VSEBINA – CONTENTS Editor’s Preface / Predgovor urednika M. Torkar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 181 IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES Pitting corrosion of TiN-coated stainless steel in 3 % NaCl solution Jami~asta korozija nerjavnega jekla s prevleko TiN v 3-odstotni raztopini NaCl I. Kucuk, C. Sarioglu . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 183 Design of a wideband planar antenna on an epoxy-resin-reinforced woven-glass material [irokopasovna ploskovna antena na epoksi smoli, oja~ani s steklenimi vlakni R. Azim, M. T. Islam . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 193 Influence of the strain rate on the PLC effect and acoustic emission in single crystals of the CuZn30 alloy compressed at an elevated temperature Vpliv hitrosti deformacije na pojav PLC in akusti~no emisijo monokristalov zlitine CuZn30, stiskane pri povi{ani temperaturi W. Ozgowicz, B. Grzegorczyk, A. Pawe³ek, A. Pi¹tkowski, Z. Ranachowski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 197 Determination of elastic-plastic properties of Alporas foam at the cell-wall level using microscale-cantilever bending tests Dolo~anje elasti~nih in plasti~nih lastnosti pene Alporas na ravni celi~ne stene z upogibnimi preizkusi z mikroskopsko iglo T. Doktor, D. Kytýø, P. Koudelka, P. Zlámal, T. Fíla, O. Jirou{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 203 Electrochemical behavior of biocompatible alloys Elektrokemijsko vedenje biokompatibilnih zlitin I. Petrá{ová, M. Losertová. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 207 Preparation and thermal stability of ultra-fine and nano-grained commercially pure titanium wires using CONFORM equipment Priprava komercialne ultradrobne in nanozrnate Ti-`ice z opremo CONFORM in njena termi~na stabilnost T. Kubina, J. Dlouhý, M. Köver, M. Dománková, J. Hodek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 213 Estimation of the thermal contact conductance from unsteady temperature measurements Dolo~anje kontaktne toplotne prevodnosti iz neravnote`nega merjenja temperature J. Kvapil, M. Pohanka, J. Horský . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 219 Potentiodynamic and XPS studies of X10CrNi18-8 steel after ethylene oxide sterilization Potenciodinami~ne in XPS analize jekla X10CrNi18-8 po sterilizaciji z etilen oksidom W. Walke, J. Przondziono . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 223 TEM replica of a fluoride-miserite glass-ceramic glaze microstructure TEM-replike mikrostrukture steklokerami~ne fluor-mizeritne glazure J. Ma. Rincón, R. Casasola . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 229 Experimental analysis and modeling of the buckling of a loaded honeycomb sandwich composite Eksperimentalna analiza in modeliranje upogibanja obremenjenega satastega sendvi~nega kompozita A. Bentouhami, B. Keskes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 235 Fatigue behaviour of X70 steel in crude oil Vedenje jekla X70 pri utrujanju v surovi nafti ¼. Gajdo{, M. [perl, J. Bystrianský . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 243 Use of Larson-Miller parameter for modeling the progress of isothermal solidification during transient-liquid-phase bonding of IN718 superalloy Uporaba Larson-Millerjevega parametra za modeliranje izotermnega strjevanja pri spajanju z vmesno teko~o fazo superzlitine IN718 M. Pouranvari, S. M. Mousavizadeh . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 247 Effect of severe air-blast shot peening on the wear characteristics of CP titanium Vpliv intenzivnega povr{inskega kovanja s peskanjem z zrakom na obrabne lastnosti CP-titana A. C. Karaoglanli . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 253 ISSN 1580-2949 UDK 669+666+678+53 MTAEC9, 49(2)179–309(2015) MATER. TEHNOL. LETNIK VOLUME 49 [TEV. NO. 2 STR. P. 179–309 LJUBLJANA SLOVENIJA MAR.–APR. 2015 STROKOVNI ^LANKI – PROFESSIONAL ARTICLES Finite-element minimization of the welding distortion of dissimilar joints of carbon steel and stainless steel Uporaba kon~nih elementov za zmanj{anje popa~enja oblike pri varjenju ogljikovega in nerjavnega jekla E. Ranjbarnodeh, M. Pouranvari, M. Farajpour . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 259 Magnesium-alloy die forgings for automotive applications Izkovki iz magnezijevih zlitin za avtomobilsko industrijo M. Madaj, M. Greger, V. Karas . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 267 Resistance to electrochemical corrosion of the extruded magnesium alloy AZ80 in NaCl solutions Odpornost ekstrudirane magnezijeve zlitine AZ80 proti elektrokemijski koroziji v raztopini NaCl J. Przondziono, E. Hadasik, W. Walke, J. Szala, J. Michalska, J. Wieczorek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 275 Microwave-assisted hydrothermal synthesis of Ag/ZnO sub-microparticles Hidrotermi~na sinteza podmikrometrskih delcev Ag/ZnO z mikrovalovi L. Münster, P. Ba`ant, M. Machovský, I. Kuøitka. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 281 Determination of the cause of the formation of transverse internal cracks on a continuously cast slab Ugotavljanje vzrokov za nastanek notranjih pre~nih razpok v kontinuirno ulitem slabu Z. Franìk, M. Masarik, J. [míd . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 285 Effective preparation of non-linear material models using a programmed optimization script for a nurimerical simulation of sheet-metal processing U~inkovita priprava nelinearnih modelov materiala s programiranim optimizacijskim zapisom za numeri~no simulacijo obdelave plo~evine M. Urbánek, F. Tikal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 291 Neutralization of waste filter dust with CO2 Nevtralizacija odpadnega filtrskega prahu s CO2 A. Kra~un, I. An`el, L. Fras Zemlji~, A. Stergar{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 297 The influence of the morphology of iron powder particles on their compaction in an automatic die Vpliv morfologije delcev `elezovega prahu na njegovo sposobnost za avtomatsko enoosno stiskanje B. [u{tar{i~, M. Godec, ^. Donik, I. Paulin, S. Glode`, M. [ori, M. Ratej, N. Javornik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 303 EDITOR’S PREFACE / PREDGOVOR UREDNIKA Each issue of the journal Materiali in tehnologije/ Materials and Technology (MIT) gets larger from year to year. In Volume 48 of MIT we published 6 Issues with a total of 1030 pages. There was a total of 159 papers published: 2 as review papers, 117 as original scientific papers and 40 as professional papers. All papers were published in English (with a translation of the titles, abstracts, keywords, figure and table captions into Slo- vene). In total there were 581 authors from all over the world. The average number of authors per paper was 3.65. However, some of these authors appeared several times. The increasing number of authors from different countries is encouraging, as it demonstrates the quality and visibility of the journal in the international scientific community. Volume 48 contained 105 papers from the field of metallic materials, 11 papers from the field of inorganic materials, 10 papers from the field of poly- mers, 1 paper from the field of vacuum techniques, 6 papers from the field of chemical technology, 3 papers from the fields of nanomaterials and nanotechnologies, 6 papers from the field of building materials and 17 papers from the field of numerical methods. In comparison with previous years, we saw an in- crease in the number of papers from the field of metallic materials. The interest of authors to publish in MIT is growing, which we suppose is due to the journal being involved in the system of impact factor (IF). The IF for MIT was 0.555 in the year 2014. The impact factor of the journal depends on the papers from MIT being cited in other journals. Of course, MIT is indexed in various journal databases: the Science Citation Index Expanded, the Materials Science Citation Index, and Journal Cita- tion Reports/Science Edition. However, the papers from MIT are also indexed in numerous international secon- dary sources. The journal is available online at http://mit.imt.si. In 2014 we started with our activities to include the articles published in MIT into the system of Digital Object Identifier (DOI). We hope that the activation of the DOI system will make MIT even more interesting for authors and readers. Publication in MIT remains free of charge. However, to increase of the quality of the journal we expect from authors, papers with original scientific results, written in good English. I would like to thank all the authors, editors, techni- cal staff, lectors and reviewers for all their efforts that enable the regular publication of the journal Materiali in tehnologije/Materials and Technology. Editor-in-Chief A/Prof. Dr. Matja` Torkar Obseg posamezne {tevilke revije Materiali in tehno- logije/Materials and Technology (MIT) se iz leta v leto pove~uje. Tako je bilo v letniku 48 natisnjenih 6 rednih {tevilk v obsegu 1030 strani. Objavljenih je bilo 159 ~lankov, od tega 2 pregledna, 117 izvirnih znanstvenih in 40 strokovnih. Vseh 159 ~lankov je bilo objavljenih v angle{kem jeziku s prevodi naslovov, povzetkov, klju~nih besed in podnaslovov slik in tabel v slovenski jezik. Skupaj je bilo 581 avtorjev iz {tevilnih dr`av. Povpre~no {tevilo avtorjev na ~lanek je 3,65. Nekateri avtorji se pojavljajo tudi ve~krat. Rasto~e {tevilo avtorjev iz razli~nih dr`av je razveseljivo, kar ka`e na kakovost in odmevnost revije v mednarodnem prostoru. Po vsebini je bilo {tevilo ~lankov z naslednjih podro~ij: kovinski materiali 105, anorganski materiali 11, polimeri 10, vakuumska tehnika 1, kemijske tehnologije 6, nano- materiali in nanotehnologije 3, gradbeni materiali 6 in numeri~ne metode 17. V primerjavi s prej{njimi leti se pove~uje predvsem {tevilo ~lankov s podro~ja kovinskih materialov. V primerjavi s prej{njimi leti se zanimanje za objavo v MIT vsako leto pove~uje, domnevno zato, ker je revija vklju~ena v sistem dejavnika vpliva (Impact Factor), ki je bil v letu 2014 0,555. Ta dejavnik je odvisen nepo- sredno od ~lankov iz MIT, ki so citirani v drugih revijah. Revija je indeksirana tudi v naslednjih bazah podatkov: Science Citation Index Expanded, Materials Science Citation Index, Journal Citation Reports/Science Edition, ~lanki iz revije Materiali in tehnologije pa so indeksirani tudi v {tevilnih mednarodnih sekundarnih virih. Revija je v celotnem obsegu dostopna v elektronski obliki na http://mit.imt.si. V letu 2014 smo za~eli aktivnosti za vklju~itev revije MIT v sistem Digital Object Identifier (DOI). Upamo, da bo aktivacija sistema DOI napravila revijo Materiali in tehnologije/Materials and Technology {e bolj zanimivo za nove avtorje. Objava v MIT ostaja brez- pla~na. Za pove~anje kvalitete revije pri~akujemo, da bodo avtorji pripravili ~lanke z izvirnimi znanstvenimi spoznanji in v dobri angle{~ini. Zahvaljujem se vsem avtorjem, urednikom, tehni- ~nim sodelavcem, lektorjem in recenzentom, ki omogo- ~ajo redno izhajanje revije MIT. Glavni in odgovorni urednik doc. dr. Matja` Torkar Materiali in tehnologije / Materials and technology 49 (2015) 2, 181 181 I. KUCUK, C. SARIOGLU: PITTING CORROSION OF TiN-COATED STAINLESS STEEL IN 3 % NaCl SOLUTION PITTING CORROSION OF TiN-COATED STAINLESS STEEL IN 3 % NaCl SOLUTION JAMI^ASTA KOROZIJA NERJAVNEGA JEKLA S PREVLEKO TiN V 3-ODSTOTNI RAZTOPINI NaCl Israfil Kucuk, Cevat Sarioglu Marmara University, Dept. of Metallurgical and Materials Engineering, Göztepe kampusu, 34722 Kadiköy-Istanbul, Turkey cevat.sarioglu@marmara.edu.tr Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2014-02-12 doi:10.17222/mit.2013.176 TiN coatings deposited by arc PVD were characterized by XRD and SEM. In-situ measurements of the corrosion of the substrate and the TiN-coated substrate were made using the corrosion potential (Cor.Pot.), the polarization resistance (PR) method and electrochemical impedance spectroscopy (EIS) in a 3 % NaCl solution as a function of the immersion time. The semiconductor scale formed on the TiN was identified using a Mott-Shottky analysis as an n-type semiconductor with a flat band potential of –0.83 V vs. SCE. The TiN coating (0.5 μm thick) consisted of cubic TiN exhibiting columnar grains, pin holes, voids and porosities. The pitting corrosion of the TiN, observed visually between 1 h and 2 h, was captured by EIS and PR. The electrical circuit (EC) model used for the EIS data supported the degradation of the coating through pitting corrosion, in agreement with the visual observations. The corrosion resistance (polarization resistance) determined by the polarization resistance method (Rp) and the EIS (Rtotal) decreased suddenly during the pitting corrosion. The corrosion resistance of the TiN-coated substrate was greater than the corrosion resistance of the substrate during the approximately 24 h of exposure. Keywords: stainless steel, TiN, coating, EIS, polarization resistance, pitting corrosion Prevleka TiN, nanesena z oblo~nim PVD-postopkom, je bila pregledana z XRD in SEM. In-situ meritve korozije podlage in podlage s prevleko iz TiN so bile izvr{ene z metodo korozijskega potenciala (Cor.Pot.), polarizacijske upornosti (PR) in z elektrokemijsko impedan~no spektroskopijo (EIS) v 3-odstotni raztopini NaCl v odvisnosti od ~asa namakanja. Mott-Shottkyjeva analiza je odkrila nastanek n-tipa polprevodne {kaje na TiN s potencialom ravnih nivojev –0.83 V proti SCE. Prevleko TiN (debeline 0,5 μm) sestavljajo kubi~ni TiN s stebrastimi zrni, luknjami, prazninami in poroznostjo. Vidno odkrita jami~asta korozija TiN, opa`ena med 1 h in 2 h, je bila posneta z EIS in PR. Model elektri~nega tokokroga (EC), ki je bil uporabljen za EIS-podatke, podpira degradacijo prevleke z jami~asto korozijo, skladno z vizualnimi opa`anji. Korozijska upor- nost (polarizacijska upornost), dolo~ena z metodo polarizacijske upornosti (Rp) in EIS (Rtotal) se je nenadno zmanj{ala med jami~asto korozijo. Korozijska upornost podlage z nanosom TiN je bila ve~ja kot korozijska obstojnost podlage med izpostavitvijo okrog 24 h. Klju~ne besede: nerjavno jeklo, TiN, prevleka, EIS, polarizacijska upornost, jami~asta korozija 1 INTRODUCTION Hard ceramic coatings such as TiN have been used mainly for tribological applications, such as cutting tools. The tribological properties of single-layer and multi-layer TiN coatings were extensively studied in the literature and published in a handbook.1–8 On the other hand, the corrosion of TiN in tribological applications was often overlooked, mainly due to the shorter life time of the cutting tools. TiN, with its golden colors, has been used for decorative applications, such as watches, archi- tectural materials and ornaments. The corrosion resi- stance of the TiN coating is required for these decorative applications in addition to the wear resistance. The pitting corrosion of the TiN coating deposited on the metallic substrates AISI 304, 430 and steel was observed for different coating thicknesses, exposure times and coating techniques.9–16 Even though the widely accepted pitting corrosion mechanism of the TiN-coated substrate in the literature was local galvanic corrosion through the galvanic coupling of the TiN coating and the substrate, there remain questions about the mechanism of pitting formation and the growth of the pitting corrosion. One of the key parameters for galvanic corrosion is defects in the coatings, such as pin holes and micro and macro porosities in the coatings. The defects in the coat- ings provide the electrolyte with a path to the coating/ substrate interface.9–16 Martensitic stainless steels (EN 1.4034 was used in this work) that are generally used for the blades in kitchen appliances were coated for both decorative and wear-resistance requirements. The pitting corrosion of the TiN coating deposited by arc PVD on a stainless- steel substrate was studied in detail with the corrosion potential (Cor.Pot), the polarization resistance (PR) and electrochemical impedance spectroscopy (EIS) tech- niques. The mechanism of the pitting corrosion was evaluated with respect to the microstructure of the TiN coating. 2 EXPERIMENTAL PROCEDURES The substrate material obtained from ThyssenKrupp was EN 1.4034 (X46Cr13) stainless steel. The TiN Materiali in tehnologije / Materials and technology 49 (2015) 2, 183–192 183 UDK 669.14.018.8:620.193:620.197:621.793 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)183(2015) coating of the substrates was performed in an industrially sized arc PVD coating chamber (AFS ltd. Cop., Turkey). The details about the specimen preparation prior to the coating and the coating procedure are given in detail in17. The substrate was coated with a Ti interlayer for 1 min to improve the adhesion of the coating and later with a TiN layer for 20 min at 1.1 · 10–3 mbar of nitrogen pressure and a total pressure of 10–2 mbar with a bias voltage of –200 V. The final deposition temperature was 250 °C. The electrochemical corrosion units used to perform the EIS, the polarization resistance and the Mott-Shottky scan were a Gamry PC14/750 Poteniostat/Galvanostat/ ZRA System. Details about the polarization resistance and the EIS techniques are given in17. All the tests were performed in 0.5 M (w = 3 %) NaCl aerated water solu- tion at 25 °C using a three-electrode system (working (sample), auxiliary (graphite) and reference (standard calomel electrode (SCE)) using a Gamry paint cell unit. The Mott-Shottky analyses at a frequency of 1 Hz bet- ween +430 mV and –570 mV were performed on a TiN-coated substrate after the corrosion potential was stable. Microstructural analyses were performed before and after the corrosion using a scanning electron microscope (SEM, Jeol, JSM-5910LV) and energy-dispersive spec- troscopy (EDS). The cross-section of the coatings was observed after it was fractured in liquid nitrogen. X-ray diffraction (Rigaku, D-MAX 2200, Cu K radiation) was used to identify the structure of the coatings deposited on the substrate. 3 RESULTS 3.1 Microstructural characterization of the substrate and the TiN-coated substrate An SEM micrograph of the microstructure of the substrate material is presented in Figure 1. The carbide phases (Cr and C-rich phase identified by EDS), etched I. KUCUK, C. SARIOGLU: PITTING CORROSION OF TiN-COATED STAINLESS STEEL IN 3 % NaCl SOLUTION 184 Materiali in tehnologije / Materials and technology 49 (2015) 2, 183–192 Figure 3: SEM (SEI) micrographs taken from TiN-coating surface at: a) low magnification and b) high magnification. TiN coating deposited over carbide phases (dark areas) were distinguished with depression over the surface at a) and b). Embedded droplets deposited throughout the surface were marked at b). Bright particles on coating surface were un-embedded droplets (spherical particles). Pin holes were marked on surface b). Slika 3: SEM (SEI)-posnetka povr{ine prevleke TiN pri: a) majhni po- ve~avi in b) veliki pove~avi. Prevleka TiN je nanesena preko karbidnih faz (temna podro~ja), ki se razpoznajo po vdolbini na povr{ini a) in b). Vgnezdene kapljice, nanesene na povr{ino, so ozna~ene na b). Svetli delci na povr{ini nanosa so nevgnezdene kapljice (sferi~ni delci). Luknjice so ozna~ene na povr{ini b). Figure 1: SEM (BEI) micrograph of EP 4034 substrate where dark-grey particles were carbides and bright phases were Fe-rich Fe-Cr particles in a grey matrix Slika 1: SEM (BEI)-posnetek podlage iz EP 4034, kjer so temnosive pike karbidi, svetle pa faze z Fe bogati Fe-Cr-delci v sivi osnovi Figure 2: 2 scan obtained from the EP 4034 steel substrate and TiN-coated substrates using the Bragg-Brentano symmetric X-ray diffraction Slika 2: 2-posnetek Bragg-Brentano simetri~ne rentgenske difrakcije EP 4034 podlage in prevleke TiN na podlagi slightly more than the matrix during the electropolishing treatment, were distributed homogenously throughout the matrix. The 4034 EP stainless-steel substrate possessed a ferritic structure (-Fe), as confirmed by the XRD (Figure 2). After coating with TiN at 1.1 · 10–3 mbar of N2 partial pressure, the surface morphology of the coating reflected the morphology of the EP substrate surface, where the etched carbide phases were covered by a TiN coating (Figure 3). The TiN coating was identified as a cubic TiN phase (Figure 2). Due to the arc PVD process, the droplets formed on the surface of the TiN-coated sub- strates. There were two different types of droplets found on the surface of the TiN coatings (Figures 3 and 4). One of them was the droplets embedded to the scale, i.e., the coating. These droplets were deposited and incorpo- rated into the coatings during the coating process (Figures 3 and 4). The other droplets were un-embedded macro-particles (bright, spherical particles) (Figure 3). They were thought to be deposited on the surface through vapour-phase precipitation after the coating was finished (when the bias was interrupted). These droplets analysed by EDS contained mainly Ti and N (Ti-rich particles). Based on a detailed surface and cross-section investigation (Figures 3 and 4), it was found that the TiN coating exhibited columnar grains (50 nm diameter) that were aligned perpendicular to the substrate surface (Figure 4) and possessed a significant amount of pin holes and porosity at the surface (Figure 3). The thick- ness of the TiN coatings, measured from the cross-sec- tion (Figure 4), was about 0.5 μm. 3.2 Corrosion of the substrate and the TiN-coated sub- strate The corrosion of the substrate and the TiN-coated substrate was followed by corrosion potential, PR and EIS measurements during about 24 h of exposure. The visually observed state of the surface was noted during the corrosion evaluation. The corrosion potentials were plotted as a function of the exposure time in Figure 5. In general, the corrosion potential on the surface of the TiN decreased with time to a level close to the corrosion potential of the substrate material (Figure 5). At all times, the corrosion potential of the substrate was lower compared to the TiN coating during the approximately 24 h of exposure. The polarization resistance measurements of the substrate and the TiN-coated substrate were performed for a period of 160 s (2.7 min) as a function of the expo- sure time (Figures 6 and 7). The polarization resistance values (Rp) were determined using the polarization resistance method and are plotted in Figure 5. The pola- rization resistance value (Rp) of the substrate determined from the plots in Figure 6 increased gradually until it I. KUCUK, C. SARIOGLU: PITTING CORROSION OF TiN-COATED STAINLESS STEEL IN 3 % NaCl SOLUTION Materiali in tehnologije / Materials and technology 49 (2015) 2, 183–192 185 Figure 5: Corrosion potentials and polarization resistances (Rp) of the substrate and TiN-coated substrates determined from the polarization resistance scan (Figures 6 and 7) in 3 % NaCl solution as a function of the immersion time Slika 5: Korozijski potencial in polarizacijska upornost (Rp) podlage in podlage s prevleko TiN so dolo~ene s posnetka polarizacijske upornosti (sliki 6 in 7) v 3-odstotni raztopini NaCl v odvisnosti od ~asa namakanja Figure 4: SEM (SEI) of micrographs of fractured surface (in liquid nitrogen) of TiN coating from different areas a) and b). Columnar grains of TiN through fractured coating seen at a) and b). Embedded droplets at surface b) and at coating/substrate interface a) were marked. Slika 4: SEM (SEI)-posnetka povr{ine preloma (v teko~em du{iku) prevleke TiN na razli~nih podro~jih a) in b). Stebrasta zrna TiN skozi prelom prevleke se vidijo na a) in b). Na povr{ini b) in na stiku pre- vleka – podlaga a) so ozna~ene vgnezdene kapljice. reached a constant value of 27 k cm2 after 819 min and remained at the same level until 1431 min of exposure, Figure 5. The polarization resistance (Rp) of the TiN coating (Figure 5) determined from the polarization resi- stance plots (Figure 7), exhibited a different evolution during the exposures 24 h. At the start, the polarization resistance (Rp) after 40 min decreased from 70.5 k cm2 to 33 k cm2 within 50 min and then gradually increased to 48 k cm2 in 300 min of exposure and stayed almost constant until 1456 min. The evolution of the polariza- tion (Rp) is clear from the PR measurement in Figure 7 through the expansion and contraction of the PR data. The steep drop in the polarization-resistance value dur- ing early exposure between 40 min and 90 min coincided with the visual observation of four pits formed on the surface between 80 min and 95 min of immersion time. These pits and others formed latter grew during the exposure time. During 24 hours of exposure, the corro- sion resistance (polarization resistance, Rp) of the TiN- coated substrate (Figure 5), was greater than the corro- sion resistance of the substrate. The EIS measurements of the substrate and the TiN-coated substrate were also performed during about 24 h of exposure in a salt solution (Figures 8 and 9). The Bode and Nyquist plots for the substrate material were shown in Figures 8a and 8b. In the Nyquist plots (Fig- ure 8a), the real and imaginary impedance values gradually increased with the exposure time during the first 759 min and then stayed almost constant until 1439 min of exposure. The Bode plots (Figure 8b) exhibited a one-time constant with a minimum in the phase shift close to –80° and at a characteristic frequency of about 10 Hz. The magnitude of the impedance Z also increased with time, as shown for the selected exposure times in the Bode plot (Figure 8b). The Bode and Nyquist plots of the TiN coatings were presented in Figures 9a and 9b. The Nyquist plot (Figure 9a) clearly demonstrated the evolution of the real and imaginary impedance during the early exposure in the salt solution. The first measurement was per- formed after 40 min. In the Nyquist plot (Figure 9a) the impedance values dropped (shrinkage of curves) until 90 min and then started to increase up to 376 min. The evolution of the Nyquist plot took place during a visual observation of the pits (4 pits observed between 80 min and 95 min). Clearly, the pit formation and the growth of the pits at the early stage were captured by EIS measure- ments (particularly by the Nyquist plot, more sensitive to I. KUCUK, C. SARIOGLU: PITTING CORROSION OF TiN-COATED STAINLESS STEEL IN 3 % NaCl SOLUTION 186 Materiali in tehnologije / Materials and technology 49 (2015) 2, 183–192 Figure 7: Polarization resistance scan of TiN coating in 3 % NaCl solution as a function of immersion time Slika 7: Zapis polarizacijske upornosti prevleke TiN v 3-odstotni raztopini NaCl v odvisnosti od ~asa namakanja Figure 8: EIS data of the substrate for selected immersion time: a) Nyquist and b) the Bode plots Slika 8: EIS-podatki podlage pri izbranih ~asih namakanja: a) Nyqui- stovi in b) Bodejevi diagrami Figure 6: Polarization resistance scan of the substrate in 3 % NaCl solution as a function of immersion time Slika 6: Zapis polarizacijske upornosti podlage v 3-odstotni raztopini NaCl v odvisnosti od ~asa namakanja evolution during pitting). After 376 min of exposure, the Bode and Nyquist plots did not vary significantly and the evolution of the impedance data (Figure 9a) resembled a polarization resistance scan (Figure 7). 3.3 Light and SEM surface examination after the cor- rosion As the corrosion potentials, PR and EIS measure- ments were performed, the surfaces of the samples were visually observed during the exposure in the salt solution. There was no change in the colour or the pit formation on surface of the substrate. As mentioned before, at an early stage four pits were observed on the surface of the TiN-coated substrate and they grew, leaving behind circler brownish colour residues that were around these four pits. The other pits (up to 5) appeared at a later stage of the exposure and with a smaller size (lesser growth of pits). Figure 10 shows one of the four pits formed on the surface of the TiN coating at an early stage and which grew during 1464 min of exposure. The brownish colour observed visually for the surrounding of the pits corresponded to dark-grey circles in the SEM micrograph (Figure 10a). A large amount of oxygen and iron elements were found in these areas by EDS analysis, indicating the dissolution of the substrate and the forma- tion of Fe oxide on the coating surface. Figures 10b and 10c showed the same pit surface at a high magnification. In some areas, the coating was spalled off and in some areas they were detached from the substrate (Figure 10b). At the periphery of the pit the interior of the coating was cracked and detached from the surface (Figure 10c). Figure 11 presented the substrate surface for bare areas inside the pit at a high magnification. In I. KUCUK, C. SARIOGLU: PITTING CORROSION OF TiN-COATED STAINLESS STEEL IN 3 % NaCl SOLUTION Materiali in tehnologije / Materials and technology 49 (2015) 2, 183–192 187 Figure 10: SEM (SEI) micrographs taken from the one of the large pits formed during early exposure: a) the micrograph of the pit marked at low magnification, b) the micrograph from the interior of the pit and c) the micrograph from the periphery of the pit interior at high magni- fication Slika 10: SEM (SEI)-posnetki velike jamice, nastale v za~etku nama- kanja: a) posnetek jamice pri majhni pove~avi, b) posnetek notranjosti jamice in c) posnetek okolice jamice pri ve~ji pove~avi Figure 9: EIS data of TiN coating for selected immersion time: a) Nyquist plots and b) the Bode plots Slika 9: EIS-podatki prevleke TiN pri izbranih ~asih namakanja: a) Nyquistovi in b) Bodejevi diagrami the bare areas, there were approximately diameter 1 μm round grains of pure Cr-Fe particles. These particles (Figure 11b), identified by EDS as Cr-rich particles containing Fe but no oxygen and Ti, were thought to be deposited on the bare, exposed substrate surface after the experiment during the drying of the surface. The same morphological observations also were made for other pits. 3.4 The Mott-Shottky measurements The Mott-Shottky measurements were made for the TiN coatings and plotted in Figure 12. Before the Mott- Shottky measurement was made, there was no pit and no colour change on the surface, the corrosion potentials were stable and the EIS data indicated strong capacitive responses at an early stage of the exposure. The Mott- Shottky plot for the TiN coatings (Figure 12) exhibited a linear segment between –195 mV and –570 mV vs. SCE with a positive slope. The positive slope indicated an n-type semiconductor oxide layer on the TiN surface. 4 DISCUSSION 4.1 Structure of the TiN coatings The cubic crystal structure of the TiN was identified by XRD (Figure 2). In the literature, the single phase of the TiN coating was generally obtained with various N2 pressures, since TiN was stable across a wide stoichio- metric range.11–19 The grain morphology of the TiN coat- ing was columnar (Figure 4). The columnar grain boun- daries in the TiN coatings, aligned perpendicularly from the topmost surface down to the substrate/coating inter- face (Figure 4), were considered to be an easy path for the penetration of the electrolyte.19 The droplets, which were considered as a preferential site for pitting in,9 were found on the TiN coating (Figures 3 and 4). The TiN coatings possessed a less uniform coverage over the etched carbide phases and droplets (Figure 3), resulting in a large quantity of porosity and pin holes. All these defects (columnar grain boundary, droplets, pin holes and porosities) in the coating are preferential sites for the penetration of the electrolyte during the pitting corrosion of the TiN coating. 4.2 Mott-Shottky analysis of the TiN coating In the literature, the Mott-Shottky analysis has been used to characterize the semiconductor layer formed on the surfaces of materials and coatings.11,16–21 Based on Figure 12, it was concluded that the semiconductor layer formed on the TiN-coated substrate was n-type (presu- mably TiO2), in agreement with the literature.11,21 The Mott-Shottky equation on page 127 in20 was used to determine the flat band potential and the density of the charge (density of donors for n-type semiconductor) in the space-charge region. After taking the dielectric con- stant of TiO2 as 60, cited in21 as20, the flat band potentials and the density of the charges were determined from the linear portion of the plot in Figure 12 using the eq. in20. I. KUCUK, C. SARIOGLU: PITTING CORROSION OF TiN-COATED STAINLESS STEEL IN 3 % NaCl SOLUTION 188 Materiali in tehnologije / Materials and technology 49 (2015) 2, 183–192 Figure 12: The Mott-Shottky measurements made for TiN coating at a frequency of 1 Hz in 3 % NaCl solution Slika 12: Mott-Shottkyjeve meritve na prevleki TiN pri frekvenci 1 Hz v 3-odstotni raztopini NaCl Figure 11: SEM (SEI) micrographs taken from the interior of the pit in Figure 10 where the scale was removed: a) low magnification and b) high magnification from the same area in a). Equiaxed particles were observed on substrate surface at a) and b). Slika 11: SEM (SEI)-posnetka okolice jamice, prikazane na sliki 10, kjer je bila povr{ina odstranjena: a) majhna pove~ava in b) velika pove~ava istega podro~ja na a). Enakoosne delce se opazi na povr{ini podlage a) in b). The density of the donors charge in the n-type TiO2 was 2.01 · 1025 cm–3. Rudenja21 found similar values for a TiN coating deposited on 304 stainless steel as 2.4 · 1024 cm–3 in a solution of 0.1 M H2SO4 and 0.05 M HCl. The cal- culated flat band potentials from the intercept of the plots (Figure 12), were –0.83 V vs. SCE for the n-type TiO2. 4.3 Corrosion of the substrate and the TiN-coated sub- strate and EIS modelling The thickness of the TiN was relatively small, 0.5 μm, compared to the literature, where the thinnest coat- ing thickness studied usually about 2 μm. For decorative applications, the coating thickness was kept as small as possible for reasons of cost (0.5 μm in this work could not be the optimum thickness). Because of this small coating thickness, a coating failure of the TiN coating (pitting) as early as about 1 h was observed. The pitting corrosion of the TiN coating at an early stage of immer- sion was captured by the Cor.Pot., PR and EIS measure- ments. In order to explain the corrosion mechanism of the TiN coating, the corrosion of the substrate and then the corrosion of the TiN coating at the early stage and later within 24 h were evaluated together with the EIS data and the EC modelling, Figures 8 and 9 in the next section. The EIS data of the Bode and Nyquist plots of the substrate clearly exhibited a one-time constant (particu- larly the phase angle vs the frequency plot in the Bode plots) during 24 h of exposure in a salt solution (Figure 8). It has been well known that any parallel RC circuit found in the EC represents a time constant (), corres- ponding to the characteristic frequency (c).18,19,22 Be- cause of one time constant observation in the Bode and Nyquist plots and the absence of pitting corrosion, the EIS model (Figure 8a) proposed for uncoated substrates to simulate the interacting of the electrolyte with the surface consisting of a solution resistance (Rsol.) and in parallel the total resistance of the passive layer (capaci- tive layer), Rpassive with constant phase elements (CPE) of the passive layer. Rpassive was the resistance of the passive layer. The fit parameters, i.e., Y0, n, Rsol. and Rpassive, were determined from the best non-linear least-square fit to the electrical circuit model with a goodness-of-fit value and they are given in Table 1 for selected times of exposure. Also, the fitted EIS plots were given in Figures 8a and 8b. The goodness of fit (2) for all the samples was in the range 1–7.6 · 10–3 (Table 1). The error in the Rpassive and Y0 was about 1 %. The value of Rpassive (Figure 13) calculated from the EIS data was the same as Rp calculated from the PR method (Figure 5) as expected.22 Y0, the admittance constant of the CPE, is a measure of the capacity (C).23 Y0 equals the capacity (C) when n = 1 for an ideal capacitor. The n value is related to the roughness and the inhomogeneity of the passive (capacitive) film and is less than 1 when the surface is rough.17,18,23 The n value was about 0.84 (Table 1), and constant during the exposure, indicating that the surface roughness (surface area) did not change during the immersion. The polarization resistance of the EP substrate (Rp and Rpassive) (Figure 13), indicated that the resistance of the passive layer increased from 7.5 k cm2 to 27 k cm2 with the exposure time. The Y0 was also calculated from the EC model and plotted in Figure 14. The Y0 decreased with time and the variation was logarithmic with time as Rp and Rpassive. The capacity values of the passive film determined from Y0 (presumably Cr2O3) (Figure 14) decreased logarithmically, indicating that the passive film thickened with the exposure time under assumption that the dielectric property of the passive I. KUCUK, C. SARIOGLU: PITTING CORROSION OF TiN-COATED STAINLESS STEEL IN 3 % NaCl SOLUTION Materiali in tehnologije / Materials and technology 49 (2015) 2, 183–192 189 Table 1: The EIS fit parameters of the substrate from the electrical circuit model (EC model): Y0, n, Rsol. and Rpassive were determined by the best non-linear least-square fit with the goodness of the fit value (2) Tabela 1: EIS-parametri podlage, pridobljeni iz modela elektri~nega tokokroga (EC-model): Y0, n, Rsol. in Rpassive, so bili dolo~eni z najbolj{im ujemanjem z nelinerano metodo najmanj{ih kvadratov in vrednostjo ujemanja (2) Immersion time (min) Rsol./k cm2 Rpassive/k cm2 Y0/μcm–2 sn  n Goodness of fit (2 · 103) 39 7.1 7.71 213.4 0.79 2.47 147 7.1 13.32 160.7 0.81 6.883 283 7.1 17.16 132.5 0.83 0.981 623 7.2 24.09 107.3 0.84 6.423 1439 7.2 26.48 82.4 0.84 7.556 Figure 13: The corrosion potential (Cor.Pot.), polarization resistance (Rp and Rpassive calculated from EIS data) for the substrate in 3 % NaCl solution as a function of immersion time Slika 13: Korozijski potencial (Cor.Pot.), polarizacijska upornost (Rp in Rpassive izra~unani iz EIS-podatkov) za podlago v 3-odstotni raztopini NaCl v odvisnosti od ~asa namakanja film did not change with time. The logarithmically increase in the polarization resistance (Rp and Rpassive) with time verified this assumption (Figure 13) since (Rp and Rpassive) and the current density must relate to the rate of the thickening (dx/dt).22 The capacitive (dielectric) layer of TiO2 formed on TiN as implied by EIS data was identified with the Mott-Shottky analysis as n-type semiconductive layer and that was represented by Y0. At an early stage of exposure, the transition from strong capacitive to less capacitive behaviour was observed in the EIS measure- ments (Figure 9). The transition for the TiN-coated sub- strate was a result of the pitting formation and growth. The transition was reproduced with three samples and took place between 1 h and 2 h of exposure for all the samples. Because of the pitting formation and growth at the substrate/TiN coating interface between 80 min and 95 min, after 40 min the EIS model used for the data in Figure 9 was changed to the model employed in18,19,24 for porous coatings and paints in order to include the pitting formation and the growth at the substrate/electrolyte interface. This model included two time constants (two RC) (Figure 9a).19,24 The goodness of the EIS data fit was in the range 1.4–19 · 10–3 (Table 2). It was higher at an early exposure time due to the dynamic change of the corrosion state (Table 2). As an example, three fits are shown in Figures 9a and 9b. The representative data especially during pitting were tabulated in Table 2. During the pitting formation and growth periods Rtotal (Rpore + Rcor.) plotted in Figure 15 as well as Rpore decreased while Y0 sub., (admittance at substrate/electro- lyte interface) increased (Table 2). Rcor. was the electron charge-transfer resistance at the substrate/electrolyte interface. The decrease in the pore resistance (Rpore) and the total resistance, and the increase in the capacity (through Y0 sub.) at the substrate interface implied the degradation of the coating at the interface. As observed visually, these degradations took place through the pitting formation and growth. After the pits grew to some extent, the total resistance increased slightly and stayed constant for some periods up to 12 h and then fluctuated due to the new pits being formed (five new pits were ob- served after 24 h) (Figures 5 and 15). The total polariza- tion resistance (Rtotal) and Rp were very similar and I. KUCUK, C. SARIOGLU: PITTING CORROSION OF TiN-COATED STAINLESS STEEL IN 3 % NaCl SOLUTION 190 Materiali in tehnologije / Materials and technology 49 (2015) 2, 183–192 Table 2: The EIS fit parameters of the TiN coating from the electrical circuit model (EC model): Y0 coat., Y0 sub., n, m, Rsol., Rpore and Rcor. were determined by the best non-linear least-square fit with goodness of the fit value (2) Tabela 2: EIS-parametri, pridobljeni za prevleko TiN iz modela elektri~nega tokokroga (EC model): Y0 coat., Y0 sub., n, m, Rsol., Rpore, in Rcor., so bili dolo~eni z najbolj{im ujemanjem z metodo nelineranih najmanj{ih kvadratov in vrednostjo ujemanja (2) Immersion time (min) Rsol./ cm 2 Rpore/k cm2 Rcor./k cm2 n m Y0 (coat.)/μcm–2sn  Y0 (sub.)/ μcm–2sn  Goodness of fit (2 · 103) 40 4.9 53.93 – – 0.83 90.7 – 19.14 90 5.1 0.34 33.65 0.75 0.9 60.3 45.47 13.25 140 5.1 0.17 43.40 0.72 0.93 54.6 59.73 13.82 240 5.1 0.06 53.75 0.72 0.96 40.6 83.13 1.418 648 5.0 0.07 58.41 0.71 0.97 38.9 83.60 1.633 1124 5.2 0.05 47.78 0.73 1 30.2 84.40 1.444 1464 5.1 0.05 48.06 0.73 1 31.5 86.53 1.440 Figure 15: Corrosion potential (Cor.Pot.), polarization resistance (Rp and Rtotal calculated from EIS data) for the TiN coating in 3 % NaCl solution as a function of immersion time Slika 15: Korozijski potencial (Cor.Pot.), polarizacijska upornost (Rp in Rtotal izra~unani iz EIS podatkov) za prevleko TiN v 3-odstotni raztopini NaCl v odvisnosti od ~asa namakanja Figure 14: The admittance (Y0) and polarization resistance (Rpassive calculated from EIS data) for the substrate in 3 % NaCl solution as a function of immersion time Slika 14: Admitanca (Y0) in polarizacijska upornost (Rpassive izra~u- nana iz podatkov EIS) za podlago v 3-odstotni raztopini NaCl v odvisnosti od ~asa namakanja followed the same trend (Figure 15), indicating that during the EIS measurement the corrosion state did not change significantly and caused a significant error (the PR measurement took 2.7 min, shorter than the 8 min EIS measurement). The Y0 coat. value, (admittance constant for TiO2 layer) decreased continuously, while the Y0 sub., for the substrate/solution interface stayed relatively constant for longer times (Table 2). After the pitting corrosion, the n value for the capacitive layer at substrate/electrolyte interface remained constant at low level in the range 0.72–0.76 (Table 2). The value of m for the capacitive coating on TiN was always at a high level in the range 0.9–1.0 after pitting formation, while at the beginning during the pitting formation it was 0.83, indicating that surface roughness was significant during the early expo- sure (during pitting) and at longer exposure times the surface became smoother. Recently, He25 studied the in-situ AFM of exposed TiN (1 μm layer) deposited by DC reactive magnetron sputtering on 304 stainless steel in a 3.5 % NaCl solution and observed decreasing in roughness with time in first exposure 60 min. They ex- plained this result with an in-situ observation of closing the pin holes and small pores, presumably by corrosion products. Even though the 2 values (Table 2), were high for the early exposure up to 140 min, these EIS data could be used to bring about the evolution of the corrosion of the TiN during the pitting formation and growth. Before visual observation of the pits based on their brownish colour on the surface (Figure 10) (formation of pitting state), the capacitive response of the surface layer (TiO2) was believed to be degraded by the penetration of the electrolyte through the defects in the coatings, pre- ferentially along the columnar grain boundaries, the pin holes, the large openings or the voids and droplets to the substrate/coating interface (Figures 3 and 4). The work of Cai26 supported this conclusion. Cai26 studied the effect of a post-deposition treatment of the TiN-coated steel and stainless-steel substrates with polymethyl methacrylate (PMMA) on corrosion in a 3.5 % NaCl solution. They found a significant improvement in the corrosion resistance because of effectively sealing the open voids or pores associated with the coatings. The pits were believed to form at the defect sites (columnar grain boundaries, pin holes, large openings or voids and droplets) by galvanic coupling between the substrate surface and the TiN coating surface. There was a driving force for the galvanic corrosion since the corrosion potential on the substrate surface was more active than the corrosion potential on the TiN surface (Figure 5) in agreement with the data reported by Mendibide,11 who measured a more noble potential of TiN on the glass surface compared to the steel substrate. The growth or propagation of pits is clearly docu- mented in Figures 10 and 11 and both were captured by the PR and EIS data (Figures 5, 7, 9 and 15). At the pit areas there was no Ti, indicating that the TiN coating during the pitting did not dissolve. The growth stage of the pit involved the growth of the pit area as a result of the detachment, cracking and spallation of the TiN coating due to the dissolution of the substrate at the TiN/substrate interface with time. Even though a number of pits formed on the TiN coated surface, the corrosion resistance after about 24 h was greater than the corrosion resistance of the substrate. This results indicated that the TiN was inherently resistant to corrosion due to the formation of the n-type semiconductor passive film (presumably TiO2), as identified by the Mott-Shottky analysis (Figure 12). 5 CONCLUSIONS 1. The pitting corrosion of the TiN is directly related to the coating defects and the coating structure. The TiN coatings deposited on the substrate consisted of cubic TiN and a passive n-type oxide (presumably TiO2) film with flat band potentials of –0.83 V vs. SCE, de- termined by the Mott-Shottky analysis. The coatings defects were columnar grain boundaries extending to the coating/substrate interface, the droplets, the pin holes and the porosities. These defects were prefe- rential sites for the pitting corrosion. 2. PR and EIS measurement captured the formation and growth of the pitting corrosion. The EIS model used supported the degradation of the coating through pitting, in agreement with visual observations. The corrosion resistance (Rp and Rtotal) decreased suddenly during the pitting corrosion. The pitting corrosion was believed to take place at the defect sites by galvanic corrosion, driven by the corrosion potential differences between the TiN surface and the substrate surface. 3. The corrosion resistance of the TiN-coated substrate was greater than the corrosion resistance of the substrate for about 24 h, even though the corrosion resistance of the substrate (Rp and Rpassive) increased logarithmically as the passive layer grew with time, as indicated by the decreasing Y0. This result indicated that the TiN coating possessed a significant inherent resistance to corrosion. 4. The PR and EIS measurements used together gave similar polarization resistance results, supporting the accuracy of the EIS data and the EC model used for the substrate and the TiN-coated substrate. Acknowledgement Marmara University is greatly acknowledged for its financial support through the Contract No: FEN-KPS- 080808-0178. I. KUCUK, C. SARIOGLU: PITTING CORROSION OF TiN-COATED STAINLESS STEEL IN 3 % NaCl SOLUTION Materiali in tehnologije / Materials and technology 49 (2015) 2, 183–192 191 6 REFERENCES 1 B. Warcholinski, A. Gilewicz, Surface Engineering, 27 (2011) 7, 491–497, doi:10.1179/026708410X12786785573355 2 Y. H. Cheng, T. Browne, B. Heckerman, C. Bowman, V. Gorokhov- sky, E. I. Meletis, Surface and Coatings Technology, 205 (2010) 1, 146–151, doi:10.1016/j.surfcoat.2010.06.023 3 J. Bujak, J. Walkowicz, J. Kusinski, Surface and Coatings Techno- logy, 180 (2004), 150–157, doi:10.1016/j.surfcoat.2003.10.058 4 A. Horling, L. Hultman, M. Oden, J. Sjolen, L. Karlsson, Surface and Coatings Technology, 191 (2005), 384–392, doi:10.1016/ j.surfcoat.2004.04.056 5 C. Ducros, V. Benevent, F. Sanchette, Surface and Coatings Tech- nology, 163-164 (2003), 681-688, doi:10.1016/S0257-8972(02) 00656-4 6 C. J. Tavares, L. Rebouta, B. Almeida, J. Bessa e Sousa, Surface and Coatings Technology, 100–101 (1998), 65–71, doi:10.1016/S0257- 8972(97)00589-6 7 C. J. Tavares, L. Rebouta, M. Andritschky, S. Ramos, Journal of Ma- terials Processing Technology, 92-93 (1999), 177–183, doi:10.1016/ S0924-0136(99)00126-0 8 L. Hultman, J. E. Sundgren, Structure/property relationships for hard coatings, In: R. F. Bunshah (ed.), Handbook of Hard Coatings, William Andrew Publishing, New York 2001, 108–180 9 H. A. Jehn, Surface and Coatings Technology, 125 (2000), 212–217, doi:10.1016/S0257-8972(99)00551-4 10 M. Fenker, M. Balzer, H. Kappl, Thin Solid Films, 515 (2006) 1, 27–32, doi:10.1016/j.tsf.2005.12.020 11 C. Mendibide, P. Steyer, J. P. Millet, Surface and Coatings Techno- logy, 200 (2005), 109–112, doi:10.1016/j.surfcoat.2005.02.060 12 M. Urgen, A. F. Cakir, Surface and Coatings Technology, 96 (1997), 236–244, doi:10.1016/S0257-8972(97)00123-0 13 M. A. M. Ibrahim, S. F. Korablov, M. Yoshimura, Corrosion Science, 44 (2002), 815–828, doi:10.1016/S0010-938X(01)00102-0 14 W. J. Chou, G. P. Yu, J. H. Huang, Corrosion Science, 43 (2001), 2023–2035, doi:10.1016/S0010-938X(01)00010-5 15 V. K. William Grips, H. C. Barshilia, V. Ezhil Selvi, Kalavati, K. S. Rajam, Thin Solid Films, 514 (2006), 204–211, doi:10.1016/j.tsf. 2006.03.008 16 L. Cunha, M. Andritschky, L. Rebouta, R. Silva, Thin Solid Films, 317 (1998), 351–355, doi:10.1016/S0040-6090(97)00624-X 17 I. Kucuk, C. Sarioglu, Mater. Tehnol., 49 (2015) 1, 19–26 18 C. Liu, Q. Bi, A. Leyland, A. Matthews, Corrosion Science, 45 (2003), 1243–1256, doi:10.1016/S0010-938X(02)00213-5 19 C. Liu, Q. Bi, A. Leyland, A. Matthews, Corrosion Science, 45 (2003), 1257–1273, doi:10.1016/S0010-938X(02)00214-7 20 S. Roy Morrison, Electrochemistry at Semiconductor and Oxidized Metal Electrodes, Plenum Press, New York 1980, doi:10.1007/ 978-1-4613-3144-5 21 S. Rudenja, C. Leygraf, J. Pan, P. Kulu, E. Talimets, V. Mikli, Sur- face and Coatings Technology, 114 (1999), 129–136, doi:10.1016/ S0257-8972(99)00033-X 22 D. A. Jones, Principles and Prevention of Corrosion, Prentice hall, Inc., 1992 23 C. H. Hsu, F. Mansfeld, Corrosion, 57 (2001) 9, 747–748, doi:10. 5006/1.3280607 24 Y. H. Yoo, D. P. Le, J. G. Kim, S. K. Kim, P. V. Vinh, Thin Solid Films, 516 (2008), 3544–3548, doi:10.1016/j.tsf.2007.08.069 25 C. He, J. Zhang, J. Wang, G. Ma, D. Zhao, Q. Cai, Applied Surface Science, 276 (2013), 667–671, doi:10.1016/j.apsusc.2013.03.151 26 F. Cai, Q. Yang, X. Huang, L. R. Zhao, Corrosion Engineering, Sci- ence and Technology, 46 (2011) 4, 368–374, doi:10.1179/ 147842209X12489567719626 I. KUCUK, C. SARIOGLU: PITTING CORROSION OF TiN-COATED STAINLESS STEEL IN 3 % NaCl SOLUTION 192 Materiali in tehnologije / Materials and technology 49 (2015) 2, 183–192 R. AZIM, M. T. ISLAM: DESIGN OF A WIDEBAND PLANAR ANTENNA ON AN EPOXY-RESIN- ... DESIGN OF A WIDEBAND PLANAR ANTENNA ON AN EPOXY-RESIN-REINFORCED WOVEN-GLASS MATERIAL [IROKOPASOVNA PLOSKOVNA ANTENA NA EPOKSI SMOLI, OJA^ANI S STEKLENIMI VLAKNI Rezaul Azim1, Mohammad Tariqul Islam2 1University of Chittagong, Chittagong 4331, Bangladesh 2Department of Electrical, Electronic & Systems Engineering, Universiti Kebangsaan Malaysia, 43600 UKM Bangi, Malaysia razim71@gmail.com Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2014-03-28 doi:10.17222/mit.2013.169 In this study, the design and prototyping of a compact planar antenna is presented for wideband applications. The designed transmission-line-fed antenna is composed of a rectangular radiating patch and a partial ground plane and is printed on both sides of a 1.6 mm thick, epoxy-resin-reinforced, woven-glass dielectric material. The antenna structure is planar, and its design is simple and easy to fabricate. Compared to the other substrate materials, the proposed antenna on epoxy-resin, woven-glass material could achieve a wider and less-expansive bandwidth. Experimental results show that the designed antenna could achieve an impedance bandwidth (return loss  –10 dB) from 2.99 GHz to 18.31 GHz (6.12: 144.33 %). Moreover, the antenna has a good gain and exhibits stable radiation patterns within the operating band. Details of the proposed antenna are presented and discussed. Keywords: wideband, microstrip patch, planar antenna, epoxy resin, fiberglass V tej {tudiji je predstavljena zgradba in izdelava kompaktne ploskovne antene za {irokopasovno uporabo. Zasnovana oddajno-sprejemna antena je sestavljena iz pravokotne sevalne in delno ozemljitvene ploskve ter natisnjena na obeh straneh 1,6 mm velike plo{~ice iz epoksidne smole, oja~ane z dielektri~nimi steklenimi vlakni. Antena je ploskovna, enostavne zgradbe in enostavna za izdelavo. V primerjavi z drugimi materiali podlage predlagana zasnova na epoksidni smoli s tkanino iz steklenih vlaken omogo~a ve~jo pasovno {irino in manj{o ekspanzivnost. Eksperimentalni rezultati ka`ejo, da predlagana antena lahko dose`e impedan~no pasovno {irino (povratna izguba  –10 dB) od 2,99 GHz do 18,31 GHz (6,12: 144,33 %). Poleg tega ima antena dober izkoristek in ka`e stabilne vzorce sevanja v delovnem pasu. Predstavljene in obravnavane so podrobnosti o predlo`enem oblikovanju antene. Klju~ne besede: {irokopasovni, mikrotrakasta ploskev, ploskovna antena, epoksidna smola, steklena vlakna 1 INTRODUCTION In wireless communications technology the necessity for wide and multi-band antennas is increasing rapidly due to the need to support more users and to provide in- formation with higher data-transmission rates. Microstrip antennas are one the most suitable structures due their low profile and ease of fabrication. Compared to conven- tional, three-dimensional types of antennas, planar microstrip antennas printed on a piece of printed-circuit board have become very popular in modern wireless communications because they can be easily embedded into wireless devices or integrated with other RF cir- cuitry. Usually, a planar design can be used to reduce the volumetric size of a wideband antenna by replacing the three-dimensional radiating elements with their planar versions.1–3 Different types of planar antennas have already been proposed for wideband applications. A variety of dielec- tric materials have been used for the design and proto- typing of these antennas. A dielectric material used for the design of wideband antennas is required to feature a higher permittivity and lower dissipation factor.4 Mate- rials with a lower permittivity are good insulators for lower-frequency signals requiring high isolation in den- sely packed circuits, such as mobile communications.5 On the other hand, materials with a higher dielectric constant have a greater capability to store charge and produce larger electromagnetic fields, but limited isola- tion between the conductors.6 Moreover, by using a material with a higher permittivity a compact antenna can be designed that is capable of achieving a very wide operating band.7 For example, in8 a miniaturized, modi- fied, circular patch antenna was designed on a ceramic- polytetrafluroethylene (PTFE) composite material. With an overall size of 0.22 × 0.29 × 0.03 , the proposed antenna achieved multi-band characteristics. However, the antenna failed to fulfill the requirement for a wideband antenna having triple operating bands of 5–6.3 GHz, 9.1–9.6 GHz and 10.7–11 GHz. In9, a wideband, pentagon-shaped, planar microstrip slot antenna was designed on an epoxy-resin composite material. Combining the pentagon-shaped slot, feed line, and pentagon stub, the antenna obtained an impedance bandwidth of 124 %. However, its use in portable communication devices was limited due to its large ground plane. Ullah et al.10 proposed a double L-shaped multi-band patch antenna on a polymer-resin substrate Materiali in tehnologije / Materials and technology 49 (2015) 2, 193–196 193 UDK 621.396.67:621.375.5 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)193(2015) material. By introducing two L-shaped slots in the rectangular patch, the designed antenna could achieve a dual operating band centered at 4.85 GHz and 8.1 GHz, and thus not suitable for broadband wireless communi- cations. In this study, a simple planar monopole antenna that achieves a physically compact planar profile having sufficient impedance bandwidth and a stable radiation pattern is proposed for wideband applications. The antenna consists of a rectangular radiating patch and a partial ground plane. The transmission line-fed radiating patch is printed on one side of an epoxy-matrix-rein- forced, woven-glass material, while the partial ground plane is printed on the other side. The epoxy-matrix- reinforced, woven-glass material is a popular and versatile high-pressure thermo set plastic laminate grade with a low dissipation factor; it is the most commonly used electrical insulator that possesses a good mecha- nical strength and a nearly zero water-absorption coefficient. It is observed that the radiating patch of the proposed design has a strong coupling with the ground plane and the antenna, designed on polymer-resin com- posite material, is capable of supporting multiple reso- nance modes. The overlapping of these multiple reso- nance modes leads to the characterization of a wideband ranging from 2.99 GHz to 18.31 GHz. The simple structure, ease of fabrication, low cost, wide operating band and stable radiation patterns make the proposed antenna suitable for use in WiMAX, WLAN, C-band, UWB and X-band applications. 2 ANTENNA DESIGN The design layout of the proposed wideband antenna is illustrated in Figure 1. The antenna is designed and analysed using the method of the moment based, full- wave electromagnetic field solver IE3D. The antenna is printed on both sides of a 1.6 mm thick double-layer dielectric substrate. The dielectric material consists of an epoxy-matrix-reinforced woven glass. The fiberglass in the composition is 60 %, while the epoxy resin contri- butes 40 % of the composition. This composition of epoxy resin and fiberglass varies in the thickness and is direction dependent. One of the attractive characteristics of polymer-resin composites is that they can be shaped and reshaped repeatedly without losing their material characteristics.11 Due to the ease of fabrication, design flexibility, low manufacturing costs and market avail- ability, the epoxy-matrix-reinforced, woven-glass mate- rial has become popular in the design of communication devices. A microstrip transmission line fed rectangular patch of 13.5 mm × 14.5 mm is printed on one side of the dielectric material, while the partial ground plane with a side length of 5.5 mm is printed on the opposite side. The value of the width and length of the transmission line is set at 2.75 mm and 6 mm, respectively, so that the input impedance of the feeding of the antenna becomes equivalent to 50 . A copper cladding 35 μm is used to metalise the patch, feed line and the ground plane of the proposed antenna. The overall dimension of the designed wideband antenna is optimized to a compact size of 29 mm × 20.5 mm, which is much smaller than the antennas proposed in9,12–14 and very suitable to be integrated into portable communication devices. To investigate the effect of different dielectric mate- rials on the performance of the proposed antenna, a parametric study was conducted. The properties of the dielectric materials are tabulated in Table 1, while their effect on the return-loss characteristics is depicted in Figure 2. It is clear from the plot that the proposed antenna with an epoxy-matrix-reinforced, woven-glass material exhibits a wider operating band than the glass and ceramic PTFE. Although the antenna with the cera- mic PTFE composite material achieved a lower operating frequency because of the high dielectric constant, its bandwidth is narrower compared to glass PTFE and epoxy resin and is extremely expensive compared to the epoxy-resin dielectric material. Table 1: Properties of the dielectric materials Tabela 1: Lastnosti izolacijskih materialov Dielectric material Permittivity Loss tangent Glass PTFE 2.2 0.0009 Ceramic PTFE 10.2 0.002 Epoxy resin 4.6 0.02 The return-loss characteristics in Figure 2 show that the proposed antenna with epoxy resin material exhibits six resonances across the operating band, of which the first appears at 3.4 GHz, the second at about 6.45 GHz, the third at 9.7 GHz, the fourth at 11.0 GHz, the fifth at 12.75 GHz and the final resonance is at 16.17 GHz. The figure clearly indicates that the overlapping of these R. AZIM, M. T. ISLAM: DESIGN OF A WIDEBAND PLANAR ANTENNA ON AN EPOXY-RESIN- ... 194 Materiali in tehnologije / Materials and technology 49 (2015) 2, 193–196 Figure 1: a) Top and b) bottom views of the proposed design Slika 1: Videz predlagane zgradbe: a) zgoraj in b) spodaj resonances that are closely spaced across the spectrum leads to a wide operating bandwidth, ranging from 2.96 GHz to 18.31 GHz. At the first resonance frequency, when the antenna size is smaller than the corresponding wavelength, the electromagnetic signal can couple to the antenna size and can operate in the stationary wave mode. As the frequency increases, the antenna starts to operate in the mixed stationary wave and the travelling wave modes. At the higher edge frequency of the ope- rating band, the antenna size becomes larger, corres- ponding to the respective wavelength, and the electro- magnetic signals have to travel a long distance, resulting in a dominating travelling wave mode, as shown in Figure 3, which depicts the input impedance of the pro- posed antenna. 