VSEBINA – CONTENTS
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
Comparative study of the synthesis and photo-physical characteristics of a new blue-emitting nanocrystal for NUV-excited
LEDs
Primerjalna {tudija sinteze in opti~no-fizikalnih zna~ilnosti novih modro sevajo~ih nanokristalov za z NUV vzbujanih LED
H. Sameie, R. Salimi, A. A. Sarabi, A. A. Sabbagh Alvani, S. Nargesian, H. E. Mohammadloo, Y. Ebrahimi . . . . . . . . . . . . . . . . . . . . . 685
Development of foundry cores based on inorganic salts
Razvoj livarskih jeder na podlagi anorganskih soli
P. Jelínek, F. Mik{ovský, J. Beòo, E. Adámková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 689
Effect of the abrasive grit size on the wear behavior of ceramic coatings during a micro-abrasion test
Vpliv velikosti brusnih zrn na vedenje kerami~nih prevlek pri mikroabrazijskem preizkusu obrabe
M. Horýk Korkut, Y. Küçük, A. C. Karaoðlanli, A. Erdoðan, Y. Er, M. S. Gök . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 695
Accelerated carbide spheroidisation and refinement (ASR) of the C45 steel during induction heating
Pospe{ena sferoidizacija in udrobnenje karbidov (ASR) med indukcijskim segrevanjem jekla C45
D. Hauserova, J. Dlouhy, Z. Novy, J. Zrnik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 701
The effect of a styrene-acrylic copolymer synthesized in a high-pressure reactor on the improved corrosion protection of
a two-component polyurethane coating
Vpliv stiren-akrilnega kopolimera, sintetiziranega v visokotla~nem reaktorju na izbolj{anje korozijske za{~ite dvokomponentnega
poliuretanskega premaza
U. [egedin, K. Burja, F. Malin, S. Skale, B. Znoj, B. [ket, P. Venturini . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 707
Processing poly(ether etherketone) on a 3D printer for thermoplastic modelling
Obdelava polyether etherketonea na 3D-tiskalniku za termoplasti~no modeliranje
B. Valentan, @. Kadivnik, T. Brajlih, A. Anderson, I. Drstven{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 715
Finite-element model of the indentation for YBCO-based superconductor thin films
Model kon~nih elementov za vtiskovanje tankih plasti superprevodnika YBCO
O. Culha, I. Turkmen, M. Toparlý, E. Celik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 723
Effect of viscosity on the production of alumina borate nanofibers via electrospinning
Vpliv viskoznosti na izdelavo aluminijevooksidnoboratnih nanovlaken z elektrospiningom
M. Ozdemir, E. Celik, U. Cocen . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 735
Effects of solidification parameters on the micro- and macrostructure of the X19CrMoVNbN11-1 steel
Vplivi parametrov strjevanja na makro in mikrostrukturo jekla X19CrMoVNbN11-1
M. Kole`nik, A. Nagode, G. Klan~nik, J. Medved, B. Janet, M. Bizjak, L. Kosec. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 739
Influence of cutting speed on intensity of the plastic deformation during hard cutting
Vpliv hitrosti rezanja na intenziteto plasti~ne deformacije med odrezovanjem
M. Neslu{an, I. Mrkvica, R. ^ep, M. ^illiková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 745
Effect of iron and cerium additions on rapidly solidified Al-TM-Ce alloys
Vpliv dodatkov `eleza in cerija na hitro strjene zlitine Al-TM-Ce
A. Michalcová, D. Vojtìch, G. Schumacher, P. Novák, E. Pli`ingrová . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 757
Shrinkage behavior of a self-compacting concrete
Vedenje samozgo{~evalnega betona pri kr~enju
N. Bouhamou, N. Belas, K. Bendani, A. Mebrouki . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 763
Failure mode of M130 martensitic steel resistance-spot welds
Na~in poru{itve uporovnih to~kastih varov pri martenzitnem jeklu M130
M. Pouranvari, S. M. Mousavizadeh . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 771
Synthesizing a new type of mullite lining
Sinteza nove vrste obloge iz mulita
Z. A}imovi}, A. Terzi}, L. Andri}, L. Pavlovi}, M. Pavlovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 777
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 47(6)683–875(2013)
MATER.
TEHNOL.
LETNIK
VOLUME 47
[TEV.
NO. 6
STR.
P. 683–875
LJUBLJANA
SLOVENIJA
NOV.–DEC.
2013
Influence of processing factors on the tensile strength of 3D-printed models
Vpliv procesnih dejavnikov na natezno trdnost tridimenzionalno tiskanih modelov
T. Galeta, I. Kladari}, M. Karaka{i}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 781
Non-singular method of fundamental solutions for the deformation of two-dimensional elastic bodies in contact
Nesingularna metoda fundamentalnih re{itev za deformacijo dvo-dimenzionalnih stikajo~ih se elasti~nih teles
Q. Liu, B. [arler . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 789
Removal of surface impurities from QCM substrates with the low-pressure oxygen-plasma treatment
Odstranjevanje povr{inskih ne~isto~ s QCM-podlag z obdelavo v nizkotla~ni kisikovi plazmi
R. Zaplotnik, D. Kreuh, A. Vesel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 795
Structural, optical and electrical studies on Si-doped polymer electrolytes
[tudij strukture, opti~ne in elektri~ne lastnosti elektrolitov polimerov, dopiranih s Si
A. Saxena, P. K. Singh, B. Bhattacharya . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 799
Electroless Ni-B-W coatings for improving hardness, wear and corrosion resistance
Nana{anje Ni-B-W s cementacijskim galvaniziranjem za izbolj{anje trdote, obrabe in odpornosti proti koroziji
A. I. Aydeniz, A. Göksenli, G. Dil, F. Muhaffel, C. Calli, B. Yüksel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 803
Physical and numerical modelling of a non-stationary steel flow through a subentry shroud with an inner metering nozzle
Fizikalno in numeri~no modeliranje nestacionarnega toka jekla skozi potopljen izlivek z notranjo {obo
K. Michalek, M. Tkadle~ková, K. Gryc, J. Cupek, M. Macura . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 807
Research of thermal processes for the continuous casting of steel
Raziskave termi~nih dogajanj pri kontinuirnem ulivanju jekla
M. Velicka, D. Dittel, R. Pyszko, M. Prihoda, M. Vaculik, P. Fojtik, J. Burda. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 815
Stationary-creep-rate equations for tempered martensite
Ena~be za hitrost stacionarnega lezenja `arjenega martenzita
F. Vodopivec, B. @u`ek, D. A. Skobir Balanti~, M. Jenko, M. Godec, D. Steiner Petrovi~. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 819
Numerical calculations of stress-intensity factors for a WWER-1000 reactor pressure vessel under a pressurized thermal shock
Numeri~ni izra~uni faktorja intenzitete napetosti v reaktorski tla~ni posodi WWER-1000 pri toplotnem {oku pod tlakom
S. V. Kobelsky . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 825
Dissolution of a copper wire during a hot-dipping process using a SnCu1 lead-free solder
Raztapljanje bakrene `ice med vro~im omakanjem pri uporabi spajke brez svinca SnCu1
D. Steiner Petrovi~, J. Medved, G. Klan~nik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 831
Structural characterization of platinum foil for neural stimulating electrodes using a four-point resistivity-measuring device
Strukturne lastnosti platinske folije za elektrode pri stimulaciji `ivcev, ugotovljene z napravo za {tirito~kovno merjenje upornosti
P. Pe~lin, M. Bizjak, M. Pribo{ek, M. Godec, S. Ribari~, J. Rozman . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 837
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Variation in magnetic properties of sol-gel-synthesized cobalt ferrites
Spreminjanje magnetnih lastnosti kobaltovih feritov, sintetiziranih po sol-gel postopku
A. Hunyek, C. Sirisathitkul . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 845
KRATKE NOVICE – SHORT NOTES
Assessing the quality of corrosion education in the United Arab Emirates
Presoja kvalitete izobra`evanja na podro~ju korozije v Zdru`enih arabskih emiratih
H. L. Lim, L. Al Tayeb Ait Tayeb, M. A. Othman . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 849
DOGODKI – EVENTS
[olski center Ravne na Koro{kem kon~al projekta energetske sanacije in gradnje sodobnega izobra`evalnega centra v skupni vrednosti
pribli`no tri milijone evrov . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 853
H. SAMEIE et al.: COMPARATIVE STUDY OF THE SYNTHESIS AND PHOTO-PHYSICAL CHARACTERISTICS ...
COMPARATIVE STUDY OF THE SYNTHESIS AND
PHOTO-PHYSICAL CHARACTERISTICS OF A NEW
BLUE-EMITTING NANOCRYSTAL FOR NUV-EXCITED LEDS
PRIMERJALNA [TUDIJA SINTEZE IN OPTI^NO-FIZIKALNIH
ZNA^ILNOSTI NOVIH MODRO SEVAJO^IH NANOKRISTALOV
ZA Z NUV VZBUJANIH LED
Hassan Sameie1,2, Reza Salimi1,2, Ali Asghar Sarabi1, Ali Asghar Sabbagh Alvani2,
Saleheh Nargesian3, Hossein Eivaz Mohammadloo1,2, Yalda Ebrahimi1,2
1Amirkabir University of Technology, Faculty of Polymer Engineering & Color Tech., 424 Hafez Ave., 15875-4413 Tehran, Iran
2Amirkabir University of Technology, Color and Polymer Research Center (CPRC), 424 Hafez Ave., 15875-4413 Tehran, Iran
3University of Ottawa, Telfer School of Management, 55 Laurier Ave., East Ottawa, K1N 6N5 Ontario, Canada,
h-sameie@aut.ac.ir
Prejem rokopisa – received: 2012-07-27; sprejem za objavo – accepted for publication: 2013-04-02
In this research, a new blue-emitting nanocrystalline phosphor, SrZn2Si2O7: Eu2+, utilizable in InGaN LEDs, was successfully
synthesized via two routes: the sol-gel method (SG) and solid-state reaction (SS). The effects of the preparation processes on the
crystallization, morphology and thermal properties were analyzed by appropriate techniques, such as thermogravi-
metric-differential thermal analysis (TG-DTA), X-ray diffraction (XRD) and scanning electron microscopy (SEM). The sample
synthesized by the wet chemical method has a relatively regular morphology, a higher phase purity and a crystallite size of
approximately 30 nm. Furthermore, luminescence spectrophotometry was performed for the investigation of the optical
characteristics. The obtained phosphors emit blue light due to the 4f65d1(2D)  4f7(8S7/2) transition of the Eu2+ ions, which act as
luminescence centers in the host lattice. After the excitation in the near-UV region, the phosphors prepared by SG have a higher
emission intensity with a color coordination of x = 0.176, y = 0.193.
Keywords: synthesis, nanocrystalline, luminescence, functional materials, photonic devices
V tej raziskavi je bil uspe{no sintetiziran nov, modro sevajo~ nanokristalni fosfor SrZn2Si2O7: Eu2+, uporaben v InGaN LED, po
dveh metodah: po sol-gel-metodi (SG) in reakciji v trdnem (SS). U~inek postopka priprave na kristalizacijo, morfologijo in
termi~ne lastnosti je bil analiziran s primernimi tehnikami, kot je termogravimetri~na diferen~na termi~na analiza (TG-DTA),
rentgenska difrakcija (XRD) in vrsti~na elektronska mikroskopija (SEM). Vzorec, sintetiziran z mokro kemijsko metodo, ima
relativno pravilno morfologijo, veliko ~istost faze in velikost kristalov okrog 30 nm. Izvr{ena je bila tudi luminiscen~na spektro-
fotometrija za preiskavo opti~nih zna~ilnosti. Dobljeni fosforji so emitirali modro svetlobo zaradi 4f65d1(2D)  4f7(8S7/2)-prehoda
Eu2+ ionov, ki u~inkujejo kot centri luminiscence v gostujo~i re{etki. Po vzbujanju blizu podro~ja UV so imeli fosforji, priprav-
ljeni po SG-metodi, vi{jo intenziteto emisije s koordinacijo barve x = 0,176, y = 0,193.
Klju~ne besede: sinteza, nanokristalini~nost, luminiscenca, funkcijski materiali, fotonske naprave
1 INTRODUCTION
In recent years, light-emitting diodes (LEDs) have
emerged as a prominent class of lighting devices and the
study of RGB phosphors suitable for near-ultraviolet
(NUV) excitation has been attracting more and more
attention for fabricating white LEDs.1,2 LEDs have a
great potential to replace conventional lighting sources,
like incandescent and fluorescent lamps, due to their
many favorable characteristics, such as a long lifetime
and environment-friendly properties.3,4 The new develop-
ments in the field of optical materials are the search for
ideal/suitable phosphors for the conversion of the NUV
emission from InGaN chips into visible light. Among the
different sorts of these materials, silicate phosphors have
attracted researchers’ attention because of the advantages
of a stable crystal structure, stability to high irradiation
powers, etc.5,6 In general, various preparation methods
may greatly affect the crystallization, morphology, parti-
cle size and optical characteristics of phosphor materials.
Compared with samples obtained by the conventional SS
route, the phosphor materials synthesized by the wet
chemical method have advantages such as a low calci-
nation temperature, good mixing of the starting materials
and a higher uniformity of the particle size distribu-
tion.7–9 But from another point of view, the solid state as
the most convenient method has industrial possibilities.
In order to optimize the characteristics of SrZn2Si2O7:
Eu2+, in this study, two experimental methods, SS and
SG, were used to prepare the nanocrystalline phosphors
and the effects of the preparation processes on the
crystallization, morphologies, and optical properties
were investigated.
2 EXPERIMENTAL METHOD
Sr0.96Zn2Si2O7:0.04Eu2+ samples were synthesized
using the SG and SS methods. TEOS and nitrate salts in
SG and metal oxides and acid boric as a flux in the SS
were used to prepare the precursors. Finally, the
Materiali in tehnologije / Materials and technology 47 (2013) 6, 685–687 685
UDK 66.017:535.37:542.9 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)685(2013)
precursors were calcined at 1100 °C for 1 h in a weak
reductive atmosphere of flowing 5 % H2–95 % N2 gas.
Also, in order to characterize the final phosphors, X-ray
diffraction (XRD), scanning and transition electron
microscopy (SEM/TEM), thermogravimetric-differential
thermal analysis (TG-DTA), and fluorescence spectro-
scopy were used.
3 RESULTS AND DISCUSSION
The XRD patterns and the SEM micrographs of the
samples synthesized with SG and SS are shown in
Figure 1. The main phase can be indexed to the phase of
SrZn2Si2O7 for both samples (JCPDS 10-0051). From the
diffraction intensity, it can be seen that the order of the
crystallization and the microstructural regularity for the
phosphor which were prepared via SG are higher than
those of the SS due to the uniformity of the starting
reactants, and thus this method is more favorable for the
formation of superfine phosphors. Moreover, the TEM
image depicts that the average size of the crystallites for
the SG sample is about 30 nm.
The TG-DTA curves of the SrZn2Si2O7 precursors
were studied, as presented in Figure 2, to understand
their pyrolysis behavior and crystallization process. The
following chemical reactions could be inferred, during
the synthesis of the two phosphors:10,11
Sol-gel sample:
Reaction A:
2(Zn(NO3)2.6H2O) (cryst)  2ZnO (amorph) + 12H2O
+ 4NO2 + O2
Reaction B:
Si(OH)4 (amorph)  SiO2 (amorph) + 2H2O
Reaction C:
ZnO (amorph)  ZnO (Cryst)
Reaction D:
2(Sr(NO3)2) (cryst)  2SrO (amorph) + 4NO2 + O2
Reaction E:
SrO (amorph) + SiO2 (amorph)  SrSiO3 (cryst.)
SrSiO3 (cryst) + 2ZnO (cryst) + SiO2 (amorph) 
SrZn2Si2O7 (cryst)
Solid state sample:
Reaction A:
2H3BO3  B2O3 +3H2O
Reaction B:
SrCO3  SrO (amorph) + CO2
Reaction C:
SrO (amorph) + 2ZnO (amorph) + 2SiO2 (amorph) 
SrZn2Si2O7 (cryst)
H. SAMEIE et al.: COMPARATIVE STUDY OF THE SYNTHESIS AND PHOTO-PHYSICAL CHARACTERISTICS ...
686 Materiali in tehnologije / Materials and technology 47 (2013) 6, 685–687
Figure 2: DTA and TG curves of SrZn2Si2O7 gels from 25 °C up to
1150 °C for: a) SG and b) SS samples
Slika 2: DTA- in TG-krivulje gela SrZn2Si2O7 od 25 °C do 1150 °C
za: a) vzorce SG in b) vzorce SS
Figure 1: XRD patterns, SEM images and TEM micrograph of SrZn2Si2O7: Eu2+ prepared via different methods
Slika 1: XRD-posnetka, SEM-posnetka in TEM-posnetek SrZn2Si2O7: Eu2+, pripravljenega z razli~nimi metodami
The effects of the different synthesis methods on the
optical properties were also investigated. Figure 3 shows
the emission spectra of the SrZn2Si2O7: Eu
2+ phosphors
prepared by SG and SS. Under near-UV excitation the
phosphors emit an intense blue light, peaking at 481 nm,
with similar profiles because of the same composition
and the same crystalline lattice, while the intensity is
different, which is consistent with the degree of
crystallization of the phosphors. The emission peak is
attributed to a typical 4f65d1(2D)  4f7(8S7/2) transition of
Eu2+ and for SG sample, the color coordination is (x =
0.176, y = 0.193). However, there is no special emission
of Eu3+ in these spectra, which implies that Eu3+ ions have
been reduced to Eu2+ completely. Double or triple Eu
ions can be present in ionic solids. For the case of the
triple charged, all the 5d and 6s orbitals are empty and
the 4f is partially occupied. The optically active 4f
electrons are shielded from the crystalline electric field
by the outer 5s and 5p shells. The resulting effect is that
the neighboring ligands have very little affect on the 4f
electrons. But for the case of the divalent Eu ions, the
energy separation between the 4f7 and 4f65d1 configura-
tions will be large and these transitions are dipole-
allowed, which are about 106 times stronger than the very
frequently observed 4f4f transitions in trivalent Eu
ions.12 Therefore, a reducing atmosphere helps to reduce
Eu3+ to Eu2+ ions for better optical properties.
4 CONCLUSION
In summary, blue-emitting phosphor SrZn2Si2O7:
Eu2+ was synthesized via two methods, SS and SG, for
LED applications. The reducing atmosphere helped the
Eu3+ ions to reduce to Eu2+ and the 4f65d1(2D)  4f7(8S7/2)
transition of Eu2+ caused the strong emission peak at
about 480 nm for the sample prepared via the SG
method. The results reveal that the sample synthesized
by wet chemical method has a relatively regular morpho-
logy, a small particle size and a higher luminescence
intensity.
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H. SAMEIE et al.: COMPARATIVE STUDY OF THE SYNTHESIS AND PHOTO-PHYSICAL CHARACTERISTICS ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 685–687 687
Figure 3: Emission spectra of SrZn2Si2O7: Eu2+ prepared via different
methods
Slika 3: Emisijski spekter SrZn2Si2O7: Eu2+, pripravljen z razli~nima
metodama

P. JELÍNEK et al.: DEVELOPMENT OF FOUNDRY CORES BASED ON INORGANIC SALTS
DEVELOPMENT OF FOUNDRY CORES BASED ON
INORGANIC SALTS
RAZVOJ LIVARSKIH JEDER NA PODLAGI ANORGANSKIH SOLI
Petr Jelínek, Franti{ek Mik{ovský, Jaroslav Beòo, Eli{ka Adámková
V[B-Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, Department of Metallurgy and Foundry, Ostrava,
Czech Republic
jaroslav.beno@vsb.cz
Prejem rokopisa – received: 2012-08-30; sprejem za objavo – accepted for publication: 2013-03-20
The aim of this study is to describe the possibilities of using salt cores for gravity, low-pressure or high-pressure die-casting
technology. Determinations of the primary, secondary and final residual strengths were carried out in order to evaluate the
possibilities of the salt-core utilization. Furthermore, this contribution is focused on developing composite salts with better
mechanical properties. The solubility of these cores and a possibility of their reclamation in a closed cycle with positive impacts
on the environment were also studied.
Keywords: salt cores, inorganic salts, die casting, PUR Cold-Box, Warm-Box, core solubility and stability
Namen te {tudije je opisati mo`nosti uporabe slanih jeder pri gravitacijski, nizkotla~ni in visokotla~ni livarski tehnologiji.
Izvr{eno je bilo dolo~anje primarnih, sekundarnih in kon~nih zaostalih napetosti, da bi ocenili uporabnost slanih jeder. Poleg
tega je prispevek usmerjen na razvoj kompozitnih soli z bolj{imi mehanskimi lastnostmi. Preu~evana je bila tudi topnost jeder in
mo`nosti njihove predelave v zaprti zanki s pozitivnim u~inkom na okolje.
Klju~ne besede: slana jedra, anorganske soli, tla~no litje, PUR Cold-Box, Warm-Box, topnost jeder in stabilnost
1 INTRODUCTION
The first salt cores appeared in the foundry industry
in the 1970s. An extensive expansion took place in the
1990s in the mass production of Diesel engine pistons.
The cores of simple forms (rings) are made from cook-
ing salt (NaCl) by high-pressure compacting and they
serve for blank casting of the holes (channels) hardly
accessible to mechanical cleaning. They are dissolved in
water, which is a precondition for the use in a techno-
logically closed production cycle. Pistons are made by
gravity or low-pressure casting in dies. The strength
characteristics of the compacted cooking salt meet the
requirements regarding the primary strength (the cold
strength) and hot strength (650–700 °C) of the cores
heated before they are inserted in the mould, which is
also necessary for improving the fluidity of Al-alloys.
The main advantages of this technology are as follows:
• dimensional accuracy and smoothness of the castings
without the use of protective coats;
• solubility of the cores in water and possible recycla-
bility of the salts and water;
• environmentally friendly core production;
• sufficient storage ability of the cores under common
climatic conditions.
When changing to more complex and shape-demand-
ing cores with a possibility of an additional functioning,
using an increased press and dynamics of the metal
injected into the mould (the high-pressure casting), the
research focuses, above all, on improving the mechanical
properties, not only the primary strengths under high
temperatures, but also the residual (secondary) strengths
after a thermal exposure, which have recently drawn a lot
of attention. More efficient ways of a core production,
different from the compacting processes, are searched
for.1,2
For the blank casting of simple holes the high-pres-
sure casting technology uses metal cores, in addition to
the sand ones, made with the PUR COLD-BOX techno-
logy or the Combicore system,3 and the aluminium
channels filled with a salt mixture (cleaned with water).
Though sand cores (PUR COLD-BOX) allow blank cast-
ing of the holes with highly intricate shapes, low
strengths and the possibilities of penetrating the base
sand grains in the alloy and, above all, a bad collapsibi-
lity after the casting, are their disadvantages. Especially
with the new applications of die-casting technologies
(e.g., the master squeeze) lowering the temperature of
the injected alloy, the perfect thermal destruction of the
core polyurethane binder is not followed by difficult
cleaning.
2 PRESENT TRENDS IN THE SALT-CORE
MANUFACTURE
2.1 Cast cores from molten salts
The salt melt is cast in core boxes and to prevent the
moistening, the crystallized cores are stored in an oven
(a minimum of 200 °C). A contraction and volume
shrinkage occur during solidification. High density
(minimum porosity) prevents the dissolution in the water
and, therefore, the cast cores are difficult to remove. To
Materiali in tehnologije / Materials and technology 47 (2013) 6, 689–693 689
UDK 621.74.043:621.743 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)689(2013)
eliminate the shrinkage, a mixture of salts and sand is
recommended. Then the melt (820 °C) is cast in a core
box under pressure.
2.2 High-pressure compacting of salts
A mildly moistened salt is compacted under high
pressure. The cores are then strengthened by way of a
mechanical deformation of the grains (conglomeration)
and a recrystallization along the grain boundaries.
Compacting takes place either under low pressures
(30.0–50.0 MPa) and at the heating temperatures of
500–750 °C or under high compacting pressures
(136.0–362.8 MPa) and lower sintering temperatures
(180–300 °C), allowing a stress release.
2.3 Shooting in core boxes using the binders
After the core shooting (6 bars), the mixture with
inorganic binders (alkali silicates) is hardened either
with the aid of CO2 or with thermal dehydration
(180–210 °C). In the case of organic binders (synthetic
resins) the cores are hardened with the core-box heat
(Warm-Box, 150–300 °C). A possibility of applying
common processes and equipment to the production of
sand cores is an advantage. A higher porosity of salt
cores (25–35 %) enables even dissolution in water
though the bending strength is 2–3 times lower than that
of the compacted cores.
All the mentioned processes, using cooking salt
(NaCl), show a relatively low strength and they do not
meet the requirements for the cores for high-pressure
castings. The research was, therefore, aimed at applying
the high-pressure compacting and improving the strength
characteristics of the cores (the bending strength) by
studying the compacting conditions, the choice of inor-
ganic salts and their different mixtures, the influence of
the salt-crystal shapes, granulometry and, above all, the
composite salts with additives, as well as the hydration
and kinetics during dissolution of the cores in water.
3 INFLUENCING THE STRENGTHS OF THE
SALT CORES COMPACTED UNDER HIGH
PRESSURES
The main criterion was the bending strength mea-
sured with an adapted universal apparatus LRu-2e
(MULTISERVIS MOREK, PL). We measured both pri-
mary strengths, the cold strength and the strength under
high temperatures (650 °C), and the residual (secondary)
cold strength after the thermal exposure (650 °C, 1 h).
Primary strengths were evaluated after 48 h, after the
compacting had been completed and the second phase of
hardening – recrystallization of salt-grain boundaries –
began (Figure 1).
The moisture content of the salts plays a role in the
initial phases of the strengthening (a lower moisture
provides for a higher primary strength) but after a longer
storing time the influence of different moisture contents
is balanced (Figure 2).
The basic tests were done on the chemically pure
salts of KCl and NaCl with the granulometry as follows:
D10 = 72.6–73.4 μm, D50 = 188–189 μm, D90 =
367–370 μm, using a FRITSCH-ANALYSETTE 22
MICRO Tcplus – isopropanol medium.
P. JELÍNEK et al.: DEVELOPMENT OF FOUNDRY CORES BASED ON INORGANIC SALTS
690 Materiali in tehnologije / Materials and technology 47 (2013) 6, 689–693
Figure 3: Influence of compression power on density and bending
strength of KCl samples (d50 = 190 μm; humidity of 0.87 %)
Slika 3: Vpliv sile stiskanja na gostoto in upogibno trdnost vzorcev iz
KCl (d50 = 190 μm; vlaga 0,87 %)
Figure 2: Course of the KCl compacting with different moisture
contents after squeezing with a compression power of 100 · 103 N (23
°C, RV 65 %)
Slika 2: Potek kompaktiranja KCl z razli~no vsebnostjo vlage, po
stiskanju s silo 100 · 103 N (23 °C, RV 65 %)
Figure 1: Grain boundaries after the KCl squeezing ( = 1.85 g cm–3,
bending strength = 8.8 MPa)
Slika 1: Meje zrn po stiskanju KCl ( = 1,85 g cm–3, upogibna trdnost
= 8,8 MPa)
The influence of the compacting force
(30 · 103– 160 · 103 N) on the bending strength and
density of the cores were evaluated, too (Figure 3).
The bending strength grows with the compacting
force up to the value of 6.5 MPa, with a simultaneous
growth of density (Table 1). If the density of the
crystalline KCl salt (1.981 g cm–3) is taken as the base,
then the so-called seeming porosity m can be determined
as follows:
m =
−
⋅
 

KCl core
KCl
100% (1)
where:
m – porosity (%)
KCl – density of KCl (g cm–3)
core – density of the core (g cm–3)
Table 1: Comparison of the primary strengths of KCl after squeezing
with different compression tensions
Tabela 1: Primerjava primarne trdnosti KCl po stiskanju z razli~nimi
tla~nimi napetostmi
Compression
power 50 · 10
3 N 100 · 103 N 160 · 103 N
Compression
tension 28 MPa 56 MPa 87 MPa
Salt BS D BS D BS D
KCl 4.65 1.68 5.53 1.90 6.5 1.94
Note: BS – bending strength (N mm–2); D – density (g cm–3)
or the compacting force of 160 · 103 N m = 2.06 %. The
growth from 100 · 103 N up to 160 · 103 N does not
significantly reflect on the primary strength nor on the
density. A decrease in the seeming porosity from 4 % to
2 % is not significant either. The achieved strengths
exceed multiple times the strength of the sand cores
made with the PUR COLD-BOX technology.
4 HOT STRENGTH AND THE RESIDUAL
STRENGTH
The chosen temperature (650 °C, 0.5 h and 1 h)
corresponded to the temperatures of heating the cores
when inserting them in the dies for gravity casting of the
pistons. It is also the temperature that strengthens the
cores for possible mechanical work. A comparison of the
results of primary and residual strengths of the cores
compacted with different forces is given in Table 2. The
cores from NaCl, under the same compacting conditions,
always had substantially lower primary strengths than
the KCl cores (Table 3), even with the same densities.
Table 2: Comparison of primary strengths of KCl after squeezing with
different compression tensions
Tabela 2: Primerjava primarne trdnosti KCl po stiskanju z razli~nimi
tla~nimi napetostmi
Salt
M A ST CP BS D
% h h 103 N MPa g cm–3
KCl
1.2 0.5 24
50
3.53 1.58–1.64
3.44* 1.72
65
5.41 1.71–1.75
4.73* 1.71
0.87 1.0 72 100
6.35 1.83–1.86
6.34* 1.86–1.89
Note: M – moisture; A – annealing at 650 °C; ST – storage time; BS –
bending strength,* – primary BS; D – density
Table 3: Comparison of the primary strengths of KCl and NaCl cores
made at a constant compression power
Tabela 3: Primerjava primarne trdnosti KCl in NaCl s konstantno
tla~no silo
Salt
M A ST CP PBS D
% h h 103 N MPa g cm–3
KCl 0.87 1 72 100 6.77 1.86–1.88
NaCl 0.83 1 72 100 3.69 1.87–1.89
Note: M – moisture; A – annealing at 650 °C; ST – storage time; PBS –
primary bending strength; D – density
Different behaviour was especially evident during the
evaluation of hot strengths. While the cores from KCl
showed a plastic state at 650 °C, cracks were formed in
the cores from NaCl (Figure 4).
The reason for different behaviours of the cores from
KCl and NaCl was not found in the physical or chemical
properties of both salts (Table 4).
Table 4: General properties of used salts
Tabela 4: Splo{ne lastnosti uporabljenih soli
Properties of salts KCl NaCl
Molar mass g mol–1 74.551 58.443
Melting point °C 770.3 801
Boiling point °C 1411 1413
Density g cm–3 1.981 2.163
Water solubility 20 °C, g/100 ml 34.19 35.86
Crystal lattice edge pm a = 693 a = 562.7
Crystal structure – cubic cubic
The principal difference is found between the crystal
shapes of both salts. While KCl had a homogeneous gra-
nularity of a regular cubic form (Figure 5), NaCl typic-
ally consisted of oval crystals (Figure 6).
P. JELÍNEK et al.: DEVELOPMENT OF FOUNDRY CORES BASED ON INORGANIC SALTS
Materiali in tehnologije / Materials and technology 47 (2013) 6, 689–693 691
Figure 4: Plastic deformation of salt samples
Slika 4: Plasti~na deformacija slanih vzorcev
Differences in the salt-grain shape were also con-
firmed with the measurement of the mean roundness
related to the volume of the particles, for NaCl: SPHT3 =
0.906; for KCl: SPHT3 = 0.778. A similar dependence of
the salt-core strengths on the crystal shape was con-
firmed by C. R. Loper.4
The integrity failure (the crack) in the NaCl sample
runs through the salt crystals under the influence of a
high grain equivalent (Figure 7). Different properties of
both samples can also be identified with the aid of
morphology of the fracture surfaces. The surface of KCl
is compact without any pores or cracks (Figure 8).
5 COMPOSITE SALTS
Up to this time we worked with pure salts (KCl,
NaCl) and the mixtures with other inorganic salts
(K2CO3, Na2CO3, MgSO3, Na2SO4). However, even with
the highest compacting forces (160 · 103 N) we could not
obtain the cores with the mechanical properties that
would meet the requirements for the applications of
high-pressure casting.
For this reason we proceeded to develop composite
salts. Very fine particles with defined granulometry of
different quartz- and non-quartz-based sands were dis-
persed in pure salts (KCl, NaCl), with the aim of failing
the long dislocation lines present in the salt matrix.
Three types of additives (marked as A, B, C) in the
amounts of up to 30 % of the weight of a salt base were
tested. Some partial results are summarized in Table 5.
The results show that with the aid of the composite
salts and some additives (base sands) the core strengths
can be considerably improved, especially the hot
strength (650 °C) – the amounts of the A and B additives
were up to 10 %. The composites with the NaCl matrix
again show considerably lower strengths than those of
KCl. The achieved residual bending strengths in the
range of 8–9 MPa also make it possible to use the me-
P. JELÍNEK et al.: DEVELOPMENT OF FOUNDRY CORES BASED ON INORGANIC SALTS
692 Materiali in tehnologije / Materials and technology 47 (2013) 6, 689–693
Figure 8: Compact structure of KCl
Slika 8: Zgo{~ena struktura KCl
Figure 6: NaCl crystals
Slika 6: Kristali NaCl
Figure 5: KCl crystals
Slika 5: Kristali KCl
Figure 7: Crack in the NaCl structure
Slika 7: Razpoka v strukturi NaCl
Table 5: Strengths of composite salts
Tabela 5: Trdnost kompozitnih soli
KCl
100 · 103 N
A/% B/% C/% NaCl
100 · 103 N
C/%
0 10 30 0 10 30 0 10 30 0 10 30
after 48 h (MPa) 7.53 7.93 3.72 7.63 8.11 1.93 6.47 7.21 6.79 after 48 h (MPa) 3.63 4.26 3.77
Hot strength
650 °C, 1 h (MPa) >8.9 13.16 4.44 8.89 10.14 1.62 8.05 8.92 9.01
Hot strength
650 °C, 1h (MPa) 5.33 2.89 1.35
Residual strength
650 °C, 1 h (MPa) 8.49 8.66 3.91 6.84 9.94 6.02 7.88 7.78 8.06
Residual strength
650 °C, 1h (MPa) 2.55 6.80 1.15
chanical working of the cores, while the hot strengths of
10–13 MPa give hopes that these cores can be used for
high-pressure castings.5
6 CONCLUSION
In the technology of salt-water soluble cores both
new salt mixtures and their manufacturing processes
meeting highly demanding conditions of high-pressure
casting, especially of the automotive castings, are
searched for. High primary strengths and the strengths
under high temperatures (substantially higher than that
of PUR COLD-BOX) give possibilities of treating the
cores by mechanical working. It turns out that utilization
under high pressure, especially when using composite
salt mixtures, can be taken into account, too.
Shooting the salt cores with inorganic or organic
binders in hot-core boxes (WARM-BOX) is a topic for
the next research that could contribute to more extensive
applications of wasteless foundry technology.
Acknowledgement
The research was done with the financial support of
the Technological Agency of the Czech Republic within
the Alfa Programme, TA 02011314.
7 REFERENCES
1 P. Stingl, G. Schiller, Leichte und rückstandfreie Entkernung, Giesse-
rei Erfahrungsaustausch, (2009) 6, 4–8
2 M. Dobosz st. et. al., Development tendencies of moulding and core
sands, China Foundry, 8 (2011) 4, 438–446
3 F. Heppes, Combicore, Giesserei, 98 (2011) 8, 69
4 C. R. Loper et al., The use of salt in Foundry Cores, AFS Transac-
tion, 85–82 (1985), 545–560
5 P. Jelínek et. al., Ovlivòování pevnostních charakteristik solných, ve
vodì rozpustných jader, Slévárenství, 55 (2012) 3–4, 85–89
P. JELÍNEK et al.: DEVELOPMENT OF FOUNDRY CORES BASED ON INORGANIC SALTS
Materiali in tehnologije / Materials and technology 47 (2013) 6, 689–693 693

M. H. KORKUT et al.: EFFECT OF THE ABRASIVE GRIT SIZE ON THE WEAR BEHAVIOR ...
EFFECT OF THE ABRASIVE GRIT SIZE ON THE WEAR
BEHAVIOR OF CERAMIC COATINGS DURING A
MICRO-ABRASION TEST
VPLIV VELIKOSTI BRUSNIH ZRN NA VEDENJE KERAMI^NIH
PREVLEK PRI MIKROABRAZIJSKEM PREIZKUSU OBRABE
Mehmet Horýk Korkut1, Yýlmaz Küçük2, Abdulllah Cahýt Karaoðlanli3, Azmý Erdoðan3,
Yusuf Er4, Mustafa Sabrý Gök3
1Department of Metallurgy Education, Firat University, 23100 Elazýg, Turkey
2Department of Mechanical Engineering, Bartin University, 74100 Bartin, Turkey
3Department of Metallurgy and Materials Engineering, Bartin University, 74100 Bartin, Turkey
4Gazi Vocational High School, Section of Machine Technologies, 06500 Ankara, Turkey
aerdogan@bartin.edu.tr
Prejem rokopisa – received: 2012-08-31; sprejem za objavo – accepted for publication: 2013-03-25
Micro-abrasion tests are commonly used to perform the wear tests on hard coatings. In this study, AISI 1040 mild-steel
specimens were coated with ceramic coatings including oxides with different hardness values. The free-ball micro-abrasion
testing method was used to examine the wear behavior of the ceramic surfaces over different test durations. The
coating-hardness measurements were carried out with a micro-hardness tester. In the experimental studies, the SiC abrasive
particles with three different grain sizes (800, 1000 and 1200 mesh), were used to explore the effect of the abrasive particle size
on the wear ratio of ceramic coatings. The abrasive effect of the SiC particles in the abrasive slurry was evaluated by examining
the SEM micrographs of the coating surfaces. According to the test results, the sample surface coated with Cr2O3 exhibited a
higher wear resistance than those covered with Al2O3 and its compositions with TiO2. It can also be concluded from the results
that an increase in the percentage of the TiO2 powders in the Al2O3 mixture leads to a decrease in the wear resistance.
Keywords: micro-abrasion, ceramic coating, abrasive grain size
Za preizku{anje obrabe trdih prevlek se navadno uporablja mikroabrazijski preizkus. V tej {tudiji so bili vzorci mehkega jekla
AISI 1040 prevle~eni s kerami~no prevleko iz razli~no trdih oksidov. Za ugotavljanje vedenja kerami~nih povr{in pri obrabi z
razli~nim trajanjem preizkusa je bila uporabljena metoda mikroabrazijskega preizkusa s tremi prostimi kroglami. Trdota
prevleke je bila izmerjena z merilnikom mikrotrdote. Pri eksperimentih so bili za ugotavljanje vpliva velikosti abrazijskih delcev
na obrabo kerami~ne prevleke uporabljeni abrazijski delci SiC s tremi razli~nimi velikostmi (zrnatost 800, 1000 in 1200).
Abrazijski u~inek SiC-delcev in abrazijskih odpadkov je bil ocenjen in preiskan s SEM-posnetki povr{ine prevleke. Glede na
rezultate imajo vzorci, prekriti s Cr2O3, ve~jo odpornost proti obrabi kot vzorci z Al2O3 in njihova kombinacja s TiO2. Iz
rezultatov se lahko tudi ugotovi, da pove~anje dele`a TiO2 v me{anici z Al2O3 povzro~i zmanj{anje odpornosti proti obrabi.
Klju~ne besede: mikroabrazija, kerami~na prevleka, velikost abrazijskih zrn
1 INTRODUCTION
Generally, the coatings used for improving the
performance of industrial parts are available in various
types, ranging from hard coatings with high abrasion
resistance to soft, lubricating coatings and applications
requiring low friction coefficients.1
Ceramics stand out in the friction and abrasion appli-
cations in the industry on account of their high hardness,
high chemical stabilities, high oxidation-resistance
values, high temperature- and thermal-barrier features.
However, the high costs involved in their manufacturing
and their brittle characteristics limit the use of ceramics.
For these reasons, ceramics are preferred for the forma-
tion of anti-abrasive layers, applied with thermal-spray
methods rather than being used as bulk materials.2 For
years, various metallic and ceramic coatings have been
applied on materials in order to form abrasion-resistant
surface layers using the powders sized between 10–100
μm with thermal-spray methods.3 The essence of the
atmospheric-plasma-spray (APS) method, which is a
thermal-spray-coating method, is to form a plasma jet for
melting the powder material. Powder particles are
injected by means of a protective gas and the powders
derive their speed and temperature from the plasma jet
via the thermal and kinematic transfer. The particulates
form abrasion-resistant, rapidly stiffening and thin layers
on the surface to be coated.4–7
The micro-scale abrasion test is implemented
successfully in assessing the abrasive performances of
various materials. This technique is applied on metallic
and non-metallic bulk and coating materials along with
various abrasive slurry media.8–11 The abrasion crater
obtained in this way is measured with optical or profilo-
metric methods and thus the abrasion results are eva-
luated in terms of the volume and abrasion mecha-
nisms.12–14 The ceramic coatings such as Cr2O3, Al2O3
and TiO2 applied with the thermal-spray method were
examined tribologically at high temperatures as well as
at room temperature, in dry and lubricant sliding condi-
tions.2 However, there are not enough studies for an
Materiali in tehnologije / Materials and technology 47 (2013) 6, 695–699 695
UDK 621.793:620.178 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)695(2013)
abrasion analysis of the mentioned ceramic coatings by
means of the micro-abrasion technique. For this reason,
in this study, the abrasion behaviours of the Cr2O3,
Al2O3, Al2O3 + 13 % TiO2 and Al2O3 + 40 % TiO2 coat-
ings were examined using the micro-abrasion test.
2 MATERIALS AND METHOD
In this study, the AISI 1040 steel with a diameter of
20 mm and length of 50 mm was selected as the sub-
strate material. Firstly, the surface was cleaned of the
unwanted residues (oil, dust and residual metals) with a
blasting process to achieve a certain roughness value.
After this process, the powders with a composite rate of
80 % Cr and 20 % Ni and baked at 100 °C were applied
to the substrate material so as to form an intermediate
surface with a thickness of 30 μm. The aim of this pro-
cess is to ensure that a stronger bond is established
between the substrate surface material and the ceramic
coating material. Finally, the ceramic-powder coatings
were applied, in two stages, on the substrate surfaces
using the plasma method. The coating thickness was
about 150 μm and the process parameters used are shown
in Table 1. In accordance with the mixture composition
of the ceramic powders, different hardness values for the
ceramic coatings were obtained (Table 2). The SEM
micrographs regarding the Cr2O3 and Al2O3 coatings that
include 40 % of TiO2 are given in Figures 1 and 2,
respectively.
Table 1: Atmospheric-plasma-spraying conditions
Tabela 1: Pogoji pri napr{evanju v atmosferski plazmi
Arc flow rate 80–100 L min–1
Arc pressure 0.689 MPa
Auxiliary gas flow rate 5–15 L min–1
Auxiliary gas pressure 0.345 MPa
Spray rate 2.7–6.8 kg/h
Arc voltage 61–68 V
Arc current 400–600 A
Spray distance 70–100 mm
Table 2: Hardness values of the coatings
Tabela 2: Trdota prevlek
Coating material Hardness (Hv) Roughness (Ra/μm)
Cr2O3 960 0.353
Al2O3 820 0.326
Al2O3 + 13 % TiO2 730 0.237
Al2O3 + 40 % TiO2 630 0.250
The aim of the micro-abrasive wear test is to generate
"the wear craters" on the specimen. So, one can calculate
the wear volume (V), using the crater depth (h) and the
crater diameter (b)15,16 after the wear tests. The volu-
metric mass-loss value can be calculated using both the
crater radius and the crater height. The formulas required
for the calculations are clearly specified in17.
After the coating process, the free-ball micro-abra-
sion wear test was applied to each sample. The free-ball
micro-abrasion method is a simple test technique for
determining the wear behavior of coatings and is
explained in detail in some studies.15,18–20 The ball used
in this test was made from the AISI 52100 steel and had
a diameter of 25.4 mm. Silicon carbide (SiC) was used
as the abrasive with three particle sizes (800, 1000 and
1200 mesh). The abrasive slurry that was created was
composed of 25 % of SiC and 75 % of distilled water.
Each test was repeated three times. During the tests, the
abrasive slurry droplet was implemented as one drop per
20 seconds. To determine the experimental test combina-
tions, the abrasive grit size, the test duration, the coating
material and the spindle speed were selected as the test
factors (Table 3). Each factor has three levels except for
the coating material that has four levels. A total of forty-
five test combinations were carried out in accordance
with the factorial design. Conventional characterization
techniques, such as scanning electron microscopy
(SEM), micro hardness and X-ray diffraction, were
M. H. KORKUT et al.: EFFECT OF THE ABRASIVE GRIT SIZE ON THE WEAR BEHAVIOR ...
696 Materiali in tehnologije / Materials and technology 47 (2013) 6, 695–699
Figure 2: Cross-section of the Al2O3 + 40 % TiO2 coating layer
Slika 2: Pre~ni prerez prevleke iz Al2O3 + 40 % TiO2
Figure 1: Cross-section of the Cr2O3 coating layer
Slika 1: Pre~ni prerez prevleke iz Cr2O3
employed to study the microstructure of the coating
zone.
Table 3: Factors and their levels
Tabela 3: Dejavniki in njihova veli~ina
Factor Level 1 Level 2 Level 3 Level 4
Coating material Cr2O3 Al2O3 Al2O3 +13 % TiO2
Al2O3 +
40 % TiO2
Abrasive grit size
(mesh) 800 1000 1200
Test duration
(min) 1 2 3
Spindle speed
(r/min) 115 160 230
3 RESULTS AND DISCUSSION
After the micro-abrasion tests, the volumetric wear
values were calculated and so, the influence of each test
factor on the results could be determined using an
ANOVA table (Table 4).
The wear loss values for all parameters are given in
Table 5. The factors that have significant effects on the
mass-loss value of the coatings were determined by
analyzing the ANOVA table. Hence, from the ANOVA
table, the significance of the factors, according to their
effects on the mass-loss values, can be presented in a
descending order such as the grit size, the test duration
and the coating material. However, the spindle speed
does not have a significant effect on the wear behavior, if
the selected confidence interval is 95 %. This case can be
attributed to the reduced effect of free-mass weight of
steel ball with increasing the tangential velocity.
Increasing the tangential velocity also leads to a decrease
in abrasive effect by applying lower pressure to the
surface of coating. The effects of the abrasive grit size
and the coating-material factors on the mass loss are
given in Figure 3.
The volumetric mass-loss values increased in all the
samples as a function of the increasing abrasive particle
size. Additionally, it was observed that the Cr2O3 coating
exhibited the lowest degree of abrasion due to its high
hardness values contrary to the expectation. This case
can be explained by taking into account both the surface
roughness and hardness values of coated samples (Table
2), it was observed, contrary to the expectations, that the
Cr2O3 coating exhibited the lowest degree of abrasion
due to its high hardness values. As it is well-known,
hardness is one of the most important parameters in
increasing the resistance to abrasion. On the other hand,
the surfaces with high roughness and hardness values
need to be abraded more as they have low rates of
toughness. However, the abrasive particles fed into the
medium in the form of slurry soon filled in the rough
surfaces, thus acting as a bed between the free ball and
the material, preventing a rapid abrasion of a sample.
M. H. KORKUT et al.: EFFECT OF THE ABRASIVE GRIT SIZE ON THE WEAR BEHAVIOR ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 695–699 697
Table 4: ANOVA table for the main factors
Tabela 4: Tabela ANOVA glavnih dejavnikov
Factor Degree of freedom Sum of squares Mean square F value Prob>F
Model 9 4.084E-03 4.538E-04 9.24 <0.0001
Material 3 9.217E-04 3.072E-04 6.26 0.0016*
Spindle speed 2 3.012E-04 1.506E-04 3.07 0.0592
Duration 2 9.419E-04 4.710E-04 9.59 0.0005*
Grit size 2 2.031E-03 1.016E-03 20.69 <0.0001*
Residual 35 1.719E-03 4.911E-05
Corrected total 44 5.802E-03
*significant factor
Figure 4: Effect of the abrasive grit size vs. the coating material on
the mass loss
Slika 4: Vpliv velikosti zrn abrazijskega sredstva in materiala prevle-
ke na izgubo mase
Figure 3: Effect of the abrasive grit size vs. the coating material on
the mass loss
Slika 3: Vpliv velikosti zrn abrazijskega sredstva in materiala prevle-
ke na izgubo mase
The increase in the mass losses as a function of the
increasing grit size can be explained as follows: An
increase in the size of the abrasive grit causes an increase
in the kinetic energy and the contact area of the grit as a
function of speed; thereby, the abrasive capacity is
increased in a directly proportional way. It is observed
that the factor with the second biggest importance is the
test duration. An increase in the rate of fresh abraders
along with the extended duration, in accordance with the
abrasive slurry dropping frequency (one drop per 20 s),
caused a further loss in the mass (Figure 4). Further-
more, the abrasive grits came into contact with a wider
area of the sample surface, along with the sliding dist-
ance that increased in direct proportion with the test
duration.
Thirdly, taking the coating-material factor into
account, Cr2O3 demonstrated the highest abrasion resi-
stance when compared to the hardness values. Addition-
ally, the hardness value was reduced, along with the
M. H. KORKUT et al.: EFFECT OF THE ABRASIVE GRIT SIZE ON THE WEAR BEHAVIOR ...
698 Materiali in tehnologije / Materials and technology 47 (2013) 6, 695–699
Table 5: Volume loss values for all parameters
Tabela 5: Zmanj{anje volumna za vse parametre
Run Material Spindle speed /(r/min) Duration /min Grit size mesh Volumetric mass loss (mm3)
1 Cr2O3 160 3 1000 0.004794
2 Al2O3+13TiO2 230 3 1000 0.028231
3 Cr2O3 115 1 1000 0.003041
4 Al2O3 160 1 800 0.015952
5 Cr2O3 160 2 800 0.007424
6 Al2O3+40TiO2 160 1 1000 0.012665
7 Cr2O3 115 3 800 0.008543
8 Al2O3+13TiO2 160 2 1000 0.015395
9 Al2O3+13TiO2 115 2 800 0.019840
10 Al2O3+40TiO2 160 1 800 0.016524
11 Al2O3+13TiO2 230 1 1000 0.008543
12 Al2O3+40TiO2 115 1 1200 0.005107
13 Al2O3+13TiO2 115 1 1200 0.003621
14 Al2O3+40TiO2 230 1 1200 0.007531
15 Al2O3+40TiO2 115 3 1000 0.017919
16 Cr2O3 115 2 1200 0.003747
17 Al2O3+40TiO2 160 2 1200 0.005605
18 Al2O3+40TiO2 230 2 800 0.033143
19 Al2O3 115 2 1200 0.002218
20 Al2O3+13TiO2 160 1 1000 0.007748
21 Al2O3+40TiO2 115 1 800 0.009525
22 Al2O3 115 1 1000 0.010870
23 Al2O3+13TiO2 115 3 1000 0.015952
24 Al2O3 160 3 1200 0.001384
25 Al2O3 230 2 800 0.033470
26 Cr2O3 160 1 1200 0.003380
27 Al2O3+40TiO2 160 3 1000 0.017112
28 Cr2O3 230 1 800 0.007009
29 Al2O3+13TiO2 160 1 1200 0.005188
30 Cr2O3 230 2 1000 0.006231
31 Al2O3+13TiO2 160 3 800 0.039037
32 Al2O3+13TiO2 230 2 1200 0.008543
33 Al2O3 230 1 1200 0.002934
34 Al2O3+13TiO2 230 1 800 0.019840
35 Al2O3+13TiO2 115 3 1200 0.009915
36 Al2O3+40TiO2 230 3 1200 0.012197
37 Al2O3 115 3 800 0.032819
38 Al2O3 160 2 1000 0.018125
39 Al2O3 230 3 1000 0.017112
40 Al2O3+13TiO2 115 1 800 0.009784
41 Cr2O3 230 3 1200 0.004281
42 Al2O3+40TiO2 230 3 800 0.059443
43 Al2O3+13TiO2 115 2 1000 0.010729
44 Al2O3+40TiO2 230 1 1000 0.012508
45 Al2O3+40TiO2 115 2 1000 0.015765
increase in the TiO2 percentage in the Al2O3 content,
resulting in an increase in the abrasion amount.
4 CONCLUSIONS
In this study, the wear behaviors of different types of
ceramic coatings were evaluated via the micro-abrasion
tests. The conclusions can be summarized as follows:
Among the specimens, the ceramic coating including
Cr2O3 has the highest value of hardness. The other coat-
ing-hardness quantities, according to their order of mag-
nitude, are pure Al2O3, Al2O3 + 13 % TiO2 and Al2O3 +
40 % TiO2.
The most influential factor for the mass loss was the
abrasive grit size. The test duration and the coating
material were the other effective factors. An increase in
the abrasive grit size led to certain differences in the
abrasive surface topography. The influence of the plastic
deformation of the sample surface with a mesh grit size
of 800 is greater than that of 1 200 mesh.
It was observed that the abrasive mechanism changed
with the change in the coating composition. The plastic
deformation on the surfaces of the Cr2O3 and Al2O3 coat-
ings was seen to be more prominent, whereas the
increase in the rate of TiO2 in Al2O3 brought about a
decrease in the hardness, resulting in a reduced plastic-
deformation severity and a smoother surface topography.
The spindle speed has not much effect on the value of
mass loss due to reducing effect of tangential velocity on
the pressure applied on surface of coating by the steel
ball. The micro-abrasion technique is a method that can
be used for comparative assessments, examining the
abrasive behaviours of ceramic coatings.
5 REFERENCES
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Gancedo, A. Rincon, Friction and wear behaviour of plasma-sprayed
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normal loads, Tribology International, 29 (1996) 4, 333–343
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process parameters on wear behaviour of alumina–titania coatings,
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M. H. KORKUT et al.: EFFECT OF THE ABRASIVE GRIT SIZE ON THE WEAR BEHAVIOR ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 695–699 699

D. HAUSEROVA et al.: ACCELERATED CARBIDE SPHEROIDISATION AND REFINEMENT (ASR) ...
ACCELERATED CARBIDE SPHEROIDISATION AND
REFINEMENT (ASR) OF THE C45 STEEL DURING
INDUCTION HEATING
POSPE[ENA SFEROIDIZACIJA IN UDROBNENJE KARBIDOV
(ASR) MED INDUKCIJSKIM SEGREVANJEM JEKLA C45
Daniela Hauserova, Jaromir Dlouhy, Zbysek Novy, Jozef Zrnik
COMTES FHT, Prumyslova 995, 334 41 Dobrany, Czech Republic
daniela.hauserova@comtesfht.cz
Prejem rokopisa – received: 2012-09-14; sprejem za objavo – accepted for publication: 2013-04-08
The present research focused on improving the mechanical properties of the C45 structural carbon steel. The microstructure of
this material can be transformed, with the accelerated carbide spheroidisation and refinement (ASR) process, to increase the
proof stress and ultimate tensile strength, as well as plasticity and toughness. In this study, the heat-treatment cycle was carried
out using induction heating, which offers high heating rates. The process consists of rapid heating of the feedstock to the
austenitizing temperature, a short hold, cooling down and rapid cycling at around the transformation temperature Ac1. Within a
short interval of time, of the order of no more than a few minutes, this cycle produces a microstructure of fine ferrite grains with
globular carbides. The present paper describes the impact of ASR process parameters on the mechanical properties and micro-
structure of the C45 steel.
Keywords: induction heating, grain refinement, carbide spheroidisation
Ta raziskava je bila usmerjena v izbolj{anje mehanskih lastnosti konstrukcijskega oglji~nega jekla C45. Mikrostruktura tega
materiala se lahko spremeni s postopkom pospe{ene sferoidizacije in udrobnenja karbidov (ASR), kar povi{a mejo plasti~nosti,
natezno trdnost, kot tudi plasti~nost in `ilavost. V tej {tudiji je bil izveden cikel toplotne obdelave z indukcijskim segrevanjem,
ki omogo~a velike hitrosti segrevanja. Postopek sestavlja hitro segrevanje ogrevanca na temperaturo avstenitizacije, kratko
zadr`anje, ohlajanje in hitro cikliranje okrog temperature transformacije Ac1. V ~asovnem intervalu – ne ve~ kot nekaj minut –
to cikliranje povzro~i nastanek mikrostrukture iz drobnih feritnih zrn z globularnimi karbidi. ^lanek opisuje vpliv procesnih
parametrov ASR na mehanske lastnosti in mikrostrukturo jekla C45.
Klju~ne besede: indukcijsko segrevanje, udrobnenje zrn, sferoidizacija karbida
1 INTRODUCTION
Grain refinement in steels is a traditional process for
enhancing their strength. It is often used in plain carbon
steels where the appropriate treatment increases the yield
strength and ultimate strength without additional alloy-
ing or hardening, thus saving costs.1 In addition to the
grain size, carbide morphology is important with regard
to toughness. Where high toughness is required, globular
carbides are the right choice, whereas lamellar pearlite is
not a suitable type of microstructure. When obtaining the
optimum combination of strength and toughness, the
ideal microstructure consists of fine ferrite grains and
globular carbides.
ASR (Accelerated Spheroidisation and Refinement)
is a process producing fine grains and globular carbides
throughout the workpiece in a very short time: of the
order of no more than a few minutes.2 This paper
describes the results of an optimization of the ASR
process using induction heating. This heating method
allows the workpiece temperature to be changed very
rapidly in contrast with conventional annealing in a
furnace.3
2 EXPERIMENTAL WORK
The experimental material was hot rolled and nor-
malized C45 structural steel. The chemical composition
of the steel is listed in Table 1. The initial microstructure
consisted of pearlite and ferrite with areas of lamellar
pearlite (Figures 1a and b). The ferrite grain size was
about 20 μm.
Heat treatment was carried out using a medium-
frequency converter. The number of coils of the inductor
was selected in such a way as to provide as homo-
geneous an electromagnetic field as possible in the coil
centre, leading to a uniform temperature field. The
heating was controlled by a PLC unit and the tempera-
ture was measured by means of a pyrometer and thermo-
couples (Figure 2). The heat-treated specimens were
15-mm-diameter bars with a length of 400 mm.
Materiali in tehnologije / Materials and technology 47 (2013) 6, 701–705 701
UDK 669.14.018.291:621.78.012.5 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)701(2013)
Table 1: Chemical composition of the C45 steel in mass fractions
(w/%)
Tabela 1: Kemijska sestava jekla C45 v masnih dele`ih (w/%)
C Mn Si P S Cr Ni Mo
0.47 0.70 0.23 0.026 0.017 0.07 0.02 0.007
For the purpose of converting lamellar carbides to a
globular morphology, schedules were designed with the
temperature cycling at around or slightly above Ac1 (see
Table 2). The Ac1 temperatures for both heating rates
were determined from the temperature-time records of
induction heating. Ac1 for the heating rate of 80 °C/s was
762 °C and for 20 °C/s it was 747 °C. All the schedules
comprised rapid induction heating, subsequent rapid
temperature cycling and air cooling. Each of schedules 1,
2 and 3 had four temperature cycles. The maximum and
minimum cycle temperatures varied. The difference
between the maximum and minimum cycle temperatures
was 60 °C. In schedule 1, the maximum cycle tempera-
ture was 740 °C; in schedule 2, it was 760 °C and in
schedule 3, it was 780 °C. These schedules took 200 s,
starting with heating and ending with cooling the speci-
men to 550 °C.
Table 2: List of schedules
Tabela 2: Seznam zaporedja ukrepov
Schedule
Heating
rate
(°C/s)
Austeni-
tizing
(°C)
Th/°C Tl/°C Numberof cycles
1 80 – 740 680 4
2 80 – 760 700 4
3 80 – 780 720 4
4 80 840/25 s – – –
5 80 840/25 s 740 670 4
6 80 840/25 s 780 720 4
7 20 – 760 700 4
8 20 – 780 720 4
9 20 840/25 s 780 720 4
Schedule 4 was designed for a microstructure refine-
ment. It comprised austenitizing and the subsequent air
cooling, i.e., essentially, a normalizing process. The
combined normalizing and carbide spheroidising were
intended to create fine ferrite grains with globular car-
bides (schedules 5 and 6). The austenitizing temperature
was 840 °C. Schedule 5 included austenitizing at 840 °C,
cooling in air to 600 °C, heating to 740 °C and additional
three cycles at around this temperature. Schedule 6 was
only different in the cycling temperature that was 780 °C.
An overview of the schedules is given in Table 1.
Schedule 4 took 250 s, while schedules 5 and 6 took
about 300 s between the start of heating and cooling to
the temperature of 550 °C.
Schedules 7, 8 and 9 were proposed with the aim of
finding the influence of the heating rate on the material
properties and microstructure. The heating rate was
reduced from 80 °C/s to 20 °C/s. Apart from the heating
rates, the pairs of schedules 7 and 2, schedules 8 and 3,
and schedules 9 and 6 were identical.
3 RESULTS AND DISCUSSION
3.1 Metallography
The temperatures achieved with schedules 1 and 2
were not sufficient for an extensive austenitization of the
microstructure. When compared with the initial micro-
structure, the grain size remained almost unchanged; part
of the cementite lamellae in pearlite transformed to
rod-like particles.
D. HAUSEROVA et al.: ACCELERATED CARBIDE SPHEROIDISATION AND REFINEMENT (ASR) ...
702 Materiali in tehnologije / Materials and technology 47 (2013) 6, 701–705
Figure 2: Induction heating of a sample
Slika 2: Indukcijsko segrevanje vzorca
Figure 1: a) Optical micrograph of the initial microstructure, b)
scanning electron micrograph of the initial microstructure
Slika 1: a) Svetlobni posnetek za~etne mikrostrukture, b) SEM-posne-
tek za~etne mikrostrukture
The temperature record for schedule 3 showed a tem-
perature hold during cooling, indicating a pearlitic trans-
formation. Temperature cycling, therefore, led to a
partial austenitization accompanied by cementite dissolu-
tion. Cementite lamellae underwent a partial fragmentation
(Figures 3a and b). The micrographs on these figures
show a partial spheroidisation of cementite lamellae and
a certain expansion of cementite into the ferritic regions.
Schedule 4 led to an almost complete microstructure
austenitization. The resulting microstructure consists of
ferrite and pearlite. Part of the pearlitic cementite is
lamellar and part is spheroidised. Lamellar cementite
was probably newly formed. Globular cementite appears
to be the original spheroidised cementite that did not
dissolve during the schedule. Cooling in air resulted in a
rapid eutectoid transformation and the undissolved
cementite lamellae acted as a nucleation agent.4 It was
mainly this schedule that produced considerably refined
ferrite grains: from 20 μm to less than 10 μm.
Schedules 5 and 6 were designed as a combination of
normalisation and the temperature cycling at around
critical temperature Ac1. The initial stage of the sche-
D. HAUSEROVA et al.: ACCELERATED CARBIDE SPHEROIDISATION AND REFINEMENT (ASR) ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 701–705 703
Figure 5: a) Scanning electron micrograph of the initial state, b) scann-
ing electron micrograph after schedule 9
Slika 5: a) SEM-posnetek za~etne mikrostrukture, b) SEM-posnetek
mikrostrukture po ukrepu 9
Figure 3: a) Optical micrograph after schedule 3, b) scanning electron
micrograph after schedule 3
Slika 3: a) Svetlobni posnetek mikrostrukture po ukrepu 3, b) SEM-
posnetek mikrostrukture po ukrepu 3
Figure 4: EBSD map with grain boundaries after schedule 5
Slika 4: EBSD-posnetek zrn z mejami zrn po ukrepu 5
dules was intended to refine the grains, whereas the
cycles at a lower temperature were to spheroidize the
carbides. The microstructures after schedules 5 and 6
were very similar. These schedules led to a grain refine-
ment in the steel producing new carbide regions between
the new finer grains, as did schedule 4. The cycling at
around Ac1 fragmented and partially spheroidised the
cementite lamellae. The EBSD map (Figure 4) shows a
substructure of the ferrite grains both in pearlite and
ferrite regions. The thick lines represent high-angle
boundaries (>15°) and the thin lines represent low-angle
boundaries (from 2° to 15°).
In the schedules with lower heating rates, the tempe-
rature of 760 °C (schedule 7) was sufficient for a partial
austenitization and a certain expansion of the cementite
into ferritic regions to take place. It was due to the fact
that, at lower heating rates, transformation start tempera-
tures become lower. A partial austenitization at 760 °C is
evidenced by the temperature record as well. The tem-
perature records for all the schedules with lower heating
rates (schedules 7, 8 and 9) showed a temperature pla-
teau on the cooling curve, indicating a pearlitic transfor-
mation. Temperature cycling, thus, leads to a partial
austenitization accompanied by cementite dissolution.
The cementite lamellae underwent a partial fragmenta-
tion and the structure underwent a grain refinement.
Figures 5a and b illustrate a comparison between the
scanning electron micrographs of the initial state and the
heat-treated sample 9. The heat-treated specimen under-
went a grain refinement and a partial carbide spheroidi-
sation (ASR).
3.2 Mechanical Properties
Tension tests were performed on the specimens with
the diameter of 8 mm and length of 50 mm. The
specimens used for impact testing were of the miniature
Charpy type, with the size of 3 mm × 4 mm × 27 mm
and a 1 mm deep V-notch. Hardness values were measu-
red on the Vickers HV30 scale.
In Table 3 you can see the results of proof stress
(PS), ultimate tensile strength (UTS), elongation (A5),
area reduction (RA), impact toughness (KCV) and
Vickers hardness (HV). The results for the initial state
were: PS = 309 MPa, UTS = 642 MPa, elongation A5 =
27 %, RA = 47 %, impact toughness KCV = 37 J/cm2 and
hardness reached 179 HV.
For schedules 1, 2 and 3, the proof stress increased
with the growing temperature of the cycles. Schedules 1
and 2 led to a significant increase in the toughness that
could have been caused by a recovery of ferrite grains.
Schedule 3 did not cause a significant increase in the
notch toughness, probably due to a formation of new fine
lamellar pearlite.
The specimens treated according to schedule 4
showed a higher proof stress, a slightly higher ultimate
strength and a higher toughness. This was due to the
grain refinement. Schedule 5 led to a similar enhance-
ment of the mechanical properties as schedule 4 but
resulting in the highest proof stress and in a higher
toughness. The specimens of schedule 6 did not have a
higher toughness than the initial material. This might be
attributed to the formation of a large amount of new
lamellar pearlite during the final cycling at around 780
°C and 720 °C.
A trend in the proof-stress values was apparent in the
specimens treated with the schedules involving the
heating rate of 20 °C/s or higher heating rates. The proof
stress rose upon cycling at around a higher temperature,
i.e., 780 °C. The specimens that were austenitized and
subjected to cycling at around 780 °C exhibited an even
higher proof stress: 423 MPa. Their ultimate strength
was slightly lower or equal to that of the initial material.
However, elongation as well as impact toughness
increased for all three specimens. The highest impact-
toughness and hardness values were found in the speci-
mens for schedule 9. In this case, the impact toughness is
higher than that of the specimen treated with the
schedule with the highest heating rate. This might be
attributed to the lower amount of new cementite
lamellae. Despite that the difference in the mechanical
properties of the specimens heated at higher and lower
rates was not substantial. Hence, no significant impact of
the heating rate on the mechanical properties of the
material was found.
4 CONCLUSION
Grain refinement and partial carbide spheroidisation
were achieved in the AISI C45 steel using induction
heating. The temperature cycling at around the Ac1 tem-
perature was responsible for the change in the cementite
morphology. Ferritic grain size was reduced by the
austenitization and the following cooling in air. These
changes in the microstructure were reflected in an
increase in the proof stress, ultimate tensile strength and
impact toughness, whereas the elongation value
remained unchanged. The proof stress was increased by
120 MPa, the ultimate tensile strength by 20 MPa and
D. HAUSEROVA et al.: ACCELERATED CARBIDE SPHEROIDISATION AND REFINEMENT (ASR) ...
704 Materiali in tehnologije / Materials and technology 47 (2013) 6, 701–705
Table 3: Mechanical properties
Tabela 3: Mehanske lastnosti
Sche-
dule
PS/
MPa
UTS/
MPa A5/% RA/%
KCV/
(J cm–2) HV
OS 309 642 27 47 37 179
1 381 645 29 48 99 165
2 387 645 30 49 93 167
3 390 633 32 57 41 169
4 417 653 30 56 52 174
5 431 660 29 55 58 178
6 426 643 32 61 38 174
7 356 613 32 60 61 169
8 403 635 31 58 58 178
9 423 638 32 60 62 180
the impact toughness by 20 J/cm2 in comparison with the
initial state. The schedule that led to this improvement in
the mechanical properties took about 300 s and consisted
of the normalizing and subsequent cycling at around the
Ac1 transformation temperature.
Acknowledgment
The results presented in this paper were obtained
within the project West-Bohemian Centre of Materials
and Metallurgy CZ.1.05/2.1.00/03.0077, co-funded by
the European Regional Development Fund.
5 REFERENCES
1 L. Kraus, S. Nemecek, J. Kasl, Mechanical properties of cast Mn
steel after intercritical heat treatment and microalloying, Trans-
ferability of fracture mechanical characteristics, 78 (2002), 15–32
2 B. Masek, H. Jirkova, L. Kucerova, Rapid Spheroidization and Grain
Refinement Caused by Thermomechanical Treatment for Plain Struc-
tural Steel, Materials Science Forum, 706–709 (2012), 2770–2775
3 A. Kamygabi-Gol, M. Seikh-Amiri, Spheroidising Kinetics and Opti-
mization of Heat Treatment Parameters in CK60 Steel Using Taguchi
Robust Design, Journal of Iron and Steel, 17 (2010), 45–52
4 P. R. Howell, The Pearlite Reaction in Steels: Mechanisms and Cry-
stallography, Materials Characterization, 40 (1998), 227–260
D. HAUSEROVA et al.: ACCELERATED CARBIDE SPHEROIDISATION AND REFINEMENT (ASR) ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 701–705 705

