VSEBINA – CONTENTS IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES The microstructure of metastable austenite in X5CrNi18-10 steel after its strain-induced martensitic transformation Mikrostruktura metastabilnega avstenita po pretvorbi v napetostno inducirani martenzit v jeklu X5CrNi18-10 A. Kurc-Lisiecka, W. Ozgowicz, E. Kalinowska-Ozgowicz, W. Maziarz . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 837 The structure and morphology of the surface of duplex layers after saturation of the base layer with carbon Struktura in morfologija povr{ine dupleks plasti po nasi~enju osnovne plasti z ogljikom W. Skoneczny . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 845 Modeling of shot-peening effects on the surface properties of a (TiB + TiC)/Ti–6Al–4V composite employing artificial neural networks Modeliranje vpliva hladnega povr{inskega kovanja na lastnosti povr{ine (TiB + TiC)/Ti-6Al-4V kompozita s pomo~jo umetnih nevronskih mre` E. Maleki, A. Zabihollah . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 851 Analysis of twin-roll casting AA8079 alloy 6.35-μm foil rolling process Analiza procesa valjanja 6,35 μm folije iz zlitine AA8079 ulite med dvema valjema A. Can, H. Arikan, K. Çýnar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 861 Antimicrobial modification of polypropylene with silver nanoparticles immobilized on zinc stearate Protimikrobno spreminjanje polipropilena z nanodelci srebra, imobiliziranih na cinkovem stearatu G. Jandikova, P. Holcapkova, M. Hrabalikova, M. Machovsky, V. Sedlarik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 869 A new wideband negative-refractive-index metamaterial Novi {irokopasovni metamaterial z negativnim lomnim koli~nikom S. S. Islam, M. R. Iqbal Faruque, M. J. Hossain, M. T. Islam. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 873 Evaluation of the degree of degradation using the impact-echo method in civil engineering Ocena stopnje degradacije v gradbeni{tvu z uporabo metode odmeva zvo~nih valov D. [tefková, K. Tim~aková, L. Topoláø, P. Cikrle. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 879 Non-traditional whiteware based on calcium aluminate cement Netradicionalni porcelan na osnovi kalcij aluminatnega cementa R. Sokolar. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 885 Behaviour of new ODS alloys under single and multiple deformation Obna{anje novih ODS zlitin pri enojni in ve~kratni deformaciji B. Ma{ek, O. Khalaj, Z. Nový, T. Kubina, H. Jirkova, J. Svoboda, C. [tádler . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 891 Electromagnetic-shielding effectiveness and fracture behavior of laminated (Ni–NiAl3) composites U~inkovitost elektromagnetne za{~ite in obna{anje pri lomu laminiranega kompozita (Ni-NiAl3) T. Yener, S. C. Yener, S. Zeytýn . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 899 Effect of thermomechanical treatment on the intergranular corrosion of Al-Mg-Si-Type alloy bars Vpliv termomehanske predelave na interkristalno korozijo palic iz zlitin Al-Mg-Si P. Sláma, J. Nacházel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 903 Valorization of brick wastes in the fabrication of concrete blocks Ocena odpadkov iz opeke pri proizvodnji betonskih zidakov Y. Ghernouti, B. Rabehi, T. Bouziani, R. Chaid . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 911 Porous magnesium alloys prepared by powder metallurgy Porozne magnezijeve zlitine, izdelane s pomo~jo metalurgije prahov P. Salvetr, P. Novák, D. Vojtìch . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 917 Influence of nano-sized cobalt oxide additions on the structural and electrical properties of nickel-manganite-based NTC thermistors Vpliv dodatka nanodelcev kobaltovega oksida na zgradbo in elektri~ne lastnosti NTC termistorjev na osnovi nikljevega manganita G. Hardal, B. Y. Price . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 923 Durability of alumina silicate concrete based on slag/fly-ash blends against acid and chloride environments Zdr`ljivost betona na osnovi glinice in silikatov iz me{anice `lindra/lete~i pepel na kislo in kloridno okolje R. Gopalakrishnan, K. Chinnaraju . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 929 ISSN 1580-2949 UDK 669+666+678+53 MTAEC9, 50(6)835–1040(2016) MATER. TEHNOL. LETNIK VOLUME 50 [TEV. NO. 6 STR. P. 835–1040 LJUBLJANA SLOVENIJA NOV.–DEC. 2016 The size effect of heat-transfer surfaces on boiling Vpliv velikosti povr{in, ki prena{ajo toploto na vrenje P. Kracík, M. Balas, M. Lisy, J. Pospí{il . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 939 Effect of gas atmosphere on the non-metallic inclusions in laser-welded trip steel with Al and Si additions Vpliv plinske atmosfere na nekovinske vklju~ke v lasersko varjenem trip jeklu z dodatkom Al in Si A. Grajcar, M. Ró¿añski, M. Kamiñska, B. Grzegorczyk . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 945 Machining parameters influencing in electro chemical machining on AA6061 MMC Parametri strojne obdelave, ki vplivajo na elektrokemijsko strojno obdelavo AA6061 MMC C. J. Thankaraj Mariapushpam, D. Ravindran, M. D. Anand . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 951 Modeling the deep drawing of an AISI 304 stainless-steel rectangular cup using the finite-element method and an experimental validation Modeliranje globokega vleka pravokotne ~a{e iz AISI 304 nerjavnega jekla z metodo kon~nih elementov in z eksperimentalnim preverjanjem B. Sener, H. Kurtaran . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 961 Surface and anticorrosion properties of hydrophobic and hydrophilic TiO2 coatings on a stainless-steel substrate Povr{inske in protikorozijske lastnosti hidrofobnih in hidrofilnih TiO2 prevlek na jekleni podlagi M. Conradi, A. Kocijan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 967 Electroslag remelting: A process overview Elektropretaljevanje pod `lindro – pregled procesa B. Arh, B. Podgornik, J. Burja. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 971 Continuous vertical casting of a NiTi alloy Vertikalno kontinuirno litje NiTi zlitine A. Stamboli}, I. An`el, G. Lojen, A. Kocijan, M. Jenko, R. Rudolf . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 981 Hot tensile testing of SAF 2205 duplex stainless steel Vro~i natezni preskusi dupleks nerjavnega jekla SAF 2205 F. Tehovnik, B. @u`ek, J. Burja . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 989 STROKOVNI ^LANKI – PROFESSIONAL ARTICLES A high-efficiency automatic de-bubbling system for liquid silicone rubber Visokozmogljiv sistem za odpravljanje mehur~kov v teko~i silikonski gumi C.-C. Kuo, C.-M. Huang . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 995 Impact toughness of WMD after MAG welding with micro-jet cooling Udarna `ilavost WMD po MAG varjenju z mikro-jet hlajenjem T. Wegrzyn, J. Piwnik, A. Borek, A. Kurc-Lisiecka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1001 Forming-limit diagrams and strain-rate-dependent mechanical properties of AA6019-T4 and AA6061-T4 aluminium sheet materials Mejni diagrami preoblikovanja in odvisnost mehanskih lastnosti od hitrosti preoblikovanja aluminijevih plo~evin iz AA6019-T4 in AA6061-T4 O. Çavuºoðlu, A. G. Leacock, H. Gürün . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1005 Effect of alternative heat-treatment parameters on the aging behavior of short-fiber-reinforced 2124 Al composites Vpliv alternativnih parametrov toplotne obdelave na staranje 2124 Al kompozita, oja~anega s kratkimi vlakni Y. Altunpak, S. Aslan, M. Oðuz Güler, H. Akbulut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1011 A. KURC-LISIECKA et al.: THE MICROSTRUCTURE OF METASTABLE AUSTENITE IN X5CrNi18-10 STEEL ... 837–843 THE MICROSTRUCTURE OF METASTABLE AUSTENITE IN X5CrNi18-10 STEEL AFTER ITS STRAIN-INDUCED MARTENSITIC TRANSFORMATION MIKROSTRUKTURA METASTABILNEGA AVSTENITA PO PRETVORBI V NAPETOSTNO INDUCIRANI MARTENZIT V JEKLU X5CrNi18-10 Agnieszka Kurc-Lisiecka1, Wojciech Ozgowicz2, El¿bieta Kalinowska-Ozgowicz3, Wojciech Maziarz4 1Rail Transport Department, University of Dabrowa Gornicza, Cieplaka Str. 1C, 41-300 Dabrowa Gornicza, Poland 2Institute of Engineering Materials and Biomaterials, Silesian University of Technology, Gliwice, Poland 3Fundamentals of Technology Faculty, Lublin University of Technology, Nadbystrzycka Str. 38, 20-618 Lublin, Poland 4Institute of Metallurgy and Materials Science of the Polish Academy of Sciences, Reymonta Str. 25, 30-059 Krakow, Poland akurc@wp.pl Prejem rokopisa – received: 2015-05-19; sprejem za objavo – accepted for publication: 2015-11-05 doi:10.17222/mit.2015.102 The performed investigations concerned the influence of the degree and temperature of deformation on the microstructure of metastable austenite in the stainless steel X5CrNi18-10 after its strain-induced martensitic transformation. Samples of steel strip were cold rolled within a degree of deformation from 20 % to 70 % and stretched at a low temperature of -196 °C. The microstructure was observed by means of scanning electron microscopy (SEM) and transmission electron microscopy (TEM, HREM). It wasen found that after cold rolling with a small degree of deformation (20 %) in the tested steel, generally a single-phase microstructure of the matrix  is found with a high density of dislocations and numerous deformation bands morphologically characteristic of stainless steel with a low stacking-fault energy. After rolling with a 50 % thickness reduction, however, the microstructure displayed deformation twins as well as refined morphologic formations of the phase ’, mostly localized in the vicinity of the grain boundaries of the metastable matrix , and also trace amounts of carbide precipitates. In samples stretched at a temperature of -196 °C the microstructure of the matrix displayed a considerable density of dislocations with lath areas of the martensite ’ and precipitations of the carbides M23C6. Moreover, the tested steel revealed a crystallographic dependence of the planes and directions on the identified phases  and ’, corresponding to dependences of the Kurdjumov-Sachs type, independent of the method and temperature of the plastic deformation. Tests carried out in the TEM proved that the typical sites of nucleation induced by the plastic deformation of martensite are the shear bands, particularly their intersection. The preferred mechanism of transformation, observed in the conditions of cold rolling is, however, a direct transformation of the type  (fcc)  ’ (bcc). Keywords: austenitic stainless steels, cold rolling, microstructure, phase transformation, strain- induced martensite Izvedene so bile preiskave vpliva temperature in stopnje deformacije na mikrostrukturo metastabilnega avstenita po njegovi pretvorbi v napetostno inducirani martenzit v jeklu X5CrNi18-10. Vzorci v obliki trakov so bili hladno valjani s stopnjo deformacije od 20 % do 70 % in natezani pri nizki temperaturi – 196 °C. Mikrostruktura je bila opazovana s pomo~jo vrsti~ne elektronske mikroskopije (SEM) in s presevno elektronsko mikroskopijo (TEM, HREM). Ugotovljeno je, da je po hladnem valjanju z majhno stopnjo deformacije (20 %) v preizku{anem jeklu dobljena enofazna mikrostruktura z osnovo , z visoko gostoto dislokacij in {tevilnimi deformacijskimi pasovi, ki so morfolo{ka zna~ilnost nerjavnega jekla z nizko energijo napake zloga. Po valjanju s 50 % zmanj{anjem debeline, se v mikrostrukturi poka`ejo deformacijski dvoj~ki, kot tudi drobni nastanki faze ’, ve~inoma v bli`ini mej zrn metastabilne osnove  in tudi sledi izlo~kov karbidov. V vzorcih natezno obremenjenih pri temperaturi –196 °C je mikrostruktura osnove pokazala precej{njo gostoto dislokacij z latastimi podro~ji martenzita ’ in izlo~ki karbidov M23C6. Poleg tega je preiskovano jeklo pokazalo kristalografsko odvisnost usmerjenosti ravnin in ploskev v identificiranih fazah  in ’, ustrezno odvisnosti vrste Kurdjumov-Sachs, neodvisno od metode in temperature plasti~ne deformacije. Preiskave izvedene na TEM so potrdile, da so zna~ilna mesta nukleacije martenzita, inducirane s plasti~no deformacijo, stri`ni pasovi, posebno {e njihova kri`anja. Prednostni mehanizem premene, opa`ene pri hladnem valjanju, je neposredna premena vrste  (fcc)  ’ (bcc). Klju~ne besede: avstenitna nerjavna jekla, hladno valjanje, mikrostruktura, fazna premena, napetostno inducirani martenzit 1 INTRODUCTION Austenitic stainless steels are widely used in many engineering applications, such as in the chemical, machi- nery, food, automotive, nuclear and shipbuilding indus- tries, due to their excellent corrosion resistance, weld- ability, and mechanical properties. However, some of these austenitic steels with a lower content of Ni can undergo a transformation to martensite during cold working.1 A different martensite morphology can be formed due to these processes, mainly strain-induced or stress-induced martensite.2 In austenitic stainless steels two types of martensite can form spontaneously, i.e., body-centered cubic (bcc) martensite ’ and hexagonal close-packed (hcp) martensite . The amount of  and/or ’ martensite depends on the chemical composition, stacking-fault energy, phase stability and processing parameters, such as stress state, temperature, strain rate. Materiali in tehnologije / Materials and technology 50 (2016) 6, 837–843 837 UDK 669.112.227.34:669.15-194.5:669.112.227.1 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)837(2016) During the deformation process different transformation sequences take place, such as:     ’ or   ’. In the transformation mode     ’,  martensite acts as the precursor phase of ’. The formation of ’ is closely related to the shear bands, which are planar defects associated with the overlapping of stacking faults on 111. Depending on the nature of the overlapping, twins,  martensite or stacking-fault bundles may be formed. Twins are formed when stacking faults overlap on successive {111} planes, whereas  martensite is generated if the overlapping of the stacking faults occurs on alternate 111 planes. Stacking-fault bundles arise from the irregular overlapping of stacking faults.1–6 The presence of deformation-induced martensite may be a harmful phenomenon and may cause a delayed cracking of deep-drawn austenitic stainless-steel compo- nents. On the other hand, the formation of martensite resulting from the plastic deformation of metastable austenite is of great interest for the production of high- strength and ductile austenitic stainless steels.2,3 The aim of the present study was to analyze the morphological details of strain-induced martensite in cold-rolled Cr-Ni steel. 