Cu-atomov sodeluje v postopku obarvanje/beljenje. Vrsti~na elektronska mikroskopija
(SEM) je potrdila ob~utne spremembe v morfologiji povr{ine zaradi prehodov obarvanja in obeljenja, posebno velikost zrn in
poroznost plasti CuxO. Morfologija povr{ine plasti je bila pregledana z mikroskopijo na atomsko silo (AFM). Ta tehnika
omogo~a pomembne sklepe glede povr{inske hrapavosti, kot tudi homogenosti sestave plasti pred elektrokemijsko obdelavo in
po njej. Ti rezultati se ujemajo s tistimi, dobljenimi z XRD-posnetkov. Lahko sklepamo, da so centri obarvanja locirani v nekaj
monoplasteh na stiku z elektrolitom.
Klju~ne besede: elektrokromizem, Cu2O, CuO, XPS, XRD, SEM, AFM
1 INTRODUCTION
Electrochromic materials are known as modulators of
the reflection/transmission of the incident illumination1.
There are many applications of electrochromic films, one
of the most important being the production of large-area
electrochromic displays2. Rear-view mirrors with a vari-
able reflectance, based on electrochromic oxide films,
are commercially available for many kinds of vehicles2.
The most attractive, very useful and environmental
friendly applications of electrochromic materials are the
so-called "smart windows"3 that are able to automatically
modulate the incoming solar illumination in the interior.
Smart windows may also be alternatively powered by
photovoltaic cells, thus, operating as energetically inde-
pendent devices.
A copper (I) oxide (Cu2O) thin film has been a sub-
ject of research in numerous studies, as a candidate for a
solar cell application4. It is known that copper oxide thin
films exhibit cathode electrochromism5–8, i.e., they are
transparent for visible light in their oxidized state and
almost black in their reduced state.
Thin CuxO films can be deposited with different tech-
niques: sputtering8, electrochemical deposition9,10, sol-
gel-like dip technique5,11, thermal oxidation12, anode oxi-
dation13 or chemical-deposition method14–17. Chemically
deposited CuxO films were the subject of our previous
research17 and are the topic of our present interest. Our
former publication18 showed some of the advantages and
disadvantages of this electrochromic material for solar-
light-modulation applications, implying the necessity for
a more profound surface examination of their light-
switching mechanism.
Materiali in tehnologije / Materials and technology 49 (2015) 3, 387–393 387
UDK 532.6:669.058 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)387(2015)
2 EXPERIMENTAL WORK
2.1 Electrochromic CuxO film preparation
Fluorine-doped tin oxide films on glass (FTO) with a
sheet resistance of 10–20  and transmittance for the
visible light of about 80–85 %, products of Solar-Split,
Croatia, were used as the substrates. Electrochromic
CuxO films were deposited with the electroless-chemi-
cal-bath deposition method, described elsewhere15–17.
The thickness of CuxO was estimated from the SEM
image of the cross-section of the glass/film interface.
2.2 Surface analysis of the Cu2O electrochromic films
The X-ray diffraction patterns of the CuxO films on
the SnO2 (FTO) substrates in their as-prepared, bleached
and colored states were recorded with a Rigaku Ultima
IV X-ray diffractometer, using Cu-K radiation of
0.15418 nm. XRD scans were taken within an angular
interval 2 from 20 ° to 70 ° via a step method applying
0.02 ° steps (2) and a counting time of 0.6 s per scan.
Scattered X-ray radiation was detected using a DteX
detector. The XRD pattern of the FTO (SnO2) substrate
was recorded as well. In addition to the identification,
the XRD patterns were also used for estimating the ave-
rage crystal size with the Scherrer method.
The X-ray photoelectron spectroscopy (XPS)
instrument consisted of a non-chromatized X-ray source
with an Mg-Al dual anode. A hemispherical electrostatic
electron analyzer VSW HA100 was employed under the
following operating conditions: a fixed-analyzer-transi-
tion (FAT) mode of 50 eV and a resolution of 1.1 eV. The
XPS spectra were taken from the CuxO electrochromic
films deposited onto the conductive FTO surface, in their
as-prepared, bleached and colored states. The Mg-K
XPS spectra were taken from the film surfaces with no
previous Ar+ ion-beam cleaning, taking into account that
the electrochromic modifications possibly take place on
the immediate surface of the interface with the elec-
trolytes. The binding energy of the Cu2p electrons was
subjected to an analysis. Taking into account that the
Cu2p electron binding energy shifts due to the presence
of various copper-oxygen compounds each of the com-
pounds was quantified.
Scanning electron microscopy (SEM) was used for
an estimation of the CuxO film thickness by taking the
scans from the film/glass substrate cross-section profile.
Surface SEM scans were taken from the surface of the
CuxO/substrate as well as from the sole substrate. The
SEM analyses of the scans taken from the CuxO/FTO
samples in their as-prepared, bleached and colored states
were used for depicting the variation in the crystal grain
size and the porosity. A digitalized system of JEOL
JSM-T220A SEM was used.
The surface morphology and phase composition of
the investigated CuxO films, in their as-prepared, colored
and bleached state, were studied with atomic force
microscopy (AFM). The height, amplitude error and
phase images were measured using a scanning probe
Shimadzu microscope SPM 9600 operating in a dynamic
and phase mode. The measurements were performed
using silicon SPM probes with a resonance frequency of
320 kHz and a force constant of 42 N/m. Several
different regions on the investigated sample surface were
explored. The scan rate was 1 Hz or 2 Hz depending on
the size of the scanned area (1 μm or 5 μm). The image
resolution was 512 lines per each scan direction. The
measured images were only flattened, without any
further processing. The surface roughness was calculated
using the SMP Manager data processing software.
3 RESULTS AND DISCUSSION
3.1 Film bleaching and coloration
Three CuxO films were deposited under equal condi-
tions onto the FTO substrate. One of the films was kept
in the as-prepared condition for further examination. The
second film was brought into the bleached state as it was
biased with +1 V against the FTO adjacent electrode in a
0.2 M KNO3 solution for 10 seconds. The third film was
fully colored upon biasing with –1 V against the FTO
adjacent electrode in the 0.2 M KNO3 solution for 10 s.
The three films were subjected to a surface analysis,
described in section 3.2. It is worth mentioning that the
films retained their electrochromic colors for a long time
(over one year) upon the completion of the surface anal-
ysis.
3.2 Surface analyses of the CuxO films (XRD, XPS,
SEM and AFM)
The XRD patterns of the CuxO electrochromic films,
deposited onto the FTO substrate in their as-prepared,
bleached and colored state, are presented in Figure 1.
The SnO2 (FTO) substrate alone was also analyzed. The
corresponding CPDS files for CuO, Cu2O and SnO2 are
given in19–21.
As can be seen from Figure 1, only three peaks were
identified in addition to the one of the SnO2 substrate.
The analysis showed that all three detectable peaks, low
in intensity, pertained to the Cu2O cuprite phase. The
small Cu2O peak intensity and the absence of the iden-
tifiable peaks typical for the tenorite crystalline oxide
state of copper (CuO) did not allow us to draw any con-
clusions about the electrochromic reversible switching
between the two Cu-oxide states. Thus, one may specu-
late that the transition from the bleached to the colored
state and its reverse process occur through the known
transition of only a minor (sub-detectable) portion of the
dominant oxide of Cu2O into CuO.
The diffraction pattern of the as-prepared CuxO
sample shows two diffraction peaks between the 2
angles of 42 ° and 36 °, which are characteristic for cop-
per (I) oxide. Following the fitting procedure, the angular
position and the full width at half maximum (FWHM) of
M. M. RISTOVA et al.: SURFACE ANALYSIS OF ELECTROCHROMIC CuxO FILMS ...
388 Materiali in tehnologije / Materials and technology 49 (2015) 3, 387–393
detectable diffraction peaks were found. The average
crystal diameter was estimated to be about 24 nm using
the Scherrer equation.
Figure 2 presents broad XPS spectra of the CuxO
film in its as-prepared, bleached and colored states. The
distinct peaks of the binding energy were found to
originate from all the elements of the film compound,
such as Cu (the 2p3/2 main peak centered at around 932
eV and a strong satellite centered at 943 eV), O (O1s)
and C (C1s). The expected Auger peak of Cu L3M4.5M4.5
is not sufficiently distinct on either of the three sample
broad spectra). The three Cu2p3 spectra were corrected
by fixing the C1s peak to 285 eV. Each of the corrected
XPS Cu2p3 spectra was then deconvoluted into two
Gaussian peaks corresponding to CuO, known to appear
between 934.2 eV 22–24 and 935.2 eV 21, and to Cu2O,
known to appear between 932.3 eV and 933.8 eV 22–25.
Since the Cu0 and Cu2O Cu2p3 peaks can not be resolved
with deconvolution (the binding energy shift between the
two is only about 0.1 eV)24, Cu(OH)2 can not be clearly
distinguished from CuO at 935.3 eV 22 due to their over-
lap. Disregarding the presence of the copper metal and
hydroxide, we can assume that the entire Cu2p signal
originates from the CuO and Cu2O compounds. Each of
the two Gaussian-peak areas (A1 and A2) of the Cu2p3/2
zero-leveled XPS spectra on Figure 3 is proportional to
M. M. RISTOVA et al.: SURFACE ANALYSIS OF ELECTROCHROMIC CuxO FILMS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 387–393 389
Figure 1: XRD patterns of the CuxO electrochromic films deposited
onto the FTO substrate in their as-prepared, bleached and colored
state. The pattern of the substrate and the corresponding CPDS files
are given on the graph.
Slika 1: XRD-posnetki elektrokromiznih plasti, nanesenih na FTO-
podlago v njihovem pripravljenem, obeljenem in obarvanem stanju.
Na sliki so prikazani vzorec podlage in ustrezne CPDS-datoteke.
Figure 3: Deconvolution of the XPS spectra of Cu2p3 electrons on
two Gaussian peaks: a) as-prepared, b) bleached and c) colored elec-
trochromic CuxO film on FTO substrates. A1 and A2 are the corres-
ponding peak areas.
Slika 3: Upadanje XPS-spektra elektronov Cu2p3 na dveh Gaussovih
vrhovih: a) pripravljene, b) obeljene in c) obarvane elektrokromizne
plasti CuxO na FTO-podlagi. A1 in A2 sta odgovarjajo~i podro~ji
vrhov.
Figure 2: Broad XPS spectra of the CuxO film in its as-prepared,
bleached and colored states
Slika 2: [irokopasovni XPS-spekter plasti CuxO v pripravljenem,
obeljenem in obarvanem stanju
the signal yield from the Cu atoms bonded in either CuO
or Cu2O. Hence, the relative peak area of either A1 or A2
versus the total peak area (A1+A2) in percent can be
considered a measure for the Cu+1 (as in CuO) and Cu+2
(as in Cu2O) quantities in the film. A possible presence
of the elemental copper and hydroxide was neglected
because these are not active participants in the reversible
electrochromic redox reactions.
The results for the Cu2p binding energies of the three
states of the electrochromic CuxO film from Figure 3 are
M. M. RISTOVA et al.: SURFACE ANALYSIS OF ELECTROCHROMIC CuxO FILMS ...
390 Materiali in tehnologije / Materials and technology 49 (2015) 3, 387–393
Figure 5: a) SEM of the CuxO film grown on an amorphous substrate
(glass), b) SEM of the cross-section of the CuxO film on a glass inter-
face. The estimated film thickness is about 250 nm.
Slika 5: a) SEM-posnetek CuxO filma, nastalega na amorfni podlagi
(steklo), b) SEM-posnetek prereza plasti CuxO na steklu. Debelina
plasti je okrog 250 nm.
Table 1: Binding energies of Cu2p3/2 electrons, obtained from XPS spectra with deconvolution on two Gaussian peaks with areas A1 and A2.
Reference values are given for comparison.
Tabela 1: Energije vezav elektronov Cu2p3/2, dobljenih iz XPS-spektra z dekonvolucijo v dva Gaussova vrhova A1 in A2. Za primerjavo so
prikazane referen~ne vrednosti.
B. E. Cu2p in Cu+2
bonding (as in CuO)
(eV)
B. E. Cu2p in Cu+1
bonding (as in Cu2O)
(eV)
ΔB. E.
(eV)
Fraction of Cu+2 (as in
CuO) of the total Cu
atoms A2/(A1+A2)
(%)
Fraction of Cu+1 (as in
Cu2O) of the total Cu
atoms =A1/(A1+A2)
(%)
XPS handbook22 933.6
Literature –XPS
native oxides23 934.2 932.5 1.7 22.1 77.9
T. Ghodselahi et al.24 934.5–935.2 932.3–933.8 1.4–2.2
I. G. Casella et al.25 932.2–932.8
As-prepared 935.5 933.7 1.8 9.8 90.2
Colored 935.5 933.7 1.8 6.8 93.2
Bleached 935.3 933.7 1.6 10.2 89.8
Figure 4: SEM micrographs taken of the surfaces of the: a) as-pre-
pared, b) bleached and c) colored films on FTO substrates
Slika 4: SEM-posnetki povr{ine plasti na FTO-podlagi: a) priprav-
ljeno, b) obeljeno in c) obarvano
summarized in Table 1. It is evident that the experimen-
tal values revealed somewhat greater binding energies
(about 1.2 eV) than those reported for the bulk samples23,
but they are in good agreement with those reported for
the films24,25. The results showed that the binding ener-
gies pertaining to the two Cu2p3/2 peaks retained the
known difference of about 1.7 eV 23,24.
From Table 1 it is clear that the as-prepared sample
contained 90.2 % Cu+1 (bonded as in Cu2O, the yellowish
transparent oxide). From this point, it is evident that the
amount of Cu+1 slightly increased in the bleached sam-
ple, up to 93.2 %, but decreased in the colored sample, to
89.8 %. By the same time, the amount of Cu+2 atoms (as
in the CuO black-colored oxide) decreased in the
bleached sample and increased in the colored sample,
correspondingly. Hence, it can be assumed that the elec-
trochromic switching occurs due to the transition of only
about 3.4 % Cu atoms (93.2–89.8 %) from Cu+ to Cu+2
and vice versa. This result is in agreement with the XRD
patterns, revealing no detectable peaks of the CuO phase
due to their relatively small quantity compared to those
of the Cu2O phase. Furthermore, the characteristic O1s
peaks of the XPS spectrum were practically useless for
the CuO and Cu2O quantification because the O1s peak
at 530.6 eV 23 can be associated with both the CuO and
Cu2O phases, while the peak at 531.6 eV can only be
associated with the electrochromic inactive Cu(OH)2
compound.
If one takes into account the crystallography data for
the unit cells of cuprite-Cu2O and tenorite-CuO, one can
make an estimation of the possible variations in the vol-
ume that may induce the strain in the CuxO structure and,
hence, invoke film degradation due to repeatable cycling.
Under the assumption that the CuxO electrochromic
behavior relies on only 3.4 % of the Cu atoms, we can
estimate the relative changes in the crystallite volume, in
which about 3.4 % Cu atoms from tenorite (a monoclinic
structure with a unit cell volume Vc (CuO) = 81.03 · 10–3
nm3 19) recrystallize into cuprite (a cubic structure with a
unit cell volume Vc (Cu2O) = 77.83 · 10–3 nm3 20). One
can assume that the crystallites within the film undergo a
negligible relative volume reduction during the colora-
tion and a volume expansion during the bleaching.
Hence, the limited cycling lifetime that was previously
observed in these films17,18 cannot be ascribed to the
volume friction within the crystallites.
Figures 4a to 4c present the SEM micrographs of the
surfaces of the as-prepared, bleached and colored films
on the FTO substrates, correspondingly. The micrograph
on Figure 4a reveals round grains of the as-prepared
film with a diameter of about 250 nm. The grain size of
the bleached films on Figure 4b seems similar. However,
the micrograph of the colored-sample surface from
Figure 4c reveals notably smaller round crystal grains.
M. M. RISTOVA et al.: SURFACE ANALYSIS OF ELECTROCHROMIC CuxO FILMS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 387–393 391
Figure 7: 2D and 3D surface images of FTO substrates
Slika 7: 2D- in 3D-posnetek povr{ine FTO-podlage
Figure 6: Two-dimensional surface images of: a) as-prepared, b) colored and c) bleached CuxO thin films
Slika 6: Dvodimenzionalni posnetek povr{ine tanke plasti CuxO: a) pripravljeno, b) obarvano in c) obeljeno
Furthermore, as a consequence of the grain shrinkage,
the porosity of the colored film seems to have increased
notably (an obviously larger empty space among the
grains upon the coloration). In addition, Figure 5 shows
the SEMs of: (a) the CuxO films deposited onto the glass
(amorphous) substrate and (b) a profile of about 250 nm
CuxO film on the glass substrate (the interface
cross-section). From Figure 5a it is obvious that the
CuxO film grows irregularly and non-homogeneously on
an amorphous substrate, such as glass. However, the
SEMs on Figure 4, along with the XRD patterns,
showed a mixed amorphous-crystalline growth on the
crystalline substrate (FTO).
Two-dimensional AFM surface images of the as-pre-
pared, bleached and colored CuxO thin films are pre-
sented in Figure 6. For a comparison, Figure 7 includes
2D and 3D surface images of the FTO substrate. The
scale of all the three images is 5 μm. As can be seen
from Figure 6, there are significant differences between
the surfaces of the as-prepared, bleached and colored
samples. In addition, from Figure 8 one can see that the
electrochemical voltage cycling notably affects the
surface roughness. In other words, the oxidation process
(coloration), taking place on a sample surface, deteri-
orates the surface smoothness, increasing the root-
mean-square roughness. On the other hand, the reduction
process (bleaching) is followed by an improvement of
the surface smoothness. Furthermore, the roughness
parameter (Rq) derived from the AFM scans of the
observed surfaces (an area of 5 μm × 5 μm) decreased
from 480 nm for the colored to 108 nm for the bleached
sample.
In order to compare the compositional homogeneity
of the CuxO thin films in their as-prepared, colored and
bleached state, their phase surface images are presented
on Figure 9, correspondingly. Besides Cu2O and CuO,
some minor amounts of Cu(OH)2 were found in all three
samples. Furthermore, the results showed that the cuprite
(Cu2O) phase is the most abundant in the bleached state,
which is in agreement with the XPS results of this study.
Comparing the phase images of the as-prepared and
colored states, it could be concluded that the amount of
the tenorite (CuO) phase is comparable in both the
as-prepared and colored states which is again in line with
the XPS results.
4 CONCLUSION
The coloration process of the chemically deposited
CuxO films can be considered as a reversible transfer of
the film 3–4 % of the total 90 % Cu+1 into Cu+2. The
coloration can thus be attributed to a decrease of 3–4 %
M. M. RISTOVA et al.: SURFACE ANALYSIS OF ELECTROCHROMIC CuxO FILMS ...
392 Materiali in tehnologije / Materials and technology 49 (2015) 3, 387–393
Figure 9: Phase surface images of CuxO thin films: a) as-prepared, b) colored and c) bleached. The scale is 5 μm in both directions and 90 nm in
height.
Slika 9: Povr{inska slika faz v tanki plasti CuxO: a) pripravljeno, b) obarvano in c) obeljeno. Merilo je 5 μm v obeh smereh in 90 nm po vi{ini.
Figure 8: Three-dimensional surface images of CuxO thin films: a) as-prepared, b) colored and c) bleached
Slika 8: Tridimenzionalen posnetek povr{ine tanke plasti CuxO: a) pripravljeno, b) obarvano in c) obeljeno
of the interstitial oxygen, creating oxygen vacancies in
the CuO film. From all the above findings, it appears that
only 3–4 % of the copper atoms in the film surface
represent the coloration centers that are driven to switch
the transmittance between 30 % and 80 % 17,18. The SEM
micrographs showed that the grain size of the CuxO
crystallites notably shrank upon the coloration of the
film, whereas the porosity grew up. Significant changes
in the surface morphology among the as-prepared and
treated samples were also detected by AFM. While the
oxidation process (the coloring) deteriorated the surface
smoothness, the reduction (the bleaching) invoked the
surface smoothening. The AFM phase imaging revealed
the existence of three different phases in the as-depo-
sited, colored and bleached films: Cu2O, CuO and a mi-
nor amount of the hydroxide phase – Cu(OH)2. In
agreement with the XPS results, the amount of the
dominant cuprite phase was found to be the highest for
the bleached state, while the tenorite phase (CuO)
appeared in comparable amounts in the as-prepared and
the colored states. The limited quantity of the electro-
chromic active Cu-atoms within the thin CuxO films on
the transparent conductive oxide (TCO) limits its perfor-
mance in the white-light-transmittance modulation bet-
ween 20 % and 80 %. If the CuxO films are sensitized as
nanocrystals within a TCO matrix, it can be expected
that a more efficient coloration can be achieved.
5 REFERENCES
1 C. G. Granqvist, Handbook of Inorganic Electrochromic Materials,
Elsevier, Amsterdam 1995
2 J. I. Pankove, Display Devices, Springer, Berlin 1980
3 C. M. Lampert, T. R Omstead, P. C. Yu, Sol. Energy Mater., 14
(1986), 161–174, doi:10.1016/0165-1633(86)90043-2
4 A. E. Rakhshani, J. Appl. Phys., 62 (1987), 1528–1529, doi:10.
1063/1.339619
5 N. Ozer, F. Tepehan, Sol. Energy Mater. Sol. Cells, 30 (1993), 1–26,
doi:10.1016/0927-0248(93)90027-Z
6 H. Demiryont, US Patent No. 4830471, 16 May 1989
7 F. I. Brown, S. C. Schulz, US Patent No. 5585959, 17 Dec. 1996
8 T. J. Richardson, J. L. Slack, M. R. Rubin, Electrochim. Acta, 46
(2001), 2281–2284, doi:10.1016/S0013-4686(01)00397-8
9 M. E. Abu-Zeid, A. E. Rakhshani, A. A. Al-Jassar, Y. A. Youssef,
Phys. Stat. Solidi (A), 93 (1986), 613-620, doi:10.1002/pssa.
2210930226
10 V. Georgieva, M. Ristov, Sol. Energy Mater. Sol. Cells, 73 (2002),
67–73, doi:10.1016/S0927-0248(01)00112-X
11 S. C. Ray, Sol. Energy Mater. Sol. Cells, 68 (2001), 307–312,
doi:10.1016/S0927-0248(00)00364-0
12 W. M. Sears, E. Fortin, Sol. Energy Mater., 10 (1984), 93–103,
doi:10.1016/0165-1633(84)90011-X
13 E. Fortin, D. Masson, Solid-State Electron., 25 (1982), 281–283,
doi:10.1016/0038-1101(82)90136-8
14 A. Roos, T. Chibuye, B. Karlsson, Sol. Energy Mater., 7 (1983),
453–465, doi:10.1016/0165-1633(83)90018-7
15 M. Ristov, Gj. Sinadinovski, I. Grozdanov, Thin Solid Films, 123
(1985), 63–67, doi:10.1016/0040-6090(85)90041-0
16 M. Ristova, J. Velevska, M. Ristov, Sol. Energy Mater. Sol. Cells, 71
(2002), 219–230, doi:10.1016/S0927-0248(01)00061-7
17 R. Neskovska, M. Ristova, J. Velevska, M. Ristov, Thin Solid Films,
515 (2007), 4717–4721, doi:10.1016/j.tsf.2006.12.121
18 M. Ristova, R. Neskovska, V. Mirceski, Sol. Energy Mater. Sol.
Cells, 91 (2007) 14, 1361–1365, doi:10.1016/j.solmat.2007.05.018
19 Powder Diffraction File No. 4680, CPDS International Center for
Diffraction Data, Newtown Square, 2006
20 Powder Diffraction File No. 1104, CPDS International Center for
Diffraction Data, Newtown Square, 1985
21 Powder Diffraction File No. 778, CPDS International Center for
Diffraction Data, Newtown Square, 1975
22 G. E. Muileenberg (Ed.), Handbook of X-ray photoelectron spectro-
scopy, Perkin – Elmer Corporation, 1979
23 XPS Handbook of elements and native oxides, XPS International
Incorporated, 1999
24 T. Ghodselahi, M. A. Vesaghi, A. A. Shafiekhani, A. Baghizadeh, M.
Lameii, Applied Surface Science, 255 (2008), 2730–2734,
doi:10.1016/j.apsusc.2008.08.110
25 I. G. Casella, M. Gatta, Journal of Electroanalytical Chemistry, 494
(2000), 12–20, doi:10.1016/S0022-0728(00)00375-2
M. M. RISTOVA et al.: SURFACE ANALYSIS OF ELECTROCHROMIC CuxO FILMS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 387–393 393

M. OPIELA: THERMODYNAMIC ANALYSIS OF THE PRECIPITATION OF CARBONITRIDES ...
THERMODYNAMIC ANALYSIS OF THE PRECIPITATION OF
CARBONITRIDES IN MICROALLOYED STEELS
TERMODINAMSKA ANALIZA IZLO^ANJA KARBONITRIDOV V
MIKROLEGIRANIH JEKLIH
Marek Opiela
Silesian University of Technology, Institute of Engineering Materials and Biomaterials, Konarskiego Street 18a, 44-100 Gliwice, Poland
marek.opiela@polsl.pl
Prejem rokopisa – received: 2014-07-01; sprejem za objavo – accepted for publication: 2014-07-08
doi:10.17222/mit.2014.096
The production of mass-products of microalloyed steels with a high strength requires a proper adjustment of the thermo-
mechanical processing and the kinetics of the precipitation of the MX-type (M – microalloying element, X – metalloid) phases
in austenite. The understanding of the effect of carbonitrides on the processes of hot-working and cooling from the finishing
deformation temperature requires the knowledge of the mechanism of their formation and their stability in austenite. The
research was carried out on newly manufactured Ti-V and Ti-Nb-V steels for forged machine parts with thermomechanical
processing. The analysis of the precipitation of the carbonitrides with the stoichiometric compositions of TixV1–xCyN1–y and
TixNbvV1–x–vCyN1–y was based on the Hillert-Staffanson model improved by Adrian. The effect of the austenitizing temperature
in the range from 900 °C to 1200 °C on the grain size of the original austenite was investigated to verify the results of the
calculation. The result provides the basis for a suitable design of the manufacturing process of the thermomechanical treatment
to obtain high-strength forged elements of microalloyed steels.
Keywords: microalloyed steels, MX-type phases, thermomechanical processing
Proizvodnja masovnih izdelkov iz mikrolegiranih jekel z visoko trdnostjo zahteva pravilno prilagoditev termomehanske
predelave in kinetike izlo~anja faze MX (M – mikrolegirni element, X – metaloid) v avstenitu. Razumevanje vpliva karbo-
nitridov na procese vro~e predelave in ohlajanja iz temperature kon~ne deformacije zahteva poznanje mehanizma njihovega
nastanka in stabilnosti v avstenitu. Raziskava je bila izvr{ena na novo izdelanih jeklih za odkovke Ti-V in Ti-Nb-V po postopku
termomehanske obdelave. Analiza izlo~anja karbonitridov s stehiometri~no sestavo TixV1–xCyN1–y in TixNbvV1–x–vCyN1–y temelji
na Hillertovem in Staffansonovem modelu, ki ga je izbolj{al Adrian. Preiskovan je bil vpliv temperature avstenitizacije v
obmo~ju 900 °C do 1200 °C na prvotno velikost avstenitnih zrn, da bi preverili rezultate izra~unov. Rezultati omogo~ajo prido-
bitev osnove za pravilno na~rtovanje postopka izdelave s termomehansko obdelavo odkovkov z veliko trdnostjo iz mikro-
legiranih jekel.
Klju~ne besede: mikrolegirana jekla, faze vrste MX, termomehanska obdelava
1 INTRODUCTION
The production of metallurgical products with high
mechanical properties from microalloyed steels requires
the conditions of plastic working to be adjusted to the
kinetics of the dissolution (precipitation) of the MX type
of micro-additions introduced into steel. The solubility
of the MX phases, i.e., nitrides, carbides and carbo-
nitrides of the alloying elements such as Ti, Nb, V, Zr or
B in austenite is determined with the logarithm of the
solubility product expressed with the equation:1–5
lg [M] · [X] = B – A/T (1)
where [M] and [X] are the mass fractions of the metallic
micro-addition and the metalloid dissolved in austenite
at temperature T, respectively, A and B are the constants
associated with the free enthalpy of the MX-phase
formation. It should be noted that constants A and B
depend on the method of determination; for this reason
slightly different values for the same phase can be found
in6,7.
High technical usability belongs to the diagrams of
different MX phases dissolved in austenite, determined
for different steels on the basis of Equation (1) as well as
to the computer programs calculating the chemical com-
positions and fractions of these phases as a function of
temperature. The chemical composition of austenite
([M], [X]) and the portion of the undissolved compound
can be determined using the solubility product and the
mass balance of the elements included in the reaction. A
simplified thermodynamic model based on the laws of
thermodynamic equilibrium, concerning the precipitation
of the compound, is described with a system of equat-
ions:8
[M] · [X] = kMX (2)
M M
(M)
(MX)
< MX >= +[ ] (3)
X X
(M)
(MX)
< MX >= +[ ] (4)
where kMX = [M] · [X] indicates the solubility product of
MX, M and X indicate the total concentrations of these
elements in steel in mass fractions (w/%), the atomic
mass of the elements, as the symbols in parentheses, ( ).
Materiali in tehnologije / Materials and technology 49 (2015) 3, 395–401 395
UDK 621.77:536.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)395(2015)
The solution of this system of equations is a
quadratic function that allows us to calculate the amount
of the interstitial element dissolved in austenite:
[ ] +[ ]2X X
(X)
(M)
M X
(X)
(M) MX
⋅ ⋅
⎡
⎣⎢
⎤
⎦⎥
− =– k 0 (5)
By avoiding the negative solution and  > 0, the
equation has two solutions. However, in the case of  =
0, the austenitizing temperature is higher than, or equal
to, the temperature of the dissolution of the compound.
The part of the interstitial element dissolved in austenite
with  > 0 is determined from the following equation:
[ ] =
X
X
(X)
(M)
M +
(X)
(M)
M X
(X)
(M)
2
MX– –⋅ ⋅
⎛
⎝
⎜ ⎞
⎠
⎟ + ⋅
2
4 k
(6)
The models of the precipitation processes in micro-
alloyed steels, based on the laws of thermodynamic
equilibrium, also include the formation of complex car-
bonitrides during hot working. The carbides and nitrides
of the Nb, Ti and V micro-additions reveal mutual solu-
bility and, as a result of this process, carbonitrides are
formed, with the chemical composition and dissolution
temperature dependent on the chemistry of the steel.9
The precipitation of the complex carbonitrides with a
MCyN1–y stoichiometric constitution is determined with a
thermodynamic model based on the assumptions of the
sublattice model created by Hillert and Staffanson10, with
the basic simplifying assumption that the metallic ele-
ment M and the interstitial elements (C, N) in the steel
form dilute solutions in austenite and that their activities
meet Henry’s law. The final form of the equations of this
model describes the state of thermodynamic equilibrium
in the Fe-M-C-N system:
ln ( )
yk
y
L
RT
MC CN
M
M C[ ] [ ]a a⋅
⎛
⎝
⎜
⎞
⎠
⎟ + − =1 02 (7)
ln
( )1
02
−
⋅
⎛
⎝
⎜
⎞
⎠
⎟ + =
y k
y
L
RT
MN CN
M
M N[ ] [ ]a a
(8)
where [Ma], [Na] and [Ca] indicate the atomic fractions
of the metallic element [M] and interstitial elements [N]
and [C] dissolved in austenite, kMN and kMC indicate the
reaction equilibrium constants, LMCN is the interaction
parameter of the M element affecting C-N, and the y and
(1 – y) solutions indicate the MC and MN moles, res-
pectively. The system of Equations (7) and (8) contains
four unknowns: [Ma], [Na], [Ca] and y. In the solutions
based on the mass balance of carbonitrides, the
following reactions are used:
M Ma a[ ]= + −
f
f
2
1( ) (9)
C Ca a[ ]= + −
yf
f
2
1( ) (10)
N Na a[ ]=
−
+ −
( )
( )
1
2
1
y f
f (11)
with the amount fraction of the MCyN1–y precipitations.
Complex M’xM”1–xCyN1–y or M’xM”vM’”1–x–vCyN1–y
type carbonitrides can be formed in microalloyed steels
containing two or three micro-additions (M’, M’’, M’’’)
at the same time. A thermodynamic model describing the
state of thermodynamic equilibrium in steel containing
up to three micro-additions and Al was elaborated by
Adrian11,12. It can calculate the chemical composition of
austenite and the concentration of carbonitride in a
Fe-Nb-V-Ti-Al-C-N multicomponent system. The final
form of the equations determining the thermodynamic
equilibrium of the Fe-M’-M”-M’”Al-C-N system is as
follows:13
y
xyk
y
x y k
y
ln ( ) ln
( )M'C M'N
M' C M' N[ ] [ ] [ ] [ ]a a a a⋅
+ −
−
⋅
+
+
1
1
( )1 0− =y
L
RT
CN
(12)
y
vyk
y
v y k
y
ln ( ) ln
( )M"C M"N
M" C M" N[ ] [ ] [ ] [ ]a a a a⋅
+ −
−
⋅
+
+
1
1
( )1 0− =y
L
RT
CN
(13)
y
x v yk
y
x v y k
ln
( )
( ) ln
( )( )1
1
1 1− −
⋅
+ −
− − −M"'C M"
M'" C[ ] [ ]a a
'N
CN
M"' N[ ] [ ]a a⋅
+
+ − =y y
L
RT
( )1 0
(14)
vy
x k
v k
x v y
x v
ln ( )( ) ln
( )[ ]
[ ]
[a
a
aM"'
M
M'M'C
M'C
+ − − −
− −
1 1
1 ]
[ ]a
k
X k
M'N
M'NM"'
+ −
−
⋅
+ − =( ) ln
( )
( )1
1
1 02y
x y k
y y
L
RT
M'N CN
M' N[ ] [ ]a a
(15)
[ l ] [ ] =a a AlNA N⋅ k (16)
The system of equations consists of eleven unknowns
describing the chemical compositions of [M’a], [M”a],
[M’”a], [Ala], [Ca], [Na] of austenite and (x, v, y)
carbonitrides. The subsequent equations, describing the
law of conservation of mass during the reaction of the
carbonitride and nitride AlN formations are necessary
for the solution:
M' M'a a a[ ]= + − −
x
f f f
2
1( ) (17)
M" M"a a a[ ]= + − −
v
f f f
2
1( ) (18)
M"' M"'a a a[ ]= + − −
z
f f f
2
1( ) (19)
C Ca a a[ ]= + − −
y
f f f
2
1( ) (20)
M. OPIELA: THERMODYNAMIC ANALYSIS OF THE PRECIPITATION OF CARBONITRIDES ...
396 Materiali in tehnologije / Materials and technology 49 (2015) 3, 395–401
N Na a a[ ]=
−
+ − −
1
2
1
y
f f f( ) (21)
Al Aa
a
a a[ l ]= + − −
f
f f f
2
1( ) (22)
where f is the amount fraction of carbonitride, fa is the
amount fraction of nitride AlN, kMX is the product of the
MX compound solubility converted into amount frac-
tions (xMX/%), the values in square brackets [ ] indicate
the concentrations of the elements in the solution in
amount fractions (%), the values in parentheses ( )
indicate the atomic mass of the elements and the values
without any brackets indicate the total concentrations of
the elements in steel in mount fractions (%).
The input data for the solution of the above system of
equations are: the chemical composition of steel, the
austenitizing temperature, the kMX product of the solubi-
lity of simple carbides and nitrides, MX, the parameter
of the impact of the M element on the C-N solution, LCN
= –4260 J/mol, and the R gas constant.
The presented model can be used for the solutions of
the technological problems associated with the produc-
tion of microalloyed steels, by designing and modifying
chemical compositions. The calculation involves the
following parameters: the chemical composition of
austenite, the chemical composition and the fraction of
carbonitrides and the dissolving temperature of carbo-
nitrides.
2 EXPERIMENTAL PROCEDURE
The research was performed on newly manufactured
microalloyed steels (Table 1). Steels with the weight of
100 kg were molten in a VSG-100 type laboratory
vacuum-induction PVA TePla AG furnace. The steels
were cast in argon, forming square ingots with the
dimensions of 160/140 mm × 640 mm and hot worked to
32 mm × 160 mm flat bars, by open die forging in a
high-speed hydraulic press, applying a force of 300 MN.
The range of the forging temperature was 1200–900 °C.
The thermodynamic analysis of the equilibrium of
the structural constituents in the stable austenite of the
steels mainly focused on the analytic calculations of the
austenite chemical composition, and the amounts and
chemical compositions of the potential interstitial phases
of carbides, nitrides or complex carbonitrides, performed
as a function of the heating or cooling temperature. Also,
the calculations based on the Hillert-Staffanson thermo-
dynamic model developed by Adrian for the analysis of
interstitial complex phases were conducted. For the
calculation of the chemical composition of austenite and
the concentration of carbonitride based on the chemistry
of the analyzed microalloyed steels, CarbNit11, the
computer program operating in a Delphi environment
was used.
To verify the performed analysis, the effect of the
austenitizing temperature on the austenite grain size was
investigated. The samples of 25 mm × 20 mm × 32 mm
were austenitized at the temperatures of (900, 1000,
1100 and 1200) °C for 30 min and water-quenched. To
reveal the grain boundaries, the samples were etched in a
saturated water solution of picric acid with an addition of
CuCl2 at a temperature of 60 °C. The metallographic
observations of the etched specimens were carried out
using a Leica MEF 4A light microscope applying the
magnifications of 200–800-times.
3 RESULTS AND DISCUSSION
The micro-additions of Nb, Ti and V in the investi-
gated steels form carbides, nitrides or simple and com-
plex carbonitrides, most often with a NaCl cubic lattice
and complete intersolubility. These phases nucleated
both heterogeneously, on the existing precipitates such as
TiN and TiC created at an increased temperature, and in
a homogeneous, independent way.
The concentration of Ti in the examined steel was
selected to bond with all the nitrogen. For this reason,
VN, NbN and AlN were not included in the thermo-
dynamic analysis of the single interstitial phases. In
addition, AlN does not dissolve in carbonitrides as it has
a different, hexagonal lattice. The solubility of particular
phases of the austenite of microalloyed steels is deter-
mined with the logarithmic dependence (Equation 1).
The calculated solubility temperatures of the investigated
phases, formed in the examined A and B microalloyed
steels are listed in Table 2 with respect to the solubility
products of the TiN, TiC and VC type phases precipi-
tated in steel A and also of NbC in steel B. The maxi-
M. OPIELA: THERMODYNAMIC ANALYSIS OF THE PRECIPITATION OF CARBONITRIDES ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 395–401 397
Table 1: Chemical compositions of the investigated steels in mass fractions, w/%
Tabela 1: Kemijska sestava preiskovanih jekel v masnih dele`ih, w/%
Steel
Mass fractions, w/%
C Mn Si P S Cr Ni Mo Nb Ti V B Al
A 0.31 1.45 0.30 0.006 0.004 0.26 0.11 0.22 – 0.033 0.008 0.003 0.040
B 0.28 1.41 0.29 0.008 0.004 0.26 0.11 0.22 0.027 0.028 0.019 0.003 0.025
Table 2: Summary of the solubility temperatures of individual
interstitial phases
Tabela 2: Pregled temperature topnosti posameznih intersticijskih faz
No. Type ofMX phase
Constants in Equation
(1)
Solubility
temperature, °C
A B Steel A Steel B
1. TiN 15490 5.19 1350 1331
2. TiC 10745 5.33 1188 1167
3. NbC 7900 3.42 – 1137
4. VC 9500 6.72 734 776
mum solubility temperatures of the TiN type interstitial
phases in the  base were 1350 °C and 1331 °C for the A
and B steels, respectively. Thus, it is assumed that a high
austenitizing temperature – close to 1200 °C – does not
cause a significant grain growth of the austenite through
the undissolved TiN fraction and the high heating and
forging temperatures increase the durability of forging
tools.
The NbC carbide in the austenite of the B steel with
the maximum solubility temperature of 1137 °C contri-
butes essentially to the steel strengthening through a
grain refinement and precipitation hardening with a
simultaneous decrease in the ductile-to-brittle-transition
temperature. In turn, the VC carbide is completely
dissolved in the austenite of the examined steels in the
range of hot-working temperature.
Due to the mutual solubility of the interstitial phases,
a complex TixV1–xCyN1–y carbonitride can form in steel A
with the micro-additions of Ti and V, under the
conditions of thermodynamic equilibrium, and a
TixNbvV1–x–vCyN1–y complex carbonitride forms in the B
steel with the micro-additions of Ti, Nb and V (Table 3).
The calculation results of the chemical composition of
the austenite of the examined microalloyed steels, name-
ly, the determination of the temperature dependence of
the amounts of the metallic element [M] = f(T) and
non-metallic element [N] = f(T), as well as the assumed
chemical constitution of the TixV1–xCyN1–y and
TixNbvV1–x–vCyN1–y type complex carbonitrides as y = f(T)
with the defined volume fractions of the analyzed VV =
f(T) phases are depicted in Figures 1 and 2.
Table 3: Calculated stoichiometric thermodynamic-equilibrium com-
positions of complex carbonitrides in the investigated steels
Tabela 3: Izra~unane stehiometri~ne termodinamsko ravnote`ne
sestave kompleksnih karbonitridov v preiskovanih jeklih
No.
Tempe-
rature
°C
Type of carbonitrides
Steel A Steel B
1 850 Ti0.854V0.146C0.645N0.355 Ti0.600Nb0.292V0.108C0.748N0.252
2 900 Ti0.874V0.126C0.610N0.390 Ti0.646Nb0.301V0.053C0.697N0.303
3 950 Ti0.890V0.110C0.598N0.402 Ti0.668Nb0.308V0.024C0.680N0.320
4 1000 Ti0.904V0.096C0.580N0.420 Ti0.681Nb0.300V0.019C0.662N0.338
5 1050 Ti0.917V0.083C0.550N0.450 Ti0.705Nb0.282V0.013C0.628N0.372
6 1100 Ti0.930V0.070C0.502N0.498 Ti0.736Nb0.258V0.006C0.578N0.422
7 1150 Ti0.942V0.058C0.432N0.568 Ti0.786Nb0.214C0.481N0.519
8 1200 Ti0.954V0.046C0.349N0.651 Ti0.851Nb0.149C0.350N0.650
9 1250 Ti0.963V0.037C0.272N0.728 Ti0.897Nb0.103C0.244N0.756
10 1300 Ti0.970V0.030C0.210N0.790 Ti0.915Nb0.085C0.180N0.820
11 1350 Ti0.975V0.025C0.165N0.835 Ti0.923Nb0.077C0.146N0.854
The analysis of thermodynamic equilibrium leads to
the conclusion that the volume fraction Vv(T) of
Ti0.978V0.022C0.146N0.854 increases with a decrease in the
temperature to approximately amount fraction 0.15 % at
850 °C in steel A, cooled from 1378 °C (Figure 1). The
M. OPIELA: THERMODYNAMIC ANALYSIS OF THE PRECIPITATION OF CARBONITRIDES ...
398 Materiali in tehnologije / Materials and technology 49 (2015) 3, 395–401
Figure 1: Calculation results for: a) TixV1–xCyN1–y carbonitride, y = f(T), b) volume fraction of (Ti, V)(C, N) carbonitride and BN, VV = f(T), c,
d) elements dissolved in austenite: [M] = f(T) and [N] = f(T); steel A
Slika 1: Rezultati izra~unov: a) TixV1–xCyN1–y karbonitrid, y = f(T), b) volumenski dele` (Ti, V)(C, N) karbonitrida in BN, VV = f(T), c, d)
elementi, raztopljeni v avstenitu: [M] = f(T) in [N] = f(T); jeklo A
amount of V in the carbonitride increases from amount
fractions 0.02 % to 0.15 % cooled from 1378 °C to 850
°C. In the temperature range from 1378 °C to 856 °C, the
amount of carbon in the carbonitride increases from
amount fractions 0.14 % to 0.65 %, while in the same
temperature range the amount of nitrogen decreases from
amount fractions 0.85 % to 0.36 %. At 1100 °C, the
amounts of C and N are similar, being 0.502 % and
0.498 %, respectively. The calculated stoichiometric
compositions, in the conditions of thermodynamic
equilibrium, for the analyzed (Ti, V)(C, N) carbonitrides
are listed in Table 3.
The thermodynamic-equilibrium solubility tempera-
ture of Ti0.928Nb0.072C0.140N0.860 in steel B is 1361 °C (Fig-
ure 2). The thermodynamic analysis of the precipitation
of the investigated carbonitrides allows us to state that its
volume fraction in austenite VV = f(T) increased up to
about amount fraction 0.14 % due to a decrease in the
temperature to 850 °C. The decrease in the temperature
from 1300 °C to 1100 °C causes a relative decrease in
the Ti concentration and an increase in the Nb concen-
tration. By further lowering the temperature to 850 °C a
mild decrease in the Ti amount and a slight increase in
the Nb amount occur; an increase in the V amount
fraction to 0.11 % in the analyzed carbonitride occurs in
the temperature range from 1100 °C to 850 °C. By
analyzing the amounts of the interstitial elements in the
investigated carbonitride, it was found that the carbon
concentration at 1361 °C is 0.14 % and it increases up to
0.75 % at 850 °C. In the examined temperature range the
M. OPIELA: THERMODYNAMIC ANALYSIS OF THE PRECIPITATION OF CARBONITRIDES ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 395–401 399
Figure 3: Austenite grains and MX particles in steel A; austenitizing
temperatures of: a) 900 °C and b) 1200 °C
Slika 3: Avstenitna zrna in MX-delci v jeklu A; temperatura avste-
nitizacije: a) 900 °C in b) 1200 °C
Figure 2: Calculation results of the chemical compositions of: a) TixNbvV1–x–vCyN1–y carbonitride, y = f(T), b) volume fraction of (Ti, Nb, V)(C,
N) carbonitride and BN, VV = f(T), c, d) elements dissolved in austenite: [M] = f(T) and [N] = f(T); steel B
Slika 2: Rezultati izra~unov kemijske sestave: a) TixNbvV1–x–vCyN1–y karbonitrid, y = f(T), b) volumenski dele` (Ti, Nb, V)(C, N) karbonitrida in
BN, VV = f(T), c, d) elementi, raztopljeni v avstenitu: [M] = f(T) in [N] = f(T); jeklo B
changes in the nitrogen amount are opposite to the
changes in the amount of carbon.
The effect of the temperature from 900 °C to 1200 °C
on the austenite grain size was investigated to verify the
results of the calculation of the precipitation of MX
particles. Austenite grains and particles are presented in
Figures 3 and 4, while Figure 5 shows the effect of the
temperature on the average size of the austenite grains at
a given temperature. The size of the austenite grains at
900 °C to 1000 °C is very small, from 11 μm to 18 μm in
the Ti-V steel and 8 μm to 12 μm in the Ti-Nb-V steel.
This fine-grained microstructure is a result of the pre-
sence of significant fractions of the (Ti, V)(C, N)
particles and complex (Ti, Nb, V)(C, N) carbonitrides
(respectively for the A and B steels) that effectively
inhibit the growth of the austenite grains. An increase in
the temperature causes a gradual growth in the austenite
grains, more distinct for the Nb-free steel (Figure 5).
The average diameters of the austenite grains after auste-
nitizing the steels at the temperature of 1200 °C are 66
μm and 62 μm for the A and B steel, respectively. These
values are several times lower compared to the C-Mn
steels, where a typical austenite grain size at 1200 °C is
reported to be from 200 μm to 300 μm.1,4 The observed
grain sizes of the austenite for the investigated steels are
also similar to those reported for the other Ti-V or Ti-Nb
steels.4,14–17
The austenite grain growth in steel A is mainly con-
trolled with the (Ti, V)(C, N) particles generally precipi-
tating above 1200 °C. The curve in Figure 5 for the Nb
steel is characteristic for microalloyed steels and the
grain growth can be combined with the precipitation
process of complex carbonitrides. Due to the gradual
dissolution of the Nb-rich MX-type phases, the austenite
grain size increases gradually with the increasing
austenitizing temperature and the growth is much faster
above 1100 °C. At 1200 °C the fractions of the particles
in both steels are comparable, consisting only of the
Ti-rich MX-type phases that lead to similar austenite
grain sizes.
4 CONCLUSIONS
The simplified-thermodynamic-model analysis of the
precipitation of the MX-type phases in both steels found
the highest thermal stability in the austenite of TiN. The
TiN precipitation starts at around 1350 °C. The TiC
precipitation occurs in the temperature range from 1150
°C to 1200 °C. It was shown in1,4,14,17,18 that the MX-type
phases show mutual solubility. Hence, the austenite grain
growth in steel A should be controlled with complex (Ti,
V)(C, N) carbonitrides. More complex carbonitrides are
found in the Ti-Nb steel and the grain growth in steel B
is controlled with the precipitation of the (Ti, Nb, V)(C,
N) particles.
The analysis of the precipitation of the MX-type
phases in austenite allows us to select a proper forging
temperature range, which should correspond to the tem-
perature range of the precipitation of these phases. High
mechanical properties of forged parts can be achieved
with an appropriate selection of the forging conditions,
i.e., the temperature of charge heating and the plastic
deformation range since the distribution of the strain and
strain rate during the production of die forgings with a
M. OPIELA: THERMODYNAMIC ANALYSIS OF THE PRECIPITATION OF CARBONITRIDES ...
400 Materiali in tehnologije / Materials and technology 49 (2015) 3, 395–401
Figure 5: Influence of the austenitizing temperature on the austenite
grain size
Slika 5: Vpliv temperature avstenitizacije na velikost avstenitnih zrn
Figure 4: Austenite grains and MX particles in steel B; austenitizing
temperatures of: a) 900 °C and b) 1200 °C
Slika 4: Avstenitna zrna in MX delci v jeklu B; temperatura avste-
nitizacije: a) 900 °C in b) 1200 °C
complex shape is difficult to adjust. Forge heating should
not lead to a total dissolution of the interstitial elements
in the solid solution because it causes a grain growth.
Hence, the A and B steels may be heated up to the
forging temperature of 1200 °C, without an excessive
growth of the austenite grains.
The investigation of the influence of the austenitizing
temperature on the austenite grain size confirms the
correctness of the precipitation analysis of the MX
phases in the investigated steels. The steels have a fine-
grained austenite over the whole investigated austeni-
tizing temperature range. In both steels the austenite
grain size grows with the increasing temperature and it is
several times lower in comparison to the C-Mn steels. A
faster grain growth is observed for the Ti-V steel. In the
Nb steel, the grain growth is delayed due to Ti and Nb
combined microalloying, resulting in a slower dissolu-
tion of the complex (Ti, Nb, V)(C, N) particles.
The performed analysis of the precipitation process
of the MX-type phases provides the basis for a suitable
design of the thermomechanical processing of micro-
alloyed steels for high-strength forged machine parts.
5 REFERENCES
1 T. Gladman, The physical metallurgy of microalloyed steels, 1st ed.,
The University Press, Cambridge 1997
2 J. Adamczyk, Journal of Achievements in Materials and Manu-
facturing, 20 (2007), 399–402
3 M. Opiela, A. Grajcar, Arch. Civ. Mech. Eng., 12 (2012), 327–333,
doi:10.1016/j.acme.2012.06.003
4 R. Kuziak, T. Bold, Y. Cheng, J. Mater. Proc. Tech., 53 (1995),
255–262, doi:10.1016/0924-0136(95)01986-L
5 J. Adamczyk, E. Kalinowska-Ozgowicz, W. Ozgowicz, R. Wusatow-
ski, J. Mater. Proc. Tech., 53 (1995), 23–32, doi:10.1016/0924-0136
(95)01958-H
6 H. Adrian, Proc. of the Inter. Conf. Microalloying ’95, Iron and Steel
Soc., Pittsburgh, 1995, 285–305
7 E. J. Palmiere, Proc. of the Inter. Conf. Microalloying ’95, Iron and
Steel Soc., Pittsburgh, 1995, 307–320
8 D. A. Skobir, M. Godec, M. Balcar, M. Jenko, Mater. Tehnol., 44
(2010) 6, 343–347
9 A. Grajcar, S. Lesz, Mater. Sci. Forum, 706–709 (2012), 2124–2129,
doi:10.4028/www.scientific.net/MSF.706-709.2124
10 M. Hillert, L. I. Staffanson, Acta Chem. Scand., 24 (1970),
3618–3636, doi:10.3891/acta.chem.scand.24-3618
11 H. Adrian, Scientific Editions of University of Mining and Metal-
lurgy, Kraków, 1995
12 H. Adrian, Mater. Sci. Technol., 8 (1992), 406–415
13 M. Opiela, Journal of Achievements in Materials and Manufacturing,
47 (2011), 7–18
14 D. K. Matlock, G. Krauss, J. G. Speer, J. Mater. Proc. Tech., 117
(2001), 324–328, doi:10.1016/S0924-0136(01)00792-0
15 M. MacKenzie, A. J. Craven, C. J. Collins, Scri. Mater., 54 (2006),
1–5, doi:10.1016/j.scriptamat.2005.09.018
16 P. S. Bandyopadhyay, S. K. Ghosh, S. Kundu, S. Chatterjee, Metall.
Mater. Trans. A, 42 (2011), 2742–2752, doi:10.1007/s11661-011-
0711-2
17 A. Nowotnik, T. Siwecki, J. Microsc., 237 (2008), 258–262,
doi:10.1111/j.1365-2818.2009.03238.x
18 M. Opiela, A. Grajcar, Arch. Civ. Mech. Eng., 12 (2012), 427–435,
doi:10.1016/j.acme.2012.06.013
M. OPIELA: THERMODYNAMIC ANALYSIS OF THE PRECIPITATION OF CARBONITRIDES ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 395–401 401