U. [EGEDIN et al.: THE EFFECT OF A STYRENE-ACRYLIC COPOLYMER SYNTHESIZED ...
THE EFFECT OF A STYRENE-ACRYLIC COPOLYMER
SYNTHESIZED IN A HIGH-PRESSURE REACTOR ON
THE IMPROVED CORROSION PROTECTION OF A
TWO-COMPONENT POLYURETHANE COATING
VPLIV STIREN-AKRILNEGA KOPOLIMERA, SINTETIZIRANEGA
V VISOKOTLA^NEM REAKTORJU NA IZBOLJ[ANJE
KOROZIJSKE ZA[^ITE DVOKOMPONENTNEGA
POLIURETANSKEGA PREMAZA
Urban [egedin1, Klemen Burja2, Franci Malin1, Sa{a Skale1, Bogdan Znoj1,
Boris [ket3, Peter Venturini1
1Central R&D of Helios Group, Helios Dom`ale, d. d., Koli~evo 2, 1230 Dom`ale, Slovenia
2Centre of Excellence for Polymeric Materials and Technologies, Tehnolo{ki park 24, 1000 Ljubljana, Slovenia
3Faculty of Chemistry and Chemical Technology, University of Ljubljana, A{ker~eva 5, 1000 Ljubljana, Slovenia
urban.segedin@helios.si
Prejem rokopisa – received: 2012-09-28; sprejem za objavo – accepted for publication: 2013-04-02
In the present research the synthesis of a high-solids styrene-acrylic copolymer in a high-pressure reactor is described and a
comparison of its properties with those of a copolymer synthesized at atmospheric pressure are given. The described procedure
enables the synthesis of styrene-acrylic copolymers with a higher solids content under elevated pressure in a significantly
shorter time, when compared to the process conducted at atmospheric pressure. The elevated pressure and higher temperature of
the synthesis result in a copolymer with a lower average molecular mass and a narrower molecular mass distribution as well as a
lower viscosity of the solution with an equal dry weight. Synthesized copolymers were used for the preparation of the two-
component polyurethane coatings for corrosion protection. Cross-linked films were formed with a hexamethylene diisocianate
homopolymer. With electrochemical impedance spectroscopy measurements a higher level of cross-linking and a lower porosity
of the two-component polyurethane coating based on styrene-acrylic copolymer synthesized in a high-pressure reactor have
been determined. This implies a better anti-corrosion protection of the metal substrates.
Keywords: styrene-acrylic copolymer, high-pressure synthesis, 2K-PUR coatings, electrochemical impedance spectroscopy
V raziskavi je opisan postopek sinteze stiren-akrilnega kopolimera s pove~ano vsebnostjo suhe snovi pri povi{anem tlaku, kar
vodi do skraj{anja ~asa sinteze v primerjavi s postopkom pri atmosferskem tlaku. Podana je primerjava lastnosti kopolimera,
sintetiziranega pri povi{anem tlaku, in kopolimera, sintetiziranega pri atmosferskem tlaku. Sinteza pri povi{anem tlaku in vi{ji
temperaturi vodi do manj{e povpre~ne molekulske mase, o`je porazdelitve molekulskih mas ter do ni`je viskoznosti raztopine
kopolimera pri enaki suhi snovi. Sintetizirana produkta sta bila uporabljena za pripravo dvokomponentnih poliuretanskih
premazov za protikorozijsko za{~ito. Z uporabo homopolimera heksametilen diizocianata so bili pripravljeni zamre`eni
premazi. Z elektrokemijsko impedan~no spektroskopijo je bilo ugotovljeno, da je stopnja zamre`enosti premaza s
stiren-akrilnim kopolimerom, sintetiziranim pri povi{anem tlaku, vi{ja, poroznost pa ni`ja. To ka`e na bolj{o protikorozijsko
za{~ito kovinskih podlag.
Klju~ne besede: stiren-akrilni kopolimer, sinteza pri visokem tlaku, 2K-PUR premazi, elektrokemijska impedan~na spektro-
skopija
1 INTRODUCTION
Styrene acrylic copolymers (SACs) are widely used
in the coatings industry as one of the main components
of coating formulations. SACs are part of one-com-
ponent acrylic as well as two-component polyurethane
coatings (2K-PUR). In the latter the SACs represent
component A (polyol) that reacts with component B
(polyisocyanate) to form polyurethane.1 The main
reaction occurs between the hydroxyl groups of the
polyol and the isocyanate groups of the polyisocyanate
to give urethanes. Other reactions, leading to the forma-
tion of allophanate, amine, urea, biuret and amide link-
ages1–6 are shown in Figure 1. The amount of hydroxyl
functional (meth)acrylic monomers in the SAC formula-
tion determines the degree of cross-linking, the porosity
and the hardness of the cross-linked 2K-PUR films.
When formulating the SAC, at least one hydroxyl func-
tional (meth)acrylic monomer is used. The rest of the
monomers are used to control the glass-transition tem-
perature (Tg) and the viscosity or to give additional
functionality to the SAC. Styrene is commonly used in
formulations for the adjustment of Tg, hardness and the
rate of physical drying. In some cases, special monomers
like monoglycidyl modified esters, ethylene, propylene
and alphaolefins with 7–20 carbon atoms are used to
achieve higher solids and lower viscosities for the SAC
solutions.2–10
Syntheses are usually performed at the reflux tempe-
rature of the organic solvent or the solvent mixture. The
possibility of conducting reactions at elevated pressure is
Materiali in tehnologije / Materials and technology 47 (2013) 6, 707–714 707
UDK 620.197:678.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)707(2013)
mentioned in literature, but the syntheses are limited to
special, highly volatile or gaseous monomers and no
comparison with atmospheric-pressure synthesis has
been reported.8,9,11,12
The effects of the process parameters on the pro-
perties of SAC are also well known.13 The increase of the
initiator concentration in the reaction mixture results in a
lower average molecular mass (Mw) and faster polymeri-
zation rates. The increase in the temperature increases
the polymerization rate as well. Higher reaction tempe-
ratures reduce the Mw, the molecular-mass distribution
(Mw/Mn) and the viscosity.13
Solutions of SAC in organic solvents are being
replaced by high-solids SACs (HS SACs) as a result of
European legislation on volatile organic compounds
(VOCs) in coatings (Directive 2004/42/CE). One way is
to synthesize the SACs with a lower and less-dispersed
Mw. This can be achieved by increasing the initiator
concentration, slower dosing, raising the reaction tempe-
rature and using a chain-transfer reagent. In most cases
the synthesis is performed in the excess solvent and
distillation is needed to reach the desired dry weight. In
some cases the solvent is distilled off and replaced with a
more suitable solvent, and as a result lower solution
viscosities are obtained.1,2,8–13
Electrochemical impedance spectroscopy (EIS) is a
powerful tool, not only to monitor the degradation of
coatings, but also a rapid non-destructive method to
assess the quality of protective performance prior to the
development of corrosion.14–16 In the early stages of the
coating degradation the coating’s capacitance, the War-
burg coefficient and the pore resistance are present. The
measurements are fitted with equivalent circuits (trans-
mission line model), which describe the charge-transfer
mechanisms through the coatings (Figure 2). The coat-
ing capacitance (Cc) represents the water uptake in the
coatings and corresponds to the charge separation on the
protective film. The water accumulation on the phase
boundaries between the coating and the steel substrate
enables ion diffusion through the film matrix. This
mechanism of charge transfer is represented by the War-
burg coefficient (c). Classical ion conduction through
pores filled with an electrolyte solution is represented by
the pore resistance (Rpo).14,15,17–27 EIS measurements are
usually conducted in the frequency range from 10–2 Hz to
105 Hz. The results are presented as Bode and Nyquist
diagrams. The high quality of 2K-PUR coatings is repre-
sented with phase values of over 80° over a wide range
of frequencies and the impedance values are mainly
determined by the coating’s capacitance. The pure capa-
citor has impedance values over 1010  cm2, displaying
superb barrier properties. With the development of pores
phase and absolute impedance values decrease at lower
frequencies. Typical impedance values in this case are
around 107–108  cm2. The impedance values of the
coatings displaying entire loss of corrosion protection
are around 104  cm2.16,27
The aim of the present work is to give a comparison
of the polymerization processes at atmospheric and ele-
vated pressure. Random copolymers were synthesized
via free-radical polymerization in a semi-batch process.
A comparison of the properties of both end-products is
presented as well. It is shown that besides a shorter reac-
tion time the quality of cross-linked 2K-PUR films of
U. [EGEDIN et al.: THE EFFECT OF A STYRENE-ACRYLIC COPOLYMER SYNTHESIZED ...
708 Materiali in tehnologije / Materials and technology 47 (2013) 6, 707–714
Figure 2: Charge-transfer mechanisms through the coating film
during EIS measurements for the porous film prior to the development
of corrosion17
Slika 2: Mehanizmi prenosa naboja preko premaza pri EIS-meritvah
za plast z razvitimi porami pred nastankom korozije17
Figure 1: Formation of urethane group and other reactions present
during the cross-linking of SAC and polyisocyanate1,3
Slika 1: Nastanek uretanske vezi ter druge reakcije pri zamre`evanju
SAC in poliizocianata1,3
SAC synthesized in a high-pressure reactor is also im-
proved.
2 EXPERIMENTAL
2.1 Materials
Acrylic acid (AA, Clariant), butyl acrylate (BA,
BASF), 2-hydroxypropyl methacrylate (HPMA, Pro-
chema) and styrene (S, Enichem) monomers as well as
the heat-sensitive free-radical initiator t-butyl perben-
zoate (t-BPB, t1/2 = 10 h, T = 104 °C, Dow) were used for
the syntheses of the SACs. Methoxypropyl acetate
(MPA, BASF) served as an organic solvent. Hexame-
thylene diisocyanate homopolymer (HDI homopolymer;
Desmodur N75 (D-N75, Bayer)) was used for the pre-
paration of the 2K-PUR coatings.
2.2 Preparation of a monomer mixture for the synthe-
ses of SACs
For the purpose of the present study two styrene-
acrylic copolymers suitable for the preparation of the
2K-PUR coatings were synthesized. Both were prepared
from the same monomer mixture containing AA, BA,
HPMA and S. All the monomers were industrial grade
with hydroquinone monomethyl ether (MEHQ) as a
stabilizer and were used as delivered without purifi-
cation. The initiator (t-BPB) was added and the mixture
was stirred under mild shear conditions. The compo-
sition of the final mixture prepared for dosing is given in
Table 1.
Table 1: Monomer mixture prepared for syntheses of SACs (mass
fraction: w/%; mol fraction: x/%)
Tabela 1: Me{anica monomera, pripravljena za sintezi SAC (masni
dele`: w/%; molski dele`: x/%)
w/% x/%
AA 0.8 1.31
BA 30.7 28.36
S 45.5 51.72
HPMA 21.7 17.82
t-BPB 1.3 0.79
 100.0 100.00
2.3 Syntheses of SACs
Both syntheses were carried out in a semi-batch
mode on a reaction calorimeter RC1eTM (Mettler Toledo)
with iControl software for monitoring and controlling
the process parameters.
2.3.1 The synthesis of SAC 1
SAC 1 was synthesized in a 2-L triple-wall glass
reactor AP01-2-RTC with temperature sensors, calibra-
tion heater, a 4-bladed glass pitch blade stirrer, glass
condenser, N2 purge and a ReactIR 45 m Probe A with
an MCT Detector using HappGenzel apodization, a
DiComp (diamond) probe connected via a K6 Conduit
(16 mm probe) with a sampling area of 4000 cm–1 to 650
cm–1 at 8 cm–1 wavelength resolution and a sampling
interval of 30 s during the synthesis. Online data from
FTIR (Fourier Transform Infrared Spectroscopy) measu-
rements was collected and analyzed with iC IR software
(Mettler Toledo) in real time. The reactor was filled with
400 g of MPA and heated to 135 °C at atmospheric
pressure (below the boiling point of MPA at atmospheric
pressure, 145 °C) with a stirring rate of 350 r/min. When
the working temperature was reached, 900 g of monomer
mixture with an initiator was continuously dosed into the
reactor for 3 h using a Prominent Beta/4 solenoid dia-
phragm metering pump. After the dosing was finished
the reaction mixture was kept at the working temperature
for an additional 2.5 h.
2.3.2 The synthesis of SAC 2
SAC 2 was synthesized in a 1.8 L stainless-steel high-
pressure reactor HP60 at a temperature of 175 °C and
elevated pressure. The reactor was equipped with tem-
perature sensors, a pressure sensor, a calibration heater, a
stainless-steel anchor stirrer, a N2 purge and an online
FTIR measuring system (as described in the synthesis of
SAC 1). The tightly closed reactor was filled with 300 g
of MPA and 78 g of monomer mixture with an initiator
and heated to 175 °C at 280 kPa of absolute pressure
with a stirring rate of 200 r/min. When the working tem-
perature was reached, 668 g of monomer mixture with an
initiator was continuously dosed into the reactor for 90
min using a Prominent Beta/4 pump. After the dosing
was finished the absolute pressure reached 550 kPa and
the reaction mixture was kept at the working temperature
for an additional 20 min, after which 15 g of the mass
fraction 10 % t-BPB in MPA was dosed in 15 min for
residual monomer depletion. After an additional 20 min
heating at the working temperature and 570 kPa of
absolute pressure the reaction mixture was cooled down
and the pressure released.
2.4 Characterization of SACs
The chemical and physical properties of the products
were evaluated through standard test methods. Acid-
value (AV) measurements were carried out in com-
pliance with the standard EN-ISO-2114, hydroxyl value
(OH, EN-ISO-4629), dry weight of copolymer solution
in MPA (DW, ISO-3251) and viscosity (, EN-ISO-
3219). The viscosity measurements were performed at
23 °C on a rotational viscometer (Haake Viscotester 550)
with a coaxial cylinder sensor system.
2.4.1 Nuclear magnetic resonance (NMR)
measurements
The composition of both copolymers and coatings
was evaluated via 1H and 13C NMR measurements. 1D
1H and 13C NMR spectra of the SACs, D-N75 and
2K-PURs were recorded at 25 °C on a Unity Inova 300
MHz NMR spectrometer (Agilent Technologies). The 1H
U. [EGEDIN et al.: THE EFFECT OF A STYRENE-ACRYLIC COPOLYMER SYNTHESIZED ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 707–714 709
and 13C chemical shifts were referred to the residual
signal of DMSO-d6 at  = 2.50 · 10–6 for 1H and  =
39.50 · 10–6 for 13C.
2.4.2 Rheological properties
The rheological properties of the SAC solutions (as
synthesized and diluted to the mass fraction 50 %) were
assessed using a Physica MCR301 rheometer (Anton
Paar GmbH). A cone-plate geometry sensor system (R =
25 mm, gap = 0.207 mm) was used to perform rotational
tests where the samples were exposed to increasing shear
rates (0.1 s–1 to 670 s–1) in the first run and to decreasing
shear rates (670 s–1 to 0.1 s–1) in the second. In both
cases, 29 measuring points on a logarithmic scale were
taken.
2.4.3 DSC evaluations
The Tg of dried SACs was determined by differential
scanning calorimetry (DSC) on a DSC 1 STARe System
calorimeter (Mettler Toledo). The SAC samples were
heated twice at 20 K/min in the temperature range from
–50 °C to 200 °C with a nitrogen purge of 20 mL/min.
The second heating run was used for the Tg determina-
tion. For each sample at least three distinct measure-
ments were performed. Average values were then
calculated and given as a result.
2.4.4 Gas chromatography – mass spectroscopy
measurements
The amount of residual monomers in the SACs was
determined with coupled gas chromatography – mass
spectroscopy (GC-MS). A 6890N GC Network System
with a HP-5 column with polysiloxane-based column
packing (30 m in length, I. D. 0.32 mm, film 0.25 μm)
and 5973 Network Mass Selective Detector (Agilent
Technologies) were used. The sample solution concen-
tration in methylenechloride was 0.1 g g–1 and heptane
was added as an internal standard. Helium was used as
an inert carrier gas with a constant flow rate of 0.7 mL
min–1. Standard monomer solutions with known concen-
trations were used for the quantitative analyses. The area
under each peak was used to determine the quantity of
the specific compound.
2.4.5 Size-exclusion chromatography
Mw and Mw/Mn were measured using size-exclusion
chromatography (SEC) on a Waters 2690 chromatograph
with a Waters 410 refractive-index detector (Waters
S.A.S.). Tetrahydrofuran (THF) was used as an eluent
with a flow rate of 0.7 mL min–1. Three 7.8 × 300 mm
single pore-sized columns with a styrene divinylbenzene
copolymer column packing and a particle size of 5 μm
were used in a sequence - Styragel HR 0.5 column with
an effective molecular mass range of 0–103, Styragel HR
4E (50–105) and Styragel HR 5E (2 · 103 – 4 · 106). A
sample solution concentration in THF was 2 mg mL–1.
Before the analysis the samples were filtered through a
0.45 μm PTFE filter and a 100 μL of prepared sample
was injected onto the column. As a result, the average
molecular masses relative to polystyrene standards are
given.
2.5 Preparation of cross-linked 2K-PUR coating films
Cross-linked 2K-PUR coating films were prepared
with a HDI-homopolymer (D-N75). Diluted SAC (50 %)
and D-N75 were mixed with a propeller stirrer (dry
weight mass ratio 3.2 : 1) at 1500 r/min for 15 min,
applied on 20 cm × 7.5 cm steel plates with an applicator
and left to dry for 48 h. The dry film thickness of each
2K-PUR film was measured using the Elcometer 456
with the F1 magnetic induction probe (Elcometer Ltd.).
For each film at least 20 measurements were performed,
in compliance with the standard ISO 2178.
2.6 Electrochemical impedance spectroscopy
The electrochemical properties of the cross-linked
coating films were measured via EIS at 20 °C. The
measuring (Tait) cell for the EIS contained a working
electrode (coated sample plate) with a 34.21 cm2 area,
immersed in a NaCl solution 0.1M, a hastelloy counter
electrode and a standard calomel reference electrode
(SCE). The potential difference between the working and
the counter electrode versus SCE was -0.6 V and the
amplitude of the signal was 30 mV. The measurements
were performed with a Parstat 2273 Advanced Electro-
chemical System (Advanced Measurement Technology
Inc.) and the data analyzed using the Electrochemistry
Power Suite and PowerSINE software (Princeton
Applied Research). The value of the frequency was in
the range from 65536 Hz to 0.1 Hz.
3 RESULTS AND DISCUSSION
3.1 Syntheses of the SACs
Both syntheses described above were previously
optimized in terms of dosing rate, initiator concentration
and working temperature to give products suitable for the
application of 2K-PUR coatings (evaluated by the
amount of residual monomers in the final product, Table
2). The synthesis of SAC 1 was carried out at 135 °C to
ensure constant conditions. The temperature for the syn-
thesis of SAC 2 was carefully selected at 175 °C.
According to the literature data, the coloration of the
product progresses with temperature. In some cases 180
°C can already be too high13. The pressure in the syn-
thesis of the SAC 2 was raised due to the evaporation of
the solvent and the monomers as well as inflow of a
fresh monomer mixture. This caused the elevation of the
boiling point of the reaction mixture. Previous syntheses
(not shown) indicated that an additional initiator is
needed to reach a sufficient monomer conversion – a
lower final residual monomer content (Table 2). Per-
forming the synthesis at 175 °C significantly reduces the
U. [EGEDIN et al.: THE EFFECT OF A STYRENE-ACRYLIC COPOLYMER SYNTHESIZED ...
710 Materiali in tehnologije / Materials and technology 47 (2013) 6, 707–714
reaction times in comparison to a synthesis at 135 °C
(2 h 25 min versus 5 h 30 min). Higher reaction tempera-
tures led to a lower Mw and a narrower Mw/Mn 13. This
reduces the viscosity of the system and allows for the
preparation of high-solids SACs (in our case SAC 2,
Table 2).
Table 2: Properties of synthesized SACs
Tabela 2: Lastnosti sintetiziranih SAC
SAC 1 SAC 2
DW [w/%] 62.4 71.3
/(mPa s) (23 °C, 0.1 s–1) 8200 15500
Mw 23000 9000
Mw /Mn * 2.6 2.3
Tg/°C ** 32 ± 2 23 ± 2
AV [mg KOH / g dry copolymer] 9.3 9.3
OH [mg KOH / g dry copolymer] 110 110
Residual monomers [w/%] 0.5 0.4
* ratio between mass average (Mw) and number average molecular
mass (Mn)
** Tg, Fox = 33.3 °C
Figure 3 shows a comparison of both syntheses
through thermal conversion. The reaction calorimeter
measured the heat released due to the polymerization (qr)
in real time. From this data the thermal conversion was
calculated as the ratio between the heat release until a
certain time (tr) and the total heat release from the
reaction (ttot stands for the total reaction time):
X
q t t
q t t
t
tth
r
r
d
d
r
tot
=
∫
∫
( )
( )
0
0
(1)
It should be stressed that a 100 % thermal conversion
does not imply total monomer consumption. For a better
comparison the x-axis in Figure 3 is not absolute time,
but rather the reaction coordinate (	), defined as the
progress of syntheses (	 = 0 meaning start of dosing,
	 = 1 meaning the end of synthesis – the beginning of
cooling the reaction mixture down). Figure 3 shows that
in the synthesis of SAC 2, a 98.2 % thermal conversion
was reached after the dosing was completed (at 	 =
0.56), whereas in the synthesis of SAC 1 only a 74.4 %
thermal conversion was achieved (at 	 = 0.58). In terms
of heat removal, only 1.8 % of all the reaction heat
needed to be removed after the dosing in the synthesis of
SAC 2. On the other hand, in the synthesis of SAC 1,
25.6 % of all the reaction heat needed to be removed
after the dosing was completed. In terms of heat release
controlled by dosing, the synthesis in a high-pressure
reactor turns out to be safer than the synthesis at a
temperature below the boiling point of the solvent.
3.2 Characterization of the SACs
The described syntheses performed on the RC1eTM
reaction calorimeter were made in parallel. There were
no discrepancies between the parallels in terms of
U. [EGEDIN et al.: THE EFFECT OF A STYRENE-ACRYLIC COPOLYMER SYNTHESIZED ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 707–714 711
Figure 4: Comparison of 1H NMR spectra of SACs and 2K-PURs; * denotes signals of solvents MPA and DMSO at 2.5 · 10–6
Slika 4: Primerjava 1H-spektrov SAC in 2K-PUR; * ozna~uje signale topil MPA in DMSO pri 2,5 · 10–6
Figure 3: Thermal conversion for polymerizations of SACs
Slika 3: Termi~na konverzija pri polimerizacijah SAC
process parameters due to the excellent process control.
Therefore, the syntheses were highly reproducible. This
resulted in matching properties of the products such as
Mw, Mw/Mn, residual monomers and viscosity. The results
are not shown due to the negligible differences.
Furthermore, analyses of the acid and hydroxyl
values (Table 2) confirmed the same functionality of
both SACs. Since the on-line FTIR data provided insuffi-
cient information on the SACs’ composition, NMR
measurements were performed. The 1H and 13C NMR
spectra of the dry SACs and 2K-PURs show no signifi-
cant differences in functionality. As can be seen in Fig-
ure 4 there are no peaks in the 5–6 · 10–6 range of the 1H
NMR spectra of the SACs that are assigned to the
protons around the double bonds of the (meth)acrylic
monomers. Also, there is no peak at 112 · 10–6 in the 13C
NMR spectra of the SACs assigned to vinyl =CH2 of
styrene (Figure 5). This supports the GC-MS measure-
ments of the residual monomers and implies that the
monomer mixture was sufficiently converted to the copo-
lymer.
The HDI homopolymer 1H NMR spectrum (Figure
6) shows peaks at 8.3 · 10–6 (-NH-CO-, biuret H), 3–3.5 ·
10–6 (-CH2-N=, HDI) and 1–2 · 10–6 (-CH2-, HDI). The
typical peaks in the 13C NMR spectrum are at 155 · 10–6
(-NH-CO-N, biuret), 121 · 10–6 (-NCO, isocyanate) and
25–45 · 10–6 (-CH2-, HDI). This confirms the structure of
the HDI homopolymer (D-N75). The HDI monomers
react with water to form trimeres with a biuret linkage.
These trimeres further react to form the HDI-homopoly-
mer.
In the preparation of the 2K-PURs the depletion of
the –NCO group in the 13C NMR spectra can be seen
(Figure 5, SAC 1 PUR and SAC 2 PUR), whereas the
peaks at 8.3 (1H NMR) and 155 · 10–6 (13C NMR)
assigned to the biuret remain visible. The urethane
groups are overlapping with the styrene aromatic ring
(7 · 10–6 and 7–7.8 · 10–6 in 1H NMR spectra, Figure 4).
Small amounts of urea produced by side reactions can be
seen from the 1H NMR spectra at 5.7 · 10–6 28–31. The
U. [EGEDIN et al.: THE EFFECT OF A STYRENE-ACRYLIC COPOLYMER SYNTHESIZED ...
712 Materiali in tehnologije / Materials and technology 47 (2013) 6, 707–714
Figure 5: Comparison of 13C NMR spectra of SACs and 2K-PURs; * denotes signals of solvent MPA
Slika 5: Primerjava 13C-spektrov SAC in 2K-PUR; * ozna~uje signale topila MPA
Figure 6: 1H and 13C NMR spectra of HDI homopolymer with trimer
structure
Slika 6: 1H- in 13C NMR-spektra HDI homopolimera s strukturo
trimera
NMR results show the similar composition of both SACs
and of both 2K-PUR films.
As Table 2 shows, a high solids SAC was synthe-
sized (DW of SAC 2 > 70 %). A higher temperature in
the synthesis of the SAC 2 led to a lower Mw (9000
versus 23000) and a narrower Mw/Mn (2.3 versus 2.6).
The shorter copolymer chains result in a decrease of the
SAC 2 Tg in comparison to the SAC 1 (23 °C versus 32
°C), which complies with the literature data13. Both
SACs have a Tg lower than that predicted by the Fox
equation (33 °C). The viscosity of the SAC 2 is higher
due to the higher solids content (15500 mPa s versus
8200 mPa s, Table 2, Figure 7). After the dilution of
both SACs with additional solvent to 50 % the effect of
shorter copolymer chains can be observed as lower
viscosity values (160 mPa s versus 570 mPa s, Figure 7).
Both products were suitable for the preparation of
2K-PUR coatings.
The results of the EIS measurements and the
evaluations for cross-linked 2K-PUR films of SAC 1 and
SAC 2 are shown in Figure 8 (Bode diagrams) and Fig-
ure 9 (Nyquist diagram) and the fitting model para-
meters are given in Table 3. The experimental data were
fitted by equivalent circuits – the transmission-line
model16,19,27 – indicating a better quality of cross-linked
2K-PUR film prepared with the SAC 2. The trans-
mission-line model for the SAC 2 film has no pore
resistance (Rpo > 109  cm–2, Figure 9), indicating that
no pores are developed. The lower Warburg coefficient
(c) of the SAC 2 represents a better cross-linking of the
shorter SAC chains and a lower coating capacitance (Cc)
represents a lower water uptake of the cross-linked
2K-PUR film prepared with the SAC 2 in comparison to
the one prepared with SAC 1.12,16,17,27 The impedance
values shown in the Bode and Nyquist diagrams (Fig-
ures 8 and 9) for cross-linked 2K-PUR films prepared
with the SAC 1 and SAC 2 in combination with compa-
rable film thicknesses (Table 3) of both films directly
indicate the better quality of the SAC 2.16,17,27 The impe-
dance values at low frequencies are above 107  for both
films and the improved quality of the 2K-PUR film
prepared with the SAC 2 can be seen. The generally
accepted absolute value of the impedance for the
U. [EGEDIN et al.: THE EFFECT OF A STYRENE-ACRYLIC COPOLYMER SYNTHESIZED ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 707–714 713
Figure 9: Nyquist diagram for cross-linked 2K-PUR films (SAC 1 and
SAC 2) with equivalent circuits used for fitting model (RS not shown);

 arrow denotes the increase in angular frequency
Slika 9: Nyquistov diagram zamre`enih 2K-PUR premazov (SAC 1 in
SAC 2) z odgovarjajo~ima nadomestnima tokokrogoma (RS ni
prikazan); 
 pu{~ica ka`e v smeri nara{~anja kotne frekvence
Figure 8: Bode diagrams for cross-linked 2K-PUR films; lines repre-
sent the fitting model
Slika 8: Bodejeva diagrama za zamre`ene 2K-PUR premaze; ~rte
prikazujejo prileganje modela
Table 3: Electrochemical properties of cross-linked 2K-PUR films
Tabela 3: Elektrokemijske lastnosti zamre`enih 2K-PUR premazov
SAC 1 SAC 2
dc/μm 75 ± 3 65 ± 3
c/(–1 rad–1/2 s1/2 cm2) 4.2 × 10–8 1.3 × 10–8
Rpo/( cm–2) 3.7 × 107 ND
Cc/(F cm–2) 6.1 × 10–9 1.7 × 10–9
dc = thickness of cross-linked 2K-PUR film
c = Warburg coefficient of film
Rpo = pore resistance
Cc = capacitance of cross-linked 2K-PUR film
Figure 7: Flow curves of original (closed symbol) and diluted SACs
(open symbol)
Slika 7: Tokovne krivulje osnovnih (~rne oznake) in razred~enih SAC
(bele oznake)
protective coatings is at least 106 .32 This means that
both SACs are suitable for the preparation of 2K-PUR
coatings.
4 CONCLUSIONS
Synthesis in a high-pressure reactor under high tem-
perature and elevated pressure proved to be a better
alternative to the synthesis of high-solids SACs in a
solution of an organic solvent at atmospheric pressure.
The proposed direct synthesis leads to a product with no
need of subsequent solvent removal by distillation and
compliant with the European legislation on volatile orga-
nic compounds in coatings. Apart from the significant
cutoff in the reaction times a much higher thermal
conversion is reached while the monomer mixture is still
being dosed when compared to the synthesis below the
reflux temperature of the solvent. The cross-linked
2K-PUR coating prepared with the SAC 2 shows better
anti-corrosion properties in comparison to the SAC 1,
making it suitable for the protection of metal substrates.
Acknowledgements
Operation was part financed by the European Union,
European Social Fund. Operation implemented in the
framework of the Operational Program for Human
Resources Development for the Period 2007-2013, Pri-
ority axis 1: Promoting entrepreneurship and adaptabi-
lity, Main type of activity 1.1.: Experts and researchers
for competitive enterprises.
The authors acknowledge the support from CE Poli-
MaT (financially supported by the Ministry of Educa-
tion, Science, Culture and Sport of the Republic of Slo-
venia through the contract No. 3211-10-000057 (Centre
of Excellence for Polymer Materials and Technologies))
where the syntheses on RC1eTM were performed.
5 REFERENCES
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1045–1055
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Corrosion Science, 52 (2010), 37–44
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Materials Letters, 3 (2012), 161–171
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1287–1309
U. [EGEDIN et al.: THE EFFECT OF A STYRENE-ACRYLIC COPOLYMER SYNTHESIZED ...
714 Materiali in tehnologije / Materials and technology 47 (2013) 6, 707–714
B. VALENTAN et al.: PROCESSING POLY(ETHER ETHERKETONE) ON A 3D PRINTER ...
PROCESSING POLY(ETHER ETHERKETONE) ON A 3D
PRINTER FOR THERMOPLASTIC MODELLING
OBDELAVA POLYETHER ETHERKETONEA NA 3D-TISKALNIKU
ZA TERMOPLASTI^NO MODELIRANJE
Bogdan Valentan1, @iga Kadivnik2, Toma` Brajlih2, Andy Anderson3, Igor Drstven{ek1,2
1Ortotip, d. o. o., Ulica [kofa Maksimiljana Dr`e~nika 6, 2000 Maribor, Slovenia
2University of Maribor, Faculty of Mechanical Engineering, Smetanova ulica 17, 2000 Maribor, Slovenia
3Invibio Ltd.,Technology Centre-Hillhouse International, Thornton Cleveleys, Lancashire FY5 4QD, United Kingdom
valentan@gmail.com
Prejem rokopisa – received: 2012-10-12; sprejem za objavo – accepted for publication: 2013-03-25
PEEK, poly(ether etherketone), is one of the high-quality industrial polymers. It is widely used in extremely demanding areas
like automotive, aircraft and space industries. Because of the fact that it is bio-compatible, PEEK is also used for medical
implants that are usually made by milling a block of the material. The article presents the results of an investigation of
processing PEEK on a 3D printer for thermoplastic modelling. The used procedure is one of the additive manufacturing
procedures and, as such, it builds a product by adding material layer by layer to get the finished product. Commercially
available machines are unable to achieve the required melting and environment temperatures, so a new machine was developed.
The machine was designed and built at the company Ortotip d.o.o. and it is able to produce the parts of up to 130 mm × 130 mm
× 150 mm. After the initial testing, test specimens, according to standards EN ISO 527-2: 2012 and EN ISO 178: 2011, were
produced and tested at the facilities of the PEEK manufacturer Invibio (from the UK). The article presents the steps taken when
developing the PEEK modelling machine, the test methods to verify the mechanical properties of manufactured products and the
results of the material testing. The machine was developed to produce medical implants (specific maxillofacial prosthesis), but
with additional testing (that will help to improve the mechanical properties of produced parts) practically all bone-replacement
implants can be made.
Keywords: PEEK, 3D printer, implant, medical application, thermoplastic, FDM, biocompatible, additive technology
PEEK, polyether etherketon, je eden izmed visoko kvalitetnih industrijskih polimerov. Uporablja se na zahtevnih podro~jih, kot
so avtomobilska, letalska in vesoljska industrija. Zaradi dejstva, da je biokompatibilen, je uporaben tudi za medicinske vsadke,
ki so navadno narejeni s frezanjem iz bloka materiala. ^lanek predstavlja izsledke raziskave oblikovanja PEEK s
3D-tiskalnikom za termoplasti~no modeliranje. Uporabljen postopek spada med tako imenovane dodajalne tehnologije in kot tak
gradi izdelek po slojih. Komercialno dostopni stroji ne zmorejo dose~i zahtevanih procesnih in okoljskih temperatur, zato je bila
razvita nova naprava. Ta je bila oblikovana in izdelana v podjetju Ortotip, d. o. o., in je primerna za izdelke do velikosti 130 mm
× 130 mm × 150 mm. Po uvodnih preizkusih so bile izdelane preizkusne epruvete v skladu s standardoma EN ISO 527-2: 2012
in EN ISO 178: 2011 in preizku{ene od proizvajalca PEEK-materiala Invibio (iz VB). Predstavljeni so klju~ni koraki pri razvoju
naprave za direktno izdelavo modelov, preizkusne metode za verifikacijo mehanskih lastnosti izdelanih kosov in rezultati
meritev. Naprava je bila razvita za izdelavo medicinskih vsadkov (posebno lobanjske vsadke), vendar bo ob dodatnih preizkusih
(za izbolj{anje mehanskih lastnosti izdelanih kosov) primerna za vse vrste medicinskih vsadkov.
Klju~ne besede: PEEK, 3D-tiskalnik, vsadek, medicinska aplikacija, termoplast, FDM, biokompatibilnost, dodajalne tehnologije
1 INTRODUCTION
Today the majority of long-lasting human-body
implants are made from three materials. The first is the
bone cement, to technicians more known as poly(methyl
methacrylate) – PMMA. Its great advantage is that it can
be processed during the operation since it can be made of
two components (liquid and powder) that are mixed
together and directly processed (by hand or with some
basic tools) in a short time (the curing time can be as
short as 2 min). The production of the implants made
from solid PMMA is used only in special cases like the
lens for trabeculum. The second material is titanium
(actually one of the titanium alloys) that is processed
with a conventional machining process (usually the CNC
milling or turning) or with one of the additive manu-
facturing procedures like SLM (selective laser melting),1
MLSS (metal-laser-sintering system)2 or EBM (elec-
tron-beam melting).3 After a long-term use titanium
shows some problems,4 but it is irreplaceable when
excellent mechanical properties are needed. The material
that promises the most is PEEK – poly(ether etherketo-
ne). The implants from PEEK5–7 are usually made with
the conventional machining process or the method, pre-
sented in 2010, of direct manufacturing with an SLS
(selective laser sintering) machine.8 Since PEEK is ther-
moplastic it can be formed also with the other proce-
dures that are widely used for other thermoplastics.9 The
challenges are the specific requirements relating to a
higher viscosity (Figure 1) that need to be taken into
consideration, so the idea of developing a dedicated
device (a 3D printer for thermoplastic modelling) was
born.
2 BASIC PRINCIPLES OF THE 3D PRINTER
FOR THERMOPLASTIC MODELLING
The 3D printer for thermoplastic modelling is basi-
cally an FDM (fused-deposition modelling) machine. As
Materiali in tehnologije / Materials and technology 47 (2013) 6, 715–721 715
UDK 621.7:004.89 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)715(2013)
any additive manufacturing procedure (also known as the
rapid prototyping)10–14 it starts with a 3D CAD model
that is usually exported as an STL file from a CAD
(engineering or dedicated medical) program. The STL
file is then sliced by the computer software into
horizontal layers that are as high as the layer in the
machine. The rod-shaped filament is supplied to the
machine through a nozzle. The nozzle is computer
controlled in the X, Y plane and, in one layer, it forms a
raster of the item in the respective layer. The material is
liquefied in the nozzle and it hardens quickly when
applied to the layer at a lower temperature. The entire
system must be preheated to a higher surrounding
temperature. After the layer is finished, the working bed
in the Z direction is lowered by the thickness of one layer
and the new layer is extruded. When the geometry of a
part is more complex15 (overhangs), a support structure
must be added.
3 MACHINE DESIGN
The machine design was based on the 3D printer
Ciciprinter16 presented in 2011 and was made with the
CAD program Solid Works. In the prototypes the open-
source control electronics17 and the software made for
RepRap printers18 were used, with a simple adaptation,
to achieve the desired application. The basic principle of
thermoplastic modelling stays the same, but there are
completely different requirements for industrial or
medical machines, comparable with the distinction
between home-used or educational tools. For medical
equipment, standard ISO 13485 19 is valid and for
electronic circuits, several other guidelines (like
89/392/EEC, 73/23/EEC and 89/336/EEC) need to be
taken into consideration. PEEK, like most of the widely
used thermoplastic materials (ABS, HDPE, LDPE, etc.),
is not easy to process and in order to get the best results,
the manufacturer’s instructions9 must be followed. The
basic material properties were studied and discussed
with the material supplier to get the data of the pro-
cessing parameters for a small-nozzle extrusion.
At this point the decision was made to develop a
PEEK printer that uses a rod-shaped material rather than
a more classical approach with the grains and an
extrusion screw because of the required screw length, a
need for an expensive regulation equipment (pressure
measurements and electronic valves) and a requirement
for cleaning after the completion of each batch (which,
in the case of a unique or small-series production, means
practically every day).9
For the testing purposes the first machine with a
heating chamber was produced (Figure 2). The heaters
with a combined power of 1000 W were used to preheat
the working plate to around 240–300 °C. The first
problem was the implementation of a mobile platform
inside the chamber that required a mechanism for an
accurate movement in the Z direction with a minimum
loss of heat. The soft isolation and the central holder for
the platform were eventually used. The second problem
was how to assure the movement in the XY direction of
the nozzle and prevent the heat loss at the top of the
chamber. The known industrial solution with protective
bellows would be an elegant solution, but the working
B. VALENTAN et al.: PROCESSING POLY(ETHER ETHERKETONE) ON A 3D PRINTER ...
716 Materiali in tehnologije / Materials and technology 47 (2013) 6, 715–721
Figure 2: First laboratory PEEK machine. On the top there is an
extruder mechanism with the 1 : 5 thrust ratio from the small NEMA
17 step motor to the specially shaped stainless-steel thrust screw. The
PEEK rod diameter of 2.2 mm (± 0.1 mm) was routed from the nearby
roll. The maximum chamber temperature was only around 300 °C
because of the 1000 W heat source and a poorly applied insulation.
Slika 2: Prva preizkusna naprava za predelavo PEEK-a. Na vrhu je
mehanizem za podajanje materiala z razmerjem 1 : 5 in pogonskim
kora~nim motorjem NEMA 17. Pogon je izveden preko nazob~anega
pogonskega vijaka iz nerjavnega jekla. PEEK-`ica, premera 2,2 mm
(±0,1 mm), je bila speljana iz koluta v bli`ini. Zaradi grelnikov 1000
W in slabe izvedbe izolacije je bila najvi{ja dose`ena temperatura
komore samo okrog 300 °C.
Figure 1: Shear viscosity versus the temperature for some common,
industrial, engineering thermoplastics, comparable to PEEK Optima®
LT1 9
Slika 1: Viskoznost v odvisnosti od temperature za nekaj termoplastov
v {iroki uporabi v primerjavi z PEEK Optima® LT1 9
temperature was simply too high, so a movable insulated
panel was used. The design of the hot-end part, together
with the extrusion nozzle, required another special
solution. The hot end is the part where the melting
process takes place, so it needs to be insulated from the
hot-end holder to prevent the heat transfer through the
leading pipe that could cause the material melting in the
leading pipe and the subsequent blockage of the material
at the next start. Some tests were made with ceramic or
composite materials,20–22 but, at the end, the nozzle was
thermally separated from the holder with a thermal
barrier made on the leading pipe in the shape of narrow-
ing/ thinning walls.
The second machine (Figure 3) was designed with
all the updates that had been developed in practice after
the previous tests. The working chamber was insulated
with a thicker insulation and was more carefully closed
so that the heat losses were, subsequently, lower than
before. The building chamber was increased to allow a
more equal air heating and a bit larger building volume,
especially in the Z direction. With the use of more
powerful heaters, the preheating time needed to reach the
working temperature (now set at 280 °C to 300 °C) was
shortened from more than one hour to 20 min. The
extrusion tests also showed a need for a new hot-end
design, so the hot end with a bayonet mechanism for a
fast and simple exchange was made.
4 MATERIAL PROPERTIES
The PEEK supplier Invibio prepared two different
PEEK Optima® materials – LT3 and LT1. The properties
of both materials,23 for the case of injection moulding,
can be found in Table 1 and the shear viscosity versus
the shear stress at different temperatures is presented in
Figure 4. When used in the injection-moulding process,
LT3 has slightly higher mechanical properties and a
lower viscosity and as such it should be ideal for our
application.
Table 1: Mechanical properties of the PEEK Optima® LT1 and LT3
materials processed with injection moulding23
Tabela 1: Mehanske lastnosti PEEK Optima® LT1 in LT3 materialov,
izdelanih z brizganjem v forme23
Property Units LT1 LT3
Melt viscosity (kN s)/m2 0.44 0.16
Density g cm–3 1.3 1.3
Tensile strength (Yield) MPa 100 108
Tensile elongation (Break) % 40 25
Flexural strength MPa 165 170
Water absorption (24h) % 0.5 0.5
Melt temperature °C 340 340
Mould shrinkage % 1.2 1.3
5 PRELIMINARY TESTS
The first tests were made with PEEK LT3 on a simple
extrusion device. Those tests just confirmed the possi-
bilities of processing PEEK in the desired way (extrusion
through a small nozzle). Figure 5 presents the results of
the first extrusion test that, due to a short delivery time,
was made with an industrial PEEK form another
manufacturer.24 The first extrusion was successful, so
during the next steps the first PEEK 3D printer (shown
previously on Figure 3) was constructed.
B. VALENTAN et al.: PROCESSING POLY(ETHER ETHERKETONE) ON A 3D PRINTER ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 715–721 717
Figure 4: Shear viscosity versus shear stress at different temperatures
shows that PEEK Optima® LT3 has a significantly lower viscosity
when compared to LT1 or LT2 9
Slika 4: Viskoznost v odvisnosti od stri`ne napetosti in temperature
ka`e na ni`jo viskoznost PEEK Optima® LT3 v primerjavi z LT1 ali
LT2 9
Figure 3: PEEK2 machine with a better insulation, more powerful
heaters, a preheated chamber for the material and a new and improved
electric controller
Slika 3: PEEK2-naprava z izbolj{ano izolacijo, mo~nej{imi grelniki,
komoro za predgretje materiala in novim, izbolj{anim elektri~nim
krmiljem
After the approval of the possibility of using the
extrusion principle, the first PEEK processing machine
was made and the first parts were produced. After seve- ral production cycles some problems with the material
thermal decomposition inside the extrusion nozzle
appeared (Figure 6).
To verify the cause of a degraded polymer, additional
tests were made. The presence of a degraded polymer
was performed using a TA Instruments Q2000 differen-
tial scanning calorimeter with the settings presented in
Table 2.
The results (Figure 7) of a multi-cycle DSC analysis
did not indicate the presence of a degraded polymer. The
difference between the first and the last recrystallization
cycles was approximately 1.7 °C. If a degraded polymer
had been present, the value (difference) would have been
significantly larger, at least 5 °C to 10 °C. It is, therefore,
likely that the bubbles are due to the presence of the
moisture in the polymer when the sample specimens
were made.
To get the best possible results, the material was
dried again (according to the manufacturer,9 the LT3
material should be dried for at least 3 h at 150 °C or 12 h
at 120 °C) and a new, clean extruder was used for the
sample production.
6 PROCESSING DATA AND TEST PROCEDURES
According to the standards EN ISO 527-2: 2012 pla-
stics – determination of tensile properties and EN ISO
178: 2011 plastics – determination of flexural properties,
several sets of test specimens were made (Figure 8).
Optical microscopy was performed using an Olympus
SZX7 optical microscope and the Spot imaging software.
The tensile-strength testing was performed on an Instron
3367 tensometer using the Invibio laboratory test method
INV-LAB-TM12 revision 2. SEM was performed on a
Hitachi TM3000 scanning electron microscope.
The SEM analyses of the gaps and voids (Figure 9)
showed a presence of moisture even after the raw ma-
terial was dried according to the material manufacturer’s
B. VALENTAN et al.: PROCESSING POLY(ETHER ETHERKETONE) ON A 3D PRINTER ...
718 Materiali in tehnologije / Materials and technology 47 (2013) 6, 715–721
Figure 5: First extrusion tests showed that an extrusion of PEEK
through a small-diameter nozzle is possible
Slika 5: Prvo poizkusno ekstrudiranje je pokazalo, da je ekstrudiranje
PEEK-a skozi {obo majhnega premera mogo~e
Figure 6: Degradation of PEEK inside the nozzle as a result of several
heating cycles. The left figure shows the inside of the nozzle after
being used for 15 cycles (15 working days). The right figure shows
small particles that are caused by the material degradation inside the
nozzle and are eventually deposited onto the product.
Slika 6: Razgradnja PEEK-a v {obi kot rezultat ve~kratnega segreva-
nja. Leva slika prikazuje notranjost {obe po 15 ciklih (15 delovnih
dneh). Na desni sliki so vidni majhni ~rni delci, ki so se pojavili v
izdelku kot posledica razgrajenega materiala, ki je iz notranjosti {obe
prehajal v izdelek.
Table 2: Test settings of the TA Instruments Q2000 differential scan-
ning calorimeter
Tabela 2: Nastavitve TA-naprave Q2000 za diferen~no vrsti~no kalo-
rimetrijo
Parameter Value
Number of cycles 5
Sampling interval 0.20 s/pt
Equilibrate temperature 30.00 °C
Ramp 1 20.00 °C/min to 400.00 °C
Isothermal 2.00 min
Mark end of cycle 0
Ramp 2 20.00 °C/min to 30.00 °C
Isothermal 5.00 min
Mark end of cycle 0
Ramp 3 20.00 °C/min to 400.00 °C
Mark end of cycle 0
Figure 7: Results of the differential scanning colorimeter. Set-A
traces show the first and the fifth crystallisation cycles.
Slika 7: Rezultati diferen~nega vrsti~nega kolorimetra. Sledi seta A
prikazujeta prvi in peti cikel kristalizacije.
data. The problem is that it takes several hours to pro-
duce a set of samples, so the material gains the moisture
back from the surrounding air. The results lead us to
improve the design of the PEEK printer and include a
drying chamber, so that the material can be in dry
condition all the time.
The results of the tensile stress showed that the
tensile strength of the parts made with the presented
technology (thermoplastic modelling or FDM) is signi-
ficantly lower (Figure 10) than the results for the parts
made with injection moulding (Table 1). However, the
tensile strength is still high enough for designated appli-
cations (higher than for the implants made of PMMA).
To test the impact of heat treatment (present due to a
long production time) on the product, the test specimens
were produced and left in the chamber at a high tempera-
ture for additional 12 h. The results (Figure 11) showed
it has a negative impact on the mechanical properties that
should be avoided, so the parts should be taken out of the
chamber immediately after the production.
The last test presents the problem of a non-uniform
raw material that was delivered (Figure 12). These test
specimens showed that a change in the rod diameter
from 2.3 mm to 2 mm (in the case of the first delivered
charge, resulting in a material deficit in volume of 30 %)
can cause a drop in mechanical properties of up to 50 %
(Figure 13).
B. VALENTAN et al.: PROCESSING POLY(ETHER ETHERKETONE) ON A 3D PRINTER ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 715–721 719
Figure 11: Response of the test specimens after being left in a heated
environment for 12 h. Mechanical properties of the products under the
obligatory high temperature are deteriorating with time. The colour
change from light brown to dark brown was also detected. Specimens
2 and 9 have unusually high values, so they are excluded from the
average-value calculations.
Slika 11: Diagram preizkusnih vzorcev, zadr`anih na visoki tempe-
raturi okoli 12 h. Mehanske lastnosti izdelkov se zaradi vpliva visokih
temperatur hitro slab{ajo. Opaziti je tudi spremembo v barvi izdelkov,
od svetlo rjave na temno rjavo. Vzorca 2 in 9 imata nenavadno visoke
vrednosti, zato pri izra~unih srednje vrednosti nista bila upo{tevana.
Figure 9: Gaps and bubbles in the model are a result of the moisture
in the material
Slika 9: Vrzeli in mehur~ki v izdelku so posledica vlage v materialu
Figure 10: Typical response of a PEEK specimen. After trimming the
second PEEK printer, the average tensile stress of approximately 60
MPa can be achieved for the large parts that are taken out of the pro-
duction chamber immediately after the end of the production.
Slika 10: Zna~ilen diagram preizku{anja mehanskih lastnosti. Po na-
stavitvi parametrov na drugem PEEK-tiskalniku je mogo~e dose~i
natezno trdnost okrog 60 MPa, ~e so deli odstranjeni iz delovne ko-
more takoj po kon~ani izdelavi.
Figure 8: PEEK test specimens made on the PEEK plate. A nicely
fulfilled structure is seen and so are the layers in the Z direction.
Slika 8: PEEK-preizkusni vzorci, izdelani na delovni plo{~i iz
PEEK-a. Vidi se lepo izpopolnjena notranja struktura in sloji v
Z-smeri.
7 PRODUCTS
For additional tests and calibration three basic shapes
were used: a cube to get the right amount of the extruded
material, a cylinder to check the X and Y movement
synchronisation and a test specimen according to the
standards (ISO) for mechanical properties. The goal of
the 3D printer is to produce more complex shapes, so
real-life products (also in the shapes of implants) were
also produced. Figure 14 presents one of the first
successfully produced bigger parts (a height of 100 mm)
with a really small wall thickness, due to poor mecha-
nical properties later replaced by the LT3 PEEK Optima®
material. Figure 15 shows one of the latest parts in the
shape of a real-size (a width of approx. 120 mm) skull
implant with a support structure. The implant was made
in the vertical position since we found that in this
position smoother products can be made.
8 CONCLUSION
The article presents only a small amount of the
research done in the project to make a 3D printer capable
of directly producing PEEK parts. In the process of
developing the appropriate device, several other products
were made. The PEEK printer is currently in its final
prototype phase, it is functional but to produce a part,
additional knowledge on the part orientation, machine
tricks and G-code generation is still needed. Because of
the bio-compatibility requirement, the whole hot end (the
part where PEEK changes the aggregate state from solid
to liquid) is exchangeable and must be exchanged after a
part has been produced. The hot end is then renewed at
the manufacturer’s facility.
Future development is planned to improve the me-
chanical properties, so some more tests with appropriate
building temperatures and speeds need to be done. The
next step is also to replace the control electronics with
B. VALENTAN et al.: PROCESSING POLY(ETHER ETHERKETONE) ON A 3D PRINTER ...
720 Materiali in tehnologije / Materials and technology 47 (2013) 6, 715–721
Figure 15: Part of a real-shape implant (a skull implant) made with
the support material in the vertical position
Slika 15: Del lobanjskega vsadka, izdelanega z dodatnimi podporami
v pokon~nem polo`aju
Figure 14: Thin-wall tornado (the wall thickness of only 0.55 mm)
produced from LT3 PEEK
Slika 14: Tankostenski tornado (debelina stene samo 0,55 mm), izde-
lan iz materiala PEEK LT3
Figure 13: Drop in the mechanical properties as a result of the change
in the diameter of the rod from 2.3 mm to 2 mm
Slika 13: Poslab{anje mehanskih lastnosti kot posledica spremembe
premera `ice iz 2,3 mm na 2 mm
Figure 12: Non-uniform raw rod varying from 2 mm to 2.3 mm in
diameter presents a big problem, since it leads to a lack of the material
inside the produced part and, subsequently, to a drop in the mecha-
nical properties.
Slika 12: Neenakomeren prerez osnovnega materiala, ki variira med
2 mm in 2,3 mm, je velik problem, saj povzro~i pomanjkanje mate-
riala v izdelku in posledi~no slab{e mehanske lastnosti.
robust, accurate and reliable home-made electronics
developed specially for this application (so additional
parameters will be monitored and, if needed, more
powerful motors can be used).
Acknowledgement
This work has been partly financed from the Euro-
pean Union – European Social Funds.
9 REFERENCES
1 Producer of Selective Laser Melting machines [cited 2012-10-10].
Available from World Wide Web: http://www.realizer.com/en/
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3 Producer of the EBM® machines that are builds parts layer-by-layer
from metal powder using a powerful electron beam [cited 2012-
10-10]. Available from World Wide Web: http://www.arcam.com/
4 M. Joseph, Titanium Implants Can Cause Cancer [cited 2012-10-10].
Available from World Wide Web: http://www.anaturalcure.com/
mambo/index2.php?option=com_content&do_pdf=1&id=482
5 P. Scolozzi, A. Martinez, B. Jaques, Complex Orbito-fronto-temporal
reconstruction using computer designed PEEK implant, The journal
of craniofacial surgery, 18 (2007) 1, 224–227
6 M. M. Kim, D. O. K. Boahene, P. J. Byrne, Use of customized Poly-
etheretherketone (PEEK) implants in the reconstruction of complex
maxillofacialdefect, Arch. facial. plast. surg., 11 (2009) 1, 53–57
7 M. M. Hanosano, N. Goel, F. DeMonte, Calvarial reconstruction
with Polyetheretherketone implants, Annals of plastic surgery, 62
(2009) 6, 653–655
8 M. Schmidt, D. Pohle, T. Rechtenwald, Selective Laser Sintering of
PEEK, CIRP Annals – Manufacturing Technology, 56 (2007) 1,
205–208
9 PEEK Optima®, Unfilled PEEK-OPTIMA Polymer Processing
Guide [cited 2012-10-10]. Available from World Wide Web: http://
www.invibio.com/resource-library/guidelines-preview.php?id=19
10 I. Drstven{ek, Slojevite dodajalne tehnologije, In: K. Kuzman, (ed.),
Moderno proizvodno in`enirstvo, priro~nik, Grafis trade, Grosuplje
2010, 777–812
11 I. Drstven{ek, T. Brajlih, B. Valentan, Hitra izdelava prototipov, In:
J. Bali~, I. Ve`a, F. ^u{ (eds.), Napredne proizvodne tehnologije,
Univerza v Mariboru, Fakulteta za strojni{tvo, Sveu~ili{te u Splitu,
Fakultet elektrotehnike, Maribor, Split 2007, 98–120
12 A. Pilipovic, P. Raos, M. Sercer, Experimental analysis of properties
of materials for rapid prototyping, International journal of advanced
manufacturing technology, 40 (2009) 1–2, 105–115
13 T. Brajlih, B. Valentan, J. Bali~, I. Drstven{ek, Speed and accuracy
evaluation of additive manufacturing machines, Rapid prototyping j.,
17 (2010) 1, 64–75
14 K. Lokesh, P. K. Jain, Selection of rapid prototyping technology, Ad-
vances in Production Engineering & Management, 5 (2010) 2, 75–84
15 B. Valentan, T. Brajlih, I. Drstven{ek, J. Bali~, Development of a
part-complexity evaluation model for application in additive
fabrication technologies, Stroj. vestn., 57 (2011) 10, 709–718
16 Web page of 3D printer producer Ortotip, d. o. o. [cited 2012-10-10].
Available from World Wide Web: http://www.ortotip.com/soustvar-
jalec.html
17 Generation 6 Electronic - web page with resources for GEN 6 elec-
tronic [cited 2012-10-10]. Available from World Wide Web:
http://reprap.org/wiki/Generation_6_Electronics
18 Web page of RepRap printers [cited 2012-10-10]. Available from
World Wide Web: http://www.reprap.org/wiki/RepRap
19 ISO 13485:2003 standard [cited 2012-10-10]. Available from World
Wide Web: http://www.iso.org/iso/catalogue_detail?csnumber=
36786
20 A. Maglica, K. Krnel, M. Ambro`i~, Ceramic composites based on
silicon nitride, Mater. Tehnol., 43 (2009) 3, 165–169
21 T. Kroupa, R. Zem~ík, J. Klepá~ek, The temperature dependence of
the parameters of non-linear stress-strain relations for carbon-epoxy
composites, Mater. Tehnol., 43 (2009) 2, 69–72
22 I. N. Orbulov, K. Májlinger, Microstructure of metal-matrix com-
posites reinforced by ceramic microballoons, Mater. Tehnol., 46
(2012) 4, 375–382
23 PEEK-OPTIMA Typical Material Properties [cited 2012-10-10].
Available from World Wide Web: http://www.invibio.com/resource-
library/brochures-preview.php?id=30
24 Properties of Quadraplastics Ketron PEEK 1000 [cited 2012-10-10].
Available from World Wide Web: http://www.quadrantplastics.com/
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/AEP/KetronPEEK_1000_E_PDS_0907.pdf
B. VALENTAN et al.: PROCESSING POLY(ETHER ETHERKETONE) ON A 3D PRINTER ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 715–721 721

O. CULHA et al.: FINITE-ELEMENT MODEL OF THE INDENTATION FOR YBCO-BASED SUPERCONDUCTOR ...
FINITE-ELEMENT MODEL OF THE INDENTATION FOR
YBCO-BASED SUPERCONDUCTOR THIN FILMS
MODEL KON^NIH ELEMENTOV ZA VTISKOVANJE TANKIH
PLASTI SUPERPREVODNIKA YBCO
Osman Culha1, Ilyas Turkmen1, Mustafa Toparlý2, Erdal Celik2
1Celal Bayar University, Department of Materials Engineering, Muradiye Campus, Manisa, Turkey
2Dokuz Eylul University, Department of Metallurgical and Materials Engineering, Tinaztepe Campus, Buca, Izmir, Turkey
osman.culha@cbu.edu.tr
Prejem rokopisa – received: 2012-10-16; sprejem za objavo – accepted for publication: 2013-03-25
Superconducting films with poor mechanical properties are useless even if they possess good transport and flux-pinning proper-
ties. Since additive particles as pinning centers are important changes in a microstructure, their effect on the micromechanical
properties such as Young’s modulus and hardness have to be investigated with respect to the additional-particle type and quan-
tity, using experimental and numerical methods. In this study, films were dip-coated onto (001) SrTiO3 (STO) single-crystal
substrates with metalorganic deposition using the trifluoroacetate (TFA-MOD) technique. The phase analysis and the micro-
structure of the superconducting thin films were determined with an X-ray diffractometer (XRD) and a scanning electron micro-
scope (SEM). The mechanical-property variations of the pure YBCO and the YBCO thin films with Mn (reacting as BaMnO3)
were experimentally obtained with nanoindentation techniques. Thus, the BaMnO3 nanoparticle effects on the structural and
mechanical properties of the films were observed. According to the nanoindentation results, the Young’s modulus and
indentation hardness of the films decreased from 88.54 GPa to 76.47 GPa and from 12.51 GPa to 3.88 GPa, respectively,
depending on the additive particles. In addition, the finite-element modeling (FEM) of the indentation was applied to estimate
the failure stress/stress distribution relation at the contact region between the indenter and the surface of a YBCO-based thin
film, obtaining the same force/penetration depth curve as with the indentation experiment. According to these main aims of
FEM, the mesh-design effect, material properties and the boundary condition of the axisymmetric model were chosen and
optimized to obtain the mechanical results of the instrumented indentation.
Keywords: superconducting films, mechanical properties, nanoindentation, finite-element modelling, stress distribution
Superprevodne tanke plasti s slabimi mehanskimi lastnostmi so neuporabne, ~eprav imajo dobre transportne lastnosti in
zmo`nost zadr`evanja toka. Ker so dodani delci kot mesta zadr`evanja pomembna sprememba v mikrostrukturi, je treba z
eksperimentalnimi in numeri~nimi metodami raziskati njihov vpliv na mikromehanske lastnosti, kot sta Youngov modul in
trdota, v odvisnosti od vrste in koli~ine dodanih delcev. V tej {tudiji so bile mono- kristalne tanke plasti (001) SrTiO3 (STO) s
potapljanjem oblo`ene s kovinsko-organskim nanosom z uporabo tehnike s trifluoroacetatom (TFA-MOD). Fazna analiza in
mikrostruktura superprevodnih tankih plasti sta bili dolo~eni z rentgenskim difraktometrom (XRD) oziroma z vrsti~nim elek-
tronskim mikroskopom (SEM). Spreminjanje mehanskih lastnosti ~istega YBCO in tankih plasti YBCO z Mn (reagira kot
BaMnO3) je bilo eksperimentalno dolo~eno z nanovtiskovanjem. Opa`eni so bili u~inki nanodelcev BaMnO3 na strukturne in
mehanske lastnosti plasti. Na podlagi rezultatov nanovtiskovanja sta se Youngov modul in trdota plasti zmanj{ala iz 88,54 GPa
na 76,47 GPa oziroma iz 12,51 GPa na 3,88 GPa, odvisno od dodanih delcev. Dodatno je bilo uporabljeno tudi modeliranje
vtiskovanja s kon~nimi elementi (FEM) za dolo~anje poru{itvene razporeditve napetosti v podro~ju stika vtisnega telesa in
povr{ine tanke plasti na osnovi YBCO, da bi dobili krivulje sila – globina vtiskovanja pri preizkusu vtiskovanja. Skladno s
ciljem FEM-modela, je bil izbran u~inek oblike mre`e, lastnosti materiala in robni pogoji za osnosimetri~ni model in izvr{eno je
bilo optimiranje glede na dose`ene rezultate pri mehanskem instrumentiranem vtiskovanju.
Klju~ne besede: superprevodne tanke plasti, mehanske lastnosti, nanovtiskovanje, modeliranje s kon~nimi elementi, razporedi-
tev napetosti
1 INTRODUCTION
Since the discovery of high-Tc superconductors,
numerous researchers focused on their electric proper-
ties. The production effect and the critical-current den-
sity (Jc) dependence on the external magnetic field and
on the temperature in high-Tc superconductor (HTS)
YBCO films have been widely investigated.1–7 The
results demonstrate a high potential for an application of
YBCO films in power devices such as fault current limi-
ters (FCL) and motors. Jc depends on the vortex-pinning
capability of a film, the optimization of which is
currently achieved by controlling the microstructure and
the defect distribution (BaMeO3 as a nanorod and dot) in
a film. However, increasing the pinning-center density
up to the value, at which the distance between the nearest
centers reaches the coherence length will severely
deteriorate the mechanical properties of the films. Thus,
it may occur that for many of the HTS thin-film appli-
cations requiring a high critical-current density, the
limiting factor will no longer be the critical-current
density magnitude but the mechanical strength against
the body forces caused by flux pinning. The mechanical
strength makes a crucial contribution to the stability and
credibility of any device based on this form of materials.8
In addition, the mechanical properties are also very
important for the final industrial application.
In the literature, there have been several reports on
their hardness, bending strength and internal friction. A
precise knowledge of the elastic and deformation beha-
Materiali in tehnologije / Materials and technology 47 (2013) 6, 723–733 723
UDK 532.6:538.945 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)723(2013)
viour is a prerequisite for a successful production and
operation of such devices. However, very little informa-
tion on the elastic-to-plastic transition is available in the
literature.9–11 On the other hand, nanoindentation has
been used routinely in the mechanical characterization of
thin films and thin-surface layers in recent years.12 The
technique applies a programmed function of an in-
creasing and decreasing load to the surface of interest
with an indenter of a well-defined shape, continuously
measuring the indenter displacement (Figure 1). The
advantage of this method is that the mechanical
information, such as the elastic modulus, can be obtained
through an analysis of the load-displacement behavior
using only the coating of the material to be tested on a
substrate made from a different material. This makes it
an ideal tool. The technique has been used to assess the
elastic and plastic properties of coatings on a range of
substrates, but there are limitations in measuring the
properties of much thinner coatings, particularly when
elastic properties are required.13,14 For the coatings with a
thickness of a few hundred nanometers, it has been
suggested that extrapolating the properties determined
for a range of peak loads or indenter displacements to
zero load/depth can be used to determine the
coating-only properties. With high-quality sharp
indenters it is possible to assess coatings down to the
thickness of 100 nm using this method.15
The indentation-system frame stiffness and the
diamond tip shape were carefully calibrated with a
fused-silica test sample, using the standard Oliver and
Pharr method (1992), before and after the measurements
with no change for either of them. Nano-indentation load
(P) vs. displacement (h) curves were then recorded for
each indent and only the ones showing evidence of a
plastic deformation (i.e., the loading and unloading
curves are different) were used in the analysis of Young’s
modulus using the above methods.15–18 In addition, most
of the previous studies about the depth-sensing nano-
indentation technique demonstrate that: to extract the
mechanical properties of the coatings or films from the
indentation test, we should limit the indentation depth to
less than 10–20 % of the film thickness.19,20 As we know,
during the nanoindentation experiments, especially when
the indentation depth is about 100–1000 nm, size-scale-
dependent indentation effects become important. These
possible size-scale-dependent effects on the hardness
have been modeled using higher-order theories21–23 and
these effects can also be related to the bluntness of the
nominally sharp indenters.24,25 For the cases when the tip
bluntness is on the same order as the indentation depth,
Borodich and Keer24 proposed fundamental relations for
depth of indentation, size of contact region, load, hard-
ness, and contact area for various boundary conditions. If
the indentation is sufficiently deep (typically deeper than
1 μm), then the scale-dependent effects become small
and negligible. Therefore, the limitation of the inden-
tation depth to less than 10–20 % of the film thickness
for the films with the thickness measured in micrometers
is not a good rule. However, when the indentation depth
is deeper than 1 μm, the deformation of the substrate has
an effect on extracting the mechanical properties of the
film from the indentation test for very thin films. Saha
and Nix26 examined the effects of the substrate on
determining the hardness and elastic modulus of thin
films with nanoindentation. They found that the effect of
the substrate hardness on the film hardness was negli-
gible in the case of soft films and the substrate hardness
affected the measured film hardness in the case of a hard
film on a soft substrate.
Compared to hardness, the nanoindentation measure-
ment of the elastic modulus of thin films was more
strongly affected by the substrate. Recently, with respect
to an elastic-plastic film applied to an elastic substrate
system, some researchers23,27,28 have proposed the tech-
niques that allow us to extract a film’s elastic-plastic
properties, taking into account the influence of the
elastic substrate from a conical indentation test using the
finite-element method (FEM). In this study, elastic
properties and the limits (E and f) of YBCO-based films
with/without BaMnO3 defects as pinning centers were
designated with the finite-element method (FEM) and
nanohardness experiments according to the theory of
indentation. Furthermore, failure (f), contact-deforma-
tion characteristics, stress distribution near the contacts
of the films depending on the amount of formed
BaMnO3 particles were obtained using an algorithm with
FEM and a comparison between the experimental and
simulation results of the nanoindentation.
2 EXPERIMENTAL PROCEDURE
Initially, (100) STO single-crystalline substrates with
dimension of 10 mm × 10 mm × 0.75 mm were rinsed in
acetone using a standard ultrasonic cleaner. After that,
the solutions were deposited on the substrates using a
dip-coating process with a withdrawal speed of 0.3 cm/s
in a vacuum atmosphere. The deposited gel films were
converted to an epitaxial pure YBCO and a YBCO film
with BaMnO3 nanoparticles through a combination of
calcining and heat-treatment procedures. After that, the
structural development of the produced thin films was
investigated with the X-ray diffraction (XRD-Rigaku
D/MAX-2200/PC) patterns that were recorded using the
Co K irradiation (the wavelength,  = 0.178897 nm),
with the scanning range between 2
 = 10° and 90°. The
surface topographies and additional-particle effects on
the microstructure of the films were examined with a
high-resolution scanning electron microscope SEM
(JEOL JSM 6060). The hardness and Young’s modulus
of the produced thin films were measured under the peak
load of 300 μN three times with the CSM Berkovich
nanoindentation tester (the loading-unloading test mode)
to determine the additional-particle effects on the mecha-
nical properties.
O. CULHA et al.: FINITE-ELEMENT MODEL OF THE INDENTATION FOR YBCO-BASED SUPERCONDUCTOR ...
724 Materiali in tehnologije / Materials and technology 47 (2013) 6, 723–733
3 INDENTATION MODELING OF YBCO-BASED
THIN FILMS
A simulation of YBCO-based thin-film indentation
analysis exposed very important mechanical properties
such as the maximum (fracture-fracture) strength that
influenced the mechanical stability of the coated super-
conductor under service conditions. The elastic modulus
and hardness of YBCO-based thin films could be deter-
mined from the instrumented indentation with the load-
ing-unloading curves. However, the maximum stress, the
stress-strain relationship, the substrate and the thickness
effect on the characteristic loading-unloading curves
could not be obtained from the experimental results. At
this time, researchers used a simulation and finite-ele-
ment modeling with the developed algorithms. Normally,
until the 2000 s, a finite-element analysis of an indenta-
tion problem was applied to bulk ceramic and metallic-
based materials. However, after the development of
thin-film technologies and materials had increased, the
determination of mechanical properties of films and
coatings became more focused. For instance, thin-film-
based technologies provide for a higher performance, a
higher density and a smaller overall size of a micro-
system package. Although thin films are increasingly
used for the sake of performance, functionality and size,
it is important to understand the mechanical behavior of
thin films to address the reliability concerns.
The indentation analysis of YBCO-based thin films
was started with a design of an axisymmetric model. A
diamond conical surface with a half-apex angle of 70.3°
was used to model the widely used Berkovich indenter.
In addition, the thicknesses of YBCO-based thin films
and substrates were 300 nm and 5000 nm, respectively,
and the ratio of the film thickness to the substrate was
O. CULHA et al.: FINITE-ELEMENT MODEL OF THE INDENTATION FOR YBCO-BASED SUPERCONDUCTOR ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 723–733 725
Figure 2: Finite-element model of YBCO-based thin films on an STO
single-crystal substrate
Slika 2: Model kon~nih elementov tanke plasti na osnovi YBCO na
monokristalni podlagi STO
Figure 1: Characteristic indentation curve and representation of the
parameters
Slika 1: Zna~ilna krivulja vtiskovanja in predstavljeni parametri
6 % in this model. The substrate dimensions (both width
and height) are taken to be at least 16 times larger than
those of the film, so as to simulate a semi-infinite
substrate (Figure 2a). On the other hand, four different
material inputs were loaded to the model as a pure
YBCO thin film, a YBCO thin film with 0.05 g Mn, a
YBCO thin film with 0.10 g Mn and a YBCO thin film
with 0.15 g Mn (BaMnO3). The input data of the material
properties of the axisymmetric thin-film model with a
substrate and an equivalent cone are listed in Tables 1
and 2. Elastic modules were experimentally determined
from the indentation unloading curves for four different
YBCO-based thin-film models. So, the determined
elastic modules were loaded with Poisson’s ratio as the
material elastic properties in the finite-element simu-
lation. While fixing the elastic properties presented in
Table 2 for each sample, the maximum stresses (failures)
of YBCO-based thin films were changed to determine
the same penetration force/depth curve with the experi-
mental indentation analysis.
The properties of the contact region between the
indenter and the surface of a thin film/ substrate system
could be designed in the assembly and interaction
modulus of the Abaqus package program. The contact
areas of YBCO-based thin films were divided into parti-
tions. As known from the indentation-analysis theory, a
fine-mesh design could be applied to the mesh part so as
to create the partitions near the contact region. The con-
tact regions of YBCO-based thin films were calculated
as the dimensions of 300 mm × 400 nm (thickness ×
length). After this area determination (depending on the
indentation depth), the whole model began to partition
from the contact region to the substrate with dimensions
of 500 nm and 400 nm, respectively (Figure 2b). This
partitioned model can be used for different element-type
and region determinations for a fine-mesh design of a
contact area. At the same time, interaction properties for
the indenter and the film surface are established as the
finite-sliding formulation and surface-to-surface con-
straint-enforcement method. The contact property of the
entire model was assumed to be tangential, being fric-
tionless between the indenter and the surface of a film.
As the main aim of the study was to obtain the stress
distributions and force-depth curves of YBCO thin films,
the mesh design of the entire model became important,
especially when a problematic region included a contact
between two separated surfaces. As we already know, the
instrumented indentation analysis was carried out with a
diamond indenter that penetrated the surface of a thin
film at the nanoscale, deforming the contact region, not
the whole surface, plastically under the applied load. The
importance of a divided surface near a contact region
required very correct and sensitive results for simulated
loading-unloading curves of YBCO-based thin films. For
that reason, the elements were the finest in the central
contact area and became coarser outwards for the entire
model and a magnified view of the contact region can be
seen in Figure 2. According to Figure 2c, the contact
region was meshed as a structured CAX4R element type,
A-4 node bilinear axisymmetric quadrilateral, reduced
integration and hourglass control with 2, 4, 8, 16, 32 and
40 numbers for a determination of the indentation depth
and load dependency on the mesh-element number. In
this model, as presented in Figure 2c, the contact-region
size of YBCO-based thin films was 300 nm × 400 nm.
When the element number increased from 2 to 40, the
element size was decreased from approximately 200 nm
to 10 nm. The smallest element size and the total ele-
ment number were 10 nm and 3764, respectively, which
enabled an accurate determination of the real-impression
size.
O. CULHA et al.: FINITE-ELEMENT MODEL OF THE INDENTATION FOR YBCO-BASED SUPERCONDUCTOR ...
726 Materiali in tehnologije / Materials and technology 47 (2013) 6, 723–733
Table 1: Indentation-experiment results for YBCO-based superconducting thin films
Tabela 1: Rezultati eksperimentalnega vtiskovanja v superprevodno plast na osnovi YBCO
Material Force(μN)
Max. depth
(nm)
Res. depth
(nm)
Film thick-
ness (nm)
Hardness
(HV)
Indentation
hardness
(GPa)
Young’s
modulus
(GPa)
Pure YBCO thin film
300
40.24 ± 6.4 30.12 ± 7.8 292 ± 9 695 ± 28 12.51 ± 4.8 88.54 ± 3.1
YBCO thin film with 0.05 g Mn 42.32 ± 9.4 32.11 ± 8.7 297 ± 5 525 ±16 8.21 ± 1.2 83.41 ± 1.8
YBCO thin film with 0.10 g Mn 42.45 ± 5.9 33.53 ± 4.4 294 ± 6 495 ± 21 5.75 ± 1.1 79.11 ± 1.9
YBCO thin film with 0.15 g Mn 47.04 ± 6.1 37.48 ± 5.7 305 ± 8 454 ± 18 3.88 ± 0.8 76.47 ± 1.3
Table 2: Material properties of the entire model
Tabela 2: Lastnosti materiala za celoten model
Material type Elastic modulus(GPa)
Poisson ratio
(v)
Maximum stress
(GPa)
Ratio of maximum
stress to elastic
modulus
Pure YBCO thin film 88.54
0.261
7.50–9.50 0.0847–0.1073
YBCO thin film with 0.05 g Mn 83.41 6.50–9.00 0.0779–0.1079
YBCO thin film with 0.10 g Mn 79.11 5.00–6.75 0.0632–0.0853
YBCO thin film with 0.15 g Mn 76.47 2.75–4.25 0.0360–0.0556
STO substrate 278 0.238 – –
Equivalent cone 1040 0.07 – –
4 RESULTS AND DISCUSSION
4.1 Characterization of YBCO-based thin films
Details of the phase analysis of YBCO-based thin
films are shown in Figure 3 relating to the pure YBCO,
and the YBCO with 0.05 g, 0.10 g and 0.15 g of
BaMnO3 nanoparticles on the STO single-crystal sub-
strate. XRD patterns illustrate that YBCO films have
(001) and parallel plane reflections for the pure YBCO
thin film. The major diffraction peaks corresponding to
the (00l) parallel plane was developed. In addition, no
second phases such as Y2Cu2O5, BaF2 and CuO were
found, and only the pure YBCO phase was formed.
Apart from that, BaMnO3 perovskite peaks with a low
intensity were determined on account of the Mn doping
effect. On the other hand, the surface analysis of
YBCO-based thin films indicated that structural defects
could react as nanodots or nanoparticles (due to the Mn
addition) along the c-axis of a YBCO film as presented
in Figure 4. Since Mn reacts with Ba and a BaMnO3
perovskite structure forms in the YBCO film during the
heat-treatment process, the microstructures of supercon-
ducting thin films were changed as expected.
4.2 Instrumented indentation results of YBCO-based
thin films
Since the general purpose is to determine the mecha-
nical properties such as the Young’s modulus and hard-
ness of the pure YBCO thin film and the YBCO thin
films with a Mn addition (Mn reacts as BaMnO3), an
instrumented nanoindentation test was used. At this time,
the importance of the applied load became apparent. The
indentation response of YBCO-based thin films on the
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Figure 4: Surface analysis of YBCO-based superconductor thin films with additional particles
Slika 4: Analiza povr{ine superprevodne tanke plasti na osnovi YBCO z dodanimi delci
Figure 3: Phase analysis of YBCO-based superconductor thin films
with additional particles
Slika 3: Fazna analiza superprevodne tanke plasti na osnovi YBCO z
dodanimi delci
STO substrate is a complex function of the elastic and
plastic properties of both the film and the substrate and it
is crucial to understand how the intrinsic mechanical
properties of the film can be determined from the overall
mechanical response of the film/substrate system. As the
values of the elastic modulus and hardness, determined
from indentations, do not depend on the value of h
(indentation depth) and, therefore, on the value of the
maximum load, the indentation depth should not exceed
10–20 % of the film thickness, otherwise the results will
be affected by the properties of the substrate. Another
parameter that can influence the indentation test is the
surface roughness. It has a very active role in the inden-
tation experiment at the nanoscale. If the surface rough-
ness is bigger than the maximum indentation depth for
the applied load, the curve has a very scattered loading
and unloading part. According to the surface analysis,
the surface roughness of the pure YBCO thin film is
(21 ± 5) nm.
As presented in Table 1, the ratio of the indentation
depth to the film thicknesses is applicable for the instru-
mented indentation with the 300 μN applied load. The
smoothing procedure was applied to all of the instrumen-
ted indentation results of the samples. The loading-un-
loading (load-displacement) curves of YBCO thin films
under the 300 μN applied peak load are shown in Figure
5, at the nanometer scale, three times and the total test
time was 60 s. The maximum and residual indentation
depths of the YBCO thin film were increased from
(40.24 ± 6.4) nm to (47.04 ± 6.1) nm and from (30.12 ±
7.8) nm to (37.48 ± 5.7) nm, respectively, by increasing
the Mn content in the structure (Table 1). The ratios of
the indentation depth to the film thickness of the YBCO
thin film, the YBCO with 0.05 g Mn, the YBCO with
0.10 g Mn and the YBCO with 0.15 g Mn were 13.78 %,
14.14 %, 14.43 % and 15.04 %, respectively. Further-
more, the elastic modulus and instrumented hardness of
YBCO thin films were decreased from (88.54 ± 3.1) GPa
to (76.47 ± 1.3) GPa and from (12.51 ± 5.1) GPa to (3.88
± 0.8) GPa under the 300 μN applied load as listed in
Table 1. The loading and unloading parts of the curves
have some fluctuations due to the surface roughness and
porosity of YBCO-based thin films. For that reason the
standard deviations were added to all the indentation-test
characteristics, such as the maximum indentation depth,
the residual depth, the elastic modulus and the instru-
mented hardness of the films.
According to the results, the indentation hardness of
the films decreased with the increasing Mn content in the
film structure. Depending on the Mn content, the maxi-
mum indentation depth, the residual depth and the
indentation hardness were changed scientifically. Since
the Mn content increased in the microstructures of
YBCO-based films, the intensity of BaMnO3 phases in
the X-ray analysis increased. Furthermore, the SEM ana-
lysis showed that the BaMnO3 phase structure dispersed
and agglomerated in the YBCO thin-film structure. So,
the characteristic indentation curves of YBCO-based thin
films changed depending on the Mn content and the
BaMnO3 phase structure. As expected, BaMnO3 reacted
as a defect, becoming the pinning center in the structure.
When the defect concentration increased due to the
increased Mn content, the stability of the structure and
interatomic bonding mechanism changed to particle-
doped YBCO-based thin films. Although the lattice
parameter was very close to YBCO, the BaMnO3 based
structure in YBCO thin films caused some distortion (the
residual stress) and irregularities. Although the indenta-
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Figure 5: Indentation curves of YBCO-based thin films on the STO single-crystal substrate
Slika 5: Krivulje vtiskovanja tanke plasti na osnovi YBCO na monokristalni podlagi STO
tion load was fixed at 300 μN, the indentation hardness
was decreased as if the indentation size effect had
influenced the mechanical properties. Since the hardness
is accepted as an inherent material property, it should not
vary with the indentation load and size but can change
with different phase formations. A decrease in the hard-
ness due to the increasing BaMnO3 content in the YBCO
thin-film structure caused differences in the indentation
depth. According to this explanation, it can be concluded
that the pure YBCO thin film is harder and more brittle
than the BaMnO3 films. As presented in Figure 6, the
indentation hardness of YBCO-based thin films de-
creased from 12.51 GPa to 3.88 GPa depending on the
Mn content in the structure. The elastic-modulus vari-
ation of YBCO-based thin films is also demonstrated in
Figure 6. Although hardness was very sensitive to the
maximum indentation depth and thickness ratio of the
samples, changing from 12.51 GPa to 3.88 GPa, the ela-
stic modulus of YBCO-based thin films did not show
any sharp decrease. However, as listed in Table 1, the
elastic modulus of the pure YBCO, the YBCO with 0.05
g Mn, the YBCO with 0.10 g Mn and the YBCO with
0.15 g Mn was 88.54 GPa, 83.41 GPa, 79.11 GPa and
76.47 GPa, respectively. However, the references about
the mechanical properties of this material, particularly
the yield strength and the stress-strain curve, are scarce.
The mechanical properties (the hardness, the Young’s
modulus and the fracture toughness) of YBCO samples
were examined with the techniques such as ultrasound29
X-ray diffraction30 and nanoindentation.31 The reported
values of the Young’s modulus for Y-123 are within the
range of E = 40–200 GPa. This large scatter may be due
to the residual porosity and poor contact between the
grains.32 Other authors,33 also using nanoindentation,
reported a value of E = 171–181 GPa for the YBCO
samples textured with the Bridgman technique, which is
in agreement with Johansen, who applied between
30 mN and 100 mN. Nanohardness values in the range of
7.8–8.0 GPa at the maximum loads of 30 mN were
recently reported by Lim and Chaudhri34 for a bulk,
single-crystal YBCO. Roa et al.33 found a hardness value
of (8.9 ± 0.1) GPa using nanoindentation on the YBCO
samples textured with the Bridgman technique.32–37
4.3 FEM results of indentation
The depth-sensing instrumented nanoindentation
technique provided a continuous record of the variation
in the indentation load and the penetration depth into the
specimens. This technique has attracted considerable
attention in recent years due to its high resolution at a
low-load scale with respect to determining the elastic
modulus and hardness of thin films. The importance of
the mechanical stability of superconductor thin films
with/without additional particles under service condition
was exposed during the experimental work and finite-
element modeling.
Additional particles behave as defect centers increas-
ing the superconducting properties under a given tempe-
rature as explained above. While increasing supercon-
ducting properties with additional particles, in this study
they are Mn (BaMnO3) particles, the elastic modulus and
the hardness of thin films decreased from 88.54 GPa to
76.47 GPa, 12.51 GPa and 3.88 GPa. However, the vari-
ations in the plastic properties of YBCO-based thin films
with additional particles could not be determined with
the instrumented indentation test. For that reason, the
finite-element modeling of YBCO-based thin films with
and without BaMnO3 particles was carried out using the
Abaqus 6.10-1 package program, taking into account the
experimentally determined elastic properties of the films.
According to the depth analysis, indentation with a
different element number and size for FEM, 40 elements
with the element size of 10 nm, was chosen, due to its
sensitivity at the nanometer scale, for the simulated
loading and unloading curves of the pure YBCO thin
film and the YBCO thin films with 0.05 g, 0.10 g and
0.15 g Mn. Figure 7 shows that the simulated maximum
and residual indentation depth variations for the four
different YBCO-based thin films depend on the maxi-
mum stress rises under the 300 μN indentation load.
According to the FEM indentation analysis, the maxi-
mum depth values decreased from 40.74 nm to 39.58
nm, 43.13 nm to 41.36 nm, 44.89 nm to 42.84 nm and
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Figure 6: Indentation-hardness and elastic-modulus variations of
YBCO-based thin films
Slika 6: Spreminjanje trdote in modula elasti~nosti tanke plasti na
osnovi YBCO
51.15 nm to 46.71 nm due to an increase in the maxi-
mum stress value from 2.75 GPa to 9 GPa, respectively,
for the material input data of the pure YBCO thin film
and the films with Mn additions. On the other hand, the
residual indentation depth of thin films decreased from
22.17 nm to 19.18 nm, 23.28 nm to 17.91 nm, 28.28 nm
to 23.58 nm and 38.50 nm to 30.81 nm under the con-
ditions of the same indentation-analysis input data
(Figures 7a, b, c and d). However, the increases in the
Mn content in the structures of YBCO-based thin films
caused significant variations in mechanical properties.
When the indentation results for the pure YBCO-based
thin film and the film with the addition of 0.15 g Mn are
compared under the same simulation conditions, it is
found that the maximum and residual depths were
increased from 40.74 nm to 51.15 nm and from 19.18 nm
to 30.81 nm (Figure 7), respectively. It means that the
ductility or the plastic-deformation properties of
YBCO-based thin films increased due to the increased
Mn content in the thin-film structure or the BaMnO3 for-
mation during the production procedure of the instru-
mented indentation.
Depending on the experimental set-up used to per-
form the indentation, the loading and unloading charac-
teristics of the P – h curve, and thus the material proper-
ties estimated from the analysis of the P – h curves,
could be obtained at varying levels of accuracy. Thus, the
sensitivity of the estimated elasto-plastic properties to
the variations in the input parameters obtained from the
P – h curves is an important issue. In the FEM study, the
structural model of the instrumented indentation test was
performed as an axisymmetric and, therefore, two-
dimensional model. The indentation process is complex,
so simplifications and assumptions have to be made to
achieve a low numerical cost but, of course, sufficient
accuracy. For the time being, the problem is idealized: it
was assumed that the surface of the specimen was ideally
smooth, whereas in reality a certain roughness may be
present as well, e.g., thin films of four different YBCOs.
The indenter is an equivalent cone and, furthermore,
isothermal conditions are adopted. The model set-up is
described below. First, the appropriate boundary condi-
tions were selected and the contact to be modeled was
defined. All the simulations were carried out under the
load-controlled conditions. Two reasons can be given for
this: First, a real-indentation experiment is hard to exe-
cute under displacement-controlled conditions, i.e., up to
the maximum displacement. Second, regarding the nu-
merical cost and accuracy, neither displacement-con-
trolled nor load-controlled experiments show an import-
ant advantage. A force-depth (P – h) curve is a
straightforward characterization of indentation results
that can be precisely obtained with the instrumented
indentation equipment. As described above, the surfaces
of YBCO-based thin films were ideally smooth. How-
ever, the instrumented indentation results are affected by
the surface roughness of materials. Both loading and
unloading parts of indentation curves can have more
scattered data depending on the ratio of the indentation
depth and roughness to the film thickness. Figure 8a
presents both indentation and simulation results of pure
YBCO-based thin films under the same condition of
instrumented indentation. The same maximum penetra-
tion depth was obtained by setting E at 88.54 GPa and
the maximum stress at 8.00 GPa for the experimental and
simulation analyses (Table 2) for the pure YBCO thin
film. The surface-roughness effect on the indentation
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730 Materiali in tehnologije / Materials and technology 47 (2013) 6, 723–733
Figure 7: Depth analysis of the FEM model for YBCO-based thin film
Slika 7: Analiza globine s FEM-modelom pri tanki plasti na osnovi YBCO
simulation could not be seen due to the choice of an
ideally smooth surface. According to the experimental
indentation results, hmax and hr were (40.24 ± 6.4) nm and
(30.12 ± 7.8) nm, respectively, for the 300 μN applied
load. Similarly, the finite-element modeling of the pure
YBCO-based thin-film indentation had the maximum
and the residual indentation depths of 40.41 nm and
32.27 nm. Good agreement was found between the maxi-
mum indentation depth and the loading curves of pure
YBCO-based thin films, having an error of 0.4 %. In
addition, when the average experimental-indentation
loading/unloading was considered, this error value was
graphically obtained as 7 %, as presented in Figure 8a.
This may be due to the differences in the maximum
strength or due to the use of a constitutive model of both
YBCO-based thin films, like the elastic/perfect plastic
one without any strain-hardening exponent. On the other
hand, Figures 8b, c and d show instrumented (the ave-
rage for three different regions) and simulation indenta-
tion results for the YBCO-based thin films with 0.05,
0.10 and 0.15 g Mn under the 300 μN applied load. The
experimentally determined maximum and residual
indentation depths for YBCO-based thin films were
(42.31 ± 9.4) nm and (32.11 ± 8.7) nm, (42.45 ± 5.9) nm
and (33.53 ± 4.4) nm, and (47.04 ± 6.1) nm and (37.48 ±
5.7) nm, respectively. According to the FEM indentation
results, the maximum and residual indentation depths for
the YBCO-based thin film with Mn were 42.69 nm and
38.61 nm, 42.84 nm and 39.42 nm, and 47.17 nm and
40.43 nm, in line with the material input data from Table
2. When the average curves of the indentation were con-
sidered, the loading parts of the curves were much closer
to the numerically determined ones.
With respect to the pure YBCO thin-film indentation
results, the unloading part of the curves was separated as
it was modeled without a strain hardening exponent. The
strain hardening exponent of ceramics and, especially,
oxide thin films was close to zero. As expected, cera-
mics- and oxide-based materials did not plastically
deform and harden with deformation like the metals.
Additionally, they showed brittle failures at the maxi-
mum stress. The main aim of the FEM study of inden-
tation was to obtain the same experimental and numeric
loading/unloading curves for YBCO-based thin films,
while the elastic and plastic properties of the model were
iterated according to Table 2. Depending on the
material-property variation listed in Table 2 for the FEM
model, the maximum stress of YBCO-based films was
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Figure 9: Stress-strain distributions on the contact regions between
the indenter and the surfaces of the films
Slika 9: Razporeditev napetost – raztezek na podro~ju kontakta med
vtisnim telesom in povr{ino tanke plasti
Figure 8: Comparison of experimental and simulation indentation
curves
Slika 8: Primerjava eksperimentalnih in simuliranih krivulj vtisko-
vanja
changed from 2.75 GPa to 9.50 GPa for different Mn
additions. When the Mn content in the structure
increased from 0 g to 0.15 g, allowing a BaMnO3 forma-
tion, the maximum stress of the films decreased from
8.00 GPa to 4.00 GPa.
Stress distributions near the contact region of the
indented material and the indenter give very important
information about the simulation of the experimental
indentation. The effects of the maximum depth, the resi-
dual depth and the film thickness on the substrate-film
system under the applied load could be observed when
the field-output request of the indentation analysis was
undertaken depending on the steps of the simulations.
Figure 9 shows the Von Misses equivalent stress distri-
butions for the pure YBCO-based thin film and the
YBCO with 0.05 g, 0.10 g and 0.15 g Mn under the 300
μN applied load at the contact regions. According to the
analysis of the indentation, firstly, the indenter app-
roached and made a contact with the surface. Secondly,
the indenter was penetrating the surface until a load of
300 μN was reached. As seen in Figure 9, the stress
distribution shows yielded elements near the contact
region for the indenter and the surface. It is understood
that if an increase in the stress formation approaches the
set maximum stress during the loading step, the plasti-
cally deformed part of the contact region can be analyzed
easily when the unloading step of the simulation has
finished. The maximum stresses for the pure YBCO and
the ones with additional Mn were applied to the FEM
model, described in Table 2 as 7.50–9.50 GPa,
6.50–9.00 GPa, 5.00–6.75 GPa and 2.75–4.25 GPa,
respectively. The same penetration force/depth curves
were obtained with the instrumented indentation and
simulation for the above films when the maximum
stresses were loaded into the Abaqus 6.10-1 program in
the following order: 8.00 GPa, 7.00 GPa, 6.75 GPa and
4.00 GPa.
5 CONCLUSION
In this study, finite-element modeling (FEM) of the
indentation was applied to estimate the failure stress/
stress distribution relation at the contact region between
an indenter and a surface of YBCO-based thin films on
an STO single-crystal substrate, and obtain the same
force/penetration depth curve with an indentation expe-
riment to determine the additional-particle effects on
both structural and mechanical properties. So, the
following results were obtained:
– XRD patterns show that YBCO films have (001) and
parallel plane reflections for both the pure YBCO
and the YBCO with a BaMnO3 thin film.
– SEM micrographs indicate structural defects com-
prised of the nanodots or nanoparticles of BaMnO3
along the c-axis of a YBCO film. These properties
result in an enhanced pinning over the pure YBCO
film.
– The calculated Young’s modulus of YBCO-based
thin films with/without BaMnO3 decreased down to
the range of 88.54–76.41 GPa, with the increased
Mn content in the microstructure. In addition, the in-
dentation-hardness values of the films were de-
creased from 12.51 GPa to 3.88 GPa depending on
the Mn content.
– The finite-element-analysis results showed that the
maximum, or failure, stress of YBCO-based thin
films decreased due to the increased Mn content in
the microstructure. The yield stresses of the pure
YBCO thin film, the YBCO thin film with 0.05 g
Mn, the YBCO thin film with 0.10 g Mn and the
YBCO thin film with 0.15 g Mn were found to be
8.00 GPa, 7.00 GPa, 6.75 GPa and 4.00 GPa, respec-
tively, by comparing the experimental work and
numerical load-penetration-depth analysis under the
applied load of 300 μN.
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O. CULHA et al.: FINITE-ELEMENT MODEL OF THE INDENTATION FOR YBCO-BASED SUPERCONDUCTOR ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 723–733 733