2 MATERIAL AND EXPERIMENTAL PROCEDURE The investigations concerned austenitic stainless steel of the type X5CrNi18-10 in compliance with PN-EN 10088:1-2007 7 with the chemical composition quoted in Table 1. The input material in the form of steel strip, 2 mm thick, 40 mm in width and 700 mm long was supersaturated in water after its austenitizing for 1 h at a temperature of 1100 °C and cold rolled with a 20 %, 50 % and 70 % thickness reduction. After rolling with a draft of 70 %, samples of the tested steel were subjected to a tensile test at a low temperature of –196 °C with a strain rate  of about 10–5 s–1. Table 1: Chemical composition of the investigated steel Tabela 1: Kemijska sestava preiskovanih jekel Elements content, in mass fractions (w/%) C Mn Si P S Cr Ni Ti Al Fe 0.024 1.32 0.43 0.028 0.005 18.53 7.8 0.010 0.01 bal. The hardness measurements of the investigated cold-rolled steel were carried out with a microhardness tester FM 700 produced by Future-Tech (Japan), according to the standard PN-EN ISO 6507-1:2007.8 The hardness was also determined in the case of the sample after 70 % degree and stretched at a low temperature of –196 °C with a strain rate  of about 10–5 s–1. The measurements were made using the Vickers method on metallographic samples with a load of 50 N for a time of 30 s. The microstructural investigations were performed with scanning (SEM) and transmission electron micro- scopy, as well as high-resolution electron microscopy (HREM). Applying SEM, a metallographic polished sec- tion after cold rolling with a draft of 20 % and stretching at the temperature of liquid nitrogen was detected. These observations were made by means of SEM of the SUPRA type from Zeiss (Germany) with a magnification of 15.000×. The section that was mechanically polished was etched in the reagent Mi17Fe.9 TEM observations were carried out using thin foils on the samples of strip after cold rolling with a draft of 50 %, and on samples stretched at a temperature of –196 °C. The preparation of the foils comprised a cutting out of disks, 3 mm in dia- meter, from a strip with a thickness of 1.0 and 0.6 mm, grinding with abrasive paper until the samples reached a thickness of 0.1 mm. The Tenupol-5 double jet electro- polisher was used for thin foil preparation from the samples in an electrolyte containing nitric acid and methanol (1:3). The microstructure was observed by means of TEM of the type Technai G2 F20 applying an accelerating voltage of 200 kV equipped with high-angle annular dark-field (HAADF) and energy-dispersive (EDS) detector. The phases were identified based on electron diffraction. The procedure was aided by the A. KURC-LISIECKA et al.: THE MICROSTRUCTURE OF METASTABLE AUSTENITE IN X5CrNi18-10 STEEL ... 838 Materiali in tehnologije / Materials and technology 50 (2016) 6, 837–843 Figure 1: Microstructure of investigated steel X5CrNi18-10: a) after 20 % of deformation, b) after cold-rolling with 70 % and tensile test at –196 °C, etching- Mi17Fe Slika 1: Mikrostruktura preiskovanega jekla X5CrNi18-10: a) po 20 % deformaciji, b) po hladnem valjanju s 70 % in nateznim preizkusom pri –196 °C, jedkano z Mi17Fe computer software Gatan and a crystallographic data- base. 3 RESULTS AND DISCUSSION In the supersaturated state the investigated steel dis- plays a single-phase austenite structure with a diameter of the average grains in the matrix  amounting to about 75 ìm and a hardness of about 125 HV0.5, containing many annealed twins and single clusters of non-metallic inclusions. After cold rolling in the range 20–30 % metallographically distinctly elongated grains of the matrix  with a hardness of 323 HV5 (Table 2) could be detected with numerous effects of work hardening in the form of fine parallel and intersected lines and slip bands, as well as shear bands, which are probably sites of mar- tensite ’ nucleation. Table 2: Results of the hardness measurement of the investigated cold-rolled and stretched steel Tabela 2: Meritve trdote preiskovanih hladno valjanih in natezanih jekel No. Material condition Hardness, HV Hard- ness, HV 5 Number of measurement 1 2 3 1 supersaturated 144.7 148.5 145.8 146.3 2 cold rolled zh=20% 321.5 322.7 325.9 323.4 3 cold rolled zh=50% 411.5 410.8 408.7 410.3 4 cold rolled zh=70% 418.5 417.4 418.6 418.1 5 cold rolled with zh=70% and stretched at –196C 460.1 461.3 459.2 460.2 The results of the observation of the microstructure of the investigated steel after cold-rolling with a degree of deformation of 20 % and 70 % and after stretching at a temperature of –196 °C carried out on a scanning elec- tron microscope (SEM) are presented in the micrographs A. KURC-LISIECKA et al.: THE MICROSTRUCTURE OF METASTABLE AUSTENITE IN X5CrNi18-10 STEEL ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 837–843 839 Figure 3: TEM micrograph structure of X5CrNi18-10 steel after 50 % of deformation: a) microstructure of the matrix  containing microtwins and martensite ’, b) dark field taken of reflection (200) , c), d) diffraction pattern Slika 3: TEM-posnetek strukture jekla X5CrNi18-10 po 50 % deformaciji: a) mikrostruktura osnove , ki vsebuje mikrodvoj~ke in martenzit ’, b) temno polje pri odsevu (200) , c), d) uklonska slika Figure 2: TEM micrograph of X5CrNi18-10 steel after 50 % of deformation: a) a band of austenite containing microtwins, b) diffraction pattern Slika 2: TEM-posnetek jekla X5CrNi18-10 po 5 % deformaciji: a) pas avstenita, ki vsebuje mikrodvoj~ke, b) uklonska slika A. KURC-LISIECKA et al.: THE MICROSTRUCTURE OF METASTABLE AUSTENITE IN X5CrNi18-10 STEEL ... 840 Materiali in tehnologije / Materials and technology 50 (2016) 6, 837–843 Figure 5: TEM micrograph structure of X5CrNi18-10 steel after 50 % of deformation: a) cell microstructure of austenite containing a variable dislocation density and ultra-fine lath of martensite ’, b) diffraction pattern Slika 5: TEM-posnetek strukture jekla X5CrNi18-10 po 50 % defor- maciji: a) celi~na mikrostruktura avstenita vsebuje razli~no gostoto dislokacij in ultra drobni latasti martenzit ’, b) uklonska slika Figure 4: TEM micrograph structure of X5CrNi18-10 steel after 50 % of deformation: a) subgrain of austenite containing a high density of dislocations and ’, bright field, b) dark field taken of the reflection (110)’, c) diffraction pattern Slika 4: TEM-posnetek strukture jekla X5CrNi18-10 po 50 % deformaciji: a) podzrna avstenita vsebujejo veliko gostoto dislokacij in ’, svetlo polje, b) temno polje pri odsevu (110) ’, c) uklonska slika Figure 6: High-resolution (HREM) micrograph: a) dislocation structure of the matrix  of steel B after cold rolling (zh=50 %) and Fourier transform (FFT), b) detail A of Figure 6a – modulated structure (IFFT) with microtwins bands and Fourier transform (FFT), c) solution of Fourier transform in Figures 6a and 6b Slika 6: Visokolo~ljivi posnetek (HREM): a) struktura dislokacij v osnovi  jekla B po hladnem valjanju (zh=50 %) in Fourierjeva pretvorba (FFT), b) detajl A na Sliki 6a modulirana struktura (IFFT) s pasovi mikro tvoj~kov in Fourierjeva pretvorba (FFT), c) re{itev Foirierjeve pretvorbe na Slikah 6a in 6b in Figure 1. In the structure of the steel, complex effects of deformation inside the grains  and at the boundaries are revealed (Figure 1a). Plastic deformation leads to a distinct elongation of the grains in the direction of roll- ing and to the formation of numerous slide bands and shear bands, in which probably the martensite ’ is loca- lized (Figure 1b). The hardness of the examined steel increases with an increasing degree of deformation. With the increase of the cold rolling degree from 50 % to 70 % the hardness of the investigated steel increases from 410 to 418 HV5, respectively (Table 2). As suggested in10,11 the twins, the dislocation density, the nucleation of martensite ’ and the increase of the volume fraction of martensite ’ phase during the transformation are the major factors influencing the hardness of the investigated steel. Heterogeneities in the plastic deformation in the form of shear bands were found mainly in the case of larger, cold plastic working and tensile tests at reduced tempe- ratures up to –196 °C. Thin foils in the TEM revealed in steel X5CrNi18-10, cold rolled with a degree of defor- mation of 50 %, a cellular structure of dislocations of an austenitic matrix with a considerable density of disloca- tions with local twins (Figure 2). Also, single reflexes of the type (112)’ and (123)’ were observed, resulting from the martensitic phase ’ (Figure 3d). Based on electron diffraction and the dark-field method, the localisation of the deformation twins could be identified and the direction of twinning (TD) <111> o was deter- mined (Figures 2b and 3d). In the microstructure of the investigated steel, highly elongated subgrain  and shearing bands dominate, and also an ultra-fine lath of martensite ’ with a characteristic dislocation forest (Figures 4 and 5). After cold rolling, observed in high- resolution microscopy (HREM), the structure of the investigated steel reveals significant morphological details – on the nanometer scale – microbands of mecha- nical twinning as well as in the range of periodicity of the structure and its modulated character (Figure 6). The disclosed structural periodicity is reflected in the A. KURC-LISIECKA et al.: THE MICROSTRUCTURE OF METASTABLE AUSTENITE IN X5CrNi18-10 STEEL ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 837–843 841 Figure 8: TEM micrograph structure of the X5CrNi18-10 steel: a) mechanical twins after 70 % of deformation and tensile test at tem- perature of –196 °C with  =10–5 s–1, b) dark field in (002), c), d) diffraction pattern Slika 8: TEM-posnetek strukture jekla X5CrNi18-10: a) mehanski dvoj~ki po 70 % deformaciji in nateznem preizkusu pri –196 °C, z  =10–5 s–1, b) temno polje pri (002) , c), d) uklonska slika Figure 7: TEM micrograph structure of X5CrNi18-10 steel after cold-rolling with 70 % and tensile test at temperature of -196 °C with a strain rate of 10–5 s–1: a) microstructure of elongated subgrain  with shear bands and heavily deformed lath of martensite ’, b) diffraction pattern from a, c) detail in Figure a, d) diffraction pattern from c Slika 7: TEM-posnetek strukture jekla X5CrNi18-10 po hladnem valjanju, 70 % in nateznem preizkusu pri –196 °C, s hitrostjo obre- menjevanja 10–5 s–1: a) mikrostruktura razpotegnjenih podzrn , s stri`nimi pasovi in mo~no deformiranimi latami martenzita-’, b) uklonska slika iz a, c) detajl iz slike a, d) uklonska slika iz c distribution of the atomic cores in the inverse Fourier transform (IFFT) (Figure 6c). High-resolution analysis of the sequence of microtwin bands (Figure 6a) and the corresponding Fourier transforms (FFT), comprising the entire area of HREM (Figure 6b) and the marked area of the microstructure (Figure 6c) justify the statement that the visible and most intense reflections result mainly from the matrix  oriented as a zone axis of orientation [011] (Figure 6b). The additional weak reflections between the reflections of the planes (002) and (111)  can be attributed to the deformation twins (Figure 6b). However, the presence of two weak reflections between the beam passing (000) and the planes (111) and (111), dividing these distances into three equal parts with a length of 1/3 (111), require an explanation. The presence of these reflections is not justified, however, in the case of twin orientations.12 In the transform (FFT) concerning the area of the band of microtwins (Figure 6c) there is a twin orientation with strong defocusing reflections in the planes (111). The inverse Fourier transform resulting from a transform (FFT) (Figure 6c) after filtering out the noise reveals that in the matrix  (M-matrix) a microtwin (T-twin) is located with a width of about 5 nm, inside which the modulation effects are visible. Modulations are caused by periodic sequences of stacking faults occurring on the following planes (111) (Figure 6d). The investigated cold-rolled steel with the degree of deformation of 70 % and then subjected to a tensile test at strain rate () of about 10–5 s–1 at cryogenic tem- peratures –196 °C displayed – similar to the cold-rolling – a subgrain structure elongated in the rolling direction with a high density of dislocations (Figure 7) and a considerably higher density of microtwinning (Figure 8). The hardness in these areas reaches about 460 HV5 (Table 2). The subgrain boundaries and the microtwins constitute potential locations for the phase ’, in the form of elongated lamellar areas with a width of approximately 0.1 μm (Figure 7b). It can be assumed that the nucleation of the phase ’ occurs preferentially in microtwins areas, mainly at their borders. It is sig- nificantly associated with the accumulation of the stress in a dislocation field, as suggested in11. Electron-diffrac- tion analysis of the investigated steel not only provides evidence for the presence in its structure of martensitic phase ’ (Figures 3d, 4b, 5b, 7a, 7b and 8d) and M23C6 type carbides (Figure 8d), but also the occurrence of a crystallographic relationship between the matrix  and phase ’ type K-S, namely: (111)  II (011)’ and <011>  II <111> ’ (Figure 4b), also quoted with res- pect to similar grades of Cr-Ni steel.11,13 4 CONCLUDING REMARKS The structural investigations of the steel X5CrNi18-10 conducted in a TEM and the analysis of the obtained results allows us to draw the following conclusions: The plastic deformation of the investigated steel X5CrNi18-10 induces the direct transformation of metastable austenite to the deformation martensite ’ of the (bcc) lattice during both the cold-rolling process, as well as the tensile test at temperatures lowered to –196 °C. The microstructure of the investigated steel after cold rolling with a degree of deformation in the range from 50 % to 70 % observed in the TEM, displays a high dis- location density in the matrix  and the presence of mechanical twins, as well as shearing bands in the area where the lamellar formations of the martensite ’ phase nucleate. The cold rolling and stretching at low temperature of the austenitic stainless-steel sheets resulted in the occur- rence of the strain-induced   ’ phase transformation. During plastic deformation the volume fraction of mar- tensite ’ phase increases, which causes the hardening of the investigated steel. The hardness of the cold-rolled steel within the draft from 20–70 % is from the range 323–418 HV5, whereas in the case of samples after 70 % degree of cold rolling and stretching at –196 °C it is about 460 HV5. High-resolution electron microscopy (HREM) of the microstructure revealed essential morphological details of the resulting microbands of twins on the nanometer scale. The application of Fourier’s reverse transform (IFFT) indicated a periodicity of the analyzed structure and its modulated character in the range of appearing sequentially, the local disorder of the crystalline lattice. The transformations   ’ of the investigated steel induced by plastic deformation indicate a typical crystal- lographic relationship between austenite and martensite ’ given by Kurdjumov-Sachs, in the form: (111)II(011)’ and <011>II<111>’. Acknowledgements The authors gratefully acknowledge financial support from the research project: Innovative sanitary sewage system DEMONSTRATOR + NCBR under the contract No. UOD-DEM-1-591/001. 5 REFERENCES 1 K. H. Lo, C. H. Shek, J. K. L. Lai, Morphologies and characteristics of deformation induced martensite during tensile deformation of 304 LN stainless steel, Materials Science and Engineering, 65R (2009) 4–6, 39–104, doi:10.1016/j.msea.2007.09.005 2 A. Das, S. Sivaprasad, M. Ghosh, P. C. Chakraborti, S. Tarafder, Morphologies and characteristics of deformation induced martensite during tensile deformation of 304 LN stainless steel, Materials Science and Engineering, 486A (2008), 283–286, doi:10.1016/ j.msea.2007.09.005 3 J. Huang, X. Ye, Z. Xu, Effect of Cold Rolling on Microstructure and Mechanical Properties of AISI 301LN Metastable Austenitic Stain- A. KURC-LISIECKA et al.: THE MICROSTRUCTURE OF METASTABLE AUSTENITE IN X5CrNi18-10 STEEL ... 842 Materiali in tehnologije / Materials and technology 50 (2016) 6, 837–843 less Steels, Journal of Iron and Steel Research International, 19 (2012) 10, 59–63, doi:10.1016/S1006-706X(12)60153-8 4 C. J. Gunter, R. P. Reed, Transaction of American Society for Metals, 55 (1962) 3, 399–419 5 S. Rajasekhara, L. P. Karjalainen, A. Kyröläinen, P. J. Ferreira, Microstructure evolution in nano/submicron grained AISI 301LN stainless steel, Materials Science and Engineering, 527A (2010), 1986–1996, doi:10.1016/j.msea.2009.11.037 6 T. Angel, Journal of the Iron and Steel Institute, 177 (1954), 165–174 7 European Standard, Stainless Steels - Part 1: List of stainless steels, Polish version PN-EN 10088:1-2007 8 European Standard, Metals. Hardness measurements made by Vickers’s method, Polish version PN-EN ISO 6507-1:2007 9 ASTM E407, Standard Practice for Microetching Metals and Alloys 10 W. S. Lee, C. F. Lin, The morphologies and characteristics of impact-induced martensite in 304L stainless steel, Scripta Materialia, 43 (2000) 8, 777–782, doi:10.1016/S1359-6462(00)00487-5 11 J. A. Venables, The martensite transformation in stainless steel, Phi- losophical Magazine, 73 (1962) 7, 35–44, doi:10.1080/ 14786436208201856 12 W. Maziarz, Structure changes of Co–Ni–Al ferromagnetic shape memory alloys after vacuum annealing and hot rolling, Journal of Alloys and Compounds, 448 (2008) 1–2, 223–226, doi:10.1016/ j.jallcom.2006.12.044 13 M. Blicharski, S. Gorczyca, Structural inhomogeneity of deformed austenitic stainless steel, Metal Science, 12 (1978) 7, 303–312, doi:10.1179/msc.1978.12.7.303 A. KURC-LISIECKA et al.: THE MICROSTRUCTURE OF METASTABLE AUSTENITE IN X5CrNi18-10 STEEL ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 837–843 843 W. SKONECZNY: THE STRUCTURE AND MORPHOLOGY OF THE SURFACE OF DUPLEX LAYERS ... 