V. KHARCHENKO et al.: EXPERIMENTAL INVESTIGATION OF THE CRACK-INITIATION MOMENT ...
EXPERIMENTAL INVESTIGATION OF THE CRACK-INITIATION
MOMENT OF CHARPY SPECIMENS UNDER IMPACT LOADING
EKSPERIMENTALNA PREISKAVA TRENUTKA INICIACIJE
RAZPOKE PRI UDARNI OBREMENITVI CHARPYJEVIH
VZORCEV
Valeriy Kharchenko, Evgeniy Kondryakov, Alexandr Panasenko
Institute for Problems of Strength, National Academy of Sciences, Tymiryazevska str. 2, Kyiv, Ukraine
alaenonn@ya.ru
Prejem rokopisa – received: 2014-07-09; sprejem za objavo – accepted for publication: 2014-09-05
doi:10.17222/mit.2014.103
A new experimental method for the investigation of crack initiation and propagation was developed. Strain gauges on a
specimen surface made it possible to obtain surface-deformation data. These data were compared with the force-time curve
obtained using a vertical instrumented impact tester. An analysis of the results of this comparison enabled us to determine the
moment of crack initiation on the force-time curve. Investigations of crack propagation were conducted in two orthogonal
directions.
Keywords: Charpy specimens, crack, moment of crack initiation, specific zones of fracture
Razvita je bila nova eksperimentalna metoda za preiskavo iniciacije in rasti razpoke. Merilni listi~i na povr{ini vzorca
omogo~ajo pridobivanje podatkov o deformaciji povr{ine. Ti podatki so bili primerjani s krivuljo sila – ~as, dobljeno iz
instrumentirane vertikalne udarne naprave. Analiza rezultatov te primerjave omogo~a dolo~itev trenutka iniciacije razpoke na
krivulji sila – ~as. Preiskave rasti razpoke so bile izvr{ene v dveh ortogonalnih smereh.
Klju~ne besede: Charpyjevi vzorci, razpoka, trenutek iniciacije razpoke, posebna podro~ja na prelomu
1 INTRODUCTION
At present, impact bending tests are one of the
simplest and cheapest methods to determine the material
properties that describe its tendency to brittle fracture.
The simplicity and efficiency of this method, a relative
ease of calibration and adjustment of the equipment
allow this test method to be used in many areas of
science and technology and industrial sectors, particu-
larly, in the programs aimed to identify and predict the
properties of reactor-vessel materials based on sur-
veillance-specimen tests.1–3 Obtaining load diagrams for
the contact between a specimen and a striker under
impact loading with the use of strain gauges and modern
recording systems implemented in a vertical instru-
mented drop-weight impact-testing machine, as well as a
further comparison of the data-analysis results with the
results of fractographic investigations, make it possible
to get more important information about the crack-
propagation mechanisms in Charpy specimens.4,5
2 MATERIALS
The present paper describes the results of an investi-
gation of the moment of crack initiation and the features
of crack propagation in Charpy V-notch specimens in the
course of impact-bending testing on a vertical instru-
mented drop-weight impact-testing machine. The sche-
me, the methods and the test procedures are described in
detail.6
Standard Charpy V-notch specimens were used. The
specimen material is hot-rolled sheet of steel 45, with a
microstructure of a ferrite-pearlite mixture (Figure 1).
The mean diameter of ferrite grains is 35 μm, and the
mean diameter of pearlite grains is 54 μm.
Materiali in tehnologije / Materials and technology 49 (2015) 3, 403–408 403
UDK 620.1:539.42 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)403(2015)
Figure 1: Microstructure of a specimen of steel 45, 500-times, 1 –
ferrite; 2 – pearlite
Slika 1: Mikrostruktura vzorca jekla 45, pove~ava 500-kratna, 1 –
ferit; 2 – perlit
3 METHODS AND EXPERIMENTAL WORK
To investigate the moment of crack initiation and the
features of its propagation, two directions of the crack
propagation were chosen – the direction of the rolling
plane and the one perpendicular to it.
Specimens were produced according to the State
Standard of Ukraine GOST-9454-78. The dimensions of
the specimens are 55 cm × 10 cm × 10 cm. The notch
parameters are: the notch depth 2 mm, the convergence
angle 45 °, the curvature radius near the notch tip 0.25
mm. The specimens were produced from a metal sheet
with collinear directions of the principal axis of a
specimen and the direction of the metal-sheet rolling
plane. The plane of crack propagation was perpendicular
to the direction of the rolling plane. Specimens with two
different orientations of the crack-propagation direction
were produced (Figure 2). The specimen marked with
"T" has a notch orientation (and direction of crack pro-
pagation) perpendicular to the rolling plane. The speci-
men marked with "S" has a notch orientation (and
direction of crack propagation) parallel to the rolling
plane.
The crack propagation during the fracture of the
Charpy specimens under impact bending has some speci-
fic features.7,8 With instrumented impact tests, diagrams
of the impact contact force with the specimen force time
P(t) are obtained and their analysis enables us to
describe every stage of the crack propagation in detail. A
comparison of the diagram specific zones with the spe-
cific zones of the fracture makes it possible to calculate
the specific energy for the crack propagation in a given
zone and relate the energy to the fracture mechanism.5,9
To determine the moment of crack initiation (brittle
or ductile) in the P(t) diagram, additional gauges for
measuring lateral deformation were used. The gauges
were located on the specimen surface near the V-notch
concentrator along the assumed crack-propagation front
(Figure 3). The signal recording the channels for the
gauge and the P(t) diagram were synchronized in time.
The discretization of the signal in time was 2e-7 s (the
sampling frequency was 5 · 106 Hz). The temperature on
the specimen surface during the testing was measured
with a chromel/alumel thermocouple.
Fractographic investigations of the specimen fracture
at the macro-level were performed using an Axiotech-
Vario microscope, while the micro-zones of the speci-
men fracture were studied using a SEM-100U micro-
scope.
4 RESULTS AND DISCUSSION
The height of the blade fall for the tested group
varied from 0.05 m to 0.5 m. A height (h) of 0.1 m was
used for a ductile/brittle crack initiation inside a
specimen without an expansion on the lateral surface of
the specimen. A height of 0.05 m was used for a
shear-lip formation on the lateral surface without a crack
initiation inside the specimen. A height of 0.5 m was
used for the final rupture of the specimens with a brittle
crack initiation followed by ductile crack growing.
All the experimental data with the parameters of the
specimens and the test conditions are shown in Table 1.
Figures 4 and 5 show the diagrams of the gauge
signal and P(t) and Figures 6 and 7 show the enlarged
zones (in time) in the vicinity of the diagram maximum
values for specimens ShK7-1 (T = 25.8 °C, h = 1 m, the
concentrator orientation is S) and ShK7-2 (T = 35 °C, h
= 1 m, the concentrator orientation is S). The maximum
V. KHARCHENKO et al.: EXPERIMENTAL INVESTIGATION OF THE CRACK-INITIATION MOMENT ...
404 Materiali in tehnologije / Materials and technology 49 (2015) 3, 403–408
Figure 4: Specimen ShK7-1, h = 50 cm: 1 – diagram of the striker
contact with the specimen, 2 – diagram of the gauge signal
Slika 4: Vzorec ShK7-1, h = 50 cm: 1 – diagram stika kladiva z
vzorcem, 2 – diagram signala iz merilnega listi~a
Figure 3: Half of the specimen with a strain gauge
Slika 3: Polovica vzorca z merilnim listi~em
Figure 2: Scheme of the produced specimens with two different notch
orientations (marked as S and T) relative to the direction of the rolling
plane
Slika 2: Shematski prikaz izdelave vzorcev z dvema orientacijama
razpoke (oznaka S in T) glede na smer ravnine valjanja
force is situated close to the beginning of drastic force
reduction.
The diagrams for specimen ShK7-2 have similar
shapes (Figure 5). The results of the specimen testing at
various temperatures revealed that the maximum value of
the gauge diagram has a tendency to shift mainly to the
left, with an increase in the temperature, in the direction
of the maximum value of the P(t) diagram.
The fractographic analysis revealed several typical
zones on the fracture surfaces: a stable crack-growth
area, an unstable crack-jump area, a rupture area and a
shear-lip area. Figure 8 presents a macro-fracture of
specimen ShK7-1. An analysis of the physical interpre-
tation of the strain-gauge diagram suggests that the
essential feature of the fracture of these specimens is the
form of a stable crack-growth area, which is of an
elongated triangle or trapezoid with the maximum length
(in the direction of the main crack propagation) in the
central section of the specimen and with the minimum
length (frequently close to zero) at the fracture edges,
i.e., close to the lateral surfaces of the specimen (the
shear-lip zone).
Thus, a crack initiation is most likely to occur in the
middle of a specimen with the subsequent extension to
the surface, which relates to the formation of lateral
necking (shear lips).
To confirm the given assumption, additional strain
tests on the Charpy specimens with and without a crack
initiation were conducted. In the light of this fact it is
obvious that the strain gauge responds to the strain of a
specimen surface during the brittle-crack propagation (an
unstable crack-jump area). It is improbable that the
gauge will register an initiation of a crack of a stable
growth considering its current shape and dimensions.
V. KHARCHENKO et al.: EXPERIMENTAL INVESTIGATION OF THE CRACK-INITIATION MOMENT ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 403–408 405
Table 1: Experimental data and test conditions
Tabela 1: Eksperimentalni podatki in razmere pri preizku{anju
Specimen Notchorientation
Temperature,
°C
Height of the
blade fall, h/m
Impact
velocity, m/s Experimental result
ShK7-1 S 25.8 0.5 3.13 Rupture with a brittle crack initiation
ShK7-2 S 35 0.5 3.13 Rupture with a brittle crack initiation
ShK7-3 T 26.6 0.05 0.99 No crack initiation, a shear-lip formation
ShK7-3 T 26.6 0.1 1.4
Brittle-crack initiation in the center without
an expansion on the lateral surface of the
specimen
ShK7-3 T 26.6 0.5 3.13 Final rupture after an initiated brittle crack
ShK7-4 T 23 0.1 1.4 Ductile-crack initiation and shear-lipformation
ShK7-4 T 23 0.5 3.13 Final rupture after an initiated ductile crack
Figure 7: Specimen ShK7-2, h = 50 cm. An enlarged section of the
diagrams: 1 – diagram of the striker contact with the specimen, 2 –
diagram of the gauge signal.
Slika 7: Vzorec ShK7-2, h = 50 cm. Pove~ano podro~je: 1 – diagram
stika kladiva z vzorcem, 2 – diagram signala iz merilnega listi~a.
Figure 5: Specimen ShK7-2, h = 50 cm: 1 – diagram of the striker
contact with the specimen, 2 – diagram of the gauge signal
Slika 5: Vzorec ShK7-2, h = 50 cm: 1 – diagram stika kladiva z
vzorcem, 2 – diagram signala iz merilnega listi~a
Figure 6: Specimen ShK7-1, h = 50 cm. An enlarged section: 1 –
diagram of the striker contact with the specimen, 2 – diagram of the
gauge signal.
Slika 6: Vzorec ShK7-1, h = 50 cm. Pove~ano podro~je: 1 – diagram
stika kladiva z vzorcem, 2 – diagram signala iz merilnega listi~a.
Using these considerations one can explain the fact that
the maximum value in the gauge diagram does not coin-
cide with the maximum value in the contact diagram, as
it is shifted to the right in time with respect to the contact
diagram.
The values for the striker velocity at the moment of
its contact with a specimen, at which the specimen
underwent deformation without a crack initiation, were
determined by varying the values of the height (h) of the
blade fall.
Figure 9 shows the diagrams of the strain gauge and
P(t) for specimen ShK7-3 (T = 26.6 °C, the concentrator
orientation is T) at the impact velocity of 1 m/s. At this
impact velocity no crack initiation is observed (during a
visual inspection using optical methods both on the
surfaces and inside the concentrator). Moreover, the
gauge registered the strain on the specimen surface (a
shear-lip initiation).
It should be mentioned that the maximum value of
the gauge signal was significantly lower (0.32 mV) com-
pared with the maximum value of the gauge signal at the
crack initiation (0.55 mV).
At the impact velocity of 1.4 m/s (Figures 10 and 11)
a crack initiated in the central section of specimen
ShK7-3 (Figure 12). It is obvious that the crack
nucleates in the central section of the specimen (the
V. KHARCHENKO et al.: EXPERIMENTAL INVESTIGATION OF THE CRACK-INITIATION MOMENT ...
406 Materiali in tehnologije / Materials and technology 49 (2015) 3, 403–408
Figure 12: Specimen ShK7-3 (initiated crack), h = 10 cm
Slika 12: Vzorec ShK7-3 (za~etna razpoka), h = 10 cm
Figure 9: Specimen ShK7-3, h = 5 cm: 1 – diagram of the striker
contact with the specimen, 2 – diagram of the gauge signal
Slika 9: Vzorec ShK7-3, h = 5 cm: 1 – diagram stika kladiva z vzor-
cem, 2 – diagram signala iz merilnega listi~a
Figure 10: Specimen ShK7-3, h = 10 cm: 1 – diagram of the striker
contact with the specimen, 2 – diagram of the gauge signal
Slika 10: Vzorec ShK7-3, h = 10 cm: 1 – diagram stika kladiva z
vzorcem, 2 – diagram signala iz merilnega listi~a
Figure 8: Fractograph of a macro-fracture of a Charpy specimen: 1 –
stable crack-growth area, 2 – unstable crack-jump area, 3 – rupture
area, 4 – shear lips
Slika 8: Makroposnetek preloma Charpy vzorca: 1 – podro~je stabilne
rasti razpoke, 2 – nestabilno podro~je skoka razpoke, 3 – prelom pri
upogibu, 4 – stri`ne ustnice
Figure 11: Specimen ShK7-3, h = 10 cm. An enlarged section of the
diagrams: 1 – diagram of the striker contact with the specimen, 2 –
diagram of the gauge signal.
Slika 11: Vzorec ShK7-3, h = 10 cm. Pove~ano podro~je: 1 – diagram
stika kladiva z vzorcem, 2 – diagram signala iz merilnega listi~a.
crack opening has the maximum value – of approxima-
tely 1 mm in the center).
In specimens ShK7-3 and ShK7-4, at the impact
velocity of 1.4 m/s, the initiated cracks did not extend to
the lateral surfaces of the specimens (Figure 12). The
P(t) diagram clearly shows the region of an abrupt
decrease in the force at the moment of brittle-crack
initiation and propagation in the central section of the
specimen.
Then, the specimen was fractured at the impact
velocity of 3 m/s. Figure 13 provides a diagram of the
fracture. Like in the P(t) diagram there is a lack of the
area of deformation/crack initiation; the initial section
has the form of a straight line (elastic deformation) with
its fluctuations. This is related to the presence of the
crack initiated in specimen ShK7-3.
A large area of a stable crack growth and an area of
an unstable crack jump (brittle fracture) formed at the
impact velocity of 1.4 m/s are shown on Figure 14.
Figure 15 shows the diagram of the gauge signal and
the P(t) diagram for specimen ShK7-4. The crack
initiation was similar to the one in specimen ShK7-3
after the impact with the velocity of 1.4 m/s. The crack
initiated inside the specimen and it demonstrated its
maximum opening in the specimen central section, but
did not extend to the lateral surfaces of the specimen
(Figure 16). The gauge-signal diagram has its maximum
value shifted to the left, relative to the maximum value in
the P(t) diagram.
In this case the gauge measuring the strain on the
specimen surface did not register the initiation of a stable
V. KHARCHENKO et al.: EXPERIMENTAL INVESTIGATION OF THE CRACK-INITIATION MOMENT ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 403–408 407
Figure 17: Specimen ShK7-4, h = 50 cm: 1 – diagram of the striker
contact with the specimen, 2 – diagram of the gauge signal
Slika 17: Vzorec ShK7-4, h = 50 cm: 1 – diagram stika kladiva z
vzorcem, 2 – diagram signala iz merilnega listi~a
Figure 14: Specimen ShK7-3 (fracture)
Slika 14: Vzorec ShK7-3 (prelom)
Figure 15: Specimen ShK7-4, h = 10 cm: 1 – diagram of the striker
contact with the specimen, 2 – diagram of the gauge signal
Slika 15: Vzorec ShK7-4, h = 10 cm: 1 – diagram stika kladiva z
vzorcem, 2 – diagram signala iz merilnega listi~a
Figure 13: Specimen ShK7-3, h = 50 cm: 1 – diagram of the striker
contact with the specimen, 2 – diagram of the gauge signal
Slika 13: Vzorec ShK7-3, h = 50 cm: 1 – diagram stika kladiva z
vzorcem, 2 – diagram signala iz merilnega listi~a
Figure 16: Crack nucleation, specimen ShK7-4, h = 10 cm
Slika 16: Nukleacija razpoke, vzorec ShK7-4, h = 10 cm
ductile crack inside the specimen. The crack did not
extend to the specimen lateral surfaces.
The P(t) diagram does not include the area of an
abrupt force decrease, which implies the propagation of
an unstable brittle crack in the specimen.
Specimen ShK7-4 was fractured at the impact velo-
city of 3 m/cm. The P(t) diagram for specimen ShK7-4
(Figure 17) is similar to the diagram for specimen
ShK7-3 (Figure 13).
The deformation/crack initiation site is not observed
and the force increases linearly with the time in the
initial part of the diagram. The maximum value in the
gauge-signal diagram is more to the left compared with
the one for specimen ShK7-3. The signal level is very
low (0.028 mV), which evidences a significant deforma-
tion of the gauge tracks after the impact with the velocity
of 1.4 m/s. At the fracture there is no unstable (brittle)
crack-jump site (Figure 18); however, the gauge regi-
stered the moment of a ductile crack extension to the
specimen lateral surface as demonstrated by its diagram.
The conducted investigations showed that the used
procedure for the determination of the crack-initiation
moment exhibits a number of drawbacks associated
primarily with the formation of lateral necking. To obtain
more accurate results, specimens with deep side grooves
should be used to eliminate the influence of the speci-
men lateral-surface deformation.
5 CONCLUSIONS
The investigations of the moment of crack initiation
in Charpy V-notch specimens were performed. The
procedure of recording the crack initiation during the
impact-bending testing using a vertical instrumented
impact tester was developed and tested. A comparison of
the gauge diagrams with the striker-contact diagrams
made it possible to determine the moment of crack
initiation in the contact diagram. The comparison of the
gauge/deformation diagrams with the results of the
fractographic investigations enabled us to find the
relation between the moment of crack initiation in the
striker contact diagram and the specimen with specific
fracture zones.
It was determined that a crack in a Charpy specimen
formed during the impact-bending testing is initiated in
the specimen central section. The applicability of the
given procedure for recording both a brittle-crack
initiation in the specimen central section and a propaga-
tion of a ductile crack to the specimen lateral surface was
shown.
There are large differences between the material
properties of the specimens with the T-orientation and
S-orientation. This difference can be revealed during an
impact-loading test. The shapes of the diagrams of the
specimens with different orientations of the notch are
significantly different.
6 REFERENCES
1 PNAE-G-007-86 Norms based on the strength of equipment and
pipelines of nuclear power plants, Energoatomizdat, Moscow 1989
2 VERLIFE - Unified Procedure for Lifetime Assessment of Compo-
nents and Piping in WWER NPPs, European Commission under the
Euratom Research and Training Programme on Nuclear Energy,
Version 8, 2008, 275 p
3 RD EO 1.1.2.09.0789-2009 – Method of determining fracture tough-
ness according to test results witness samples underlying strength
and life of WWER - 1000 NPPs, Energoatomizdat, Moscow 2009
4 Standard ISO EN DIN 14 556 Instrumented Impact Test, Deuth
Verlag, Berlin 2000
5 R. Chaouadi, A. Fabry, On the utilization of the instrumented Char-
py impact test for characterization the flow and fracture behavior of
reactor vessel steels, European Structural Integrity Society, 30
(2002), 103–117, doi:10.1016/S1566-1369(02)80011-5
6 V. Kharchenko, E. Kondryakov, V. Babutskii, V. Zhmaka, An instru-
mented testing machine for impact tests: basic elements, operation
analysis, Reliability and Life of Machines and Structures, 27 (2006),
121–130
7 V. Kharchenko, E. Kondryakov, A. Panasenko, Crack propagation
peculiarities in steels at Charpy and disc-shaped specimens tests,
Questions of nuclear science and technology, 84 (2013) 2, 31–38
8 L. Botvina, Fracture, Kinetics, Mechanisms, General Laws, Nauka,
Moscow 2008 (in Russian)
9 V. Goritskii, Thermal embrittleness of steel, Metalurizdat., Moscow
2007
V. KHARCHENKO et al.: EXPERIMENTAL INVESTIGATION OF THE CRACK-INITIATION MOMENT ...
408 Materiali in tehnologije / Materials and technology 49 (2015) 3, 403–408
Figure 18: Fracture of specimen ShK7-4
Slika 18: Prelom vzorca ShK7-4
S. LESZ et al.: STRUCTURAL, THERMAL AND MAGNETIC PROPERTIES OF Fe-Co-Ni-B-Si-Nb ...
STRUCTURAL, THERMAL AND MAGNETIC PROPERTIES OF
Fe-Co-Ni-B-Si-Nb BULK AMORPHOUS ALLOY
STRUKTURNE, TERMI^NE IN MAGNETNE LASTNOSTI
MASIVNE AMORFNE ZLITINE Fe-Co-Ni-B-Si-Nb
Sabina Lesz1, Marcin Nabia³ek2, Ryszard Nowosielski1
1Silesian University of Technology, Institute of Engineering Materials and Biomaterials, Konarskiego Street 18a, 44-100 Gliwice, Poland
2Institute of Physics, Czestochowa University of Technology, Av. Armii Krajowej 19, 42-200 Czestochowa, Poland
sabina.lesz@polsl.pl
Prejem rokopisa – received: 2014-07-15; sprejem za objavo – accepted for publication: 2014-09-05
doi:10.17222/mit.2014.108
In the present paper the structure, thermal stability and magnetic properties of the Fe43Co22Ni7B19Si5Nb4 bulk amorphous alloy
were investigated. The investigated alloy was cast as rods with three different diameters. The thermal stability associated with
the glass transition temperature (Tg), crystallization temperature (Tx) and supercooled-liquid region (Tx = Tx – Tg) was
examined with differential scanning calorimetry (DSC). The Curie temperature of the investigated glassy rods was determined
from the results obtained with the DSC method. The magnetic properties and microstructure of the rods were examined with the
vibrating-sample magnetometer (VSM) and X-ray diffraction (XRD) methods, respectively. The crystallization temperature (Tx)
and the glass transition temperature (Tg) as well as the parameter of Tx = Tx – Tg as the criterion of the glass-forming ability
(GFA) of the investigated alloy were determined. The investigated alloys have good soft-magnetic properties.
Keywords: bulk amorphous alloy, structure, thermal and magnetic properties
Predstavljena je preiskava strukture, toplotne stabilnosti in magnetnih lastnosti masivne amorfne zlitine Fe43Co22Ni7B19Si5Nb4.
Preiskovana zlitina je bila ulita kot palice s tremi razli~nimi premeri. Toplotna stabilnost, povezana s prehodom v steklasto
stanje (Tg), temperaturo kristalizacije (Tx) in s podhlajenim podro~jem taline (Tx = Tx – Tg), je bila preiskovana z diferen~no
vrsti~no kalorimetrijo (DSC). Curiejeva temperatura preiskovanih steklastih palic je bila dolo~ena iz rezultatov, dobljenih pri
DSC-metodi. Magnetne lastnosti in mikrostruktura palic so bile preiskane z magnetometrom z vibrirajo~imi vzorci (VSM) in z
metodo rentgenske difrakcije (XRD). Dolo~eni so bili temperatura kristalizacije (Tx) in temperatura prehoda v steklasto stanje
(Tg) ter tudi parameter Tx = Tx – Tg kot merilo sposobnosti tvorbe steklastega stanja (GFA) preiskovanih zlitin. Preiskovane
zlitine imajo dobre mehkomagnetne lastnosti.
Klju~ne besede: masivna amorfna zlitina, struktura, termi~ne in magnetne lastnosti
1 INTRODUCTION
A large number of studies on the development of
soft-magnetic metallic glasses have been carried out over
the last 20 years. It is well recognized that the low
glass-forming ability (GFA) of Fe-based alloys has
limited the potential of using them as engineering ma-
terials. For this reason extensive efforts have been carried
out to improve the GFA of metallic materials and the
understanding of the mechanism of the effects of various
factors on the formation, crystallization, thermal stability
and property of bulk metallic glass (BMG). Bulk
metallic glasses (BMGs) represent a new class of amor-
phous metallic alloys. BMGs are valuable materials for
environmental applications (e.g., solar cells, hydrogen
production, the systems for retention and purification of
dangerous pollutants, the nuclear industry, etc.) and for
industrial applications in different areas (e.g., aerospace,
automotive, electronics, computer, telecommunication
areas, etc.).1–12
These multi-component metallic alloys can be ob-
tained at low cooling rates of 1 K/s to 100 K/s, which
allow an increase in the time (from milliseconds to minu-
tes) before the crystallization, enabling a greater critical
casting thickness (> 1 cm) by conventional mould-
ing.1,2,4–7
Among BMGs, the Fe-based BMGs are more
attractive for application since they do not exhibit only
good properties, such as excellent soft-magnetic proper-
ties, a high strength and a good corrosion resistance, but
are also cheaper in comparison to the other BMGs.1–10
For the preparation of a Fe-based BMG, Fe80B20 is
often used as the starting alloy. Later the Nb metal with a
high melting temperature is added. The additions of
small amounts of Nb to (Fe,Co,Ni)-(B,Si) alloys are
effective for the increase in the GFA through the increase
in the stability of the supercooled liquid against crystalli-
zation.3 A temperature interval of the supercooled-liquid
region Tx has been suggested to evaluate the
glass-forming ability (GFA) of bulk amorphous alloys.
An addition of amount fraction of Nb 4 % was found to
be very effective in improving the GFA of Fe- and
Co-based glassy alloys.7
As BMGs can be produced by adding four and five
elements to the basic ternary alloys, small amounts of the
elements such Ni, Co and Si were added. A partial sub-
stitution of Fe with the other magnetic elements, Ni or
Materiali in tehnologije / Materials and technology 49 (2015) 3, 409–412 409
UDK 544.23:537.622 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)409(2015)
Co, may significantly enhance the GFA and soft-magne-
tic properties of the Fe-based glass-forming alloys. The
metalloid elements of Si and B play a crucial role in the
formation of BMGs. They also affect the GFA, the
thermal stability, the crystallization and the properties of
BMGs. These materials have a strong affinity with the
conventional BMG base elements such Fe and rare-earth
elements, i.e., they have a large, negative heat of mixing
with these base elements. The metalloid elements result
in crystallization, degrading the GFA of the BMGs, but,
on the other hand, due to a small atomic size of the Si
and B atoms, a proper addition can tighten the alloy
structure, stabilizing the alloy against crystallization.3
The Fe-Co based glassy alloys exhibit good soft-
magnetic properties, i.e., a high saturation magnetization
(0.8–1.3 T) and a low coercivity (1–2.5 A/m).3 Magnetic
properties of these alloys are dependent on the Ni and Fe
contents. A decrease in the coercivity (Hc) with the
increasing Co content was found to originate in the
reduction of saturation magnetostriction.3 Coercivity Hc
is proportional to the ratio of saturation magnetostriction
(s) to saturation magnetization (Js), i.e.:8
H V d
Jc
s
s
≈ ⋅Δ 