M. OZDEMIR et al.: EFFECT OF VISCOSITY ON THE PRODUCTION OF ALUMINA BORATE NANOFIBERS ...
EFFECT OF VISCOSITY ON THE PRODUCTION OF
ALUMINA BORATE NANOFIBERS VIA
ELECTROSPINNING
VPLIV VISKOZNOSTI NA IZDELAVO
ALUMINIJEVOOKSIDNOBORATNIH NANOVLAKEN Z
ELEKTROSPININGOM
Mehtap Ozdemir1,2, Erdal Celik1, Umit Cocen1
1Dokuz Eylul University, Faculty of Engineering, Dept. of Metallurgical and Materials Engineering, 35160, Buca, Izmir, Turkey
2Gediz University, Faculty of Engineering, Dept. of Electrical and Electronics Engineering, 35665, Seyrek, Izmir, Turkey
mehtap.koklu@gediz.edu.tr
Prejem rokopisa – received: 2012-11-06; sprejem za objavo – accepted for publication: 2013-04-02
Alumina borate (Al18B4O33) nanofibers have been successfully fabricated with the electrospinning method using a solution that
contains polyvinyl alcohol (PVA) and aluminium acetate stabilized with boric acid. The effect of the viscosity on the production
of alumina borate nanofibers was investigated and the optimum parameters were determined to obtain the best fibers. DTA/TG
analyses were performed to determine the heat-treatment regime. The morphology of the fibers was evaluated by SEM analyses,
the phase structure of the fibers was obtained using XRD and the chemical bonding structure was determined by FTIR. On the
whole, the results indicated that the viscosity of a solution is an important parameter for producing a homogeneous and thin
fiber, while the calcination temperature has a significant effect on the phase structure of the fibers.
Keywords: nanomaterials, nanofiber, electrospinning, Al18B4O33
Aluminijevooksidnoboratna (Al18B4O33) nanovlakna so bila uspe{no izdelana z elektrospining metodo z uporabo raztopine, ki
vsebuje polivinil alkohol (PVA) in aluminijev acetat, stabiliziran z borovo kislino. Preiskovan je bil vpliv viskoznosti na
izdelavo aluminijevooksidnoboratnih nanovlaken in dolo~eni so bili optimalni parametri za pridobivanje najbolj{ih vlaken. Za
dolo~itev re`ima toplotne obdelave so bile izvr{ene DTA/TG-analize. Morfologija vlaken je bila ocenjena s SEM-analizo, fazna
struktura vlaken je bila dobljena z rentgensko difrakcijo (XRD), strukture kemijskih vezav pa so bile ugotovljene s FTIR.
Rezultati so pokazali, da je viskoznost raztopine pomemben parameter za proizvodnjo homogenih in tankih vlaken, medtem ko
ima temperatura kalcinacije pomemben vpliv na strukturo faz vlaken.
Klju~ne besede: nanomateriali, nanovlakna, elektrospining, Al18B4O33
1 INTRODUCTION
One-dimensional materials, such as nanowires,
nanorods, nanowhiskers and nanofibers, have stimulated
great interest due to their importance in basic scientific
research and their potential in technological applications.
They are expected to play an important role as both inter-
connecter and functional components in the fabrication
of nanoscale devices. Many unique properties have
already been proposed or demonstrated for this class of
materials, such as a high elastic modulus and tensile
strength, chemical inertness, excellent resistance to
oxidation/corrosion, a low thermal expansion coefficient,
a high thermal conductivity, a good stability at high tem-
perature and low-cost production.1–5 Of these properties,
aluminium borate is a famous material as it can be used
in a variety of applications, such as high-temperature
structural components, nonlinear optical and tribological
materials, electronic ceramics and reinforced composite
materials.2,6
Due to the superior properties of alumina borates,
nanostructures such as filaments, powders, fibers, whi-
skers and wires, they have been synthesized using vari-
ous techniques. These are vapor-solid reaction,7 powder
technology,8 the electrospinning technique,3,9 the flux
method with microwaves,10 calcination techniques.1,6,11–14
Dai et al.3 and Tuttle et al.9 have produced approximately
alumina borate nanofibers 550 nm using the electro-
spinning technique. Because it is difficult to fabricate the
extremely fine nanofibers using electrospinning, we have
investigated the effect of viscosity on the process to pro-
duce thin nanofibers.
Electrospinning is the most preferred technique to
produce nanofibers with a diameter ranging from 20 nm
to 1000 nm. The diameter of the fibers depends on the
process parameters, such as the viscosity of the solution,
the applied electric field, the distance between the collec-
tor and the needle, and the feeding rate of the solution.15
The viscosity of the solution and its electrical properties
determine the extent of the elongation of the solution that
affects the diameter of the resultant electrospun fibers.16
Moreover, the viscosity of the solution is very important
for its spinability. When the viscosity is too high, pump-
ing of the solution through the syringe needle becomes
difficult, which may result in drying of the solution at the
tip of the needle before the electrospinning initiates.
Additionally, the applied electric field affects the mor-
phology of the fibers that are obtained. Generally, a
Materiali in tehnologije / Materials and technology 47 (2013) 6, 735–738 735
UDK 620.3:66.017 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)735(2013)
higher voltage leads to greater stretching of the solution
due to the greater columbic forces in the jet as well as
the stronger electric field. These have the effect of reduc-
ing the fiber diameter.16 On the other hand, the distance
between the tip and the collector plate must be high for
thin fibers. The feed rate determines the amount of solu-
tion available for electrospinning. When the feed rate is
increased, the fiber diameter increases.16
In this study we attempt to form alumina borate
nanofibers and investigate the effect of viscosity during
the electrospinning process on the nanofiber morphology
and spinability, and the effect of temperature on the
phase structure.
2 EXPERIMENTAL
PVA solutions were prepared in different concen-
trations in mass fractions, i.e., 6 %, 8 %, and 10 %, by
dissolving the PVA in 80 °C distilled water and cooling
down to room temperature while stirring for 24 h. Next,
1 g of aluminium acetate stabilized with boric acid
(CH3CO2Al(OH)2·1/3H3BO3) was added to the PVA
solution and the stirring was continued until a transpa-
rent and homogeneous solution was obtained. The visco-
sities of solutions were determined by CVO 100 Digital
Rheometer (Bohlin Instrument).
In order to produce the nanofibers, transparent
solutions were immediately loaded into a plastic syringe.
A 22-gauge stainless-steel needle was used as a nozzle.
The emitting electrode from a power supply was attached
to the needle. The grounding electrode from the same
power supply was attached to a piece of 316L stainless
steel, which was used as the collector plate and was
placed approximately 10 cm below the tip of the needle.
A high voltage ranging between 20 kV and 30 kV was
applied across the needle and a non-woven mat of fine
fibers were fabricated. The alumina borate/PVA fibers
were heat treated at 800–1200 °C for 2 h in air using an
electrical tube furnace to obtain the alumina borate
fibers.
The prepared solution used in the electrospinning
was dried at 100 °C for 1 h and the obtained powder was
subjected to thermogravimetric differential thermal
analysis (DTA/TG) (Shimadzu DTG-60H) to define the
reaction type of the intermediate temperature products
and to use a suitable process regime. The chemical
bonding structures of the fibers before and after the heat
treatment were determined with Fourier Transform Infra-
red Spectroscopy (FTIR) (Perkin Elmer Spectrum BX).
The X-ray diffraction (XRD) measurements were per-
formed for the crystal-phase identification (Rigaku
D/Max-2200/PC) with CuK radiation. The morphology
and the average fiber diameter of the nanofibers were
characterized using a scanning electron microscope
(JSM-6060 JEOL).
3 RESULTS AND DISCUSSION
The solution viscosity plays a major role in pro-
ducing a uniform nanofiber. For a low viscosity it is
common to find beads along the fibers. As the viscosity
increases, a gradual change in the shape of the fibers
takes place until smooth fibers are obtained.16 For
spinability the viscosity must be neither very high nor
very low. The viscosity of the prepared solutions at
various concentrations are determined in the range
0.17–2.34 Pa s (Figure 1). In the Figure 1 it is clear that
as the concentration of the solution increases, the visco-
sity of the solution increases. Furthermore, the viscosity
does not change with time, indicating the stability of the
solution, which is an important property for obtaining
uniform fibers. While the viscosity of the solution pre-
pared by dissolving 6 % PVA is very low for spinability,
it is very high for the solution prepared by dissolving 10
% PVA. However, the viscosity of the solution prepared
by dissolving 8 % PVA is convenient for spinability.
Figure 2 shows SEM images of heat-treated
Al18B4O33 nanofibers that were prepared with solutions
that have different concentrations. It is clear that the
fibers that were obtained from the 6 % PVA solutions are
very thin, but the amount of them is very small. On the
other hand, the diameter of the fibers produced from the
10 % PVA solution is very large. These results were con-
sistent with the viscosity analyses. As a result, the solu-
tion prepared with 8 % PVA was determined to be the
optimum solution and used in the following process.
The DTA/TG analysis of he prepared alumina-bo-
rate/PVA solution is shown in Figure 3. The DTA curve
shows endothermic peaks between 100 °C and 300 °C
that indicates the loss of citric acid, which gets stabilized
with boric acid in the aluminium acetate. Degradation of
the PVA occurs between 300 °C and 600 °C in four
steps. A broad endothermic peak in the range 900 °C to
M. OZDEMIR et al.: EFFECT OF VISCOSITY ON THE PRODUCTION OF ALUMINA BORATE NANOFIBERS ...
736 Materiali in tehnologije / Materials and technology 47 (2013) 6, 735–738
Figure 1: Viscosity of the solutions prepared with mass fractions 6 %,
8 % and 10 % PVA solutions
Slika 1: Viskoznost raztopine, pripravljene z masnimi dele`i 6 %, 8 %
in 10 % PVA
1200 °C indicates the crystallization of the Al18B4O33.
According to the studied TG curve, the weight loss of the
alumina-borate/PVA solution was determined to be
approximately 63 % and it was followed by the removal
of the PVA from the sol structure.
Figure 4 shows the FTIR absorption peaks of the
PVA, the alumina-borate/PVA composite fibers and the
alumina borate fibers heat treated at 800 °C, 1000 °C and
1200 °C. A broad peak at about 3400 cm–1 corresponds
to the H-OH stretch. The peaks between 1300 cm–1 and
1700 cm–1 correspond to the characteristic PVA peaks in
Figure 4a.3 When the PVA is mixed with Al acetate that
is stabilized with boric acid the absorption peaks show a
change (Figure 4b). This indicates that there is an inte-
raction between them. After a heat treatment at 800 °C,
there are no PVA peaks because of the degradation of the
polymer (Figure 4c). When the calcination temperature
is increased to 1000 °C and 1200 °C (Figures 4d and
4e), a class of peaks appeared at 650–950 cm–1 and
1200– 1500 cm–1. These peaks can be attributed to the
formation of crystalline alumina borate.3 Furthermore,
all the FTIR graphs show a sharp peak at about 2400
cm–1, which corresponds to the CO2 peaks that appear
because of the ambient atmosphere during the analyses.
Figure 5 presents a typical XRD pattern of the
products before and after the heat treatment at different
temperatures and the pure PVA powder. The characte-
ristic peak of the PVA at 2 = 20° is shown in Figures 5a
and 5b. When the alumina borate/PVA composite fibers
M. OZDEMIR et al.: EFFECT OF VISCOSITY ON THE PRODUCTION OF ALUMINA BORATE NANOFIBERS ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 735–738 737
Figure 2: SEM images of nanofibers prepared at: a) 6 %, b) 8 % and c) 10 % concentrations
Slika 2: SEM-posnetki nanovlaken, pripravljenih s koncentracijo: a) 6 %, b) 8 % in c) 10 %
Figure 4: FTIR spectra of: a) PVA, b) alumina-borate/PVA composite
fibers and alumina borate fibers heat treated at: c) 800 °C, d) 1000 °C
and e) 1200 °C
Slika 4: FTIR-spektri: a) PVA, b) aluminijevooksidni borat/PVA kom-
pozitna vlakna in aluminijevooksidnoboratna vlakna, toplotno obde-
lana pri: c) 800 °C, d) 1000 °C in e) 1200 °C
Figure 3: DTA/TG analyses of the prepared solutions prepared with
the mass fraction 8 % PVA solution
Slika 3: DTA/TG-analiza pripravljenih raztopin, pripravljenih z
masnim dele`em PVA raztopine 8 %
Figure 5: XRD pattern of: a) PVA, b) alumina-borate/PVA composite
fibers and alumina borate fibers heat treated at: c) 800 °C, d) 1000 °C,
e) 1200 °C
Slika 5: XRD-posnetek: a) PVA, b) aluminijevooksidni borat/PVA
kompozitna vlakna in aluminijevooksidnoboratna vlakna, toplotno
obdelana pri: c) 800 °C, d) 1000 °C, e) 1200 °C
were heat treated at 800 °C, this peak disappeared and
the crystalline Al4B2O9 peaks occurred. The peak posi-
tions and the relative intensities match well with those of
the fiber Al4B2O9 (JCPDS 00-047-0319). Nevertheless,
as the intensity of the peaks increased with the increase
in the calcination temperature from 800 °C to 1000 °C,
the crystalline Al18B4O33 formed. At a calcination tempe-
rature of 1200 °C, all the XRD peaks belong to
crystalline Al18B4O33, which matches with JCPDS
00-053-1233 and 00-032-0003.
4 CONCLUSION
In summary, this study was structured to use the elec-
trospinning technique to produce alumina borate/PVA
composite nanofibers. In addition, the effect of viscosity
on the spinability and the morphology of the alumina
borate/PVA nanofibers were investigated. After a heat
treatment at 1200 °C, alumina borate (Al18B4O33) fibers
were produced.
On the whole, we found that the viscosity of the pre-
pared solution affects the fiber morphology. The increase
in the viscosity depends on the increase of the con-
centration of the solution, while the spinability of the
solution adopts an opposing relation with the dense
solution, which leads to clogging of the needle tip. As a
conclusion, the viscosity of the PVA polymer solution
has a very important effect on the diameters of the nano-
fibers. In addition, the phase structure of the aluminium
borate depends on the calcination temperature.
Acknowledgment
This work was supported by BOREN (National
Boron Research Institute) and TUBITAK (Scientific and
Technical Research Council of Turkey) project number
105M363. Also, Mehtap Ozdemir acknowledges support
from TUBITAK in the framework of the National
Scholarship Programmes for PhD Students.
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Qi, Materials Chemistry and Physics, 101 (2007), 499–504
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Qi, Materials Research Bulletin, 42 (2007), 482–486
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duction to Electrospinning and Nanofibers, World Scientific
Publishing Company, United States of America 2005, 69
M. OZDEMIR et al.: EFFECT OF VISCOSITY ON THE PRODUCTION OF ALUMINA BORATE NANOFIBERS ...
738 Materiali in tehnologije / Materials and technology 47 (2013) 6, 735–738
M. KOLE@NIK et al.: EFFECTS OF SOLIDIFICATION PARAMETERS ON THE MICRO- ...
EFFECTS OF SOLIDIFICATION PARAMETERS ON THE
MICRO- AND MACROSTRUCTURE OF THE
X19CrMoVNbN11-1 STEEL
VPLIVI PARAMETROV STRJEVANJA NA MAKRO IN
MIKROSTRUKTURO JEKLA X19CrMoVNbN11-1
Mitja Kole`nik1, Ale{ Nagode2, Grega Klan~nik2, Jo`e Medved2, Bernarda Janet1,
Milan Bizjak2, Ladislav Kosec2
1Metal Ravne, d. o. o., Koro{ka cesta 14, 2390 Ravne na Koro{kem, Slovenia
2Faculty of Natural Sciences and Engineering, University of Ljubljana, A{ker~eva cesta 12, 1000 Ljubljana, Slovenia
mitja.koleznik@metalravne.com
Prejem rokopisa – received: 2012-11-27; sprejem za objavo – accepted for publication: 2013-04-15
Differential scanning calorimetry (DSC) of the X19CrMoVNbN11-1 steel was used for determining phase-transformation
temperatures. To study the effect of the cooling rate on the shape and chemical composition of primary precipitates with SEM
and EDXS, an investigation was made on DSC samples. With a decreasing cooling rate precipitates became shorter and thicker
and with a higher cooling rate the contents of Nb and V increased.
The effects of solidification parameters on micro-segregations were investigated on the re-melted steel solidified at different
parameters (experimental ingots). The microstructure was examined using SEM and EDXS. A higher volume fraction of
segregation is present in the more rapidly cooled steel. In positive segregations the largest increase of the Cr content and a minor
increase in the concentrations of Mo and V were observed. The effect of melt superheating on the columnar and equiaxed
crystallization zone shares was also investigated. At higher melt temperatures a higher fraction of columnar crystals was
present.
Keywords: solidification parameters, transformation temperatures, precipitate shape, segregation, melt superheating,
crystallization zones
Za dolo~itev temperatur faznih pretvorb v jeklu X19CrMoVNbN11-1 smo izvedli diferen~no vrsti~no kalorimetrijo (DSC). Pri
vzorcih smo z vrsti~nim elektronskim mikroskopom (SEM) in energijsko disperzijsko spektroskopijo rentgenskih `arkov
(EDXS) preiskovali vpliv razli~nih hitrosti ohlajanja na obliko in kemijsko sestavo primarnih karbonitridov. Z manj{o hitrostjo
ohlajanja so bili izlo~ki kraj{i in debelej{i. Vsebnost Nb in V v izlo~kih pa je nara{~ala z ve~jo hitrostjo ohlajanja.
Izdelani so bili eksperimentalni ingoti, ki so bili strjeni v razli~nih razmerah. Preiskovali smo vpliv razmer pri strjevanju na
razse`nost izcejanja, ki smo ga ugotavljali z vrsti~nim elektronskim mikroskopom in EDXS-analizo. Obseg izcejanja je bil ve~ji
pri jeklu z ve~jo hitrostjo ohlajanja. V pozitivnih izcejah se je najbolj pove~ala koncentracija Cr, v manj{em obsegu pa se je
pove~ala tudi koncentracija Mo in V. Preiskali smo tudi vpliv pregretja taline na dele` kristalizacijskih podro~ij s stebrastimi in
enakoosnimi zrni. Z vi{jo temperaturo pregretja taline se je pove~al dele` stebrastih kristalnih zrn.
Klju~ne besede: vpliv parametrov strjevanja, temperature premen, oblika izlo~kov, izcejanje, pregretje taline, kristalizacijska
podro~ja
1 INTRODUCTION
The vital parts (like the components for steam and
gas turbines, steam pipes, boilers, etc.) for thermal power
plants must fulfil very high quality standards. The most
widely used creep-resistant steels in power plants are
martensitic steels with 9 % to 12 % Cr. They have the
best combination of physical and elastic properties, a
high creep strength, a high resistance against thermal
fatigue and a high steam-oxidation resistance.1–5 To
ensure the adequate properties the minimum structure
homogeneity right from the start of a manufacturing
process is necessary.
During solidification in industrial conditions,
non-equilibrium conditions prevail resulting in chemical
inhomogeneities on the macro- and micro-scale (macro-
and micro-segregations).6–8 Micro- or crystal segrega-
tions represent a change in the concentration of alloying
elements across a grain cross-section with low melting
elements on the grain boundaries. The smallest concen-
tration of alloying elements is expected to be found in
the core of dendrites. Since the mobility of atoms in the
solid phase is rather small and the diffusion times are
short, chemical inhomogeneities remain to room tempe-
rature. The intensity of segregations and the microstruc-
tural evolution in an as-cast sample depends on the
solidification rate. Casting parameters also affect the
volume fraction of individual crystallization zones in
solidified ingots. The anisotropy of columnar zones can
decrease the ductility of steel during a plastic
deformation.9–11
2 EXPERIMENTAL WORK
The alloy used in this study was the creep-resistance
steel X19CrMoVNbN11-1 with a typical chemical
composition shown in Table 1. Differential scanning
calorimetry was used for determining the characteristic
Materiali in tehnologije / Materials and technology 47 (2013) 6, 739–744 739
UDK 669.14:669.112.221:536.421.4 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)739(2013)
temperatures of solidification, using a Netzsch STA 449
C Jupiter. In Table 2 the parameters of DSC analyses are
given. DSC analyses were performed in a protective
atmosphere of volume fraction 99.999 % N2 and, as a
reference, an empty corundum crucible was used. DSC
specimens were also used to determine the effects of
solidification parameters on the microstructure.
Thermodynamic calculations were performed using the
ThermoCalc program with an internal thermodynamic
database, TCFE7.
Table 2: Parameters of an DSC analysis
Tabela 2: Parametri DSC-analize
sample Tmax/°C vheating /(K min–1)
vcooling/
(K min–1)
S1 1550 1 1
S25 1550 25 25
S6P 1600 6 6
S6 1550 6 6
Table 3: Casting parameters of experimental ingots
Tabela 3: Parametri litja eksperimentalnih ingotov
No. mload/kg Tmold/°C Tcasting/°C Moldisolation
1 11 318 1547 no
2 10.97 280 1560 yes
3 11.02 330 1690 yes
4 8.75 30 1550 no
The experimental ingots of the X19CrMoVNbN11-1
creep-resistant steel were prepared by melting the steel
in an inductive vacuum furnace. The steel was solidified
under different cooling conditions. To achieve different
cooling rates of the castings, different preheated moulds
were used. The casting parameters are shown in Table 3
and the cooling rates of the ingots are shown in Figure 1.
The samples for the investigation of the macro- and
microstructure were cut from the ingots without heat
treatment. For the metallographic examination of the
macrostructure, the steel plates were grinded and etched
in a hot 50 % HCl solution at 50 °C. The samples for the
microstructural investigation were grinded, polished and
etched with the Villela etchant.
The metallographic investigation of the steel was
carried out on a scanning electron microscope (SEM),
JEOL JSM-5610 with an energy-dispersive x-ray spec-
trometer (EDXS), GRESHAM Scientific Instruments,
model SIRIUS 1YSUTW. The results are given as a
quantitative spot (EDXS spectrum) and qualitative
EDXS line profile and as mapping analyses.
2.1 Thermodynamic model for M(C, N)
In the CALPHAD approach each phase should be
described with a proper thermodynamic model. In this
paper emphasis is placed only on the thermodynamic
modelling of M(C, N). M stands for the metals mixing
on the first sublattice, sublattice #1. Ma1(C, N)a2 is
modelled with two sublattice models with the second
sublattice representing the interstitial atoms of C and N.
Va stands for vacancy. The carbonitrides present in steels
can be modelled with a combination of Fe, Nb, Ti and
others on the first sublattice, and C, N and Va on the
second one; an example is given for (Fe, Nb, Ti)a1(C, N,
Va)a2. The Gibbs free energy, Gm, of such a formula for
one formula unit is described with the following
equation:12
Gm = yFeyc0GFe:C + yNbyc0GNb:C + yTiyc0GTi:C +
+ yFeyN0GFe:N + yNbyN0GNb:N + yTiyN0GTi:N +
+ yFeyVa0GFe:Va + yNbyVa0GNb:Va + yTiyVa0GTi:Va +
+ a1RT(yFe lnyFe + yNb lnyNb + yTi lnyTi) +
+ a2 RT(yC lnyC + yN lnyN + yVa lnyVa) + EGm
where yi represents the site fractions of component i on
the first or second sublattice (I or II). G i j:
0 is the Gibbs
free energy where only i and j are occupying the sub-
lattices. Gm
E represents the excess Gibbs energy. In
addition, the term for magnetic contribution can also be
added. For a more complex formula with more elements
on the first sublattice the Gibbs energy is described
below; here a1 or a2 represent the site occupancy:
G y y G
RT a y y a y y
i
ji
j i j
i
i
i j
j
m = +
+ +
∑∑
∑ ∑
ª ªª
:
ª ª ªªln( ) ln(
0
1 2 j EG
ªª )
⎛
⎝
⎜
⎞
⎠
⎟ + m
M. KOLE@NIK et al.: EFFECTS OF SOLIDIFICATION PARAMETERS ON THE MICRO- ...
740 Materiali in tehnologije / Materials and technology 47 (2013) 6, 739–744
Table 1: Typical chemical composition of the X19CrMoVNbN11-1 steel in mass fractions, w/%
Tabela 1: Tipi~na kemijska sestava jekla X19CrMoVNbN11-1 v masnih dele`ih, w/%
C Si Mn Cr Mo Ni V Nb N B
0.19 0.30 0.55 10.75 0.75 0.5 0.20 0.33 0.08 max. 0.0015
Figure 1: Cooling rates of ingots
Slika 1: Hitrost ohlajanja posameznega ingota
In this study the modelling was done using the
following formula for carbonitride with an fcc structure:
(Nb, Cr, Fe, Mn, Mo, Ni, Si, V)1(C, N, Va)1
where in sublattice #1 the major constituents are Nb,
Cr, V and in sublattice #2 the major constituents are C
and N, with Va as a minor constituent.
3 RESULTS AND DISCUSSION
Differential scanning calorimetry was used for
determining transformation temperatures. The increasing
heating rate influenced the phase transformation at
higher temperatures and with the increasing cooling rate
these transformations were moved to lower temperatures.
Table 4 shows the Ar1 temperatures of the solidification
curves.
Table 4: Phase-transformation temperatures
Tabela 4: Temperature faznih premen
sample S1 S6 S25 S6P
Tmax/°C 1550 1550 1550 1600
vs/vo 1/1 6/6 25/25 6/6
Ar1/°C 812 823 835 821
Tliquidus/°C 1497 1510 1462 1495
Tsolidus/°C) 1492 1493 1443 1480
The metallographic investigation of DSC samples
revealed that the cooling rate has a great effect on the
shape and chemical composition of the precipitates.
Combining the Thermo-Calc software and EDXS ana-
lysis, the precipitates were identified as Nb(C, N) with
low contents of V and Mo. At the highest cooling rate of
25 K min–1, the precipitates were needle shaped with a
diameter of approximately 0.5 μm and a length of up to
35 μm with a faceted surface. At a lower cooling rate, the
precipitates were shorter and thicker, and their surface
was less faceted. At the cooling rate of 6 K min–1, the
diameter of the precipitates was approximately 1 μm and
the length was approximately 15 μm. The precipitates
formed at different cooling rates are shown in Figures 2
and 3. A micro-chemical analysis of the precipitates
indicated that at higher cooling rates the contents of Nb
and V increased Figure 4 shows the concentration ratio
of Nb and V w(Nb)/w(V) in the precipitates as a function
of the cooling rate. At higher cooling rates the
w(Nb)/w(V) ratio was increased indicating that, at a
higher cooling rate, the content of Nb in the precipitates
increased faster than the content of V.
A comparison of the macro- and microstructures of
the experimental ingots was made. In all the ingots three
distinct crystallization zones were found. In the central
region there was a zone of fine, randomly oriented,
equiaxed grains and, in the outer region, a zone of
columnar grains was observed. On the ingot surfaces, a
zone of very fine equiaxed grains (chill crystals) was
observed. The degree of melt superheating has the
strongest influence on the shares of the columnar and
equiaxed zones in the experimental ingots (in our case
the temperatures of melt superheating and the casting
temperatures from Table 3 are equal). The smallest share
of the columnar zone was found in ingot 1 with the melt
heated to the lowest temperature (1547 °C). The highest
share of the columnar zone was found in ingot 3 with the
overheating of up to 1690 °C. Table 5 shows the shares
M. KOLE@NIK et al.: EFFECTS OF SOLIDIFICATION PARAMETERS ON THE MICRO- ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 739–744 741
Figure 4: Concentration ratio of Nb and V in precipitates as a
function of the cooling rate
Slika 4: Razmerje koncentracije Nb in V v izlo~kih v odvisnosti od
hitrosti ohlajanja
Figure 2: Shape of precipitates at 25 K min–1
Slika 2: Oblika izlo~kov pri hitrosti ohlajanja 25 K min–1
Figure 3: Shape of precipitates at 6 K min–1
Slika 3: Oblika izlo~kov pri hitrosti ohlajanja 6 K min–1
of the columnar zones (wtrans/%) in each ingot. In Figure
5 a comparison of the macrostructures of ingots 2 and 3
can be observed together with the corresponding dia-
meters of the equiaxed zones. At a higher degree of melt
superheating more heterogeneous nuclei dissolve. The
smaller is the concentration of heterogeneous nuclei in a
melt, the higher the degree of supercooling that is needed
for the solidification to start. A higher degree of super-
cooling increased the share of the columnar zone in a
solidified ingot.
Table 5: Share of the columnar zone in ingots
Tabela 5: Dele` podro~ja stebrastih zrn v posameznem ingotu
ingot 1 ingot 2 ingot 3 ingot 4
Tcasting/°C 1547 1560 1690 1550
wtrans/% 88.9 90.7 95.5 89.9
The cooling curves of experimental ingots are shown
in Figure 1. It can be seen that every curve shows three
distinct regions. Region I with the highest slope repre-
sents a very intense cooling of the ingots immediately
after casting when the melt comes into contact with the
wall of the colder metallic mold. However, the steepest
slope shows ingot 3 with the highest casting temperature.
In region II the cooling rates slow down due to the
generated crystallisation heat. In this region the cooling
rates depend on the casting parameters (the mold tem-
perature and the mold isolation) (Table 3). In this region,
the highest drop in the temperature is observed for ingot
4 with the coldest mold (not preheated) and without
isolation. In region III with the stationary cooling having
the lowest cooling rates, a small difference in the slopes
of the cooling curves are found except for ingot 4 with
an unpreheated mold and without isolation. In all the
investigated ingots crystal segregations were observed
with the segregation intensity (the volume fraction)
depending on solidification parameters. The most intense
segregation was found in ingot 4, cast in a cold mold (the
highest cooling rate). The least intense segregation was
found in ingot 2, cooled with the lowest rate. In the
region with the maximum and minimum contents of
alloying elements, a micro-chemical analysis was carried
out. In Figure 6 the EDXS micro-chemical analysis of
ingot 4 is presented and the chemical composition of the
selected areas is given in Table 6. Differences in the
chemical composition affect the intensity of etching.
Darker areas on Figure 6 are positive segregations with
higher chromium contents. The concavity of positive
segregations indicates that they were solidified last with
the formation of one carbide eutectic with the lowest
melting point (Figure 7). The largest variation between
the alloying elements was measured for the concentra-
tion of Cr. The difference between the maximum and
minimum Cr contents was much more explicit for the
rapidly cooled ingot. For ingot 2 the difference in the
concentration of Cr was approximately 2.5 % and for
M. KOLE@NIK et al.: EFFECTS OF SOLIDIFICATION PARAMETERS ON THE MICRO- ...
742 Materiali in tehnologije / Materials and technology 47 (2013) 6, 739–744
Figure 6: EDXS analyis of the microstructure of ingot 4. Chemical
compositions of areas 1 and 2 are in Table 6. The darker area – higher
contents of alloying elements.
Slika 6: Mikrokemijska analiza (EDXS) ingota 4. Kemijska sestava
podro~ja 1 in podro~ja 2 je podana v tabeli 6. Temnej{a obmo~ja –
pove~ana koncentracija zlitinskih elementov.
Table 6: Chemical composition (EDXS analysis) of segregations in
mass fractions, w/%; see also Figure 6
Tabela 6: Kemijska sestava (EDXS-analiza) izcej v masnih dele`ih,
w/%; glej sliko 6
Fe Si Cr Mn Mo Ni V Nb
Area 1 84.45 0.29 13.10 0.32 0.66 0.62 0.29 0.27
Area 2 89.24 0.26 9.03 0.36 0.42 0.48 0.18 0.03
Figure 7: Positive segregations with carbide eutectic
Slika 7: Pozitivne izceje s karbidnim evtektikom
Figure 5: Equiaxed and columnar zones in ingots 2 and 3
Slika 5: Primerjava dele`a podro~ij s stebrastimi in enakoosnimi zrni
v ingotih 2 in 3
ingot 4 the difference was approximately 4 % (mass
fractions). The segregation of Cr in each ingot was
characterised with the segregation index Is (1) and
segregation rate Isg (2):
I
C
Cs
= max
min
(1)
I
C C
Csg
=
−max min
0
(2)
where Cmax is the maximum, and Cmin is the minimum
content, while C0 is the average content of Cr in each
ingot. The results are shown in Table 7. The concen-
trations of V and Mo were also increased in positive
segregations.
Table 7: Mass index and level of the Cr segregation in ingots
Tabela 7: Masni indeks ter stopnja izcejanja Cr za posamezen ingot
Cmax/% Cmin/% C0/% Is Isg/%
ingot 1 11.12 9.47 9.98 1.17 16.54
ingot 2 12.88 10.41 10.24 1.24 24.44
ingot 3 11.83 10.10 10.10 1.17 17.12
ingot 4 13.10 9.03 10.02 1.45 40.63
The change in the concentration of the alloying
elements in the dendrite cross-section was determined
with the qualitative EDXS line analysis of ingot 4 in
Figure 8. The line intersects one positive segregation in
two areas (between 20–40 μm and between 56–68 μm
along with the carbide at approximately 62). There is a
significant increase in the concentrations of Cr and Mn
and a minor increase in the concentrations of V in the
areas of positive segregation. At the carbide intersection,
a large increase in the carbide-forming elements, like
Nb, Mo and V, is present. In the microstructure in Figure
8 carbide eutectic is also presented. In Figure 9 a
qualitative mapping analysis of carbide eutectic is
shown, indicating increased contents of Mo, Nb and V
like for the carbide in the line analysis from Figure 8.
4 CONCLUSIONS
To determine the influence of solidification para-
meters, the macro- and microstructure as well as the
chemical homogeneity of the creep-resistant steel of
grade X19CrMoVNbN11-1 were investigated.
– The heating and cooling rates affect the solidifica-
tion interval. By increasing the heating rate, the
solidification interval is shifted to a higher tem-
perature, while with a higher cooling rate the solidi-
fication interval is shifted to a lower temperature.
– The cooling rate affects the shape and chemical
composition of precipitates. At the higher cooling
rates the precipitate shapes were long faceted
needles, and at a lower cooling rate the precipitates
were shorter and thicker. They were identified as the
Nb(C, N)-type precipitates.
– At the higher cooling rates the contents of Nb and V
were higher.
– The degree of melt overheating influenced the
respective proportions of columnar and equiaxed
zones. With higher melt superheating the share of
the columnar zones increases.
– At a higher cooling rate the intensity of segregation
and the intensity of positive segregations of Cr were
greater.
– In the investigated steel the last solidified melt had a
eutectic composition.
M. KOLE@NIK et al.: EFFECTS OF SOLIDIFICATION PARAMETERS ON THE MICRO- ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 739–744 743
Figure 8: Qualitative EDXS line analysis of ingot 4
Slika 8: Kvalitativna ~rtna EDXS-analiza ingota 4
Figure 9: Qualitative EDXS mapping analysis of the segregations
with carbide eutectic
Slika 9: Kvalitativna ploskovna EDXS-analiza izcej s karbidnim
evtektikom
5 REFERENCES
1 A. Nagode, L. Kosec, B. Ule, G. Kosec, Review of creep resistant
alloys for power plant applications, Metalurgija, 50 (2011) 1, 45–48
2 F. Vodopivec, D. A. Skobir, M. Jenko, B. @u`ek, M. Godec, Nucle-
ation and growth of M23C6 particles in high-chromium creep resistant
steel, Mater. Tehnol., 46 (2012) 6, 633–636
3 D. Rojas Jara, 9–12 % Cr heat resistant steels: Alloy design, TEM
characterisation of microstructure evolution and creep response at
650 °C, Fakultät für Maschinenbau der Ruhr, Universität Bochum,
2011
4 R. L. Klueh, R. H. Donald, High-Chromium Ferritic and Martensitic
Steels for Nuclear Applications, ASTM International, (2001), 5–27
5 F. Abe, T. Kern, R. Viswanathan, Creep-resistant steels, Woodhead
Pub. Ltd, Cambrige, England 2008
6 W. Kurz, D. J. Fisher, Fundamentals of solidification, Trans Tech
Publications Ltd, Aedermannsdorf 1986
7 University of Cambridge, Casting: Microstructure and segregation in
castings [online]. 2004, [cited 2012-07-26]. Available from World
Wide Web: http://www.doitpoms.ac.uk/tlplib/casting/microsegre-
gation.php
8 S. Spai}, Fizikalna metalurgija: zgradba kovinskih materialov, strje-
vanje kovinskih zlitin, Naravoslovnotehni{ka fakulteta, Univerza v
Ljubljani, Ljubljana 2002
9 M. Durand-Charre, Microstructure of Steels and Cast Irons, Springer,
Paris 2003
10 J. W. Elmer, S. M. Allen, T. W. Eagar, Microstructural development
during solidification of stainless steel alloys, Metallurgical transac-
tions, 20 (1989), 2117–2131
11 H. Biloni, W. J. Boettinger, Solidification, In: R. W. Chan, P. Haasen
(eds.), Physical Metallurgy, Elsevier Science BV, Amsterdam 1996,
670–842
12 B. J. Lee, Thermodynamic Assessment of the Fe-Nb-Ti-C-N System,
Metallurgical and Materials Transactions A, 32 (2001), 2423–2439
M. KOLE@NIK et al.: EFFECTS OF SOLIDIFICATION PARAMETERS ON THE MICRO- ...
744 Materiali in tehnologije / Materials and technology 47 (2013) 6, 739–744
M. NESLU[AN et al.: INFLUENCE OF CUTTING SPEED ON INTENSITY OF THE PLASTIC DEFORMATION ...
INFLUENCE OF CUTTING SPEED ON INTENSITY OF
THE PLASTIC DEFORMATION DURING HARD
CUTTING
VPLIV HITROSTI REZANJA NA INTENZITETO PLASTI^NE
DEFORMACIJE MED ODREZOVANJEM
Miroslav Neslu{an1, Ivan Mrkvica2, Robert ^ep2, Mária ^illiková1
1University of @ilina, Faculty of Mechanical Engineering, @ilina, Slovakia
2V[B-Technical University of Ostrava, Ostrava, Czech Republic
miroslav.neslusan@fstroj.utc.sk, robert.cep@vsb.cz
Prejem rokopisa – received: 2013-01-29; sprejem za objavo – accepted for publication: 2013-03-27
The paper deals with an analysis of deformation processes and the aspects related to a chip formation such as the chip thickness,
the chip ratio, the shear angle and the chip segmentation during turning the hardened steel 100Cr6. This paper investigates the
influence of the cutting speed through a metallographic analysis, a calculation of significant aspects of deformation processes
and the resulting experimental study. This experimental study is based on an application of acoustic emission. The results of this
study indicate that the cutting speed significantly affects the parameters such as the chip ratio, deformation angle or the speeds
in the cutting zone. On the other hand, the experimental study allows an analysis of the specific characters of deformation
processes and their real intensity.
Keywords: turning, hardened steel, chip segmentation, acoustic emission
^lanek obravnava analizo deformacijskih procesov in vidike pri nastanku odrezka, kot so debelina odrezka, odvzem, rezalni kot
in segmentacija odrezkov med stru`enjem kaljenega jekla 100Cr6. V ~lanku je obravnavan vpliv hitrosti odrezovanja z
metalografsko analizo, izra~un pomembnih vidikov deformacijskih procesov in temu slede~a eksperimentalna {tudija. Le-ta
temelji na uporabi akusti~ne emisije. Rezultati te {tudije ka`ejo pomemben vpliv hitrosti rezanja na parametre, kot so odvzem,
deformacije in hitrosti v coni rezanja. Po drugi strani eksperimentalna {tudija omogo~a analizo posebnosti deformacijskih
procesov in njihove realne intenzivnosti.
Klju~ne besede: stru`enje, kaljeno jeklo, segmentacija ostru`kov, akusti~na emisija
1 INTRODUCTION
With the development of super-hard cutting materials
such as ceramics, CBN, PCBN and PCD, etc., the tech-
nology of hard turning has attracted considerable interest
from several leading manufacturers. In order to specify
the potential of this new production technology, several
issues of hard turning such as cutting mechanism, tool
wear, machined surface integrity, etc., have been recently
investigated.1,2 One of the significant observations during
hard turning, which is different from machining ductile
materials, is that consistently cyclic, segmented chips are
produced, essentially without any deformation at the
shear plane.3–5
In metal cutting, the principal chip morphologies
classified by Komanduri and Brown,6 are continuous,
serrated chips. In machining, hard and difficult-to-
machine materials tend to localize the heat generated due
to the strain localization in a narrow band called an
adiabatic shear band represented in Figures 1 and 2.
Adiabatic shear banding investigated by Recht3 is used to
describe a localization phenomenon that occurs in the
high-strain-rate plastic-deformation processes such as
cutting. This phenomenon can be explained in the
following way: during a deformation the rate of heat
generation is determined by the strain rate, and the heat-
dissipation rate is controlled by the heat conductivity of
the material. When the heat-generation rate is larger than
the heat-dissipation rate, the temperature increases and
the softening associated with such temperature increases
as well, exceeding the strain hardening and the catastro-
phic propagation of the shear occurs. Thermoplastic
instability is a significant phenomenon in the segmented
chip formation. The thermoplastic instability is caused
by a decrease in the flow stress due to the thermal soften-
ing associated with an increase in the strain, offsetting
Materiali in tehnologije / Materials and technology 47 (2013) 6, 745–755 745
UDK 621.941:539.37 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)745(2013)
Figure 1: Illustration of a chip formation during hard turning
Slika 1: Prikaz nastanka odrezka pri stru`enju
the strain hardening.4 The materials sensitive to a
formation of a shear-localized chip can be characterized
with poor thermal properties (titanium alloys, maraging
steels) and a limited ductility (hardened steels).5
A proper understanding of the material-removal
mechanisms taking place during hard cutting is essential
for a process evaluation. An analysis of the work area is
necessary to describe the chip generation in hardened
materials. Depending on the cutting parameters and
workpiece material properties, cutting may either lead to
continuous or discontinuous chip formation.2,4,7 Conti-
nuous chips are formed during turning conventional soft
steels (Figure 3 – a decrease in the flow stress due to the
thermal softening associated with an increase in the
strain is less than the associated strain hardening4), while
hard turning can lead to a formation of segmented chips.
Figures 1 and 2 illustrate a segmented chip during turn-
ing the hardened steel 100Cr6. Figure 2 shows that the
plastic deformation inside the segment is low and the
material in this area stays nearly untouched. Deformation
processes are concentrated in the shear zone, tool-chip
contact and tool-workpiece contact. On the other hand,
the formation of a continuous chip leads to a more
homogenous dissipation of deformation processes across
the whole chip (Figure 3).
It was found that the total energies entering the cut-
ting processes of conventional and hard turning (derived
from the cutting speed and tangential component of the
cutting force) are nearly the same.8,9 The difference can
be found in the specific character related to the dissipa-
tion of this energy. An evaluation of the real intensity of
deformation processes should be carried out with respect
to the specific character of the chip formation. Intensity
of deformation processes can be evaluated from the ma-
croscopic point of view through metallographic obser-
vations. This conventional approach allows a calculation
of the conventional parameters based on the measure-
ments of chip thickness, undeformed chip thickness and
derived parameters (deformation angle, chip speed, chip
deformation and others). In this way, the average inten-
sity of deformation processes in the cutting zone (across
the whole chip) can be investigated.5,8
An investigation of the real intensity of deformation
processes related to the specific cutting zones is quite
difficult and can be realized through acoustic emission
(AE). The AE non-destructive technique is based on a
detection and conversion of these high-frequency elastic
waves to electrical signals. The major AE sources10,11 in
a metal-cutting process are:
• deformation and fracture of work materials in the
shear zone, tool-chip and tool-workpiece contacts;
• deformation and fracture of the cutting tool;
• collision, entangling and breakage of chips.
The main advantage of using AE to monitor a
machining operation is that the frequency range of an AE
signal is much higher than that of the machine vibrations
and environmental noises, and does not interfere with the
cutting operation.10 Due to its sensitivity to various con-
tact areas and deformation regions during cutting, the AE
signal is the basic tool for process monitoring.
Figures 2 and 3 show that the characters of defor-
mation processes during the conventional soft and hard
turnings differ. And so, the aim of this paper is to per-
form an analysis of the deformation processes in the
cutting zone through the conventional metallographic
observations and with the AE signal. Moreover, the
differences between the conventional soft turning and
hard turning should be investigated.
2 METALLOGRAPHIC OBSERVATIONS
The general aim of taking metallographic chip sam-
ples was to analyze the significant parameters of defor-
mation processes in the cutting zone during hard turning
and investigate the segmentation frequency in the chip
(and comparing it with the frequency analysis from the
accelerometer). The same analysis (except for the seg-
mentation frequency) was carried out for soft steel with a
view to comparing deformation processes in the cutting
zone with respect to different properties of the machined
material. The experimental setup and cutting conditions
are listed in Table 1. Figure 4 shows the chips, on which
periodic cracks can be observed. These series of seg-
ments are measured all over the chip, and from more
than 20 values the mean values are calculated. In the
extensive cutting tests, the segmentation length and chip
M. NESLU[AN et al.: INFLUENCE OF CUTTING SPEED ON INTENSITY OF THE PLASTIC DEFORMATION ...
746 Materiali in tehnologije / Materials and technology 47 (2013) 6, 745–755
Figure 3: Chips after turning the 100Cr6 (annealed), SEM, vc = 100
m min–1, f = 0.21 mm
Slika 3: Odrezki po stru`enju 100Cr6 (`arjeno), SEM, vc = 100 m min–1,
f = 0,21 mm
Figure 2: Chips after turning the 100Cr6 (hardened, 62 HRC), SEM,
vc = 100 m min–1, f = 0.21 mm
Slika 2: Odrezki po stru`enju 100Cr6 (kaljeno, 62 HRC), SEM, vc =
100 m min–1, f = 0,21 mm
thickness were statistically established. The parameters
such as the chip ratio, the chip deformation, the chip
speed, the shear speed, or the deformation angle were
calculated. Moreover, the segmentation frequency could
be evaluated. To obtain the segmentation frequency, the
segmentation length had to be measured with a metallo-
graphic microscope and, knowing the cutting speed, the
shear plane speed or the chip speed, and the frequency
could be calculated.
Table 1: Experimental conditions during hard turning
Tabela 1: Eksperimentalni pogoji pri stru`enju
Cutting tool: TiC reinforced Al2O3 ceramic inserts
DNGA150408 (TiN coating),
rake angle n= –7°
Work material: 100Cr6 (hardened, 62 HRC and
annealed, 27 HRC), external diameter of
56 mm, internal diameter of 40 mm, 125
mm long
Cutting
condition:
vc = 25–250 m min
–1, f = 0.09 mm,
ap = 0.25 mm (constant), dry cutting
Machine tool: CNC Lathe Hurco TM8
The chip thickness was measured with a light microscope
Intensity of a plastic deformation in the cutting zone
can be expressed in terms of various parameters. Except
for the chip ratio, there are parameters such as the degree
of segmentation (G) or chip deformation (sh) and other
derived parameters such as the chip and shear speeds.
The measurement of the chip thickness was carried
out with an optical microscope and through the measure-
ment of the etched chips shown on Figure 4.
The measurement of the chip thickness (hc) allows us
to calculate the chip ratio (equation (1)) and the other
related parameters such as the shear angle 1, the chip
speed (vch), the shear speed (vsh) and the chip deforma-
tion (sh):
K
h
h
= c (1)
where h is the undeformed-chip thickness and hc is the
chip thickness. The deformation angle can be calculated
through equation (2):
tg
sin
n
n
Φ1 = −
cos 
K
(2)
The chip speed and shear speed can be derived from
the cutting speed and shear angle (equations (3) and (4)):
v vch c
1
1 n
=
−
sin
cos( )

 
(3)
A chip deformation can be expressed in a similar
way, through equation (5):
v vsh c
n
1 n
=
−
cos
cos( )