3 RESULTS AND DISCUSSION The return-loss characteristic of the realised antenna was measured in an anechoic chamber using an Agilent E8362C vector network analyser. Figure 4 plots the measured and simulated return-loss curves. The simu- lated –10 dB return-loss bandwidth ranges from 2.96 GHz to 18.31 GHz, equivalent to a fractional bandwidth of 144.33 %. This wideband characteristic of the pro- posed planar antenna is confirmed in measurements, with only a small shift of the lower edge frequency to 2.99 GHz. Despite a very small size, the proposed antenna achieved a sufficient operating band to cover the WiMAX, WLAN, C-band, UWB and X-frequency bands. Although there is a disparity between the measured and simulated resonances that can possibly be attributed to the manufacturing tolerances and imperfect soldering effects of the SMA connector, the measured resonance frequencies are nearly identical to the simulated fre- quency. This mismatch may also be due to the effect of the RF feeding cable, which is used in the measure- ments, but not considered during the simulation. The ripples observed in the measured result may be caused by the current drain from the conducting ground plane to the outer shield of the RF feeding cable, which is not presumed during the simulation. The peak gain of an epoxy-resin-based, wide-band antenna is illustrated in Figure 5. Despite a very com- pact dimension, the designed antenna exhibits a good gain. As the antenna would be used for short-distance wireless communication, the achieved peak gain is within the acceptable limits. The gain of the proposed R. AZIM, M. T. ISLAM: DESIGN OF A WIDEBAND PLANAR ANTENNA ON AN EPOXY-RESIN- ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 193–196 195 Figure 4: Measured and simulated return losses Slika 4: Izmerjene in simulirane povratne izgube Figure 2: Return-loss characteristics for different materials Slika 2: Zna~ilnosti povratnih izgub za razli~ne materiale Figure 5: Gain of the proposed wideband antenna Slika 5: Izkoristek predlagane {irokopasovne antene Figure 3: Real and imaginary parts of the input impedance Slika 3: Realni in navidezni del vhodne impedance antenna could be improved using ceramic PTFE micro- wave substrate material rather than epoxy-resin material. However, the use of the ceramic PTFE increases the cost of the antenna and hence this option is not considered in the design. The E- and H-plane radiation patterns of the pro- posed antenna at (3.2, 6.1 and 12.7) GHz are depicted in Figure 6. At lower frequencies, the antenna exhibits a bidirectional radiation pattern for both the E and H planes, and the patterns are approximately the same as the pattern of a typical monopole antenna. As the fre- quency increases, a higher-order harmonic is introduced to the patterns, and both the H and E planes become mo- re directional, but still retain their bidirectionality. Some dips are observed, mainly at higher frequencies, and might be because the microstrip feed line is printed directly below the partial ground plane and might also be caused by the feed connector. However, the radiation patterns are remarkably stable throughout the operating band. 4 CONCLUSIONS A compact planar microstrip antenna is designed and fabricated for a wideband application. The proposed antenna consists of a partial ground plane and a trans- mission-line-fed rectangular radiating element. The antenna design is analysed and optimised by the method of moment-based software IE3D and is verified by means of a prototype. The antenna is designed and fabri- cated on epoxy-matrix-reinforced, woven-glass material. Compared to the other dielectric materials, the proposed antenna with an epoxy matrix reinforced woven-glass material exhibits better performance in terms of band- width, return-loss, gain and radiation patterns. Experi- mental results show that the proposed antenna could achieve an impedance bandwidth from 2.99 GHz to 18.31 GHz (6.12: 1, 15.32 GHz) for a return loss of  –10 dB. Moreover, the antenna achieved a good gain and stable radiation patterns. All these features of the proposed epoxy-resin-based antenna make it a worthy candidate for wideband applications in portable communication devices. 5 REFERENCES 1 M. J. Ammann, Z. N. Chen, Wideband monopole antennas for multi- band wireless systems, IEEE Antennas and Propagation Magazine, 45 (2003), 146–150, doi:10.1109/MAP.2003.1203133 2 X. H. Wu, Z. N. Chen, Comparison of planar dipoles in UWB applications, IEEE Transactions on Antennas and Propagations, 53 (2005), 1973–1983, doi:10.1109/TAP.2005.848471 3 M. Samsuzzaman, M. T. Islam, J. S. Mandeep, N. Misran, Printed wide-slot antenna design with bandwidth and gain enhancement on low-cost substrate, The Scientific World Journal 2014 (2014), Article ID 804068, 10 pages, doi:10.1155/2014/804068 4 K. Oohira, Development of an antenna material based on rubber that has flexibility and high impact resistance, NTN Technical Review, 76 (2008), 58–63 5 M. Samsuzzaman, M. T. Islam, J. S. Mandeep, Parametric analysis of a glass-micro fibre-reinforced PTFE material, multiband, patch- structure antenna for satellite applications, Optoelectronics and Advanced Materials–Rapid Communications, 7 (2013), 760–769 6 A. Aguayo, Analyzing advances in antenna materials, Antenna Sys- tems & Technology, 12 (2010), 14–15 7 M. H. Ullah, M. T. Islam, A compact square loop patch antenna on high dielectric ceramic–PTFE composite material, Applied Physics A, 113 (2013), 185–193, doi:10.1007/s00339-012-7511-4 8 M. H. Ullah, M. T. Islam, Miniaturized modified circular patch mo- nopole antenna on ceramic-polytetrafluroethylene composite mate- rial substrate, Journal of Computational Electronics, 13 (2014), 211–216, doi:10.1007/s10825-013-0501-8 9 S. K. Rajgopal, S. K. Sharma, Investigations on ultrawideband pen- tagon shape microstrip slot antenna for wireless communications, IEEE Transactions on Antennas and Propagations, 57 (2009), 1353–1359, doi:10.1109/TAP.2009.2016694 10 M. H. Ullah, M. T. Islam, J. S. Mandeep, N. Misran, A new double L-shape multiband patch antenna on polymer resin material substrate, Applied Physics A: Materials Science & Processing, 110 (2012), 199–205, doi:10.1007/s00339-012-7114-0 11 I. Yarovskya, E. Evansb, Computer simulation of structure and pro- perties of crosslinked polymers: application to epoxy resins, Poly- mer, 43 (2002), 963–969, doi:10.1016/S0032-3861(01)00634-6 12 R. Azim, M. T. Islam, N. Misran, Microstrip line-fed printed planar monopole antenna for UWB applications, Arab Journal for Science and Engineering, 38 (2013), 2415–2422, doi:10.1007/s13369-013- 0553-x 13 L. Liu, S. W. Cheung, R. Azim, M. T. Islam, A compact circular-ring antenna for ultra-wideband applications, Microwave and Optical Technology Letters, 53 (2011), 2283–2288, doi:10.1002/mop 14 R. Azim, M. T. Islam, N. Misran, A. T. Mobashsher, Compact UWB planar antenna for broadband applications, Informacije MIDEM, 41 (2011), 37–40 R. AZIM, M. T. ISLAM: DESIGN OF A WIDEBAND PLANAR ANTENNA ON AN EPOXY-RESIN- ... 196 Materiali in tehnologije / Materials and technology 49 (2015) 2, 193–196 Figure 6: Radiation patterns at different frequencies (Solid line: co-polarised field, crossed line: cross-polarised field) Slika 6: Vzorci sevanja pri razli~nih frekvencah (polna ~rta: kopolari- zirano polje, pre~no: pre~no polarizirano polje) W. OZGOWICZ et al.: INFLUENCE OF THE STRAIN RATE ON THE PLC EFFECT AND ACOUSTIC EMISSION ... INFLUENCE OF THE STRAIN RATE ON THE PLC EFFECT AND ACOUSTIC EMISSION IN SINGLE CRYSTALS OF THE CuZn30 ALLOY COMPRESSED AT AN ELEVATED TEMPERATURE VPLIV HITROSTI DEFORMACIJE NA POJAV PLC IN AKUSTI^NO EMISIJO MONOKRISTALOV ZLITINE CuZn30, STISKANE PRI POVI[ANI TEMPERATURI Wojciech Ozgowicz1, Barbara Grzegorczyk1, Andrzej Pawe³ek2, Andrzej Pi¹tkowski2, Zbigniew Ranachowski3 1Institute of Engineering Materials and Biomaterials of the Silesian University of Technology, Konarskiego St. 18A, 44-100 Gliwice, Poland 2Institute of Metallurgy and Materials Science of Polish Academy of Sciences, Reymonta St. 25, 30-059 Cracow, Poland 3Institute of Fundamental Technological Research of the Polish Academy of Sciences, Pawinskiego St. 5B, 02-106 Warsaw, Poland barbara.grzegorczyk@polsl.pl Prejem rokopisa – received: 2013-10-01; sprejem za objavo – accepted for publication: 2014-02-25 doi:10.17222/mit.2013.195 The purpose of these investigations was to determine the effect of the strain rate on the phenomenon of a heterogeneous plastic deformation of the Portevin–Le Chatelier type while testing free compression of CuZn30 single crystals with a crystallographic orientation of [139] at 300 °C. Moreover, the relations between the work-hardening curve – displaying the PLC effect and the characteristics of the signals of the acoustic emission generated in the uniaxial-compression test were determined. It was found that the process of plastic deformation of the tested single crystals in the analyzed range of the frequencies up to 35 kHz generates differentiated sources of acoustic-energy emission, mainly the impulsive emission generated by signal events, correlated with the oscillations of the stresses on the work-hardening curves – . The strain rate mainly causes the changes in the intensity of the oscillation typical for the PLC effect. Keywords: plastic strain, Portevin–Le Chatelier effect (PLC), single crystals, copper alloys, compression test, acoustic emission (AE) Namen raziskav je bil ugotoviti vpliv hitrosti deformiranja pri preizkusu stiskanja monokristalov CuZn30 s kristalografsko orientacijo [139] pri 300 °C na pojav heterogene plasti~ne deformacije vrste Portevin-Le Chatelier. Poleg tega so bile ugotovljene {e odvisnosti med krivuljo deformacijskega utrjevanja – , ki ka`e pojav PLC, in zna~ilnostmi signalov akusti~ne emisije pri enoosnem tla~nem preizkusu. Ugotovljeno je, da proces plasti~ne deformacije preizku{anih kristalov v analiziranem podro~ju frekvenc do 35 kHz proizvaja diferencirane vire emisije akusti~ne energije, predvsem impulzivne emisije signalov, ki se skladajo z oscilacijo napetosti na krivulji deformacijskega utrjevanja – . Hitrost deformacije se ka`e predvsem v spremembi intenzitete oscilacij, zna~ilnih za PLC-pojav. Klju~ne besede: plasti~na deformacija, Portevin-Le Chatelierov pojav (PLC), monokristali, zlitine bakra, tla~ni preizkus, akusti~na emisija (AE) 1 INTRODUCTION In order to determine the kinetic relations and struc- tural changes, conditioning the mechanisms of plastic deformation it is necessary to know the process and understand the essence of various phenomena occurring in the course of plastic deformations. In many alloys a heterogeneous plastic deformation in the form of irregu- larities on the work-hardening curve is observed during tensile (or compression) tests. The earliest investigations concerning this phenomenon in medium-carbon steel and aluminium were published by Portevin and Le Chatelier in 1923, hence, it is called the PLC effect. Although the effect of the instability of plastic deformations of the PLC type has been known and investigated for nearly one hundred years, it is still not fully recognized and explained.1–7 Generally, this effect was investigated on the basis of material factors, taking into account the microstructural conditions of an initiation of a localized plastic deformation and the rheological factors connected with the mechanics of plastic deformations at various thermodynamic and physico-chemical conditions. One of the more recent methods of analyzing the phenomena occurring in the course of a plastic deforma- tion is the acoustic emission (AE). This method consists of a detection and analysis of the acoustic signal emitted by the material while it is being mechanically loaded. The acoustic signal is a result of the propagation of elastic waves generated in the material due to a fast release of the energy accumulated in this material. The shape of an AE signal is affected by many factors such as the chemical composition and macrostructure of the tested material, the kind of the applied heat treatment,8,9 the temperature and strain rate ( ),10 the grain size,11 the texture and the state of precipitations.12,13 The inve- stigations based on AE measurements are characterized by non-invasiveness and an incomparably high sensiti- vity in recording physical phenomena in comparison with the other methods of investigations. Materiali in tehnologije / Materials and technology 49 (2015) 2, 197–202 197 UDK 669.35’5:620.173:548.55 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)197(2015) The aim of this paper is to determine the influence of the strain rate in the range of 10–5–10–1 s–1 on a non-ho- mogeneous plastic deformation, the so-called Porte- vin-Le Chatelier effect, in a single-crystalline CuZn30 alloy with the crystallographic orientation of [1 39], compressed at a temperature of 300 °C. The additional aim is to determine the relation between the behaviour of the signals of acoustic emissions generated in the course of a plastic deformation at elevated temperature and the shape of the work-hardening curves – in the range of the PLC effect. 2 EXPERIMENTAL PROCEDURE The investigated material was a single-crystalline CuZn30 alloy in the form of a rod with a diameter of 3.8 mm, a length of about 200 mm and crystallographic orientation of [1 39], the chemical composition of which is shown in Table 1. Single crystals of this alloy were obtained with the Bridgmen’s method in a vertical, electrical tubular heating furnace by displacing the temperature gradient of the zone in the furnace in relation to the crucible (a quartz tube), with a charge in an atmosphere of inert gas. In order to realize the purpose of these investigations the following procedures had to be carried out: • mechanical tests of free compression of single crystals at 300 °C and a strain rate of 10–5–10–1 s–1, • investigations concerning the plastic deformation of single crystals by means of acoustic emission. Compression tests of single crystals were performed at a temperature of 300 °C and at  amounting to 10–5–10–1 s–1 on a universal testing machine INSTRON 3382 equipped with a duct die including heating ele- ments and a quartz outlet of the waveguide with an AE sensor (Figure 1). For the AE measurements the follow- ing setting was used: the acoustic-emission wideband sensor of type WD (PAC) provided by a company called Physical Acoustic Corp., a low-noise amplifier set to 66 dB of gain, background-noise discrimination at 0.75 V and 12-bit digital-to-analogue conversion at a rate of samples 88.2 · 103 s–1. The settings used for the AE acquisition had been adjusted experimentally during the W. OZGOWICZ et al.: INFLUENCE OF THE STRAIN RATE ON THE PLC EFFECT AND ACOUSTIC EMISSION ... 198 Materiali in tehnologije / Materials and technology 49 (2015) 2, 197–202 Figure 2: Simplified block diagram of the measurement and AE recording system Slika 2: Poenostavljeni sestav merilnega in zapisovalnega sistema AE Figure 1: Duct-die block for deforming the samples at an elevated temperature with an AE probe: 1 – die-block body, 2 – guide bar, 3 – sample, 4 – punch, 5 – wave guide, 6 – AE probe, 7 – duct, 8 – heating elements, 9 – foamed polystyrene, 10 – tack bolt Slika 1: Orodje s kanalom za deformiranje vzorcev pri povi{anih temperaturah z uporabo AE-sonde: 1 – telo orodja, 2 – pali~asto vodilo, 3 – vzorec, 4 – tla~ilna palica, 5 – valovni vodnik, 6 – AE-sonda, 7 – vodilni kanal, 8 – grelni element, 9 – penjeni polistiren, 10 – spenjalni vijak Table 1: Chemical composition of the monocrystalline alloy applied in the investigations Tabela 1: Kemijska sestava monokristalne zlitine, uporabljene pri preiskavah No. Determination of the alloy and the kind ofthe applied analysis of CuZn30 Chemical composition, mass fractions, w/% Zn Fe Al Ni Sn Pb Cu 1 smelting analysis of CuZn30 ingot 30.3 0.024 0.039 0.024 0.003 0.01 bal. 2 CuZn30 PN-EN 12163:2002 28.3–30.3 max 0.05 max 0.02 max 0.3 max 0.1 max 0.05 bal. previous tests dealing with the tension and compression of metal samples.14,15 The final deformation of the sample after the compression test amounted to about 50 %. The values of the forces within the entire range of the measurements were recorded with an accuracy of up to 0.5 %. The acoustic emission was measured during the compression tests of microcrystalline samples. In order to reduce the coefficient of friction between the butting face of the compressed sample and the steel punch, the samples were coated with a strip of Teflon. The block diagram of the measurement, the system and the recording of AE are illustrated in Figure 2. The AE measurement system was connected with the recording system of the testing machine. The tests were carried out in the Accredited Laboratory of the Strength of Materials, the Polish Academy of Science, Cracow. 3 RESULTS AND DISCUSSION The influence of the strain rate in the range of 10–5–10–1 s–1 on the mechanical characteristics – of the CuZn30 single crystals with the initial crystallo- graphic orientation of [1 39] at the constant temperature of the compression amounting to 300 °C is presented in Figure 3 and Table 2. It was found that the strain rate does not cause any changes in the general shape of the work-hardening curves of the investigated single crystals, but causing a considerable effect on the level of the stresses at the yield point and intensive oscillation of the stresses on these curves. When the strain rate grows, a tendency toward decreased values of the actual stresses could be observed, particularly during the initial stage of the deformation (15–20 %). Moreover, a tendency toward decreased oscillation of the stresses on the analytical work-hardening curves was observed. At the investigated temperature of the compression, the strain rate considerably influences the intensity of the oscillation of the stresses characterized by an occurrence of the PLC effect (Figure 3). The conditions of the plastic instability of the deformed single crystals are distinct in the case of low values of  (10–5–10–4 s–1) but they fade at a medium strain rate of about 10–3 s–1 and do not occur when  >10–2–10–1 s–1. It was also found that the strain rate only slightly affects the type of the oscillation of the stress in the described conditions of deformation. According to the classification by Brindley and Worthington16,17 three fundamental types of the oscil- lation of stresses are to be distinguished. Type A occurs at a low temperature of a tensile test, characterized by an increase, followed by a sudden drop in the force, such as the oscillation of the forces occurring periodically. At elevated temperatures type-B deformation oscillations occur that are smaller and irregular. These oscillations are symmetrical in relation to the level of the work- hardening curve. Type C occurs at the highest tempe- rature of a plastic deformation. These oscillations are characterized by a decrease of the force in relation to the level of the work-hardening curve and they are often separated from the area of homogeneous plastic defor- mation. According to this system, on the compression curves ( – ) of the tested single crystals in the range of deformation (5–15 %), type-B oscillations were detected. The initiation of the PLC effect, conditioned by the critical strain ( c), was observed on the work-harden- ing curves in the range of (1.2–2 %). In most cases the initiation of the oscillations on the curves of the com- pression coincides with the yield point of the material. The value of c proved to be rather independent of the strain rate. In order to quantitatively describe the serra- tion on the stress-strain curves, the analytic parameters used in quantitative fractography were applied.11 One of them, the so-called coefficient of development of the RL line, was derived from the relation of RL = L/L’, where L is the length of a given segment of the – line and L’ is the length of its projection. An approximately normal projection on the stress-strain curves was used. The geometrical analysis also indicates that for the serration occurring with a particular frequency (f), RL can be calculated with equation RL ≅ −( )A/f 1 , where A is the average amplitude of serration. The parameters calcu- lated in compliance with the MATLAB software are gathered in Table 2. It was found that the higher the values of coefficient RL (at a higher amplitude and Materiali in tehnologije / Materials and technology 49 (2015) 2, 197–202 199 W. OZGOWICZ et al.: INFLUENCE OF THE STRAIN RATE ON THE PLC EFFECT AND ACOUSTIC EMISSION ... Figure 3: Influence of the strain rate on the shape of curves – of CuZn30 single crystals with the initial orientation of [139], com- pressed at 300 °C Slika 3: Vpliv stopnje deformacije na obliko krivulje – monokri- stalov CuZn30 z za~etno orientacijo [139], stisnjenih pri 300 °C Table 2: Specification of the coefficients of the shape of – curves in CuZn30 single crystals with the orientation of [139] compressed in the range of deformations (5–15 %) Tabela 2: Pregled koeficientov oblike krivulje – v monokristalih CuZn30 z orientacijo [139], stisnjenih v obmo~ju deformacij (5–15 %) No. Temperature of defor- mation (°C)  /s–1 c/% Coefficient of the shape of – curves with the PLC effect RL A f 1 300 10–2 2.7 6.0 1.5 9.2 2 10–3 2.6 4.1 2.9 12.2 3 10–4 1.6 43.7 3.8 166.2 4 10–5 1.2 71.6 5.5 395.2 frequency of serration), the larger is the instability of the PLC plastic deformation (Table 2). The highest values of coefficient RL (about 72) in the range of deformations (5–15 %) as well as the highest value of the amplitude of stresses (calculated as 2 · A) amounting to about 11 MPa were determined for the samples compressed at the temperature of 300 °C and with  amounting to about 10–5 s–1. The obtained values confirm the qualitative description of the compression curves. The obtained results allow us to maintain that in the tested single crystals the PLC effect is a result of dyna- mic strain aging (DSA), which is an interaction between the sliding dislocations and free atoms. The sources of these interactions, complying with the dynamic-disloca- tion model of the PLC effect, are the multiplications of dislocations in the course of being affected by Frank- Read sources. The occurrence of DSA is conditioned by the rate of the migration of foreign atoms constituting the Cottrell atmosphere. The effects of the atoms of the alloy and the impurities with dislocations are responsible for retarding the dislocations, and thus also for hardening the alloy. If, therefore, a given strain rate is "imposed", we must also apply a stress that exceeds the resistance to the motion of dislocations, including the resistance resulting from the effect of the atoms of the alloy with dislocations. In order to explain, in a simple way, the qualitative model of the changes in the stress in the hardening diagram, we must assume that the break-away of a dislocation from the alloy atoms reduces the stress. This is indispensable for a further displacement of the dislocation by the value of this effect. The moment of a break-off of a dislocation from the blocking atoms involves a sudden drop in the force exciting the deforma- tion.18–20 According to the dislocation-dynamic model of the PLC effect, every local drop in the loading force recorded on the – curves is connected with unlocking the dislocation sources in a certain localized area of the sample. If there is a high concentration of internal stresses, the adjacent sources of dislocation will, succes- sively, become unlocked due to these stresses. Therefore, the unlocked sources of dislocation operate in a state of being considerably overcharged, so that the dynamically generated dislocation increases. Consequently, this involves a formation of sliding bands, which then propa- gate until the time of waiting (tw) again reaches the value of the time of aging (ta). Then, all the sources of disloca- tion are effectively locked by the Cottrell atmosphere,21 after which the process of unlocking starts again. The results of the measurements of the acoustic emis- sion (AE) during the compression testing of CuZn30 single crystals with the orientation of [1 39] at the temperature of 300 °C and the strain rate of 10–5–10–1 s–1 are gathered in Figures 4 to 6 and Table 3. In the inve- stigations of AE the descriptors of the recorded signals are usually added, such as: • the count rate over the preset noise-discrimination level • the signal amplitude within an AE event • the energy of an AE event, understood as a half of the product of the AE-event duration and the squared amplitude of the AE event. Table 3: Specification of the descriptors of the AE signals in the Re range concerning the investigated parameters of compression Tabela 3: Pregled opisovalcev AE-signalov v obmo~ju Re preiskova- nih parametrov pri tla~nem preizkusu No. Temperature of deforma- tion (°C) Strain rate  /s–1 Sum of AE events in the Re range Average energy of the AE events in the Re range (pJ) 1 300 10–4 14539 20.3 2 10–3 8903 42 3 10–2 6714 68 4 10–1 816 30 The last parameter was applied in the presented research because the squared signal amplitude combined with the event duration applied in the formula for data processing distinguishes well the signals caused by material effects from the unwanted signals caused by the drive of the loading machine and other system noises. This is because the unwanted AE sources have different spectral characteristics than the desired ones. In most compression tests of the investigated single crystals a distinct growth of the activity of the AE energy was recorded during the initial stage of their compres- sion and in the area of passing from the elastic to the plastic range (Figures 4 and 5). This growth depends on the value of  when deformations amount to about 2–5 %. An increase in AE is in both cases characterized by a more or less wide maximum change in the energy of the signal, after which AE reaches its minimum. The observed increase in the energy of the AE activity in the initial stages of the – curves may, however, be caused, among others, by mechanical factors due to the friction and the matching of the sample to the compress- W. OZGOWICZ et al.: INFLUENCE OF THE STRAIN RATE ON THE PLC EFFECT AND ACOUSTIC EMISSION ... 200 Materiali in tehnologije / Materials and technology 49 (2015) 2, 197–202 Figure 4: Dependence of the AE energy on the PLC effect recorded on the work-hardening curve during the compression testing of mono- crystalline CuZn30 alloy at 300 °C and  up to about 10–4 s–1 Slika 4: Odvisnost energije AE od PLC-u~inka, zabele`enega na kri- vulji preoblikovalnega utrjevanja pri tla~nem preizkusu monokristala zlitine CuZn30 pri 300 °C in  do 10–4 s–1 ing punch of the testing machine. In the range of the yield point, however, where AE displays a merely physi- cal aspect, the increasing AE activity is undoubtedly connected with the processes of dislocation. The AE level in this stage of the hardening of the alloy was found to be much more intensive than in the advanced stages of deformation occurring in the cases of greater drafts. The sum and mean energy of AE depending on the strain rate during the compression test are gathered in Table 3. It was found that in the  range from about 10–4 s–1 to about 10–2 s–1 the sum of events in the range of the yield point (Re) increases with the slowing down of the strain rate. The opposite is the case when the energy of a single event reaches its mean value. A higher average value of the energy of the event, except for  amounting to about 10–1 s–1, corresponds to a smaller sum of events. The minimum sum of events of about 816 also corresponds to a low value (30 pJ) of the average energy of AE events. The highest sum of events occurred in the case of the single crystals deformed at the temperature of 300 °C with  amounting to about 10–4 s–1 and the mean value of the energy of event of about 20 pJ. It was found that in the first stage of the hardening (the stage of easy sliding) of the single crystals that do not display any PLC effect (Figure 5), the AE level is lower, growing with the increase in the rate of deforma- tion, whereas in the single crystals displaying the PLC effect, the AE level is higher and its characteristics are more complex. A change in the strain rate at a given temperature does not involve any essential changes in the level of the frequency of AE. In the second stage of the hardening within the  range of (10–5–10–1 s–1) the AE activity is greater, particularly during the initial stage. These are mostly single samples of the AE signal cha- racterized by differentiated energy or a continuously increasing level of the AE energy, which grows with the increasing rate of deformation, independent of the occurrence of the PLC effect in the course of the com- pression of single crystals. When a work-hardening curve is of a parabolic cha- racter (the third stage of hardening), AE appears in the form of strong cumulative maximum values of the energy changes in the signal in time, particularly at the beginning. Irrespective of the appearance of the PLC effect in the course of the compression testing of single crystals, it was found that in the final stage of the defor- mation there is also a range, in which the AE energy is intensified, although less intensive with respect to the number of pulses than in the Re range. The observed correlations between the AE behaviour and the exerted compressive force and the evolution of the microstruc- ture may be satisfactorily explained with the qualitative level based on the dynamic processes of dislocation connected with the motion of dislocation.8–10,19 The processes of the formation of sliding lines in the range of the yield point were proved to be acoustically most effective. Every collective dislocation motion in the sliding systems probably leads to a release of elastic energy, generating the recorded signals. In most compression tests of the CuZn30 single crystals displaying the PLC effect sudden reductions of the compressive forces recorded on the work-hardening curve were found to be distinctly correlated with the AE peaks (Figure 6). It is supposed that every increase in the oscillation of the stress on the – curves is connected with a locking of dislocations, and that every violent reduction represents a motion of a large group of dislocations (avalanches of dislocations). Similarly, it is assumed that only regular oscillations of the stresses of certain given types are connected with the generation of slip bands of the PLC effect. W. OZGOWICZ et al.: INFLUENCE OF THE STRAIN RATE ON THE PLC EFFECT AND ACOUSTIC EMISSION ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 197–202 201 Figure 6: a) Dependence of the AE energy on the serration along the stress-strain curve (within the fragment of the diagram presented in Figure 4) of a CuZn30 single crystal with the orientation of [139] at 300 °C and  up to about 10–4 s–1 and b) an acoustogram correspond- ing to this set of EA events Slika 6: a) Odvisnost energije AE od nazob~anosti vzdol` krivulje napetost – deformacija (z izrezkom iz diagrama, predstavljenega na sliki 4) pri CuZn30-monokristalu z orientacijo [139] pri 300 °C in  do 10–4 s–1 in b) akustogram, ki ustreza temu sklopu AE-dogodkov Figure 5: Dependence of the AE energy on the shape of the work- hardening curve during the compression testing of a CuZn single crystal at 300 °C and  amounting to about 10–1 s–1 Slika 5: Odvisnost energije AE od oblike krivulje napetostnega utrje- vanja med tla~nim preizkusom monokristala CuZn pri 300 °C in vrednostjo  do 10–1 s–1 4 CONCLUSIONS The investigations dealt with in this paper lead to the following conclusions: • The phenomenon of instability of the plastic deformation of the Portevin-Le Chatelier type occurs in the tested single crystals of alloy CuZn30 in the course of free compression at 300 °C within a limited range of the strain rate from 10–5 s–1 to 10–2 s–1. • In the conditions of free compression applied in the investigations, CuZn30 single crystals display charac- teristic oscillations of the stresses on the work- hardening curves, corresponding to the oscillations classified in the literature as the B type. • The increasing strain rate in the compression tests leads to a distinct decrease in the stress-oscillation intensity, typical for the PLC effect. • The plastic deformation of the investigated single crystals at elevated temperature generates in the analyzed range of frequencies (up to 35 kHz) diversified source of AE energy, mainly an impulsing emission from single events, i.e., pulsating acoustic signals with a high energy in the frequency band from 4 kHz to 8 kHz. • The correlation between the PLC and AE effects in the tested single crystals can be explicitly explained on the basis of a dislocation-dynamic model of the PLC effect. • In the process of plastic deformation of the tested single crystals, the applied AE method exhibits the dependence of the activity of acoustic emissions on the given stage of the hardening of the analyzed alloy. • The intensity of AE increases mainly in the range of the yield point (Re) on the – curves, and also in the case of considerable deformations. When the maximum intensity of the AE signal has been reached, it fades, in most cases, to the minimum value. 5 REFERENCES 1 A. Portevin, F. Le Chatelier, Comptes Rendus de l’Académie des Sciences Paris, 176 (1923), 507–510 2 V. Scott, F. Franklin, F. Mertens, M. Marder, Physical Review E, 62 (2000), 8195–8206, doi:10.1103/PhysRevE.62.8195 3 P. Hähner, E. Rizzi, Acta Materialia, 51 (2003), 3385–3397, doi:10.1016/S1359-6454(03)00122-8 4 J. Balik, Materials Science and Engineering, A316 (2001), 102–108, doi:10.1016/S0921-5093(01)01223-0 5 Z. Jiang, Q. Zhang, H. Jiang, Z. Chen, X. Wu, Materials Science and Engineering A, 403 (2005) 1, 154–164, doi:10.1016/j.msea.2005. 05.059 6 W. Ozgowicz, B. Grzegorczyk, E. Kalinowska-Ozgowicz, Journal of Achievements in Materials and Manufacturing Engineering, 29 (2008) 2, 123–136 7 W. Ozgowicz, B. Grzegorczyk, Archives of Materials Science and Engineering, 39 (2009) 1, 5–12 8 I. Malecki, J. Ranachowski, Acoustic emission: Sources, Methods Applications, PASCAL Publications, Warsaw 1994 (in Polish) 9 Z. Ranachowski, Methods of measurement and analysis of acoustic emission signal, Institute of Fundamental Technological Research Polish Academy of Sciences, Warsaw, 1997 (in Polish) 10 A. Pawelek, J. Kusnierz, Z. Ranachowski, Z. Jasienski, Archives of Acoustics, 31 (2006), 102–122 11 J. Zdunek, J. P³owieæ, J. Mizera, W. Spychalski, K. J. Kurzyd³owski, In¿ynieria Materia³owa, 3 (2010), 577–581 12 J. M. Arimi, E. Duggan, M. O’Sullivan, J. G. Lyng, E. D. O’Riordan, Journal of Texture Studies, 41 (2010), 320–340, doi:10.1111/ j.1745-4603.2010.00224.x 13 J. S. Chen, C. Karlsson, M. Povey, Journal of Texture Studies, 36 (2005), 139–156, doi:10.1111/j.1745-4603.2005.00008.x 14 S. Kudela, A. Pawe³ek, Z. Ranachowski, A. Pi¹tkowski, Metallic Materials, 49 (2011), 271–277, doi:10.4149/km_2011_4_271 15 A. Pawelek, J. Kusnierz, Z. Ranachowski, Z. Jasienski, J. Bogucka, Archives of Metallurgy and Materials, 54 (2009), 83–88 16 B. J. Brindley, P. J. Worthington, Scripta Metallurgica, 4 (1970) 4, 295–297, doi:10.1016/0036-9748(70)90124-9 17 P. Rodriguez, Material Science, 6 (1984), 653–663, doi:10.1007/ BF02743993 18 A. H. Cottrell, Philosophical Magazine, 44 (1953), 829–832, doi:10.1080/14786440808520347 19 P. G. McCormick, Acta Metallurgica, 22 (1974) 4, 489–493, doi:10.1016/0001-6160(74)90102-3 20 A. Van den Beukel, Phys. Stat. Sol. (a), 30 (1975) 1, 197–206, doi:10.1002/pssa.2210300120 21 A. Pawelek, Dislocation aspects of acoustic emission in plastic deformation processes of metals, Aleksander Krupkowski Institute of Metallurgy and Materials Science, Polish Academy of Sciences, Ed. OREKOP s.c., Cracow, 2006, 1–136 (in Polish) W. OZGOWICZ et al.: INFLUENCE OF THE STRAIN RATE ON THE PLC EFFECT AND ACOUSTIC EMISSION ... 202 Materiali in tehnologije / Materials and technology 49 (2015) 2, 197–202 T. DOKTOR et al.: DETERMINATION OF ELASTIC-PLASTIC PROPERTIES OF ALPORAS FOAM ... DETERMINATION OF ELASTIC-PLASTIC PROPERTIES OF ALPORAS FOAM AT THE CELL-WALL LEVEL USING MICROSCALE-CANTILEVER BENDING TESTS DOLO^ANJE ELASTI^NIH IN PLASTI^NIH LASTNOSTI PENE ALPORAS NA RAVNI CELI^NE STENE Z UPOGIBNIMI PREIZKUSI Z MIKROSKOPSKO IGLO Tomá{ Doktor, Daniel Kytýø, Petr Koudelka, Petr Zlámal, Tomá{ Fíla, Ondøej Jirou{ek Institute of Theoretical and Applied Mechanics, Academy of Sciences of the Czech Republic, Prosecká 76, 190 00 Prague, Czech Republic doktor@itam.cas.cz Prejem rokopisa – received: 2013-10-01; sprejem za objavo – accepted for publication: 2014-05-06 doi:10.17222/mit.2013.207 The presented paper is focused on determining the mechanical properties of the Alporas closed-cell aluminium foam. To utilise the favourable properties of cellular metals (e.g., high strength-weight ratio, energy absorption or insulation capabilities) a detailed description of the mechanical properties is required. Cellular metals exhibit heterogeneity at several scale levels. The contribution of the internal structure to the overall mechanical properties may not be in detail evaluated utilizing only the macroscopic testing. On the other hand, the compact material of cell walls is influenced by its composition (titanium- and calcium-rich regions are present in aluminium). Therefore, localised testing techniques with a small region of interest (e.g., indentation methods) may neglect the inhomogeneity along the cell walls. Hence, a testing of isolated cell walls was performed. A custom-developed modular loading device (based on precise linear bearing stages) was assembled to enable cantilever bending tests. The load was applied with a stepper motor and the loading force was measured with a micro-scale load cell with a loading capacity of 2.25 N. Displacements of the samples were measured optically. Several points along the longitudinal axis of a sample were tracked using the Lucas-Kanade tracking algorithm and the obtained displacements were compared to the analytically prescribed deflection curve. Based on the obtained deflections and measured forces, a stress-strain diagram was constructed and the constants of the elastic-plastic material model were evaluated. Keywords: aluminium foam, cantilever bending, micromechanics, optical strain measurement ^lanek obravnava dolo~anje mehanskih lastnosti aluminijaste pene Alporas z zaprtimi porami. Za uporabo ugodnih lastnosti celi~nih kovinskih materialov (npr. visoko razmerje med trdnostjo in maso, absorpcija energije ali izolacijske zmo`nosti), je potrebeno poznati mehanske lastnosti. Celi~ni materiali ka`ejo heterogenost, ~e jih opazujemo pri razli~nih merilih velikosti. Podrobna ocena notranje zgradbe in njen prispevek k mehanskim lastnostim ni mogo~ zgolj z uporabo makroskopskega preizku{anja. Po drugi strani pa zgradba vpliva na kompakten material celi~nih sten (v aluminiju je mogo~e najti podro~ja, ki so bogata s titanom in kalcijem). Zato je mogo~e, da preve~ lokalizirano preizku{anje (npr. metoda za dolo~anje trdote) zanemari nehomogena podro~ja na stenah celic. Zato je bilo izvr{eno preizku{anje posamezne celi~ne stene. Sestavljena je bila modularna obremenitvena naprava, ki omogo~a natan~ne stopnje linearne obremenitve, da so mogo~i upogibni preizkusi z mikroskopsko iglo. Vzorec je bil obremenjen s kora~nim motorjem, pri ~emer je bila sila obremenitve izmerjena z mikroskopsko merilno celico z nosilnostjo 2,25 N. Premik vzorcev je bil izmerjen opti~no. Nekaj to~k ob vzdol`ni osi vzorca je bilo spremljano s sledilnim algoritmom Lucas-Kanade, ugotovljeni premiki pa so bili primerjani z analiti~no dolo~eno krivuljo deformacije. Na podlagi ugotovljenih deformacij in izmerjenih sil je bil postavljen diagram odvisnosti med obremenitvijo in deformacijo. Ocenjene so bile konstante elasti~no-plasti~nega modela materiala. Klju~ne besede: aluminijska pena, upogibanje igle, mikromehanika, opti~no merjenje deformacije 1 INTRODUCTION Aluminum foams are lightweight materials with a favourable combination of mechanical properties, e. g., a high strength-to-weight ratio. This material, widely used in damping applications as an impact-energy absorbent and as a thermal or acoustic insulation, was studied in detail at the macroscopic level1,2. Moreover, models for predicting the macroscopic deformation behaviour were developed3. To enable the exploitation of the favourable proper- ties of aluminium foams for structural applications, a detailed description of their deformation behaviour is required. Due to complex inner structures of closed-cell foams, exhibiting a significant inhomogeneity at the cellular level4 (in terms of the size and shape of the cells as well as the thickness of cell walls), an analysis of the deformation response of a cellular structure requires a complex mathematical description and modelling. The numerical models of metal foams have to reflect both the geometry of the cellular structure and the mechanical properties of the base material. Since the cell-wall material contains residuals of the foaming agents (calcium and titanium) and micropores are present after the foaming process, the mechanical properties should not be measured solely with direct localized measurement techniques (e.g., the nanoinden- Materiali in tehnologije / Materials and technology 49 (2015) 2, 203–206 203 UDK 669.71:620.17:532.6 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)203(2015) tation where the measured properties have to be pro- cessed with a homogenization procedure5). To provide a more complex description of the foam’s deformation response a microstructural numerical model of a representative volume element was developed in our recent study based on X-ray imaging and tomographic reconstruction6. However, the parameters of the consti- tutive material model that was necessary in the finite- element simulation have to take into consideration also the material properties at the level of cell walls. In this study, a micromechanical testing methodology was employed to determine the mechanical properties of the base material of the Alporas aluminium foam. A cantilever arrangement of the loading test was used to ensure a simple description of the boundary conditions for both analytical and numerical evaluations of the experiments. 2 MATERIALS AND METHODS 2.1 Specimen preparation The specimens were prepared from the Alporas closed-cell aluminium foam (Shinko Wire, Inc, Japan). Flat cell walls were identified in the cellular structure. The cells containing such walls were carefully extracted using a hand-operated micro-lathe (Dremel FortiFlex, Bosch GmbH, Germany). To protect the base cell-wall material against a plastic deformation the specimens were embedded into rosin (with a high purity and transparency and a melting point not exceeding 80 °C) during every distinct step of the preparation procedure. A specimen’s shape was finalized by precise grinding and polishing using an oscillation grinding machine, TegraPol (Struers, A/S, Denmark). 2.2 Sample-volume description To determine the geometrical characteristics required for evaluating the bending tests (i.e., the distance between the clamp and the loading point and the cross- sectional characteristics) a set of projections of each specimen was acquired with a scanning electron micro- scope (SEM) (Mira II, Tescan, s. r. o., Czech Republic). Since the specimens had a metallic nature, the secondary electron probe was employed for the scanning. The obtained projections provided a resolution of 2 μm per pixel. For each specimen three projections were acquired, a floor plan (normal to the loading direction, Figure 1a) and two longitudinal faces (parallel to the in-plane loading, Figure 1b). A volumetric model of each specimen was developed using a custom-assembled image-processing procedure developed in the language of computational environment MatLab (Mathworks, Inc., USA). 2.3 Experimental set-up For the loading test an in-house loading set-up was designed and assembled. The set-up was based on linear bearing stages (Standa Ltd, Lithuania) with a resolution of 1 μm and a travel range of up to 50 mm. The motion of the loading point was provided with a precise linear bearing stage, UMR 9.0 (Newport, Ltd, USA) with a resolution of 0.1 μm and a travel range of 5 mm, which was driven by a stepper motor (Microcon, s. r. o., Czech Republic). The reaction force at the loading point was measured with a miniature load cell in the cantilever arrangement (FPB350, Futek Inc., USA) with a loading capacity of 2.25 N and a read-out unit, OM911 (Orbit Merret, s. r. o., Czech Republic) with a sampling rate of T. DOKTOR et al.: DETERMINATION OF ELASTIC-PLASTIC PROPERTIES OF ALPORAS FOAM ... 204 Materiali in tehnologije / Materials and technology 49 (2015) 2, 203–206 Figure 2: Experimental set-up for micro-scale loading tests: a) spe- cimen, b) load cell, c) stepper motor, d) linear bearing stages Slika 2: Eksperimentalni sestav za preizkuse obremenitve na mikro- skopski ravni: a) vzorec, b) merilna celica, c) kora~ni motor, d) na- stavljanje linearnih obremenitev Figure 1: a) Specimen extracted from the cellular structure of Alporas aluminium foam, b) cross-section (SEM) Slika 1: a) Vzorec, izrezan iz celi~ne stene aluminijske pene Alporas, b) pre~ni prerez (SEM) 100 Hz. A detailed description of the loading set-up is depicted in Figure 2. 2.4 Loading procedure The loading was performed in the cantilever arrange- ment in order to overcome the issues of low stiffness of the supports and the eccentricity of the loading point that occurred during our previous studies where the three- point bending arrangement was used. The clamp was implemented by placing a specimen between two prisms connected together with a pair of screws (Figure 2). The loading tests were displacement controlled. The synchronization of the loading and force logs was en- sured with the custom software based on an open-source LinuxCNC project and real-time Linux kernel. The loading rate was set to 1 μm s–1. 2.5 Strain measurement Due to a high compliance of the load cell (even higher than the deflection of the loaded point of a specimen) the strain could not be determined directly from the displacement prescribed by the linear bearing stage. Instead, the displacements were measured opti- cally from a set of projections of the loading scene captured with a CCD camera (Manta G504B, Allied Vision Technologies, GmbH, Germany) with a resolution of 2452 px × 2056 px attached to a light microscope (Navitar Imaging, Inc., USA) that provided a magni- fication of up to 24×. The acquisition of the projections was controlled by in-house-built software based on the OpenCV7 (Open Source Computer Vision) library and Python programming language. The acquisition rate of the camera was 2 fps, which enabled an identification of a sufficient number of points on the force-displacement curve. A selected loading scene is depicted in Figure 3. 2.6 Strain-stress curve evaluation From the sets of projections, displacements were determined using a digital-image-correlation (DIC) soft- ware tool8 based on the Lucas-Kanade tracking algo- rithm9. The points on the specimens’ surfaces were selected along their longitudinal axes and the displace- ments of each point were tracked by searching for the highest correlation coefficient between the two conse- quent projections. The engineering stress ( ) and strain ( ) values were determined from the geometrical properties of a tested specimen and optically measured displacements of the loading point using Equations (1) and (2): = 3 2 2 uh l (1) = Flh I Z2 (2) Here u denotes the displacement of the loading tip, h is the height of a specimen, F is the loading force, l is the distance between the clamp and the loading point and IZ is the axial quadratic moment of inertia of the loaded cross-section. From the slope of the unloading phase, the Young’s modulus was determined and the yield point was estimated using the offset method at a 0.2 % strain level. 3 RESULTS AND DISCUSSION Deformation behaviour of the isolated cell walls was described using micro-scale cantilever measurements. The obtained stress-strain curves are presented in Figure 4 (selected curves only). The curves exhibit similar slopes in their initial parts as well as in their unloading phases. Different portions of the plastic strain are caused by the non-uniformity in the specimens’ dimensions due to a highly irregular geometrical arrangement of the cellular structure. The obtained values of the Young’s modulus are (36.7 ± 5.23) GPa, the yield stress reached (39.0 ± 9.7) MPa and the yield strain reached (0.279 ± 0.044) %. 4 CONCLUSIONS Based on a series of the cantilever bending tests of the isolated cell walls of the Alporas aluminium foam elastic and plastic mechanical properties were deter- T. DOKTOR et al.: DETERMINATION OF ELASTIC-PLASTIC PROPERTIES OF ALPORAS FOAM ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 203–206 205 Figure 4: Stress-strain curves of the tested specimens Slika 4: Krivulje napetost – raztezek preizku{anih vzorcev Figure 3: Loading scene of a clamped specimen (image acquired with a CCD camera and light microscope, the bar corresponds to 1 mm) Slika 3: Obremenjen vzorec (posnetek narejen s CCD-kamero, dol`ina traku ustreza 1 mm) mined. The obtained results scatter slightly among the measured sets, which might be caused by a highly irre- gular nature of the tested material in terms of its geo- metrical arrangement. This issue may be overcome using an inverse numerical evaluation of the measured results in conjunction with the precise geometrical models obtained with the X-ray tomography. The obtained results may serve as the parameters for the constitutive material models for microstructural numerical models. The conjunction of the material pro- perties of the foam’s base material and the microstruc- tural models (e. g., developed by X-ray computed micro-tomography) will provide a more detailed descrip- tion of the mechanical behaviour of the metal foam. Acknowledgements The authors would like to express their gratitude to the Czech Science Foundation (research project No. P105/12/0824). The institutional support of RVO: 68378297 is also gratefully acknowledged for the finan- cial support of this research. We were delighted to cooperate with MDDr. Eva Kamenická, an excellent ope- rator of the micro-lathe. 5 REFERENCES 1 R. Pippan, Deformation behaviour of closed-cell aluminium foams in tension, Acta Materialia, 49 (2001), 2463–2470, doi:10.1016/S1359- 6454(01)00152-5 2 M. I. Idris, T. Vodenitcharova, M. Hoffman, Mechanical behaviour and energy absorption of closed-cell aluminium foam panels in uniaxial compression, Materials Science and Engineering A, 517 (2009) 1–2, 37–45, doi:10.1016/j.msea.2009.03.067 3 Y. Mu, G. Yao, L. Liang, H. Luo, G. Zu, Scripta Materialia, 63 (2010), 629–632, doi:10.1016/j.scriptamat.2010.05.041 4 V. Králík, J. Nìme~ek, Two-scale model for prediction of macrosco- pic elastic properties of aluminium foam, Chemicke Listy, 106 (2012) 3, 458–461 5 J. Nìme~ek, P. Zlámal, V. Králík, J. Nìme~ková, Modeling Inelastic Properties of Metal Foam Based on Results from Nanoindentation, 14th Int. Conf. on Civil, Struct. & Env. Eng. Comp., Cagliari, 2013, 103 6 O. Jirou{ek, T. Doktor, D. Kytýø, P. Zlámal, T. Fíla, P. Koudelka, I. Jandejsek, D. Vavøík, X-ray and finite element analysis of deforma- tion response of closed-cell metal foam subjected to compressive loading, Journal of Instrumentation, 8 (2013), C02012, doi:10.1088/ 1748-0221/8/02/C02012 7 G. Bradski, The OpenCV Library, Dr. Dobb’s Journal of Software Tools [online] (2000), article id 2236121, [posted 2008-01-15 19:21:54], Available from World Wide Web: http://www.drdobbs. com/open-source/the-opencv-library/184404319 8 I. Jandejsek, J. Valach, D. Vavrik, Optimization and Calibration of Digital Image Correlation Method, Proceedings of Experimental Stress Analysis 2010, Velke Losiny, Czech Republic, 2010, 121–126 9 B. D. Lucas, T. Kanade, An Iterative Image Registration Technique with an Application to Stereo Vision, Proceedings of Imaging Under- standing Workshop, Vancouver, 1981, 121–130 T. DOKTOR et al.: DETERMINATION OF ELASTIC-PLASTIC PROPERTIES OF ALPORAS FOAM ... 206 Materiali in tehnologije / Materials and technology 49 (2015) 2, 203–206 I. PETRÁ[OVÁ, M. LOSERTOVÁ: ELECTROCHEMICAL BEHAVIOR OF BIOCOMPATIBLE ALLOYS ELECTROCHEMICAL BEHAVIOR OF BIOCOMPATIBLE ALLOYS ELEKTROKEMIJSKO VEDENJE BIOKOMPATIBILNIH ZLITIN Ivana Petrá{ová, Monika Losertová V[B-Technical University of Ostrava, Faculty of Metallurgy and Materials Engineering, Department of Non-Ferrous Metals, Rafining and Recycling, 17. listopadu 15/2172, 70833 Ostrava, Czech Republic ivana.petrasova@vsb.cz Prejem rokopisa – received: 2013-10-02; sprejem za objavo – accepted for publication: 2014-05-09 doi:10.17222/mit.2013.218 The electrochemical behavior of Ti6Al4V and Ti22Nb alloys was studied in a 0.15 M (0.9 %) physiological sodium chloride solution at room temperature (22 ± 1) °C. The experimental samples of the Ti6Al4V alloy were in different states of the thermomechanical treatment: as-received and hot rolled. The samples of Ti22Nb were studied in the as-cast, heat-treated and aged stages. The microstructural features of the Ti6Al4V alloy in three different states influenced not only the passivation but also the rate of corrosion that was (0.12, 0.07 and 0.10) mm per year for the equiaxed ( + ), acicular  in the transformed grains and coarse lamellar ( + ) phases, respectively. The as-cast TiNb sample with a dendritic microstructure and very fine martensite showed the lowest corrosion rate of 0.26 mm per year unlike the specimens after the heat treatment and aging with the rate of 0.34 mm and 0.33 mm per year, respectively. Keywords: Ti6Al4V, Ti22Nb, electrochemical behavior Preu~evano je bilo elektrokemijsko vedenje zlitin Ti6Al4V in Ti22Nb v fiziolo{ki raztopini natrijevega klorida 0,15 M (0,9 %) pri sobni temperaturi (22 ± 1) °C. Vzorci zlitine Ti6Al4V za preizkuse so bili v razli~nih termomehanskih stanjih: v dobavljenem stanju in vro~e valjani. Vzorci Ti22Nb so bili v litem stanju, toplotno obdelani in starani. Razlike v mikrostrukturi zlitine Ti6Al4V v treh razli~nih stanjih vplivajo na pasivacijo in na hitrost korozije, ki je bila (0,12, 0,07 in 0,10) mm na leto pri enakoosnih ( + ), igli~astih - in transformiranih -zrnih ter grobo lamelarnih fazah ( + ). Vzorci TiNb z dendritno mikrostrukturo in drobnim martenzitom so pokazali najmanj{o hitrost korozije 0,26 mm na leto, v primerjavi z vzorci po toplotni obdelavi in staranju pa z 0,34 mm in 0,33 mm na leto. Klju~ne besede: Ti6Al4V, Ti22Nb, elektrokemijsko vedenje 1 INTRODUCTION The 316L stainless steel, cobalt-chromium alloys and Ti-alloys are three important classes of the most used metallic biomaterials, especially for orthopedic implants, even in highly loaded areas such as artificial joints1,2. Nowadays, titanium and its alloys are the most attractive biomaterials for orthopedic implants and other devices for dental applications due to their relatively good fatigue resistance, excellent biocompatibility, better corrosion resistance in body fluids and lower elasticity modulus compared to the other metallic biomaterials1,3. Pure titanium and ( + ) Ti-alloys were originally used as structural materials, especially for aerospace struc- tures, and only afterwards they were adopted for biomedical applications. In the recent decades, many titanium alloys have been developed. The Ti6Al4V became the standard alloy and it was the first titanium alloy used as a biomaterial1. Recently, the development of titanium alloys has drawn considerable attention in the biomedical area for their much lower Young’s modulus compared to the  (105 GPa for pure Ti)4 or ( + ) Ti alloys (110 GPa for Ti6Al4V)4, thus exhibiting a better biomechanical compatibility. A large modulus mismatch between a metallic implant and the adjacent bone (the Young’s modulus of the human bone is 10–30 GPa) will cause the stress-shielding effect, leading to an excessive bone resorption and implant loosening4. Recently developed alloys based on TiNb showing superelasticity5 have a high potential to serve as alter- natives for the NiTi shape-memory alloys in biomedical applications. The reason for the substitution of Ni with Nb is a very good cytocompatibility of Nb as indicated in the studies in vitro and in vivo6. The corrosion resi- stance of the TiNb alloys has also been shown to be similar or superior to that of Ti alone. Despite the poten- tial benefits of the TiNb alloys, their development for biomedical applications is still at the beginning7. The aim of the presented work is to compare the microstructure effects on the electrochemical behavior of two titanium alloys with different thermal and mecha- nical treatments. The electrochemical experiments were performed on Ti6Al4V and TiNb with amount fraction x = 22 % Nb (Ti22Nb) in the NaCl solution. Most of the recent works focused on the biocompa- tible titanium alloys in different stages of the microstruc- ture deal with the mechanical behavior8. The effects of the microstructure on the corrosion properties of tita- nium alloys have not been studied extensively9. 2 EXPERIMENTAL The study of electrochemical behavior was per- formed on two titanium alloys in various states of the thermal and thermomechanical processing. Materiali in tehnologije / Materials and technology 49 (2015) 2, 207–211 207 UDK 544.6:57:669.295 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)207(2015) The Ti6Al4V alloy with the composition given in Table 1 was studied in three microstructure stages: in the as-received stage and after being rolled above the transus at two diverse temperatures of 900 °C or 1100 °C. The Ti22Nb alloy was tested in following states: in the as-cast state, after being solution annealed at 900 °C for 1 h and water quenched or after being solution annealed at 900 °C for 1 h, aged at 400 °C for 1 h and water quenched10. The denotation of the specimens in relation to the treatment is listed in Table 2. Table 1: Chemical composition of the studied Ti6Al4V alloy in mass fractions, w/% Tabela 1: Kemijska sestava uporabljene zlitine Ti6Al4V v masnih dele`ih, w/% Al V C Fe H O N Y O + N 5.50– 6.75 3.50– 4.50 max. 0.08 max. 0.30 max. 0.0125 max. 0.20 max. 0.05 max. 0.005 max. 0.25 Table 2: Denotation of experimental specimens Tabela 2: Oznake vzorcev Stage/alloy as-recived rolled at900 °C rolled at 1100 °C Ti6Al4V A B C Stage/alloy as-cast solutionannealed aged Ti22Nb D E F For light microscopy the specimen were prepared with the standard metallographic techniques involving grinding, mechanical polishing and etching in Kroll’s reagent containing HF, HNO3 and distilled water (8 : 15 : 77). The microstructures were observed in the etched state using a light microscope GX51. Electrochemical studies were carried out using linear polarization at room temperature (22 ± 1) °C. The corro- sion tests were performed using an Autolab PGSTAT 128n apparatus and a personal computer with software Nova 1.7. The electrochemical measurements were performed using a three-electrode cell containing a working electrode (a specimen) with the exposed area of 16 mm2 or 10 mm2, an Ag/AgCl (SSC) electrode in saturated KCl as the reference electrode and a platinum mesh as the counter electrode. A solution of 0.15 M (0.9 %) NaCl was used as the electrolyte. The specimens were grinded before the tests with the wheel paper of up to 1000 mesh and then polarized from –2 V to +1.5 V with a scan rate of 2 mV/s. The scan rate was selected according to11. 