845–850 THE STRUCTURE AND MORPHOLOGY OF THE SURFACE OF DUPLEX LAYERS AFTER SATURATION OF THE BASE LAYER WITH CARBON STRUKTURA IN MORFOLOGIJA POVR[INE DUPLEKS PLASTI PO NASI^ENJU OSNOVNE PLASTI Z OGLJIKOM W³adys³aw Skoneczny University of Silesia, Ul. ¯ytnia 10, 41-200 Sosnowiec, Poland Wladyslaw.Skoneczny@us.edu.pl Prejem rokopisa – received: 2015-05-29; sprejem za objavo – accepted for publication: 2015-12-15 doi:10.17222/mit.2015.108 The paper presents a new method of obtaining aluminium-oxide-based, duplex-type layers. In the first stage, the base layer is produced via the hard anodising of aluminium alloys in order to obtain the optimal structural and morphological properties. Following anodising, the samples with an Al2O3 layer are rinsed in distilled water in order to remove the electrolyte. Graphite was introduced into the structure of aluminium oxide during a thermal treatment in a solid medium consisting of graphite dust. Afterwards, the properties of the obtained layers were determined using a scanning electron microscope, a transmission electron microscope and an atomic force microscope (AFM), as well as X-ray diffraction. The structure of the duplex-type layers con- tains carbon and other precipitates, which are typical for an alloy with additions of Fe, Mn, Cr and other elements. Carbon precipitates have a relatively weak connection with the matrix, as an envelope with numerous discontinuities forms around each carbon precipitate. Carbon precipitates are considerably larger than alloy precipitates, have micrometre dimensions, occur in groups and are composed of small grouped nanometric particles that form larger agglomerates. This means that there are nano- metric particles inside the micrometric ones. Keywords: aluminium alloys, nano-layers, SEM, AFM, EDS ^lanek predstavlja novo metodo za izdelavo dupleksne plasti aluminijevega oksida. V prvi stopnji se izdela osnovna plast s trdim anodiziranjem aluminijevih zlitin, da se dobi optimalno strukturo in morfolo{ke lastnosti. Po anodizaciji so vzorci z Al2O3 plastjo potopljeni v destilirano vodo, da se izpere elektrolit. V strukturo aluminijevega oksida se uvede grafit med toplotno obdelavo v trdem mediju, ki je vseboval grafitni prah. Potem so bile dolo~ene lastnosti dobljenih plasti, s pomo~jo vrsti~nega elektronskega mikroskopa, presevnega mikroskopa in mikroskopa na atomsko silo (AFM), kot tudi z rentgensko difrakcijo. Struktura dupleksnih plasti je vsebovala izlo~ke ogljika in druge izlo~ke, ki so zna~ilni za zlitine, z dodatkom Fe, Mn, Cr in drugih elementov. Izlo~ki ogljika imajo relativno {ibko povezavo z osnovo, saj nastaja okrog vsakega izlo~ka ogljika ovojnica s {tevilnimi diskontinuitetami. Izlo~ki ogljika so mnogo ve~ji kot izlo~ki zlitin, imajo mikrometrske dimenzije, se pojavljajo v skupinah in so sestavljeni iz malih gru~ nanometri~nih delcev, ki tvorijo ve~je skupke. To pomeni, da so znotraj mikrometrskih delcev prisotni nanometrski delci. Klju~ne besede: aluminijeve zlitine, nanoplasti, SEM, AFM, EDS 1 INTRODUCTION Aluminium alloys and the composite layers formed on them are widely used in engineering (components of heat engines and power machines, in aircraft and space industries) owing to their very good thermal conducti- vity, low density and high strength. Oxide layers obtained via hard anodising may, to a large degree, change their properties, depending on the process conditions. Thanks to their characteristic porous structure, Al2O3 layers can be used in a number of tech- nology fields. Oxide coatings have been used in surface engineering as protective-decorative or electro-isolating layers for many years. In recent years, Al2O3 layers have been used for the sliding couples in heat engines:1–6 • cylinder bearing surfaces in lubricant-free compres- sors, • combustion-engine pistons, • cylinder bearing surfaces in lubricant-free pneumatic servo-motors, • shock-absorber components. One of the most recent applications of oxide layers are the moulds used for producing nano-elements with a diameter of 4–200 nm.7 Machine components produced from aluminium or its alloys with a specially prepared surface layer are used more and more frequently. The main arguments for broadening the scope of applications for aluminium- oxide-based, duplex-type layers in surface engineering are as follows: • easy access to devices and very inexpensive oper- ation, • the possibility of conducting a surface treatment on all Al groups and its alloys, • high efficiency of the duplex-type treatment, • the possibility of obtaining electro-isolating and pro- tective-decorative coatings, as well as hard, abrasion- resistant layers, • satisfactory quality of the surface after processing, Materiali in tehnologije / Materials and technology 50 (2016) 6, 845–850 845 UDK 669.058:620.198:669.715:537.533.35 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)845(2016) • the possibility of forming surface layers with a gradient, composite and, first of all, an amorphous structure. The wear of materials used in machine building takes place mostly on the surface and determines the ope- rational durability of a given appliance. The past few years have been a period of intensive development in state-of-the-art, duplex, hybrid technologies. It seems that only a combination of different surface-engineering technologies, which is the foundation of the duplex technology can lead to surfaces with versatile qualities that are to meet the requirements of modern technology. As the name indicates, the duplex technology involves a sequential application of two defined surface-engineer- ing technologies to produce a surface composed of com- bined properties, unattainable by any other individual form of surface engineering. This paper presents a wide range of possibilities to shape the properties of aluminium-oxide-based surface layers using the duplex method. For technological reasons, hard anodising is applied for lubricant-free sliding couples. The most widespread method so far has been anodising in sulphuric or oxalic acids at lowered temperatures, from 263 K to 278 K, depending on the type of electrolyte. Therefore, research has been conducted for years on anodizing at elevated temperatures. The purpose of the research has been to apply such an electrolyte that would make it possible to obtain hard oxide layers at room temperature. Elimina- tion of the electrolyte cooling stage would considerably reduce the cost of producing oxide coatings. It would be possible if hard anodising could be conducted at tempe- ratures of 293–313 K and higher. At the same time, coatings with better properties could be obtained owing to the Al2O3 oxide phase transition at a temperature of 293 K. An increase in temperature accompanying the oxidation process is conducive to etching aluminium oxide fibres. As a consequence, an oxide cell of a more regular (ideal) structure is formed.1,8,9 Increasing the electrolyte temperature would also have a considerable effect on the porosity of the oxide coating.1–4 The porosity of the obtained oxide layers is of major import- ance to their utilisation for a sliding interaction with plastic materials. Adding organic substances with surface-active properties to the electrolyte has a large influence on the mechanism of the formation of oxide coatings on aluminium. The mechanism of the influence of organic substances depends on the properties of the addition and on the composition and properties of the electrolyte. A supposition can be made that under proper conditions, surface-active substances fully or partly cover the surface of the anode (on active places), as a result of which the oxidation of aluminium is considerably hindered. On the other hand, the adsorption of organic substances at the anode – the electrolyte interface causes the secondary dissolution of the layer by the electrolyte to stop. The role of this mechanism is performed precisely by the addition of the above-mentioned organic (dicarboxylic) acids. The method developed in the Division of Upper Layer Technologies, University of Silesia, does not require cooling and the process heat is used to control the properties of the obtained oxide coatings. Controlling anodising parameters allows, within some limits, programming the selected functional properties of future upper layers.3–12 The above-mentioned method consists of oxidising aluminium and its alloys in three-component electro- lytes. Dicarboxylic acid is added to the mixture of sulphuric and oxalic acids. These acids have an aliphatic chain of various lengths in their structure and are arranged in a row: 1) malonic acid - CH2(COOH)2 2) succinic acid - (CH2)2(COOH)2 3) glutaric acid - (CH2)3(COOH)2 4) adipic acid - (CH2)4(COOH)2 5) pimelic acid - (CH2)5(COOH)2 6) suberic acid - (CH2)6(COOH)2 7) azelaic acid - (CH2)7(COOH)2 8) sebacic acid - (CH2)8(COOH)2 9) phthalic acid - C6H4(COOH)2. 2 EXPERIMENTAL PART The hard anodising method developed in the Division of Upper Layer Technologies, University of Silesia, is the basis for the production of surface layers from aluminium oxide by means of duplex methods. It enables the control of process parameters, which allows, within some limits, programming the selected functional pro- perties of the obtained upper layers. According to the method proposed, anodising is conducted in a three-com- ponent water electrolyte that consists, of among others, sulphuric and oxalic acids as well as an addition of succinic acid. Research was carried out on the alumi- nium alloy AlMg2. The electrolyte temperature during anodising falls within the range 293–313 K, whereas the anodic density of current was between 2 A/dm2 and 4 A/dm2. Such a range of temperatures means that the anodising process can be initiated at an ambient temperature. Tests of the structure and morphology of the surface of Al2O3 layers were conducted using a Phi- lips XL30 scanning microscope (SEM). An atomic force microscope (AFM) of the VEECO company, MULTOMODE model, operating in the standard contact mode, was used for the examination of the micro-unevenesses of the obtained oxide layers via the electrochemical method in three-component electro- lytes after graphite saturation. A DRON-2 diffractometer was used for an X-ray phase analysis of the obtained Al2O3 and duplex layers. Al2O3 layers have a columnar (fibrous) structure, which is shown in Figure 1. The technology of obtaining duplex layers based on aluminium oxide consists of a two-stage production W. SKONECZNY: THE STRUCTURE AND MORPHOLOGY OF THE SURFACE OF DUPLEX LAYERS ... 846 Materiali in tehnologije / Materials and technology 50 (2016) 6, 845–850 process. In the first stage, the base layer is produced via hard anodising in order to obtain the optimal structural and morphological properties (gradient structure and high porosity). Prior to the oxidation process, the samples are etched for 40 min in a 5 % solution of KOH and next, in order to reverse the etching reaction, tinned for 10 min in a 10 % HNO3 solution. The etching and tinning processes are followed by rinsing in distilled water. In the first phase, the surfaces are subjected to anodic oxidation in an electrolyte composed of an aqueous solution of sulphuric, succinic and oxalic acid. Following anodising, the samples with an Al2O3 layer are rinsed in distilled water in order to rinse out the electrolyte. Graphite was introduced into the structure of the aluminium oxide during the thermal treatment in a solid medium consisting of graphite dust. Samples sprinkled with graphite dust are tightly closed in boxes and heated in an electric oven at the following tempe- ratures: (343, 363, 383 and 403) K for (24, 36 and 48) h for each temperature. When the process ends, the sam- ples were cleaned with compressed air. 3 RESULTS AND DISCUSSION A fibrous structure causes micro- and nanoporosity of the Al2O3 layer. An example of the morphology of the surface of the Al2O3 layers, obtained via hard anodising, is shown in Figure 2. The measurement results of the diameter of fibres and the number of fibres and nanopores are juxtaposed in Table 1. Tests of the structure of the duplex-type layers after graphite infiltration were conducted using a transmission electron microscope. The results of the tests of the structure of duplex-type layers after carbonisation of the Al2O3 layers via saturation are presented in Figures 3 and 4. An analysis of the chemical composition was also made using a Philips XL30 scanning electron micro- scope with an EDS attachment. The results of the tests of W. SKONECZNY: THE STRUCTURE AND MORPHOLOGY OF THE SURFACE OF DUPLEX LAYERS ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 845–850 847 Figure 2: SEM image of nanopores in the Al2O3 layer Slika 2: SEM-posnetek nanopor v plasti Al2O3 Table 1: The dimensions, the number of fibres and the number of pores Tabela 1: Dimenzije, {tevilo vlaken in {tevilo por Production parameters Fibre diameter (nm) Number of fibres per mm2 Number of pores per mm2 j= 2 A/dm2 t = 293 K 100 100 × 10 6 200 × 106 j= 2 A/dm2 t = 303 K 120 64 × 10 6 128 × 106 j= 2 A/dm2 t = 313 K 200 25 × 10 6 50 × 106 j= 3 A/dm2 t = 293 K 110 81 × 10 6 162 × 106 j= 3 A/dm2 t = 302 K 140 49 × 10 6 98 × 106 j= 3 A/dm2 t = 303 K 160 36 × 10 6 72 × 106 Figure 1: Cross-sectional SEM image a columnar-fibrous structure of an oxide layer Slika 1: SEM-posnetek preseka stebrasto-vlaknaste zgradbe oksidne plasti Figure 3: Top view of the surface of a duplex-type layer after satu- ration with graphite, with spectra (transmission microscope) Slika 3: Slika povr{ine dupleksne vrste plasti po nasi~enju z grafitom, s spektri (presevni mikroskop) the change in the chemical composition of these layers obtained during carbonisation at the temperatures of 383 K are presented in Figure 5. The structure of the duplex-type layers contains carbon and other precipitates, which are typical for the alloy, with additions of Fe, Mn, Cr and other elements. Carbon precipitates have a relatively weak connection with the matrix, as an envelope with numerous discon- tinuities forms around each carbon precipitate. Carbon precipitates are considerably larger than alloy precipi- tates, have micrometric dimensions, occur in groups and are composed of small grouped nanometric particles that form larger agglomerates. This means that there are nanometric particles inside micrometric ones. Therefore, the structure of carbon particles is complex. The results of the tests of duplex-type layers obtained via the electrochemical method in three-component electrolytes after graphite saturation are presented in Figure 6. The following is shown: a) a topographic three-dimensional image, b) the geometric microstruc- ture and a histogram. X-ray diffractograms of the Al2O3 layer obtained on a crystalline alloy of AlMg2 aluminium and aluminium oxide, as well as a duplex layer are presented in Figure 7. An X-ray phase analysis has shown that the obtained Al2O3 layers are amorphous and also contain two peaks characteristic for the crystalline form of Kappa Al2O3, a phase of an indeterminate structure type, entered into the crystallographic data catalogue under the number 04-0878. A very strong typical peak from the graphite phase called chaolite (22-1069) is also present. W. SKONECZNY: THE STRUCTURE AND MORPHOLOGY OF THE SURFACE OF DUPLEX LAYERS ... 