 
(4)


  sh
n
1 n 1
=
− ⋅
cos
cos( ) sin
(5)
A degree of segmentation is calculated through
equation (6) and illustrated in Figures 4 and 5:
G
h h
h
=
−c 0
c
(6)
Figure 6 illustrates that the chip thickness during
turning hardened steel is much lower than that during
turning annealed steel. The metallographic observations
verify the previous investigations performed under simi-
lar conditions.12–14 The formation of the segments during
turning hardened steel causes an elongation and a
decrease in the chip thickness. As a result of the for-
mation of thin and long chips, the chip ratio is smaller
than 1, contrary to the turning of annealed steel (thick
and short continuous chips causing the chip ratio to be
M. NESLU[AN et al.: INFLUENCE OF CUTTING SPEED ON INTENSITY OF THE PLASTIC DEFORMATION ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 745–755 747
Figure 4: Chips after hard turning: a) vc = 25 m min–1, f = 0.09 mm,
b) vc = 50 m min–1, f = 0.09 mm, c) vc = 100 m min–1, f = 0.09 mm, d)
vc = 200 m min–1, f = 0.09 mm
Slika 4: Odrezki po stru`enju: a) vc = 25 m min–1, f = 0,09 mm, b) vc
= 50 m min–1, f = 0,09 mm, c) vc = 100 m min–1, f = 0,09 mm, d) vc =
200 m min–1, f = 0,09 mm
more than 1, Figure 7). The average intensity of plastic
deformation expressed with the chip ratio is much lower
during hard turning than during the turning of soft steel.
The low intensity of plastic deformation must be
attributed, first of all, to the material inside the segment.
The plastic deformation inside the segment is low and
the material in this area stays untouched. Although the
plastic deformation in the localized areas of the segmen-
ted chip is extremely high (white areas), the total defor-
mation of the segmented chip is much lower than for a
continuous chip (during turning the annealed steel), as
seen on Figure 8. On the other hand, the intensity of
plastic deformation significantly changes with the cut-
ting speed in the case of hard turning. The segmented
chip becomes more continuous with a decreasing cutting
speed. This aspect is verified with the degree of segmen-
tation illustrated in Figure 9. The continuous chip can be
expressed as a segmented chip with the zero degree of
segmentation. Figure 9 illustrates that the degree of
segmentation strongly decreases with a decreasing
cutting speed. The specific character of the chip for-
mation is related to the high shear angle, higher than the
shear angle for turning annealed steel (Figure 10). The
chip thickness, the chip ratio and the chip deformation
increase with a decreasing cutting speed because the
average intensity of plastic deformation increases (the
chip becomes more and more continuous).
M. NESLU[AN et al.: INFLUENCE OF CUTTING SPEED ON INTENSITY OF THE PLASTIC DEFORMATION ...
748 Materiali in tehnologije / Materials and technology 47 (2013) 6, 745–755
Figure 9: Influence of cutting speed on degree of segmentation
Slika 9: Vpliv hitrosti rezanja na stopnjo segmentacije
Figure 7: Influence of cutting speed on chip ratio
Slika 7: Vpliv hitrosti rezanja na razmerje odrezkov
Figure 5: Illustration for a calculation of chip segmentation
Slika 5: Prikaz segmentacije odrezkov za izra~un
Figure 6: Influence of cutting speed on chip thickness
Slika 6: Vpliv hitrosti rezanja na debelino odrezka
Figure 8: Influence of cutting speed on chip deformation
Slika 8: Vpliv hitrosti rezanja na deformacijo odrezkov
Figure 10: Influence of cutting speed on shear angle
Slika 10: Vpliv hitrosti rezanja na stri`ni kot
As a result of the formation of the thin and long chips
(in the case of turning hardened steel) the shear speed
and the chip speed are much higher than in the case of
turning annealed steel (Figures 11 and 12). This aspect
is attributed to the balance in the cutting zone between
the incoming and outgoing volumes of the material. The
increase in the chip ratio is strongly associated with a
decrease in the chip speed.
It can be easily observed that the specific process of a
chip segmentation leads to the thinning of the plastically
deformed region as the chip moves up the tool face.12,13
The significance of the thinning of the plastically de-
formed region as the chip moves up the tool face is that
this gives rise to a chip ratio less than one. This is
usually the case when hard steel is turned with a negative
rake tool. An important consequence is that the chip
speed will be greater than the cutting speed and the shear
angle will be greater than 45°.
3 EXPERIMENTAL RESULTS
The term segmented chip is often used to describe all
of the cyclic types. This is unfortunate since these types
of chips are distinctly different. For example, the cycle
frequency for a wavy chip is typically about 100 Hz,
while the one for a segmented chip is by 2 to 4 orders of
magnitude greater.13–15 Dynamic forces that fluctuate at a
frequency over 10 kHz are difficult to measure.
A conventional piezoelectric dynamometer limits the
frequency response to about 3.5 kHz. However, an esti-
mation of the relative changes in the force components
and the frequency of force fluctuation may be obtained
by using wire-resistance strain gauges15 or accelerome-
ters. A conventional accelerometer limits the frequency
response to about 20 kHz (special accelerometers limit
the frequency response to about 50 kHz).
On the other hand, acoustic-emission techniques
allow us to investigate the processes that fluctuate at a
frequency over several MHz. AE signals can be classi-
fied into two types,16 as either continuous-type AE
signals or burst-type AE signals. Continuous signals are
associated with the shearing in the primary zone and the
wear on the tool flank, while the burst type signals are
observed during a crack growth in a material, tool, mate-
rial fracture or chip breakage.
Apart from the conventional process monitoring, AE
can be applied to an analysis of a chip form and a chip
flow. The specific character of an AE signal can be ob-
served under the specific cutting conditions related to a
specific chip formation. Hard turning represents the case
of a specific character of deformation processes mixed
with a fracture of a machined material. There is a strong
relaxation character of the related AE signal during hard
turning. Uehara17 reported on remarkable patterns of the
AE related to the segmented-chip formation. The AE sig-
nals accompanying the formation of a segmented chip
exhibit remarkable patterns; the tool side signal shows a
periodic bursting. The amplitude of AE varies correspon-
dingly to the periodic change of the cutting force.
The elastic waves arising from different locations in
the cutting zone can be detected through an AE
system.10,17 The friction between the tool and the work-
piece in the primary shear zone generates a continuous
AE signal (during a formation of a continuous chip),
providing comprehensive information on the cutting
process. On the other hand, the chip formation during
machining hardened steel allows us to work out the
conditions for crack initiation and propagation. The sur-
faces that need to be machined are not perfectly smooth
but rough and composed of microscopic ridges, cracks,
voids, etc. Machining hardened materials using high
compressive stress creates a subsurface material flow
leading to the formation of cracks in the free surface.
Strong elastic waves related to the crack formation dur-
ing the segmented-chip formation can be detected
through the AE systems, being inherent to the AE signals
and related to the deformation processes in the cutting
zone. The character of the AE signals during hard turn-
ing is complicated because mixed types of AE signals
are generated in different zones. The aim of this experi-
mental study is to analyze these aspects.
The experimental setup is shown in Figures 13 and
14. Commercial piezoelectric AE sensors (D9241A – the
frequency range of 20 kHz to 180 kHz, WD – the fre-
quency range of 100 kHz to 1000 kHz) by Physical
M. NESLU[AN et al.: INFLUENCE OF CUTTING SPEED ON INTENSITY OF THE PLASTIC DEFORMATION ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 745–755 749
Figure 12: Influence of cutting speed on shear speed
Slika 12: Vpliv hitrosti rezanja na hitrost stri`enja
Figure 11: Influence of cutting speed on chip speed
Slika 11: Vpliv hitrosti rezanja na hitrost odrezka
Acoustics Corporation (PSC) were mounted on the top
of the tool holder. To maintain a good propagation of the
signals from the tool holder to the sensor, a semi-solid,
high-vacuum grease was used. During the experiment,
the AE signals were amplified, high passed at 20 kHz,
low passed at 1000 kHz, and then sent through a pre-
amplifier, at a gain of 40 dB, to the signal-processing
software package. All the cutting tests were performed
on a CNC lathe. The signals were real-time sampled,
amplified, digitized and then fed to the signal processing
unit. The AE signals were post-processed using AEwin.
3.1 Frequency analysis
When analyzing AE signals during hard turning, the
segmentation frequency must be taken into consideration
in relation to the frequency response of the applied AE
sensors. Two sensors with different frequency ranges
were applied. While the D9241A AE sensor allows us to
analyze low frequencies (from 20 kHz), the low
frequency limit of the WD AE sensor is 100 kHz.
The mean cycle frequency of the chip segmentation
can be determined by dividing the speed of the chip vch
by the mean spacing of the points of the maximum chip
thickness pc (the segment length in Figure 1) using
equation 7:13
segmentation frequency =
v
p
ch
c
(7)
To obtain the segmentation frequency, the segmen-
tation length had to be measured with the metallographic
microscope. Figure 4 illustrates that the segment length
is not constant; it varies and so the segmentation fre-
quency is not established as an exact value, but should be
determined within a certain interval. As shown on Fig-
ure 15, the calculated segmentation frequencies lie in the
frequency range of 10 kHz to 85 kHz. The verification of
these frequencies can be performed through an appli-
cation of the D9241A AE sensor, as the calculated
frequencies are in the frequency range (or close to the
low frequency limit at the low cutting speed) of this
sensor.
An AE analysis could be limited under specific
conditions because it could be difficult to find a relation
between a physical process and the character of the AE
signals. Because of this, many parameters of the AE sig-
nals implemented into the software package, such as the
amplitude, the number of hits and counts, the signal
strength, the energy, the average frequency and others
could be analyzed. These parameters are derived from an
AE wave with relation to the threshold value and they do
not have to reflect the dynamic character of the process
in real time. For example, Figure 15 illustrates that the
segmentation frequency increases with the cutting speed.
Figure 16 shows the relaxing and periodic character of
the signal at different cutting speeds.
The time period between two consecutive amplitudes
of an AE signal should reflect the segmentation frequen-
cy and should change in relation to different frequencies
under different cutting conditions. Figure 16 illustrates
that the time period between two consecutive pulses of
an AE signal does not change and stays nearly constant.
This is associated with the AE waveform transformation.
The transformation of an AE waveform is compounded
by the sensor response. When a resonant sensor is exci-
ted by a broadband transient pulse, it rings like a bell at
M. NESLU[AN et al.: INFLUENCE OF CUTTING SPEED ON INTENSITY OF THE PLASTIC DEFORMATION ...
750 Materiali in tehnologije / Materials and technology 47 (2013) 6, 745–755
Figure 15: Influence of cutting speed on the peak frequency (PF) of
the D9241A sensor and calculated segmentation frequency (CA)
Slika 15: Vpliv hitrosti rezanja na maksimalno frekvenco (PF) sen-
zorja D9241A in izra~unana frekvenca segmentacije (CA)
Figure 14: Detail of sensor placement
Slika 14: Detajl namestitve senzorjev
Figure 13: Schematic of experimental setup
Slika 13: Shematski prikaz eksperimentalnega sestava
its own frequency of oscillation. Therefore, the electrical
signal at the sensor output is the product of this ringing,
thus compounding the effects of multiple paths and
multiple wave modes, with which the wave travels from
the source to the sensor. The transformation of the AE
waveform leads to a mismatch between the real dyna-
mics of the signal and the derived AE parameters.
For example, the average frequency of the AE signal
is determined as the average frequency over the entire
AE hit and is derived from the AE-signal duration and
the number of the counts (a count is a signal excursion
over the AE threshold). This derived parameter (during
this experiment) is in the range of 21 kHz to 24 kHz (for
the D9241A sensor) and does not match the calculated
values of the segmentation frequency in Figure 16. On
the other hand, the applied AE system has the capability
of calculating and processing the frequency-derived AE
features in real time. Up to six frequency-based features
can be processed by the AE system. Each of these
features requires that a real time FFT is performed on the
received AE-hit waveform. Except for four partial power
frequency features, including the AE frequency features,
they include frequency centroid and peak frequency
(PF). The investigations into the centroid and peak fre-
quency show that only the peak frequency (illustrated in
Figure 15) matches the calculated values. The peak
frequency is defined as a point in the power spectrum, at
which the peak magnitude occurs. A real-time FFT (fast
Fourier transformation) is performed on the waveform
associated with the AE hit. The frequency that contains
the largest magnitude is reported. Like the calculated
values, the peak frequencies do not represent the exact
values, but the AE Win software also processes this
feature as an interval. Figure 16 shows a good correla-
tion between the peak frequencies and the calculated
frequencies as an evidence of the AE system’s capability
of truly reflecting the specific processes in the cutting
zone during hard turning.
3.2 Intensity of deformation processes
It should be established that the AE sensors’
responses depend on the frequency range of the applied
sensor and the cutting speed. Figures 16 and 17 illustrate
the AE signals for different cutting speeds and the
related FFT spectrums. It can be easily observed that the
course of an AE signal (for the low-frequency sensor
D9241A) reflects the relaxation character of the chip
formation. The relaxation course of the signal is related
to the relaxation process of the stress ahead of the cutting
edge and the crack propagation in the shear region.
According to the theory of the crack propagation and
segment formation, the change in the amplitude of AE
indicates a change in the sliding velocity during the
tool-chip interface. Many pulse-like signals are observed,
corresponding to the periodic fluctuation (the relaxation
character) of the cutting process. The signal level of the
AE between these pulses is quite small. During the
segmented chip formation, the chip slides over the rake
face with a varying speed corresponding to the period of
the fracture in the shear plane.
Figure 16 shows the periodic peaks in the FFT spec-
trum. This character of the FFT spectrum confirms the
dominant periodic character of the recorded signal and
an ability of the D9241A AE sensor (the frequency range
of 20 kHz to 180 kHz) to detect the periodic process
typical for a segmented-chip formation. The segmen-
tation frequencies of the investigated cutting speeds lie in
the frequency range of the D9241A AE sensor (or close
to the low-frequency limit). On the other hand, Figure
17 illustrates that these peaks are missing during the
formation of a continuous chip. Moreover, the excited
amplitudes in the FFT spectrum during hard turning are
M. NESLU[AN et al.: INFLUENCE OF CUTTING SPEED ON INTENSITY OF THE PLASTIC DEFORMATION ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 745–755 751
Figure 16: Signal of AE and its FFT spectrum for the 9241A sensor, hard turning: a) vc = 25 m min–1, b) vc = 100 m min–1, c) vc = 200 m min–1
Slika 16: Signal AE in njegov FFT-spekter za senzor 9241A, stru`enje: a) vc = 25 m min–1, b) vc = 100 m min–1, c) vc = 200 m min–1
significantly higher (above 50 dB, up to the 200 kHz
frequency) in comparison with those occurring during
the formation of continuous chips. It should be stated
that the amplitudes of the AE signals, together with the
amplitudes in the FFT spectrums, are nearly on the same
level for all the cutting speeds. This indicates that the
low-frequency AE sensor D9241A can be applied when
making dynamic analyses (frequency responses) of hard
turning, but this sensor is not sensitive to the variation of
the deformation-process intensity in the cutting zone.
Considering the AE WD sensor (the frequency range
of 100 kHz to 1000 kHz), all the segmentation fre-
quencies lie outside the frequency range of this sensor
and the periodic character of the AE signal is missing.
The FFT spectrum of the AE signal for the WD sensor is
without the periodic peaks in this spectrum (Figure 17)
as well as FFT spectrum in the case of turning annealing
steel (Figure 18).
It was reported in the previous sections that AE
signals can be classified into two types as either conti-
nuous-type AE signals or burst-type AE signals. Conti-
nuous signals are associated with the shearing in the
primary zone, in the tool-chip and tool-workpiece con-
tacts (Figure 19). These processes can be investigated
and detected with the WD AE sensor, because the signals
related to the chip segmentation (the crack propagation)
are out of the segmentation-frequency range. Figure 17
shows that the amplitude of the AE signal for the WD
sensor increases with the cutting speed. This aspect is
associated with the increasing intensity of the friction
M. NESLU[AN et al.: INFLUENCE OF CUTTING SPEED ON INTENSITY OF THE PLASTIC DEFORMATION ...
752 Materiali in tehnologije / Materials and technology 47 (2013) 6, 745–755
Figure 18: Signal of AE and its FFT spectrum, turning of annealed steel: a) vc = 25 m min–1, 9241A sensor, b) vc = 100 m min–1, 9241A sensor,
c) vc = 200 m min–1, WD sensor
Slika 18: Signal AE in njegov FFT-spekter, stru`enje `arjenega jekla: a) vc = 25 m min–1, 9241A senzor, b) vc = 100 m min–1, 9241A senzor, c) vc
= 200 m min–1, WD senzor
Figure 17: Signal of AE and its FFT spectrum for the WD sensor, hard turning
Slika 17: Signal AE in njegov FFT-spekter za WD-senzor, stru`enje
processes in the cutting zone in relation to the increasing
chip and shear speeds. Moreover, the amplitudes of the
related frequencies in the FFT spectrum are excited
above 50 dB, across the whole frequency range and at
the cutting speed of 200 m min–1, while the amplitudes in
the frequency spectrum at the cutting speed of 25
m min–1 are lower (Figure 17).
An application of the D9241A sensor is limited for
these analyses. The burst-type AE signal (related to the
crack propagation) superposes with the signal from the
tool-chip and tool-workpiece interface and so it causes
the difficulties in the investigation of the processes in
these regions. The segmentation frequencies lie in the
frequency range of the D9241A sensor. These fre-
quencies are associated with the crack formation and its
propagation in the shear zone. The signal associated with
a crack formation and its prolongation in the shear zone
is very strong (the amplitude of the AE signals for all the
cutting speeds is on the maximum level of 100 dB)
generating a massive low-frequency noise. This noise
negatively influences the detection of the friction process
in the shear zone, the tool-chip and tool-workpiece con-
tacts (Figure 19). The AE signals recorded by the
D9241A sensor are not sensitive to the friction processes
in the cutting zone. Figures 20, 21 and 22 show that the
RMS value, the signal strength and the absolute energy
of the AE signal do not vary significantly with the
increasing cutting speed. (The signal strength is defined
as the integral of the rectified voltage signal over the
duration of the AE waveform packet. The absolute ener-
gy is a true energy measure of the AE hit.)
On the other hand, the frequency range of the WD
sensor lies above the segmentation frequencies and so
the low-frequency limit of 100 kHz of the WD sensor
represents a high-pass filter. The RMS values, the
strength and the absolute energy of the AE signal
increase with the increasing cutting speed (Figures 20,
21 and 22). This aspect is associated with an increasing
intensity of the deformation processes in the cutting
zone. These processes are dissipated to the narrow white
areas in the shear zone (the white area is caused by the
plastic deformation in this zone, contrary to the crack
formation in the area without an occurrence of a
structure transformation) as well as to the tool-chip and
tool-workpiece contacts. The increase of the values in
Figures 20, 21 and 22 cannot be attributed to the tool
wear because all the tests were carried out in the normal
phase of the tool wear (its VB being from 0.1 mm to
M. NESLU[AN et al.: INFLUENCE OF CUTTING SPEED ON INTENSITY OF THE PLASTIC DEFORMATION ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 745–755 753
Figure 22: Influence of cutting speed on absolute energy of AE
Slika 22: Vpliv hitrosti rezanja na absolutno energijo AE
Figure 20: Influence of cutting speed on RMS values of AE
Slika 20: Vpliv hitrosti rezanja na RMS-vrednosti AE
Figure 19: Cutting zone for hard cutting: Ia) microcracked and
plastically deformed shear zone, Ib) cracked region, II) tool-chip
contact, III) tool-workpiece contact
Slika 19: Cona odrezovanja pri rezanju: Ia) plasti~no deformirano
stri`no podro~je z mikrorazpokami, Ib) razpokano podro~je, II) stik
orodje-odrezek, III) stik orodje-obdelovanec
Figure 21: Influence of cutting speed on AE signal strength
Slika 21: Vpliv hitrosti rezanja na mo~ AE-signala
0.12 mm). The increasing intensity of deformation pro-
cesses is attributed to the following aspects:
the increasing chip segmentation with the increasing
cutting speed; a more intensive thinning of the micro-
cracked region (the plastically deformed shear zone) as
the chip moves up the face of the tool;
• the increasing non-homogeneity in the stress and
temperature distribution in the cutting zone with the
increasing cutting speed, together with the increasing
chip and shear speeds;
• the temperature in the cutting zone increases with the
increased cutting speed8 (forming suitable conditions
for the processes of plastic deformation and the
related structure transformation and reducing the
zones of an undeformed chip).
The sensitivity of the AE signals concerning the
intensity of the real-deformation processes in the cutting
zone can be illustrated through a comparison of the AE
signal and FFT spectrums during conventional (annealed
steel) turning and hard turning. The analytic approach
shows that the chip deformation and chip ratio are higher
for the conventional continuous chips (Figures 7 and 8),
but the amplitude of the AE signals and the excited
amplitudes of the FFT spectrums in Figure 18 are
significantly lower. While the amplitudes of the AE
signal during hard turning oscillate between ±10000 mV,
the formation of a continuous chip leads to an oscillation
on the scale of ± 100 mV. Moreover, the excited ampli-
tudes of the FFT spectrum in Figure 18 are close to zero
(except for the resonance frequency represented by the
maximum amplitude in the spectrum), while, in contrast,
those amplitudes excited above 50 dB cover the whole
frequency spectrum during hard turning (Figure 17).
Some other derived parameters such as the RMS value of
the AE signal, the signal strength and the absolute
energy are closely associated with the emission signals,
correlating with the amplitude range of the responded
emission waves. Figures 20, 21 and 22 illustrate that all
the values are significantly lower during the formation of
conventional chips.
The metallographic and experimental studies give
contradictory results because of their different methodo-
logies. For example, the values of the chip ratio or the
chip deformation for the hard-turning process in Figures
7 and 8 decrease with the increasing cutting speed. This
indicates a decreasing intensity of deformation processes
in the cutting zone. However, Figures 17, 18, 21 and 22
show that the absolute energy, the signal strength and the
amplitudes of the AE signals increase with the increasing
cutting speeds.
The conventional analytical approach does not facili-
tate an evaluation of the real intensity of deformation
processes in the specific zones of the cutting process.
This approach takes the whole chip into the conside-
ration. The calculated values do not reflect the non-
homogeneity of deformation processes during hard
cutting, but provide information about the average value
across the whole chip. On the other hand, the experi-
mental study (based on AE) can identify the real
intensity of this specific process.
4 CONCLUSIONS
The findings of this study show that the AE signals
can be used to monitor the dynamic character and inten-
sity of plastic deformation in the cutting zone during
hard turning in the following ways:
• metallographic analyses and the related calculations
allow an analysis of the real intensity of deformation
processes during the formation of conventional conti-
nuous chips, but this approach is less sensitive to the
identification of the specific-character chip formation
during hard turning;
• experimental analysis of the chip formation during
hard turning should be investigated with respect to
the specific dynamic character of the cutting process
and the related frequencies;
• burst type of the AE signal is associated with a crack
propagation in the shear zone because of the relaxing
character of the chip formation and the related seg-
mentation frequencies;
• analysis of the chip segmentation requires a match of
the segmentation frequency and the frequency range
of the applied AE sensor;
• continuous type of the AE signals is associated with
the plastic deformation in the cutting zone related to
the plastic deformation in the shear zone, the tool-
chip and tool-workpiece sliding contacts;
• continuous type of the AE signal is more sensitive to
the real intensity of deformation processes during
hard turning, contrary to the metallographic analyses
and the calculated parameters.
The dynamic character of a cutting process in hard
turning and the specific character of a chip formation
significantly affect the parameters such as the shear and
chip speeds, the friction processes in the cutting zone,
the related heat generation and high temperatures in this
zone. These parameters have an impact on the surface
quality represented by the residual stresses, surface
hardness, structural changes and other attributes. And so,
the studies about the dynamic character of the hard-turn-
ing process should be carried out.
Acknowledgement
This paper was supported by the Students Grant
Competition of VSB-TU Ostrava, "SP2012/68 Effective
machining of progressive materials and integrity surface
evaluation", and also by the VEGA agency.
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A. MICHALCOVÁ et al.: EFFECT OF IRON AND CERIUM ADDITIONS ON RAPIDLY SOLIDIFIED Al-TM-Ce ALLOYS
EFFECT OF IRON AND CERIUM ADDITIONS ON
RAPIDLY SOLIDIFIED Al-TM-Ce ALLOYS
VPLIV DODATKOV @ELEZA IN CERIJA NA HITRO STRJENE
ZLITINE Al-TM-Ce
Alena Michalcová1, Dalibor Vojtìch1, Gerhard Schumacher2, Pavel Novák1, Eva Pli`ingrová3
1Institute of Chemical Technology, Department of Metals and Corrosion Engineering, Technicka 5, 166 28 Prague, Czech Republic
2Helmholtz-Zentrum Berlin für Materialien und Energie GmbH, Hahn-Meitner-Platz 1, 14109 Berlin, Germany
3Institute of Inorganic Chemistry AS CR, v.v.i., Husinec-Rez 1001, 258 68 Rez, Czech Republic
alena.michalcova@vscht.cz
Prejem rokopisa – received: 2013-02-23; sprejem za objavo – accepted for publication: 2013-03-25
The influence of the additions of the mass fractions 3 % of Fe and 1 % of Ce to an AlCr4Ti0.4 alloy is described in this paper.
The iron addition modifies the phase composition of the material leading to the formation of the Al13Fe4 crystalline and
Al80(Cr,Fe)20 quasicrystaline phases. It also causes a grain refinement that was proved with transmission electron microscopy
and X-ray diffraction. The microhardness of the Fe-alloyed material is changed by about a factor of two in comparison to the
material without Fe. On the other hand, the Fe addition has no influence on the thermal stability of the alloy. The effect of 1 %
of Ce on the AlCr4Ti0.4 alloy hardly affects the microhardness but increases the thermal stability during long-term annealing at
400 °C.
Keywords: rapid solidification, aluminium, quasicrystals
V prispevku je opisan vpliv dodatka z masnim dele`em 3 % `eleza in 1 % cerija zlitini AlCr4Ti0,4. Dodatek `eleza spremeni
sestavo faz v materialu in povzro~i nastanek Al13Fe4 kristalne in Al80(Cr,Fe)20 kvazikristalne faze. Povzro~i tudi zmanj{anje zrn,
kar je bilo dokazano s presevno elektronsko mikroskopijo in z rentgensko difrakcijo. Mikrotrdota z Fe legiranega materiala se
spremeni za faktor dve glede na material brez Fe. Po drugi strani pa dodatek Fe ne vpliva na termi~no stabilnost zlitine. U~inek
1 % Ce v AlCr4Ti0,4 neznatno vpliva na mikrotrdoto, vendar pa pove~a termi~no stabilnost med dolgotrajnim `arjenjem pri 400
°C.
Klju~ne besede: hitro strjevanje, aluminij, kvazikristali
1 INTRODUCTION
Rapidly solidified (RS) aluminium alloys are promis-
ing materials for structural applications. Their main
advantage is the superior strain-to-weight ratio. The best
mechanical properties and thermal stability are achieved
in an aluminium rare-earth transition-metal (RE-TM) alloy
system with aluminium contents higher than 80 %.1,2
These alloys are usually prepared under the maximum
available cooling rates that, together with the appropriate
chemical compositions, lead to a formation of amor-
phous or partially amorphous alloys.1 Such metastable
materials are very interesting from the investigation point
of view, but their potential for the applications at
elevated temperatures is limited. The amorphous phase
decomposes under the heat treatment, forming inter-
metallic phases in the fcc-Al matrix.3 Depending on the
heat-treatment regime, annealed materials can exhibit
higher hardness and other mechanical properties such as
the yield strength and the ultimate tensile strength.1 The
production of such materials is highly demanding with
respect to processing and it would be considerably com-
plicated to manufacture them in industrial conditions.
This paper is focused on an investigation of the rapidly
solidified alloys, prepared with the cooling rates achiev-
able in industry.
The influences of single alloying elements on
extremely rapidly solidified aluminium alloys are
described in detail.1,4 The authors usually focus on the
effect of each element on the glass-forming ability
(GFA), which qualifies the tendency of an alloy to form a
glassy structure. Three empirical rules for a GFA deter-
mination were established:
1. The higher the number of elements in a system, the
higher is GFA.
2. GFA is also high due to a negative mixing enthalpy
of the components in a system.
3. A GFA increase is generally observed together with
great differences in the atomic ratios of the elements.1
The definition of the GFA described for glassy alloys
can be generalized as the ability to form a metastable
phase (MPFA), indicating the tendency of a system to
form not only the glassy state, but, generally, to form any
metastable phase, e.g., amorphous and quasicrystaline
phases and supersaturated solid-solution phases.
In the case of amorphous alloys, the functions of all
the REs are approximately the same with respect to the
increasing GFA.5 The influences of single TMs on Al-
based alloys cannot be simplified only in terms of GFA,
because they vary significantly5 on the basis of the inter-
atomic interactions. The Al-based alloys containing Fe
are very exceptional. It was reported that there is a strong
Materiali in tehnologije / Materials and technology 47 (2013) 6, 757–761 757
UDK 669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)757(2013)
interaction between the Al and Fe atoms in the amor-
phous Al-RE-Fe alloys6 and even in the alloy melts.7 It
was found, in the amorphous Al-Ce-Fe alloys, that the
Ce atoms are basically randomly distributed and that the
Fe atoms are surrounded by many Al atoms.6 This is in
agreement with the idea that the icosahedral Al-Fe
clusters are present already in the melts.2,7 It is generally
supposed that the melt clusters are formed by the
elements with the lowest negative mixing enthalpies.
Interestingly, the behavior of the Al-Ce-Fe alloys is in
disagreement with this rule because the mixing enthalpy
of Al-Fe is –11 kJ/mol 8 and the mixing enthalpy of
Al-Ce is -38 kJ/mol.9 It was also shown that an Fe
addition changes the crystallization mechanism of an
Al86Mm4Ni10 alloy (Mm means mischmetal – an alloy
of Ce, La, Nd and Pr), increasing its thermal stability.6
Ce precipitates from the Al-TM based amorphous
alloys in the form of an Al4Ce phase.7,10 This phase is
formed due to a decomposition of the Al92Ce8 metastable
phase11 and is one of the main strengthening phases. It
was also reported12 that REs can be present in four
different phases in the Al-RE alloys or Al-TM-RE alloys
rich in REs. This possibility of nucleating various phases
can represent a "confusion principle" for an alloy. On the
other hand, when Fe is added to an alloy, a formation of
intermetallic phases is preferred.12 The question is in
which form Ce will be present in the crystalline RS
alloys and whether it will significantly influence the
alloys.
In this paper, the influence of a mass fraction 3 % Fe
addition on the structure, hardness and thermal stability
of the nanocrystalline AlCr4Ti0.4Ce1 alloy will be
shown. This influence will also be compared to the influ-
ence of a 1 % Ce addition to the AlCr4 Fe3Ti0.4 alloy.
2 EXPERIMENT
The alloys with the compositions given in Table 1
were studied in this paper. Ingots of the alloys were
prepared by melting the appropriate amounts of master
alloys and pure metals (AlCr11, AlTi4, Al, Fe and Ce)
in a vacuum induction furnace. Consequently, RS rib-
bons were prepared by melt spinning. The cooling wheel
diameter of the melt-spinning equipment was 14 cm. The
cooling wheel velocity was 1800 r/min, corresponding to
the circumferential speeds of 28 m/s. Melt-spinning pro-
cedures were performed in air. The rapidly solidified
alloys were produced in the form of thin foils with a
thickness of approximately 50 μm and a width of 4–5
mm. The chemical composition was determined with
X-ray fluorescence spectroscopy (XRF) using a spectro-
meter ARL 9400 XP for Ce, and with atomic absorption
spectroscopy (AAS) using GBC 932plus for the other
elements.
The microstructure of the cross-sectioned ribbons
was examined using a scanning electron microscope
(SEM) Hitachi S 4700 (30 kV) and a light microscope
Neophot 2 in the immersion mode. The samples were
etched in a solution of 0.5 % HF for 20 s. The phase
composition of the materials was determined with X-ray
diffraction (XRD) (PAN analytical X’Pert PRO + High
Score Plus, Cu anode). All diffraction measurements
were performed on the wheel side of the RS ribbons. The
grain size was estimated from XRD diffraction patterns
as the size of the coherently diffracting domain using the
Scherrer formula (1),13 attributing the peak width
expansion only to the influence of the average grain size:
Crystallite size (average) =
K
B


cos
(1)
In equation (1)  describes the wavelength of the
incident X-rays, K is the constant that varies with the
method of taking the line width (0.89 < K < 1) and B de-
scribes the structural broadening, given by the difference
in the integral profile width between the standard and the
unknown samples:13
B = Bobs. – Bstd. (2)
The peak at the 2 position of about 78° corres-
ponding to the diffraction on the Al (311) plane was used
for the estimation. The Rietveld structure refinement was
performed with the program Topas 4. The samples were
examined using a transmission electron microscope
(TEM) Jeol 3010 (an accelerating voltage of 300 kV,
LaB6). Thin foils for TEM were prepared by grinding the
free sides of the RS ribbons until the final thickness
reached approximately 50 μm, and by subsequently
electropolishing the mixture of C2H5OH:HNO3 with the
ratio of 3:1 at 10 V and –20 °C. The structural investiga-
tion was completed with the EDS (Oxford Instruments)
measurements. The Vickers hardness HV 0.02 was also
measured on the cross-sectioned ribbons. The load of 20
g was chosen to be the optimum one for measuring the
average hardness of the ribbons. The indentor always
pointed to the middle of the ribbons.
A long-term annealing of the alloys was carried out
in an electric resistance furnace. Selected samples were
annealed at 400 °C to determine the thermal stability of
the materials. After the long-term annealing (25, 50, 75
and 100) h, the room-temperature hardness HV 0.02 was
measured.
3 RESULTS AND DISCUSSION
All the RS ribbons showed structural gradients in the
cross-sections, as illustrated in Figures 1 and 2. The
cooling wheel sides of the ribbons, shown at the bottom
A. MICHALCOVÁ et al.: EFFECT OF IRON AND CERIUM ADDITIONS ON RAPIDLY SOLIDIFIED Al-TM-Ce ALLOYS
758 Materiali in tehnologije / Materials and technology 47 (2013) 6, 757–761
Table 1: Chemical compositions of the studied alloys in mass frac-
tions (w/%) determined with an XRF analysis
Tabela 1: Kemijska sestava preiskovanih zlitin v masnih dele`ih
(w/%), dolo~enih z XRF-analizo
Alloy Al Cr Ti Fe Ce
AlCr2 97.24 2.13 0.38 0.25 0
AlCr4Ce1 93.96 3.84 0.40 0.96 0.84
AlCr4Fe3 91.70 4.16 0.63 3.51 0
AlCr5Fe3Ce1 90.33 4.93 0.29 3.43 1.02
of Figures 1 and 2 and marked with arrows, exhibit a
very fine microstructure, called an ultra-rapidly-soli-
dified area (URSA). The thickness of URSA depends on
two main factors: 1) experimental parameters and 2)
alloy properties. Experimental parameters include the
cooling wheel speed, the atmosphere of the experiment
and the geometry of the experimental device. These cha-
racteristics were identical during the preparation of all
the rapidly solidified alloys. Therefore, all the differen-
ces among the rapidly solidified ribbons can be assumed
to result from the chemical compositions of the alloys.
Figure 1 shows the structure of a cross-sectioned
ribbon of the AlCr4Ce1 alloy. In comparison with the
AlCr5Fe3Ce1 alloy, documented in Figure 2, its URSA
is thinner and the microstructure is not so fine, because
intermetallic particles (the dark areas) are observed on
the free (top) side of the ribbon. In the URSA, a high-
entropy alloy is formed. The ability to form such an alloy
is increased with the number of elements in the system,
especially with the number of transition metals. For this
reason, the URSA of the AlCr5Fe3Ce1 (quaternary)
alloy is much larger than that of the AlCr4Ce1 (ternary)
alloy.
According to XRD (Figure 3), the AlCr5Fe3Ce1
alloy is composed of the fcc-Al, Al13Cr2, Al13Fe4, Al3Ti
and Al3Ce crystalline intermetallic phases and the
Al80(Cr,Fe)20 decagonal quasicrystalline phase. The
AlCr4Ce1 alloy, on the other hand, contains neither
Al13Fe4 nor the quasicrystalline phase. This suggests that
the Fe addition leads to a formation of two very different
intermetallic phases – Al13Fe4 and quasicrystalline
Al80(Cr,Fe)20 – therefore, strongly increasing the MPFA
due to the so called "confusion principle".14,15
The fcc-Al matrix grain size determined with the
Scherrer formula and Rietveld analysis is given in Table
2. The values are different because of the used methods
and their limitations, but the grain-size trend is in
sufficient agreement. The total amount of the alloying
elements (Cr, Ti, Fe, Ce) in the AlCr2 alloy is signifi-
cantly lower than in the other alloys. This results in a
lower number of crystallization nuclei of the interme-
tallic phases during the solidification and in a huge grain
size of the AlCr2 alloy.
Table 2: Grain size of the fcc-Al matrix in nm determined with the
Scherrer formula and Rietveld analysis
Tabela 2: Velikost zrn fcc-Al osnove v nm, dolo~ena po Scherrerjevi
ena~bi in z Rietveldovo analizo
Alloy Scherrer formula Rietveld analysis
AlCr2 232 118
AlCr4Ce1 123 76
AlCr4Fe3 49 50
AlCr5Fe3Ce1 46 32
The addition of the mass fraction 1 % Ce causes a
mild grain refinement. The reason for this is not only a
higher amount of the alloying elements. The possibility
A. MICHALCOVÁ et al.: EFFECT OF IRON AND CERIUM ADDITIONS ON RAPIDLY SOLIDIFIED Al-TM-Ce ALLOYS
Materiali in tehnologije / Materials and technology 47 (2013) 6, 757–761 759
Figure 2: LM micrograph of a cross-section of a rapidly solidified
ribbon of the AlCr5Fe3Ce1 alloy
Slika 2: Posnetek prereza hitro strjenega traku zlitine AlCr5Fe3Ce1
Figure 3: XRD patterns of: a) AlCr2, b) AlCr4Ce1, c) AlCr4Fe3 and
d) AlCr5Fe3Ce1 alloys
Slika 3: Rentgenski posnetki zlitin: a) AlCr2, b) AlCr4Ce1, c)
AlCr4Fe3 in d) AlCr5Fe3Ce1
Figure 1: LM micrograph of a cross-section of a rapidly solidified
ribbon of the AlCr4Ce1 alloy
Slika 1: Posnetek prereza hitro strjenega traku zlitine AlCr4Ce1
of an intermetallic-phase formation increases with the Ce
addition and so this addition supports the alloy metasta-
bility on the basis of the confusion principle. The Ce
addition also enables the existence of an alloy melt at
lower temperatures, because the Al-Ce system exhibits a
eutectic. The Fe addition has a strong effect on the
grain-size refinement, as it can be seen from the diffe-
rence between the grain sizes of the AlCr4Ce1 and
AlCr5Fe3Ce1 alloys. A very small grain size of the RS
materials is caused by a suppression of the diffusion rate
due to a very high quenching rate. Surprisingly, the
AlCr4Fe3 and AlCr5Fe3Ce1 alloys exhibited almost the
same grain sizes. This was the reason for determining the
grain sizes of these alloys also with an image analysis
and the results of the grain-size distribution are given in
Figures 4 and 5. The values of the grain size determined
from the XRD pattern are about half of the values deter-
mined from the TEM image. This is usual because what
is seen, in a TEM micrograph, as single grains, can be
more coherent diffracting domains for the X-ray me-
thods. However, all the methods showed that the grain
sizes of the AlCr4Fe3 and AlCr5Fe3Ce1 alloys are
approximately the same. In the Fe-containing alloys the
grain refinement is predominantly caused by the Fe
addition and the influence of Ce on the grain size is
almost negligible. The interaction of Al and Fe in a melt
and during the solidification is so strong that it probably
leads to a primary formation of very fine intermetallic
phases. These phases can act as solidification nuclei for
the fcc-Al matrix. These phases are then located at the
grain boundaries or at the triple points of the matrix
grains, protecting the alloys from a consequent grain
coarsening.16
Figure 6 presents the structure of the AlCr4Ce1 alloy
composed of the Al grains and intermetallic particles
containing mainly Al and Cr. These intermetallics are
most likely Al13Cr12, see the XRD patterns in Figure 3.
The microstructure of the AlCr5Fe3Ce1 alloy is shown
A. MICHALCOVÁ et al.: EFFECT OF IRON AND CERIUM ADDITIONS ON RAPIDLY SOLIDIFIED Al-TM-Ce ALLOYS
760 Materiali in tehnologije / Materials and technology 47 (2013) 6, 757–761
Figure 7: TEM micrograph of a cross-section of a rapidly solidified
ribbon of the AlCr5Fe3Ce1 alloy (1 – Al grains, 2 – Al13Cr2, 4 –
Al80(Cr,Fe)20, 5 – Al13Fe4, 6 – Al3Ce)
Slika 7: TEM-posnetek prereza hitro strjenega traku zlitine
AlCr5Fe3Ce1 (1 – Al zrna, 2 – Al13Cr2, 4 – Al80(Cr,Fe)20, 5 –
Al13Fe4, 6 – Al3Ce)
Figure 5: Grain-size distribution of the RS AlCr5Fe3Ce1 alloy
Slika 5: Razporeditev velikosti zrn hitro strjene zlitine AlCr5Fe3Ce1
Figure 6: TEM micrograph of a cross-section of a rapidly solidified
ribbon of the AlCr4Ce1 alloy (1 – Al grains, 2 – Al13Cr2)
Slika 6: TEM-posnetek prereza hitro strjenega traku zlitine AlCr4Ce1
(1 – Al zrna, 2 – Al13Cr2)
Figure 4: Grain-size distribution of the RS AlCr4Fe3 alloy
Slika 4: Razporeditev velikosti zrn hitro strjene zlitine AlCr4Fe3
in Figure 7. Beside the Al grains and Al13Cr2 interme-
tallics in Figure 7, there are also spherical particles of
the Al80(Cr,Fe)20 quasicrystalline phase as well as the
Al13Fe4 and Al3Ce phase particles.
Figure 8 shows the dependence of the room-tem-
perature hardness on the annealing time after annealing
at 400 °C. The microstructure of the AlCr5Fe3Ce1 alloy
is supposed to be very stable because its hardness value
did not change even after 100 h of annealing. The
hardness of the other alloys slightly decreased. The Fe
addition also caused a significant increase in the initial
hardness of the rapidly solidified ribbons. In comparison
with the AlCr5Fe3Ce1 alloy the hardness of the
AlCr4Fe3 alloy slightly decreased with time. This
suggests that the Ce addition stabilizes the alloy. The
Al3Ce particles, present in the AlCr5Fe3Ce1 alloy, are
placed mainly at the grain boundaries suppressing the
grain coarsening during the thermal exposure.
4 CONCLUSION
An addition of a few mass fractions of Fe signifi-
cantly modifies the microstructure and the hardness of
the rapidly solidified Al-Cr-Ti alloys. This results in an
increase in the thickness of the ultra-rapidly-solidified
areas of ribbons. It also causes a decrease in the fcc-Al
matrix grain size and a significant increase in the hard-
ness of the RS ribbons. An influence of the Fe addition
on the thermal stability of the alloy during the heat treat-
ment was not observed.
An addition of a very small amount (1 %) of Ce leads
to the thermal stabilization of the alloy at the elevated
temperatures of about 400 °C. The influence of Ce on
the matrix grain size was not observed in the Al-Cr-Fe-Ti
alloy. In the Al-Cr-Ti alloy, the Ce addition causes a mild
decrease in the grain size.
Acknowledgement
This work was financially supported by the Czech
Science Foundation (project No. P108/12/G043).
5 REFERENCES
1 A. Inoue et al., NanoStructured Materials, 11 (1999), 221
2 K. Saksl et al., Journal of applied physics, 97 (2005), 113507
3 M. Galano et al., Acta Materialia, 57 (2009), 5107
4 P. Jur~i et al., Mater. Tehnol., 41 (2007) 6, 283–287
5 H. Yang et al., Journal of Non-Crystalline Solids, 354 (2008), 3473
6 H. W. Jin et al., Journal of Materials Science, 36 (2001), 2089
7 C. Zhang et al., Materials Science and Engineering A, 323 (2002),
226
8 C. Dong et al., Journal of Physics: Conference Series, 98 (2008),
012015
9 K. Song et al., Journal of Alloys and Compounds, 440 (2007), L8
10 O. D. Neikov et al., Materials Science and Engineering A, 447
(2008), 80
11 S. H. Wang et al., Journal of Alloys Compounds, 441 (2007), 135
12 P. Rizzi et al., Materials Science and Engineering A, 304–306
(2001), 574
13 High Score Plus 2, user manual
14 A. L. Greer, Science, 267 (1995), 1947
15 M. P. Miller et al., Bulk Metallic Glasses, Springer-Verlag, 2007, 51
16 C. J. Ehrstrom et al., Materials Science and Engineering A, 186
(1994), 55–64
A. MICHALCOVÁ et al.: EFFECT OF IRON AND CERIUM ADDITIONS ON RAPIDLY SOLIDIFIED Al-TM-Ce ALLOYS
Materiali in tehnologije / Materials and technology 47 (2013) 6, 757–761 761
Figure 8: Room-temperature hardness dependence on the annealing
time at 400 °C for: a) AlCr5Fe3Ce1, b) AlCr4Fe3, c) AlCr4Ce1 and
d) AlCr2 alloys
Slika 8: Odvisnost trdote pri sobni temperaturi od ~asa `arjenja pri
400 oC za zlitine: a) AlCr5Fe3Ce1, b) AlCr4Fe3, c) AlCr4Ce1 in d)
AlCr2

N. BOUHAMOU et al.: SHRINKAGE BEHAVIOR OF A SELF-COMPACTING CONCRETE
SHRINKAGE BEHAVIOR OF A SELF-COMPACTING
CONCRETE
VEDENJE SAMOZGO[^EVALNEGA BETONA PRI KR^ENJU
Nasr-Eddine Bouhamou, Nadia Belas, Karim Bendani, Abdelkader Mebrouki
LCTPE Laboratory, Civil Engineering Department, University of Mostaganem, Algeria
bendanik@yahoo.fr
Prejem rokopisa – received: 2013-03-01; sprejem za objavo – accepted for publication: 2013-03-26
This paper presents the influence of the mixing constituents on the behavior of self-compacting concretes (SCCs), especially,
the effect of the paste volume in their fresh and hardened state. It explains the roles of the pore network and the microstructure
of the hydrates in drying and autogenous shrinkage of SCCs. Several tests such as slump flow, L-box, sieve stability, bleeding,
mechanical strength, free shrinkage (drying and autogenous shrinkage) and microstructural tests (mineralogical
characterization, porosimetric distributions) were carried out in order to understand the roles played by various components
likely to influence the formulation of an SCC. The results obtained offer interesting prospects to optimize a SCC using local
materials in Algeria. This study has allowed the development of various SCC formulations that fulfill the rheological criteria
such as good deformability, low bleeding, low segregation and a better mechanical performance.
Keywords: self-compacting concrete, limestone fillers, porosimeter, drying shrinkage, autogenous shrinkage, mechanical
strength
^lanek opisuje vpliv prime{anih dodatkov na vedenje samozgo{~evalnega betona (SCC) na volumen v sve`em in v strjenem
stanju. Razlo`ena je vloga mre`e por in mikrostrukture hidratov pri su{enju in avtogenem kr~enju SCC. Da bi razumeli vlogo
posameznih delov in sestavin, ki opredeljujejo SCC, je bilo izvr{enih ve~ preizkusov, kot pojemanje toka, L-Box, stabilnost pri
sejanju, izcejanje, mehanska trdnost, prosto kr~enje (su{enje in avtogeno kr~enje), mikrostrukturni preizkusi (mineralo{ka
karakterizacija, razporeditev por). Dobljeni rezultati dajejo zanimive mo`nosti za optimiranje SCC z uporabo lokalnih mine-
ralov v Al`iriji. Ta {tudija je omogo~ila razli~ne formulacije SCC, ki izpolnjujejo reolo{ka merila, kot je dobra deformabilnost,
manj izcejanja, manj segregacij in bolj{e mehanske zmogljivosti.
Klju~ne besede: samozgo{~evalni beton, polnila iz apnenca, porozimeter, kr~enje pri su{enju, avtogeno kr~enje, mehanska
trdnost
1 INTRODUCTION
Self-compacting concrete is a fluid mixture suitable
for being placed in the structures with a congested
reinforcement without any vibration. Such concrete
should have a low viscosity during the pouring to ensure
the high-flow ability, and moderate viscosity to resist
segregation and bleeding. Moreover, it must maintain its
homogeneity during the transportation, placing and
curing to ensure an adequate structural performance and
long-term durability.
However, despite the interesting features that they
offer, in particular, in their fresh state, some drawbacks
of their long-term behavior can hinder their use. Indeed,
the physicochemical phenomena inherent to the free
shrinkage of SCCs have not yet been clearly specified. If
it seems, from the formulation point of view, that one
controls the proportioning of various components, to
obtain a suitable fluidity, by preserving a good homoge-
neity and even a high mechanical strength,1,2 it should
also be noted that the problem of the strains remains to
be treated.
This research was undertaken in this context and
must give answers to several questions, some of them
concerning the progression of mechanical and physical
properties of concrete with time. The objective of the
current work, using the local materials, is to acquire
some knowledge explaining the shrinkage behavior of
SCCs with variations in the relevant constituents, in
particular, the volume of paste. The results obtained in
this study will help us to have a more precise idea about
the values of shrinkage and to determine the difference
between the SCC and the ordinary concrete (OC),
adopting an optimum formulation.
2 MATERIALS AND METHODS
2.1 Materials used
2.1.1 Cement
The cement used is CPJ CEM II/A 32.5, obtained
from the Zahana factory (in the west of Algeria). Its
physical and chemical properties in weight percent are
given in Table 1.
2.1.2 Limestone fillers
The limestone used is from the Kristel quarry (in the
west of Algeria). The sample analyzed is essentially
constituted of limestone (w(CaCO3) = 85.45 %)
containing also a considerable quantity of silica (w(SiO2)
= 10.81 %). Its physical and chemical properties in mass
fractions are given in Table 1.
Materiali in tehnologije / Materials and technology 47 (2013) 6, 763–769 763
UDK 691:620.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)763(2013)
Table 1: Physical properties and chemical analysis of the cement and
limestone fillers
Tabela 1: Fizikalne lastnosti in kemijska analiza cementa ter polnil iz
apnenca
Properties Cement Limestonefillers
Physical properties
Bulk density (g/cm3) 1.09 0.87
Specific gravity (g/cm3) 3.00 2.66
Fineness (Blaine) (cm2/g) 3100 2880
Chemical analysis
(mass fractions, w/%)
SiO2 21.93 10.81
CaO 63.87 47.51
MgO 0.21 0.21
Fe2O3 4.26 0.76
Al2O3 6.81 0.31
SO3 1.31 –
Loss on ignition 1.83 40.69
Free CaO 0.13 –
Carbonates – 85.45
CO2 – 37.60
H2O – 3.09
2.1.3 Aggregates
Table 2 gives the aggregate properties for all the
mixes used in this study.
Table 2: Physical properties of the aggregates
Tabela 2: Fizikalne lastnosti agregatov
Marine
sand (Sm)
Quarry
sand (Sc) Gravel (G)
Grade 0/2 0/3 3/8 8/15
Composition siliceous limestone limestone limestone
Specific gravity
(g/cm3) 2.64 2.66 2.67 2.67
Absorption (%) – – 1.28 0.93
2.1.4 Admixture
A Viscocrete 20 HE superplasticizer, non-colored,
containing acrylic copolymer, developed by the company
of Sika, France, and complying with NF EN 934-2, was
used for all the SCC mixes. This high-range water-
reducing admixture has a dry extract content of 40 % and
a unit mass of 1.085.
For the ordinary concrete (OC), the water-reducing
superplasticizer used was a Plastiment BV 40. This
admixture has a dry extract content of 36.6–40.4 % and a
unit mass of 1.185.
2.2 Concrete mixtures
Three self-compacting concretes and one ordinary
vibrated concrete were designed to study the effect of the
paste volume on the SCC behavior in its fresh and
hardened states.
The superplasticizer proportioning and the water/
cement and filler/cement ratios were maintained constant
for all the SCC mixes, i.e., Sp = 1 %, W/C = 0.5 and F/C
= 0.25.
For the ordinary vibrated concrete, the mix was
obtained using the Dreux-Gorisse method and the same
W/C ratio as for the SCC was used for comparison. The
compositions of different mixes are given in Table 3.
2.3 Tests on the fresh concrete
The fresh-state characterization of the SCC was limi-
ted to the tests recommended by the French Association
of Civil Engineering (AFGC),3 i.e, the mini slump flow,
L-box, sieve stability and bleeding.
2.4 Tests on the hardened concrete
2.4.1 Mechanical strengths
The samples used to determine the compressive
mechanical strength for different concretes studied, are
cylindrical test tubes with an diameter 11 cm and height
22 cm. Once removed from the mold, they were pre-
served in water for (1, 7, 28 and 90) d.
2.4.2 Free shrinkage
The shrinkage strains were measured using a contrac-
tometer on the prismatic test tube with the dimensions of
7 cm × 7 cm × 28 cm, fulfilling two requirements:
• to obtain the total shrinkage with hydrous exchange
of the material with the environment;
• to measure the autogenous shrinkage, without
hydrous exchange with the environment, by wrapping
the test tubes in one or two aluminum sticker sheets.
After the removal from the mold (after 24 h), the
total- and autogenous-shrinkage measurements are, at
the beginning, worked out at very short times, followed
by a periodic increase.
N. BOUHAMOU et al.: SHRINKAGE BEHAVIOR OF A SELF-COMPACTING CONCRETE
764 Materiali in tehnologije / Materials and technology 47 (2013) 6, 763–769
Table 3: Mix proportioning
Tabela 3: Me{alna razmerja
Description
Vpaste
(%)
Mix proportioning (kg/m3)
Cement Limestonefillers
Efficient
water
Viscocrete
20 HE
Plastiment
BV 40 Sm Sc G (3/8) G (8/15)
SCC1 35 375.5 94 188 3.76 – 601 259 346 519
SCC2 37.5 400 100 200 4 – 578 249 333 499
SCC3 40 429 107 215 4.29 – 554 239 319 479
OC 34 350 – 175 – 5.25 129 576 204 947
2.4.3 Measure of porosity
With the mercury intrusion porosimetry (MIP), the
samples are introduced into a chamber, the chamber is
evacuated, the samples are surrounded by mercury and
the pressure on mercury is gradually increased. As the
pressure increases, mercury is forced into the pores on
the surface of the sample. By tracking the pressure and
intrusion volumes during the experiment, it is possible to
measure the connecting pore necks of a continuous
system or the breakthrough pressure in a discontinuous
system. The pore width corresponding to the highest rate
of the mercury intrusion per change in the pressure is
known as the "threshold", "critical" or "percolation" pore
width. Using this technique, one also measures the total
porosity of a sample as that corresponding to the volume
of the mercury intruding at the maximum experimental
pressure, divided by the bulk volume of the unintruded
sample.
The humidity exchanges of the concrete with the
external media are then linked to the porous structure.
The knowledge of the concrete porosity can thus prove
to be useful when estimating the relative differences
between the microstructural formulations of the SCCs
that are likely to explain the results obtained for the
shrinkage with respect to its composition. In this context,
the porosities of different SCC formulations were cha-
racterized (after about 300 d, i.e., the last period of the
shrinkage measurement). The measurements were
realized in the civil engineering laboratory of La
Rochelle (France) by means of a mercury porosimeter
Autopore III from Micrometrics (ASTM D4404-10). The
range of the measurement was from 3 nm to 200 μm.
3 RESULTS AND ANALYSIS
3.1 Fresh state
The response of the characterization tests performed
on the prepared concrete is given in Table 4.4
Table 4: Influence of the W/C and F/C ratios on the behavior of SCC
in the fresh state
Tabela 4: Vpliv razmerij W/C in F/C na vedenje SCC v sve`em stanju
Tests
SCC1 SCC2 SCC3
Vpaste =
35 %
Vpaste =
37.5 %
Vpaste =
40 %
Slump flow (mm) 570 700 750
L-box H2/H1 0.73 0.87 0.91
Sieve stability  (%) 5.66 7.68 15.01
Bleeding (‰) 0.93 1.15 3.15
The volume of paste is supposed to play two roles in
SCC. Initially, it fluxes the material by limiting the
contacts between the aggregates. Then, it splits off the
gravels sufficiently to avoid a formation of clusters
against the reinforcements, responsible for the flow
blocking. It is supposed that a minimum volume of paste
is needed to fulfill the two functions.
Table 4 shows that among the tested concretes, only
one presents a slump shorter than 60 cm. It is clear that
an increase in the volume of paste contributes to a signi-
ficant improvement in workability. This improvement is
principally due to the reduction of the coarse aggregate
content and, moreover, an increase in the volume, in
particular, in the crushed aggregates, induces an im-
portant friction.5 This is the case of the SCC with a
volume of paste of 35 %.
The results obtained from the L-box test given in
Table 4 show a good proportioning between the volume
of paste and the h2/h1 ratio of the concretes. The SCC
with a volume of paste of 35 % showed a poor flow in
the L-box test with the blocking h2/h1 ratio not exceeding
0.80.
For this purpose, it is noted that a paste content of
approximately 35 % does not allow an SCC to fulfill the
requirements of the AFGC.
Contrary to the slump and L-box tests, the concrete
with the 40 % volume of paste was at the extreme limit
to the domain of self-compactness since the percentage
of its latency is slightly higher than 15 %.
This concrete also developed a poor stability with
respect to bleeding, as its percentage exceeded 3 ‰. This
segregation is principally due to the increase in the
volume of paste compared to that of the aggregates.
3.2 Hardened state
3.2.1 Evolution of the compressive mechanical strength
An SCC presenting a good workability must contain
enough volume of paste to cover the aggregate surface in
order to minimize the frictions between the particles, on
the one hand, and, on the other hand, an additional quan-
tity of paste is necessary to obtain a better workability.6
When injected into the pores of the aggregates, this paste
is called the matrix and its apparent properties are
affected by the geometric arrangement of the skeleton.
N. BOUHAMOU et al.: SHRINKAGE BEHAVIOR OF A SELF-COMPACTING CONCRETE
Materiali in tehnologije / Materials and technology 47 (2013) 6, 763–769 765
Figure 1: Influence of the volume of paste on the compressive strengths
of SCCs
Slika 1: Vpliv koli~ine paste na tla~no trdnost SCC
More precisely, it is the average distance between two
coarse adjacent aggregates, called the maximum thick-
ness of the paste, which reflects the topology influence
of the skeleton: the shorter the distance, the higher is the
concrete strength. It is supposed that there is a minimum
volume of paste supporting this condition.
Figure 1 shows the effect of the volume of paste on
the compressive strength after (1, 7, 28 and 90) d. At the
first approximation, it is the density of the paste that
varies with respect to its volume percentage. It is noted
that the volume of paste within the interval (37.5–40 %)
contributes to the increase in the strengths for two aging
periods before causing their decrease when its rate
exceeds the critical value of about 37.5 %.
There is only one explanation for that: within the
volume-of-paste interval (37.5–40 %), the quantity of
paste is important and the distance between the aggre-
gates is larger, causing the frictions between them and,
consequently, engendering the weaker strengths.
The compressive strength deviation for SCC/OC
observed in Figure 1 can be explained with a higher
proportion of the paste used for SCC (376 l/m3 against
341 l/m3 for OC) and the absence of the limestone fillers
in OC which may have had a negative effect upon the
density.7
3.2.2 Free shrinkage
The concrete starts to undergo geometric strains as
soon as its installation in the framework has been
completed. These dimensional changes develop in diffe-
rent directions and are governed by various physical and
chemical phenomena. They take place in the material
free from stresses, creating a need for a loaded material.
3.2.2.1 Autogenous shrinkage
The development of autogenous shrinkage begins as
soon as the concrete begins to settle and it will stabilize
after a few months. This is linked to the auto-drying, i.e.,
the internal-humidity decrease due to the water con-
sumption by the hydrates.5
The volume of paste in the concrete is generally
much lower than in the mortar. And yet, it is the paste
which retracts and not the aggregate skeleton, which, on
the contrary, obstructs the shrinkage. Therefore, it can be
concluded that it is the volume-of-paste proportioning
that mainly governs the shrinkage values.
It is observed in Figure 2 that the autogenous shrin-
kage increases with the volume of paste. The shrinkage
progress, with respect to the volume of paste, does not
seem to be linear. This is perhaps due to the skeleton-
structure variation.
The autogenous shrinkage after 300 d is about 317
μm/m for the SCC with the Vpaste of 40 % and 275 μm/m
for the SCC with the Vpaste of 37.5 %, showing a
difference of 42 μm/m. The SCC with the Vpaste of 35 %
presents a shrinkage of 248 μm/m, i.e., a reduction of 69
μm/m compared to the first formula. These results reveal
that the recorded autogenous strains on the SCC with the
Vpaste of 40 % are much higher than those measured on
the two other SCCs. This highlights the importance of
the aggregate content. In fact, the strains measured on
the concrete are the effective strains of this hetero-
geneous material, consisting of the aggregates and the
paste (cement + fillers + superplasticizer + water). And
yet, the aggregates restrict the straining of the paste due
to the physical/chemical process linked to hydration.
Therefore, an aggregate-content reduction and an
increase in the volume of paste allow much higher
effective strains.
Indeed, the effective strains are determined with the
respective parameters (especially the elastic properties,
but also the dry density of the granular mixture).
At last, it can be said that the autogenous shrinkage
of an SCC increases with an increase in the volume of
paste. The reason is that only the paste creeps. The
volume of paste is also a parameter that influences the
shrinkage. The same result was reported in the litera-
ture.8,9
The results show that, on the whole, the autogenous
shrinkage of an SCC is greater than that of the ordinary
concrete (OC). After the first three months, the observed
difference between the two types of concrete varied from
10 % to 40 %, i.e., having a strain range of 1.1 to 1.4
(Figure 2).
3.2.2.2 Shrinkage due to drying and weight loss
Figure 3 shows a variation in the shrinkage time due
to drying, for an SCC with respect to the volume of
paste. This shrinkage can be considered to be due only to
the evaporation of the water contained in the hydrated
cement paste that develops from the surfaces when
exposed to external ambiance.
The drying shrinkage increases with an increase in
the volume of paste. Two distinct phases appear in the
strains developed: during the first phase of six months
N. BOUHAMOU et al.: SHRINKAGE BEHAVIOR OF A SELF-COMPACTING CONCRETE
766 Materiali in tehnologije / Materials and technology 47 (2013) 6, 763–769
Figure 2: Influence of the volume of paste on the autogenous shrink-
age of SCC
Slika 2: Vpliv koli~ine paste na avtogeno kr~enje SCC
the evolution speed is high and in the second phase it
slows down significantly to become constant. These
results are very compatible with those obtained from
reference.9
The strain analysis of the SCCs shows that the
formulation with the Vpaste of 40 % is characterized with a
higher drying shrinkage than that of the two other for-
mulations (Vpaste = 37.5 % and Vpaste = 35 %) after 300 d;
the differences are 60 μm/m when compared to the SCC
with the Vpaste of 37.5 % and 96 μm/m compared to the
SCC with the Vpaste of 35 %, corresponding to the in-
crease range of 12 % and 21 %, respectively.
The differences in drying shrinkage between the SCC
and the OC shown in Figure 4 are explained with the
paste proportioning that is very important for the SCCs
(the minimum proportioning for the SCCs is higher by
35 l/m3), leading to a much higher volume of hydrates in
the SCCs than in the OC and, therefore, creating a very
important volumetric strain.
On the other hand, Figure 4 shows that the water-
quantity variation for all the concretes is approximately
linear as long as the water evaporation continues with a
logarithmic time scale. All the concretes continue to
hydrate and the evaporated quantity of water is then the
function of the volume of paste.
Figure 5 brings an additional proof of the influence
of the volume of paste on the shrinkage due to drying. In
fact, different evolution curves of the shrinkage due to
drying, with respect to weight loss, show that after the
first evaporation of water with no consequences on the
shrinkage, there is the second stage, where the shrinkage
progresses with the water consumption.
It is observed that the weight losses for the SCC and
OC are different, but also that the drying-shrinkage evo-
lution with respect to the weight loss can be considered
as linear.
3.2.2.3 Total shrinkage
The relative results for the SCC total shrinkage
during 300 d are presented in Figure 6. From the
quantitative point of view, the experimental value for the
total shrinkage of the SCC containing a higher quantity
of paste (40 %), after 300 d, is 873 μm/m. The values,
for the same time, for the two other formulations are as
follows: 771 μm/m for the SCC containing 37.5 % of
paste and 708 μm/m for the SCC containing 35 % of
paste.
The shrinkage of the SCC with the Vpaste of 35 % is
slightly lower than that of the SCC with the Vpaste of 37.5
%, but it shows a defect in the stability, observed during
the L-box test, causing a problem according to the
AFGC recommendations.
The observed experimental data reveal significant
differences between the total shrinkages of the SCCs.
This deviation can be seen as a consequence of the
volume-of-paste increase.
The SCCs are more susceptible to deformability than
the OC (Figure 6) because of a higher quantity of paste.
N. BOUHAMOU et al.: SHRINKAGE BEHAVIOR OF A SELF-COMPACTING CONCRETE
Materiali in tehnologije / Materials and technology 47 (2013) 6, 763–769 767
Figure 4: Loss of weight with respect to the volume of paste – a loga-
rithmic scale
Slika 4: Zmanj{anje mase glede na koli~ino paste – logaritemsko
merilo
Figure 5: Drying shrinkage versus the loss of mass
Slika 5: Kr~enje pri su{enju v primerjavi z zmanj{anjem mase
Figure 3: Influence of the volume of paste on the drying shrinkage of
SCC
Slika 3: Vpliv koli~ine paste na kr~enje pri su{enju SCC
While making this conclusion, we neglect, however, the
fact that the SCC paste does not have the same compo-
sition as the OC paste (additions, superplasticizer, water
quantity). Besides, E. Proust8 has found an opposed ten-
dency of the other concretes. B. Persson10 has shown
that, with the same strength, there will be no behavioral
difference between SCC and OC, in spite of the volume-
of-paste deviation.
On the other hand, Y. Klug and K. Holschemacher11
found repeatedly in a large database that the total SCC
shrinkage was higher, by 10 % to 50 %, than that of the
OC. The SCC shrinkage was also found to be higher than
that of the OC by other researchers.12–14
3.3 Study of porosity
The results obtained from the porosimetry test for the
SCCs with respect to different values of paste propor-
tioning (Vpaste = 35 %, Vpaste = 37.5 % and Vpaste = 40 %)
are shown in Figures 7 and 8.
In Figure 7, it is observed that the cumulative volu-
me of the SCC pores containing a higher paste propor-
tioning (40 %) becomes more important than those of the
two other SCCs (the Vpaste of 35 % and the Vpaste of 37. 5 %)
as the pore size decreases. This seems to be in accordan-
ce with the autogenous shrinkage obtained for these
concretes.
As far as the shrinkage due to drying is concerned,
the smallest pores (<0.01 μm) in the SCC with the Vpaste
of 40 % have a greater proportioning than those of the
SCC with the Vpaste of 37.5 % and the SCC with the Vpaste
of 35 %. The experimental results given in section 3.2.2
reveal, effectively, that the drying shrinkage of the SCC
with a high quantity of paste (Vpaste = 40 %) is greater
than in the case of the two other formulations (Vpaste =
37.5 % and Vpaste = 35 %).
On the other hand, it is noted on the graph of Figure
8, for all the SCCs, that the porosities are closer in a very
marked peak zone, between 0.02 μm and 0.2 μm, and,
leaving aside the presence of the remainder of the peak
in the zone (10–100 μm) for the SCC with the Vpaste of
40 % and the SCC with the Vpaste of 35 %, the other for-
mulation shows a low porosity for higher pore sizes. The
paste corresponding to this SCC (Vpaste = 37.5 %) must
appear on this scale as a homogenous and, consequently,
tougher material.
Figure 8 shows, in fact, that a lower macroporosity
difference (0.1 %) seems to explain a much higher me-
chanical strength of the corresponding SCCs (section
3.2.1).
4 CONCLUSION
In this study, it is confirmed that it is possible to
produce SCCs with the Algerian local materials having
the same properties as those known internationally.
Therefore, it is advisable to pay attention to the import-
N. BOUHAMOU et al.: SHRINKAGE BEHAVIOR OF A SELF-COMPACTING CONCRETE
768 Materiali in tehnologije / Materials and technology 47 (2013) 6, 763–769
Figure 7: Cumulative volume of the mercury intrusion into the sample
versus the pore diameter
Slika 7: Kumulativni volumen vdora `ivega srebra v vzorec v primer-
javi s premerom por
Figure 8: Differential volume of the mercury intrusion into the sample
versus the pore diameter
Slika 8: Diferen~ni volumen vdora `ivega srebra v vzorec v primerjavi
s porami
Figure 6: Influence of the volume of paste on the total shrinkage
Slika 6: Vpliv koli~ine paste na celotno kr~enje
ance of the mix proportioning, which is of great influ-
ence, in particular, the W/C and F/C ratios, to achieve the
best required properties of an SCC.
Based on the obtained results, it has been shown that,
in the tested range (35–37.5 %), the volume of paste
contributes to an increase in the SCC strength. In the
range of 37.5–40 %, the paste proportioning has a signi-
ficant influence on the concrete behavior resulting in a
remarkable fall in the strength.
The SCC shrinkage is directly proportional to the
paste proportioning. The presence of a net difference bet-
ween the values for the tested SCC shrinkage seems to
indicate that the volume of paste is also a principal para-
meter affecting the shrinkage.
The results obtained from the mercury-intrusion
porosimetry test support certain assumptions put forward
and allow us to explain, in a pertinent manner, the beha-
vior of the SCCs in their hardened state. In fact, the
microstructural study of the samples has confirmed that
the hydration process is not the same for different SCCs,
depending on the paste quantity. The volumetric porosity
also shows relative differences in the SCC microstructure
formulations, confirming the results obtained experimen-
tally.
Finally, taking into account the relative measurement
uncertainties, it is believed that the stabilization of the
SCC strains will be faster than that for the OC, for which
the tendency to stabilization seems to differ with time.
Therefore, certain reduced deviations can be observed
over a long term. More tests made over several years will
be necessary for a better quantification of this pheno-
menon.
5 REFERENCES
1 H. Okamura, K. Ozawa, M. Ouchi, Self-compacting concrete, Struc-
tural Concrete, J. adv. concrete technology, 1 (2003) 1, 5–15
2 M. Shahul Hameed, A. S. S. Sekar, V. Sarswathy, Strength and per-
meability characteristics study of self-compacting concrete using
crusher rock dust and marble sludge powder, Arab J. Sci. Eng., 37
(2012), 561–574
3 AFGC, Les bétons autoplaçants, Recommandations provisoires,
French Association of Civil Engineering, 2000
4 N. Bouhamou, N. Belas, H. Mesbah, A. Mebrouki, Y. Ammar, Etude
du comportement à l’état frais des bétons autoplaçants à base de
matériaux locaux, Canadian Journal of Civil Engineering, 35 (2008)
7, 653–662
5 A. Neville, Propriétés des bétons, Centre de Recherche International
du Béton, Editions Eyrolles, 2000
6 S. G. Oh, T. Nogushi, F. Tomosawa, Toward mix design for rheology
of self-compacting concrete, Proc. First International Rilem Sym-
posium on Self-Compacting Concrete, Rilem, 1999
7 K. H. Khayat, Utility of statistical models in proportioning self-con-
solidating concrete, Proc. First International Rilem Symposium on
Self-Compacting Concrete, Rilem, 1999
8 E. Proust, Retrait et fluage des bétons autoplaçants – vers une com-
préhension des comportements différés, Doctorat Thesis, Institut
National des Sciences Appliquées, Toulouse, 2002
9 P. Turcry, Retrait et fissuration des bétons autoplaçants – Influence
de la formulation, Doctorat thesis, Ecole Centrale de Nantes et
l’Université de Nantes, Nantes, 2004
10 B. Persson, Creep, shrinkage and elastic modulus of self-compacting
concrete, Technical Report, Japan Society for Promotion of Science,
Lund Institute of Technology, Division of Building Materials, Lund
University, Lund, 1997
11 Y. Klug, K. Holschemacher, Comparison of hardened properties of
self-compacting and normal vibrated concrete, Proc. Third Interna-
tional Rilem Symposium on Self-Compacting Concrete, Rilem, 2003
12 J. Ambroise, J. Pera, S. Rols, Les bétons autonivellants, Annales du
Bâtiment et des Travaux Publics, 1 (1997), 37–41
13 C. Hu, B. Barbieri, Comparaison des retraits des bétons autonivelants
et d’un béton fluide traditionnel, Sciences des Matériaux et Pro-
priétés des bétons 1ères rencontres internationales, Toulouse, 1998,
265–272
14 S. Rols, J. Ambroise, J. Pera, Development of admixture for self-
leveling concrete, 5th CANMET/ACI Int. Conf. on Superplasticizer
and Other Chemical Admixtures in Concrete, SP 173, Rome, 1997,
493–509
N. BOUHAMOU et al.: SHRINKAGE BEHAVIOR OF A SELF-COMPACTING CONCRETE
Materiali in tehnologije / Materials and technology 47 (2013) 6, 763–769 769