3 RESULTS AND DISCUSSION The microstructures of the electrochemically tested specimens are shown on the micrographs in Figures 1 and 2. The fine-grained microstructure of the Ti6Al4V alloy presented in Figure 1a consisted of equiaxed  grains (light) in the transformed matrix (dark) con- taining coarse acicular . After the hot rolling at 900 °C, the coarse grains transformed into fine acicular , as shown in Figure 1b. Figure 1c shows coarse lamellar  at the prior -grain boundaries, grains with fine  precipitates in the centre of the specimen and coarse plate-like  at the surface of the specimen after the hot rolling at 1100 °C. The as-cast dendritic microstructure of the Ti22Nb alloy was formed of ( + ) and very fine martensite needles, as presented in Figure 2a. In spite of the annealing at 900 °C for one hour in argon, the dendritic structure remained conserved and the water quenching from 900 °C (above the transus) retained the phase I. PETRÁ[OVÁ, M. LOSERTOVÁ: ELECTROCHEMICAL BEHAVIOR OF BIOCOMPATIBLE ALLOYS 208 Materiali in tehnologije / Materials and technology 49 (2015) 2, 207–211 Figure 1: Light micrographs of Ti6Al4V: a) as-received (A), b) after hot rolling at 900 °C (B), c) after hot rolling at 1100 °C (C) Slika 1: Mikrostruktura Ti6Al4V: a) dobavljeno stanje (A), b) po vro~em valjanju na 900 °C (B), c) po vro~em valjanju na 1100 °C (C) that transformed into martensite at lower temperatures10. Indeed, very fine martensite needles were observed in the annealed and quenched microstructure at a high magnification of light microscopy, as shown in Figure 2b. After the aging at 400 °C and water quenching the dendritic microstructure was formed of the ( + ) phases (Figure 2c). No martensite was observed. The linear-polarization results for the Ti6Al4V and Ti22Nb alloys in the 0.15 M NaCl solution are shown in Figures 3 and 4. The polarization parameters, including the corrosion potential (Ecorr), the corrosion current density (jcorr) and the polarization resistence (Rp) obtained using the Tafel extrapolation method, are listed in Table 3. Comparing the linear polarization curves of the three specimens of the Ti6Al4V alloy shown in Figure 3, it can be seen that there are no significant diferences between the cathodic polarization curves. The nature of the anodic polarization curve for specimen C (Figure 3, curve C) indicates an almost stable passive behavior over the entire potential range. The corrosion potential (Ecorr) estimated from the Tafel region is –0.361 V (SSC). The passive current first continuously increases with the potential, but at around 0 V a slight decrease is observed; then, from about 0.5 V the passive current slightly increases with the higher potentials. The Ecorr of samples I. PETRÁ[OVÁ, M. LOSERTOVÁ: ELECTROCHEMICAL BEHAVIOR OF BIOCOMPATIBLE ALLOYS Materiali in tehnologije / Materials and technology 49 (2015) 2, 207–211 209 Figure 3: Linear polarization curves obtained in the 0.15 M NaCl solution at room temperature for different Ti6Al4V stages: A – as-received, B – after hot rolling at 900 °C, C – after hot rolling at 1100 °C Slika 3: Linearne polarizacijske krivulje, dobljene v 0,15 M raztopini NaCl pri sobni temperaturi za Ti6Al4V, za stanja: A – dobavljeno stanje, B – po vro~em valjanju na 900 °C, C – po vro~em valjanju na 1100 °C Figure 2: Light micrographs of Ti22Nb: a) as-cast (D), b) after solution annealing (E), c) after solution annealing and aging (F) Slika 2: Mikrostruktura Ti22Nb: a) lito stanje (D), b) po raztopnem `arjenju (E), c) po raztopnem `arjenju in staranju (F) Table 3: Values of the corrosion parameters determined with the Tafel plot analysis for Ti6Al4V and Ti22Nb alloys in 0.15 M NaCl solution Tabela 3: Korozijski parametri, dolo~eni s Taflovo analizo za zlitini Ti6Al4V in Ti22Nb v 0,15 M raztopini NaCl Alloy State of treatment Ti6Al4V A Ti6Al4V B Ti6Al4V C Ti22Nb D Ti22Nb E Ti22Nb F jcorr/(μA/cm2) 7.23 4.03 6.19 14.86 20.53 20.35 Ecorr/V –0.467 –0.444 –0.361 –0.330 –0.536 –0.545 Rp/kΩ 11.03 27.32 23.80 9.80 14.60 13.60 A and B (curves A, B, Figure 3) shifts to more negative values than in the case of specimen C. For specimen A, the passive film is broken down at more positive potentials than for sample B and then the current sharply increases (Figure 3, curve A). Specimen B (Figure 3, curve B) shows several active-passive transitions followed by the passivation domain at the higher potentials. This behavior may be related to the effect of the presence of the fine acicular  phase in the micro- structure that can act as effective galvanic couples leading to a higher rate of corrosion12. Indeed, the micro- structural features of all three states influenced not only the passivation but also the rate of corrosion that was (0.12, 0.07 and 0.10) mm per year for the equiaxed ( + ) (A), acicular  in the transformed grains (B) and coarse lamellar ( + ) (C), respectively. The samples of the Ti22Nb alloy showed quite simi- lar polarization behaviors with the increasing potential. The as-cast specimen of Ti22Nb (Figure 4, curve D) exhibited a corrosion potential Ecorr of –0.324 V (SSC), which was slightly more noble than the corrosion poten- tials observed for samples E (–0.499 V (SSC)) and F (–0.537 V (SSC)). The polarization behavior can be re- garded as stable passivity. Comparing the microstructural features of the three specimens, it can be concluded that the more is retained, the higher stability of the passivation region can be expected. Sample D having a dendritic microstructure with very fine martensite had the lowest corrosion rate (0.26 mm per year), unlike specimens E and F after the heat treatment (0.34 mm and 0.33 mm per year, respectively). The electrochemical behaviors of both alloys in the 0.15 M NaCl solution are dissimilar, as seen in Figures 3 and 4. For the Ti6Al4V alloy, a passive region followed by a breakdown and repassivation was observed on the anodic polarization diagrams in all three cases. On the contrary, the Ti22Nb alloy showed a relatively stable passivation region over the entire potential range. However, the corrosion rate for Ti22Nb was about three times higher. No visible surface changes were observed on any specimens after the anodic polarization. 4 CONCLUSION Based on the experimental results it can be concluded that the corrosion resistance of titanium alloys is determined not only by the alloy composition but also by the microstructure formed after different thermal or thermomechanical treatments. The electrochemical behavior of the Ti6Al4V and Ti22Nb alloys was studied in the 0.15 M (0.9 %) physio- logical NaCl solution at room temperature. The experi- mental samples of the Ti6Al4V alloy were tested in different states of the thermomechanical treatment: the as-received state with the ( + ) microstructure and after being hot rolled at 900 °C and 1100 °C with a fine acicular  phase and a coarse lamellar ( + ) phase, respectively. The samples of Ti22Nb were studied in the following stages: the as-cast stage with very fine marten- site needles in the dendritic ( + ) microstructure, the stage after being heat treated at 900 °C or quenched, having martensite needles in the retained phase and after being aged at 400 °C, having a dendritic ( + ) microstructure. With respect to the corrosion related to the microstructural features, Ti6Al4V in all three investigated stages displayed lower values of the corrosion rate than measured for the Ti22Nb alloys. Nevertheless, the samples of the Ti22Nb alloys showed a more stable passivation behavior than the Ti6Al4V alloy. No visible surface changes were observed on any specimens after the anodic polarization. Acknowledgments The experimental work was performed with the support of projects No. TA03010804 financed by the Technology Agency of the Czech Republic, No. CZ.1.05/2.1.00/01.0040 "Regional Materials Science and Technology Centre" within the frame of the operational program "Research and Development for Innovations" financed by the Structural Funds and from the state budget of the Czech Republic, and the Grant of SGS, No. SP2013/64 – "Specific research in metallurgy, material and processing engineering" financed by V[B – Technical University of Ostrava, Czech Republic. 5 REFERENCES 1 T. C. Niemeyer, C. R. Grandini, L. M. C. Pinto, A. C. D. Angelo, S. G. Schneider, Journal of Alloys and Compounds, 476 (2009) 1–2, 172–175, doi:10.1016/j.jallcom.2008.09.026 I. PETRÁ[OVÁ, M. LOSERTOVÁ: ELECTROCHEMICAL BEHAVIOR OF BIOCOMPATIBLE ALLOYS 210 Materiali in tehnologije / Materials and technology 49 (2015) 2, 207–211 Figure 4: Linear polarization curves obtained in the 0.15 M NaCl solution at room temperature for different Ti22Nb stages: D – as-cast, E – after solution annealing at 900 °C, F – after aging at 400 °C Slika 4: Linearne polarizacijske krivulje, dobljene v 0,15 M raztopini NaCl pri sobni temperaturi za Ti22Nb, za stanja: D – lito stanje, E – po raztopnem `arjenju na 900 °C, F – po staranju na 400 °C 2 H. Zohdi, H. R. Shahverdi, S. M. M. Hadavi, Electrochemistry Communications, 13 (2011) 8, 840–843, doi:10.1016/j.elecom.2011. 05.017 3 Y. J. Bai, Y. B. Wang, Y. Cheng, F. Deng, Y. F. Zheng, S. C. Wei, Materials Science and Engineering C, 31 (2011) 3, 702–711, doi:10.1016/j.msec.2010.12.010 4 F. X. Xie, X. B. He, S. L. Cao, X. Lu, X. H. Qu, Corrosion Science, 67 (2013), 217–224, doi:10.1016/j.corsci.2012.10.036 5 H. Y. Kim, Y. Ikehara, J. I. Kim, H. Hosoda, S. Miyazaki, Acta Materialia, 54 (2006) 9, 2419–2429, doi:10.1016/j.actamat.2006. 01.019 6 H. Matsuno, A. Yokoyama, F. Watari, M. Uo, T. Kawasaki, Bioma- terials, 22 (2001) 11, 1253–1262, doi:10.1016/S0142-9612(00) 00275-1 7 R. E. McMahon, J. Ma, S. V. Verkhoturov, D. Munoz-Pinto, I. Kara- man, F. Rubitschek, H. J. Maier, M. S. Hahn, Acta Biomaterialia, 8 (2012) 7, 2863–2870, doi:10.1016/j.actbio.2012.03.034 8 M. Niinomi, M. Nakai, J. Hieda, Acta Biomaterialia, 8 (2012) 11, 3888–3903, doi:10.1016/j.actbio.2012.06.037 9 M. Atapour, A. Pilchak, G. S. Frankel, J. C. Williams, M. H. Fathi, M. Shamanian, Corrosion, 66 (2010) 6, 065004-1–065004-9, doi:10. 5006/1.3452400 10 P. [tìpán, M. Losertová, Proc. of the 21st International Conference on Metallurgy and Materials, Metal 2012, Brno, 2012, 1581 11 S. Barril, S. Mischler, D. Landolt, Wear, 259 (2005) 1–6, 282–291, doi:10.1016/j.wear.2004.12.012 12 F. Karimzadeh, M. Heidarbeigy, A. Saatchi, Journal of Materials Processing Technology, 206 (2008) 1–3, 388–394, doi:10.1016/ j.jmatprotec.2007.12.065 I. PETRÁ[OVÁ, M. LOSERTOVÁ: ELECTROCHEMICAL BEHAVIOR OF BIOCOMPATIBLE ALLOYS Materiali in tehnologije / Materials and technology 49 (2015) 2, 207–211 211 T. KUBINA et al.: PREPARATION AND THERMAL STABILITY OF ULTRA-FINE AND NANO-GRAINED ... PREPARATION AND THERMAL STABILITY OF ULTRA-FINE AND NANO-GRAINED COMMERCIALLY PURE TITANIUM WIRES USING CONFORM EQUIPMENT PRIPRAVA KOMERCIALNE ULTRADROBNE IN NANOZRNATE Ti-@ICE Z OPREMO CONFORM IN NJENA TERMI^NA STABILNOST Tomá{ Kubina1, Jaromír Dlouhý1, Michal Köver1, Mária Dománková2, Josef Hodek1 1COMTES FHT, Prumyslova 995, 334 41 Dobrany, Czech Republic 2Institute of Materials Science, Faculty of Materials Science and Technology in Trnava, STU in Bratislava, Bottova 25, 917 24 Trnava, Slovakia tomas.kubina@comtesfht.cz Prejem rokopisa – received: 2013-10-03; sprejem za objavo – accepted for publication: 2014-04-01 doi:10.17222/mit.2013.226 Processes based on severe plastic deformation (SPD) capable of producing microstructures with sizes of the order of nanometres are gaining increasing importance. One of the available ways to make production more efficient is to combine the CONFORM continuous extrusion process with the ECAP method. This paper describes our initial experience with this combined process in a CONFORM 315i machine, equipped with a specially designed die chamber. Trials were performed to explore the impact of the CONFORM equipment’s settings on the microstructure of the Ti wire. The feedstock consisted of CP-Ti grade 2 bar with a diameter of 10 mm. The decisive parameter for the entire process, i.e., the die-chamber temperature, was varied and controlled. Specimens with grain sizes of 1.4 μm and 420 nm were obtained. Using these specimens, the temperatures at which the recovery processes began to take effect were determined by thermal analysis. Keywords: CONFORM-ECAP, titanium wire, ultrafine microstructure, nanostructure, thermal stability Postopki, ki temeljijo na veliki plasti~ni deformaciji (SPD), pri katerih nastaja mikrostruktura z nanometrskimi dimenzijami, pridobivajo na pomenu. Ena od mogo~ih poti za bolj u~inkovito proizvodnjo je kombinacija kontinuirnega postopka ekstruzije CONFORM in metode ECAP. Ta ~lanek opisuje za~etne izku{nje s tem kombiniranim postopkom z napravo CONFORM 315i, opremljeno s posebno oblikovano komoro. Preizkusi so bili opravljeni zato, da bi ugotovili vpliv nastavitve naprave CONFORM na mikrostrukturo Ti-`ice. Vlo`ek je bila palica CP-Titan 2 premera 10 mm. Spreminjana in kontrolirana je bila temperatura komore kot klju~ni parameter celotnega postopka. Dobljeni so bili vzorci z velikostjo zrn 1,4 μm in 420 nm. Pri teh vzorcih je bila s termi~no analizo dolo~ena temperatura, pri kateri se za~ne proces poprave. Klju~ne besede: CONFORM-ECAP, `ica iz titana, ultradrobna mikrostruktura, nanostruktura, termi~na stabilnost 1 INTRODUCTION In the past 15 years, numerous SPD processes (Se- vere Plastic Deformation) have been developed. These processes are used for achieving grain refinement in materials, typically to a grain size between 100 nm and 400 nm. Their efficiency in processing a large volume of material is, however, still insufficient for industrial-scale applications. This drawback has now been overcome by the CONFORM method, which is known for a long time and is used for the continuous, industrial-scale produc- tion of sections, mostly from aluminium. By merging two processes, where the feedstock is forced through a die by the friction force of a wheel, ECAP becomes a continuous process1,2. The first results of limited-scope process experiments conducted by COMTES FHT were presented in3. The ultrafine titanium obtained was annealed isothermally to explore the growth of the mean grain size. In this case, the grain-growth kinetics can be described as grain coarsening from the very beginning of the annealing process. The isothermal grain coarsening is described with the equation: d d tK Q RT n n1 0 1 0 / / exp− = −⎛⎝ ⎜ ⎞ ⎠ ⎟ (1) where d is the mean grain size at the annealing time t, T is the temperature, n is the time exponent, d0 is the initial grain size, K0 is a constant, R is the universal gas constant and Q is the grain-growth activation energy. Equation (1) was used for expressing the distorted grain-growth rate in CP-Ti in terms of the time exponent (n) and the activation energy (Q). The scatter in the measured mean grain size also caused problems in finding the activation energy Q. Its value was deter- mined using graphical methods: 248 kJ mol–1 (at the mean n = 0.19)4. This value is substantially higher than the activation energy of Ti self-diffusion (169.1 kJ mol–1)5 and the grain-growth activation energy in CP-Ti after eight ECAP passes (176 kJ mol–1)6. Materiali in tehnologije / Materials and technology 49 (2015) 2, 213–217 213 UDK 621.777:669.295 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)213(2015) This paper expands the knowledge of the production of nanostructured titanium and describes its thermal stability by thermo-physical tests with a contrast to the stability of an ultrafine-grained variant. 2 EXPERIMENTAL The feedstock consisted of CP-Ti grade 2 bar with a diameter of 10 mm. The chemical composition of the feedstock is shown in Table 1. It was measured using a Bruker Q4 Tasman optical emission spectrometer and a Bruker G8 Galileo gas analyser. Table 1: Chemical composition of feedstock in mass fractions, w/% Tabela 1: Kemijska sestava uporabljenega materiala v masnih dele`ih, w/% Fe O C H N Ti 0.046 0.12 0.023 0.0026 0.0076 99.822 Titanium bars were converted into an "endless" bar of the same diameter as the feedstock using the CON- FORM 315i machine. The material’s mechanical properties at room temperature were determined with cylindrical tensile test specimens with a gauge length of 25 mm and a diameter of 5 mm. In addition, V-notch impact tests were conducted using specimens with a 3 mm × 4 mm cross-section. Two processing experi- ments were conducted with different sets of extrusion parameters. The microstructures of the specimens were observed in a JEOL JSM 7400 FEG scanning electron microscope with a field-emission gun and with a Nordlys (Oxford Instruments) EBSD (Electron Backscatter Diffraction) detector. The specimens for the EBSD analysis were prepared by ion polishing in a JEOL SM-09010 Cross Section Polisher. The EBSD observation conditions were as follows: 25 kV voltage, working distance of 15.5 mm, 500 point × 500 point lattice and a step size of 0.1 μm. The EBSD maps were displayed and edited using HKL Channel 5 software. The intercept grain size was assessed by measuring the lengths between points of intersection between grain boundaries and a square grid, according to the Czech Standard ^SN EN ISO 643. For the purposes of observation in the transmission electron microscope (TEM), thin foils were prepared with final electrolytic thinning in a Tenupol 5 device, using a solution of 300 mL CH3OH + 175 mL 2-butanol + 30 mL HClO4 at –10 °C and a voltage of 40 V. The TEM analysis was performed in a JEOL 200CX instru- ment with an acceleration voltage of 200 kV. Selective electron diffraction was used for the determination of the phases. The effect of deformation on the thermal expansion of titanium and the temperature range for recovery were explored using a Linseis L75 Platinum horizontal dilato- meter with an Al2O3 specimen chamber and a pull bar. The temperature changes were monitored with a thermo- couple on specimens of diameter 5 mm and lengths of approximately 20 mm. Nitrogen (N2) was used as the protective gas. After heating to 950 °C at a rate of 3 K/min the specimen was cooled at 20 K/min to 600 °C. Then the specimen was left to cool in air to the ambient temperature. The recovery processes were monitored in a Linseis PT-1600 heat-flux calorimeter equipped with an S-type thermocouple. The measurement was conducted in Ar at a flow rate of 600 mL/min on specimens with a mass of approximately 41 mg, cut off from diameter bars 5 mm and placed in Al2O3 crucibles with lids. 3 MICROSTRUCTURE AFTER EXPERIMENTAL PROCESSING The feedstock microstructure contained equiaxed grains with scarce twins (Figure 1). After the deforma- tion it consisted of two types: recrystallized equiaxed grains and a small proportion (less than 15 %) of dis- torted grains divided into sub-grains by low-angle boun- daries. The as-received and distorted microstructures exhibited an identical pronounced texture with (1000) planes aligned with the specimen axis. The experimental extrusion was based on varying the key parameters: the speed of the wheel, the die-chamber temperatures and the cooling downstream of the die chamber. The pro- cess-controlling parameter was the die-chamber tempe- rature, and it was varied from an initial 500 °C to a final 350 °C. The specimens represent runs at various die- chamber temperatures. The microstructure of the specimen processed in the die chamber at a temperature of 500 °C consisted of recrystallized equiaxed grains with a bimodal size with an average size of 1.9 μm. Neither the small nor the large grains exhibit distorted structures. The die-chamber temperature of 450 °C is sufficient for the microstructure to recover/recrystallize. No effects of the cooling were detected. The reco- very/recrystallization and potential grain growth finished before the specimen was cooled. The specimen pro- T. KUBINA et al.: PREPARATION AND THERMAL STABILITY OF ULTRA-FINE AND NANO-GRAINED ... 214 Materiali in tehnologije / Materials and technology 49 (2015) 2, 213–217 Figure 1: Microstructural state of the feedstock material (length of scale 200 μm) Slika 1: Stanje mikrostrukture izhodnega materiala (dol`ina skale je 200 μm) cessed at the die chamber temperature of 400 °C began to show changes as it contained a small amount (10–15 %) of deformed, un-recrystallized grains. Dis- torted grains with a size of no more than 5 μm × 10 μm were divided by low-angle boundaries into sub-grains. The average grain size was 1.9 μm. Figure 2 shows the microstructure upon water cooling after the processing at a chamber temperature of 350 °C, consisting of slightly elongated, deformed grains. The EBSD analysis was focused on the centre of the circular cross-section of the extruded product. The analysed surface is in a plane that is parallel to the bending plane/flow plane in the CONFORM chamber. On the EBSD maps the axis of the extruded section is vertical and the specimen grain size was 1.4 μm. The titanium wire processed in this way was then used for the annealing experiment. The mechanical properties of the feedstock and of the Ti wire after a single pass at the die chamber temperature of 350 °C are listed in Table 2. As expected, the yield stress and the ultimate strength of the product are higher than those of the feedstock. On the other hand, its contraction, elongation and the impact toughness upon the first pass are lower than those of the feedstock. In the second processing experiment, the feedstock preheating device of the CONFORM machine was turned off, the chamber temperature was set to 200 °C and the wheel speed was 0.5 m s–1. In this experiment, three CONFORM passes of Ti wire were used. As the amount of strain introduced into the material was sub- stantial, the EBSD examination did not give any result. All three specimens were also examined using the TEM. 3.1 TEM analysis of the microstructure From samples of Ti wire upon each pass, foils for the transmission electron microscopy observations, oriented in longitudinal and transverse directions, were prepared and in this work the results of the examination of the lon- gitudinal cross-sections are presented. In this direction, the change in grain size with the introduced strain was critically assessed. Upon the first CONFORM pass with a chamber tem- perature of 200 °C, the microstructure on the longitudi- nal section consisted of polyhedral grains with an ele- vated dislocation density, distorted grains and even disc-shaped grains, as shown in Figure 3. T. KUBINA et al.: PREPARATION AND THERMAL STABILITY OF ULTRA-FINE AND NANO-GRAINED ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 213–217 215 Figure 4: General view of Ti wire substructure after second pass at 200 °C Slika 4: Videz podstrukture v Ti-`ici po drugem prehodu pri 200 °C Figure 2: Microstructural state of CP-Ti subjected to one pass through the CONFORM machine at 350 °C (length of scale 200 μm) Slika 2: Stanje mikrostrukture CP-Ti po 1 prehodu skozi napravo CONFORM pri 350 °C (dol`ina skale je 200 μm) Table 2: Mechanical properties and average grain size for various states of CP-Ti Tabela 2: Mehanske lastnosti in povpre~na velikost zrn pri razli~nih stanjih CP-Ti Proof stress Rp0.2/MPa UTS Rm/MPa Ag/% A5/% Z/% KCV/(J cm–2) d/μm Feedstock 354 470 9.3 32.3 64.2 64.2 5.39 Single pass (350 °C) 620 694 12.0 26.3 55.7 27.5 1.4 Three passes (200 °C) 637 698 1.77 17.8 66.2 –- 0.42 Figure 3: Region with grains with deformation texture within Ti wire after the first pass at 200 °C Slika 3: Podro~je z zrni z deformacijsko teksturo v Ti-`ici po prvem prehodu pri 200 °C The nature of the substructure of the samples taken upon the second pass is similar to the condition seen on the transverse cross-section and the very fine-grained polyhedral microstructure is shown in Figure 4. The grain size measured was dstr (310 ± 30) nm, with a difference compared to the previous specimen. On the longitudinal section there were areas of distorted grains or grains with a disc morphology, while the previous specimen contained no such grains. Figure 5 characterizes the substructure of the speci- mens on the longitudinal section through the wire upon three CONFORM passes. The nature of the substructure is very similar to that on the transverse cross-section. The grains are polygonal, with a mean size of dstr (420 ± 30) nm. On this section too, areas can be found where the dislocation density is low, but locations with a higher dislocation density are present as well. 4 RESULTS OF THE THERMAL ANALYSIS The results of the dilatometer measurement were processed with Linseis Data Evaluation software (Figure 6). The specimen after one pass through the CONFORM machine at 350 °C showed a reduced elongation, which can be explained by the annihilation of dislocations and the elimination of the lattice stress. This was not ob- served in the annealed feedstock, which received no deformation and so eventual effects of the phase trans- formation can be ruled out. The calculated difference between the changes in the lengths of the feedstock and the processed specimen suggest that in the temperature range 432–576 °C the recovery caused a substantial change in the length. The same method was used for measuring the variation in the thermal expansion for the specimen upon three passes at 200 °C. The lower limit of the interval of the initial recovery of deformed Ti wire is shifted towards notably lower temperatures, to 300–560 °C, according to Figure 7. The results of the differential scanning calorimeter (DSC) measurement on a specimen after a single pass at 350 °C are shown in Figure 8. Above 300 °C, the tem- perature increase in both specimens is in accordance with the thermal schedule (constant heating rate of T. KUBINA et al.: PREPARATION AND THERMAL STABILITY OF ULTRA-FINE AND NANO-GRAINED ... 216 Materiali in tehnologije / Materials and technology 49 (2015) 2, 213–217 Figure 7: Dependence of the change of length on the temperature for different states of CP-Ti processed at 200 °C Slika 7: Odvisnost spremembe dol`ine od temperature za razli~na stanja CP-Ti, predelanega pri 200 °C Figure 5: General view of Ti wire substructure after third pass at 200 °C Slika 5: Videz podstrukture v Ti-`ici po tretjem prehodu pri 200 °C Figure 8: Heat flux vs. temperature dependence in a specimen after a single pass at 350 °C Slika 8: Odvisnost toplotnega toka od temperature v vzorcu po enem prehodu pri 350 °C Figure 6: Dependence of the change in length on the temperature for different states of CP-Ti processed at 350 °C Slika 6: Odvisnost spremembe dol`ine od temperature za razli~na stanja CP-Ti, predelanega pri 350 °C 10 K/min). Despite that, the curves for both specimens (the processed one and the one annealed in the DSC) show differences. As the calculation shows, there is a steeper increase in the stress in the processed specimen at temperatures above 440 °C (Figure 8). Apparently, an exothermic reaction, generating heat, takes place. This region matches the recovery zone identified in the dilatometric measurement. (The higher onset tempera- ture found in the DSC may be due to the higher heating rate, the larger specimen, and the resulting response). It may be attributed to the release of heat by recovery stress relaxation. In the processed specimens, the recrystalliza- tion peak was detected at lower temperatures, in accordance with the theoretical and empirical knowledge of the dependence of recrystallization on the temperature and the strain. In the specimen processed with three passes through the CONFORM machine at 200 °C, the DSC method showed a more notable decrease in the temperature for the onset of the exothermic reaction, being approximately 320 °C, as shown in Figure 9. The recrystallization temperature in the undeformed specimen was Trx = 615 °C, and 594 °C for a single-pass "conformed" specimen. The decrease in the recrystalliza- tion temperature down to 527 °C in a CP-Ti specimen after eight ECAP passes has been reported in7. 5 CONCLUSION The effect of CONFORM straining on the micro- structure of a Ti wire was investigated. It was found that the die-chamber temperature had the strongest influence. The smallest mean grain size found by EBSD was 1.4 μm, which was achieved at a die-chamber tempe- rature of 350 °C. The microstructure of the processed material after three passes at a die chamber temperature of 200 °C was examined. After the first pass, the micro- structure contained a notable number of slip bands, and after the second and third passes, the character of micro- structure was identical, with fine polyhedral grains with low and high dislocation densities. After the third CON- FORM pass, the mean grain size of dstr (420 ± 30) nm was obtained. The thermal analysis showed that the increasing amount of strain introduced temperatures for the onset of the recovery decrease. In the Ti wire with the mean grain size of 1.4 μm, the recovery began at approximately 440 °C, and for the Ti wire with the grain size of 420 nm, this temperature is even lower, i.e., 310 °C. As polyhedral grains with varying dislocation densi- ties are present, is becomes clear that the recovery phenomena take place even during the forming process. It also suggests that the forming temperature is lower than the recorded die-chamber temperature. 6 REFERENCES 1 G. J. Raab, R. Z. Valiev, T. C. Lowe, Y. T. Zhu, Continuous pro- cessing of ultrafine grained Al by ECAP-Conform, Materials Science and Engineering: A, 382 (2004) 1/2, 30–34, doi:10.1016/j.msea. 2004.04.021 2 G. Raab et al., Long-Length Ultrafine-Grained Titanium Rods Pro- duced by ECAP-Conform, Materials Science Forum, 584–586 (2008), 80–85, doi:10.4028/www.scientific.net/MSF.584-586.80 3 M. Duchek, T. Kubina, J. Hodek, J. Dlouhy, Development of the pro- duction of ultrafine-grained titanium with the conform equipment, Mater. Tehnol., 47 (2013) 4, 515–518 4 T. Kubina, J. Dlouhy, M. Köver, J. Hodek, Preparation and thermal stability of ultra fine-grained commercially pure titanium wire, Proc. of Recent trends in structural materials COMAT 2012, Pilsen, 2012, [cited 2013-09-30]. Available from world Wide Web: http://www. comat.cz/files/proceedings/11/reports/1301.pdf 5 E. A. Brandes, G. B. Brook, Smithels Metals Reference Book, 7th ed., Butterworth-Hernemann, Oxford 1992 6 M. Hoseini et al., Thermal stability and annealing behaviour of ultrafine grained commercially pure titanium, Materials Science and Engineering A, 532 (2012), 58–63, doi:10.1016/j.msea.2011.10.062 7 J. Gubicza et al., Microstructure of severely deformed metals determined by X-ray peak profile analysis, Journal of Alloys and Compounds, 378 (2004) 1/2, 248–252, doi:10.1016/j.jallcom.2003. 11.162 T. KUBINA et al.: PREPARATION AND THERMAL STABILITY OF ULTRA-FINE AND NANO-GRAINED ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 213–217 217 Figure 9: Heat flux vs. temperature dependence in a specimen after three passes at 200 °C Slika 9: Odvisnost toplotnega toka od temperature v vzorcu po treh prehodih pri 200 °C J. KVAPIL et al.: ESTIMATION OF THE THERMAL CONTACT CONDUCTANCE ... ESTIMATION OF THE THERMAL CONTACT CONDUCTANCE FROM UNSTEADY TEMPERATURE MEASUREMENTS DOLO^ANJE KONTAKTNE TOPLOTNE PREVODNOSTI IZ NERAVNOTE@NEGA MERJENJA TEMPERATURE Jiøí Kvapil, Michal Pohanka, Jaroslav Horský Heat Transfer and Fluid Flow Laboratory, Faculty of Mechanical Engineering, Brno University of Technology, Technická 2, 616 69 Brno, Czech Republic kvapil@fme.vutbr.cz Prejem rokopisa – received: 2013-10-08; sprejem za objavo – accepted for publication: 2014-03-28 doi:10.17222/mit.2013.238 Thermal contact conductance is an important parameter for describing the heat transfer between two bodies. When two solids are put in contact and heat transfer occurs, a temperature drop is observed at the interface between the solids. This is caused by an imperfect joint, which occurs because the real surfaces are not perfectly smooth and flat. This paper describes an experimental device for the evaluation of the thermal contact conductance, which was designed and fabricated in the Heat Transfer and Fluid Flow Laboratory. This device was built mainly for simulating metal-forming conditions, which include high pressures (up to 360 MPa) and high temperatures (up to 1200 °C) in the contact of two solids. The principle of this investigation is the unsteady measurement of the temperatures of two solids that are put in contact under different conditions. The surface temperature and thermal contact conductance can be calculated from the measured temperatures by an inverse heat-transfer task. The measured temperature history and the calculated values of the thermal contact conductance for pilot tests are presented in this paper. Keywords: thermal contact conductance, inverse heat conduction problem, heat-transfer coefficient Kontaktna toplotna prevodnost je pomemben parameter za opisovanje prehoda toplote med dvema telesoma. Ko sta dve trdni snovi v stiku, se pri prenosu toplote opazi zni`anje temperature na stiku med dvema trdnima snovema. To nastane zaradi nepopolnega stika, ker realne povr{ine niso popolnoma gladke in ravne. Ta ~lanek opisuje eksperimentalno napravo za oceno kontaktnega prevajanja toplote, konstruirane v laboratoriju za prenos toplote in toka teko~in. Naprava je bila zgrajena predvsem za simulacijo razmer pri preoblikovanju materialov, ki vklju~ujejo velik tlak (do 360 MPa) in visoke temperature (do 1200 °C) na stiku med dvema trdnima materialoma. Princip teh raziskav je neravnote`no merjenje temperature dveh trdnih snovi, ki sta v kontaktu v razli~nih razmerah. Temperatura povr{ine in kontaktna prevodnost toplote se lahko izra~unata iz izmerjenih temperatur z upo{tevanjem inverznega prenosa toplote. V tem ~lanku sta predstavljeni zgodovina merjenja temperature in izra~unane vrednosti kontaktne toplotne prevodnosti pri opravljenih preizkusih. Klju~ne besede: toplotna prevodnost kontakta, problem inverznega prenosa toplote, koeficient toplotne prevodnosti 1 INTRODUCTION Heat transfer over an interface has been the subject of research for decades, as it plays an important role in applications such as nuclear-reactor cooling, the aero- dynamic heating of supersonic aircraft and missiles, satellite thermal control, the packaging of electronics, turbine and internal combustion engine design, etc. When two solids with different temperatures come into contact, heat transfer occurs. A temperature drop is observed at the interface between the solids because of the surface imperfections. No truly smooth surface really exists. In reality, the contact is created at only a few discrete points. The contact point size and the density depend on the surface roughness, the physical properties of asperities and the contact pressure. The real contact- area fraction is only 0.01–0.1 % without pressure. Pressure is one of the most important parameters that can affect the real contact area. Nevertheless, for metals the direct contact takes an area of around 1–2 % for several tens of MPa in the contact.1 The heat transfer is de- creased due to the (air) layer partially filling the voids between the surfaces, although heat flows or radiates through these voids. Many models and empirical and semi-empirical correlations to predict the thermal contact conductance have been published.2 However, these mo- dels have restrictions in terms of the maximum contact pressure (up to 7 MPa) and temperature and are not con- venient for an estimation of the thermal contact con- ductance in conditions simulating metal forming and hot rolling. 2 METHODS OF MEASUREMENT Usually, thermal contact conductance is measured using steady-state experiments (Figure 1).3 The two bo- dies are in contact and their ends are cooled and heated. The temperature distribution is measured using thermo- couples inside the bodies. After a couple of hours, the constant heat flux q is obtained and the temperature drop T at the interface is estimated by extrapolation. From this, the thermal contact conductance hc can be cal- culated: h q Tc =  (1) Materiali in tehnologije / Materials and technology 49 (2015) 2, 219–222 219 UDK 536.21:536.28 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)219(2015) The next procedure to estimate the thermal contact conductance is an unsteady measurement. This proce- dure is described here. Two thermocouples are embedded close to the surface of the bodies in contact. The measured temperature history is used for an inverse cal- culation and the thermal contact conductance is derived. This method is much faster but more difficult to calculate in comparison with steady-state experiments. 3 EXPERIMENTAL PROCESS An experimental device (Figure 2) for estimating the thermal contact conductance in various conditions was built in the Heat Transfer and Fluid Flow Laboratory. The main part of the device is a steel body in the shape of a hollow cylinder. Two smaller cylinders, top and bottom, and a spring are located inside the hollow cylin- der (Figure 3). A temperature sensor is embedded in the bottom cylinder and has a diameter of 12 mm. The top cylinder is used for compressing the spring to the required force, which is measured by a force sensor. The trigger mechanism is used to prevent the bottom cylinder and the sensor from moving. The first thermocouple (type K, diameter of 0.5 mm) is built in the temperature sensor at a depth of 0.9 mm from the surface and the second thermocouple (type K, diameter of 1.5 mm) is in the test plate at a depth of 2 mm from the surface. This device can be used for measuring a variety of initial conditions, like contact pressure (up to 360 MPa), temperature (up to 1200 °C), different types of materials in contact, surface roughness and scales on the surface. 3.1 Experimental procedure The experiment begins by heating the test plate to the required temperature. Independent of the heating, the J. KVAPIL et al.: ESTIMATION OF THE THERMAL CONTACT CONDUCTANCE ... 220 Materiali in tehnologije / Materials and technology 49 (2015) 2, 219–222 Figure 3: Cross-section of the experimental device Slika 3: Prerez naprave za preizkuse Figure 1: Experimental device for steady-state measurement of the thermal contact conductance3 Slika 1: Eksperimentalna naprava za ravnote`ne meritve kontaktne toplotne prevodnosti3 Figure 4: Contact of the temperature sensor and the heated test plate in detail Slika 4: Detajl stika senzorja temperature in ogrevane preizkusne plo{~e Figure 2: Experimental device Slika 2: Naprava za preizkuse spring is pressed to the required load. Then the test plate is inserted into the experimental device and the trigger mechanism is released. When the surfaces of the sensor and the test plate are in contact (see the detail in Figure 4), heat transfer occurs because of the different tempera- tures of the bodies in contact. The temperatures are measured and stored in the data logger. Three experiments (Table 1) with different initial pa- rameters were performed and the measured temperature history from Exp2 is shown in Figure 5. The time 0 s marks the time of contact. Table 1: List of experiments Tabela 1: Seznam preizkusov Test name Contactpressure Test plate Initial temperature Surface (roughness) Exp1 25 MPa 330 °C grinded (Ra 0.8) Exp2 25 MPa 530 °C grinded (Ra 0.8) Exp3 70 MPa 820 °C grinded (Ra 0.8) The heat flux and the surface temperatures of the temperature sensor and the test plate during the experi- ment are calculated by the inverse heat-conduction task. 4 INVERSE HEAT-CONDUCTION TASK Two 2D models, one for the temperature sensor and the second for the test plate, were used for the numerical computation. The models also include the thermocouples inside because the homogeneity of the material is dis- rupted by the inserted thermocouples, and thus the temperature profile is also disrupted. A one-dimensional sequential Beck’s approach4–6 is used to compute the heat fluxes and the surface temperatures of the temperature sensor. The main feature of this method is the sequential estimation of the time-varying heat fluxes and surface temperatures and using future time-step data to stabilize the ill-posed problem. The measured temperature history from the temperature sensor is used as the input T* in the minimizing equation: SSE T Ti i i m m f = − = + + ∑ ( )* 2 1 (2) where m is the current time, f is the number of future time steps and Ti is the computed temperatures from the forward solver7. The SSE denotes the sum of square errors. The value of the surface heat flux q at time m is: q q T T m m i i q i i m m f i i m m f m = + − ⋅ − == + + = + + ∑ ∑ 1 0 1 2 1 ( ) ( ) *   (3)  i i m T q = ∂ ∂ (4) where i is a sensitivity coefficient at time index i to the heat-flux pulse at time m. The temperatures Ti q m =0 at the thermocouple location of the temperature sensor computed from the forward solver use all the previously computed heat fluxes without the current one qm. When the heat flux is found for time m, the corresponding sur- face temperatures of the temperature sensor Tsurf m 1 and the test plate Tsurf m 2 are computed from the forward solver using qm as the boundary condition in the contact area for the temperature sensor and the test plate. Using this procedure, the whole heat-flux history and the surface-temperature history are computed. When the surface heat flux qm and the surface temperatures Tsurf m 1 and Tsurf m 2 are known, the thermal contact conductance hc is computed from: ( ) ( )h q T T T T c m m surf m surf m surf m surf m = + − +− −2 2 1 1 1 1 2 2 (5) 5 RESULTS The computed heat-flux history and surface-tempera- ture histories of the bodies in contact from Exp2 are shown in Figures 6 and 7. J. KVAPIL et al.: ESTIMATION OF THE THERMAL CONTACT CONDUCTANCE ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 219–222 221 Figure 6: Computed heat-flux distribution of Exp2 Slika 6: Izra~unana razporeditev toka toplote po Exp2 Figure 5: Temperature history of Exp2 Slika 5: Potek temperature pri Exp2 The inverse calculations, made for all the experi- ments and results in the form of thermal contact conduc- tance, are shown in Figure 8. Fieberg3 dealt with the unsteady measurement of thermal contact conductance using a method where the temperature is measured using a high-speed infrared camera. During the experiments, he focused on condi- tions similar to those in a combustion engine, where the contact pressure can be up to 250 MPa and the tempe- rature can be up to 630 °C. He mainly tested steel-alu- minium alloys, but he also published an experiment with steel. His distribution of the thermal contact conductance with steel bodies for 34 MPa in the contact and a starting temperature at 270 °C is shown in Figure 8. Theoretically, the distribution of the thermal contact conductance with time should be constant. This corres- ponds quite well with our experiments, and also with Fieberg3, but in our case, directly after the contact, it takes some time (around 0.3–0.5 s) to stabilize the level of thermal contact conductance. This is caused by vibra- tions after the sudden contact of the temperature sensor and the heated test plate. It is obvious that an important influence on the thermal contact conductance comes from the contact pressure and the initial temperature of the test plate. 6 CONCLUSION In this paper, an experimental device for the unsteady measurement of the temperatures of two solids that are put in contact has been described. New numerical mo- dels for computing the thermal contact conductance from the temperature history were developed, three pilot expe- riments were made and the results were presented. The computed distributions of the thermal contact conduc- tance show a strong dependence on the contact pressure and the initial temperature. The estimated thermal con- tact conductance for a contact pressure of 25 MPa and an initial temperature of 330 °C is 9100 W/(m2 K), while for 25 MPa and 530 °C it is 12200 W/(m2 K) and for 70 MPa and 820 °C it is 35500 W/(m2 K). These results could be used in numerical simulations of metal forming and hot rolling where the values of thermal contact conductance in the interface between two solids are still missing. Additionally, there is a need for similar experi- ments focusing on the influence of the thermal contact conductance by surface roughness, type of material and thickness of the scales. Acknowledgement The research in the presented paper has been sup- ported within the project No. CZ.1.07/2.3.00/ 20.0188, HEATEAM – Multidisciplinary Team for Research and Development of Heat Proceeding. 7 REFERENCES 1 F. P. Bowden, D. Tabor, The friction and lubrication of solids, Oxford University Press, London 1950, 391 p. 2 A. Wang, J. Zhao, Review of prediction for thermal contact resistance, Science China Technological Sciences, 53 (2010) 7, 1798–1808, doi:10.1007/s11431-009-3190-6 3 C. Fieberg, R. Kneer, Determination of thermal contact resistance from transient temperature measurements, International Journal of Heat and Mass Transfer, 51 (2008) 5–6, 1017–1023, doi:10.1016/ j.ijheatmasstransfer.2007.05.004 4 J. Beck, B. Blackwell, C. R. Clair, Inverse heat conduction: ill-posed problems, Wiley, New York 1985, 308 p. 5 M. Pohanka, K. A. Woodbury, A Downhill Simplex method for computation of interfacial heat transfer coefficients in alloy casting, Inverse Problems in Engineering, 11 (2003), 409–424, doi:10.1080/ 1068276031000109899 6 M. Raudensky, Heat Transfer Coefficient Estimation by Inverse Conduction Algorithm, International Journal of Numerical Methods for Heat and Fluid Flow, 3 (1993) 3, 257–266, doi:10.1108/ eb017530 7 W. J. Minkowycz, E. M. Sparrow, J. Y. Murthy, Handbook of Nume- rical Heat Transfer, 2nd edition, John Willey & Sons, New Jersey 2006, 968 p. J. KVAPIL et al.: ESTIMATION OF THE THERMAL CONTACT CONDUCTANCE ... 222 Materiali in tehnologije / Materials and technology 49 (2015) 2, 219–222 Figure 8: Distributions of thermal heat conduction for various initial conditions Slika 8: Razporeditev toplotne prevodnosti za razli~ne za~etne razmere Figure 7: Computed surface temperatures of Exp2 Slika 7: Izra~unane temperature povr{ine pri Exp2 W. WALKE, J. PRZONDZIONO: POTENTIODYNAMIC AND XPS STUDIES OF X10CrNi18-8 STEEL ... POTENTIODYNAMIC AND XPS STUDIES OF X10CrNi18-8 STEEL AFTER ETHYLENE OXIDE STERILIZATION POTENCIODINAMI^NE IN XPS ANALIZE JEKLA X10CrNi18-8 PO STERILIZACIJI Z ETILEN OKSIDOM Witold Walke1, Joanna Przondziono2 1Silesian University of Technology, Faculty of Biomedical Engineering, Ch. de Gaulle’a 66, 41-800 Zabrze, Poland 2Silesian University of Technology, Faculty of Materials Engineering and Metallurgy, Krasiñskiego 8, 40-019 Katowice, Poland witold.walke@polsl.pl Prejem rokopisa – received: 2013-11-14; sprejem za objavo – accepted for publication: 2014-05-09 doi:10.17222/mit.2013.281 An innovative development in the treatment of heart, blood and vascular-system diseases using low-invasive techniques led to a development of new forms of tools (among other devices: cardiologic guide wires) made of the X10CrNi18-8 steel. The authors of this study made an attempt to evaluate the impact of one of the medical sterilisation methods, i.e., the ethylene oxide sterilisation of a wire made of the X10CrNi18-8 steel after electrochemical polishing and chemical passivation. One of the basic criteria for deciding about the suitability of a specific material for a vascular tool is the proper corrosion resistance in a blood environment. It is directly connected with the chemical composition of the surface layer. Therefore, an evaluation was made on the basis of pitting-corrosion tests and the tests of the chemical compositions of the surface layers by means of the XPS method. The samples were subjected to the tests before and after the ethylene oxide sterilisation. The obtained results explicitly prove that the chemical-passivation process of the X10CrNi18-8 steel improves its resistance to corrosion in a blood environment. The resistance is higher due to the creation of a thin passive layer, mainly built of Fe2O3 and Cr2O3 on the surface, which was proved during the XPS tests. Next, the sterilisation in ethylene oxide had a favourable influence on the electrochemical properties of the X10CrNi18-8 steel, irrespective of the way of surface preparation. The presence of the alloying elements in the oxidised form was also detected in the surface layer, contributing to the improvement of the corrosion resistance in contact with the blood. Keywords: X10CrNi18-8 steel, EO sterilization, XPS, pitting corrosion Razvoj inovativnih metod pri zdravljenju srca, kot tudi krvno-`ilnih bolezni, z uporabo malo invazivnih tehnik je pripeljal do novih oblik orodij (med drugim tudi kardiolo{kih uvajalnih `ic), izdelanih iz jekla X10CrNi18-8. Avtorji te {tudije si prizadevajo oceniti vpliv ene od sterilizacijskih metod, to je sterilizacija `ice, izdelane iz jekla X10CrNi18-8 z etilen oksidom, po elektrokemijskem poliranju in po kemijski pasivaciji. Eden od osnovnih meril, ki odlo~ajo o primernosti dolo~enega materiala za orodja za o`ilja, je primerna korozijska obstojnost v krvi. Ta je neposredno povezana s kemijsko sestavo povr{inske plasti. Zato je bila izdelana ocena na podlagi preizkusov jami~aste korozije in analize sestave povr{ine z XPS-metodo. Vzorci so bili preizku{eni pred sterilizacijo z etilen oksidom in po njej. Dobljeni rezultati so neposredno dokazali, da proces kemijske pasivacije jekla X10CrNi18-8 izbolj{a njegovo odpornost proti koroziji v okolju s krvjo. To je povezano z nastankom tanke pasivne plasti iz Fe2O3 in Cr2O3 na povr{ini, kar je bilo dokazano z XPS-analizami. Sterilizacija in etilen oksid imata ugoden vpliv na elektrokemijske lastnosti jekla X10CrNi18-8, ne glede na na~in priprave povr{ine. Na povr{ini je bila odkrita prisotnost oksidiranih legirnih elementov, ki prispevajo k izbolj{anju korozijske odpornosti pri stiku s krvjo. Klju~ne besede: jeklo X10CrNi18-8, EO-sterilizacija, XPS, jami~asta korozija 1 INTRODUCTION Cardiologic guide wires have an extremely important role in the success of a coronary intervention. It is believed that the success of such a treatment depends on the type and quality of a guide wire. The guide wire is the first to reach the area of atherosclerotic lesion and it passes through a narrowing, previously passing along a complicated route inside the vessels, e.g., coronary arteries. Quite frequently, it also has to pass through a lesion that is totally closed. Guide wires differ as far as their profile, flexibility and tip configuration are con- cerned. The market offers guide wires with a J-shaped tip or a straight tip for individual forming. The diameter of modern guide wires is within the range of 0.35–0.45 mm, and their length is within the range of 180–300 mm. These wires, due to their direct contact with the blood, must feature a proper set of electrochemical properties also determined by the conditions of medical sterilisation.1,2 An important issue in the process of forming guide-wire functional characteristics is the selection of the characteristics of the metallic material they are made of.3–5 Mechanical characteristics of the material are selected on the ground of biomechanical characteristics determined for individual forms of instru- ments, taking the anatomical load for each type of the system into consideration. A crucial issue for the cardio- logic guide-wire functional characteristics is the proper method of the surface preparation. The restrictions arising from implementing a tool made of a metallic material in the blood and vascular system have made numerous researchers look for the ways of efficiently diminishing the influences of the body fluids on the metal corrosion. It results from the analysis of the subject-matter data that one of the main factors influencing the efficiency of a treatment is the guide- Materiali in tehnologije / Materials and technology 49 (2015) 2, 223–227 223 UDK 669.14:620.193:620.197.2:577 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)223(2015) wire surface roughness. The significance of the surface-layer chemical composition cannot be neglected either.6,7 One of the basic treatments increasing the corrosion resistance of steel in the environment of human blood is the chemical-passivation process. Therefore, the authors of this study attempted to make an evaluation of the influence of the surface treatment (electrochemical polishing and chemical passivation) of steel X10CrNi18-8 on the biotolerance in the environment of human blood under the conditions of sterilisation with ethylene oxide. 2 EXPERIMENTAL WORK Samples made of steel X10CrNi18-8 in the form of a wire with a diameter of 1 mm were selected for the tests. The chemical composition of the steel is presented in Table 1. Table 1: Chemical composition of X10CrNi18-8 steel in mass fractions, w/% Tabela 1: Kemijska sestava jekla X10CrNi18-8 v masnih dele`ih, w/% Steel C Mn Si P S Cr X10CrNi18-8 0.08 0.91 0.68 0.028 0.001 17.96 The differentiation of the surface roughness was achieved by means of mechanical treatment – grinding (Ra = 0.60 μm) and mechanical polishing (Ra = 0.14 μm). Chemical passivation was made in HNO3 40 %. Next, the samples were subjected to sterilisation with ethylene oxide. The sterilisation was performed at the temperature of 55 °C in a Steri-Vac 5 XL steriliser. Before the tests, all the samples were cleaned in ethanol 96 % in an ultrasonic disintegrator. Next, the samples corresponding to the consecutive stages of the surface preparation were tested for the resistance to pitting corrosion. The tests were carried out in accordance with ASTM.8,9 Polari- sation curves were registered by means of potentiostat PGP-201 by Radiometer. A saturated calomel electrode (SCE) of the KP-113 type was used as the reference electrode. A platinum electrode of the PtP-201 type served as the auxiliary electrode. The change in the potential in the anodic direction took place at the rate of 1 mV/s. Once the current density reached the value of 1 mA/cm2, the polarisation direction was changed and, at the same time, a return curve was registered. The tests were made in an alternative solution simulating a human-blood environment – in an artificial-blood plasma. The chemical composition of the artificial-blood plasma is presented in Table 2. The temperature of the artificial-blood plasma during the tests was (37 ± 1) °C and pH = 7.2.10,11 The Stern method was applied for determining the parameters typical of the corrosion resistance of the tested alloys. Table 2: Chemical composition of artificial plasma Tabela 2: Kemijska sestava umetne plazme Component Amount in distilled water g/L NaCl 6.8 CaCl2 0.2 KCl 0.4 MgSO4 0.1 NaHCO3 2.2 Na2HPO4 0.126 NaH2PO4 0.026 XPS (X-ray photoelectron spectroscopy) was applied for identifying the chemical composition of the sample surface layer before and after the chemical-passivation process and sterilization in ethylene oxide. A photo- electron spectrometer made by Prevac with an analyser and monochromatic X-ray source by VG Scienta were used. Photoelectrons were activated with an X-ray tube with an aluminium anode and a quartz monochromator that ensured the Al K radiation with an energy of 1486.6 eV. The spectra in a wide range of electron bond energy as well as detailed spectra with high resolution made as in-depth profiles were tested. Ion etching was performed with Ar+ ions with an energy of 4 keV. The analysed area was rectangular with the following dimensions: 260 μm × 1000 μm, with the longer side parallel to the axis of the samples. The ion-etching area was 2000 μm × 2000 μm, which ensured a safe distance from the edge of the crater created as a result of the argon-ion bombardment of the surface. The chemical composition was determined by integrating the respective photoemission lines with the application of programme MULTIPAK by Physical Electronics. W. WALKE, J. PRZONDZIONO: POTENTIODYNAMIC AND XPS STUDIES OF X10CrNi18-8 STEEL ... 224 Materiali in tehnologije / Materials and technology 49 (2015) 2, 223–227 Table 3: Potentiodynamic-test results for X10CrNi18-8 steel Tabela 3: Rezultati potenciodinami~nih preizkusov na jeklu X10CrNi18-8 Surface type Corrosion potential Ecorr/mV Breakdown potential Eb/mV Polarisation resistance (average) Rp/(k cm2) Corrosion current density (average) icorr./(μA/cm2) Before EO Polished –63 to –55 +377 to +383 229 0.114 Passivated +15 to +25 +690 to +700 1070 0.024 After EO Polished –81 to –71 +520 to +560 1060 0.025 Passivated –72 to –63 +790 to +830 2470 0.010 3 RESULTS AND DISCUSSION The performed potentiodynamic tests in the artifi- cial-blood plasma supplied the information regarding the corrosion resistance of austenitic steel X10CrNi18-8 with its surface prepared in different ways and subjected to the sterilisation process. Potentiodynamic-test results are presented in Table 3. Anodic-polarisation curves are shown in Figure 1. It was determined that the average corrosion-poten- tial value for the wire subjected to electrochemical polishing was lower (Ecorr = –59 mV) than that obtained for the wire subjected to chemical passivation (Ecorr = +20 mV). The mean values of the perforation potential and possible repassivation, determined on the ground of the registered polarisation curves, were also lower for the wire subjected to polishing. It was also proved that there were significant differences between the determined values of corrosion-current density icorr and polarisation resistance Rp (Figure 1 and Table 3). In order to identify the chemical composition of the surface layer created as the result of the electrochemical treatment and sterilisation process, XPS tests were performed. An analysis of the spectral lines for the respective elements enabled us to draw the conclusions regarding their chemical states and their changes depending on the depth of the analysed layer (Figure 2). The bond energy of electron states depends on the chemical compound, in which an element is present. A favourable decrease in the main alloying elements of Fe, Cr and Ni in relation to the substrate, as well as an increase in the participation of oxide compounds of those W. WALKE, J. PRZONDZIONO: POTENTIODYNAMIC AND XPS STUDIES OF X10CrNi18-8 STEEL ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 223–227 225 Figure 3: XPS spectra of X10CrNi18-8 steel for lines (electrochemi- cal polishing): a) O1s, b) Ni2p3/2, c) Fe2p3/2, d) Cr2p Slika 3: XPS-spektri za jeklo X10CrNi18-8 za linije (elektrokemijsko poliranje): a) O1s, b) Ni2p3/2, c) Fe2p3/2, d) Cr2p Figure 1: Polarisation curves determined for X10CrNi18-8 steel after electrochemical surface treatment: a) before sterilisation in ethylene oxide, b) after sterilisation in ethylene oxide Slika 1: Polarizacijske krivulje, ugotovljene pri jeklu X10CrNi18-8 po elektrokemijski obdelavi povr{ine: a) pred sterilizacijo v etilen oksidu, b) po sterilizaciji v etilen oksidu Figure 2: In-depth profiles registered for X10CrNi18-8 steel after the processes of: a) electrochemical polishing, b) electrochemical polishing and sterilisation in ethylene oxide, c) chemical passivation, d) chemical passivation and sterilisation in ethylene oxide Slika 2: Profil koncentracije v globino pri jeklu X10CrNi18-8 po: a) elektrokemijskem poliranju, b) elektrokemijskem poliranju in steri- lizaciji z etilen oksidom, c) kemijski pasivaciji, d) kemijski pasivaciji in sterilizaciji z etilen oksidom elements were detected in the composition of the passive layer obtained after the electrochemical treatment and sterilisation in ethylene oxide (Figures 3 to 6). Oxide compounds feature better haemocompatibility. Measured line O1s may be subordinate to several chemical states. The main line with the energy of 285.29 eV could be assigned to hydrocarbons, ever present on the surface. A weak line with the maximum energy of 283.50 eV that came from carbides, e.g., Cr3C2, was also visible. Other weak lines at their maximum levels, with their energy slightly higher than 285.30 eV could be attributed to the organic compounds containing oxygen or the carbonates created during the surface treatment of the sample. With respect to analysing the Fe spectrum, two chemical states are visible: metallic Fe (707 eV) and the form of oxide, Fe2O3 (711 eV), which may be determined through a comparison with the spectra of the respective iron oxides and an analysis of the satellites, e.g., 719 eV – the typi- cal position of Fe2O3.12 For chromium, the oxidized condition was dominant; line Cr2p3/2 with an energy of 577.0 eV was emitted by Cr2O3, whereas the weak line with an energy of 574.4 eV was emitted by metallic chromium. The line of oxygen included various chemical states. The line with an energy of about 531.6 eV, which could be attributed to Cr2O3, was dominant.13 For the transition metals, such as chromium or nickel, the effect of the reduction due to the ion bombardment was taken into consideration. This effect is particularly important when interpreting the nickel spectra. Nickel oxides are largely reduced due to the ions with an energy of 4 keV. The absence of nickel oxides in the spectra may just partially result from this effect (Table 4). Table 4: Results of XPS analyses Tabela 4: Rezultati XPS-analiz Surface type Elements in amount fractions, x/% O C Fe Cr Ni Before EO Polished 35.16 52.05 7.61 5.17 - Passivated 55.11 41.21 2.23 1.03 0.41 After EO Polished 34.55 50.23 9.29 5.92 - Passivated 60.76 33.25 3.25 3.52 0.21 W. WALKE, J. PRZONDZIONO: POTENTIODYNAMIC AND XPS STUDIES OF X10CrNi18-8 STEEL ... 226 Materiali in tehnologije / Materials and technology 49 (2015) 2, 223–227 Figure 4: XPS spectra of X10CrNi18-8 steel for lines (electroche- mical polishing and sterilisation in ethylene oxide): a) O1s, b) Ni2p3/2, c) Fe2p3/2, d) Cr2p Slika 4: XPS-spektri za jeklo X10CrNi18-8 za linije (elektrokemijsko poliranje in sterilizacija z etilen oksidom): a) O1s, b) Ni2p3/2, c) Fe2p3/2, d) Cr2p Figure 6: XPS spectra of X10CrNi18-8 steel for lines (chemical passivation and sterilisation in ethylene oxide): a) O1s, b) Ni2p3/2, c) Fe2p3/2, d) Cr2p Slika 6: XPS-spektri za jeklo X10CrNi18-8 za linije (kemijska pasivacija in sterilizacija z etilen oksidom): a) O1s, b) Ni2p3/2, c) Fe2p3/2, d) Cr2p Figure 5: XPS spectra of X10CrNi18-8 steel for lines (chemical passivation): a) O1s, b) Ni2p3/2, c) Fe2p3/2, d) Cr2p Slika 5: XPS-spektri za jeklo X10CrNi18-8 za linije (kemijska pasivacija): a) O1s, b) Ni2p3/2, c) Fe2p3/2, d) Cr2p 4 CONCLUSIONS An innovative progress in treating cardiovascular- system diseases using low-invasive methods led to the production of new forms of tools, such as guide wires. Those tools, introduced as the first ones to a narrowed section of a blood vessel, many times increase its cross- section in a situation of its total closure. Some limita- tions, observed in the clinical practice, resulting from the introduction of the tools made of metallic materials to blood vessels, are mainly connected with the blood coagulation on their surfaces.14 Therefore, the study contains an evaluation of the impact of the surface treatment that enables us to increase the biotolerance of the X10CrNi18-8 steel in a blood environment under the conditions of sterilisation with ethylene oxide. The analysis of the results of the electrochemical- corrosion tests proved differentiated corrosion resistance of the wires made of the X10CrNi18-8 steel. Less favourable values of the determined parameters that cha- racterise the corrosion resistance result from restricting their surface treatment only to polishing (Table 3). Next, the in-depth profiles obtained for individual variants of the surface treatment showed that electroche- mical polishing, chemical passivation and sterilisation in ethylene oxide caused an increase in the oxygen in the surface layer in relation to the substrate, mainly resulting in the oxides such as: Cr2O3 and Fe2O3. Consequently, the concentrations of individual alloying elements (Fe, Cr, Ni) in relation to the chemical composition of the substrate were substantially decreased, which is found to be extremely favourable for the haemocompatibility of guide wires. All the elements occurred mainly in the oxi- dised condition. To sum up, it must be stated that the chemical-passi- vation process performed after electrochemical polishing of the X10CrNi18-8 steel has a favourable impact on the physical and chemical characteristics, improving the haemocompatibility of the cardiologic guide wires made of this type of steel. Acknowledgements This project was financed from the funds of the National Science Centre in Cracow. 