848 Materiali in tehnologije / Materials and technology 50 (2016) 6, 845–850 Figure 5: Changes in the chemical composition of duplex-type layers obtained after saturation at the temperature of 383 K: a) change in the oxygen and aluminium content, b) change in the graphite content Slika 5: Spreminjanje kemijske sestave dupleksne plasti, dobljene po nasi~enju na temperaturi 383 K: a) spreminjanje vsebnosti kisika in aluminija, b) spreminjanje vsebnosti grafita Figure 6: a) AFM images of the surface (3D) of duplex-type layers and b) micro-unevennesses of the surface, and a histogram; anodising process parameters: j = 3 A/dm2, T = 303 K, t = 1 h; carburisation parameters: T = 343 K, t = 24 h Slika 6: a) AFM-posnetek povr{ine (3D) dupleksnih plasti in b) mikro-neravnine na povr{ini in histogram s parametri procesa anodizacije : j = 3 A/dm2, T = 303 K, t = 1 h; parametri naoglji~enja: T = 343 K, t = 24 h Figure 4: Cross-sectional of a duplex-type layer after saturation with graphite, with spectra (transmission microscope) Slika 4: Presek dupleksne plasti po nasi~enju z grafitom, s spektri (presevni mikroskop) The porosity of the Al2O3 layers has a significant in- fluence on their properties, including their wear resis- tance, the capacity for the sorption of lubricants, the possibility of a sliding film forming from a plastic (containing graphite) and first of all the susceptibility to further modification, which is used for obtaining layers via the duplex method. It is the micro- and nanoporosity of the Al2O3 layer that was used during the second stage of the dual technology to form duplex-type surface layers. The second technology used to produce duplex- type layers consisted of infiltration with graphite. Owing to this, it will become possible to fill the above-described structure of the Al2O3 layer that is obtained via hard anodising with graphite. 4 CONCLUSIONS The technology of obtaining aluminium-oxide-based duplex layers consists involves a two-stage production process. In the first stage, the base layer is produced via hard anodising in order to obtain optimal structural and morphological properties (gradient structure and high porosity). Following anodising, the samples with an Al2O3 layer are rinsed in distilled water in order to rinse out the electrolyte. Graphite was introduced into the structure of aluminium oxide during the thermal treat- ment in a solid medium consisting of graphite dust. The structure of the duplex-type layers contains carbon precipitates and other precipitates, which are typical for the aluminium alloy, with additions of Fe, Mn, Cr and other elements. Carbon precipitates have a relatively weak connection with the matrix, as an envelope with numerous discontinuities forms around each carbon precipitate. Surface layers obtained via the duplex method can have better wear resistance, a lower friction coefficient used to cover cylinders of lubricant-free compressors, pneumatic servo-motors or shock absorber components due to the increase of mechanical properties and the increased graphite content in the structure of the layers (which causes a decrease in the motion resistance of the kinematic nodes). The obtained aluminium-oxide-based, duplex-type surface layers fully confirm the usefulness of the new surface-treatment technologies in increasing the operational durability of the sliding couples of piston machines. 5 REFERENCES 1 W. Skoneczny, Model of structure of Al203 layer obtained via hard anodizing method, Surface Engineering, 17 (2001), 389–392, doi:10.1179/sureng.026708401101518060 2 W. Skoneczny, J. Jurasik, A. Burian, Investigations of the surfaces morphology of Al2O3 layers by atomic force microscopy, Materials Science Poland, 3 (2004), 265–278 3 T. Kmita, W. Skoneczny, Gradient layers on aluminum alloys created electrolytically, Chemical and Process Engineering, 26 (2005), 735–744 4 T. Kmita, J. Szade, W. Skoneczny, Gradient oxide layers with an in- creased carbon content on en AW-5251 alloy, Chemical and Process Engineering, 29 (2008), 375–387 5 W. Skoneczny, M. Bara, Aluminum oxide composite layers obtained by the electrochemical method in the presence of graphite, Material Science Poland, 25 (2007) 4, 1053–1062 6 M. Bara, W. Skoneczny, M. Hajduga, Ceramic-graphite surface layers obtained by the duplex method on an aluminum alloy sub- strate, Chemical and Process Engineering, 30 (2009), 431–442 7 S. Jeong, H. Hwang, S. Hwang, K. Lee, Carbon nanotubes based on anodic aluminum oxide nano-template, Carbon, 42 (2004), 2073–2080, doi:10.1016/carbon 2004.04.015 8 K. Wada, T. Shimohina, M. Yamada, N. Baba, Microstructure of porous anodic oxide films on aluminum, Journal of Materials Science, 21 (1986), 3810–3816, doi:10.1007/J.MSBF00553432 9 J. Mikulskas, S. Joudkazis, S. Jagminas, S. Meskins, J. G. Dumas, J. Vaitkus, R. Tomasiunas, Aluminum oxide film for 2D photonic structure; room temperature formation, Optical Materials, 17 (2001), 343–346, doi:10.1016/Opt.MatS0925-3467(01)00100-8 10 J. De Leat, H. Terryn, J. Vereecken, Development of an optical model for steady state porous anodic films on aluminium formed in phosphoric acid, Thin Solid Films, 320 (1997), 241–252, doi:10.1016/S0040-6090(97)00741-4 W. SKONECZNY: THE STRUCTURE AND MORPHOLOGY OF THE SURFACE OF DUPLEX LAYERS ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 845–850 849 Figure 7: XRD spectra of the: a) substrate, b) aluminium oxide and c) a duplex layer Slika 7: Rentgenogram: a) podlaga, b) aluminijev oksid in c) dupleks- na plast 11 M. J. Bartolome, V. Lopez, E. Escudero, G. Caruana, J. A. Gonzales, Changes in the specific surface area of porous aluminium oxide films during sealing, Surface & Coatings Technology, 200 (2006), 4530–4537, doi:10.1016/j.surfcoat.2005.03.019 12 V. Lopez, E. Otero, A. Bautista, J. Gonzales, Sealing of anodic films obtained in oxalic acid baths, Surface and Coatings Technology, 124 (2000), 76–84, doi:10.1016/S0257-8972(99)00626-X 850 Materiali in tehnologije / Materials and technology 50 (2016) 6, 845–850 W. SKONECZNY: THE STRUCTURE AND MORPHOLOGY OF THE SURFACE OF DUPLEX LAYERS ... E. MALEKI, A. ZABIHOLLAH: MODELING OF SHOT-PEENING EFFECTS ON THE SURFACE PROPERTIES ... 851–860 MODELING OF SHOT-PEENING EFFECTS ON THE SURFACE PROPERTIES OF A (TiB + TiC)/Ti–6Al–4V COMPOSITE EMPLOYING ARTIFICIAL NEURAL NETWORKS MODELIRANJE VPLIVA HLADNEGA POVR[INSKEGA KOVANJA NA LASTNOSTI POVR[INE (TiB + TiC)/Ti-6Al-4V KOMPOZITA S POMO^JO UMETNIH NEVRONSKIH MRE@ Erfan Maleki, Abolghassem Zabihollah Sharif University of Technology, International Campus, Department of Mechanical Engineering, Kish Island, 7941776655, Iran maleki_erfan@kish.sharif.edu, maleky.erfan@gmail.com Prejem rokopisa – received: 2015-06-29; sprejem za objavo – accepted for publication: 2015-11-13 doi:10.17222/mit.2015.140 Titanium matrix composites (TMCs) have wide application prospects in the field of aerospace, automobile and other industries because of their good properties, such as high specific strength, good ductility, and excellent fatigue properties. However, in order to improve their fatigue strength and life, crack initiation and growth at the surface layers must be suppressed using surface treatments. Shot peening (SP) is an effective surface mechanical treatment method widely used in industry which can improve the mechanical properties of a surface. However, artificial neural networks (ANNs) have been used as an efficient approach to predict and optimize the science and engineering problems. In the present study the effects of SP on TMC were modeled by means of ANN and the capability of the ANN in predicting the output parameters is investigated. A back-pro- pagation (BP) error algorithm is developed for the network training. Data of experimental tests on the (TiB + TiC)/Ti–6Al–4V composite are employed in order to train the network. The volume fractions of the reinforcements (TiB + TiC) were 5 % and 8 %. ANN testing is accomplished using different experimental data thaat were not used during the network training. The distance from the surface (depth) and SP intensity are regarded as input parameters and residual stress and hardness of the Ti–6Al–4V before and after the SP and adding reinforcements are gathered as the output parameters of the network. A comparison was made between experimental and predicted data. The predicted results were in good agreement with experi- mental ones, which indicates that developed neural network can be used for modeling the SP process on TMCs. Keywords: titanium matrix composites, surface treatment, shot peening, artificial neural networks, residual stress, hardness Kompoziti na osnovi titana (TMCs) imajo {iroko mo`nost uporabe na podro~ju letalstva, avtomobilske in druge industrije zaradi njihovih dobrih lastnosti, kot so: velika specifi~na trdnost, dobra duktilnost in odli~na odpornost na utrujanje. Vseeno pa je za pove~anje odpornosti na utrujanje in `ivljenjsko dobo, potrebna povr{inska obdelava, da se zavre nastanek razpok in njihova rast na povr{ini. Hladno povr{insko kovanje (SP) je u~inkovita mehanska metoda, ki se v industriji pogosto uporablja za izbolj{anje mehanskih lastnosti povr{ine. Umetne nevronske mre`e (ANNs) se uporabljajo kot u~inkovit pribli`ek za napovedovanje in optimiranje znanstvenih osnov in in`eniringa tega problema. V {tudiji so bili modelirani vplivi SP na TMC s pomo~jo ANN in preiskovana je bila zmo`nost napovedovanja izhodnih parametrov z ANN. Za usposabljanje mre`e je bil razvit algoritem vzvrat- nega {irjenja napak (BP). Podatki iz eksperimentalnih preizkusov na (TiB + TiC)/Ti–6Al–4V kompozitu so uporabljeni za usposabljanje mre`e. Volumska dele`a delcev (TiB + TiC) za oja~anje sta bila 5 % in 8 %. ANN preizku{anje je bilo izvedeno z uporabo razli~nih eksperimentalnih podatkov, ki niso bili uporabljeni pri usposabljanju mre`e. Razdalja od povr{ine (globina) in intenziteta SP sta uporabljeni kot vhodna parametra, preostala, napetost in trdota Ti–6Al–4V, pred in po SP, in dodatku delcev za oja~anje, sta izbrana kot izhodna parametra mre`e. Izvedena je bila primerjava med eksperimentalnimi in predvidenimi podatki. Predvideni rezultati so se dobro ujemali z eksperimentalnimi, kar ka`e na to, da se razvito nevronsko mre`o lahko uporabi pri modeliranju SP postopka na TMC. Klju~ne besede: kompoziti na osnovi titana, obdelava povr{ine, hladno kovanje povr{ine, umetna nevronska mre`a, zaostale napetosti, trdota 1 INTRODUCTION Titanium matrix composites (TMCs) have attracted considerable interest due to their attractive properties over titanium alloys, such as high elastic modulus, high strength, superior creep and fatigue resistances at ambient and elevated temperatures.1–4 The fabrication of TMCs using in-situ technology is simple and does not result in the pollution of an interface.5,6 TMCs can be reinforced with continuous fibers, whiskers or particles.7 As is well known, the mechanical properties of the com- posites depend on matrix, reinforcement and reinforce- ment/matrix interface, which bonds the formers to- gether.8 Compared with continuous fibers, TMCs rein- forced with whiskers or particles exhibit more isotropic behaviors. The fabrication of these materials is more convenient and cost effective; therefore, they have drawn extensive attention recently.9 Titanium monoboride (TiB) whiskers and titanium carbide (TiC) particles offer high modulus, relative chemical stability, and high thermal stability, while maintaining similar density and thermal expansion coefficient to those of the titanium matrix, as well as clean interfaces without any unfavorable reaction between the precipitates and the titanium matrix.10–12 The Materiali in tehnologije / Materials and technology 50 (2016) 6, 851–860 851 UDK 62-4:669.018.25:004.032.26 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)851(2016) reinforcements were obtained according to the high-tem- perature reactions as follows in Equations (1) and (2):13,14 5Ti+B4C = 4TiB+TiC (1) Ti+C = TiC (2) TMCs co-reinforced with TiB whiskers and TiC particles have been fabricated and investigated for their mechanical properties15–17 and have been extensively demonstrated to possess superior mechanical properties over the single TiB or TiC reinforced discontinuously reinforced titanium matrix composites (DRTMCs).18–21 As an effective and important surface-treatment method, shot peening (SP) can introduce high residual compressive stress (RCS) and microstructure variation at near surface layers, which can enhance their fatigue pro- perties compared to non-peened materials. The process of SP involves the bombardment of spherical balls of a hard material against the surface of components, which induces the strong elastic–plastic deformation at the sur- face and sub-surface regions. In the deformation layers, high RCS and microstructure refinements are introduced after SP. The residual stresses and hardness are very important properties of materials after the SP treatment. In the field of science and engineering, artificial neural networks (ANNs) are some of the most important research areas. ANN is a modeling tool to solve linear and nonlinear multivariate regression problems.22 Recently, ANN models were widely utilized to interpret and correlate the variable relationships in complex non- linear data sets. The present study proposes a new approach based on ANNs to investigate the effects of SP process on mechanical and metallurgical properties (TiB + TiC)/Ti–6Al–4V composite. Residual stress and hard- ness were modeled by ANN. 20 data of experimental tests results from the total of 30, are used to train the networks, while in the networks testing 10 different experimental data which were not used during training are used. Since the experimental results did not include the training sets the performance of the ANN is eva- luated in a fine way. 2 EXPERIMENTAL PART The experimental data are obtained from Xie et al.23 The materials of (TiB + TiC)/Ti–6Al–4V (TiB:TiC = 1:1 (vol.%)) were fabricated via in-situ technology. Two types of theoretical total volume fraction of reinforce- ments (TiB + TiC) were 5 % and 8 %. The SP treatment was performed using an air-blast machine. The related information of used SP process is demonstrated in Table 1. The Almen specimens are A type and the diameter of peening nozzle was 15 mm and the distance between nozzle and sample was 100 mm. In order to obtain the depth distribution of the residual stress and hardness, the thin top surface layers were removed one by one via the method of chemical etch with a solution of water, nitric acid, and hydrofluoric acid in the ratio 31:12:7. All the measurements were carried out at room temperature. The method of residual stress and hardness measurements are X ray stress analysis and digital microhardness test res- pectively.23 Table 1: Parameters of the SP process treatments23 Tabela 1: Parametri uporabljenega procesa SP23 SP intensity (mm A) Shot material Shot diameter (mm) Shot hardness (HV) Jet pressure (MPa) SP time (min) Cover- age (%) 0.15 Caststeel 0.6 610 0.2 0.50 100 0.30 Caststeel 0.6 610 0.3 0.50 100 The results indicate that the increased reinforcements and SP intensities enhance the surface roughness after SP. Both the compressive residual stresses and hardness increase with the increase of the SP intensity, which is mainly due to the plastic deformation and high disloca- tion density in the near surface layer. Moreover, the rein- forcement particles can act as the block sources during dislocation movements. After an appropriate SP treat- ment, the increased CRS and hardness are beneficial to industrial applications A table shows the obtained values of the experimental results on (TiB + TiC)/Ti–6Al–4V for 30 different samples. The SP intensity for non- peened specimens has been shown by zero in Table 2. 