M. POURANVARI, S. M. MOUSAVIZADEH: FAILURE MODE OF M130 MARTENSITIC STEEL RESISTANCE-SPOT WELDS
FAILURE MODE OF M130 MARTENSITIC STEEL
RESISTANCE-SPOT WELDS
NA^IN PORU[ITVE UPOROVNIH TO^KASTIH VAROV PRI
MARTENZITNEM JEKLU M130
Majid Pouranvari1, Seyed Mostafa Mousavizadeh2
1Materials and Metallurgical Engineering Department, Dezful Branch, Islamic Azad University, Dezful, Iran
2Department of Materials Engineering, Faculty of Engineering, Hakim Sabzevari University, Sabzevar, Iran
mpouranvari@yahoo.com
Prejem rokopisa – received: 2013-03-04; sprejem za objavo – accepted for publication: 2013-03-25
This paper discusses the transition from the interfacial to the pullout failure mode for the M130 martensitic, advanced
high-strength steel (AHSS) during a quasi-static tensile-shear test. It was studied whether the conventional/industrial weld-size
criteria can produce pullout failure modes for M130 spot welds. The minimum fusion-zone size required to ensure the pullout
failure mode during the tensile-shear test is estimated using an analytical model, taking into account the weld metallurgical
characteristics, with good accuracy. Based on the theoretical analysis it was concluded that the failure-mode transition in M130
resistance-spot welds is influenced by the sheet thickness, the fusion-zone hardness and the intercritical heat-affected-zone
hardness.
Keywords: martensitic steel, resistance-spot welds, failure mode, HAZ softening
Ta ~lanek obravnava prehod od medploskovnega do iztrganega na~ina poru{itve pri M130 martenzitnem naprednem
visokotrdnostnem jeklu (AHSS) med kvazistati~nim natezno-stri`nim preizkusom. Raziskano je bilo, ali je pri obi~ajnih
kriterijih industrijskih zvarov mogo~e dobiti poru{itev z iztrganjem pri to~kastih zvarih M130. Z analiti~nim modelom, ki
upo{teva metalur{ke zna~ilnosti zvara, je bilo mogo~e z dobro zanesljivostjo dolo~iti minimalno podro~je zlivanja za poru{itev z
iztrganjem med natezno-stri`nim preizkusom. Na podlagi teoreti~nih analiz je bilo ugotovljeno, da na prehod na~ina poru{itve
pri M130 uporovno varjenih zvarih vplivajo debelina plo~evine, trdota podro~ja zlivanja in interkriti~na trdota toplotno
vplivanega podro~ja.
Klju~ne besede: martenzitno jeklo, uporovno varjeni zvari, na~in poru{itve, meh~anje HAZ
1 INTRODUCTION
Advanced high-strength steels (AHSSs) have been
introduced to vehicle designs in an effort to improve the
collision-energy management and passenger safety,
while maintaining or reducing the vehicle weight that, in
turn, creates a better fuel economy.1 Due to a very high
strength and low formability, martensitic advanced high-
strength steels (AHSSs) are good candidates for high-
stiffness, load-transferring barriers and anti-intrusion
barriers (e.g., A- and B-pillar reinforcements, rocker
reinforcements, roof rails, front and rear bumpers, side-
wall members, cross beams and door beams) improving
crash management and protection of passengers during
side-impact collisions.2
Resistance-spot welding is the predominant joining
process in the automotive industry.3–5 The failure mode
of resistance-spot welds (RSWs) is a qualitative measure
of mechanical properties.4,5 Figure 1 shows a schematic
representation of the fracture surfaces in the main failure
modes during a mechanical testing of the spot welds.
Basically, spot welds can fail in two distinct modes
described as follows:6–11
1. Interfacial failure (IF) mode, in which a fracture pro-
pagates through the fusion zone. It is believed that
this failure mode has a detrimental effect on the
crashworthiness of the vehicles.
2. Pullout failure (PF) mode, in which a failure occurs
due to a withdrawal of the weld nugget from one
sheet. In this mode, a fracture may be initiated in the
BM, HAZ or HAZ/FZ depending on the metallur-
gical and geometrical characteristics of the weld zone
Materiali in tehnologije / Materials and technology 47 (2013) 6, 771–776 771
UDK 621.791.053:620.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)771(2013)
Figure 1: Schematic representation of: a) interfacial, b) pullout failure
modes
Slika 1: Shematski prikaz na~ina poru{itve: a) medploskovna, b) z
iztrganjem
and the loading conditions. Generally, the PF mode
exhibits the most satisfactory mechanical properties.
Thus, vehicle crashworthiness, as the main concern
in the automotive design, can be dramatically reduced if
spot welds fail due to the interfacial mode. The pullout
failure mode indicates, during a quality control, that the
same weld would have indeed been able to transmit a
high level of force, thus, causing a severe plastic defor-
mation in its adjacent components and an increased
strain-energy dissipation in crash conditions. Therefore,
it is necessary to adjust the welding parameters so that
the pullout failure mode is guaranteed.
The RSWs of AHSSs exhibit a higher tendency to
fail in the interfacial failure mode than those of tradi-
tional steels (i.e., low-carbon and HSLA steels).5–10
There are several size criteria for resistance-spot welds
included in industrial standards. "The effectiveness of
these criteria for evaluating AHSS spot welds, however,
has not been adequately addressed in the automotive
welding community; it was simply adopted from the
mild steel practice and applied to AHSS spot welds".12
The failure mode of the AHSS spot welds is a complex
phenomenon causing an interaction among the geometri-
cal factors, weld metallurgical properties and the loading
mode. Consequently, it is necessary to develop a spot-
weld sizing criterion based on the failure mode to reach a
deeper understanding of the factors governing the failure
mode of spot welds.
In this paper, the failure-mode transition of M130
martensitic steel resistance-spot welds under the tensile-
shear loading condition is studied. The objectives of this
study are:
1. to examine whether the existing weld-size criteria can
produce the pullout failure mode for M130 spot
welds;
2. to determine the critical fusion-zone size ensuring the
pullout failure mode of M130 spot welds.
2 EXPERIMENTAL PROCEDURE
An M130 martensitic sheet with a thickness of 2 mm
was used as the base metal in this study. The chemical
composition and mechanical properties of the M130
used in this study are given in Table 1. Spot welding was
performed using a 120 kVA AC pedestal-type resi-
stance-spot welding machine, operating at 50 Hz and
controlled by PLC. Welding was conducted using a 45°,
truncated-cone, RWMA class 2 electrode with a face
diameter 8 mm. To study the effects of the weld FZ size
on the failure mode and mechanical properties, spot
welding was performed in various welding conditions.
The welding time, electrode force and electrode holding
time after current-off were selected on the basis of the
thickness of the base material and were kept constant at
0.5 s, 4.5 kN and 0.2 cycles. The welding current was
changed step by step from 5 kA to 12.5 kA. The step size
was 0.5 kA.
The quasi-static tensile-shear test samples were pre-
pared according to the ANSI/AWS/SAE/D8.9-97 stan-
dard.13 Figure 2a shows the sample dimensions during
the tensile-shear test. Since the tensile-shear specimen is
asymmetrical, two shims with the same thickness were
added at the grip sections of the specimen to ensure an
alignment and to reduce the sheet bending and nugget
rotation. The tensile-shear tests were performed at the
cross-head speed of 2 mm/min with an Instron universal
testing machine. The peak load and failure energy were
extracted from the load-displacement curve (Figure 2b).
The failure modes were determined by observing the
weld fracture surfaces.
Microhardness test, a technique that has proven to be
useful for quantifying the microstructure/mechanical
property relationship, was used to determine the hard-
ness profile. The hardness profile in the diagonal direc-
tion was obtained using the Vickers microhardness
testing with an indenter load of 10 g and a speed of 10
M. POURANVARI, S. M. MOUSAVIZADEH: FAILURE MODE OF M130 MARTENSITIC STEEL RESISTANCE-SPOT WELDS
772 Materiali in tehnologije / Materials and technology 47 (2013) 6, 771–776
Figure 2: a) Tensile-shear specimen dimensions, b) a typical load-dis-
placement curve with the extracted parameters; Pmax: peak load, Wmax:
energy absorption
Slika 2: a) Dimenzije vzorcev za natezno-stri`ni preizkus, b) zna~ilna
krivulja obremenitev – raztezek s parametri; Pmax: najve~ja obreme-
nitev, Wmax: absorpcija energije
Table 1: Chemical composition and mechanical properties of the
investigated M130 steel
Tabela 1: Kemijska sestava in mehanske lastnosti preiskovanega jekla
M130
Chemical composition (w/%) Tensile properties
C Mn Si S P YS* (MPa) UTS** (MPa)
0.11 0.53 0.07 0.02 0.02 851 960
*Yield strength
** Ultimate tensile strength
indentations/min. The indentations were made on three
paths with a 0.3 mm spacing.
3 RESULTS AND DISCUSSION
3.1 Hardness characteristics
Macrostructural characteristics of the RSWs, particu-
larly the fusion-zone size, the microstructural and hard-
ness characteristics, play important roles in their failure
behavior and failure mode.14–16 Rapid heating and cool-
ing induced by the resistance-spot-welding thermal
cycles significantly alter the microstructure in the joint
zone. A typical macrostructure of M130 RSW is shown
in Figure 3a indicating that there are distinct zones in
the weldment including:
1. the fusion zone (FZ) or weld nugget, melted during
the welding process and later resolidified, showing a
cast structure. The macrostructure of the weld nugget
consists of columnar grains;
2. the heat-affected zone (HAZ) that is not melted but it
undergoes microstructural changes during welding;
3. the base metal (BM).
Figure 3b shows a typical hardness map of the M130
spot welds. The average hardness of the BM is about 300
HV. The average hardness of the FZ is 360 HV, which is
about 1.2 times higher than the hardness of the BM. The
microstructure of both BM and FZ is almost martensitic.
Therefore, the higher hardness of the FZ, compared to
the BM, can be related to the higher cooling rate during
RSW. It has been reported that an increasing cooling rate
can increase the dislocation density in the martensitic
laths increasing its hardness.
A reduction in the hardness (softening), with respect
to the BM, was observed in the HAZ. The minimum
hardness of the HAZ is about 215 HV (i.e., the maxi-
mum hardness reduction of 85 HV). The observed
softening can be related to the following reasons:
1. In the intercritical HAZ (ICHAZ) region, the peak
temperature is ranging between Ac1 and Ac3 and the
BM microstructure transforms into ferrite plus auste-
nite during heating. Due to the fast welding cooling
rates, austenite can subsequently transform into mar-
tensite. Therefore, the formation of an allotriomor-
phic ferrite phase in the ICHAZ leads to a hardness
reduction with respect to the fully martensitic BM.
2. In the subcritical HAZ (SCHAZ), the peak tempera-
ture is below Ac1 resulting in tempering of the meta-
stable martensite. This issue causes a reduction in the
hardness with respect to the BM.
3.2 Failure-mode analysis
Figure 4 shows the effect of the FZ size on the fail-
ure mode of M130 welds. As can be seen, the increasing
FZ size alters the failure mode from the interfacial to the
pullout mode. There is a critical FZ size, above which
the pullout failure mode is guaranteed. According to
Figure 4, the minimum FZ size of 9.2 mm is required to
ensure the PF mode.
In order to study the effect of the failure mode on the
mechanical performance of M130 resistance-spot welds,
box plots of peak load and energy absorption are shown
in Figures 5a and 5b, respectively. As can be seen in
Figure 5, the average peak load of those spot welds that
failed in the IF is lower than the peak load of the spot
welds that failed in the PF mode. Therefore, due to its
significant impact on the joint reliability, the failure
mode has been an interesting issue of some recent stu-
dies. The transition from the IF mode to the PF mode is
generally related to the increase in the FZ size above the
M. POURANVARI, S. M. MOUSAVIZADEH: FAILURE MODE OF M130 MARTENSITIC STEEL RESISTANCE-SPOT WELDS
Materiali in tehnologije / Materials and technology 47 (2013) 6, 771–776 773
Figure 4: Effect of the FZ size on the failure mode of M130 marten-
sitic steel resistance-spot welds
Slika 4: Vpliv velikosti FZ na na~in poru{itve uporovnega to~kastega
zvara pri martenzitnem jeklu M130
Figure 3: a) Typical macrostructure of M130 martensitic steel resi-
stance-spot weld. Failure paths in the interfacial failure (IF) mode and
pullout failure (PF) mode are also shown in the macrograph with the
arrows. b) Typical hardness map
Slika 3: a) Zna~ilna makrostruktura uporovnega to~kastega zvara pri
martenzitnem jeklu M130. S pu{~icami so prikazane poti poru{itve pri
medploskovni poru{itvi (IF) in pri poru{itvi z iztrganjem (PF). b) Zna-
~ilna razporeditev trdote
minimum value. In the following sections, the failure-
mode transition is compared with the existing criterion
for the weld-nugget sizing of a spot weld. For practical
purposes, it is interesting to compare the experimentally
determined minimum FZ size required to obtain the PF
mode with the existing industrial standards for weld-
nugget sizing. Various industrial standards have recom-
mended a minimum weld size for a given sheet thick-
ness:
1. According to AWS/ANSI/AISI,12 the weld-button
sizing used to ensure that the weld size was large
enough to carry the desired load, is based on equation
(1):
D = 4t0.5 (1)
where D is the weld-nugget size and t is the sheet
thickness (mm).
2. According to the Japanese JIS Z314017 and German
DVS292318 standards the required weld size is spe-
cified according to equation (2):
D = 5t0.5 (2)
3. The minimum weld-nugget size (equation 3) and the
nominal weld size (equation 4) are also frequently
used in certain industries:19
D = 0.69 (1.65t – 0.007)0.5 (3)
D = 0.86 (1.65t – 0.007)0.5 (4)
Figure 6 compares the experimentally determined
critical FZ size for the spot welds made on M130 steel 2
mm and the weld size recommended by the industry. As
can be seen, the most common weld-sizing criteria of
4t0.5 and 5t0.5 are not sufficient to produce a weld with the
PF mode. Also, equation (3) and equation (4) are not
sufficient to avoid the interfacial failure. Despite the fact
that the industrial recommendations work well for
obtaining the pullout mode of low-carbon steels, the
sizing based on these recommendations does not guaran-
tee the PF mode during the tensile-shear testing of the
AHSS spot welds. This is due to the fact that these
recommendations ignore the effect of metallurgical
factors on the failure mode so that these models include
only the sheet thickness for the sake of simplicity. Rada-
kovic and Tumuluru20 used finite-element modeling to
predict the resistance-spot-weld failure mode and the
loads in the shear-tension tests of advanced high-strength
steels (AHSS). In the finite-element model, the base
material, the heat-affected zone and the fusion-zone
properties were assumed to be homogeneous. According
to their finite-element modeling results, the critical
fusion-zone size can be expressed as follows:
DC = 4t (5)
According to their results, the critical weld size for a
thick sheet 2 mm is 8 mm. However, the experimental
value for a thick M130 2 mm is 9.2 mm, which is above
the recommended value of this mode. Indeed, this model
assumes the mechanical properties across the weldment
to be homogenous and does not take into account the
HAZ softening. In the following section an analytical
M. POURANVARI, S. M. MOUSAVIZADEH: FAILURE MODE OF M130 MARTENSITIC STEEL RESISTANCE-SPOT WELDS
774 Materiali in tehnologije / Materials and technology 47 (2013) 6, 771–776
Figure 6: Comparison of the experimentally determined critical FZ
zone (DC) for M130 steel 2 mm with the thickness-based recommen-
dations. The predicated value of DC obtained with the analytical
model is also shown.
Slika 6: Primerjava eksperimentalno dolo~enega kriti~nega podro~ja
FZ (DC) za jeklo M130 s priporo~ljivo debelino. Prikazana je tudi
napovedana vrednost za DC, dobljena z analiti~nim modelom.
Figure 5: Box plots of: a) peak load (Pmax) and b) failure energy
(Wmax) versus failure mode for M130 martensitic steel resistance-spot
welds
Slika 5: Prikaz: a) najve~je obremenitve (Pmax) in b) energije
poru{itve (Wmax) glede na na~in poru{itve to~kasto uporovno varjenih
zvarov martenzitnega jekla M130
model considering the weld hardness characteristics is
used to determine the critical FZ size to ensure the pull-
out failure mode.
3.3 Prediction of the failure mode taking into account
metallurgical factors
The failure of the resistance-spot welds during the
tensile-shear test can be described as a competition
between the shear plastic deformation of the fusion zone
(i. e., the IF mode) and the necking of the base metal or
HAZ (i. e, the PF mode).11,16 Spot welds usually fail in
the mode requiring less force during a fracture. It is well
documented that the driving force for the IF is the shear
stress along the sheet/sheet interface, while the driving
force for the PF is the tensile stress around the weld
nugget.11,16 In order to develop a model that predicts the
failure mode, the first necessary step is to develop the
equations for calculating the required force for each
failure mode.
First, we have to consider the peak load of the spot
welds in the interfacial mode. Considering the nugget as
a cylinder with a specific (D) diameter and (2t) height,
the failure load of the interfacial failure mode (FIF) can
be expressed with equation (6):
FIF = /4 D
2FZ (6)
where FZ is the shear strength of the weld nugget.
Now, the peak load of the spot weld in the pullout
failure mode is considered. It is assumed that, in the
pullout failure mode, the failure is initiated when the
maximum experienced radial tensile stress at the nugget
circumference reaches the ultimate tensile strength of the
failure location. Therefore, the failure load in the PF
mode can be expressed as:
FPF =  DtPFL (7)
where PFL is the ultimate tensile strength of the pullout
failure location. In the case of M130, where there is
significant softening in the ICHAZ, the peak load in the
PF mode can be expressed as follows:
FPF =  DtICHAZ (8)
where ICHAZ is the tensile strength of the ICHAZ. To
obtain the critical nugget diameter, DC, equations (6)
and (8) are intersected resulting in equation (9):
DC = 4tICHAZ/FZ (9)
The spot welds with D < DC tend to fail via the inter-
facial mode, contrary to the welds with D > DC that tend
to fail via the preferred pullout mode.
A direct measurement of the mechanical properties of
different regions of a spot weld is difficult. It is well
known that there is a direct relationship between the
tensile strength of a material and its hardness. The shear
strength of a material can be related linearly to its tensile
strength using a constant coefficient, f. Therefore,
equation 9 can be rewritten as follows:
DC = 4tHICHAZ/fHFZ (10)
Now, to validate the model, the experimental results
are compared with the analytical results. In the case of
the M130 martensitic steel, the average FZ hardness is
approximately 380 HV and the hardness of the softened
zone in the ICHAZ is about 225 HV. Therefore, the hard-
ness ratio of the FZ to the failure location is about 1.68.
According to the Tresca criteria, the ratio of the ultimate
shear strength to the ultimate tensile strength, f, is 0.5.
By substituting these values in equation (10), the critical
fusion-zone size is calculated to be 9.4 mm. Figure 6
shows that this value is a reasonable estimation of the
critical FZ size.
Figure 7 compares the analytical model and the
existing criteria for a sheet 2 mm. As can be seen, the
sizing based on 4t0.5, Eq.3, 5t0.5, Eq.4 and 4t criteria is not
appropriate for obtaining the PF mode when the hardness
ratios are smaller than 2.85, 2.65, 2.3, 2.1 and 2, respec-
tively. As the hardness ratio is reduced, the discrepancy
between the analytical model and the industrial sizing
criteria is increased.
4 CONCLUSIONS
In this work, the failure-mode transition of M130
martensitic resistance-spot welds is investigated. The
following conclusions can be drawn from this work:
1. A reduction in the hardness (softening), with respect
to the BM, was observed in the HAZ. The heat-
affected-zone softening reduced the strength of the
HAZ, resulting in the strain localization and, hence,
causing the PF mode.
2. There is a critical FZ size ensuring the pullout failure
mode during the tensile-shear test. According to the
M. POURANVARI, S. M. MOUSAVIZADEH: FAILURE MODE OF M130 MARTENSITIC STEEL RESISTANCE-SPOT WELDS
Materiali in tehnologije / Materials and technology 47 (2013) 6, 771–776 775
Figure 7: Comparison of the mode prediction and thickness-based
recommendations when the sheet thickness (t) is 2 mm. The hardness
ratio (K) is defined as the ratio of the FZ hardness to the pullout-fail-
ure-location hardness.
Slika 7: Primerjava napovedane vrste poru{itve in priporo~il debeline,
~e je debelina plo~evine (t) 2 mm. Razmerje trdote (K) je dolo~eno kot
razmerje trdote FZ in lokalne trdote pri iztrganju.
theoretical analysis presented in this study, the
failure-mode transition of the M130 spot welds is
governed by the sheet thickness, FZ hardness and
ICHAZ hardness.
3. The existing industrial weld-nugget sizing criteria are
not sufficient to ensure the pullout failure mode dur-
ing the tensile-shear testing of the M130 resistance-
spot welds. The proposed model can predict the fail-
ure mode with good accuracy.
5 REFERENCES
1 M. Pouranvari, Influence of Welding Parameters on Peak Load and
Energy Absorption of Dissimilar Resistance Spot Welds of DP600
and AISI1008 Steels, Canadian Metallurgical Quarterly, 50 (2011),
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lines, Version 4.1, International Iron and Steel Institute, Washington
DC, 2009
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187–194
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and failure mode on peak load and energy absorption of advanced
high strength steel spot welds under lap shear loading conditions,
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on the welding variables in the resistance spot welding of ferrite-
martensite DP980 advanced high-strength steel, Mater. Tehnol., 46
(2012) 6, 665–671
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Automotive Steels: Process, Structure and Properties, Sci. Technol.
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performance of dual phase steel resistance spot welds, Sci. Technol.
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ence of weld size and microstructure of dissimilar AHSS resistance
spot welds, Sci. Technol. Weld. Joining, 13 (2008), 769–776
10 M. S. Khan, S. D. Bhole, D. L. Chen, G. Boudreal, E. Biro, J. V.
Deventer, Canadian Metallurgical Quarterly, 48 (2009), 303–310
11 M. Pouranvari, H. R. Asgari, S. M. Mosavizadeh, P. H. Marashi, M.
Goodarzi, Effect of weld nugget size on overload failure mode of
resistance spot welds, Sci. Technol. Weld. Joining, 12 (2007),
217–225
12 X. Sun, E. V. Stephens, M. A. Khaleel, Effects of Fusion Zone Size
and Failure Mode on Peak Load and Energy Absorption of Advanced
High-Strength Steel Spot Welds, Welding Journal, 86 (2007),
18s–25s
13 Recommended Practices for Test Methods and Evaluation the Resi-
stance Spot Welding Behavior of Automotive Sheet Steels, ANSI/
AWS/SAE D8.9-97, 1997
14 M. Pouranvari, S. P. H. Marashi, Similar and dissimilar RSW of low
carbon and austenitic stainless steels: effect of weld microstructure
and hardness profile on failure mode, Mater. Sci. Technol., 25
(2009), 1411–1416
15 M. Pouranvari, Failure mode transition in similar and dissimilar resi-
stance spot welds of HSLA and low carbon steels, Canadian Metal-
lurgical Quarterly, 51 (2012), 67–74
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sition in AHSS resistance spot welds. Part II: Experimental investi-
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17 Japanese Industrial Standard, Method of Inspection for Spot Welds,
JIS Z 3140, 1989
18 German Standard, Resistance Spot Welding, DVS 2923, 1986
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impact strengths of spot welded HSLA and low carbon steel joints,
SAE 820281, Society of Automotive Engineers, Warrendale, PA,
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776 Materiali in tehnologije / Materials and technology 47 (2013) 6, 771–776
Z. A]IMOVI] et al.: SYNTHESIZING A NEW TYPE OF MULLITE LINING
SYNTHESIZING A NEW TYPE OF MULLITE LINING
SINTEZA NOVE VRSTE OBLOGE IZ MULITA
Zagorka A}imovi}1, Anja Terzi}2, Ljubi{a Andri}3, Ljubica Pavlovi}3,
Marko Pavlovi}1
1University of Belgrade, Faculty of Technology and Metallurgy, Belgrade, Serbia
2Institute for Materials Testing, Belgrade, Serbia
3Institute for Technology of Nuclear and Other Mineral Raw Materials, Belgrade, Serbia
anja.terzic@institutims.rs
Prejem rokopisa – received: 2013-03-07; sprejem za objavo – accepted for publication: 2013-03-27
Various possibilities for developing new mullite-based refractory linings that can be applied in a casting process were
investigated and are presented in this paper. An optimization of the refractory-lining composition design with the controlled
rheological properties was achieved by applying different lining components and altering the lining-production procedure.
Mullite was used as a high-temperature filler. A mullite sample was tested with the following methods: X-ray diffraction
analysis, differential thermal analysis and scanning-electron microscopy. The particle shape and particle size were analyzed with
the program package for an image analysis called OZARIA 2.5. It was proved that an application of this type of lining has a
positive effect on the surface quality, structural and mechanical properties of the castings of Fe-C alloys obtained by casting into
sand molds, according to the method of expandable patterns (the EPC casting process).
Keywords: refractory lining, mullite, quality of casting, EPC casting process
V ~lanku so predstavljene razli~ne mo`nosti za razvoj nove mulitne ognjevzdr`ne obloge, uporabne pri postopku ulivanja.
Optimiranje sestave ognjevzdr`ne obloge s kontroliranimi reolo{kimi lastnostmi je bilo dose`eno z uporabo razli~nih sestavin
obloge in s spremembo postopka izdelave obloge. Mulit je bil uporabljen kot visokotemperaturno polnilo. Vzorec mulita je bil
preiskovan z naslednjimi metodami: z rentgensko difrakcijo, diferen~no termi~no analizo in vrsti~no elektronsko mikroskopijo.
Oblika in velikost zrn sta bili dolo~eni s programsko opremo za analizo slik OZARIA 2.5. Dokazano je bilo, da uporaba te vrste
obloge ugodno vpliva na kvaliteto povr{ine, strukturo in mehanske lastnosti ulitkov iz Fe-C zlitin pri ulivanju v pe{~ene forme
po metodi ekspandirane pene (EPC-postopek ulivanja).
Klju~ne besede: ognjevzdr`na obloga, mulit, kvaliteta ulitkov, EPC-postopek ulivanja
1 INTRODUCTION
The main role of a lining is the formation of an
efficient, unbreakable and firm refractory barrier which
separates the sandy substrate from the liquid metal flow.
For such a role, certain lining properties are required:
high refractoriness, suitable gas permeability, simple
application, good adhesion to a sandy mold and polymer
model, simple adjustment of the lining-layer thickness
and high drying rate. These requirements can be success-
fully fulfilled with an optimization of the lining compo-
sition and production technology.1–3
The basic characteristic of the EPC casting process is
that the patterns and gating of molds made of polymers
stay in the cast until the liquid-metal inflow occurs. The
pattern-decomposition kinetics is the function of the
liquid-metal temperature, with which the pattern comes
in contact. The important factors influencing the pattern
decomposition and, consequently, the evaporation pro-
cess, besides temperature and pattern density, are: type
of refractory lining, thickness of the lining layers cover-
ing the evaporable pattern, type and size of the sand
grains and their granulation, permeability of the sandy
model, gating of the mold construction, etc.4–7 Manufac-
turing the castings with the projected application quality
by means of the EPC process has not been investigated
enough and, thus, there is a need for a systematic
research of the štriad’ including structure/properties/
technology, to which special attention was paid in this
paper.
The mullite was chosen as the refractory-lining filler
due to the following properties: low thermal conductivity
(6 W m–1 K–1);8 low coefficient of the linear thermal ex-
pansion (5.4 · 10–6 °C–1 at 25 °C);9 high thermal-shock
resistance (quenching/500) and high maximum-use tem-
perature of 1650 °C;10 flexural strength of 180 MPa;
elastic modulus of 151 GPa; compresive strength of 1310
MPa; hardness of 1070 kg mm–2; extreme resistance to
liquid-metal absorption;6 no gas production when in
contact with liquid metal. Different additive types and
various quantities were tested in order to enable the best
possible absorption between the additives and the
refractory filler particles and, thus, the maintenance of
the filler in a dispersed state and prevention of the filler
build-up or segregation.
2 EXPERIMENTAL WORK
For the synthesis of mullite (3Al2O32SiO2) the
mixture of kaolin and alumina was used with an addition
of a mineralizer (1 % of NaF). Alumina was added in
order to achieve the mullite stoichiometric ratio of 3 : 2.
The crushing of the reactive components and the homo-
Materiali in tehnologije / Materials and technology 47 (2013) 6, 777–780 777
UDK 621.74:666.76:666.762.14 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)777(2013)
genization were performed in a planetary ball mill PM
100 with sintered, aluminium-oxide grinding balls. After
the homogenization and an addition of the mineralizer,
the mixture was wetted with water and, subsequently,
pressed in the mold with the 100 N/mm2 pressure and,
afterwards, dried. The synthesis of mullite was per-
formed by means of isothermal heating in the laboratory
high-temperature šNetzsch’ furnace at the temperature of
1450 °C, with the heating rate of 10 °C/min in air atmo-
sphere.
The lining compositions were defined (Table 1) and
the lining-component preparation methods were deter-
mined.
Table 1: Composition of refractory mullite-based linings
Tabela 1: Sestava ognjevzdr`ne obloge na osnovi mulita
Component Refractory liningbased on alcohol
Refractory lining
based on water
ararsid14156519
Refractory filler
Mullite with the
grain size of
35–40 μm,
90–94 %
Mullite with the
grain size of
40 μm, 93–95 %
Binding agent Colophonium
(C20H30O2), 2.5 %
Bentonite, 2.5%;
Bindal H, 0.5%;
Na3P3O3, 1–3 %
Additive/
suspension
Bentone 25,
0.8–1 %
Carboxymethyl-
cellulose, 0.5–1 %
Solvent Alcohol Water
Refractory linings were applied to the sandy molds
with a brush. During the application of the refractory
lining on the polymer model using the technique of
immersion into the tank with a lining, the process
parameters were: the suspension density of 2 g/cm3, and
the suspension temperature of 25 °C. The drying proce-
dure was as follows: for the water-based linings, the
duration of drying the first layer was 2 hours and the
final layer was dried for 24 h; for the alcohol-based
linings burning was used. The thickness of the lining
layer on the model after drying was 0.5–1 mm.
For casting, Fe-C alloys were used. The casting tem-
perature was 1350 °C. For the production of sandy
molds, the mold mixture based on quartz sand was used,
with the mean grain size being 0.17 mm, with an addi-
tion of bentonite (3 %) and dextrin (0.5 %). To produce
the EPC-casting-process molds, dry quartz sand with the
mean grain size of 0.25 mm was used and evaporable
models were made of polystyrene with the density of
20 kg/m3.
The mineral-phase composition of mullite was ana-
lyzed by means of X-ray powder diffraction (an XRD-
Philips PW-1710 diffractometer). DTA was performed
with a Shimadzu DTA-50 apparatus. The microstructure
of the samples was characterized with the scanning-elec-
tron-microscopy method (SEM) using a JEOL JSM-6390
Lv microscope. Distribution of the refractory filler and
bonding agent in a lining suspension was conducted with
a polarized-light optical microscope of the passing-light
JENAPOL type (Carl Zeiss – Jena). The analysis of the
particle size and shape factor was conducted with the PC
software OZARIA 2.5.
3 RESULTS AND DISCUSSION
In Figure 1 the results of the X-ray structural ana-
lysis of mullite powder are shown. The mean grain size
of the refractory filler was between 35–40 μm, the
grain-shape mean factor was 0.63, which means that the
grains are round and suitable for the production of
homogeneous linings.
A DTA curve for the mullite sample is presented in
Figure 2. It can be concluded that mullite has a high
refractoriness and, thus, it is suitable for casting Fe-C
alloys.
In Figure 3, the results of the qualitative mineralo-
gical analysis of the filler based on mullite are shown.
The analysis shows that the mullite particles are princi-
pally of equal size and morphology, but there are also
some differences in the particle size. This is favorable
Z. A]IMOVI] et al.: SYNTHESIZING A NEW TYPE OF MULLITE LINING
778 Materiali in tehnologije / Materials and technology 47 (2013) 6, 777–780
Figure 2: DTA curve of mullite
Slika 2: DTA-krivulja mulita
Figure 1: X-ray diffractogram of mullite
Slika 1: Rentgenski difraktogram mulita
since the particles of diverse granulations contribute to
forming an equalized, continuous lining layer on the
polymer pattern, due to a better harmony between the
particles.
A histogram of the mullite-powder particle size is
given in Figure 4, while Figure 5 represents a histogram
of the mullite-particle shape factor.
Microstructural analyses of the lining filler and the
suspension samples proved that the filler particles have
predominantly a uniform size and morphology. The filler
particles were mainly oval (Figure 5). The filler particles
with the sizes of 20–40 μm were predominantly present,
making 50 % of the total mullite-particle granulate (Fig-
ure 4). It was estimated that the oval-shaped particles of
various grain sizes contribute to the forming of a uniform
and consistent lining layer on the mold and model sur-
faces due to the stronger interrelations among the parti-
cles, which was proved with the results of the lining-
property tests. The sediment-stability test performed on
all three types of refractory linings showed that solid
matters build up to the amount of 5–8 %, which is in
accordance with the lining-quality requirements.
The homogeneity of the refractory-filler distribution
also depends on the suspension preparation and appli-
cation technology. The filler-particle concentration and
the adhesion have a significant influence on the suspen-
sion rheological properties. In the case of an increase in
the concentration of the filler in suspension, it was esta-
blished that the adhesion forces among mullite particles
were also increased and that, under the influence of
rheological additives and binding agents, constant and
uniform coating layers might be formed on the applied
surfaces. The lining adheres easily to the surfaces
applied and does not get cracked or wiped out after
drying, which was proved after a visual examination of
several test samples (Figure 6).
An application of diluted linings or the linings whose
components are not homogenized enough with careful
stirring does not create a good surface adhesion. Further-
more, in this case dried lining layers are not uniform,
while the linings applied in thicker layers often crack
after drying (Figure 7).
Z. A]IMOVI] et al.: SYNTHESIZING A NEW TYPE OF MULLITE LINING
Materiali in tehnologije / Materials and technology 47 (2013) 6, 777–780 779
Figure 6: Appearance of the mullite-based lining on a test sample
without visible flaws
Slika 6: Videz obloge na osnovi mulita na preizkusnem vzorcu brez
vidnih napak
Figure 4: Mullite-particle size frequency
Slika 4: Velikostna razporeditev mulitnih zrn
Figure 5: Distribution of the particle-shape factor of mullite
Slika 5: Razporeditev faktorja oblike zrn mulita
Figure 3: Mullite particles (SEM)
Slika 3: SEM-posnetek zrn mulita
Due to the results of preliminary testing and visual
examination of the samples, further tests were carried out
using a lining with the density of 2000 kg/m3. The lining
components were carefully stirred during the application
in order to achieve a homogeneous filler distribution in
the suspension.
The experiment showed that alcohol-based linings
are suitable for an application in sandy molds and cores.
Both alcohol-based and water-based linings applied did
not penetrate the test-tube surfaces made of polystyrene.
Applications of the three types of linings, during the
processes of casting, in the sandy molds, and the EPC
casting procedure enable a production of the castings
with advanced quality settings. The prepared linings with
a 2000 kg/m3 density were applied in two layers and,
after drying, it was noted that constant, thin lining layers
were formed on the molds and polymer models. This
method provided a good gas permeability of the lining
layers, enabling a faster liquid-metal cooling in the mold
and a formation of a tiny-grain cast structure, confirmed
with the results of testing the structural and mechanical
properties of the castings. The application of thin lining
layers positively influences the porosity reduction of the
castings.
4 CONCLUSION
The investigation showed that the linings based on
synthesized mullite met the casting-application require-
ments creating an acceptable casting-surface finish. This
justifies the applied sintering regime at 1450 °C for the
mullite synthesization. The investigated linings showed a
possibility of an easy application on the sandy molds and
polymer models: the linings evenly flew down during the
pouring and they were easily applied with a brush with-
out leaving traces, a leakage or a formation of coat drops
or lumps. After drying, the lining surfaces were smooth,
the layers had an even thickness and the models did not
show any bubbling, cracks or peeling. The linings added
stiffness to the cluster and allowed the foam-decompo-
sition products to escape. The application of the linings
in thinner layers, approximately 0.5–1 mm, due to im-
proved permeability, showed that the cast quality was
higher, having no visible flaws (porosity, bubbles or
cracks). Although all the investigated refractory linings
gave satisfying properties, water-based linings are more
ecologically and economically sustainable than alcohol-
based linings.
Further research on improving this type of refractory
linings will be done in terms of improved properties of
the mullite-based filler using a mechanical activation
process enhancing an improvement of the lining rheolo-
gical properties and the lining suspension stability. Also,
further research of the types of refractory linings inve-
stigated in this paper will concentrate on the mechanical
properties of metal castings in order to confirm the con-
nection between the mechanical performances and the
advanced morphological characteristics of the linings.
Acknowledgements
This investigation was supported and funded by the
Ministry of Education, Science and Technological Deve-
lopment of the Republic of Serbia and it was conducted
within projects 33007, 172057 and 45008.
5 REFERENCES
1 F. G. Martins, C. A. Silva de Oliveira, Study of full-mold casting
process for Al–Si hipoeuthetic alloys, Journal of Materials Process-
ing Technology, 179 (2006), 196–201
2 M. R. Barone, D. A. Caulk, A foam ablation model for lost foam
casting of aluminum, International Journal of Heat and Mass
Transfer, 48 (2005), 4132–4149
3 H. Gökce, D. Agaogullari, M. Lütfi Övecoglu, I. Duman, T. Boyraz,
Characterization of microstructural and thermal properties of
steatite/cordierite ceramics prepared by using natural raw materials,
Journal of the European Ceramic Society, 31 (2011), 2741–2747
4 R. E. Moore, Refractories, Structure and Properties, Encyclopedia of
Materials: Science and Technology, Press, Oxford 2001, 8079–8099
5 A. Prsti}, Z. A}imovi} - Pavlovi}, S. Gruji}, M. \uri~i}, Lj. Andri},
Different ceramic linings for application in foundry, Proceedings of
the 43rd October Conference on Mining and Metallurgy, Kladovo,
2011, 79–82
6 J. Stjernberg, B. Lindblom, J. Wikström, M. L. Antti, M. Odén,
Microstructural characterization of alkali metal mediated high tem-
perature reactions in mullite based refractories, Ceramics Interna-
tional, 36 (2010) 2, 733–740
7 U. Steinhauser, W. Braue, J. Göring, B. Kanka, H. Schneider, A new
concept for thermal protection of all-mullite composites in com-
bustion chambers, Journal of the European Ceramic Society, 20
(2000) 5, 651–658
8 G. Brunauer, F. Frey, H. Boysen, H. Schneider, High temperature
thermal expansion of mullite: an in situ neutron diffraction study up
to 1600 °C, Journal of the European Ceramic Society, 21 (2001),
2563–2567
9 N. M. Rendtorff, L. B. Garrido, E. F. Aglietti, Thermal shock resi-
stance and fatigue of Zircon–Mullite composite materials, Ceramics
International, 37 (2011) 4, 1427–1434
10 N. Rendtorff, L. Garrido, E. Aglietti, Mullite–zirconia–zircon com-
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tional, 35 (2009), 779–786
Z. A]IMOVI] et al.: SYNTHESIZING A NEW TYPE OF MULLITE LINING
780 Materiali in tehnologije / Materials and technology 47 (2013) 6, 777–780
Figure 7: Appearance of the mullite-based lining on a test sample
with visible flaws
Slika 7: Videz obloge na osnovi mulita na preizkusnem vzorcu z
vidnimi napakami
T. GALETA et al.: INFLUENCE OF PROCESSING FACTORS ON THE TENSILE STRENGTH OF 3D-PRINTED MODELS
INFLUENCE OF PROCESSING FACTORS ON THE
TENSILE STRENGTH OF 3D-PRINTED MODELS
VPLIV PROCESNIH DEJAVNIKOV NA NATEZNO TRDNOST
TRIDIMENZIONALNO TISKANIH MODELOV
Tomislav Galeta, Ivica Kladari}, Mirko Karaka{i}
J. J. Strossmayer University of Osijek, Mechanical Engineering Faculty in Slavonski Brod, Trg Ivane Brli}-Ma`urani} 2, 35000 Slavonski
Brod, Croatia
tgaleta@sfsb.hr
Prejem rokopisa – received: 2013-03-08; sprejem za objavo – accepted for publication: 2013-05-21
Since the models produced with three-dimensional printing are not as strong as the ones made with the other rapid prototyping
technologies, the main objective of this research was to determine the influence of selected processing factors on the tensile
strength and to determine the factor combination that provides the highest strength. Test samples were prepared on a 3D printer
with variations in the layer thickness, building orientation and infiltrant type. The secondary objective was to evaluate the
application of more affordable alternative infiltrants used instead of genuine infiltrants. The results of the tensile test revealed
that the strength of 3D-printed samples comes mainly from the infiltrants, but it may be additionally increased by selecting the
best combination of the other two processing factors. The strength of the samples infiltrated with alternative infiltrants was
equivalent to that obtained with genuine infiltrants, thus confirming the use of alternative infiltrants.
Keywords: rapid prototyping, three-dimensional printing, tensile strength
Modeli, ki so proizvedeni s tridimenzionalnim tiskanjem, niso tako mo~ni v primerjavi z drugimi tehnologijami hitrega
prototipiranja, zato je bil glavni cilj raziskave ugotoviti, kak{en vpliv imajo izbrani procesni dejavniki na natezno trdnost, in
kombinacijo, ki zagotavlja najve~jo trdnost. Vzorci za preizkuse so bili pripravljeni na tridimenzionalnem tiskalniku s
spreminjanjem debeline plasti, smeri nalaganja in vrste veziva. Dodaten cilj je bil oceniti uporabo dostopnej{ih nadomestnih
vezivnih sredstev, namesto originalnih. Rezultati preizkusa natezne trdnosti so pokazali, da je trdnost 3D tiskanih vzorcev
najbolj odvisna od vezivnega sredstva, mogo~e pa jo je pove~ati z izbiro najbolj{e kombinacije drugih dveh procesnih
dejavnikov. Trdnost vzorcev, ki so bili infiltrirani z alternativnimi vezivnimi sredstvi, je bila enakovredna tisti, pridobljeni z
originalnimi vezivnimi sredstvi, kar potrjuje uporabo alternativnih vezivnih sredstev.
Klju~ne besede: hitro prototipiranje, tridimenzionalno tiskanje, natezna trdnost
1 INTRODUCTION
The models produced with three-dimensional print-
ing (3DP) are not as strong as the ones made with the
other rapid prototyping (RP) technologies. Several
authors emphasized the issue of the 3DP strength in their
researches. Pilipovi} et al.1 compared the properties of
the samples made with two similar RP procedures: the
samples produced with "the stipulated standard 3DP
procedure" showed mechanical properties that were
inferior to those of the samples produced with the hybrid
Polyjet procedure. The results obtained in that paper and
frequent customer demands for better tensile properties
motivated us to perform further research in order to
improve the strength of the models produced on the con-
sidered model of 3DP machines. An initiative research
was performed and its results were published in a
master’s degree thesis2 and in our papers published in
conferences.3,4
There is a constant effort of the researchers to impro-
ve the strength and other mechanical properties of 3DP
models. This is evident from the number of recently
published papers. Patirupanusara et al. performed studies
to enhance the mechanical properties of 3DP specimens
on the basis of polymethyl methacrylate.5,6 Suwanprateeb
found that the use of light-cured acrylate resin as an
infiltrant can enhance the flexural modulus and the
flexural strength of natural-polymer-based 3DP parts to
be close to the general use of polymethyl methacrylate
resin.7 Chumnanklang et al. revealed that pre-coated
particles would yield a stronger 3DP hydroxyapatite
part.8 Hydroxyapatite is widely used as a medical, highly
biocompatible bone-substitute material. Furthermore,
several researchers gave their contributions on the
subjects closely related to the mechanical properties of
3DP models.9–13
The overall work on the enhancement of 3DP mecha-
nical properties can also be noticed on the market. The
manufacturers of 3DP systems frequently deliver new
enhanced models of printers, together with improved
materials, system software or improved alternative spare
parts for the existing models. These efforts will reduce
the damages to 3DP models that occur during explora-
tion. However, from the 3DP owner’s point of view, it
should be stressed that most damages occur on green
models, i.e., before post-processing and infiltration
(Figure 1), although customers are usually not informed
about such damages. In the previous research2, the
obtained average tensile strength of green samples was
0.95 MPa. However, if the samples were dried in the
Materiali in tehnologije / Materials and technology 47 (2013) 6, 781–788 781
UDK 621.7:004.896 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)781(2013)
oven for at least two hours at 55 °C, the average tensile
strength of green samples increased to 1.52 MPa.
Original equipment manufacturers (OEM) constantly
make additional efforts to enhance the mechanical pro-
perties of the green 3DP model. OEM efforts result in
new base materials, printer upgrades and a development
of new 3DP machines.
The main objective of our research was to determine
the influence of the selected 3DP processing factors on
the tensile strength and to determine the factor combi-
nation that provides the highest strength. The secondary
objective was to evaluate the application of the alterna-
tive infiltrants that we used instead of the genuine manu-
facturer’s infiltrants for the considered 3DP system. If
the obtained properties are equivalent, the use of the
alternative infiltrants can help to reduce the printing
costs and to acquire infiltrants from the more available
alternative suppliers. For this purpose we carried out a
set of experiments on the considered 3DP system.
2 TESTING THE EQUIPMENT AND SAMPLE
MATERIALS
The 3D printer, used for these experiments, was
model Z310, a product of Z Corporation. It is a low-cost,
monochrome 3D printer suitable for the RP education or
for small and medium-sized companies. The printer
firmware version was 10.158 and the test samples were
prepared in printer software ZPrint version 7.5.2314,15
(Figure 2).
The considered 3D printer combines the layered
approach of RP technologies and conventional ink-jet
printing. It prints a binder fluid through a conventional
ink-jet print head into the powder, one layer onto ano-
ther, from the lowest model cross-section to the highest
one (Figure 3). After printing, the printed models are
dried in the building box (Figure 4), then removed from
the powder bed, de-powdered with compressed air, dried
in the oven and infiltrated for the maximum strength.
There are several base materials, i.e., powder types
available for the above mentioned 3D printer. For our
T. GALETA et al.: INFLUENCE OF PROCESSING FACTORS ON THE TENSILE STRENGTH OF 3D-PRINTED MODELS
782 Materiali in tehnologije / Materials and technology 47 (2013) 6, 781–788
Figure 3: 3D printer main components – a section view
Slika 3: Glavni sestavni deli 3D-tiskalnika – pregled komponent
Figure 1: 3DP model of a spherical gas tank – damaged (left) and repaired with infiltration (right)
Slika 1: 3DP-model okroglega rezervoarja plina – po{kodovan (levo) in popravljen z infiltriranjem (desno)
Figure 2: Test-sample build setup in software ZPrint
Slika 2: Razporeditev preizkusnih vzorcev s programsko opremo
ZPrint
experiment, we used the plaster-based zp130 powder
with an appropriate binder, zb56. The zp130 powder is
recommended for the accuracy and for delicate models.
It is a mixture of plaster, vinyl polymer and sulphate
salt.16
All the test samples were dried twice before the infil-
tration, as recommended in15,17: first in the printer’s
building box for one hour and then, after de-powdering,
in the oven for at least two hours at 55 °C.
After drying, the samples were infiltrated with an
appropriate infiltrant, taking into account the appropriate
combination of the experiment. The applied infiltrants
were: cyanoacrylate-based Loctite 40618; epoxy-resin-
based Loctite Hysol 948319 and normal-wax Cera Alba.
We applied the alternative infiltrants that are the most
similar to the corresponding genuine manufacturer’s
infiltrants: Loctite 406 instead of Z-Bond20; Loctite
Hysol 9483 instead of Z-Max; and normal-wax Cera
Alba instead of Paraplast X-TRA Wax.21
Subsequent to the infiltration, i.e., prior to the ten-
sile-strength test, all the samples rested in room condi-
tions for a minimum of 24 h to obtain the final strength,
as recommended for the epoxy-resin infiltrant.22 We
measured the dimensions of the test samples with a
digital caliper Lux Profi, model 572587, with a measu-
rement range of 0–150 mm and an accuracy of 0.01 mm.
The tensile test was performed at room temperature
on the tensile-testing machine ZMGi 250 made by VEB
Thuringer Industriewerk with the jaw-motion speed of 1
mm/min.
3 DESIGN OF THE EXPERIMENT
The models for the tensile tests used in our experi-
ments were those defined with standard ISO 527:1993.
The nominal dimensions of the test sample are presented
in Figure 5. The dimensions measured for the tensile test
were: the neck width (W1) and the height (H). The other
linear dimensions like the total length (L) and the width
at the end (W) were controlled according to the require-
ments of the tensile test.
We considered the combinations of the following
processing factors: the layer thickness, the building
orientation and the infiltrant type. Every combination of
the processing factors was denoted with an appropriate
unique label made of three characters (Table 1).
In the considered ZPrint software version, the layer
thickness can be selected from two possible values: 0.1
mm or 0.0875 mm. The first thickness is the default for
the printer and it is, therefore, marked with number 1 at
the beginning of the label for a particular combination of
the factors. Congruently, the second thickness is marked
with number 2 in a combination label. The lower layer
thickness provides for finer printed models than the
higher or coarser layer thickness, with the lower rough-
ness notable especially on the beveled edges and faces.
The sample model can be oriented in any possible
direction inside the printer building box. We considered
two main directions: the 1st direction denotes the orienta-
tion of a sample’s largest L dimension towards the build-
ing X axis and the 2nd direction denotes its orientation
towards the Y axis. Thereby, the samples were aligned at
the bottom plane of the building box. In the experiment
label for a particular combination of the factors, the mark
for the sample orientation is in the second place of the
label, expressed with letters X or Y, respectively.
The orientation towards the building Z axis was
omitted since it significantly prolonged the printing time:
for example, to print five samples oriented with their
largest L dimensions towards the building X axis, it takes
only 15 min for 39 printed layers; the same time and the
same number of layers are necessary when the samples
are oriented towards the building Y axis; but when they
are oriented towards the Z axis, the process is prolonged
to 3 hours and 51 min for 1476 layers.
Although all the other orientations are also available
to be combined, additional combinations were omitted in
order to reduce the total number of the experiments.
However, if the results of the experiments had
indicated a significant influence of the orientation on the
tensile strength, additional experiments could have been
easily performed.
T. GALETA et al.: INFLUENCE OF PROCESSING FACTORS ON THE TENSILE STRENGTH OF 3D-PRINTED MODELS
Materiali in tehnologije / Materials and technology 47 (2013) 6, 781–788 783
Figure 5: Test sample
Slika 5: Preizkusni vzorec
Figure 4: Test-sample 3D printing
Slika 4: 3D-tiskanje preizkusnih vzorcev
The type of the infiltrant is marked with the letter at
the end of the label for a particular combination of the
factors. The letters used for a particular infiltrant type
are: W for wax, E for epoxy resin and C for cyanoacry-
late.
For every particular experiment, i.e., particular
combination of factors in each set, six test samples were
printed. During the first tests with the probationary
3D-printed samples, most of the samples tended to slip
out of the tensile-testing machine jaws. Therefore, addi-
tional rubber pads were inserted between the clamping
surfaces in order to ensure a sufficient grip of the
machine jaws holding a test sample.
4 RESULTS
The dimensions that have the biggest influence on the
test results were measured and analyzed prior to the
tests: the test-sample neck width (W1) and the height (H).
The results of measuring the neck height revealed
that most of the samples exceed the limits of tolerance
specified with the standard ((10 ± 0.2) mm). A correc-
tion to the exceeding dimension can be performed with
the afterward processing (e.g., sanding) or this excess
can be prevented prior to 3D printing. The prevention
can be performed during the model preparation in the
printer software with an appropriate anisotropic scale
factor as presented in our published research.23 The
measured values of the neck height or thickness of the
samples, H, are mainly within the limits of tolerance
specified with the standard ((4 ± 0.2) mm). The neck
width and the sample height multiplied together
determine the cross-sectional area A0 of the test samples
(Table 2). The cross-sectional area A0 of the test samples
is used later to calculate the tensile strength.
The last two rows in the table contain the calculated
values: the arithmetic mean and the standard deviation.
The results of the tensile test are expressed with the
values of the breaking force, Fm, and presented in Table
3. In the previous researches it was proved that the
breaking force of 3DP samples is very close or, often,
equal to the maximum testing force1,2; consequently, the
breaking force values are considered to narrow the focus
of the research.
T. GALETA et al.: INFLUENCE OF PROCESSING FACTORS ON THE TENSILE STRENGTH OF 3D-PRINTED MODELS
784 Materiali in tehnologije / Materials and technology 47 (2013) 6, 781–788
Table 1: Combinations of the processing factors and experiment labels
Tabela 1: Kombinacije procesnih dejavnikov in oznaka preizkusa
Layer thickness 0.1 mm 0.0875 mm
Infiltrant Wax Epoxy resin Cyanoacrylate Wax Epoxy resin Cyanoacrylate
Orientation X Y X Y X Y X Y X Y X Y
Experiment label 1XW 1YW 1XE 1YE 1XC 1YC 2XW 2YW 2XE 2YE 2XC 2YC
Table 2: Neck cross-sectional area (A0) of the test samples
Tabela 2: Prerez (A0) preizkusnih vzorcev
Area A0/mm2
Experiment
label 1XW 1YW 1XE 1YE 1XC 1YC 2XW 2YW 2XE 2YE 2XC 2YC
1 41.94 43.60 42.32 43.28 43.42 44.49 43.23 44.13 42.84 43.08 44.27 44.24
2 41.35 43.59 41.69 42.95 44.07 43.52 43.19 43.48 42.70 42.89 43.58 44.95
3 42.28 42.44 41.00 40.59 44.33 42.95 43.73 43.45 43.40 44.61 44.27 44.24
4 41.47 42.50 40.85 41.13 44.51 43.20 43.22 43.58 42.72 44.25 43.58 44.95
5 44.96 42.67 42.42 41.56 43.20 43.66 47.60 44.90 46.46 45.33 49.85 50.80
6 43.93 43.93 42.22 41.71 42.92 43.14 45.63 46.63 45.70 45.39 49.90 50.11
x 42.66 43.12 41.75 41.87 43.74 43.49 44.43 44.36 43.97 44.26 45.91 46.55
S 1.46 0.66 0.69 1.05 0.65 0.55 1.81 1.24 1.67 1.08 3.09 3.05
Table 3: Breaking force
Tabela 3: Sila ob poru{itvi
Force Fm/N
Experiment
label 1XW 1YW 1XE 1YE 1XC 1YC 2XW 2YW 2XE 2YE 2XC 2YC
1 119 104 237 327 174 177 115 137 366 421 163 223
2 134 154 273 321 154 185 116 126 358 299 214 206
3 100 138 278 256 121 160 121 146 338 483 225 156
4 121 120 230 276 161 143 131 151 380 502 233 123
5 93 112 273 279 137 163 126 105 251 303 208 246
6 133 128 290 295 165 167 101 126 295 332 252 192
x 116.67 126.00 263.50 292.33 152.00 165.83 118.33 131.83 331.33 390.00 215.83 191.00
S 16.91 18.15 24.16 27.55 19.62 14.54 10.42 16.63 49.19 90.96 30.14 45.00
The typical breakage outlook presents a fragile frac-
ture without the previously observable plastic deforma-
tion (Figure 6). Moreover, the locations of the breakage
are not concentrated in the middle of the sample-neck
length but accidentally distributed along the neck length.
Such accidental distribution of the breakage location
indicates an internal structural inconsistency and its
significant influence on the mechanical properties of
3DP models.
The maximum tensile stress (Rm), i.e., the tensile
strength is calculated from the results with the following
formula:
R
F
Am
m=
0
(1)
5 ANALYSIS AND DISCUSSION
The maximum particular tensile strength of
11.35 MPa was achieved for the 4th sample from the set
labeled as 2YE, i.e., the sample was printed with a finer
layer thickness of 0.0875 mm, infiltrated with epoxy
resin and oriented towards axis Y. The particular value is
shaded with gray and underlined continuously in Table
4.
The minimum particular tensile strength of 2.07 MPa
was achieved for the 5th sample from the set labeled as
1XW, i.e., the sample was printed with a coarser layer
thickness of 0.1 mm, infiltrated with wax and oriented
towards axis X. This overall lowest value is shaded with
light gray and has a dotted underline in Table 4.
An overview of the sorted average tensile strengths,
presented in Figure 7, reveals the strongest and the
weakest experiment sets. The strongest experiment set is
the one that includes the strongest particular sample,
labeled as 2YE. Although this set also showed the
highest standard deviation referring to the significant
diversity of the strength values, it has the highest average
strength. However, the weakest experiment set was not
the one including the weakest particular sample, labeled
as 1XW. The weakest average strength was revealed by
the set labeled as 2XW, having also the lowest standard
deviation.
The samples infiltrated with epoxy resin obtained the
highest strength in comparison with the other infiltrant
types. Among the epoxy-infiltrated samples, the ones
having a finer layer thickness (2YE and 2XE) showed a
higher strength than those with a coarser thickness (1YE
and 1XE). However, it should be noted that the results
for the finer thickness are significantly more dispersed
than the results for the coarser thickness: the standard
deviation for the 2YE set is 2.08 MPa, for 2XE it is 1.37
T. GALETA et al.: INFLUENCE OF PROCESSING FACTORS ON THE TENSILE STRENGTH OF 3D-PRINTED MODELS
Materiali in tehnologije / Materials and technology 47 (2013) 6, 781–788 785
Table 4: Tensile strength
Tabela 4: Natezna trdnost
Tensile strength Rm/MPa
Experiment
label 1XW 1YW 1XE 1YE 1XC 1YC 2XW 2YW 2XE 2YE 2XC 2YC
1 2.84 2.39 5.60 7.55 4.01 3.98 2.66 3.10 8.54 9.77 3.68 5.04
2 3.24 3.53 6.55 7.47 3.49 4.25 2.69 2.90 8.38 6.97 4.91 4.58
3 2.37 3.25 6.78 6.31 2.73 3.73 2.77 3.36 7.79 10.83 5.08 3.53
4 2.92 2.82 5.63 6.71 3.62 3.31 3.03 3.47 8.90 11.35 5.35 2.74
5 2.07 2.62 6.44 6.71 3.17 3.73 2.65 2.34 5.40 6.68 4.17 4.84
6 3.03 2.91 6.87 7.07 3.84 3.87 2.21 2.70 6.46 7.31 5.05 3.83
x 2.74 2.92 6.31 6.97 3.48 3.81 2.67 2.98 7.58 8.82 4.71 4.09
S 0.44 0.42 0.56 0.49 0.47 0.31 0.26 0.42 1.37 2.08 0.64 0.89
Figure 7: Sorted average tensile strength
Slika 7: Razporeditev povpre~ne natezne trdnosti
Figure 6: Typical breakage outlook
Slika 6: Zna~ilni videz poru{enega vzorca
MPa, while for 1YE it is only 0.49 and for 1XE only
0.56. The orientation of the epoxied samples towards
building axis Y provided a somewhat better strength than
those oriented towards axis X. The samples infiltrated
with cyanoacrylate obtained a medium strength. The
cyanoacrylated samples with a finer layer thickness also
showed a higher strength than those with a coarser thick-
ness, being similar to the epoxied samples. However, a
different orientation of these samples did not cause any
obvious difference in the strength.
The waxed samples acquired the lowest strength. The
orientation of the waxed samples towards building axis Y
gave a slightly better strength than the orientation
towards axis X. The variation in the layer thickness did
not show any discrepancy in the strength.
In order to verify the observed principles and rela-
tionships between the sample strength and the processing
factors, we carried out a factorial analysis of the variance
(ANOVA) and summarized the results in Table 5. The
abbreviations used for the processing factors are: LT –
layer thickness; O – orientation; I – infiltrant. ANOVA
confirms a major influence of the infiltrant type and a
minor influence of the other observed processing factors
on the resulting strength of the 3DP samples. The varian-
ces in the processing-factor combinations (LT*O, LT*I,
O*I, LT*O*I) also indicate a minor influence of the con-
sidered combinations on the tensile strength.
Table 5: ANOVA of the tensile strength and the processing factors
Tabela 5: ANOVA natezne trdnosti in procesni dejavniki
Factors SS DOF MS Var P F(0.95))
LT 10.61 1 10.61 14.38 0.00 4
O 2.23 1 2.23 3.02 0.09 4
I 272.47 2 136.24 184.55 0.00 4
LT*O 0.03 1 0.03 0.04 0.85 4
LT*I 7.37 2 3.68 4.99 0.01 4
O*I 3.68 2 1.84 2.49 0.09 4
LT*O*I 1.85 2 0.93 1.25 0.29 4
Error 44.29 60 0.74
To evaluate the application of alternative infiltrants
on our samples instead of the genuine, manufacturer’s
infiltrants, the results are compared with those obtained
for the genuine infiltrants2 and presented in Figure 8. In
both researches, the samples were printed with the same
model of 3D printer but with different printers. The
tensile tests were performed with the same tensile-testing
machines.
The genuine epoxied samples show a significantly
better average strength when compared to the overall
average strength of the alternative epoxied samples.
However, it should be noted that only five genuine
epoxied samples were tested in the comparative research
where the samples were printed with a finer layer thick-
ness, directed towards building orientation Y and laid
with the narrow side on the bottom of the building box.
Therefore, it should be compared with the test involving
six appropriate alternative epoxied samples labeled as
2YE, with the average strength of 8.82 MPa (Table 4).
The alternative cyanoacrylated samples show a much
better average strength when compared to the overall
average strength of the genuine cyanoacrylated samples.
There were sufficient numbers of samples and combina-
tions of the factors in the two researches for them to be
compared: 40 genuine cyanoacrylated samples distri-
buted over 5 distinctive combinations of the processing
factors2 and 24 alternative cyanoacrylated samples used
with 4 factor combinations. Since the strength of the
cyanoacrylated samples can also be influenced by the
residual moisture24 and the moisture was not measured,
we must express some reservation about the results of
the comparison between the cyanoacrylated samples.
The genuine and alternative waxed samples reveal
almost equal average strengths after a comparison bet-
ween sufficient numbers of samples (genuine: 50, alter-
native: 24) and combinations of factors in both
researches.
Furthermore, if the results are compared with the
previous research published in1, it can be noticed that the
average values of the measured strength correspond to
the published values with an evident improvement for
the epoxied samples and a smaller decrease for the
cyanoacrylated samples than in the previous research.
The improvement for the epoxied samples is most likely
due to an improved combination of the processing fac-
tors used for more samples.
6 CONCLUSIONS
The strength of 3D-printed samples mainly comes
from the infiltrants. The results obtained from the
conducted experiments clearly confirm a major role of
the infiltration type, putting ahead epoxy resin as the
strongest one, followed by cyanoacrylate and wax. If the
maximum strength is required, as it is common for
functional prototypes or molding models, then an addi-
tional increase in the strength can be obtained by sele-
cting the best combination of the other two processing
factors.
T. GALETA et al.: INFLUENCE OF PROCESSING FACTORS ON THE TENSILE STRENGTH OF 3D-PRINTED MODELS
786 Materiali in tehnologije / Materials and technology 47 (2013) 6, 781–788
Figure 8: Comparison of the average tensile strengths for the alter-
native and genuine infiltrants
Slika 8: Primerjava povpre~ne natezne trdnosti med alternativnim in
originalnim vezivnim sredstvom
In the case of the infiltration with epoxy resin, to
obtain the maximum strength, the model should have a
finer layer thickness, i.e., 0.0875 mm. The infiltration
with epoxy resin is the slowest one since additional time
is necessary for the resin to obtain full strength, usually
one day at room temperature or two hours in a heated
oven at 70 °C.24
Also, the most important dimension of the printed
model should be oriented towards building direction Y, if
possible taking into account the size of the model and the
size of the building box. The most probable reason for
this is a coincidence between the orientations of axis Y
and the movement of the printing head – the direction of
the binding-material application, so that the direction of
the binding-material application coincides with the
longitudinal direction of the test-tube orientation along
axis Y. On the other hand, the orientation of test tubes
along axis X coincides with the powdering direction, so
the application of the binding material is performed
transversely to the direction of the orientation of the test
tube, so it is expected that the value of the tensile
strength is lower. The only exceptions are the cases of
2XC and 2YC, where the average tensile strength is
higher in the direction of axis X. However, a higher
standard deviation of set 2YC implies possible impurities
in the powder and post-processing differences.
Although the infiltration with cyanoacrylate provides
a clear white color of a 3D-printed model, different from
a dim yellow one obtained with epoxy or wax, the expe-
riments did not confirm cyanoacrylate as the best choice
when the strength of a model is required. The cyanoacry-
late infiltration is the most expensive solution,25 but it
still remains to be the only solution if two or more
models have to be joined or mated together after 3D
printing, when a large model is divided due to a limi-
tation of the building-box size. A cyanoacrylated model
should be 3D printed with a finer layer thickness of
0.0875 mm and the most important dimension of the
printed model should be oriented towards building direc-
tion X to obtain the highest possible strength.
The infiltration with wax, the cheapest available
solution, is also the weakest solution. However, such a
low strength can still satisfy customer demands, because
many 3D-printed models are produced only for represen-
tative purposes without special demands on the mecha-
nical properties. Neither the orientation nor the layer
thickness has a significant influence on the resulting
strength of a 3D-printed model infiltrated with wax.
The evaluation of the alternative infiltrants, applied
instead of the genuine manufacturer’s infiltrants,
revealed that the obtained levels of the strength are
equivalent to those obtained with genuine infiltrants for
the considered 3DP system. Therefore, the use of alter-
native infiltration solutions has been confirmed with
respect to the tensile strength.
Although it is possible that some of the presented
conclusions are valid for similar machines or even rapid
prototyping techniques, all the conclusions should be
considered only for the selected 3D printer and the
selected materials. New materials and new equipment for
3D printing are developed constantly and may demand
new analyses.
Acknowledgements
The work presented in this paper was financially
supported by the Ministry of Science, Education and
Sports, Republic of Croatia, through the scientific
project No. 152-1521473-3111. Special thanks to
Professor Pero Raos, Professor Milan Kljajin and system
engineer Miroslav Mazurek for supporting the work pre-
sented in this paper; to Vedran Galeta for the assistance
in the experiments; and to Professor @eljka Rosandi} for
proofreading the manuscript.
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788 Materiali in tehnologije / Materials and technology 47 (2013) 6, 781–788
Q. LIU, B. [ARLER: NON-SINGULAR METHOD OF FUNDAMENTAL SOLUTIONS FOR THE DEFORMATION ...
NON-SINGULAR METHOD OF FUNDAMENTAL
SOLUTIONS FOR THE DEFORMATION OF
TWO-DIMENSIONAL ELASTIC BODIES IN CONTACT
NESINGULARNA METODA FUNDAMENTALNIH RE[ITEV ZA
DEFORMACIJO DVO-DIMENZIONALNIH STIKAJO^IH SE
ELASTI^NIH TELES
Qingguo Liu1, Bo`idar [arler1,2,3
1University of Nova Gorica, Nova Gorica, Slovenia
2IMT, Ljubljana, Slovenia
3Center of Excellence BIK, Solkan, Slovenia
qingguo.liu@ung.si, bozidar.sarler@imt.si
Prejem rokopisa – received: 2013-03-11; sprejem za objavo – accepted for publication: 2013-10-10
The development of an effective new numerical method for the simulation of the micromechanics of multi-grain systems in
contact is developed in the present paper. The method is based on the Method of Fundamental Solutions (MFS) for two-dimen-
sional plane strain isotropic elasticity and employs the Kelvin Fundamental Solution (FS). The main drawback of MFS is the
presence of an artificial boundary, outside the physical boundary, for positioning the source points of the FS, which is difficult
or impossible in multi-body problems. In order to remove the singularities of the FS the point sources are replaced by the
distributed sources over circular disks. The values of the distributed sources are calculated in a closed form in the case of the
Dirichlet boundary conditions. In the case of the Neumann boundary conditions the respective values of the derivatives of the FS
are calculated indirectly from the considerations of the solution of simple displacement fields. A problem of two, four and nine
bodies in contact is tackled. The newly developed method is verified based on a comparison with the classic MFS. The nume-
rical method will form a part of the microstructure-deformation model, coupled with the macroscopic thermo-mechanics simu-
lation system for continuous casting, hot rolling and heat treatment.
Keywords: isotropic elasticity, plane strain, Navier’s equation, displacement and traction boundary conditions, non-singular
method of fundamental solutions, Kelvin’s fundamental solution
V ~lanku opisujemo u~inkovito novo numeri~no metodo za simulacijo mikromehanike sistemov z ve~ zrni v stiku. Ta temelji na
metodi fundamentalnih re{itev (MFR) za dvo-dimenzionalne probleme izotropnih elasti~nih ravninskih deformacij in uporablja
Kelvinovo fundamentalno re{itev (FR). Bistvena slabost MFR je prisotnost fiktivnega roba zunaj fizikalnega roba za postavitev
izvirnih to~k FR, kar je te`avno ali nemogo~e pri problemih z ve~ telesi. Z odstranitvijo nesingularnosti FR smo to~kovne izvire
nadomestili s porazdeljenimi izviri po krogih. Vrednosti porazdeljenih izvirov so izra~unane v analiti~ni obliki za Dirichletove
robne pogoje. Pri Neumann-ovih robnih pogojih so vrednosti odvodov FR izra~unane posredno pri upo{tevanju re{itev
preprostih polj premika. Obravnavani so sistemi z dvema, {tirimi in devetimi telesi v stiku. Nova metoda je verificirana na
podlagi primerjave s klasi~no MFR. Uporabljena bo v deformacijskem modelu mikrostukture, povezanim z makroskopskim
termomehanskim sistemom za simulacijo kontinuirnega ulivanja, vro~ega valjanja in toplotne obdelave.
Klju~ne besede: izotropna elasti~nost, ravninska deformacja, Navierove ena~be, robni pogoji premika in vle~enja, nesingularna
metoda fundamentalnih re{itev, Kelvinova fundamentalna re{itev
1 INTRODUCTION
The physical modelling of metallurgical processes1
consists of modelling the relations between the process
parameters and the macroscopic velocity, temperature,
concentration, and stress fields, and the relations bet-
ween these fields and the evolution of the microstructure.
In such multi-scale modelling, the transport phenomena
and solid mechanics on the level of microstructure play
an important role and have to be properly computatio-
nally modelled.2,3 We have, in the recent years, deve-
loped a completely new generation of meshless methods,
based on local collocation with radial basis functions, for
solving these models on different scales. The main
advantage of the method is in their similar structure in 2
and 3D, no need for polygonisation, ease of coding, high
accuracy, robustness and flexibility. The method has
already been developed for modelling very complex
phenomena such as macro-segregation on the macrosco-
pic level4 as well as dendritic growth on the micro-level.5
An extension of the meshless method, based on a
collocation with a fundamental solution (Method of
Fundamental Solutions (MFS)), for the simulation of the
deformation of the multiple grains in an ideal mechanical
contact is presented in the present paper. An extension of
the represented method with anisotropic and plastic
deformation capabilities will be used in our future ther-
mo-mechanical calculations for the microstructure evo-
lution in metallurgical processes.6
The main idea of MFS consists of approximating the
solution of the partial differential equation by a linear
combination of fundamental solutions, defined in source
points. The expansion coefficients are calculated by
collocation or least-squares fit of the boundary condi-
tions. The fundamental solution is usually singular in the
Materiali in tehnologije / Materials and technology 47 (2013) 6, 789–793 789
UDK 519.61/.64:539.3 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)789(2013)
source points and this is the reason why the source points
are located outside the domain in the MFS. Then, the
original problem is reduced to determining the unknown
coefficients of the fundamental solutions and the coordi-
nates of the source points by requiring the approximation
to satisfy the boundary conditions and hence solving a
non-linear problem. If the source points are a priori
fixed, then the coefficients of the MFS approximation are
determined by solving a linear problem. The MFS has
become very popular in recent years because of its sim-
plicity. Clearly, it is applicable when the fundamental
solution of the partial differential operator of the govern-
ing equation (or of the system of governing equations) of
the problem under consideration is known. The MFS has
been widely used7 for the solution of problems in linear
elasticity.
In the traditional MFS, the fictitious boundary, posi-
tioned outside the problem domain, is required to place
the source points. This is very impractical or even impo-
ssible, particularly when solving muti-body problems. In
recent years, various efforts have been made, aiming to
remove this barrier in the MFS, so that the source points
can be placed directly on the real boundary.8–12 In the
present paper, we use a non-singular MFS based on8 to
deal with the 2D multi-body isotropic elasticity pro-
blems. The application of a non-singular method of
fundamental solutions (NMFS) in two-dimensional
isotropic linear elasticity has been originally developed
in.13 We extend these developments to multi-body pro-
blems in the present paper. We respectively use area-
distributed sources covering the source points to replace
the concentrated point sources. This NMFS approach
also does not require information about the neighbouring
points for each source point, thus it is a truly mesh-free
boundary method. The derivatives of the fundamental
solution in the distributed source points are calculated by
adopting the methodology in9 from the Laplace to Kelvin
fundamental solution.
The rest of the paper is structured as follows. The
solution procedure is given for MFS and NMFS. Nume-
ral results with 1, 4 and 9 bodies in contact are given,
followed by conclusions and further research.
2 GOVERNING EQUATIONS OF ELASTICITY
FOR THE MULTI-BODY PROBLEM
We consider a two-dimensional domain  with the
boundary , divided into M sub-domains  = I  II
 ...  M with boundaries  = (I  II  ...  M) –
I-II – ... – I-M – ... – (M-I)-M as shown in Figure 1. Each
of the domains is occupied by an isotropic, ideally elastic
material with different material properties, in general.
Let us introduce a two-dimensional Cartesian coordinate
system with orthonormal base vectors ix and iy, and coor-
dinates px and py of point P with the position vector p =
pxix + pyiy. The solid is governed by Navier’s equations
for plane strain problems, which are the conditions for
equilibrium, expressed with the displacement u. The
following governing equations are valid in the sub-
domain m,m = I, II, ..., M, p  m:
2 1
1 2
1
1 2
2
2
2
2
2
( ) ( ) ( ) ( )−
−
+ +
−
v
v
u
p
u
p v
u
m
m
x
x
x
y m
y∂
∂
∂
∂
∂
∂
p p p
p p
v
v
u
p
u
p v
x y
m
m
y
y
y
x m
∂
∂
∂
∂
∂
∂
=
−
−
+ +
−
0
2 1
1 2
1
1 2
2
2
2
2
( ) ( ) ( )p p 2
0
u
p p
x
x y
( )p
∂ ∂
=
(1)
where vm represents the Poisson ratio in the subdomain
m. The boundary  is divided into two not necessarily
connected parts  = D + T. On the part D the displa-
cement (Dirichlet) boundary conditions are given, and
on the part T the traction (Neumann) boundary condi-
tions are given:
  
 

   I I II II M M
I
u u u u
x y D
( ) ( ) ... ( ) ( )
, ,
p p p p
p
+ + + =
= ∈
t t t t
x y T
    
 
I II II M M( ) ( ) ... ( ) ( )
, ,
p p p p
p
+ + + =
= ∈
(2)
where:


m
m
m
=
∈
∉
⎧
⎨
⎩
1
0
,
,
p
p
(3)
On the interface between different regions, displace-
ment continuity and traction equilibrium conditions have
been assumed:
u u x y
t t
m k 
 
  m k
m k
( ) ( ) , ,
( ) ( )
p p p
p p
− = = ∈ ∩
+ =
0
0
I, II, ..., M
  = ∈ ∩
=
x y
m k
m k, ,
,
p (4)
The strains 	; , 	 = x, y are related to the displace-
ment gradients by:
 

	

	
	

∂
∂
∂
∂
u
p
u
p
+
⎛
⎝
⎜
⎞
⎠
⎟ (5)
Q. LIU, B. [ARLER: NON-SINGULAR METHOD OF FUNDAMENTAL SOLUTIONS FOR THE DEFORMATION ...
790 Materiali in tehnologije / Materials and technology 47 (2013) 6, 789–793
Figure 1: A scheme of the multi-region problem. Each of the sub-
domains can have different elastic properties.
Slika 1: Shema problema z ve~ obmo~ji. Vsako podobmo~je ima
lahko razli~ne elasti~ne lastnosti.
The stress components 	; , 	 = x, y are for the plane-
strain cases related to the strains through Hooke’s law:
       	 	 	m xx yy m( )+ +2 (6)
where μm = Em / 2(1 + vm) is the shear modulus of elasti-
city, Em is a constant, known as the modulus of elasti-
city, or Young’s modulus, m = 2 vmμm / (1 – 2vm) is the
Lamé constant, and 	 is the Kronecker delta:

 	
 		
=
=
∉
⎧
⎨
⎩
1
0
,
,
(7)
3 SOLUTION PROCEDURE
The fields on each of the sub-domains are repre-
sented by collocation with fundamental solutions in the
boundary points. The collocation needs to satisfy the
boundary conditions between different regions and outer
boundaries. In a numerical implementation of MFS and
NMFS, we assume that one boundary collocation point
belongs to only two regions at once. In order to keep the
formulation simple we do not put the discretisation
points on the corners that might belong to three or more
regions at once. Explicit expressions for Kelvin’s funda-
mental solution of elastostatics, used in the collocation,
are given14 in a two-dimensional plane-strain situation by:
U
v
v
r
p s p s
	 	
  	

( , )
( )
( )lg
( )(
p s =
−
− ⎛
⎝
⎜ ⎞
⎠
⎟ +
− −1
8 1
3 4
1
π
	 )
r 2
⎧
⎨
⎩
⎫
⎬
⎭
 	 = x, y (8)
where the material properties depend on the position in
a subdomain. U	 (p,s) represents the displacement in
the  direction at point p due to a unit point force acting
in the 	 direction at point s. r p s p sx x y y= − −( ) +( )
2 2
is the distance between the collocation point p and the
source point s.
It can be shown that the following ux and uy satisfy
the governing Eq. (1):
u U Ux xx
n
N
n n xy
n
N
n n( )
~
( , )
~
( , )p p p p p= +
= =
∑ ∑
1 1
 
u U Uy yx
n
N
n n yy
n
N
n n( )
~
( , )
~
( , )p p p p p= +
= =
∑ ∑
1 1
  (9)
p p∉
=
∑ A Rn
n
N
( , )
1
where n and n represent arbitrary constants and n =
nm, n = nm, N = Nm when p  m. Nm is the number
of p  m, m = I, II, ..., M. A(pn,R) (Figure 2)
represents a circle with radius R, centred around pn · pn
= s represents points on the physical boundary. p and pn
belong to the same sub-domain. The quantity U	 (p,pn)
is singular when p = pn. We use the following de-sin-
gularization technique, proposed by6, for evaluating the
singular values:
~
( , )
( , )
U
U
R
Un
n n
	
	

p p
p p p p
=
≠
1
2π 	
( , )
( , )
p p p p
s
n n
A R
Ad =
⎧
⎨
⎪
⎩⎪
∫ (10)
The tractions can be expressed as:
t T Tx xx
n
N
n n xy
n
N
n n( ) ( , ) ( , )p p p p p= +
= =
∑ ∑
1 1
 
t T Ty yx
n
N
n n yy
n
N
n n( ) ( , ) ( , )p p p p p= +
= =
∑ ∑
1 1
  (11)
where:
T
v
v
U
p
v
v
U
xx n
xx n
x
yx( , )
( ) ( , ) ( ,
p p
p p p p= −
−
+
−
2 1
1 2
2
1 2
 ∂
∂
∂ n
y
nx
xx n
y
yx n
x
p
n
U
p
U
p
)
( , ) ( , )
∂
∂
∂
∂
∂
⎡
⎣
⎢
⎢
⎤
⎦
⎥
⎥
+
+
⎡
⎣
⎢
⎢
 
p p p p ⎤
⎦
⎥
⎥
nny
T
v
v
U
p
v
v
U
xy n
xy n
x
yy( , )
( ) ( , ) ( ,
p p
p p p p
= −
−
+
−
2 1
1 2
2
1 2
 ∂
∂
∂ n
y
nx
xy n
y
yy n
x
p
n
U
p
U
p
)
( , ) ( , )
∂
∂
∂
∂
∂
⎡
⎣
⎢
⎢
⎤
⎦
⎥
⎥
+
+
⎡
⎣
⎢
⎢
 
p p p p ⎤
⎦
⎥
⎥
nny
T
U
p
U
p
nyx n
yx n
x
xx n
y
nx( , )
( , ) ( , )
p p
p p p p= +
⎡
⎣
⎢
⎢
⎤
⎦
⎥
⎥
 
∂
∂
∂
∂
+
+ −
−
+
−
⎡
⎣
⎢
2 1
1 2
2
1 2
 ( ) ( , ) ( , )v
v
U
p
v
v
U
p
yx n
y
xx n
x
∂
∂
∂
∂
p p p p
⎢
⎤
⎦
⎥
⎥
nnx
T
U
p
U
p
nyy n
yy n
x
xy n
y
nx( , )
( , ) ( , )
p p
p p p p
= +
⎡
⎣
⎢
⎢
⎤
⎦
⎥
⎥
 