(1)
and the slope is related to the volume (V) and density
(d) of internal defects in the glassy structure.8
Due to their unique properties, the Fe-Co based
glassy alloys have been commercialized in the following
application fields: precision-mould material, precision-
imprint material, precision-sensor material, precision-
machinery material, surface-coating material, cutting-
tool material, shot penning material, fuel-cell separator
material and so forth.1,2,9,10
In the present paper the structural, thermal and mag-
netic properties of a Fe-Co-Ni-B-Si-Nb bulk amorphous
alloy with a selected chemical composition was investi-
gated.
2 EXPERIMENTAL PROCEDURE
Investigations were carried out on amorphous rods
with a composition of [(Fe0.6Co0.3Ni0.1)0.75B0.2Si0.05]96Nb4.
Fe-based master-alloy ingots with a composition of
[(Fe0.6Co0.3Ni0.1)0.75B0.2Si0.05]96Nb4 were prepared by
induction melting of pure Fe, Co, Ni, Nb and pure B and
Si crystals in an argon atmosphere. The
Fe43Co22Ni7B19Si5Nb4 alloy composition represents the
nominal atomic percentages. The master alloy was
melted in a quartz crucible using an induction coil. Rods
with (1.5, 2.5 and 3) mm diameters were prepared with
the pressure copper-mould casting method.11
The microstructure of the rods was examined with
the X-ray diffraction (XRD) method. The X-ray method
was performed using a Seifert-FPM XRD 7 diffracto-
meter with filtered Co-K radiation.
The thermal stability associated with the glass tran-
sition temperature (Tg), crystallization temperature (Tx)
and supercooled-liquid region (Tx = Tx – Tg) was
examined with differential scanning calorimetry (DSC)
at a heating rate of 0.1 K/s. The Curie temperature of the
investigated glassy rods was determined from the results
obtained with the DSC method.
High-field magnetization curves were measured with
a vibrating-sample magnetometer (VSM) in a magnetic
field up to 2 T. The magnetizing field was parallel to the
sample length to minimize the demagnetization effect.
The magnetization curves were analyzed using the least-
squares method.
3 RESULTS AND DISCUSSION
It was found from the obtained results of the struc-
tural studies performed with X-ray diffraction that the
diffraction patterns of the surface rods with (1.5, 2.5 and
3.0) mm diameters of the Fe43Co22Ni7B19Si5Nb4 alloy
consist of a broad-angle peak, indicating the existence of
an amorphous phase (Figure 1).
The DSC curves determined on the
Fe43Co22Ni7B19Si5Nb4 rods with the diameters of (1.5, 2.5
and 3) mm in the as-cast state for the studied alloy are
shown in Figures 2 to 4, and summarized in Table 1.
Table 1 also gives information about the thermal proper-
ties of the studied amorphous-alloy rods. The onset
crystallization temperatures Tx for the glassy rod samples
with the diameters of (1.5, 2.5 and 3) mm are slightly
different and equal to (828, 827 and 826) K (Figures 2 to
4), respectively. It is seen that Tx decreases from 828 K
to 826 K with an increase in the diameter of the rods. On
the basis of an analysis of DSC curves the glass
transition temperature Tg and supercooled-liquid region
S. LESZ et al.: STRUCTURAL, THERMAL AND MAGNETIC PROPERTIES OF Fe-Co-Ni-B-Si-Nb ...
410 Materiali in tehnologije / Materials and technology 49 (2015) 3, 409–412
Figure 1: X-ray diffraction patterns of the bulk amorphous
Fe43Co22Ni7B19Si5Nb4 rods
Slika 1: Rentgenogrami masivnih amorfnih palic
Fe43Co22Ni7B19Si5Nb4
Tx = Tx – Tg for the glassy rod samples with the dia-
meters of 1.5 mm to 3 mm are determined, too.
The value of the supercooled-liquid region is an experi-
mental parameter that determines the glass-forming
ability of the tested alloy. The glass transition tempera-
ture Tg and supercooled-liquid region Tx for the glassy
rod samples with the diameters of (1.5, 2.5 and 3) mm
are: Tg = 794 K, Tx = 34 K (Figure 2), Tg = 790 K, Tx
= 37 K (Figure 3), Tg = 797 K, Tx = 29 K (Figure 4),
respectively. The value of the Curie temperature TC for
the Fe43Co22Ni7B19Si5Nb4 rods with the diameters of (1.5,
2.5 and 3.0) mm is (652, 650 and 655) K, respectively.
Similar values of Tg and TC were obtained in1,9, where the
results are Tg = 813 K and TC = 643 K for the
[(Fe0.6Co0.3Ni0.1)0.75B0.2Si0.05]96Nb4 alloy in the form of a
rod with a diameter of 4 mm.
S. LESZ et al.: STRUCTURAL, THERMAL AND MAGNETIC PROPERTIES OF Fe-Co-Ni-B-Si-Nb ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 409–412 411
Figure 5: Magnetic hysteresis loops of the bulk amorphous
Fe43Co22Ni7B19Si5Nb4 rods
Slika 5: Magnetne histerezne zanke masivnih steklastih palic
Fe43Co22Ni7B19Si5Nb4
Figure 3: DSC curve of the Fe43Co22Ni7B19Si5Nb4 glassy-alloy rod
with a diameter of 2.5 mm
Slika 3: DSC-krivulja palice premera 2,5 mm iz steklaste zlitine
Fe43Co22Ni7B19Si5Nb4
Figure 4: DSC curve of the Fe43Co22Ni7B19Si5Nb4 glassy-alloy rod
with a diameter of 3.0 mm
Slika 4: DSC-krivulja palice premera 3,0 mm iz steklaste zlitine
Fe43Co22Ni7B19Si5Nb4
Table 1: Thermal (Tg – the glass transition temperature, Tx – the
crystallization temperature, Tx – the temperature interval of the
supercooled-liquid region) and magnetic (TC – the Curie temperature,
Ms – the saturation induction) properties of the bulk glassy
Fe43Co22Ni7B19Si5Nb4 rods with the diameters of (1.5, 2.5 and 3.0)
mm
Tabela 1: Termi~ne (Tg – temperatura prehoda v steklasto stanje, Tx –
temperatura kristalizacije, Tx – temperaturni interval podro~ja
superpodhlajene taline) in magnetne (TC – Curiejeva temperatura, Ms
– nasi~enje indukcije) lastnosti masivnih steklastih palic
Fe43Co22Ni7B19Si5Nb4 s premeri (1,5, 2,5 in 3,0) mm
Diameter Thermal properties Magnetic properties
!/mm Tg/K Tx/K Tx/K =Tx-Tg
TC/K Ms/T
1.5 794 828 34 652 1.07
2.5 790 827 37 650 1.22
3.0 797 826 29 655 1.18
Figure 2: DSC curve of the Fe43Co22Ni7B19Si5Nb4 glassy-alloy rod
with a diameter of 1.5 mm
Slika 2: DSC-krivulja palice premera 1,5 mm iz steklaste zlitine
Fe43Co22Ni7B19Si5Nb4
The saturation induction (Ms) of the studied glassy
rods is (1.07, 1.22 and 1.18) T for the samples with the
diameters of (1.5, 2.5 and 3) mm, respectively (Figure
5).
The obtained magnetic properties allow us to classify
the studied bulk amorphous alloy in the as-cast state as a
soft-magnetic material. These excellent magnetic proper-
ties lead us to believe that the Fe-based amorphous alloy
could be used as a new engineering and functional mate-
rial intended for the parts of inductive components.
4 CONCLUSIONS
Bulk metallic glass rods with the diameters of (1.5,
2.5 and 3) mm and a composition of
Fe43Co22Ni7B19Si5Nb4 were made by pressure
copper-mould casting. The glassy rods show good soft-
magnetic properties and thermal stability.
A high magnetization of 1.07 T to 1.22 T of the
Fe43Co22Ni7B19Si5Nb4 rods leads us to believe that the
Fe-based bulk glassy alloy with a Ni addition will be
used as a new engineering material for the parts of
micro-motors, force sensors and other applications.
Moreover, force sensors based on the newly developed
amorphous alloys may operate in a high-temperature
range. The temperature of the operation of such a sensor
is limited mainly by the Curie temperature and the value
of TC for the Fe43Co22Ni7B19Si5Nb4 alloy is in the range
from 650 K to 655 K.
5 REFERENCES
1 A. Inoue, B. L. Shen, C. T. Chan, Fe- and Co-based bulk glassy
alloys with ultrahigh strength of over 4000 MPa, Intermetallics, 14
(2006), 936–944, doi:10.1016/j.intermet.2006.01.038
2 A. Inoue, B. Shen, A. Takeuchi, Properties and applications of bulk
glassy alloys in late transition metal-based systems, Materials
Science and Engineering, A 441 (2006), 18–25, doi:10.1016/j.msea.
2006.02.416
3 W. H. Wang, Roles of minor additions in formation and properties of
bulk metallic glasses, Progress in Materials Science, 52 (2007),
540–596, doi:10.1016/j.pmatsci.2006.07.003
4 A. Inoue, Stabilization of metallic supercooled liquid and bulk amor-
phous alloys, Acta Materialia, 48 (2000), 279–306, doi:10.1016/
s1359-6454(99)00300-6
5 B. Shen, A. Inoue, C. Chang, Superhigh strength and good soft-mag-
netic properties of (Fe,Co)–B–Si–Nb bulk glassy alloys with high
glass-forming ability, Applied Physics Letters, 85 (2004) 21,
4911–4913, doi:10.1063/1.1827349
6 K. F. Yao, C. Q. Zhang, Fe-based bulk metallic glass with high pla-
sticity, Applied Physics Letters, 90 (2007), 061901, doi:10.1063/
1.2437722
7 C. Chang, B. Shen, A. Inoue, FeNi-based bulk glassy alloys with
superhigh mechanical strength and excellent soft magnetic pro-
perties, Applied Physics Letters, 89 (2006) 5, 051912, doi:10.1063/
1.2266702
8 T. Bitoh, A. Makino, A. Inoue, Origin of Low Coercivity of Fe-(Al,
Ga)-(P, C, B, Si, Ge) Bulk Glassy Alloys, Materials Transactions, 44
(2003) 10, 2020–2024
9 B. Shen, C. Chang, A. Inoue, Formation, ductile deformation beha-
vior and soft-magnetic properties of (Fe,Co,Ni)–B–Si–Nb bulk
glassy alloys, Intermetallics, 15 (2007), 9–16, doi:10.1016/j.intermet.
2005.11.037
10 S. Lesz, R. Babilas, M. Nabialek, M. Szota, M. Dospial, R. Nowo-
sielski, The characterization of structure, thermal stability and mag-
netic properties of Fe–Co–B–Si–Nb bulk amorphous and nanocry-
stalline alloys, Journal of Alloys and Compounds, 509 (2011),
197–201, doi:10.1016/j.jallcom.2010.12.146
11 R. Nowosielski, R. Babilas, G. Dercz, L. Pajak, Structure of Fe-
based metallic glass after crystallization process, Solid State Pheno-
mena, 163 (2010), 165–168, doi:10.4028/www.scientific.net/SSP.
163.165
12 L. A. Dobrzañski, A. Drygala, Influence of Laser Processing on
Polycrystalline Silicon Surface, Materials Science Forum, 706–709
(2012), 829–834, doi:10.4028/www.scientific.net/MSF.706-709.829
S. LESZ et al.: STRUCTURAL, THERMAL AND MAGNETIC PROPERTIES OF Fe-Co-Ni-B-Si-Nb ...
412 Materiali in tehnologije / Materials and technology 49 (2015) 3, 409–412
M. MIHAILOVI] et al.: THE NANO-WETTING ASPECT AT THE LIQUID-METAL/SiC INTERFACE
THE NANO-WETTING ASPECT AT THE LIQUID-METAL/SiC
INTERFACE
VIDIK NANOOMAKANJA NA STIKU STALJENA KOVINA-SiC
Marija Mihailovi}1, Karlo Rai}2, Aleksandra Patari}1, Tatjana Volkov - Husovi}2
1Institute for Technology of Nuclear and other Mineral Raw Materials, Franchet d'Esperey St. 86, 11000 Belgrade, Serbia
2Faculty of Technology and Metallurgy, University of Belgrade, Karnegijeva 4, 11000 Belgrade, Serbia
m.mihailovic@itnms.ac.rs
Prejem rokopisa – received: 2014-07-18; sprejem za objavo – accepted for publication: 2014-07-31
doi:10.17222/mit.2014.111
The wetting process on the nano-scale, as an initial and essential step in liquid metal/ceramic joining, is discussed here. Thanks
to recent breakthroughs in experimental techniques with nanometre resolution, questions posed several decades ago are being
looked at again. Despite recorded facts on acting mechanisms, the published results are very diverse due to the variety of
materials and their structures, as well as experimental conditions, so the modeling is inevitable for process development and to
overcome the multi-scale influencing parameters issues. A nano-scale wetting model have been proposed and tested on results
obtained in a liquid-metal/SiC system that was published in the literature.
Keywords: wetting, modelling, nano-scale, metal/ceramic interface
Razlo`en je postopek omakanja na nanonivoju kot za~etni in bistveni del pri spajanju staljena kovina-keramika. Zaradi sedanjih,
prelomnih eksperimentalnih tehnik do nanometrske resolucije se ponovno pojavljajo vpra{anja, stara ve~ desetletij. Kljub
dejstvom glede delujo~ih mehanizmov so objavljeni rezultati zelo razli~ni zaradi razli~nosti materialov in njihovih struktur, kot
tudi eksperimentalnih razmer. Zato je neizogibno modeliranje razvoja procesov, da se prese`e {tevilne ve~dimenzijske vplivne
parametre. Na podlagi literaturnih podatkov iz sistema staljena kovina-SiC je bil predlo`en in preizku{en model omakanja na
nanonivoju.
Klju~ne besede: omakanje, modeliranje, nanopodro~je, stik kovina-keramika
1 INTRODUCTION
Although such different materials in terms of heat
and electrical conductivity, as well as hardness, ductility,
wear or corrosion resistance, metals and ceramics have
integrated the advantages of their differences in many
modern applications when operating together. Wetting as
the initial and inevitable phenomenon of the liquid metal
to ceramic joining process have been investigated, both
experimentally and theoretically for more than three
decades.1–5 In the most recent decade, experimental tech-
niques enabled an insight into wetting phenomena at
high resolution, i.e., at the nanoscale.6 Nevertheless,
modeling is still a required method for wetting-process
prediction. Despite, or simply because of modern experi-
mental techniques, any investigation of metal/ceramic
wetting mechanisms acting on the micro and nano levels
is still demanding, both experimentally and theoreti-
cally.5,7,8 These mechanisms are important for under-
standing metal/ceramics interfacial bonding and further
process development.3 The trend in several fields of
science today is miniaturization and developing towards
the nano-scale, both for the sake of sensitive processes
and for lowering the costs; whereas for metal/ceramic
systems the intention is the miniaturization of electronic
devices.8,9 This pushed SiC to the center of investigations
again. Besides superconductivity, it has a very low
coefficient of thermal expansion and an absence of phase
transformations at operating temperatures, enabling the
observed crystal structure to be stable. Together with
new technical developments, new theoretical concepts
are also required.
There are several theories used to describe interface
bonding in the two adjacent components, i.e., liquid
metal and solid ceramic, materials that are diverse due to
the different types of atomic bonding.5 One of the first
theories was confined to reactive wetting, known as the
šreaction product control’ theory, claimed that the inter-
face reactions take control over the wetting mechanism
at the interface,1,10 while the opposite theory claimed that
chemical reactions are not crucial for controlling the
wetting phenomenon, but the capillary effects and
adsorption of metals onto the ceramic substrate, with
triple line ridging11,12. During years of research, theories
were supported by experimental data, some statements
were reconsidered,13,14 and plenty of influential para-
meters were investigated. For the sake of plenty of inve-
stigated metal/ceramic systems and different approaches,
there were variations in the experimental work, as well
as in the suggested theoretical models. For the compre-
hensive modeling, both the physical and chemical
approaches should be taken into account, as well as the
mechanisms acting at the nano scale.
Materiali in tehnologije / Materials and technology 49 (2015) 3, 413–416 413
UDK 532.64:541.183 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)413(2015)
2 A NANO-SCALE MECHANISM – THE
FRACTIONAL SURFACE APPROCAH
The macroscopically measured contact angle, ,
defined by Young’s equation (1), is usually interpreted as
the bonding quality of the metal/ceramic interface, but
this criteria is valid only up to the micron scale. The
atomic structure of the liquid metal and of the substrate
becomes important on the nano-scale, and nano-wetting
properties have an important influence on the macro-
scopic wetting behavior of liquids on solid surfaces.4,6,8
The wetting properties of the liquid metal/ceramic
interface are strongly affected by the composition of the
solid and liquid components, the roughness and other
irregularities on the macro level, or grain-boundary
grooves and lattice pits at the micro-scale, the surface
pattern of the ceramic substrate, and either the reactive
or non-reactive wetting, by the mechanisms occurring at
the nano scale, including thermal influences.4–7,15–17
Although the geometrically structured substrates
were a matter of interest for years18 at macro and micro
scale, they can be investigated in a new light, since the
recent development of experimental techniques at the
nano level, as well as of accompanied theories.4,6,19
It is shown that a change in the type of liquid metal/
substrate interface, in the same system (Ni-Si/C system
with a formed reaction layer of SiC and hence an inter-
face change), leads to a remarkable change in the wett-
ing: from contact angles much higher than 90 ° to contact
angles in the range 20–40 °.20 This refers to the macro-
scopically measured contact angle .
There is a wide range of reported measured contact
angles for the pure liquid metals in a contact with SiC,
for different temperatures, in a literature review.3 A dra-
matic change in the contact-angle values for the majority
of reactive metals is normally used in metal/ceramic
brazing is registered, even reflecting in a wetting-non-
wetting transition for the same metal with the tempera-
ture change. The situation is more complicated when the
liquid metal is in fact an alloy.
Recent in-situ experiments at the nano level brought
into connection the high traveling angle of the molten
metal over the SiC basal plane with a high interfacial
energy between the molten metal (Ti) and the SiC.6 So,
the nano-level analysis can also start from the interfacial
energies relation in the Young’s equation (1):
cos 
 

=
−SA LS
LA
(1)
where SA, LS and LA are solid-air (i.e., the corres-
ponding atmosphere), liquid-solid and liquid-air inter-
facial energies, respectively. Combining this with the
Cassie and Baxter relationship for rough surfaces,18
modified with two coefficients concerning the surface
area, one can see that the net energy in the system is
expressed according to Equation (2):
E f f f fN LA LA LA= − −2 1 2 1    cos ( cos ) (2)
where f1 is the total area of the solid-liquid interface and
f2 is the total area of the liquid-air interface (fractional
surface factors).
Using the energy Equation (2), the cosine of the
apparent contact angle, A, for the geometrically struc-
tured surface, can be expressed as:
cos cos 

A
N
LA
E
f f= −1 2 (3)
The atomistic arrangement of atoms in a liquid metal
is not entirely random. During melting the crystal struc-
tures are being broken down, and the average packing
density becomes smaller than in the solid state. The
interatomic forces keep trying to establish the original
arrangement, at the same time being disturbed by the
thermal motion of the atoms. So, the melt is much more
like the crystal than the completely random state of a
gas. Figure 1 shows the liquid metal atoms in contact
with the ceramic substrate atoms, with the apparent
contact angle " and advancing contact angle between the
liquid metal atoms and the ceramic substrate (). Several
authors21 found that in the liquid Cu/SiC system, the
liquid Cu spreads over the hexagonal crystal structure of
-SiC, maintaining a hexagonal shape. Since SiC has
over 250 crystalline forms, it is necessary to develop a
model taking into account the crystal lattice and the
planes’ orientation.
M. MIHAILOVI] et al.: THE NANO-WETTING ASPECT AT THE LIQUID-METAL/SiC INTERFACE
414 Materiali in tehnologije / Materials and technology 49 (2015) 3, 413–416
Figure 1: The liquid-metal atoms in contact with the ceramic-sub-
strate atoms
Slika 1: Atomi staljene kovine v stiku z atomi keramike v podlagi
Figure 2: Angles describing the position of liquid metal atoms (LMA)
in contact with the ceramic substrate atoms (CA)
Slika 2: Koti, ki opisujejo polo`aj atomov v staljeni kovini (LMA) v
stiku z atomi iz keramike v podlagi (CA)
At the atomistic level, the new dimensions of f1 and f2
can be introduced. Areas f1 and f2 can be derived from
the value of the angle , the contact angle between the
atoms of the metal and the ceramic, as denoted in Figure
2, and the atom packing density s in the plane of the
crystal lattice along which the contact with the liquid
metal is established.
If the reciprocal atom packing density is denoted
with s, and expressed through the atomic radius in the
plane along which the wetting process occurs in a
monocrystal, then the total plane area (P) will be:
P = s ·  r2 (4)
P1 is the ceramic atom surface area in contact with
the metal atoms:
[ ]P r1 22 1 90= ⋅ − − °π sin( ) (5)
while P2 is the area of the liquid metal at the metal/cera-
mic interface, which is not in contact with the crystal
lattice atoms:
P P r2 1
2= −π (6)
The contact point of the liquid metal atoms in contact
with the solid substrate atoms is at the distance r1, shown
in Figure 2. According to Figure 2, it can be written:
r r1 90= ⋅ − °cos( ) (7)
and hence the f1 and f2 are:
[ ]
f
P
P s1
1 2 1 90= =
⋅ − − °sin( )
(8)
f
P
P
s
s s2
2
2
290 1
1
90= =
− − °
= − − °
cos ( )
cos ( )

 (9)
Similar to our calculations, there is an expression for
the apparent contact-angle cosine on a super-hydropho-
bic surface, Equation (10).22 With the same Cassie-Bax-
ter approach, it is postulated that the measured contact
angle is a sum calculated for n surfaces:
cos ( )

 A
LA
n n, SA n, LS
nwhere
= −
=
=
=
∑
∑
1
1
1
1
f
f
n
N
n
N
(10)
where SA, LS and LA are the interfacial energies, as
described previously, and fn is the fractional coverage of
the nth chemical species.
There is a coincidence in the wetting approaches at
the nano level, although contradictory processes have
been observed: the wetting and super-non-wetting. It has
been observed that both wetting processes can be
enhanced, i.e., modified by introducing the geometrically
structured surface approach.
3 RESULTS AND DISCUSSION
Knowing the f1 and f2 factors in their new, nano-scale
meaning, the apparent contact angle on a nano-structured
surface can be calculated according to Equation (3), for
contact angles larger than 90 ° (the non-wetting case). For
the wetting case, angles smaller than 90 °, this equation is
modified by using supplementary angles to the apparent
contact angles, due to the different substrate (CA) and
the position of the liquid metal atoms (LMAs), and
hence the different geometry. So, Equation (3),
according to Figures 2 and 3, becomes (11):
cos " = f1 cos(180°– ) – f2 (11)
Most researchers deal with the -SiC(0001) and
-SiC(111) crystallographic planes, so this calculation is
based on the determined atomic packing factors for these
types of crystal lattice, although there are more than 250
different polytypes of SiC.23,24 The obtained results are
summarized in Figure 3.
The influence of the substrate planes’ orientation on
the calculated values of the apparent contact angles is
clear in Figure 3. A deviation from the straight line,
representing the characteristic advancing contact angles,
differs more between the two different SiC planes for the
wetting case, i.e., for angles smaller than 90 °, in com-
parison to the non-wetting case, where the lines are
almost overlapped. It is clear that the minimum discre-
pancy between the advancing and the apparent calculated
contact angle is around 90 °, i.e., where the transition
wetting–non-wetting occurs. Besides, the trend is
somewhat similar, but shifted, for both the -SiC(0001)
and -SiC(111) crystallographic planes. This is due to
the similarity in the configuration between those two
crystallographic planes, already reported by others.3
It must be emphasized that this consideration is
restricted to the non-reactive wetting case. During the
formation of the new reaction product at interface, the
distortion of the substrate lattice at the interface is inevi-
table, so other specific lattice factor calculations would
have to be performed.
M. MIHAILOVI] et al.: THE NANO-WETTING ASPECT AT THE LIQUID-METAL/SiC INTERFACE
Materiali in tehnologije / Materials and technology 49 (2015) 3, 413–416 415
Figure 3: Comparison between calculated values of apparent contact
angles " and the characteristic advancing contact angles  for
different orientations of the substrate planes
Slika 3: Primerjave med izra~unanimi vrednostmi navideznih kotov
stika " in zna~ilnih napredujo~ih kotov kontakta  pri razli~nih
orientacijah ravnin podlage
4 CONCLUSIONS
The wetting effect on the nano-scale depends on the
structure of the crystal lattice and the planes’ orientation
according to the proposed model tested with measured
apparent contact angles.
This is an aspect of possible liquid-metal/ceramic
interface phenomena and this approach should contribute
to a better understanding of the complex interface wett-
ing behavior and should help in predicting the wetting
modification on the nano-scale. But there is still the need
for further investigations and modeling of the wetting at
the liquid metal/ceramic interface, since quite different
mechanisms take place at the nano level, compared to the
well-established theories based on macroscopic contact-
angle measurements.
Acknowledgement
The authors wish to acknowledge the financial
support from the Ministry of Education, Science and
Technological Development of the Republic of Serbia
through the projects TR 34002 and P-172005.
5 REFERENCES
1 B. Drevet, K. Landry, P. Vikner, N. Eustathopoulos, Scripta Mate-
rialia, 35 (1996) 11, 1265–1270, doi:10.1016/1359-6462(96)00305-3
2 P. Wynblatt, Acta Mater., 48 (2000), 4439–4447, doi:10.1016/S1359-
6454(00)00230-5
3 G. W. Liu, M. L. Muolo, F. Valenza, A. Passerone, Ceramics Inter-
national, 36 (2010), 1177–1188, doi:10.1007/s10853-009-3858-0
4 T. Hofmann, M. Tasinkevych, A. Checco, E. Dobisz, S. Dietrich, B.
M. Ocko, Physical Review Letters, 104 (2010) 10, 106102,
doi:10.1103/PhysRevLett.104.106102
5 M. W. Finnis, J. Phys.: Condens. Matter., 8 (1996) 32, 5811–5836,
doi:10.1088/0953-8984/8/32/003
6 S. I. Tanaka, C. Iwamoto, Materials Science and Engineering A, 495
(2008), 168–173, doi:10.1016/j.msea.2007.11.096
7 K. T. Raic, Ceramics International, 26 (2000) 1, 19–24, doi:10.1016/
S0272-8842(99)00013-9
8 S. Dietrich, M. N. Popescu, M. Rauscher, J. Phys.: Condens. Matter.,
17 (2005) 9, S577–S593, doi:10.1088/0953-8984/17/9/017
9 S. Stopic, R. Rudolf, J. Bogovic, P. Majeri~, M. ^oli}, S. Tomi}, M.
Jenko, B. Friedrich, Mater. Tehnol., 47 (2013) 5, 577–583
10 C. Rado, B. Drevet, N. Eustathopoulos, Acta Mater., 48 (2000),
4483–4491, doi:10.1016/S1359-6454(00)00235-4
11 E. Saiz, R. M. Cannon, A. P. Tomsia, Acta Mater., 48 (2000),
4449–4462, doi:10.1016/S1359-6454(00)00231-7
12 E. Saiz, A. P. Tomsia, R. M. Cannon, Scripta Mater., 44 (2001),
159–164, doi:10.1016/S1359-6462(00)00549-2
13 N. Eustathopoulos, Curr. Opin. Solid State Mater. Sci., 9 (2005),
152–160, doi:10.1016/j.cossms.2006.04.004
14 E. Saiz, A. P. Tomsia, Curr. Opin. Solid State Mater. Sci., 9 (2005),
167–173, doi:10.1016/j.cossms.2006.04.005
15 M. Mihailovi}, T. Volkov-Husovi}, K. Rai}, Adv. Sci. Tec., 45
(2006), 1526–1531, doi:10.4028/www.scientific.net/AST.45
16 M. Tasinkevyich, S. Dietrich, Eur. Phys. J. E, 23 (2007), 117–128,
doi:10.1140/epje/i2007-10184-5
17 K. T. Rai}, R. Rudolf, A. Todorovi}, D. Stamenkovi}, I. An`el,
Mater. Tehnol., 44 (2010) 2, 59–66
18 A. B. D. Cassie, S. Baxter, Trans. Faraday. Soc., 40 (1944), 546–551,
doi:10.1039/TF9444000546
19 K. T. Raic, Adv. Sci. Tec., 32 (2003), 725–733
20 V. Bougiouri, R. Voytovych, O. Dezellus, N. Eustathopoulos, J.
Mater. Sci., 42 (2007), 2016–2023, doi:10.1007/s10853-006-1483-8
21 K. Nogi, Y. Hirata, T. Matsumoto. H. Fuji, Journal of Physics:
Conference Series, 165 (2009) 1, 012073, doi:10.1088/1742-6596/
165/1/012073
22 A. H. F. Wu, K. Nakanishi, K. L. Cho, R. Lamb, Biointerphases, 8
(2013) 5, 1–10, doi:10.1186/1559-4106-8-5
23 M. E. Levinshtein, S. L. Rumyantsev, M. S. Shur, Properties of Ad-
vanced Semiconductor Materials, John Wiley & Sons, 2001
24 A. Gasse, G. Chaumat, C. Rado, N. Eustathopoulos, J. Mat. Sci. Let.,
15 (1996) 18, 1630–1632, doi:10.1007/BF00278110
M. MIHAILOVI] et al.: THE NANO-WETTING ASPECT AT THE LIQUID-METAL/SiC INTERFACE
416 Materiali in tehnologije / Materials and technology 49 (2015) 3, 413–416
H. [IMONOVÁ et al.: THE EFFECT OF A SUPERPLASTICIZER ADMIXTURE ON THE MECHANICAL ...
THE EFFECT OF A SUPERPLASTICIZER ADMIXTURE ON THE
MECHANICAL FRACTURE PARAMETERS OF CONCRETE
VPLIV DODATKA SUPERPLASTIFIKATORJA NA PARAMETRE
MEHANSKEGA ZLOMA BETONA
Hana [imonová, Ivana Havlíková, Petr Danìk, Zbynìk Ker{ner, Tomá{ Vymazal
Brno University of Technology, Faculty of Civil Engineering, Veveøí 331/95, 602 00 Brno, Czech Republic
simonova.h@fce.vutbr.cz
Prejem rokopisa – received: 2014-07-23; sprejem za objavo – accepted for publication: 2014-09-03
doi:10.17222/mit.2014.114
This paper focuses on the mechanical fracture parameters obtained from records of three-point bending tests on concrete
specimens with a central edge notch. The tests were conducted on four sets of specimens made of different materials. The
concrete of the specimens was different in the dosage of Portland cement CEM I 42.5 R (305 kg/m3 or 355 kg/m3) and
superplasticizer (none or 0.25 % of the mass of the cement). The consistency class of fresh concrete determined using a slump
test and a flow-table test was the same for all the mixtures. Three specimens in each set were tested after aging for 28 d.
Increasing the dosage of cement and superplasticizer admixture influences the mechanical fracture properties of the concrete in
both positive and negative ways. It follows that it is proper to monitor not only the effect of the superplasticizer admixture on
the compressive strength values, but also focused attention on the fracture parameter values. The resistance to stable and
unstable crack propagation, which is evidently connected with the durability of material, was quantified using a double-K
fracture model.
Keywords: concrete, mechanical fracture parameters, superplasticizer, fracture test, double-K model
V tem ~lanku se avtorji osredinjajo na parametre mehanskega zloma, ugotovljene pri trito~kovnem upogibnem preizkusu
vzorcev iz betona s sredinsko zarezo na robu. Preizkusi so bili izvr{eni na {tirih serijah vzorcev z razliko v materialu. Beton
vzorcev se je razlikoval v dodanem Portland cementu CEM I 42,5 R (305 kg/m3 ali 355 kg/m3) in superplastifikatorju (z masnim
dele`em 0,25 % cementa ali brez njega). Razred konsistence sve`ega betona, dolo~en s preizkusom posedanja in preizkusom
te~enja na mizi, je bil enak za vse me{anice. Iz vsake serije so bili preizku{eni trije vzorci, stari 28 d. Nara{~anje dodatka
cementa in superplastifikatorja je vplivalo na lastnosti pri mehanskem lomljenju betona v pozitivno in negativno smer. Iz tega
izhaja, da je bilo treba nadzirati ne samo vpliv dodatka superplastifikatorja na tla~no trdnost, temve~ tudi nameniti pozornost
vrednostim parametrov pri prelomu. Odpornost proti stabilnem in nestabilnem napredovanju razpoke, ki je nedvoumno
povezano z zdr`ljivostjo materiala, je bila ocenjena koli~insko z uporabo dvojnega K-modela preloma.
Klju~ne besede: beton, parametri mehanskega preloma, superplastifikator, preizkus preloma, dvojni K-model
1 INTRODUCTION
The project of the Grant Agency of the Czech Re-
public "Assessment and Prediction of the Concrete Co-
ver Layers Durability" deals with the study of problems
of concrete cover layers’ durability and contributes to the
development of knowledge in the field of durability
assessment and evaluation. The project is aimed espe-
cially at the determination of the transport characteristics
of the concrete cover layers using the water- and gas-per-
meability methods. These so-called "durability parame-
ters" are completed, especially with the fracture para-
meters (e.g., fracture toughness and fracture energy) and
the basic physical and mechanical properties of fresh and
hardened concrete1.
The quality of the surface layer is significantly
related to the permeability of the material, which defines
its transport properties, and the thermal and electrical
conductivity. The permeability of the concrete with the
dense aggregate depends mainly on the porosity of the
cement stone structure and is affected, besides other
things, by cracks with a width greater than 10–4 m,
resulting from the hardening of the concrete. Cracks (or
microcracks) may combine together during the external
loading of the concrete structure and create cracks,
causing a significant reduction of durability or even
serious failure of the structure. The connection of dura-
bility and content of the microcracks in the concrete is
therefore obvious. The content of microcracks, and their
resistance to stable and unstable propagation can be
quantified by a number of fracture parameters.
In this paper the authors are focused on the mecha-
nical fracture parameters obtained from records of
three-point bending tests on concrete specimens with a
central edge notch. The tests were conducted on four sets
of specimens made of different materials. The concrete
of the specimens was different in terms of the dosage of
Portland cement CEM I 42.5 R and superplasticizer.
The resistance to stable and unstable crack propa-
gation was quantified using the double-K fracture
model2. Using this double-K model it is possible to
determine the critical crack-tip opening and the fracture
toughness of the investigated concrete, and quantify – as
indicated model name – two different levels of crack pro-
Materiali in tehnologije / Materials and technology 49 (2015) 3, 417–421 417
UDK 691:620.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)417(2015)
pagation: initiation, which corresponds to the beginning
of the stable crack growth and the level of unstable crack
propagation.
The results obtained using the double-K model are
completed by values of the compressive strength, the
modulus of elasticity, the effective fracture toughness
and the specific fracture energy.
2 EXPERIMENTAL PART
2.1 Material
The tests were conducted on four sets of specimens
differing in terms of the material. The Portland cement
CEM I 42.5 R was used as the binder. The aggregate
consisted of three grain size fractions (0–4, 4–8 and
8–16) mm. The superplasticizer Sika ViscoCrete 4035
was used in two mixtures in the amount of 0.25 % of the
weight of the cement. Table 1 introduces the details
regarding the theoretical concrete mixtures’ composition
and their designation. Table 2 introduces the real con-
crete mixtures’ composition. Table 3 introduces the
properties of the fresh concrete, which were determined
according Czech Standards3–6 in the laboratory at the
concrete-mixing plant, when all four mixtures were
prepared. The consistency class of the fresh concrete
determined using the slump test and the flow-table test
was the same for all the mixtures (Table 3).
2.2 Testing procedure
Beam specimens (of nominal dimensions 100 mm ×
100 mm × 400 mm) with a central edge notch were sub-
jected to three-point bending fracture tests (Figure 1).
The notches were made before testing with a diamond
blade saw. The notch depth was approximately equal to
1/3 of the depth of the specimen. The span length was
equal to 300 mm. Three specimens from each set were
tested after aging for 28 d. The fracture tests were
carried out using a Heckert FPZ 100/1 testing machine
(Figure 2) with the load range of 0–10 kN.
Load versus deflection diagrams (P – d-diagrams)
and load versus crack mouth opening diagrams (P –
CMOD-diagrams) were recorded using induction sensors
and an extensometer (crack-opening displacement trans-
ducer) connected in a HBM SPIDER 8 device during the
fracture experiments.
H. [IMONOVÁ et al.: THE EFFECT OF A SUPERPLASTICIZER ADMIXTURE ON THE MECHANICAL ...
418 Materiali in tehnologije / Materials and technology 49 (2015) 3, 417–421
Figure 2: Fracture-test configuration in testing machine
Slika 2: Priprava preizkusa preloma na preizkusni napravi
Table 1: Theoretical composition of the mixtures
Tabela 1: Teoreti~na sestava me{anic
Component
Mixture
0/1 1/1 0/2 1/2
CEM I 42.5 R (kg) 305 305 355 355
Sand 0–4 (kg) 929 951 886 923
Aggregate 4–8 (kg) 182 186 182 186
Aggregate 8–16 (kg) 690 706 690 706
Water (kg) 200 185 209 180
Superplasticizer (kg) 0 0.76 0 0.89
Table 2: Real composition of the mixtures
Tabela 2: Realna sestava me{anic
Component
Mixture
0/1 1/1 0/2 1/2
CEM I 42.5 R (kg) 309 303 358 359
Sand 0–4 (kg) 927 952 892 921
Aggregate 4–8 (kg) 182 190 175 190
Aggregate 8–16 (kg) 698 707 695 712
Water (kg) 202 149 192 170
Superplasticizer (kg) 0 0.73 0 0.93
Table 3: Properties of fresh concrete
Tabela 3: Lastnosti sve`ega betona
Property
Mixture
0/1 1/1 0/2 1/2
Density (kg/m3) 2315 2275 2315 2300
Slump test (mm) 60 60 60 50
class S2 S2 S2 S2
Flow test (mm) 410 360 385 350
class F2 F2 F2 F2
Air content (%) 2.7 3.6 2.5 2.8
Figure 1: Three-point bending fracture test geometry
Slika 1: Geometrija vzorca za trito~kovni upogibni preizkus
Cubes with edge lengths of 150 mm were used for
the determination of the compressive strength values.
The compressive tests were carried out using a
FORM+TEST ALPHA 3-3000 testing machine with the
load range 0–3000 kN. The compressive strength tests
were performed according to the Czech standard ^SN
EN 12 390-37; the loading rate was 0.6 MPa/s.
2.3 Methods
The modulus of elasticity values were obtained from
the first (almost) linear part of the corrected P – d-dia-
grams. The effective fracture toughness was determined
using the Effective Crack Model8, which combines linear
elastic fracture mechanics and the crack-length approach.
Estimations of the fracture energy values according to
the RILEM method were calculated using a "work of
fracture" value9. Note that, especially in the case of the
plain concrete specimens, a stability loss during the
displacement-controlled loading can occur due to the low
rigidity of the loading set-up. This stability loss appears
as a jump in the measured parameters. A procedure was
developed to recognise this problem and correct it in the
case of fracture tests conducted on concrete specimens10.
The measured P – CMOD-diagrams were used to de-
termine the fracture parameters of the double-K model.
Two levels of crack propagation were quantified: the
initiation of stable crack growth and the level of unstable
crack propagation.
The unstable fracture toughness KIcun was numeri-
cally determined first, followed by the cohesive fracture
toughness KIcc. When both of these values are known, the
following formula can be used to calculate the initiation
fracture toughness KIcini:
K K KIc
ini
Ic
un
Ic
c= − (1)
Details regarding the calculation of both the unstable
and cohesive fracture toughness can be found, e.g., in the
following articles11,12.
Finally, in accordance with Equation (2) the value of
the load Pini is determined. This value can be defined as
the load level at the beginning of stable-crack propa-
gation from an initial crack/notch and can be obtained
using the expression1:
P
WK
SF aini
Ic
ini
=
4
1 0 0( )
(2)
where W = 1/6 BD2 is the section modulus, B is the spe-
cimen width, D is the specimen depth, S is the load
span, a0 is the initial notch length according to Figure 1
and F1(0) is the geometry function given by the
following equation:
F1 0
0 0 0
0
199 1 215 3 93 2 7
1 2 1
( )
. ( )( . . . )
( )(