5 REFERENCES 1 M. Kaczmarek, W. Walke, W. Kajzer, Arch. Mater. Sci. Eng., 28 (2007) 5, 273–276 2 J. Marciniak, J. Tyrlik-Held, W. Walke, Z. Paszenda, Eng. of Biomat., X (2007) 69–72, 90–93 3 D. Baim, W. Grossman, Angiography and Intervention, 6th edition, LWW, Philadelphia 2000 4 D. Baim, W. Grossman, Angiography and Intervention, 7th edition, LWW, Philadelphia 2006 5 M. Dewey, F. Teige, D. Schnapauff, Ann. of Int. Med., 145 (2006) 6, 407–415, doi:10.7326/0003-4819-145-6-200609190-00004 6 W. Walke, J. Przondziono, Metalurgija, 50 (2011) 3, 201–204 7 W. Walke, J. Przondziono, Arch. of Metall. and Mat., 58 (2013) 2, 625–630, doi:10.2478/amm-2013-0048 8 ASTM F2129-08 Standard Practice for Selecting Generic Biological Test Methods for Materials and Devices, 2008 9 ASTM F746-04(2009): Standard Test Method for Pitting or Crevice Corrosion of Metallic Surgical Implant Materials, 2009 10 J. Marciniak, J. Tyrlik-Held, W. Walke, Z. Paszenda, Arch. Mater. Sci. Eng., 28 (2007) 5, 289–292 11 W. Kajzer, A. Krauze, W. Walke, J. Marciniak, J. Achiev. Mat. Manuf. Eng., 18 (2006), 115–118 12 C. R. Clayton, Y. C. Lu, J. Electrochem. Soc., 133 (1986) 12, 2465–2473, doi:10.1149/1.2108451 13 J. M. Grimal, P. Marcus, Corr. Sci., 5 (1992), 805–814, doi:10.1016/ 0010-938X(92)90113-H 14 A. Tortoriello, G. Pedrizzetti, J. of Biomech., 37 (2004), 1–11, doi:10.1016/S0021-9290(03)00259-8 W. WALKE, J. PRZONDZIONO: POTENTIODYNAMIC AND XPS STUDIES OF X10CrNi18-8 STEEL ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 223–227 227 J. Ma. RINCÓN, R. CASASOLA: TEM REPLICA OF A FLUORIDE-MISERITE GLASS-CERAMIC GLAZE MICROSTRUCTURE TEM REPLICA OF A FLUORIDE-MISERITE GLASS-CERAMIC GLAZE MICROSTRUCTURE TEM-REPLIKE MIKROSTRUKTURE STEKLOKERAMI^NE FLUOR-MIZERITNE GLAZURE Jesus Ma. Rincón, Raquel Casasola Vitreous and Ceramics Lab/Group, Instituto CC Construcción E. Torroja, CSIC, Madrid jrincon@ietcc.csic.es Prejem rokopisa – received: 2014-01-27; sprejem za objavo – accepted for publication: 2014-04-10 doi:10.17222/mit.2014.020 Several glazes were obtained for an application onto the surface of clay-based ceramic tiles that are usually produced by fast firing. Original glasses within the F-K2O-CaO-SiO2 system were obtained and during the sinter- crystallization of glass ceramic, fluoride-miserite (KCa5Si8O22F2) nanocrystals with an amount of 4.5 % F2 were precipitated in the glassy matrix. These glassy tiles can be used for floor and wall coverings in buildings, for certain decorations in civil engineering and for pavements exhibiting good wear and anti-slip properties. The TEM-replica and extraction-replica methods were revisited for observing the microstructures of these glasses under the conventional TEM to elucidate clearly the phase separation that gives rise to the opalescence of these glasses and is strongly dependent on the fluorine included in the structures of these glasses. Keywords: TEM-replica method, glass ceramics, glazes, tiles, miserite Pridobljenih je bilo ve~ glazur za uporabo na povr{ini kerami~nih plo{~ic na osnovi gline, ki se navadno proizvajajo s hitrim `ganjem. V originalnem steklu iz sistema F-K2O-CaO-SiO2 so bili s postopkom sintranja in kristalizacije steklokeramike izlo~eni nanokristali fluor-mizerita (KCa5Si8O22F2) z vsebnostjo 4,5 % F2. Te steklaste plo{~ice se lahko uporabijo na tleh ali stenah poslopij, ali kot dekoracija v gradbeni{tvu in za plo~nike, ker izkazujejo dobre obrabne in protizdrsne lastnosti. Uporabljena je bila metoda TEM-replik in ekstrakcijskih replik za opazovanje mikrostrukture teh stekel v konvencionalnem TEM za pojasnitev lo~evanja faz, ki povzro~ijo opalescenco teh stekel, kar je mo~no odvisno od fluora, ki je vklju~en v strukturi teh stekel. Klju~ne besede: metoda TEM-replik, steklokeramike, glazura, plo{~ice, mizerit 1 INTRODUCTION Glass ceramics developed widely in the previous century and found new applications in the first decade of the 21st century1. These applications include the use of the glass-ceramic processing (nucleation + crystal growth) in the production of improved glazes for ceramic-tile coverings2. These new glazes produced from frits give rise to a wide range of crystallites precipitated in the vitreous matrix of a glaze and are compatible with the fast-firing process and, even more importantly, with the traditional clay-ceramic and/or porcelainized substrates. The liquid-liquid phase separation or immiscibility in the glazes susceptible to a transformation in microcrystalline coatings has not yet been well described for ceramic tiles. With respect to the above, the glass ceramics, in which the fluoride phases are crystallized, are of great interest for the production of the materials with high toughness and flexural-strength values. The aim of this work is to show the microstructures obtained in several original K2O-CaO-SiO2 glasses with different fluorine amounts, close to the composition of the crystalline phase called miserite K(Ca,Ce)6Si8O22(OH,F)2.3 Though TEM/EDS allow us to carry out a microstructure determination of ceramics4, due to the difficulties of the ion-thinning preparation, it was the objective of this research to apply the old traditional extraction carbon-triafol replica method to demonstrate the capability of observing the glass-in- glass or liquid-liquid phase separation5,6, which addition- ally allowed us to simultaneously perform an analysis of the particle extraction, thus saving the time and cost required for the thinning methods. 2 MATERIALS AND METHODS The original or starting glasses were obtained from some batch compositions with a miserite stoichiometric composition, K(Ca,Ce)6Si8O22(OH,F)2, without cerium, where the hydroxyl ions were first substituted with fluorine and later, in the second series, corrected with the consecutive additions of mass fractions (5, 8, 9 and 10) % fluorine added as a pure chemical CaF2 powder. Potassium and calcium were added as carbonates and alumina, as well as silica, as pure oxides. The melting was carried out in a super-kanthal furnace at 1450 °C for one hour using alumina-silica crucibles from Lomba SL, Galicia, Spain. Table 1 shows the initial, basic compo- sitions and Table 2 shows the corrected and final compositions determined with SEM/EDS for the M-4, M-8 and M-10 glasses. The M-5 composition is the Materiali in tehnologije / Materials and technology 49 (2015) 2, 229–233 229 UDK 666.3/.7:620.187 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)229(2015) closest to the stoichiometric miserite, being similar to M-10 but including Na2O to facilitate the melting and the pouring operation after the melting. The M-4 glass includes some alumina stabilizing this glass and a higher amount of K2O, while M-8 and M-9 allow a comparison of the increased calcium additions. As can be seen in the original composition, the basicity or the CaO/SiO2 ratio, which is well known, is a factor controlling the viscosity variation at high temperatures in the 0.29–0.58 range. Table 1: Initial or theoretical compositions of the “miserite glasses”, studied with TEM, in mass fractions, w/% Tabela 1: Za~etna ali teoreti~na sestava “mizeritnih stekel”, preiskovanih s TEM, v masnih dele`ih, w/% Oxide (w/%) M-4 M-5 M-8 M-9 M-10 Na2O – – – – 5.00 K2O 7.35 5.56 6.66 6.24 5.29 CaO 26.64 33.14 20.00 25.00 31.48 Al2O3 8.00 – – – – SiO2 55.00 56.81 67.97 63.72 53.97 F (expressed as F2) 3.00 4.49 5.37 5.04 4.27 Basicity (CaO/SiO2) 0.48 0.58 0.29 0.39 0.58 After the melting of these original glasses, it was necessary to prepare new glasses with the increasing and sequential fluorine additions in order to determine the corrosion of the crucibles and the volatilization of fluo- rine in this type of glasses. Figure 1 shows that fluorine volatilization growths linearly when the fluorine amount in the miserite glasses is increased. In the same way Figure 2 shows the alumina amounts in the final glasses prepared with the increasing fluorine amount. It can be seen that the alumina amount added to these glasses on the basis of the corrosion of the alumina-silica crucibles also grows linearly, with the silica being in the same range for all the glasses considered, as can be seen in Table 2, affecting the actual compositions of the glasses analysed with SEM/EDS. Table 2: Compositions, obtained with SEM/EDS, of the glass samples corrected with the increasing fluorine amount (fluorine and alumina are deduced from the crucible corrosion and the volatilization curves for M-5 and M-9 glasses) Tabela 2: SEM/EDS-sestava v popravljenih vzorcih stekla s pove~ano vsebnostjo fluora (fluor in glinica izvirata iz korozije lonca in krivulje izhlapevanja za stekla M-5 in M-9) Oxide (w/%) M-E M-5 M-8 M-9 M-10 K2O 4.79 4.61 4.56 4.52 4.48 CaO 34.02 33.85 31.37 31.73 32.21 Al2O3 4.63 4.98 7.66 7.70 7.87 SiO2 54.05 53.58 52.13 51.37 50.82 F (expressed as F2) 2.51 2.80 4.27 4.90 5.63 Basicity (CaO/SiO2) 0.62 0.63 0.60 0.60 0.63 Then, the final compositions of the sequential fluo- ride-miserite glasses show that the potassium amount is stable in different compositions, being volatile to the same extent in all the glasses studied. Calcium is also stable in all the glasses and close to the stoichiometric miserite composition. Even the basicities of these glasses are very similar, with the increasing fluorine amount being in a very narrow range of 0.60–0.63. However, the alumina based on the crucible corrosion grows with the increasing fluorine amount. In order to investigate the microstructures of the original or starting glasses formulated around the miserite composition, the TEM-replica method was used to avoid the expensive and time-consuming preparations of the ion-thinning methods. Thus, the double triafol- carbon replica method used for years for observing very J. Ma. RINCÓN, R. CASASOLA: TEM REPLICA OF A FLUORIDE-MISERITE GLASS-CERAMIC GLAZE MICROSTRUCTURE 230 Materiali in tehnologije / Materials and technology 49 (2015) 2, 229–233 Figure 1: Absolute values of fluorine losses depending on the fluorine added to the original glasses Slika 1: Absolutna vrednost izgub fluora v odvisnosti od dodanega fluora originalnim steklom Figure 2: Alumina amounts determined with SEM/EDS versus the fluoride added to the sequential original glasses Slika 2: Vsebnost glinice, ugotovljene s SEM/EDS, v odvisnosti od dodajanja fluorida zaporednim steklom fragile samples with TEM was revisited as shown by Rincón et al.7 Figure 3 shows a simple drawing of the method steps from a fresh fracture and HF 2 %, 10 s, to the etched glass surfaces. The final carbon-replica skins were deposited on 3 mm Cu grids after transfering, smooth drying in an acetone bath and dissolving the first triafol plastic replica. The TEM equiment for the final observation of the replicas was a Philips CM-10 instru- ment working at 100 kV. 3 RESULTS AND DISCUSSION Figures 4 and 5 show a summary of the main obser- vations of the microstructures of the fluoride-miserite composition glasses. In the M-4 glass (Figure 4) a clear liquid-liquid phase separation of dispersed droplets with some networks in some areas can be seen. In the case of glass M-5, the glass-in-glass phase separation is not present due to the crystallization of very small crystallites with an elongated shape in some cases. According to the binary CaO–SiO2 phase diagram these crystals in the M-5 glass must be of wollastonite, coexisting with the liquid phase. For the other glasses investigated with TEM, the representative micrographs obtained with the replica observations are shown in Figures 5a to 5d. It can be seen that the M-8, M-9 and M-10 glasses show different microstructures. Thus, in the M-8 glass, there are clusters of crystals coexisting with opaque nanocrystals, possibly of cuspidine and/or miserite. The microstructure of glasses M-9 and M-10 is completely different, there are rounded and pseudo-polyonal habit crystallites that must be undissolved Ca2F in the glassy matrix. Droplets are shown at higher magnifications. In these glasses the Ca2F crystals coexist with very small nanodroplets of the residual phase separation or, more probably, the nano- crystals of miserite (dark nanocrystals) and cuspidine (white nanocrystals). Table 3 shows volume fraction of immiscibility and size of droplets and nanocrystals. When the compositions of these miserite glasses are located in the binary CaSiO3–SiO2 diagram taken from8 (Figure 6), showing the over-liquidus glass-in-glass J. Ma. RINCÓN, R. CASASOLA: TEM REPLICA OF A FLUORIDE-MISERITE GLASS-CERAMIC GLAZE MICROSTRUCTURE Materiali in tehnologije / Materials and technology 49 (2015) 2, 229–233 231 Figure 5: TEM double-replica micrographs for: a) M-5, b) M-8, c) M-9 and d) M-10 glasses: the inserted bars correspond to these mag- nifications: a) 5000 nm, b) 2000 nm, c) 5000 nm and d) 2000 nm Slika 5: Mikroposnetki TEM dvojne replike iz stekel: a) M-5, b) M-8, c) M-9 in d) M-10; vrisano merilo pomeni ustrezno pove~avo: a) 5000 nm, b) 2000 nm, c) 5000 nm in d) 2000 nm Table 3: Quantitative evaluation of the glass-in-glass phase separation and the crystallites precipitated in the miserite glasses from the TEM replica observations Tabela 3: Kvantitativna ocena izlo~anja stekla v steklu in izlo~kov kristalitov v mizeritnih steklih iz opazovanj TEM-replik Original glasses Vf (esti- mated)/% Droplet average diameter (nm ± 10 nm) Crystal average diameter (nm ± 20 nm) M-4 50 120 No crystallites M-5 30 No droplets 800–1000 M-8 80 130(interconnected) clusters = 1800 M-9 80 100 830 M-10 10 100 (very fewdroplets) 480 Figure 3: Succesive steps for the preparation of direct and double triafol-carbon replicas7 Slika 3: Zaporedje stopenj priprave neposrednih in dvojnih replik triafol – ogljik7 Figure 4: TEM micrographs of double replicas of M-4 glass: the inserted bars correspond to these magnifications: a) 5000 nm, b) 1000 nm, c) 1000 nm and d) 500 nm Slika 4: TEM-posnetki dvojnih replik iz stekla M-4: prikazano merilo ustreza naslednjim pove~avam: a) 5000 nm, b) 1000 nm, c) 1000 nm in d) 500 nm phase separation, it can be seen that all the considered compositions are far from this zone and that droplets of the phase separation are present in the M-4, M-8 and M-9 glasses. The M-8 and M-9 compositions are located in the two phases of silica + liquid area below the liqui- dus. The M-4 composition is very close to the eutectic, between the above-mentioned area and the wollastonite + liquid area below the liquidus. The M-5 and M-10 glasses with very similar compositions are located in the wollastonite + liquid area. Therefore, these glasses were produced during the cooling as a precipitation of very small crystals of wollastonite (CaO–SiO2) with no glass-in-glass phase separation. Conversely, the compo- sitions of M-8 and M-9 show a phase separation during the TEM replica observations of the droplets that must be enriched with SiO2 according to the theory of the phase separation in glasses5,6. This fact made us think that there is a dome below the liquidus in this zone, con- sisting of silica + liquid zone and that this is the only explanation for the immiscibility in these glasses, con- sidering the simplification of this composition system. This proposed cupula or dome of immiscibility is congruent with the inclusion of the third modifier cations in these glasses (Tables 1 and 2) and the relative location of immiscibility in the binary systems, as can be seen in Figure 6, where an extension of the immiscibility in this area and/or, as proposed, the second immiscibility dome occurring below the liquidus can be found in this area of the simplified phase diagram. Even more, if we consider the ternary systems of Na2O–CaO–SiO2, K2O–CaO–SiO2 and Al2O3–CaO–SiO2 the phase-separation areas are very J. Ma. RINCÓN, R. CASASOLA: TEM REPLICA OF A FLUORIDE-MISERITE GLASS-CERAMIC GLAZE MICROSTRUCTURE 232 Materiali in tehnologije / Materials and technology 49 (2015) 2, 229–233 Figure 9: Microstructure of the M-8 miserite glass showing a spino- dal-like decomposition Slika 9: Mikrostruktura mizeritnega stekla M-8, ki ka`e razgradnjo, podobno spinodalni reakciji Figure 7: a) Liquid-phase separation below the solidus line in a binary system and b) liquid-phase separation below the liquidus in a binary system9 Slika 7: a) Izlo~anje staljene faze pod solidusno linijo v binarnem sistemu in b) izlo~anje staljene faze pod likvidusno linijo v binarnem sistemu9 Figure 8: Ternary compositions of glasses showing the narrow glass-in-glass and/or liquid-phase separation zones for the ternary systems: RO–K2O–SiO2, RO–Na2O–SiO2, RO–Al2O3–SiO2 and K2O–CaO–SiO2 6 Slika 8: Ternarna sestava stekel, ki prikazuje ozko podro~je steklo v steklu in/ali podro~je izlo~anja taline v ternarnih sistemih: RO–K2O– SiO2, RO–Na2O–SiO2 , RO–Al2O3–SiO2 in K2O–CaO–SiO2 6 Figure 6: CaSiO3–SiO2 binary system with the liquid-liquid phase- separation zone (a phase-separation dome below the liquidus and near the eutectic and/or an extension to the third component are proposed) (modification of the figure from8) Slika 6: Binarni sistem CaSiO3–SiO2 z lo~enim podro~jem izlo~anja talina-talina (predlagano je izlo~anje kupolaste faze pod likvidusom blizu evtektika in/ali raz{iritev do tretje komponente; prirejeno po viru8) narrow being the miserite glasses within the limits of these immiscibility zones. Therefore, the presence of liquid immiscibility in these glasses is very critical but, in any case, favored by the presence of fluorine ions in a glass network structure. As shown in Figure 6, this area of the metastable immiscibility in miserite glasses can be produced at the 1436 °C solidus limit and extended below such a solidus line in the metastable conditions, which are the usual liquid-freezing conditions for obtaining glasses. Another possible hypothesis is that the known area of the meta- stable miscibility at the solidus where the eutectic and the liquidus line of the monotectic reaction (1705 °C) are located can be extended or widened to the lower silica amount in this diagram. This situation is well-known in the binary R2O–SiO2 systems as shown in Figure 79 and even in the corresponding ternary diagrams (Figure 8)5,6. The miserite M-8 glass shows, in some zones, an interconnected-droplet microstructure, while this type of immiscibility microstructure was not observed in the other miserite glasses (Figure 9). In principle, this may correspond to the spinodal decomposition as demon- strated by the Cahn theory for the compositions in the center of the phase-separation cupolas5,6. In this case, while observing the relative location of the M-8 glass in a simplified pseudobinary system involving CaO (Figure 6) it can be seen that this composition is close to the centre of the proposed second dome, close to Tc, the temperature of the maximum immiscibility for the binary domes. Therefore, in spite of the very narrow areas of the phase separation in the basic composition of the K2O–CaO–Al2O3–SiO2 glasses and the similarly that was demonstrated by Chiang and Kingery10 it is possible to obtain the glasses with a liquid-liquid immiscibility according to the several microstructures of the phase- separation phenomena. Finally, it is evident that this immiscibility can be observed with scanning electron microscopy (SEM). However, our experience in observing this effect in glasses showed that a clear definition of immiscibility is only possible when revisiting the traditional TEM replica method, as was the case in this investigation. More re- search is now in progress involving different composi- tions of mica glass-ceramics such as fluorphlogopite and others11. 4 CONCLUSIONS The TEM replica method was revisited to study the microstructures of the original fluoride-miserite glasses, using the conventional TEM, and to elucidate the phase separation inherent to these opalescent glasses. These liquid-liquid and/or glass-in-glass phase separations strongly depend on the fluorine level in a vitreous structure, reaching the maximum values with the lowest fluorine amounts. At higher fluorine concentrations there are precipitations of Ca2F and the nanocrystals of wollastonite, cuspidine and/or miserite, depending on the glass composition. Acknowledgement The authors wish to thank for the facilities provided by the Polytechnic University of Valencia (UPV), Servei de Microscopia Electronica, and the valuable help from Manuel Planes when using the TEM equipment. They also thank Eduardo Cabrero, IETcc, CSIC, for helping them draw the original figures and Dr. M. Romero from the IETcc-CSIC for giving valuable advice to R. Casa- sola regarding her Ph. D. Thesis. 5 REFERENCES 1 A. G. Guy, Essentials of Materials Science, McGraw-Hill, 1976, 97 2 J. Ma. Rincón, Principles of nucleation and controlled crystallization of glasses, Polym. Plast. Technol. Engineering, 31 (1992) 3–4, 309–357, doi:10.1080/03602559208017751 3 R. Casasola, J. M. Pérez, J. Ma. Rincón, M. Romero, 50th Congress of the Spanish Glass and Ceramics Society VI-P-01, Bol. Soc. Esp. Ceram. Vidr., 49 (2010) 5, 35 4 J. Ma. Rincón, Microstructural characterization by electron micro- scopy of ceramics and glasses, Microscopy and Analysis, (1996), 23–25 5 J. Ma. Rincón, Separación de fases en el vidrio, Bol. Soc. Esp. Ceram. Vidr., 11 (1972) 1, 111–125 6 J. Ma. Rincón, A. Durán, Separación de fases en Vidrios, El Sistema Na2O-B2O3-SiO2, Edited by the Spanish Glass and Ceramic Society, SECV, Arganda del Rey, Madrid, 1982 7 J. Ma. Rincón, P. Callejas, F. Capel, Fractografía de vidrios y materiales vitrocristalinos, Bol. Soc. Esp. Ceram. Vidr., 28 (1989) 4, 257–267 8 J. R. Taylor, A. T. Dinsdale, Thermodynamic and phase diagram data for the CaO-SiO2 system, Calphad, 14 (1990) 1, 71–88, doi:10.1016/ 0364-5916(90)90041-W 9 E. Plumat, La formation de systèmes pseudo-vitreux et pseudo-cri- stallines, Silicates Industriels, 1 (1967), 5–13 10 Y. M. Chiang, W. D. Kingery, Spinodal decomposition in a K2O–Al2O3–CaO–SiO2 glass, Journal of the Amer. Ceram. Soc., 66 (1983) 9, c171–c172, doi:10.1111/j.1151-2916.1983.tb10632.x 11 R. Casasola, Ph. D. Thesis, Autonoma University of Madrid, 2013 J. Ma. RINCÓN, R. CASASOLA: TEM REPLICA OF A FLUORIDE-MISERITE GLASS-CERAMIC GLAZE MICROSTRUCTURE Materiali in tehnologije / Materials and technology 49 (2015) 2, 229–233 233 A. BENTOUHAMI, B. KESKES: EXPERIMENTAL ANALYSIS AND MODELING OF THE BUCKLING ... EXPERIMENTAL ANALYSIS AND MODELING OF THE BUCKLING OF A LOADED HONEYCOMB SANDWICH COMPOSITE EKSPERIMENTALNA ANALIZA IN MODELIRANJE UPOGIBANJA OBREMENJENEGA SATASTEGA SENDVI^NEGA KOMPOZITA Abderrahmane Bentouhami, Boualem Keskes Laboratoire de mécanique de précision appliquée, Institut d’Optique et Mécanique de Précision, Université Ferhat Abbas de Sétif, Algérie bentouhamidahman@yahoo.fr, bkeskes2012@gmail.com Prejem rokopisa – received: 2014-02-22; sprejem za objavo – accepted for publication: 2014-03-28 doi:10.17222/mit.2014.039 Sandwich panels have the best stiffness-to-lightness ratio, which is what makes them very useful in industrial applications. This paper is focused on a study of the buckling capacities of the core components under uniaxial compression. The critical buckling loads for various core densities and materials of honeycomb panels were experimentally and numerically investigated. The specimens under lateral loading showed three zones: zone 1 is the initial elastic state, followed by the plateau region in zone 2, while zone 3 shows a monotonically stiffening region, associated with the densification of the material. The effect of the core density and its materials on the behavior and the damage was highlighted. From the experiment it is clear that the buckling load of the specimens increases as the core density is increasing. In terms of stiffness and load at failure, the honeycomb sandwich panel had better mechanical characteristics than its components. The study also calculated the numerical buckling loads of the panels using the ABAQUS finite-element analysis program. The achieved experimental and numerical results were compared with each other. In conclusion, a good correlation between theory and experiment was found. Keywords: honeycomb sandwich panel, buckling analysis, compression, finite-element method, collapse Sendvi~ni paneli imajo najbolj{e razmerje med togostjo in maso. To jih dela primerne za industrijsko uporabo. Ta ~lanek je usmerjen v {tudij zdr`ljivosti za upogibanje klju~nih komponent pri enoosni tla~ni obremenitvi. Eksperimentalno in numeri~no so bile preiskane kriti~ne upogibne obremenitve za razli~ne klju~ne gostote in material satastih plo{~. Vzorci, stransko obre- menjeni, so pokazali tri podro~ja: podro~je 1 je za~etno elasti~no stanje, ki mu sledi podro~je platoja, tj. podro~je 2. Podro~je 3 prikazuje monotono upogibanje, povezano z zgo{~evanjem materiala. Ocenjen je bil vpliv klju~ne gostote in materialov glede vedenja in po{kodb. Iz preizkusov je razvidno, da z nara{~anjem klju~ne gostote nara{~a tudi odpornost proti upogibni obremenitvi vzorcev. Glede na upogibanje in obremenitev pri poru{itvi ima satasta sendvi~na plo{~a bolj{e mehanske lastnosti v primerjavi z njenimi komponentami. V {tudiji je tudi izra~unana numeri~na upogibna obremenitev panelov z analizo kon~nih elementov s programom ABAQUS. Primerjani so dobljeni eksperimentalni in numeri~ni rezultati. Dobljena je bila dobra korelacija med teorijo in eksperimentalnimi rezultati. Klju~ne besede: satasta sendvi~na plo{~a, analiza upogibanja, tla~enje, metoda kon~nih elementov, poru{itev 1 INTRODUCTION Honeycomb sandwich panels are increasingly used in engineering applications (aviation, astronautics and navigation, automotive, etc.), where a high rigidity as well as lightness is important. A number of core mate- rials and core configurations have been proposed recently. The most commonly used core materials are honeycombs and foams. Honeycomb sandwich panels are obtained by covering the upper and lower surfaces of the honeycomb with sheets. Metal or non-metal mate- rials can be used as the lower and upper surface face sheet materials of honeycomb sandwich panels. The widespread use of honeycomb in practice generated a need to establish their mechanical properties. Previous studies on the crushing behavior of honeycomb struc- tures included the early work reported by McFarland,1 who developed a semi-empirical model to predict the crushing stress of hexagonal cell structures subjected to axial loading. This model as later improved to incor- porate both the bending and extensional deformation of such cellular structures.2 Meanwhile, the mechanical properties of honeycomb structures in the lateral directions were investigated both analytically and experimentally by Gibson and Ashby,3 and Gibson et al.4 In their works, phenomenological models were proposed. More detailed analyses of the buckling of tubes with various geometries have been reported, and some of these results correlated well with the experimental data.5–9 A more complete description of the buckling mechanisms for thin-walled tubular struc- tures subjected to quasi-static and dynamic loading can be found in10,11. On the other hand, the mechanical beha- vior of sandwich composite panels made of honeycomb cellular structures has been studied extensively. For example, the force-indentation relationship for beams and plates made of sandwich polymer composites was investigated by Wu and Sun.12 This relationship was later adopted by Lee and Tsotsis13 to develop an impact model to predict the transient responses of sandwich composite Materiali in tehnologije / Materials and technology 49 (2015) 2, 235–242 235 UDK 519.61/.64:620.174:620.173 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)235(2015) parcels. As for sandwich plates and shells whose faces were made of metal, experimental results have been reported by Goldsmith and Sackman.14 A plasticity model has been reported by Jamjian et al.15 for metallic sandwich panels subjected to impact, in which the honeycomb was treated as a continuum with a discon- tinuous density. Kaman et al.16 have experimentally and numerically investigated the behavior and failure mechanisms of honeycomb core panels. They have determined that the critical buckling load of Nomex is higher than that of aluminum comb composite panels for all cell sizes. Castanié et al.17 showed that the compression load is essentially taken by the vertical edges of the hexagonal cell. From the axial compression collapse tests on the aluminum honeycomb sandwich panel specimen, various potential influential parameters, i.e., the core height, core-cell thickness and panel aspect ratio, it was ob- served that the core height would be a crucial parameter affecting the sandwich panel’s ultimate compressive strength.18 Tsang and Lagace19 reported the different failure mechanisms in impact-damaged sandwich panels subjected to uniaxial compressive loads. They observed that the damage propagation and final failure modes were dependent on the relative extents of the core and face-sheet damage. They reported dimple propagation across the width of the panel, which occurred in the presence of core damage, with the final failure mode being a face-sheet fracture. Zhou and Mayer20 studied the shear, tearing, and compression tests over honeycomb aluminum, which showed different failure modes in- volved in a general crash accident. Mohr and Do- yoyo,21,22 Hong et al.23 performed multi-axial loading tests of honeycomb materials and derived the macrosco- pic yield functions for the honeycomb materials. Wilbert et al.24 proved that following an initial linear response, the cell walls buckle elastically. The post-buckling response is initially stiff and stable, but the inelastic action progressively softens it, leading to a limit load instability. The deformation localizes first at mid-height in the form of a sharp buckle, which with the load con- tinuing to drop more into folding. When the walls of the fold come into contact, the local collapse is arrested, the load begins to recover, and a second fold develops on one side of the first one. The second fold in turn col- lapses, forming a new load peak and a second trough. This progressive folding keeps repeating until the whole panel is consumed and the structure returns to a stiff response. The present study is concerned with the more tra- ditional problem of transverse compression. In particular, we aim to establish all aspects of the compressive res- ponse and explain the buckling phenomena of honey- comb sandwich structures. The critical buckling loads for various core densities and the materials of honey- comb panels are experimentally and numerically investi- gated. The different specimens exhibited similar load/ displacement curves and the differences observed were only due to the behavior of the different materials. The study also calculates the numeric buckling loads of the panels using the ABAQUS finite-element analysis pro- gram. The achieved experimental and numerical results are compared with each other and the results are presented in curves. In conclusion, a good correlation between theory and experiment was found. 2 EXPERIMENTAL PROCEDURES The critical buckling loads and crushing behavior of the honeycomb sandwich panels were determined by running through compression tests using a computerized universal testing machine Zwick/Roell (100 kN) (Figure 1). The test procedure for the compressive properties was A. BENTOUHAMI, B. KESKES: EXPERIMENTAL ANALYSIS AND MODELING OF THE BUCKLING ... 236 Materiali in tehnologije / Materials and technology 49 (2015) 2, 235–242 Figure 2: Typologies of investigated honeycomb sandwich: a) aluminum core, b) Nomex core Slika 2: Vrste preiskovanih satastih sendvi~nih plo{~: a) jedro iz aluminija, b) jedro nomex Figure 1: Experimental set-up of the compressive test Slika 1: Eksperimentalni sestav za tla~ni preizkus as per the ASTM C 365 standards. Aluminum honey- comb with cell sizes of (3.2, 6.4, 9.6 and 19.2) mm and densities of (29, 41, 82 and 130) kg/m3, Nomex honey- comb with a cell size of 3.2 mm and densities of (48, 80, 128 and 144) kg/m3 were used to complete the tests series (Figure 2). The dimension of the compressive specimen was 50 mm × 50 mm × 10 mm. The tests were performed at a constant cross-head speed of 0.5 mm/min. 2.1 Modeling of the cell walls of honeycomb composi- tes When the honeycomb composite was loaded in com- pressive mode, it was assumed that a uniform com- pression was achieved on the two edges parallel to the compressive loading direction of each wall, as shown in Figure 3. It was also assumed that the cell walls of the honeycomb composite were rigidly constrained by the neighboring cell walls and that all the cell walls were deformed to the same strain. Therefore, the compressive stress of the honeycomb composite is the sum of the stresses carried by the individual cell walls. The assump- tions, due to its geometrical symmetry of the cross-sec- tion, are as follows:25 When the thickness hc of a honeycomb core is not large compared with the length of side a, the buckling mode of a cell shell is based on a cell wall, every wall has a similar buckling mode, and the phases among the cell walls are the same or reversed. At the same time, every prismatic edge remains a straight line, and each cell wall looks like a rectangular thin-walled plate simply supported on all four edges. If a sandwich panel with a hexagonal cell core is thick, the buckling mode of the cell shell is a deflection of the axial centerline, and the deformation of every prismatic edge is similar to that of he axial centerline. According to Equation (1), every wall becomes a rectan- gular, thin-walled, nd simply supported on all four edges, as described in Figure 3. The buckling investigation on a ell wall can be substituted for that of the entire cell shell. The governing differential equation of the cell wall can be expressed as follows: D W x W x y W y F W xx ∂ ∂ ∂ ∂ ∂ ∂ ∂ ∂ ∂ 4 4 4 2 2 4 4 2 22+ + ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ + (1) The boundary conditions are often written as: W = 0, ∂ ∂ 2 2 W y = 0 when y = 0, a (2) W = 0, ∂ ∂ 2 2 W x = 0 when x = 0, hc (3) Under the restriction of boundary conditions, the solution (local buckling) of Equation (1) can be written as: W x y A m x h n y amnnm ( , ) sin sin= = ∞ = ∞ ∑∑ π π c11 (4) The formula for a rectangular cell wall under equal uniform compression on two opposite edges, hc, was shown as in Equation (5). The theoretical compressive stress on a cell wall used in this study was based on Zhang and Ashby’s model and can be expressed as follows: F K E v t h = − ⎛ ⎝ ⎜ ⎞ ⎠ ⎟C c c1 2 3 (5) where KC is the end constraint factor in the compression mode and its value26,27 is 5.73, E is the elastic modulus of the cell walls, v is the Poisson’s ratio of the cell walls, tc is the thickness of the cell wall and hc is the length of the free wall. Equation (1) is expressed as the load F on a cell wall. The compressive load of the individual hexagonal cell of the honeycomb core is the sum of the loads carried by the individual cell walls. The total compressive load is 10F, which is the sum of the compressive load, 2F, carried out by the free walls with single thickness and the compressive load, 8F, and carried by the ribbon with double thickness because the load is proportional to the cube of the thickness, as shown in Equation (5). The area, Ahex, of the individual hexagonal cell in a honey- comb core is calculated as 2hc cos  × hc sin  + 2hc cos  × hc, where  is the angle of the inclined cell wall. The compressive strength, C, carried by the unit hexagonal cell, is expressed as in Equation (2): c hex C= = − + 10 5 1 12 3 3 F A K E v t hc( ) cos( sin ) (6) 3 RESULTS AND DISCUSSION Figure 4 shows a typical stress-strain curve obtained from a compressive test of the honeycomb composite. The compressive deformation process can be catego- rized into three regions (1, 2 and 3) based on the com- pressive stress-strain behavior. The figure shows that the stress-strain relationship is linear in Region 1 up to the bare compressive strength. The honeycomb cell walls are in the elastic buckling condition in Region 1 (Figure 5). A. BENTOUHAMI, B. KESKES: EXPERIMENTAL ANALYSIS AND MODELING OF THE BUCKLING ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 235–242 237 Figure 3: Modeling of honeycomb core with hexagonal cell under uniform compression loading Slika 3: Modeliranje satastega jedra s heksagonalnimi celicami pri enakomerni tla~ni obremenitvi Later, a sudden decrease in the compressive stress occurs in Region 2. In this region, the core walls are in the plastic buckling condition and, as a result, wall fold- ing occurs (Figure 5). The compressive stress remains approximately stable in Region 3 until the densification of the folds in the honeycomb core. This stable stress value is defined as the crushing strength. In this region the crushing and fracture of the cores start (Figure 5). Depending on the core densification at the end of Region 3, an increase in the compressive stress is observed, as was reported in16–24. Initial collapse occurs at a load that is about twice that of the average steady load causing progressive crushing. The amplitudes of the little peaks, which sig- nify progressive folding collapse, are higher initially and gradually decrease, as shown in Figure 4. Plastic collapse always occurred at one (usually the top) end and the deformation front gradually progressed with conti- nued crushing until the plastic folding deformation approached the lower end of the specimen. Then the load increased very rapidly, indicating the densification of the specimen. The load-displacement graphics of the alumi- num and Nomex honeycomb sandwich panels for diffe- rent densities, resulting from the experiment, are given in A. BENTOUHAMI, B. KESKES: EXPERIMENTAL ANALYSIS AND MODELING OF THE BUCKLING ... 238 Materiali in tehnologije / Materials and technology 49 (2015) 2, 235–242 Figure 7: Evolution of the critical maximum load with the core density Slika 7: Odvisnost kriti~ne maksimalne obremenitve od gostote jedra Figure 5: Stages of quasi-static compression test of aluminum honey- comb: (1) initial state, (2) buckling initiation, (3) progressive folding and (4) densification Slika 5: Stopnje kvazistati~nega tla~nega preizkusa satja iz aluminija: (1) za~etno stanje, (2) za~etek upogibanja, (3) napredovanje zlaganja, (4) zgo{~evanje Figure 6: Load-displacement curves for honeycomb sandwich panels for different core densities: a) aluminium and b) Nomex Slika 6: Krivulje obremenitev – raztezek za sataste sendvi~ne panelne plo{~e z razli~no gostoto jedra: a) aluminij, b) nomex Figure 4: Load-displacement curve for an aluminum honeycomb sandwich Slika 4: Krivulja obremenitev – raztezek satasto sendvi~ne plo{~e iz aluminija Figure 6. It is clear that the maximum critical buckling load values of the aluminum core panels are higher than those of the Nomex core panels (Figure 7). Also, as the density increased, the compressive strength of both the Nomex and aluminum honeycomb panels increased. The honeycomb compressive behavior intrinsically relates to the cell-wall buckling behavior under com- pression, because in reality the vertical cell walls can never be compressed along the length direction until a pure compressive failure due to the instability of the thin structure occurs. Therefore, the dominant mode of damage in these structures is the buckling of the cells’ walls. Besides the compressive strength, the establishment of the incurred failure modes during the experiment is also important. As the load increased, the initial honey- comb wall buckling and the later regional cell wall fold- ing and core crushing were observed in the aluminum core sandwich panels (Figure 8). The failure modes of the Nomex panels under compression load show a simi- lar behavior to that of the aluminum honeycomb. But for the Nomex core panels, which are more than aluminum, prior to the core crushing failure, crack generation occurred (Figure 9). The failure, which started as a cell- wall buckling, caused cracks under greater compression loads (Figures 8 and 9) than were reported in16,21–24. For the honeycomb core compressed in the axial direction, the localization occurs in the well-defined plastic collapse bands at the interface between the crushed and uncrushed structural regions. Figure 10 shows the evolution of the load for two different cell-wall thicknesses. From the obtained results, we see that the stiffness of the honeycomb sandwich panels increases with the cell-wall thickness. To account for the influence of the thickness of the cell wall, the compression test has been made and the results of the variation of the load for two different wall thicknesses of the honeycomb core for a 6.4 mm cell size, the obtained results are given in the figure. Accord- ing to the results of the experiment, we find that the wall thickness of the cell also has a significant impact on the A. BENTOUHAMI, B. KESKES: EXPERIMENTAL ANALYSIS AND MODELING OF THE BUCKLING ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 235–242 239 Figure 9: Failure modes of Nomex honeycomb panels for a cell size of 3.2 mm Slika 9: Na~in poru{itve sataste plo{~e nomex z velikostjo celice 3,2 mm Figure 10: Effect of the cell-wall thickness on the buckling load for an aluminum honeycomb core Slika 10: Vpliv debeline sten celic na obremenitev pri upogibu za satasto celico iz aluminija Figure 8: Failure mode of aluminum honeycomb sandwich Slika 8: Na~in poru{itve satastega sendvi~a iz aluminija critical buckling load. It is clear that the critical buckling load increased with an increase of the cell-wall thick- ness. 4 NUMERICAL STUDY The ABAQUS package program, based on the finite- element method, is used as the solution method to numerically establish the critical buckling loads and the complex compressive response and crushing of honey- comb sandwich panels after the buckling observed in the experiments. The shell elements are appropriate because the thickness is the smallest planar dimension (Figures 11 and 12). The finite-element mesh uses the fully inte- grated S4R shell element as a regular mesh with nearly square elements, while the number of elements used was selected from convergence studies. The panels, for which the compression test was performed, are three-dimen- sional and the mechanical properties of the utilized materials are given in Table 1. Table 1: Sizes cell and mechanical properties of used core materials for modeling Tabela 1: Velikost celic in mehanske lastnosti uporabljenih materialov za jedra za modeliranje Density (kg/cm3) hc/mm S/mm tc/mm Ec/MPa vc 29 8.8 19.2 0.080 69000 0.3 To study the sensitivity of the critical buckling load with various parameters, such as the cell’s number, the cell’s size and the thickness of the cell wall, we con- A. BENTOUHAMI, B. KESKES: EXPERIMENTAL ANALYSIS AND MODELING OF THE BUCKLING ... 240 Materiali in tehnologije / Materials and technology 49 (2015) 2, 235–242 Figure 12: Finite-element mesh (S4R) model for buckling simulation of honeycomb sandwich panel Slika 12: Model mre`e kon~nih elementov (S4R) za simulacijo upogibanja sataste sendvi~ne plo{~e Figure 11: Finite-element model and boundary conditions of honey- comb sandwich panels Slika 11: Model kon~nih elementov in robni pogoji za satasto send- vi~no plo{~o Figure 15: Buckling shapes for different values of hc Slika 15: Oblike upogibanja pri razli~nih vrednostih hc Figure 14: Variation of the numerical critical buckling load (Fcn) with the core thickness (hc) Slika 14: Spreminjanje numeri~ne kriti~ne obremenitve (Fcn) pri upogibu z debelino jedra (hc) Figure 13: Variation of buckling load with cell-wall thickness for different honeycomb core materials for a cell size 19.2 mm Slika 13: Spreminjanje obremenitve pri upogibanju v odvisnosti od debeline stene celice za razli~ne materiale satastega jedra pri velikosti celice 19,2 mm ducted the following simulation runs on the finite-ele- ment code ABAQUS. To account for the influence of the cell-wall thick- ness, simulations using the finite element code were made and the results of the variation of the numerical critical buckling load and the cell-wall thickness of the honeycomb sandwich panels for 19.2 mm cell sizes obtained are given in Figure 13. According to the results of the simulation, we find that the wall thicknesses of the cell have a significant impact on the critical buckling load and the resistance of the honeycomb sandwich panels. It is established that the critical buckling load increased with an increase of the cell-wall thickness. To explain the effect of the core’s thickness (hc) on the numerical critical buckling (Fcn), the differences applications on the finite-elements code ABAQUS are realized. We can simply understand from Figure 14 that by increasing the height of the core (hc), the value of the numerical critical buckling load (Fcn) decreases. The different values of (hc), (Fcn) and the correspond- ing analytical buckling equations are summarized in Table 2. The buckling modes obtained for each (hc) are shown in Figure 15. Figure 16 represents a comparison between failure modes of one cell obtained from expe- riment and numerical results for aluminum honeycomb core with cell’s size of 19.2 mm. Figure 17 shows the evolution of the critical buckling load with the cell number of the honeycomb core for a 19.2 mm cell size. From the obtained results it is clear that the stiffness of the honeycomb sandwich panels increases with the number of cells. It is seen that the experimental and numerical deformation conditions are relatively coherent. 5 CONCLUSIONS The buckling load of sandwich plates with a honeycomb core subjected to compression damage has been investigated experimentally and numerically. The cell size and the wall thickness, and the materials, are parameters that have to be determined with respect to the usage area of the honeycomb sandwich structures opti- mally. The honeycomb compressive behavior intrinsi- cally relates to the cell-wall buckling behavior under in-plane compression, because in reality the vertical cell walls can never be compressed along the length direction until a pure compressive failure due to the instability of the thin structure occurs. The following are the obtained results of the study: • The critical buckling load of aluminum panels is determined to be higher than that of the Nomex honeycomb panels. • The failure modes of the Nomex honeycomb sand- wich panels under compression load show similar A. BENTOUHAMI, B. KESKES: EXPERIMENTAL ANALYSIS AND MODELING OF THE BUCKLING ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 235–242 241 Figure 17: Variation of buckling load with the cell number for honey- comb core materials for a 19.2 mm cell size Slika 17: Spreminjanje upogibne obremenitve s {tevilom celic za material s satastim jedrom pri velikosti celice 19,2 mm Table 2: Different height of core (hc) with correspondent buckling load (Fcn) and analytical equations of buckling Tabela 2: Razli~ne vi{ine jeder (hc) s pripadajo~o obremenitvijo pri upogibu (Fcn) in analitsko ena~bo upogibanja hc R = hc/a Fcn/N analytical equations ofbuckling 8.8 2 ≥ R 928.8 W(x,y)= A x h y a11 sin sin π π c 20 2 6≤ ≤R 562.67 W(x,y)= A x h y a21 2 sin sin π π c 30 6 12≤ ≤R 507 W(x,y)= A x h y a31 3 sin sin π π c 40 12 2 5≤ ≤R 469 W(x,y)= A x h y a41 4 sin sin π π c Figure 16: a) Experimental and b) numerical deformation of alumi- num honeycomb for a cell size of 19.2 mm Slika 16: a) Eksperimentalna in b) numeri~na deformacija satja iz alu- minija za velikost celice 19,2 mm behavior as that of the aluminum honeycombs. However, for the Nomex core panels, which are much more brittle than aluminum, prior to the core crush- ing failure, crack generation incurred. • As the core’s density increased, the maximum critical buckling load increased, both for the Nomex and the aluminum comb panels. • It was observed that the core’s height is a crucial parameter affecting the ultimate compressive strength of the sandwich panel. • Besides the core’s densities and the core height, the core-wall thickness also has an important effect on the critical buckling load. The critical buckling load increased as the cell-wall thickness increased. 6 REFERENCES 1 R. K. McFarland Jr, Hexagonal cell structures under post-buckling axial load, AIAA Journal, 1 (1963) 6, 1380–1385, doi:10.2514/ 3.1798 2 T. Wierzbicki, Crushing analysis of metal honeycombs, International Journal of lmpact Engineering, 1 (1983) 2, 157–174, doi:10.1016/ 0734-743X(83)90004-0 3 L. J. Gibson, M. F. Ashby, Cellular Solids, Structure and Properties, Pergamon Press, Oxford 1988 4 L. J. Gibson, M. F. Ashby, G. S. Schajer, C. I. Robertson, The mechanics of twodimensional cellular materials, Proc. R. Soc. Lond. A, 382 (1982), 43–59, doi:10.1098/rspa.1982.0087 5 D. Mohr, T. Wierzbicki, Crushing of soft-core sandwich profiles: experiments and analysis, International Journal of Mechanical Sciences, 45 (2003) 2, 253–271, doi:10.1016/S0020-7403(03) 00053-5 6 W. Abramowicz, T. Wierzbicki, Axial crushing of multicorner sheet metal columns, Journal of Applied Mechanics, 56 (1989) 1, 113–120, doi:10.1115/1.3176030 7 S. Li, S. R. Reid, Relationship between the elastic buckling of square tubes and rectangular plates, Journal of Applied Mechanics, 57 (1990) 4, 969–973, doi:10.1115/1.2897669 8 W. Abramowicz, N. Jones, Dynamic axial crushing of square tubes, International Journal of lmpact Engineering, 2 (1984) 2, 179–208, doi:10.1016/0734-743X(84)90005-8 9 W. Abramowicz, N. Jones, Dynamic axial crushing of circular tubes, International Journal of Impact Engineering, 2 (1984) 3, 263–281, doi:10.1016/0734-743X(84)90010-1 10 N. Jones, Structural Impact, Cambridge University Press, Cambridge 1989 11 W. Johnson, S. R. Reid, Metallic energy dissipating systems, and the update, Applied Mechanics Reviews, 31 (1978), 277–288, Updated in: 39 (1986), 315-319 12 C. L. Wu, C. T. Sun, Low velocity impact damage in composite sandwich beams, Composite Structures, 34 (1996) 1, 21-27, doi:10.1016/0263-8223(95)00127-1 13 S. M. Lee, T. K. Tsotsis, Indentation failure behaviour of honeycomb sandwich panels, Composites Science and Technology, 60 (2000) 8, 1147–1159, doi:10.106/S0266-3538(00)00023-3 14 W. Goldsmith, J. L. Sackman, An experimental study of energy absorption in impact on sandwich plates, International Journal of Impact Engineering, 12 (1992) 2, 241–262, doi:10.1016/0734- 743X(92)90447-2 15 M. Jamjian, J. L. Sackman, W. Goldsmith, Response of an infinite plate on a honeycomb foundation to a rigid cylindrical impactor, International Journal of Impact Engineering, 15 (1994) 3, 183–200, doi:10.1016/S 0734-743X(05)80012-0 16 M. O. Kaman, M. Y. Solmaz, K. Turan, Experimental and Numerical analysis of Critical Buckling Load of Honeycomb Sandwich panels, Journal of Composite Materials, 44 (2010) 24, 2819-2831, doi:10.1177/0021998310371541 17 B. Castanié, C. Bouvet, Y. Aminanda, J. J. Barrau, P. Thevenet, Modelling of low-energy/low-velocity impact on Nomex honeycomb sandwich structures with metallic skins, International Journal of Impact Engineering, 35 (2008) 7, 620–634, doi:10.1016/j.ijimpeng. 2007.02.008 18 M. Giglio, A. Manes, A. Gilioli, Investigations on sandwich core properties through an experimental – numerical approach, Compo- sites: Part B: Engineering, 43 (2012) 2, 361–374, doi:10.1016/ j.compositesb.2011.08.016 19 P. H. W. Tsang, P. A. Lagace, Failure Mechanisms of Impact- damaged Sandwich Panels Under Axial Compression, AIAA, (1994) 1396, 745–754, doi:10.2514/6.1994-1396 20 Q. Zhou, R. R. Mayer, Characterization of Aluminum Honeycomb Material Failure in Large Deformation Compression, Shear and Tearing, Journal of Engineering Materials and Technology, 124 (2002) 4, 412-420, doi:101115/1.1491575 21 D. Mohr, M. Doyoyo, Large plastic deformation of metallic honey- comb: orthotropic rate-independent constitutive model, International Journal of Solids and Structures, 41 (2004) 16-17, 4435–4456, doi:10.1016/j.ijsolstr.2004.02.062 22 D. Mohr, M. Doyoyo, Deformation-induced folding systems in thin-walled monolithic hexagonal metallic honeycomb, Int. J. Solids Struct., 41 (2004) 41, 3353–3377, doi:10.1016/j.ijsolstr.2004.01.014 23 S. T. Hong, J. Pan, T. Tyan, P. Prasad, Quasi-static Crush Behaviors of Aluminum Honeycomb Specimens under Compression Dominant Combined Loads, International Journal of Plasticity, 22 (2006) 1, 73–109, doi:10.1016/j.ijplas.2005.02.002 24 A. Wilbert, W. J. Jang, S. Kyriakides, J. F. Floccari, Buckling and progressive crushing of laterally loaded honeycomb, International Journal of Solids and Structures, 48 (2011) 5, 803–816, doi:10.1016/ j.ijsolstr.2010.11.014 25 S. Liang, H. L. Chen, Investigation on the square cell honeycomb structures under axial loading, Composite Structures, 72 (2006) 4, 446–454, doi:10.1016/j.compstruct.2005.01.022 26 H. S. Lee, S. H. Hong, J. R. Lee, Y. K. Kim, Mechanical behavior and failure process during compressive and shear deformation of honeycomb composite at elevated temperatures, Journal of Materials Science, 37 (2002) 6, 1265-1272, doi:10.1023/A:1014344228141 27 X. Fan, I. Verpoest, D. Vandepitte, Finite Element Analysis of Out-of-plane Compressive Properties of Thermoplastic Honeycomb, Journal of Sandwich Structures and Materials, 8 (2006) 5, 437-458, doi:10.1177/1099636206065862 A. BENTOUHAMI, B. KESKES: EXPERIMENTAL ANALYSIS AND MODELING OF THE BUCKLING ... 242 Materiali in tehnologije / Materials and technology 49 (2015) 2, 235–242 ¼. GAJDO[ et al.: FATIGUE BEHAVIOUR OF X70 STEEL IN CRUDE OIL FATIGUE BEHAVIOUR OF X70 STEEL IN CRUDE OIL VEDENJE JEKLA X70 PRI UTRUJANJU V SUROVI NAFTI ¼ubomír Gajdo{1, Martin [perl1, Jaroslav Bystrianský2 1Institute of Theoretical & Applied Mechanics, Academy of Sciences of the Czech Republic, v.v.i., Department of Thin-Walled Structures, Prosecká 25, Prague 9, 190 00, Czech Republic 2Institute of Chemical Technology, Technická 5, Prague 6, 166 28, Czech Republic gajdos@itam.cas.cz Prejem rokopisa – received: 2014-02-25; sprejem za objavo – accepted for publication: 2014-04-10 doi:10.17222/mit.2014.041 The fatigue properties of the X70 steel in crude oil and in the water separated from the crude-oil phase were tested, the stress-cycle-asymmetry ratio being R = 0. As a reference, fatigue tests in air were also carried out. The fatigue specimens used were prepared from the Vohburg an der Donau – Nelahozeves (IKL) crude-oil pipeline after 13 years of exploitation. The tests were aimed at assessing the degree of degradation of the fatigue properties of this steel due to the environmental impacts typical for the crude-oil processing and transport. For added value, fatigue tests were also carried out in an alkaline solution on specimens prepared from a new pipe. The results of the tests showed that the fatigue properties of the steel in crude oil were slightly better than in air, but much poorer in the water separated from the crude-oil phase. Keywords: corrosion fatigue, S-N curve, X70 steel, crude oil, separated water Preizku{eno je bilo vedenje jekla X70 pri utrujanju v surovi nafti in v vodi, izlo~eni iz surove nafte, pri ~emer je bilo razmerje asimetrije napetostnih ciklov R = 0. Za primerjavo so bili izvr{eni tudi preizkusi utrujanja na zraku. Uporabljeni vzorci za utrujanje so bili izrezani iz naftovoda Vohburg an der Donau – Nelahozeves (IKL) za surovo nafto po 13 letih obratovanja. Namen preizkusov je bila ocena stopnje degradacije zaradi vpliva okolja, zna~ilnega za pridobivanje in transport surove nafte. Kot zanimivost so bili izvr{eni tudi preizkusi utrujanja v alkalni raztopini pri vzorcih iz novih cevi. Rezultati preizkusov so pokazali, da je vedenje jekla pri utrujanju v nafti nekaj bolj{e kot na zraku, vendar ob~utno slab{e v vodi, izlo~eni iz surove nafte. Klju~ne besede: korozijsko utrujanje, S-N-krivulja, jeklo X70, surova nafta, izlo~ena voda 1 INTRODUCTION The transportation of crude oil from a production site to a refinery and from the refinery to the oil consumers needs to be very safe because, in the event of an acci- dent, an oil-spillage disaster endangers the natural envi- ronment and causes financial losses. Crude oil is transported by sea in tankers and on land by trucks, trains and pipelines. Tankers are generally subjected to the corrosive action of seawater, and studies of the effect of seawater on the fatigue properties of the steels used in ship structures are therefore very important1. No less important is the transportation of crude oil on land through the pipelines. In comparison with the trucks and trains used for transporting oil, in recent decades the pipelines have proved to be highly economical and very safe. A number of pipelines are currently under construc- tion all over the world. The crude-oil flow rate is not constant as it varies according to the demand for crude oil. The crude-oil pressure increases when the flow volume increases and vice versa. Based on the records obtained, the pressure fluctuation is about 7 bar. On average, there are ten pressure changes in the month. The pressure fluctuation may result in the corrosion fatigue of a pipeline wall and some of its parts become corrosive2. It is therefore useful and has become a matter of interest to examine the fatigue behaviour of the steel in contact with the trans- ported crude oil. Naturally, it would be desirable to conduct the fatigue tests with the frequency of the stress changes of the same order as the frequency of the crude-oil pressure changes encountered in practice (10–4 Hz), but this would require tests of unrealistic duration. It is more convenient to carry out ordinary resonant-fre- quency fatigue tests which, of course, cannot represent the actual corrosion-fatigue resistance of a crude-oil pipeline. Nevertheless, the results can be taken as infor- mative data indicating relative fatigue-resistance values of the pipelines for various crude oils. 2 FATIGUE TESTS The material used in the tests (X70) was taken from a real IKL oil pipeline, with an outside diameter of 711 mm and a wall thickness of 8.5 mm, after the pipeline had been in operation for about 13 years. The steel contained mass fractions (w/%) C 0.18, Mn 1.3, Si 0.23, P 0.009, S 0.021, Ni 0.03, and V 0.01 and its structure was bainitic with the bandings in the longitudinal direc- tion. The static tensile properties in the hoop direction were as follows: the ultimate tensile strength Rm = 610 MPa, the yield stress Rp0.2 = 495 MPa, and the elongation at failure A5 = 31 %. This steel is designated as X70-A. Materiali in tehnologije / Materials and technology 49 (2015) 2, 243–246 243 UDK 620.178.3 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)243(2015) For added value, fatigue tests were also carried out in an alkaline solution. They were performed on another steel sample of the same grade (X70), taken from a new pipe with a diameter of 508 mm and a wall thickness of 10 mm. The static mechanical properties of this steel were as follows: Rm = 625 MPa, Rp0.2 = 500 MPa, and A5 = 24 %. In order to distinguish it from the steel taken from the pipeline, this steel was designated as X70-B. The fatigue tests were performed in zero-to-tension loading with the resonance frequency of 130–133 Hz, using an Instron electromagnetic resonance fatigue ma- chine, model 1603, with an automatic system for the mean-load control and with the maximum force level F = 100 kN. Flat specimens of a varying width and a constant thickness were used in the tests (Figure 1). The reference fatigue tests of X70-A were made in air, while the main fatigue tests were made in liquid environments: in crude oil and in the water separated from the crude-oil phase. The fatigue tests of the X70-B steel were made in a 1 N water alkaline solution of Na2CO3 and NaHCO3 in the ratio of 1 : 1 with pH = 9.325.3 In cooperation with the RCP Prague, a special sealing cell was developed for the fatigue tests in a liquid medium (Figure 2). The number of cycles to fracture Nf was determined as the number of the stress cycles that the resonance machine was able to exert without changing the set-up parameters. The relative sizes of the cracks at the instant when the machine was automatically stopped were approximately 0.25–0.4 of the minimum width of a specimen. The fatigue tests in the water phase separated from the crude oil (Vohburg–Nelahozeves) were, to some extent, specific since this environment is, in fact, a resi- due of the deposit water that is present, in small amounts, in the transmitted crude oil, either in an emul- sified form or a dissolved form (in very small amounts)4. The chemical composition of separated water is shown in Table 1. Table 1: Results of an analysis of the water phase Tabela 1: Rezultati analize vodne faze pH – 7.0 specific conductivity mS m–1 2000 dissolved substances g dm–3 14 neutralization capacity (KNK4.5) mmol dm–3 11  (Ca + Mg) mmol dm–3 22 chlorides ( Cl– ) g dm–3 8.1 sulphates ( SO42– ) mg dm–3 31 oxygen (15–25 °C) mg dm–3 9.5–8.5 Before the actual fatigue tests, the water was left in contact with the air in order to saturate with oxygen (car- bon dioxide). The oxygen amount corresponded to the saturation at the temperatures between 18 °C and 23 °C. 3 RESULTS OF THE FATIGUE TESTS All the results of the fatigue tests are presented in Figure 3. In addition to the curves pertaining to the X70-A steel specimens there is also an S-N curve ob- tained for the specimens from the X70-B steel. The level of the fatigue limit in zero-to-tension loading of the X70-A steel specimens is represented by a horizontal dashed line. It is taken as 0.6 Rm 5 which, in this case, yields a magnitude of 366 MPa. The figure shows that at the upper margin of the fa- mily of all the experimental points (indicating the best fatigue properties) there are triangles corresponding to crude oil and dark diamonds corresponding to the alkaline solution applied to the X70-B steel. In the bottom part of the family of the experimental points there are, with a big downward shift, light diamonds that relate to separated water. Slightly below the triangles ¼. GAJDO[ et al.: FATIGUE BEHAVIOUR OF X70 STEEL IN CRUDE OIL 244 Materiali in tehnologije / Materials and technology 49 (2015) 2, 243–246 Figure 1: Fatigue specimen of varying width Slika 1: Vzorec za utrujanje z razli~no {irino Figure 2: View of a specimen in a sealing cell, in the grips of the fatigue machine Slika 2: Vzorec v zatesnjeni celici, vpet v ~eljusti naprave za utrujanje there are squares representing the results obtained in air. This shows that, in relation to the fatigue properties of the X70-A steel in air, the influence of crude oil appears to be non-aggressive, and it also follows that the effect of the alkaline solution on the X70-B steel appears to be very similar to the effect of crude oil on the X70-A steel. However, separated water has a very negative influence on the corrosion-fatigue properties of X70-A. Some of the results obtained for the X70-A steel are presented in6. If the S-N curve obtained in air is considered as a reference curve, then, according to this diagram, the crude-oil environment does not worsen the fatigue-corro- sion resistance of the steel up to the fatigue limit region, i.e., up to a life of 2 · 106 cycles. The decrease in the maximum stress in a cycle with the number of cycles to fracture is, however, more rapid in crude oil than in air. As indicated by the position of the S-N curve for separated water in the diagram, the presence of separated water in crude oil leads to a considerable deterioration of the fatigue properties of the tested steel. If we express the S-N curves in the following form: max ( ) = A N f b (1) constants A and b take the magnitudes shown in Table 2. Table 2: Constants of the S-N curves for X70-A and X70-B steels Tabela 2: Konstanti krivulje S-N za jekli X70-A in X70-B environment steel A b air X70-A 1147.4 0.0813 crude oil X70-A 1803.1 0.1129 separated water X70-A 1188.3 0.0971 alkaline solution X70-B 2098.3 0.1239 As already specified, by identifying the fatigue limit with the maximum stress in a cycle for a life of Nf = 2 · 106 cycles, we can obtain probable fatigue limits in zero-to-tension loading. For the X70-A steel, these are: f = 353 MPa in air, f = 350 MPa in crude oil, and f = 290 MPa in separated water; and for the X70-B steel f = 348 MPa in the alkaline solution. 