∂
∂
∂
∂
+
+ −
−
+
−
⎡
⎣
⎢
2 1
1 2
2
1 2
 ( ) ( , ) ( , )v
v
U
p
v
v
U
p
yy n
y
xy n
x
∂
∂
∂
∂
p p p p
⎢
⎤
⎦
⎥
⎥
nny
(12)
t = tm, T	 = T	m, nn = nnm, , 	 = x, y when p  m.
The coefficients n and n are calculated from a sys-
tem of (2NI + 2NII + ... + 2NM) × + (2NI + 2NII + ... +
2NM) algebraic equations, obtained by collocating the
boundary conditions:
Ax = b (13)
Q. LIU, B. [ARLER: NON-SINGULAR METHOD OF FUNDAMENTAL SOLUTIONS FOR THE DEFORMATION ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 789–793 791
Figure 2: Distributed source on a disk A(pn)
Slika 2: Porazdeljeni izviri na disku A(pn)
where A is composed of
~
( , )U n p p and
~
( , )T n p p , x is
composed of n and n, and b is composed of u, t and
0. The explicit form of the elements of the algebraic
equation system (13) can be found in.11
The diagonal terms
~
( , )T l l p p , ,  = x, y, l = 1, ...,
NI + ... + NM, in Eq. (13) are determined indirectly for the
collocation points on T. For this purpose, the method
proposed in7 for potential problems is applied to deter-
mine the diagonal coefficients of Eq. (13). In the
approach, we first assume two simple solutions. The first
simple solution is u p cx x x( )p = + , uy ( )p = 0 everywhere.
The second simple solution is ux ( )p = 0, u p cy y y( )p = +
everywhere. We solve them for the corresponding n
( )1 ,
n
( )1 and n
( )2 , n
( )2 using only the displacement boundary
condition. From these two solutions, we can also know:
∂
∂
∂
∂
∂
∂
∂u
p
u
p
u
p
u
x
x
x
y
y
x
y
( ) ( ) ( ) ( )( ) ( ) ( ) (1 1
1 1
p p p
= = =1
p
p p p
)
( ) ( ) ( )( ) ( )
( )
∂
∂
∂
∂
∂
∂
∂
p
u
p
u
p
u
p
y
y y
y
y
=
= = =
0
2
2
2
2 2
0
∂
∂
u
p
y
y
( ) ( )2
1
p
=
(14)
By substituting n
( )1 , n
( )1 and n
( )2 , n
( )2 and Eq. (9)
into Eq. (14), we can obtain the diagonal terms
~
( , )T l p pl , ,  = x, y, l = 1, ..., NI + ... + NM. The
constant c should be selected in such a way that all the
points in the upper cases do not move the same distance.
So that the denominators in the upper derivations are not
zero.
By knowing all the elements Aij and bk of the system
(13), we can determine all the values of n and n.
Subsequently, we can calculate the displacement for all
the domain points using Eq. (9).
4 NUMERICAL EXAMPLES
We consider a square with side a = 3m centred
around px = 0m, py = 0m for testing the performance of
the method. We distinguish three sub-examples. In the
first one, the whole square is occupied by one material,
with the material properties E = 1N/m2, v = 0.3, in the
second one, the square is split into four parts with the
same material properties as in the first example EI = EII =
... = EIV = 1N/m2 vI = vII = ... = vIV = 0.3, and in the third
one, the square is split into nine parts with the same
Q. LIU, B. [ARLER: NON-SINGULAR METHOD OF FUNDAMENTAL SOLUTIONS FOR THE DEFORMATION ...
792 Materiali in tehnologije / Materials and technology 47 (2013) 6, 789–793
Figure 5: The deformation calculated with MFS and NMFS for a
nine-domain case EI = EII = ... = EIV = 1N/m2, vI = vII = ... = vIV = 0.3
and N = 360 (: collocation points, : source points, ×: MFS solution,
: NMFS solution). The position of the source points in MFS is on
squares around each of the square physical domains.
Slika 5: Deformacija, izra~unana z MFR in NMFR, za primer z
devetimi obmo~ji EI = EII = ... = EIV = 1N/m2, vI = vII = ... = vIV = 0,3
in N = 360 (: kolokacijske to~ke, : izvirne to~ke, ×: MFR-re{itev,
: NMFR-re{itev). Pozicija izvirnih to~k je na kvadratih okoli
vsakega od {tirih kvadratnih domen.
Figure 3: The deformation calculated with MFS and NMFS for a
one-domain case with E = 1N/m2, v = 0.3 and N = 120 (: collocation
points, : source points, ×: MFS solution, : NMFS solution)
Slika 3: Deformacija, izra~unana z MFR in NMFR, za primer z enim
obmo~jem z E = 1N/m2, v = 0,3 in N = 120 (: kolokacijska to~ka, :
izvirna to~ka, ×: MFR-re{itev, : NMFR-re{itev)
Figure 4: The deformation calculated with MFS and NMFS for a
four-domain case EI = EII = ... = EIV = 1N/m2, vI = vII = ... = vIV = 0.3
and N = 240 (: collocation points, : source points, ×: MFS solution,
: NMFS solution). The position of the source points in MFS is on
squares around each of the square physical domains.
Slika 4: Deformacja, izra~unanana z MFR in NMFR, za primer s
{tirimi obmo~ji EI = EII = ... = EIV = 1N/m2, vI = vII = ... = vIV = 0,3 in
N = 240 (: kolokacijska to~ka, : izvirna to~ka, ×: MFR-re{itev, :
NMFR-re{itev). Pozicija izvirnih to~k je na kvadratih okoli vsakega
od {tirih kvadratnih domen.
material properties as in the first example EI = EII = ... =
EIX = 1N/m2 vI = vII = ... = vIX = 0.3. We considered the
solution of the Navier equations in this square subject to
the boundary conditions ux = 0m, uy = −01. m on the points
of the north side with py = 1.5m; ux = 0m, uy = 01. m on
the south side with py = –1.5m and on the east px = 1.5m
and west px = –1.5m sides t /x = 0N m
2 , t /y = 0N m
2 is set.
A plot of the deformation, calculated with the defined
three sub-examples is shown in Figures 3, 4 and 5,
respectively. The following parameters were used R =
d/5, RI = dI/5, ..., RIX = dIX/5, where d, dm, m = I, II, ..., IX
are the smallest distances between two nodes on the
boundary, Rm is the radius of the circle centred the point
pn  m, cx = cy = cxI = cyI = L = cxIX = cyIX = 4. The dist-
ance of the fictitious boundary from the true boundary
for the MFS is set to RM = 5d, RMI = 5dI, ..., RMIX = 5dIX.
Figures 3, 4 and 5 show good agreement between the
solution for a one-domain region and a solution recalcu-
lated with the four and nine regions in ideal mechanical
contact and with the same material properties. The
maximum absolute difference in displacements between
the values in Figures 3 and 4 at the outer boundary are
ux = 0.0417m, uy = 0.0886m, and between Figures 3
and 5 ux = 0.0017m, uy = 0.0012m, respectively.
5 CONCLUSION
A new, non-singular method of fundamental
solutions13 is extended in the present paper to solve
multi-dimensional linear elasticity problems. In this
approach, the singular values of the fundamental solution
are integrated over small circular discs, so that the
coefficients in the system of equations can be evaluated
analytically and consistently, leading to an extremely
simple computer implementation of this method. The
method essentially gives the same results as the classic
MFS. It has the advantage that the artificial boundary is
not present; however, at the expense of solving three
times the systems of algebraic equations in comparison
with only one solution in MFS. The main advantage of
the method is that the discretisation is performed only on
the boundary of the domain and no polygonisation is
needed, like in the finite-element method. The NMFS
method, presented in this paper, can be adapted or
extended to handle many related problems, such as
three-dimensional elasticity, anisotropic elasticity, and
multi-body problems, which all represent directions for
our further investigations. The advantage of not having
to generate the artificial boundary is particularly
welcome in these types of problems. The method will be
used in the future for the calculation of multigrain
deformation problems in steel and aluminium alloys,
with realistic grain shapes, obtained from the microscope
images. The developed method is believed to represent a
most simple state-of-the-art way to numerically cope
with these types of problems.
Acknowledgement
This paper forms a part of the project L2-3651
Simulation and Optimisation of Casting, Rolling and
Heat Treatment Processes for Competitive Production of
Topmost Steels, supported by the Slovenian Research
Agency (ARRS) and [tore Steel company. The Centre of
Excellence for Biosensors, Instrumentation and Process
Control (COBIK) is an operation financed by the
European Union, European Regional Development Fund
and the Republic of Slovenia.
6 REFERENCES
1 B. [arler, R. Vertnik, S. [aleti}, G. Manojlovi}, J. Cesar, Berg-
Huettenmaenn. Monatsh., 150 (2005), 300–306
2 T. Mura, Micromechanics of Defects in Solids, 2ed, Martinus Nijhoff
Publishers, Netherlands 1987
3 M. Braccini, M. Dupeux, Mechanics of Solid Interfaces, Wiley, New
York 2012
4 G. Kosec, M. Zalo`nik, B. [arler, H. Ccombeau, CMC: Computers,
Materials & Continua, 22 (2011), 169–195
5 A. Z. Lorbiecka, B. [arler, CMC: Computers, Materials & Continua,
18 (2010), 69–103
6 U. Hanoglu, S. Islam, B. [arler, Mater. Tehnol., 45 (2011) 6,
545–547
7 Y. S. Smyrlis, Mathematics of Comptation, 78 (2009), 1399–1434
8 Y. J. Liu, Engineering Analysis with Boundary Elements, 34 (2010),
914–919
9 B. [arler, Engineering Analysis with Boundary Elements, 33 (2009),
1374–1382
10 D. L. Young, K. H. Chen, J. T. Chen, J. H. Kao, CMES: Computer
Modeling in Engineering & Sciences, 19 (2007), 197–222
11 D. L. Young, K. H. Chen, C. W. Lee, Journal of Computational
Physics, 209 (2005), 290–321
12 W. Chen, F. Z. Wang, Engineering Analysis with Boundary Ele-
ments, 34 (2010), 530–532
13 Q. G. Liu, B. [arler, CMES: Computer Modeling in Engineering and
Sciences, 91 (2013), 235–266
14 D. E. Beskos, Boundary Element Methods in Mechanics, Elsevier,
Amsterdam 1987, 33
Q. LIU, B. [ARLER: NON-SINGULAR METHOD OF FUNDAMENTAL SOLUTIONS FOR THE DEFORMATION ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 789–793 793

R. ZAPLOTNIK et al.: REMOVAL OF SURFACE IMPURITIES FROM QCM SUBSTRATES ...
REMOVAL OF SURFACE IMPURITIES FROM QCM
SUBSTRATES WITH THE LOW-PRESSURE
OXYGEN-PLASMA TREATMENT
ODSTRANJEVANJE POVR[INSKIH NE^ISTO^ S QCM-PODLAG Z
OBDELAVO V NIZKOTLA^NI KISIKOVI PLAZMI
Rok Zaplotnik1, Darij Kreuh2, Alenka Vesel1
1Jo`ef Stefan Institute, Department of Surface Engineering and Optoelectronics, Jamova cesta 39, 1000 Ljubljana, Slovenia
2Ekliptik, d. o. o., Teslova 30, 1000 Ljubljana, Slovenia
rok.zaplotnik@ijs.si
Prejem rokopisa – received: 2013-03-12; sprejem za objavo – accepted for publication: 2013-03-25
A novel method for cleaning quartz crystals used in QCM is presented. Crystals were covered with a thin film of impurities
containing protein fibrinogen, sodium stearate and sodium chloride. Samples were exposed to the oxygen plasma created by an
electrodeless radiofrequency discharge in the H-mode. The oxygen pressure was fixed to 10 Pa, the nominal power of the
RF-generator was 200 W, the reflected power was about 24 W, and the treatment times were up to 40 s. The plasma was
characterized with optical emission spectroscopy. During the treatment of the samples the spectral features indicating a removal
of impurities appeared in the optical spectra. The temporal evolution of characteristic optical features allowed for a real-time
monitoring of the cleaning procedure. All the spectral features have been normalized to the temporal evolution of the atomic
oxygen line at 615 nm. The CO features appeared in the first few seconds of the plasma treatment indicating a preferential
removal of organic components. The Na atomic line at 589 nm became measurable after a few seconds of the plasma treatment,
increasing for about 10 s and decreasing thereafter until it became very low after about 35 s. The results clearly indicate the
applicability of oxygen plasma for a rapid removal of impurities from the surface of quartz crystals. Organic impurities are
removed in a very short time due to an extremely high affinity of oxidation, while inorganic impurities can be removed only
with a prolonged treatment when the thermal effects become important.
Keywords: protein fibrinogen, oxygen plasma, surface cleaning, etching
V ~lanku predstavljamo novo metodo ~i{~enja QCM-podlag iz kremenovega kristala. Kristale smo prekrili s tanko plastjo
ne~isto~, ki so vsebovale protein fibrinogen, natrijev stearat in natrijev klorid. Vzorce smo izpostavili kisikovi plazmi, ustvarjeni
z brezelektrodno radiofrekven~no razelektritvijo v H-na~inu. Obdelovali smo jih pri konstantnem tlaku 10 Pa, ~asi obdelave so
bili do 40 s, mo~ RF-generatorja je bila 200 W, odbita mo~ pa okoli 24 W. Zna~ilnosti plazme smo merili z opti~no emisijsko
spektroskopijo. Med obdelavo vzorcev so se na opti~nem spektru pojavile emisijske spektralne ~rte, ki jasno nakazujejo
odstranjevanje ne~isto~. ^asovni potek teh ~rt nam je omogo~il spremljanje postopka ~i{~enja v realnem ~asu. Vse emisijske
~rte smo normalizirali s ~asovnim potekom kisikove atomske ~rte pri 615 nm. V prvih nekaj sekundah so se v spektrih pojavile
emisijske ~rte molekule CO, kar nakazuje odstranjevanje organskih komponent. Intenziteta natrijeve emisijske ~rte pri 589 nm
postane merljiva po nekaj sekundah plazemske obdelave, se ve~a pribli`no 10 s, dose`e maksimum in se nato manj{a, dokler ne
dose`e nizkih vrednosti pri 35 s. Rezultati jasno prikazujejo uporabnost kisikove plazme pri hitrem odstranjevanju ne~isto~ s
povr{in kremenovih kristalov. Organske ne~isto~e o~istimo v zelo kratkem ~asu zaradi visoke afinitete oksidacije, medtem ko
anorganske ne~isto~e lahko odstranimo le z dalj{imi obdelavami, ko termi~ni efekti postanejo pomembni.
Klju~ne besede: protein fibrinogen, kisikova plazma, ~i{~enje povr{in, jedkanje
1 INTRODUCTION
Quartz crystal microbalance (QCM) is one of the
most accurate devices for determining thin-film thick-
ness.1–6 The method is based on measuring the resonance
frequency of the oscillating crystals, depending on the
mass of the material deposited on the quartz crystal.7
Quartz crystals are rarely used several times if the
deposited material is not removed properly, making this
technique rather expensive. In the vacuum practice
crystals are often cleaned and reused in order to reduce
costs. Several techniques for removing the surface impu-
rities have been applied and most of them use wet chemi-
cal cleaning. Not only is this technology unfriendly to
the environment, it often does not ensure an atomically
clean surface. One can use different solvents for re-
moving specific impurities such as organic solvents for
removing organic materials and acids or bases for re-
moving inorganic impurities. The latter may destroy the
samples if not used properly. In many practical appli-
cations hundreds of samples should be measured and the
cleaning of crystals may represent a major concern
especially because the atomic cleanliness of a surface is
not always guaranteed. In order to overcome this pro-
blem we have developed an alternative technique for
removing the surface impurities that is based on an
application of highly aggressive oxygen plasma.8
2 EXPERIMENTAL WORK
Commercially available QCM crystals have been
used for the experiments. The crystals were covered with
about a 100-nm thick film containing a mixture of fibri-
nogen, sodium citrate and sodium chloride. A commer-
cially available lyophilized mixture of these materials
Materiali in tehnologije / Materials and technology 47 (2013) 6, 795–797 795
UDK 533.9:669.787 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)795(2013)
(Fibrinogen, Fraction I, type III from human plasma)
was purchased from Sigma-Aldrich, dissolved in MiliQ
water in order to obtain a 0.1 % solution, applied on a
quartz crystal and distributed evenly by spinning it at
500 r/min. The samples were mounted into a vacuum
system presented schematically in Figure 1. The system
is pumped with a two-stage rotary pump with the nomi-
nal pumping speed of 80 m3 h–1. Commercially available
oxygen is leaked into the system on the other side, so the
oxygen pressure during continuous pumping is 10 Pa.
The pressure is measured with a precise pressure gauge
MKS Baratron 722A. Plasma was excited inside a glass
tube with a diameter of 4 cm. A copper coil was con-
nected to an RF-generator via a matching network. The
generator operated at the standard frequency of 13.56
MHz and the nominal power variable up to 1000 W.
During the current experiment the applied power was
200 W and the reflected power was 24 W, so the plasma
was well matched with the generator. Luminous plasma
was concentrated within the coil indicating almost a pure
H-mode of the electrical discharge.
The plasma was characterized with optical emission
spectroscopy. We used a spectrometer AVASPEC 3648
and the integration time was fixed to 50 ms. Such an
integration time, rather unusual for a characterization of
luminous plasma in the H-mode, was applied in order to
let the major peaks oversaturate and detect the minor
peaks and bands.
3 RESULTS
A typical spectrum of gaseous plasma at the absence
of any sample is presented in Figure 2, while a spectrum
of plasma during the treatment of a quartz crystal con-
taminated with a film of impurities is shown in Figure 3.
Both spectra show oversaturated major atomic lines, but
the spectral features corresponding to the etching of
R. ZAPLOTNIK et al.: REMOVAL OF SURFACE IMPURITIES FROM QCM SUBSTRATES ...
796 Materiali in tehnologije / Materials and technology 47 (2013) 6, 795–797
Figure 3: Typical optical spectrum of oxygen plasma at the presence
of a sample
Slika 3: Zna~ilni opti~ni spekter kisikove plazme pri obdelavi vzorca
Figure 2: Typical optical spectrum of oxygen plasma at the absence of
a sample
Slika 2: Zna~ilni opti~ni spekter kisikove plazme brez vzorca
Figure 1: Schematic of the experimental set-up
Slika 1: Shema eksperimentalne postavitve
Figure 4: Temporal evolution of specific spectral features during the
plasma cleaning of the quartz crystal contaminated with a 100-nm
thick film of impurities
Slika 4: ^asovni potek dolo~enih emisijskih spektralnih ~rt med
plazemskim ~i{~enjem kremenovega kristala z naneseno debelo
plastjo ne~isto~ 100 nm
surface impurities are clearly visible in Figure 3.
Selected spectral features during the plasma treatment of
a sample are presented in Figure 4. These spectral
features have been normalized to the temporal evolution
of the atomic oxygen line at 615 nm.
4 DISCUSSION
Optical spectra of oxygen plasma during the treat-
ment of quartz crystals, contaminated with surface
impurities, have been collected continuously for 40 s. A
spectrum without a sample was also measured and it is
presented in Figure 2. As expected, neutral oxygen-atom
atomic lines prevail. The absence of any molecular bands
just indicates a well-known characteristic of plasma,
created by the RF-discharge in the H-mode – an almost
100 % dissociation rate. Such a very high dissociation
rate is due to a very low coefficient for a heterogeneous
surface recombination of O-atoms on the glass surface as
well as the absence of three body collisions at 10 Pa that
could otherwise allow for a gas-phase recombination.
Plasma is therefore almost 100 % dissociated and many
oxygen atoms are probably found in the excited states
that are metastable. Such atoms are chemically extre-
mely reactive and will interact with organic as well as
some inorganic impurities even at room temperature.
Well pronounced spectral features peaking at 309 nm
correspond to OH radicals. This peak is due to the
presence of water vapour in our vacuum system that has
never been baked. Water molecules dissociate to OH and
H radicals in a rather powerful discharge. The spectral
features presented in Figure 3 are different. Not only are
the hydrogen atomic lines corresponding to Balmer
series much higher, but other features are observed as
well. The most significant one is the broad continuum
that has been assigned to various transitions of CO
radicals.9,10 Also, the Na atomic line at 589 nm is clearly
visible. The existence of these two features indicates a
substantial etching of the surface impurities deposited on
quartz crystals.
The time evolution of some spectral features is shown
in Figure 4. Let us first discuss the extremely sharp peak
of the OH radical found just after turning on the dis-
charge. As mentioned earlier, our system was frequently
vented and never baked so that the surfaces are covered
with adsorbed water molecules. As soon as the discharge
is turned on the molecules desorb and quickly dissociate
in gaseous plasma and are excited by the inelastic colli-
sions with free electrons. Since the amount of water
molecules on the surfaces is rather low, practically all the
molecules desorb in a very short time. This is the reason
of the sharp peak observed in Figure 4. For the next few
seconds the intensity of the OH peak slowly decreases
until it becomes constant. Such a decrease cannot be
attributed to the desorption of water molecules from the
discharge tube, as it is probably due to the oxidation of
organic impurities containing hydrogen. This hypothesis
is sound with the behaviour of the CO peak. This peak
appears immediately after turning on the discharge.
Since it is absent in the plasma created in an empty tube,
it can be attributed to the oxidation of organic materials.
The fact that the peak appears only in the first few se-
conds of the plasma treatment indicates a rapid oxidation
of organic impurities. Such a rapid oxidation is probably
due to a rather high density of excited oxygen atoms in
plasma. As mentioned earlier such metastable atoms are
chemically extremely reactive. The sodium peak is also
presented in Figure 4. The temporal evolution of this
peak is completely different from the OH and CO peaks.
The sodium peak becomes measurably high after about 5
s and its intensity is the highest after about 15 s. Oxygen
plasma is obviously aggressive enough to react also with
this sort of impurities.
5 CONCLUSIONS
A novel method for cleaning quartz crystals was pre-
sented. The method is ecologically benign since it does
not produce any wastes. The method allows for a rapid
removal of organic impurities since about a thick film
100 nm of fibrinogen with stearates is effectively
removed in a few seconds. As expected, inorganic
impurities such as sodium chloride are more resistant
and it, therefore, takes a longer time to remove them. It is
worth mentioning that oxygen plasma does not interact
with silicon dioxide so that the properties of quartz
crystals are preserved.
Acknowledgements
The authors acknowledge the financial support from
the Slovenian Research Agency through project L7-4035
(Toward ecologically benign alternative for cleaning of
delicate biomedical instruments).
6 REFERENCES
1 B. D. Vogt, E. K. Lin, W. I. Wu, C. C. White, Journal of Physical
Chemistry B, 108 (2004) 34, 12685–12690
2 F. Höök, B. Kasemo, T. Nylander, C. Fant, K. Sott, H. Elwing,
Analytical Chemistry, 73 (2001) 24, 5796–5804
3 T. Indest, J. Laine, L. S. Johansson, K. Stana-Kleinschek, S. Strnad,
R. Dworczak, V. Ribitsch, Biomacromolecules, 10 (2009) 3,
630–637
4 F. Caruso, K. Niikura, D. N. Furlong, Y. Okahata, Langmuir, 13
(1997) 13, 3422–3426
5 M. M. Ayad, M. A. Shenashin, European Polymer Journal, 39 (2003)
7, 1319–1324
6 T. Indest, J. Laine, K. Stana-Kleinschek, L. Fras-Zemljic, Colloids
Surfaces A, 360 (2010) 1–3, 210–219
7 A. Doliska, A. Vesel, M. Kolar, K. Stana-Kleinschek, M. Mozetic,
Surface and Interface Analysis, 44 (2012) 13, 56–61
8 M. Mozetic, Mater. Tehnol., 44 (2010) 4, 165–171
9 N. Glavan Vukelic, M. Biscan, S. Milosevic, Spectral simulations of
CO molecular emission observed in inductively coupled radio
frequency plasmas, Proceedings of the 3rd International Conference
on Advanced Plasma Technologies, Bohinj, 2010, 38–43
10 A. Vesel, M. Mozetic, A. Drenik, M. Balat-Pichelin, Chemical
Physics, 382 (2011) 1–3, 127–131
R. ZAPLOTNIK et al.: REMOVAL OF SURFACE IMPURITIES FROM QCM SUBSTRATES ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 795–797 797

A. SAXENA et al.: STRUCTURAL, OPTICAL AND ELECTRICAL STUDIES ON Si-DOPED POLYMER ELECTROLYTES
STRUCTURAL, OPTICAL AND ELECTRICAL STUDIES
ON Si-DOPED POLYMER ELECTROLYTES
[TUDIJ STRUKTURE, OPTI^NE IN ELEKTRI^NE LASTNOSTI
ELEKTROLITOV POLIMEROV, DOPIRANIH S Si
Amit Saxena1,2, Pramod Kumar Singh2, Bhaskar Bhattacharya2
1Department of Physics, Hindustan College of Science and Technology, Farah, Mathura- 281122, India
2Material Research Laboratory, School of Engineering and Technology, Sharda University, G. Noida 201310, India
saxena.electronics@gmail.com
Prejem rokopisa – received: 2013-03-25; sprejem za objavo – accepted for publication: 2013-04-02
In this paper we have studied the modification of the structural, optical and electrical properties of a polymer electrolyte film of
polyethylene oxide complexed with sodium iodide (PEO:NaI) by doping with Si particles. Structural studies were carried out
using Fourier-transform infrared spectroscopy (FTIR) and X-ray diffraction (XRD) measurements. The surface morphology of
the polymer electrolyte film doped with Si particles was studied using polarized optical microscopy (POM). For the electrical
properties, the electrical conductivity of the films was measured using complex impedance spectroscopy. The ionic transference
number measurement (tion) was carried out and explained in detail to support the electrical conductivity data.
Keywords: polymer electrolyte, FTIR, XRD, POM, ionic conductivity
V tem ~lanku smo preu~evali spremembe strukture, opti~nih in elektri~nih lastnosti polimernih elektrolitskih plasti polietilen
oksida, kompleksiranega z natrijevim iodidom (PEO:NaI) po dopiranju z delci Si. [tudij strukture je bil opravljen z uporabo
infrarde~e spektroskopije s Fourierjevo transformacijo (FTIR) in z rentgensko difrakcijo (XRD). Morfologija povr{ine tanke
plasti polimernega elektrolita, dopiranega z delci Si, je bila preu~evana s svetlobno mikroskopijo s polarizirano svetlobo (POM).
Od elektri~nih lastnosti je bila izmerjena elektri~na prevodnost plasti s kompleksno impedan~no spektroskopijo. Izvr{eno in
razlo`eno je bilo merjenje {tevilke ionske transference (tion), ki prispeva k podatkom o elektri~ni prevodnosti.
Klju~ne besede: elektrolit iz polimera, FTIR, XRD, POM, ionska prevodnost
1 INTRODUCTION
Ion-conducting solid polymer electrolytes are argu-
ably the most attractive materials today because of their
potential applications in various electrochemical devices,
such as dye-sensitized solar cells, supercapacitors,
rechargeable batteries, sensors, etc.1–7 For electroche-
mical applications, the electrolytes must be ironically
conductive, and at the same time electrode must be
compatible with it. Good quality, mechanically strong,
semi-crystalline films could be obtained, but unfortuna-
tely in most cases the obtainable conductivity is too low
(10–7–10–8 S/cm), while for an efficient device we need at
least three-to-four orders of magnitude higher conduc-
tivity (10–3–10–4 S/cm). To achieve this conductivity
value, different approaches have been proposed in the
literature, like the use of plasticizers, copolymerization,
the formation of blends, changing polymer chain length,
composite formation, etc. Of these, composite formation
is quite attractive and many composites/blends with vari-
ous types of dispersoid have been reported in the
literature.8–10 In this communication, we report an
important observation that if semiconducting Si particles
are used in forming a composite, a significant conduc-
tivity enhancement as well as a structural change can be
obtained.
2 EXPERIMENTAL
Before forming the composite, the weight amount of
polymer (PEO) and salt (NaI) are taken and films are
prepared using a standard solution cast technique. The
cation-to-monomer ratio was kept fixed at !0.065 for all
the samples. The different stoichiometric ratios of the
semiconducting Si powder (different size) were taken as
the dopant to form the composite electrolyte and added
to the ion-conducting polymer electrolyte solution. The
resultant Si-doped polymer electrolyte solutions were
stirred continuously. The final viscous solutions thus
obtained were poured into a polypropylene Petri dish.
The films, so prepared, were subjected to different
characterizations.
To confirm the interaction between the polymer and
the semiconductor, Fourier-transform infrared (FTIR)
spectra were recorded in the attenuated total reflectance
(ATR) mode using a Bruker Tensor 27 spectrometer with
a resolution of 4 cm–1 in the vibrational frequency range
600–3600 cm–1. The X-ray diffraction (XRD, Rigaku
D/MAX-RC 12 kW) patterns of the samples were taken
with 2
 values ranging from 15° to 65° at a scanning rate
of 10°/min. The surface morphologies of the composite
polymer film were studied with an optical microscope
(OM, Leica DM LB) under a cross polarizer.
In order to evaluate the ionic conductivity of the
polymer electrolyte film, impedance spectroscopic tech-
Materiali in tehnologije / Materials and technology 47 (2013) 6, 799–802 799
UDK 678.84 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)799(2013)
niques were used. The conductivities of the polymer
films were evaluated from the bulk resistance by ac
complex impedance analysis over a frequency range of
100 Hz to 1 MHz. The transference numbers (tion) of the
Si-doped polymer electrolyte composite sample are
measured using Wagner’s polarization method.2 In this
method the sample is placed between two electrodes in
such a way that it minimizes the contact resistance. In an
ideal case one of the electrodes is blocking, whereas the
other is non-blocking. In the present case we have used
stainless-steel electrodes. A dc potential is applied across
the sample in such a way that mobile ion species move
towards the non-blocking electrode and a small polari-
zation current flows, and finally the sample becomes
polarized. The initial total current Ii and the final residual
current If are use to evaluate the ionic transference
number:
T
I I
Iion
initial final
initial
=
−
3 RESULT AND DISCUSSION
3.1 FTIR study
The composite natures of the Si-doped polymer
electrolyte and the host polymer electrolyte (PEO:NaI)
were confirmed using the FTIR spectra of different
samples. Figure 1 shows the recorded infrared spectra of
the PEO:NaI (line 1) and PEO:NaI doped with Si
polymer electrolyte films (line 2). Peaks corresponding
to various conformations of the polymer chain could be
identified.11–13 No additional peaks were observed in the
FTIR spectra of the Si-doped polymer electrolyte (line 2)
other than the host polymer electrolyte (line 1).
This affirms the composite nature of the films. Addi-
tionally, the absence of any other peak re-confirms that
no chemical reaction takes place between the host and
the doped semiconductor. The recorded X-ray diffraction
patterns of the Si-doped polymer electrolyte as well as
the host polymer electrolyte (without Si-doping) are
shown in Figure 2. It is well known that the PEO shows
a partially crystalline nature, containing a broad hollow
between 2 = 15° and 20°. The Si-doped polymer elec-
trolyte films seems superimposed on this broad hollow.10
No other peak than the host and the semiconductors
clearly confirms the composite nature, as we observed in
our FTIR study, discussed earlier.
3.2 Surface study using polarized optical microscopy
(POM)
The surface features of polymer electrolyte films
with and without Si are further characterized by pola-
rized optical microscopy (POM). Figure 3 shows the
POM photographs of different samples. It was clear that
the pure PEO film (Figure 3a) shows well the semi-
crystalline nature in which large size spherulites are
tightly interconnected with each other. Adding NaI into
the PEO matrix (Figure 3b), the spherulite size becomes
small, while the amorphous region (black portion) is
A. SAXENA et al.: STRUCTURAL, OPTICAL AND ELECTRICAL STUDIES ON Si-DOPED POLYMER ELECTROLYTES
800 Materiali in tehnologije / Materials and technology 47 (2013) 6, 799–802
Figure 2: XRD pattern of pure PEO:NaI (line 2) and PEO:NaI doped
with Si particles (line 1) polymer electrolyte films
Slika 2: XRD-posnetek ~istega PEO:NaI (linija 2) in PEO:NaI, dopi-
ranega z delci Si (linija 1) polimerne elektrolitske plasti
Figure 1: The FTIR spectra of the Si-doped polymer film (line 2) and
polymer electrolyte film (PEO:NaI) without Si-doping (line 1)
Slika 1: FTIR-spektra s Si dopirane polimerne plasti (linija 2) in
polimerne elektrolitske plasti (PEO:NaI) brez dopiranja s Si (linija 1)
Figure 3: Polarized optical micrographs of: a) pure PEO, b) PEO:NaI
and c) Si-doped PEO:NaI polymer electrolyte films
Slika 3: Opti~ni posnetki v polarizirani svetlobi: a) ~isti PEO, b)
PEO:NaI in c) s Si dopirana PEO:NaI polimerna elektrolitska plast
increased. It was also noticed that the amorphous region
further increases after the Si doping (Figure 3c) and Si is
preferred to obtain the assembled amorphous region.
This shows good agreement with our FTIR and XRD
measurements.
3.3 Electrical studies
The complex impedance spectroscopy was used to
calculate the bulk electrical conductivity of the polymer
electrolyte films.14–18 We used pressure-contacted stain-
less-steel electrodes and connected it with the CH instru-
ments electrochemical workstation (Model CHI604D)
with the frequency range 100 Hz to 1 MHz. The electri-
cal conductivity () was evaluated using the formula:
 =G
l
A
where  is the ionic conductivity, G is the conductance
(G = 1/Rb where Rb is the bulk resistance), l is the
thickness of the film and A is the area of the polymer
sample. The cole-cole plot (complex impedance plot) of
a typical sample of PEO + NaI doped with the mass
fraction 10 % of Si is shown in Figure 4.
The calculated ionic conductivity values are plotted
in Figure 5. Figure 5 shows the conductivity vs. com-
position plot of Si-doped PEO:NaI polymer electrolyte
films calculated from various impedance plots. Since the
PEO shows a very low conductivity at room temperature
(10–9 S cm–1), we have added a small amount of NaI to
improve its conductivity. It is clear from the figure that
the conductivity increases with increasing Si concen-
tration and the conductivity maxima obtained at the mass
fraction of 10 % Si concentration with conductivity value
of 4.4 · 10–6 S cm–1. This enhancement in conductivity is
attributed to the enhancement in the amorphous region
by doping with Si particles, which is clearly affirmed by
our POM data, discussed earlier.
3.4 Ionic transference number measurement
The total ionic transference numbers of the Si-doped
polymer electrolyte samples are calculated using
Wagner’s polarization method. Figure 6 shows the
variation of the current with the time plot for PEO:NaI +
8 % Si composite polymer electrolyte film. It can be
observed that the current initially decreases rapidly as
the time increases. The initial current is due to ionic
behaviour, but after a short time the sample becomes
polarized and the current is decreased due to the lack of
ions. Finally, there are no ions, but a very small number
of electrons are still being present inside the sample,
which are responsible for the residual current. The
transference number calculated in the present system is
0.99, which shows that the system is mostly due to ionic
behaviour and is a good candidate for a solid electrolyte.
A. SAXENA et al.: STRUCTURAL, OPTICAL AND ELECTRICAL STUDIES ON Si-DOPED POLYMER ELECTROLYTES
Materiali in tehnologije / Materials and technology 47 (2013) 6, 799–802 801
Figure 6: Time vs. current plot for PEO:NaI + 8 % Si composite poly-
mer electrolyte
Slika 6: ^as v odvisnosti od toka za kompozitni polimerni elektrolit
PEO:NaI z masnim dele`em 8 % Si
Figure 4: Cole-cole plot of PEO:NaI + 10 % Si polymer electrolyte
system
Slika 4: Cole-cole diagram PEO:NaI + 10 % Si sistema elektrolitske-
ga polimera
Figure 5: Ionic conductivity () vs. concentration of NaI plot in
Si-doped PEO:NaI polymer electrolyte system
Slika 5: Ionska prevodnost () v odvisnosti od koncentracije NaI v
sistemu s Si dopiranega PEO:NaI elektrolitskega polimera
4 CONCLUSIONS
A composite polymer electrolyte based on Si-doped
PEO:NaI has been developed and characterized. Elec-
trical conductivity measurements show the enhancement
in the electrical conductivity by Si doping. IR as well as
XRD affirmed the composite nature of the polymer elec-
trolyte samples. Polarized optical microscopy (POM)
confirmed the enhancement in the amorphous reason by
Si doping, which is a well-known condition for conduc-
tivity enhancement. Ionic transference number measure-
ments (Tion) suggest that the developed composite poly-
mer electrolytes are ionic in nature.
Acknowledgement
This work was supported by DST project (SR/S2/
CMP-0065/2010) government of India.
5 REFERENCES
1 J. R. MacCallum, C. A. Vincent (eds.), Polymer Electrolyte Reviews,
1, 2, Elsevier Science Publishers Ltd., New York 1989, 1991
2 P. V. Wright, Br. Polym. J., 7 (1975), 319
3 D. E. Fenton, J. M. Parker, P. V. Wright, Polymer, 14 (1973), 589
4 S. Chandra, Superionic solids: principles and applications, North
Holland, Amsterdam 1981
5 A. F. Nogueira, J. R. Durrant, M. A. De Paoli, Advanced Materials,
13 (2001), 826
6 A. Chandra, P. K. Singh, S. Chandra, Solid State Ionics, 154–155
(2002), 15
7 P. Kumar, P. K. Singh, B. Bhattacharya, Ionics, 17 (2011), 721–725
8 P. V. Braun, P. Osenar, S. I. Stupp, Nature, 380 (1996), 327
9 R. Singh, N. A. Jadhav, S. Majumder, B. Bhattacharya, P. K. Singh,
Carbohydrate Polymers, 91 (2013), 682
10 B. Bhattacharya, J. Y. Leem, J. Geng, H. T. Jung, J. K. Park, Lang-
muir, 25 (2009), 3276
11 B. L. Papke, M. A. Ratner, D. F. J. Shriver, Electrochem. Soc., 129
(1982), 1434
12 J. Y. Lee, B. Bhattacharya, D. W. Kim, J. K. Park, J. Phys. Chem. C,
112 (2008), 12576
13 G. Socrates, Infrared and Raman Characteristic Group Frequencies,
3rd ed., John Wiley & Sons, Ltd., New York 2005
14 P. K. Singh, B. Bhattacharya, R. K. Nagarale, Journal of Applied
Polymer Science, 118 (2010), 2976
15 P. K. Varshney, S. Gupta, Ionics, 17 (2011), 479
16 A. Chandra, S. Chandra, Phys Rev B, 49 (1994), 630
17 P. K. Singh, K. W. Kim, K. I. Kim, N. G. Park, H. W. Rhee, Journal
of Nanoscience and Nanotechnology, 8 (2008), 5271
18 P. K. Singh, B. Bhattacharya, Optoelectronics and Advanced Mate-
rials-Rapid communications, 7 (2013), 157
A. SAXENA et al.: STRUCTURAL, OPTICAL AND ELECTRICAL STUDIES ON Si-DOPED POLYMER ELECTROLYTES
802 Materiali in tehnologije / Materials and technology 47 (2013) 6, 799–802
A. I. AYDENIZ et al.: ELECTROLESS Ni-B-W COATINGS FOR IMPROVING HARDNESS, WEAR ...
ELECTROLESS Ni-B-W COATINGS FOR IMPROVING
HARDNESS, WEAR AND CORROSION RESISTANCE
NANA[ANJE Ni-B-W S CEMENTACIJSKIM GALVANIZIRANJEM
ZA IZBOLJ[ANJE TRDOTE, OBRABE IN ODPORNOSTI PROTI
KOROZIJI
Ali Imre Aydeniz1, Ali Göksenli1, Gökce Dil1, Faiz Muhaffel2, Cagdas Calli2,
Behiye Yüksel3
1TU Istanbul, Mechanical Engineering Department, 34437 Istanbul, Turkey
2TU Istanbul, Mechanical Metallurgical and Materials Eng. Dep., 34469 Maslak, Istanbul, Turkey
3Istanbul Aydin University, Dep. of Mechanical Engineering, 34668 Istanbul, Turkey
aydenizal@itu.edu.tr
Prejem rokopisa – received: 2013-03-26; sprejem za objavo – accepted for publication: 2013-04-08
In this study the formation of a Ni-B-W coating on steel using an electroless plating process and evaluation of the hardness,
wear and corrosion resistance was analyzed. The Ni-B-W coating was prepared using an alkaline borohydride-reduced electro-
less nickel bath. Scanning electron microscopy (SEM) of the cross-sectional view of the Ni-B-W coating was analyzed and the
layer characteristics were investigated. The Ni-B-W coating was characterized using XRD. The study reveals that the Ni-B-W
coating is amorphous in the as-plated condition and after heat treatment at 450 °C for 1 h the Ni-B-W coatings crystallize and
produce nickel and nickel borides. The hardness of the as-plated and heat-treated Ni-B-W coatings were compared with Ni-B
coatings. In both coatings the hardness increased with heat treatment. The wear resistance of Ni-B-W coatings annealed at
different temperatures was analyzed using a ball-on-flat test. The wear resistance of the coating increased with heat treatment
and reached a maximum value at 400 °C. An optical microscope and EDS analysis were applied on the wear tracks to determine
the wear mechanism. Ductile failure and an abrasive wear mechanism were dominant for the coating heat treated up to 400 °C.
Network cracks and shearing marks were detected for the coatings annealed at 450 °C and above, indicating brittle failure. For
investigating the corrosion characteristics of the coatings, immersion tests in 5 % H2SO4 and 10 % HCl acidic solutions for
seven days were applied. According to the corrosion test results, the Ni-B-W coatings showed good corrosion resistance in the
H2SO4 solutions.
Keywords: electroless plating, hardness, wear resistance, corrosion resistance, Ni-B-W coating
V tej {tudiji je bil analiziran postopek cementacijskega galvaniziranja za nana{anje Ni-B-W na jeklo in analizirana je bila trdota,
obraba in korozijska odpornost. Nanos Ni-B-W je bil pripravljen z uporabo alkalne z borhidridom reducirane cementacijske
galvanske kopeli niklja. Z vrsti~nim elektronskim mikroskopom (SEM) je bil analiziran pre~ni prerez nanosa Ni-B-W in
preiskane so bile zna~ilnosti plasti. Nanos Ni-B-W je bil karakteriziran z XRD. [tudija je odkrila, da je nanos Ni-B-W amorfen
v nanesenem stanju, po toplotni obdelavi pri 450 °C, 1 h je bil nanos Ni-B-W kristaliziran in nastali so nikelj in nikljevi boridi.
Primerjani sta bili trdota nanesenega in toplotno obdelanega nanosa Ni-B-W. V obeh primerih je trdota narastla s toplotno
obdelavo. S preizkusom s kroglico na ploskvi je bila preizku{ena obrabna odpornost nanosa Ni-B-W, `arjenega pri razli~nih
temperaturah. Odpornost proti obrabi se je pove~evala s toplotno obdelavo in je dosegla najvi{jo vrednost pri 400 °C. Svetlobni
mikroskop in EDS-analiza sta bila uporabljena na obrabljeni te~ini, da bi ugotovili mehanizem obrabe. @ilav prelom in meha-
nizem abrazijske obrabe sta prevladovala pri nanosu, segretem do 400 °C. Pri nanosu, segretem na 450 °C in vi{je, se je pojavila
mre`a razpok in znaki stri`enja, kar ka`e na krhek prelom. Za ugotavljanje korozijskih lastnosti so bili uporabljeni preizkusi s
potapljanjem sedem dni v raztopino 5 % H2SO4 in 10 % HCl. Pri korozijskem preizkusu je nanos Ni-B-W pokazal dobro
korozijsko odpornost v raztopini H2SO4.
Klju~ne besede: cementacijsko galvaniziranje, trdota, odpornost proti obrabi, odpornost proti koroziji, Ni-B-W nanos
1 INTRODUCTION
Electroless deposition is a chemical deposition
technology, involving the deposition of metals from a
solution onto surfaces without applying an external
electric voltage.1 The coating forms a uniform layer,
independent of the substrate geometry, with good
adhesion.2 As reducing agents for the electroless nickel,
hypophosphite and borohydrite have been used.3 The
main advantages of Ni-P coatings are their low cost, ease
of control and high corrosion resistance.4 The Ni-B
coatings have exhibited interesting properties, like
improved hardness and wear resistance.5 The as-plated
Ni-B coatings’ hardness can be further increased by
applying a heat treatment up to 1200 VHN, which could
make it an alternative for chromium coatings.6 Also, the
columnar structure of Ni-B coatings is useful in adhesive
wear due to the good lubrication characteristics.7 The
major limitations of Ni-B coatings are their high costs,
poor control of the coating conditions and a low corro-
sion resistance compared to Ni-P coatings.8 To improve
the properties of Ni-B coatings, a second metal salt into
nickel solutions, such as W, Mo or Co, is added.9–11
Among these metals, tungsten is an attractive metal due
to its excellent, electrical, wear and corrosion resistance.
Up to now, only a few analyses were carried out on the
preparation and tribological–corrosion analyses of
Ni-B-W coatings.9,12
Materiali in tehnologije / Materials and technology 47 (2013) 6, 803–806 803
UDK 621.793:620.178.1 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)803(2013)
2 EXPERIMENTAL PROCEDURES
Plain carbon steel samples with dimensions of 10
mm × 10 mm × 30 mm were used as the substrate
material. The plain carbon steel substrates were grinded,
polished and then soaked in trichloroethylene. After
which they were cleaned with a detergent at 70 °C.
Lastly, the samples were placed in a 30 % HCl solution
for 2 min, followed by a rinse in distilled water. The
cleaned substrates were subsequently placed in a bath.
For the electroless Ni-B-W bath, sodium borohydride
was preferred. The stabilizer used was lead tungstate. To
prevent the precipitation of nickel hydroxide, complex-
ing agents such as ethlenediamine are required in the
bath. The bath composition and the operating conditions
used for preparing the Ni–B-W coatings are given in
Table 1.
Table 1: Composition of the bath and the experimental parameters
used for electroless Ni-B-W coatings
Tabela 1: Sestava kopeli in eksperimentalni parametri, uporabljeni pri
cementacijskem galvaniziranju Ni-B-W nanosa
Operation Plating bath composition Condition
Electroless
Ni-B-W
NiCl2 24 g/L
EDA 60 mL/L
KOH 26.5 g/L
NaBH4 120 g/L
NaOH 263 g/L
PbWO4 2.6 g/L
EDTA 13 g/L
Na2WO4 2H2O 40 g/L
pH = 13.5 ± 0.2
T = (88 ± 2) °C
t = 2 h
A 2.6 mL reducing solution and a 2.6 mL stabilizing
solution were added to the heated bath every 30 min
during the coating procedure. A scanning electron
microscope (SEM) was used to obtain the surface
morphology and the cross-section of the Ni-B-W layer.
A microhardness tester with a Vickers indenter load of
100 g and an indentation time of 20 s was used. The dry
sliding wear performances of the selected coatings were
evaluated in atmospheric conditions using a ball-on-flat
wear tester. The wear tests were carried out by applying
a normal load of 5 N with a diameter 6 mm aluminum
oxide (Al2O3) ball. The stroke and sliding speed of the
balls on the coatings were 5 mm and 10 mm/s. To
evaluate the wear mechanisms, wear tracks were studied
using light microscopy, SEM and energy-dispersive
spectrometry (EDS). To evaluate the corrosion resistance
and the possible passivation behavior of the samples,
immersion tests in different media were carried out.
3 RESULTS AND DISCUSSION
The chemical composition of the coating was anal-
ysed using EDS. The amount of tungsten in the coating
was in mass fraction 1.1 %. The XRD patterns of the
Ni-B-W coatings on a steel substrate in the as-plated and
heat-treated (300 °C and 450 °C) conditions are analyzed
in Figure 1. For the as-deposited condition, the coating
was an amorphous phase. At 300 °C, crystallization
started to occur. After the examination, the existence of
Ni3B peaks after the heat treatment at 450 °C proved the
crystallization of the Ni-B-W layer.
The surface morphology of the Ni-B-W coating
(Figure 2a) after heat treatment shows the formation of
the developed crystallites.
A. I. AYDENIZ et al.: ELECTROLESS Ni-B-W COATINGS FOR IMPROVING HARDNESS, WEAR ...
804 Materiali in tehnologije / Materials and technology 47 (2013) 6, 803–806
Figure 2: a) Surface morphology of heat-treated Ni-B-W coating and
b) the cross-section morphology of the Ni-B-W coating detected by
SEM
Slika 2: a) Morfologija povr{ine toplotno obdelanega nanosa Ni-B-W
in b) morfologija prereza nanosa Ni-B-W (SEM)
Figure 1: XRD pattern of the Ni-B-W coating
Slika 1: XRD-posnetek nanosa Ni-B-W
The spherical nodules, which can be seen in Figure
2a, are formed because of the precipitation of crystalline
nickel and nickel borides (Ni3B) in the coatings. From
the cross-section morphology of the as-plated Ni-B-W
coatings detected by SEM it is clear that the coating is
uniform and is closely connected to the substrate,
indicating good adhesion to the substrate. The thickness
of the coating is 28 μm.
The microhardness of the Ni-B and Ni-B-W coatings
after different heat-treatment temperatures are given in
Table 2. The microhardness of the coatings increased
with an increase of the heat-treatment temperature. This
is due to the formation of micro-grains and a hard nickel
boride phase (Ni3B) in the Ni-B-W coatings, which was
confirmed by the XRD measurements. The Ni-B coating
reached its maximum hardness after the heat treatment at
350 °C, with 1131 VHN, and the Ni-B-W coating at 400
°C, with 1278 VHN.
Table 2: Micro-Vickers hardness of Ni-B and Ni-B-W coatings after
different heat-treatment temperatures
Tabela 2: Trdota po mikrovikersu nanosa Ni-B in Ni-B-W po toplotni
obdelavi pri razli~nih temperaturah
Temperature (°C) Ni-B Ni-B-W
300 979 995
350 1131 1147
400 1109 1278
450 987 1247
500 599 1116
The evaluation of the friction coefficients and the
wear resistance of the steel, as-plated and heat-treated
Ni-B-W coatings, are presented in Table 3. It is clear
that applying the heat treatment increases slightly the
friction coefficient. The results of the wear tests on
different electroless coatings were analyzed using a
profilometer, by measuring the wear track area after 100
m of sliding distance (Table 3).
Table 3: Coefficient of friction, wear track area of as-plated and
heat-treated Ni-B-W coatings
Tabela 3: Koeficient trenja, velikost podro~ja s sledovi obrabe pri
nanesenem in toplotno obdelanem nanosu Ni-B-W
Annealing
temp.
Fric.
coeff.
Wear track area
(μm2)
Rel. wear
resist.
Steel 0.48 859 1.00
As-plated 0.53 231.2 3.71
350 °C 0.62 134.3 6.39
400 °C 0.54 90.1 9.53
450 °C 0.54 399.5 2.15
550 °C 0.58 670.8 1.28
The specific wear rates of all the coatings indicate a
better wear resistance than the steel. Although the
as-deposited coatings show the lowest value of friction
coefficient, the wear rate of the as-deposited sample is
higher than the steel. The best wear resistance was
achieved for the samples annealed at 400 °C. This is
because of the maximum hardness of the coatings
achieved by heating at this temperature. When heated
beyond this temperature, the specific wear resistance
decreased due to the softening of the coatings by grain
coarsening. The EDS analyses of the wear tracks on the
coated samples show that the wear-track depth does not
reach the substrate.
To understand the wear mechanism of the different
coatings, wear-track patterns produced by the Al2O3 ball
were studied (Figure 3).
The wear tracks of the as-plated and the coating
annealed at 400 °C show a similarity. In both coatings,
grooves along the sliding direction are observed. Also,
high plasticity is detected, indicating micro-cutting,
which is a typical ductile failure mechanism. No pits are
detected. Therefore, it can be concluded that an abrasive
wear mechanism is dominant. A network of cracks and
shearing marks are observed for the coatings annealed at
450 °C (Figure 3b) and 550 °C. For a detailed investi-
gation of the wear mechanism, an EDS analysis was
carried out on the wear tracks to detect aluminum, which
would originate from the Al2O3 friction ball, indicating
the presence of adhesive wear. The EDS analysis indi-
cates that in the case of the as-plated and 400 °C
A. I. AYDENIZ et al.: ELECTROLESS Ni-B-W COATINGS FOR IMPROVING HARDNESS, WEAR ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 803–806 805
Figure 3: Wear tracks after the wear tests on the Ni-B-W coatings
annealed at: a) 400 °C, b) 550 °C
Slika 3: Sledovi obrabe po preizkusu obrabe pri nanosu Ni-B-W po
`arjenju na: a) 450 °C in b) 550 °C
annealed coatings, the layers are dominated by nickel
and contain aluminum of only 0.8 % in terms of weight.
In the annealed coatings at 450 °C and 550 °C, the
amount of aluminum increased slightly (to 1.7 % Al).
This indicates that with the annealed coatings at 450 °C
and above, the adhesive wear mechanism is also active in
addition to the abrasive wear mechanism. Another
important factor is the effect of tungsten on the wear
mechanism. As mentioned before, by annealing the
coating beyond 400 °C, a phase transformation from an
amorphous to a crystalline structure is detected. With the
formation of the crystalline structure, grains start to
occur, resulting in the formation of grain boundaries.
These boundaries act as diffusion channels for the
tungsten to move from the inside to the coating surface,
causing the formation of a dense and brittle tungsten
oxide film on the surface. This brittle tungsten oxide film
is also an additional reason for the formation of a
network of cracks after the wear tests for the coatings
annealed at 450 °C and above.
For analyzing the corrosion resistance of the coat-
ings, immersion tests were carried out. Steel, Ni-B and
Ni-B-W coatings were immersed in 10 % HCl and 5 %
H2SO4 acid solutions for 7 d and the weight loss was
measured each day (Table 4).
Table 4: Weight loss of steel and as-plated Ni-B and Ni-B-W coatings
in different acids, mg
Tabela 4: Zmanj{anje mase jekla, nanesenega Ni-B in nanosa Ni-B-W
v razli~nih kislinah, mg
Day
Steel Ni-B Ni-B-W
HCl H2SO4 HCl H2SO4 HCl H2SO4
1 11.1 390.6 0.1 12.9 17.8 15.4
2 14.2 791.4 1.2 76.7 19.7 21.9
3 18.1 1201.5 4.6 209.1 36.0 30.6
4 23.7 1564.3 9.6 472.4 56.9 58.3
5 26.6 1890.5 13.8 729.8 92.7 102.4
6 35.7 2219.5 14.7 1049.6 145.5 112.5
7 40.8 2357.5 26.5 1460.0 219.1 241.4
The corrosion rate of the as-plated Ni-B coatings in
both the HCl and H2SO4 solutions is lower compared to
the steel. The corrosion rate of the Ni-B-W coating in
HCl solutions is the highest; however, the corrosion rate
in the H2SO4 solution is the lowest. Therefore, it can be
concluded that Ni-B-W coatings can be used in a H2SO4
environment.
4 CONCLUSION
Ni-B and Ni-B-W coatings were prepared on steel by
electroless deposition using dual baths. The Ni–B-W
coatings are amorphous in their as-plated condition, but
after a heat treatment at 450 °C for 1 h, the Ni–B-W
coatings crystallize and produce nickel and nickel
borides. SEM of the cross-section view of the coatings
show that the coatings are uniform and the compatibility
is good. The hardness and specific wear resistance of the
Ni–B-W coating increased with the heat treatment and
reached maximum values at 400 °C. Above 400 °C, the
hardness and wear resistance decreased due to grain
coarsening. After analyses of the wear profiles it was
concluded that ductile failure occurred in the coatings up
to 400 °C and that abrasive wear is the dominant wear
mechanism. With coatings at 450 °C and above, surface
cracks are determined and the adhesive wear mechanism
is also active in addition to the abrasive wear mecha-
nism. The formation of a tungsten oxide layer due to the
crystallization was also a reason for the formation of
network cracks. The results of the immersion test in acid
solutions show that a Ni-B-W coating could provide
electrochemical protection in a H2SO4 environment. It
can be concluded that electroless Ni-B-W coatings will
be a useful replacement for Ni–B coatings.
5 REFERENCES
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1412–1418
2 A. Contreras, C. Leon, O. Jimenez, E. Sosa, R. Perez, Applied
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chemistry, 28 (1998), 559–563
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Coating Tech., 190 (2005), 115–121
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M. Polukarov, Protection of metals, 41 (2005), 55–62
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(2011), 1–6
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Tech., 179 (2004), 179–185
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Materials Res., 651 (2013), 263–268
A. I. AYDENIZ et al.: ELECTROLESS Ni-B-W COATINGS FOR IMPROVING HARDNESS, WEAR ...
806 Materiali in tehnologije / Materials and technology 47 (2013) 6, 803–806
K. MICHALEK et al.: PHYSICAL AND NUMERICAL MODELLING OF A NON-STATIONARY STEEL FLOW ...
PHYSICAL AND NUMERICAL MODELLING OF
A NON-STATIONARY STEEL FLOW THROUGH
A SUBENTRY SHROUD WITH AN INNER METERING
NOZZLE
FIZIKALNO IN NUMERI^NO MODELIRANJE
NESTACIONARNEGA TOKA JEKLA SKOZI POTOPLJEN IZLIVEK
Z NOTRANJO [OBO
Karel Michalek1, Markéta Tkadle~ková1, Karel Gryc1, Jiøí Cupek2,
Michal Macura3
1V[B-Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, Department of Metallurgy and Foundry and Regional
Materials Science and Technology Centre, Ostrava, Czech Republic
2Tøinecké @elezárny, Inc., Tøinec, Czech Republic
3Vesuvius ^eská republika, a.s., Tøinec, Czech Republic
karel.michalek@vsb.cz
Prejem rokopisa – received: 2013-04-02; sprejem za objavo – accepted for publication: 2013-05-23
The paper presents new knowledge about physical and numerical modelling of a non-stationary steel flow into a mould through
a subentry shroud with an inner, pressed metering nozzle. The physical and numerical modelling was realized under the
conditions of the Department of Metallurgy and Foundry at VSB-Technical University of Ostrava. A special type of the subentry
shroud is used during continuous casting of steel in Tøinecké `elezárny, a.s. During continuous casting of steel, two
unfavourable phenomena were observed. In the first case, it was not possible to increase the casting speed, though the diameter
of the metering nozzle was extended. In the second case, a fluctuation of the casting speed among individual casting strands was
detected. These two problems did not allow an improvement of the performance of the casting machine. Therefore, the physical
and numerical modelling was performed. Attention was focused on the verification of the effect of the inner diameter of the
nozzle body and internal diameter of the metering nozzle on the resulting volume flow rates. Four diameters of the metering
nozzle – (16; 17; 17.5; 18) mm – were tested. The physical modelling was done on a 1 : 1 model constructed from Plexiglas.
The numerical modelling was realized in the ANSYS FLUENT software. On the basis of the results of the modelling study and
in cooperation with the VESUVIUS company, a new type of the profile of the subentry shroud with a metering nozzle was
designed. The first experimental results in the steel plant led to an increase in the productivity of the continuous-casting
machine.
Keywords: steel, continuous-casting machine, subentry shroud, metering nozzle, physical and numerical modelling
^lanek predstavlja novo znanje o fizikalnem in numeri~nem modeliranju nestacionarnega toka jekla v kokilo skozi potopljen
izlivek z notranjo {obo. Fizikalno in numeri~no modeliranje je bilo izvr{eno v Oddelku za metalurgijo in livarstvo na VSB
Tehni{ki univerzi, Ostrava. V Tøinecké `elezárny, a. s., se uporabljajo posebne izvedbe potopljenih izlivkov med kontinuirnim
ulivanjem jekla. Pri tem sta bila opa`ena dva ne`elena pojava. Pri prvem ni bilo mogo~e pove~ati hitrosti litja, ~eprav je bil
pove~an premer notranjega izlivka, pri drugem pa je bilo opa`eno spreminjanje hitrosti litja med posameznimi `ilami. Ta dva
pojava nista omogo~ila izbolj{anja zmogljivosti livne naprave. Zato je bilo izvr{eno fizikalno in numeri~no modeliranje.
Pozornost je bila usmerjena na preverjanje vpliva notranjega premera telesa izlivka in notranjega premera izlivne {obe na
stopnje pretoka. Preizku{eni so bili {tirje premeri (16; 17; 17,5 in 18) mm notranjega izlivka. Fizikalni model je bil izdelan na
modelu 1 : 1, izdelanem iz pleksi stekla. Numeri~ni model je bil izdelan s programsko opremo ANSYS FLUENT. Na osnovi
rezultatov in s sodelovanjem podjetja VESUVIUS je bila konstruirana nova vrsta potopnega izlivka z notranjo izlivno {obo. Prvi
preizkusi v jeklarni so omogo~ili pove~anje produktivnosti naprave za kontinuirno ulivanje.
Klju~ne besede: jeklo, naprava za kontinuirno ulivanje, potopni izlivek, izlivna {oba, fizikalno in numeri~no modeliranje
1 INTRODUCTION
A special type of the subentry shroud with an inner,
pressed metering nozzle is used during continuous
casting of steel on the CCM No.2 in Tøinecké `elezárny,
a.s. (T@, a.s.). A metering nozzle serves here as a
stabilizing element adjusting the volume flow rate of the
steel. However, while using this type of the nozzle two
adverse effects were observed:
1) It was not possible to increase the casting speed when
increasing the diameter of the metering nozzle.
2) Fluctuations of the casting speed in individual casting
streams were detected.
Both of these factors eliminated the possibility of
increasing the performance of the entire casting machine.
Due to a long-term mutual cooperation, the T@, a. s.,
team contacted the research team at the Department of
Metallurgy and Foundry at VSB-TU Ostrava, which has
extensive experience in the optimization of metallurgical
processes using modelling methods.1–4 The sophisticated
methodology of the experiments and a well-selected
system of physical and numerical modelling results,
transformed into the operational experience with verified
use of the similarity theory, provide an opportunity to
address the other fundamental technological problems.5–7
Physical and numerical modelling was used to study the
Materiali in tehnologije / Materials and technology 47 (2013) 6, 807–814 807
UDK 621.74.047:519.61/.64 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)807(2013)
behaviour of a steel flow through an inner metering
nozzle. Its goal was to assess the behaviour of the steel
flow through the nozzle, analyse its causes and propose
solutions in the form of a new geometry of the inner,
pressed metering nozzle of the subentry shroud. The
parameters such as the influences of the inner diameter
of the body of the subentry shroud and the inner
diameter of the metering nozzle on the resulting volume
flow rates at a constant hydrostatic (ferrostatic) height
above the subentry shroud (a constant steel height in the
tundish) were tested. The efforts were aimed at finding
the limiting factor that restricts the increase in the rate of
the steel flow through the subentry shroud even when the
diameter of the inner metering nozzle is increased.
2 ANALYSIS OF THE PROBLEM
Fluctuations of the casting speed during the casting
of steel on the CCM No. 2, equipped with the shrouds
with inner metering nozzles, are perfectly evident from
Figures 1 and 2. Figure 1 shows the behaviour of the
casting speed on individual casting strands (CS) during a
sequence of heats cast using the subentry shrouds
equipped with inner metering nozzles. See Figure 2 for
an example of the behaviour of the casting speed at a
constant steel weight in the tundish. The following was
determined:
• When using the subentry shrouds equipped with
metering nozzles, the variability of the casting speed
is much greater than when regulating it with the
stopper rods only.
• The casting speed (or the steel flow through the
nozzle) is slightly dependent on the weight of the
steel in the tundish (or on the steel height in the
tundish) – Figure 1.
• Despite the nearly constant weight of the steel in the
tundish, the casting speed increased in some cases –
Figure 2.
• Fluctuations in the volume flow rate indicate the
presence of a non-stationary flow in the nozzle.
• An inability to further increase the casting speed
while increasing the metering-nozzle diameter shows
the existence of an undetermined limiting factor
affecting the flow, which limits the volume flow of
the steel through the subentry shroud.
On the basis of the calculation of the volume flow
rate for the casting speed of 2.5 m min–1 and the cast
dimensions of 150 mm × 150 mm, it was found that:
• The rate of the volume flow of the liquid steel
through the nozzle with the inner diameter of 20 mm
is approximately 60 L min–1.
• The calculated Re criterion reaches approximately
89,000.
• The value noted above is specified for circular
cross-sections far beyond the critical value of the
transition from the laminar to the turbulent flow. In
other words, we can expect that, within the nozzle,
the steel flow is of a strong turbulent character,
which, of course, bears the risk of large-scale non-
stationarities of the steel flow, especially if the
diameter of the nozzle is, at some spot/location, made
even narrower (in the area of the inner, pressed
metering nozzle).
• It is likely that the hydraulic resistance of the
subentry shroud below the metering nozzle was so
high that an increase in the metering-nozzle diameter
could no longer allow a further increase in the steel
flow rate (the casting speed).
To confirm these partial, theoretical conclusions, the
researchers decided to assess the nature of the steel flow
using physical and numerical modelling.
3 PHYSICAL MODELLING
Physical modelling was performed with a nozzle
constructed from Plexiglas on the geometric scale of 1 : 1
(Figure 3a). The following diameters of the metering
K. MICHALEK et al.: PHYSICAL AND NUMERICAL MODELLING OF A NON-STATIONARY STEEL FLOW ...
808 Materiali in tehnologije / Materials and technology 47 (2013) 6, 807–814
Figure 2: Example of the behaviour of the casting speed on casting
strand No.1 at a constant weight of the steel in the tundish
Slika 2: Primer vedenja hitrosti litja na livni napravi {t. 1 pri konstant-
ni masi jekla v vmesni ponovci
Figure 1: Example of the behaviour of the casting speed on individual
casting strands during a sequence of heats cast using subentry nozzles
equipped with metering nozzles
Slika 1: Primer vedenja hitrosti ulivanja na posamezni livni `ili med
sekven~nim ulivanjem talin z uporabo potopljenega izlivka, opremlje-
nega z notranjo izlivno {obo
nozzles were tested: (16; 17; 17.5 and 18) mm, having
the outlets (the internal diameters of the nozzles below
the metering nozzles) of 20 mm and 24 mm, hereafter
referred to as nozzle A and nozzle B. The submersion of
a nozzle below the surface of the steel in the mould was
in all the cases identical – 120 mm. This value
corresponds to the central part of the bottom zirconium
protection.
Given that when applying the geometric scale of 1 : 1,
both identity criteria, Re and Fr, are maintained, we can
assume there is a significant agreement between the
phenomena occurring in the flow in the actual nozzle and
in its model in terms of the character of the flow, swirl
and stationarity (stability). It is evident that the thermal
processes that may affect the steel flow cannot be
detected with this type of modelling. In particular, it is
the influence of the temperature on the steel viscosity,
erosion and corrosion of the material of a nozzle (that is
contingent upon the flow rate of steel and its tempera-
ture) and the associated change in the internal profile of
the nozzle, which may be essential for the stability of the
flow and its fluctuations.
4 NUMERICAL MODELLING
Physical modelling was also complemented with
numerical modelling of the steel-flow behaviour in the
original nozzle with the metering nozzles having the
diameters of (17; 17.5 and 18) mm and with an outlet of
only 20 mm in diameter. Numerical modelling of the
steel flow through a subentry shroud with inner metering
nozzles was performed using the simulation package
ANSYS WORKBENCH, providing a 3D modeller called
DesignModeler, a generator of a computational mesh and
the CFD software FLUENT. The nozzle geometry was
modelled on the 1 : 1 scale. The 3D CAD geometry of
the nozzle, including a detailed view of the area of the
metering nozzle is shown in Figure 3b.
The volume mesh was automatically generated for
the nozzle. The entrance into the nozzle was defined as
the PRESSURE INLET and the exit from the nozzle was
defined as the PRESSURE OUTLET. These boundary
conditions allow a definition of the input parameters
when the input speed is not known exactly, but the
pressure indicated by the height of the steel in the tun-
dish is known. The total pressure defined the pressure-
input condition, which is the sum of the hydrostatic and
dynamic pressure.8
The hydrostatic pressure is defined by the product of
the steel height in the tundish with the steel density and
gravitational acceleration. The dynamic pressure is
defined as the square of the flow speed multiplied by the
density of steel and one-half.8
The pressure outlet was defined in a similar way to
the inlet pressure. However, only the hydrostatic/ferro-
static output pressure, the turbulence intensity and the
hydraulic diameter need to be specified. The pressure
outlet was the same for all simulated variants.
The steel flow was simulated to be incompressible
and turbulent with the turbulence-model k- standard
having a standard wall function. The gravity of 9.81 m s–2
was considered in the z axis. In the simulation, the
operating conditions and material properties listed in
Table 1 were defined.
Table 1: Operating conditions and material properties of the steel at
the temperature of 1843 K
Tabela 1: Delovne razmere in materialne lastnosti jekla pri tempera-
turi 1843 K
Condition Value Unit
Operating pressure 101325 Pa
Operating temperature 1843 K
Operating density 7030 kg m–3
Heat capacity 750 J kg–1 K–1
Heat conductivity 41 W m–1 K–1
Viscosity 0.006995 Pa s
The study focused on the effect of the inner geometry
of the nozzle on the rate of the flow through the nozzle.
Attention was also focused on the character of the
velocity and pressure field inside the nozzle. The results
were also complemented with the distribution of the
shear stress on the walls of the nozzle, allowing a pre-
diction of the risk of erosion of the nozzle material.
5 DISCUSSION OF RESULTS
5.1 Effect of the metering-nozzle diameter and nozzle-
outlet diameter on volume flow rates
The effect of the diameter of the metering nozzle and
the diameter of the nozzle outlet on the achieved maxi-
mum volume flow rates at a constant hydrostatic (ferro-
K. MICHALEK et al.: PHYSICAL AND NUMERICAL MODELLING OF A NON-STATIONARY STEEL FLOW ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 807–814 809
Figure 3: a) Physical model, b) 3D CAD model of the metering
nozzle of the subentry shroud
Slika 3: a) Fizikalni model, b) 3D CAD model notranje izlivne {obe
potopljenega izlivka
static) height above the nozzle (the constant steel height
in the tundish) obtained with physical (PM) and nume-
rical modelling (NM) is shown in Figure 4.
The main findings that can be defined from the graph
above are as follows:
• An increase in the metering-nozzle diameter results
in the increases in the volume flow rates for both
types of nozzle, A and B (with the outlet diameters of
20 mm and 24 mm).
• The nozzle with the outlet diameter of 20 mm shows
an increase in the volume flow rate of only 1.4 L min–1
(or 1.3 L min–1 according to the results of the nume-
rical simulation) when switching from the metering-
nozzle diameter 17 mm to 18 mm. This increase
corresponds to a very small increase in the casting
speed, i.e., 0.06 m min–1.
• The nozzle with an outlet diameter of 24 mm shows a
more noticeable increase in the volume flow rate, i.e.,
3.8 L min–1, representing an increase in the casting
speed by 0.16 m min–1.
• A very good agreement of the results of the physical
and numerical modelling for the nozzle outlet of 20
mm is evident.
To sum up the results stated above, we can conclude
that when we use the nozzle with a 20-mm outlet dia-
meter, the flow rate can be increased only very slightly
by simply increasing the diameter of the metering
nozzle. This is due to a great hydraulic resistance of the
nozzle located below the metering nozzle, which pre-
vents a bigger increase in the volume flow rate. An
increase in the flow rate can only be achieved by
increasing the input ferrostatic pressure (the height of the
steel in the tundish), but only to a very limited extent. A
more significant increase in the volume flow rate and,
thus, the casting speed can be expected after increasing
the diameter of the nozzle below the metering nozzle
from 20 mm to 24 mm, when the diameter of the
metering nozzle has already started to perform the actual
role of the flow-rate restrictor.
5.2 Fluctuation of the rate of the volume flow through
the subentry shroud/nozzle
The fluctuation of the flow rate (and, in practice, of
the casting speed) is an interesting phenomenon that may
be associated with the changes in ferrostatic conditions,
especially with a change in the flow character in the
nozzle itself.
A change in ferrostatic conditions may be related to a
change in the steel weight in the tundish, i.e., with a
change in the height of the steel in the tundish. As the
flow rate is not only a function of the height of the steel
in the tundish, but, more accurately, a function of the
distance between the height of steel in the tundish and
the height of steel in the mould, it is also necessary to
take into account the change in the elevation position of
the tundish, which is sometimes made to allow an even
wear of the outer surface of the nozzle. A shift in the
steel-level distances by 10 cm will change the volume
flow rate of the model by 1.33 L min–1 and, in the
operating conditions, it will change the casting speed by
0.055 m min–1.
However, a far more significant factor is the character
of the steel flow rate in the nozzle. Figure 5 shows the
flow character displayed with the velocity vectors in the
cross-section of a nozzle. Figure 5 shows that an
increase in the metering-nozzle diameter (17 mm  17.5
mm  18 mm) leads to a slight decrease in the
maximum velocity in the area of the metering nozzle,
while the area of the medium speed of 3.2 m s–1 slightly
increases below the metering nozzle and in the
nozzle-outlet area. On the other hand, the walls of the
nozzle show a slowdown of the casting strand due to the
presence of the so-called laminar sublayer9, which is a
typical phenomenon in a highly turbulent flow of a liquid
in the pipelines. The velocity of the steel flow is, thus,
not the same in the entire section of the nozzle, as shown
in the detailed views of the distribution of the velocity
vectors at the nozzle outlets.
The velocity distribution also determines the
distribution of the pressure of the flowing steel in a
nozzle. As confirmed with the detailed views of the
pressure profiles at the metering nozzles (Figure 6) an
enlarged diameter of a metering nozzle results in a
reduction of the pressure (or, as noted above, of the
maximum speed). On the other hand, below the metering
nozzles, as shown by the details in Figure 6, the bigger
the metering-nozzle diameter, resulting in a slight
increase in the volume flow rate, the bigger is the area
with the maximum pressure field near the nozzle outlet.
However, if one looks closely, an increase in the area
with the maximum pressure field is very subtle, which
could be related, as stated above, to a large hydraulic
resistance, which also prevents a higher increase in the
volume flow rate.
K. MICHALEK et al.: PHYSICAL AND NUMERICAL MODELLING OF A NON-STATIONARY STEEL FLOW ...
810 Materiali in tehnologije / Materials and technology 47 (2013) 6, 807–814
Figure 4: Effect of the metering-nozzle diameter and nozzle-outlet
diameter on the volume flow through the subentry shroud for a 120
mm submersion in the mould
Slika 4: Vpliv premera notranje izlivne {obe in izstopni premer izlivka
na volumenski tok skozi potopljen izlivek pri potopitvi 120 mm v
kokilo
K. MICHALEK et al.: PHYSICAL AND NUMERICAL MODELLING OF A NON-STATIONARY STEEL FLOW ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 807–814 811
Figure 7: a) Character of the steel flows in the nozzles shown with the velocity vectors (m s–1) at the cross-sections of the nozzles – in the
metering nozzles and b) details just below the metering nozzles
Slika 7: a) Vedenje toka jekla v {obi, prikazano z vektorji hitrosti (m s–1) na prerezu {obe – notranje {obe in b) detajl tik pod izlivno {obo
Figure 6: Detailed views of the pressure profiles (Pa): a) in the metering nozzles and b) in the areas of the nozzle outlets
Slika 6: Podroben prikaz profilov tlaka (Pa): a) v notranji izlivni {obi in b) na podro~ju izstopa iz {obe
Figure 5: Steel-flow characteristics in the nozzles displayed with the velocity vectors (m s–1) in the cross-sections of the nozzles with detailed
views at the nozzle outlets for individual nozzles with the metering-nozzle diameters
Slika 5: Zna~ilnosti toka jekla v {obi, prikazane z vektorji hitrosti (m s–1) na prerezu {obe z detajlnim pogledom pri izstopu iz {obe za posamezne
{obe s premeri notranjih {ob
As explained above, the flow fluctuations may be
caused by a high Re number of the steel flow in the
nozzle. A particularly critical point can be considered to
be the narrowing in the metering nozzle. See Figure 7a
for a detailed view of the behaviour of the flow in the
metering nozzle and see Figure 7b for the way of the
velocity distribution just below the metering nozzle,
where the nozzle is expanded to a diameter 20 mm. By
observing the velocity vectors, a velocity profile can be
noted, from which it is obvious that the strand is being
strongly held back on the metering-nozzle walls. In the
area of the burst expansion just below the metering
nozzle, a "weak spot" is visible, where the velocity is
minimum, and where the minimum contact of the nozzle
walls with the flowing steel can be expected. An intense
steel strand begins to flow around the wall approximately
3 cm below the metering-nozzle extension. This area
may be, thus, prone to a greater erosion of the internal
profile of the nozzle.
It is understandable that during the casting a profile
may change (due to erosion, corrosion, fouling, etc.),
which greatly and adversely affects the flow character
and, thus, the stationarity of the flow fluctuations and the
emergence of steel-flow fluctuations. A more significant
impact of the metering-nozzle wear will more likely
apply to a smaller diameter of the metering nozzle, as
evidenced with the distribution of the shear stress on the
nozzle walls in Figure 8. The size of the shear stress
signals a danger of erosion. The most significant risk of
erosion is at the inlet to the metering nozzle. In the area
of the burst expansion below the metering nozzle, a
"weak spot", as mentioned above, with a length of
approximately 2 cm can be seen as well as the impact of
the strand on the walls of the nozzle and, thus, there is a
potential danger of a further significant erosion at a di-
stance of approximately 3 cm below the metering nozzle.
The shear-stress results confirmed the earlier finding that
an increase in the diameter of the metering nozzle
influences the length of the "weak spot", which is
reduced at the extension below the metering nozzle.
When designing the metering nozzles, the principles
resulting from the design of the nozzles must be
followed. In general, in the recommended design the
angle of the leading edges at the narrowing should be
less than 10° and, in the critical cases, it should be 5°. In
other words, at the connection of the metering nozzle to
the body of the nozzle there should be no sharp
transition; neither should there be one at the end of the
metering nozzle (the drainage edge).
These conclusions were also confirmed experimen-
tally on the physical model. It was also found that in the
case of a sharp transition at a metering nozzle, it was not
possible to provide a stable flow rate; the fluctuations
amounted to 10 % of the volume flow rate values.
6 DESIGN OF A NEW NOZZLE PROFILE
Based on the above results, a design of a new
structure (the inner profile) of the subentry nozzle was
created in cooperation with the engineers at T@, a. s., and
the staff of the VESUVIUS company, which should
provide a greater volume flow of the steel (and, thus,
higher casting speeds) and also eliminate the pulsation.
In the new construction design of the nozzle (the C
nozzle) the lower end of the metering nozzle was moved
510 mm away from the nozzle outlet and the inner
diameter of the nozzle above the metering nozzle was
enlarged to 30 mm. The inner diameter of the nozzle
below the metering nozzle remained the same
(20/22 mm), as well as the metering-nozzle profile with
a 17.5 mm diameter. A comparison of the original and
modified geometry is shown in Figure 9.
Prior to manufacturing an operational variant and
mounting the nozzle for the first test, a verification of the
impact of the diameters of the metering nozzles on the
volume flow rates, as well as the occurrence of
non-stationarity of the flow were carried out using
K. MICHALEK et al.: PHYSICAL AND NUMERICAL MODELLING OF A NON-STATIONARY STEEL FLOW ...
812 Materiali in tehnologije / Materials and technology 47 (2013) 6, 807–814
Figure 8: Distribution of the shear-stress profile (a risk of erosion) on the walls of the outlets for individual diameters of the metering nozzles
Slika 8: Razporeditev stri`nih napetosti (rizik erozije) na stenah pri izhodu za posamezni premer notranjega izlivka
physical and numerical modelling. Physical modelling
was performed using the same methodology with the
nozzles on a geometric scale of 1 : 1, and with the inter-
changeable metering nozzles of the diameters of (16, 17,
17.5 and 18) mm. Numerical modelling was performed
for the geometry of the new nozzle on the scale of 1 : 1
with only the metering nozzle of the 17.5 mm diameter.
The exact same input and output parameters of the
calculation were used in the simulations, including the
final step of the calculation. Figure 10 shows the effect
of the diameter of the metering nozzle on the volume
flow rates. To enable a comparison, the previously
discussed dependences on the metering nozzle in the
inlet part of a nozzle are also shown.
As evident from Figure 10, there was an increase in
the flow rate for all the tested diameters of the metering
nozzles at nozzle C during both physical and numerical
modelling in comparison with nozzles A and B. It can be
noted that the curve behaviour for nozzle C is shifted in
relation to the curve of nozzle A by approximately
11 L min–1, which is an increase by 18 % to 20 %. For
the casting speed of 2.5 m min–1 this would mean an
increase to 2.9 m min–1 or 3 m min–1.
Similarly to nozzle A, a relatively low dependence of
the volume flow rate on the diameter of the metering
nozzle was registered in nozzle C. The cause may be the
same: a poorly proportioned diameter of the nozzle
below the metering nozzle. Increasing the diameter to
approximately 25 mm would result in a greater incline of
nozzle C and, thus, to a further increase in the flow
through a subentry nozzle. These nozzles are signifi-
cantly more sensitive to potential erosion in the metering
nozzle.
Operational tests of the new subentry nozzle of type
C resulted in an increase in the casting speed of the first
casting sequence from 2.4 m min–1 to 2.8 m min–1,
representing an increase by 17 %. The situation was
similar with the second and third test sequences of the
heats.
7 CONCLUSIONS
The current type of the subentry shroud equipped
with metering nozzles used in the continuous casting of
steel on the CCM No. 2 in Tøinecké `elezárny, a.s., did
not allow a performance improvement of the casting
machine even if the diameter of the metering nozzle was
increased. For this reason, physical and numerical
modelling was conducted, the goal of which was to
assess and justify the behaviour of the steel flow and
suggest a possible solution. On the basis of theoretical
calculations and performed physical and numerical
modelling, the following was found:
An inability to increase the casting speeds by further
increasing the diameters of the metering nozzles in the
subentry nozzles is due to a too small diameter of the
nozzle-A outlet (20/22 mm) and, in particular, due to its
length (approximately 1000 mm). Both parameters
increase the hydraulic resistance and limit a further
increase in the volume flow rate of the steel.
Fluctuations in the volume flow rate (and, thus, the
casting speed) may be related to the turbulent nature of
the flow; a metering nozzle and its implementation can
have a significant impact; for this reason it should not
have sharp leading and output edges – it is recommended
that an angle is up to 5° and, at the maximum, up to 10°.
On the basis of the knowledge obtained with the
physical and numerical modelling, a new type of nozzle,
K. MICHALEK et al.: PHYSICAL AND NUMERICAL MODELLING OF A NON-STATIONARY STEEL FLOW ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 807–814 813
Figure 10: Effect of the metering-nozzle diameter on the rates of the
volume flows through subentry nozzles A, B, C at a 120 mm submer-
sion
Slika 10: Vpliv premera izlivka na hitrosti pretoka volumna skozi
potopljeno {obo A, B, C pri potopitvi 120 mm
Figure 9: Comparison of the inner profiles of the original and new
nozzles
Slika 9: Primerjava notranjega profila originalnega izlivka in novega
izlivka
known as type-C nozzle, has been proposed. Here a shift
of the lower end of the metering nozzle to a distance of
510 mm from the nozzle outlet, and an enlargement of
the inner diameter of the nozzle above the metering
nozzle to 30 mm were carried out.
Operational testing of the subentry nozzle led, during
the first sequence, to an increase in the casting speed
from 2.4 m min–1 to 2.8 m min–1, an increase by 17 %,
and similar results were achieved during the second and
third test sequences of the heats.
Acknowledgements
This paper could not have been realized without the
financial support from Tøinecke `elezárny, a. s., and
close cooperation with Vesuvius ^eská republika, a. s.
Further acknowledgement is owed to the project No.
CZ.1.05/2.1.00/01.0040 "Regional Materials Science and
Technology Centre", within the frame of the operation
programme "Research and Development for Innova-
tions" financed from the Structural Funds and from the
state budget of the Czech Republic.
8 REFERENCES
1 K. Michalek, K. Gryc, M. Tkadle~ková, D. Bocek, Model study of
tundish steel intermixing and operational verification, Archives of
metallurgy and materials, 57 (2012) 1, 291–296
2 K. Michalek, M. Tkadle~ková, K. Gryc, P. Klus, Z. Hudzieczek, V.
Sikora, P. Støasák, Optimization of argon blowing conditions for the
steel homogenization in ladle by numerical modelling, Proc. of the
20th Anniversary International Conference on Metallurgy and Mate-
rials, Metal 2011, Brno, 2011, 143–149
3 K. Gryc, K. Michalek, M. Tkadle~ková, Z. Hudzieczek, Physical
Modelling of Flow Pattern in 5-strand Asymmetrical Tundish with
Baffles, Proc. of the 19th International Conference of Metallurgy and
Materials, Metal 2010, Ro`nov p. Radho{tìm, 2010, 42–46
4 M. Tkadle~ková, K. Michalek, K. Gryc, Z. Hudzieczek, J. Pindor, P.
Støasák, J. Morávka, Comparison of extent of the intermixed zone
achieved under different boundary conditions of continuous billets
casting, Proc. of the 19th International Conference of Metallurgy and
Materials, Metal 2010, Ro`nov p. Radho{tìm, 2010, 47–52
5 J. Stetina, L. Klimes, T. Mauder, F. Kavicka, Final-structure precon-
dition of continuously cast billets, Mater. Tehnol., 46 (2012) 2,
155–160
6 J. Stetina, F. Kavicka, The influence of the chemical composition of
steel on the numerical simulation of a continuously cast slab, Mater.
Tehnol., 45 (2011) 4, 363–367
7 J. Pieprzyca, Z. Kudlinski, Determining transition zone size in conti-
nuous casting by the method of modelling, Stahl und Eisen, 124
(2004), 51–53
8 ANSYS FLUENT, Users Guide
9 M. Kozubková, Modelling of fluid flow, V[B-Ostrava, Ostrava, 2008
(in Czech), www.338.vsb.cz/PDF/Kozubkova-Fluent.pdf
K. MICHALEK et al.: PHYSICAL AND NUMERICAL MODELLING OF A NON-STATIONARY STEEL FLOW ...
814 Materiali in tehnologije / Materials and technology 47 (2013) 6, 807–814
M. VELICKA et al.: RESEARCH OF THERMAL PROCESSES FOR THE CONTINUOUS CASTING OF STEEL
RESEARCH OF THERMAL PROCESSES FOR THE
CONTINUOUS CASTING OF STEEL
RAZISKAVE TERMI^NIH DOGAJANJ PRI KONTINUIRNEM
ULIVANJU JEKLA
Marek Velicka, David Dittel, Rene Pyszko, Miroslav Prihoda, Miroslav Vaculik,
Pavel Fojtik, Jiri Burda
VSB – Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, 17. listopadu 15, 708 33 Ostrava, Czech Republic
marek.velicka@vsb.cz
Prejem rokopisa – received: 2013-04-03; sprejem za objavo – accepted for publication: 2013-04-08
The article deals with the determination of the basic indicators of heat transfer in the continuous casting of steel, which can be
described as an unsteady process with complicated boundary conditions for the solution. An analytical solution of this problem
is practically impossible and, therefore, mathematical modelling is applied with a certain simplification of the real conditions
and with a description of those criteria that influence the most the process of solidification and cooling. Using a simulation
program and the knowledge of input parameters, it was possible to predict the distribution of the thermal field of a continuously
cast blank in the course of its casting. Simulations also allowed us to deal with the issues of the inner structure, surface quality,
mechanical properties of a continuously cast blank, metallurgical length, change in the thickness of a strand shell and over-
heating of steel. Some results obtained with numerical simulations are documented for concrete examples.
Keywords: continuous casting of steel, modelling, heat transfer, shell, mould
^lanek obravnava dolo~anje osnovnih pokazateljev prenosa toplote pri kontinuirnem ulivanju jekla, ki se lahko opi{ejo kot
nestabilen proces s kompliciranimi robnimi pogoji za re{itev. Analitska re{itev tega problema je prakti~no nemogo~a, zato je
bilo uporabljeno matemati~no modeliranje z nekaterimi poenostavitvami realnih pogojev in z opisom tistih meril, ki najbolj
vplivajo na proces strjevanja in ohlajanja. S programom za simulacijo in poznanjem vhodnih parametrov je bilo mogo~e pred-
videti razporeditev temperaturnega polja kontinuirno ulite gredice med njenim ulivanjem. Uporaba simulacije je omogo~ila opis
notranje strukture, kvalitete povr{ine, mehanskih lastnosti kontinuirno ulite gredice, metalur{ke dol`ine, spreminjanja debeline
strjene skorje in pregretja jekla. Nekateri rezultati, dobljeni z numeri~no simulacijo, so prikazani za konkretne primere.
Klju~ne besede: kontinuirno ulivanje jekla, modeliranje, prenos toplote, skorja, kokila
1 INTRODUCTION
Mathematical precise description of thermal pro-
cesses during the continuous casting of steel is very
difficult, since the process of cooling and solidification
of a continuously cast blank is influenced by many
factors.
For this reason it is necessary to find, with the help of
mathematical and physical modelling, the criteria that
have the biggest influence on the solidification and cool-
ing of a continuously cast blank. Understanding the
thermal processes taking place during the continuous
casting of steel is of crucial importance because it ena-
bles a prediction of a formation of defects, an enhance-
ment of the thermal processes during continuous casting,
the optimum locations for cooling nozzles or a minimisa-
tion of breakout risks, etc.1
It is evident, that it is impossible to optimise the pro-
cess of continuous casting of steel only by modelling. A
very close interaction with the results of experimental
measurements is always necessary since these results
introduce, into the system, characteristic features of a
concrete continuous-casting machine. At present, it is
possible to perform, with the use of sophisticated
software, not only the thermal calculations but also the
calculations of stress conditions, and also to predict
segregation during the continuous casting of steel. The
finite-element method or finite-difference method are
used most frequently as the calculation algorithm in this
software.
This paper deals with the solution of thermal pro-
cesses during the continuous casting of round steel
blanks with a diameter of 410 mm, using the simulation
software based on the explicit finite-element method.
2 THERMAL FIELD OF A CONTINUOUSLY
CAST BLANK
The kinetics of an unsteady thermal field is described
with the Fourier partial differential equation.2 If we want
to describe the thermal field of a moving, continuously
cast blank, it is necessary to take into account, in the
transverse direction, also the rate of casting, by trans-
forming the classical Fourier equation to the Fourier-
Kirchhoff equation:
∂
∂
t
a t
q
c
v
 