3 ARTIFICIAL NEURAL NETWORKS Artificial intelligence (AI) systems such as artificial neural networks (ANNs) have found many applications in science and engineering problems in the past decade. The concept of an ANN has emerged with the idea that it E. MALEKI, A. ZABIHOLLAH: MODELING OF SHOT-PEENING EFFECTS ON THE SURFACE PROPERTIES ... 852 Materiali in tehnologije / Materials and technology 50 (2016) 6, 851–860 Figure 1: Schematic of neuron: a) a biological neuron, b) an artificial neuron Slika 1: Shematski prikaz nevrona: a) biolo{ki nevron, b) umetni ne- vron simulates the operating principles of a human brain. The first studies were made with mathematical modeling of biological neurons that make up the brain cells.24 Basi- cally, the brain functions with a very dense network of neurons. The brain contains a lot of neurons connected to each other by many interconnections. A neuron consists mainly of the following parts: dendrite, cell body and axon.25 Dendrite gets the signals from various other neurons to the neuron and carries them to the cell body for processing, after that an axon carries the signal from the cell body to various other neurons. Similarly, the neural units in the artificial neural network are developed as a very approximate model of the natural biological neurons.26 Figure 1 shows a natural biological neuron (Figure 1a) and an artificial neuron (Figure 1b) that is a computational and mathematical model of the biological neuron. A single neuron computes the sum of its inputs, which are multiplying with a variant called the weight, adds a bias term, and drives the result through a generally nonlinear transfer function to produce a single output termed the activation level of the neuron. An ANN model is created by interconnection of many of the neurons in a known configuration. The primary elements characterizing the neural network are the distributed representation of information, local operations and non-linear processing. Structurally, every ANN is made up of three sections: input, hidden and output layers.27 The structure of an ANN model is deter- mined by the number of its layers and the number of nodes in each layer and the nature of the transfer function.28,29 Selecting the optimum architecture of the network is one of the challenging steps in ANN mo- deling. The term “architecture” refers to the number of layers in the network and the number of neurons in each layer. However, there is no straightforward method to estimate the optimal number of hidden layers and neu- rons in each layer.30,31 Thus trial-and-error methods have been used by many researchers to determine such case- dependent parameters for studies involving ANN-based models.32 Figure 2 represents the architecture of the neural network. In this network, each input consists of r parameters and each output comprises s parameters, E. MALEKI, A. ZABIHOLLAH: MODELING OF SHOT-PEENING EFFECTS ON THE SURFACE PROPERTIES ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 851–860 853 Table 2: Values of the SP process effects on residual stress and hardness of (TiB + TiC)/Ti–6Al–4V Composite specimens23 Tabela 2: Vrednosti SP postopka, ki vplivajo na zaostale napetosti in trdoto (TiB + TiC) / Ti-6Al-4V kompozitnih vzorcev23 Sample No. Depth SP intensity(mm A) Residual Stress (MPa) Hardness (HV) matrix 5 % (TIB+TIC) 8 % (TIB+TIC) matrix 5 % (TIB+TIC) 8 % (TIB+TIC) 1 0 0.00 10.42 18.04 17.93 334.72 380.37 417.57 2 0 0.30 -522.83 -524.25 -575.51 524.87 560.38 637.31 3 0 0.15 -375.97 -434.55 -481.94 484.31 512.07 584.43 4 15 0.00 25.11 5.022 6.84 328.67 393.61 436.62 5 15 0.30 -613.76 -648.62 -657.79 436.11 512.19 628.86 6 15 0.15 -417.53 -499.54 -539.63 418.45 492.85 523.28 7 25 0.00 -9.57 -20.26 -12.97 325.48 381.28 420.16 8 25 0.30 -608.67 -650.57 -672.54 414.97 468.23 557.85 9 25 0.15 -408.92 -465.67 -574.80 387.62 451.56 507.09 10 50 0.00 14.35 -29.48 -9.19 315.39 403.32 411.77 11 50 0.30 -581.27 -608.46 -626.77 378.62 455.55 530.79 12 50 0.15 -397.63 -419.90 -545.84 372.53 431.43 475.18 13 75 0.00 14.48 14.02 -12.72 338.28 381.40 423.67 14 75 0.30 -564.84 -570.02 -586.49 385.38 438.64 538.40 15 75 0.15 -354.00 -386.95 -516.88 358.29 420.55 446.63 16 100 0.00 -12.84 -17.01 -14.41 340.03 396.67 414.43 17 100 0.30 -557.53 -524.25 -515.10 368.47 433.57 509.66 18 100 0.15 -324.71 -337.52 -264.28 349.94 411.35 459.31 19 150 0.00 -5.26 -11.81 -8.67 330.01 375.66 409.48 20 150 0.30 -422.37 -335.69 -299.08 350.72 388.76 438.64 21 150 0.15 -302.74 -249.65 -120.90 340.80 386.23 435.87 22 200 0.00 16.96 -13.89 -4.75 319.14 386.77 429.04 23 200 0.30 -323.74 -220.36 -101.37 354.10 402.29 434.42 24 200 0.15 -236.84 -194.73 -61.16 344.29 382.99 424.21 25 250 0.00 28.21 14.93 17.41 337.86 375.06 424.09 26 250 0.30 -46.11 -20.82 -28.14 343.11 380.31 428.50 27 250 0.15 -31.80 -18.99 -15.97 336.83 378.90 410.87 28 300 0.00 -13.62 -23.50 21.33 344.74 369.26 414.91 29 300 0.30 -35.15 -18.99 -33.63 324.51 381.15 430.88 30 300 0.15 -18.99 -11.67 -16.23 326.86 376.50 419.41 while p, w, b, f and a represent the inputs, weight matrixes, bias vectors, transfer function in neurons, and outputs, respectively.33 Mathematically, a layer n may be described by Equations (3) and (4):34 u w ps n s n i r i= = ∑ 1 (3) a f v f u bs s n s n s n= = +( ) ( ) (4) where p1, p2, . . ., pr are the input signals, wk1, wk2, . . ., wns are the synaptic weights of neuron n, un is the linear combiner output due to input signals, bn is the bias, f is the transfer function and a1, a2, . . ., as are the output signals of the neuron. The tangent sigmoid (Tansig) (x), logarithmic sigmoid (Logsig) (x) and linear (x) transfer function are described as follows in Equations (5), (6) and (7):35 ( )x e x = + −− 2 1 12 (5) ( )x e x = + − 2 1 (6) ( ) ( )x linear x= (7) 3.1 Training of ANN The training of the ANN is performed by adjusting the connection weights. It is performed by iteratively adjusting the weights (w) of the connections and biases (b) in the network in order to minimize a predefined cost function.36 The ANNs are trained with a training set of input and known output data. An ANN is better trained as more input data are used. The performance of an ANN is generally based on the parameters’ architecture and the setting. As was mentioned, one of the most difficult tasks in studying ANNs is finding an appropriate architecture. This task is performed via trial and error and the number of middle layers and neuron presented in each layer is being identified. Appropriate designation of the initial amounts of weights and biases is very effective on the performance of network and the time of calculation. But there is not a reasonable law and process to identify a suitable architecture. The only step which is very time consuming is the trial and error. One can get an idea by looking at a problem and decide to start with simple networks; going on to complex ones until the solution is within acceptable limits of uncertainty. Furthermore, the point that must be considered in training of the network is the rate of input and output data scattering. In this study all values of each input and output data parameters are divided to maximum absolute value of them and normalized also the used data are dimensionless. The normalized data are in range of [–1, +1]. In the present study, a feed forward ANN based on back propagation (BP) error algorithm, which is the most popular one in training of ANNs is used. BP is a descent algorithm, which attempts to minimize the error during iterations. The weights of the network are adjusted by the algorithm such that the error is decreased along a descent direction. In the back-propagation learning, the actual outputs are compared with the target values to derive the error signals, which are propagated backward layer by layer for the updating of the synaptic weights in all the lower layers. E. MALEKI, A. ZABIHOLLAH: MODELING OF SHOT-PEENING EFFECTS ON THE SURFACE PROPERTIES ... 854 Materiali in tehnologije / Materials and technology 50 (2016) 6, 851–860 Figure 3: A Conceptual structure of network with four layers Slika 3: Konceptualna zgradba mre`e s {tirimi nivoji Figure 2: Architecture of neural network33 Slika 2: Zgradba nevronske mre`e33 3.2 Implementation of ANN In this paper the effects of the SP process on surface properties including of (TiB + TiC) / Ti–6Al–4V com- posite were modeled by means of an ANN. In implemen- tation of the ANN distance from the surface (depth) and SP intensity are regarded as inputs and the residual stress and hardness are gathered as outputs of the networks. Different networks with different architecture and net- work parameters were trained for the prediction of resi- dual stress and hardness. Figure 3, for an example, re- presents the schematic architecture of ANN for modeling of the mentioned output parameters: a four-layer feed forward with BP algorithm with full interconnection. This neural network model has a powerful input-output mapping capability. With the use of enough hidden neurons, it can effectively approximate any continuous nonlinear function. In the considered network, two inputs are logged into the input layer to determine the two outputs. In the ANN methodology, the sample data is often subdivided into training and testing sets. The distinctions among these subsets are crucial.37 Ripley defines the following: Training set: a set of examples used for learning that is to fit the parameters of the classifier. Testing set: a set of examples used only to assess the performance of a fully-specified classifier. 3.3 Performance evaluation of ANN The performance of the ANN models in predicting the shot-peening effects on residual stress and hardness of (TiB + TiC)/Ti–6Al–4V composite were statistically evaluated using four prediction score metrics calculated from the test dataset: Pearson coefficient of correlation (PCC), root mean square error (RMSE), mean relative error (MRE) and mean absolute error (MAE). These parameters were determined using the following Equations (8), (9), (10) and (11): PCC f F f F f F EXP i EXP ANN i ANN i n EXP i EXP = − − − − = ∑ ( ) ( ) ( ) , , , 1 2 − − = ∑ ( ),f FANN i ANN i n 2 1 (8) RMSE f f n EXP i ANN i n = − = ∑ ( ), 2 1 (9) MRE n f f f EXP i ANN EXP ii n = − × = ∑1 1001 1 , , , (10) MRE n f fEXP i ANN i n = − = ∑1 1 1 , , (11) where n is the number of used sample for modeling, fEXP is the experimental value and fANN is the networks predicted value. Also, the values of FEXP and FANN are calculating as follows in Equations (12) and (13): F n fEXP EXP i i n = = ∑1 1 , (12) F n fANN ANN i i n = = ∑1 1 , (13) 3.4 Generating model function After the neural network is trained successfully with four layers, the values of the four parameters of the net- work (p, b, w and f) can be obtained. The function that correlates the inputs to the corresponding output can be calculated by applying the aforementioned parameters. Finally, the model function can determined in Equations (14) and (15): a1= f1(w1p+b1) (14a) a2= f2(w2p1+b2) (14b) a3= f3(w3p2+b3) (14c) a4= f4(w4p3+b4) (14d) G (g(1), g(2))= a4= = f4(w4f3 (w3 2(w2f1(w1p+b1)+ b2) +b3)+ b4) (15) where a1, a2 and a3 are the outputs of the first, second and third layer, respectively; a4 is the fourth layer out- put, which is equal to the function G (g(1), g(2)). The function G gets the values of the input parameters. The function of g(1)and g(2) represent the residual stress and hardness, respectively. The methodology used for neural network application in this study is as follows: 1. Start; 2. Normalize the data (inputs & outputs); 3. Feed the data to artificial neural network; 4. Find network optimum parameters; 5. Execute network training; 6. Obtain Pearson correlation coefficient; 7. If PCC = 0.99 go to 8, if not go back to 4 with re- vising the parameters of network; 8. Continue processing until obtaining desired conver- gence between experimental and predicted values; 9. Obtain weights & biases values; 10. Create the model function; 11. Conduct analysis based on model function; 12. Verify the results using experimental values; 13. Calculate the error for each answer; 14. End. 4 RESULTS AND DISCUSSION In order to train the ANNs in this study, the obtained experimental test results on shot peened (TiB + TiC)/Ti–6Al–4V composite specimens are employed. Different networks were trained to achieve the optimum structure (OS) in order to generate a model function (MF). After the OS is selected and the MF is generated, operation of the network is tested with the use of them (OS & MF). Twenty sample data (data of samples 1-20) E. MALEKI, A. ZABIHOLLAH: MODELING OF SHOT-PEENING EFFECTS ON THE SURFACE PROPERTIES ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 851–860 855 were used from the total of 30, as data sets for network training. Table 3 shows the normalized sample data used for networks training. In the network testing, 10 different sample data (data of samples 21–30) which were not used during training are employed. Therefore, the whole experimental results did not comprise in the training. For investigative purposes, out of 30 samples data, 67 % data had taken for training and 33 % data for testing. Several networks have been trained with different architecture to find the OS of ANNs, to predict the regarded outputs, with the best performance and the highest PCC, the least RMSE, MRE and MAE. Related information of some the different trained networks for modelling of matrix hardness are shown in Table 4. The ordinal numbers shown in the "Layer Structure" were used to indicate the total number of neurons in the input, hidden and output layers, respectively. Results of the networks were investi- gated and the ANN modelling number 11 with 2×8×16×2 structure is selected for modelling and simu- lation. Figures 4 and 5 show the obtained values of the ANN response in comparison with experimental values for each 20 training samples (samples 1–20) for residual stress and hardness, respectively, using the selected network. After the network was trained, the selected network is tested. Figures 6 and 7 have been demonstrated the predicted and experimental values of residual stress and hardness for 10 different testing samples (samples 21–30) respectively. E. MALEKI, A. ZABIHOLLAH: MODELING OF SHOT-PEENING EFFECTS ON THE SURFACE PROPERTIES ... 