   




=
− − − +
+ −0
3 2) /
(3)
where 0 is the ratio a0/D.
3 RESULTS
The arithmetic mean and the standard deviation
values of the selected parameters are introduced in
following figures: compressive strength fc (Figure 3),
elasticity modulus Ec (Figure 4), specific fracture energy
GF (Figure 5), effective fracture toughness KIce (Figure
6), initiation level of the stress intensity factor KIcini
(Figure 7), ratio KIcini/KIcun (Figure 8), i.e., the ratio of
the initiation fracture toughness to unstable fracture
toughness, and the ratio Pini/Pmax (Figure 9), i.e., the ratio
between the force at the beginning of stable-crack pro-
pagation from an initial stress concentrator and the maxi-
mum force corresponding to the peak of the P –
CMOD-diagram.
The relative mean values of these properties are intro-
duced in Table 4 – the 100 % value for each material
parameter represents the values of this parameter for the
concrete with same amount of cement without superpla-
sticizer.
H. [IMONOVÁ et al.: THE EFFECT OF A SUPERPLASTICIZER ADMIXTURE ON THE MECHANICAL ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 417–421 419
Figure 4: Comparison of the modulus elasticity values Ec of the
concrete sets
Slika 4: Primerjava vrednosti modula elasti~nosti Ec za skupine
betonov
Figure 3: Comparison of the compressive strength values fc of the
concrete sets
Slika 3: Primerjava vrednosti tla~ne trdnosti fc za skupine betonov
The concrete with the superplasticizer admixture
achieved higher compressive strength values for both
dosages of cement – concrete 1/1 had about 30 % higher
value of compressive strength than concrete 0/1 without
superplasticizer, and the concrete 1/2 was about 15 %
higher than concrete 0/2.
The values of the modulus of elasticity also increased
with a superplasticizer admixture in the case of a lower
dosage of cement – concrete 1/1 achieved an about 20 %
higher value of this parameter than the concrete 0/1. On
the other hand, the modulus of elasticity values de-
creased in the case of concrete 1/2 in comparison with
the concrete 0/2 by about 4 % – but the variability is
higher in the case of concrete 1/2, so we could say that
the superplasticizer admixture had no effect on the elasti-
H. [IMONOVÁ et al.: THE EFFECT OF A SUPERPLASTICIZER ADMIXTURE ON THE MECHANICAL ...
420 Materiali in tehnologije / Materials and technology 49 (2015) 3, 417–421
Figure 9: Comparison of the ratio Pini/Pmax of the concrete sets
Slika 9: Primerjava razmerja Pini/Pmax za skupine betonov
Table 4: The relative mean values of selected parameters of concrete
sets in %
Tabela 4: Relativne glavne vrednosti izbranih parametrov vrst betona
v %
Parameter
Concrete
0/1 1/1 0/2 1/2
fc/MPa 100 129.1 100 114.0
Ec/GPa 100 120.1 100 96.2
GF/J/m2 100 113.5 100 103.9
KIce/(MPa m1/2) 100 109.1 100 104.0
KIcini/(MPa m1/2) 100 117.9 100 75.8
KIcini/KIcun 100 101.3 100 79.0
Pini/Pmax 100 98.2 100 84.9
Figure 7: Comparison of the initiation level of stress intensity factor
KIcini of the concrete sets
Slika 7: Primerjava iniciacijskega nivoja napetostnega intenzitetnega
faktorja KIcini za skupine betonov
Figure 6: Comparison of the effective fracture toughness KIce of the
concrete sets
Slika 6: Primerjava efektivne lomne `ilavosti KIce za skupine betonov
Figure 5: Comparison of the specific fracture energy GF of the
concrete sets
Slika 5: Primerjava specifi~ne prelomne energije GF za skupine be-
tonov
Figure 8: Comparison of the ratio KIcini/KIcun of the concrete sets
Slika 8: Primerjava razmerij KIcini/KIcun za skupine betonov
city modulus value in the case of a higher dosage of
cement.
The fracture energy quantifies the brittleness/tough-
ness of the material through the evaluation of the whole
P – d-diagram. This parameter value increased with the
superplasticizer admixture in both cases of cement
dosage. In addition, these values decreased with an
increase of the cement dosage in both cases of concrete,
with or without the superplasticizer. The variability of
the results is relatively high.
The effective fracture toughness takes into conside-
ration the brittleness/toughness of the materials through
the encompassing nonlinearity of the P – d-diagram
before reaching the peak load. This parameter value has
probably slightly increased with the superplasticizer
admixture in both cases of the cement dosage, but the
variability of the results is quite high – concrete 1/1 had
an about 10 % higher value than the concrete 0/1 without
the superplasticizer, and the concrete 1/2 had an about 5
% higher value than concrete 0/2.
The initiation fracture-toughness value increased
with the superplasticizer admixture in the case of the
lower dosage of cement – concrete 1/1 achieved an about
20 % higher value for this parameter than concrete 0/1.
On the other hand, this parameter decreased in the case
of concrete 1/2 in comparison with concrete 0/2 by about
25 %.
The relative value of the stress-intensity factor for the
the initiation of stable crack growth from the initial notch
quantifies the initiation brittleness/toughness of the
material corresponding to the loss of linearity of the P –
CMOD-diagram before reaching the peak load. This
parameter value did not change with the superplasticizer
admixture presence for a lower dosage of cement. On the
other hand, this parameter decreased in the case of
concrete 1/2 in comparison with concrete 0/2 by about
20 %. Similar results were achieved for the ratio between
the force at the beginning of the stable-crack propagation
from an initial stress concentrator and the maximum
force corresponding to the peak of the P – CMOD-dia-
gram.
4 CONCLUSIONS
The authors focused their attention on the mechanical
fracture parameters determined via the evaluation of the
records of the experiments performed on four sets of
concrete specimens with the stress concentrator. The
concrete used in each set differed in terms of the dosage
of Portland cement and superplasticizer. Increasing the
dosage of cement and superplasticizer admixture influ-
ences the mechanical fracture properties of the concrete
in both positive and negative ways. It follows that it is
proper to monitor not only the effect of superplasticizer
admixture on the compressive strength values13,14, but
also focus attention on the fracture parameter values.
Particularly in the case of a higher dosage of cement the
superplasticizer admixture presence had a negative effect
on the values of the fracture parameters that quantify the
resistance to stable and unstable crack propagation.
Acknowledgement
This research was carried out with the financial
support of the Czech Science Foundation, project GA
CR 13-18870S and the European Union’s "Operational
Programme Research and Development for Innovations",
No. CZ.1.05/2.1.00/03.0097 as an activity of the regional
Centre AdMaS "Advanced Materials, Structures and
Technologies".
5 REFERENCES
1 A. M. Nevile, Properties of Concrete, Pearson Education Limited,
Harlow 2011, 846
2 S. Kumar, S. V. Barai, Concrete Fracture Models and Applications,
Springer, Berlin 2011, 406
3 ^SN EN 12350-2 Testing fresh concrete – Part 2: Slump-test, Czech
Standards Institute, 2009. (This standard is the Czech version of the
European Standard EN 12350 -2: 2009)
4 ^SN EN 12350-5 Testing fresh concrete – Part 5: Flow table test,
Czech Standards Institute, 2009, (This standard is the Czech version
of the European Standard EN 12350 -5: 2009)
5 ^SN EN 12350-6 Testing fresh concrete – Part 6: Density, Czech
Standards Institute, 2009 (This standard is the Czech version of the
European Standard EN 12350 -6: 2009)
6 ^SN EN 12350-6 Testing fresh concrete – Part 7: Air content –
Pressure methods, Czech Standards Institute, 2009, (This standard is
the Czech version of the European Standard EN 12350-7: 2009)
7 ^SN EN 12390-3 Testing hardened concrete – Part 3: Compressive
strength of test specimens, Czech Standards Institute, 2009, (This
standard is the Czech version of the European Standard EN 12390-3:
2009)
8 B. L. Karihaloo, Fracture Mechanics and Structural Concrete, Long-
man Scientific & Technical, New York 1995, 346
9 RILEM TC-50 FMC (Recommendation): Determination of the
fracture energy of mortar and concrete by means of three-point bend
test on notched beams, Materials & Structures, 18 (1985), 285–290
10 P. Frantík, J. Prù{a, Z. Ker{ner, J. Macur, About stability loss during
displacement-controlled loading, Proc. of Inter. Conf. on Fibre Con-
crete 2007, Prague, 2007
11 S. Xu, H. W. Reinhardt, Z. Wu, Y. Zhao, Otto-Graf-Journal, 14
(2003), 131–157
12 X. Zhang, S. Xu, Engineering Fracture Mechanics, 78 (2011),
2115–2138, doi:10.1016/j.engfracmech.2011.03.014
13 M. Mittal, S. Basu, A. Sofi, International Journal of Civil Engi-
neering, 2 (2013) 4, 61–66
14 S. M. A. El-Gamal, F. M. Al-Nowaiserb, A. O. Al-Baityb, Journal of
Advanced Research, 3 (2012) 2, 119–124, doi:10.1016/j.jare.
2011.05.008
H. [IMONOVÁ et al.: THE EFFECT OF A SUPERPLASTICIZER ADMIXTURE ON THE MECHANICAL ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 417–421 421

A. GUWER et al.: PROPERTIES AND STRUCTURE OF Cu-Ti-Zr-Ni AMORPHOUS POWDERS ...
PROPERTIES AND STRUCTURE OF Cu-Ti-Zr-Ni AMORPHOUS
POWDERS PREPARED BY MECHANICAL ALLOYING
LASTNOSTI IN STRUKTURA AMORFNIH PRAHOV Cu-Ti-Zr-Ni,
PRIPRAVLJENIH Z MEHANSKIM LEGIRANJEM
Aleksandra Guwer, Ryszard Nowosielski, Anna Lebuda
Silesian University of Technology, Faculty of Mechanical Engineering, Institute of Engineering Materials and Biomaterials, Konarskiego Street
18A, 44-100 Gliwice, Poland
aleksandra.guwer@polsl.pl
Prejem rokopisa – received: 2014-07-26; sprejem za objavo – accepted for publication: 2014-09-02
doi:10.17222/mit.2014.119
The method of fabrication, an investigation and a comparison of the structure, size and shape of grains of a quaternary
Cu-Ti-Zr-Ni alloy were investigated. Cu-based amorphous alloys have a high strength, ductility, fracture toughness, fatigue
strength and excellent corrosion resistance in solutions such as H2SO4, NaOH, NaCl and HNO3. Samples of powders were
prepared by mechanical alloying in a high-energy ball mill SPEX 8000. To obtain the amorphous structure of the Cu47Ti34Zr11Ni8
powder, various milling times were used. Finally, four samples for testing were obtained with pure Cu, Ti, Ni, Zr (99.99 %). The
structure of the Cu47Ti34Zr11Ni8 powders was examined by X-ray diffraction (XRD) after 7 h, 8 h, 9 h and 10 h of milling time.
The chemical composition, particle size and shape of the prepared powders were investigated by scanning electron microscopy
(SEM). The microhardness was measured by using a Vickers hardness-testing machine with automatic track measurement. The
fully amorphous powders were obtained after 10 h of milling. The prolonged time of milling resulted in an increased particle
size and a changed shape of the powders. The highest microhardness was obtained for the amorphous samples.
In further work the studied amorphous powders will be consolidated using spark-plasma sintering, which is an innovative
method for the production of amorphous alloys.
Keywords: mechanical alloying, Cu-based amorphous alloys, SEM, XRD, microhardness
Preiskovan je bil na~in izdelave, preiskava in primerjava strukture, velikosti in oblike zrn kvaternerne zlitine Cu-Ti-Zr-Ni.
Amorfne zlitine na osnovi Cu imajo visoko trdnost, duktilnost, lomno `ilavost, odpornost proti utrujanju in odli~no odpornost
proti koroziji v raztopinah H2SO4, NaOH, NaCl in HNO3. Vzorci prahov so bili pripravljeni z mehanskim legiranjem v
visokoenergijskem krogli~nem mlinu SPEX 8000. Za zagotovitev amorfne strukture prahu Cu47Ti34Zr11Ni8 so bili uporabljeni
razli~ni ~asi mletja. Iz ~istega Cu, Ti, Ni, Zr (99,99 %) so bili izdelani {tirje preizku{anci. Struktura prahov Cu47Ti34Zr11Ni8 je bila
pregledana z rentgensko difrakcijo (XRD) po 7 h, 8 h, 9 h in 10 h mletja. Kemijska sestava, velikost in oblika delcev priprav-
ljenih prahov je bila preiskana z vrsti~nim elektronskim mikroskopom (SEM). Mikrotrdota je bila izmerjena z avtomatsko
napravo za merjenje trdote po Vickersu. Popolnoma amorfni prahovi so bili dobljeni po 10 h mletja. Pri podalj{anju ~asa mletja
je narasla velikost in spremenila se je oblika delcev prahov. Najvi{jo mikrotrdoto so imeli amorfni vzorci.
V nadaljevanju dela bodo preiskovani amorfni prahovi, sintrani z uporabo iskrilnega plazemskega sintranja, ki je inovativna
metoda za izdelavo amorfnih zlitin.
Klju~ne besede: mehansko legiranje, amorfne zlitine na osnovi Cu, SEM, XRD, mikrotrdota
1 INTRODUCTION
Bulk amorphous metallic alloys exhibit many supe-
rior properties compared to crystalline alloys. Lately, it
has been noted that rods and ribbons of Cu-based alloys
demonstrate a high tensile strength, fatigue strength,
fracture strength, ductility, relatively low cost of pro-
ducts, a good glass-forming ability and excellent corro-
sion resistance in solutions such as H2SO4, NaOH, NaCl
and HNO3 1–5.
The most frequently encountered methods for the
preparation of amorphous materials are casting methods.
An alternative process to prepare amorphous alloys is
mechanical alloying combined with the method of
spark-plasma sintering. Using this production method
Cu-based amorphous alloys were produced by, e.g., Kim
et al.6 and Chu et al.7
Mechanical alloying (MA) is defined as a high-
energy milling process during which the particles are
subjected to multiple cold welding, cracking and
re-welding. With rapid cold deformation the specimen’s
temperature is increased because of the transformation of
the mechanical work into heat. The MA process allows
the alloying of elements that are difficult or impossible
to combine by conventional casting methods. The pro-
ducts of MA are advanced materials, including equili-
brium, non-equilibrium (amorphous, quasicrystals, nano-
crystalline) and composite materials. The final material
properties depend on the MA process parameters (kind
of mill, size and amount of grinding media, temperature
and atmosphere of milling, ratio of grinding media mass
to powder mass, etc.)8,9.
In this paper we report on the fabrication and an
investigation of Cu47Ti34Zr11Ni8 alloy powder prepared
by mechanical alloying. The purpose of the present work
was to obtain amorphous powders that could be sintered
in the future.
Materiali in tehnologije / Materials and technology 49 (2015) 3, 423–427 423
UDK 544.022.51:669.017.13:621.762 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)423(2015)
2 EXPERIMENTAL
2.1 Materials
Four samples with the composition Cu47Ti34Zr11Ni8
were prepared using elemental powders of copper, tita-
nium, zirconium and nickel (99.99 % purity, < 325
mesh). Each sample was prepared with 8 g of properly
weighed powders. The masses and melting points10 of
the individual elements (Cu, Ti, Zr, Ni) are shown in
Table 1. The powder composition was weighed on an
analytical high-precision balance AS/X.
Table 1: Characteristics of used elements (Cu, Ti, Zr, Ni)
Tabela 1: Zna~ilnosti uporabljenih elementov (Cu, Ti, Zr, Ni)
Powder x/% m(8 g)/g Tm/°C
Copper 47 3.9252 1085 10
Titanium 34 2.1389 1670 10
Zirconium 11 1.3187 1854 10
Nickel 8 0.6170 1453 10
x/% – amount fraction
x/% – mno`inski dele`
2.2 Research methodology
Four different milling times were applied: (7, 8, 9,
10) h. The process of mechanical alloying was inter-
rupted every 30 min for 30 min to lower the temperature
of the crucible and the powders. Cr steel balls of 13 mm
diameter were used and the ball-to-powder weight ratio
was 5 : 1. The powder mixture and the Cr steel balls
were placed in an austenitic crucible in an argon atmo-
sphere inside a glove bag, as shown in Figure 1.
A high-energy ball mill SPEX 8000 CertiPrep Mixer/
Mill "shaker" type was used, which generated vibrations
of the balls and the powder inside the container11,12.
An X-ray diffractometer X’Pert Pro Panalytical and
radiation ( Co-K) of 0.178897 nm were used to study
the structure of the obtained powders. The data of the
diffraction lines were recorded using the "step-scanning"
method in the 2 range from 30 ° to 70 ° and with a
0.013 ° step. The time of the step was 40 s and the scann-
ing speed was 0.084 ° s–1.
The particles size and shape of the Cu47Ti34Zr11Ni8
powders were assessed using the microscope SEM
SUPRA 25 ZEISS with a magnification up to 500-times
The chemical compositions of the samples were
measured with energy-dispersive X-ray spectroscopy
(EDS) with an EDS analyzer as part of the SEM. The
values of the characteristic radiation energy allow a
qualitative analysis in the test sample, and the intensity
(peaks height) allows for a quantitative analysis.
The microhardnesses of the particles were measured
by the Vickers tester with automatic track measurement
using image analysis FUTURETECH FM-ARS 9000.
The microhardness measurements were made under a
load of 0.97 N. In each of the prepared samples, seven
particles were tested.
3 RESULTS AND DISCUSSION
3.1 XRD analysis
Figure 2 demonstrates the XRD patterns of the
Cu47Ti34Zr11Ni8 powders after different milling times (7
h, 8 h, 9 h, 10 h). After 7 h of mechanical alloying there
is no significant change in the position of the diffraction
peaks and the slightly diminished intensity of those
peaks is observed. After 8 h and 9 h of processing the
broadening and intensity reduction of the crystalline
diffraction lines were observed and a maximum broad
diffuse diffraction started to form, and after 10 h of
milling the samples were amorphous. The diffraction
pattern shows a single broad diffraction halo with the 2
range of 43–54 ° from the amorphous phase without
simple peaks (Figure 2d).
The same alloy was tested by Shengzhong et al.13 The
team of researchers used different process parameters for
a QM-1SP planetary high-energy ball miller and pure
elemental powders, i.e., 99.9 %. The process of
mechanical alloying was interrupted every hour for 30
min. They obtained an amorphous phase after 8 h, 9 h,
10 h and 12.5 h of milling time.
The amorphous structure of the Cu50Ti50 powders was
obtained after 8 h of mechanical alloying by using
identical parameters to those indicated in this article14.
3.2 Microstructure
Figure 3 shows the powders after: a) 7 h, b) 8 h, c) 9
h, d) 10 h of milling time. The initial size of the powders
A. GUWER et al.: PROPERTIES AND STRUCTURE OF Cu-Ti-Zr-Ni AMORPHOUS POWDERS ...
424 Materiali in tehnologije / Materials and technology 49 (2015) 3, 423–427
Figure 1: Schematic illustration of the cylindrical steel vessel placed
in the holder inside the SPEX 8000 mill
Slika 1: Shematski prikaz cilindri~ne jeklene posode, postavljene v
mlin SPEX 8000
was about 44 μm. As a result of the mechanical synthesis
the powders changed their size and shape. The largest
particles were found after 7 h of milling time (238 μm ×
143 μm). During this milling time, the particles were
stuck to large agglomerates, then after 8 h of milling
time the particles disintegrated, because after 8 h of mill-
ing the particles were crushed to a smaller average of 47
μm × 25 μm. By using longer milling times (9 h, 10 h),
the particles size was increased and their shape became
more homogeneous and spherical. However, their size
was below that after 7 h of milling time. The average
size of the particles after the milling time is listed in
Table 2.
Table 2: Average particle size (μm) of the MA powders
Tabela 2: Povpre~na velikost delcev (μm) MA-prahov
Time of mechani-
cal alloying (h) 7 8 9 10
Average particle
size (μm) 238 × 143 47 × 25 63 × 41 87 × 62
A. GUWER et al.: PROPERTIES AND STRUCTURE OF Cu-Ti-Zr-Ni AMORPHOUS POWDERS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 423–427 425
Figure 2: X-ray diffraction pattern of Cu47Ti34Zr11Ni8 powders after:
a) 7 h, b) 8 h, c) 9 h, d) 10 h of mechanical alloying
Slika 2: Posnetek rentgenske difrakcije prahov Cu47Ti34Zr11Ni8 po: a)
7 h, b) 8 h, c) 9 h, d) 10 h mehanskega legiranja
Figure 4: a) EDS spectrum with marked EDS X-ray lines and b) SEM
micrographs of Cu47Ti34Zr11Ni8 powders after 10 h of mechanical
alloying with 30 min interruption
Slika 4: a) EDS-spekter z ozna~enimi EDS rentgenskimi linijami in b)
SEM-posnetek prahov Cu47Ti34Zr11Ni8 po 10 h mehanskega legiranja
s prekinitvijo 30 min
Figure 3: Shape and size of Cu47Ti34Zr11Ni8 powder after: a) 7 h, b) 8
h, c) 9 h, d) 10 h of mechanical alloying, (SEM, magnifications
500-times)
Slika 3: Oblika in velikost prahu Cu47Ti34Zr11Ni8 po: a) 7 h, b) 8 h, c)
9 h, d) 10 h mehanskega legiranja, (SEM, pove~ava 500-kratna)
Figure 4 depicts the XRD spectrum and the analyzed
area of the Cu47Ti34Zr11Ni8 powder after 10 h of milling.
Energy-dispersive X-ray analysis (EDS) shows the X-ray
lines of copper, titanium, zirconium and nickel elements
in the sample. The amount of Cu, Zr, Ni and Ti depends
on the time of milling. Table 3 presents the detailed
results of the chemical analysis for every sample. The
particles contain the basic components (Ti, Cu, Zr and
Ni). The initial atomic percentage of Cu equals 47 %, for
Ti it is 34 %, for Zr it is 11 % and for Ni it is 8 %. The
results indicate that the obtained powder particles after
the alloying process have a very similar atomic compo-
sition compared to the initial weighed composition. The
chemical composition of the milled powders confirms
the existence of the metals identified from the XRD
spectra.
Table 3: Chemical composition of the powders surface
Tabela 3: Kemijska analiza povr{ine prahov
Milling Time (h) Element x/%
0
Cu 47
Ti 34
Zr 11
Ni 8
7
Cu 50.61
Ti 32.89
Zr 09.23
Ni 07.27
8
Cu 49.58
Ti 33.02
Zr 9.82
Ni 7.58
9
Cu 48.73
Ti 33.43
Zr 10.02
Ni 7.82
10
Cu 51.50
Ti 30.94
Zr 08.98
Ni 08.57
3.3 Microhardness
The microhardness was measured on pressed pow-
ders with ten indentations for each sample and are shown
in Figure 5. The deduced average microhardness after
milling times (7 h, 8 h, 9 h, 10 h) is shown in Table 4.
The highest average microhardness was obtained for the
powders after 10 h of milling time (553 HV), i.e., for the
powders with the fully amorphous structure. The average
microhardness increases with the milling time. The
difference between the lowest 334 HV, after 7 h of
milling, and the highest (518 HV), after 10 h of milling,
was 184 HV. This indicates the great heterogeneity of the
obtained particles. The average microhardness of the
amorphous powder Cu47Ti34Zr11Ni8 (553 HV) is higher
than that of the amorphous powders Cu50Ti50 (542 HV).14
4 CONCLUSIONS
The result of the tests and the examination of the
Cu47Ti34Zr11Ni8 powders lead to the following conclu-
sions:
• It is possible to obtain an amorphous structure for a
four-component alloy Cu47Ti34Zr11Ni8 by using me-
chanical synthesis in a SPEX 8000 mill.
• An amorphous structure was obtained for the 10 h
milling-time sample.
• The largest particles are obtained after 7 h milling
and the smallest after 8 h milling. The largest shape
and the best size regularity were obtained for the
amorphous powders.
• The presence of the initial elements Cu, Ti, Zr, Ni in
the milled particles was confirmed. The content of
elements in the milled powders corresponds to the
initial weighed composition.
• The average microhardness value increases with the
milling time and the highest hardness is achieved in
the amorphous sample (553 HV).
Acknowledgments
The work was partially supported by the National
Science Centre under research Project No.: 2012/07/N/
ST8/03437.
A. GUWER et al.: PROPERTIES AND STRUCTURE OF Cu-Ti-Zr-Ni AMORPHOUS POWDERS ...
426 Materiali in tehnologije / Materials and technology 49 (2015) 3, 423–427
Figure 5: Powders microhardness after different milling times
Slika 5: Mikrotrdote prahov po razli~nih ~asih mletja
Table 4: Average microhardness after different mechanical-alloying times
Tabela 4: Spreminjanje povpre~ne mikrotrdote pri razli~nem trajanju mehanskega legiranja
Samples Cu47Ti34Zr11Ni8 (7 h) Cu47Ti34Zr11Ni8 (8 h) Cu47Ti34Zr11Ni8 (9 h) Cu47Ti34Zr11Ni8 (10 h)
The average microhardness (HV) 428 496 545 553
5 REFERENCES
1 P. Lee, C. Yao, J. Chen, L. Wang. R. Jeng, Y. Lin, Preparation and
thermal stability of mechanically alloyed Cu–Zr–Ti–Y amorphous
powders, Materials Science and Engineering A, 375–377 (2004),
829–833, doi:10.1016/j.msea.2003.10.107
2 C. Suryanarayana, A. Inoue, Bulk Metallic Glasses, CRC Press,
Boca Raton, London, New York 2011, 313–322
3 H. Kim, K. Sumiyama, K. Suzuki, Formation and thermal stability of
nanocrystalline Cu-Ti-Ni prepared by mechanical alloying, Journal
of Alloys and Compounds, 239 (1996), 88–93, doi:10.1016/0925-
8388(96)02274-8
4 C. Hu, H. Wu, Formation of Cu–Zr–Ti amorphous powders with
Band Si additions by mechanical alloying technique, Journal of
Alloys and Compounds, 434–435 (2007), 390–393, doi:10.1016/
j.jallcom.2006.08.219
5 A. Inoue, B. Shen, A. Takeuchi, Fabrication, properties and appli-
cations of bulk glassy alloys in late transition metal-based systems,
Materials Science and Engineering A, 441 (2006), 18–25,
doi:10.1016/j.msea.2006.02.416
6 C. K. Kim, H. S. Lee, S. Y. Shin, J. C. Lee, D. H. Kim, S. Lee, Mi-
crostructure and mechanical properties of Cu – based bulk amor-
phous alloy billets fabricated by spark plasma sintering, Materials
Science and Engineering A, 406 (2005), 293–299, doi:10.1016/
j.msea.2005.06.043
7 Z. H. Chu, H. Kato, G. Q. Xie, G. Y. Yuan, W. J. Ding, A. Inoue,
Consolidation and mechanical properties of Cu46Zr42Al7Y5 metallic
glass by spark plasma sintering, Journal of Non-Crystalline Solids,
358 (2012), 1263–1267, doi:10.1016/j.jnoncrysol.2012.02.027
8 C. Suryanarayana, Recent developments in mechanical alloying, Re-
views on Advanced Materials Science, 18 (2008), 203–211
9 M. Adamiak, Mechanical alloying for fabrication of aluminium
matrix composite powders with Ti-Al intermetallics reinforcement,
Journal of Achievements in Materials and Manufacturing Engineer-
ing, 31 (2008), 191–196
10 The Engineering ToolBox /cited 2014-07-10/. Available from World
Wide Web: http://www.engineeringtoolbox.com/melting-tempera-
ture-metals-d_860.html
11 R. Nowosielski, R. Babilas, G. Dercz, L. Pajak, J. Wrona, Structure
and properties of barium ferrite powders prepared by milling and
annealing, Journal of Achievements in Materials and Manufacturing
Engineering, 22 (2007), 45–48
12 R. Nowosielski, R. Babilas, G. Dercz, L. Paj¹k, Microstructure of
composite material with powders of barium ferrite, Journal of
Achievements in Materials and Manufacturing Engineering, 17
(2006), 117–120
13 K. Shengzhong, F. Liu, D. Yutian, X. Guangji, Synthesis and mag-
netic properties of Cu-based amorphous alloys made by mechanical
alloying, Intermetallics, 12 (2004), 1115–1118, doi:10.1016/j.inter-
met.2004.04.007
14 A. Guwer, R. Nowosielski, A. Borowski, R. Babilas, Fabrication of
copper-titanium powders prepared by mechanical alloying, Indian
Journal of Engineering & Materials Sciences, 21 (2014), 265–271
A. GUWER et al.: PROPERTIES AND STRUCTURE OF Cu-Ti-Zr-Ni AMORPHOUS POWDERS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 423–427 427