4 DISCUSSION It was experimentally proved that crude oil has no negative effect on the fatigue properties of the X70-A steel. This contrasts with the effect of separated water, which exhibited corrosion aggressivity towards this steel. The corrosion aggressivity of separated water occurs due to its saturation with oxygen and a high amount of chlo- rides7; however, the presence of acid carbonates in high concentrations reduces its aggressivity8. Under the con- ditions of a crude-oil pipeline, separated water can con- tain only a limited amount of oxygen, but under the conditions of long-term storage separated water can be- come saturated with air9. This environment can then be considered as the most aggressive possible environ- ment10,11. In other words, due to a low pH value and a high [c(Cl) + c(SO4)]/c(HCO3) ratio, separated water is an environment characterized by the highest chemical corrosion activity. The effect of a crude-oil environment on the fatigue properties of steel is comparable to that of an inert envi- ronment since it manifests itself by preventing the surface from having contact with the oxygen in the air (fatigue tests in air). The composition of the alkaline solution (its alkali- nity and the presence of the HCO3– / CO32– system) leads to a passivation of the steel that, subsequently, results in a lower aggressivity of the solution. 5 CONCLUSIONS 1. A comparison of the fatigue properties of the X70 steel in air and in crude oil has shown that crude oil had no aggressive effect on the steel in the sense of reducing its fatigue characteristics. The S-N curve for crude oil was above the reference S-N curve for air. A similar fatigue behaviour of the steel was exhibited when an alkaline solution was used, i.e., the corresponding S-N curves were very close to each other. However, quite a different situation was observed in the fatigue tests performed in separated water, which had an aggressive effect on the steel. The corresponding S-N curve was below the other curves, the decrease in terms of the cyclic load being 15–20 % with regard to the reference S-N curve. 2. It can be stated that a crude-oil environment im- proves the fatigue properties of the steel up to the region of the fatigue limit, i.e., up to a life of Nf = 2 · 106 cycles. The decrease in the maximum stress in a cycle with the increasing fatigue life is much more rapid in crude oil than in air. By contrast, separated water markedly wor- sens the fatigue properties. Thus, for example, when the specimens are cycled in crude oil with max = 380 MPa, their mean life is Nf = 976500 cycles; however, their life is reduced to Nf = 125700 cycles when they are cycled in separated water, i.e., their life is reduced by a factor of ¼. GAJDO[ et al.: FATIGUE BEHAVIOUR OF X70 STEEL IN CRUDE OIL Materiali in tehnologije / Materials and technology 49 (2015) 2, 243–246 245 Figure 3: Aggregate presentation of the fatigue-test results Slika 3: Skupna predstavitev rezultatov preizkusov utrujanja approximately 7.8. In this case the slope of the decrease in the cyclic stress is roughly the same as in air. 3. If the fatigue limit is considered as the maximum stress in a zero-to-tension cycle (the stress range) corres- ponding to a life of Nf = 2 · 106 cycles, the fatigue limit for crude oil (350 MPa) is practically the same as those for air (353 MPa) and alkaline solution (348 MPa). The smallest magnitude of the fatigue limit (290 MPa) is found for separated water. Acknowledgements This work was supported by RVO: 68378297 and by the projects P105/10/2052 of the Grant Agency of the Czech Republic, and TE 02000162 of the Technological Agency of the Czech Republic. 6 REFERENCES 1 V. Kasemi, A. Haxhiraj, Mater. Tehnol., 42 (2008) 6, 169–170 2 ¼. Gajdo{ et al., Structural Integrity of Pressure Pipelines, Transgas, a.s., Prague 2004, 137–160 3 M. [perl, Vliv korozního po{kození na provozní spolehlivost plyno- vodních potrubí, (The Effect of Corrosion Damage on Operation Reliability of Gas Linepipes), Ph.D. Dissertation, Czech Technical University, Faculty of Transportation Sciences, Prague, 2008 4 Z. A. Foroulis, Werkstoffe und Korrosion (Materials and Corrosion), 33 (1982) 3, 121–131, doi:10.1002/maco.19820330302 5 P. Hopkins, Defect Assessment in Pipelines, APA Course, Prague, Czech Republic, 2001 6 ¼. Gajdo{, M. [perl, J. Bystrianský, J. Pressure Vessel Technol., 137 (2015) 5, 051401, doi:10.1115/1.4029659 7 L. Garverick, Corrosion in the Petrochemical Industry, ASM International, 1994 8 H. H. Uhlig, R. W. Revie, Corrosion and Corrosion Control, 3rd ed., Wiley-Interscience, New York, USA 1985, 415–441 9 DIN 50 930 (1993): Korrosion Metallischer Werkstoffe im Innern von Rohrleitungen, Behältern und Apparaten bei Korrosionsbela- stung durch Wässer, Teile 1 bis 5, (Corrosion of Metallic Materials inside Pipelines, Vessels and Apparatuses Caused by Water, Parts 1–5), 1993 10 R. Martínez-Palou et al., Journal of Petroleum Science and Engineer- ing, 75 (2011) 3–4, 274–282, doi:10.1016/j.petrol.2010.11.020 11 N. H. Abdurahman, Y. M. Rosli, N. H. Azhari, B. A. Hyder, Journal of Petroleum Science and Engineering, 90–91 (2012), 139–144, doi:10.1016/j.petrol.2012.04.025 ¼. GAJDO[ et al.: FATIGUE BEHAVIOUR OF X70 STEEL IN CRUDE OIL 246 Materiali in tehnologije / Materials and technology 49 (2015) 2, 243–246 M. POURANVARI, S. M. MOUSAVIZADEH: USE OF LARSON-MILLER PARAMETER FOR MODELING THE PROGRESS ... USE OF LARSON-MILLER PARAMETER FOR MODELING THE PROGRESS OF ISOTHERMAL SOLIDIFICATION DURING TRANSIENT-LIQUID-PHASE BONDING OF IN718 SUPERALLOY UPORABA LARSON-MILLERJEVEGA PARAMETRA ZA MODELIRANJE IZOTERMNEGA STRJEVANJA PRI SPAJANJU Z VMESNO TEKO^O FAZO SUPERZLITINE IN718 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: 2014-03-10; sprejem za objavo – accepted for publication: 2014-05-09 doi:10.17222/mit.2014.048 The progress of diffusion-induced isothermal solidification, which is a vital issue in transient-liquid-phase bonding, is modeled using the Larson-Miller parameter (LMP). The solidification mechanisms of the liquid phase during the TLP bonding of a wrought IN718 nickel-based superalloy with different bonding times/temperatures and the data were analyzed using the Larson-Miller parameter. Based on the Larson-Miller parameter (LMP), the TLP bonding parameter was defined in the form of LMP = TB[70 + ln(tB)] to explore the effects of the bonding temperature (TB) and the bonding time (tB) on the isothermal solidification progress. It was found that there is a direct linear relationship between the LMP and the isothermal-solidification- zone size. Keywords: transient-liquid-phase bonding, isothermal solidification, superalloy, Larson-Miller parameter Napredovanje z difuzijo spodbujenega izotermnega strjevanja, ki je vitalen del spajanja z vmesno teko~o fazo, je modelirano z uporabo Larson-Millerjevega parametra (LMP). Mehanizem strjevanja teko~e faze med TLP-spajanjem superzlitine IN718 na osnovi niklja pri razli~nih ~asih/temperaturah ter podatki so bili analizirani z uporabo Larson-Millerjevega parametra. Na tej podlagi je bil dolo~en parameter TLP-spajanja v obliki LMP = TB[70 + ln(tB)], da bi ugotovili vpliv temperature spajanja (TB) in ~asa spajanja (tB) na napredovanje izotermnega strjevanja. Ugotovljeno je bilo, da obstaja neposredna linearna odvisnost med LMP in velikostjo podro~ja izotermnega strjevanja. Klju~ne besede: spajanje s prehodno teko~o fazo, izotermno strjevanje, superzlitina, Larson-Millerjev parameter 1 INTRODUCTION Nickel-based superalloys are extensively used in the modern industry due to their excellent high-temperature tensile strength, stress rupture and creep properties, fatigue strength, oxidation and corrosion resistance and microstructural stability at elevated temperatures1–3. Transient-liquid-phase (TLP) bonding is considered as an interesting repairing/joining process for nickel- based superalloys due to its ability to produce near-ideal joints4. In the TLP bonding, by placing a thin filler alloy (i.e., interlayer) of an alloying metal containing a melting-point-depressant (MPD) element between the two pieces of the base metal to be joined and heating the entire assembly, a liquid phase is formed5. The diffusion of the melting-point depressants (MPD) from the molten interlayer into the base metal causes significant compo- sitional changes in the liquid phase and increases the liquidus temperature of the liquid phase. Once the liquidus temperature reaches the bonding temperature, isothermal solidification starts. As a result of the absence of a solute rejection at the solid/liquid interface during the isothermal solidification, the formation of the second phase is basically prevented and a single-phase micro- structure is formed in the ISZ. If a sample is cooled down before the completion of isothermal solidification, the residual liquid will be solidified during the cooling. A non-equilibrium segregation of the MPD elements results in a formation of eutectic-type solidification pro- ducts consisting of hard and brittle intermetallic pha- ses6–9. Therefore, there is a critical bonding time (tIS) beyond which the formation of a eutectic-type micro- structure in the joint centerline is precluded (i.e., the entire liquid phase is solidified isothermally). In their studies of the effects of the stress and tem- perature on the creep-rupture time of gas-turbine alloys, Larson and Miller10 showed that their results could be collapsed onto a common curve using a normalization of the following form: LMP = T[C + Ln(t)]. This norma- lization has been widely used since as a phenomenolo- gical method to relate the effects of the time and temperature on various thermally activated processes (for example11,12). Since the isothermal solidification process is the key stage of the TLP bonding, this paper aims at investigating the application of the Larson-Miller para- meter to modeling the progress of isothermal solidifi- Materiali in tehnologije / Materials and technology 49 (2015) 2, 247–251 247 UDK 621.791/.792:669.24 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)247(2015) cation during the transient-liquid-phase bonding of a superalloy. 2 EXPERIMENTAL PROCEDURE The wrought IN718 nickel-based superalloy (Ni-20.26Fe-18.23Cr-4.85Nb-3.10Mo-0.93Ti-0.61 Al-0.19Si (w/%)) was TLP bonded using a 50 μm thick BNi-2 (Ni-7Cr-4.5Si-3.2B-3Fe). 10 mm × 5 mm × 5 mm coupons were sectioned from the base metal using an electro-discharge machine. Thereafter, in order to remove the oxide layer, the contacting surfaces were ground using 600-grade SiC paper and then ultra- sonically cleaned in an acetone bath. The interlayer was then inserted between the two base-metal coupons. A Cr-Mo steel fixture (Figure 1) was used to fix the coupons in order to hold this sandwich assembly and reduce the metal flow during the TLP operation. Figure 1 shows the thermal history during the bonding of the samples. The bonding was carried out in a furnace at the temperatures of (1273, 1323, 1373 and 1423) K under a vacuum of approximately 1.33 · 10-5 mbar. The bonding time was varied from 1–60 min. The bonded specimens were sectioned perpendicular to the bond and then microstructural observations of the cross-sections of the specimens were carried out using a light microscope and a field-emission scanning electron microscope (FESEM). For the microstructural exami- nations, the specimens were etched using a solution con- taining 10 mL of HNO3, 10 mL of C2H4O2 and 15 mL of HCl. Semi-quantitative chemical analyses of the phases formed in the centerline of the bond region and adjacent to the base metal were conducted with a JEOL 5900 FESEM equipped with an ultra-thin-window Oxford energy-dispersive X-ray spectrometer (EDS). 3 RESULTS AND DISCUSSION 3.1 Microstructure evolution during isothermal solidi- fication Figure 2a depicts a typical microstructure of TLP bonded IN718 using a Ni-7Cr-4.5Si-3.2B-3Fe filler alloy M. POURANVARI, S. M. MOUSAVIZADEH: USE OF LARSON-MILLER PARAMETER FOR MODELING THE PROGRESS ... 248 Materiali in tehnologije / Materials and technology 49 (2015) 2, 247–251 Figure 2: a) Typical overview of TLP bonded wrought IN718 with the partial-isothermal-solidification joint region, b) detailed view of ASZ: the microconstituents marked as P, Q and R are Ni-rich boride, Cr-rich boride and Ni-rich silicide. The zone marked as S is the eutectic gamma solid solution containing fine Ni3Si precipitates. Slika 2: a) Zna~ilen videz spojenih TLP-vzorcev IN718 s podro~jem delnega izotermnega strjevanja, b) detajl videza ASZ: mikrogradniki, ozna~eni s P, Q in R so z Ni bogati borid, s Cr bogati borid in z Ni bogati silicid. Podro~je, ozna~eno s S, je evtekti~na gama-raztopina, ki vsebuje drobne izlo~ke Ni3Si. Figure 1: Temperature-time history during the TLP bonding of the IN718 alloy. The fixture used to hold the samples during the bonding is also superimposed in the figure. Slika 1: Temperaturna in ~asovna odvisnost med TLP-spajanjem zliti- ne IN718. Prikazano je tudi uporabljeno dr`alo vzorcev med spaja- njem. Table 1: Chemical compositions (amount fraction, x/% ) of different metallic constituents for various phases observed in the athermal solidi- fication zone Tabela 1: Kemijska sestava (mno`inski dele`, x/%) razli~nih kovinskih gradnikov v razli~nih fazah, opa`enih v obmo~ju neizotermnega strjevanja Microconstituent/Element Ni Cr Fe Si Mo Nb Al Ti P 84.05 6.53 3.53 0.95 – 3.55 1.38 – Q 13.42 79.32 1.40 0.10 2.63 3.00 0.13 – R 79.10 3.15 2.23 15.05 0.12 0.12 0.2 0.04 S 77.04 7.61 4.64 9.62 0.21 0.52 0.14 0.22 at 1050 °C for 10 min indicating a steep microstructural gradient in the joint area. Three distinct microstructural zones can be distinguished in the affected bonding area, namely: • the isothermal solidification zone (ISZ), • the athermal solidification zone (ASZ), • the diffusion-affected zone (DAZ). The ISZ is composed of a thin layer of the  solid solution formed at the solid/liquid interface towards the joint centerline due to isothermal solidification. As can be seen in Figure 2b, the ASZ is composed of four dis- tinct phases marked with P, Q, R and S. Table 1 shows the EDS-SEM analysis of the phases formed in the ASZ. According to Table 1: • P is a Ni-rich boride intermetallic formed during the cooling of the residual liquid. • Q is a Cr-rich boride intermetallic formed during the cooling of the residual liquid. • R represents Ni-rich the silicide eutectic microcon- stituents formed during the cooling of the residual liquid. • S is the eutectic- formed during the non-isothermal solidification of the residual liquid as part of the eutectic solidification product. This region contains extensive fine Ni3Si precipitates formed during the solid-state precipitation reaction during the cooling, not directly from the solidification of the remaining liquid. According to the morphology of the phases in the ASZ, it can be concluded that they are formed during the eutectic-type reaction during the cooling of the liquid phase. Considering the brittleness of the intermetallic phases in the bonding condition when the isothermal solidification is not accomplished completely, the stress concentration associated with the eutectic microconsti- tuents plays the key role in determining the joint strength. It was shown that the size of the athermal soli- dification zone (ASZ) is the controlling factor for the shear strength of the partially isothermally solidified TLP bonded IN7186. 3.2 Effect of process parameters on isothermal solidi- fication progress The time required to obtain a eutectic-free joint centerline is the key parameter in designing a TLP bonding cycle. Studying the effects of the bonding parameters on the time required to ensure the completion of isothermal solidification can serve as the first step toward the optimization of the process parameters to obtain an isothermally solidified joint free of deleterious intermetallic formations in the joint centerline. There- fore, to investigate the effects of the bonding parameters on the isothermal solidification progress, bonding was carried out at the temperatures of (1273, 1323, 1373 and 1423) K with various durations. Taking the -solid solution/eutectic boundary as the solid/residual liquid interface at the end of each holding time, the average ASZ size was measured. Figure 3a shows the variation in the ASZ size with respect to the bonding time at va- rious bonding temperatures. Figure 3b shows a contour plot of the ASZ versus the bonding time and tem- perature. As can be seen, the ASZ size decreases with the increasing bonding temperature and bonding time due to a larger diffusion of the melting-point-depressant M. POURANVARI, S. M. MOUSAVIZADEH: USE OF LARSON-MILLER PARAMETER FOR MODELING THE PROGRESS ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 247–251 249 Figure 4: Typical eutectic-free joint indicating the completion of iso- thermal solidification Slika 4: Zna~ilen spoj brez evtektika, ki prikazuje dokon~anje izo- termnega strjevanja Figure 3: a) Effect of the bonding time and temperature on the ASZ size, b) contour plot of the ASZ size versus the bonding time and temperature Slika 3: a) Vpliv ~asa in temperature spajanja na velikost ASZ, b) pri- kaz obrisa velikosti ASZ, glede na ~as in temperaturo spajanja (MPD) elements into the base metal. When the bonding time was sufficient for the completion of isothermal solidification, a eutectic-free joint centerline was obtained (Figure 4). According to Figure 3, the time required to ensure the completion of isothermal solidification decreases as the bonding temperature increases. In this section the progress of isothermal solidifica- tion is analyzed using the Larson-Miller normalization. The development of the Larson-Miller normalization starts by expressing the rate for a thermally activated process, r, as an Arrhenius-type equation10. The rate of isothermal solidification is equal to the rate of the solid/liquid movement during the bonding process that can be expressed as follows13: r X t t D C C C t x X t = = − ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ = d d L S ( ) ( ) ( ) ∂ ∂ (1) where D is the diffusivity of the MPD element in the BM, CL and CS are the equilibrium liquidus and solidus, respectively, ( / ) ( )∂ ∂C t x X t= is the gradient of the MPD element at the solid/liquid interface and X(t) is the liquid/solid interface position. Since the diffusion coefficient is related to the tem- perature using an Arrhenius-type equation, the rate of isothermal solidification in a given time can be expressed as an Arrhenius-type equation: r A Q RT = −⎛⎝ ⎜ ⎞ ⎠ ⎟exp (2) where A is the pre-exponential factor, Q is the activation energy, R is the universal gas constant and T is the tem- perature. Rearranging Equation (2) leads to the following equation: [ ]Q R T A r= −ln( ) ln( ) (3) Since the rate r is inversely proportional to time t, Equation (3) can be further written as: [ ]Q R T A t= +ln( ) ln( ) (4) Equation (4) serves as the basis for the Larson-Miller parameter, the LMP: [ ]LMP Q R T C t= = + ln( ) (5) with the Larson-Miller constant defined as C = ln A. The width of the ISZ was measured for each bonding time and temperature and the variation in the ISZ size versus the LMP with various Larson-Miller constants (C) was plotted. A linear trend line was used to fit the data to the LMP. The coefficient of determination (R2) was determined at each LM constant (C). Figure 5a shows the variation in the coefficient of determination (R2) versus the LM constant. According to Figure 5a the best fitting was obtained when C was 70. Figure 5b shows the variation in the ISZ size versus LMP = T[70 + ln(t)] confirming the following linear relationship between the ISZ size and the LMP: Size ISZ/μm = 0.0041 TB[70 + ln(tB)] – 324.07 (6) where TB (K) is the bonding temperature and tB (h) is the bonding time. Therefore, this normalization allowed the progress of isothermal solidification for different bonding times and at different bonding temperatures to be described with a common curve. 4 CONCLUSIONS The understanding of isothermal solidification is of key importance in determining the mechanical perfor- mance of a transient-liquid-phase bonded joint. In this paper, the isothermal solidification during the TLP bond- ing of a wrought IN718 superalloy is studied and the results are analyzed using the Larson-Miller parameter. The following conclusions can be drawn from this study: M. POURANVARI, S. M. MOUSAVIZADEH: USE OF LARSON-MILLER PARAMETER FOR MODELING THE PROGRESS ... 250 Materiali in tehnologije / Materials and technology 49 (2015) 2, 247–251 Figure 5: a) Coefficient of determination for fitting versus the Larson-Miller constant, C, b) ISZ size versus LMP [T(70 + ln (t))] in TLP bonding of IN718/Ni-Cr-Fe-Si-B/IN718 Slika 5: a) Koeficient ujemanja z Larson-Millerjevo konstanto C, b) velikost ISZ v odvisnosti od LMP [T(70 + ln (t))] in pri TLP-spajanju IN718/Ni-Cr-Fe-Si-B/ IN718 1) When the bonding time in not sufficient for the iso- thermal solidification completion, the joint centerline contains a eutectic-type microconstituent that can degrade the mechanical properties of the joint. 2) The isothermal solidification time required to obtain a eutectic-free joint centerline is reduced as the bond- ing temperature increases from 1273 K to 1423 K. 3) Based on the Larson-Miller parameters (LMP), the TLP bonding parameter is defined as TB [70 + Ln(tB)] to explore the effects of the bonding temperature (TB) and the bonding time (tB) on the isothermal solidi- fication progress. It was found that there is a direct linear relationship between the LMP and the isother- mal-solidification-zone size. Acknowledgement The authors would like to thank the Dezful Branch, Islamic Azad University for providing the financial support of this work under the research project entitled: "Investigation on the effect of chemical composition of filler metal on the phase transformations and mechanical properties of diffusion brazed IN718 nickel based superalloy". 5 REFERENCES 1 A. Milosavljevic, S. Petronic, S. Polic-Radovanovic, J. Babic, D. Bajic, Mater. Tehnol., 46 (2012) 4, 411–417 2 R. Sunulahpa{i}, M. Oru~, M. Had`ali}, M. Rimac, Mater. Tehnol., 46 (2012) 3, 263–267 3 A. Semenov, S. Semenov, A. Nazarenko, L. Getsov, Mater. Tehnol., 46 (2012) 3, 197–203 4 W. F. Gale, D. A. Butts, Sci. Technol. Weld. Joining, 9 (2004), 283–300, doi:10.1179/136217104225021724 5 I. Tuah-Poku, M. Dollar, T. B. Massalski, Metall. Trans. A, 9 (1988), 675–686, doi:10.1007/bf02649282 6 M. Pouranvari, A. Ekrami, A. H. Kokabi, Mater. Sci. Eng. A, 568 (2013), 76–82, doi:10.1016/j.msea.2013.01.029 7 D. S. Duvall, W. A. Owczarski, D. F. Paulonis, Weld. J., 53 (1974), 203–214 8 M. Pouranvari, A. Ekrami, A. H. Kokabi, Mater. Tehnol., 47 (2013) 5, 593–599 9 M. Pouranvari, Mater. Tehnol., 48 (2014) 1, 113–118 10 F. R. Larson, J. A. Miller, Trans. ASME, 74 (1952), 765–775 11 S. Venkadesan, A. K. Bhaduri, P. Rodriguez, K. A. Padmanabhan, J. Nucl. Mater., 186 (1992), 177–184, doi:10.1016/0022-3115(92) 90332-f 12 A. M. Limarga, J. Iveland, M. Gentleman, D. M. Lipkin, D. R. Clarke, Acta Mater., 59 (2011), 1162–1167, doi:10.1016/j.actamat. 2010.10.049 13 Y. Zhou, W. F. Gale, T. H. North, Int. Mater. Rev., 40 (1995), 181–196, doi:10.1179/imr.1995.40.5.181 M. POURANVARI, S. M. MOUSAVIZADEH: USE OF LARSON-MILLER PARAMETER FOR MODELING THE PROGRESS ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 247–251 251 A. C. KARAOGLANLI: EFFECT OF SEVERE AIR-BLAST SHOT PEENING ON THE WEAR CHARACTERISTICS ... EFFECT OF SEVERE AIR-BLAST SHOT PEENING ON THE WEAR CHARACTERISTICS OF CP TITANIUM VPLIV INTENZIVNEGA POVR[INSKEGA KOVANJA S PESKANJEM Z ZRAKOM NA OBRABNE LASTNOSTI CP-TITANA Abdullah Cahit Karaoglanli Department of Metallurgical and Materials Engineering, Bartin University, 74100 Bartin, Turkey karaoglanli@bartin.edu.tr, cahitkaraoglanli@gmail.com Prejem rokopisa – received: 2014-03-24; sprejem za objavo – accepted for publication: 2014-06-05 doi:10.17222/mit.2014.055 In this study, air-blast shot peening was applied to analyze the wear characteristics of CP titanium (Grade II). The specimens were exposed to different plastic-deformation rates via different severe shot-peening conditions in order to determine the wear behaviour of CP titanium. A free-ball micro-abrasion test was performed on the specimens shot peened with different Almen intensities. Nanohardness measurements were also performed to investigate the work-hardened layer – the coarse-grained- structure transition zone. Light microscopy and scanning electron microscopy (SEM) were used to analyze both the wear tracks and the severely deformed layer. As a result, the plastically deformed layer thickness reaches approximately 100 μm beneath the surface. Moreover, the hardness and wear durability after severe shot peening is increased. Keywords: air-blast shot peening, wear, nanoindentation, plastic deformation, ultra-fine-grain durability V tej {tudiji je bilo uporabljeno peskanje z zrakom za {tudij lastnosti pri obrabi CP-titana (Grade II). Vzorci so bili izpostavljeni razli~nim stopnjam deformacije v razli~nih razmerah peskanja, da bi ugotovili vedenje CP-titana pri obrabi. Izvr{en je bil abrazijski preizkus z mikrokroglicami na vzorcih po peskanju z razli~no Almen-intenziteto. Izvr{ene so bile meritve nanotrdote za preiskavo utrjenega sloja in grobozrnate strukture prehodne cone. Za analizo sledov obrabe in mo~no deformiranih slojev sta bili uporabljeni svetlobna in vrsti~na elektronska mikroskopija (SEM). Debelina plasti~no deformiranega sloja dose`e globino okrog 100 μm pod povr{ino. Trdota in odpornost proti obrabi se po mo~nem peskanju pove~ata. Klju~ne besede: hladno kovanje povr{ine s peskanjem z zrakom, obraba, nanoodtisek, plasti~na deformacija, odpornost ultra drobnih zrn 1 INTRODUCTION Surface treatments are generally applied to metallic materials, particularly machine parts before service con- ditions to increase service life and efficiency.1 Mechani- cal properties such as wear, fatigue, fretting fatigue and corrosion are influenced by surface treatments.2,3 To analyze these mechanical and physical effects on the sur- face-treated materials, nanoindentation, scratch, hardness and thermal tests are performed.4–9 Surface treatments are investigated within the branch of mechanical and thermal surface treatments. Mechanical surface treat- ments cover a wide variety of processes and shot peen- ing, laser peening, deep drawing, burnishing, sand blast- ing, brush cleaning are given as example processes.10,11 Nitriding, carburising, nitrocarburising, plasma nitriding and boriding processes are beneath the thermal-surface treatments.12 Shot peening has been widely used as a mechanical surface treatment to improve the fatigue resistance of critical machine parts.13 Also, the shot-peening effect on the oxidation, corrosion and fretting-fatigue properties of materials has been studied.13–15 Oxidation and the above mentioned properties are also important for coatings and other surface treat- ments.16,17 With respect to enhancing the mechanical properties without altering the chemical compositions of materials, severe plastic deformations attract a lot of attention.18 Severe plastic deformation, just as ECAP (equal-channel angular pressing), HPT (high-pressure torsion) or ARB (accumulative roll bonding) is performed to increase mechanical properties by decreasing the grain size of the whole bulk materials.18,19 Nevertheless, these methods are applied restrictedly due to the high-pressure require- ments. Shot peening may be more influential if per- formed as a severe plastic deformation.19 It applies a high plastic deformation only to the material surfaces made of a wide variety of materials due to the easiness of the application.13,20 By raising the conditions of the conven- tional shot peening, severe shot peening which applies a very high plastic deformation to a material surface was conceived.20 In recent years, studies depicted that severe shot peening made positive contributions to the increased wear characteristics, creating fine-grained-surface bulk materials.21 Severe shot peening is applied in the ways of high-energy shot peening, ultrasonic shot peening, etc.22–24 Most of the severe-plastic-deformation studies show a highly deformed layer with ultra-fine grains that has superior mechanical properties in comparison with the interior parts.25 From this point of view, obtaining a highly deformed surface with better properties is a way of improving the wear properties of metallic mate- Materiali in tehnologije / Materials and technology 49 (2015) 2, 253–258 253 UDK 669.295:539.538:621.793/.795 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 49(2)253(2015) rials.26–28 Several techniques are used for determining the wear behaviour of materials.29 Micro-abrasion is one of these test techniques due to its simplicity with respect to the abrasive-wear behaviour of materials.30 In this study, attention is focused on the mechanical properties and the wear characteristics of the bulk CP titanium materials with severely deformed surfaces. The specimens are subjected to a free-ball micro-abrasion test subsequent to a severe shot-peening process. The distinction between the severe shot-peened and as-received titanium specimens are performed via mechanical and physical investigations. 2 EXPERIMENTAL STUDIES Commercially pure titanium (Grade II) material with the dimensions of 20 mm × 20 mm × 8 mm was ma- chined and a normalizing heat treatment was performed to release the machining and manufacturing effects. The surfaces of the specimens were ground with 200, 400, 800 and 1200 grade emery papers and then mechanically polished with 6 μm and 1 μm pastes. The specimens were subjected to the air-blast severe shot peening with different Almen intensities (Table 1). In addition, the air pressures were (750, 800 and 870) kPa for 31A, 35A and 9C, respectively. The Almen intensities were selected due to the high-plastic-deformation exposure.20 The intensities were 31A, 35A and 9C. The C Almen strip is used as the highest plastic deformation in comparison to the A and N strips.31 Table 1: Severe-shot-peening conditions for CP-titanium (Grade II) specimens Tabela 1: Razmere pri mo~nem peskanju vzorca CP-titan (Grade II) Specimen No. Almen intensity Shot type Shot size Coverage (%) 1 31A SAE-J2175 S230 200 2 35A SAE-J2175 S230 200 3 9C SAE-J2175 S230 200 The specimens were etched with a 3 % Nital solution following the severe shot-peening process. Afterwards light and SEM (Vega Tescan) microstructure images were obtained on the cross-sections and the peened surfaces. A Schimadzu DUH-W201S ultra-micro-hard- ness tester was used to determine the hardness alteration from the surface to the interior. The experimental load applied was 50 mN and the duration was 10 s. The fixed-ball micro-abrasion test method was used for determining the wear performance of the shot-peened surfaces with different Almen intensities. The diameter, the material and the hardness of the ball were 25.4 mm, AISI 52100 steel and 65 Rc, respectively. The wear volume was calculated using Equation (1):32 V b R ≈ π 4 64 for b << R (1) where V is the volume of the material removed by wear, b is the diameter of the wear crater, and R is the radius of the ball. The fixed-ball micro-abrasion tests were performed on the peened surfaces of the CP-titanium (Grade II) specimens for 2 min, at 120 r/min, under the normal loads of (0.5, 1 and 1.5) N to determine the wear-volume loss on the peened-specimen surfaces. 800-mesh SiC particles were used as the abrasive and distilled water including 25 % SiC particles was used as the abrasive solution. 3 RESULTS AND DISCUSSION The specimens peened with the 31A, 35A and 9C Almen intensities are shown in Figure 1. The images show the cross-sections of the peened specimens. A severe plastic deformation with a severe shot-peening A. C. KARAOGLANLI: EFFECT OF SEVERE AIR-BLAST SHOT PEENING ON THE WEAR CHARACTERISTICS ... 254 Materiali in tehnologije / Materials and technology 49 (2015) 2, 253–258 Figure 1: Light-microscope images of severely shot-peened titanium specimens peened with: a) 31A, b) 35A, c) 9C Almen intensities Slika 1: Mikrostruktura mo~no peskanega vzorca iz titana z razli~no Almen-intenziteto: a) 31A, b) 35A, c) 9C exposure just below the surface is observed. Approxi- mately 100 μm beneath the surface, grain boundaries become dense and lose their homogeneity and visibility. With the increasing Almen intensity, especially at the intensity of 9C, a very dense, severely deformed surface can be noticed; it is separated from the interior structure because fine grains appear due to a high plastic defor- mation. Also, shot tracks and waves on the surface peened with 9C are denser compared with 31A and 35A. In the literature the reported surface roughness also increases with the increasing plastic deformation.20 The severely deformed structure of the surface was investigated using a SEM analysis (Figure 2). In line with the studies made before, the images of the spe- cimens peened with 31A, 35A and 9C show that a high plastic deformation ruined the homogenous micro- structure and created quite a dense and heterogenous ultra-fine-grained structure.22,33–35 Although the magnifi- cations are high, the grain boundaries cannot be seen. Due to a dislocation-density increase and piling up around the grains, the boundaries are invisible just like reported in36,37. Nanohardness measurements were performed to determine the effect of the plastic deformation on the surface and the effect release to the interior. Figure 3 depicts the hardness variation from the surface to the bulk interior. As seen on the figure, after approximately 150 μm a large part of the plastic-deformation impact was released, being very similar to the ones presented in21,38. The hardness increase is the highest on the sur- faces of the specimens. Nevertheless, for the specimen peened with the 9C Almen intensity, the hardness decrease does not occur as abruptly as in the cases of 31A and 35A. As seen from Figure 1, an ultra-fine- grain, highly deformed layer provides a higher and more stabilized hardness down to 50 μm. A. C. KARAOGLANLI: EFFECT OF SEVERE AIR-BLAST SHOT PEENING ON THE WEAR CHARACTERISTICS ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 253–258 255 Figure 4: Volume loss versus Almen intensity for the normal load Slika 4: Izguba volumna v odvisnosti od Almen-intenzitete pri nor- malni obremenitvi Figure 2: SEM images of severely shot-peened titanium specimens peened with: a) 31A, b) 35A, c) 9C Almen intensities Slika 2: SEM-posnetki mo~no peskanega vzorca iz titana z razli~no Almen-intenziteto: a) 31A, b) 35A, c) 9C Figure 3: Hardness variation versus depth according to the Almen intensities Slika 3: Spreminjanje trdote z globino glede na razli~no Almen-inten- ziteto The volume loss is evaluated on the basis of the crater dimensions from the light and SEM images. The mathematical approach taken from the literature was used in this study.21 The volume loss decreases with the increasing Almen intensity. The Almen intensity causes an exposure to a severe plastic deformation, increasing the hardness of the contact surfaces of the materials. The hardness increase induces a reduction in the wear-vo- lume loss. Figure 4 shows the graphs indicating the volume loss of the specimens, subjected to the wear process using the 800-mesh SiC abrasives under the (0.5, 1 and 1.5) N loads and shot peened at different Almen intensities. As seen in Figure 4, the volume loss of the specimens increased in parallel with the increasing load. This is due to the fact that the force applied on the particles, stuck between the subsurface and the ball, increases resulting in a higher shear force with a deeper plunge of the abrasive particles into the specimen. The lowest volume losses are observed on the speci- men shot peened at the 9C Almen intensity, followed by the specimens shot-peened with the 35A and 31A Almen intensities, respectively. This can be primarily attributed to the surface hardness of the specimens. It would be appropriate to correlate the surface hardness of a speci- men with the plastic deformation occurring on the surface area of the specimen depending on the increasing Almen intensity, since the stresses generated on the surface result in an increased dislocation density, hence, an increased hardness in this area of the specimen. Additionally, the density of the compressive stresses generated in this area varies depending on the Almen intensity and has a positive effect on the wear resistance. Figure 5 shows a SEM image of a crater formed as a result of the micro-abrasion wear test made on a speci- men that was shot-peened at the 9C Almen intensity. The wear tracks, obtained as a result of the tests carried out in compliance with the ASTM G77 standard using different Almen values, exhibited a circular geometry as expected. Figures 6a and 6b show the wear-surface images of the specimens subjected to the micro-scale abrasion pro- cess under the loads of 0.5 N and 1.5 N and shot-peened at the 31A Almen intensity. A three-body wear-abrasion mechanism was encountered on the specimens that were subjected to the wear process under 0.5 N. This is due to the rolling abrasion of the abrasive particles on the surface. A two-body abrasion-wear mechanism, thereby, a groove formation was observed on the specimens, subjected to the wear process under 1.5 N. This is a con- sequence of the plunging of the abrasive SiC particles into the ball surface with the increasing load, which results in a cut-off titanium alloy. The wear-surface images of the specimens shot peened at the 35A Almen intensity under the 0.5 N and 1.5 N loads are given in Figures 7a and 7b. Here, the wear mechanisms occurring at the 31A Almen intensity are observed as well; however, the groove depths, obtained under the load of 1.5 N, happened to be lower due to a higher material hardness. The wear-surface images of the specimens shot peened at the 9C Almen intensity are shown in Figure 8. The wear mechanism obtained under the 1.5 N load differs from that of the other two specimens. The plung- ing of the abrasive particles into the material surface was obstructed due to an increased surface hardness and a three-body abrasion mechanism was observed due to the rolling of these particles. A. C. KARAOGLANLI: EFFECT OF SEVERE AIR-BLAST SHOT PEENING ON THE WEAR CHARACTERISTICS ... 256 Materiali in tehnologije / Materials and technology 49 (2015) 2, 253–258 Figure 5: SEM image of a wear crater obtained on the titanium alloy shot peened with the 9C Almen intensity Slika 5: SEM-posnetek kraterja na titanovi zlitini, peskani z 9C Almen-intenziteto Figure 6: SEM images of the worn surfaces obtained with the 31A Almen intensity at the normal loads of: a) 0.5 N and b) 1.5 N Slika 6: SEM-posnetka obrabljene povr{ine, obdelane z Almen-inten- ziteto 31A pri obte`bi: a) 0,5 N in b) 1,5 N 4 CONCLUSIONS In this study, commercially pure titanium (Grade II) specimens were subjected to severe air-blast shot peening with different Almen intensities. Different Almen intensities exposed the specimen surfaces to different plastic-deformation rates. Severe plastic-defor- mation rates caused a microstructural and mechanical behaviour alteration. With the increasing Almen inten- sity a finer and deeper grain structure occurred and this structure become more complex. However, according to the SEM images, although the magnifications were so high, the grain boundaries cannot be seen and the struc- ture has a very dense layer and a plastically deformed surface. The hardness variation is also compatible with the optical and SEM images. With the decreasing Almen intensity the hardness was reduced and after approxi- mately 100 μm a large part of the plastic deformation was released. The wear-volume loss was reduced with the increasing Almen intensity. The highest wear resistance was obtained with the specimens shot peened at the 9C Almen intensity, followed by the 35A and 31A Almen intensities. On the specimens shot peened at the 9C Almen intensity a three-body abrasion-wear mechanism was observed at all normal loads. However, on the specimens shot peened at the 35A and 31A Almen inten- sities, three-body abrasion was observed at the low load (0.5 N) and two-body abrasion was observed at the high load (1.5 N). These results show that the Almen intensity and, thereby, the surface hardness are the key parameters for the control of the wear behaviour of a Ti alloy. Acknowledgments The author gratefully acknowledges the Turkish Air Force Air Supply and Maintenance Command for the shot-peening applications. 5 REFERENCES 1 K. Z. Shepelyakovskiil, Bulk-surface hardening as a method for increasing strength, reliability, and service life of machine parts, Metal Science and Heat Treatment, 37 (1995) 11, 433–440, doi: 10.1007/BF01154216 2 G. Q. Wu, Z. Li, W. Sha, H. H. Li, L. J. Huang, Effect of fretting on fatigue performance of Ti-1023 titanium alloy, Wear, 309 (2014) 1–2, 74–81, doi:10.1016/j.wear.2013.10.010 3 W. Brostow, K. Czechowski, W. Polowski, P. Rusek, D. Tobo³a, I. Wronska, Slide diamond burnishing of tool steels with adhesive coatings and diffusion layers, Materials Research Innovations, 17 (2013) 4, 269–277, doi:10.1179/1433075X12Y.0000000060 4 W. M. Ke, F. C. Zhang, Z. N. Yang, M. Zhang, Micro-characte- rization of macro-sliding wear for steel, Materials Characterization, 82 (2013), 120–129, doi:10.1016/j.matchar.2013.05.009 5 A. Khellouki, J. Rech, H. Zahouani, Micro-scale investigation on belt finishing cutting mechanisms by scratch tests, Wear, 308 (2013), 17–28, doi:10.1016/j.wear.2013.09.016 A. C. KARAOGLANLI: EFFECT OF SEVERE AIR-BLAST SHOT PEENING ON THE WEAR CHARACTERISTICS ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 253–258 257 Figure 8: SEM images of the worn surfaces obtained with the 9C Almen intensity at the normal loads of: a) 0.5 N and b) 1.5 N Slika 8: SEM-posnetka obrabljene povr{ine, obdelane z Almen-inten- ziteto 9C pri obte`bi: a) 0,5 N in b) 1,5 N Figure 7: SEM images of the worn surfaces obtained with the 35A Almen intensity at the normal loads of: a) 0.5 N and b) 1.5 N Slika 7: SEM-posnetka obrabljene povr{ine, obdelane z Almen-inten- ziteto 35A pri obte`bi: a) 0,5 N in b) 1,5 N 6 B. D. Beakea, T. W. Liskiewicz, Comparison of nano-fretting and nano-scratch tests on biomedical materials, Tribology International, 63 (2013), 123–131, doi:10.1016/j.triboint.2012.08.007 7 L. Jiangliang, X. Dangsheng, W. Hongyan, H. Zhongjia, D. Jihui, T. Rajnesh, Tribological Properties of Laser Surface Texturing and Molybdenizing Duplex-Treated Ni-Base Alloy, Tribology Transac- tions, 53 (2010), 195–202, doi:10.1080/10402000903097478 8 A. C. Karaoglanli, G. Erdogan, A. Turk, I. Ozdemir, F. Ustel, Study of the microstructural and oxidation behavior of YSZ and YSZ/ Al2O3 TBCs with HVOF bond coatings, Mater. Tehnol., 46 (2012) 5, 439–444 9 C. Edoardo, G. Luca, P. Barbara, Modelling of the transient thermal field in laser surface treatment test, The International Journal of Advanced Manufacturing Technology, 40 (2009), 307–315, doi: 10.1007/s00170-007-1350-z 10 A. Igor, K. N. Ravi, S. Yuji, W. Lothar, O. R. Robert, On the effect of deep-rolling and laser-peening on the stress-controlled low- and high-cycle fatigue behavior of Ti–6Al–4V at elevated temperatures up to 550 °C, International Journal of Fatigue, 44 (2012), 292–302, doi:10.1016/j.ijfatigue.2012.03.008 11 S. Anand Kumar, R. Sundar, S. Ganesh Sundara Raman, H. Kumar, R. Gnanamoorthy, R. Kaul, K. Ranganathan, S. M. Oak, L. M. Kuk- reja, Fretting Wear Behavior of Laser Peened Ti-6Al-4V, Tribology Transactions, 55 (2012), 615–623, doi:10.1080/10402004.2012. 686087 12 B. Edenhofer, W. Grafen, J. Müller-Ziller, Plasma-carburising a surface heat treatment process for the new century, Surface and Coatings Technology, 142–144 (2001), 225–234, doi:10.1016/ S0257-8972(01)01136-7 13 M. A. S. Torres, H. J. C. Voorwald, An evaluation of shot peening, residual stress and stress relaxation on the fatigue life of AISI 4340 steel, International Journal of Fatigue, 24 (2002), 877–886, doi:10.1016/S0142-1123(01)00205-5 14 V. Azar, B. Hashemi, Y. M. Rezaee, The effect of shot peening on fatigue and corrosion behavior of 316L stainless steel in Ringer’s solution, Surface and Coatings Technology, 204 (2010), 3546–3551, doi:10.1016/j.surfcoat.2010.04.015 15 X. H. Zhang, D. X. Liu, H. B. Tan, X. F. Wang, Effect of TiN/Ti composite coating and shot peening on fretting fatigue behavior of TC17 alloy at 350 °C, Surface and Coatings Technology, 203 (2009), 2315–2321, doi:10.1016/j.surfcoat.2009.02.058 16 A. C. Karaoglanli, E. Altuncu, I. Ozdemir, A. Turk, F. Ustel, Struc- ture and durability evaluation of YSZ + Al2O3 composite TBCs with APS and HVOF bond coats under thermal cycling conditions, Surface and Coatings Technology, 205 (2011) 2, 369–373, doi:10.1016/j.surfcoat.2011.04.081 17 A. C. Karaoglanli, H. Dikici, Y. Kucuk, Effects of Heat Treatment on Adhesion Strength of Thermal Barrier Coating Systems, Engineering Failure Analysis, 32 (2013), 16–22, doi:10.1016/j.engfailanal.2013. 02.029 18 K. Kyungjin, Y. Jonghun, Evolution of the microstructure and me- chanical properties of AZ61 alloy processed by half channel angular extrusion (HCAE), a novel severe plastic deformation process, Materials Science and Engineering A, 578 (2013), 160–166, doi:10.1016/j.msea.2013.04.073 19 E. Kaveh, I. Kazutaka, K. Takanobu, H. Zenji, Equal-Channel Angu- lar Pressing and High-Pressure Torsion of Pure Copper: Evolution of Electrical Conductivity and Hardness with Strain, Materials Transac- tions, 53 (2012) 1, 123–127, doi:10.2320/matertrans.MD201109 20 O. Unal, R. Varol, Almen intensity effect on microstructure and mechanical properties of low carbon steel subjected to severe shot peening, Applied Surface Science, 290 (2014), 40–47, doi:10.1016/ j.apsusc.2013.10.184 21 O. Unal, R. Varol, A. Erdogan, M. S. Gok, Wear behaviour of low carbon steel after severe shot peening, Materials Research Innova- tions, 17 (2013), 519–523, doi:10.1179/1433075X13Y.0000000106 22 S. Bagherifard, P. I. Fernández, R. Ghelichi, M. Guagliano, Fatigue properties of nanocrystallized surfaces obtained by high energy shot peening, Procedia Engineering, 2 (2010), 1683–1690, doi:10.1016/ j.proeng.2010.03.181 23 Y. K. Gao, Improvement of fatigue property in 7050–T7451 alumi- num alloy by laser peening and shot peening, Materials Science and Engineering A, 528 (2011), 3823–3828, doi:10.1016/j.msea.2011. 01.077 24 T. Chaise, J. Li, D. Nélias, R. Kubler, S. Taheri, G. Douchet, V. Robin, P. Gilles, Modelling of multiple impacts for the prediction of distortions and residual stresses induced by ultrasonic shot peening (USP), Journal of Materials Processing Technology, 212 (2012), 2080–2090, doi:10.1016/j.jmatprotec.2012.05.005 25 N. A. Prakash, R. Gnanamoorthy, M. Kamaraj, Surface nanocrystal- lization of aluminium alloy by controlled ball impact technique, Surface and Coatings Technology, 210 (2012), 78–89, doi:10.1016/j.surfcoat.2012.08.069 26 K. S. Anand, R. S. Ganesh Sundara, T. S. N. Sankara Narayanan, R. Gnanamoorthy, Fretting wear behaviour of surface mechanical attrition treated alloy 718, Surface and Coatings Technology, 206 (2012), 4425–4432, doi:10.1016/j.surfcoat.2012.04.085 27 T. Fu, Z. F. Zhou, Y. M. Zhou, X. D. Zhu, Q. F. Zeng, C. P. Wang, K. Y. Li, J. Lu, Mechanical properties of DLC coating sputter deposited on surface nanocrystallized 304 stainless steel, Surface and Coatings Technology, 207 (2012), 555–564, doi:10.1016/j.surfcoat.2012. 07.076 28 M. Mubarak Ali, S. Ganesh Sundara Raman, S. D. Pathak, R. Gna- namoorthy, Influence of plasma nitriding on fretting wear behaviour of Ti–6Al–4V, Tribology International, 43 (2010), 152–160, doi:10.1016/j.triboint.2009.05.020 29 H. Caliskan, A. Erdogan, P. Panjan, M. S. Gök, A. C. Karaoglanli, Micro-abrasion wear testing of multilayer nanocomposite TiAlSiN/ TiSiN/TiAlN hard coatings deposited on the AISI H11 steel, Mater. Tehnol., 47 (2013) 5, 563–568 30 A. C. Karaoglanli, H. Caliskan, M. Gok, A. Erdogan, A. Turk, A comparative study of the microabrasion wear behavior of CoNiCrAlY coatings fabricated by APS, HVOF and CGDS techni- ques, Tribology Transactions, 57 (2014) 1, 11–17, doi:10.1080/ 10402004.2013.820372 31 S. Bagherifard, M. Guagliano, Influence of mesh parameters on FE simulation of severe shot peening (SSP) aimed at generating nano- crystallized surface layer, Procedia Engineering, 10 (2011), 2923–2930, doi:10.1016/j.proeng.2011.04.485 32 D. Sun, J. A. Wharton, R. J. K. Wood, Micro-abrasion mechanisms of cast CoCrMo in simulated body fluids, Wear, 26 (2009), 1845–1855, doi:10.1016/j.wear.2009.03.005 33 S. Bagherifard, M. Guagliano, Fatigue behavior of a low-alloy steel with nanostructured surface obtained by severe shot peening, Engi- neering Fracture Mechanics, 81 (2012), 56–68, doi:10.1016/ j.engfracmech.2011.06.011 34 K. Dai, L. Shaw, Comparison between shot peening and surface nanocrystallization and hardening processes, Materials Science and Engineering A, 463 (2007), 46–53, doi:10.1016/j.msea.2006.07.159 35 G. Liu, J. Lu, K. Lu, Surface nanocrystallization of 316L stainless steel induced by ultrasonic shot peening, Materials Science and Engineering A, 286 (2000), 91–95, doi:10.1016/S0921-5093(00) 00686-9 36 M. A. Terres, N. Laalai, H. Sidhom, Effect of nitriding and shot- peening on the fatigue behavior of 42CrMo4 steel: Experimental analysis and predictive approach, Materials & Design, 35 (2012), 741–748, doi:10.1016/j.matdes.2011.09.055 37 K. Farokhzadeh, J. Qian, A. Edrisy, Effect of SPD surface layer on plasma nitriding of Ti–6Al–4V alloy, Materials Science and Engi- neering A, 589 (2014), 199–208, doi:10.1016/j.msea.2013.09.077 38 G. Li, J. Chen, D. Guan, Friction and wear behaviors of nanocry- stalline surface layer of medium carbon steel, Tribology Interna- tional, 43 (2010), 2216–2221, doi:10.1016/j.triboint.2010.07.004 A. C. KARAOGLANLI: EFFECT OF SEVERE AIR-BLAST SHOT PEENING ON THE WEAR CHARACTERISTICS ... 258 Materiali in tehnologije / Materials and technology 49 (2015) 2, 253–258 E. RANJBARNODEH et al.: FINITE-ELEMENT MINIMIZATION OF THE WELDING DISTORTION ... FINITE-ELEMENT MINIMIZATION OF THE WELDING DISTORTION OF DISSIMILAR JOINTS OF CARBON STEEL AND STAINLESS STEEL UPORABA KON^NIH ELEMENTOV ZA ZMANJ[ANJE POPA^ENJA OBLIKE PRI VARJENJU OGLJIKOVEGA IN NERJAVNEGA JEKLA Eslam Ranjbarnodeh1, Majid Pouranvari2, Mehdi Farajpour3 1Mining and Metallurgical Engineering Department, Amirkabir University of Technology, Tehran, Iran 2Young Researchers Club, Dezful Branch, Islamic Azad University, Dezful, Iran 3Department of Mechanical Engineering, Islamic Azad University, East Tehran Branch, Tehran, Iran islam_ranjbar@yahoo.com, islam_ranjbar@aut.ac.ir Prejem rokopisa – received: 2012-08-07; sprejem za objavo – accepted for publication: 2014-05-27 doi:10.17222/mit.2012.057 In the present study, on the basis of a verified model, the effects of geometrical and operational variables on the welding distortion of dissimilar joints were investigated. Then, considering the welding current, the fixing time and the sequence of the welding operation, the magnitude of the welding deformation was minimized. The results showed that in the studied dissimilar joint between carbon steel and stainless steel, the minimum distortion occurred when the welding current was about 95 A, the fixing time was 120 s and the symmetric layout was used. Keywords: minimization, welding distortion, dissimilar joint, welding variables V tej {tudiji je bil preiskovan s preizku{enim modelom vpliv geometrijskih in procesnih spremenljivk na popa~enje oblike pri varjenju razli~nih materialov. Obseg deformacije pri varjenju je bil zmanj{an z upo{tevanjem varilnega toka, dolo~itvi ~asa in zaporedja varilskih operacij. Rezultati pri {tudiju varjenja ogljikovega in nerjavnega jekla so pokazali, da je najmanj{e popa~enje oblike dose`eno pri varilnem toku okrog 95 A, ~asu 120 s in simetri~ni izvedbi varjenja. Klju~ne besede: zmanj{anje, popa~enje oblike pri varjenju, spajanje razli~nih materialov, spremenljivke pri varjenju 1 INTRODUCTION Dissimilar joints are made of two materials that are considerably different in mechanical and/or chemical senses. Alloying between the base metals and filler metals is of great importance when dissimilar joints are made using fusion-welding processes. The resulting weld metal can show completely different mechanical properties as well as the distribution of the residual stress and the distortion within the weldment compared to the applied base metal(s) and filler metal. In this respect, proper designing of the welding procedure for these welds is of importance to engineers and scientists1. In welding operations, residual stresses and distortions are formed in the welded structures due to severe tempera- ture gradients, as they can cause brittle fractures re- ducing the fatigue life as well as stress corrosion crack- ing. There are many published studies about the effect of welding residual stresses on the performance of welded structures. For example, Francis et al.2 studied the influence of the residual stress in the vicinity of a weld on the structural integrity. They found that the martensite start temperature of the weld filler metal can be adjusted to mitigate the residual-stress distributions in ferritic steel welds. Klob~ar et al.3 developed a finite-element model to predict the deformation and residual stresses and detect the areas critical to cracking during the repair welding of complex-geometry tooling. Francis et al.4 reviewed the metallurgical issues that arise in ferritic steel welds, relating them to the difficul- ties of calculating residual stresses and highlighting some stimulating areas for future research. It should be mentioned that welding residual stresses and distortions are affected by many factors: the thermo-physical and mechanical properties of base metal(s), the geometry of the weldment, the heat input, the welding sequence, etc.5 So far, a few studies about the residual stresses in similar as well as dissimilar arc-welding operations have been published. For instance, Sahin et al.6 used a two-dimen- sional finite-element model to predict the residual stresses in a brazed joint between 1.0402 and brass (BS CZ107). They found higher residual stresses in the steel part of the joint with the higher yield strength. Katsareas et al.7 developed two-dimensional and three-dimensional models to predict the residual-stress distribution in a dissimilar joint between A508 and 1.4301. They used the element-death-and-birth technique to simulate an addi- tion of the filler metal to the weld pool. Ranjbarnodeh et al.8 simulated the heat transfer in dissimilar arc welding of stainless steel to carbon steel. They used the finite- element software ANSYS to investigate the effects of process parameters on the temperature distribution and residual stresses of a dissimilar joint. According to the Materiali in tehnologije / Materials and technology 49 (2015) 2, 259–265 259 UDK 519.61/.64:621.791/.792 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 49(2)259(2015) literature, there are several studies on the effects of welding parameters on the magnitude and distribution of welding residual stresses. Teng et al.9 developed a model to evaluate the effects of welding speed, specimen size, mechanical constraint and preheating on residual stresses. In another study, Teng et al.10 evaluated the residual stresses with various types of welding sequences in the single-pass welding. But, few researches examined the finite-element simulation of the effect of the welding parameters on the residual distortion. Tsirkas et al.11 developed a model to study the effect of the welding speed on the residual stresses and distortion in butt welds of similar butt joints, in which the AH36 steel was used as the base metal. Their results showed that increasing the welding speed greatly affected the welding residual distortion. Considering the published works on dissimilar joints, further studies are necessary to investigate the welding residual distortions of dissimilar welds, particularly under different welding conditions, as well as the weld- ing sequences and the fixing time in a fixture. Moreover, to the best knowledge of the authors, there is no compre- hensive research on the minimization of the welding distortion of dissimilar joints. Therefore, in the present study, a thermo-mechanical model was utilized to evaluate the effects of the welding current, the welding sequence, the fixing time, the similarity and the joint geometry on the residual distortions of dissimilar butt joints of carbon and stainless steels made with the TIG welding process. A finite-element software, ANSYS, was used to solve the governing equations of the heat transfer and elastic-plastic distortion. Finally, consider- ing the fixing time, the welding current and the sequence, the magnitude of the welding distortion was minimized. 2 MATHEMATICAL MODEL AND EXPERIMENTAL PROCEDURES In this work, the finite-element method is employed to solve the heat-conduction problem during and after dissimilar welding. Equation (1) can be employed to describe the temperature variations inside the parts being welded: ∂ ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ + ∂ ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ + ∂ ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ = ∂ ∂ ∂ ∂ ∂ ∂ ∂ ∂ ∂ ∂ ∂x k T x y k T y z k T z C T  t (1) where T denotes the temperature, k is the thermal con- ductivity, C is the specific heat,  is the density, t represents the welding time, z represents the welding direction, x is the transverse direction and y is the thickness direction. In addition, convection-conduction boundary conditions presented in Equation (2) are assumed, except for the region affected by the welding arc where the Gaussian heat source is employed as the boundary condition as illustrated in Equation (3): − = −k T n h T T ∂ ∂ ( )a (2) k T y q r VI r r∂ ∂ π = = − ⎛⎝ ⎜ ⎞ ⎠ ⎟ ⎡ ⎣⎢ ⎤ ⎦⎥ ( ) exp  2 1 22 2 (3) Here, n denotes the normal direction to the surface boundary, Ta is the ambient temperature, q(r) is the welding input energy, r is the distance from the center of the heat source and is the Gaussian distribution parameter, which is assumed as the radius of the area to which 95 % of the energy is entered8, while r is assumed to be 1.5 mm and  = 0.6.8 It should be mentioned that the finite-element soft- ware, ANYSYS, is employed to solve the above heat- conduction problem. Regarding the thermal response of the material being welded, it is expected to produce a severe temperature gradient close to the welding arc and, therefore, very fine elements are required in the region of the weld pool to archive accurate results. Thus, the mesh is generated in such a way that the size of the elements increases exponentially in the transverse direction, i.e., the x-axis. The mesh system used in the model is illustrated in Figure 1. At the same time, the mechanical response of the weldment should be determined by solving the equilibrium problems as shown in Equations (4) and (5): ij j jb, + = 0 (4) ij ji= (5) E. RANJBARNODEH et al.: FINITE-ELEMENT MINIMIZATION OF THE WELDING DISTORTION ... 