856 Materiali in tehnologije / Materials and technology 50 (2016) 6, 851–860 Table 3: Normalized sample data used for networks training Tabela 3: Normalizirani podatki vzorca, uporabljenega pri usposabljanju mre`e Sample No. Depth SP intensity Residual stress Hardness matrix 5 %(TIB+TIC) 8 % (TIB+TIC) matrix 5 % (TIB+TIC) 8 % (TIB+TIC) 1 0.0000 0.0 0.0170 0.0277 0.0267 0.6377 0.6788 0.6552 2 0.0000 1.0 -0.8518 -0.8058 -0.8557 1.0000 1.0000 1.0000 3 0.0000 0.5 -0.6126 -0.6680 -0.7166 0.9227 0.9138 0.9170 4 0.0500 0.0 0.0409 0.0077 0.0102 0.6262 0.7024 0.6851 5 0.0500 1.0 -1.0000 -0.9970 -0.9781 0.8309 0.9140 0.9867 6 0.0500 0.5 -0.6803 -0.7678 -0.8024 0.7972 0.8795 0.8211 7 0.0833 0.0 -0.0156 -0.0311 -0.0193 0.6201 0.6804 0.6593 8 0.0833 1.0 -0.9917 -1.0000 -1.0000 0.7906 0.8356 0.8753 9 0.0833 0.5 -0.6663 -0.7158 -0.8547 0.7385 0.8058 0.7957 10 0.1667 0.0 0.0234 -0.0453 -0.0137 0.6009 0.7197 0.6461 11 0.1667 1.0 -0.9471 -0.9353 -0.9319 0.7214 0.8129 0.8329 12 0.1667 0.5 -0.6479 -0.6454 -0.8116 0.7098 0.7699 0.7456 13 0.2500 0.0 0.0236 0.0216 -0.0189 0.6445 0.6806 0.6648 14 0.2500 1.0 -0.9203 -0.8762 -0.8721 0.7342 0.7828 0.8448 15 0.2500 0.5 -0.5768 -0.5948 -0.7685 0.6826 0.7505 0.7008 16 0.3333 0.0 -0.0209 -0.0261 -0.0214 0.6478 0.7079 0.6503 17 0.3333 1.0 -0.9084 -0.8058 -0.7659 0.702 0.7737 0.7997 18 0.3333 0.5 -0.5291 -0.5188 -0.3930 0.6667 0.7341 0.7207 19 0.5000 0.0 -0.0086 -0.0182 -0.0129 0.6287 0.6704 0.6425 20 0.5000 1.0 -0.6882 -0.5160 -0.4447 0.6682 0.6937 0.6883 Table 4: Relevant information of 12 different networks for modeling of matrix hardness Tabela 4: Pomembne informacije o 12 razli~nih mre`ah pri modeliranju trdote osnove ANN Modeling no. Rate of training Layers structure Hidden transfer function Output transfer function PCC RMSE MRE (%) MAE 1 0.090 2×2×4×2 Logsig Linear 0.97035 0.7677 0.1552 0.6680 2 0.095 2×2×6×2 Tansig Linear 0.97421 0.7018 0.1399 0.5742 3 0.110 2×2×8×2 Logsig Tansig 0.98460 0.6671 0.1007 0.5018 4 0.100 2×4×4×2 Tansig Linear 0.98662 0.4163 0.0938 0.4261 5 0.115 2×4×6×2 Logsig Linear 0.99003 0.2397 0.0875 0.3459 6 0.120 2×4×10×2 Logsig Linear 0.99150 0.2078 0.0617 0.2822 7 0.115 2×6×10×2 Tansig Tansig 0.99877 0.3400 0.0587 0.2401 8 0.130 2×6×12×2 Tansig Linear 0.99901 0.2229 0.0461 0.1886 9 0.145 2×6×16×2 Logsig Tansig 0.99936 0.1997 0.0384 0.1597 10 0.160 2×8×10×2 Logsig Logsig 0.99911 0.1609 0.0331 0.1529 11 0.165 2×8×16×2 Logsig Logsig 0.99979 0.0985 0.0194 0.0853 12 0.165 2×8×20×2 Tansig Linear 0.99963 0.1265 0.0247 0.1173 E. MALEKI, A. ZABIHOLLAH: MODELING OF SHOT-PEENING EFFECTS ON THE SURFACE PROPERTIES ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 851–860 857 Figure 7: Comparison of predicted values (ANN response) with experimental values for each 20 testing samples (samples 21–30) for hardness: a) matrix, b) 5 % TIB+TIC and c) 8 % TIB+TIC Slika 7: Primerjava napovedanih vrednosti (odgovor ANN) z eksperimentalnimi vrednostmi za vsakega od 20 preizku{enih vzorcev (vzorci 21-30) za trdoto: a) osnova, b) 5 % TiB+TiC in c) 8 % TiB+TiC Figure 6: Comparison of predicted values (ANN response) with experimental values for each 20 testing samples (samples 21–30) for residual stress: a) matrix, b) 5 % TIB+TIC and c) 8 % TIB+TIC Slika 6: Primerjava napovedanih vrednosti (odgovor ANN) z eksperimentalnimi vrednostmi za vsakega od 20 preizku{enih vzorcev (vzorci 21-30) za zaostale napetosti: a) osnova, b) 5 % TiB+TiC in c) 8 % TiB+TiC Figure 5: Comparison of predicted values (ANN response) with experimental values for each 20 training samples (samples 1–20) for hardness: a) matrix, b) 5 % TiB+TiC and c) 8 % TiB+TiC Slika 5: Primerjava napovedanih vrednosti (odgovor ANN) z eksperimentalnimi vrednostmi za vsakega od 20 vzorcev usposabljanja (vzorci 1-20) za trdoto: a) osnova, b) 5 % TiB+TiC in c) 8 % TiB+TiC Figure 4: Comparison of predicted values (ANN response) with experimental values for each 20 training samples (samples 1–20) for residual stress: a) matrix, b) 5 % TiB+TiC and c) 8 % TiB+TiC Figure 4: Primerjava napovedanih vrednosti (odgovor ANN) z eksperimentalnimi vrednostmi za vsakega od 20 vzorcev usposabljanja (vzorci 1-20) za zaostale napetosti: a) osnova, b) 5 % TiB+TiC in c) 8 % TiB+TiC Figure 8 illustrates the obtained relative error (RE) values of the residual stress and hardness for the testing samples. In modeling of residual stress, according to the Figure 8a, the minimum and maximum (min., max.) determined relative errors for matrix, 5 % reinforcement and 8 % reinforcement are (0.000,0.2633), (0.0103, 0.1714) and (0.0099, 0.2616), respectively. Based on Figure 8b, similarly minimum and maximum calculated REs for matrix, 5 % reinforcement and 8 % reinforce- ment in modeling of hardness are (0.0015, 0.0587), (0.0026, 0.0870) and (0.0023, 0.0919), respectively. According to the obtained values of the ANN for training and testing samples, data corresponding to the used network are shown in Table 5. In network training it is observed that the values of PCC for each considered output parameters are more than 99.7 %. The values of training RMSE, MRE and E. MALEKI, A. ZABIHOLLAH: MODELING OF SHOT-PEENING EFFECTS ON THE SURFACE PROPERTIES ... 858 Materiali in tehnologije / Materials and technology 50 (2016) 6, 851–860 Figure 9: Distribution of residual stresses along the depth from the surface obtained by OS and MF of ANN for different SP intensities of: a) 0 mm A, b) 0.15 mm A and c) 0.3 mm A Slika 9: Razporeditev spreminjanja zaostalih napetosti v globino od povr{ine, dobljene pri OS in MF z ANN pri razli~nih intenzitetah SP: a) 0 mm A, b) 0,15 mm A in c) 0,3 mm A Figure 8: Values of obtained relative error for testing samples (sam- ples 21–30) for considered output parameters: a) residual stress, b) hardness Slika 8: Vrednosti dobljene relativne napake preizku{anih vzorcev (vzorci 21-30) pri upo{tevanih izhodnih parametrih: a) zaostale nape- tosti, b) trdota Table 5: Obtained values of PCC, RMSE, MRE and MAE for trained and tested network Tabela 5: Dobljene vrednosti PCC, RMSE, MRE in MAE pri usposobljeni in pri preizku{eni mre`i Output parameter Training Testing PCC RMSE MRE (%) MAE PCC RMSE MRE (%) MAE Residual stress Matrix 0.99914 0.1781 0.1104 0.1099 0.99846 0.2529 0.1223 0.1315 5 % rein. 0.99875 0.0651 0.0872 0.0471 0.99816 0.0906 0.0910 0.0521 8 % rein. 0.99766 0.0597 0.1228 0.0382 0.99683 0.0647 0.1342 0.0440 Hardness Matrix 0.99979 0.0985 0.0194 0.0853 0.99912 0.1154 0.0276 0.0937 5 % rein. 0.99901 0.1544 0.0313 0.1372 0.99858 0.1783 0.0368 0.1420 8 % rein. 0.99837 0.1976 0.0346 0.1475 0.99721 0.2071 0.0419 0.1780 MAE are very close to 0 and they are in little intervals and their ranges are [0.0597, 0.1976], [0.0194, 0.1228] and [0.0382, 0.14715], respectively. So, it is concluded that networks are trained finely and adjusted carefully. Likewise in network testing the values of PCC are more than 99.6 % and it is observed that values of testing PCC have a negligible reduction in comparison with the train- ing. Moreover, the values of testing RMSE, MRE and MAE are in a tiny span as well and their ranges are [0.0647, 0.2071], [0.0276, 0.1342] and [0.0440, 0.1780], respectively. Based on the achieved values for the statistical errors for both training and testing samples it is concluded that the error values are acceptable and implementation of ANN is accomplished in a good way. In residual stress modeling for each case of network training and testing, the obtained values for 5 % TiB + TiC., matrix and 8 % TiB + TiC and in modeling of hardness, achieved values for matrix, 5 % TiB + TiC and 8 % TiB + TiC have the smallest errors, respectively. Distributions of the residual stress and hardness from the shot-peened surface to the bulk material (25-300 μm) for SP intensity of (0.00, 0.15 and 0.30) are shown in Figures 9 and 10, which are achieved by OS and MF of the used ANN modeling in this paper. 5 CONCLUSION In present study the application of ANNs was investigated, aiming to create models to predict and optimize the SP process effects with different intensities on the residual stress and hardness of a (TiB + TiC)/Ti–6Al–4V composite. Experimental data show that both the residual stress and the hardness were increased with an improvement of the SP intensities. Residual stress and hardness were modeled using ANN for three cases: matrix, 5 % TiB + TiC and 8 % TiB + TiC. The obtained results indicate that statistical errors for RSME, MRE and MAE are in very small range and so close to 0. Moreover, the values of PCC for all of the regarded output parameters in implemented networks are more than 99 %. 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Mohagheghtabar, Application of artificial neural network (ANN) for the prediction of size of silver nanoparticles prepared by green method, Digest Journal of Nanomaterials and Biostructures, 8 (2013), 541–549 E. MALEKI, A. ZABIHOLLAH: MODELING OF SHOT-PEENING EFFECTS ON THE SURFACE PROPERTIES ... 860 Materiali in tehnologije / Materials and technology 50 (2016) 6, 851–860 A. CAN et al.: ANALYSIS OF TWIN-ROLL CASTING AA8079 ALLOY 6.35-μm FOIL ROLLING PROCESS 861–868 ANALYSIS OF TWIN-ROLL CASTING AA8079 ALLOY 6.35-μm FOIL ROLLING PROCESS ANALIZA PROCESA VALJANJA 6,35 μm FOLIJE IZ ZLITINE AA8079 ULITE MED DVEMA VALJEMA Ahmet Can1, Hüseyin Arikan2, Kadir Çýnar2 1Necmettin Erbakan University, Faculty of Engineering, Department of Industrial Design, Konya, Turkey 2Necmettin Erbakan University, Faculty of Engineering, Department of Mechanical Engineering, Seydiºehir, Konya, Turkey ahmetcan@konya.edu.tr Prejem rokopisa – received: 2015-06-29; sprejem za objavo – accepted for publication: 2015-12-16 doi:10.17222/mit.2015.134 In this work the rolling process and properties of a 6.35-μm twin-roll casting AA8079 aluminum alloy foil was analyzed. First, the 8-mm-thick sheets were produced with a twin-roll casting technology. This product was annealed and cold rolled to a 6.35-μm foil with suitable processing conditions. The mechanical tests and microhardness measurement was applied to specimens derived from all the foil-rolling process stages. On the other hand, the specimens’ surface roughness and the surface structure are visualized with an atomic force microscope and an SEM. The microstructural investigation is realized with an optical microscope and XRD. The von-Misses total effective strain was calculated by determining the incremental work for all of the cold-rolling cycle. The alloy showed very low ductility in the tensile tests because of the second-phase metastable intermetallic particles such as Al3Fe. The maximum elongation at the breaking value was measured for 256-μm-foil as 4.5 %. On the other hand, the alloy did not show any significant strain hardening after the cold rolling during the plastic-deformation stages. Keywords: aluminum foil, cold rolling, twin roll casting V delu je bil analiziran postopek valjanja in latnosti 6,35 folije iz AA8079 aluminijeve zlitine, ulite med dvema valjema. Najprej je bil izdelan 8 mm debel trak po postopku ulivanja med dvema valjema. Trak je bil primerno `arjen in hladno zvaljan v 6.35 μm folijo. Iz vseh stopenj procesa valjanja so bili vzeti vzorci na katerih so bile dolo~ene mehanske lastnosti in izmerjena mikro trdota. Poleg tega je bila hrapavost povr{ine in struktura povr{ine vizualizirana z mikroskopom na atomsko silo (AFM) in iz SEM. Preiskava mikrostrukture je bila izvr{ena s svetlobnim mikroskopom in z rentgensko difrakcijo (XRD). Za dolo~anje stopnjujo~ega dela, med celotnim ciklom hladnega valjanja, je bila izra~unana celotna von-Misses efektivna napetost. Zlitina je pokazala zelo nizko duktilnost pri nateznih preizkusih zaradi vsebnosti delcev sekundarne metastabilne intermetalne faze Al3Fe. Maksimalni raztezek pri poru{itvi je bil pri 256 μm debeli foliji 4.5 %. Po drugi plati pa zlitina ni kazala nobenega ob~utnega napetostnega utrjevanja med posameznimi fazami plasti~ne deformacije. Klju~ne besede: aluminijeva folija, hladno valjanje, ulivanje med dvema valjema 1 INTRODUCTION The production of aluminum alloys with twin roll casting (TRC) technology has been introduced to the industry about 50 years ago. It was claimed that TRC would offer significant reduction in the cost of aluminum sheet and foil production, compared to the conventional production technique, i.e. DC casting and hot rolling. Major evolution in the TRC technology has been attained in the last 5 years. It has been widely accepted due to its low investment cost, operational cost and flexibility provided to the production planning.1 The strengthening of metals due to increase in lattice defects during cold deformation makes a thermodyna- mically unstable structure and promotes subsequent restoration phenomena. The restoration processes can change microstructures as well as mechanical and physi- cal properties of metals and alloys while required me- chanical and physical properties may be achieved by adjusting the deformation and annealing variables.2 Some of the researchers3–6 studied about cold rolling of various aluminum alloys. D. Wang et al.3 studied about severe cold rolling (CR) deformation properties of AA 7050. The strength of the 7050 samples increased with increasing the CR reduction. The yield and ultimate strengths of the CR sample with a reduction of 67 % in- creased by 16.5 % and 9.2 %, respectively. Wang re- ported that both the residual dislocations and hetero- geneously nucleated fine-phase particles in the matrix increased the strength of the CR samples. Z. Liang et al.4 studied the evolution of texture as well as microstructures in an AA 7055 aluminum alloy during cold rolling. Author reported that more micro-bands are formed in the center of the plate with the rolling reduction, while the spacing between two bands decreases. S. X. Zhou et al.5 studied about cold Rolling of AA 1050 alloys which are produced by hot finishing rolling and twin roll casting. Authors studied the microstructure mechanical properties such as tensile strength, yield stress elongation area reduction, elastic modulus hardness and impact energy and formability. Materiali in tehnologije / Materials and technology 50 (2016) 6, 861–868 861 UDK 662.2.036:621.771.8:669.715 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)861(2016) J. G. Lenard6 studied the effect of roll roughness on the rolling parameters during cold rolling of 6061-T6 alloys. The effects of the roughness of the work roll on the roll force roll torque and the forward slip and additionally the frictional mechanisms were identified. Author reported that high roughness appeared to increase the possibility of insufficient lubrication at the interfaces. While both adhesive and ploughing forces were present in all instances, the ploughing forces became dominant at higher rolling speeds. The contribution of ploughing to frictional resistance increased as the roll roughness in- creased to a certain value and beyond that its behavior depended on the rolling speed. J. G. Lenard and S. Zhang7 studied a similar work with commercially pure alumi- num. Using lighter oil, boundary or mixed lubrication is produced. With higher viscosity oil, negative forward slip is observed, indicating the onset of hydrodynamic lubrication. The coefficients of friction are found to increase with increasing reduction and decreasing rolling speeds.7 K. S. S. Sathees et al.8 studied about with purity alu- minum sheets which were subjected to intense plastic straining by constrained Groove pressing method. The tensile behavior evolution with increased straining indi- cates substantial improvement of yield strength by 5.3 times from 17 MPa to 90 MPa during first pass corrobo- rated to grain refinement observed. Quantitative assess- ment of degree of deformation homogeneity using micro hardness profiles reveal relatively better strain homo- geneity at higher number of passes. G. Liv9 studied the development of surface structure forming properties and corrosion resistance during cold rolling of twin-roll cast of AA 3003. It is found that the as-cast surface determines the development of the sur- face topography during cold rolling. This is due to the large roughness associated with the groove/shingle con- figuration of the as-cast surface. Author pointed out that, the initial surface topography and the cold rolling are of great importance to the quality of the end product of the cold rolling. Since there are differences in the initial topography induced by the surface of the casting rolls, there will be differences in the development of the sur- face during the deformation sequence. The rough pattern of the as-cast surface results in large gorges in the first pass. In the second pass shingles are smeared out on top of the gorges. Patches of the shingle are not welded to the bulk sheet. It must be focused to metallurgical principles of Alu- minum alloys for determining the mechanical and forma- bility behavior of Al-Fe alloys. The main concern is pro- pagation of second phase intermetallic particles which is a function of both the cooling rate and the