P. BÍLEK et al.: INTERACTION OF Cr2N AND Cr2N/Ag THIN FILMS WITH CuZn-BRASS COUNTERPART ...
INTERACTION OF Cr2N AND Cr2N/Ag THIN FILMS WITH
CuZn-BRASS COUNTERPART DURING BALL-ON-DISC TESTING
INTERAKCIJA Cr2N IN Cr2N/Ag TANKIH PLASTI V PARU S
CuZn-MEDENINO MED PREIZKUSOM KROGLA NA DISK
Pavel Bílek, Peter Jur~i, Petra Dulová, Mária Hudáková, Jana Pta~inová,
Matej Pa{ák
Institute of Materials Science, Faculty of Materials Science and Technology in Trnava, Slovak University of Technology in Bratislava,
Paulínská 16, 917 24 Trnava, Slovak Republic
pavel-bilek@email.cz
Prejem rokopisa – received: 2014-07-28; sprejem za objavo – accepted for publication: 2014-09-04
doi:10.17222/mit.2014.121
Cr2N- and Cr2N/Ag-nanocomposite thin films were deposited on a substrate made of Cr-V ledeburitic tool steel Vanadis 6 by
reactive magnetron sputtering, at a deposition temperature of 500 °C, using pure Cr and Ag targets, in a composite, low-pressure
N2/Ar atmosphere. The additions of silver to the Cr2N/Ag coatings were w = (3, 7, 11 and 15) %. Tribological testing using a
ball-on-disc apparatus was realized at ambient temperature and, for the Cr2N with the additions of mass fractions of Ag 7 % and
11 %, at the elevated temperatures of (300, 400 and 500) °C, respectively. Balls made of binary CuZn-brass (55 % Cu, 45 % Zn)
were used as the counterparts. The wear tracks after the ball-on-disc testing and the worn balls were analyzed with scanning
electron microscopy (SEM) and a microanalysis (EDX), and the wear rates were calculated. The adhesive wear was derived
from a quantitative-point metallographic analysis. The obtained results infer that a considerable material transfer from the
counterpart onto the surface of the coatings takes place during dry sliding. The material transfer (and the adhesive wear of the
counterpart) is mainly due to the low shear strength of the brass used. Two main trends were observed. The first one shows that
the adhesive material transfer decreases with the increasing silver content when tested at ambient temperature. The second trend
indicates that the use of a higher testing temperature leads to a higher adhesive wear of the counterpart.
Keywords: Cr2N/Ag-nanocomposite PVD coatings, ball-on-disc, adhesion, friction coefficient, wear rate
Cr2N- in Cr2N/Ag-nanokompozitne tanke plasti so bile nanesene na podlago iz Cr-V ledeburitnega orodnega jekla Vanadis 6 z
reaktivnim nana{anjem z magnetronom pri temperaturi nana{anja 500 °C z uporabo tar~ iz ~istega Cr in Ag v kompozit v
nizkotla~ni atmosferi N2/Ar. Dodatki srebra v nanose Cr2N/Ag so bili w = (3, 7, 11 in 15) %. Tribolo{ki preizkusi z uporabo
naprave krogla na disk so bili izvr{eni pri sobni temperaturi in pri Cr2N tudi z masnim dele`em dodatka 7 % in 11 % Ag pri
povi{anih temperaturah (300, 400 in 500) °C. Krogle, izdelane iz binarne CuZn-medenine (55 % Cu, 45 % Zn), so bile izbrane
kot par. Sledovi obrabe po preizkusu krogla na disk in obrabljene krogle so bili analizirani z vrsti~no elektronsko mikroskopijo
(SEM) in mikroanalizo (EDX), izra~unane pa so bile tudi hitrosti obrabe. Adhezivna obraba je bila prikazana s kvantitativno
to~kasto metalografsko analizo. Dobljeni rezultati ka`ejo, da se med suhim drsenjem pojavi ob~uten prenos materiala iz krogle
na povr{ino nanosa. Prenos materiala (in adhezijska obraba krogle) je nastal ve~inoma zaradi majhne stri`ne trdnosti medenine.
Opa`eni sta bili dve glavni usmeritvi. Prva je bila, da se pri preizku{anju pri sobni temperaturi adhezivni prenos materiala
zmanj{uje z nara{~ajo~o vsebnostjo srebra. Druga pa, da uporaba vi{je temperature pri preizkusu povzro~i ve~jo adhezivno
obrabo krogle.
Klju~ne besede: Cr2N/Ag-nanokompozitni PVD-nanosi, krogla na disk, adhezija, koeficient trenja, hitrost obrabe
1 INTRODUCTION
Hard ceramic coatings like CrN and TiN have been
used for the last three decades due to their high hardness
and chemical stability, high oxidation resistance and low
wear rate1,2. They have gained great scientific interest
and industrial popularity due to these properties in
copper machining3, alumina die casting and forming, and
wood processing4. However, the friction coefficient of
most transition-metal nitride coatings is fairly high
(0.6–0.8) and the tribological effectiveness, especially at
elevated temperatures, is insufficient5,6. Therefore, a lot
of effort has been made in recent years to decrease the
friction coefficient at room as well as elevated tempera-
tures. Multi-functional coatings combining soft lubri-
cating phases within a hard wear-resistance matrix offer
good properties7,8. These coatings generally include one
or more nanocrystalline phases in a functional matrix to
provide improved mechanical and tribological properties
and/or corrosion resistance. Some coatings are designed
to be adaptive, that is, their properties follow the changes
in the operating conditions. An example of an adaptive
coating is a hard wear-resistance matrix with an incorpo-
ration of soft metals like Cu, Ag or Au. This method can
improve the lubricating in specific tribological condi-
tions9. Chromium nitrides combined with noble metals
are relatively easy to co-deposit by reactive magnetron
sputtering and they form nanocomposite structures due
to a lack of miscibility between the matrix and the lubri-
cant10,11.
Silver is most commonly used as an addition to the
TM-nitride thin films. It exhibits a stable chemical beha-
viour over a wide temperature range as well as in a
variety of aggressive environments. Ag is capable to mi-
Materiali in tehnologije / Materials and technology 49 (2015) 3, 429–433 429
UDK 621.793:539.92:539.375.6 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)429(2015)
grate to a free surface to form Ag particles providing
lubrication above 300 °C and, thereby, distinctively
reducing the friction coefficient12.
Our recent investigations of the magnetron-sputtered
Cr2N-films with w = (3, 7, 11 and 15) % of silver
amounts, deposited on the Cr-V ledeburitic steel Vanadis
6, can be summarized as follows13–17: the incorporation
of silver in the Cr2N matrix led to an improvement in the
tribological properties at elevated temperatures, espe-
cially at 400 °C and 500 °C. Nevertheless, the addition
of 15 % of Ag made the film too soft and sensitive to the
wear, which resulted in a worsening of the tribological
performance. On the other hand, the films with w = (7
and 11) % of Ag additions seem to be very promising.
In this paper, the tribological performance against a
CuZn-brass counterpart of the nanocomposite coatings
consisting of a hard Cr2N matrix co-deposited with diffe-
rent Ag additions is investigated. The films were depo-
sited onto the Cr-V ledeburitic steel Vanadis 6 using the
magnetron-sputter technique.
2 EXPERIMENTAL WORK
The substrate material was the PM ledeburitic tool
steel Vanadis 6 with mass fractions 2.1 % C, 1.0 % Si,
0.4 % Mn, 6.8 % Cr, 1.5 % Mo, 5.4 % V and Fe as the
balance element.
The samples used for the investigation and the con-
ditions for depositing Cr2N- and Cr2N/Ag- coatings were
reported elsewhere11. In the case of the Cr2N coatings,
during the deposition, the power was 2.9 kW per cathode
(both Cr). For the production of the Ag-containing coat-
ings, the power of the Cr cathode was kept at 5.8 kW,
while the power of the Ag cathode was varied (0.10,
0.21, 0.34 and 0.45) kW in order to prepare the films
with different Ag concentrations (3, 7, 11 and 15) %.
The tribological properties of the coatings were
measured using a CSM ball-on-disc tribometer at room
temperature and, for the coatings with 7 % and 11 % of
Ag, also at the elevated temperatures up to 500 °C. The
balls of 6 mm in diameter, made from CuZn-brass
(55 % Cu, 45 % Zn), were used for the tests. No external
lubricant was added during the measurements. The
normal load used for the investigation was 1 N and the
total sliding distance for each measurement was 100 m.
The volume losses V of the worn balls were calculated
on the basis of the sketch shown on Figure 1 using the
following formula:
V = (( · h) · (32+h2))/6 (1)
where R is the ball radius, h and  are the height and the
radius of the worn spherical segment of the ball.
Relating the volume loss to the normal load and sliding
distance, the wear rates W were calculated.
After the testing, the wear tracks and the worn balls
were examined with a scanning electron microscope
(SEM) JEOL JSM-7600F and an energy dispersive X-ray
analysis (EDX).
The adhesive wear was derived from the quantita-
tive-point metallographic analysis carried out on the
SEM micrographs of the tracks after the ball-on-disc
testing.
3 RESULTS AND DISCUSSIONS
Figure 2 shows a detail of the track after the ball-
on-disc test of the Cr2N/11Ag coating tested at room
temperature against the CuZn-brass counterpart. The
surface of the coating did not show any indications of
damage. On the other hand, thanks to the corresponding
EDX maps, a considerable material transfer from the
counterpart to the surface of the coating was detected; it
was mainly due to the low shear strength of the brass
used.
P. BÍLEK et al.: INTERACTION OF Cr2N AND Cr2N/Ag THIN FILMS WITH CuZn-BRASS COUNTERPART ...
430 Materiali in tehnologije / Materials and technology 49 (2015) 3, 429–433
Figure 2: Transferred material of the CuZn-brass counterpart to the
surface of the Cr2N/11Ag coating after the ball-on-disc test at room
temperature: a) overview, b) EDX of chromium, c) EDX of silver, d)
EDX of copper, e) EDX of zinc
Slika 2: Prenesen material CuZn-medenine iz krogle na povr{ino
Cr2N/11Ag-nanosa po preizkusu krogla na disk; preizku{eno pri sobni
temperaturi: a) videz, b) EDX-kroma, c) EDX-srebra, d) EDX-bakra,
e) EDX cinka
Figure 1: Sketch of a worn ball for the volume-loss calculation
Slika 1: Skica prikaza obrabe krogle za ra~unanje volumenske izgube
Figure 3 depicts the worn CuZn-brass ball after the
ball-on-disc test of the Cr2N/11Ag coating at room tem-
perature. The parallel grooves oriented along the sliding
direction are well visible and the diameter of the worn
spherical segment is easily measurable. The surfaces of
the worn balls were investigated with an EDX analysis to
confirm the presence of silver. However, no signal of
silver was obtained and one could assume that the silver
content on the surface was too low to be detected with
the EDX analysis.
The mean value of the friction coefficient examined
at room temperature decreased with the increasing silver
content in the Cr2N matrix until 7 % (Figure 4). How-
ever, the influence of a higher silver content was not as
positive for the friction coefficient as for the coating with
7 % Ag. The wear rates of the balls decreased by about
50 % for the coatings with the silver addition in com-
parison with the pure Cr2N, and in the case of the coating
with 15 % of Ag the decrease was about 75 % (Figure 4).
These results were expected since silver can act as a
solid lubricant, facilitating the sliding of the balls12. In
our previous works11,17 silver particles were well visible
on the surface of the coatings after the deposition. On the
other hand, the experiments carried out at the elevated
temperature showed the opposite tendency. An increase
in the testing temperature led to a slight increase in the
mean value of the friction coefficient (Figure 5). As
reported in the previous works16,17, a decrease in the
friction coefficient was observed with the increasing test-
ing temperature, but there an alumina counterpart was
used. In the case of the CuZn-brass counterpart the me-
chanism of the wear was different and a softening of the
CuZn-brass alloy also took place at high temperatures.
Nevertheless, the wear rates of the CuZn-brass balls
during the tribological testing of the coatings with mass
fractions 7 % and 11 % of Ag at elevated temperatures
were lower in comparison with the measurement at room
temperature (Figure 5). The minimum values for both
coatings were found at the temperature of 300 °C.
Figure 6 depicts the dependence of the friction
coefficient on the sliding distance for the Cr2N + 11Ag
coating tested at ambient and elevated temperatures. All
the measurements show a considerable instability. The
friction coefficient is oscillating around the mean value
in the range of 0.2. This can be explained with the cre-
P. BÍLEK et al.: INTERACTION OF Cr2N AND Cr2N/Ag THIN FILMS WITH CuZn-BRASS COUNTERPART ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 429–433 431
Figure 5: Wear rates W and friction coefficient μ of Cr2N + 7Ag* and
Cr2N + 11Ag coatings tested at room and elevated temperatures, *pre-
vious result16
Slika 5: Hitrost obrabe W in koeficient trenja μ pri Cr2N + 7Ag* in
Cr2N + 11Ag-nanosih, preizku{anih pri sobni in povi{anih tempera-
turah, *predhodni rezultati16
Figure 3: Worn CuZn-brass ball after the ball-on-disc test of the
Cr2N/11Ag coating at room temperature
Slika 3: Obrabljena krogla iz CuZn-medenine po preizkusu krogla na
disku na Cr2N/11Ag-nanosu, preizku{ano pri sobni temperaturi
Figure 6: Dependence of friction coefficient μ on sliding distance L of
the Cr2N + 11Ag coating tested at different temperatures against the
CuZn-brass counterpart
Slika 6: Odvisnost koeficienta trenja μ od dol`ine drsenja L pri Cr2N
+ 11Ag-nanosu, preizku{anem pri razli~nih temperaturah, v paru z
CuZn-medeninasto kroglo
Figure 4: Wear rates W and friction coefficient μ of Cr2N and Cr2N/
Ag coatings tested at room temperature
Slika 4: Hitrost obrabe W in koeficient trenja μ pri Cr2N- in Cr2N/
Ag-prevlekah, preizku{anih pri sobni temperaturi
ation of adhesion joins between the surface of the coat-
ing and the ball during the sliding and their subsequent
release, accompanied with the oscillating of the friction
coefficient.
A quantitative-point metallographic analysis was
used to describe the adhesion interaction during the
ball-on-disc testing. Figure 7 shows the results of the
testing at room temperature for the uncoated steel Vana-
dis 6, the pure Cr2N and Cr2N/Ag coatings. The highest
adhesion wear was found for the uncoated substrate,
where the area of adhesion interaction was A = (81 ±
5) %. For the coated substrate it is clearly evident that
the adhesion wear decreases with the increasing silver
content; the minimum A = (55 ± 8) % was found for the
Cr2N coating with the highest silver addition of 15 %.
One can assume that the incorporation of silver into the
Cr2N matrix improves the wear of the CuZn-brass ball
front view and also the area of adhesion interaction at
room temperature.
On the other hand, the material transfer is more re-
markable in the conditions of higher testing temperatures
(Figure 8). Both coatings, Cr2N with mass fractions of
Ag w = (7 and 11) %, showed the same tendency,
although the coating with 11 % of Ag exhibited a slower
increase in the adhesion wear with the increasing testing
temperature than the coating with 7 % of Ag. Most
likely, the higher adhesion interaction at higher tempera-
tures is caused by the softening of the CuZn-brass mate-
rial, when the atoms of the ball material easily create the
adhesion joins.
4 CONCLUSIONS
The friction and wear characteristics of the Cr2N and
Cr2N/Ag coatings prepared with the magnetron-sputter-
deposition method were investigated at room tempera-
ture and elevated temperatures during a ball-on-disc
testing against a CuZn-brass counterpart. The results can
be summarized as follows:
• During the tribological testing, a considerable mate-
rial transfer from the CuZn-brass counterpart to the
surfaces of all the tested coatings was observed.
• For the Cr2N/Ag coatings, a lower friction coefficient
and also lower wear rates of the balls were found
during the testing at room temperature in comparison
with the pure Cr2N. This phenomenon is attributed to
the silver incorporated into the Cr2N matrix.
• Higher testing temperatures led to a slight increase in
the friction coefficient. On the other hand, the wear
rates of the balls were further decreasing, with the
minimum values at 300 °C.
• During the testing at room temperature, the highest
material transfer was found for the uncoated steel
Vanadis 6. The adhesion wear was lower when the
Cr2N coating was tested. The decrease is more re-
markable with the increasing silver content incorpo-
rated into the Cr2N matrix.
• Higher testing temperatures led to increased adhesion
interaction of both Cr2N coatings, with mass fractions
of Ag 7 % and 11 %.
Acknowledgements
This publication is the result of implementing the
project: CE for development and application of advanced
diagnostic methods in processing of metallic and
non-metallic materials, ITMS: 26220120048, supported
by the Research & Development Operational Programme
funded by the ERDF. This research was supported by
grant project VEGA 1/1035/12.
5 REFERENCES
1 S. K. Pradhan, C. Nouveau, A. Vasin, M. A. Djouadi, Surface and
Coatings Technology, 200 (2005), 141–145, doi:10.1016/j.surfcoat.
2005.02.038
2 P. Panjan, B. Navin{ek, A. Cvelbar, A. Zalar, I. Milo{ev, Thin Solid
Films, 281–282 (1996), 298-301, doi:10.1016/0040-6090(96)
08663-4
P. BÍLEK et al.: INTERACTION OF Cr2N AND Cr2N/Ag THIN FILMS WITH CuZn-BRASS COUNTERPART ...
432 Materiali in tehnologije / Materials and technology 49 (2015) 3, 429–433
Figure 8: Area A of adhesion interactions between CuZn-brass and
Cr2N coatings with mass fractions of Ag 7 % and 11 % during the
ball-on-disc testing at different temperatures
Slika 8: Podro~je A adhezijske interakcije med CuZn-medenino in
Cr2N-nanosi z masnim dele`em Ag 7 % in 11 % med preizkusi krogla
na disk pri razli~nih temperaturah
Figure 7: Area A of adhesion interactions between CuZn-brass and
uncoated steel Vanadis 6, Cr2N and Cr2N/Ag coatings during ball-on-
disc testing, at room temperature
Slika 7: Podro~je A adhezijske interakcije med CuZn-medenino in
neprekritim jeklom Vanadis 6 ter Cr2N- in Cr2N/Ag-nanosi med
preizkusom krogla na disk pri sobni temperaturi
3 A. Kondo, T. Oogami, K. Sato, Y. Tanaka, Surface and Coatings
Technology, 177–178 (2004), 238–244, doi:10.1016/j.surfcoat.
2003.09.039
4 R. Gahlin, M. Bronmark, P. Hedenqvist, S. Hogmark, G. Hakansson,
Surface and Coatings Technology, 76–77 (1995), 174–180,
doi:10.1016/0257-8972(95)02597-9
5 T. Polcar, T. Kubart, R. Novak, L. Kopecky, P. Siroky, Surface and
Coatings Technology, 193 (2005), 192–199, doi:10.1016/j.surfcoat.
2004.07.098
6 M. Brizuela, A. Garcia-Luis, I. Braceras, J. I. Onate, J. C. San-
chez-Lopez, D. Martinez-Martinez, C. Lopez-Cartes, A. Fernandez,
Surface and Coatings Technology, 200 (2005), 192–197, doi:10.1016/
j.surfcoat.2005.02.105
7 S. M. Aouadi, B. Luster, P. Kohli, C. Muratore, A. A. Voevodin, Sur-
face and Coatings Technology, 204 (2009), 962–968, doi:10.1016/
j.surfcoat.2009.04.010
8 C. Muratore, J. J. Hu, A. A. Voevodin, Thin Solid Films, 515 (2007),
3638–3643, doi:10.1016/j.tsf.2006.09.051
9 H. C. Barshilia, M. S. Prakash, A. Jain, K. S. Rajam, Vacuum, 77
(2005), 169–179, doi:10.1016/j.vacuum.2004.08.020
10 C. P. Mulligan, T. A. Blanchet, D. Gall, Wear, 269 (2010), 125–131,
doi:10.1016/j.wear.2010.03.015
11 P. Bílek, P. Jur~i, M. Hudáková, M. Pa{ák, M. Kusý, J. Bohovi~ová,
Applied Surface Science, 307 (2014), 13–19, doi:10.1016/j.apsusc.
2014.03.044
12 C. P. Mulligan, P. A. Papi, D. Gall, Thin Solid Films, 520 (2012),
6774–6779, doi:10.1016/j.tsf.2012.06.082
13 P. Jur~i, J. Bohovi~ová, M. Hudáková, P. Bílek, Mater. Tehnol., 48
(2014) 2, 159–170
14 P. Jur~i, I. Dlouhý, Applied Surface Science, 257 (2011),
10581–10589, doi:10.1016/j.apsusc.2011.07.054
15 P. Jur~i, S. Krum, Materials Engineering, 19 (2012), 64–70
16 P. Bílek, P. Jur~i, M. Hudáková, L. ^aplovi~, M. Novák, Mater.
Tehnol., 48 (2014) 5, 669-673
17 P. Bílek, P. Jur~i, M. Hudáková, L. ^aplovi~, M. Novák, Wear,
(2015), doi:10.1016/j.wear.2015.03.019
P. BÍLEK et al.: INTERACTION OF Cr2N AND Cr2N/Ag THIN FILMS WITH CuZn-BRASS COUNTERPART ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 429–433 433

B. PONIKU et al.: USING SIMULATED SPECTRA TO TEST THE EFFICIENCY ...
USING SIMULATED SPECTRA TO TEST THE EFFICIENCY OF
SPECTRAL PROCESSING SOFTWARE IN REDUCING THE NOISE
IN AUGER ELECTRON SPECTRA
UPORABA SIMULIRANEGA SPEKTRA ZA PREIZKUS U^INKOVITOSTI
PROGRAMSKE OPREME PREDELAVE SPEKTRA PRI ZMANJ[ANJU [UMA
SPEKTRA AUGERJEVIH ELEKTRONOV
Besnik Poniku1,2, Igor Beli~1, Monika Jenko1
1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2Jo`ef Stefan International Postgraduate School, Jamova 39, 1000 Ljubljana, Slovenia
besnik.poniku@imt.si
Prejem rokopisa – received: 2015-01-04; sprejem za objavo – accepted for publication: 2015-01-21
doi:10.17222/mit.2015.013
When attempting to automate Auger spectra analyses it becomes necessary to have a deeper knowledge of the constituent
elements of the spectra. In order to obtain a reliable analysis, the unavoidable spectral noise must be reduced, thus giving a
clearer view to the spectral peaks and the spectral background. Therefore, the necessary step is to analyze the spectral noise and
to find a way to evaluate the noise-reduction algorithms. A method in which simulated Auger electron spectra are used for
testing the efficiency of noise-reduction routines has been proposed. The performance of noise-reduction procedures on
measured spectra cannot be evaluated since the intrinsic noiseless spectra is never available for reference; therefore, the spectra
were simulated and the noise-reduction routines were used on the simulated spectra. After the processing, the simulated
noiseless spectrum is subtracted from the complete spectrum, leaving the remaining noise for further analysis and a comparison
with the exactly known simulated noise. For each spectrum data point the noise ratios are calculated by dividing the remaining
noise levels by the initial noise. When plotting the noise ratios for each respective processing route, it was found that most of the
noise ratios lie in the interval –1 to +1, indicating an improvement in regard to the initial noise. Such a plot of the noise ratios
offers a convenient way for assessing the efficiency of the noise-reduction routine at a glance.
Keywords: Auger electron spectroscopy, spectra simulator, spectral noise, noise reduction
Avtomatizacija postopka analize Augerjevih spektrov zahteva dobro poznanje posameznih sestavnih elementov spektra.
Zanesljivost avtomatske analize je v prvi vrsti odvisna od tega, v kolik{ni meri nam uspe zmanj{ati spektru prime{an {um, ki
sicer zamegli tako spektralne vrhove kot tudi spektralno ozadje. Zato moramo najprej analizirati lastnosti {uma, prime{anega
spektrom, in poiskati na~ine za ovrednotenje delovanja orodij, ki {um zmanj{ujejo. V ~lanku predlagamo uporabo simulatorja
Augerjevih spektrov, ker sicer pri izmerjenih spektrih nikoli ne poznamo oblike prime{anega {uma in torej nimamo osnove za
dobro ovrednotenje delovanja uporabljenih orodij. Po uporabi orodja za zmanj{evanje {uma, ki deluje na simuliranem spektru,
odstranimo natan~no poznano spektralno ozadje in spektralne vrhove. Tako dobimo preostali {um, ki ga primerjamo z znanim
za~etnim {umom. V vsaki to~ki spektra so izra~unana razmerja med za~etnim in kon~nim {umom. Razmerja so nato prikazana v
grafu in v veliki ve~ini spektralnih to~k le`ijo v intervalu –1, 1. Tako dobimo vizualno predstavitev delovanja orodij za
zmanj{evanje {uma, ki omogo~a hitro oceno u~inkovitosti preizku{anega orodja.
Klju~ne besede: Augerjeva elektronska spektroskopija, simulator spektra, spektralni {um, zmanj{evanje {uma
1 INTRODUCTION
Auger electron spectroscopy is a technique often
used for the elemental characterization of the surface of
conductive samples.1–8 Apart from a high surface sensiti-
vity,9 due to the fact that the primary electron beam can
be focused down to approximately 10 nm in diameter,10
analyses with very good spatial resolution can also be
performed. This fact makes it possible to analyze
features on a nanometer scale on the surface through this
technique.
To interpret the measured spectra the measured data
have to be manipulated by software for signal pro-
cessing. This manipulation inevitably leaves its mark on
the results obtained.11
Smoothing is one of the methods that are used for the
purpose of reducing the noise in Auger electron spec-
tra.12 Very little can be said about the efficiency of such
procedures in reducing the noise when applying them in
measured spectra, because noise in both the input and the
output spectra is at unknown levels. The aim of this work
is to show a simple way in which the performance of the
noise-reduction techniques can be assessed using
simulated spectra. Using simulated spectra to assess the
efficiency of noise-reduction routines is very appro-
priate. This comes about due to the fact that the values
for the different components of the simulated AE spectra
(including the noise) are known before processing, and
thus any change due to the processing route may be
found and then compared to the initial preprocessed
values.
Materiali in tehnologije / Materials and technology 49 (2015) 3, 435–439 435
UDK 543.428.2:544.171.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(3)435(2015)
2 EXPERIMENT
The construction of the simulator for gathering the
simulated spectra used for this assessment is described in
detail in11. For the construction of this simulator a num-
ber of measured AE spectra obtained from spring-steel
samples were closely inspected. The neural network was
used to model the primary background by selecting a
number of representative points for the background and
including them in the training data set for the neural
network. After carefully observing the behavior of the
background in the measured spectra an equation was
derived, which then would be used for generating various
primary backgrounds that would resemble those
observed in the measured spectra. After removing the
primary background defined in this way the peak base
and the peaks remained. The peak base was also
modeled in the same way using the neural network, and
the removal of the defined peak base left only the charac-
teristic peaks. The peak base and the peaks of various
elements were saved in the database. Combining the
generated primary backgrounds, on the one hand, and the
peak base and characteristic peaks from the database, on
the other, produced the simulated spectrum. The gene-
rated noise that was then added to such a spectrum was
also made to resemble the noise observed in measured
spectra. It is important to note that while the components
of the simulated spectra such as the background and
noise are made to resemble those of the measured
spectra, their exact values are simulated and therefore
known and stored in the computer (Figure 1).
Through the modeling of the background, which was
performed using the neural network, we have found that
the AE spectra consist of three main components: the
primary background, the peak base, and the peaks (Fig-
ure 2).
From the set of standard AE spectra that were ob-
tained using COMPRO10, a freely available online spec-
tral database, the peak base and the peaks of elements
such as Al, C, Co, Cu, Fe, Au, Ni, O, Si, Ag, Ti, and V
were extracted (as shown in Figure 3 for the case of
iron) and were stored separately.
The AE spectra simulator combines the extracted
peak base and the peaks from various standard elements,
and it combines them with the randomly defined primary
background (Figure 4) to form the complete simulated
spectrum without the noise.
At the end of the simulation process the random
noise is added and also stored separately for further use.
We have ensured that the properties of the simulated
noise resemble the properties of the noise in the
measured AE spectra.
Other AE spectra simulators can be used for this pur-
pose as well. One such simulator is SESSA (Simulation
of Electron Spectra for Surface Analysis). SESSA is
B. PONIKU et al.: USING SIMULATED SPECTRA TO TEST THE EFFICIENCY ...
436 Materiali in tehnologije / Materials and technology 49 (2015) 3, 435–439
Figure 3: a) The AE spectra peak base and b) spectral peaks
Slika 3: a) Podlaga spektralnih vrhov in b) spektralni vrhovi AE-spek-
tra
Figure 1: Simulated Auger electron spectrum
Slika 1: Simuliran spekter Augerjeve elektronske spektroskopije
Figure 2: The AE spectra constituent elements: the primary back-
ground, the peaks base, and the peaks
Slika 2: Sestavni elementi AE-spektra: primarno ozadje, podlaga
spektralnih vrhov in spektralni vrhovi
intended for facilitating the quantitative interpretation of
electron spectra (Auger and XPS spectra), and therefore
a lot of attention is paid to the detailed physical pheno-
mena related to the excitation and emission of the Auger
electron or photoelectron. The database of SESSA
contains the data of many physical parameters needed in
quantitative electron spectroscopy (AES and XPS).13 The
simulations needed for the purpose discussed in this
paper do not require such detailed simulations. The key
factor here is that the spectra resemble the real measured
ones, and that the values of the different components of
the spectra are known before the processing starts. This
fact is of utmost importance for the comparison of values
of any of the spectral components before and after the
processing.
As mentioned in the introduction, smoothing is often
used for reducing the noise in Auger electron spectra.
For spectra measured with a energy step size 1 eV a
5-point averaging window is recommended.14 Since most
of the measured spectra that were used when building the
simulator were of energy step size 1 eV, the same
averaging window was used for processing the spectrum
shown in Figure 1. The processing was performed using
CasaXPS. The resulting spectrum is given in Figure 5.
The same simulated spectrum (Figure 1) was pro-
cessed by applying a notch filter. The threshold fre-
quency was selected arbitrarily for this case, just for a
comparison. This procedure was completed using
Audacity, where this procedure is used to reduce the
noise in sound files. The resulting spectrum is given in
Figure 6.
3 RESULTS AND DISCUSSION
After applying the smoothing procedures to the
simulated spectrum, the noiseless simulated signal
(Snoiseless) was subtracted from the processed spectrum
(Sprocessed), thus obtaining the remaining noise after pro-
cessing (Nremaining):
Nremaining = Sprocessed – Snoiseless (1)
Such a procedure was used for obtaining the values
of the remaining noise for each data point. These values
were then compared to the initial simulated noise, Ninitial,
which is the random noise added to the simulated spec-
trum, thus obtaining the noise ratios that serve as a
measure of the efficiency of the processing software in
reducing the noise and bringing the signal closer to the
noiseless one:
Nratio = Nremaining /Ninitial (2)
Figure 7 illustrates the concept behind the use of the
noise ratios for this kind of evaluation of the efficiency in
noise reduction.
As may be inferred from Figure 7, when manipula-
ting the signal for reducing the noise the obtained signal
will have a new value with a different deviation from the
target point, the noiseless signal. This new difference
from the noiseless signal, the remaining noise, will be
smaller than, equal to, or greater than that of the initial
noise. Thus, for one specific data point a noise ratio of
B. PONIKU et al.: USING SIMULATED SPECTRA TO TEST THE EFFICIENCY ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 435–439 437
Figure 6: Processed spectrum from Figure 1 using a notch filter in
Audacity
Slika 6: AE-spekter s slike 1 po uporabi ozkopasovnega filtra (pro-
gram Audacity)
Figure 4: The randomly defined primary backgrounds
Slika 4: Naklju~no dolo~ena primarna ozadja
Figure 5: Processed spectrum from Figure 1 using a 5-point window
smoothing in CasaXPS
Slika 5: AE-spekter s slike 1 po uporabljenem glajenju z oknom {irine
5 to~k (program CasaXPS)
less than one means that the processing was successful in
reducing the noise, a noise ratio of one means that the
noise level is kept the same, and a noise ratio of more
than one means that the noise level is actually increased
due to the processing of the spectrum. A graphical plot
of the noise ratios would serve as a quick assessment at a
glance with respect to the success of the noise-reduction
routine, as will be shown later on.
By using Equation (1), first the values for the remain-
ing noise were found for each data point, and then by
dividing at the respective data points according to
Equation (2), the noise ratios were found and recorded in
the data sheet, as shown in Table 1 for the spectrum pro-
cessed in CasaXPS, and in Table 2 for the spectrum
processed in Audacity.
By plotting the obtained noise ratios for each data
point according to Equation (2), the graph obtained will
give, at a glance, an indication of the improvement with
regards to the noise. Figure 8 shows such graphs for the
two processing routes discussed in this paper.
As can be seen in Figures 8a and 8b, in both cases
most of the points representing the noise ratios occupy
the region between –1 and 1, while some of the peaks lie
outside these boundaries. If the ratio of the remaining
noise to the initial noise is less than 1, this indicates that
the new signal after processing is actually closer to the
real signal than the one before processing, thus indi-
cating an improvement with respect to the noise. The
noise ratios whose values lie outside the [–1,1] interval
B. PONIKU et al.: USING SIMULATED SPECTRA TO TEST THE EFFICIENCY ...
438 Materiali in tehnologije / Materials and technology 49 (2015) 3, 435–439
Figure 7: Concept of noise ratios as a measure of the efficiency in
noise reduction
Slika 7: Koncept razmerja {umov kot merilo u~inkovitosti zmanj{e-
vanja {uma
Figure 8: Noise ratios from: a) the 5-point smoothed spectrum and b)
the spectrum smoothed using a notch filter
Slika 8: Razmerja amplitud {umov pri: a) filtriranju spektra s 5-to~-
kovnim povpre~enjem in b) pri filtriranju spektra z ozkopasovnim
filtrom
Table 1: Data sheet with the noise ratios from the 5-point smoothed
Auger spectrum
Tabela 1: Razmerja amplitud {umov pri glajenju spektra s 5-to~-
kovnim povpre~enjem
Table 2: Data sheet with the noise ratios from the AE spectrum pro-
cessed using a notch filter
Tabela 2: Razmerja amplitud {umov pri filtriranju spektra z ozkopa-
sovnim filtrom
indicate that for those specific data points the processing
has actually worsened the situation with respect to the
noise. Plotting these noise ratios for all the data points in
the spectrum offers a convenient way to see at a glance
whether there is an improvement in terms of the noise or
not, as well as for comparing different processing tech-
niques for this purpose. Again, the usefulness of the
simulated spectra must be stressed in this regard, because
no such comparison can be made if the values for the
noise are not known at the beginning, as it is in the case
of the measured spectra.
4 CONCLUSIONS
Simulated spectra have been used to assess the
performance of two noise-reduction techniques. A sim-
ple idea of using the noise ratios as a measure of the
efficiency of the noise-reduction routines was presented.
By applying this idea on a simulated spectrum, which
was processed using two different procedures, the values
for the noise ratios at each data point were found for the
respective procedures. Plotting the noise ratios provided
a convenient way to assess, at a glance, the efficiency of
the processing route in reducing the noise. The noise
ratios for the majority of the data points lie in the [–1,1]
interval, indicating an improvement with respect to the
noise.
The only way in which the values of the noise ratios
can be obtained is if the values of the noise and the
noiseless signal are known before and after the process-
ing. Such a condition can be fulfilled if simulated spectra
are used in the assessment stage, as shown in this paper.
5 REFERENCES
1 D. R. Baer et al., Challenges in Applying Surface Analysis Methods
to Nanoparticles and Nanostructured Materials, Journal of Surface
Analysis, 12 (2005) 2, 101–108
2 D. R. Baer et al., Characterization challenges for nanomaterials, Sur-
face and Interface Analysis, 40 (2008), 529–537, doi:10.1002/sia.
2726
3 D. R. Baer, D. J. Gaspar, P. Nachimuthu, S. D. Techane, D. G. Cast-
ner, Application of Surface Chemical Analysis Tools for Charac-
terization of Nanoparticles, Analytical and Bioanalytical Chemistry,
396 (2010) 3, 983–1002, doi:10.1007/s00216-009-3360-1
4 A. S. Karakoti et al., Preparation and characterization challenges to
understanding environmental and biological impacts of ceria
nanoparticles, Surface and Interface Analysis, 44 (2012), 882–889,
doi:10.1002/sia.5006
5 D. R. Baer, Application of surface analysis methods to nanomate-
rials: summary of ISO/TC 201 technical report: ISO 14187:2011 –
surface chemical analysis – characterization of nanomaterials, Sur-
face and Interface Analysis 44 (2012), 1305–1308, doi:10.1002/sia.
4938
6 M. P. Seah, Summary of ISO/TC 201 Standard: XXIII, ISO
24236:2005 – Surface chemical analysis – Auger electron spectro-
scopy – Repeatability and constancy of intensity scale, Surface and
Interface Analysis 39 (2007), 86–88, doi:10.1002/sia.2493
7 M. P. Seah, Summary of ISO/TC 201 Standard XI. ISO 17974: 2002
– Surface chemical analysis – High-resolution Auger electron
spectrometers – Calibration of energy scales for elemental and
chemical-state analysis, Surface and Interface Analysis, 35 (2003),
327–328, doi:10.1002/sia.1529
8 M. P. Seah, Summary of ISO/TC 201 Standard XII. ISO 17973:2002
– Surface chemical analysis – Medium – resolution Auger electron
spectrometers – Calibration of energy scales for elemental analysis,
Surface and Interface Analysis, 35 (2003), 329–330, doi:10.1002/
sia.1530
9 N/A. Auger Electron Spectroscopy (AES): What is AES?
http://www.phi.com/techniques/aes.html (accessed November 2014).
Physical Electronics, Inc. Chanhassen, MN, 2006 – 2014.
10 VG Scientific, Microlab 310 – F: Operators Manual, VG Scientific,
East Sussex, 1997
11 B. Poniku, I. Beli~, M. Jenko, The Auger spectra recognition and
modeling: Modeling Auger spectra for effective background removal
and noise reduction, Lambert Academic Publishing, Saarbrücken
2011, doi:10.13140/2.1.3620.7045
12 I. S. Gilmore, M. P. Seah, Savitzky and Golay differentiation in AES,
Applied Surface Science, 93 (1996) 3, 273–280, doi:10.1016/0169-
4332(95)00345-2
13 W. S. M. Werner, W. Smekal, C. J. Powell, NIST Database for the
Simulation of Electron Spectra for Surface Analysis (SESSA) –
User’s Guide, Version 1.3, National Institute of Standards and
Technology, Maryland, 2011
14 N. Fairley, CasaXPS Manual 2.3.15: Introduction to XPS and AES,
Casa Software Ltd., Devon, 2009
B. PONIKU et al.: USING SIMULATED SPECTRA TO TEST THE EFFICIENCY ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 435–439 439

Y. Y. ÖZBEK, M. DURMAN: SURFACE BEHAVIOR OF AISI 4140 MODIFIED WITH THE PULSED-PLASMA TECHNIQUE
SURFACE BEHAVIOR OF AISI 4140 MODIFIED WITH THE
PULSED-PLASMA TECHNIQUE
LASTNOSTI POVR[INE AISI 4140, SPREMENJENE S TEHNIKO
PULZIRAJO^E PLAZME
Yýldýz Yaralý Özbek, Mehmet Durman
Sakarya University, Engineering Faculty, Department of Metallurgical and Metarials Engineering, Esentepe Campus, 54187 Sakarya, Turkey
yyarali@sakarya.edu.tr
Prejem rokopisa – received: 2013-10-02; sprejem za objavo – accepted for publication: 2014-07-17
doi:10.17222/mit.2013.219
In this study, the microstructure and surface properties of a low-alloy steel (AISI 4140) treated with pulsed plasma were
investigated. Three different plasma-gun nozzle distances of (60, 70 and 80) mm and one battery capacity were chosen for a
surface modification. The modified surface layers were examined using a light microscope and X-ray analyses were carried out
for all the samples. The X-ray diffraction confirmed a development of new phases after the surface treatment. The samples were
subjected to micro-hardness measurements and it was found that the hardness values of the modified surfaces were four times
higher than those of the untreated samples. The samples were immersed into liquid nitrogen and then broken in a Charpy
machine. The fractured surfaces were exposed to SEM and EDS analyses. At the end of the study, thin grains originating from
the consumable electrode were detected. After the pulsed-plasma treatment, new structures were obtained.
Keywords: pulsed plasma, fracture, consumable electrode, modification
V tej {tudiji so bile preiskovane mikrostruktura in lastnosti povr{ine malolegiranega jekla (AISI 4140) po obdelavi s pulzirajo~o
plazmo. Za spremembo povr{ine so bile izbrane tri razli~ne razdalje plazemske {obe od povr{ine: 60 mm, 70 mm in 80 mm, in
kapaciteta ene baterije. Spremenjene povr{inske plasti so bile preiskovane s svetlobnim mikroskopom in izvr{ene so bile rent-
genske analize. Rentgenska difrakcija je potrdila nastanek novih faz po obdelavi povr{ine. Na vzorcih so bile izvr{ene meritve
mikrotrdote in ugotovljeno je bilo, da je trdota spremenjene povr{ine {tirikrat ve~ja od tiste pri neobdelanih vzorcih. Le-ti so bili
potopljeni v teko~i du{ik in nato prelomljeni na napravi Charpy. Povr{ina preloma je bila analizirana s SEM in EDS. Odkrita so
bila drobna zrna, ki izvirajo iz elektrode. Po obdelavi s pulzirajo~o plazmo je nastala nova mikrostruktura.
Klju~ne besede: pulzirajo~a plazma, prelom, porabljiva elektroda, sprememba
1 INTRODUCTION
High-intensity pulse or plasma irradiation has
recently gained a growing interest as a potential tool in
surface engineering. As in the cases of a laser or electron
beam, ions from a pulsed beam rapidly heat the surface
of the irradiated material. The surface remains at a high
temperature (up to the melting point or higher) for a
period in the nanosecond to microsecond range, and then
rapidly cools through conduction into the bulk at the
rates of the order of l07–1010 K/s. Obviously, the details
of the heat evolution in a substrate depend on its thermal
properties and dimensions as well as on the beam para-
meters. Heat-induced processes result in several non-
equilibrium phenomena such as the mixing of metallic
overlayers on various (even hardly miscible) substrates,
the formation of metastable crystalline alloys and so on.
Besides purely thermal effects, ion or plasma beams are
used as well. It is also possible, under appropriate con-
ditions, to modify the surface properties of solids via
thermal effects in conjunction with the mass transport.1–3
The pulsed-plasma process is used to improve the
surface properties of the workpieces of tool steels.1–6 The
pulsed-plasma system has high rates of heating and
cooling. These lead to the formation of a nano/microcry-
stalline structure, a high dislocation density and a growth
of the concentration of the alloying elements and, thus,
an intensification of the diffusion mechanisms.3–5
In general, a modified surface consists of a com-
pound layer (the white layer) that is a few micrometers
deep. In the diffusion zone the nitrogen atoms can be
interstitially dissolved or precipitated as iron nitrides,
tungsten and/or tungsten alloys from the consumable
electrode used. As a result, metastable states may appear
in the surface layers, which are the origin of the
improved physical, chemical and mechanical properties
unattainable with the conventional surface-treatment
techniques.5
The process parameters have significant effects in
determining the final structure and mechanical properties
of the surfaces. Among these parameters the controlling
gas diffusion through the nozzle distance and the number
of pulses, and the plasma composition are the most im-
portant ones. A proper combination of these parameters
would provide the best surface properties and set out the
duration of the process time as an important economic
factor.1 The pulsed plasma is the most advantageous one;
its process time is very short (1 min), it is more eco-
nomical and it can produce superior mechanical proper-
ties.
Materiali in tehnologije / Materials and technology 49 (2015) 3, 441–445 441
UDK 669.14:621.78.015 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(3)441(2015)
The aim of this study is to improve the material
performance and functional properties of the surface of a
commonly used steel without long and expensive heat-
treatment operations. AISI 4140 was used in this study.
The AISI 4140 steel is the most common type of steel
dealt with in the studies discussing different methods.
Different parameters of pulsed plasma were tried on the
samples.
2 EXPERIMENTAL PROCEDURE
The standard, medium-carbon, low-alloy AISI 4140
steel was used in the study and the chemical composition
of the steel is given in Table 1. The samples were ex-
posed to the pulsed-plasma-modification technique. A
schematic illustration of the pulsed-plasma technique
used for the modification of the samples is shown in
Figure 1. The parameters of pulsed plasma are given in
Table 2.
Table 1: Chemical composition of the AISI 4140 steel used in the
study in mass fractions, w/%
Tabela 1: Kemijska sestava jekla AISI 4140, uporabljenega za
preizkuse v masnih dele`ih, w/%
AISI
4140 C Si Mn P S Cr Mo
w/% 0.40 0.30 0.70 0.035 0.035 0.98 0.27
Table 2: Parameters of the pulsed-plasma process
Tabela 2: Parametri postopka pulzirajo~e plazme
No. Nozzledistance mm
Number of
pulses
Battery
capacity, μF
Consumable
electrode, W
1 70 15 800 Tungsten
2 70 10 800 Tungsten
3 70 5 800 Tungsten
4 80 15 800 Tungsten
5 80 10 800 Tungsten
6 80 5 800 Tungsten
7 60 15 800 Tungsten
8 60 10 800 Tungsten
9 60 5 800 Tungsten
The specimens used in the pulsed-plasma experi-
ments were cut from the center of the modified surfaces
with a cutting machine (Discotom-6) and then put in
Bakelite. After that, they were grinded with the emery
paper and polished. The samples were studied with a
light microscope using different magnifications. The
hardness values of the specimens were measured with a
Future-Tech test apparatus for 15 s under a load 5 g. The
phase compositions of the modified surfaces were
investigated with X-ray diffractometry (XRD) using a
Rigaku diffractometer employing monochromatic Cu-K
radiation. After the pulsed-plasma process, the notched
specimens were immersed into liquid nitrogen for two
minutes and then they were cracked with a Charpy test
machine from the notched regions. The fractured regions
were analyses with a scanning electron microscope
JEOL 6600 (SEM) and an energy dispersive spectro-
meter (EDS).
3 RESULTS AND DISCUSSION
In Figure 2, a light microstructure of Sample 4 is
shown. It was obvious from the microstructural exami-
nation that the modified layer and the substrate could be
easily seen due to the light contrast. A white layer was
formed on the surface layer.
A homogeneous and ordered layer was obtained by
increasing the pulse number as seen in the light
Y. Y. ÖZBEK, M. DURMAN: SURFACE BEHAVIOR OF AISI 4140 MODIFIED WITH THE PULSED-PLASMA TECHNIQUE
442 Materiali in tehnologije / Materials and technology 49 (2015) 3, 441–445
Figure 2: Microphotograph of the cross-section of Sample 4 modified
with pulsed plasma
Slika 2: Mikrostruktura prereza vzorca 4, obdelanega s pulzirajo~o
plazmo
Figure 1: Schematic presentation of the pulsed-plasma modification
system: 1-detonation chamber, 2-central electrode, anode, 3-conical
electrode, cathode, 4-interelectrode gap, 5-consumable electrode,
6-power supply, 7-gap between the electrodes, 8-pulsed plasma
forming, 9-work surface
Slika 1: Shematski prikaz sistema s pulzirajo~o plazmo: 1-detona-
cijska komora, 2-centralna elektroda anoda, 3-koni~na elektroda
katoda, 4-vrzel med elektrodama, 5-porabljiva elektroda, 6-izvir
energije, 7-vrzel med elektrodama, 8-nastanek pulzirajo~e plazme, 9-
delovna povr{ina
photographs. When the pulse number was increased, the
thickness of the modified layer increased as due to the
increased pulse number the process time and the amount
of ionized gases were increased.6,7
Rapid heating and solidification induced a heavy
plastic deformation, which caused a formation of dislo-
cation cells due to one pulse bombardment. After the
multi-pulse bombardments, both austenite and carbide
types of the nanostructure particles were formed from
the supersaturated solution.
The pulsed-plasma system affected the grain size of
the modified layer.4 The grain structure of the outer
surface of the modified layer was very fine and dense.
The previous studies showed that the thermal stability,
mechanical and tribological properties and the corrosion
behavior of the materials are greatly influenced by the
grain size.2,8,9
Figure 3 shows the EPMA results for the modified
layer. The amounts of nitrogen and tungsten changed
from the surface to the inner space. Firstly, the nitrogen
was increased, then it was decreased and after that point
the amount of tungsten was increased. These phases
were very important for the surface properties.8
The XRD results for the sample groups are given in
Figure 4. The figure clearly shows that the pulsed-
plasma treatment changed the diffraction profiles of the
samples. The results of the X-ray analysis indicate that
the new phases (such as -Fe, W and carbides) are
formed on the steel surface after the plasma processing.
Tungsten oxide was also determined from the analysis,
due to the tungsten consumable electrode used during the
pulsed-plasma treatment.10
While there was only the -Fe phase in the untreated
AISI 4140 steel (Figure 4a), Fe3N, -Fe and tungsten
were observed in the modified layer after the treatment
(Figure 4b). As seen in Figure 4b, the pulse number is
the only difference between the three specimens. The
increased number of pulses leads to the growth of some
existing phases, such as ’-Fe, and the formation of new
phases since the increasing pulse number also increases
the pulsed-plasma-treatment time. This results in an in-
creased amount of the ionized products doped into the
surface. In addition, crystalline phases can be formed
more easily with the increasing treatment time.6,7
Sample 7, having the highest pulse number, has the
largest peak spacing (increased FWHM values). The
increasing FWHM values of these phases cause a
decrease in the grain size in accordance with the Scherer
equation.8 Besides, a high cooling rate results in a small
size of the precipitates, making them hard to be found
with XRD.8,9
Y. Y. ÖZBEK, M. DURMAN: SURFACE BEHAVIOR OF AISI 4140 MODIFIED WITH THE PULSED-PLASMA TECHNIQUE
Materiali in tehnologije / Materials and technology 49 (2015) 3, 441–445 443
Figure 4: XRD analysis of: a) un-treated AISI 4140 steel, b) Sample
7, Sample 8 and Sample 9
Slika 4: Rentgenogram: a) neobdelano AISI 4140 jeklo, b) vzorec 7,
vzorec 8 in vzorec 9
Figure 3: EPMA results for the modified layer of Sample 12
Slika 3: Rezultati EPMA spremenjene plasti vzorca 12
A comparison of the X-ray analysis results for the
sample surfaces processed with different numbers of
pulses showed a profound effect on the creation of new
phases in the surface structure, hence, enhancing the
surface properties.10,11
Figure 5 shows the results of the microhardness
measurements for all the samples. The figure clearly
illustrates that the pulsed-plasma treatment has a pro-
found effect on increasing the microhardness values of
the AISI 4140 steel, depending on the treatment parame-
ters used. The new phases and the structure transforma-
tions improved the surface mechanical characteristics.
It was also found that there were some tungsten,
nitride and austenite phases on the surfaces of the
modified samples. The amount of these strong phases
precipitated in the matrix phase was one of the reasons
leading to an increase in the hardness values of the
treated samples.4,12–14
Prior to the pulsed-plasma treatment, the initial hard-
ness value of the specimens was recorded as 180 HV and
later its value was increased up to 700–950 HV. The
hardness value measured for Sample 1 (15 pulses) was
910 HV; for Sample 2 (10 pulses) it was 900 HV and for
Sample 3 (5 pulses) it was 820 HV. The number of
pulses affected the hardness values. An increase in the
pulse number, or the intensity of the energy density ab-
sorbed by the surface, leads to the growth of the micro-
hardness value.2,4,9
In the experiments, the maximum value of the surface
microhardness was achieved for the samples treated with
the maximum number of pulses. However, in these treat-
ment regimes, a partial melting of the surface layer was
observed, being due to the high temperatures caused by
the frequently repeated pulses.
The structural defects, or disorder trappings, are
easily formed at the high solidification rate induced by
pulsed plasma, increasing the material hardness, known
as the established point-defect strengthening models.13
Another reason for the increase in the hardness
values of the samples was the decreased grain size due to
Y. Y. ÖZBEK, M. DURMAN: SURFACE BEHAVIOR OF AISI 4140 MODIFIED WITH THE PULSED-PLASMA TECHNIQUE
444 Materiali in tehnologije / Materials and technology 49 (2015) 3, 441–445
No. 1 2 3 4 5 6 7 8 9
W
w/% 2.610 3.200 1.198 2.067 1.154 1.190 2.091 1.889 1.190
Figure 6: EDS analyses of the fracture surface of Sample 3
Slika 6: EDS-analize na povr{ini preloma vzorca 3
Figure 5: Microhardness values of the samples
Slika 5: Mikrotrdota vzorcev
Figure 7: SEM micrographs of fracture surfaces: a) Sample 5, b)
Sample 6, c) Sample 7
Slika 7: SEM-posnetki povr{ine preloma: a) vzorec 5, b) vzorec 6, c)
vzorec 7
the increased cooling rate within the cycles of pulses. In
the present experiment, the samples subjected to a high
number of pulses during the pulsed-plasma treatment
exhibit a low grain size and a high amount of the
phase.11,13
The EDS analyses of the modified layer of Sample 3
are given in Figure 6. The modified layer and the resul-
tant structure were easily detected and clearly seen in the
fractured regions of the specimens. The EDS analysis
was performed and the tungsten element was observed.
In addition to this, the consumable-electrode tungsten
amount in the matrix phase was calculated on the basis
of the EDS analysis. Since there is no tungsten ingre-
dient in the base metal, the existence of the tungsten
from the electrode after the analysis indicates that the
pulsed-plasma treatment was performed successfully.
The process parameters affect the structure and
thickness of the modified layer (Figures 7a to 7c). From
the SEM micrographs, it is observed that the fracture
mechanism was brittle as characterized by the cleavage
facets in the bulk material.10,11 Different fracture mecha-
nisms occurred on the modified layer and the substrate.
A fine-grained structure observed in the modified zone
where a ductile fracture occurred is an indication of a
hard structure. But ultra-fine intermetallic compounds
were found in the matrix of this zone.4,14
4 CONCLUSION
In the light of the results of the experimental studies
carried out on the surfaces of AISI 4140 steels modified
with the pulsed-plasma technique, the findings given
below were obtained:
1) The pulsed-plasma technique is used for a surface
modification. Due to this technique, the thickness of a
modified layer increases with the increasing pulse num-
ber, and the resultant structure becomes homogeneous.
2) In addition, a decrease in the grain size is observed
as a result of fast heating and cooling of the modified
layer. The grains outside of the modified zone are larger
than the ones in the inside regions.
3) The phase and structure transformations occurred
during the modification of the surfaces. The -Fe, Fe3N
and W compounds were determined with the XRD
studies.
4) The new phases and structure transformations
improved the surface mechanical characteristics. The
hardness values increased four times.
5) A ductile fracture occurred on the modified layer.
A cleavage fracture was seen on the substrate.
5 REFERENCES
1 K. M. Zhang, J. X. Zou, B. Bolle, T. Grosdidier, Vacuum, 87 (2013),
60–68, doi:10.1016/j.vacuum.2012.03.061
2 N. Y. Tyurin, O. V. Kolisnichenko, N. G. Tsygankov, The Paton
Welding Journal, (2004), 38–43
3 C. Kwietniewski, W. Fontana, C. Moraes, A. S. Rocha, T. Hirsch, A.
Reguly, Surface and Coatings Technology, 179 (2004) 1, 27–32,
doi:10.1016/S0257-8972(03)00795-3
4 Q. F. Guan, H. Zou, G. T. Zou, A. M. Wu, S. Z. Hao, J. X. Zou, Y.
Qin, C. Dong, Q. Y. Zhang, Surface & Coatings Technology, 196
(2005) 1–3, 145–149, doi:10.1016/j.surfcoat.2004.08.104
5 J. L. Fan, T. Liu, H. C. Cheng, D. L. Wang, Journal of Materials
Processing Technology, 208 (2008) 1–3, 463–469, doi:10.1016/
j.jmatprotec.2008.01.010
6 Y. Y. Özbek, M. Durman, H. Akbulut, Tribology Transactions, 52
(2009) 2, 213–222, doi:10.1080/10402000802369721
7 V. V. Uglov, V. M. Anishchik, N. N. Cherenda, Y. V. Sveshnikov, V.
M. Astashynski, E. A. Kostyukevich, A. M. Kuzmitski, V. V. Asker-
ko, Surface & Coatings Technology, 202 (2008) 11, 2439–2442,
doi:10.1016/j.surfcoat.2007.08.045
8 Y. Yarali, M. Durman, H. Akbulut, Key Engineering Materials,
264–268 (2004), 561–564, doi:10.4028/www.scientific.net/KEM.
264-268.561
9 Z. Zhou, J. Linke, G. Pintsuk, J. Du, S. Song, C. Ge, Journal of Nu-
clear Materials, 386–388 (2009), 733–735, doi:10.1016/j.jnucmat.
2008.12.306
10 A. Buranawong, N. Witit-anun, S. Chaiyakun, A. Pokaipisit, P. Lim-
suwan, Thin Solid Films, 519 (2011) 15, 4963–4968, doi:10.1016/
j.tsf.2011.01.062
11 Y. Hao, B. Gao, G. F. Tu, S. W. Li, C. Dong, Z. G. Zhang, Nuclear
Instruments and Methods in Physics Research B, 269 (2011) 13,
1499–1505, doi:10.1016/j.nimb.2011.04.010
12 Y. Y. Özbek, H. Akbulut, A. Ozel, M. Durman, Surface Modification
of M2 Steel by Pulsed Plasma Technique, ICAMM, 2010, 46–50
13 J. X. Zou, T. Grosdidier, K. M. Zhang, C. Dong, Applied Surface
Science, 255 (2009) 9, 4758–4764, doi:10.1016/j.apsusc.2008.10.123
14 X. Cheng, S. Hu, W. Song, X. Xiong, Applied Surface Science, 286
(2013), 334–343, doi:10.1016/j.apsusc.2013.09.083
Y. Y. ÖZBEK, M. DURMAN: SURFACE BEHAVIOR OF AISI 4140 MODIFIED WITH THE PULSED-PLASMA TECHNIQUE
Materiali in tehnologije / Materials and technology 49 (2015) 3, 441–445 445