260 Materiali in tehnologije / Materials and technology 49 (2015) 2, 259–265 Figure 1: a) Model geometry and b) the used finite-element mesh Slika 1: a) Geometrija modela in b) uporabljena mre`a kon~nih ele- mentov where ij is the Cauchy stress tensor and bi is the body force vector, while the thermo-elastic-plastic behavior, based on the Von Mises yield criterion and the isotropic strain-hardening rule, is considered in the model. Accordingly, the constitutive equations can be written as follows:6 [ ] [ ][ ] [ ]d d dep th = −D C T (6) where Dep is the elastic-plastic matrix, Cth is the thermal strain matrix, d is the stress increment, d is the strain increment and dT is the temperature increment. Since the thermal elastic-plastic analysis is a non-linear and path-dependent problem, the incremental calculations, together with the iterative solution techniques, are employed in the model. It is worth noting that, in the finite-element analysis, for the short samples a total of 5670 elements and 7130 nodes were employed, while for the long samples a total of 11340 elements and 14105 nodes were used in the analysis. The temperature-dependent material properties are assumed for both the stainless-steel and low-carbon- steel parts. The convection heat-transfer coefficient was taken as 25 W/(m2 K) for the surfaces in contact. The total duration of the joining process consisted of two main parts in the model. The first part was used to complete the welding stage and the remaining time was used to perform the cooling in and out of the fixture. During and after the welding stage the model was con- strained in order to prevent a rigid body motion. It should be noted that for the mechanical part of the simulation the results of the thermal part were applied as the body force and thermal stresses were calculated at each step, in other words, the thermo-mechanical prob- lem was handled as a sequentially coupled one.8 Using the validated model8, the simulations were repeated for different lengths, thicknesses, welding currents, welding sequences and holding times of the fixture and similar joint. Table 1 shows the chemical compositions of the steels used in the welding experiments. The dimensions of the samples were L × 80 × t (mm3) where L is the sample length and t is the sample thickness, but the filler was not used. The welding parameters used for preparing the experimental samples are shown in Table 2. The welding speed was 3.56 mm/s for all the samples. After the completion of the joints, the maximum magnitude of the welding distortion of all the samples was measured using a height gauge in the Y-direction as shown in Fig- ure 2 and the resulting distortions were compared with the computed results. Table 1: Chemical compositions of the applied base metals in mass fractions, w/% Tabela 1: Kemijska sestava uporabljenih jekel v masnih dele`ih, w/% Material C Si Mn Ni Cr Ti Stainless steel (AISI409) 0.015 0.59 0.27 0.13 11.28 0.17 Carbon steel (CK4) 0.025 0.013 0.19 0.04 0.01 – Table 2: Applied welding parameters Tabela 2: Uporabljeni parametri varjenja Joint type Weldingsequence Welding current (A) Thick- ness (mm) Length (mm) Sample dissimilar progressive 120 2 450 120LPD dissimilar progressive 120 2 225 120SPD dissimilar progressive 150 3 225 150Thick dissimilar progressive 150 2 225 150Thin dissimilar progressive 105 2 450 105LPD dissimilar progressive 135 2 450 135LPD dissimilar progressive 95 2 450 95LPD dissimilar symmetric 95 2 450 95LSD similar progressive 120 2 225 120SS-SS similar progressive 120 2 225 120CS-CS dissimilar back step 120 2 225 120SBD dissimilar symmetric 120 2 225 120SSD 3 RESULTS AND DISCUSSION The comparison of the experimental and simulation results of the distortion measurements are presented in Table 3. As can be seen, there is a logical consistency between the experimental and simulated data. The effect of the sample length on the welding residual distortion was investigated using the samples with two different E. RANJBARNODEH et al.: FINITE-ELEMENT MINIMIZATION OF THE WELDING DISTORTION ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 259–265 261 Figure 2: Measurement setup for the welding-induced distortion Slika 2: Merjenje popa~enja oblike zaradi varjenja Table 3: Comparison between experimental and simulated distortion data Tabela 3: Primerjava med eksperimentalnimi in simuliranimi podatki o popa~enju oblike Simulated distortion (mm) Experimental distortion (mm) Sample 5.47 8.1 70SPD 16.4 22.1 70LPD 17.6 18.5 105LPD 4.43 6.1 120SPD 21.9 24.3 120LPD 18.2 23.1 105LPS 19.4 24.2 120LPS 19.7 25.5 135LPS 10.2 9.8 120LPDS lengths, 120SPD and 120LPD. The distortion of the longer sample, with twice the length, is over five times larger than that of the shorter sample as shown in Figure 3. This can be related to the lower stiffness of the longer sample (120LPD). Stiffness is the resistance of a body to distortion due to an applied force. Contrary to elastic modulus, stiffness is not an intrinsic material property. In other words, it depends on the geometry and material property (elastic modulus). For an element in tension or compression, the axial stiffness is defined with Equation (7): K A E L W t E L = × = × × (7) where L denotes the sample length, A is the cross-sec- tional area, E is the Young’s modulus, W is the width of the specimen and t shows the thickness. The double length of the longer sample causes the half stiffness, i.e., a lower resistance to distortion and, accordingly, a higher welding distortion. The next point is the effect of the thickness on the welding distortion. Figure 4 com- pares the distortions for two different thicknesses. Again, according to Equation (7), a higher thickness means a higher stiffness. Therefore, the thicker structure is more resistant to distortion. Increasing the thickness from 2 mm to 3 mm decreased the maximum magnitude of distortion from 6.7 mm to 0.3 mm. The maximum distortions for different welding currents are compared in Table 4. As expected, a higher welding current induced a higher welding distortion due E. RANJBARNODEH et al.: FINITE-ELEMENT MINIMIZATION OF THE WELDING DISTORTION ... 262 Materiali in tehnologije / Materials and technology 49 (2015) 2, 259–265 Figure 4: Effect of the thickness on the welding distortion: a) thicker sample (3 mm), b) thinner sample (2 mm) Slika 4: Vpliv debeline na popa~enje oblike pri varjenju: a) debelej{i vzorec (3 mm), b) tanj{i vzorec (2 mm) Figure 3: Effect of the length on the welding distortion: a) longer sample, b) shorter sample Slika 3: Vpliv dol`ine na popa~enje oblike pri varjenju: a) dalj{i vzorec, b) kraj{i vzorec Figure 5: Effect of the welding current on the joint penetration Slika 5: Vpliv varilnega toka na penetracijo v spoju to a higher welding-heat input. Therefore, it can be con- cluded that the minimum distortion would be attained for the minimum applicable welding heat input keeping a full penetration. Using the simulation results, the mini- mum welding current to reach a full penetration was found to be about 95 A, as illustrated in Figure 5. Using this welding current, the simulation was repeated and the result showed that the minimum possible distortion for a progressive layout was about 6.8 mm (Figure 6). Table 4: Maximum distortions for different welding currents Tabela 4: Maksimalna popa~enja pri razli~nih tokovih varjenja 135LPD 120LPD 105LPD Sample 135 120 105 Welding current (A) 23.7 21.9 17.6 Distortion (mm) The effect of the welding sequence on the welding distortion was investigated for three different sequences. Figure 7 shows a schematic illustration of the used welding sequences and Figure 8 indicates the effects of the used welding sequences on distortion. As it is de- monstrated, the symmetric layout causes the minimum residual distortion. This can be due to the movement of the welding heat source from the middle of the plate with the minimum degree of freedom to the endpoint with the highest degree of freedom. It should be noted that this result is in agreement with the other researchers.9,12 According to the authors’ knowledge there is no detailed numerical study on the effect of the holding time in a E. RANJBARNODEH et al.: FINITE-ELEMENT MINIMIZATION OF THE WELDING DISTORTION ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 259–265 263 Figure 8: Effect of the welding layout on distortion: a) progressive, b) back-step and c) symmetric welding layout Slika 8: Vpliv postavitve varjenja na popa~enje oblike pri varjenju: a) progresivno, b) s koraki nazaj in c) simetri~na postavitev varjenja Figure 7: Schematic illustration of the used welding sequences Slika 7: Shematski prikaz uporabljenih sekvenc varjenja Figure 6: Distortion of sample 95LPD Slika 6: Popa~enje oblike pri vzorcu 95LPD fixture on the welding distortion. The question is what the effect of the holding time in a fixture on the mag- nitude of distortion is when a structure is welded under mechanical restraint. In order to answer this question the final distortion of sample 95LPD was determined for different holding times in the fixture and afterwards the structure was relaxed due to the removal of the applied mechanical constrains. As depicted in Figure 9, in- creasing the holding time decreases the residual welding distortion. After 120 s, the distortion-holding time curve enters a plateau region. This means that the joint in the fixture was cooled down and that, after this time, the holding time did not have a significant effect on the final magnitude of distortion. This can be due to the high yield strength of the base metal at room temperature, preventing additional distortion. Figure 10 compares the distortions of similar and dissimilar joints. It is seen that a similar joint of stainless steel will cause a higher mag- nitude of distortion in comparison with a dissimilar joint. This may be attributed to the low thermal conductivity, low yield strength and the high thermal-expansion coeffi- cient of stainless steel. Using the previous findings of this study, a simulation was done for different welding currents to reach the minimum welding distortion. The result is depicted in Figure 11. As can be seen, using the minimum welding current (95 A), the symmetric sequen- ce and the 120-second holding time, the final magnitude of the welding distortion can be reduced by only about 0.4 mm which is very small in comparison with 23.7 mm for sample 135LPD. 4 CONCLUSIONS In this work, a verified 3D finite-element model8 was utilized to evaluate the effects of the geometry, the welding current, the welding sequence, the holding time in the fixture and the similarity of the joint on the welding distortion of a dissimilar joint between carbon steel and stainless steel. Also, considering the welding current, the sequence and the fixing time, the magnitude of the welding distortion was minimized. The results showed that: E. RANJBARNODEH et al.: FINITE-ELEMENT MINIMIZATION OF THE WELDING DISTORTION ... 264 Materiali in tehnologije / Materials and technology 49 (2015) 2, 259–265 Figure 11: Distortion of sample 95LSD Slika 11: Sprememba oblike vzorca 95LSD Figure 10: Comparison of the distortions of: a) similar and b) dissimi- lar joints Slika 10: Primerjava popa~enja oblike pri spajanju: a) enakih in b) razli~nih materialov Figure 9: Effect of the holding time in the fixture on distortion Slika 9: Vpliv ~asa zadr`anja v dr`alu na popa~enje oblike • A higher welding current, i.e., a higher heat input produces a larger distortion. • Increasing the length and decreasing the thickness of the welded plates reduce the stiffness of the structure and, consequently, increase the welding distortion. • Due to the low thermal conductivity and yield strength and the high thermal-expansion coefficient, the welding distortion of a similar joint of stainless steel is larger than in the case of a dissimilar joint of carbon steel and stainless steel. • Increasing the holding time in the fixture first reduces the distortion but after 120 seconds, it no longer has a significant effect. • The symmetric welding sequence was found to be an effective way to reduce the welding distortion of a joint compared to the other welding layouts applied in the current study. • The minimum welding distortion in the studied joint can be obtained using the symmetric sequence, the welding current of 95 A and the holding time of 120 s. 5 REFERENCES 1 Welding handbook, vol. 4, AWS, Miami 1997, 514–515 2 J. A. Francis, H. J. Stone, S. Kundu, R. B. Rogge, H. K. D. H. Bha- deshia, P. J. Withers, L. Karlsson, Transformation Temperatures and Welding Residual Stresses in Ferritic Steels, Proc. of ASME Pressure Vessels and Piping Conference, San Antonio, 2007, 949–956, doi: 10.1115/PVP2007-26544 3 D. Klob~ar, J. Tu{ek, B. Taljat, Finite element modeling of GTA weld surfacing applied to hot-work tooling, Comput. Mater. Sci., 31 (2004), 368–378, doi:10.1016/j.commatsci.2004.03.022 4 J. A. Francis, H. K. D. H. Bhadeshia, P. J. Withers, Mater. Sci. Tech- nol., 23 (2007), 1009–1020, doi:10.1179/174328407X213116 5 S. Kou, Welding Metallurgy, John Wiley & Sons, New Jersey 2003, 122 6 S. Sahin, M. Toparli, I. Ozdemir, S. Sasaki, J. Mater. Process. Technol., 132 (2003), 235–241, doi:10.1016/S0924-0136(02) 00932-9 7 D. E. Katsareas, A. G. Yostous, Mater. Sci. Forum, 490–491 (2005), 53–61, doi:10.4028/www.scientific.net/MSF.490-491.53 8 E. Ranjbarnodeh, S. Serajzadeh, A. H. Kokabi, A. Fisher, J. Mater. Sci., l46 (2010), 3225–3232, doi:10.1007/s10853-010-5207-8 9 T. Teng, C. Lin, Int. J. Press. Vessels Pip., 75 (1998), 857–864, doi: 10.1016/S0308-0161(98)00084-2 10 T. Teng, T. Chang, W. Tseng, Comput. Struct., l81 (2003), 273–286, doi:10.1016/S0045-7949(02)00447-9 11 S. A. Tsirkas, P. Papanikos, T. Kermanidis, J. Mater. Process. Tech- nol., 134 (2003), 59–69, doi:10.1016/S0924-0136(02)00921-4 12 L. Gannon, Y. Liu, N. Pegg, M. Smith, Mar. Struct., 23 (2010), 385–404, doi:10.1016/j.marstruc.2010.05.002 E. RANJBARNODEH et al.: FINITE-ELEMENT MINIMIZATION OF THE WELDING DISTORTION ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 259–265 265 M. MADAJ et al.: MAGNESIUM-ALLOY DIE FORGINGS FOR AUTOMOTIVE APPLICATIONS MAGNESIUM-ALLOY DIE FORGINGS FOR AUTOMOTIVE APPLICATIONS IZKOVKI IZ MAGNEZIJEVIH ZLITIN ZA AVTOMOBILSKO INDUSTRIJO Michal Madaj1, Miroslav Greger1, Vlastimil Karas2 1V[B-Technical University of Ostrava, Regional Materials Science and Technology Centre, 17. listopadu 15/2172, 708 33 Ostrava-Poruba, Czech Republic 2KOVOLIT, a.s., Nádra`ní 344, 664 42 Modøice, Czech Republic michal.madaj@vsb.cz, vlastimil.karas@kovolit.cz Prejem rokopisa – received: 2013-09-30; sprejem za objavo – accepted for publication: 2014-06-03 doi:10.17222/mit.2013.174 The paper presents an investigation of the effect of process variables and material condition on the forgeability of magnesium wrought alloys of the Mg-Al-Zn group. The experimental work included the studies of forging capabilities of the alloys in open-die forging at hot- and warm-working temperatures. Forging tests were performed for the material in both the as-cast and as-worked conditions, for two variants of the work-piece geometry. Different variants of the work piece indicated fracture- related problems in forging magnesium alloys in the warm-working temperature mode, which involved an interaction between the material composition and process variables, and the state of stress. By means of numerical calculations it was concluded that, in addition to the material condition, the favourable state of stress, provided by a closed die, could greatly improve the formability of magnesium alloys in the warm-working range. Keywords: forging, magnesium alloys, automotive applications ^lanek predstavlja preiskavo vplivov procesnih spremenljivk in materiala na kovnost magnezijevih zlitin iz skupine Mg-Al-Zn. Eksperimentalno delo je vklju~evalo {tudij sposobnosti za kovanje zlitin v odprtem orodju pri temperaturah vro~ega in toplega preoblikovanja. Preizkusi kovanja so bili izvr{eni za material v litem stanju in za `e predelan material za dve vrsti izkovkov z razli~no geometrijo. Razli~ne variacije izkovkov so pokazale pri kovanju magnezijevih zlitin te`avo z razpokami pri toplem preoblikovanju ter interakcijo med sestavo materiala, procesnimi spremenljivkami in stanjem napetosti. Z numeri~nimi izra~uni je bilo ugotovljeno, da dodatno k razmeram materiala lahko ugodno stanje napetosti, ki se ga dose`e z zaprtim orodjem, mo~no izbolj{a preoblikovalnost magnezijevih zlitin v obmo~ju toplega preoblikovanja. Klju~ne besede: kovanje, magnezijeve zlitine, uporaba v avtomobilski industriji 1 INTRODUCTION The automotive industry, characterised by having the largest potential for development, is becoming an important user of magnesium materials. The use of magnesium in vehicles was, for decades, limited to the castings of complicated shapes for engines and wheels. Traditional die casting dominated for economic reasons. A possibility of using the components from magnesium materials including chassis and drives is now being considered. It turns out that it is suitable to replace the so-far used parts made of steel and aluminium with magnesium alloys. The use of magnesium alloys for the components of the chassis leads to high requirements for their strength, toughness and service life. Most of these properties are achieved by forging. The importance of using the forgings from magnesium alloys in passenger vehicles in comparison with the currently used die-cast castings is continuously increasing1. The use of magne- sium alloys in cars depends on the price relation between aluminium and magnesium alloys. Table 1 compares the current economic possibilities of replacing aluminium alloys with magnesium alloys, as well as the price rela- tions expected in the years to come. Figure 1 shows that the main market for forged com- ponents is the automotive industry. The forging industry Materiali in tehnologije / Materials and technology 49 (2015) 2, 267–273 267 UDK 669.721.5:539.511:621.73 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 49(2)267(2015) Table 1: Price relations between the forgings made of aluminium and magnesium alloys1 Tabela 1: Primerjava cen izkovkov iz aluminija in magnezijevih zlitin1 Price relation aluminium - magnesium Aluminium Magnesium – current price Magnesium – target price /kg /dm3 /kg /dm3 /kg /dm3 Basic metal 2.4 6.5 4.3 7.7 3.6 6.5 Initial blank 0.7 1.9 2.9 to 4.3 5.2 to 7.7 1.4 to 2.1 2.5 to 3.7 Forging and finishing 5–7 14.3–19.8 10–20 18–36 5–10 9–18 Total costs 8–10 23–28 17–29 31–51 10–16 1–28 Comparison with Al alloys 100 % 100 % 210–280 % 140–180 % 120–160 % 80–100 % is thus faced with some particular trends that relate to the developments within this sector2. It is increasingly recognised that aluminium (with a density of 2,700 kg/m3) and magnesium (1,800 kg/m3) are attractive alternatives to steel (7,800 kg/m3). Notably magnesium is the lightest available engineering metal, being by 75 % lighter than steel and by 35 % lighter than aluminium3. To further specify this aspect, Figure 2 gives an overview of the intrinsic weight-saving potential for some magnesium wrought alloys in comparison with the aluminium reference alloy which is in use for forgings. Figure 2 distinguishes between some distinct modes of loading, taking into account the relevant material properties: modulus of elasticity E, yield stress YS and density  (for the other modes of loading, other design parameters apply). Although these data depend some- what on the assumptions and the values of the specific property, this basic approach clearly demonstrates that benefits are anticipated for the strength-related and, in particular, for the bending-relevant parts, with a potential gain for magnesium of even 37 % over aluminium3. The mechanical properties of Mg can be substantially increased by alloying it with aluminium (up to w = 10 %), zinc (up to w = 6 %), manganese (up to w = 2.5 %) and zirconium (up to w = 1.5 %). Aluminium and zinc form a solid solution with magnesium. Intermetallic phases of types Mg17Al12 and MgZn2 are formed when their amounts are higher. In both cases the quantity of the alloying element increases the basic mechanical pro- perties. Manganese and magnesium form a solid solu- tion, -Mg. The solubility of manganese in magnesium decreases with the decreasing temperature and the -Mn phase precipitates from the -Mg solid solution. An addition of manganese does not influence the achieved strength characteristics, but it favourably influences the resistance to corrosion. An increase in the level of the resistance to corrosion can be explained with the fact that a thin layer of Mg-Mn oxides is formed on the surface. An addition of manganese decreases the effect of iron in magnesium. Manganese and iron form a compound of a high density, which settles at the bottom of the bath during the melting. Apart from the basic additions of these elements, an addition of tin is also used in magne- sium alloys. Tin is soluble in magnesium at the tempe- rature of 645 °C up to the amount of approximately w = 10 %. Its solubility decreases with the temperature, with a simultaneous precipitation of the phase (Mg2Sn). Complex Mg-Al-Mn alloys alloyed additionally with w = 5 % of Sn have good hot formability. Silicon is insoluble in magnesium. They form an intermetallic phase of the Mg2Si type, which significantly strengthens the basic matrix4. Due to a significant increase in the brittleness, the amount of silicon in the alloys is under w = 0.3 %. As the alloying of magnesium alloys with zirconium refines the grains, the achieved level of mechanical properties increases and, at the same time, the resistance to corrosion decreases. The elements of rare-earth metals or thorium increase the refractoriness of magnesium alloys. Beryllium, in the amount of w = 0.005–0.012 %, decreases the oxidation of alloys during melting, casting and heat treatment5. One of the limitations of using the formed magne- sium alloys for the production of forged parts is their low formability. For this reason, most of these materials are processed by forming at elevated temperatures, which is reflected in their strength properties. A lower forging temperature increases the precision of the forged parts, but it greatly deteriorates their formability and the mate- rial factors – the grain size, the limited number of slip systems at low temperatures and the resistance of a metal to formation of cracks, which is one of the important fac- tors of formability. Magnesium and the majority of its alloys crystallise in a hexagonal system. This system is characterised by a reduced formability, which is caused by a small number of slip mechanisms. Slips of dislocations take place in selected crystallographic planes and directions, and they are controlled by three known laws. Up to the tempera- ture of 220 °C the only slip plane in magnesium is in the basal plane (0001) and direction [1120]. At higher temperatures the slip begins in plane (1110) in direction [1120]. These are the planes and directions in HCP lattices that are most densely occupied by atoms. The formability increases significantly with an increase in the M. MADAJ et al.: MAGNESIUM-ALLOY DIE FORGINGS FOR AUTOMOTIVE APPLICATIONS 268 Materiali in tehnologije / Materials and technology 49 (2015) 2, 267–273 Figure 1: Customer profile of the European forging industry2 Slika 1: Profil porabnikov evropske industrije izkovkov2 Figure 2: Mass-saving potential of magnesium over aluminium for some typical loading situations3 Slika 2: Mo`nost prihranka mase magnezija v primerjavi z alumini- jem za nekatere zna~ilne primere obremenitve3 slip systems. The values of the critical slip stress (kr) for pure magnesium are low. The value of the critical slip stress depends on the purity of the metal, the structure and thermodynamic conditions of deformation. The higher the purity of the metal, the lower is the magnitude of the critical slip stress. The impurities forming solid solutions with the basic metal increase kr more intensely than the impurities that are insoluble in the basic metal6. If a metal and admixture form a solid solution, then the value of the critical stress increases in dependence of the difference between the magnitudes of atoms of both metals, and the difference between the electrochemical properties of both metals. The admixture elements in magnesium interact with the dislocations, increasing the critical slip stress. The influence of the admixture elements on kr can be determined with the following equation:  kr = c n (1) where c is the concentration of the admixture element and n is the exponent (n 0.5–0.66). The values of the critical slip stress decrease for the majority of metals with the increasing temperature. The influence of the temperature is not unequivocal in the case of magnesium and its alloys. Various slip planes can act at various temperatures. For example, at room tem- perature Mg alloys have only one system of slip planes. The number of active slip planes increases with an increase in the temperature, which is manifested with a rapid decrease in the slip stress. The yield strength of magnesium alloys can be approximately determined with the following equation:  k kr= m (2) where m is the Schmid factor (mmax 0.5). Table 2 presents the basic parameters of the technical procedure of forging magnesium alloys, as well as their mechanical and technological properties. The basic properties of magnesium alloys depend on the achieved structural state, which is a function of the chemical composition, applied deformation and heat treatment. Recrystallisation annealing is performed at the temperature of around 350 °C. The recrystallisation of magnesium alloys strengthened by deformation starts in the temperature interval of 250–280 °C. This tempera- ture interval depends on the degree of strain hardening. Most of the magnesium alloys alloyed with manganese or aluminium are used in the heat-treated conditions, i.e., after quenching and aging. The achieved higher strength is connected with the changed solubility of the admixture elements – Al, Zn and Zr – in dependence of the tem- perature. The heating before quenching is selected in such a way that the segregated intermetallic phases of types MgZn2, Mg17Al12, Mg3Al2Zn2 are dissolved in a solid solution. A homogenous oversaturated solid solu- tion is obtained after the quenching. During aging the strengthening phases precipitate. A characteristic pro- perty of magnesium alloys is a low rate of diffusion pro- cesses and that is why the processes of the phase trans- formation run very slowly. During the heating before quenching the dwell times of 4–24 h are applied. Arti- ficial aging in magnesium alloys runs within the interval from 16–24 h. The selected magnesium alloys can also be quenched by cooling them in air from the finish- forging temperature. The consequential aging performed directly from the finish-forging temperature is used without the inclusion of the previous solution annealing and quenching. The temperatures of the solution annealing of magnesium alloys vary from 380–420 °C. Controlled aging is performed at the temperatures from 200–300 °C. This procedure of heat treatment is marked as T1 and T4. To achieve the maximum level of strength- ening, it is necessary to apply an aging temperature from 175–200 °C. The changes in the properties achieved by aging are smaller for magnesium alloys in comparison with aluminium alloys. An increase in the strength pro- perties after aging is not higher than 20–35 %. However, the plastic properties of alloys decrease after aging. For these reasons the most frequently used heat treatment is the homogenisation annealing. The mechanical proper- ties are enhanced as a result of a more homogeneous structure. An application of natural aging does not prac- tically lead to more significant changes in the strength properties6,7. 2 EXPERIMENTAL WORK We experimentally verified the forging procedure on the piston-rod and plate forgings, the final shapes of which are illustrated in Figure 3. Bars with the diameter of 30 mm and the length of 178 mm were used as initial blanks for the piston-rod forgings. The flat blank for the plate forging had the following dimensions: 130 mm × 150 mm × 13 mm. The forged materials were made of magnesium alloys of types AZ31, AZ61 and AZ91. The M. MADAJ et al.: MAGNESIUM-ALLOY DIE FORGINGS FOR AUTOMOTIVE APPLICATIONS Materiali in tehnologije / Materials and technology 49 (2015) 2, 267–273 269 Table 2: Forging temperatures, mechanical and technological properties of the forgings from magnesium alloys6 Tabela 2: Temperature kovanja, mehanske in tehnolo{ke lastnosti izkovkov iz magnezijevih zlitin6 Alloy Forging temperatures (°C) Mechanical properties Technological properties for forgings for die forgings YS/MPa UTS/MPa Elong. l/% Weldability Resistance tocorrosion AZ31 290–345 260–315 195 260 9.0 O G AZ61 315–370 290–345 180 295 12.0 G G AZ91 300–385 205–290 250 345 5.0 G G Note: O – outstanding, G – good chemical compositions of the forged alloys are presented in Table 3. Forgings were used in the heat-treated as well as in the non-treated states. Before forming, the input blanks were subjected to homogenisation annealing at the tem- peratures of 380–420 °C. The duration of annealing was 15 h. After the forging the samples were subjected to the heat treatment (recrystallisation annealing), which con- sisted of a gradual reheating of the forgings in the fur- nace at a rate of 20 °C per minute up to a temperature of approximately 420 °C. The forgings were left at this temperature for three hours and then cooled in water. Approximately four hours after the completion of the annealing the surfaces of the forgings were blasted with Cr-balls. The forging of the piston-rod and plate forgings was performed in an open die on a hydraulic press MW PA 200. The samples were forged with a single strike at a temperature of 300–350 °C depending on the type of the alloy (Table 4). In the case of the piston-rod forging, the bleed was cut-off in the hot state, immediately after the completion of the forging. The temperature of the die tool was approximately 150–170 °C. The Acheson Dag 554/20 lubricant diluted with water at a ratio of 1 : 20 was used as a lubricating medium. The power of the stamping machine was set to 45 % and its stroke to 220 mm (the lowest possible value on the machine for such forgings). After the forging the test samples for a metallogra- phic analysis of the structure were taken from the forgings. The piston-rod samples were cut along the axis of symmetry in order to examine the change in the struc- ture at individual places of the cut. The sample prepara- tion also included their grinding, polishing and sub- sequent etching. The polishing was performed in two phases. In the first phase the samples were polished on a cloth with a soft nap using a polishing suspension based on Al2O3. However, after the completion of the first phase of the polishing, the sample surfaces still contained a large amount of scratches and it was, therefore, necessary to start the second phase of polishing on a very fine velvet cloth with short hair. Thus polished samples were cleaned with water, rinsed in alcohol and dried by warm air. The prepared surfaces were etched in 4 % HNO3 (Nital) in order to remove the deformed layer for its identification. After the heat treatment, the test specimens for the determination of Brinell hardness were taken from the forgings. The hardness test was performed 10 d after the forging, prior to the heat treatment and also after the heat treatment. The load during the hardness test was 306.5 N and the diameter of the indenting ball was 2.5 mm. Three indents were made on each sample while keep- ing the distance between individual indents in accord- ance with the ISO 6506 standard recommending at least 5 mm in order to avoid the results to be influenced by strain hardening. The samples were subjected to a load for approximately 25 seconds. 3 RESULTS AND DISCUSSIONS The deformation behaviour and development of the structures of six alloys and two shapes of the products were verified experimentally. All the forgings were forged without any problem and with respect to techno- logy no problem occurred during the forging of magne- sium alloys. After the forging the flow stress was assessed quantitatively (Table 5). Table 5: Flow stress of the formed alloys Tabela 5: Napetost te~enja preoblikovanih zlitin Material – shape Average residualenergy (kN) AZ31- piston rod 5166 AZ61- piston rod 4864 AZ91- piston rod 4858 AZ31- plate 4496 AZ61- plate 4369 AZ91- plate 4285 M. MADAJ et al.: MAGNESIUM-ALLOY DIE FORGINGS FOR AUTOMOTIVE APPLICATIONS 270 Materiali in tehnologije / Materials and technology 49 (2015) 2, 267–273 Figure 3: a) Shape of the flat-plate forging and b) shape of the piston-rod forging Slika 3: a) Oblika izkovka plo{~e in b) oblika izkovka ojnice Table 3: Chemical composition of magnesium alloys for forgings Tabela 3: Kemijska sestava magnezijevih zlitin za kovanje Alloy Amounts of alloying elements in mass fractions, w/% Al Zn Mn Si Cu Fe Ni AZ31 2.50–3.50 0.20– 0.80  0.200  0.100  0.05  0.005  0.005 AZ61 6.76 0.38 0.13 0.05 0.006 0.011 – AZ91 8.76 0.73 0.22 0.05 0.010 0.011 – Table 4: Initial parameters for forging the piston rods Tabela 4: Za~etni parametri pri kovanju ojnice Alloy Tempe- rature (°C) Mass of the blank (g) Dimensions and shape of forgings were satisfactory (%) Flow stress AZ31 350 30 100 low AZ61 320 30 97 medium AZ91 300 30 95 high During the forging after a strike the energy was dis- charged onto the contact surfaces – it means that the forgings with the lowest residual energy had the highest flow stress. It follows from the obtained results that among the magnesium alloys the AZ91 alloy has the highest flow stress, followed progressively by AZ61 and AZ31. Some of the piston-rod forgings had a scratch on the larger diameter, which might have been related to the crack formation and such parts were investigated. The shapes of the forgings are shown in Figure 4. For the preparation of the flat forging it was abso- lutely necessary to adapt the contact surfaces, which had to be parallel and smooth since any possible deep scratch could cause a formation of cracks. The shapes of the forgings are shown in Figure 5. The metallographic investigation of the samples was performed in the initial state, after the heat treatment and, finally, after the forging and heat treatment. The analysis of the microstructure was focused on the forma- tion of the cracks that were highlighted during the tech- nological production of the forgings and the individual phenomena associated with the forming of magnesium alloys (dynamic recrystallisation, twinning, growth and shape of grains). In the initial state, the microstructures of magnesium alloys AZ31, AZ61 and AZ91 contained the majority phase (a solid solution of aluminium in magnesium) and two types of the minority phase. The first type of the minority phase consisted of relatively massive particles of Mg17Al12, while the second type consisted of fine, needle-like particles of the same phase present in the vicinity of the grain boundaries (Figure 6). The objective of the homogenisation annealing was to remove the segregation heterogeneities of the admix- ture elements. During the homogenisation annealing the M. MADAJ et al.: MAGNESIUM-ALLOY DIE FORGINGS FOR AUTOMOTIVE APPLICATIONS Materiali in tehnologije / Materials and technology 49 (2015) 2, 267–273 271 Figure 6: Microstructure of alloy AZ91 in the as-cast state, a cross- section Slika 6: Strjevalna mikrostruktura zlitine AZ91; pre~ni prerez Figure 4: Shapes of the forgings from the magnesium alloys: a) AZ31, b) AZ61, c) AZ91 Slika 4: Videz izkovkov iz magnezijevih zlitin: a) AZ31, b) AZ61, c) AZ91 Figure 7: Microstructure of the AZ91 alloy after the homogenisation annealing, a cross-section Slika 7: Mikrostruktura zlitine AZ91 po homogenizacijskem `arjenju; pre~ni prerez Figure 5: Shapes of the forgings from the magnesium alloys: a) AZ31, b) AZ61 Slika 5: Videz izkovkov iz magnezijevih zlitin: a) AZ31, b) AZ61 segregated phases on the grain boundaries dissolved in the basic matrix and the chemical composition of the alloy was more homogenous (Figure 7). This improved the formability and enhanced the level of mechanical properties. The structures of the forged piston rods made of alloys AZ31 and AZ61 did not show any defects that could cause a subsequent failure of the components. The grains of both alloys were stretched in the direction of the intensive flow of the material, i.e., in the longitudinal direction of the forging. In alloy AZ61 a dynamic recry- stallisation took place, which started at the boundaries of the original grains, being supported by a sufficiently large number of precipitates of Mg17Al12 and by the created twins. The recrystallisation gradually expanded toward the centre of the basic grains. The piston rod made of alloy AZ91 had its central part without any sig- nificant failures; however, its peripheral parts were characterised by the cracks that penetrated deeper into the sample. In this alloy a dynamic recrystallisation on the grain boundaries took place, but there was insuffi- cient time for its propagation into the entire volume (Figure 8). With respect to the flat blank for plate forging, the best forging over the entire cross-section was achieved with the AZ31 alloy that showed no cracks. The AZ61 alloy contained no cracks in the central part of the com- ponent, but under the surface some cavities were formed, which could cause a failure of the given component during the subsequent use. The most damaged microstructure was found for alloy AZ91, where cracks were present right under the surface in all the areas, penetrating deeper into the com- ponent and, in some cases, they appeared on the surface as well. It was evident from the metallographic investi- gation that the cracks were preferentially formed on the grain boundaries, particularly in the presence of phase Mg17Al12 or in the places where this phase was dissolved. The grains of all the forged alloys were considerably stretched in the direction of the bleed groove and in alloy AZ91 they led to a formation of a crack over the entire component. After the annealing, a complete recrystallisation of the grains occurred in the sample made of alloy AZ91. The grains contained acicular particles, spread over the entire grains. In alloy AZ61, the recrystallised grains were present only on the borders of the original grains, but they did not spread throughout the entire volume. The chemical composition, i.e., the amount of alumi- nium, had a great influence on these processes. The secondary phases, and zinc- and aluminium-based M. MADAJ et al.: MAGNESIUM-ALLOY DIE FORGINGS FOR AUTOMOTIVE APPLICATIONS 272 Materiali in tehnologije / Materials and technology 49 (2015) 2, 267–273 Figure 9: Hardness values of the piston-rod forgings in the initial state, after the forming and after the heat treatment Slika 9: Trdota izkovka ojnice v za~etnem stanju, po kovanju in po toplotni obdelavi Figure 8: Microstructures of the alloys: a) AZ31, b) AZ61 and c) AZ91 after the forging, a cross-section Slika 8: Mikrostrukture zlitin: a) AZ31, b) AZ61 in c) AZ91, po kovanju; pre~ni prerez precipitates dissolved during the heat treatment in the basic matrix. The mechanical properties of the forgings and their evolution in dependence of the heat treatment were verified with a hardness test. For all the investigated magnesium alloys in the initial state, the Brinell hardness values were measured after the forming and after the heat treatment, and they were averaged for individual alloys. The resulting hardness values are represented in Figures 9 and 10. It follows from the results of the hardness tests that the heat treatment and forming had considerable influences on the hardness of magnesium alloys. The aluminium amount in the material is another factor, influencing the hardness. The material hardness de- creases with the decreasing aluminium amount in the material (Table 3, Figures 9 and 10). During the forming the hardness of alloy AZ91 increased consider- ably – even by 19 HB. The hardness of alloy AZ61 increased by 10 HB. The weakest influence of the form- ing on the hardness value was found for alloy AZ31. After the heat treatment the hardness of alloy AZ91 dropped considerably, which was caused by the recrystallisation, taking place during the heat treatment. The heat treatment had no significant impact on the resulting hardness of alloys AZ31 and AZ61. 4 CONCLUSIONS The deformation behaviour of alloys AZ31, AZ61 and AZ91 during die forging was experimentally veri- fied. The influences of the forging technology and homo- genisation annealing on the structures and properties of the forgings were compared. The influences of the heat treatment and the forming temperature on the final structure and mechanical properties were evaluated. It follows from the obtained results that the aluminium amount in the material as well as the heat treatment and forming had considerable influences on the hardness of the magnesium alloys. The initial structures used in the tests were in the as-cast form, which probably caused the formation of cracks in the alloys with higher aluminium amounts (AZ61, AZ91). The results confirmed the suitability of applying heat treatment before forging. This procedure enabled us to obtain the forgings with a more homogeneous structure. After the application of forming some micro-cracks and voids were detected in the magnesium alloys of AZ61 and AZ91. The cracks were located just under the surface and they penetrated deeper into the material. In alloy AZ91 the micro-cracks were formed throughout the entire volume and the initiations of these micro-cracks were preferentially on the grain boundaries, mainly in the area of particles Mg17Al12. From the structural point of view, alloy AZ31, in which no structural defects were detected, was found satisfactory. The highest strength and hardness were obtained with alloy AZ91. It follows from the obtained results that with the increasing alumi- nium amount the hardness of forged magnesium alloys increases as well. Acknowledgement This paper was created within project No. CZ.1.05/ 2.1.00/01.0040 "Regional Materials Science and Techno- logy Centre" within the frame of operational programme "Research and Development for Innovations" financed by the Structural Funds and the state budget of the Czech Republic. 5 REFERENCES 1 R. Matsumoto, K. Osakada, Development of warm forging method for magnesium alloy, Materials Transactions, 45 (2004) 9, 2838–2844, doi:10.2320/matertrans.45.2838 2 R. Asakawa, K. Hirukawa, Technology for manufacturing magne- sium alloy components with excellent heat resistance, Kobelco technology review, 31 (2013), 76–81 3 B. P³onka, M. L. Grega, K. Remsak et al., Die forging of high- strength magnesium alloys – the structure and mechanical properties in different heat treatment conditions, Archives of Metallurgy and Materials, 58 (2013) 1, 127–132, doi:10. 2478/v10172-012-0162-9 4 M. Greger, L. ^í`ek, I. Juøi~ka et al., Possibilities of mechanical pro- perties and microstructure improvement of magnesium alloys, Archives of Materials Science and Engineering, 28 (2007) 2, 83–90 5 M. Greger, M. Widomská, V. Karas, Properties of forging from mag- nesium alloys and their use in industry, Conference proceedings, Metal 2012, Ostrava, 2012, 440–445 6 M. Greger, M. Widomská, Structural characteristics of magnesium alloys along the equal channel angular pressing, Advances in Engi- neering Plasticity and its Applications, Shanghai Jiaong University, Shanghai, 2004, 1083–1088, doi:10.4028/www.scientific.net/KEM. 274-276.1083 7 D. Kobold, T. Pepelnjak, G. Gantar et al., Analysis of deformation characteristics of magnesium AZ80 wrought alloy under hot condi- tions, Journal of Mechanical Engineering, 56 (2010) 12, 823–832 M. MADAJ et al.: MAGNESIUM-ALLOY DIE FORGINGS FOR AUTOMOTIVE APPLICATIONS Materiali in tehnologije / Materials and technology 49 (2015) 2, 267–273 273 Figure 10: Hardness values of the plate forgings in the initial state, after the forming and after the heat treatment Slika 10: Trdota izkovka plo{~e v za~etnem stanju, po kovanju in po toplotni obdelavi J. PRZONDZIONO et al.: RESISTANCE TO ELECTROCHEMICAL CORROSION OF THE EXTRUDED MAGNESIUM ... RESISTANCE TO ELECTROCHEMICAL CORROSION OF THE EXTRUDED MAGNESIUM ALLOY AZ80 IN NaCl SOLUTIONS ODPORNOST EKSTRUDIRANE MAGNEZIJEVE ZLITINE AZ80 PROTI ELEKTROKEMIJSKI KOROZIJI V RAZTOPINI NaCl Joanna Przondziono1, Eugeniusz Hadasik1, Witold Walke2, Janusz Szala1, Joanna Michalska1, Jakub Wieczorek1 1Silesian University of Technology, Faculty of Materials Engineering and Metallurgy, Krasiñskiego 8, 40-019 Katowice, Poland 2Silesian University of Technology, Faculty of Biomedical Engineering, Ch. de Gaulle’a 66, 41-800 Zabrze, Poland joanna.przondziono@polsl.pl Prejem rokopisa – received: 2013-10-01; sprejem za objavo – accepted for publication: 2014-04-03 doi:10.17222/mit.2013.208 The purpose of this study was to evaluate the electrochemical corrosion resistance of the extruded magnesium alloy AZ80 in NaCl solutions. The resistance to electrochemical corrosion was evaluated on the grounds of registered anodic polarisation curves. Potentiodynamic tests were performed in solution with a concentration of 0.01–2.00 M NaCl. In addition, immersion tests were performed, and they allowed us to determine the corrosion rate. Scanning electron microscopy was applied to observe the microstructure after the immersion tests (after removing the corrosion products). Phenomena that happen on the surface of the alloy were evaluated with the application of electrochemical impedance spectroscopy. The tests enabled us to determine the impedance spectra of the system and the data obtained during the measurement was matched to the equivalent system. An optical profilometer was used for the measurement of the geometrical features of the surface of the alloy. The results of the performed tests prove explicitly the deterioration of the corrosion characteristics of the alloy with an increase in the molar concentration of the NaCl solution. A decrease of the corrosion potential and the polarisation resistance was observed, as well as an increase of the corrosion current density. It was proved that irrespective of the concentration, pitting corrosion can be found on the surface of the alloy. The potential to use the extruded magnesium alloy AZ80 in the aircraft and automotive industries is connected with the necessity to apply protective layers on elements made from the tested alloy. Keywords: extruded magnesium alloy AZ80, electrochemical corrosion, potentiodynamic and immersion tests, SEM, EIS Namen te {tudije je bil oceniti odpornost ekstrudirane magnezijeve zlitine AZ80 proti elektrokemijski koroziji v raztopini NaCl. Odpornost proti elektrokemijski koroziji je bila ugotovljena na podlagi anodnih polarizacijskih krivulj. Izvr{eni so bili poten- ciodinamski preizkusi v raztopini s koncentracijo od 0,01–2 M NaCl. Dodatno so bili izvr{eni tudi preizkusi s potapljanjem, ki so omogo~ili dolo~itev hitrosti korozije. Vrsti~na elektronska mikroskopija je bila uporabljena za slikanje mikrostrukture po potapljanju (po odstranitvi korozijskih produktov). Pojavi na povr{ini zlitine so bili ocenjeni z elektrokemijsko impedan~no spektroskopijo. Preizkusi so omogo~ili dolo~itev impedan~nega spektra sistema in z meritvami dobljeni podatki so se ujemali z ekvivalentnim sistemom. Opti~ni profilometer je bil uporabljen za merjenje geometrijskih pojavov na povr{ini zlitine. Rezultati izvr{enih preizkusov so potrdili poslab{anje korozijskih lastnosti zlitine pri pove~anju molske koncentracije raztopine NaCl. Opa`eno je bilo zmanj{anje korozijskega potenciala in polarizacijske odpornosti, kot tudi pove~anje gostote korozijskega toka. Dokazano je, da se ne glede na koncentracijo jami~asta korozija pojavi na povr{ini zlitine. Mo`nosti uporabe ekstrudirane magnezijeve zlitine AZ80 v letalstvu in avtomobilski industriji je povezana z nujnostjo uporabe za{~itnih plasti na komponentah iz preizku{ane zlitine. Klju~ne besede: ekstrudirana magnezijeva zlitina AZ80, elektrokemijska korozija, potenciodinamski preizkus in preizkus s potapljanjem, SEM, EIS 1 INTRODUCTION The application of magnesium alloys that can be subject to plastic strain is far less popular than for alloys obtained through casting. This results from a number of technological difficulties during plastic working (which are connected with their low formability at ambient temperature), as well as from high production costs. One of the methods of magnesium-alloy forming is extrusion, which is usually realised within the temperature range 320–450 °C at a rate of 1–25 m/min. Hot isostatic pressing (HIP) has been intensively developing for the last couple of years. Thanks to favourable thermal and mechanical conditions, hot isostatic pressing can be exe- cuted at lower temperatures and a larger grain size reduc- tion for the magnesium alloys can be obtained. Equal channel angular pressing (ECAP) is also increasing in popularity1–5. The main problem when using magnesium alloys in the aircraft and automotive industries is their suscepti- bility to electrochemical corrosion. Magnesium as a highly electronegative element that features extreme susceptibility to passing into electrolyte solutions. The standard electrochemical potential of magnesium Eo is –2.37 V, whereas in a 3 % solution of sodium chloride it is –1.63 V (SCE). Magnesium alloys feature good corro- sion resistance in weather conditions and when they are put to the reaction of alkaline, chromate and water- fluoric solutions of acids as well as to the majority of organic chemical compounds, e.g., hydrocarbons, alde- hydes, alcohols (with the exception of methanol), phenols, amines, esters and most oils. Magnesium is not Materiali in tehnologije / Materials and technology 49 (2015) 2, 275–280 275 UDK 669.721.5:620.193 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 49(2)275(2015) resistant to the influence of water containing trace elements of heavy-metals ions, sea water, inorganic and organic acids and acid salts (e.g., ammonium), anhy- drous methanol, gasoline containing lead (and its com- pounds), and freon containing water6–12. It is extremely prone to electrochemical and chemical corrosion, in particular in an environment that contains chloride ions, which substantially limits the area of this alloy’s application. The reason for the low corrosion resi- stance of magnesium is the insufficient protective proper- ties of the layer of oxides that is formed on the surface in an oxidising atmosphere or the layer of hydroxides in water solutions. Electrochemical corrosion is most often displayed by metal defects on the surface (spots and pits) or by the deterioration of the material’s strength13–19. The purpose of this study was to evaluate the resi- stance to electrochemical corrosion of the magnesium alloy AZ80 after extrusion. Corrosion tests were made in NaCl solutions featuring a concentration of chloride ions within the range of 0.01–2.00 M NaCl. Potentiodynamic tests allowed us to register anodic polarisation curves. Immersion tests in NaCl solutions were performed in the time period 1–5 d. A scanning electron microscope served to make images of the AZ80 alloy surface after the corrosion tests. Phenomena that happen on the sur- face of the alloy were evaluated with the application of electrochemical impedance spectroscopy. The surface morphology after the corrosion tests was evaluated by means of a surface analyser. 2 EXPERIMENTAL Samples of AZ80 alloy after extrusion served as stock material for the tests. The chemical composition and the mechanical properties of the alloy are presented in Tables 1 and 2. Table 1: Chemical composition of magnesium alloy AZ80 in mass fractions, w/% Tabela 1: Kemijska sestava magnezijeve zlitine AZ80 v masnih dele`ih, w/% Al Zn Si Mn Cu Fe Mg 8.2 0.34 0.02 0.13 <0.03 0.005 bal. Table 2: Mechanical characteristics of magnesium alloy AZ80 after extrusion Tabela 2: Mehanske lastnosti magnezijeve zlitine AZ80 po ekstruziji Rm/MPa Rp0.2/MPa A/% 343 258 13.5 The resistance to electrochemical corrosion was evaluated on the grounds of registered anodic polarisa- tion curves with the application of the VoltaLabPGP201 testing system by Radiometer. A saturated calomel elec- trode (NEK) of the KP-113 type served as a reference electrode. A platinum electrode of the PtP-201 type served as an auxiliary electrode. The tests were per- formed in solutions featuring various molar concentra- tions of NaCl solution (0.01–2.00 M NaCl). The tempe- rature of the solutions during the test was (21 ± 1) °C. The immersion tests were performed at ambient temperature with the application of the immersion method in 0.01–2.00 M NaCl solution for 1–5 d. After grinding of the surface of the samples, they were weighed and the mass m0 was determined. After immersion of the alloy in the NaCl solution for 1–5 d, samples were taken out and the corrosion products were removed in the reagent containing 200 g/L CrO3 and 10 g/L AgNO3. Next, they were washed with distilled water, degreased with acetone, dried and weighed again with a determination of the mass m1. The performed tests enabled us to determine the corrosion rate. Potentiody- namic tests and immersion tests were performed for three samples of AZ80 alloy for each concentration of NaCl solution. A scanning electron microscope with field emission FE SEM S-4200 Hitachi in cooperation with a spectro- meter EDS Voyager 3500 Noran Instruments was used to make qualitative and quantitative analyses of the chemical composition in micro-areas. The analysis of the morphology of the AZ80 surface was presented in diagrams and profilographs made with the application of an optical surface analyser MicroProf by FRT. In order to obtain information about the physical and chemical properties of the surface of the samples made from the AZ61 alloy, tests with the application of elec- trochemical impedance spectroscopy (EIS) were per- formed. The measurements were made with the applica- tion of an AutoLab PGSTAT 302N measuring system equipped with a FRA2 (Frequency Response Analyser) module. The tests of electrochemical impedance spectro- scopy are a linear measurement of the electrical response of the tested metallic material to stimulation with an electromagnetic signal over a wide range of frequencies. The performed tests enabled a direct comparison of the real behaviour of the object with its equivalent system, which is a model that relates to the physically realized impedance. 3 RESULTS AND DISCUSSION The potentiodynamic tests results are presented in Table 3. The anodic polarisation curves are shown in Figure 1. Table 3: Potentiodynamic tests results of AZ80 alloy – mean values Tabela 3: Rezultati potenciodinami~nih preizkusov zlitine AZ80 – srednje vrednosti Molar concentration /M* Ecorr/mV Icorr/(A/cm 2) Rp/(Ω cm2) 0.01 –1434 0.010 2610 0.2 –1540 0.062 427 0.6 –1573 0.093 280 1.0 –1576 0.128 203 2.0 –1593 0.252 104 * M = mol/L (ISO 80000) J. PRZONDZIONO et al.: RESISTANCE TO ELECTROCHEMICAL CORROSION OF THE EXTRUDED MAGNESIUM ... 276 Materiali in tehnologije / Materials and technology 49 (2015) 2, 275–280 The tests proved that the corrosion characteristics of the alloy decrease with the increase of the chloride ion concentration. The corrosion potential decreased from Ecorr = –1434 mV (0.01 M NaCl) to Ecorr = –1593 mV (2 M NaCl). It was observed that the polarisation resi- stance decreased from Rp = 610  cm2 (0.01 M NaCl) to Rp = 104  cm2 (2 M NaCl). The corrosion current den- sity increased from icorr = 0.01 μA/cm2 (0.01 M NaCl) to icorr = 0.252 μA/cm2 (2 M NaCl). Table 4 presents selected results of the immersion test. The corrosion rate V in the immersion test was determined from Equation (1): V m m St = −0 1 (1) where V is the corrosion rate (mg/(cm2 d)), m0 is the ini- tial mass of the sample (mg), m1 is the sample mass after removal of the corrosion products (mg), S is the area (cm2), and t is the exposure time (d). Table 4: Immersion test results Tabela 4: Rezultati preizkusov s potapljanjem Concentration NaCl M Corrosion rate, V/(mg cm–2 d–1) 1 d 5 d 0.01 0.256 0.343 0.6 1.103 1.295 2.0 2.332 3.850 The results of the immersion tests for the tested alloy confirmed, just as the potentiodynamic tests did, that the alloy was more prone to electrochemical corrosion when the molar concentration of the solution increased. The corrosion rate in a solution with a concentration of 0.01 M NaCl increased from 0.256 mg/(cm2 d) after the 1st day to 0.343 mg/(cm2 d) after 5 d. The tests performed in the time period 1 d to 5 d in the 2 M NaCl solution showed that corrosion rate increased from 2.332 mg/(cm2 d) to 3.85 mg/(cm2 d). The results of the tests performed with the scanning microscope FE SEM S-4200 Hitachi are presented in Figures 2 and 3. J. PRZONDZIONO et al.: RESISTANCE TO ELECTROCHEMICAL CORROSION OF THE EXTRUDED MAGNESIUM ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 275–280 277 Figure 1: Anodic polarisation curves of AZ80 alloy Slika 1: Anodna polarizacijska krivulja zlitine AZ80 Figure 2: The surface of the alloy after 1 d exposure in the solution with a concentration of: a) 0.01 M, b) 0.6 M and c) 2 M NaCl Slika 2: Povr{ina zlitine po izpostavi 1 d v raztopini s koncentracijo: a) 0,01 M, b) 0,6 M in c) 2 M NaCl Quantitative and qualitative analyses enabled us to identify the intermetallic phases present in the magne- sium alloy AZ80. The presence of phases of the MgAl-, MgMnAl-, and MgAlSi-type was detected. Figures 4 and 5 present the results of the tests per- formed with the optical surface analyser MicroProf by FRT, that are illustrated by selected 3D images of the AZ80 alloy and the distribution of the roughness. J. PRZONDZIONO et al.: RESISTANCE TO ELECTROCHEMICAL CORROSION OF THE EXTRUDED MAGNESIUM ... 278 Materiali in tehnologije / Materials and technology 49 (2015) 2, 275–280 Figure 5: a) 3D image of the surface of the alloy and b) roughness distribution after 5 d exposure in 2 M NaCl solution Slika 5: a) 3D posnetek povr{ine zlitine in b) razporeditev hrapavosti po izpostavitvi 5 d v raztopini 2 M NaCl Figure 4: a) 3D image of the surface of the alloy and b) roughness distribution after 5 d exposure in 0.01 M NaCl solution Slika 4: a) 3D posnetek povr{ine zlitine in b) razporeditev hrapavosti po izpostavitvi 5 d v raztopini 0,01 M NaCl Table 5: EIS analysis results Tabela 5: Rezultati EIS-analize NaCl concen- tration, M Rs/ (k cm2) Rf/ (k cm2) Cf/ (μF cm–2) CPEf Cdl/ (μF cm–2) L/ (H cm–2) Rct/ (k cm2) RL/ (k cm2)Y01/ (–1 cm-2 s-n) n2 0.01M 2.46 0.55 14.36 – – 13.51 0.37e–7 0.21 0.20 0.2 M 0.49 1.10 19.03 – – 7.99 0.14e–7 0.69 1.28 0.6 M 0.21 0.17 – 0.1572e–4 0.90 480.0 7.13 0.05 0.02 1 M 0.14 0.41 – 0.1475e–4 0.89 129.8 35.13 0.16 0.03 2 M 0.08 0.01 – 0.1652e–4 0.91 526.0 5.19 0.03 0.01 Figure 3: The surface of the alloy after 5 d exposure in the solution with a concentration of: a) 0.01 M, b) 0.6 M and c) 2 M NaCl Slika 3: Povr{ina zlitine po izpostavi 5 d v raztopini s koncentracijo: a) 0,01 M, b) 0,6 M in c) 2 M NaCl It was proved that with an increase of both the exposure time and the NaCl solution concentration the roughness parameters of the AZ80 deteriorated substan- tially. For instance, the average arithmetic deviation of the roughness profile Ra increases when the concen- tration is 0.01 M NaCl, from 0.745 μm (1 d) to 1.25 μm (5 d), and when the concentration is 2 M NaCl, from 0.944 μm (1 d) to 3.3 μm (5 d). The maximum height of the roughness profile Rz for the same concentrations 0.01 and 2 M increases from 4.89 μm (1 d) to 8.41 μm (5 d) and from 6.4 μm (1 d) even to 37 μm (5 d). The results of the electrochemical impedance tests of the AZ80 alloy are presented in Table 5. The obtained diagrams enabled us to match equiva- lent systems that are physical models depicting phenomena taking place in the respective object. It was proved that the best matching of the experimental im- pedance spectra is obtained with the application of an equivalent electrical system consisting of: • For samples exposed to 0.01 M and 0.2 M NaCl solu- tion from two consecutive parallel systems: within the range of low frequencies from the parallel capaci- tance system connected with the resistance of ion transition through phase boundary: metal – solution Rct, resistance RL together with coil (metallic conduc- tor with electromagnetic induction) L representing the corrosion processes; within the range of medium frequencies from the parallel capacitance system Cf connected with the resistance of transition of the Rf ions placed on the surface of the alloy in the result of corrosion (the layer consisting of corrosion products) and resistance at high frequencies, which may be attributed to the resistance of the electrolyte Rs (Table 5). In Figure 6 Rf and Cf designate, respectively, the resi- stance and capacitance of the layer created as the result of corrosion (corrosion product layer). Rct indicates the resistance of the charge transition and Cdl the capacitance of the double (porous) layer, then RL and L – induction loop, implicating the initiation and development of the pitting corrosion process. The mathematical model of the impedance for the system: AZ80 alloy – double layer –NaCl solution (0.01 M and 0.2 M) is presented by Equation (2): Z R R j C R j C R j L S= + + + + + + − 1 1 1 1 1 / ( ) / ( ) / ( ) f f ct dl L    (2) • For samples exposed to a solution from 0.6 M to 2 M NaCl within the range of low frequencies from a parallel capacitance system connected with the re- sistance of ion transition through the phase-boundary metal – solution Rct, resistance RL with the coil (metallic conductor with electromagnetic induction) L representing corrosion processes; within the range of medium frequencies from the parallel system of CPEf connected with the resistance of transition of the Rf ions placed on the surface of the alloy as the result of corrosion (corrosion product layer) and the resistance at high frequencies that may be attributed to the resistance of electrolyte Rs (Table 5). In Figure 7 CPEf indicates CPE depicting the cha- racter of the layer created as the result of the corrosion (corrosion product layer), Rf indicates, respectively, the resistance of that layer, whereas Rct resistance of ion transition, and Cdl capacitance of double (porous) layer, when RL and L – induction loop, implicating the ini- tiation and development of the pitting corrosion process. The mathematical model of the impedance for the system: AZ80 alloy – double layer –NaCl solution (0.6 M, 1 M and 2 M) is presented by Equation (3): Z R R Y j R j C R j L S n= + + + + + + − 1 1 1 1 1 0/ ( ) / ( ) / ( ) f ct dl L    (3) 4 CONCLUSIONS It is anticipated that the application of magnesium alloys in future years will be systematically increasing and more and more machine parts and units will be made from that group of materials. Their advantage is the fact that they can be formed with the application of casting as well as plastic working. The application of the magnesium alloy AZ80 after plastic working is dependent to a large extent on its resi- stance to electrochemical corrosion. The results of the performed tests prove explicitly the deterioration of the corrosion characteristics of the alloy with an increase of the molar concentration of the NaCl solution. Poten- tiodynamic tests performed in solutions with concentra- tions of 0.01–2.00 M NaCl showed that with an increase J. PRZONDZIONO et al.: RESISTANCE TO ELECTROCHEMICAL CORROSION OF THE EXTRUDED MAGNESIUM ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 275–280 279 Figure 7: Physical model of equivalent electrical system of corrosion system metal – solution Slika 7: Fizikalni model elektri~nega sistema, enakovrednega korozij- skemu sistemu kovina – raztopina Figure 6: Physical model of equivalent electrical system Slika 6: Fizikalni model enakovrednega elektri~nega sistema in the chloride ions concentration, a decrease of the corrosion potential and polarisation resistance, as well as an increase of the corrosion current density of the alloy can be observed. The deterioration of the corrosion characteristics with an increase of the NaCl solution concentration was also confirmed by immersion tests and during the metallographic tests. Microscopic tests of the samples made of the AZ80 alloy enable us to observe corrosion pits at each stage of the test. Within the early stage of pitting, pits were selec- tively located in the areas where non-metallic precipi- tates or inclusions were present (Figures 2a and 3a). The effect of the internal galvanic corrosion was evidenced. The secondary-phase particles were preferentially and uniformly corroded, while the  phase was being obviously unattacked. During the development of corro- sion, galvanic corrosion should thus have been less important and the increased corrosion attack of the whole structure was noticed. Significant degradation of the grain boundaries (Figure 3c) and the existence of crevices (Figures 2c and 3b) were observed. Deep anodic etching of the micrstructure and corres- ponding roughness of the samples were observed with an increase of the exposure time. The performed EIS tests enabled a direct comparison of the behaviour of the real object with its equivalent system, i.e., an electrical model related to physically realised impedance. To sum up, it must be highlighted that pitting corro- sion is present in the tested magnesium alloy. This proves that the extruded magnesium alloy AZ80 is not resistant to that type of corrosion. The prospects for its application in the aircraft industry trigger the need for covering elements made of the tested alloy with protec- tive coatings. Acknowledgements Financial support of Structural Funds in the Operatio- nal Programme–Innovative Economy (IE OP) financed from the European Regional Development Fund – Project "Modern material technologies in aerospace industry", No POIG.0101.02-00-015/08 is gratefully acknowledged. 