chemical com- position of Al–Fe alloys and their effects on mechanical properties. Some of the researchers focused the metallurgical behavior of Al-Fe powders and alloys.10–13 On the other hand some of the researchers focused the rapid cooling rate and severe deformation effects on the different Al-Fe alloys.14–16 M. Aghaie-Khafri and R. Mah- mudi17 have investigated the plastic instability and necking behavior of AA8079 aluminum alloy sheet in temper-annealed and fully annealed conditions. In this work, the rolling process and properties of 6.35 mm twin roll casting AA8079 aluminum alloy foil was analyzed. Foils production were tested on the level of industrial trials with no rupture by the time rolling process. Firstly the 8 mm thickness sheets were pro- duced by twin roll casting technology. This product was annealed and cold rolled to 6.35 μm foil with suitable process conditions. The mechanical tests and micro hard- ness measurement is applied to specimens derived from whole foil rolling process stages. On the other hand the specimens’ surface roughness and the surface structure is visualized with Atomic Force Microscope and SEM. 2 MATERIAL AND EXPERIMENTAL PART The material used in this experimental investigation was an aluminum-rich eutectic alloy AA8079. As de- rived from XRF analyses the alloy containing (in mass fractions, w/%) 99 % Al, 0.8 % Fe and 0.12 % Si with minor constituents of 0.02 % Cu, 0.02 % Mn, 0.009 % Zn, and 0.022 % Ti. Notice that this work is not only an experimental work. The whole experimental outputs were derived from a real industrial production process. So, nearly 2000 kg raw material of AA8079 was melted and roll-casted to 8 mm thickness and cold rolled to 4 mm at one step in CR (Cold Rolling Machine). Then the material homogenized at 580 °C for 8 h, furnace cooled, and then cold rolled to the thickness 0.53 mm with four steps. Then the material annealed in a furnace at 450 °C and a holding time of 4 h. Then cold rolled to the initial foil thickness 250 μm with one step. The 250 μm sheet was cold rolled to 100 and then 56 μm with two steps in FR-I rolling machine. The 56 μm foil was cold rolled to 14 μm with two steps in FR-II rolling machine. Then the final thickness foil was derived by cold rolling with twofold (14+14 μm foil on foil) foil in FR-III rolling machine. The diameter of the (CR) cold rolling pin was 400 mm, and the foil rolling (FRI-III)) pin was 240 mm. The Rolling speed of CR, FRI,FRII,FRIII were, 12-175- 325-560 m/min respectively. After final (6.35 μm) cold rolling, the annealing in a furnace at 270 °C and a hold- ing time of 11 h is applied. The experimental samples were derived from all rolling stages. All the experimental numeric outputs such as tensile-yield strength, elongation at break and surface roughness values verified with minimum 3 times re- peated tests results. These outputs were averaged and given with derived error band in the graphs. The micro hardness test was realized in NDT MH-140 under 5 g loading forces. The surface roughnesses of the foils were measured with Mitutoyo Surf Test SJ301 roughness measurement device. The 3D surface topography was determined with Park XE-100 Atomic Force Microscope A. CAN et al.: ANALYSIS OF TWIN-ROLL CASTING AA8079 ALLOY 6.35-μm FOIL ROLLING PROCESS 862 Materiali in tehnologije / Materials and technology 50 (2016) 6, 861–868 (AFM). The samples were polished with Struers polish- ing devices with suitable abrasives. The specimens thinner than 250 μm etched with 0.5 % HF solutions in 30 s and then cleaned with alcohol and dried with air. 8-4-2 mm specimens were etched electroliticaly under the 24 °C, 18 V electric current in a 5 % Tetra Fluoro- boric acid solution in 120 s. And Nikon stereo light microscope is used for determining the microstructure with X10-500 magnification. The grain dimensions are derived with ASTM E112. The tensile tests were realized with Testometrik DBBMTCL-250 Kgf device. And the pinhole counting is realized with special lighting table with BS-EN 546-4 procedures. 3 RESULTS AND DISCUSSION 3.1 Mechanical properties A tensile test procedure is applied after all foil cold rolling stages and annealing procedures. The ultimate engineering stress before rupture of specimen, Engineer- ing Tensile Stress ( u), Yield Stress ( 0.2) and Strain at Failure (f) values were determined with 100 mm initial length samples (A100). Nearly same tendency and gra- phics were observed in all tensile tests for cold rolled specimens. A specific stress-strain graphics for AA8079 were given in Figure 1 as representative. Aluminum on the other hand having a FCC crystal structure does not show the definite yield point in comparison to those of the BCC structure materials, but shows a smooth engi- neering stress-strain curve. The yield strength ( 0.2) therefore has to be calculated from the load at 0.2 % strain. It can be observed from Figure 1 that the material does not show any significant strain hardening after yield point. It causes the develop plastic instability and, there- by, very low ductilities. When this situation compared with the previous researchs the same tendency is ob- served. This undesirable phenomenon and the associated strain localization can be avoided by employing anneal- ing process.17 The von-misses total effective strain (vm) was calculated for determining the incremental work for all cold rolling cycle. Figure 2a shows the vm versus (Ten- sile Stress) u, Figure 2b shows the vm versus (Yield Stress) 0.2. As can be seen in Figures 2a and 2b an increase can be observed in after first cold Rolling (vm = 1.5). It can be explained by strain hardening of cold worked alloy. But nearly no change was observed on tensile and the yield stress between vm= 1.5 and 4.2. It can be explained by very low and saturated strain hardening and very long post uniform elongations as depicted in stress-strain curve in Figure 1. Notice that the AA8079 material includes 99 % pure aluminum. When the tensile and yield strength graphs are compared with the previous researcher result the same tendency can be observed. K. S. S. Satheesh and T. Raghu8 reported that the yield strength ( 0.2) increases signifi- cantly after first pass of cold working, whereas the tensile strength ( u) is nearly showed same behavior. Considerable increase in strength observed after first pass is mainly attributed to the significant decrease in A. CAN et al.: ANALYSIS OF TWIN-ROLL CASTING AA8079 ALLOY 6.35-μm FOIL ROLLING PROCESS Materiali in tehnologije / Materials and technology 50 (2016) 6, 861–868 863 Figure 2: The effect of equal strain on mechanical properties: a) tensile strength, b) yield strength, c) elongation at break Slika 2: Vpliv enake napetosti na mehanske lastnosti: a) natezna trdnost, b) meja plasti~nosti, c) raztezek ob poru{itvi Figure 1: The stress-strain graphic for 256 μm specimen Slika 1: Diagram napetost-raztezek za vzorec debeline 256 μm grain sizes and the increased dislocation density which necessitates higher applied stress for dislocation motion by slip. A horizontal trend in 0.2 & u is observed in subsequent passes which is attributed to the increased recovery/annihilation of dislocations with increasing accumulated strain and formation of micro cracks. Propagation of intermetallic phases in structure has great affect on mechanical properties. Previous researchers indicate that the intermetallic phases reduce the strength and elongation capability of the aluminum alloys. When tensile fracture surfaces of aluminum alloy samples investigated with SEM, it is observed that fracture occurs secondary intermetallic phases and inclusion concentration regions.18 P. J. Appsa et al.15 investigated the effect of coarse second-phase particles on the rate of grain refinement and material properties during severe deformation pro- cessing. Authors indicate that the hardness development of the AA8079, with increasing deformation process strain and the hardness of the AA8079 was slightly higher than that of the single-phase alloy, due to the pre- sence of the second-phase particles. During deformation, the work hardening of the AA8079 alloy saturated much more rapidly than the single-phase alloy and reached a plateau at a strain of only vm~3 showing little further increase even after a strain of 10. This behavior would be expected to correspond to a continued fast micro struc- tural refinement with increasing strain. The restoration processes can change microstructures as well as mechanical and physical properties of metals and alloys while required mechanical and physical pro- perties may be achieved by adjusting the deformation and annealing variables.2 The annealing process was decreased the tensile and the yield strength significantly. Reversely the annealing process was increased the elon- gation at break values. The maximum elongation at break f value is observed as 4.4 from 250 μm to 15 μm foils. This value was decreased to 1.2 % for final 6.5 μm foils. Then the last foil annealing treatment increased the elongation value to 2.3 %. It can be concluded that the recovery phenomenon is the major reason of decrease in flow stress. The steep increase in ductility implies that a complete softening due to recrystallization and grain refinement has taken place, and there covered structure, which is expected to be the main cause of the observed premature failure, has been removed.8,17 Figure 3 shows the micro-hardness results of the rolling stages from TRC to last foil annealing. The first cold rolling process after TRC is decreased to micro- hardness from 38 HV to 47 HV. During the deformation process, after the 1st annealing process the hardness of the alloy saturated rapidly and reached a plateau at a 250 μm thickness. This behavior would be expected to correspond to a continued fast micro structural refine- ment with increasing strain. The annealing at 450 °C and a holding time of four hours decreased the micro-hard- ness from 47 HV to 29 HV. After annealing the cold rolling was decreased to 29 HV to 42 HV and the micro hardness has no significant change during the cold roll- ing process until the last annealing process. The micro hardness was decreased to the least value after the last foil annealing process. As compared the micro hardness behavior with previous works; the results are coinciding with each other. Salehi reported the variation of micro- hardness after cold rolling with 20 %, 30 % and 40 % reduction in thickness. As a result, increased cold work- ing increases the micro-hardness of the structure by increased dislocation density and deformed grains.2,8,15 There is significant increase in hardness after first pass, after 1st annealing followed by marginal increase after second pass. During subsequent passes the hardness drops slightly and tends to remain fairly uniform fairly coinciding with earlier findings.8 After final cold rolling 14 μm to 6.5 μm, (vm =4.2) the annealing in a furnace at 270 °C and a holding time of 11 h is applied. This treatment reduced the tensile and the yield strength of the materials 71 MPa and 55 MPa respectively. The cold rolled and annealed 6.5 μm foils stress-strain graphics were depicted in Figures 4a and 4b.The cold rolled 6.5 μm foil has showed maximum elongation  = 1.2. This value was increased to 2.3 % after the last annealing process. These elongations at break values show that this material can be called a brittle material because of low ductility. This undesirable phenomenon and the associated strain localization can be avoided by employing suitable annealing procedures. The annealed specimen’s tensile graphics has some diffe- rences from cold rolled specimens. The annealed speci- mens showed a dynamic deformation aging behavior or The Portevin–Le Chatelier effect (PLC) effect. PLC describes a serrated stress-strain curve or jerky flow, which some materials exhibit as they undergo plastic deformation. This behavior is an expected behavior on annealed aluminum alloys only in limited regimes of strain rate. In a uniaxial tension test for instance, this A. CAN et al.: ANALYSIS OF TWIN-ROLL CASTING AA8079 ALLOY 6.35-μm FOIL ROLLING PROCESS 864 Materiali in tehnologije / Materials and technology 50 (2016) 6, 861–868 Figure 3: The variation of Micro hardness of the foils during cold rolling Slika 3: Spreminjanje mikrotrdote folij med hladnim valjanjem irregular flow results in inhomogeneous deformation with various localization bands. These bands can be static, hopping and sometimes propagating along the specimen. It is also observed in presence of this irregular flow that some materials fail by a shear localization mode prior to any diffuse necking in uniaxial tension and even under more complex states of stress.19 3.2 Microstructure properties The samples were characterized by XRD with a Bruker D8 advance diffractometer (40 kV, 40 mA), in Bragg-Brentano reflection geometry with Cu-K radi- ation ( = 0.154 nm). The data were obtained between 10° and 90° in steps of 0.1 with counting time of 3. When the previous research are investigated various author focused the intermetallic phase formation. On the other hand some of the researchers focused the TRC process which is very important the on the effect on cooling rate. Al-rich portion of the Al–Fe binary phase diagram shows many intermetallic formed by the peri- tectic reactions.13 The Al13Fe4, Al5Fe2, Al3Fe (AlnFem) etc. are possible intermetallic phases in Aluminum alloys.1,10 Figure 5 shows the XRD diffractogram pattern of the 6.35 μm foil of AA8079 alloy. The peak in 2 78° indicates the -Al (311) with miller index. The remain- ing peaks indicate Al3Fe intermetallic phases. Alloying of Al with Fe increases the high temperature strength due to a dispersion of second phase particles. Some of the researchers15 introduces this Al3Fe intermetallic peak as Al13Fe4. On the other hand some of the researcher reported that, these two monoclinic structural submicron intermetallics are very close to each other.12,20 The rapid cooling and solidification have great effect on formation of the intermetallic phase. However, the development of solidification microstructures along the (Twin Roll Casting) TRC process has some particular characteristics. The TRC makes the solidification with water cooled cylinders and it causes high cooling rate solidification and deforming near the surface.1 High cooling rates near the surface cause the formation of metastable intermetallic Al6Fe and AlmFe compounds in addition to the stable Al3Fe. It is often considered that at high cooling rates, due to kinetic restrictions there is not enough time for the atoms to arrange themselves in a stable structure.1,13 During the cooling stage no nuclea- tion is involved and an epitaxial solidification occurs. Moreover, the onset of solidification at the molten sub- strate interface is characterized by a solidification velo- city that approaches zero, favoring the initial formation of the stable Al–Al3Fe eutectic. This structure probably continues to grow in spite of the sudden increase in the solidification velocity over the surface, i.e., the equili- brium Fe aluminide is not displaced with increasing solidification velocity by a metastable aluminide.10,14,15 Previous researchers reported that the Si content is lower than 0.15 % of mass fractions of Si, which is the limiting Si content permitting to avoid AlFeSi to be the dominant intermetallic phase.16 The material used in the experiment (AA8079) has the Si ratio of 0.12 and no AlFeSi intermetallic was observed in the structure. A microstructure samples were derived after all rolling process. The procedure for etching and electro polishing for thick specimens was described in the pre- vious sections. Figures 6a and 6b shows the etched and electro polished optical microscope images respectively. As depicted in Figures 6a and 6b the intermetallic phases needles oriented along the nonhomogeneous grains. Orientation along the casting directions between the grains boundaries were not observed on 8 mm TRC samples. The 8 to 4 mm cold rolled specimen’s micro- structure was depicted in Figures 6c and 6d. The effect A. CAN et al.: ANALYSIS OF TWIN-ROLL CASTING AA8079 ALLOY 6.35-μm FOIL ROLLING PROCESS Materiali in tehnologije / Materials and technology 50 (2016) 6, 861–868 865 Figure 5: X-ray diffractogram of the Twin Roll Casting 6.35 μm AA8079 Foil Slika 5: Rentgenogram 6,35 μm folije iz traku AA8079, ulitega med dvema valjema Figure 4: The stress-strain graphic for 6.5 μm specimen: a) cold rolled, b) annealed-270 °C, 11 h Slika 4: Diagrami napetost-raztezek pri vzorcu debeline 6,5 μm: a) hladno valjano, b) `arjeno 270 oC, 11 h of the 50 % plastic deformation on cold rolling can be observed on these figures. The narrowing affect along the grain boundary can be observed as comparing the Figures 6a and 6c. After 4 mm cold rolling process the annealing (580 °CC – 8 h) is applied before the foil roll- ing process. Figures 6e and 6f show the annealed sam- ples microstructure. The needle shape of the interme- tallic particles transforms to bulk shape by the recovery effect. The microstructures of the 250 μm to 6.5 μm foil specimens’ light microscope images were determined. The 250 μm foil microstructure is depicted in Figure 6g as representative. As depicted in Figure 6g the grains are elongated along the rolling direction and the narrowed vertically to rolling direction by comparing the previous stage microstructure as given in Figures 6a and 6f. The distance between intermetallic particles are decreased with increased plastic deformation ratio. The grain size vertically to rolling direction is measured with image processing and illustrated with grain size vs rolling stages from TRC to 6.5 μm foil in Figure 7. As depicted in Figure 7 the grain size was decreased from 150 μm to 0.5 μm. After vm = 2.4 strain, the submicron grain size were observed in the microstructure. 