A. KHOLODKOVA et al.: PREPARATION AND DIELECTRIC PROPERTIES OF THERMO-VAPOROUS BaTiO3 CERAMICS
PREPARATION AND DIELECTRIC PROPERTIES OF
THERMO-VAPOROUS BaTiO3 CERAMICS
PRIPRAVA IN DIELEKTRI^NE LASTNOSTI TERMO-PARNO
POROZNE KERAMIKE BaTiO3
Anastasia Kholodkova1, Marina Danchevskaya1, Nellya Popova2,
Liana Pavlyukova2, Alexandr Fionov3
1Chemistry Department, Moscow State University, GSP-1, Leninskie Gory 1-3, 119991 Moscow, Russia
2D. Mendeleev University of Chemical Technology of Russia, Geroev Panfilovtsev 20, 125047 Moscow, Russia
3Kotel’nikov Institute of Radio Engineering and Electronics of RAS, Mokhovaya 11-7, 125009 Moscow, Russia
anastasia.kholodkova@gmail.com
Prejem rokopisa – received: 2013-11-10; sprejem za objavo – accepted for publication: 2014-07-17
doi:10.17222/mit.2013.276
A crystalline BaTiO3 powder was synthesized at 350 °C for 0–20 h from TiO2 (> 99 % purity) and BaO (> 98 % purity) with
water vapour acting as the reaction media. According to the XRD and SEM results, the BaTiO3 synthesized for 3 h proved to be
the most adequate raw material for ferroelectric ceramics among the obtained samples as long as it consisted of pure crystalline
sphere-shaped BaTiO3 particles with the average size of 156 nm. Pellets were pressed at (100, 150 and 200) MPa and sintered at
1300 °C for 1 h. The influence of the compacting pressure on the dielectric characteristics of BaTiO3 ceramics was studied by
monitoring the permittivity and loss-tangent values of the pellets at 20 Hz–2 MHz.
Keywords: barium titanate, thermo-vaporous synthesis, microstructure, dielectric ceramics
Kristalni prah BaTiO3 je bil sintetiziran pri 350 °C od 0 do 20 h iz TiO2 (~istost: > 99 %) in BaO (~istost: > 98 %) z vodno paro
kot reakcijski medij. Iz rentgenogramov in SEM-posnetkov izhaja, da je BaTiO3, sintetiziran 3 h, najprimernej{a surovina med
vsemi vzorci za feroelektri~no keramiko, dokler sestoji iz ~istih kristalnih okroglih delcev BaTiO3 s povpre~no velikostjo 156
nm. Peleti so bili stiskani pri (100, 150 in 200) MPa in sintrani 1 h pri temperaturi 1300 °C. Vpliv tlaka pri stiskanju na
dielektri~ne lastnosti keramike BaTiO3 je bil preiskovana s spremljanjem permitivnosti in velikosti izgube tg  peletov pri 20
Hz–2 MHz.
Klju~ne besede: barijev titanat, termo-parno porozna sinteza, mikrostruktura, dielektri~na keramika
1 INTRODUCTION
Since the 1940s barium titanate, BaTiO3, has been
known for its extremely high values of ferroelectric cha-
racteristics which make it widely used in the production
of multilayer ceramic capacitors (MLCCs), resistors with
a positive temperature coefficient of resistivity (PTCR),
temperature/humidity/gas sensors, piezoelectric transdu-
cers and actuators, ultrasonic and electro-optic devices,
IR-detectors, etc.1–4 Generally, the BaTiO3 powder is
obtained as the raw material for the bulk-ceramic manu-
facturing, as well as thin-film and composite-material
production.4 For this purpose a homogeneous, well-dis-
persed pure BaTiO3 powder, consisting of spherical
particles up to 200 nm in size, is required.1,5,6 Various
synthesis routes for the as-characterized powder have
been developed over several decades. In addition to the
conventional solid-state method, the techniques such as
pyrolysis (Pechini, citrate processes), dispersion (cate-
cholate synthesis, spray pyrolysis, sol-gel), precipitation
(oxalate, hydrothermal and solvothermal synthesis) are
widely used for the fine-crystalline BaTiO3 process-
ing.1,7–14 But, as some of them are multistage and require
the use of auxiliary substances, mostly in the solid state,
hydrothermal and oxalate techniques are industrially
suitable. The development of a simple and low-cost
method for the industrial BaTiO3 production remains to
be a pending problem.
In the present work water vapour at 350 °C was used
as the medium for a BaTiO3 synthesis from simple
oxides. A similar technique combined with the treatment
in supercritical water fluid was previously successfully
used for the production of MgAl2O4, ZnAl2O4, Y3Al5O12,
BaFe12O19, LiNbO315,16 and also BaTiO3.17,18 As ceramic
manufacturing is one of the main application areas for
the BaTiO3 powder, steps were taken to develop this
technique for obtaining the BaTiO3 with the thermo-
vaporous process. We prepared ceramic samples in the
same conditions, but varied the compressing pressure
and studied the phase content, the microstructure and
dielectric properties of the pellets in order to determine
the most appropriate value of the pressure for the cera-
mic-manufacturing route.
2 EXPERIMENTAL WORK
The synthesis of BaTiO3 was performed in laboratory
stainless-steel autoclaves using BaO (> 98 % purity) and
TiO2 (> 99 % purity) as the starting reagents. As BaO
Materiali in tehnologije / Materials and technology 49 (2015) 3, 447–451 447
UDK 666.3/.7:621.762.5 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(3)447(2015)
interacts with CO2 in the air, forming BaCO3, in order to
avoid a lack of Ba2+ ions in the reaction mixture, an
amount fraction of excess of BaO 5 % was used. After a
thorough mixing by means of grinding in an agate
mortar with a pestle, the reagents were placed into a
special container inside the autoclave, separated from the
distilled water. Hermetically closed autoclaves were
heated up to 350 °C (the water-vapour pressure of 16
MPa), kept at this temperature for (0, 0.5, 1, 2, 3, 4, 5,
20) h and then cooled so that the water vapour condensed
at the bottom of the autoclaves separated from the pro-
duct. The product was first washed with acetic acid
solution to avoid a BaCO3 contamination and then with
distilled water.
To produce a ceramic powder, the sample synthesized
for 3 h (named BT-3h) was mixed with 1 % PVA and
uniaxially pressed into pellets at (100, 150 and 200) MPa
at room temperature. The pellets were sintered at 1300
°C for 1 h.
The phase contents of the powders and ceramics were
identified with an X-ray diffraction analysis (STOE
STADI P) using the Cu-K radiation in a range of 20 ° #
2 # 80 °. The crystallite size was calculated with the
Scherrer equation. The morphologies of the powder and
ceramic samples were studied with scanning electron
microscopy (JSM-6390 LA). The dielectric permittivity
and loss tangent of the pellets were calculated from the
capacity and the conductivity, respectively, of the plane
condenser, in which each pellet was used as a dielectric
layer. Dielectric characteristics were obtained with a pre-
cision LCR meter Agilent E4980a with a frequency
range of 20 Hz–2 MHz.
3 RESULTS AND DISCUSSION
During the thermo-vaporous BaTiO3 synthesis, H2O
molecules from the vapour became incorporated into the
TiO2 structure due to the dissociative absorption mani-
fested in the breaking of the Ti-O bonds and the creation
of the Ti-OH bonds. In these conditions the TiO2 struc-
ture becomes more flexible, interacting with the Ba2+
ions and reorganising into BaTiO3. The XRD analysis of
the powders prepared at 350 °C in the water-vapour
atmosphere over the periods of 0–20 h showed that the
powders consisted of crystalline BaTiO3 (Figure 1). The
formation of BaTiO3 from TiO2 and Ba(OH)2 occurred
already during the heating, thus, the sample synthesized
for 0 h contained only crystalline BaTiO3. The interac-
tion of the newly formed BaTiO3 phase with water
vapour led to an elimination of lattice defects and to a
perfection of the crystalline structure. Figure 2 shows
the BaTiO3 crystallite-size dependence on the duration of
the thermo-vaporous synthesis, calculated from the
Scherrer equation. The crystallite size of the samples
synthesized for 0–4 h fluctuates in a range of 35–45 nm,
while, in the case of a longer synthesis, the crystallite
A. KHOLODKOVA et al.: PREPARATION AND DIELECTRIC PROPERTIES OF THERMO-VAPOROUS BaTiO3 CERAMICS
448 Materiali in tehnologije / Materials and technology 49 (2015) 3, 447–451
Figure 3: Box charts of the crystal-size distribution of BaTiO3
synthesized in water vapour at 350 °C and 16 MPa for 0–20 h
Slika 3: [katlasti diagram razporeditve velikosti kristalnih zrn
BaTiO3, sintetiziranih od 0 do 20 ur, v vodni pari pri 350 oC in 16
Figure 1: XRD patterns of the BaTiO3 powders synthesized in water
vapour at 350 °C and 16 MPa for 0–20 h
Slika 1: Rentgenogram prahov BaTiO3 sintetiziranih od 0 do 20 ur, v
vodni pari pri 350 °C in 16 MPa
Figure 2: Crystal-size distributions of BaTiO3 synthesized in water
vapour at 350 °C and 16 MPa for 0–20 h
Slika 2: Razporeditev velikosti kristalov BaTiO3, sintetiziranih od
0 do 20 ur, v vodni pari pri 350 oC in 16 MPa
size is reduced. This effect can be explained with the
interaction of the excessive amounts of Ba(OH)2 in the
reacting mixture with the already formed BaTiO3. It is
known that the crystallite size from the Scherrer equation
is sensitive to phase inhomogeneities.
In the SEM images of the synthesized samples the
crystals of BaTiO3 exhibit a narrow size distribution. The
average crystal size slightly varies in a range of 150–188
nm without a distinct relation to the duration of the
synthesis (Figure 3). There is a clear effect of the
reaction time on the shape of the crystals. The samples
processed for 0–3 h consist of sphere-shaped particles
(Figures 4a and 4b). A longer processing leads to a
A. KHOLODKOVA et al.: PREPARATION AND DIELECTRIC PROPERTIES OF THERMO-VAPOROUS BaTiO3 CERAMICS
Materiali in tehnologije / Materials and technology 49 (2015) 3, 447–451 449
Figure 4: SEM images of BaTiO3 synthesized in water vapour at 350 °C and 16 MPa for: a) 0 h, b) 3 h, c) 4 h, d) 20 h
Slika 4: SEM posnetki BaTiO3, sintetiziranega v vodni pari pri 350 oC in 16 MPa po: a) 0 urah, b) 3 urah, c) 4 urah, d) 20 urah
Figure 6: SEM images of the BaTiO3 pellets pressed at: a) 100 MPa,
b) 150 MPa, c) 200 MPa, and sintered at 1300 °C for 1 h
Slika 6: SEM posnetki BaTiO3 peletov, stisnjenih pri: a) 100 MPa, b)
150 MPa, c) 200 MPa in sintranih 1 uro na 1300 oC
Figure 5: XRD patterns of the BaTiO3 pellets pressed at 100–200
MPa and sintered at 1300 °C for 1 h and BaTiO3 powder BT 3 h used
as the raw material
Slika 5: Rentgenogram peletov BaTiO3, stisnjenih pri 100–200 MPa
in sintranih 1 uro na 1300 oC in prah BaTiO3 po 3 urah, uporabljen kot
surovina
formation of crystal facets. In the sample synthesized for
4 h cube-shaped particles could be observed among the
sphere-shaped ones and, after 20 h, the sample mainly
consisted of faceted crystals (Figures 4c and 4d). The
formation of the faceted crystals is related to the BaTiO3
structure perfection due to its interaction with water
vapour.
Ceramic pellets were manufactured from the BaTiO3
powder sample synthesized in water vapour for 3 h, so
that it consisted of pure, crystalline and sphere-shaped
BaTiO3 particles. Three pellets were shaped at different
compacting pressures, while the temperature and dura-
tion of the sintering were taken from the literature.19–23
The geometric density of the pellets was 84–86 % of the
theoretical value of 6.01 g/cm3 (Table 1) and it naturally
rose with the increase in the compaction pressure. In
addition to the BaTiO3, the XRD analysis showed the
presence of an impurity phase in every pellet at 2 = 29 °
(Figure 5), which is a complex barium aluminate titanate
originating from the milling process.
Table 1: Ceramic-processing conditions and density of the pellets
Tabela 1: Pogoji pri izdelavi keramike in gostota peletov
Pellet
Compacting
pressure,
MPa
Sintering
temperature,
°C
Density,
g/cm3
Relative
density, %
Pel-100 100
1300
5.07 84
Pel-150 150 5.14 86
Pel-200 200 5.15 86
In the SEM images the pellets consist of the grains of
160–180 nm in size. Typical features of the pellet micro-
structure are sub- and micron-sized pores and plates
(Figures 6a to 6c). The presence of the plates shows that
the recrystallization occurred during the sintering and
suggests that the temperature of 1300 °C chosen on the
basis of the reference literature is higher than the appro-
priate sintering temperature for the thermo-vaporous
BaTiO3 powder.
The frequency dependence of the dielectric permitti-
vity and loss tangent is shown in Figure 7. The values of
both parameters decrease with the increase in the
frequency. This phenomenon is common for all dielec-
trics. The higher the frequency the less polarization can
be realized in a dielectric.24 As the permittivity is a para-
meter that shows the polarizability of a dielectric,25 it is
reduced with a frequency increase. Notably, the per-
mittivity of the pellets shows a strong dependence on the
compacting pressure. The permittivity of the pellet
manufactured at 150 MPa exhibits the smallest variation
in the range of 20 kHz–2 MHz in comparison with the
other two pellets. Both the pellets made at 200 MPa and
100 MPa show a more pronounced variation in the
permittivity (Figure 7). The loss tangent shows a similar
tendency as the permittivity, depending on the frequency,
and the pellet prepared at 150 MPa exhibits the lowest
values of tg  among the three examined pellets. Com-
paring these results, it can be concluded that 150 MPa is
the most appropriate compacting pressure for thermo-
vaporous BaTiO3 ceramics.
4 CONCLUSION
The present work reports on the results of a
thermo-vaporous synthesis and ceramic processing of
crystalline BaTiO3. The optimum duration of the
synthesis in water vapour at 350 °C and 16 MPa is 3 h in
order to obtain a raw material for ceramics. The study of
the ceramic microstructure showed that the sintering
temperature for the thermo-vaporous BaTiO3 powder
should be lower than 1300 °C. For the ceramics with the
permittivity weakly dependent on the frequency in the
range of 20 Hz–2 MHz and a low loss tangent, the
compacting pressure of 150 MPa involving 1 % PVA as
the binder is the most applicable.
Acknowledgements
This work was partly supported by the M. V.
Lomonosov State University Program of Development.
The authors are thankful to Dr. G. P. Muravieva for the
help with the XRD analysis and S. A. Chernyak for the
SEM study.
5 REFERENCES
1 W. Maison, R. Kleeberg, R. B. Heimann, S. Phanichphant, Phase
content, tetragonality and crystallite size of nanoscaled barium
titanate synthesized by catecholate process: effect of calcination
temperature, J. Eur. Ceram. Soc., 23 (2003) 1, 127–132, doi:10.1016/
S0955-2219(02)00071-7
2 D. H. Yoon, B. I. Lee, Processing of barium titanate tapes with diffe-
rent binders for MLCC applications – Part I: Optimization using
design of experiments, J. Eur. Ceram. Soc., 24 (2004) 5, 739–752,
doi:10.1016/S0955-2219(03)00333-9
A. KHOLODKOVA et al.: PREPARATION AND DIELECTRIC PROPERTIES OF THERMO-VAPOROUS BaTiO3 CERAMICS
450 Materiali in tehnologije / Materials and technology 49 (2015) 3, 447–451
Figure 7: Frequency dependencies of the permittivity and loss tangent
of the BaTiO3 ceramic pellets pressed at 100–200 MPa and sintered at
1300 °C for 1 h
Slika 7: Frekven~na odvisnost permitivnosti in tangenta izgub BaTiO3
kerami~nih peletov, stisnjenih pri 100-200 MPA in sintranih 1 uro na
1300 oC
3 U. A. Joshi, S. Yoon, S. Baik, J. S. Lee, Surfactant-free hydrothermal
synthesis of highly tetragonal barium titanate nanowires: a structural
investigation, J. Phys. Chem. B., 110 (2006), 12249–12256,
doi:10.1021/jp0600110
4 D. Padalia, G. Bisht, U. C. Johri, K. Asokan, Fabrication and cha-
racterization of cerium doped barium titanate/PMMA nanocompo-
sites, Solid State Sciences, 19 (2013), 122–129, doi:10.1016/
j.solidstatesciences.2013.02.002
5 Ch. Pithan, D. Hennings, R. Waser, Progress in synthesis of nano-
crystalline BaTiO3 powders for MLCC, Int. J. Appl. Ceram.
Technol., 2 (2005) 1, 1–14, doi:10.1111/j.1744-7402.2005.02008.x
6 D. H. Yoon, Tetragonality of barium titanate powder for ceramic
capacitor application, Journal of Ceramic Processing Research, 7
(2006) 4, 343–354
7 W. S. Jung, J. H. Kim, H. T. Kim, D. H. Yoon, Effect of temperature
schedule on the particle size of barium titanate during solid-state
reaction, Materials Letters, 64 (2010) 2, 170–172, doi:10.1016/
j.matlet.2009.10.035
8 V. Vinithini, P. Singh, M. Balasubramanian, Synthesis of barium
titanate nanopowder using polymeric precursor method, Ceramics
International, 32 (2006) 2, 99–103, doi:10.1016/j.ceramint.2004.
12.012
9 P. Duran, F. Capel, J. Tartaj, D. Gutierrez, C. Moure, Heating-rate
effect on the BaTiO3 formation by thermal decomposition of metal
citrate polymeric precursors, Solid State Ionics, 141–142 (2001),
529–539, doi:10.1016/S0167-2738(01)00742-1
10 A. Purwanto, W. N. Wang, I. W. Lenggoro, K. Okuyama, Formation
of BaTiO3 nanoparticles from an aqueous precursor by flame-
assisted spray pyrolysis, J. Eur. Ceram. Soc., 27 (2007) 16,
4489–4497, doi:10.1016/j.jeurceramsoc.2007.04.009
11 M. R. Mohammadi, D. J. Fray, Sol-gel derived nanocrystalline and
mesoporous barium strontium titanate prepared at room temperature,
Particuology, 9 (2011) 3, 235–242, doi:10.1016/j.partic.2010.08.012
12 Y. B. Khollam, A. S. Deshpande, H. S. Potdar, S. B. Deshpande, S.
K. Date, A. J. Patil, A self-sustaining acid-base reaction in semi-
aqueous media for synthesis of barium titanyl oxalate leading to
BaTiO3 powders, Materials Letters, 55 (2002) 3, 175–181,
doi:10.1016/S0167-577X(01)00642-5
13 K. Y. Chen, Y. W. Chen, Preparation of barium titanate ultrafine
particles from rutile titania by a hydrothermal conversion, Powder
Technology, 141 (2004) 1–2, 69–74, doi:10.1016/j.powtec.2004.
03.002
14 S. G. Kwon, K. Choi, B. I. Kim, Solvothermal synthesis of nano-
sized tetragonal barium titanate powders, Materials Letters, 60
(2006) 7, 979–982, doi:10.1016/j.matlet.2005.10.089
15 M. N. Danchevskaya, Yu. D. Ivakin, S. N. Torbin, G. P. Muravieva,
O. G. Ovchinnikova, Thermovaporous synthesis of complicated
oxides, J. Mater. Sci., 41 (2006) 5, 1385–1390, doi:10.1007/s10853-
006-7411-0
16 Yu. D. Ivakin, M. N. Danchevskaya, G. P. Muravieva, Kinetic model
and mechanism of Y3Al5O12 formation in hydrothermal and thermo-
vaporous synthesis, High Pressure Research, 20 (2001) 1–6, 87–98,
doi:10.1080/08957950108206156
17 A. A. Kholodkova, M. N. Danchevskaya, A. S. Fionov, Investigation
of fine-crystalline barium titanate synthesized in water fluid, Abstr.
of the XIV Inter. Scientific Conf. High-Tech in Chemical Engineer-
ing, Tula, 2012, 547
18 A. A. Kholodkova, M. N. Danchevskaya, A. S. Fionov, Study of
nanocrystalline barium titanate formation in water vapour conditions,
Proc. of the 4th Inter. Conf. NANOCON 2012, Brno, 2012, 66–71
19 W. Cai, C. Fu, Z. Lin, X. Deng, Vanadium doping effect on micro-
structure and dielectric properties of barium titanate ceramics,
Ceramics International, 37 (2011), 3643–3650, doi:10.1016/
j.ceramint.2011.06.024
20 M. M. Vijatovic Petrovic, J. D. Bobic, T. Ramoska, B. D. Stojanovic,
Electrical properties of lanthanum doped barium titanate ceramics,
Materials Characterization, 62 (2011), 1000–1006, doi:10.1016/
j.matchar.2011.07.013
21 P. Kumar, S. Singh, M. Spah, J. K. Juneja, Ch. Pracash, K. K. Raina,
Synthesis and dielectric properties of substituted barium titanate
ceramics, J. of Alloys and Compounds, 489 (2010), 59–63,
doi:10.1016/j.jallcom.2009.08.024
22 S. B. Narang, D. Kaur, Sh. Bahel, Dielectric properties of lanthanum
substituted barium titanate microwave ceramics, Materials Letters,
60 (2006), 3179–3182, doi:10.1016/j.matlet.2006.02.079
23 L. J. Gao, X. L. Liu, J. Q. Zhang, S. Q. Wang, J. F. Chen, Grain-
controlled barium titanate ceramics prepared from high-gravity reac-
tive precipitation process powder, Materials Chemistry and Physics,
88 (2004), 27–31, doi:10.1016/j.matchemphys.2004.03.023
24 T. Lee, I. A. Aksay, Hierarchical structure-ferroelectricity rela-
tionship of barium titanate particles, Crystal Growth & Design, 1
(2001) 5, 401–419, doi:10.1021/cg010012b
25 B. M. Yavorskiy, A. A. Detlaf, Spravochnik po fizike, Nauka, Mos-
cow 1974
A. KHOLODKOVA et al.: PREPARATION AND DIELECTRIC PROPERTIES OF THERMO-VAPOROUS BaTiO3 CERAMICS
Materiali in tehnologije / Materials and technology 49 (2015) 3, 447–451 451