5 REFERENCES 1 A. Kie³bus, D. Kuc, T. Rzychoñ, Magnesium alloys – microstructure, properties and application, Modern metallic materials – presence and future, Department of Materials Engineering and Metallurgy, Katowice, 2009 (in Polish) 2 K. Bry³a, J. Dutkiewicz, L. Lityñska-Dobrzyñska, L. L. Rokhlin, P. Kurtyka, Influence of number of ECAP passes on microstructure and mechanical properties of AZ31 magnesium alloy, Archives of Metallurgy and Materials, 57 (2012) 3, 711–717, doi:10.2478/ v10172-012-0077-5 3 Z. Cyganek, M. Tkocz, The effect of AZ31 alloy flow stress des- cription on the accuracy of forward extrusion FE simulation results, Archives of Metallurgy and Materials, 57 (2012) 1, 199–204, doi:10.2478/v10172-012-0010-y 4 R. Kawalla, G. Lehmann, M. Ullmann, H. P. Vogt, Magnesium semi-finished products for vehicle construction, Archives of Civil and Mechanical Engineering, 8 (2008) 2, 93–101, doi:10.1016/ S1644-9665(12)60196-4 5 M. Gandara, Recent growing demand for magnesium in the automo- tive industry, Mater. Tehnol., 45 (2011) 6, 633–637 6 G. L. Makar, J. Kruger, Corrosion of magnesium, International Mate- rials Reviews, 38 (1993) 3, 138–153, doi:10.1179/095066093790 326320 7 G. Song, A. Atrens, X. Wu, B. Zhang, Corrosion behaviour of AZ21, AZ501 and AZ91 in sodium chloride, Corrosion Science, 40 (1998) 10, 1769–1791, doi:10.1016/S0010-938X(98)00078-X 8 G. Song, A. Atrens, Corrosion mechanisms of magnesium alloys, Advanced Engineering Materials, 1 (1999) 1, 11–33, doi:10.1002/ (SICI)1527-2648(199909)1:1 11::AID-ADEM11 3.0.CO;2-N 9 J. Przondziono, W. Walke, E. Hadasik, Galvanic corrosion test of magnesium alloys after plastic forming, Solid State Phenomena, 191 (2012), 169–176, doi:10.4028/www.scientific.net/SSP.191.169 10 H. Altun, S. Sen, Studies on the influence of chloride ion concen- tration and pH on the corrosion and electrochemical behaviour of AZ63 magnesium alloy, Materials and Design, 25 (2004) 7, 637–643, doi:10.1016/j.matdes.2004.02.002 11 R. Ambat, N. N. Aung, W. Zhou, Evaluation of microstructural effects on corrosion behaviour of AZ91D magnesium alloy, Corro- sion Science, 42 (2000) 8, 1433–1455, doi:10.1016/S0010-938X (99)00143-2 12 J. Przondziono, W. Walke, E. Hadasik, J. Szala, J. Wieczorek, Corrosion resistance tests of magnesium alloy WE43 after extrusion, Metalurgija, 52 (2013) 2, 242–246 13 S. Amira, D. Dubé, R. Tremblay, E. Ghali, Influence of the micro- structure on the corrosion behavior of AXJ530 magnesium alloy in 3.5 % NaCl solution, Materials Characterization, 59 (2008) 10, 1508–1517, doi:10.1016/j.matchar.2008.01.018 14 Z. Yu, D. Ju, H. Zhao, Effect of Stress on the Electrochemical Corro- sion Behavior of Mg-Zn-In-Sn Alloy, International Journal of Electrochemical Science, 7 (2012) 8, 7098–7110 15 T. Zhang, G. Meng, Y. Shao, Z. Cui, F. Wang, Corrosion of hot extrusion AZ91 magnesium alloy, Corrosion Science, 53 (2011) 9, 2934–2942, doi:10.1016/j.corsci.2011/05.035 16 Z. Yu, D. Ju, N. Takashi, Effect of Stresses for Electrochemical Calculations of Mg-Zn-In-Sn Alloy, International Journal of Electro- chemical Science, 7 (2012) 10, 10164–10174 17 P. Lichy, J. Beòo, M. Cagala, J. Hampl, Thermophysical and thermo- mechanical properties of selected alloys based on magnesium, Metalurgija, 52 (2013) 4, 473–476 18 P. Pal~ek, M. Chalupová, I. Hlavá~ová, Efect of microstructure on the development of plastic deformation around the propagating cracks, Acta Metallurgica Slovaca – Conference, 3 (2013), 202–208, doi:10.12776/amsc.v3.128 19 M. Kappes, M. Iannuzzi, R. M. Carranza, Pre-exposure embrittle- ment and stress corrosion cracking of magnesium alloy AZ31B in chloride solutions, Corrosion, 70 (2014) 7, 667–677, doi:10.5006/ 1172 J. PRZONDZIONO et al.: RESISTANCE TO ELECTROCHEMICAL CORROSION OF THE EXTRUDED MAGNESIUM ... 280 Materiali in tehnologije / Materials and technology 49 (2015) 2, 275–280 L. MÜNSTER et al.: MICROWAVE-ASSISTED HYDROTHERMAL SYNTHESIS OF Ag/ZnO SUB-MICROPARTICLES MICROWAVE-ASSISTED HYDROTHERMAL SYNTHESIS OF Ag/ZnO SUB-MICROPARTICLES HIDROTERMI^NA SINTEZA PODMIKROMETRSKIH DELCEV Ag/ZnO Z MIKROVALOVI Luká{ Münster1,2, Pavel Ba`ant1,2, Michal Machovský2, Ivo Kuøitka1,2 1Polymer Centre, Faculty of Technology, Tomas Bata University in Zlin, Nam. T. G. Masaryka 275, 762 72 Zlin, Czech Republic 2Centre of Polymer Systems, University Institute, Tomas Bata University in Zlin, Nad Ovcirnou 3685, 760 01 Zlin, Czech Republic l_munster@ft.utb.cz Prejem rokopisa – received: 2013-10-02; sprejem za objavo – accepted for publication: 2014-04-03 doi:10.17222/mit.2013.223 A fast and environmentally friendly, microwave-assisted, hydrothermal synthesis was utilized for the preparation of a Ag/ZnO hybrid system by using zinc acetate dihydrate and silver nitrate as the sources of Zn2+ and Ag+, and hexamethylentetramine as the reducing and precipitating agent. The influence of the concentration was investigated by X-ray diffraction analysis, scanning electron microscopy and energy-dispersive analysis. It was found that the concentration has a strong effect on the morphology and proportion between the Ag and ZnO components of the prepared particulate materials. With a decreasing concentration, the morphology of the ZnO changed from twinned or single frustums to rod-like microparticles, whereas the silver morphology changed from large polygon-shaped microparticles to very small, spherical nanoparticles. Keywords: zinc oxide, silver, nanoparticle, hydrothermal, microwave synthesis Uporabljena je bila hitra in okolju prijazna mikrovalovna hidrotermi~na sinteza hibridnega sistema Ag/ZnO z uporabo cink-acetat dihidrata in srebrovega nitrata kot vira Zn2+ in Ag+ ter heksametilentetramina kot reducenta in sredstva za izlo~anje. Z rentgensko difrakcijo, vrsti~no elektronsko mikroskopijo in energijsko disperzijsko analizo je bil preu~evan vpliv koncentracije. Ugotovljeno je bilo, da ima koncentracija mo~an vpliv na morfologijo in razmerje med komponentama Ag in ZnO pripravljenega zrnatega materiala. Z zmanj{anjem koncentracije se je morfologija ZnO spremenila iz dvoj~i~nih ali sto`~astih mikrodelcev v pali~aste, medtem ko se je morfologija srebra iz velikih poligonalnih mikrodelcev spremenila v majhne sferi~ne nanodelce. Klju~ne besede: cinkov oksid, srebro, nanodelec, hidrotermi~en, sinteza z mikrovalovi 1 INTRODUCTION Various metals (Au, Ag) and metal oxides (ZnO, TiO2, SnO2, CuO) have been the subjects of investigation for a long time because of their unique optical, electrical, and mechanical properties that can be modulated by varying the size and the morphology of the particles on the nano- and sub-micro scales1,2. Recently, hybrid metal-semiconductor materials have attracted great attention because their metal and semiconductor inter- face possesses a unique electronic band structure result- ing in a specific chemical activity3. Therefore, great efforts have been devoted to the synthesis of hybrid materials by combining various metals and semiconduc- tors such as Ag/TiO2, Ag/SiO2, Pt/SnO2 and Ag/ZnO.4–7 Among them, Ag/ZnO is of great interest because of potential applications in many fields, such as catalysis and medicine.7,8 This particulate material has been successfully demonstrated as a suitable filler for an inorganic antibacterial polymer system9. ZnO is an im- portant semiconducting material because of its unique optoelectronic and piezoelectric properties.10,11 More- over, it possesses strong antibacterial activity towards gram-negative bacteria and it is both non-toxic and biocompatible10. Metallic silver, its salts and complexes have been exploited for their antibacterial properties for millennia before the realization that bacteria are the agents of the infection2. Silver nanoparticles are com- mercially available on the antibacterial-additives market, exhibiting a broad spectrum of antibacterial activity against gram-positive, gram-negative bacteria, fungi, cer- tain viruses, including antibacterial-resistant strains.12–14 Therefore, with the combination of Ag and ZnO in a hybrid material on the nano- or sub-micro scales could achieve a synergic effect of antibacterial and photocata- lytic activity, biosensing, electrochemical properties, etc.15–17 Hence, we developed a simple, fast and efficient preparation of a Ag/ZnO hybrid via a microwave (MW) assisted hydrothermal synthesis using a modified open-vessel laboratory microwave-oven system18. Intro- ducing microwaves instead of the conventional heating allowed us to shorten the reaction time to 12 min, com- pared to hours reported in references for conventional heating17. The concentration of reactants is one of the key factors affecting the properties of the prepared hybrid material. Therefore, a study of the influence of the reactants’ concentrations on the morphology and particles size was performed. Materiali in tehnologije / Materials and technology 49 (2015) 2, 281–284 281 UDK 54.057:542.9 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 49(2)281(2015) 2 EXPERIMENTAL WORK 2.1 Sample preparation Zinc acetate dihydrate Zn(Ac)2 · 2H2O (PENTA) and silver nitrate AgNO3 (PENTA) were used as the sources of Zn2+ and Ag+, respectively. Hexamethylentetramine (HMTA), (CH2)6N4 (LACHNER) was used as the source of OH–, acting as a precipitating agent and CH2O as a reducing agent. All the chemicals were of analytical grade and used as received without any further purifica- tion. Demineralised water with a conductivity 0.1 μS cm–1 was used throughout the experiment. The synthesis was carried out using an open-vessel microwave-oven laboratory system MWG1K-10 ope- rated at 800 W and 2.45 GHz (RADAN) with an external reflux cooler. In a typical experimental procedure a given amount of precursor (Zn2+, Ag+ or both) and re- ducing/precipitating agent were dissolved in 100 mL and 50 mL of demineralised water, respectively. The precur- sor solutions were first preheated in the microwave oven for 2 min, then the solution of HMTA was added via a dropping funnel, and then microwave heating continued for another 10 min. Hence, the total time of the micro- wave synthesis was 12 min for each prepared sample. The reaction mixture was always left to cool to approxi- mately 50 °C and then the material was collected and washed by microfiltration. Samples were dried at 40 °C in a laboratory oven until a constant weight was achieved. In all the experiments, the molar ratio of Zn2+/Ag+ was set to 10/1. In order to investigate the influence of concentration, the initial solution mixture concentration, which is close to the saturation limit for zinc acetate (set I), was lowered ten (set II) and fifty times (set III), giving rise to three sets of samples. The concentration of HMTA was kept at a constant ratio to the precursors as well. For the reaction-mixture compo- sitions (Table 1). 2.2 Characterization methods Crystal-phase structure identification was conducted using the multifunctional X-ray diffractometer PANaly- tical X’Pert Pro MPD (PANalytical) and Cu-K X-ray source ( = 0.15418 nm), at 40 kV and 30 mA and the 2 diffraction angle range 10–90 °. The morphology was observed in a scanning electron microscope (SEM) Vega II LMU (TECSAN) at an acceleration voltage of 10 kV. Semi-quantitative elemental analysis was performed with the energy-dispersive X-ray (EDX) analyzer (Oxford Instruments, INCA, UK) mounted on the microscope. 3 RESULTS AND DISCUSSION All the samples were examined by powder X-ray diffraction (XRD) analyses. As an example of the crystal-phase structure investigation, X-ray patterns of the samples prepared at the lowest concentration, i.e., pure ZnO, Ag and hybrid Ag-ZnO (set III) are shown in Figure 1. Other sets of the experiment with higher concentrations of precursors and reducing/precipitating agents (set I and II) showed similar XRD patterns. Specific diffraction lines (labeled by #) were observed at approximately 2 = (31.7, 34.4, 36.2, 47.5, 56.5, 62.8, 66.2, 67.8, 72.5, 76.8, 81.3 and 89.7) °. According to the L. MÜNSTER et al.: MICROWAVE-ASSISTED HYDROTHERMAL SYNTHESIS OF Ag/ZnO SUB-MICROPARTICLES 282 Materiali in tehnologije / Materials and technology 49 (2015) 2, 281–284 Table 1: Sets of samples and amounts of precursors and reducing/precipitating agent used in MW synthesis and yields of the powder product Tabela 1: Sklopi vzorcev in koli~ina predhodnikov in reducirnega/izlo~evalnega sredstva, uporabljenih v MW-sintezi, in izkoristki proizvodov v obliki prahu Set Sample Zn(Ac)2·2H2Omol AgNO3 mol HMTA mol Yield g Theoretical Yield g I I-ZnO 0.0500 – 0.0500 0.740 4.070 I-Ag – 0.0050 0.0500 0.501 0.539 I-AgZnO 0.0500 0.0050 0.0500 1.160 4.609 II II-ZnO 0.0050 – 0.0050 0.190 0.407 II-Ag – 0.0005 0.0050 0.048 0.054 II-AgZnO 0.0050 0.0005 0.0050 0.182 0.461 III III-ZnO 0.0010 – 0.0010 0.025 0.041 III-Ag – 0.0001 0.0010 0.003 0.005 III-AgZnO 0.0010 0.0001 0.0010 0.026 0.046 Figure 1: X-ray diffractogram of samples from set III, the 2 range is cropped at lower values in the graph as no peaks were detected bet- ween 10 ° and 30 ° Slika 1: Rentgenski difraktogram vzorcev iz sklopa III, podro~je 2 je odrezano pri ni`jih vrednostih, ker v obmo~ju 10–30 ° ni bilo vrhov JCDD PDF-2 entry 01-079-0207, these peaks can be unambiguously assigned to the ZnO wurtzite hexagonal crystal phase. Other observed diffraction peaks (labeled by *) at approximately 2 = (38.2, 44.4, 64.6, 77.6 and 81.8) ° were assigned to a metallic silver cubic crystal phase according to JCDD PDF-2 entry 01-087-0720. Moreover, the sharpness and the narrowness of the diffraction peaks imply well-developed crystalline struc- tures for both the Ag and ZnO, with no level traces of crystalline impurities. The morphology of the prepared particles is shown in Figure 2. A back-scattered electron (BSE) detector was utilized to enhance the contrast between the Ag and ZnO particles. It is clear that the concentration strongly influenced the particles size and the morphology. For pure ZnO synthesized at the highest concentration (set I), the particles form typical twinned frustums joined by their apical bases, in many cases hollow if single frustums. The dimensions of these particles are diameter up to 1.5 μm and in length up to 2 μm. The particles with a hollow core have a wall thickness starting from tens of nanometers. Lowering the concentration to one tenth of the ZnO precursor (set II) resulted in an increased aspect ratio from approximately 1 : 1 observed for pure ZnO at the highest concentration (set I) to 6 : 1. The approxi- mate diameter and length were 500 nm and 3 μm. At the lowest concentration (set III), rod-shaped ZnO particles have the same aspect ratio; however, there is observable a slight increase of the particles’ length and diameter. The morphology of the Ag particles prepared by using the highest concentration of the precursor (set I), dis- played polygon-shaped particles with a diameter up to 1.5 μm. The lowering of the concentration of the Ag precursor (set II) results in spherically shaped Ag particles with an approximate diameter of 250 nm. The lowest concentration of Ag (set III) displayed the smallest particles with a size in the range of tens of L. MÜNSTER et al.: MICROWAVE-ASSISTED HYDROTHERMAL SYNTHESIS OF Ag/ZnO SUB-MICROPARTICLES Materiali in tehnologije / Materials and technology 49 (2015) 2, 281–284 283 Figure 3: EDX spectra of samples from set III Slika 3: EDX-spektri vzorcev iz sklopa III Figure 2: SEM micrographs of prepared samples of ZnO, Ag and AgZnO particles arranged in dependence of the concentration of precursors and HMTA in the reaction mixture Slika 2: SEM-posnetki pripravljenih vzorcev ZnO, Ag in AgZnO delcev, razvr{~enih v odvisnosti od koncentracije predhodnikov in HMTA v reakcijski zmesi nanometers. For the hybrid AgZnO system, the trend is very similar as that observed for the pure Ag and ZnO components in the prepared material. Figure 3 shows typical spectra obtained from the EDX analyses. For the sake of brevity, only spectra for the lowest concentration (set III) are presented. The presence of elemental zinc (peaks at (1.0, 8.6 and 9.5) keV) and oxygen (peak at 0.5 keV) from the ZnO parti- cles was confirmed in the samples ZnO, Ag/ZnO. Ele- mental silver (peaks at 2.6 keV and 3.0 keV) was confirmed even within the lowest concentration used in the samples Ag and Ag/ZnO. Elemental carbon with a peak at 0.3 keV can be assigned to the carbon tape used for fixation of the sample to the holder inside the micro- scope chamber or to other organic impurities. The semi-quantitative results from the EDX in the weight percentage of Ag, Zn and O content are listed in Table 2. It is clear that the weight percentage of Zn increases with lowering of the concentration, whereas the content of Ag decreases with lowering the concentration, which is in accordance with observations based on the SEM micro- graphs. The large observed content of O points towards the presence of water adsorbed at the surface or in the pores of the particles. Table 2: EDX results of prepared samples, content of Zn, Ag and O in mass fractions, w/% Tabela 2: EDX-rezultati pripravljenih vzorcev, vsebnost Zn, Ag in O v masnih dele`ih, w/% EDX (w/%) Set ZnO Ag AgZnO I Zn: 59.1; O: 40.9 Ag: 100 Ag: 39.5; Zn: 31.7;O: 28.8 II Zn: 60.9; O: 39.1 Ag: 100 Ag: 28.2; Zn: 36.4;O: 35.5 III Zn: 65.2; O: 34.8 Ag: 100 Ag: 3.3; Zn: 62.2;O: 34.4 4 CONCLUSION A simple one-pot open-vessel microwave-assisted hydrothermal synthesis of a ZnO, Ag and Ag/ZnO hybrid combination of both types of particles was intro- duced. The influence of the concentration of precursors and the precipitating/reducing agent on the properties and on the morphology was investigated. Powder XRD analysis showed wurtzite crystalline phases of zinc oxide in all the samples prepared using Zn(Ac)2 · 2H2O as a precursor and cubic crystalline phases of silver prepared using the precursor AgNO3, even at the lowest concen- trations. The alternation of the concentration has a marked effect on the particle morphology: (i) the de- crease of the concentration of the zinc precursor results in prolongation and an increased size of the ZnO parti- cles, whereas (ii) the size of the silver particles decreases directly with the concentration decrease, from microme- ter down to tens of nanometers, and the highest percen- tage yield was almost 93 % of pure silver particles. Acknowledgment The authors wish to thank the internal grant of TBU in Zlin No. IGA/FT/2013/014 funded from the resources of specific university research for financial support. This article was written with the support of the Operational Program "Research and Development for Innovations" co-funded by the European Regional Deve- lopment Fund (ERDF) and the national budget of the Czech Republic, within the Centre of Polymer Systems project (reg. number: CZ.1.05/2.1.00/03.0111). This article was written with the support of the Operational Program "Education for Competitiveness" co-funded by the European Social Fund (ESF) and the national budget of the Czech Republic, within the "Ad- vanced Theoretical and Experimental Studies of Polymer Systems" project (reg. number: CZ.1.07/2.3.00/20.0104). 5 REFERENCES 1 D. C. Look, Materials Science and Engineering: B, 80 (2001), 383–387, doi:10.1016/S0921-5107(00)00604-8 2 C. Marambio-Jones, E. M. V. Hoek, Journal of Nanoparticle Re- search, 12 (2010), 1531–1551, doi:10.1007/s11051-010-9900-y 3 Y. H. Zheng, L. R. Zheng, Y. Y. Zhan, X. Y. Lin, Q. Zheng, K. M. Wei, Inorganic Chemistry, 46 (2007), 6980–6986, doi:10.1021/ ic700688f 4 W. Su, S. S. Wei, S. Q. Hu, J. X. Tang, Journal of Hazardous Ma- terials, 172 (2009), 716–720, doi:10.1016/j.jhazmat.2009.07.056 5 G. Gu, J. Xu, Y. Wu, M. Chen, L. Wu, Journal of Colloid and Inter- face Science, 359 (2011), 327–33, doi:10.1016/j.jcis.2011.04.002 6 M. H. Madhusudhana Reddy, A. N. Chandorkar, Thin Solid Films, 349 (1999), 260–265, doi:10.1016/S0040-6090(99)00194-7 7 C. Ren, B. Yang, M. Wu, J. Xu, Z. Fu, Y. lv, T. Guo, Y. Zhao, C. Zhu, Journal of Hazardous Materials, 182 (2010), 123–129, doi:10.1016/ j.jhazmat.2010.05.141 8 A. Meng, S. Sun, Z. Li, J. Han, Advanced Powder Technology, 24 (2013), 224–228, doi:10.1016/j.apt.2012.06.006 9 P. Bazant, I. Kuritka, O. Hudecek, M. Machovsky, M. Mrlik, T. Sedlacek, Polymer Composites, 35 (2014), 19–26, doi:10.1002/ pc.22629 10 C. Jagadish, S. Pearton, Zinc Oxide Bulk, Thin Films and Nano- structures, Elsevier, Amsterdam 2006 11 A. Corso, M. Posternak, R. Resta, Physical Review: B, 50 (1994), 10715–10721, doi:10.1103/PhysRevB.50.10715 12 R. M. Slawson, M. I. Van Dyke, H. Lee, J. T. Trevors, Plasmid, 27 (1992), 72–79, doi:10.1016/0147-619X(92)90008-X 13 D. J. Balazs, K. Triandafillu, P. Wood, Y. Chevolot, C. Van Delden, H. Harms, C. Hollestein, H. J. Mathieu, Biomaterials, 25 (2004), 2139–2151, doi:10.1016/j.biomaterials.2003.08.053 14 N. Stobie, B. Duffy, D. E. McCormack, J. Colreavy, M. Hildalgo, P. Mchale, S. J. Hinder, Biomaterials, 29 (2008), 963–969, doi:10.1016/ j.biomaterials.2007.10.057 15 W. Lu, G. Liu, S. Gao, S. Xing, J. Wang, Nanotechnology, 19 (2008), 1–10, doi:10.1088/0957-4484/19/44/445711 16 C. Tian, Q. Zhang, B. Jiang, G. Tian, H. Fu, Journal of Alloys and Compounds, 509 (2011), 6935–6941, doi:10.1016/j.jallcom.2011. 04.005 17 X. Feng, Y. Cheng, C. Ye, J. Ye, J. Peng, J. Hu, Materials Letters, 79 (2012), 205–208, doi:10.1016/j.matlet.2012.03.098 18 M. Machovsky, P. Bazant, Z. Kozakova, M. Pastorek, P. Zlebek, I. Kuritka, Open Vessel Microwave-Assisted Synthesis of Ag/Zno Hybrid Fillers with Antibacterial Activity, Proc. of the Nanocon, Brno, 2011, 628–634 L. MÜNSTER et al.: MICROWAVE-ASSISTED HYDROTHERMAL SYNTHESIS OF Ag/ZnO SUB-MICROPARTICLES 284 Materiali in tehnologije / Materials and technology 49 (2015) 2, 281–284 Z. FRANÌK et al.: DETERMINATION OF THE CAUSE OF THE FORMATION OF TRANSVERSE INTERNAL CRACKS ... DETERMINATION OF THE CAUSE OF THE FORMATION OF TRANSVERSE INTERNAL CRACKS ON A CONTINUOUSLY CAST SLAB UGOTAVLJANJE VZROKOV ZA NASTANEK NOTRANJIH PRE^NIH RAZPOK V KONTINUIRNO ULITEM SLABU Zdenìk Franìk1, Milo{ Masarik2, Jaromír [míd3 1Silesian University in Opava, School of Business Administration in Karvina, Univerzitní nám. 1934/3, 733 40 Karviná, Czech Republic 2VITKOVICE STEEL, a.s., ^eskobratrská 3321/46, 702 00 Ostrava, Czech Republic 3TaM, Technologie a Metalurgie, Areál VÚH@, a.s., 739 51 Dobrá 120, Czech Republic franek@opf.slu.cz Prejem rokopisa – received: 2013-10-03; sprejem za objavo – accepted for publication: 2014-06-09 doi:10.17222/mit.2013.224 The paper describes a determination of the cause of transverse internal cracks on a continuously cast slab using an analytical software tool, called LITIOS. It is a complex system based on long-term monitoring of the casting parameters and their impact on the quality of the slabs. Transverse internal cracks are characterised by identifying the possible causes of their origin. Using the LITIOS system, selected constant (invariable) and variable parameters of the casting were assigned to specific cracks and on their basis the causes of the cracks were determined. Keywords: continuous casting of steel, slab, casting parameters, quality prediction, crack, software tool ^lanek opisuje ugotavljanje vzrokov notranjih pre~nih razpok v kontinuirno litem slabu z uporabo analitskega programskega orodja LITIOS. To je kompleksen sistem, ki temelji na dolgotrajnem nadzoru parametrov pri litju in njihovem vplivu na kvaliteto slabov. Ocenjene so pre~ne notranje razpoke s prikazom mogo~ih vzrokov za njihov nastanek. Z uporabo sistema LITIOS so bile dodeljene izbrane konstante (nespremenljivke) in spremenljivi parametri litja dolo~eni razpoki in na tej osnovi so bili dolo~eni vzroki za nastanek razpok. Klju~ne besede: kontinuirno litje jekla, slab, parametri litja, napovedovanje kakovosti, razpoka, programsko orodje 1 INTRODUCTION Continuous casting is a complex process, characte- rised by a considerable instability, i.e., by the so-called fluctuations. The fluctuations have a negative influence on the quality of slabs. To a large extent, they make it difficult to determine the causes of defects and, hence, also the optimum technological parameters for the casting process or their control during the events caused by the technology. The whole process of casting steel is thoroughly mapped by measuring and reading the monitored data, including the chemical-analysis results. Quality ratings are assigned to the produced continuously cast slabs and to the products rolled from them. In this way, a timeline of very extensive data, mapping the production process is created. It can be, however, concluded that the potential of these data has not yet been sufficiently used. The reason this is, among other things, the fact that it is very difficult to handle these data and that assigning the data to the appropriate place on a slab and a possible defect is not easy. In spite of that, the operation of continuous casting, especially the casting of slabs, is nowadays practically impossible without the use of the systems for monitoring the casting parameters and the software systems for optimising the casting parameters. An example of an artificial-intelligence technique for optimising the process parameters used for the conti- nuous casting of steel is given in1. The mathematical modelling and optimisation strategy, genetic algorithm and knowledge base, applied to the continuous casting of steel are given in2. A description of a properly developed algorithm solving an essential component of the above mentioned systems, including the assigning of the data to a specific place on the slab, is described in several works, e.g., in3. A multi-dimensional, statistical, model-based system for monitoring continuous casting and detecting the risk of impending breakouts is also the subject of the U.S. patent.4 Today, specialised companies offer different solutions for meeting the requirements for an ever increasing quality of steel, with simultaneous reductions in the manufacturing costs. For example, the company Siemens Metals Technologies provides the basic automation solu- tions for all the types of continuous-casting machines, i.e., for casting slabs, blooms, billets, the blanks for dog-bone sections (beam-blank casters) as well as for the equipment for producing endless strips. Sets of automa- tion programs meet the requirements for the performance of production units, dimensional flexibility, the quality of semi-products and final products, and the maximum Materiali in tehnologije / Materials and technology 49 (2015) 2, 285–290 285 UDK 621.74.047:620.192.46 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 49(2)285(2015) yield. Reference5 states that the described solutions were installed on hundreds of casting machines worldwide, in both the new and the existing plants. An extensive and comprehensive range of optimi- sation processes, optimisation models, technology-mana- gement functions and services is crucial for meeting the production and quality goals. On-line data links for a con- tinuous support provide a rapid system for modernisa- tion, maintenance and metallurgical assistance. Referen- ce5 also states that with the support of a vast experience the described packages combine excellent technology and automation solutions that perfectly control the pro- cess of continuous casting and its complex parameters. 2 USE OF THE SYSTEMS FOR MONITORING THE PARAMETERS AND QUALITY PREDICTION The data obtained from the above-mentioned systems for monitoring the casting parameters and the software systems for optimising the casting parameters call for the use of advanced statistical methods for monitoring and evaluating the quality of continuously cast blanks6–8 and for the subsequent prediction of defects. At present, it can be said that the basic technical development of the classical machines for continuous casting of steel has mostly reached such a level that a further development is often not necessary. Minor metal- lurgical, technical and technological improvements are naturally still ongoing in these areas, but, in general, it can be said that other possibilities for making the production more efficient through major changes are already limited by very narrow barriers. One of the areas where some reserves still exist is increasing and control of the quality of continuously cast blanks and, particularly, the final products of a metallur- gical plant. This can be achieved by introducing predic- tion systems, in the first stage, for predicting the quality of continuously cast blanks and then of the final products. According to the design of these systems their primary objective is to predict the quality of the products at the moment when they leave the production line, or even before it. There is a possibility of a transition from a discon- tinuous production in the segment of liquid-steel rolled products to a continuous production, with evidenced savings, by eliminating the energy-intensive reheating before the next forming of the slabs. The prediction of quality allows another direct processing of the slabs without a significant risk that the absence of their control in the cold state, with the subsequent removal of the detected defects, will lead to a rejection or a quality degradation of the final products due to their defects. Even in the case of non-removable defects of the slabs, their rejection before the useless further processing is economically advantageous. The second benefit of the prediction systems, which probably has not yet been sufficiently appreciated, is the use of the feedback between the quality of a slab and a final product, and the technological parameters for casting steel on a CCM, or the parameters for processing a continuously cast slab at a rolling mill. Archiving and statistical processing of the data describing these relations (the parameters versus the quality level) can provide, with a surprising success, the optimum techno- logical parameters for casting and for further processing. The methods for acquiring data and explaining their use in the metallurgical-industry operating conditions as a tool for improving the steelmaking processes, descrip- tion of the basic processes and various approaches to the implementation of the necessary software tools, as well as data acquisition, creation of knowledge databases and their storing, displaying of the results and diagnostics – all of these represent the essential activities and the segments of the developed software systems aimed at achieving the basic objective, i.e., an increase in the quality and in the yield of the production of steel pro- ducts.9 3 LITIOS SOFTWARE SYSTEM Within the research, under the support of various projects, in cooperation with technologists and metal- lurgists, numerous statistical analyses were made on the concrete CCM. The procedure and the obtained result for the specific defect are described below. This article describes the possibility to monitor dozens of the para- meters that are assigned to any section of a cast slab. In the event that the Baumann sulphur print made on a transverse sample taken from a cast slab in the standard manner shows any defects, it is possible to assign certain casting parameters to this sample and compare them with the documented reality. This procedure is then also the basic prerequisite for predicting the quality of cast slabs and, subsequently, of the sheets rolled from these slabs. If we assume, already at the general level, the knowledge about the qualitative relation between a casting parameter and a slab defect, this procedure is then used for confirming their relation. And, in the case of the limit values that differ between different CCMs, this procedure allows their exact deter- mination. This makes it possible to express the basic result of a prediction with a YES-NO verdict about the presence of a defect. Another necessary prerequisite for analysing the causes of defects in a slab is also the selection and assignment of the casting parameters to the place where these defects are formed. In the case of Baumann sulphur prints the above refers to the data assigned to the place, from which they were taken. The company EVRAZ VÍTKOVICE STEEL, a. s., implemented an original complex system of long-term monitoring of the casting parameters and of their impact on the quality of the slabs, called LITIOS.3 This software system is organically linked with an on-line temperature Z. FRANÌK et al.: DETERMINATION OF THE CAUSE OF THE FORMATION OF TRANSVERSE INTERNAL CRACKS ... 286 Materiali in tehnologije / Materials and technology 49 (2015) 2, 285–290 model, as well as with an on-line module of data acquisi- tion. The system deals with all the data available from the CCM process. The system comprises the recording and filtering of the data, their sorting, entering into the relational database system, as well as the data aggrega- tion and their graphical interpretation. Technological data, measured every 10 seconds, are entered from the temperature model. The software of the temperature model directly enters the data into the database of the LITIOS system. The LITIOS software system also records all the necessary data about each sequence of the superior system for the automatic control of the steel shop, called FLS, and it enters them into the database system. It enables a filtration of the data and performs the necessary data aggregation. This aggregation is necessary in order to simplify the work and handle large amounts of data. It turns out that it is appropriate and sufficient to aggregate the data per meter of the length of a casting strand. The developed software is modular, using the latest knowledge on the database technology and data-analysis methods. The system architecture is shown in Figure 1. Data storage is performed according to the hierarchy of their formation: the data about the sequence, heat, slab, measured value (the so-called channels) of the continuous casting machine, the data about the quality of the slabs and the products rolled from them. The files are divided into items with their attributes. The whole sys- tem (Figure 1) thus contains approximately 500 items, and if we take into account that the technological para- meters, the so-called channels, are measured every 10 seconds, we see that a very large amount of data is created. The system provides reports and it is possible to view the data for individual sequences, heats and slabs. The selected data are represented graphically. Various selections of the data are available for the analytical methods. A user can transform the data into a matrix of the causes – the measured values and the consequences – i.e., the quality indicators. The employees of a steel shop use these narrowed data analyses of the evolution of casting. The system allows an export of the selected data to the other statistical programs for more detailed analyses. The data from the LITIOS system were used for the analyses of the crack formation presented below. 4 TRANSVERSE INTERNAL CRACKS These cracks, also called the half-way cracks because of their position on a slab cross-section, which are caused by straightening or bending, are situated between the surface and the centre of a slab, in the plane perpen- dicular to the direction of casting. They mostly occur in the top half of a slab cross-section. They are schemati- cally illustrated in Figure 2. In some cases transverse internal cracks in conti- nuously cast slabs may cause a deterioration of the utility properties of the heavy plates rolled from them, or possibly even their rejection from the production as a result of an ultrasonic inspection. The best prevention against these undesirable phenomena during the produc- tion of heavy plates is an elimination of these cracks already during the continuous casting of slabs on a CCM. The basic causes of the formation of these cracks are high tensile deformations in the high-temperature zone of low strength and ductility. The formation of these deformations is explained with three different facts10: • intensive secondary cooling, which causes high-tem- perature reheating of the surface of a continuously cast slab Z. FRANÌK et al.: DETERMINATION OF THE CAUSE OF THE FORMATION OF TRANSVERSE INTERNAL CRACKS ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 285–290 287 Figure 2: Schematic illustration of transverse cracks and localisation of their occurrence in the transverse and longitudinal sections of a continuously cast slab Slika 2: Shematski prikaz pre~nih razpok in mesto njihove pojave na pre~nem in vzdol`nem prerezu kontinuirno litega slaba Figure 1: Architecture of the LITIOS monitoring system Slika 1: Shema nadzornega sistema LITIOS • a depression of the wide side of a slab • slab bending and straightening, particularly in the temperature interval of the reduced strength and ductility of steel, i.e., within the temperature interval from 950 °C to 700 °C In addition to the usual sensitivity of some steel grades to the formation of these cracks, there are also some other technological parameters that contribute to the formation of transverse cracks: • a high casting temperature, causing a wide zone of a columnar casting structure, sometimes taking place even after the transcrystallisation • a low casting rate, causing a reduction in the slab temperature in the area of its straightening • a high casting rate • a drastic increase in the cooling intensity after the exit of a slab from the mould • a chemical composition of steel; micro-alloyed steels are particularly susceptible to the formation of these cracks since micro-alloying elements, namely, vana- dium, titanium, niobium, etc., reduce the ductility of steel exactly at the critical temperatures of the slab straightening. When inspecting the causes of the formation of trans- verse intermediate cracks we see an obvious antagonism of the influence of the secondary cooling on the forma- tion of these defects. The so-called soft cooling is required, which causes a hotter and thinner strand shell leading to a lower resistance to depression. The compro- mise must be, therefore, the optimum intensity of the cooling. In the case of a formation of the cracks caused by slab bending and straightening, it is probable that a more important role is played by the straightening, indicated by the fact that the cracks are almost always situated in the top part of a slab, i.e., in the half of the cross-section belonging to the smaller radius of bending. This defect cannot be removed. If it is situated at a sufficient depth under the slab surface and if further processing of the slab is not too demanding, this defect need not present a cause of difficulties during such processing or a cause of a lower quality of the plate. A determination of the presence and form of trans- verse cracks as well as their quantification according to the methodology used in the given case are performed in the standard manner on macro-etches and/or on Baumann sulphur prints. Figure 3 shows an example of such manifestations of transverse cracks. 5 DETERMINATION OF THE CAUSES OF CRACK FORMATION For a determination of the cause, or causes, of the formation of internal transverse cracks on the concrete slab we first chose the parameters that might have caused the defect. With the LITIOS analytical software tool we then assigned the values of the selected casting para- meters to the evaluated sample and compared these values with the values defined, in the standard manner, for the given CCM as satisfactory, i.e., with the limits specified by the supplier of the equipment or by the CCM operator on the basis of long-term experience. For an exploration of the causes of the internal transverse cracks in the specific slab we chose the heat and the sample that was subjected to a metallographic evaluation. A transverse sample was taken, i.e., a sample from the plane perpendicular to the direction of casting. The sample was cut with oxygen-acetylene flame at the end of the cooling bed. For a determination of the possible degree of defects the sample was then polished in a metallographic laboratory, etched in order to see the nature of the macrostructure and then Baumann sulphur prints were made. The evaluation presented in the table below was carried out according to the commonly used standard methodology for evaluating cast slabs. Table 1 below presents the degrees of defects deter- mined on the Baumann print in the metallographic testing laboratory. Transverse internal cracks were visible on the sample. Using a 6-degree scale, from degree 0 denoting no defect to degree 5 denoting the largest defect, we determined the cracks of the 3rd degree, which already exceeded the limit of the harmlessness of the defect, see the Baumann sulphur print in Figure 4. Z. FRANÌK et al.: DETERMINATION OF THE CAUSE OF THE FORMATION OF TRANSVERSE INTERNAL CRACKS ... 288 Materiali in tehnologije / Materials and technology 49 (2015) 2, 285–290 Figure 3: Transverse internal cracks in the top half of a macro-etched slab sample; example from the catalogue of defects Slika 3: Notranje pre~ne razpoke v gornji polovici makro jedkanega vzorca slaba; primer iz kataloga napak Table 1: Evaluation of defects on the sulphur print Tabela 1: Stopnje napak, dolo~ene na Baumannovem odtisu v metalografskem laboratoriju Defect Pointinclusions Cluster inclusions Central segregation Lateral cracks Corner cracks Transverse internal cracks Longitudinal internal cracks Degree 2 0 2 2 0 3 NA The investigated heat and slab were produced using the following "fix" technical parameters that did not change during the casting of the given heat: • Steel grade – low-carbon, micro-alloyed with vana- dium, titanium and niobium • Chemical composition – 0.06 % C, 1.66 % Mn, 0.30 % Si, 0.016 % P, 0.005 % S, 0.02 % V, 0.03 % Nb, 0.005 % Ti • Liquidus temperature – 1516 °C • Slab dimensions – 180 mm × 1580 mm The values of the variable casting parameters pre- sented in Table 2 and assigned to the exactly defined section of the slab, from which the sample was taken for a metallographic investigation, were the following: Only one parameter, which directly specifies the degree of failure of its stipulated value, is contained in the values presented in this table. It is the non-achieve- ment of the minimum casting rate specified by the technological standard. In the investigated heat it was necessary to reduce the casting rate on the upstream units due to technological causes. For a possible compa- rison with the state before the slowdown of the casting rate and for an evaluation of a possible influence of the deviations of the values of individual parameters, we give, in the last row of the table, the average values of the monitored parameters, acquired from approximately 250 heats of the same slab format and the same steel grade. 6 RESULTS AND DISCUSSION On the basis of the presented values of the selected casting parameters and the occurrence of the internal transverse cracks on the slab it is possible to draw rather unequivocal conclusions: • The cause of the internal transverse cracks in the specific slab represented by a sample of the Baumann sulphur print and macro-etching, was a reduction in the casting rate. The value of the casting rate specified by the technological standard for the given steel grade of 1.20 m/min was reduced down to 0.74 m/min, indicating the overall drop in the casting rate by 38 %. • The reduction in the casting rate is also documented in the system of prediction by the monitored parameter of the non-achievement of the minimum casting rate – the value recorded in our case was minus 470 mm. • The reduced casting rate entailed a reduction in the surface temperature of the measured planes of the secondary cooling and, consequently, also a reduction in the density of the heat flux from the slab into the cooling media. Particularly the reduction in the surface temperature of the slab in the area of its straightening might have been the primary cause of the transverse cracks. • The steel grade also played its unfavourable role during the formation of the investigated cracks. Namely, the micro-alloying elements – vanadium, niobium and titanium – reduce the steel ductility already at approximately 950 °C. They form the so-called "trough" of the ductility reduced below this temperature value, which means that micro-alloyed steels are here more susceptible to a crack formation in continuously cast products. The paper describes one of the possible approaches of using the prediction system, or using the values of the casting parameters selected for this system, for an exact determination of the causes of the transverse internal cracks on a concrete slab. In the professional literature, we did not find any solutions or results for exactly this, or a similar, problem. On the other hand, it is possible to find studies or articles dealing with a prediction system and the use of this software as an analytical tool for an assessment of the continuous casting process and the quality of cast blanks. Z. FRANÌK et al.: DETERMINATION OF THE CAUSE OF THE FORMATION OF TRANSVERSE INTERNAL CRACKS ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 285–290 289 Figure 4: Transverse internal cracks on the sulphur print of the investigated sample Slika 4: Pre~ne notranje razpoke na Baumannovem odtisu preiskova- nega vzorca Table 2: Values of the parameters for the investigated sample and the average values for 250 heats Tabela 2: Vrednosti parametrov preiskovanih vzorcev in povpre~ne vrednosti iz 250 {ar` Parameter Casting rate(Vg)/(m min–1) Vg below the minimum limit Vg/(mm min–1) Overheating of steel above Tl in the tundish (°C) Type Seg 6 (°C) Type Seg 11 (°C) Heat-flux density (W m–2) Value 0.74 –470 35 915 842 989705 Average for 250 1.19 3 33 965 876 1086773 Their results are generally comparable with the use of the LITIOS analytical software tool, presented here. They are presented, for example, in11, which describes the strategies and methods for monitoring and controll- ing the quality during the production of flat steel. It gives a complete description of the quality, technological parameters and information technology. The data are archived and used in an appropriate manner for inter- vening at a specified place in order to improve the quality. The percentage of success is said to be 85 %. Reference12 gives an example of a complex system integrated in CSP at the ACB company, including the modules that collect and process the data, search and visualise them, draw up statistics and analyse the phenomena. Individual, special cases can be quickly solved with the use of this software. 7 CONCLUSIONS The cause of the transverse internal cracks in a continuously cast slab was determined with a complex system of a long-term monitoring of the casting para- meters using the LITIOS analytical software tool: • In a metallographic laboratory a cross-section sample taken from a continuously cast slab was evaluated. Some transverse internal cracks were found on it at a degree exceeding the harmlessness of such a defect. • The technologist determined the possible technolo- gical parameters of the casting that might have caused these cracks. • Subsequently, using the LITIOS analytical software tools, we assigned concrete variable casting values, corresponding exactly to the taken sample, to these parameters. • Only one technological parameter was found to be beyond the standard limits set for casting the relevant steel grade. This analysis clearly revealed the cause of the examined cracks. It was the forced reduction in the cast- ing rate to approximately 62 % of the value specified by the standard technological procedure. By using a similar procedure it is possible to deter- mine the technological cause of any defect found on the Baumann sulphur print or on a macro-etched sample taken from a continuously cast slab, and to perform interventions aiming at controlling the casting techno- logy, leading to a minimisation or complete elimination of the causes of the defect during the subsequent casting. By using the LITIOS software tool, together with examining the samples and determining their exact evaluation in the metallographic laboratory, we can efficiently contribute to improving the quality of cast slabs and, ultimately, to increasing the efficiency of the final products made from them. 8 REFERENCES 1 C. A. Santos, J. A. Spim, M. C. Flerardi, A. Garcia, The use of arti- ficial intelligence technique for the optimisation of process para- meters used in the continuous casting of steel, Applied Mathematical Modelling, 26 (2002) 11, 1077–1092, doi:10.1016/s0307-904x(02) 00062-8 2 C. A. Santos, J. A. Spim, A. Garcia, Mathematical modelling and optimization strategies (genetic algorithm and knowledge base) applied to the continuous casting of steel, Engineering Applications of Artificial Intelligence, 16 (2003) 5–6, 511–527, doi:10.1016/ S0952-1976(03)00072-1 3 F. Kavi~ka, Z. Franìk, J. [tìtina, Software Analytical Instrument for Assessment of the Process of Casting Slabs, Proceedings of the 10th International Conference on Numerical Methods in Industrial Forming Processes NUMIFORM 2010, Pohang, Republic of Korea, 2010, 586–592, doi:10.1063/1.3457607 4 V. Vaculik, R. B. MacCuish, R. K. Mutha, Multivariate statistical model-based system for monitoring the operation of a continuous caster and detecting the onset of impending breakouts, US Patent 6564119, 2003 5 Electrics and Automation for Continuous Casting – SIMETAL CC Control, Basic automation, Metals magazine, 1 (2014) 6 J. [tìtina, Z. Franìk, F. Kavi~ka, M. Masarik, V. Krol, Quality Opti- mization of Casting Slab via Mathematical and Statistical Method, DVD Proceedings of the 7th European Continuous Casting Confe- rence METEC InSteelCon, Düsseldorf, Germany, 2011, 193–200 7 T. Mauder, C. Sandera, J. Stetina, M. Seda, Optimization of The Quality of Continuously Cast Steel Slab Using the Firefly Algorithm, Mater. Tehnol., 45 (2011) 4, 347–350 8 T. Mauder, Z. Franek, F. Kavicka, M. Masarik, J. Stetina, A Mathe- matical & Stochastic Modelling of the Concasting of Steel Slabs, Proceedings of the 18th International Conference on Metallurgy and Materials, Hradec nad Moravici, 2009, 41–48 9 A. Ebel, J. Hackmann, N. Holzknecht, N. Link, H. Peters, Indus- trielles data mining in der stahlindustrie, Stahl und Eisen, 132 (2012) C.2, 29–37 10 J. [míd, Catalogue of defects of continuously cast slabs, Technologie und Metallurgie, Ltd., Business Studies, 2011, 125 p. 11 H. Peters, T. Heckenthaler, N. Holzknecht, Strategies and methods for quality monitoring and quality control in flat steel production, Stahl und Eisen, 126 (2005) C.7, 29–36 12 M. Reifferscheid, J. Kempken, M. Bruns, J. I. L. Garcia-Echave, J. M. Ovejero, Integrated Product Improvement by Quality Analysis and Modelling, Stahl und Eisen, 125 (2005) 12, 29–34 Z. FRANÌK et al.: DETERMINATION OF THE CAUSE OF THE FORMATION OF TRANSVERSE INTERNAL CRACKS ... 290 Materiali in tehnologije / Materials and technology 49 (2015) 2, 285–290 M. URBÁNEK, F. TIKAL: EFFECTIVE PREPARATION OF NON-LINEAR MATERIAL MODELS ... EFFECTIVE PREPARATION OF NON-LINEAR MATERIAL MODELS USING A PROGRAMMED OPTIMIZATION SCRIPT FOR A NURIMERICAL SIMULATION OF SHEET-METAL PROCESSING U^INKOVITA PRIPRAVA NELINEARNIH MODELOV MATERIALA S PROGRAMIRANIM OPTIMIZACIJSKIM ZAPISOM ZA NUMERI^NO SIMULACIJO OBDELAVE PLO^EVINE Miroslav Urbánek, Filip Tikal COMTES FHT a.s., Prumyslova 995, Dobrany, Czech Republic miroslav.urbanek@comtesfht.cz Prejem rokopisa – received: 2013-10-14; sprejem za objavo – accepted for publication: 2014-05-23 doi:10.17222/mit.2013.248 Progressive methods and technologies are the key to the dynamic development of the automotive and electrical-engineering industries. The processes in sheet-metal processing have changed fundamentally since the end of the 1990s. Previously, the operations such as cutting, punching holes, etc., were carried out separately on different press machines. These operations can now be integrated into a single tool on one press machine due to the development of progressive tools and, especially, the fine-blanking technology. The result is an already completed component that can be used for the assembly. The development of progressive tools must be supported with the FEM simulations of sheet-metal processing that are dependent on their inputs. Therefore, only a correct material model can be expected to provide the right results. For the reasons described above, an experimental program dealing with measuring and fitting the data to the models with ortho- tropic material properties such as rolled sheets was designed and implemented. The aim is to obtain the material models of rolled sheets made of selected aluminum, copper and steel alloys. One of the objectives of the solution is to provide more efficient and more accurate data fitting because the more accurate the input material data are, the more accurate are the simu- lation results for sheet-metal processing. The optimization script using the simplex method was made for fitting. The main function of the optimization script is to specify the parameters of the material model iteratively and to compare the simulation results and the mechanical-test results. The script was programmed in the Python environment for the MSC.MARC/MENTAT software using the Johnson-Cook plasti- city model. Fitting the data from the pressure tests by Rastegaev at different loading speeds is presented. The difference between the measured and simulated curves is less than 1 %. Keywords: FEM simulation, measurement, compression test, fitting, MSC.MARC, Python Napredne metode in tehnologije so klju~ne za dinami~en razvoj avtomobilske in elektro- industrije. Procesi preoblikovanja plo~evin so se bistveno spremenili od devetdesetih let zadnjega stoletja. Prej so se operacije, kot so rezanje, prebijanje lukenj in podobno, izvajale lo~eno na razli~nih strojih. Zaradi razvoja naprednih orodij in {e posebej tehnologije precizijskega {tancanja se te operacije lahko zdru`i v enem orodju na enem stroju. Rezultat je kompletna komponenta, ki je primerna za vgradnjo. Razvoj naprednih orodij mora biti podprt s FEM-simulacijami obdelave plo~evine, kar je odvisno od vhodnih veli~in. Zato samo pravilen model materiala lahko zagotavlja prave rezultate. Iz navedenih razlogov je bil postavljen in uporabljen eksperimentalni program, ki obravnava merjenje in ujemanje podatkov za modele z ortotropnimi lastnostmi materiala, kot je valjana plo~evina. Namen je dobiti modele materiala valjane plo~evine za izbrane aluminijeve in bakrove zlitine ter jekla. Eden od ciljev je zagotoviti bolj u~inkovito in bolj zanesljivo pridobivanje podatkov, kajti ~im bolj zanesljivi so vhodni podatki materiala, bolj natan~ni so rezultati simulacije pri obdelovanju plo~evine. Za prilagajanje je bila uporabljena optimizacija zapisa z uporabo simpleksne metode. Glavna vloga optimizacijskega zapisa je iterativna opredelitev parametrov modela materiala in primerjava rezultatov simulacije z rezultati mehanskih preizkusov. Zapis je bil programiran v okolju Python za programsko opremo MSC.MARC/MENTAT z uporabo Johnson-Cookovega modela pla- sti~nosti. Predstavljeno je ujemanje podatkov iz Rastegaevega tla~nega preizkusa pri razli~nih hitrostih obremenjevanja. Razlika med izmerjenimi in simuliranimi krivuljami je manj kot 1 %. Klju~ne besede: FEM-simulacija, merjenje, tla~ni preizkus, ujemanje, MSC.MARC, Python 1 INTRODUCTION The economic crisis dragging on since 2007 exerts an increasing pressure on the suppliers of the components for the automotive and electrotechnical industries. In general, the aim is to maintain a high quality of the com- ponents while reducing the inputs. The aim can only be achieved by deploying more advanced and less costly technologies that offer a higher efficiency. Such were also the reasons leading to the creation of the interna- tional EUREKA project aimed primarily at replacing the costly machining processes with lower-cost and more advanced manufacturing processes for making flat products. The existing fine-blanking process delivers precision sheet products with the tolerances of IT8-IT9. Fine-blanking combined with bulk forming (chamfering, stepping down and other operations) of metal-sheet components represents an advanced manufacturing approach. The production time for a part made with this technique is reduced to seconds, while the required accuracy and quality of the functional surfaces are maintained. A typical example of a flat product with Materiali in tehnologije / Materials and technology 49 (2015) 2, 291–295 291 UDK 519.61/.64:621.9 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 49(2)291(2015) spatial features, such as a cylindrical recess, is shown in Figure 1.1–3 Today’s professional literature around the world des- cribes various approaches to measuring material charac- teristics. They concern not only elastic-plastic materials, such as metal alloys, but also viscoelastic materials, such as rubber. Elastic-plastic models typically rely on a suit- able form of the Johnson-Cook flow-stress model, defined with Equation (1) representing the stress-strain curve. This model can provide a comprehensive descrip- tion of a material at various strain rates and tempera- tures.4,5 Previously, we performed fitting by hand when indi- vidual states were fitted, which provided us with more accurate results for a single load condition (a single temperature or strain rate), but for a comprehensive assessment of material properties this method is insuffi- cient and also time-consuming. In addition, the depen- dence of the material behavior on the states was not respected. Specifically, a material stiffens more during faster loading and it is more pliable at a higher tempe- rature. 