= ⋅∇ +
⋅
2
p
(1)
where a/(m2/s) is the temperature-conductivity coeffi-
cient, ∇ 2 /(m–2) is the Laplace's operator, cp/(J/kg K) is
Materiali in tehnologije / Materials and technology 47 (2013) 6, 815–818 815
UDK 621.74.047:519.68 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)815(2013)
the specific heat capacity, /(kg/m3) is the density and
qv/(W/m3) is the inner heat source.
This equation can only be solved with the conditions
of monovalency of the solution, which, in the classical
concept, are divided into geometric, physical, surface
and initial conditions. Physical conditions are generally
created with the chemical composition of the cast steel,
geometric conditions are created with the shape of the
cast products, and the initial condition is closely linked
to the temperature of the liquid steel in a tundish.
Numerous researches concerning the continuous casting
of steel3–6 combined with modelling physical models use
the Neumann surface condition for the mould and the
Fourier condition for the secondary and tertiary cooling.
3 SIMULATION MODELLING
For numerical modelling, the ProCAST program tool
was used, which is based on the implicit finite element
method. Alternatively, an original program code Tefis
based on the explicit finite difference method, specially
designed for numerical modelling of the continuously
cast blank temperature field, was used. The program
Tefis enables real-time simulations too. The second
method requires compliance with a numeric stability
condition, which describes mutual dependence between
fineness of the calculation mesh and calculation time
step.
An application of program ProCAST consists of five
modules, which are functionally mutually inter-
connected. The first module is used for defining the body
geometry and for selecting the fineness and the shape of
the calculation mesh. It is then necessary to determine
the physical, initial and surface conditions that may be
modified in line with the requirements of the next
module. After correctly entering the initial calculation
parameters, the main calculation module of the software
launches the program.
The program allows a graphical visualisation of the
calculated results and an export of the values, graphs and
images for further processing. The program can launch a
calculation in two modes – the thermal and flow modes.
The results of the thermal mode give the thermal fields
of a round, continuously cast blank without the velocity
vectors of liquid steel. If it is also required to visualise
the velocity vectors, then it is necessary to activate the
flow mode (Figure 1).
Distribution of the temperatures during the solidifi-
cation or cooling is calculated in the nodal points of the
whole volume of a continuously cast blank (Figure 2).
The program also comprises a vast database of the
information about cast steel grades, including their ther-
mophysical properties (density, specific heat capacity,
heat conductivity, viscosity, etc.).
4 RESULTS OF NUMERICAL SIMULATIONS
Within the research two steel grades of different
chemical compositions used for the production of con-
tinuously cast blanks with a diameter of 410 mm were
subjected to a simulation (Table 1). The following
influences were analysed: the influence of the chemical
composition, the casting rate, overheating of the steel
above the liquidus temperature, the influence of the mol-
M. VELICKA et al.: RESEARCH OF THERMAL PROCESSES FOR THE CONTINUOUS CASTING OF STEEL
816 Materiali in tehnologije / Materials and technology 47 (2013) 6, 815–818
Figure 2: Calculation mesh of a blank
Slika 2: Ra~unska mre`a gredice
Figure 1: Velocity vectors of a blank
Slika 1: Vektor hitrosti v gredici
ten-steel level in the mould on the resulting technolo-
gical parameters, generally represented with the metal-
lurgical length, the length of the liquid phase, the surface
temperatures of a continuously cast blank at the outlets
of the primary and secondary cooling zones, and the
thickness of the strand shell.
Table 1: Chemical composition of the steel in mass fractions, w/%
Tabela 1: Kemijska sestava jekla v masnih dele`ih, w/%
Brand of
the steel
Chemical composition, w/%
C Mn Si P S
Steel A 0.168 1.350 0.381 0.013 0.007
Steel B 0.913 0.342 0.246 0.009 0.004
The chemical compositions, especially the carbon
contents, have a principal influence on the heat removal
in all the zones of the continuous-casting machine. In the
case of low-carbon steels the biggest shrinkage of the
strand shell occurs as a result of a peritectic reaction as
well as a considerable deceleration of its growth, which
is manifested by a reduction of the heat-flow density
through the mould wall to its minimum. Another result
of this reaction is an increased occurrence of surface
defects of a continuously cast blank, particularly of
surface cracks.
The casting rate is related to the dimensions of a con-
tinuously cast blank, the type of steel and the type of
mould. Higher values of the casting rate cause a shorter
stay of a continuously cast blank in the mould, which
increases the surface temperatures and, simultaneously,
also the local values of the thermal-flow density.
It may be stated on the basis of the performed simu-
lations that, with respect to investigating the influence of
the casting rate on the metallurgical length and the length
of the liquid phase, their very strong linear dependence is
evident and is best characterised with a high value of the
slope in the regression equations (Figure 3).
The growth of the strand shell with the changing
casting rate (Figure 4) confirms the previous scientific
research works,7,8 which document the fact that the para-
bolic law may be used for the solidification of an ingot to
a distance of approximately 75 % of the diameter of a
continuously cast blank from the ingot mould surface.
When this critical value is exceeded, the differences are
more pronounced, as it is known that the solidification of
round, continuously cast blanks accelerates towards their
centre, in comparison to the ingot mould.
The thickness of the shell can determine the results of
the logarithmic equation for steel A:
	

= − ⋅ ⋅ + ⋅ ⋅
⋅
− −
− ⋅ ⋅−
( . ln . )
( . ln
1490 10 1683 102 1
6 378 10 3
v z
v z + ⋅
−5 874 10 1. )
(2)
And for steel B, the following applies:
	

= − ⋅ ⋅ + ⋅ ⋅
⋅
− −
− ⋅ ⋅−
( . ln . )
( . ln
13690 10 1869 102 1
9 369 10 3
v z
v z + ⋅
−5 456 10 1. )
(3)
where vz/(m/s) is the specific casting rate and /s is the
casting time.
Finally, an analysis of the influence of overheating of
steel on the investigated technological parameters was
made. Overheating of steel can be defined as the diffe-
rence between the casting temperature and the liquidus
M. VELICKA et al.: RESEARCH OF THERMAL PROCESSES FOR THE CONTINUOUS CASTING OF STEEL
Materiali in tehnologije / Materials and technology 47 (2013) 6, 815–818 817
Figure 5: Blank-temperature change depending on the overheating of
steel B
Slika 5: Sprememba temperature gredice glede na pregrevanje jekla B
Figure 3: Influence of the metallurgical length depending on the cast-
ing speed in the case of steel A
Slika 3: Vpliv metalur{ke dol`ine v odvisnosti od hitrosti litja v pri-
meru jekla A
Figure 4: Increase in the thickness of the casting shell depending on
the casting speed in the case of steel A
Slika 4: Pove~anje debeline skorje, v odvisnosti od hitrosti litja v pri-
meru jekla A
temperature. This temperature difference should be, from
the viewpoint of the operation, as small as possible, since
it reduces the thermal stress and the energy intensity of
the continuous casting of steel, as well as the scrap factor
and the formation of surface cracks (Figure 5).
The results of the simulations also show that the
greater is the overheating of steel, the later is the solidi-
fication of a continuously cast blank and the greater is
the distance of this solidification from the molten-steel
level. The overheating of steel in the interval from 0 °C to
45 °C increases the metallurgical length, on average, by
0.5 m, and a similar trend is characteristic of all steel
grades. Even a greater influence of this parameter on the
increase in the length of the liquid phase was confirmed.
An increase in the overheating temperature by 40 °C
above the liquidus temperature caused an elongation of
the metallurgical length, on average by 3 %; however,
the length of the liquid phase was increased even by
20 %, while the surface temperature of a continuously
cast blank at the output from the primary and secondary
cooling zones increased by up to 3 %.
5 CONCLUSION
The optimum configuration of the thermal mode of
the continuous casting machine substantially influences
the quality of the cast products. The processes partici-
pating in the cooling and solidification of a continuously
cast blank show their physical characters in the compli-
cated transfer phenomena of the heat and mass transfer.
In order to allow a general solvability of these tasks,
computer simulations are used making it possible to defi-
ne the optimum parameters of the continuous casting of
steel. The casting rate, which influences practically all
the casting parameters, appeared to be most important
quantity.
Acknowledgements
The work was realised within the frame of the grant
project No. SP 2013/53, under the financial support of
The Ministry of Education, Youth and Sports.
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massive ductile cast-iron casting, Mater. Tehnol., 44 (2010) 2, 93–98
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Speed on Slab Broadening in Continuous Casting, Steel Research
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M. VELICKA et al.: RESEARCH OF THERMAL PROCESSES FOR THE CONTINUOUS CASTING OF STEEL
818 Materiali in tehnologije / Materials and technology 47 (2013) 6, 815–818
F. VODOPIVEC et al.: STATIONARY-CREEP-RATE EQUATIONS FOR TEMPERED MARTENSITE
STATIONARY-CREEP-RATE EQUATIONS FOR
TEMPERED MARTENSITE
ENA^BE ZA HITROST STACIONARNEGA LEZENJA @ARJENEGA
MARTENZITA
Franc Vodopivec, Borut @u`ek, Danijela A. Skobir Balanti~, Monika Jenko,
Matja` Godec, Darja Steiner Petrovi~
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
franc.vodopivec@imt.si
Prejem rokopisa – received: 2013-07-01; sprejem za objavo – accepted for publication: 2013-08-28
The stationary-creep equations including parameters as the average particle spacing or/and particle size are examined in terms
of the effects of different equation parameters on the difference of calculated and experimental creep rate. It is concluded that
the difference between the experimental creep rate and the creep rate deduced from the recently proposed creep equation is in
the range of about ±25 % if the effects of the steel composition and the changes in the particle size and spacing during the creep
test are considered with sufficient accuracy.
Keywords: creep-resistant steel, creep equations, particle size and spacing, difference between the calculated and experimental
creep rate
Ena~be za izra~un hitrosti stacionarnega lezenja, ki vklju~ujejo kot parametre povpre~no velikost in oddaljenost izlo~kov, so
analizirane s stali{~a razlike med eksperimentalno in izra~unano hitrostjo lezenja. Sklep je: pred kratkim predlo`ena ena~ba da
razliko v razponu okoli 25 %, ~e so prav upo{tevani vplivi sestave jekla in spremembe velikosti in oddaljenosti izlo~kov pri
preizkusih lezenja.
Klju~ne besede: jeklo odporno proti lezenju, ena~be za izra~un hitrosti lezenja, velikost in oddaljenost izlo~kov, razlika med
izra~unano in eksperimentalno hitrostjo lezenja
1 INTRODUCTION
Theoretical and semi-empirical constitutive equations
were proposed for calculating the stationary (secondary)
creep rate of high-chromium creep-resistant steels using
as parameters physical constants, creep-test conditions,
steel properties, microstructure as well as the average
particle size and spacing1–5. The equations were deduced
applying different concepts of interaction between parti-
cle shape, size and spacing and mobile dislocations in
the metal matrix. Also, more than 20 semi-empirical
creep equations have been proposed6–8 so far using the
rationalisation of the stress-strain-creep test results
obtained for different steels. In these equations, the base
parameters of microstructures of creep-resistant steels,
matrix grain size and particle size and spacing are
considered as constants based on the results of the creep
tests. The recently derived creep equation is based on the
disjunctive-matrix (dm) concept of particle stress
transfer of acting stress on gliding and climbing of
dislocation segments in the matrix after crossing the
strait between two neighbour particles9. By specific
angles of acting stress, matrix grain-dislocation slip
plane and stress-transfer particle face, the dislocation
gliding stress is decreased significantly, while the climb
stress changes much less10. It was concluded that the
creep rate depended on the dislocation glide velocity in
disjunctive matrix9,10 which is proportional to the gliding
stress1. The calculations have shown that in a disjunctive
matrix with backstress due to opposite directions of the
components of acting and transfer stress10, the gliding
stress may be decreased by up to about four orders of
magnitude. The concept of a dm stationary creep of
tempered martensite allows a simple explanation of the
effect of matrix grain size on the stationary-creep rate
and of the transition from stationary to accelerating
creep10. In this article, the accuracy and reliability of
parameters of the equations, proposed so far for calcul-
ating the stationary-creep rate, such as volume share,
particle size and/or spacing are examined.
2 STATIONARY-CREEP EQUATIONS
Equation (1) was deduced for a dislocation line
advancing through a field of obstacles (particles) by
locally climbing over them1:

 
=
⋅ ⋅ ⋅
⋅ ⋅
⎛
⎝
⎜
⎞
⎠
⎟ ⋅
⎛
⎝
⎜ ⎞
⎠
⎟2 D G b
f k T G
(1)
where  is the creep rate (s–1), D is the iron-diffusion
rate in ferrite, G is the shear modulus at T, b is the
Burgers vector, f is the volume fraction of particles in
the matrix, k is the Boltzmann constant and  is the
acting stress.
It was stated that the equation is reasonable for the
limit of density of mobile dislocations:  < 1/(d/2)2. An
Materiali in tehnologije / Materials and technology 47 (2013) 6, 819–824 819
UDK 669.14.018.44:539.37 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)819(2013)
interaction of particles and mobile dislocations is
considered including the volume share of carbide, f. In
this work, experimental and calculated creep rates are
obtained for the steel with the average particle sizes of d1
= 280 nm and d2 = 420 nm; the limits 1 = 1.96 × 1014
m–2 and 2 = 4.41 × 1014 m–2 are obtained, while  =
0.978 × 1014 m–2 is deduced later for the examined steel.
Thus, the validity limit  < 1/(d/2)2 for equation (1) is
fulfilled.
Equation (2)2 was derived as an improvement of
equation (1) considering the density of mobile disloca-
tions () and particle spacing ():

   
=
⋅ ⋅ ⋅ ⋅
⋅ ⋅
b D
k T G
2
(2)
For the detachment concept of stationary creep based
on TEM observations of the detachment of the disloca-
tion line from the free surface of oxide particles in alu-
minium, the creep equations (3) were derived3–5:
 exp
 
=
⋅ ⋅⎛
⎝
⎜ ⎞
⎠
⎟ ⋅ −
6 D
b
E
E G b r k
k
= ⋅ ⋅ ⋅ −
−⎡
⎣⎢
⎤
⎦⎥
⎛
⎝
⎜
⎞
⎠
⎟2
3 2
1
1
( )
/
 d
(3)
where E is the detachment activation energy, d is the
detachment stress and k = Tp/Tm with the Tm dislocation
line energy.
It was concluded11–13 that equations (3) required
significant improvements for the calculation of the sta-
tionary-creep rate of the creep-resistant steels streng-
thened with non-coherent particles.
Equation (2) was modified introducing the constants
for the particle size and spacing, considering their depen-
dence on temperature and rationalising the acting stress
with the shear modulus or yield stress at the creep-test
temperature. Reasonable fits of the experimental and
calculated creep rates for different test temperatures were
obtained with the stress exponents in the range from n =
4.5 to n = 16 and the creep activation energy signifi-
cantly higher than the activation energy of iron diffusion
in ferrite was calculated6–8,14–16:
 exp
( / )

 
= ⋅
⎛
⎝
⎜ ⎞
⎠
⎟ ⋅
−
C
G
n
C
G
RT
1
2
(4)
Based on the experimental data11–13, with C1 = 4.3 ×
10–41 K s–1 and C2 = 2.85 × 107, the stress exponent n =
4.5 and the activation energy of 680 kJ m–1 was
deduced16–18. Recently, only a small difference was found
in the creep activation energy for the creep rates calcu-
lated considering a negligible change in the particle size
and spacing19.
The simplified equation (5) includes the parameters
of the previous equations and considers the effect of the
particles as a difference of particle spacing and size ( – d).9
The basis of the equation is the assumption that the
mobile-dislocation driving energy is consumed mainly
for the lengthening of the dislocation segment due to the
transition of the disjunctive-particle matrix. Based on the
dm stationary-creep concept and experimental data, the
stress exponent n = 3.65 was deduced and the obtained
difference between the calculated and experimental creep
rates was about 10 %.9 For the cube-shaped particles, the
exponent up to n  4 was deduced recently10:

( )

  
=
⋅ ⋅ ⋅ ⋅ −
⋅ ⋅
c b D d
k T G
2
(5)
The creep rate  in s–1 is obtained in equation (5) and
written as:

( )
 	
  
=
⋅ ⋅ ⋅ −
⋅ ⋅
c
b D d
k T G
2
and 	





= = 4.78 · 103
The density of mobile dislocations was calculated
as18:


=
⋅ ⋅ ⋅
⎛
⎝
⎜ ⎞
⎠
⎟
a M G b
2
(6)
where M = 3 (the Taylor factor).
In Table 1 for the tempered martensite with a uni-
form distribution of particles, the experimental creep
rate, the creep rates calculated using equations (1), (2)
and (5) with c = 1, b = 2.5 × 10–10 m,  = 170 MPa,  =
9.78 × 1013 m–2, D = 1.19 × 10–19 m2 s–1 (in19), k = 1.38 ×
10–23 J K–1, G = 57.3 × 103 MPa, T = 853 K and f = 0.037
(in9,20) are given.
Table 1: Creep rates calculated for two average particle sizes and
spacings. In equation (1) shear stress g = a/M = 170/3 = 56.6 GPa
and in equation (5) n = 3.65 were considered.
Tabela 1: Hitrosti lezenja, izra~unane za dve povpre~ni velikosti
izlo~kov in razdalji med njimi. V ena~bi (1) je upo{tevan stri`ni modul
g = a/M = 170/3 = 56,6 GPa in n = 3,65.
Particles Calc. creep rate, /s–1 Exp. creeprate
Size Spacing  
D  s–1 s–1
nm nm Eq. 1 Eq. 2 Eq. 5
280 1065 2.36 ·10–11 3.42 ·10–9 12.2 ·10–8 12.1 ·10–8
420 1650 5.42 ·10–9 18.9 ·10–8 21.2 ·10–8
The difference between the experimental and cal-
culated creep rate for equation (5) is reasonable. It is
greater if deduced from equation (1), probably due to the
substitution of the effects of segmentation of the dislo-
cation line, the interaction dislocation segments and the
particles with the volume fraction of particles. The lower
creep rate deduced from equation (4) is ascribed to the
assumption of an interaction between the mobile dislo-
cations and the modified Orowan mechanism. However,
by stationary creep, the applied stress is significantly
lower than the Orowan stress. The creep rates quoted in
this work were determined and calculated for 100 h tests
with static-tensile stress of 170 MPa at 853 K. At this
temperature, the Young modulus (E) for two high-
chromium creep-resistant steels is E853 = 57.3 × 103 MPa
F. VODOPIVEC et al.: STATIONARY-CREEP-RATE EQUATIONS FOR TEMPERED MARTENSITE
820 Materiali in tehnologije / Materials and technology 47 (2013) 6, 819–824
= 0.82 × E22 with the E22 Young modulus at room
temperature21. Assuming that the effects of temperature
on yield stress (y) and Young modulus are equal, the
experimental rates in Table 2 were obtained as follows:
y853  0.82 × y22 = 0.82 × 400 = 328 MPa. The ratio of
the acting-creep stress to the yield stress, /y = 170/328
= 0.52 at 853 K, makes the Orowan interaction between
mobile dislocations and non-shearing particles
improbable at the creep-test temperature.
3 ACCURACY AND RELIABILITY OF THE
PARAMETERS IN EQUATION (5)
In creep-resistant steels, the ratio of numbers of iron
atoms and atoms of the alloying elements is above 0.85.
The difference between the parameters of chromium and
iron is: a (Cr-Fe) = 2.88 × 10–10 – 2.48 × 10–10 = 0.4 ×
10–10 m. Theoretically, the maximum extension of ferrite
lattice due to chromium atoms in the solid solution is 0.4
× 0.15 = 0.06 × 10–10 m, thus 0.06/2.48 = 0.024 × 10–10
m. It is concluded that the effect of the difference in the
Burgers vector due to alloying elements in solution in
ferrite for creep-resistant steels is negligible.
The lattice parameters of ferrite and the Burgers vec-
tor are increased at the creep-test temperature because of
the thermal expansion. The coefficient of linear thermal
expansion of  iron is 11.8 × 10–6 K–1. Consequently, at
the creep-test temperature, the lattice parameter is for
2.48 × 10–10  11.6 × 10–6  853 = 2.45 × 10–12 m. Thus,
the effect of increase in the -Fe lattice at creep tempe-
rature is negligible.
The iron-diffusion rate was deduced from data in19.
The values of activation energy for steels with 0.87 % Cr
to 5 % Cr differ by less than 5 % from the average value.
Considering the difference in D0 it is concluded that the
error due to differences in the diffusion rate is within
±10 %.
The acting stress is included as a parameter in all
theoretical and semi-empirical equations, while n = 1, a
= 170 MPa, 1.1 = 284 MPa, 1.2 = 475 MPa, 1.5 = 2 116
MPa and 2 = 28 900 MPa. It is evident that already the
difference of n = 0.2 generates the difference between
the calculated and experimental creep rates of 2.8 times.
The accuracy of the stress-exponent deduction depends
on the accuracy of the assumed changes in particle con-
tent, size and spacing during the creep tests. The effect
of the stress exponent is explained with the dm concept
of the stationary creep10. It was, namely, calculated that
the gliding stress in disjunctive parts of the matrix could
be decreased by up to 104 times, with respect to the
acting stress, to the backstress generated by the angles of
acting stress, grain-dislocation slip plane and stress-
transfer particle face. Thus, it seems reasonable to
ascribe a great range of the quoted stress exponents (n =
4.5 to n = 16) to the undervaluation of the effect of
changes of particle size and spacing during the experi-
mental creep tests.
The mobile-dislocation density calculated using
equation (5)18,  = 0.978 × 1014 m–2, was checked using
the relations suggested for copper22:
 s G b= ⋅ ⋅ ⋅
1
2
and  = ⋅ ⋅v b (7)
where s is the acting shear stress (s = a/3), v is the
dislocation velocity, while the other parameters are the
same as in equation (5); it was obtained that  = 1.04 ×
1014 m–2. The difference between the two deduced
mobile-dislocation densities of 5.3 % supports the con-
clusion that the accuracy of the calculations of the
mobile-dislocation density deduced from equation (6) is
of ±10 %.
In the creep equations, the coefficient of iron diffu-
sion in ferrite determined for the equilibrium content of
vacancies is considered. The tensile stress increases the
content of vacancies23,24 and the increase in the diffusion
rate is deduced from this relation24:
D
c
c
D1
0
1
0=
⎛
⎝
⎜
⎞
⎠
⎟ ⋅ (8)
where c0 and D0 are the vacancy content and diffusion
rate at  = 0, while c1 and D1 are the vacancy content
and diffusion rate at the acting stress. It was deduced
from data in24 that D1 = 1.14D0. In the case of
calculations of creep rate at the acting stress a = 170
MPa, the effect of the increase of diffusion rate is
relatively low with respect to the effect of the stress
exponent, however, it should not be neglected in accu-
rate calculations.
For the activation energy of iron diffusion in ferrite,
Q = 	 and D0 251 kJ/mol 19 as 289 kJ/mol 22 are quoted.
The calculated diffusion rates at 853 K differ by several
orders of magnitude. Thus, it is clear that the calculated
creep-rate values are reliable only if deduced using
reliable data in the deduction of iron diffusion in ferrite.
In this case, the error in the calculation of the creep rates
would be in the range of 10 %.
In modern high-chromium creep-resistant steels, the
molar contents of carbide-forming vanadium and nio-
bium are lower than the molar content of carbon, while
the content of chromium is higher. Consequently, the
particles of carbide M23C6 with a molar composition of
Cr18Fe3Mo223–25 are stable up to the austenite tempe-
rature. The solubility of VC and NbC carbides in ferrite
is higher and at 700 °C virtually all VC and a significant
part of NbC are dissolved in ferrite. Consequently, the
effects of the creep-testing temperature on the quantity,
size and spacing of particles of different carbides may be
different. It is greater for a lower content of carbide-
forming elements in solution in ferrite and higher solid
solubility of carbide.
From the data in25, it is calculated that in the
0.18C11.4Cr steel with the particles size of 0.1 μm,
M23C6 increase after 1000 h at 650 °C by about d =
2 %, and by about d = 12 % at 750 °C. According to
F. VODOPIVEC et al.: STATIONARY-CREEP-RATE EQUATIONS FOR TEMPERED MARTENSITE
Materiali in tehnologije / Materials and technology 47 (2013) 6, 819–824 821
equation (5), the creep rate is proportional to ( – d). The
particle spacing is deduced from2:
 =
⋅
⋅
4
1 3
d
fπ /
(9)
By constant volume share of carbide, an increase in ave-
rage particle size of 12 %, the increase of particle
spacing is   = 4 d = 48 % and, according to equation
(5), the increase of difference between experimental and
calculated creep rates of (48 – 12) % = 36 %. The logi-
cal conclusion is that by stationary-creep tests, reliable
data could be obtained only by a very limited change
initial particles size and spacing. The difference bet-
ween calculated and experimental creep rates would be
even greater if the dissolution of the VC and NbC
particles was omitted.
Assuming that after the tempering at 800 °C M23C6
particles are situated in the centre of cubes with edge e,
equation (10) was deduced20. Further, it was assumed
that, in the micrograph plane, the sections of all cubes
are squares with side e and with a particle in the square
centre. The average size of the particle is than20:
d
F f
N
=
⋅ ⋅π 1 2 1 3
1 24
/ /
/
p
(10)
where F is the surface area of the assessed micrograph
and Np is the number of the particles in the micrograph,
f is the volume fraction of carbide calculated from the
content of carbon in the steel and the previously quoted
relation2, e =  = (4d)/ f1/3. In equation (5) it is
assumed that the cube edge and the square side lengths
are equal. This assumption is partly incorrect, as the
geometrical shape of a cube section depends on the
angle between the section (micrograph) plane and the
cube edge, which is, on average, a rectangle with the
average side length s = 1.205e = (e + e 2)/210. The
average spacing of the particles in the micrograph plane
is then not  = F/N but ' = 1.205 F/N and equation
(10) is inexact. However, the size of most of the parti-
cles is below the average size25. On the SEM micro-
graphs, the particles with the size of 25 nm are visible
by magnification of 104 times. The dissolution velocity
of particles, vd ! d3/d2, is inversely proportional to the
particle surface area26. As in equation (5) ( – d) is
included, it is assumed that the omission of the particles
unseen on the micrographs at the magnification of 5 ×
103 would decrease the calculated creep rate by up to
10 % if interaction particle-mobile dislocation was inde-
pendent of particle size.
The volume fraction of carbide f is a parameter in
several equations and it is deduced considering the
contents of carbon and of carbide-forming elements in
steel2,20,27. For high-chromium steels, the calculated value
of f is accurate as the content of carbon is much higher
than the content of chromium bound in M23C6. The
calculation of the carbide-volume fraction is unreliable
in case of a strong effect of temperature on the solubility
of particles of different carbides, i.e., VC and NbC
carbides in creep resistant steels. In this case, the volume
fraction of carbides is calculated with an acceptable
accuracy up to about 650 °C when vanadium in niobium
are virtually totally bound in VC and NbC and the
difference in carbon is bound to M23C6.
4 ACCURACY OF EQUATION (10)
Equation (10)20 was deduced assuming:
• a spherical shape of the particles with different sizes,
• the particles are situated in the centres of the cubes of
different sizes;
• the particle spacing is equal to the cube edge, which
is equal to the square side on the micrographs and 
= (F/N)1/2 with F – surface and N – number of parti-
cles on the examined micrograph.
In Table 2 the results of an assessment of the micro-
graphs of the same specimen of the X20 steel, quenched
and tempered for 400 h at 800 °C, are listed. The volume
fraction of carbide Cr18Fe3Mo2C6, with f = 0.045, in the
steel tempered at 800 °C 20 was deduced as:
f
C
Cr Fe Mo C
C
V
=
⋅
+ + +
×
( )
.
18 3 2 6
6
7 04 st
(11)
where C is the content of C in 100 g of steel in g; Cr, Fe,
Mo and C are the atom weights; 7.04 is the specific
weight of carbide (g/cm3) and Vst is the volume of 100 g
of steel.
Table 2: Results of assessment of micrographs of M23C6 particles in
grip and gauge parts of a specimen after 100 h of creep tests with  =
170 MPa and 853 K
Tabela 2: Rezultati analiz izlo~kov M23C6 na posnetkih iz merilnega
in vpetega dela preizku{anca po preizkusu lezenja 100 h pri 853 K in
170 MPa
Magnifi-
cation
Gauge, particles, μm Grip, particles, μm
Spacing Size Spacing Size
3500 0.68 0.19 0.72 0.20
5000 0.61 0.17 0.58 0.16
10000 0.55 0.15 0.55 0.15
The average particle spacing and size decrease with
the micrograph magnification and reflect the discerni-
bility of very small particles. By the magnification of 104
times, the particles of a size down to 25 nm are dis-
cerned. The omission of smaller particles would cause an
error of about 10 % with respect to the particle spacing
and size; however, this inaccuracy is diminished with the
assumption that the section of the cubes in the micro-
graph plane is orthogonal to the cube edge, while the
section is rectangular and the average particle spacing is,
accordingly, greater. The resulting error is opposed to the
error due to the number of very small particles undis-
cerned on the examined micrograph. The comparison of
data suggests that by sufficient magnification of the
micrographs and a sufficient number of counted parti-
F. VODOPIVEC et al.: STATIONARY-CREEP-RATE EQUATIONS FOR TEMPERED MARTENSITE
822 Materiali in tehnologije / Materials and technology 47 (2013) 6, 819–824
cles, the range of error of assessment the particle size
and spacing is ±10 %.
The error could be significantly greater by insuffi-
cient surface quality of the SEM micrograph specimen.
In Figures 1 and 2 SEM micrographs of the same steel
are shown. The surface quality of the specimen in Figure
1 was sufficient and only M23C6 particles of different
sizes are visible, while the quality of surface of the spe-
cimen in Figure 2 was insufficient, as, in strongly etched
areas, a great density of very small particles is visible
due to an unsuited specimen preparation.
5 CONCLUSIONS
The accuracy of the quoted creep-rate equations,
including as parameters particles size and spacing, de-
pends on the accuracy of consideration of the changes
due to the creep-test stress and temperature, and is
related to:
• physical parameters at test temperature with a negli-
gible change by change of test temperature;
• inaccuracies of the calculated changes in steel para-
meters: iron-diffusion rate in ferrite, shear modulus
and mobile-dislocation density are in the range of
±10 %;
• changes of average particle-size spacing during the
tests may be greater than 100 % depending on the
creep-test temperature, time and solubility of parti-
cles in ferrite;
• accuracy and reliability of the stress exponent as, due
to the exponent change of 0.2, the calculated creep
rate increases 2.8 times.
Thus, if the effects of enumerated changes of para-
meters are considered with sufficient accuracy, espe-
cially the change in the particle size and spacing, the
stress coefficient was deduced from tests enabling the
acurate changes of particle size and spacing, the diffe-
rence between the experimental creep rate and the rate
calculated for tempered martensite using equation (5),
should be in the range of ±25 %.
Acknowledgment
The authors thank the Slovenian Research Agency
for the funding of this work.
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Materiali in tehnologije / Materials and technology 47 (2013) 6, 819–824 823
Figure 1: Particles in the gauge specimen part
Slika 1: Izlo~ki v merilnem delu preizku{anca
Figure 2: Particles in the grip specimen part with areas of inappro-
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S. V. KOBELSKY: NUMERICAL CALCULATIONS OF STRESS-INTENSITY FACTORS FOR A WWER-1000 REACTOR ...
NUMERICAL CALCULATIONS OF STRESS-INTENSITY
FACTORS FOR A WWER-1000 REACTOR PRESSURE
VESSEL UNDER A PRESSURIZED THERMAL SHOCK
NUMERI^NI IZRA^UNI FAKTORJA INTENZITETE NAPETOSTI V
REAKTORSKI TLA^NI POSODI WWER-1000 PRI TOPLOTNEM
[OKU POD TLAKOM
Sergey Vladimirovich Kobelsky
G. S. Pisarenko Institute for Problems of Strength, National Academy of Science of Ukraine, Kiev, Ukraine
ksv@ipp.kiev.ua
Prejem rokopisa – received: 2013-07-24; sprejem za objavo – accepted for publication: 2013-10-03
Two methods of determining stress-intensity factors based on the calculations of J- and G-integrals are analyzed for a simulated
emergency-cooling regime of a WWER-1000 reactor pressure vessel. A finite-element analysis of the convergence of the results
determining the stress-intensity factors using different approaches and different finite-element meshes was done.
Keywords: reactor pressure vessel (RPV), finite-element method (FEM), J- and G-integral, crack, stress-intensity factor (SIF)
Analizirani sta bili dve metodi za dolo~anje faktorja intenzitete napetosti na osnovi izra~una J- in G-integrala pri simuliranem
re`imu hitrega nujnega ohlajanja tla~ne posode reaktorja WWER-1000. Narejena je bila analiza raztrosa rezultatov s kon~nimi
elementi pri dolo~anju faktorja intenzitete napetosti in razli~nih mre`ah kon~nih elementov.
Klju~ne besede: tla~na posoda reaktorja (RPV), metoda kon~nih elementov (FEM), J- in G-integral, razpoka, faktor intenzitete
napetosti (SIF)
1 INTRODUCTION
The safety control of nuclear power plants and, in
particular, of reactor pressure vessels (RPV) as their
most critical elements, is one of the important theoretical
and practical problems. One of the promising ways of
solving this problem is mathematical modeling of the ki-
netics of the thermomechanical state of a reactor pres-
sure vessel, which requires solutions of non-stationary
non-linear problems of thermomechanics. The computa-
tional justification for the strength of the RPV is carried
out on the basis of an analysis determining fracture-me-
chanics parameter values, namely, stress-intensity factors
(SIF).
The complexity of the problem under consideration, a
large amount of computations and the necessity of per-
forming multiple-choice calculations make it essentially
impossible to justify the convergence of the obtained re-
sults when solving practical problems with the finite-ele-
ment method (FEM).1 In the majority of cases, for the
structures of a complex shape, calculations are per-
formed on finite-element meshes with a spacing of 0.5
mm to 1 mm in the vicinity of crack front points. At the
same time, in the finite-element method, the convergence
of numerical-calculation results is a necessary but not a
sufficient condition for their reliability. The degree of re-
liability of the results can be enhanced by comparing the
values of the analyzed parameters obtained with the use
of various hypotheses or criteria.2 It should be noted that
the conditions for attaining the numerical convergence
cannot automatically be extended to another class of
problems. Thus, the numerical-convergence conditions
for the problem to be solved using the deformation-plas-
ticity theory, cannot be fulfilled for the problem to be
solved on the basis of the flow theory.
The goal of the present paper is to analyze the con-
vergence of the results determining the SIF calculated
using various concepts on different finite-element
meshes.
2 PROBLEM STATEMENT
For typical emergency-cooling conditions, known as
the pressurized thermal shock, two computational mod-
els of the WWER-1000 reactor pressure vessel (RPV),
with a built-in crack and without it, are analyzed (Figure
1).
A circumferential semi-elliptical crack with the depth
of a = 15 mm and the ratio a/c = 1/3 of the half-axis3 that
is located at the level of welded joint No. 4 is postulated.
The procedure of embedding the crack in the finite-ele-
ment model of a pressure-vessel fragment is used.
3 METHOD OF CALCULATION AND THE
SOFTWARE
The problem is solved in a three-dimensional formu-
lation using a mixed finite-element method scheme
Materiali in tehnologije / Materials and technology 47 (2013) 6, 825–830 825
UDK 621.039.536.2 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)825(2013)
(MFEM).4 The main advantage of the MFEM, as com-
pared to the classical FEM approach in the form of the
displacement method, is the possibility of ensuring the
continuity of the approximation not only for displace-
ments but also for stresses and strains as well as the pos-
sibility of accurately satisfying the static boundary con-
ditions on the body surface. The mixed scheme is
implemented in the SPACE&RELAX software package
created at the G. S. Pisarenko Institute for Problems of
Strength of the National Academy of Sciences of
Ukraine.
The procedure of solving the problem includes three
stages. At the first stage, the temperature fields are deter-
mined at the specified instants of time. At the second
one, the fields of displacements, stresses and strains for
the reactor-pressure-vessel (RPV) model are calculated
taking into account the thermomechanical loading his-
tory. At the third stage, a refined calculation of the fields
of displacements, stresses and strains is performed for
the RPV fragment with a built-in crack, and the values of
the stress-intensity factors (SIF) are determined.
When solving the problem of the elasto-plastic for-
mulation, the loading process is divided into separate
time steps, within each of which the elasto-plastic prob-
lem is solved considering the residual plastic strains de-
termined at the previous step. To solve the elasto-plastic
problem at the current stage of loading, a two-step itera-
tion process is used. The problem of the plasticity theory
is solved with the method of variable parameters of elas-
ticity, whereas at the internal steps pertaining to elastic-
ity, the problems are solved using the conjugate-gradient
method with the zeroth-order initial approximation.
4 METHODS OF DETERMINING THE SIF
VALUES
To determine the SIF values, the J-integral and
G-integral concepts are used. In the problems of both
the linear and nonlinear theories of elasticity as well as
the deformation-plasticity theory, the J-integral value
characterizes the energy-release rate for a virtual
infinitesimal crack extension.5,6 The applicability of the
J-integral concept is limited by the following require-
ments: the material is nonlinearly elastic and homo-
geneous; under elasto-plastic deformation, the load
increases in proportion to one parameter, that is, the
deformation-plasticity theory is applicable. The use of
the J-integral in the flow theory has no theoretical
justification.
In this paper, the method of equivalent volume inte-
gration (EVI) proposed by G. P. Nikishkov is used for
the calculation of the J-integral value.7
The G-integral of the crack closure is defined as the
work of adhesive forces acting in a body ahead of the
crack front and preventing a separation of its surfaces
when the crack propagates negligibly. For the instant of
the crack growth onset when the value of the crack prop-
agation can be considered as infinitesimal, the G-integral
is calculated with formula:8
G = ⋅
1
2
  (1)
Here,  is the stress calculated normally to the crack
plane at the finite-element node located ahead of its
front;  is the opening displacement calculated as a dif-
ference in the displacements normal to the crack plane
for two mesh nodes adjacent to the crack front and lying
on its different surfaces.
In contrast to the J-integral, the G-integral value re-
mains meaningful when used in the flow theory for an
arbitrary loading history, because it defines the specific
work necessary for the crack growth onset in an elas-
tic-plastic body. The G-integral value is likely to be
meaningful only on condition that tensile (positive)
stresses are acting at the crack tip.
Thus, it can be noted that when solving elasto-plastic
problems with the flow theory for a complex law of load-
ing, the use of the G-integral concept is more justified
than the J-integral concept.
5 FINITE-ELEMENT MESH
Figure 2 illustrates a fragment of the finite-element
mesh being constructed in the vicinity of the crack front.
The mesh extending along the front consists of three
parts – the core, the transition zone and the coarsening
zone. The core size is selected to be small in comparison
with the size of the transition zone (5–10 %). A mesh
that is regular in the plane perpendicular to the crack
front is constructed in the core. The size of the mesh
spacing along the front is adjusted by assigning the
S. V. KOBELSKY: NUMERICAL CALCULATIONS OF STRESS-INTENSITY FACTORS FOR A WWER-1000 REACTOR ...
826 Materiali in tehnologije / Materials and technology 47 (2013) 6, 825–830
Figure 1: Computational model of the WWER-1000 reactor pressure
vessel
Slika 1: Model za izra~une tla~ne posode reaktorja WWER-1000
refinement factor whose value is selected so that, in the
vicinity of the points of interest (e.g., the deepest point
of the front), the ratio of the mesh spacing values along
the front and in the plane directed to it does not exceed
10 × 1 × 1. The fragment of the finite-element mesh
constructed in the crack plane is shown in Figure 3.
The mesh shown in Figure 3a allows obtaining reli-
able data on the fracture parameters for the points lo-
cated along the crack front (point 1 is the deepest point
of the front) but does not allow evaluating their values at
point 2 on the boundary between the base metal and the
cladding. The mesh shown in Figure 3b makes it possi-
ble to obtain reliable data on the values of the stress-in-
tensity factors and the corresponding integrals at points 1
and 2.
Unless otherwise specified, the results below were
obtained using the unclosed crack model.
6 SOLUTION OF THE PROBLEM IN THE
ELASTO-PLASTIC FORMULATION TAKING
INTO ACCOUNT THE LOADING HISTORY
When solving the problem of the elasto-plastic for-
mulation, taking into account the loading history, we an-
alyzed the sensitivity of the SIF values determined with
the following factors:
• the use of the initial RPV model with a built-in crack
and without it;
• the use of different criteria for stopping the iterative
process;
• the use of the finite-element meshes of different den-
sities.
The SIF values presented in all the plots were nor-
malized with the SIF value determined using the G-inte-
gral for the crack with the mesh parameters of 21 μm ×
5.5 μm. Relative temperature T – Tk0 is set along the hor-
izontal axe, where Tk0 is the critical-brittleness tempera-
ture of the metal under the initial conditions.
6.1 Analysis of the influence of the crack presence in
the initial RPV model and the used criteria for
stopping the iteration process
In the method of variable elasticity parameters, the
use of two criteria for stopping the iteration process is
analyzed – based on the discrepancy (2) and the G-inte-
gral value (3):
 = ≤
=
∑
( , )
( , )
r r
r r
i i
j i
j
i r
1
(2)
 = ≤−
−
( , )G G
G
i k i
i k
G (3)
where ri is the discrepancy vector, Gi–k and Gi are the
G-integral values at the i-k-th and i-th step iterations
with respect to the non-linearity, k ( 1 is the parameter,
r = 1 × 10–16, G = 1 × 10–3.
Figure 4 shows the plots of the results for fracture
parameters – CTOD and relative SIF values calculated
using the G-integral for the problems that were solved on
three meshes (1000 μm × 600 μm, 400 μm × 275 μm,
47.5 μm × 27.5 μm) under the following conditions:
• the initial RPV model with a built-in crack, the stop-
ping criterion is (2),
• the initial RPV model without a built-in crack, the
stopping criterion is (3).
The analysis of the results allows one to conclude
that it is unreasonable to build a crack into the initial full
RPV model. The use of criterion (3) instead of criterion
S. V. KOBELSKY: NUMERICAL CALCULATIONS OF STRESS-INTENSITY FACTORS FOR A WWER-1000 REACTOR ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 825–830 827
Figure 3: Finite-element mesh constructed in the crack plane: a) with
an unclosed contour, b) with a closed contour
Slika 3: Mre`a kon~nih elementov, postavljena v ravnini razpoke: a) z
nezaprtim obrisom, b) z zaprtim obrisom
Figure 2: Finite-element mesh constructed in the vicinity of the crack
front
Slika 2: Mre`a kon~nih elementov, konstruirana v bli`ini ~ela razpoke
(2) did not lead to a noticeable difference in the results of
the problem solution (the maximum difference in the SIF
values was found to be less than 0.1 %) but made it pos-
sible to reduce considerably the total time of the solution
(if the time of the problem solution with the use of crite-
rion (2) can be taken to be a unity, then, with the use of
criterion (3), the problem time was 0.55 for the mesh of
1000 μm × 600 μm, 0.4 for the mesh of 400 μm × 275
μm, and 0.7 for the mesh of 47.5 μm × 27.5 μm).
The character of the SIF variation curves (Figure 4b)
shows that the effect of the mesh-spacing value is mini-
mum at the stages of active loading. However, with an
increase in the spacing, the descending branch of the
curve departs increasingly more to the right, showing a
tendency to be vertical in the limit.
Particular attention should be given to the stress be-
havior during the mesh refinement – under unloading, a
compressive stress zone occurs in the vicinity of the
crack tip. The instant of time, at which this zone begins
to be fixed, and its size are essentially dependent on the
spacing value in the vicinity of the crack front. Thus, for
the mesh of 1000 μm ×600 μm, the compressive stresses
did not occur at all, whereas for the mesh of 400 μm ×
275 μm they appeared at the crack tip only, at the time
instant of 6400 s.
For the mesh of 21 μm × 5.5 μm, the last SIF value
was obtained for the time instant of 2000 s, whereas for
the mesh of 1000 μm × 600 μm, the SIF values were ob-
tained up to the time instant of 6 400 s. The plots of the
time dependence of the compressive-stress-zone size in
the vicinity of the crack tip for the meshes of 47.5 μm ×
27.5 μm and 21 μm × 5.5 μm are given in Figure 5.
6.2 Use of meshes of different densities
The solution of the problem with the finite-element
method can be considered to be reliable only upon reach-
ing the convergence. It is common practice to perform an
analysis of the convergence by comparing the solutions
obtained on two or more meshes, whose spacings differ
by a factor of 2 to 5 times. Figure 6 compares the rela-
tive SIF values calculated on the meshes of 1000 μm ×
600 μm, 47.5 μm × 27.5 μm and 21 μm × 5.5 μm using
the J- and G-integral concepts. At the stages of active
loading, the greatest difference in the SIF values is !9 %
for the meshes of 1000 μm × 600 μm and 47.5 μm × 27.5
μm, whereas for the mesh of 21 μm × 5.5 μm this differ-
S. V. KOBELSKY: NUMERICAL CALCULATIONS OF STRESS-INTENSITY FACTORS FOR A WWER-1000 REACTOR ...
828 Materiali in tehnologije / Materials and technology 47 (2013) 6, 825–830
Figure 6: Plots of the relative SIF values calculated on different
meshes using the J- and G-integral concepts: 1 – 1000 μm × 600 μm
(J), 2 – 1000 μm × 600 μm (G), 3 – 47.5 μm × 27.5 μm (J), 4 – 47.5
μm × 27.5 μm (G), 5 – 21 μm × 5.5 μm (J), 6 – 21 μm × 5.5 μm (G)
Slika 6: Prikaz relativnih vrednosti SIF, izra~unanih z razli~nimi mre-
`ami, s konceptom J- in G-integrala: 1 – 1000 μm × 600 μm (J), 2 –
1000 μm × 600 μm (G), 3 – 47,5 μm × 27,5 μm (J), 4 – 47,5 μm × 27,5
μm (G), 5 – 21 μm × 5,5 μm (J), 6 – 21 μm × 5,5 μm (G)
Figure 4: Plots showing the results of calculating the fracture parame-
ters: a) is for the crack-tip opening displacement, b) is for the relative
SIF values: 1 – 1000 μm × 600 μm (1), 2 – 1000 μm × 600 μm (2), 3 –
400 μm × 275 μm (1), 4 – 400 μm × 275 μm (2), 5 – 47.5 μm × 27.5
μm (1), 6 – 47.5 μm × 27.5 μm (2)
Slika 4: Prikaz rezultatov izra~unov parametrov preloma: a) je za
premik konice razpoke, b) je za relativne vrednosti SIF: 1 – 1000 μm ×
600 μm (1), 2 – 1000 μm × 600 μm (2), 3 – 400 μm × 275 μm (1), 4 –
400 μm × 275 μm (2), 5 – 47,5 μm × 27,5 μm (1), 6 – 47,5 μm × 27,5
μm (2)
Figure 5: Plots of the time dependence of the compressive-stress-zone
size in the vicinity of the crack tip:  21 μm × 5.5 μm,  47.5 μm ×
27.5 μm
Slika 5: Prikaz odvisnosti velikosti tla~ne cone na ~elu konice raz-
poke:  21 μm × 5,5 μm,  47,5 μm × 27,5 μm
ence is !3 %. It can be noted that the SIF values deter-
mined using the G-integral concept turn out to be higher
than those determined using the J-integral concepts, and
also that with a decrease in the mesh-spacing value, the
difference between the SIF values calculated with vari-
ous methods decreases, too.
6.3 Determination of the allowable critical-brittleness
temperature of the metal
The data of the SIF calculation for a circumferential
under-the-cladding crack obtained using the J- and
G-integral concepts is applied when determining the
critical-brittleness temperature of the metal. A fracture-
toughness curve is used, whose equation is given in3,9.
Figure 7a shows the plots of the temperature depen-
dence of the relative SIF values KI(T)/Kmax and the
fracture toughness KIc(T)/Kmax for the three meshes (1000
μm × 600 μm, 47 μm × 27.5 μm and 21 μm × 5.5 μm)
obtained using the J-integral concept. On the other hand,
Figure 7b illustrates the plots obtained using the G-in-
tegral concept. Table 1 summarizes the values of incre-
ments in the critical-brittleness temperature of metal
 Tka compared for the meshes of 1000 μm × 600 μm, 47
μm × 27.5 μm and 47 μm × 27.5 μm , 21 μm × 5.5 μm.
On the basis of the analysis of the results it is possible to
draw a conclusion about their convergence when
successively denser meshes are used.
The critical-brittleness temperature values of the
metal obtained using the J-integral concept turned out to
be less conservative than those obtained using the G-in-
tegral concept.
The critical-brittleness temperature value of the metal
increases and becomes determined with lesser conserva-
tism when performing a refinement of a finite-element
mesh.
Table 1: Values of increments in the critical-brittleness temperature of
the metal
Tabela 1: Vrednost korakov kriti~ne temperature krhkosti kovine
Mesh, μm × μm  Tka, J-  Tka, G-
47 × 27.5  1000 × 600 14 °C 14 °C
21 × 5.5  47 × 27.5 5 °C 3 °C
6.4 Analysis of the results obtained using the closed-
crack model
Figure 8 shows the results of determining the SIF
values at a deep point of the crack and at the midpoint of
S. V. KOBELSKY: NUMERICAL CALCULATIONS OF STRESS-INTENSITY FACTORS FOR A WWER-1000 REACTOR ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 825–830 829
Figure 8: Plots of the relative SIF values calculated on different
meshes using the J- and G-integral concepts: a) is for the deep point of
the crack, b) is for the point at the boundary between the base metal
and the cladding: 1 – 1100 μm × 360 μm (J), 2 – 1100 μm × 360 μm
(G), 3 – 100 μm × 100 μm (J), 4 – 100 μm × 100 μm (G), 5 – 12 μm ×
12 μm (J), 6 – 12 μm × 12 μm (G)
Slika 8: Relativne vrednosti SIF, izra~unane z razli~nimi mre`ami z
uporabo koncepta J- in G-integrala: a) je za to~ko v globini razpoke,
b) je za to~ko na meji med osnovnim materialom in oblogo: 1 – 1100
μm × 360 μm (J), 2 – 1100 μm × 360 μm (G), 3 – 100 μm × 100 μm
(J), 4 – 100 μm × 100 μm (G), 5 – 12 μm × 12 μm (J), 6 – 12 μm × 12
μm (G)
Figure 7: Determining the critical-brittleness temperature value of the
metal: a) the J-integral concept; b) the G-integral concept: 1 – 1000
μm × 600 μm, 2 – 47.5 μm × 27.5 μm, 3 – 21 μm × 5.5 μm
Slika 7: Dolo~anje vrednosti kriti~ne temperature krhkosti kovine: a)
koncept J- integrala, b) koncept G- integrala: 1 – 1000 μm × 600 μm, 2
– 47,5 μm × 27,5 μm, 3 – 21 μm × 5,5 μm
the boundary between the base metal and the cladding
with the use of the J- and G-integral concepts. For the
deep point, the results obtained using the G-integral con-
cept, as before, turned out to be more conservative than
those obtained using the J-integral concept. At the same
time, for the point at the boundary between the base
metal and the cladding, the results obtained using the
J-integral concept turned out to be more conservative
than those obtained with the G-integral concept. For this
point, the maximum difference in the SIF values deter-
mined using both concepts was approximately 5 %.
7 CONCLUSIONS
The SIF determination methods based on the use of
the J- and G-integral concepts were analyzed for a typi-
cal regime of the WWER-1000 RPV emergency cooling.
When solving the problem of an elastic and elasto-
plastic formulation without taking into account the load-
ing history, the difference in the SIF values is negligible,
corresponding to the theoretical statements about the
equality of J and G integrals.
The results of the SIF calculation obtained using the
initial full model of the RPV – with a built-in crack and
without it – practically coincide, which testifies to the
uselessness of the crack embedding into the initial full
model of the RPV. However, according to the require-
ments of the IAEA (International Atomic Energy
Agency), a crack is to be built in the RPV fragment un-
der consideration.
It was shown that when solving the problem of the
elasto-plastic formulation on the basis of the flow theory
under non-proportional loading, the SIF values obtained
using J- and G-integral concepts are in agreement even
at the stages of unloading, although the use of the J-inte-
gral concept under these conditions contradicts the the-
ory statements.
The results of the SIF calculation using the G-integral
concept are more conservative than those obtained using
the J-integral concept.
The character of the temperature dependence of the
fracture parameters – the stress at the crack tip and SIF –
at the stages of unloading was shown to be considerably
dependent on the density of the employed finite-element
mesh. With a decrease in the size of the finite element in
the vicinity of the crack front, an increase in the size of
the compressive-stress zone is observed, and the de-
scending branch of the SIF curve tends to become practi-
cally vertical. The analysis of the behavior of the frac-
ture-parameter curves points to the convergence of the
results, which is a necessary condition for solving the
problem with the finite-element method. However, the
use of the temperature dependence of SIF constructed
from the solution of the problem on a dense finite-ele-
ment mesh for the purpose of determining the brittleness
temperature of the metal results in obtaining, as far as we
know, less conservative results than those obtained using
the curve constructed from the problem solution on a
coarse finite-element mesh.
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S. V. KOBELSKY: NUMERICAL CALCULATIONS OF STRESS-INTENSITY FACTORS FOR A WWER-1000 REACTOR ...
830 Materiali in tehnologije / Materials and technology 47 (2013) 6, 825–830
D. STEINER PETROVI^ et al.: DISSOLUTION OF A COPPER WIRE DURING A HOT-DIPPING PROCESS ...
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RAZTAPLJANJE BAKRENE @ICE MED VRO^IM OMAKANJEM
PRI UPORABI SPAJKE BREZ SVINCA SnCu1
Darja Steiner Petrovi~1, Jo`ef Medved2, Grega Klan~nik2,3
1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
2University of Ljubljana, Faculty of Natural Sciences and Engineering, A{ker~eva 12, 1000 Ljubljana, Slovenia
3Institute for Foundry and Heat Treatment, Litostrojska 60, 1000 Ljubljana, Slovenia
darja.steiner@imt.si
Prejem rokopisa – received: 2013-09-04; sprejem za objavo – accepted for publication: 2013-10-28
Thermodynamic arguments and calculations were used to describe a complete dissolution of a copper wire in a SnCu1 lead-free
solder during hot-dipping at 400 °C. For the calculation of the phase diagram the newly reviewed Gibbs energies of the phases
were used. The experimental investigation involved a visual inspection, stereomicroscopy, scanning electron microscopy
(FE-SEM/EDX) and a thermal analysis (DSC). The results showed that the dissolution of the copper wire during hot dipping at
a selected working temperature can be attributed to the increased solubility of the copper in the liquid solder and a prolonged
time of dipping. Thus, the applied temperature was too high for the geometry, the volume-to-surface ratio, of the selected fuse
element. Laboratory simulation tests performed at 303 °C showed a much slower kinetics for the Cu pick-up.
Keywords: lead-free solder, copper, dissolution, hot dipping, thermodynamics
Ne`eleno raztapljanje tanke bakrene `ice v kopeli spajke SnCu1 pri temperaturi 400 °C smo opisali s termodinami~nimi
argumenti in izra~uni. Za izra~un faznega diagrama smo uporabili na novo optimirane Gibbsove energije posameznih faz.
Eksperimentalno delo je obsegalo vizualni pregled, stereomikroskopijo, vrsti~no elektronsko mikroskopijo (FE-SEM/EDX) in
termi~no analizo (DSC). Rezultati so pokazali, da je raztapljanje bakrene `ice med omakanjem na izbrani delovni temperaturi
posledica pove~ane topnosti bakra v talini spajke in predolgega ~asa omakanja za dano geometrijo bakrene `ice, predvsem
velikega razmerja med povr{ino in prostornino. Laboratorijski simulacijski preizkusi, izvedeni pri temperaturi 303 °C, so
pokazali veliko po~asnej{o kinetiko raztapljanja tanke bakrene `ice.
Klju~ne besede: spajke brez svinca, baker, raztapljanje, omakanje, termodinamika
1 INTRODUCTION
Fuses consist of one or more metallic conductors,
known as the fuse element, and usually have a cylin-
drical or flat strip form. The surrounding media are
granular silica quartz (SiO2) in high-breaking-capacity
(HBC) fuses, and boric acid in expansion fuses. The
element and the surroundings are housed in a body of an
insulating material (ceramic, fibre, melamine, etc.). The
fuse element’s edges are usually soldered to, or in an
electric contact with, the fuse end caps.1
Another process used in the electronics industry for
the coating of copper wires is hot dipping. The process is
carried out by immersing a pre-treated (e.g., cleaned,
etched) substrate in a bath of a molten solder metal, or an
alloy, for a specific time.
In recent years a new generation of solders has been
developed in order to replace Pb-Sn solders with
lead-free solder alloys. Many of the proposed alloy
systems are Sn-rich alloys. Thermodynamic properties
play a crucial role in the development of the new solder
materials.2–5 Although a thermodynamic assessment
cannot cover the whole situation relating to kinetic
problems, the driving forces and formation energies of
intermetallic compounds, which are the most important
parameters in the growth kinetics, may be subjected to
the thermodynamics of the system. A precise thermo-
dynamic assessment of a Cu-Sn system was given by
Shim et al.2 New data is not only essential for a new
alloy design but also very important for understanding
the reactions between solder alloys and substrate mate-
rials.
The major factors affecting a solder selection are the
melting point of an alloy, the wetting characteristics, the
cost, the availability, the environmental friendliness, etc.
Reliability-related properties include the electrical and
thermal conductivities, the mechanical strength, the
shear and tensile properties, the fatigue resistance, the
corrosion and oxidation resistance, the coefficient of
thermal expansion and the formation of intermetallic
compounds.6
The three major chemical properties affecting the use
and the long-term reliability of solders6 are as follows:
– the solubility of Cu in the solder,
– the resistance to corrosion,
– the oxidation behaviour.
In the literature a lot of data about the formation of
intermetallic compounds (IMCs) is available. Investiga-
tions of interfacial interactions and the formation of
Materiali in tehnologije / Materials and technology 47 (2013) 6, 831–836 831
UDK 669.3:536.7 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)831(2013)
IMCs involving a solid substrate and liquid soldering
alloys have mainly focused on simple geometries,
holding the solder joint above the melting temperature of
the solder alloy at a single interface and then maintaining
the system under isothermal conditions.7–9
The aim of the present case study was to highlight the
thermodynamic background of an undesirable dissolu-
tion of a fuse element (in this case a copper wire) at the
temperature of interest, in the process of hot dipping
using an Sn-Cu lead-free solder. For the calculation of a
phase diagram, the reviewed Gibbs energies of phases
were modelled using substitutional and stoichiometric
models. The optimized parameters are presented in
Appendix A.
In the investigation, a visual inspection, chemical
analysis, stereomicroscopy, scanning electron micro-
scopy (FE-SEM/EDS) and thermal analysis (DSC) were
carried out. In addition, thermodynamic calculations
using ThermoCalc were performed in order to determine
the failure analysis. The copper dissolution was inve-
stigated by hot dipping the Cu wire into a SnCu1
soldering bath.
2 EXPERIMENTAL WORK
The specimens under investigation were a mass
fraction 99.9 % Cu wire and a lead-free solder declared
as SnCu1, both of which are used in industrial produc-
tion processes for special-purpose fuses (Table 1).
The cross-sections of the materials under investi-
gation were metallographically prepared by grinding and
polishing. FE-SEM/EDX analyses were performed on
these cross-sections using a JEOL JSM 6500-F electron
microscope. An X-Ray fluorescence spectrometer (Ther-
mo Scientific Niton) was used to determine the chemical
composition of the solder alloy.
The thermal analysis was performed using differen-
tial scanning calorimetry (DSC) in an STA-449 C
Jupiter, Netzsch instrument. The DSC experiments were
conducted under the static atmosphere of argon of a
volume fraction 99.999 % purity to minimize the surface
oxidation. The empty corundum crucible was taken as a
reference. The linear temperature program for the
heating was as follows: 25 °C to 280 °C to 25 °C. The
heating rate was taken to be 5 K/min to minimize the
effects relating to the apparatus and hysteresis of the
characteristic temperatures (emphasis was given to a
determination of the liquidus temperature). There was no
isothermal step at the maximum temperature in order to
avoid a possible loss of elements. The liquidus tempe-
rature was estimated assuming that the peak temperature
of the last thermal event during heating represents the
liquidus temperature.10 The solidus temperature was
taken as the onset of the melting during heating. The
construction of the tangents for the solidus temperature
was made on the DSC cooling curve with an extrapola-
tion of the peak slope down to the baseline.
The copper dissolution was investigated by hot
dipping the Cu-wire in a SnCu1 soldering bath. The tests
were performed at 400 °C and 303 °C. For the tempera-
ture control a K-type thermocouple was used, protected
with a corundum tube and immersed into the soldering
bath. The dipping was done in air atmosphere as the
common procedure of preparing the SnCu coatings.
2.1 Thermodynamic model
The thermodynamic calculations for the Cu-Sn
binary system were performed with ThermoCalc Classic
(TCC). Additionally, a computer simulation of the
solidification of the selected solder alloy was performed
with the Scheil-Gulliver model for simulating the solidi-
fication path, knowing that an equilibrium distribution of
the elements is inhibited with the relatively high cooling
rates after the hot-dipping process. The thermodynamic
calculations were done using various models for
describing the Gibbs energies of phases.
2.2 Substitutional model
The disordered solution phases and their Gibbs free
energies, Gm
P , of fcc, liquid and bcc are described with
the following equation:
G X G X G
RT X X X X Gex
m
P
Cu Cu
P
Sn Sn
P
Cu Cu Sn Sn m
P
= + +
+ + +
0 0
( ln ln )
where exGm, XCu, R and T represent the excess molar
Gibbs energy, the molar fraction of Cu, the gas constant
(8.314 J mol–1 K–1) and the temperature (K) represent
the molar Gibbs energy of pure Cu in the P phase
(liquid, fcc, bcc). The Gibbs energy of pure element Cu
in the P phase is given relative to SER (the standard
element reference of 298.15 K).The excess contribution
is modelled as:
ex P P
Z P
G X X L L X X
L X
m
P
Cu Sn CuSn CuSn Cu Sn
CuSn Cu
= + − +
+
( ( )
(
0 1
−X Sn ) )
2
representing a deviation from the ideal solution.
2.3 Stochiometric model
The stoichiometric phases (Cu3Sn, Cu41Sn11, Cu10Sn3
and Cu6Sn5) were modelled using the following formula-
tion of the Gibbs energy with reference to the enthalpies
of pure Cu and Sn in phase  (298.15 K):
G T p H q H
G T p GSERCU Tf
m
Cu Sn
Cu Sn
Cu Sn
p q
p q
( )
( ) (
− − =
= + ⋅
∅ ∅0 0
Δ ) ( )+ ⋅q GHSERSN T
where Δ f G TCu Snp q ( ) represents the standard Gibbs free
energy. Normally, this is represented with the linear
temperature dependence:
Δ f G T a b TCu Snp q ( ) = + ⋅
D. STEINER PETROVI^ et al.: DISSOLUTION OF A COPPER WIRE DURING A HOT-DIPPING PROCESS ...
832 Materiali in tehnologije / Materials and technology 47 (2013) 6, 831–836
where a and b represent the optimized parameters.
GHSERCU and GHSERSN are the Gibbs free energies
of pure elements, Cu and Sn, also relative to the stable
state at 298.15 K. The two-sublattice model can be
represented with (Cu)p(Sn)q. The p and q symbols repre-
sent the atomic ratio (the location of a phase). The D03
phase, with a larger solubility, was modelled with a
two-sublattice model: (Cu,Sn)0.75(Sn,Cu)0.25. More de-
tails are given in2.
3 RESULTS AND DISCUSSION
Copper wire is used in the production of fuses. It has
a widespread use in electronic applications due to its
superior thermal and electrical conductivities. In this
case study a commercial copper coil with a diameter d =
±0.15 mm was investigated (Figure 1).
Problems occurred when, in the industrial process of
hot dipping, the Pb-Sn solder was replaced with a Sn-Cu
lead-free solder (Figure 2). After a few seconds in the
solder bath, declared as SnCu1, heated to T = 400 °C, the
copper wire was completely dissolved.
The DSC melting curve of the solder material was
performed to determine its liquidus temperature, which
gives very important information required for determin-
ing the optimum parameters of hot dipping. The DSC
melting curve obtained at a heating rate of 5 K/min is
presented in Figure 3. The Tonset of the major endo-
thermic reaction was around 223.9 °C. The enthalpy of
this endothermic reaction was 48.15 J/g.
From the DSC response during heating, it is clear
that the chemical composition of the investigated Sn-Cu
alloy is very close to the eutectic composition. For
non-eutectic compositions the melting curve would show
a splitting into two separate peaks due to the sample
containing the Cu solid solution in Sn and the eutectic.11
Additionally, the chemical composition of the solder
material of the Sn-Cu alloy system was identified using
X-Ray fluorescence spectrometry. The chemical compo-
sition of the solder is given in Table 1.
Table 1: Chemical composition of the lead-free Sn-Cu solder in mass
fractions (w/%)
Tabela 1: Kemijska sestava spajke Sn-Cu v masnih dele`ih (w/%)
w(Sn)/% w(Cu)/%
Solder Sn-Cu 99.183 ± 0.648 0.785 ± 0.034
The eutectic Sn-Cu solder alloy is one of the most
popular lead-free alloys used for soldering in electronic
applications.12
D. STEINER PETROVI^ et al.: DISSOLUTION OF A COPPER WIRE DURING A HOT-DIPPING PROCESS ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 831–836 833
Figure 3: DSC heating curve for the solder material (the heating rate
of 5 K/min)
Slika 3: DSC segrevalna krivulja za zlitino spajke (hitrost ogrevanja
5 K/min)
Figure 1: Fuse element (a Cu-coil)
Slika 1: Bakrena `ica za varovalko
Figure 2: SE image and the corresponding X-ray elemental mapping
showing the distribution of Cu and Sn in the solder
Slika 2: SEM-posnetek spajke in porazdelitev elementov Cu in Sn
(EDX)
The ability of a solder bath to pick up copper is
directly related to the solubility of Cu in the major
constituents.6 The dissolution kinetics of Cu in soldering
reactions with Sn-Pb and Sn-Ag solders has already been
described in13,14.
In the present study the consumption of a Cu wire
was investigated during the hot dipping of the wire into
the soldering bath at 400 °C. At this temperature the
soldering bath aggressively dissolves the whole of the Cu
wire on contact. It has been reported that the suitable
temperature to perform the soldering with the eutectic
Sn-0.7Cu alloy is 303 °C, because at this temperature the
solder exhibits the lowest value of the surface tension.12
For this reason, comparison tests were also per-
formed at 303 °C, showing much slower kinetics for the
Cu pick-up. The results of the hot dipping at 303 °C are
given in Table 2.
It was confirmed that a Cu wire with a diameter of
approximately 140 μm to 153 μm is completely dissol-
ved after a few seconds of being dipped in the solder
bath at T = 400 °C. On the other hand, at T = 303 °C, the
average Cu consumption measured after 10 s (Figure 4)
and 20 s, as the difference between the initial diameter
and that after hot dipping, was 30 % and 60 %, respec-
tively.
The problem of the rapid dissolution of a Cu wire in
a liquid solder can also be described using thermo-
dynamic arguments. Using the CALPHAD method
(CALculation of PHAse Diagrams) the phase equilibria
can be calculated with the relative Gibbs free energies of
the phases present in a particular system.15 It is clear
from the Cu-Sn phase diagram in Figure 5 that the solid
solubility of Cu in Sn at room temperature is practically
zero. However, as the temperature increases, the solu-
bility of Cu in Sn increases appreciably. The optimized
parameters taken for the Cu-Sn system2 are presented in
Appendix A.
According to the Scheil-Gulliver model the liquidus
temperature is 228 °C for Sn-0.7Cu (Figure 6). The
calculated solidification process is in good agreement
with the experimental one (Figure 3). With a decreasing
temperature, the solidification of the liquid phase
proceeds with the precipitation of Sn (the bct phase).
Below 220 °C the precipitation of the fcc intermetallic
takes place. The mass-fraction of all the stable phases in
the Sn-0.7Cu system is presented in Figure 7.
D. STEINER PETROVI^ et al.: DISSOLUTION OF A COPPER WIRE DURING A HOT-DIPPING PROCESS ...
834 Materiali in tehnologije / Materials and technology 47 (2013) 6, 831–836
Figure 6: Solidification sequence of the Sn-0.7Cu alloy
Slika 6: Potek strjevanja zlitine Sn-0,7Cu
Figure 5: Binary phase diagram for Cu-Sn
Slika 5: Binarni fazni diagram Cu-Sn
Figure 4: Cu wire after hot dipping in the SnCu1 bath at 303 °C for
10 s
Slika 4: Bakrena `ica po omakanju pri temperaturi 303 °C, 10 s v
kopeli SnCu1
Table 2: Average consumption of the Cu wire  d during hot dipping
at 303 °C in mass fractions
Tabela 2: Povpre~no odtapljanje bakrene `ice  d med omakanjem pri
303 °C v masnih dele`ih
T/°C t/s  d/%
303 10 30
303 20 60
During the soldering operation, materials from the
solid substrate dissolve and mix with the solder, allowing
a formation of intermetallic compounds (IMCs). It has
been found that in Cu-Sn systems the formation of
Cu6Sn5 is usually followed by Cu3Sn.7,8 A compressive
stress is generated as a result of the volume expansion
during the growth of IMCs, especially Cu6Sn5.16,17 The
initial formation of Cu6Sn5, followed by Cu3Sn, can be
attributed to a larger driving force for the precipitation of
Cu6Sn5 during the early stages of the Cu/Sn reaction.
Furthermore, the Cu/Sn interface is typically covered by
a Cu6Sn5 layer within a few milliseconds. After a pro-
longed Cu6Sn5 layer growth, Cu3Sn begins to appear at
the Cu/Cu6Sn5 interface.8,9
The growth of IMCs as a result of the interactions
between the eutectic Sn-0.7Cu solder and the Cu-based
alloy was studied by Dariavach et al.18 A thick scallop
layer of -Cu6Sn5 and a thin layer of -Cu3Sn were
observed at the interface. The total thickness of IMCs
and the grain size of the -phase increased with the
increasing soldering time.18 As the interfacial reaction
occurs between the substrate and the molten solder, the
liquid structure of the solder significantly affects the
reaction and the formation of interfacial compounds.
4 CONCLUSIONS
The thermodynamic assessment of Shim et al.2 was
the basis for the thermodynamic description of a Sn-Cu
system. The optimization was performed using Thermo-
Calc Classic (TCC). The optimized parameters of the
phases of interest in the Sn-Cu system are listed in
Appendix A. The thermodynamic data were used to
describe a complete dissolution of a Cu wire in a SnCu1
lead-free solder during hot dipping at 400 °C.
The melting behaviour as well as the chemical anal-
ysis of the investigated solder declared as SnCu1
confirmed the chemical composition of Sn-Cu to be very
close to the eutectic one. The measured liquidus tempe-
rature was approximately 228 °C.
According to the Sn-Cu binary phase diagram the
solubility of copper in tin increases with the temperature.
A rapid dissolution of the copper wire during hot
dipping at the working temperature T = 400 °C is attri-
buted to the increased solubility of copper in a liquid
solder. The applied temperature was too high for the
geometry of the selected fuse element (0.15-mm thick
copper wire).
The kinetics of the Cu dissolution in the Sn-Cu solder
was much slower at 303 °C. The average copper con-
sumption measured after 10 s and 20 s was approxima-
tely 30 % and 60 %, respectively.
Better results for hot dipping would be obtained by
lowering the temperature of the solder bath so that the
layer of intermetallic compounds (IMCs) would form on
the interface between the solder bath and the copper fuse
element. By lowering the working temperature for hot
dipping, the solubility of copper in the solder would
decrease and the nucleation of IMCs would be ensured.
Appendix A: Optimized parameters for the Sn-Cu
phase diagram2
LIQUID
CONSTITUENTS: CU,SN
G(LIQUID,CU;0)-H298(FCC_A1,CU;0) =
298.15<T< 1358.02:
+12964.736-9.511904*T-5.849E-21*T**7+GHSERCU
1358.02<T< 3200.00: +13495.481-9.922344*T-3.642E+29*T**(-9)
+GHSERCU
G(LIQUID,SN;0)-H298(BCT_A5,SN;0) =
100.00<T< 505.08: +7103.092-14.087767*T+1.47031E-18*T**7
+GHSERSN
505.08<T< 3000.00: +6971.587-13.814382*T+1.2307E+25*T**(-9)
+GHSERSN
L(LIQUID,CU,SN;0) = -9002.8-5.8381*T
L(LIQUID,CU,SN;1) = -20100.4+3.6366*T
L(LIQUID,CU,SN;2) = -10528.4
BCC_A2
2 SUBLATTICES, SITES 1: 3
CONSTITUENTS: CU,SN : VA
G(BCC_A2,CU:VA;0)-H298(FCC_A1,CU;0) = +GCUBCC
G(BCC_A2,SN:VA;0)-H298(BCT_A5,SN;0) = +GSNBCC
L(BCC_A2,CU,SN:VA;0) = -44821.6+51.2164*T
L(BCC_A2,CU,SN:VA;1) = -6876.5-56.4271*T
BCT_A5
CONSTITUENTS: CU,SN
G(BCT_A5,CU;0)-H298(FCC_A1,CU;0) = +GCUBCT
G(BCT_A5,SN;0)-H298(BCT_A5,SN;0) = +GHSERSN
L(BCT_A5,CU,SN;0) = 21000
CU10SN3
2 SUBLATTICES, SITES .769: .231
CONSTITUENTS: CU : SN
G(CU10SN3,CU:SN;0)-0.769 H298(FCC_A1,CU;0)-0.231
H298(BCT_A5,SN;0) =
-6655-1.4483*T+.769*GHSERCU+.231*GHSERSN
CU3SN
D. STEINER PETROVI^ et al.: DISSOLUTION OF A COPPER WIRE DURING A HOT-DIPPING PROCESS ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 831–836 835
Figure 7: Mass fraction of stable phases in the Sn-0.7Cu system
Slika 7: Masni dele` stabilnih faz v sistemu Sn-0,7Cu
2 SUBLATTICES, SITES .75: .25
CONSTITUENTS: CU : SN
G(CU3SN,CU:SN;0)-0.75 H298(FCC_A1,CU;0)-0.25
H298(BCT_A5,SN;0) =
-8194.2-.2043*T+.75*GHSERCU+.25*GHSERSN
CU41SN11
2 SUBLATTICES, SITES .788: .212
CONSTITUENTS: CU : SN
G(CU41SN11,CU:SN;0)-0.788 H298(FCC_A1,CU;0)-0.212
H298(BCT_A5,SN;0) =
-6323.5-1.2808*T+.788*GHSERCU+.212*GHSERSN
CU6SN5
2 SUBLATTICES, SITES .545: .455
CONSTITUENTS: CU : SN
G(CU6SN5,CU:SN;0)-0.545 H298(FCC_A1,CU;0)-0.455
H298(BCT_A5,SN;0) =
-6869.5-.1589*T+.545*GHSERCU+.455*GHSERSN
CU6SN5_L
2 SUBLATTICES, SITES .545: .455
CONSTITUENTS: CU : SN
G(CU6SN5_L,CU:SN;0)-0.545 H298(FCC_A1,CU;0)-0.455
H298(BCT_A5,SN;0) =
-7129.7+.4059*T+.545*GHSERCU+.455*GHSERSN
DO3
2 SUBLATTICES, SITES .75: .25
CONSTITUENTS: CU,SN : CU,SN
G(DO3,CU:CU;0)-H298(FCC_A1,CU;0) = +GCUBCC
G(DO3,SN:CU;0)-0.25 H298(FCC_A1,CU;0)-0.75
H298(BCT_A5,SN;0) =
+116674.85+4.8166*T+.75*GSNBCC+.25*GCUBCC
G(DO3,CU:SN;0)-0.75 H298(FCC_A1,CU;0)-0.25
H298(BCT_A5,SN;0) =
-10029.85+.00285*T+.75*GCUBCC+.25*GSNBCC
G(DO3,SN:SN;0)-H298(BCT_A5,SN;0) = +GSNBCC
L(DO3,CU:CU,SN;0) = -1857.8-2.5311*T
L(DO3,CU:CU,SN;1) = -2.9894*T
L(DO3,CU,SN:SN;0) = +45850-42.2191*T
FCC_A1
2 SUBLATTICES, SITES 1: 1
CONSTITUENTS: CU,SN : VA
G(FCC_A1,CU:VA;0)-H298(FCC_A1,CU;0) = +GHSERCU
G(FCC_A1,SN:VA;0)-H298(BCT_A5,SN;0) = +GSNFCC
L(FCC_A1,CU,SN:VA;0) = -11106.95+2.0791*T
L(FCC_A1,CU,SN:VA;1) = -15718.02+5.92467*T
Acknowledgements
The authors would like to acknowledge ETI, d. d.,
Izlake, Slovenia, for the supply of the specimens. The
work was also supported by the Slovenian Research
Agency (Pr. No. P2-0050 and P2-0344).
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15 H. L. Lukas, S. G. Fries, B. Sundman, Computational Thermodyna-
mics, The Calphad Method, Cambridge University Press, Cambridge
2007
16 A. Skwarek, M. Pluska, J. Ratajczak, A. Czerwinski, K. Witek, D.
Szwagierczak, Analysis of tin whisker growth on lead-free alloys
with Ni presence under thermal shock stress, Mater Sci Eng B, 176
(2011), 352–357
17 A. Skwarek, M. Pluska, A. Czerwinski, K. Witek, Influence of lami-
nate type on tin whisker growth in tin-rich lead-free solder alloys,
Mat Sci Eng B, 177 (2012), 1286–1291
18 N. Dariavach, P. Callahan, J. Liang, R. Fournelle, Intermetallic
growth kinetics for Sn-Ag, Sn-Cu, and Sn-Ag-Cu lead-free solders
on Cu, Ni, and Fe-42Ni substrates, J Electron Mater, 35 (2006),
1581–1591
D. STEINER PETROVI^ et al.: DISSOLUTION OF A COPPER WIRE DURING A HOT-DIPPING PROCESS ...
836 Materiali in tehnologije / Materials and technology 47 (2013) 6, 831–836
P. PE^LIN et al.: STRUCTURAL CHARACTERIZATION OF PLATINUM FOIL ...
STRUCTURAL CHARACTERIZATION OF PLATINUM
FOIL FOR NEURAL STIMULATING ELECTRODES
USING A FOUR-POINT RESISTIVITY-MEASURING
DEVICE
STRUKTURNE LASTNOSTI PLATINSKE FOLIJE ZA ELEKTRODE
PRI STIMULACIJI @IVCEV, UGOTOVLJENE Z NAPRAVO ZA
[TIRITO^KOVNO MERJENJE UPORNOSTI
Polona Pe~lin1, Milan Bizjak2, Marko Pribo{ek3, Matja` Godec4, Samo Ribari~5,
Janez Rozman1
1Centre for Implantable Technology and Sensors, Itis, d. o. o. Ljubljana, Lepi pot 11, 1000 Ljubljana, Slovenia
2Department of Materials and Metallurgy, Faculty of Natural Sciences and Engineering, University of Ljubljana, A{ker~eva cesta 12, 1000
Ljubljana, Slovenia
3Marko Pribo{ek, zasebni raziskovalec, Pot v smre~je 11, 1000 Ljubljana, Slovenia
4Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
5Institute of Pathophysiology, Medical Faculty, University of Ljubljana, Zalo{ka 4, 1000 Ljubljana, Slovenia
janez.rozman@guest.arnes.si
Prejem rokopisa – received: 2013-09-25; sprejem za objavo – accepted for publication: 2013-10-07
In the past few decades, considerable efforts have been devoted to develop neuroprosthetics that interface selectively with the
human nervous system via multi-electrode spiral cuffs, using implantable electronic devices. The objective of this study was to
investigate the structural properties of a cold-rolled platinum foil used to manufacture multi-electrode spiral nerve cuffs. To
achieve this objective, thick cold-rolled platinum foil strips 0.03 mm with 99.99 % purity were used. For this purpose, the strips
were mounted into the sample holder within the furnace of a custom-designed set-up. The resistivity measurements were made
using a 4-point probe technique in which the strips were subjected to dynamic annealing in an argon atmosphere within the
temperature range between room temperature and 900 °C. Finally, the microstructures of the strips, prepared using standard
metallographic techniques, were investigated using light microscopy. In the resistivity measurements, a small change is
observed at !280 °C. This change could be explained as the partial recovery elicited by a decrease of the dislocation density.
Above 500 °C, a significant decrease in the resistivity was recorded, and the decrease reached a maximum at !750 °C. These
results provide a deeper insight into the fabrication of platinum foil to be used in the further development of multi-electrode
neural stimulating spiral cuffs. The most important finding is that the results of the resistivity measurements provide one
criterion for selecting materials, and that the appropriate thermal and mechanical working processes are required to fabricate
stimulating electrodes. These results may make cold-rolled platinum ribbon the best material for the long-term application of
multi-electrode spiral cuffs in SNS.
Keywords: selective nerve stimulation (SNS), platinum electrodes, annealing, recrystallization, microstructure, resistivity, light
microscopy
V zadnjih nekaj desetletjih je bilo pri razvoju elektronskih `iv~nih protez, ki prihajajo preko ve~elektrodnih spiralnih objemk
selektivno v stik z `iv~nim sistemom ~loveka, namenjenih veliko naporov. Cilj raziskave je bil preiskati strukturne lastnosti
hladno valjane platinske folije za izdelavo ve~elektrodnih `iv~nih spiralnih objemk. Za dosego tega cilja so bili uporabljeni
trakci iz 0,03 mm debele hladno valjane platinske folije ~istosti 99,99 %. V ta namen so bili trakci montirani v nosilec vzorcev v
pe~i posebej izdelanega merilnega sistema. Meritve upornosti so bile izvedene z uporabo 4-to~kovne merilne tehnike, pri kateri
so bili trakci dinami~no `arjeni v argonu v temperaturnem obmo~ju med sobno tempreraturo in 900 °C. Na koncu je bila
mikrostruktura trakcev, pripravljenih po standardnih metalografskih tehnikah, preiskana z uporabo svetlobne mikroskopije. Pri
meritvah upornosti je bilo opaziti majhne spremembe pri temperaturi pribli`no 280 °C. Te spremembe je mogo~e pripisati
delnemu okrevanju, ki ga povzro~i zmanj{anje gostote dislokacij. Pri temperaturah nad 500 °C pa je bilo opaziti znaten padec
upornosti, ki je dosegel najve~jo vrednost pri temperaturi pribli`no 750 °C. Dobljeni rezultati omogo~ajo poglobljen vpogled v
izdelavo platinske folije, uporabne pri nadaljnjem razvoju ve~elektrodnih stimulacijskih spiralnih objemk. Najpomembnej{i
sklep je, da rezultati meritev upornosti podajajo enega od meril za izbiro materiala ter primernih procesov mehanske predelave
ter toplotne obdelave folije za izdelavo stimulacijskih elektrod. Na osnovi tovrstnih rezultatov lahko preverimo, ali je ustrezno
hladno valjana platinska folija najbolj{i material za dolgotrajno uporabo pri ve~elektrodnih spiralnih objemkah za selektivno
stimulacijo `ivcev.
Klju~ne besede: selektivna stimulacija `ivca, platinske elektrode, `arjenje, rekristalizacija, mikrostruktura, upornost, svetlobna
mikroskopija
Materiali in tehnologije / Materials and technology 47 (2013) 6, 837–843 837
UDK 611.83:615.84 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 47(6)837(2013)
1 INTRODUCTION
In the past few decades, considerable scientific and
technological efforts have been devoted to developing
neuroprostheses that interface the human nervous system
with implantable electronic devices. Potential applica-
tions include limb prostheses, bladder prostheses,
cochlear and brain-stem auditory prostheses, retinal and
cortical visual prostheses, cortical recording for the
cognitive control of assistive devices, vagus nerve stimu-
lation and deep brain stimulation for essential tremor,
Parkinson’s disease, epilepsy, dystonia, and depression.
However, all the applications require electrodes charac-
terized by high spatial and nerve fibre type selectivity
and low impedance for recording and safe reversible
charge injection for stimulation at the same time. With
this in mind, an understanding of the electrochemical
mechanisms underlying the behaviour of neural stimula-
tion and recording electrodes is important for the deve-
lopment of chronically implanted devices, particularly
those employing large numbers of electrodes.1,2
Since the installation of such a multi-electrode
system onto a peripheral nerve also causes mechanical
issues, the designer must consider and optimize both the
aforementioned issues. To minimize the mechanical
issues such as abrasive or compressive injuries and
fitness difficulties in the application of multi-electrode
systems, flexible substrates and appropriate metal mate-
rials have emerged.
In the area of Functional Electrical Stimulation,
multi-electrode cuffs have been used in peripheral nerves
for stimulation as well as for the recording of an electro-
neurogram for more than 35 years.3–5
In the last two decades, particular attention is being
paid to vagus nerve stimulation, techniques that are to be
used as a method to treat a number of autonomous
nervous-system disorders.
In this regard, selective nerve stimulation (SNS) is
usually delivered from a group of three electrodes (trip-
let) within the spiral cuff installed onto the nerve. The
efficacy of SNS is dependent exclusively on localizing
charge delivery to specific populations of nerve fibres.
Charge delivery, however, is influenced by the electro-
de-tissue interface, where a transduction of charge
carriers from electrons in the metal electrode to ions in
the tissue occurs.
Electrodes for SNS face electrochemically harsh
working conditions. At the same time, it is imperative
that in contact with the nerve tissue, the electrodes must
remain non-toxic and non-reactive with the tissue. That
means while injecting a required amount of charge, the
formation of any harmful reactions should be avoided.6–8
For SNS, different materials that support charge
injection by capacitive and faradaic mechanisms are
available. These materials have certain advantages and
limitations. Among the criteria that must be considered
when choosing the material for electrodes that make
electrical contact with a neural tissue are the mechanical
characteristics of the material.9
The noble metal commonly used as a stimulating
electrode material which is capable of supplying ade-
quate electrical charge to activate neural tissue is pure
platinum.10 As with most metals with a face-centred
cubic structure, platinum is a very ductile metal.11 High-
purity platinum is non-toxic, insoluble in mineral and
organic acids and does not corrode or tarnish. However,
soluble salts are highly toxic by deposition within the
internal organ tissues.
Apart from the chemical inertness, platinum has a
number of physical properties of great value for its use in
the technology of implantable stimulating electrodes.12–18
These include general properties, mechanical properties
and physical properties:
General Properties
Density: 21.45 g/cm3
Thermal Conductivity: 0.716 W/(cm K) at 298.2 K
Electrical Resistivity: 105 n m at 293.2 K
Mechanical Properties
Young's modulus: 168000 MPa
Tensile Strength: 125–165 MPa (annealed at 426.8 K)
Elongation: 30–40 % in 50 mm (annealed at 426.8 K)
Tensile Strength: 205–240 MPa (hard drawn, 50 % cold
worked)
Elongation: 1–3 % in 50 mm (hard drawn, 50 % cold
worked)
Elastic Modulus: 171000 MPa (annealed at 426.8 K) at
293.2 K
Elastic Modulus: 156000 MPa (hard drawn, 50 % cold
worked) at 293.2 K
Yield strength: 38-180 MPa
Vickers hardness: 549 MPa
Physical Properties
Absolute electrode potential: 4.44 ± 0.02 V at 298.2 K
Melting temperature: 1495.1 K
Thermal expansion: 8.8 μm·m–1·K–1 at 298.2 K
In the electrical stimulation of nerves, platinum is
often used in the form of the pure metal. Namely,
impurities and alloying elements may adversely affect
both its mechanical characteristics and its stability
against corrosion in physiological media. This metal
injects charge by both faradaic reactions and double-
layer charging. The relative contribution of each process
depends on the current density and the pulse width,
although the faradaic processes predominate under most
neural stimulation conditions.10,19 However, in a SNS with
a high charge density, pH shifts causing irreversible
changes in tissue proteins, metallic dissolution products,
gross hydrogen and oxygen gas bubbles, and oxidized
organic and inorganic species, could occur.20 It is of a
crucial importance that the electrodes in SNS function
without facing degradation over a prolonged time period.
P. PE^LIN et al.: STRUCTURAL CHARACTERIZATION OF PLATINUM FOIL ...
838 Materiali in tehnologije / Materials and technology 47 (2013) 6, 837–843
Although the charge-injection limits of Pt are based on
avoiding the electrolysis of water, Pt dissolution can
occur at lower charge densities.21
The principal approach to controlling the interface
voltage in SNS has been the use of charge-balanced
biphasic stimuli, having cathodic and anodic phases that
contain equal but opposite charge. However, even with
charge-balanced stimulating pulses it is possible that the
interface voltage may reach levels where electrochemical
reactions can occur.22,23
Mechanical fitness is another important issue in
multi-electrode systems design. Low strength, typical of
a platinum of high purity (99.93 %), is accepted in
stimulating and recording electrodes, despite being a
significant disadvantage. To optimise the parameters for
the thermal and mechanical processing of platinum a
knowledge of its rheological characteristics, including
deformation resistance, is required.24–26 The deformation
resistance, however, is considered in terms of the uni-
axial compression or tension of the sample under condi-
tions of plastic deformation, resulting in lattice defects.
With this regard, the cold strength depends on the degree
to which the object has been deformed: the more lattice
defects the harder the metal. Grain boundaries disrupt
the motion of dislocations through a material, so reduc-
ing crystallite size is a common way to improve the
strength. It was assumed that during cold deformation,
the deformation resistance depends only on the geome-
tric parameters of the change in shape. When a metal is
annealed, however, it loses this strength via the reaction
known as recrystallization. The annealing temperature
range for 99.93 % platinum is 400–1000 °C, and
depends on the degree of cold-working.27–29
Resistivity, sometimes called the Specific Resistance,
is one of the fundamental electrical properties of a parti-
cular material which gives it electrical resistance.30
Resistance can be calculated by using a knowledge of
both the resistivity of the material of a particular object
and the shape and geometry of the object. Accordingly,
the electric resistance of a square sheet is independent of
the size of the square, but depends only on the sheet
resistivity and the thickness of the sheet.
Characterizing the resistivity of a material for the
SNS electrode could help to determine the long-term
charge delivery to the nerve without harmful effects for
both the neural tissue and the electrode.
It has been common knowledge that, in general, the
resistivity of a metal is higher when cold-worked than in
an annealed state.31,32 Namely, grain boundaries are
defects in the crystal structure and they tend to decrease
the electrical conductivity of the material. It was found
that platinum of the highest purity almost completely
recovers its electrical properties at a temperature as low
as 300 °C, although a slight further recovery occurs with
an increase of the annealing temperature up to 1450 °C.33
The result of annealing is the diffusion of dislocations
and their cancellation through encounters with disloca-
tions of the opposite sign. In this regard, the increase in
hardness and resistivity due to the strain condition
obtained during cold-rolling could be reduced to zero by
annealing within the temperature range between 400 °C
and 800 °C.
Improved, implantable, multi-electrode stimulating
systems with greater strength, desirable for an individual
electrode, may be achieved through plastic deformation
during cold-rolling and partial annealing by carefully
choosing the appropriate annealing regimes.
In order to set up the appropriate working cycles,
metallographic analysis revealing the microstructure of
platinum after each cycle should be used. The most con-
venient way to reveal the microstructure of cold-rolled
and/or thermally treated platinum foil used for stimu-
lating electrodes is by electrolytic etching.34,35 Namely,
work hardening and thermal treatments cause significant
changes to the microstructure of the platinum with
respect to the as-cast condition.
The objective of this study was to investigate the tem-
perature-dependent structural properties of a cold-rolled
platinum foil used to manufacture stimulating electrodes
within the spiral nerve cuffs for SNS using a custom-
designed four-point resistivity-measuring device.
2 METHODS
The foil was obtained by means of cold-rolling of the
sheet material. In the most labour-intensive method of
foil preparation, a small piece of the ingot (chemical
composition of the ingot: Pt 99.99 %, Rh 0.01 %) was
extracted and, after initial working cycles, the material
was annealed and subsequently rolled flat on a mill to
obtain a thick foil 0.15 mm. Because platinum has a very
high stiffness, work hardening during working cycles
made the foil appear excessively hard. In addition, the
platinum foil experienced severe local hardening arising
from the high strain rates during cold-rolling. We pre-
sumed that a gradual increase in the dislocation density
P. PE^LIN et al.: STRUCTURAL CHARACTERIZATION OF PLATINUM FOIL ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 837–843 839
Figure 1: The 99-electrode spiral cuff and an arbitrarily chosen longi-
tudinal row of electrodes
Slika 1: Prikaz 99-elektrodne spiralne man{ete in poljubno izbrana
vzdol`na vrsta elektrod
occurred, causing increased hardness and a localized loss
of ductility. To reduce the dislocation density, renew the
softness and ductility, and provide a correct orientation
of the metal grains for further cold rolling, recrystalliza-
tion was promoted by subsequent annealing at 700–800
°C. Finally, the foil was rolled to achieve the final
thickness of 0.03 mm on a mill with diameter rollers 300
mm, using varying degrees of cold-working.
The long-term electrochemical stability of platinum
stimulating electrodes within the multi-electrode spiral
nerve cuff, shown schematically in Figure 1, may be
connected with the resistivity of the grain boundaries and
its change with time.
We measured the electrical resistivity of the platinum
strips currently used to manufacture the electrodes used
in the 99-electrode spiral nerve cuff. For this purpose,
test strips were cut from the foil, with their long axes
oriented along the rolling direction. The resistivity
measurements were made with a modified 4-probe (also
called Kelvin probe) resistance-measurement technique.
For this purpose, a high-temperature resistivity
measurement device shown schematically in Figure 2,
designed by a research team at the Department of
Materials and Metallurgy, Faculty of Natural Sciences
and Engineering, University of Ljubljana and Marko Pri-
bo{ek, an individual investigator, was used.
The main parts of the measurement device are a
custom-designed furnace that could withstand a tempe-
rature up to 900 °C and the sample holder. To maintain a
reliable contact at the sample holder during the measure-
ment, the strips were contacted by four spring-loaded
platinum probes.
Direct current was supplied via two probes (source
and sink) that were connected to the current source
(Model: 6220, Keithley Instruments, Inc., Cleveland,
Ohio, USA), and the voltage was measured between the
other two probes, connected to the nanovoltmeter
(Model: 2182A, Keithley, Keithley Instruments, Inc.,
Cleveland, Ohio, USA36; Model 6220 DC Current
Source and Model 6221 AC and DC Current Source37;
Model 2182 and 2182A Nanovoltmeter38).
The complete measuring device and the sample
holder housing four spring-loaded platinum probes are
shown in Figure 3.
To maintain a linear pattern of heating, the system
included a digital proportional-integral-derivative (PID)
controller, which uses the principle of a “continuous”
furnace controller. The PID controller was interfaced
with a computer via an RS232 port. A thermocouple
(K-type) was used to measure the furnace temperature,
while a Pt-100 sensor was used to measure the sample
temperature. The data on temperature were retrieved
using a USB data-acquisition module (Model: NI-9219,
National Instruments, Corporation, Austin, Texas,
U.S.A.) and a USB data-acquisition chassis (Model: NI
cDAQ-9174, National Instruments Corporation, Austin,
Texas, U.S.A.). With this configuration, the test current
was forced through the strip resistance using one set of
test leads, while the voltage across the sample was
measured using a second set of leads, called the sense
leads. Although some small current may flow through
P. PE^LIN et al.: STRUCTURAL CHARACTERIZATION OF PLATINUM FOIL ...
840 Materiali in tehnologije / Materials and technology 47 (2013) 6, 837–843
Figure 3: Complete measuring device: a) measuring and computer
unit, b) furnace, c) four spring-loaded platinum probes (detail), d)
vacuum unit
Slika 3: Celotna merilna naprava: a) merilna in ra~unalni{ka enota, b)
pe~, c) platinasta sonda s {tirimi vzmetmi (detajl), d) vakuumska enota
Figure 2: Schematic diagram of the complete measuring device
Slika 2: Shematski prikaz kompleta merilne opreme
Legend to the figure: RE- Recipient, IS- Isolation, CR- Ceramic retort, WC-
Water cooling, HT- Heater, SA- Sample, SS- Sample support, P1, P2, P3, P4-
Spring-loaded measuring points or contacts, EF- Electric (cable) feed through,
TC- Thermocouple, RP- Rotary pump, TP- Turbomolecular pump, HV- High-
vacuum valve, AV- Argon dosing valve, PI 1- Pirani vacuum gauge, PI 2- Pirani
vacuum gauge, PE- Penning vacuum gauge, MV- Membrane vacuum gauge
the sense leads, it is usually negligible (typically pA or
less) and can generally be ignored for all practical
purposes. Because the voltage drop across the sense
leads was negligible, the voltage measured by the nano-
voltmeter was considered to be essentially the same as
the voltage across the resistance. Both instruments
worked in tandem; the intercommunication was realized
via a Trigger Link (TL) and an RS232 serial port inter-
faced with the computer using a data bus IEE488
(GPIB), an interface (Model: NI-GPIB-USB-HS, Natio-
nal Instruments Corporation, Austin, Texas, USA), and
Lab-view software (Version 10.0, National Instruments,
National Instruments Corporation, Austin, Texas, USA).
Platinum strips were subjected to annealing and, at
the same time, the electrical resistance of the grain boun-
daries was measured using a direct method.30 For the
measuring conditions, the strips underwent annealing in
an argon atmosphere using a linearly increasing tempe-
rature range from room temperature to approximately
900 °C, so as to achieve recrystallization. The heating
rate was 5 K/min. This temperature range was initially
estimated to be high enough to ensure complete recry-
stallization.24 Under these conditions the temperature
distribution during annealing was uniform along the strip
in the argon atmosphere. In the measurements, the strips
were so placed that the grain boundaries were oriented
predominantly parallel to the length of the strip. After-
wards, at arbitrarily chosen points on the strip, platinum
probes connected to the potential leads were mechani-
cally pressed to the strip. An electrical measuring current
of approximately 2–3 A/mm2 was then passed through
the strip and a potential difference of approximately 10
μV was measured using the aforementioned nanovolt-
meter, having an input resistance of 50 M.
To reveal the changes in the microstructure that
occurred after the cold-roll hardening and the recrystalli-
zation thermal treatments of the strips, optical metallo-
graphy was employed. To obtain this goal, the most
convenient way, namely electrolytic etching, was
used.34,35 The preparation of the metallographic speci-
mens consisted of the following four steps: sectioning
the platinum foil into the strips, embedding the strips
into the epoxy resin, mechanical grinding and polishing
of the specimens, electro-polishing of the exposed
metallographic section in a controlled-atmosphere glove
box, and electrochemical etching for microstructure
detection.
For the electrolytic etching, the polished specimens
were immersed into the most widely used electrolytic
solution, i.e., a saturated solution of sodium chloride in
concentrated hydrochloric acid (100 cm3 HCl (37 %) +
10 g NaCl electrolytic). Afterwards, the samples were
exposed to electrolytic etching with an AC power supply
(3–6 V). Because of the "forced corrosion", brought on
by the etching, the specimen surface showed inhomoge-
neous corrosion, and different microstructural features,
which varied from area to area, were revealed.
3 RESULTS
The results of the resistivity measurements are
plotted in Figure 4. Because the plot was difficult to
interpret, a transformation of the plot was performed, in
which the first derivative of the resistance-change plot
was used to identify the points where the resistance
change was most apparent.
It can be seen in Figure 4 that immediately upon
heating from room temperature, the electrical resistance
of the strips linearly increased with increasing tempera-
ture. The starting conditions were as follows: tempera-
ture 27 °C and resistivity 20.846 × 10–3 . At approxi-
mately 280 °C, a small decrease in the resistivity was
observed, which could have been the result of a decrease
in the dislocation density as part of the recovery.
Nevertheless, above 500 °C, a larger decrease in resisti-
vity was recorded, and a maximum decrease was attained
at approximately 750 °C. This is consistent with the
recrystallization trend (dislocation-free grains), in which
the corresponding maximum peak was determined by
differential scanning calorimetry (not shown in this
paper).
P. PE^LIN et al.: STRUCTURAL CHARACTERIZATION OF PLATINUM FOIL ...
Materiali in tehnologije / Materials and technology 47 (2013) 6, 837–843 841
Figure 5: Microstructure: a) deformed and b) recrystallized
Slika 5: Mikrostruktura: a) deformirano, b) rekristalizirano
Figure 4: A plot of the resistivity changes during the recrystallization
versus the temperature
Slika 4: Diagram spreminjanja upornosti med rekristalizacijo v odvis-
nosti od temperature
The recrystallization is shown to depend on the
stored energy of the cold-worked foil, the nature of the
nucleation sites, and the pinning of the boundaries.
Furthermore, knowing that the hardness of platinum and
platinum alloys decreases with the increasing tempera-
ture and time of annealing, a rapid decrease in the
hardness is expected in the temperature range where the
recrystallization takes place.39 According to Loginov et
al.,24 with a decrease in the initial deformation, the
annealing point shifts towards higher temperatures.
According to a comprehensive study about the work-
hardening, recovery, and recrystallization of three grades
of platinum, reported by Raub (1964),40 pure platinum
was recovered after heavy cold-working at approxima-
tely 200 °C. The temperature ranges for the recovery and
recrystallization for platinum of different grades, how-
ever, were found to increase with increasing impurity
content.
An optical micrograph of cold-rolled and recrystalli-
zed strips is shown in Figures 5a and b, respectively.
The etched specimens showed a selectively corroded
surface as a consequence of different grain orientations,
crystal defects such as dislocations and grain boundaries,
and cold-worked regions. The grains appear as elongated
particles when the differential interference-contrast
technique was used. A finer structure is observed in the
cold-rolled sample (Figure 5a). After a short period of
recrystallization, slightly larger grains are observed
(Figure 5b). The size of the grains depends not only on
the temperature used but also on the annealing time. In
this study, the relatively small increase of grain size after
annealing confirms the benefit of using a low annealing
temperature and a short annealing time.
4 DISCUSSION
The present work aims at giving some explanations
about the microstructure of platinum foil to be used in
the fabrication of electrodes for SNS, in both the cold-
rolled hardened and annealed conditions. More specifi-
cally, this paper discusses an analysis, the design criteria,
and the structural properties of the platinum ribbon to be
used for the fabrication of capacitive as well as faradaic
recording and stimulating electrodes within the multi-
electrode spiral nerve cuff for SNS.9 Namely, it is
important to understand all the problems that may occur
when using an electrode for SNS and the situations in
which a stimulating electrode poses the greatest risk to
the nerve.6 Only by combining the knowledge about
appropriate stimulation parameters and the thermal
treatments of cold-worked platinum foil, could result in
the optimized fabrication of SNS electrodes used in
clinical studies. Namely, it is well known from the litera-
ture that significant chemical reactions may occur during
stimulation that can destroy a platinum stimulating elec-
trode.41 Decomposition of the stimulating electrode is a
far more significant problem than just the loss of the
electrode itself. It can result in the deposition of metal
ions in the body of the patient.
5 CONCLUSION
To address some of the key issues currently inhibiting
the widespread adoption of implanted stimulation tech-
nologies, research in the design and characterization of
metallic materials for implantable multi-electrode
systems should be intensified. To achieve suitable final
properties of the platinum foil for the fabrication of
stimulating electrodes, it is crucial to establish the right
combination of plastic deformation and annealing treat-
ment. From these experiments, we concluded that the
deployed measurement technique is adequately sensitive
to the resistivity changes due to the cold-rolling so that
the temperature of recrystallization could be determined.
Besides the aforementioned techniques used, however, a
wide variety of analytical techniques such as Auger and
SEM, could be used alongside it to provide a far more
complete knowledge of the microstructure. In this
regard, the direction of our further work would involve
quantitative and qualitative information about the physi-
cal and chemical changes during the cold deformation of
a platinum foil that would be obtained using the thermo-
analytical technique of differential scanning calorimetry.
Furthermore, the recrystallization kinetics of cold-rolled
platinum foil containing an oriented grain structure will
be studied.
In summary, this study utilized electrical resistivity
measurements to characterize a material with the
potential for use in the production of neural stimulation
electrodes.
Acknowledgements
This work was supported by the Ministry of Higher
Education and Science, Republic of Slovenia, research
programme P3-0171. Additional support was provided
by the ITIS d. o. o. Ljubljana, Centre for Implantable
Technology and Sensors. The research was also
performed as partial fulfilment of the requirements for
the Doctor of Philosophy Degree (Polona Pe~lin, MD).
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Materiali in tehnologije / Materials and technology 47 (2013) 6, 837–843 843