N. V. MURILLO-GUTIÉRREZ et al.: HYBRID SOL-GEL COATINGS DOPED WITH CERIUM ...
HYBRID SOL-GEL COATINGS DOPED WITH CERIUM TO
PROTECT MAGNESIUM ALLOYS FROM CORROSION
HIBRIDNI SOL-GEL-NANOSI, DOPIRANI S CERIJEM, ZA
KOROZIJSKO ZA[^ITO MAGNEZIJEVIH ZLITIN
Noé Verner Murillo-Gutiérrez, Florence Ansart, Jean-Pierre Bonino,
Marie-Jöelle Menu, Marie Gressier
Université de Toulouse UPS-INP-CNRS, Institut Carnot CIRIMAT, 118 Route de Narbonne, 31062 – Toulouse CEDEX 09, France
murillo@chimie.ups-tlse.fr
Prejem rokopisa – received: 2014-05-12; sprejem za objavo – accepted for publication: 2014-09-03
doi:10.17222/mit.2014.077
Hybrid coatings produced via the sol-gel route were deposited onto an Elektron 21 magnesium alloy. The sol consisted of
tetraethyl-orthosilicate (TEOS) and 3-(trimethoxysilyl)propylmethacrylate (MAP) to which corrosion inhibitors were added.
The influence of the cerium concentration on the anti-corrosion properties of the hybrid coating is presented. Furthermore, the
morphology of the organic/inorganic coatings deposited on the magnesium alloy was determined with scanning electron
microscopy (SEM). In parallel, the electrochemical behavior during the immersion in a 0.05 M NaCl corrosive solution was
studied with electrochemical impedance spectroscopy (EIS). It was proven that the hybrid films exhibit a high impedance
modulus during the first hours of the immersion and that an addition of cerium to the sol with a concentration of 0.01 M
considerably increases the durability of the film, delaying its degradation during the immersion. In addition, this project
especially focusses on determining the critical concentration of the cerium salt at which the impedance modulus of the hybrid
coating strongly decreases during the immersion.
Keywords: magnesium, coating, sol-gel, corrosion inhibitor, EIS
Hibridni nanosi, pripravljeni s sol-gel-postopkom, so bili naneseni na magnezijevo zlitino Elektron 21. Osnova je bila
tetraetil-ortosilikat (TEOS) in 3-(trimetoksisilil)propilmetakrilat (MAP), ki so ji bili dodani inhibitorji korozije. Predstavljen je
vpliv koncentracije cerija na protikorozijske lastnosti hibridnega nanosa. Poleg tega je bila dolo~ena morfologija organskih/
neorganskih nanosov na magnezijevo zlitino z vrsti~nim elektronskim mikroskopom (SEM). Vzporedno je bilo preu~evano
elektrokemijsko vedenje med potopitvijo v korozijsko raztopino 0,05 M NaCl, z uporabo elektrokemijske impedan~ne
spektroskopije (EIS). Dokazano je bilo, da izkazujejo hibridni nanosi visok impedan~ni modul med prvimi urami namakanja in
da dodatek cerija osnovi v koncentraciji 0,01 M mo~no pove~a zdr`ljivost nanosa z zadr`anjem njegove degradacije med
namakanjem. Ta projekt je bil usmerjen v dolo~anje kriti~ne koncentracije cerijeve soli, pri kateri se impedan~ni modul
hibridnega nanosa mo~no zmanj{a med namakanjem.
Klju~ne besede: magnezij, nanos, sol-gel, inhibitor korozije, EIS
1 INTRODUCTION
With a density equivalent to 2/3 of that of aluminium,
magnesium and its alloys are interesting weight-saving
materials for the automotive and aeronautics industries.
However, compared to steel and aluminium alloys,
magnesium alloys have a very low corrosion resistance.
In order to prevent this problem, various surface
treatments and coatings have been developed with
different techniques over the last few years1,2. However,
most of these processes make use of chromium (Cr VI)
compounds, nowadays forbidden by international
regulations since these are classified as carcinogen,
mutagenic and reprotoxic compounds. The sol-gel route
is an efficient method to produce "green" coatings and
their anti-corrosion performances have been proven
successful on steel and aluminium alloys.3–5 This project
aims to evaluate the anti-corrosive properties of a hybrid
coating obtained via the sol-gel route and deposited on a
cast Elektron 21 magnesium alloy (El21) and, secondly,
to identify its mechanisms.
2 METHODOLOGY
2.1 Preparation of the materials and coatings
Samples with the dimensions of 40 mm × 20 mm ×
6 mm were obtained by making cuttings from a cast
Elektron 21 (El21) alloy. The chemical composition of
this alloy is shown in Table 1. The samples were first
mechanically polished with abrasive papers with a grit of
up to grade 4000, then with alumina paste (3 μm and
1 μm) and finally they were rinsed with ethanol and dried
under a flux of cold air.
Table 1: Chemical composition of the Elektron 21 cast alloy in mass
fractions, w/%
Tabela 1: Kemijska sestava livne zlitine Elektron 21 v masnih dele`ih,
w/%
Element Nd Gd Zr Zn Other rareearths Mg
w/% 3.1 1.7 1 0.5 < 0.4 Balance
The sols were produced by mixing the starting pre-
cursors consisting of tetraethyl-orthosilicate (TEOS) and
Materiali in tehnologije / Materials and technology 49 (2015) 3, 453–456 453
UDK 669.721.5:621.793:620.197 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(3)453(2015)
3-(trimethoxysilyl)propyl-methacrylate (MAP), deio-
nized water and ethanol with a molar ratio of 11 : 1 : 60 :
80, under constant stirring and at room temperature. In
order to adjust the pH of the sol to 4, nitric acid (HNO3)
was drop-added to the mixture when required. The
production of cerium-doped sols was performed by
adding cerium nitrate (Ce(NO3)3 · 6H2O) at four different
concentrations: (0.005, 0.01, 0.05 and 0.1) mol L–1. The
addition of the cerium salt was carried out by previously
dissolving this compound in the corresponding water
volume of the formulation. After maturing for 24 h, the
sols were deposited on the magnesium El21 substrates
using the dip-coating technique, at a controlled with-
drawal speed of 200 mm min–1. They were then dried at
60 °C for 20 min.
2.2 Characterization techniques
The microstructures of the El21 substrate and the
sol-gel coatings were analyzed with scanning electron
microscopy (SEM) using a JEOL JSM-6510LV micro-
scope, at an operating voltage of 20 kV. Electrochemical
tests of the open circuit potential (Eocp) and electroche-
mical impedance spectroscopy (EIS) were performed in
0.05 mol L–1 of a NaCl corrosive solution at room tem-
perature, using a Bio-Logic SP-150 potentiostat. The
electrochemical cell consisted of a one-chamber three-
electrode cell, the working electrode having an exposed
area of 2 cm², delimited with an insulating tape. The
reference and auxiliary electrodes included a saturated
calomel electrode (SCE) and a platinum-foil electrode,
respectively. The EIS spectra were drawn using the
potentiostatic mode and a frequency ranging from 100
mHz and 10 mHz, with an applied voltage oscillation of
10 mV vs. OCP. For each test, three samples were ana-
lyzed in order to check the reproducibility of the tests.
3 RESULTS AND DISCUSSION
3.1 Morphology of the hybrid coatings
The microstructure of a hybrid sol-gel coating was
first observed with SEM (Figure 1). The surface of the
coating (a) shows a homogeneous surface, with the
presence of some cracks and defects spotted in the
neodymium-rich zones. The origin of these defects may
be attributed to the formation of a galvanic couple
between this intermetallic phase and the alpha phase of
magnesium2, and to the internal stresses of the hybrid
coating that lead to fracture. A cross-sectional observa-
tion of the substrate in the BSE (back-scattered elec-
trons) mode (b) allows the thickness of the hybrid
coating to be measured at around 1 μm.
3.2 Electrochemical characteristics of the hybrid coat-
ings
Firstly, the open-circuit potential of the samples was
recorded during the immersion of the substrates in a
corrosive solution containing 0.05 mol L–1 of NaCl (Fig-
ure 2). All the samples exhibit a similar behavior, except
N. V. MURILLO-GUTIÉRREZ et al.: HYBRID SOL-GEL COATINGS DOPED WITH CERIUM ...
454 Materiali in tehnologije / Materials and technology 49 (2015) 3, 453–456
Figure 2: Evolution of the open-circuit potential (Eocp) of the hybrid
coatings doped with cerium of different concentrations, during the
immersion in a corrosive solution of 0.05 mol L–1
Slika 2: Razvoj potenciala odprtega kroga (Eocp) hibridnega nanosa,
dopiranega s cerijem v razli~nih koncentracijah, med namakanjem v
korozijski raztopini 0,05 mol L–1
Figure 1: SEM images of the hybrid sol-gel coating on the El21
magnesium alloy: a) surface of the coating, b) cross-section of the
substrate in the BSE mode
Slika 1: SEM-posnetka sol-gel hibridnega nanosa na magnezijevi zli-
tini El21: a) povr{ina nanosa, b) pre~ni prerez podlage v BSE-na~inu
the hybrid coatings that were non-doped or doped with
0.005 mol L–1 of Ce. These show a stabilized potential
from the beginning of the immersion, at around –1.63 V
and –1.65 V, respectively, attributed to the insulating
effect of the protective coating. In contrast, the sol-gel
coatings containing 0.01 mol L–1 or higher concentra-
tions, present a behavior similar to that of the bare El21
alloy. This is related to the growth of the passive layer
and corrosion products at the surface of the substrates,
due to the reaction with the electrolyte.6
Secondly, the hybrid coatings were tested with EIS
after 1 h of the immersion in the corrosive solution, right
after the OCP recording. The Bode plots of the EIS
spectra obtained for different protective systems are
shown in Figure 3. It is worth noting that only the hybrid
coating doped with 0.1 mol L–1 presents a behavior
similar to that of the bare El21 substrate. The Bode
phase-angle diagram (a) shows that the last group pre-
sents two time constants, at the high and low frequencies
(10 kHz and 1 Hz, respectively). The first is normally
attributed to the capacitive response of a hybrid sol-gel
coating, which indicates that this film has a physical
barrier effect.7 The second time constant is attributed to
the presence of a porous layer in the corrosion products.8
It is important to observe that the phase angle of a hybrid
coating is the highest when it is doped with 0.01 mol L–1
of cerium. On the other hand, the impedance modulus
(b) obtained at a low frequency (10 mHz) is typically
assigned to the resistance of the electrochemical system,
and so to its corrosion resistance.9 The cerium concentra-
tions lower than 0.1 mol L–1, especially 0.01 mol L–1,
exhibit higher impedance values. Figure 4 presents the
evolution of a hybrid coating doped with 0.01 mol L–1
during its immersion in the corrosive solution. The time
constant attributed to the coating gradually disappears
with the immersion time (Figure 4a, the time constant at
10 kHz), simultaneously with the shift of the time
constant at a low frequency from 1 Hz to 20 Hz. After
48 h of immersion, both curves, representing the bare
El21 substrate and the hybrid coating, are superimposed,
meaning that the coating lost its protective properties.
However, the corrosion resistance of the hybrid film,
depicted by the impedance modulus at a low frequency
(Figure 4b) shows a progressive decrease with the time.
N. V. MURILLO-GUTIÉRREZ et al.: HYBRID SOL-GEL COATINGS DOPED WITH CERIUM ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 453–456 455
Figure 4: Bode plots of the results obtained with EIS for the hybrid
coating doped with 0.01 mol L–1 of cerium, during immersion in the
corrosive solution (0.05 mol L–1): a) phase-angle diagram, b) impe-
dance modulus
Slika 4: Bodejev diagram rezultatov, dobljenih z EIS hibridnega na-
nosa, dopiranega z 0,01 mol L–1 cerija, med namakanjem v korozijski
raztopini (0,05 mol L–1): a) fazni kotni diagram, b) impedan~ni modul
Figure 3: EIS spectra of the El21 magnesium alloy covered with the
hybrid coating, doped with different concentrations of cerium. Results
obtained after 1 hour of immersion in 0.05 mol L–1 of NaCl: a)
phase-angle diagram, b) impedance modulus.
Slika 3: EIS-spektri magnezijeve zlitine El21 s hibridnim nanosom,
dopiranim z razli~nimi koncentracijami cerija. Rezultati, dobljeni po
1 h namakanja v 0,05 mol L–1 NaCl: a) fazni kotni diagram, b) impe-
dan~ni modul.
Here, it is proven that the hybrid coating doped with
0.01 mol L–1 of cerium has the best corrosion resistance.
Higher or lower concentrations of cerium decrease the
coating resistance and capacitance, which leads to a
rapid loss of the protective properties.10 The inhibiting
properties of this element are strongly determined by the
ion concentration inside the hybrid film,11 showing the
existence of the optimum cerium concentration. This is
due to the formation of insoluble compounds such as
CeO2 and Ce(OH)3 12 that temporally block the passage
of the corrosive species through the hybrid coating to the
metallic substrate.
4 CONCLUSION
A hybrid sol-gel coating offers a slight protection to a
magnesium substrate during the first hours of an
immersion. Moreover, an addition of cerium inside the
coating with the optimum concentration of 0.01 mol L–1
leads to the increase of its anti-corrosion properties. This
is due to the corrosion-inhibiting effect of cerium ions
that allows a formation of insoluble compounds, en-
hancing the resistance of the hybrid coating to corrosive
species.
Acknowledgments
The FDA and the OSEO are gratefully acknowledged
for the funding provided for this project. The authors
would like to thank the partners in the FUI CARAIBE
project and its coordinator: the SAFRAN group.
5 REFERENCES
1 J. E. Gray, B. Luan, Journal of Alloys and Compounds, 336 (2002),
88-113, doi:10.1515/bmt-2014-4495
2 R. G. Hu, S. Zhang, J. F. Bu, C. J. Lin, G. L. Song, Progress in Orga-
nic Coatings, 73 (2012), 129-141, doi:10.1016/j.porgcoat.2011.10.
011
3 E. Certhoux, F. Ansart, V. Turq, J. P. Bonino, J. M. Sobrino, J.
Garcia, J. Reby, Progress in Organic Coatings, 76 (2013), 165-172,
doi:10.1016/j.porgcoat.2012.09.002
4 H. B. Lu, Y. Hu, M. H. Gu, S. C. Tang, H. M. Lu, X. K. Meng, Sur-
face and Coatings Technology, 204 (2009), 91-98, doi:10.1016/
j.surfcoat.2009.06.035
5 N. C. Rosero-Navarro, L. Paussa, F. Andreatta, Y. Castro, A. Duran,
M. Aparicio, L. Fedrizzi, Progress in Organic Coatings, 69 (2010),
167-174, doi:10.1016/j.porgcoat.2010.04.013
6 G. Baril, N. Pebere, Corrosion Science, 43 (2001), 471-484,
doi:10.1016/S0010-938X(00)00095-0
7 S. V. Lamaka, G. Knornschild, D. V. Snihirova, M. G. Taryba, M. L.
Zheludkevich, M. G. S. Ferreira, Electrochimica Acta, 55 (2009),
131-141, doi:10.1016/j.electacta.2009.08.018
8 I. A. Kartsonakis, A. C. Balaskas, E. P. Koumoulos, C. A. Charitidis,
G. Kordas, Corrosion Science, 65 (2012), 481-493, doi:10.1016/
j.corsci.2012.08.052
9 G. W. Walter, Corrosion Science, 26 (1986), 27-38, doi:10.1016/
0010-938X(86)90120-4
10 P. S. Correa, C. F. Malfatti, D. S. Azambuja, Progress in Organic
Coatings, 72 (2011), 739-747, doi:10.1016/j.porgcoat.2011.08.005
11 W. Trabelsi, P. Cecilio, M. G. S. Ferreira, M. F. Montemor, Progress
in Organic Coatings, 54 (2005), 276-284, doi:10.1016/j.porgcoat.
2005.07.006
12 M. F. Montemor, M. G. S. Ferreira, Electrochimica Acta, 52 (2007),
7486-7495, doi:10.1016/j.electacta.2006.12.086
N. V. MURILLO-GUTIÉRREZ et al.: HYBRID SOL-GEL COATINGS DOPED WITH CERIUM ...
456 Materiali in tehnologije / Materials and technology 49 (2015) 3, 453–456
H. GÜLER: INFLUENCE OF THE TOOL GEOMETRY AND PROCESS PARAMETERS ...
INFLUENCE OF THE TOOL GEOMETRY AND PROCESS
PARAMETERS ON THE STATIC STRENGTH AND HARDNESS OF
FRICTION-STIR SPOT-WELDED ALUMINIUM-ALLOY SHEETS
VPLIV GEOMETRIJE ORODJA IN PARAMETROV PROCESA NA
STATI^NO TRDNOST IN TRDOTO PRI VRTILNO-TORNEM
TO^KASTEM VARJENJU PLO^EVIN IZ Al-ZLITINE
Hande Güler
Uludag University, Faculty of Engineering, Department of Mechanical Engineering, 16059 Gorukle-Bursa, Turkey
handeguler@uludag.edu.tr
Prejem rokopisa – received: 2014-06-07; sprejem za objavo – accepted for publication: 2014-07-22
doi:10.17222/mit.2014.087
In this study, the effects of the tool geometry and welding parameters on the friction-stir spot-welding properties of AA
5754-H111 were studied. Two different tool-pin geometries were used and tensile shear tests were carried out to compare the
weld strength. Hardness observations were also done. The optimum tool geometry for the mentioned material was determined as
the circular pin tool and the tapered pin tool gave the lowest tensile shear load.
Keywords: Al-alloys, friction-stir spot welding, mechanical properties, hardness
V tej {tudiji so bili preu~evani vplivi geometrije orodja in parametrov varjenja na torno- vrtilno to~kasto varjenje zlitine AA
5754-H111. Dve razli~ni geometriji konice orodja sta bili uporabljeni in opravljeni so bili natezni stri`ni preizkusi za primerjavo
trdnosti zvara. Opravljene so bile tudi meritve trdote. Za ta material je bilo ugotovljeno, da je optimalna geometrija orodja v
obliki okrogle konice, pri sto`~asti konici orodja pa je bila ugotovljena najmanj{a stri`na trdnost.
Klju~ne besede: Al-zlitine, torno-vrtilno to~kasto varjenje, mehanske lastnosti, trdota
1 INTRODUCTION
Friction-stir spot welding (FSSW) is an alternative
method for spot welding of lightweight alloys that was
developed by Mazda Motor Corporation and Kawasaki
Heavy Industries in 2003.1 The FSSW method consists
of three phases: plunging, stirring and retraction. In these
phases, a cylindrical rotating tool plunges at a specific
rate into the overlapping sheets to a predetermined depth.
It is then retracted at a rapid rate either immediately or
after a dwell period. The heating is achieved with the
friction between the tool and the workpieces causing a
plastic deformation of the workpieces. The localized
heating softens the material around the pin and the
forging pressure applied with the tool shoulder results in
the formation of an annular, solid-state bond around the
pin. The retraction of the pin leaves a characteristic exit
hole.2
The significant parameters that determine the
strength of FSSW joints are the rotational speed, the
dwell time and the tool plunge depth and this method has
been successfully applied to aluminum, magnesium,
advanced high-strength steel and polymers.3 There have
been several papers relating to different FSSW process
parameters and different materials.
Külekçi4 investigated the effects of the FSSW para-
meters such as the tool rotation, the dwell time and the
tool-pin height on the tensile shear strength of an EN
AW 5005 aluminium alloy and determined the optimum
parameters. Tozaki et al.5 proposed a new tool which
uses a scroll groove to displace the material in the verti-
cal direction instead of the profiled pin and compared it
with the conventional convex-shoulder tool with a cylin-
drical pin. They found that the new tool exhibited equal,
or superior, results compared with the conventional tool.
FSSW of an Al-alloy 6016-T4 sheet was evaluated by
Yuan et al.6 using the conventional pin tool and off-
center feature tool. Different parameters were investi-
gated to determine the lap-shear separation load and they
found that the tool rotational speed and plunge depth
influenced the shear strength. Badarinarayan et al.7
studied the effects of the shoulder and pin geometry on
the hook formation, and the material flow of a friction-
stir spot-welded 5754-O aluminum alloy was investi-
gated. Külekçi et al.8 explored the hardness distribution
and the tensile shear strength of FSSW welds of the EN
AW 5005 aluminum alloy and compared them with resi-
stance spot welding.
The effects of three tool shapes (threaded-pin tool:
TPT; cylindrical tool: CT; cylindrical tool with projec-
tion: CTP) and the tool penetration depth on the joint
strength of a FSSW commercial AA 5J32 alloy with a
nominal composition in mass fractions (w/%) of
Al-5.54Mg-0.03Si-0.07Fe-0.32Cu-0.03Ti-0.01Zn was
investigated by Choi et al.9 The CTP (cylindrical tool
Materiali in tehnologije / Materials and technology 49 (2015) 3, 457–460 457
UDK 669.715:621.791 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(3)457(2015)
with projection) showed the best mechanical properties
compared to the other tool shapes.
Karthikeyan and Balasubramanian10 aimed at opti-
mizing the welding parameters to attain the maximum
lap-shear tensile strength of a friction-stir spot-welded
AA2024 aluminum alloy. For this purpose, an empirical
relationship was developed to predict the tensile shear-
fracture load.
In this work, the FSSW process was used with the
objective to investigate the effects of different welding
parameters and tool geometry on the mechanical perfor-
mance of thick AA 5754-H111 1 mm. This alloy has
been widely used in shipbuilding, the vehicle, chemical
and nuclear industries.11
However, a literature review indicated that there is
only one paper about the thick AA 5754-H111 1 mm.12
Güler12 studied the FSSW process of the thick AA
5754-H111 1 mm using only circular pinned tool. In this
context, the effects of the tool design and process para-
meters on the weld strength and hardness distribution
were studied and the effects of FSSW process were com-
pared using the tools with circular and tapered pins.
2 EXPERIMENTAL DETAILS
Commercially available aluminium-alloy plates (AA
5754-H111) with a thickness of 1 mm were used to
fabricate the joints. The chemical composition of the
material is presented in Table 1. The tensile strength, the
yield strength and the elongation are 207 MPa, 117 MPa
and 21 %, respectively. The samples for the tensile shear
test were machined out with the dimensions of 105 mm
× 45 mm × 1 mm according to the ISO 14273 standard.
The specimens were lap positioned with a 45 mm × 45
mm overlap area (Figure 1). The welded samples were
loaded on a universal testing machine with a constant
crosshead of 5 mm/min and with a load capacity of 250
kN.
FSSW was performed using two different tools made
of hot-work tool steel having different pin profiles: a
circular one, designated as FSSW-C and a tapered one,
designated as FSSW-T. The tools with different pin
geometries were selected to ensure simple manufacturing
with a small amount of tool wear. As shown in Figure 2,
each tool had a shoulder with a diameter of 15 mm and a
pin length of 1.7 mm. The circular pin had a diameter of
5 mm and the tapered pin had the diameters of 5 mm at
the bottom and 7 mm at the pin shoulder.
The prepared samples were joined by FSSW using
different tool rotational speeds and dwell times. FSSW
was performed at three tool rotational speeds, ranging
from 500 r/min to 1500 r/min, while the plunge depth of
the tool pin was 1.8 mm and the dwell time ranged from
6 s to 21 s. Considering each parameter, five joints at
different tool rotational speeds and dwell times were
produced and tested. Some samples of the FSSW joints
are shown in Figure 3.
3 RESULTS AND DISCUSSION
3.1 Tensile shear strength
Figure 4 shows a comparison of the shear load ver-
sus the dwell time obtained during the shear tensile
testing at a constant tool rotational speed. According to
Figure 4a, both the dwell time and the shear load of the
FSSW-C welds increased, while the shear load of the
FSSW-T welds remained almost constant. When using
the tapered tool with a combination of a tool rotational
speed 500 r/min and dwell time 6 s, the shear load
increased by 55 %. Besides, at the other dwell times, the
shear loads of the FSSW-C welds were higher compared
to those of the FSSW-T welds.
H. GÜLER: INFLUENCE OF THE TOOL GEOMETRY AND PROCESS PARAMETERS ...
458 Materiali in tehnologije / Materials and technology 49 (2015) 3, 457–460
Table 1: Chemical composition of the investigated material in mass fractions, w/%
Tabela 1: Kemijska sestava preiskovanega materiala v masnih dele`ih, w/%
Si Fe Mn Mg Cu Ti Cr Zn Al
Chemical composition 0.071 0.251 0.19 2.834 0.022 0.003 0.024 0.049 Balance
Figure 1: Lap-shear specimen
Slika 1: Vzorec s prekrivanjem za strig
Figure 2: Two types of tool used for FSSW: FSSW-T (tapered pin,
left) and FSSW-C (circular pin, right)
Slika 2: Dve vrsti orodja, uporabljeni za FSSW: FSSW-T (konica v
obliki prisekanega sto`ca, levo) in FSSW-C (okrogla konica, desno)
Figure 4b shows the shear-load values at a tool rota-
tional speed 1000 r/min. According to the figure, the
average shear loads of the FSSW-C welds are higher
than the ones of the FSSW-T welds. With the combina-
tion of the tool rotational speed 1000 r/min and dwell
time 11 s, the maximum increase was observed when
using the circular tool which increased the shear load by
approximately 43 %.
The tensile-shear-test results for 1500 r/min are given
in Figure 4c. An increase in the dwell time from 6 s to
21 s increased the tensile shear load of the FSSW-T
welds by 26 %, while the tensile shear load of the
FSSW-C welds remained almost constant. On the other
hand, the shear-load values were higher when a sample
was welded with the FSSW-C tool compared with the
FSSW-T tool.
According to all the tensile shear results, the speci-
mens welded with the FSSW-C tool exhibited an in-
crease in the shear load. The results of the previous
investigations7,13,14 stated that the tool-pin geometry
significantly affects the hook geometry and thereby the
shear-load capacity of FSSWs. An increase in the tensile
shear load may be explained with the microstructural
changes due to the heat generation at a stirring location.4
When using a FSSW-T tool, the materials may be stirred
severely and extensively. When the amount of the
material stirred is large, the size of a bonded region also
becomes large, resulting in a higher separation load.14
3.2 Microhardness
Hardness measurements were performed on a Vickers
microhardness tester focusing at the medium-thick area
of the upper sheet with a load of 100 g, dwell time of 10
s and spacing of 1 mm. The microhardness-distribution
examples of the FSSW-C and FSSW-T welded joints are
H. GÜLER: INFLUENCE OF THE TOOL GEOMETRY AND PROCESS PARAMETERS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 457–460 459
Figure 4: Shear load as a function of dwell time for both FSSW-C and
FSSW-T welds at a tool rotational speed: a) 500 r/min, b) 1000 r/min,
c) 1500 r/min
Slika 4: Stri`na obremenitev v odvisnosti od ~asa zadr`evanja za obe
varjenji: FSSW-C in FSSW-T pri hitrosti orodja: a) 500 r/min, b) 1000
r/min, c) 1500 r/min
Figure 3: Samples of friction-stir spot-welded joints
Slika 3: Vzorci vrtilno-torno to~kasto zvarjenih spojev
given in all the hardness profiles, and a higher Vickers
microhardness was observed close to the keyhole
because of a higher plastic deformation, which also
caused very fine, dynamic recrystallized grains. On the
other hand, in the region close to the keyhole, the
FSSW-T weld had a slightly higher hardness than the
FSSW-C weld. This situation occurred because of the
presence of a finer grain structure observed in the stir
zone of the FSSW-T weld.13 In addition, on the basis of
Figures 5a, 5b and 5c, it was concluded that the hard-
ness distribution was somewhat scattered without a
noticeable trend. There was no obvious influence of the
process parameters on the hardness distribution.4
4 CONCLUSIONS
In this work, the mechanical properties of the AA
5754-H111 aluminum-alloy material were investigated
during the FSSW process. The welding process was
performed using two different types of designed tools
(FSSW-C and FSSW-T), compared to each other. After
analyzing the experimental results, the following con-
clusions can be drawn:
1. At the tool rotational speed 500 r/min, the dwell
time is an influential factor in determining the tensile
shear load of the FSSW-C welds, while the tensile shear
loads remain almost constant at the other tool rotational
speed. When using the FSSW-T tool at the rotational
speed 1500 r/min, the tensile shear load increases if the
dwell time increases.
2. Higher shear-load values are found for the sample
welded with the FSSW-C tool, compared to the FSSW-T
tool, because of a better stirring procedure.
3. There are no noticeable effects of the tool geome-
tries and welding parameters affecting the hardness of
the mentioned material.
4. As future work, different failure modes occurring
during the lap-shear test could be studied to determine
their effects on the tensile shear-load capacity.
5 REFERENCES
1 W. Ratanathavorn, Hybrid Joining of Aluminum to Thermoplastics
with Friction Stir Welding, Master of Science Thesis, Department of
Materials Science and Engineering, KTH – Royal Institute of Tech-
nology [cited 2014-07-06], available from http://kth.diva-portal.org/
smash/get/diva2:515154/FULLTEXT01.pdf
2 S. Lathabai, M. J. Painter, G. M. D. Cantin, V. K. Tyagi, Scripta Ma-
terialia, 55 (2006), 899–902, doi:10.1016/j.scriptamat.2006.07.046
3 X. Song, L. Ke, L. Xing, F. Liu, C. Huang, The International Journal
of Advanced Manufacturing Technology, 71 (2014) 9–12, 2003–
2010, doi:10.1007/s00170-014-5632-y
4 M. K. Külekçi, Archives of Metallurgy and Materials, 59 (2014) 1,
221–224, doi:10.2478/amm-2014-0035
5 Y. Tozaki, Y. Uematsu, K. Tokaji, Journal of Materials Processing
Technology, 210 (2010), 844–851, doi:10.1016/j.jmatprotec.2010.
01.015
6 W. Yuan, R. S. Mishra, S. Webb, Y. L. Chen, B. Carlson, D. R. Her-
ling, G. J. Grant, Journal of Materials Processing Technology, 211
(2011), 972–977, doi:10.1016/j.jmatprotec.2010.12.014
7 H. Badarinarayan, Y. Shi, X. Li, K. Okamoto, International Journal
of Machine Tools & Manufacture, 49 (2009), 814–823, doi:10.1016/
j.ijmachtools.2009.06.001
8 M. K. Külekci, U. Esme, O. Er, Mater. Tehnol., 45 (2011) 5,
395–399
9 D. Choi, B. Ahn, C. Lee,Y. Yeon, K. Song, S. Jung, Materials Tran-
sactions, 51 (2010) 5, 1028–1032, doi:10.2320/matertrans.
M2009405
10 R. Karthikeyan, V. Balasubramanian, Int. J. Adv. Manuf. Technol.,
51 (2010), 173–183, doi:10.1007/s00170-010-2618-2
11 T. S. Mahmoud, T. A. Khalifa, JMEPEG, 23 (2014), 898–905,
doi:10.1007/s11665-013-0828-0
12 H. Güler, JOM, 66 (2014) 10, 2156–2160, doi:10.1007/s11837-
014-1117-6
13 H. Badarinarayan, Q. Yang, S. Zhu, International Journal of Machine
Tools & Manufacture, 49 (2009) 2, 142–148, doi:10.1016/
j.ijmachtools.2008.09.004
14 N. Pathak, K. Bandyopadhyay, M. Sarangi, S. K. Panda, Journal of
Materials Engineering and Performance, 22 (2013), 131–144,
doi:10.1007/s11665-012-0244-x
H. GÜLER: INFLUENCE OF THE TOOL GEOMETRY AND PROCESS PARAMETERS ...
460 Materiali in tehnologije / Materials and technology 49 (2015) 3, 457–460
Figure 5: Hardness distribution for both FSSW-C and FSSW-T welds
with a combination of: a) 500 r/min and 21 s, b) 1000 r/min and 11 s,
c) 1500 r/min and 21 s
Slika 5: Razporeditev trdote za oba zvara FSSW-C in FSSW-T pri
kombinaciji: a) 500 r/min in zadr`anju 21 s, b) 1000 r/min in 11 s, c)
1500 r/min in 21 s
K. M. SHOJAEI et al.: THE STABILIZATION OF NANO SILVER ON POLYESTER FILAMENT FOR ...
THE STABILIZATION OF NANO SILVER ON POLYESTER
FILAMENT FOR A MACHINE-MADE CARPET
STABILIZACIJA NANODELCEV SREBRA NA POLIESTRSKEM
VLAKNU ZA STROJNO IZDELAVO PREPROG
Khashayar Mohajer Shojaei, Ali Farrahi, Hossein Farrahi, Ahmad Farrahi
Farrokh Sepehr Kashan Company (Farrahi Carpet), Kashan, Iran
Khashayar045@yahoo.com
Prejem rokopisa – received: 2014-07-20; sprejem za objavo – accepted for publication: 2014-09-02
doi:10.17222/mit.2014.113
Nowadays, polyester filament yarns with permanent anti-bacterial characteristics are known as an innovative yarn in the textile
industry for machine-made carpets and garments. Different methods, such as chemical modification, entrapment and encapsu-
lation, have been applied to stabilize the silver nanoparticles on polyester filament yarn for use in machine-made carpets. In this
article we modified the nano-silver particles by a chemical reaction in order to produce a polyester filament yarn with permanent
anti-bacterial characteristics using a spraying method. The anti-bacterial tests were carried out on the nano-silver-coated
machine-made carpet according to ATCC 27853, 25923 and 25922 some 72 h before and after the washing process. The results
showed that the nano-silver-coated polyester machine-made carpet has a permanent anti-bacterial characteristic.
Keywords: anti bacterial characteristic, polyester filament yarn, machine-made carpet, DLS method, washing processes,
bacteria, surfactant, Mac Farlen
Dandanes je preja iz poliestrskih vlaken s trajno protibakterijsko odpornostjo poznana kot inovativna preja v tekstilni industriji
za strojno izdelavo preprog in obla~il. Razli~ne metode, kot so kemijska obdelava, ujetje in enkapsulacija, so bile uporabljene za
stabilizacijo nanodelcev srebra na vlaknih poliestrske preje za strojno izdelavo preprog. ^lanek predstavlja modifikacijo
nanodelcev srebra s kemijsko reakcijo z napr{evanjem za izdelavo vlakna poliestrske preje s stalno protibakterijsko odpornostjo.
Protibakterijski preizkusi so bili opravljeni na strojno izdelani preprogi, prekriti z nanodelci, skladno z ATCC 27853, 25923 in
25922, 72 h pred pranjem in po njem. Rezultati so pokazali, da ima strojno izdelana preproga iz poliestra, prekritega z nanodelci
srebra, trajno protibakterijsko odpornost.
Klju~ne besede: protibakterijske lastnosti, preja iz poliestrskih vlaken, strojno izdelana preproga, DLS-metoda, postopek pranja,
bakterija, povr{inska aktivnost, Mac Farlen
1 INTRODUCTION
Microbial organisms, bacteria and micro-organisms
are the main reasons for sickness, infections and bad
odours, etc. In fact, increasing death rates in many
underdeveloped countries, such as many African coun-
tries, led to the legislation of global and social policies in
order to overcome this challenge. Such problems and
needs have led to a resurgence in the use of silver- and
copper-based antiseptics that may be linked to broad-
spectrum activity and a far lower propensity to induce
microbial resistance than antibiotics1.
The anti-bacterial characteristics of silver and silver
salts have been observed since antiquity2. Silver is
currently used to control bacterial growth in a variety of
applications, including dental work, catheters and burn
wounds3,4.
In fact, it is well known that Ag ions, Ag-based com-
pounds, copper and brass compounds have a strong
biotical effect on many bacteria species, such as E. coli,
P. aeruginosa and S. aureus5.
A lot of information is available about the practical
use of nanoparticles for food safety and hygiene, the
disinfection of water in swimming pools and hospitals,
wound healing, air disinfection and surface sanitation6–8.
It is believed that the germicidal property of metals,
especially heavy metals, is due to oligodynamic effect in
which the metal and metal compounds, when introduced
into the interior of bacterial cells, have the ability to
change and then kill them in a specific way. Copper and
silver are the most studied metals for oligodynamic ac-
tion9. The data from silver suggest that its ions denature
the proteins in the bacterial cells by binding to the reac-
tive groups, resulting in their inactivation10.
Different factors influence the efficiency of silver-
and copper-based compounds, such as the particle size
and particle size distribution. In these cases, a reduction
of the particle size of silver and copper nanoparticles is a
reliable solution to improve their efficiency and biocom-
patibility. In this field, nanotechnology has a direct effect
on the elimination of particle size limitations and
changing the world outlook regarding science10,11.
In this study we investigated the stabilization of nano
silver on polyester filament yarn in order to produce a
machine-made carpet with permanent antibacterial cha-
racteristics.
Materiali in tehnologije / Materials and technology 49 (2015) 3, 461–464 461
UDK 677 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(3)461(2015)
2 MATERIALS
A nano-silver colloid was obtained from US Re-
search Nano Material Company. The polyester filament
yarns as a pile, warp and weft in a machine-made carpet
were provided by the Farrokh Sepehr Kashan Textie
Company. Gluteraldehyde (cross linking agent) and
-amino propyl tri-etoxy silane were provided by the
Sigma-Aldrich Company on a laboratory scale. An am-
photeric surfactant based on amino betaine was provided
by the Carp Company.
3 METHODS
Nano-silver colloids, including a cross-linking agent
and amphoteric surfactant in different concentrations,
were applied to the polyester filament yarn and the back
of a machine-made carpet using a spraying method. The
nano-silver-coated polyester filament yarn and polyester
machine-made carpet were dried and cured in the stenter
at 130 °C for 6 min.
4 ANALYSIS
To predict the stability of the anti-bacterial characte-
ristics, the nano-silver back-coated polyester machine-
made carpets were analysed according to ATCC 27853,
25923 and 25922 some 72 h before and after a washing
process in the Pasteur Institute of Iran.
The anti-bacterial characteristic of the nano-silver-
coated machine-made carpet were measured based on the
growth of different bacteria, such as E. coli, P. aerugi-
nosa and S. aureus with respect to the reference sample.
The stability of the nano-silver colloids and their par-
ticle size distribution were analysed using the DLS
(dynamic light scattering) method at 25 °C (Malvern seri
nano (zeta sizer) model DLS).
The FTIR analysis of the obtained solution from
washing the nano-silver-coated polyester filament was
carried out using Fourier-transform spectroscopy in the
range 300 cm–1 to 4000 cm–1.
The distribution of nano-silver on the back of the
polyester-filament machine-made carpet was analysed
using the FESEM method in 1.89 KX (Philips model
FESEM).
5 PREPARATION METHOD
The Mac Farlan solutions of different bacteria such
as P. aeruginosa (ATCC 27853), S. aureus (ATCC
25923) and E. coli (ATCC 25922) at a concentration of
1.5 × 108 CFU/mL were prepared in the first stage.
During the next stage, the nano-silver-coated ma-
chine-made carpet and the reference machine-made
carpet were put in contact with the Mac Farlen solution
of different bacteria for 24 h. After the cultivation and
incubation process for 72 h at 37 °C, the growth of the
bacteria were analysed.
6 RESULTS AND DISCUSSION
6.1 Particle size distribution of nano-silver colloid
The particle size distribution of the nano-silver
colloid was measured using the DLS method. The results
showed that the average size of the nano-silver colloid
and the Pdi constant were about 49.08 nm and 0.408 nm,
respectively. Figure 1 shows the size distribution of the
nano-silver.
6.2 Stability of the nano-silver colloid
The stability of the nano-silver colloid was measured
using the DLS method according to the zeta-potential
value. The results showed that the nano-silver colloid has
K. M. SHOJAEI et al.: THE STABILIZATION OF NANO SILVER ON POLYESTER FILAMENT FOR ...
462 Materiali in tehnologije / Materials and technology 49 (2015) 3, 461–464
Figure 1: Size distribution of nano silver colloid
Slika 1: Razporeditev velikosti nanodelcev srebra v koloidu
a suitable stability under normal conditions. In Figure 2,
the zeta-potential distribution of the nano-silver is
shown. According to Figure 2, the zeta-potential for a
nano-silver colloid is about –17.2 MV. In addition, the
nano-silver particle colloid has a suitable stability
against sedimentation.
6.3 Distribution of nano-silver on the back of polyester
machine-made carpet
The distribution of the nano silver on the back of the
machine-made carpet is shown in a FESEM micrograph
in 1.89 KX. According to Figure 3, it can be concluded
that the anti-bacterial characteristics of the polyester-fila-
ment machine-made carpet is due to the nano-silver
material that was applied on the polyester filament yarn
and the back of the machine-made carpet.
6.4 Anti-bacterial test
The nano-silver back-coated machine-made carpets
before and after washing were tested according to ATCC
27853, 25923 and 25922 in 72 h. In Tables 1 and 2, the
concentration of bacteria before and after the washing
processes are presented.
The results showed that the nano-silver-coated poly-
ester-filament machine-made carpet before and after
K. M. SHOJAEI et al.: THE STABILIZATION OF NANO SILVER ON POLYESTER FILAMENT FOR ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 461–464 463
Figure 2: Zeta-potential distribution of nano-silver colloid
Slika 2: Zeta potencial razporeditve nanodelcev srebra v koloidu
Figure 3: SEM micrographs of samples: a) polyester-filament yarn, b)
nano-silver-coated polyester-filament yarn
Slika 3: SEM-posnetka vzorcev: a) vlakna poliestrske preje, b) z
nanodelci srebra prekrita vlakna poliestrske preje
Table 1: Anti-bacterial characteristics of samples before washing
Tabela 1: Protibakterijske lastnosti vzorcev pred pranjem
Name of test bacteria Bacteria concentration(CFU/mL) The growth of bacteria after 72 h (CFU/mL)
E. coli (ATCC 25922) 1.5 × 108
Sample control 1.5 × 106
Anti bacterial sample 8.5 × 10
P. aeruginosa (ATCC 27853) 1.5 × 108
Sample control 1.5 × 106
Anti bacterial sample 7.5 × 10
S. aureus (ATCC 25923) 1.5 × 108
Sample control 1.5 × 106
Anti bacterial sample 7.5 × 10
Table 2: Anti-bacterial characteristics of samples after washing
Tabela 2: Protibakterijske lastnosti vzorcev po pranju
Name of test bacteria Bacteria concentration(CFU/mL) The growth of bacteria after 72 h (CFU/mL)
E. coli (ATCC 25922) 1.5 × 108
Sample control 1.5 × 106
Anti bacterial sample 8.5 × 10
P. aeruginosa (ATCC 27853) 1.5 × 108
Sample control 1.5 × 106
Anti bacterial sample 7.5 × 10
S. aureus (ATCC 25923) 1.5 × 108
Sample control 1.5 × 106
Anti bacterial sample 7.5 × 10
washing has suitable anti-bacterial characteristics against
E. coli, P. aeruginosa and S. aureus (99.9 % anti-bacte-
rial characteristic).
6.5 FTIR analysis of nano-silver-coated polyester-fila-
ment machine-made carpet
In order to demonstrate the permanence of the
anti-bacterial characteristic in a polyester-filament ma-
chine-made carpet, FTIR analyses carried out on a
nano-silver colloid and the obtained solution from wash-
ing the nano-silver-coated polyester-filament machine-
made carpet. Figure 3 shows the FTIR analysis of the
nano-silver colloid and the obtained solution from wash-
ing the nano-silver-coated polyester-filament machine-
made carpet (Figure 4).
According to Figure 4a, the sharp peaks at 1634 cm–1
and 1384 cm–1 indicate the formation of asymmetric and
symmetric stretching modes of metal carbonyl groups.
This is due to the stabilization of the silver nano particles
by the –COO- group of amino betaine (as an amphoteric
surfactant). The peaks at 3432 cm–1 and 2922 cm–1 are
related to the C-H stretching bond of the propyl and the
amine groups of the -amino propyl tri-etoxy silane12.
According to Figure 4b, the sharp peaks at 3444
cm–1 and 2934 cm–1 are related to the O-H and C-H
groups of the soap solution, which are used for washing
the nano-silver-coated polyester-filament machine-made
carpet.
Moreover, the lack of any peak at 1634 cm–1 and
1384 cm–1 indicates that the washing process had no
effect on removing the silver nanoparticles from the
anti-bacterial polyester-filament machine-made carpet.
7 CONCLUSION
The results showed that the nano-silver-coated poly-
ester-filament machine-made carpet has a permanent
anti-bacterial characteristic. This is due to the modifica-
tion process carried out on the nano-silver before
applying it to polyester-filament yarn and the back of the
machine-made carpet. Having permanent anti-bacterial
characteristics in the machine-made carpet make it more
suitable for use in crowded places by removing the bad
smells in a machine-made carpet due to the direct con-
tact of people with machine-made carpet’s surface.
Acknowledgement
This work was done in the Farrokh Sepehr Kashan
Company. Mr. Ali Farrahi and Mr. Ahmad Farrahi are
grateful for the financial support.
8 REFERENCES
1 S. A. Jones, P. G. Bowler, M. Walker, D. Parsons, Wound Repair and
Regeneration, 12 (2004) 3, 288–294, doi:10.1111/j.1067-1927.
2004.012304.x
2 S. Silver, L. T. Phung, Annu. Rev. Microbiol., 50 (1996), 753–789,
doi:10.1146/annurev.micro.50.1.753
3 M. Catauro, M. G. Raucci, F. De Gaetano, A. Marotta, J. Mater. Sci.
Mater. Med., 15 (2004) 7, 831–837, doi:10.1023/B:JMSM.
0000032825.51052.00
4 J. H. Crabtree, R. J. Burchette, R. A. Siddiqi, I. T. Huen, L. L. Had-
nott, A. Fishman, Perit. Dial. Int., 23 (2003) 4, 368–374
5 G. Zhao, S. E. Stevens Jr., Biometals, 11 (1998) 1, 27–32,
doi:10.1023/A:1009253223055
6 C. E. Santo, N. Taudte, D. H. Nies, G. Grass, Appl. Environ. Micro-
biol., 74 (2008) 4, 977–986, doi:10.1128/AEM.01938-07
7 S. A. Wilks, H. Michels, C. W. Keevil, International Journal of Food
Microbiology, 105 (2005) 3, 445–454, doi:10.1016/j.ijfoodmicro.
2005.04.021
8 D. S. Blanc, P. Carrara, G. Zanetti, P. Francioli, J. Hosp. Infect., 60
(2005) 1, 69–72, doi:10.1016/j.jhin.2004.10.016
9 A. J. Varkey, Scientific Research and Essays, 5 (2010) 24,
3834–3839
10 M. Yamanaka, K. Hara, J. Kudo, Appl. Environ. Microbiol., 71
(2005) 11, 7589–7593, doi:10.1128/AEM.71.11.7589-7593.2005
11 J. S. Kim, E. Kuk, K. N. Yu, J. H. Kim, S. J. Park, H. J. Lee, S. H.
Kim, Y. K. Park, Y. H. Park, C. Y. Hwang, Y. K. Kim, Y. S. Lee, D.
H. Jeong, M. H. Cho, Nanomedicine: Nanotechnology, Biology, and
Medicine, 3 (2007) 1, 95–101, doi:10.1016/j.nano.2006.12.001
12 R. Augustine, K. Rajarathinam, Int. J. Nano Dim., 2 (2012) 3,
205–212
K. M. SHOJAEI et al.: THE STABILIZATION OF NANO SILVER ON POLYESTER FILAMENT FOR ...
464 Materiali in tehnologije / Materials and technology 49 (2015) 3, 461–464
Figure 4: FTIR analysis a) nano silver b) solution from washing
Slika 4: FTIR-analiza, a) nanodelci srebra, b) raztopina pri pranju
J. BEÒO et al.: EVALUATION OF THE THERMAL RESISTANCE OF SELECTED BENTONITE BINDERS
EVALUATION OF THE THERMAL RESISTANCE OF SELECTED
BENTONITE BINDERS
OCENA TOPLOTNE UPORNOSTI IZBRANIH BENTONITNIH
VEZIV
Jaroslav Beòo1, Jiøina Vontorová2, Vlastimil Matìjka3, Karel Gál1
1Department of Metallurgy and Foundry Engineering, Faculty of Metallurgy and Material Engineering, V[B-Technical University of Ostrava,
17. listopadu 15/2172, 708 33 Ostrava – Poruba, Czech Republic
2Department of Chemistry, Faculty of Metallurgy and Material Engineering, V[B-Technical University of Ostrava, 17. listopadu 15/2172,
708 33 Ostrava – Poruba, Czech Republic
3Nanotechnology Centre, V[B-Technical University of Ostrava, 17. listopadu 15/2172, 708 33 Ostrava – Poruba, Czech Republic
jaroslav.beno@vsb.cz
Prejem rokopisa – received: 2014-07-29; sprejem za objavo – accepted for publication: 2014-09-03
doi:10.17222/mit.2014.126
Bentonite is one of the most widely used clays associated with various applications. In the case of foundry technology, bentonite
is primarily used as a binder for the mold manufacture. The thermal resistance of bentonite binders, also called the thermal
stability, is a natural property of clay minerals, depending on the source, the mineralogical and chemical composition of clay
and it is also closely connected to the bentonite structure (various interlayer ions, the level of ion substitution of
montmorillonite). Generally, there are various methods for evaluating this property. This contribution describes various methods
of determining the bentonite thermal stability based on the evaluation of the technological parameters of bentonite molding
mixtures and their comparison. These methods were chosen on the basis of a background research and practical experiences. For
the experiments the bentonites commonly used in the foundries of the Czech and Slovak region were selected.
Keywords: bentonites, clay minerals, thermal stability, dehydration, dehydroxylation
Bentonit je ena najpogosteje uporabljanih glin, povezanih z razli~no uporabo. V livarstvu se bentonit uporablja predvsem kot
vezivo pri izdelavi form. Toplotno upornost bentonitnih veziv imenujemo tudi toplotna stabilnost in je naravna lastnost
mineralov glin, odvisna je od izvora, mineralo{ke in kemijske sestave gline in je tudi tesno povezana s strukturo bentonita
(razli~ni ioni med plastmi, nivo nadomestila ionov v montmorilonitu). Obstaja ve~ metod za oceno te lastnosti. Ta prispevek
opisuje razli~ne metode dolo~anja toplotne stabilnosti bentonita, ki temeljijo na oceni tehnolo{kih parametrov bentonitnih
me{anic za forme in njihovo primerjavo. Te metode so bile izbrane na podlagi raziskav ozadja in prakti~nih izku{enj. Za
preizkuse so bili izbrani bentoniti, ki se uporabljajo v livarnah na ^e{kem in Slova{kem.
Klju~ne besede: bentoniti, minerali glin, toplotna stabilnost, dehidracija, dehidroksilacija
1 INTRODUCTION
Bentonite as a natural and abundant soil is being
widely used in a large range of industrial applications. A
high CEC, the swelling ability and a high surface area of
bentonite predetermine it for utilization in ceramics,
cosmetics, nanocomposites, environmental protection,
waste-water treatment, nuclear-waste deposits or
manufacture of foundry molds and cores.1–6
As a rule, bentonites consist of the montmorillonite
clay with the amount of the montmorillonite mineral
higher than mass fractions w = 70 % to 75 %, which
means that it contains up to 30 % of other minerals,
above all aluminosilicates and also micas and carbo-
nates. Individual bentonite localities also differ with their
genesis. In principle, the bentonites do not differ in the
chemical composition, but their behaviors are quite
different (because of their mineral composition, physical
characteristics, etc.). Nowadays the required properties
are achieved by mixing the bentonites from different
localities.
The thermostability of bentonite is connected with
the temperature of the clay dehydroxylation, i.e., with
liberating the OH– groups from the octahedral network in
the form of H2O (g). The residual oxygen remains in the
structure. At the same time the binding properties are
gradually lost and the "burnt-out bentonite" is formed.
The whole process is endothermic and it can be well mo-
nitored with a thermal analysis (DTA). The thermostabi-
lity and its testing methods were studied by Jelínek et al.7
In the bentonites of a mean quality the endothermic reac-
tions are present in the temperature region from 450 °C
to 550 °C. The bentonites with a high thermostability (a
loss of crystallic water, dehydroxylation, occurs in the
range from 700 °C up to 750 °C) often even have two
peaks, at 500 °C and 700 °C. This state is explained as
follows:
• a substitution of ions in montmorillonite octahedrons
and tetrahedrons (the influence of Fe)8
• defects in the lattice (vacancies)
• a difference in the size of montmorillonite particles9.
The influence of Fe on dehydroxylation has not yet
been unambiguously explained. A low concentration of
Fe (< 8 %)7 results in a low dehydroxylation tempera-
ture; Fe-rich bentonites have higher dehydroxylation
Materiali in tehnologije / Materials and technology 49 (2015) 3, 465–469 465
UDK 621.742.48:621.742.43 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(3)465(2015)
temperatures10. On the contrary, Grefhorst11 and Kaplun12
give an opinion that with the increasing Fe2O3 content
the dehydroxylation temperature is falling (in the range
of 2–14 %) and in the presence of black coal the fall is
even more intense.
Due to the thermal exposure of the mold buildup of
bentonite during metal casting, the amount of the active
bentonite decreases and the non-reactive (burnt-out) ben-
tonite is formed. For each cycle of the mold preparation
the used mixture is reactivated with an addition of fresh
bentonite in order to maintain the portion of active ben-
tonite (replacing the burnt-out bentonite in the sand).
The burnt-out bentonite amount of the foundry sand is
affected by the thermal stability of bentonite. The ther-
mal stability of bentonite is usually determined as the
temperature of dehydroxylation or, more precisely, the
temperature of the crystal-lattice destruction and/or the
measure of the loss of its plasticity.
The objectives of this study are: i) an evaluation of
the basic physical and chemical properties of the bento-
nites from the Czech and Slovak region commonly used
in the foundry industry, and a comparison of the parame-
ters of these bentonites with natural Na+ – the bentonite
of the Wyoming type; ii) an evaluation of the impact of
the bentonite chemical composition on its thermal stabi-
lity with respect to its utilization in the foundry industry
(with an emphasis on Fe and Mg contents); iii) a deter-
mination of the thermal stability of bentonite samples
evaluated with the methods based on an evaluation of the
technological parameters of the bentonite molding mix-
tures; iv) a comparison of individual methods and an
assessment of the optimum method for determining the
bentonite thermal stability.
2 MATERIALS AND METHODS
Four commonly used binders supplied as soda-acti-
vated foundry bentonites and one natural sodium bento-
nite of the Wyoming type were used for the experiments
within this research.
The studied bentonites comprise the following groups
of bentonites: i) two bentonites mined and produced in
the Czech Republic (assigned as Sa and K, from the
West Bohemian region), ii) two bentonites from the
Central Slovakia region, assigned as B and S; and iii) a
natural sodium bentonite of the Wyoming type (USA),
assigned as P, was used as the standard material.
The following general parameters (Table 1) com-
monly used for the characterization of bentonites were
determined: a) the moisture under the temperature of 105
°C up to the constant weight; b) pH and conductivity of
water suspension (a 1 : 10 solid-liquid ratio); c) the loss
of ignition (LOI) of dried samples (105 °C up to the con-
stant weight) at 900 °C/2 h.
The chemical compositions of the studied samples
were determined using energy dispersive fluorescence
spectrometer (XRFS) SPECTRO XEPOS (SPECTRO
Analytical Instruments GmbH) equipped with a 50 W Pd
X-ray tube. The samples for the analysis were prepared
in the form of pressed tablets (wax was used as a binder)
for this measurement.
The thermal stabilities of the selected bentonite
samples were determined as the ratio of the values of the
selected technological properties (splitting strength, wet
tensile strength, determination of the methylene-blue
consumption – an active clay test) before and after the
annealing of the bentonite molding mixture.
The samples of the bentonite molding mixture were
prepared with a 5 min homogenization of the mixture of
the studied bentonite with silica sand, in the constant
weight ratio of 8 : 100 and with an appropriate amount
of water ensuring a constant compactibility of (45 ± 3) %
using an MK 00 sand mill. The prepared mixtures were
processed into standard cylinders (Ø 50, a height of 50
mm) to obtain the samples for the determination of the
technological parameters.
The splitting strengths were measured using a
WADAP testing machine of the LRU-1 type, while the
wet tensile strength was measured using a +GF+ testing
machine of the SPNF type.
3 RESULTS AND DISCUSSION
The basic bentonite-binder parameters are summar-
ized in Table 1. The natural moisture of the samples
determined at 105 °C (as the loss of weight) varied a lot,
ranging from 6.93 % (K) to 11.69 % (standard – P). The
conductivity of the elements prepared from the studied
bentonites varies significantly in the range from 457
μS/cm to 2800 μS/cm, measured for samples P and S,
respectively. The lowest conductivity as well as the
lowest pH value obtained for bentonite P are connected
with the fact that bentonite P is a natural bentonite (not
activated by soda).
Table 1: Basic parameters of the studied bentonite samples
Tabela 1: Osnovni parametri preiskanih vzorcev bentonita
Sample
moisture pH conductivity LOI
w/% (–) μS/cm w/%
Sa 7.16 10.27 1422 16.30
K 6.93 10.16 1067 13.60
B 8.21 10.43 1456 12.60
S 10.46 10.50 2800 16.30
P 11.69 9.54 457 12.20
The values of the loss of weight at the ignition (up to
900 °C) of individual samples ranged from 12.20 (P) up
to 16.30 % (Sa). The LOI values reached the values
typical for the amount of water present in the interlayer
space of montmorillonite (approximately w = 12 %).
The chemical compositions of the studied samples
were evaluated with XRFS, whereas the amounts of the
analyzed elements were recalculated to the amounts of
oxides and resumed in Table 2. The most important
J. BEÒO et al.: EVALUATION OF THE THERMAL RESISTANCE OF SELECTED BENTONITE BINDERS
466 Materiali in tehnologije / Materials and technology 49 (2015) 3, 465–469
elements associated with the bentonite-binder behavior
under high temperatures (mainly Fe and Mg) were
selected.
Table 2: Technological parameters of the basic salt-core mixtures
Tabela 2: Tehnolo{ki parametri osnovnih slanih me{anic za jedra
Composition
(w/%)/Bentonite Sa K B S P
Na2O 2.00 < 1.00 2.50 3.05 < 1.00
Al2O3 11.80 15.80 16.40 17.16 16.70
SiO2 44.30 53.90 60.20 59.01 59.10
CaO 4.69 2.40 1.49 1.34 1.22
MgO 3.20 2.50 2.20 1.72 2.20
Fe2O3 12.88 8.50 2.16 1.77 4.31
Although sample K should be rich in the Na+ content
(the sample is a soda-activated bentonite), the chemical
analysis performed with XRFS showed the amount of
Na2O to be bellow the detection limit; the same situation
was observed for bentonite P activated without soda.
The lowest amount of Na2O, measured for samples K
and P, is in good agreement with the measured conduc-
tivity of their water suspensions (Table 1).
From the background research, mentioned above, it is
evident, that the thermal resistance of bentonite binders
to higher temperatures is closely connected to the
Mg and Fe contents. According to the chemical analysis
of the studied samples there is no significant difference
between the Mg contents of the bentonite samples.
The Mg content ranged from 1.72 % (sample S) up to
3.20 % (sample Sa). On the other hand, the Fe content
differs significantly. The highest value of the Fe content
was obtained for sample Sa (12.88 %), while the lowest
value of the Fe content was detected for sample S
(1.77 %). On the basis of the results of the XRFS anal-
ysis, it can be assumed that the highest thermal resis-
tance will be obtained for samples S, B or P.
For the determination of the thermal stability of the
studied bentonite samples two methods were selected.
The first procedure includes a determination of the
thermal stability as the ratio of the mechanical properties
of the fresh and annealed molding mixtures. The
temperature of the thermal exposition of 550 °C/1 h was
selected on the basis of the background research.6–12
The values of the splitting strength of the bentonite
molding mixture with individual bentonite samples are
summarized in Figure 1.
The evaluated values of the technological parameters
significantly depend on the type of a bentonite sample as
evident from Figures 1 and 2.
The optimum mechanical properties required for
molds (high values of the splitting strength in the fresh
state) were obtained for all the samples, but the highest
values were found for sample Sa (31 kPa) and K
(30 kPa), respectively.
However, the lowest thermal resistance, calculated as
the ratio of the splitting strengths of the fresh and
annealed bentonite molding mixtures was detected for
bentonite Sa (a 100 % decrease in the splitting-strength
values).
From this point of view, even though sample Sa in its
fresh state suggested the best mechanical properties, the
highest thermal stability was obtained for samples P and
S, where no change was found.
A slight increase in the splitting-strength values de-
tected for samples S (+5.6 %) and P (+4.3 %) was pro-
bably caused by a measurement error.
Generally, a more sensitive parameter for the evalu-
ation of the bentonite-binder quality is the determination
of the wet tensile strength of the bentonite molding mix-
ture. The results of these experiments are summarized in
Figure 2.
All the samples suggest satisfactory values of the wet
tensile strength. In practice there is a rule, according to
which the value of the wet tensile strength should be
J. BEÒO et al.: EVALUATION OF THE THERMAL RESISTANCE OF SELECTED BENTONITE BINDERS
Materiali in tehnologije / Materials and technology 49 (2015) 3, 465–469 467
Figure 2: Thermal stability of bentonites determined as the ratio of
wet tensile strengths
Slika 2: Toplotna stabilnost bentonitov, dolo~ena kot razmerje natezne
trdnosti v vla`nem
Figure 1: Thermal stability of the bentonites determined as the ratio
of splitting strengths
Slika 1: Toplotna stabilnost bentonita, dolo~ena kot razmerje trdnosti
pri cepljenju
higher than 2.0 kPa. In these experiments it ranges from
2.5 kPa (S) to 4.4 kPa (K). After the annealing the values
of the wet tensile strength significantly decreased for all
the bentonite samples.
The highest decrease was also detected for sample Sa
(–100.0 %). The lowest decrease in the wet tensile
strength was obtained for samples S (–12.0 %) and P
(–22.9 %).
The second method of determining the bentonite
thermal stability includes a preparation of two molding
mixtures. The first mixture was prepared with 5 % dried
bentonite and the second with 5 % annealed bentonite.
Then active-clay tests (a determination of the methy-
lene-blue consumption) were carried out. The thermal
stability was calculated as the ratio of the bentonite
molding mixtures with fresh (dried, 105 °C) and
annealed bentonites.
The results of these experiments are summarized in
Figure 3.
The results of these experiments (the determination
of active bentonite) suggested the same trend as the
previous experiments based on the determination of the
mechanical properties of the molding mixtures with
individual bentonite samples.
The minimum (the lowest) thermal stability was also
found for sample Sa (–53.8 %) and the maximum ther-
mal stability was found for sample P (–8.5 %).
Finally, all the most important parameters based on
the results of the thermal-stability determination for the
individual bentonite samples are summarized in Table 3.
Table 3: Resume of the thermal stability of the studied bentonite
samples
Tabela 3: Pregled toplotnih stabilnosti preiskovanih bentonitnih
vzorcev
Parameter/Bentonite Sa K B S P
STS
(%)
–100.0 –36.7 –21.7 +5.6 +4.3
WTS –100.0 –46.5 42.4 –12.0 –22.9
ABT –53.8 –18.2 29.5 –18.4 –8.5
MgO 3.20 2.50 2.20 1.72 2.20
Fe2O3 12.88 8.50 2.16 1.77 4.31
Note: STS – change in the splitting strength
WTS – change in the wet tensile strength
ABT – change in the active-bentonite test
The experiments carried out in order to evaluate the
thermal stability of the selected bentonite binders
commonly applied at the Czech and Slovak foundries
show that, in all the cases, the lowest thermal resistance
was observed for sample Sa. It is probably related to the
highest amount of Fe (12.88 %). This fact is in line with
our previous research7.
Even if sample P is natural sodium bentonite (not
soda activated) this sample shows the highest thermal
stability. It also shows the minimum decrease in the
mechanical properties and methylene-blue consumption
applied for the evaluation of the thermal stability. It is in
good accordance with the theory of clay binders.4,7
4 CONCLUSION
Thermal stability, the natural property of clay mine-
rals, depends on the source, the mineralogical and
chemical composition of clay and, thus, it is always
necessary to evaluate the suitability of a given clay for a
selected purpose.
The results obtained within this work clearly show
the influence of the Fe amount on the thermal stability of
bentonites as the highest thermal stability was observed
for bentonite P which also shows almost the lowest iron
amount. A high amount of iron was observed for sample
Sa which demonstrates the lowest thermal resistance.
The optimum method for assessing the thermal
stability of the bentonite binder appears to be the method
based on a different methylene-blue consumption of the
bentonite molding mixture with fresh and annealed
bentonite. This method eliminates the potential problems
with the preparation of a molding mixture and the
natural loss of water (drying) from the mixture during
the experiments, which can negatively affect the results
of the measurements. This method is also more sensitive
and more conclusive with respect to the evaluation of the
thermal resistance of a bentonite binder.
Acknowledgement
The research was carried out within the frame of
internal projects of the V[B-Technical University of
Ostrava, SP2014/61 and SP2014/62.
5 REFERENCES
1 C. Paluskiewicz, M. Holtzer, A. Bobrowski, FTIR analysis of
bentonite in moulding sands, Journal of Molecular Structure, 880
(2008) 1–3, 109–114, doi:10.1016/j.molstruc.2008.01.028
J. BEÒO et al.: EVALUATION OF THE THERMAL RESISTANCE OF SELECTED BENTONITE BINDERS
468 Materiali in tehnologije / Materials and technology 49 (2015) 3, 465–469
Figure 3: Thermal stability of bentonites determined as the ratio of
the methylene-blue tests
Slika 3: Toplotna stabilnost bentonitov, dolo~ena kot razmerje preiz-
kusov metilen modro
2 V. J. Inglezakis et al., Removal of Pb(II) from aqueous solutions by
using clinoptilolite and bentonite as adsorbents, Desalination, 210
(2007) 1–3, 248–256, doi:10.1016/j.desal.2006.05.049
3 H. A. Patel et al., Preparation and characterization of phosphonium
montmorillonite with enhanced thermal stability, Applied Clay
Science, 35 (2007) 3–4, 194–200, doi:10.1016/j.clay.2006.09.012
4 J. Madejova et al., Behaviour of Li+ and Cu2+ in heated mont-
morillonite: Evidence from far-, mid-, and near-IR regions,
Vibrational Spectroscopy, 40 (2006) 1, 80–88, doi:10.1016/j.vibspec.
2005.07.004
5 F. Mik{ovský, P. Lichý, The oolitization rate determination of
bentonite moulding mixtures, Archives of Foundry Engineering, 8
(2008) 2, 103–106
6 I. Vasková, J. Malik, P. Futá{, Tests of moulding mixture by using
various clay binder granularity, Archives of Foundry Engineering, 9
(2009) 1, 29–32
7 P. Jelinek et al., Thermostability of montmorillonitic clays, China
Foundry, 11 (2014) 3, 201–207
8 L. Heller et al., An approximation of the position of some cations in
dehydroxylated montmorillonite, Clay Minerals Bulletin, 4 (1961)
25, 213–220
9 R. Gauglitz, H. E. Schwiete, Thermochemical investigations on
montmorillonite with regards to type, grain size, and cation loading,
Berichte Deutsche Keramik Ges, 38 (1964), 43–49
10 P. Wu, C. Ming, The relationship between acidic activation and
microstructural changes in montmorillonite from Heping, China,
Spectrochimica Acta A, 63 (2006) 1, 85–90, doi:10.1016/j.saa.2005.
04.050
11 C. Grefhorst, Investigation of bentonites, Giesserei, 93 (2006) 5,
26–31
12 V. I. Kaplun et al., New studies of ^erkas bentonite, Litìjnoje proiz-
vodstvo, 6 (2006), 12–15
J. BEÒO et al.: EVALUATION OF THE THERMAL RESISTANCE OF SELECTED BENTONITE BINDERS
Materiali in tehnologije / Materials and technology 49 (2015) 3, 465–469 469