2 PREPARATION OF THE MATERIAL MODELS Material characteristics are measured using standar- dized and other physical methods with the test machines that record the force and time or displacement (the plunger displacement). The force-versus-displacement plots are evaluated and transformed by means of formu- las to obtain the required quantities, such as stress, strain, elongation and others. In order to obtain the correct values of material cha- racteristics, it is advisable to perform a test simulating the actual manufacturing process in a simplified form: on a smaller scale, for instance. There are various factors having an impact on the quality of the measurement record, such as the measurement method, the test con- figuration, the quality of the sensors, the machine stiff- ness, the loading velocity and others.1–3 In the present project, a rapid and effective data- fitting procedure for various physical measurements was developed and implemented. Data fitting is in essence a fine tuning of the material-property data obtained with the mechanical tests. Data fitting comprises four main steps. The first one involves a physical measurement and an evaluation. The next step is smoothing where a smooth continuous polynomial curve is fitted to the jagged test plot. In the third step, parameters are estimated with an approximation. The last step consists of a numerical analysis under the conditions identical to those of the physical measurement. Using a pre-pro- grammed optimization script, the constants for the material model are sought. In the following sections, these steps for an effective material-data preparation will be described in greater detail. For most of the steps, the scripts were developed in the user-friendly Python environment. Most scripts were written with the aid of open source libraries facili- tating the preparation. The term physical measurement can denote a stan- dard compression or tension test. The COMTES FHT a. s. company generally deals with standard and non-standard measurements of specimens and functional items. For this reason, the optimization script used for finding the parameters of a material model should be versatile enough to apply to most of the measurements. The measurement data is evaluated and corrected for further processing, if required. In general, smooth curves are easier to work with, especially when using mathe- matical algorithms. The smoothing can be modified with a user-defined parameter, which alters the shape of the resulting curve. The time required is very short: at the order of seconds. Using the corrected curve obtained from the physical measurement under the quasi-static conditions at room temperature, some parameters can be estimated with the aid of the generally known formulas based on the volu- me conservation, the yield strength and the engineering strain.2,3,5 In addition, these estimated parameters are used as the input parameters for the optimization script controll- ing the numerical simulation. The optimization script is iterative (Figure 2) seeking the constants for a general material model according to Johnson-Cook. The model is defined with an equation comprising five constants derived from the material properties. Other terms of the M. URBÁNEK, F. TIKAL: EFFECTIVE PREPARATION OF NON-LINEAR MATERIAL MODELS ... 292 Materiali in tehnologije / Materials and technology 49 (2015) 2, 291–295 Figure 2: Simplex algorithm using the Johnson-Cook model of flow stress Slika 2: Algoritem simpleks, ki uporablja Johnson-Cookov model za napetost te~enja Figure 1: Example of a flat product with spatial features Slika 1: Primer plo{~atega izdelka s prostorsko mre`o equation (Equation (1)) depend on the boundary and initial conditions of the process (temperatures, strain rates) and their meanings are described below. In the previous step, constants A, B and n of the Johnson-Cook equation, which depend on the strain hardening and yield strength, were estimated. Generally, the equations of the Johnson-Cook model describe the material in all the states. The state change is defined with the value of the relevant constant. The A, B, C, m and n constants determine the material state at all the temperatures and rates:  v nA B C T T T = + ⋅ ⋅ + ⋅ ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ ⋅ − − ( ) ln   p room mel 1 1 t room− ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ T m (1) The Johnson-Cook model represents the flow-stress values describing the material behaviour under deforma- tion. Equation (1) consists of a product of the terms governed by the plastic strain, the strain rate and the temperature. The first term ( )A B n+ ⋅ p describes the beginning of the plastic-flow segment of the stress-strain curve where A is the yield strength (MPa). The plastic- flow segment depends on the B strain-hardening modu- lus (MPa), the plastic strain p and the strain exponent n. The second term ( ln(  ))1+ ⋅C  consists of a dimen- sionless coefficient of sensitivity to strain rate C, the strain-rate ration logarithm  (s–1) and the reference strain rate  (s –1), which is normally taken as  = 1 s –1. The last term of the equation 1− − − ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ ⎛ ⎝ ⎜ ⎞ ⎠ ⎟T T T T m room melt room , which describes the effect of the temperature consists of a ratio of two temperature differences where individual terms denote the current temperature T (°C), the room temperature Troom (°C) and the melting temperature Tmelt (°C).6,7 The optimization script relies on the simplex algo- rithm, which is a linear iterative method. The outcome of the simplex algorithm is the value of the objective function. The values of the previous and the current steps are compared. The objective function is the evaluation criterion for the optimization method that seeks the minimum value of the function. With the single-digit values (e.g., 5), the curves are virtually identical. By contrast, at higher values (such as 120), there are diffe- rences between them, leading to inaccurate results. The optimization involves comparing the measured data (or smoothed data) with the results of the numerical simulation (Figure 2). COMTES FHT a. s. uses the soft- ware solutions provided by MSC Software, which has been developing the finite-element-method-based soft- ware for numerical simulations for 50 years. The script relies on MSC.MENTAT as the pre- and post-processor and on MSC.MARC as the solver. The latter was deve- loped for strongly non-linear analyses.6–8 It is important to bear in mind that errors occur in both the physical measurement and the numerical simu- lation. In addition, when the material-model constants are corrected with regard to multiple measurements under varying conditions, the results cannot be identical for all the states. 3 DESCRIPTION OF THE MATERIAL The complex behaviour of the materials can be des- cribed in a simplified manner with a tension-test curve (Figure 3) which, generally, comprises three segments. The first linear segment of the tension-test plot is des- cribed with Hooke’s law (from 0 to A) given by the Poisson’s ratio and Young’s modulus. In the second seg- ment of the plot (from A to C), the plastic deformation begins to occur at the yield stress. Beyond that point, the plastic flow continues and the material strain hardens up to the ultimate tensile strength. This red segment is described with plasticity models. In this case, it is the Johnson-Cook model. Once the ultimate strength is exceeded, the plasticity of the material is used up and cracks develop, leading to a destruction of the test spe- cimen (from C to D). The last segment of the curve is modelled using damage criteria which describe the condition of the material through the widely known Cockroft-Latham model: max ∫ ≥dt C (2) 4 COMPRESSION TEST In a demonstration of the application, the Rastegaev compression test (Figure 4) is described as a test suitable from multiple standpoints. First, the test configuration allows the data to be measured correctly without any friction effects. Furthermore, the test can be simplified for the numerical simulation to a 2D axially symmetric problem (Figure 5), due to the axially symmetric test specimen. This type of analysis provides an excellent description of the specimen throughout the process using a relatively small number of elements. Thanks to the small number of elements, short computation times are M. URBÁNEK, F. TIKAL: EFFECTIVE PREPARATION OF NON-LINEAR MATERIAL MODELS ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 291–295 293 Figure 3: Graph of the tensile test Slika 3: Diagram nateznega preizkusa achieved. With the optimization of the material para- meters for the compression test, a single iteration step of the optimization script takes approximately 10 s, of which the calculation itself takes about 6 s. The remain- ing time is used for setting up and evaluating the calcul- ation. These times are, understandably, dependent on the computer-workstation hardware. With the above steps, this short data fitting becomes a comparatively rapid pro- cedure. The optimization iteration takes place involving multiple measurement-data sets in order to find the correct material constants. Naturally, the more measure- ment-data sets are available for various temperatures and speeds, the easier it is to find the material-model con- stants. The method was applied to fitting the measurement data for the 16MnCr5 manganese steel. The specimens of this material were taken along the forming direction and loaded at the rates of 2.4 m/s and 240 mm/s at room temperature. The plot (Figure 6) contains solid lines representing the measured curves and dash-and-dot lines showing the simulation curves for identical boundary conditions. Upon fine tuning, the Johnson-Cook model parameters were as follows: A = 554 MPa, B = 198 MPa, n = 0.00035, C = 0.0215 and m = 0.0335. Although the specimens were kept at room temperature, the effect of the temperature on the specimens this small (D0 = 8 mm and height H0 = 11.5 mm) and loaded at high strain rates is not negligible. 5 CONCLUSION Using the above-demonstrated data-fitting procedure for obtaining the material constants, the material models can be fine-tuned quickly and effectively. A smart choice of the test specimens simplifies the numerical simulation and allows the entire data-fitting process to be auto- mated. Thanks to the MSC.MARC/MENTAT environ- ment, the optimization script was developed in the PYTHON language and the curve-smoothing tools, so that the material data of the Johnson-Cook flow-stress model can be used in other FEM-based programs as well. The objective of the solution developed herein was to obtain as accurate the material data as possible for con- structing a numerical model of shearing and forming flat products with spatial features for orthotropic materials. The comparison graph (Figure 6) suggests that the material model is suitable for characterizing the spatial features in forming a flat product. The main result is the working optimization script for fitting, which efficiently searches for the material-model constants that will subsequently be used for the simula- tions of cold forming. With precise results, it is possible M. URBÁNEK, F. TIKAL: EFFECTIVE PREPARATION OF NON-LINEAR MATERIAL MODELS ... 294 Materiali in tehnologije / Materials and technology 49 (2015) 2, 291–295 Figure 6: Comparison of the numerical simulation and physical measurements Slika 6: Primerjava numeri~ne simulacije in fizikalnih meritev Figure 5: Axisymmetric model of the compression test in MSC.MARC/MENTAT in the initial and deformed states Slika 5: Osnosimetri~ni model tla~nega preizkusa v MSC.MARC/ MENTAT za~etnem in deformiranem stanju Figure 4: Specimen for the Rastegaev compression test Slika 4: Vzorec za Rastegaevov tla~ni preizkus to better describe the process of forming and to optimize it not only in terms of the material flow, but also in terms of the tools stress. The next step will be the preparation of a script for fitting for various mechanical tests and material-seeking models using other measurement types, such as tensile, compressive or shear tests. Acknowledgement The authors of this paper gratefully acknowledge the support from the EUREKA LF12009 project: Research and Development of a New Technology of Cold Preci- sion Forming as a Replacement for the Cutting Opera- tions. 6 REFERENCES 1 M. Urbánek, F. Tikal, Ur~ení koeficientù materiálových modelù pro tváøecí procesy, Hutnické listy, LXVI (2013) 4, 71 2 M. [paniel, A. Prantl, J. D`ugan, J. Rù`i~ka, M. Moravec, J. Ku- `elka, Calibration of fracture locus in scope of uncoupled ela- stic–plastic-ductile fracture material models, Advances in Engineer- ing Software, 72 (2014), 95–108, doi:10.1016/j.advengsoft.2013. 05.007 3 P. Kubík, F. [ebek, J. Petru{ka, J. Hùlka, J. Rù`i~ka, M. [paniel, J. D`ugan, A. Prantl, Calibration of Selected Ductile Fracture Criteria Using Two Types of Specimens, Key Engineering Materials, 592–593 (2013), 258–261, doi:10.4028/www.scientific.net/KEM. 592-593.258 4 R. A. Smidt, F. Bitzer, P. Höfer, M. Hellmann, B. Reh, P. Radema- cher, H. Hoffman, Cold Forming and Fineblanking, Druckhaus Thomas Mützen, Germany 2007 5 J. Dzugan, M. Spaniel, P. Konopík, J. Ruzicka, J. Kuzelka, Identi- fication of Ductile Damage Parameters for Austenitic Steel, World Academy of Science, Engineering and Technology, 6 (2012) 5, 1291–1296 6 J. D`ugan, M. Zemko, Input data influence on FEM simulation of steam turbine blades materials hot forming, Materials Science Forum, 773–774 (2013), 79–88, doi:10.4028/www.scientific.net/ msf.773-774.79 7 Marc® 2012, Volume A, Theory and User Information 8 Simplexová metoda, website, www.algoritmy.net/article/1416/Sim- plexova-metoda, accessed 14 Oct. 2013 M. URBÁNEK, F. TIKAL: EFFECTIVE PREPARATION OF NON-LINEAR MATERIAL MODELS ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 291–295 295 A. KRA^UN et al.: NEUTRALIZATION OF WASTE FILTER DUST WITH CO2 NEUTRALIZATION OF WASTE FILTER DUST WITH CO2 NEVTRALIZACIJA ODPADNEGA FILTRSKEGA PRAHU S CO2 Ana Kra~un1,2, Ivan An`el1, Lidija Fras Zemlji~1, Andrej Stergar{ek3 1University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, 2000 Maribor, Slovenia 2Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia 3Kemek, d. o. o., Cimpermanova ulica 3, 1000 Ljubljana, Slovenia ana.kracun@imt.si Prejem rokopisa – received: 2014-09-29; sprejem za objavo – accepted for publication: 2014-12-12 doi:10.17222/mit.2014.247 In this paper we report on the possibility of neutralizing filter dust from Talum Livarna d.o.o. The filter dust that remains after cleaning flue gas with the classification number of waste 10 10 09* is alkaline and contains heavy metals, non-metals, organic pollutants, and, therefore, has the properties of hazardous waste. The possibility of neutralizing this dust with CO2 was studied. The results showed that the treatment successfully lowered the pH value between the limits 6 and 9, which is within the legal constraints of pollution for strong acidic or alkaline waste. The contents of the hazardous substances were lowered, i.e., As, Cu, Ba, Zn, Cd, Cr, Ni, Pb, Sn, Mn and V, with percolation values that are below the level of the prescribed threshold-limit values for substances that allows their disposal in non-hazardous waste landfills. Only the percolation values of Sb, Cd, Mo and Se exceed the prescribed threshold limit values of substances that allow their disposal in inert waste landfills. The XRD analysis after the neutralization of the filter dust using CO2 showed no presence of CaO. The neutralized filter dust can be land filled as a stabilized and unreactive waste in landfills for nonhazardous wastes. Their properties also offer the possibility for incorporating them into some other material or product, such as the production of new composite materials, their use in construction products and perhaps cements or usage in backfills. Keywords: hazardous waste, filter dust, neutralization, stabilization, chemical properties Raziskali smo mo`nost nevtralizacije filtrskega prahu iz podjetja Talum Livarna, d. o. o. Filtrski prah po ~i{~enju dimnih plinov s klasifikacijsko {tevilko odpadka 10 10 09* je alkalen, vsebuje te`ke kovine, nekovine, organska onesna`evala, zato ima lastnosti nevarnega odpadka. Filtrski prah smo nevtralizirali s CO2, da je nastal prete`no amorfen produkt. Po obdelavi smo uspe{no zni`ali pH-vrednost v meje med 6 in 9, kar je v dovoljenem obmo~ju za odpadke, onesna`ene z mo~no kislino ali bazo. Prav tako je bila zmanj{ana vsebnost nevarnih snovi, in sicer As, Cu, Ba, Zn, Cd, Cr, Ni, Pb, Sn, Mn in V, tako da so izlu`evalne vrednosti pod mejo predpisanih parametrov izlu`ka in je tako dovoljeno odlaganje na odlagali{~ih za nenevarne odpadke. Samo izlu`evalne vrednosti Sb, Cd, Mo in Se {e prekora~ujejo predpisane mejne vrednosti, ki so dovoljene za odlaganje na odlaga- li{~ih za inertne odpadke. XRD-analiza po nevtralizaciji filtrskega prahu s CO2 ni pokazala prisotnosti CaO. Nevtraliziran filtrski prah se lahko odlo`i kot stabiliziran in nereaktiven odpadek na odlagali{~ih nenevarnih odpadkov. Glede na lastnosti obstaja mo`nost predelave in uporabe v koristne namene, npr. za proizvodnjo novih kompozitnih materialov, gradbenih izdelkov, morda cementa ali za zasipavanje. Klju~ne besede: nevarni odpadki, filtrski prah, nevtralizacija, stabilizacija, kemijske lastnosti 1 INTRODUCTION Hazardous wastes are a problem of modern civiliza- tion and therefore need to be handled in a prudent manner. A rapid increase in their amount, negative effects on the environment and a growing environmental awareness have led to changes in the field of waste management in recent decades. These factors have con- tributed to stricter regulations and the development of new technical and operational solutions.1,2 The strategy of waste management in Slovenia is directed towards actions that enable the overseeing, removal and reduction of the harmful effects of these wastes on the environment and humans, as well as their preparation for reuse, recycle and use as an energy source.2 The metallurgical industry is a major source of potentially hazardous waste materials, those by-products of the production of metals and alloys. These metallur- gical wastes consist mainly of slags and dust sludges that result from flue-gas filtering. These so-called waste materials can be potentially utilized as resources, for example, dust from the EAF process can be used as a source of Zn,3,4 metallurgical slags can be used in the production of building materials.5 A special aspect of the steel industry that requires particular attention is the pro- duction of stainless steel. During stainless-steel produc- tion, chromium oxidation occurs, which leads to the for- mation of CrOx phases. They not only represent a loss from the production point of view,6–8 but also represent an environmental risk, because they can oxidize to dan- gerous hexavalent chromium (Cr6+).9 During the melting of secondary aluminium contaminated with oil, paint or plastic flue gases that contain dust contaminated with heavy metals, nonmetals, dioxins, furans and fluorides.10 Waste filter dust is formed inside the exhaust gas treatment apparatus of foundry furnaces during operation with help of an additive DESOMIX HK. The additive DESOMIX HK is a mixture of calcium hydroxide [Ca(OH)2] and active chalk. Filter dust is classified as a hazardous waste with a classification number 10 10 09*, which puts it as far as regulation is concerned1 amongst the group of dusts that contain dangerous substances. Materiali in tehnologije / Materials and technology 49 (2015) 2, 297–301 297 UDK 658.567.5 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 49(2)297(2015) The filter dusts are homogenous, alkaline, contain heavy metals, nonmetals, dioxins, furans and fluorides. Be- cause the limit values of these toxic substances are exceeded during leeching, it is determined that these dusts have the properties of hazardous waste. Therefore, it has to be deposited in accordance with the relevant regulations.2 Industrial wastes are deposited in sub- terranean depots, usually prepared in closed coal and ore mines.11 There are many waste substances that react with carbon dioxide, for example, metallurgic slags, dusts from industrial thermical processes, combustion remains in which calcium oxide is a key component. The reaction between calcium hydroxide [Ca(OH)2] and carbon dio- xide (CO2) is called carbonization. It is a natural reaction between an oxide or a base and CO2.12,13 Direct binding between CO2 and calcium hydroxide is slow. However, it can be accelerated with the addition of water. In such cases the CO2 dissolves in water and dissociates and reacts with the dissolved and dissociated calcium hydro- xide, therefore forming carbonate materials.12,13 During the formation of these carbonate materials the pH of the waste dust decreases, which results in the minimal solu- bility of the metals present in the waste dust.12–14 Com- plete carbonization can decrease the leaching of metals by 80 %.14–16 Previous research14–28 has shown that metal stabilization processes are more efficient if the treatment of the waste dust is conducted over an extended period of time and at high CO2 concentrations. The main goal of this research was to determine whether it is possible to obtain a chemically stable pro- duct from waste filter dust with the process of neutra- lization using CO2. This would also enable this material to be reused in some other useful application or at least to process it enough so that long-term monitored depo- sition would be possible according to Slovenian regula- tions.2 Should the treated waste powder leaching values be inside the legal constraints for nonhazardous waste there would be a possibility that this waste could be deposited as a cost-reducing measure in dumps for nonhazardous waste. If it could achieve the constraints for inert waste it could be reused for practical purposes, for example, the production of new composite materials, use in construction materials, perhaps in cements or use in landfills. 2 EXPERIMENTAL Filter dust, a by-product of the treatment of exhaust gases, is classified as a hazardous waste with the classifi- cation number 10 10 09*. The analyzed sample comes from the release of a exhaust-gas treatment apparatus filter L9 in the company Talum Livarna, d. o. o. The granulometric composition of the filter dust and the average particle size was determined with a device for laser-diffraction particle size analysis Cilas 1064. A qualitative analysis of the sample microstructure was facilitated using a high-resolution Scanning Electron Microscope SIRION NC 400, equipped with NICA brand EDS detector. The X-ray analysis was conducted using a Philips XRD-analyzer W1800 hardware and X’pert High Score software. 2.1 Neutralization of the filter dust with CO2 The first step in this research was the determination of the initial pH of the leachate of the filter dust sample originating from Talum Livarna, d. o. o. Three samples with different moisture contents were prepared: • 11 % moisture content (sample label NFP1). • 8 % moisture content (sample label NFP2) • 14 % moisture content (sample label NFP3). The moisturized samples were stirred using a glass rod, which helped to achieve a uniform water consi- stency of the sample. A Rushton mixer and a supply tube for the CO2 were placed inside a glass beaker. The samples were exposed to a flow of 10 L/min of CO2 for a period of 1 h. The leaching and neutralization of the waste filter dust were carried out in accordance with the standard SIST EN 12457 – 4 (24 h leaching with water; ratio water/solid was 10/1). This was followed by a measu- rement of the pH and the leached inorganic parameters (metals) with help of the ICP-MS method. The ICP-MS device ionizes the sample with an inductively coupled plasma of argon gas. The ionized particles of the sample are then directed into a mass spectrometer. With the help of the ICP the presence of the following elements could be determined: Ag, Al, As, Ba, Ca, Cd, Co, Cr, Cu, Fe, Mn, Ni, Pb, Sb, Se, Sn, Sr, Zn, B, Be, Mo, Tl, V. 3 RESULTS AND DISCUSSION 3.1 Analysis of waste filter dust With the aid of laser diffraction we determined that the sizes of the filter particles range from 0.47 μm to 47.30 μm. Some 20 % of the particles are in the size range between 0.47 μm and 10 μm, while the remaining A. KRA^UN et al.: NEUTRALIZATION OF WASTE FILTER DUST WITH CO2 298 Materiali in tehnologije / Materials and technology 49 (2015) 2, 297–301 Figure 1: Size of the filter particulates Slika 1: Velikost delcev filtrskega prahu 80 % fall between 10 μm and 47.30 μm. Figure 1 shows the size of filter particulates. The SEM analysis of the sample dust showed that the morphology of the sample is not uniform. The particles were of different shapes, i.e., rod-shaped, spherical and asymmetrical (Figure 2). The particles also differed according to the chemical composition. The EDS analy- sis of the sample dust showed asymmetrical particles that had peaks for the elements: O, Al, Ca and Zn. The rod-shaped particles had peaks for the elements: O, Na, Mg, Al, Si, K and Ca. The spherical shaped had peaks for the elements: O, Si, Ti, Mg, Zn, Al, Si, Ca in K. Recurring elements with in all shapes were elements: O, Al and Ca. The results of the chemical analysis of the waste filter dust are presented in Table 1. Table 1: Levels of measured parameters in the filter dust Tabela 1: Vsebnost izmerjenih parametrov v filtrskem prahu Parameter Amount in the s. s. sample,mg/kg Al 2.1 Sb 45.1 As 10.0 Cu 540.2 Ba 3225.3 Zn 2501.6 Cd 24.7 Cr 154.7 Mo 15.4 Ni 128.0 Pb 282.3 Se 20.2 Co 2.1 Sn 130.9 Mn 130.3 Tl 0.6 V 15.4 The chemical analysis also revealed that the samples contained 3.5 % of moisture. 3.2 Filter-dust leachate analysis The filter dust leachate contains high levels of heavy metals, such as: Sb, Cu, Ba, Cd, Mo, Ni, Se. It contains chlorides and fluorides. Present are also As, Zn, Cr, Pb, Co, Sn, Mn, Tl and V (Table 2). In the presence of water these chemical elements could leach into the soil and damage the environment, if the waste dust were to be deposited freely. It can be seen by observing Table 2 that some quantities of leached heavy metals exceed the limiting values for dumping on sites equipped to handle inert waste. The leachate of filter dust has a highly alkaline pH value of 12.71. The data shows that the waste filter dust is of a heterogeneous composition and that it presents a hazard to the environment and it could have a great impact on the chemical and ecological balance, if not deposited correctly. 3.3 Analysis of filter-dust leachate neutralized with CO2 The aim of the neutralization experiment with CO2 was to evaluate the reduction of the pH values and the inorganic parameters and to attain, if it is possible, a chemically stable product that could be safely and harmlessly stored or deposited. The results show that the pH values could be reduced with this process, as shown in Figure 3. The leachate pH values for the samples NFP1, NFP2 and NFP3 were monitored for a period of 32 d. The che- mical stability of the neutralized samples was satisfac- tory and the pH values were between 6 and 9. The attained pH values confirm that the waste is not contami- nated with a strong acid or base. Because of the leachate exposure to CO2 found in the surrounding air, metal A. KRA^UN et al.: NEUTRALIZATION OF WASTE FILTER DUST WITH CO2 Materiali in tehnologije / Materials and technology 49 (2015) 2, 297–301 299 Figure 3: Filter dust leachate pH values of samples NFP1, NFP2, NFP3 Slika 3: pH-vrednosti izlu`kov vzorcev NFP1, NFP2 in NFP3 Figure 2: SEM analysis of filter dust: 1-rod-shaped particle, 2-sphe- rical particle, 3-asymmetrical particle Slika 2: SEM-analiza filtrskega prahu: 1-pali~ast delec, 2-sferi~en delec, 3-asimetri~en delec oxides bound to CO2 and there was an additional drop in the pH over a period of 32 d. In sample NFP1 the initial pH was 12.71, while after neutralization it dropped to 7.97, and after a period of 32 days it was measured at 7.30. Table 2 presents the leaching results for the chosen leachate parameters of the samples NFP1, NFP2, NFP3 and the waste filter dust, compared to environmental legal constraints. The neutralization reduced the leachate values of the As, Cu, Ba, Zn, Cr, Ni, Pb, Sn, Mn and V to a level that could enable this waste to be deposited as non-hazardous waste. The leachate values for the Sb, Cd, Mo and Se remained too high for this waste to be depo- sited as inert waste. The highest heavy-metal reduction occurred in sample NFP3 with 14 % moisture. This proves that the reaction of the carbonization is the fastest and most successful when the moisture is high and there is a lot of water present. During the neutralization waste filter dust in samples NFP1, NFP2 and NFP3 CO2 underwent a chemical reaction and bound to calcium hydroxide [Ca(OH)2] in the presence of moisture. The carbonate CaCO3 was formed. An XRD analysis showed no presence of CaO in any of the samples, which indi- cates that the carbonization was complete. The formed carbonates are less soluble and less alkaline than the oxides and hydroxides found in the filter dust. 4 CONCLUSIONS Neutralization helped to transform a hazardous waste into a non-hazardous waste, which could be deposited on sites for non-hazardous waste. CO2 neutralization pre- sents a viable technological solution for the stabilization of waste filter dust. The neutralized dust could be used as a bulk material. Acknowledgment The authors would like to thank dr. Marko Hom{ak and Talum In{titut d.o.o. for their technical support. 5 REFERENCES 1 Uredba o odlaganju odpadkov na odlagali{~ih, Uradni list RS, {t. 61/2011, str. 8857, Available from World Wide Web: http://www. uradni-list.si/1/content?id=104808 [7. 4. 2014] (in Slovene) 2 Uredba o odpadkih, Uradni list RS, {t. 103/2011, str. 13935, Available from World Wide Web: http://www.uradni-list.si/1/ content?id=106484 [27. 4. 2014] (in Slovene) 3 P. Drissen, A. Ehrenberg, M. Kuhn, D. Mudersbach, Recent Deve- lopment in Slag Treatment and Dust Recycling, Steel Research International, 80 (2009) 10, 737–745, doi:10.2374/SRI09SP055 4 C. Scharf, A. Ditze, Processing of Agglomerated Red Filter Dust in the Converter Operation from Metallurgical Point of View, Steel Research International, 84 (2013) 9, 917–925, doi:10.1002/srin. 201300126 5 V. Zalar Serjun, B. Mirti~, A. Mladenovi~, Evaluation of ladle slag as a potential material for building and civil engineering, Mater. Tehnol., 47 (2013) 5, 543–550 6 B. Arh, F. Vode, F. Tehovnik, J. Burja, Reduction of chromium oxides with calcium carbide during the stainless steelmaking process, Metalurgija, 54 (2015) 2, 368–370 7 J. Burja, F. Tehovnik, F. Vode, B. Arh, Microstructural characteriza- tion of chromium slags, Metalurgija, 54 (2015) 2, 379–382 8 J. Burja, F. Tehovnik, J. Medved, M. Godec, M. Knap, Chromite spinel formation in steelmaking slags, Mater. Tehnol., 48 (2014) 5, 753–756 9 G. Albertsson, L. Teng, B. Bjorkman, S. Seetharaman, F. Engstrom, Effect of Low Oxygen Partition in CaO-MgO-SiO2-Cr2O3-Al2O3 Synthetic Slag at Elevates Temperatures, Steel Research Inter- national, 84 (2013) 7, 670–679, doi:10.1002/srin.201200214 A. KRA^UN et al.: NEUTRALIZATION OF WASTE FILTER DUST WITH CO2 300 Materiali in tehnologije / Materials and technology 49 (2015) 2, 297–301 Table 2: Leachate results for chosen parameters in the samples NFP1, NFP2, NFP3 and filter dust – compared to limiting legal constraints Tabela 2: Vrednosti izlu`evanja nevarnih snovi vzorca NFP1, NFP2, NFP3 in filtrskega prahu - primerjava z zakonodajnimi vrednostmi Parameter Filter dust leachate result mg/kg CO2 NFP1 neutralized filter dust leachate result mg/kg CO2 NFP2 neutralized filter dust leachate result mg/kg CO2 NFP3 neutralized filter dust leachate result mg/kg Limiting legal constraint of a parameter L/S = 10 L/kg (mg/kg) Inert waste Hazardouswaste Non-hazardous waste Sb 0.26 0.34 0.26 0.31 0.06 5 0.7 As 0.01 0.01 0.01 0.01 0.5 25 2 Cu 2.63 12.99 1.95 0.38 2 100 50 Ba 73.17 5.19 4.71 4.86 20 300 100 Zn 3.41 0.77 0.38 0.37 4 200 50 Cd 0.18 0.21 0.11 0.11 0.04 5 1 Cr 0.01 0 0 0 0.5 70 10 Mo 0.49 0.77 0.73 0.65 0.5 30 10 Ni 0.44 0.20 0.15 0.11 0.4 40 10 Pb 0 0.05 0.01 0.04 0.5 50 10 Se 0.19 0.36 0.28 0.28 0.1 7 0.5 Co 0 0.01 0 0 Sn 0.05 0.01 0.01 0.01 Mn 3.73 1.92 1.57 1.51 Tl 0 0.04 0.03 0.02 V 1.00 0.31 0.23 0.10 10 A. Grochowalski, C. Lassen, M. Holtzer, M. Sadowski, T. Hudyma, Determination of PCDDs, PCDFs, PCBs and HCB Emissions from the Metallurgical Sector in Poland, Env Sci Pollut Res., 14 (2007) 5, 326–332, doi:10.1065/espr2006.05.303 11 J. A. Roether, D. J. Daniel, D. Amutha Rani, D. E. Deegan, C. R. Cheeseman, A. R. Boccaccini, Properties of sintered glass-ceramics prepared from plasma vitrified air pollution control residues, Journal of Hazardous Materials, 173 (2010) 1/3, 563–569, doi:10.1016/ j.jhazmat.2009.08.123 12 G. Montes-Hernandez, R. Perez-Lopez, F. Renard, J. M. Nieto, L. Charlet, Mineral sequestration of CO2 by aqueous carbonation of coal combustion fly-ash, Journal of Hazardous Materials, 161 (2009), 1347–1354, doi:10.1016/j.jhazmat.2008.04.104 13 J. Jianguo, C. Maozhe, Z. Yan, X. Xin, Pb stabilization in fresh fly ash from municipal solid waste incinerator using accelerated carbonation technology, Journal of Hazardous Materials, 161 (2009), 1046–1051, doi:10.1016/j.jhazmat.2008.04.051 14 H. Ecke, H. Sakanakura, T. Matsuto, N. Tanaka, A. Lagerkvist, State-of-the-art treatment processes for municipal solid waste incineration residues in Japan, Waste Management & Research, 18 (2000) 1, 41–51, doi:10.1177/0734242X0001800106 15 H. Ecke, N. Menad, A. Lagerkvist, Carbonation of municipal solid waste incineration fly ash and the impact on metal mobility, J. Environ. Eng. ASCE, 129 (2003) 5, 435–440, doi:10.1061/ (ASCE)0733-9372(2003)129:5(435) 16 H. Ecke, N. Menad, A. Lagerkvist, Treatment-oriented characteriza- tion of metal-bearing dry scrubber residue from municipal solid waste incineration (MSWI), Journal of Material Cycles and Waste Management, 4 (2002) 2, 117–126 17 H. Ecke, H. Sakanakura, T. Matsuto, N. Tanaka, A. Lagerkvist, Effect of electric arc vitrification of bottom ash on the mobility and fate of metals, Environmental Science & Technology, 35 (2001) 7, 1531–1536, doi:10.1021/es0001759 18 T. T. Eighmy, B. S. Crannell, L. G. Butler, F. K. Cartledge, E. F. Emery, D. Oblas, J. E. Krzanowski, J. D. J. Eusden, E. L. Shaw, C. A. Francis, Heavy metal stabilization in municipal solid waste com- bustion dry scrubber residue using soluble phosphate, Environmental Science & Technology, 31 (1997) 11, 3330–3338, doi:10.1021/ es970407c 19 E. Rendek, G. Ducom, P. Germain, Carbon dioxide sequestration in municipal solid waste incinerator (MSWI) bottom ash, J. Hazard. Mater., 128 (2006) 1, 73–79, doi:10.1016/j.jhazmat.2005.07.033 20 T. Van Gerven, E. Van Keer, S. Arickx, M. Jaspers, G. Wauters, C. Vandecasteele, Carbonation of MSWI-bottom ash to decrease heavy metal leaching, in view of recycling, Waste Management, 25 (2005) 3, 291–300, doi:10.1016/j.wasman.2004.07.008 21 M. Fernandez Bertos, X. Li, S. J. R. Simons, C. D. Hills, P. J. Carey, Investigation of accelerated carbonation for the stabilisation of MSW incinerator ashes and the sequestration of CO2, Green Chem., 6 (2004) 8, 428–436, doi:10.1039/B401872A 22 X. Li, M. Fernandez Bertos, C. D. Hills, P. J. Carey, S. Simon, Acce- lerated carbonation of municipal solid waste incineration fly ashes, Waste Management, 27 (2007), 1200–1206, doi:10.1016/j.wasman. 2006.06.011 23 G. Cappai, S. Carab, A. Muntoni, M. Piredda, Application of acce- lerated carbonation on MSW combustion APC residues for metal immobilization and CO2 sequestration, Journal of Hazardous Materials, 207–208 (2012), 159–164, doi:10.1016/j.jhazmat.2011. 04.013 24 T. Sicong, J. Jianguo, Z. Chang, Influence of flue gas SO2 on the toxicity of heavy metals in municipal solid waste incinerator fly ash after accelerated carbonation stabilization, Journal of Hazardous Materials, 192 (2011), 1609–1615, doi:10.1016/j.jhazmat.2011. 06.085 25 L. Wang, Y. Jin, Y. Niea, Investigation of accelerated and natural carbonation of MSWI fly ash with a high content of Ca, Journal of Hazardous Materials, 174 (2010), 334–343, doi:10.1016/j.jhazmat. 2009.09.055 26 P. Gunning, C. Hills, A. Antemir, P. Carey, Novel approaches to the valorisation of ashes using aggregation by carbonation, Proc. of the 2nd International Slag Valorisation Symposium, Leuven, 2011, 103–116 27 H. Katsuura, T. Inoue, M. Hiraoka, S. Sakai, Full-scale plant study on fly ash treatment by the acid extraction process, Waste Manage- ment, 16 (1996) 5/6, 491–499, doi:10.1016/S0956-053X(96)00091-8 28 A. J. Chandler, T. T. Eighmy, J. Hartlén, O. Hjelmar, D. S. Kosson, S. E. Sawell, H. A. van der Sloot, J. Vehlow, Municipal Solid Waste Incinerator Residue, Studies in Environmental Science 67, Elsevier Science B. V., Netherlands 1997 A. KRA^UN et al.: NEUTRALIZATION OF WASTE FILTER DUST WITH CO2 Materiali in tehnologije / Materials and technology 49 (2015) 2, 297–301 301 B. [U[TAR[I^ et al.: THE INFLUENCE OF THE MORPHOLOGY OF IRON POWDER PARTICLES ... THE INFLUENCE OF THE MORPHOLOGY OF IRON POWDER PARTICLES ON THEIR COMPACTION IN AN AUTOMATIC DIE VPLIV MORFOLOGIJE DELCEV @ELEZOVEGA PRAHU NA NJEGOVO SPOSOBNOST ZA AVTOMATSKO ENOOSNO STISKANJE Borivoj [u{tar{i~1, Matja` Godec1, ^rtomir Donik1, Irena Paulin1, Sre~ko Glode`2, Marko [ori2, Milan Ratej3, Nada Javornik3 1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia 2University of Maribor, FNM, Koro{ka cesta 160, 2000 Maribor, Slovenia 3UNIOR, Forging Industry, Kova{ka cesta 10, 3214 Zre~e, Slovenia borivoj.sustarsic@imt.si Prejem rokopisa – received: 2014-10-24; sprejem za objavo – accepted for publication: 2014-12-16 doi:10.17222/mit.2014.273 Fe- and steel-based powder metallurgy (P/M) products, such as steel gears, spurs, locking mechanisms, porous filters, sliding bearings and bushes, as well as other machine parts and structural elements, are mainly produced with the so-called conventional sintering technology. It is the most efficient technology for the mass production of small, complex, functional and structural parts. Therefore, it is the most convenient and popular among all of the P/M technologies. The most important end-user of sintered parts is the automotive industry. However, small, complex, sintered parts can also be frequently used in the furniture and household industries, precise mechanics, articles for recreation and sports. A fine, iron-based powder mixture or prealloyed powder is first automatically uniaxial-die compacted (ADC) into the final shape of the product with a mechanical or hydraulic press and then sintered in a protective atmosphere at approximately 1100 °C. The metal powder mixture must have the appropriate engineering properties given by the chemistry and particle morphology, enabling a fast and reliable die-compaction process. The most important are a high tap density, a good powder flowability and a low compressibility. All this gives the green compacts an appropriate final shape with a smooth surface, a relatively high and uniform green density, as well as a green strength without internal flaws and cracks. In the case of very small two-or-more-heights products, for example, spur gears with a low module, it is very difficult to obtain a uniform green density at acceptable compaction pressures. Often small cracks are formed at height crossings and big differences in the green density appear in smaller or thinner regions. In the frame of our investigation we analysed the influence of the selected prealloyed commercial iron powder’s morphology and its technological properties on automatic die compaction, as well as the sintering process in the case of small two-level sintered gear dimensions of 5/40–7/10 × 7 mm with module m = 0.5. The original iron powder was sieved and the finest powder particle fraction (< 45 μm) was compared with the original powder mixture considering ADC and sintering process. It was found that the selection of the finer powder mixture could not contribute to the improvement in the overall ADC process, as well as a better green compact. In the present paper the results of our investigations are presented and the reasons why a finer powder mixture cannot contribute much to an improvement of the conventional sintering process. Keywords: Fe-based alloy powders, morphology and microstructure of particles, influence on automatic die compaction, sintering Jekleni izdelki, izdelani po P/M-postopkih, kot so npr. zobniki, mehanizmi za zaklepanje, porozni filtri, drsni le`aji in pu{e, kakor tudi drugi strojni elementi ali strukturni deli, se v glavnem izdelujejo po konvencionalni tehnologiji sintranja. To je najbolj u~inkovita in zato najpopularnej{a tehnologija za masovno proizvodnjo majhnih kompliciranih funkcionalnih in strukturnih izdelkov. Avtomobilska industrija je najbolj pomemben kon~ni uporabnik sintranih delov, vendar se le-ti pogosto uporabljajo tudi v pohi{tveni industriji, beli tehniki, precizni mehaniki, v izdelkih za {port in rekreacijo itd. Fina pra{na me{anica na osnovi Fe ali predlegiranih jeklenih delcev se najprej avtomatsko enoosno stiska v kon~no obliko na mehanskih ali hidravli~nih stiskalnicah in nato sintra v za{~itni atmosferi pri pribl. 1100 °C. Kovinska pra{na me{anica mora imeti primerne tehnolo{ke lastnosti, ki omogo~ajo hiter in zanesljiv proces zgo{~evanja v surovce `elene oblike in velikosti. Te so definirane s fizikalno-kemijskimi lastnostmi in predvsem morfologijo (obliko in velikostjo) delcev. Najbolj pomembne so visoka nasipna gostota, dobra teko~nost in nizka stisljivost pra{ne me{anice. Vse to daje surovemu izdelku `eleno kon~no obliko z gladko povr{ino, relativno visoko in enakomerno zeleno gostoto in trdnostjo brez lokalnih napak in razpok. Pri zelo majhnih dvo- ali ve~vi{inskih izdelkih, kot so npr. dvojni zobniki z majhnim zobni{kim modulom, je zelo te`ko dobiti surovce (zelence) brez napak in z enakomerno gostoto pri sprejemljivih tlakih stiskanja. Pogosto na surovcih nastajajo razpoke na prehodu iz tanj{ega v debelej{i del zobnika zaradi velikih razlik v zeleni gostoti. V okviru na{ih raziskav smo analizirali vpliv morfologije izbranega predlegiranega komercialnega Fe-prahu in njegovih tehnolo{kih lastnosti na razmere pri avtomatskem enoosnem stiskanju in sintranju majhnega dvovi{inskega sintranega zobnika dimenzij 5/40–7/10 × 7 mm z modulom m = 0,5. Originalni Fe-prah smo presejali in najfinej{o frakcijo prahu (< 45 μm) uporabili za stiskanje in sintranje ter jo primerjali z originalno pra{no me{anico. Ugotovili smo, da finej{i prah ne prispeva k izbolj{anju postopka avtomatskega enoosnega stiskanja oz. k bolj{im surovcem iz ve~ razlogov. V tem prispevku predstavljamo rezultate na{ih raziskav in razloge, zakaj izbira finej{e pra{ne me{anice ne prispeva celovito k izbolj{anju konvencionalnega sintrnega postopka. Klju~ne besede: legirani prahovi na osnovi `eleza, vpliv morfologije in mikrostrukture na avtomatsko enoosno stiskanje, sintranje Materiali in tehnologije / Materials and technology 49 (2015) 2, 303–309 303 UDK 621.762:621.77 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 49(2)303(2015) 1 INTRODUCTION The conventional sintering technology is the most popular among the all P/M technologies. It enables the large-scale production of small complex parts (Figure 1) with the lowest raw-material and energy consumption.1–3 This technology is used in the Unior forging industry, Sinter workshop, Zre~e, Slovenia, mainly for the produc- tion of sintered steel parts for the automotive industry.4 Sinter technology consists of three main production phases: powder manufacture and the preparation of the mixture, automatic (uniaxial cold) die compaction (ADC) and sintering (Figure 2). The improvement of the dimensional tolerances and the mechanical properties can be obtained with te additional post-sintering ope- rations, i.e., sizing, surface and heat treatment and machining. Some geometrical limitations exist in the phase of product design and later in the phase of tool manufacture when considering uniaxial automatic die compaction. These limitations have their origin in the nature of uniaxial ADC. Namely, the forces (pressures) of com- paction are not uniformly transferred over the height and cross-section of the formed green compact because of the internal friction among the powder particles and the friction on the die walls (Figure 3). This has already been shown by the classic analysis1 introducing an equa- tion where one can see that the transferred compaction force depends not only on the internal friction and the die-wall friction, but also on the ratio between the height and the diameter of the green compact (h/D): F F ex zx D= ⋅ −0 4  / (1) In Equation (1) Fx is the resulting force at a distance x from the upper punch, F0 is the acting (compaction) force on the upper punch, μ is the die-wall friction coefficient and the factor z describes the ratio between B. [U[TAR[I^ et al.: THE INFLUENCE OF THE MORPHOLOGY OF IRON POWDER PARTICLES ... 304 Materiali in tehnologije / Materials and technology 49 (2015) 2, 303–309 Figure 2: Schematic presentation of the phases of the conventional sintering process for the production of steel parts Slika 2: Shemati~ni prikaz postopka izdelave sintranih jeklenih izdelkov Figure 3: Schematic presentation and analysis of the uniaxial cold die compaction of metal powder in a cylindrical die6 Slika 3: Shemati~ni prikaz in analiza hladnega enoosnega stiskanja prahu v cilindri~nem orodju6 Figure 1: Some typical sintered steel parts Slika 1: Nekaj tipi~nih oblik sintranih jeklenih izdelkov the normal and the powder-transferred radial stresses, which depends on the internal friction among the powder particles. From Equation (1) we can calculate the transferred force at any distance x in the formed green powder com- pact, as well as the lowest transferred force on the bottom punch (at x = h) if the die and powder charac- teristics μ and z are known. The higher is the ratio h/D, the larger are the local differences in the green density and the formed green compact has an unequal green den- sity over its volume (Figure 4). The higher is the com- paction pressure, the higher is the local and overall (average) green density of the compact. The green density distribution can be mitigated if the compaction force acts from both sides (top and bottom) when forming the green part. Therefore, modern tools for ADC consist of a large number of parts; their move- ment is programmed and controlled by a computer on hydraulic or mechanical presses (Figure 5) in order to avoid too large differences in the green density of the compact and its uniform ejection out of the die. In spite of this, it is not possible to ensure that a green compact has a completely uniform green density over the whole volume (Figure 4), especially, if in the product design phase, it is not possible to avoid larger height differences, sharp crossings and chamfers, be- cause of other functional limitations of the final sintered product.5 The results of an unsuitable geometry of the product are: non-uniform powder filling of the die, large local compaction pressures, forming large local green- density differences, and finally cracking of the green compact during ejection. However, these also lead to wear/fracture of the most loaded tool parts and overall shorten the life of the tool. Tools (dies) for the compac- tion of metal powders are very precise, made of ad- vanced tool steels and cemented carbides, and therefore, their manufacture is very complex and expensive. Tool life depends not only on its complexity but also on pow- der engineering (technological) properties. Metal pow- ders have to have a large tap density, good flowability and compressibility for the appropriate ADC. One such sintered steel product that has a difficult ADC geometry is the two-height small gear produced in the Unior factory (Figure 6). It also has a very small gear module (m = 0.5) as well as large height and dia- meter differences. B. [U[TAR[I^ et al.: THE INFLUENCE OF THE MORPHOLOGY OF IRON POWDER PARTICLES ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 303–309 305 Figure 6: a) Schematic presentation of small and large gear module and b) a photograph of the discussed small complex double spur gear with the dimensions 5/40-7/10 × 7 mm, produced in the Sinter Work- shop, Unior factory, Zre~e, Slovenia Slika 6: a) Shemati~ni prikaz ozobljenja z majhnim in velikim mo- dulom in b) posnetek obravnavanega majhnega dvojnega sintranega zobnika 5/40–7/10 × 7 mm, ki ga izdelujejo v obratu Sinter, Unior, Zre~e Figure 5: a) Modern hydraulic computer-controlled press for the auto- matic die compaction of steel powders and b) a schematic presentation of a multi-platen adapter system for the fixation of the tool into the press5 Slika 5: a) Moderna hidravli~na ra~unalni{ko vodena stiskalnica za avtomatsko enoosno stiskanje jeklenih surovcev in b) shemati~ni pri- kaz adapterja z ve~plo{~nim sistemom vpetja orodja v stiskalnico5 Figure 4: a) Schematic and b) practical presentation of the green-den- sity distribution in a cylindrical specimen after the uniaxial cold compaction of steel powder5 Slika 4: a) Shemati~ni in b) prakti~ni prikaz porazdelitve zelene gostote v cilindri~nem surovcu po enoosnem stiskanju jeklenega prahu v orodju5 This small gear is compacted with very high com- pacting pressures in order to decrease the differences in the green density in the gear teeth because of poor filling of the engraving, as well as to decrease the differences in the green density between the gear parts with a large height difference to avoid cracking at the height cross- ing. This demands high compaction pressures over 700 MPa. The result is a too short tool-life because of fre- quent fracture of the most loaded tool parts (punches and core rods, Figure 7). Different solutions (better tool materials, more precise tooling, optimization of the die and press set up) have been researched to solve this problem. But no one has found a complete result and a final solution. There- fore, we also tried with a change of the existing powder granulometry. The hypothesis was that the selection of a finer powder could offer better compressibility and filling of the die. Unfortunately, as it follows, the change of the granulometry to a finer powder also did not give an improvement of the ADC process but gave us a lot of useful and interesting information. 2 EXPERIMENTAL WORK For the production of the investigated two-height gear a standard commercial diffusion prealloyed Fe-based powder Distaloy AB, Höganäs, Sweden was used. Its average nominal bulk chemical composition in mass fractions (w/%) is: 1.7 Ni, 1.5 Cu, 0.5 Mo, and the rest is Fe. The addition of carbon (generally 0.4–0.6 % gra- phite) changes it during the sintering into a steel with the required chemical composition. The 5 kg of original Distaloy AB powder was sieved on a set of vibrating sieves in the frame of our experimental work. The finest powder fraction (< 45 μm) was selected for our sub- sequent experiments and investigations. The compressibility of the selected fine (< 45 μm) and rough (> 45 μm) mixtures was determined by instru- mented apparatus6 in standardized die dimensions of  24 mm × 16 mm. The experiment for the compressibility determination is performed at a ram speed of 10 mm/min. It is a much slower speed than the actual indu- strial ADC process. Therefore, the densification and deformation rate of the green compact in industrial conditions are different and higher (a larger number of structural defects affecting the sintering), respectively. The flowability and tap density of the selected powder mixtures were determined with a Hall flowmeter1. The prescribed amount of graphite (w = 0.5 %) and Kenolube lubricant (w = 0.9 %) were added to the original Distaloy AB and the fine sieved powder and both were then homogenized in a double-cone mixer. The experimental compaction of the gears was performed on an industrial 60 kN Dorst, Germany, mechanical press. Approxi- mately 100 gear pieces were compacted from both pow- der mixtures followed by the sintering of green compacts in an industrial continuous-belt Mahler furnace under standard sintering conditions (1120 °C/30 min) in a pro- tective atmosphere (N2 + 5/10 % H2). The sintered gears were additionally heat treated after sintering (oil quench- B. [U[TAR[I^ et al.: THE INFLUENCE OF THE MORPHOLOGY OF IRON POWDER PARTICLES ... 306 Materiali in tehnologije / Materials and technology 49 (2015) 2, 303–309 Figure 7: Schematic presentation of the cross-section of the tool for the uniaxial ADC of the double-height product5 Slika 7: Shemati~ni prikaz prereza sestave orodja za stiskanje dvo- vi{inskega izdelka5 Figure 8: SEM micrograph of: a) fine (sieved < 45 μm) and b) large (> 45 μm sieve residual) powder mixture Distaloy AB; magnification 200-times Slika 8: SEM-posnetka: a) fine (presejane na < 45 μm) in b) grobe (ostanek po sejanju > 45 μm) frakcije kovinskega prahu Distaloy AB; pove~ava 200-kratna ing from 890 °C and tempering at 200 °C/30 min) The Vickers hardness HV5 of the sintered and heat-treated gears and the mechanical moment (teeth strength) were determined. The local bulk and micro-chemical compo- sitions of the powders, green compacts and sintered gears were determined with an SEM/EDS (Jeol – JSM6500F/Oxford INCA ENERGY 450, INCA X-SIGHT LN2) and an XRF analyzer (Thermo Scienti- fic, Niton XL3t Goldd+). 3 RESULTS AND DISCUSSION Figure 8 shows scanning electron micrographs of the fine and rough fractions of the investigated powders. The powders do not have large differences in morphology (shape and surface state), with the exception of the par- ticle size. However, micro-chemical SEM/EDS analyses have shown that the local chemical composition of the fine fraction is significantly different compared to the original mixture. The most probable reason is the me- thod of powder alloying. The used Distaloy AB powder is diffusion prealloyed (Figure 9b) and segregation of the alloying elements occurred during sieving and finer powder particles have a different chemical composition than the larger ones. Figure 10 shows EDS micro-che- mical mapping analyses of the fine and large powder particles. It is clear that their local compositions are quite different. This was also confirmed by the XRF analyses, which included a much larger volume of analyzed sample. In spite of this, the local chemical compositions of all the samples differ significantly from the nominal chemical composition of the Distaloy AB powder. Table 1 shows the average chemical compositions of all the analyzed samples. It is clear that the fine powder mix- tures have a much higher content of alloying elements than the original powder mixture. As will be shown later, this over-alloying also has a significant influence on the mechanical properties of the sintered and heat-treated gears. Table 2 shows the results of the technological pro- perties of the original, fine and rough powder mixture. We can see from this table that the fine fraction mixture has a poorer flowability, a lower tap density (n) and a negligibly better compressibility (z at psr). Our hypo- thesis was that the finer powder mixture has a better ability to fill the die cavity, but obviously the experi- ments disprove this. In this way the original powder mixture has a signifi- cantly better flowability, a higher tap density and a negligibly lower compressibility, and is therefore more suitable for ADC. This was also confirmed by our indu- B. [U[TAR[I^ et al.: THE INFLUENCE OF THE MORPHOLOGY OF IRON POWDER PARTICLES ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 303–309 307 (a) Fe Ni Cu Mo Spectrum 1 92.54 4.01 1.71 1.74 Spectrum 2 98.96 0.41 0.63 0.00 (b) Fe Ni Cu Mo Spectrum 1 97.51 0.81 0.71 0.97 Spectrum 2* 87.27 0.76 9.94 2.03 * not designated in the micrograph Figure 10: Mapping SEM/EDS microanalyses of: a) fine and b) large powder particles in mass fractions, w/% Slika 10: Primer ploskovne SEM/EDS-mikroanalize: a) finega in b) grobega delca v masnih dele`ih, w/% Figure 9: Standard ways of alloying techniques in powder metallurgy5 Slika 9: Standardni na~ini legiranja v metalurgiji prahov5 Table 1: Average XRF bulk chemical compositions of analyzed samples in mass fractions, w/% Tabela 1: Povpre~ne XRF kemijske sestave analiziranih vzorcev v masnih dele`ih, w/% No. Sample Fe Ni Cu Mo Mn Cr Si 1 Fine powder 74.14 14.38 9.99 0.84 0.09 0.05 0.13 2 Green compactfine powder 83.03 9.73 5.91 0.78 0.09 0.04 0.12 3 Sintered gear finepowder 88.92 5.47 4.42 0.80 0.20 0.05 0.08 4 Green compactoriginal powder 90.57 4.72 3.53 0.60 0.10 0.06 0.17 5 Sintered compactoriginal powder 93.65 3.10 2.11 0.65 0.13 0.06 0.12 strial experiments of the gear compaction. The powder mixture better filled the die cavity and a higher average green density (approx. 7.0–7.1 g/cm3) of the gears at lower compaction pressures (approx. 180 kN) are ob- tained. The fine powder mixture did not fill the die cavity so well and a lower average green density (approx. 6.9–7.0 g/cm3) of the gears at higher compaction pressures (approx. 210 kN) were obtained. Besides this, the gears made of the fine powder mixture have poorer mechanical properties after sintering and heat-treatment (Table 3). Figures 11 and 12 show the microstructures of the sintered and heat-treated gears. Figures 11a and 11b show a typical microstructure of a polished sample in the region of the tooth-root of the sintered gear visible under a light microscope (LM). It is clear that the sintered gear made of the original mixture has a larger fraction of large pores, but it is better densified in the gear core. This could be a problem of gear resistance to wear and fatigue. On the other hand, the gear made of the finer powder mixture has well-distributed, finer pores, but it is much less densified. Figures 12a and 12b show typical microstructures of polished and etched samples of the gears after sintering B. [U[TAR[I^ et al.: THE INFLUENCE OF THE MORPHOLOGY OF IRON POWDER PARTICLES ... 308 Materiali in tehnologije / Materials and technology 49 (2015) 2, 303–309 Table 2: Technological properties of the steel powder Distaloy AB, Höganäs Tabela 2: Tehnolo{ke lastnosti kovinskega prahu Distaloy AB, Höganäs Powder type n/(g/cm3) Flowabilitys/50 g pmax/ MPa psr/MPa z/(g/cm 3) Remarks Original powder* 3.06 26 7.10 g/cm3 at 600 MPa lubricant Kenolube Fine fraction (< 45 μm) 2.75 30 722.6 595.4 7.15 lubricant stearic acid Rough fraction (> 45 μm) 2.99 26 738.5 512.4 6.94 * manufacturer’s data Figure 12: Microstructure of sintered and heat-treated gear: a) produced from original powder mixture and b) produced from fine sieved powder mixture; polished and nital-etched sample, light microscope, magnification 100-times Slika 12: Mikrostruktura sintranega in pobolj{anega zobnika: a) izdelanega iz originalne me{anice in b) iz fine me{anice; poliran in v nitalu jedkan vzorec, svetlobni mikroskop; pove~ava 100-kratna Table 3: Vickers hardness of sintered and tempered gears Tabela 3: Trdote sintranih in pobolj{anih zobni~kov Original powder Fine sieved powder Sintered gears 212 298 Tempered gears 318 293 Figure 11: Microstructure of sintered gear: a) produced from original powder mixture and b) produced from fine sieved powder mixture; po- lished sample, light microscope; magnification 50-times Slika 11: Mikrostruktura sintranega zobnika: a) izdelanega iz original- ne me{anice in b) iz fine me{anice; poliran vzorec, svetlobni mikro- skop; pove~ava 50-kratna and heat treatment, visible under a LM. The gear made of the original powder mixture has a typical and correct, but heterogeneous, martensitic/bainitic microstructure. However, the gear made of the fine powder mixture has a microstructure in accordance with the inappropriate (over-alloyed) chemical composition. 4 CONCLUSIONS The original prealloyed Fe-based powder Distaloy AB was sieved and the technological properties of the fine and rough powder fractions important for ADC were determined. The fine powder fraction has a lower tap density, worse flowability and a negligibly better com- pressibility. It was expected that the selection of the fine powder fraction can contribute to an improvement in the ADC process of a small, two-height gear, especially to better filling of the teeth engraving, as well as a more uniform green-density distribution at a lower compaction pressure. This hypothesis has been disapproved based on experimental and semi-industrial work. It has been found that the selection of the finer powder mixture also has other traps. The sieved finer fraction has a different chemical composition than the original powder mixture. This has an important influence on the sintering and heat-treatment response of the material. Therefore, the poorer mechanical properties of the gears made of the fine fraction were obtained. The open question is also the price of manufacture of the fine powder mixture with the correct chemical composition. For now, the existent ADC procedure for the selected small gears is indicated as optimal. In the future, it will be necessary to find other ways to improve the ADC of small spur gears. 5 REFERENCES 1 R. M. German, Powder Metallurgy Science, Metal Powder Industries Federation (MPIF), Second Edition, Princeton, New Jersey 1994 2 F. Thümmler, R. Oberacker, Introduction to Powder Metallurgy, The Institute of Materials, The University Press, Cambridge 1993 3 Powder Metallurgies and Application, ASM Handbook, Volume 7, 1993 4 Unior, d. d., Forging Industry, Zre~e, program Sinter, web page address: http://unior-schmiede.com/cgi-bin/cms.cgi?doc=20454 5 Powder Metallurgy, materials, processes and application, CD product of the European Commission’s Leonardo Da Vinci Program, EPMA (European Powder Metallurgy Association), 2000 6 B. [u{tar{i~ et al., An instrumented cell to analyse the behaviour of metal powders during cold uniaxial die compaction, Mater. Tehnol., 35 (2001) 6, 351–360 B. [U[TAR[I^ et al.: THE INFLUENCE OF THE MORPHOLOGY OF IRON POWDER PARTICLES ... Materiali in tehnologije / Materials and technology 49 (2015) 2, 303–309 309