A. HUNYEK, C. SIRISATHITKUL: VARIATION IN MAGNETIC PROPERTIES OF SOL-GEL-SYNTHESIZED ...
VARIATION IN MAGNETIC PROPERTIES OF
SOL-GEL-SYNTHESIZED COBALT FERRITES
SPREMINJANJE MAGNETNIH LASTNOSTI KOBALTOVIH
FERITOV, SINTETIZIRANIH PO SOL-GEL POSTOPKU
Anuchit Hunyek, Chitnarong Sirisathitkul
Molecular Technology Research Unit, School of Science, Walailak University, 80161 Nakon Si Thammarat, Thailand
schitnar@wu.ac.th
Prejem rokopisa – received: 2012-11-28; sprejem za objavo – accepted for publication: 2013-04-10
Single-phase cobalt ferrite (CoFe2O4) was synthesized with a sol-gel reaction between 10.59 g of cobalt nitrate and 29.40 g of
iron nitrate using 480 mL of 10 % polyvinyl alcohol (PVA) in water. Their magnetic squareness and coercive field were slightly
reduced with the increase in the annealing at 800 °C from 2 h to 6 h. To examine the reproducibility of ferrite products of the
PVA sol-gel method, the synthetic condition was repeated 28 times. After annealing for 4 h, CoFe2O4 samples from different
batches exhibit variations in the magnetic properties. The coercive field has a large distribution because of its sensitivity to the
particle agglomeration, whereas the squareness fluctuates in a narrower range from 0.22 to 0.29. Since the squareness is the
ratio of the remanence to the saturation magnetization, the particle agglomeration tends to increase both values while keeping
this ratio rather unchanged.
Keywords: cobalt ferrites, PVA sol-gel, annealing time, magnetic squareness, coercive field
Enofazni kobaltov ferit (CoFe2O4) je bil sintetiziran s sol-gel reakcijo med 10,59 g kobaltovega nitrata in 29,40 g `elezovega
nitrata z uporabo 480 mL 10-odstotne raztopine polivinil alkohola (PVA) v vodi. Njihova ~etverokotna magnetna oblika in
koercitivno polje sta se rahlo zmanj{ala s podalj{anjem `arjenja na 800 °C iz 2 h na 6 h. Da bi preiskali ponovljivost feritnih
proizvodov iz PVA sol-gel metode, je bila sinteza ponovljena 28-krat. Po `arjenju 4 h so kazali vzorci CoFe2O4 iz razli~nih serij
razliko v magnetnih lastnostih. Koercitivno polje ka`e velik raztros zaradi ob~utljivosti za aglomeracijo delcev, medtem ko se
~etverokotnost oblike spreminja v o`jem podro~ju od 0,22 do 0,29. Ker je ~etverokotnost razmerje med remanenco in nasi~eno
magnetizacijo, aglomeracija delcev pove~uje obe vrednosti, medtem ko razmerje ostaja skoraj nespremenjeno.
Klju~ne besede: kobaltov ferit, PVA sol-gel, ~as `arjenja, magnetna ~etverokotna oblika, koercitivno polje
1 INTRODUCTION
Cobalt ferrite (CoFe2O4) is under research and deve-
lopment for its applications in magneto-optical recording
media, magnetic refrigerants, microwave absorbers and
stress sensors.1,2 Bulk CoFe2O4 has an inverse spinel
structure consisting of a cubic close-packed (fcc)
arrangement of oxide anions, O2–. The tetrahedral and
octahedral interstitial sites in the lattice are partially
occupied by the Co2+ and Fe3+ cations, respectively.1,2
Like the other spinel ferrites (e.g., NiFe2O4, MnFe2O4,
ZnFe2O4, CuFe2O4) and barium ferrites, CoFe2O4 nano-
particles can be synthesized with the sol-gel method.3–17
In a sol-gel reaction, homogeneous CoFe2O4 nanoparti-
cles are produced in-situ with a controlled decompo-
sition of precursors, while the chelating gel is dried
during the heat treatment at a relatively low temperature.
Gatelyte et al.7 demonstrated that the characteristics of
CoFe2O4, NiFe2O4, ZnFe2O4, YFeO3 and Y3Fe5O12 pro-
duced with aqueous sol-gel reactions under the same
synthesis condition were comparable.
According to our previous sol-gel synthesis using
polyvinyl alcohol (PVA) as the chelating agent,8 the
magnetic properties of CoFe2O4 are influenced by PVA
contents for two reasons. Firstly, the single-phase
CoFe2O4 obtained in the case of sufficient PVA gives rise
to a relatively high coercive field because of its strong
magnetic anisotropy. In the case of a diluted PVA solu-
tion, the second phase (e.g., -Fe2O3) may also be pre-
sent, reducing the overall coercive field. The other
reason is the dependence of the cluster size on the PVA
contents used in the sol-gel reaction. The cluster size, on
the other hand, dictates the coercive field and magneti-
zation of CoFe2O4.9,15 The role of the PVA gel is to
cleave atoms during the reaction. An increase in PVA
improves the crystallinity of CoFe2O4 without an
excessive particle agglomeration.
The temperature of the heat treatment of the as-syn-
thesized products also affects the magnetic properties of
CoFe2O4. Our previous result agrees with the other
experimental works indicating that both the coercive
field and the magnetic squareness (the ratio of the
remanence to the saturation magnetization) of CoFe2O4
are initially increased with an increase in the annealing
temperature up to 700–800 °C and then reduced with
further increases in the temperature.8,15–19 The growth and
agglomeration of CoFe2O4 particles at high temperatures
result in a decrease in the coercive field and an increase
in the saturation magnetization. The squareness is,
therefore, reduced due to this high-temperature regime.
In addition, Toksha et al.15 reported a decrease in the
coercive field and an increase in the saturation magne-
Materiali in tehnologije / Materials and technology 47 (2013) 6, 845–848 845
UDK 537.622:549.73 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 47(6)845(2013)
tization due to an increased particle size after prolonged
annealing. Recently, Sajjia et al.14 employed the response
surface methodology to simulate the optimum heat
treatment to obtain single-phase CoFe2O4 and minimize
the cost of electricity.
To implement a sol-gel synthesis of ferrites on a
commercial scale, another aspect of concern is the
reproducibility of the products from different batches.
Although this issue is hardly addressed in research
papers, it is particularly important since ferrites of the
order of a hundred grams or less are usually obtained
with a laboratory-scale sol-gel synthesis. In this research,
the condition for preparing the single-phase CoFe2O4
from the previous work8 is reproduced in order to study
the variation in magnetic properties observed after a
repeated synthesis. The effect of the annealing time is
also investigated.
2 EXPERMENTAL WORK
The PVA solution was prepared by dissolving PVA
powders in distilled water (10 %). It was heated at 70–80
°C until the solution became clear (5–7 h). Then, 10.59 g
of cobalt nitrate, Co(NO3)2 · 6H2O, and 29.40 g of iron
nitrate, Fe(NO3)3 · 9H2O powders were mixed with the
480 mL PVA solution. The reaction mixture was stirred
for a further 3 h. After that, it was heated at 80 °C for
10–12 h or until the gel was dry.
The morphology of the ferrite product was shown
with scanning electron microscopy (SEM) and its struc-
ture was characterized with powder X-ray diffractometry
(XRD). A copper target was used as an X-ray source
(K,  = 0.154058 nm) with 40 kV between the cathode
and the copper target. The measurement was performed
in the range of 2
 angles from 10° to 80° with each
rotating step accounting for 0.02°. The crystallite size
was calculated from the peak width using the Scherrer’s
formula:
d =
0 9.
cos

 

(1)
where  is the broadening of the diffraction line
measured at half the maximum intensity and  is the
wavelength of K.
To investigate the effect of the annealing time and
reproducibility of sol-gel products, the time of annealing
at 800 °C varied between (2, 4, 6) h and the same
synthesis condition was repeated for 28 batches with the
same, annealing 4 h. The magnetic properties of all the
samples were characterized with vibrating sample
magnetometry (VSM).
3 RESULTS AND DISCUSSION
The XRD pattern in Figure 1 shows the cubic spinel
CoFe2O4 phase [JCPDS 22-1086] with the peaks at
30.0°, 35.4°, 43.0°, 56.9°, 62.6° corresponding to the
(220), (311), (400), (511) and (440) planes, respectively.
Other oxides or impurity phases are not detected and the
CoFe2O4 peaks are particularly sharp. The crystallite
size, calculated from the line broadening of the (311)
diffraction peak is approximately 45 nm. These crystalli-
tes tend to agglomerate into microscale clusters that are
quite common for sol-gel-synthesized ferrites.20 As seen
in the SEM micrograph in the inset, these aggregates
have flat surfaces and sharp edges. Compared to the
ferrite products from the previous work,8 the cluster size
tends to increase with a reduction of PVA in the sol-gel
reaction.
According to VSM hysteresis loops in Figure 2 all
the CoFe2O4 samples after varying annealing times
exhibit ferrimagnetic properties. The absence of super-
paramagnetic behaviors is due to the particle agglome-
A. HUNYEK, C. SIRISATHITKUL: VARIATION IN MAGNETIC PROPERTIES OF SOL-GEL-SYNTHESIZED ...
846 Materiali in tehnologije / Materials and technology 47 (2013) 6, 845–848
Figure 2: Hysteresis loops of CoFe2O4 from the sol-gel reaction after
annealing at 800 °C for (2, 4 and 6) h
Slika 2: Histerezna zanka CoFe2O4, izdelanega po sol-gel reakciji, po
`arjenju (2, 4 in 6) h na 800 °C
Figure 1: XRD pattern of sol-gel-synthesized ferrite products. Its
SEM micrograph is shown in the inset.
Slika 1: XRD-posnetek feritnega proizvoda, sintetiziranega po sol-gel
postopku. SEM-posnetek je vlo`en.
ration. In the applied magnetic field of 440 kA/m, the
magnetization does not reach the complete saturation.
The highest magnetic moment per mass is around 60
A m2/kg, a value comparable to that of sol-gel-synthe-
sized CoFe2O4 nanoparticles.9–11,15 Both the coercive field
and the squareness are slightly decreased with the
increase in the annealing time from 2 h to 6 h. Longer
annealing times apparently have a similar effect on the
increase in the annealing temperature beyond 800 °C and
the cluster size is, consequently, enhanced. It follows that
the coercive field and the squareness are reduced.
The variation in the coercive field of CoFe2O4 is
shown in the form of a histogram in Figure 3. Ranging
from 31.2 kA/m to 63.2 kA/m, the distribution is appro-
ximately exponential with the coercive field of less than
40 kA/m in the majority of the samples. This implies that
most particles are unevenly agglomerated, like those in
Figure 1. The average coercive field of 28 batches is
(39.8 ± 8.7) kA/m (compared to 34.4 kA/m in Figure 2).
If an enhanced coercive field is aimed for, either an
increase in the PVA contents or a decrease in the anneal-
ing temperature and time is recommended. In Figure 4,
the squareness exhibits a narrower distribution, from
0.22 to 0.29. In addition to its effect on the coercive
field, the particle agglomeration also results in an
increase in the magnetization. However, the squareness
does not exhibit a large variation, because it is the ratio
between the magnetizations. With the average value of
0.26 ± 0.03, a large fraction of the samples have a
squareness beyond 0.25. Interestingly, the minimum
squareness of 0.22 is obtained from the batch with the
smallest cluster size, signified by the maximum coercive
field.
4 CONCLUSION
The magnetic squareness and coercive field of the
single-phase CoFe2O4 synthesized with a sol-gel reaction
between Co(NO3)2 · 6H2O and Fe(NO3)3 · 9H2O exhibit a
slight decrease with the increase in the annealing time
from 2 h to 6 h at 800 °C. However, repeated syntheses
show that the magnetic properties of CoFe2O4 obtained
under the same synthetic and annealing conditions may
vary due to non-uniform particle agglomerations. This
must be taken into the account before implementing the
sol-gel-synthesized ferrites on a large scale.
Acknowledgements
This work was funded by the Industry/University
Cooperative Research Center (I/UCRC) in HDD Compo-
nent, the Faculty of Engineering, the Khon Kaen
University and the National Electronics and Computer
Technology Center, the National Science and Techno-
logy Development Agency. The characterizations using
XRD, SEM and VSM were carried out at the National
Metal and Materials Technology Center (MTEC), the
Chiang Mai University and the Kasetsart University,
respectively.
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H. L. LIM et al.: ASSESSING THE QUALITY OF CORROSION EDUCATION IN THE UNITED ARAB EMIRATES
ASSESSING THE QUALITY OF CORROSION
EDUCATION IN THE UNITED ARAB EMIRATES
PRESOJA KVALITETE IZOBRA@EVANJA NA PODRO^JU
KOROZIJE V ZDRU@ENIH ARABSKIH EMIRATIH
Hwee Ling Lim, Lyas Al Tayeb Ait Tayeb, Mohamed Abdulrahman Othman
The Petroleum Institute, P.O. Box 2533, Sas Al Nakheel Campus, Abu Dhabi, United Arab Emirates
hlim@pi.ac.ae
Prejem rokopisa – received: 2012-08-27; sprejem za objavo – accepted for publication: 2013-03-27
The economic consequences of corrosion failure can be minimized by an engineering workforce well trained in corrosion
fundamentals and management. The United Arab Emirates (UAE) incurs the second highest annual cost of corrosion after Saudi
Arabia, given its large petroleum industry. Hence, this study examined the quality of corrosion education in engineering
programs of the universities in the UAE. Using a single-embedded case-study design, academia and industry respondents were
surveyed on the competence of engineering undergraduates/graduates in corrosion. The findings showed that the dedicated
corrosion courses and engineering courses integrating corrosion into the curricula are available in the UAE universities.
Regarding the competence of engineering students/graduates, the consensus view was that there was an insufficient fundamental
knowledge of corrosion engineering. The industry respondents were highly critical, believing that graduate engineers had a
superficial understanding of corrosion in real-life design contexts. The effectiveness of engineering curricula in corrosion is
determined by both the competence in corrosion knowledge/skills and the availability of resources (qualified staff and new
knowledge from research) to support corrosion education. The findings showed that most departments would not hire new
corrosion specialists. However, the aspect of research was more encouraging with the universities reporting availability of
departmental research and industry partnerships in corrosion research. This paper gives recommendations for improving the
knowledge and skills of future engineers in corrosion management and for enhancing corrosion training to better meet the
industry needs.
Keywords: corrosion, engineering education, case study, United Arab Emirates
Ekonomske posledice korozijskih po{kodb je mogo~e zmanj{ati z in`enirji, ki so dobro izobra`eni na podro~ju osnov korozije in
upravljanja s korozijo. Zdru`eni arabski emirati (ZAE) imajo za Savdsko Arabijo drugi najve~ji letni stro{ek zaradi korozije v
svoji veliki naftni industriji. Zato ta {tudija posega na podro~je kvalitete pou~evanja korozije pri in`enirskih programih na
univerzah v ZAE. Z enotno obliko {tudija primera v obliki ankete so bile pregledane akademije in industrija o pristojnosti
in`enirjev in diplomiranih in`enirjev glede korozije. Ugotovitve ka`ejo, da so namenski te~aji o koroziji in in`enirski predmeti,
ki vklju~ujejo korozijo, vklju~eni v programih univerz v ZAE. Glede pristojnosti {tudentov tehnike in in`enirjev je bilo soglasno
potrjeno, da ni dovolj temeljnega znanja s podro~ja obvladovanja korozije. Anketiranci iz industrije so bili kriti~ni, da imajo
in`enirji pomanjkljivo razumevanje korozije pri projektih v realnem okolju. U~inkovitost in`enirskih programov glede korozije
je dolo~ena tako s kompeten~nim znanjem in izku{njami s korozijo ter z razpolo`ljivimi viri (usposobljenimi delavci in novim
znanjem iz raziskav), ki naj bi podpirale izobra`evanje o koroziji. Rezultati so pokazali, da ve~ina podjetij ne bi zaposlila novih
strokovnjakov za korozijo. Vendar pa je vidik raziskav bolj spodbuden na univerzah, ki poro~ajo o mo`nosti raziskav in
partnerstvu z industrijo na podro~ju raziskav korozije. ^lanek se kon~a s priporo~ilom po izbolj{anju znanja in izku{enj bodo~ih
in`enirjev pri upravljanju s korozijo, kot tudi spodbujanje korozijskega svetovanja za bolj{e zadovoljevanje potreb industrije.
Klju~ne besede: korozija, izobra`evanje in`enirjev, {tudij primera, Zdru`eni arabski emirati
1 INTRODUCTION
The United Arab Emirates (UAE) incurs a high annual
cost of corrosion of US$ 14.26 billion (2011), which is
about 5.2 % of the country’s GDP over three years
(2009–2011), due to its large petroleum industry1.
Corrosion is a deterioration of metals or materials due to
chemical, biological or environmental agents2,3 resulting
in serious health, safety and environmental (HSE) risks.
When poorly managed, corrosion affects the equipment
integrity and serviceability, raises the risk of discharge of
flammable fluids and gases that companies can be held
liable for4. Since it is impossible to eliminate corrosion,
engineers in the petroleum industry need to be competent
in corrosion fundamentals and management to minimize
the economic, environmental and social consequences of
a corrosion failure. Much of this competence is gained
through higher education. Yet, 54 % of the corrosion-
protection practitioners have not taken a corrosion course
in their formal education. Also, 45 % of the currently
active and experienced corrosion technologists are likely
to retire in the next 10 years4. Furthermore, there is little
research on the quality of corrosion education in the
UAE. Hence, this study aims to assess the quality of
corrosion education in engineering programs of the
universities in the UAE. The research questions are:
RQ 1a: What is the level of corrosion training available
at the universities in the UAE?
RQ 1b: Why is this level of corrosion training charac-
teristic of the UAE universities?
RQ 2a: How competent are the engineering undergra-
duates/graduates from the UAE universities in
corrosion knowledge and skills?
Materiali in tehnologije / Materials and technology 47 (2013) 6, 849–851 849
UDK 378.6:620.193(53) ISSN 1580-2949
Short note/Kratka novica MTAEC9, 47(6)849(2013)
RQ 2b: What resources are available to the UAE uni-
versities to support effective corrosion education?
2 CASE STUDY RESEARCH DESIGN
A single-embedded case design was used where,
within a single case, attention is also given to subunits.5,6
This design enabled in-depth analyses of the subunits
providing an insight into the main case. The main case is
the engineering education program of the UAE univer-
sities, while four universities were the subunits in the
primary case. Through purposive sampling, four institu-
tions were identified that met the following criteria:
location in the UAE; availability of engineering pro-
grams, for which corrosion is relevant; at least one
cohort of graduates; reliable accessibility to the partici-
pants. Due to ethical considerations, the universities
were pseudonymised.
The primary data sources were the academics (n =
58) from the four institutions and the engineers from the
oil/gas industry (n = 8). The latter were interviewed to
understand the employers’ perspective on the compe-
tence of engineering graduates in corrosion. The total
sample consisted of 66 respondents. Surveys and inter-
views (individuals, focus groups) were used to gather
quantitative and qualitative data that were statistically
and interpretively analyzed, respectively. The question-
naires provided information on the respondent demo-
graphics, the level of the existing engineering curricula
in corrosion, the effectiveness of engineering curricula in
corrosion, and recommendations for enhancing the
corrosion knowledge and skills. Table 1 summarizes the
constructs and questions that constitute the operational
measures.
The first construct, the level of the existing engineer-
ing curricula in corrosion, is defined as the extent of the
corrosion training available. The construct is operationa-
lized by Q.10 measuring the availability of corrosion
training on the scale from having a dedicated corrosion
course, teaching corrosion as part of other courses
(integrated) to not teaching corrosion at all. The second
construct, the effectiveness of the existing engineering
curricula in corrosion, is defined as the degree, to which
a course produces the desired educational outcomes
measured as (a) the level of competence in corrosion
knowledge and skills of engineering undergraduates/
graduates (Q.34-35); (b) the availability of the resources
to achieve the desired outcomes (Q.29, 32-33).
3 RESULTS
Regarding the level of corrosion training, this study
found that the dedicated corrosion courses and the
engineering courses integrating corrosion into the
curricula were available at the UAE universities.
Detailed findings on the structure of the dedicated and
integrated corrosion courses are reported in the author’s
expanded publication7. For the competence in corrosion
knowledge and skills, the consensus was that there is
insufficient fundamental knowledge of corrosion
engineering due to the limited scope of corrosion in the
curricula. Another dimension of the competence is the
ability to apply theoretical knowledge in practice.
Interestingly, while academia respondents held a positive
H. L. LIM et al.: ASSESSING THE QUALITY OF CORROSION EDUCATION IN THE UNITED ARAB EMIRATES
850 Materiali in tehnologije / Materials and technology 47 (2013) 6, 849–851
Table 1: Questions relating to constructs and measures
Tabela 1: Vpra{anja glede na~rtovanja in ukrepov
Construct – Level of the existing engineering curricula in corrosion
Q10. Does your department offer a course or courses specifically in corrosion?
Yes, my department offers course(s) specifically in corrosion
No, my department teaches corrosion as part of other courses
No, corrosion is NOT taught at all in my department
Construct – Effectiveness of the existing engineering curricula in corrosion
Level of competence
in corrosion
knowledge and skills
Q34. Do you think engineering undergraduates/graduates (from your university) have a sufficient
fundamental knowledge of corrosion engineering? Please explain.*
Q35. Do you think engineering undergraduates/graduates (from your university) have a sufficient
understanding of the importance of corrosion in engineering design? Please explain.*
Availability of
resources (manpower,
knowledge)
Q29. Would your department consider hiring a faculty member whose technical focus is corrosion?
Yes
No
Q32. Is your department doing any corrosion-specific research? If so, who is funding it?
No
Yes, funding organization(s): _____________________
Q33. Do you have any actual or potential partnerships with the industry to study corrosion or
develop continuing education for practicing engineers?
No
Yes
If so, please describe your partnerships with the industry to study corrosion.
*Industry respondents answered the variation of Q34/35 that excluded the phrase "from your university"
view that the students had a sufficient understanding of
the importance of corrosion in engineering design, the
industry respondents were more critical, believing that
the graduates had a superficial understanding of the
corrosion in real-life design contexts.
The effectiveness of engineering curricula in corro-
sion is determined by both the competence in corrosion
knowledge/skills and the availability of resources (quali-
fied staff, new knowledge from research) to support
corrosion education. Unfortunately, the engineering
departments in three out of the four universities would
not hire corrosion specialists as other topics have more
priority. However, the aspect of research is more
encouraging, as two universities reported the availability
of both departmental research and industry partnerships
in corrosion research (Table 2).
In terms of the overall quality of corrosion education,
this study concluded that while corrosion is part of most
engineering programs of the UAE universities, there are
inadequate resources to support corrosion education with
respect to hiring corrosion specialists or gaining new
knowledge from department research and industry
partnerships in corrosion research.
4 RECOMMENDATIONS AND FUTURE
RESEARCH
Respondents were asked to propose the strategies for
enhancing the corrosion knowledge and skills of
engineering undergraduates. Three main themes were
identified and the following recommendations are
offered for improving the corrosion training to better
meet industry expectations:
Set corrosion courses as mandatory courses in the
curriculum framework of engineering fields, where
corrosion is particularly relevant so that a common level
of knowledge and skills can be established.
Integrate the corrosion theory with practical expe-
rience in engineering courses and senior design projects
so that the undergraduates can be actively involved in
applying the corrosion theory to actual design processes.
Promote the awareness of the corrosion impact
through campus-wide activities to raise the appreciation
of the need for corrosion mitigation and management.
Future research in this area could include compa-
rative studies examining the differences in quality of
corrosion education in other Middle-East countries. The
major oil/gas producing countries such as Qatar, Saudi
Arabia and Kuwait have established universities offering
extensive engineering programs, for instance: Texas
A&M University in Qatar, Qatar University, King Fahd
University of Petroleum and Minerals, and Kuwait
University. Hence, such studies could capture a more
complete picture of the quality of corrosion education in
the region, where corrosion prevention and management
are crucial for its petroleum-based economies.
5 REFERENCES
1 A. Al Hashem, Corrosion in the Gulf Cooperation Council (GCC)
states: Statistics and figures, presented at the Corrosion UAE, Abu
Dhabi, UAE, 2011
2 Z. Ahmad, Principles of Corrosion Engineering and Corrosion Con-
trol, Butterworth-Heinemann, Oxford 2006
3 P. Roberge, Corrosion Engineering: Principles and Practice,
McGraw-Hill Professional, New York 2008
4 D. Rose, Current practice: The teaching of corrosion in colleges and
universities in Materials Forum 2007: Corrosion Education for the
21st Century, Washington, DC 2007, 8–12
5 R. Yin, Case Study Research: Design and Methods, 2nd ed., Sage
Publishing, Beverly Hills, CA 1994
6 R. Yin, Case Study Research: Design and Methods, 3rd ed., Sage
Publications, Inc., Thousand Oaks, CA 2003
7 H. L. Lim, Assessing level and effectiveness of corrosion education
in the United Arab Emirates, International Journal of Corrosion,
(2012), 1–10
H. L. LIM et al.: ASSESSING THE QUALITY OF CORROSION EDUCATION IN THE UNITED ARAB EMIRATES
Materiali in tehnologije / Materials and technology 47 (2013) 6, 849–851 851
Table 2: Availability of resources to support corrosion training: faculty and research
Tabela 2: Razpolo`ljivost virov za podporo pou~evanja korozije: fakultete in raziskave
University A University B University C University D
Recruitment of corrosion specialists*
Hiring corrosion specialists 0 % 8 % 6 % 55 %
Not hiring corrosion specialists 100 % 92 % 94 % 45 %
Research efforts in corrosion*
Corrosion research carried out by department 10 % 0 % 25 % 55 %
Industry research partnership available 0 % 0 % 13 % 25 %
*Based on 58 academic respondents who are instructors in mechanical, civil and chemical-engineering departments of the universities

[OLSKI CENTER RAVNE NA KORO[KEM KON^AL PROJEKTA ENER-
GETSKE SANACIJE IN GRADNJE SODOBNEGA IZOBRA@EVALNEGA
CENTRA V SKUPNI VREDNOSTI PRIBLI@NO TRI MILIJONE EVROV
Dne 25. oktobra 2013 je na Srednji {oli Ravne
potekalo uradno odprtje Medpodjetni{kega izobra-
`evalnega centra Ravne (MIC [C Ravne) za podro~je
obdelave in predelave kovin in slovesnost ob ener-
getski sanaciji objektov [olskega centra Ravne na
Koro{kem, ki naj bi za 63 % zmanj{ala porabo
toplotne energije in za 20 % elektri~ne energije.
Skupna vrednost projektov je bila pribli`no tri
milijone evrov. Odprtja se je udele`il tudi minister za
izobra`evanje, znanost in {port dr. Jernej Pikalo, ki je
poudaril pomen medpodjetni{kih izobra`evalnih cen-
trov in njihovo povezovalno vlogo med izobra`e-
vanjem in industrijo. Povedal je, da »mora sedaj
industrija jasno izraziti svoje potrebe« in tako omo-
go~iti MIC-em, da bodo »dajali tisto, kar potrebuje
slovensko gospodarstvo«. Po kon~anem uradnem delu
z nastopom Eve Boto so si udele`enci ogledali naj-
sodobnej{o tehnolo{ko opremo v skupni vrednosti
pribli`no 1 800 000 evrov in del objekta Srednje {ole
Ravne, ki je bil energetsko saniran.
[olski center Ravne na Koro{kem je septembra
kon~al energetsko sanacijo stavb Gimnazije in Srednje
{ole Ravne, sofinancirano iz Kohezijskega sklada,
oktobra pa kon~al gradnjo sodobnega in tehnolo{ko
naprednega MIC [C Ravne, ki je bil sofinanciran iz
Evropskega sklada za regionalni razvoj. Dne 25. oktobra
2013 je na Srednji {oli Ravne potekala slovesnost ob
kon~anju obeh projektov. Direktor [olskega centra in
ravnatelj Gimnazije Ravne na Koro{kem Dragomir
Benko je ob tem povedal, da »so s projektoma, vrednima
nekaj ve~ kot tri milijone evrov, opravili odli~no delo,
vzpostavili bolj{e razmere in zmanj{ali negativne vplive
na okolje«.
Energetska sanacija Gimnazije in Srednje {ole Ravne
je potekala na 6000 m2 ogrevanih povr{in in bo pred-
vidoma zmanj{ala porabo toplotne energije za 63 % ter
elektri~ne energije za 20 %. Skupna vrednost projekta je
bila pribli`no 940 000 EUR.
Starej{i del fasade Gimnazije je bil izoliran in
zaklju~en z novim ometom teranova, novej{i del pa
izoliran in zaklju~en s fasadnimi plo{~ami. Dodatno je
bila izolirana podstreha, zamenjane so bile zunanje
okenske police in dotrajana vrata. Vgrajene so bile nove
svetilke, zamenjani radiatorski termostatski ventili na
zaklep ter vgrajeni toplotna ~rpalka (zrak-voda), prezra-
~evalne naprave in razvod za prezra~evanje. Na Srednji
{oli Ravne je potekala energetska sanacija na vzhodnem
traktu objekta. Sanirana je bila fasada, zamenjana so bila
okna. Vgrajene so var~ne svetilke in senzorji gibanja. Na
vsa grelna telesa so name{~eni termostatski ventili. V
okviru sanacije je bila vgrajena prezra~evalna naprava.
Materiali in tehnologije / Materials and technology 47 (2013) 6, 853–854 853
UDK 330.322.3:37(497.12) ISSN 1580-2949
Dogodki/Events MTAEC9, 47(6)853(2013)
Na obeh objektih je vzpostavljen sistem digitalnega
obratovalnega nadzora (DOM:EU), ki omogo~a sprem-
ljanje porabe toplotne in elektri~ne energije ter vode in
stanja klimatskih razmer.
»Za vzpostavitev MIC-a je bil zgrajen prizidek in
dobavljena najsodobnej{a oprema, ki bo izbolj{ala
kakovost prakti~nega dela izobra`evanja in usposab-
ljanja dijakov, {tudentov ter odraslih udele`encev,« je na
odprtju povedala ravnateljica Srednje {ole Ravne Ivanka
Stopar. V okviru projekta so bili dobavljeni merilna,
strojna, didakti~no-tehnolo{ka, ra~unalni{ka in varilna
oprema, stroj za centriranje, mala fotovoltai~na elek-
trarna (MFE), laserski stroj, mini robot, testna steza itd.
Med najpomembnej{e pridobitve spadajo 5-osni CNC
vertikalni rezkalni obdelovalni center – DOOSAN
VC630/5AX, CNC-stru`nica z gnanim orodjem in Y-osjo
– DOOSAN PUMA 2100LY ter CNC brusilni stroj za
okroglo zunanje in notranje bru{enje – OKAMOTO
OGM-NCIII 390.
MIC tako ustvarja razmere za inovativne na~ine izo-
bra`evanja in raziskovanja v sklopu naslednjih poklicev:
oblikovalec kovin, mehatronik operater, avtoserviser,
metalurg, strojni tehnik, elektrotehnik, ra~unalni{ki in
strojni tehnik, pomo~nik pri tehnolo{kih procesih.
Obenem je MIC podporno okolje za osvajanje prakti~nih
znanj, spretnosti in sposobnosti za takoj{njo vklju~itev v
delovne procese pri delodajalcih, delodajalcem pa
omogo~a prenos znanja in sodobnih tehnologij v delovne
procese. MIC bo opravljal tudi prakti~no izpopolnjevanje
za zaposlene v obrti in gospodarstvu, za u~itelje prak-
ti~nega usposabljanja in strokovno-teoreti~nih predmetov
v strokovnih in poklicnih {olah ter za nezaposlene, ki jih
bo v MIC napotil zavod za zaposlovanje. MIC se bo
povezoval tudi z osnovnimi {olami v regiji.
Skupna vrednost investicije v gradnjo MIC-a Ravne
na podro~ju obdelave in predelave kovin je bila 2 136 000
evrov in je bila v vrednosti 1 740 000 evrov sofinancira-
na iz Evropskega sklada za regionalni razvoj.
Vir: EUTRIP
854 Materiali in tehnologije / Materials and technology 47 (2013) 6, 853–854