L. MALE^EK et al.: DEVELOPMENT OF NUMERICAL MODELS FOR THE HEAT-TREATMENT-PROCESS ...
DEVELOPMENT OF NUMERICAL MODELS FOR THE
HEAT-TREATMENT-PROCESS OPTIMISATION IN A CLOSED-DIE
FORGING PRODUCTION
RAZVOJ NUMERI^NIH MODELOV ZA OPTIMIZACIJO
POSTOPKA TOPLOTNE OBDELAVE PRI PROIZVODNJI
ODKOVKOV V ZAPRTIH UTOPNIH ORODJIH
Ladislav Male~ek1, Mikulá{ Fedorko1, Filip Van~ura2, Hana Jirková2,
Bohuslav Ma{ek2
1COMTES FHT a.s., Prùmyslová 995, 334 41 Dobøany, Czech Republic
2University of West Bohemia in Pilsen, Výzkumné centrum tváøecích technologií – FORTECH, Univerzitní 22, 306 14 Plzeò, Czech Republic
ladislav.malecek@comtesfht.cz
Prejem rokopisa – received: 2014-08-15; sprejem za objavo – accepted for publication: 2014-09-04
doi:10.17222/mit.2014.196
The paper describes a numerical simulation of the current technology of heat treatment of closed-die forgings made of the
25CrMoS4 steel. The aim of this simulation was to create a temperature model enabling a temperature analysis of closed-die
forgings during the heating to the austenitization temperature. This model would permit the heating and soaking times to be
reduced. The paper also describes a numerical simulation and material/technological modelling of the current forming techno-
logy and the subsequent still-air cooling of a selected type of closed-die forgings for the automotive industry. This numerical
simulation provides information on the material flow, the part size and the deformation rate during forming and on the
temperature conditions during handling, forming and subsequent still-air cooling. Using the material/technological modelling,
samples corresponding to the selected locations of a forging can be obtained. By combining these two techniques, controlled
cooling of closed-die steel forgings will be developed and optimized as a substitute for heat treatment. It is also possible to
optimize the process in terms of both quality and energy consumption. Both numerical simulations were applied to the
technology of forming and heat treatment of closed-die forgings of microalloyed steel, chromium-molybdenum 25CrMoS4, at
the company of Kovárna VIVA a.s.
Keywords: 25CrMoS4, MARC, DEFORM, closed-die forging
^lanek opisuje numeri~no simulacijo sedanje tehnologije toplotne obdelave odkovkov iz jekla 25CrMoS4 v zaprtih orodjih.
Namen te simulacije je bil postavitev temperaturnega modela, ki bi omogo~il temperaturno analizo odkovkov, kovanih v zaprtih
utopih, med ogrevanjem na avstenitizacijo. Ta model naj bi omogo~il skraj{anje ogrevanja in zadr`evanja na temperaturi. ^lanek
opisuje tudi numeri~no simulacijo in materialno-tehnolo{ko modeliranje sedanje tehnologije preoblikovanja in ohlajanja na
mirujo~em zraku izbranih utopnih izkovkov za avtomobilsko industrijo. Ta numeri~na simulacija omogo~a informacijo o toku
materiala, o velikosti delov in hitrosti deformacije med kovanjem in o temperaturnih razmerah med manipuliranjem, preobli-
kovanjem in ohlajanjem na mirujo~em zraku. Z materialno-tehnolo{kim modeliranjem se lahko dobijo vzorci, ki ustrezajo
izbranemu polo`aju kovanja. S kombiniranjem teh dveh tehnik bo razvito in optimirano kontrolirano ohlajanje izkovkov v
zaprtih utopih kot nadomestilo za toplotno obdelavo. Proces je mogo~e optimirati tudi s stali{~a kvalitete in porabe energije.
Obe numeri~ni simulaciji sta bili uporabljeni pri tehnologiji preoblikovanja in toplotne obdelave izkovkov iz krom-molib-
denovega mikrolegiranega jekla 25CrMoS4 v zaprtih utopnih orodjih v podjetju Kovárna VIVA, a. s.
Klju~ne besede: 25CrMoS4, MARC, DEFORM, zaprto utopno kovanje
1 INTRODUCTION
The production of closed-die steel forgings involves a
series of forming operations and the subsequent heat
treatment. The forming process typically consists of
several operations. The ones most frequently used are
upsetting, preforming, finish-forging and trimming.
Trimmed forgings are transferred with a conveyor to a
container where they cool down to the ambient tempe-
rature. In order to attain the desired mechanical proper-
ties, the cooling is followed by re-heating the concerned
parts in a continuous-tunnel furnace and by quenching
them. Today’s closed-die-forging plants strive to shorten
this cycle or even omit some of the operations.
Several approaches are available for achieving this
goal. One of them involves the use of numerical simula-
tions. The present paper focuses on two possible
applications of a numerical simulation to optimise the
production of closed-die forgings. The first one aims at
optimising the heating and soaking of forged parts prior
to quenching. The other uses a numerical simulation for
constructing a material/technological model in order to
develop a new method of the thermomechanical treat-
ment of forged parts.
The goal of the first application was to construct a
temperature model. It would be used for predicting the
temperature fields in the forged parts during heating and
soaking at the quenching temperature in the existing
heat-treatment process. Knowing the temperature distri-
bution, it is possible to adjust the process and potentially
reduce the tact time in the production.
Materiali in tehnologije / Materials and technology 49 (2015) 3, 471–475 471
UDK 519.61/.64:621.78:621.73.043 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 49(3)471(2015)
With respect to the second application, the develop-
ment of the material/technological model, the paper
describes a comprehensive numerical simulation of a
forming process, including the subsequent still-air cool-
ing. The forged part in question belongs to a larger group
of products of a similar shape. The paper also presents
the results of physical modelling of the forging process
on a thermomechanical simulator. It compares the
properties of the resulting specimens with the conditions
of the corresponding locations within the actual forged
part.
2 NUMERICAL MODELLING OF HEAT
TREATING A FORGED PART
The objective of the numerical simulation of the heat
treatment was to map the effects of the radiant heat from
the furnace lining on the forgings and the effects of the
radiant heat between the forgings themselves. The forged
parts were made of the 25CrMoS4 material (Figure 1 and
Table 1). In the process, these forgings were arranged in
a charging basket passing through a continuous heating
furnace.
The model was constructed with the use of the data
obtained from the heat-treatment lines. The computation
was carried out using the MSC.MARC/MENTAT soft-
ware. This software employs the finite-element method
and is suitable for solving multiphysical problems1.
The simulation of the heat treatment was an iterative
process. The goal was to fine-tune the simulation to
match the data obtained from the heat-treatment lines.
The iterative approach consisted of a gradual refinement
and an addition of the input data to the computational
model. The computational model comprised three types
of bodies (groups of objects forming a single entity): the
furnace, the charging basket and the forgings (Figure 2).
The CAD models provided the input data for generating
the mesh in individual bodies.
The meshes used for solving the problem consisted
of hexagonal elements for the furnace and the basket and
tetragonal elements for the complex-shaped forgings.
The element size was changing in all the bodies as the
computation was gradually made more accurate. There
were two reasons for it. One was related to the total
number of the elements and the other to the element size
ratio with respect to the view-factor setting. The view
factor is used in analysing the heat transfer by radiation.
It defines the proportion of the radiation from surface A
that reaches surface B. In the model, the view factor indi-
cates the visibility of the face elements of the individual
bodies in the furnace to one another and to the elements
of the inner surface of the furnace chamber. As a rule,
the more elements there are in a computational model,
the more accurate the results are – and the more face
surfaces of the elements there are. With these numbers
increasing, the computation time of the furnace heating
simulation increases as well. For this reason, the analysis
was first tried out using a simplified thermal model
shown in Figure 3. The goal was to examine the effect of
the view factor on the heat transfer by radiation between
two simplified objects.
The meshed objects were assigned material proper-
ties. The properties (the thermal conductivity and the
specific heat) were measured for the forgings using
thermophysical measurement methods. The material
properties of the basket and the furnace were retrieved
from the material data library of the software. The com-
putation was fine-tuned by defining a permanent thermal
contact between the basket and the forged parts. The
L. MALE^EK et al.: DEVELOPMENT OF NUMERICAL MODELS FOR THE HEAT-TREATMENT-PROCESS ...
472 Materiali in tehnologije / Materials and technology 49 (2015) 3, 471–475
Table 1: Chemical composition of 25CrMoS4 steel in volume fractions, "/%
Tabela 1: Kemijska sestava 25CrMoS4 jekla v prostorninskih dele`ih, "/%
Element C Mn Si max. P max. S Cr max. Mo max.
Content 0.22–0.29 0.60–0.90 0.40 0.035 0.02–0.04 0.90–1.20 0.15–0.30
Figure 2: Bodies used in computing a temperature model in the
MARC software environment
Slika 2: Telesa, uporabljena za izra~un temperaturnega modela v
okolju programske opreme MARC
Figure 1: Shape of a forging – a 3D view
Slika 1: Oblika izkovka – 3D-pogled
initial temperature of the forgings was 20 °C. At the start
of the simulation, the furnace temperature was 690 °C. It
changed during the simulation in accordance with the
schedule used. The furnace heating and soaking schedule
was constructed in accordance with the real-world
conditions. It was applied to the side walls and the top
wall of the furnace chamber. Heating by radiation was
first modelled using the MONTE-CARLO method
which, however, did not yield adequate results. There-
fore, the HEMI_CUBE method was employed. This
method uses a pre-defined hollow space, within which
the heat is reflected from or absorbed by the objects. The
hollow space is a numerical zone where the outer ele-
ments of the bodies constitute a working space within
which the view factor is computed.
Due to the increasing computation time, some
aspects that substantially complicated the simulation
were neglected and certain preconditions were defined.
The variation in the position of the basket inside the
furnace was neglected, as it can be taken into account by
adjusting the thermal schedule. The wire basket was
substituted with a solid metal-sheet container in order to
shorten the computation of the view factor. The tempera-
ture field inside the furnace was considered to be uni-
form, although the actual temperature field is not con-
stant. It is affected by opening the furnace door, by the
transitions between its zones with different temperatures,
the types of heating and the temperature-measurement
methods. The results of the FEM simulation were com-
pared with the temperature curves obtained in the se-
lected locations of the real-world forgings in the produc-
tion (Figure 4).
Simulation results (Figure 5) were in agreement with
the temperature curves obtained for the forged parts in
the continuous furnace. Therefore, the numerical model
is suitable for this type of analysis. It can be used for
predicting the temperatures of the forged parts during
heating and soaking before quenching.
3 NUMERICAL SIMULATION OF FORMING
AND COOLING A SPECIFIC TYPE OF FORGED
PARTS
A numerical simulation of forging a selected type of
forged part (Figure 1) was carried out using DEFORM
3D, a program developed for modelling forging pro-
cesses. The input data for the simulation was obtained by
measuring the mechanical and thermophysical properties
of the 25CrMoS4 steel, the material of the forged part.
The goal of the measurement was to obtain an accurate
description of the plastic and temperature behaviours of
the material for the numerical simulation. The plastic
behaviour of the forged material was described with the
flow stress/temperature (T), flow stress/strain (e) and
flow stress/strain rate (ë) relationships in the form of
curves. The flow-stress levels were found using the
Rastegaev test2,3. The temperature behaviour of the work-
piece, i.e. the changes in the temperature field within the
L. MALE^EK et al.: DEVELOPMENT OF NUMERICAL MODELS FOR THE HEAT-TREATMENT-PROCESS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 471–475 473
Figure 5: Comparison between the temperatures found with the
numerical simulation and the field measurement (curve 1 shows the
furnace temperature)
Slika 5: Primerjava med temperaturami, dobljenimi z numeri~no
simulacijo in z meritvami (krivulja 1 prikazuje temperaturo pe~i)
Figure 4: Layout of forged parts in the furnace with thermocouple
locations. The layout was also used for the simulation of heat treat-
ment (quenching).
Slika 4: Razporeditev izkovkov v pe~i s polo`ajem termoelementov.
Razporeditev je bila uporabljena tudi za simulacijo toplotne obdelave
(kaljenja).
Figure 3: Trial model (examination of the effect of the view factor
between two simple objects: cuboids)
Slika 3: Preizkusni model (preiskava vpliva faktorja videza med dve-
ma enostavnima kvadratastima objektoma)
forged part during forming, was described using the
measured specific-heat and thermal-conductivity values,
as in the previous simulation of heat treatment. A
kinematic model of the LMZ 2500 press, in which the
actual forged part was made, was developed. The
simulation was based on the forging-sequence
description provided by the company of Kovárna VIVA
a.s., as well as on the manufacturing-route analysis and
on the field measurement (Figure 6).
The model comprised all the forming operations.
Their sequence consisted of: upsetting – preforming –
finish-forging – trimming. After the trimming, the forg-
ing cooled in still air to the ambient temperature. All the
relevant handling times were taken into consideration,
including the duration of the transfer of the forged part
by the conveyor to the container.
The goal of the numerical modelling was to obtain
the strain and temperature versus the time plots which
were going to be used as the input data for the thermo-
mechanical simulator (for the material/technological
modelling). The material/technological modelling allows
the entire process model to be validated using real
specimens and also permits the microstructure evolution
and mechanical properties to be mapped4. The point-
tracking method was employed to determine the tempe-
rature-versus-time and strain-versus-time curves for the
selected locations during the production of the forged
part (Figure 7). A single representative point (P1) was
selected for the physical simulation. The information
obtained for this point of the forged part, i.e. the
L. MALE^EK et al.: DEVELOPMENT OF NUMERICAL MODELS FOR THE HEAT-TREATMENT-PROCESS ...
474 Materiali in tehnologije / Materials and technology 49 (2015) 3, 471–475
Figure 7: Tracked points on the cross-section of the FEM model of the forged part (left) and a micrograph of the P1 location on the real-world
forged part (right)
Slika 7: Spremljane to~ke na prerezu FEM-modela izkovka (levo) in mikrostruktura polo`aja P1 v realnem izkovku (desno)
Figure 6: Block diagram of the manufacturing process of the forged part
Slika 6: Blok-diagram izdelavnega procesa izkovkov
Figure 9: Bainite-ferrite microstructure of the forged part (HV 284)
Slika 9: Bainitno-feritna mikrostruktura izkovka (HV 284)
Figure 8: Strain and temperature plots for the tracked point
Slika 8: Diagram napetosti in temperature za preiskovano to~ko
strain-time and temperature-time curves, is shown in
Figure 8.
Using this data, a schedule for the thermomechanical
simulator was developed and applied to an actual speci-
men. The microstructure of the real-world part (Figure
9) was then compared with the specimen microstructure
upon the physical simulation (Figure 10) conducted for
the selected point (P1). In both cases, the microstructure
consisted of bainite and a portion of ferrite. For the sake
of comparison, the measured Vickers-hardness values are
shown as well.
4 CONCLUSIONS
Finite-element-method-based simulation is a power-
ful tool that can provide information about the variables
that are difficult to measure otherwise: the strain and
temperature curves for particular points of a forged part.
The knowledge of these values is the key to optimising
the existing processes and developing new procedures
and materials. This is, however, impossible without veri-
fied models, required for a reliable analysis of the pro-
cess. The present work deals with two applications of a
FEM simulation to analysing the manufacturing routes in
closed-die forging.
The first application of the numerical simulation in-
volved constructing a temperature model. It described
the temperature changes in closed-die forgings during
the heating to the austenitizing temperature before the
quenching. Using this model, the heating and soaking
times of the forgings in the furnace can be shortened, the
optimum layout of the forgings in the furnace can be
found and various types of problems solved.
In the model, all the heat-transfer modes were taken
into consideration. The most effective method of the
solution was sought, taking account of the accuracy of
the results. Due to the complexity of the problem, the
computation times of the simulation variants were on the
order of hundreds of hours. The sizes of database files
even exceeded 100 GB. For this reason, this model will
continue to be developed in an effort to shorten the com-
putation times and reduce the data storage requirements.
Gradual improvement in the accuracy of the model is a
matter of course.
The second application of the numerical simulation
involved an analysis of a closed-die forging process for a
selected forged part. This model was developed to obtain
the temperature and strain data to be used as the input
data in constructing a material/technological model.
Such a model combines the findings from the numerical
and physical simulations for assessing the feasibility of
substituting the existing hardening process. The available
alternative is the thermomechanical treatment (combin-
ing forming and the subsequent controlled cooling).
It was found that thermomechanical treatment can
produce practically identical properties of a workpiece as
conventional hardening. However, such results should be
interpreted with caution and this finding should be
supported by a larger body of statistical data. In future
efforts, the FEM simulation of forming processes will be
refined, e.g., using Johnson-Cook model for describing
the plastic behaviour of a forged part instead of the curve
plots employed so far.
Acknowledgement
This paper reports the results obtained under project
TA02010390 "Innovation and Development of New
Thermo-Mechanical and Heat Treatment Processes of
Die Forgings by the Transfer of Findings Obtained from
Material-Technological Modelling". The project runs in
the framework of the ALFA programme and is funded
from the specific resources of the state budget for re-
search and development through the Technology Agency
of the Czech Republic.
5 REFERENCES
1 Marc® 2012, Volume A: Theory and User Information, User’s
manual
2 I. Poláková, T. Kubina, Flow stress determination methods for nume-
rical modelling, 22nd International Conference on Metallurgy and
Materials, METAL 2013 , Brno, Czech Republic, 2013, 273– 278
3 S. L. Semiatin, T. Altan, Measurement and Interpretation of Flow
Stress Data for the Simulation of Metal-Forming Processes, Mate-
rials and Manufacturing Directorate, Air Force Research Laboratory,
Wright-Patterson Air Force Base, Ohio, USA, 2010, 1–57
4 V. Pile~ek, F. Van~ura, H. Jirková, B. Ma{ek, Material-Technological
Modelling of the Die Forging of 42CrMoS4 Steel, Mater. Tehnol., 48
(2014) 6, 869–873
L. MALE^EK et al.: DEVELOPMENT OF NUMERICAL MODELS FOR THE HEAT-TREATMENT-PROCESS ...
Materiali in tehnologije / Materials and technology 49 (2015) 3, 471–475 475
Figure 10: Bainite-ferrite microstructure of the physical-simulation
specimen (HV 270)
Slika 10: Bainitno-feritna mikrostruktura v fizikalno simuliranem
vzorcu (HV 270)

IN MEMORIAM
Alojz Pre{ern, dipl. in`. metalurgije (1920–2015)
Alojz Pre{ern se je rodil 23. 12. 1920 v Globokem pri
Polj~anah. Na Montanistiko na Univerzi v Ljubljani se je
vpisal leta 1939, {tudij metalurgije pa je kon~al leta 1943
v Leobnu. Konec vojne je do~akal v Prekomorski brigadi
NOV. Leta 1946 se je zaposlil kot asistent v jeklarni
@elezarne Ravne, maja 1947 je bil prestavljen v @ele-
zarno Jesenice, leta 1948 pa v @elezarno Zenica, nato pa
je pri{el ponovno v @elezarno Jesenice, kjer je ostal do
avgusta 1963. Povojni ~as je bil ~as najve~jega vzpona
jeklarstva pri nas, pri katerem je Alojz Pre{ern inten-
zivno sodeloval. Postal je vodja vseh topilnic v @elezarni
Jesenice in bil hkrati med najve~jimi strokovnjaki za
jeklarstvo v Sloveniji ter tudi {ir{e v Jugoslaviji. Takrat
so se uvajala nova jekla za gradnjo energetskih objektov,
predelovalno industrijo, strojegradnjo in ladjedelni{tvo.
Velik dose`ek v tedanjem ~asu je bil prehod iz generator-
skega plina na mazut pri kurjenju Siemens-Martinovih
pe~i. Jeklarna na Jesenicah je bila takrat {ola jeklarstva
za vso Jugoslavijo. Za svoje strokovne in organizacijske
dose`ke je bil Alojz Pre{ern odlikovan z redom dela III.
stopnje.
Leta 1963 se je zaposlil na Metalur{kem in{titutu kot
tehni~ni direktor, nato pa je bil leta 1966 imenovan za
direktorja in{tituta. Ponovno je bil imenovan za direk-
torja Metalur{kega in{tituta {e v letih 1970, 1974, 1978
in 1982. Pri svojem delu na in{titutu si je prizadeval
in{titut bolj vklju~iti v tehnolo{ki razvoj slovenske
metalurgije in je sodeloval v prizadevanjih pri zdru`eva-
nju slovenskih `elezarn. Leta 1968 je bil sklenjen
dogovor o sodelovanju med tremi slovenskimi `elezar-
nami in delovnimi organizacijami barvne metalurgije.
Kot direktor je Alojz Pre{ern ves ~as ohranjal in negoval
vezi in{tituta s slovenskimi `elezarnami in drugimi
podjetji, ker se je zavedal, da je to edina trdna garancija
za obstoj in{tituta. Njegova velika zasluga je tudi, da se
je leta 1973 Metalur{ki in{titut pridru`il SOZD-u Slo-
venskih `elezarn kot samostojna delovna organizacija,
obenem pa je in{titut ohranil status osrednje raziskovalne
organizacije za vso slovensko metalurgijo.
Kljub mestu direktorja in{tituta je Alojz Pre{ern vse
do leta 1978 intenzivno delal tudi kot raziskovalec na
podro~ju jeklarskih tehnologij in kemizma reakcij v sta-
ljenem jeklu. Podro~ja njegovih raziskovanj so bila:
rekonstrukcije pe~i, gorilniki in zgorevanje, vakuumske
tehnologije, vpihovanje argona in pra{nih snovi, emisije,
povr{inske napake na gredicah, predvsem pa korelacije
med dezoksidacijskimi postopki in nekovinskimi
vklju~ki v jeklu. Upokojil se je leta 1986.
V zasebnem `ivljenju je bil Alojz Pre{ern skrben o~e
sinu in h~eri, ki sta oba uspe{na, vsak na svojem pod-
ro~ju. V prostem ~asu se je ukvarjal tudi s slikarstvom in
zapustil obse`no zbirko svojih slikarskih stvaritev.
Direktorja Alojza Pre{erna, dipl. in`. metalurgije, je
odlikovala strokovna, raziskovalna in poslovodna spo-
sobnost, obenem pa tudi pristnost, tovari{tvo, neposred-
nost in preprostost v odnosih do sodelavcev. Zato se ga
bomo spominjali tudi kot dobrega prijatelja, tovari{a in
sodelavca.
Dr. Matja` Torkar,
Glavni in odgovorni urednik revije
Materiali in Tehnologije
Materiali in tehnologije / Materials and technology 49 (2015) 3, 477 477