3.3 Surface properties Not only the metallurgical and mechanical properties of the aluminum sheets and foils are very important, but also the surface properties of the aluminum sheets and foils are very important and it must be characterized to A. CAN et al.: ANALYSIS OF TWIN-ROLL CASTING AA8079 ALLOY 6.35-μm FOIL ROLLING PROCESS 866 Materiali in tehnologije / Materials and technology 50 (2016) 6, 861–868 Figure 7: The variation of grain size with TRC to foil rolling stages Slika 7: Spreminjanje velikosti zrn od TRC do kon~ne izvaljane folije Figure 6: Etched and electro polished light microscope images: a), b) twin roll casting, c), d) cold rolled, e), f) 580 oC, 8 h annealed and g) 250 μm foil rolling samples Slika 6: Mikrostruktura po jedkanju in elektropoliranju: a), b) ulito med dvema valjema, c), d) hladno valjano, e), f) `arjeno 8 ur na 580 oC in g) vzorec valjane 250 μm folije Figure 8: The atomic force microscope visualization of the thin foils from 250 to 6.5 μm Slika 8: Vizualizacija povr{ine folij od 250 μm do 6,5 μm, z mikro- skopom na atomsko silo desired functions as where they used. The chemical and electrochemical surface properties and the surface struc- ture are very important in lithography sheets and also in food industry. In contrary to hydrophilic lithography sheets the surface must be smooth and shiny in food in- dustry. So in this work the surface properties were deter- mined with atomic force microscope (AFM), surface roughness measurement and optical microscope. And also the pin holes and rolling tracks were determined with the SEM and optically with light source. The AFM scanning images were depicted in Figure 8 from 250 μm to 6.5 μm. Notice that the vertical scale is used as μm and nm depending on the roughness of the samples. The valley and the peaks are oriented along the rolling direc- tion. Figure 9 shows the surface roughness of the rolled samples and the roller pins surface roughness. As de- picted in Figure 9 the surface roughness is decreased with rolling stages. The surface roughness of the rolling pin and the foils are very close in almost all cases espe- cially in thinner foils. The pin hole counting is realized with a special light source in an dark ambient. The holes are counted for per 1 dm2 as detailed in ISO-EN 546-4. Figure 10a shows a sample pin hole SEM images. The pin hole diameter is measured as 1 μm and meanly 10 pin holes are counted in 1 dm2 standard area. And also the rolling tracks, po- rous and the skid effected banked up structure can be observed from Figures 10a and 10b. 4 CONCLUSION Foils production were tested on the level of industrial trials with no rupture by the time rolling process. It has been shown that selected TRC parameters result in the production of 8 mm sheet of good quality, with espe- cially: fine microstructure with adequate grain refiner addition and annealing conditions. Although the hard phases (that could lead to porosity problems) in the microstructure the 8 mm to 6.35 μm foil rolling was realized by obtaining the adequate surface properties such as sufficient pin hole, porosity The von-misses total effective strain was calculated for determining the incremental work for all cold rolling cycle. After first cold rolling (vm = 1,5) the yield and the tensile stress were increased in a limited range. This in- crease was explained by strain hardening of cold worked alloy. After continued deformation no change was ob- served on tensile and the yield stress between vm = 1.5 and 4.2. It can be explained by nearly saturated strain hardening behavior after a critical plastic deformation. On the other hand the maximum elongation at break f value is observed as 4.4 from 250 μm to 15 μm foils. These elongations at break values show that this material can be called a brittle material because of low ductility. The first cold rolling 8 mm to 4 mm process after TRC is decreased to micro hardness to the highest value of 47 HV. During the deformation process, after the 1st annealing process the hardness of the alloy saturated rapidly and reached a plateau at a 250 μm thickness. This behavior would be expected to correspond to a continued fast microstructure refinement with increasing strain. The annealed foil specimens showed a dynamic deformation aging behavior or The Portevin–Le Chate- lier effect (PLC) effect in tensile test. PLC describes a serrated stress-strain curve or jerky flow, which some materials exhibit as they undergo plastic deformation. The XRD analyses shows that the TRC casting AA8079 alloys includes -Al (311) and the Al3Fe meta- A. CAN et al.: ANALYSIS OF TWIN-ROLL CASTING AA8079 ALLOY 6.35-μm FOIL ROLLING PROCESS Materiali in tehnologije / Materials and technology 50 (2016) 6, 861–868 867 Figure 10: The surface structures: a) SEM image of a pin hole 5000 ×, b) rolling trucks and porous of the surface 200 × Slika 10: Struktura povr{ine: a) SEM posnetek luknjice 5000 ×, b) sledi valjanja in poroznost povr{ine 200 × Figure 9: The variation of the surface roughness in foil rolling stages Slika 9: Spreminjanje hrapavosti povr{ine pri valjanju folij stable inter-metallic phases because of high cooling rates, due to kinetic restrictions there is not enough time for the atoms to arrange themselves in a stable structure. The AFM scanning images indicates that the valley and the peaks are oriented along the rolling direction. The surface roughness of the rolling pin and the foils are very close in almost all cases especially in thinner foils. 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Kamio, Fabrication of Al/Al3Fe composites by plasma synthesis method, Materials Science and Engineering A, 343 (2003), 199–209, doi:10.1016/S092- 5093(02) 00380-5 A. CAN et al.: ANALYSIS OF TWIN-ROLL CASTING AA8079 ALLOY 6.35-μm FOIL ROLLING PROCESS 868 Materiali in tehnologije / Materials and technology 50 (2016) 6, 861–868 G. JANDIKOVA et al.: ANTIMICROBIAL MODIFICATION OF POLYPROPYLENE WITH SILVER NANOPARTICLES ... 869–871 ANTIMICROBIAL MODIFICATION OF POLYPROPYLENE WITH SILVER NANOPARTICLES IMMOBILIZED ON ZINC STEARATE PROTIMIKROBNO SPREMINJANJE POLIPROPILENA Z NANODELCI SREBRA, IMOBILIZIRANIH NA CINKOVEM STEARATU Gabriela Jandikova, Pavlina Holcapkova, Martina Hrabalikova, Michal Machovsky, Vladimir Sedlarik Tomas Bata University in Zlin, University Institute, Centre of Polymer Systems, tr. Tomase Bati 5678, 76001 Zlín, Czech Republic sedlarik@cps.utb.cz Prejem rokopisa – received: 2015-06-30; sprejem za objavo – accepted for publication: 2015-12-01 doi:10.17222/mit.2015.152 The microwave synthesis of Ag nanoparticles on zinc stearate (ZnSt/Ag) was performed to obtain an antimicrobial additive for a polypropylene matrix. Thermoplastically prepared polymer composites contained (1, 3, 5 and 10) % of mass fractions of ZnSt/Ag. The effect of the presence of additives on the morphology and mechanical properties of composites was studied by scanning electron microscopy and stress-strain analysis. The antimicrobial activity of the composites was studied according to the ISO 22196 standard. The results showed that sufficient antimicrobial activity of the composites against both Gram-positive and Gram-negative bacterial strains was observed in the case of the composites with the highest filling studied. Keywords: antibacterial, polypropylene composite, zinc stearate, Ag nanoparticles Izvedena je bila sinteza nanodelcev Ag na cinkovem stearatu (ZnSt/Ag) z mikrovalovi, da bi dobili protimikrobni dodatek poli- propilenski osnovi. Termoplasti~no pripravljeni kompozit polimera je vseboval (1, 3, 5 in 10) % masnega dele`a ZnSt/Ag. Vpliv prisotnosti dodatkov na morfologijo in mehanske lastnosti kompozitov je bil prou~evan z vrsti~no elektronsko mikroskopijo in analizo napetost-raztezek. Protimikrobna aktivnost kompozita je bila prou~evana skladno s standardom ISO 22196. Rezultati so pokazali, da je primerna protimikrobna aktivnost pri sevih, gram pozitivnih in gram negativnih bakterij, dose`ena v primeru kompozita z najve~jim prou~evanim dodatkom. Klju~ne besede: protibakterijsko, polipropilenski kompozit, cinkov stearat, nanodelci Ag 1 INTRODUCTION Antimicrobial modifications of polymers are used to prevent or inhibit the growth of microorganisms on its surface. Such a modification may find utilization in food packaging, medical applications and especially in hygi- enic materials or textile production. Nowadays, a commonly used method for the modification of polymers is an addition of an antimicrobial agent/additive directly into the polymer matrix. Currently, silver-based (Ag) additives have received significant attention due to the low toxicity of the active Ag ion to human cells as well as for being a long-lasting biocide with high thermal sta- bility and low volatility.1,2 Microwave (MW) synthesis is one of the well-known effective methods for the prepa- ration of Ag NPs.3,4 The immobilization of Ag NPs by MW synthesis on various organic substrates has been studied by P. Bazant et al.2 The authors successfully immobilized nano-silver, nanostructured ZnO and hybrid nanostructured Ag/ZnO on a wood flour (WF) surface by MW synthesis. Subsequently, the modified WF was compounded into a PVC matrix (5 % of mass fractions loading) and the antimicrobial activity was tested while the most efficient system was the hybrid nanostructured Ag/ZnO. N. Iqbal et al.1 described the surface modifica- tion by the MW synthesis of Ag NPs on the surface of inorganic substances. The Ag NPs were successfully bonded on hydroxyapatite and caused antimicrobial activity of the prepared system. The antimicrobial modification of polypropylene with Ag NPs prepared by MW synthesis and immobi- lized on a zinc stearate surface was studied in this work. The prepared composites were characterized by scanning electron microscopy, stress-strain analysis and antimicro- bial testing according to the ISO 22196 standard. 2 EXPERIMENTAL PART 2.1 Materials Polypropylene (PP) resin (C706-21 NA HP, density = 0.9 g.cm–3, MFR = 21 g.10 min–1) used in this work was a product of the Braskem company (Brazil). Zinc steara- te (ZnSt) was supplied by Sigma Aldrich (USA). Silver nitrate (AgNO3), Hexamethylenetetramine (HMTA) and ethanol were purchased from PENTA, Czech Republic. 2.2 Preparation of hybrid ZnSt/Ag particles and PP composites Hybrid ZnSt/Ag particles were prepared under reflux in the MW open vessel system MWG1K-10 (RADAN, Materiali in tehnologije / Materials and technology 50 (2016) 6, 869–871 869 UDK 620.1:678.742.3:669.22:669.5:661.8'075.3 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)869(2016) Czech Republic; 1.5 kW, 2.45 GHz) operating in conti- nuous mode (zero idle time) with an external cooler. First, 200 mL of AgNO3 (0.85 g) solution in water and 450 mL of ZnSt (11.02 g) dispersion in ethanol were transferred into a 1000 mL reaction bottle. The reaction mixture was heated in a MW oven for 2 min. After that, 100 mL of HMTA (7.00 g) solution in water was added and the MW heating continued for 10 min under conti- nuous stirring (250 min–1). The reaction product was collected by suction microfiltration and left to dry in a laboratory oven (50 °C) up to the constant weight. The prepared ZnSt/Ag system contained 2.4 % of mass frac- tions of Ag (determined by atomic absorption spectro- scopy, Agilent DUO 240FS/240Z/UltrAA). The prepared hybrid particles were incorporated into the PP matrix by twin screw micro-compounder DSM Xplore (15 mL chamber volume). The temperature of all three zones was 200 °C, speed 100 min–1 and time of mixing 10 min. The concentration of the filler was from 1 % to 10 % of mass fractions. 2.3 Characterization methods The structure of the prepared samples was observed by scanning electron microscopy (VEGA IILMU, TESCAN). The specimens were coated with a thin Au/Pd layer. The microscope was operated in vacuum mode at an acceleration voltage of 10 kV. The mechanical characterization of the samples (dog-bone shaped specimens) was performed with a M350-5 CT Materials Testing Machine. The cross-head speed was 10 mm/min. Each test consisted of seven re- plicate measurements. Antimicrobial testing was per- formed according to the ISO 22196:2007 international standard against Escherichia coli and Staphylococcus aureus. 3 RESULTS AND DISCUSSION SEM micrographs of the ZnSt/Ag additive and PP 10 % ZnSt/Ag composite are shown in Figure 1. Immo- bilized Ag particles with a size below 100 nm are visible in the case of the pure additive (Figure 1a) as well in the composite (Figure 1b). The size and shape of the MW prepared Ag NPs correspond to the results published by Bazant et al.2 where Ag NPs were immobilized on a wood flour. Furthermore, the SEM analysis reveals that immobilized Ag NPs have good cohesion to the ZnSt substrate, even after thermoplastic processing. The mechanical properties of the composites were noticeably influenced when the loading of the filler was above 5 % of mass fractions. However, the addition of 1 % of mass fractions of ZnSt/Ag improved the tensile strain by approximately 16 % in comparison with neat PP, while the tensile strength remained unchanged. In the case of the composites containing 10 % of mass fractions of ZnSt/Ag the Young’s modulus, tensile strength and tensile strain were reduced by approximately 35 %, 20 %, and 81 %, respectively (Table 1). The mechanical properties of the composites are most dependent on the G. JANDIKOVA et al.: ANTIMICROBIAL MODIFICATION OF POLYPROPYLENE WITH SILVER NANOPARTICLES ... 870 Materiali in tehnologije / Materials and technology 50 (2016) 6, 869–871 Table 1: Summary of stress-strain analysis results of prepared PP/ZnSt/Ag composites Tabela 1: Zbir rezultatov analiz napetost-raztezek pripravljenih kompozitov PP/ZnSt/Ag Young’s modulus (MPa) Tensile strength (MPa) Strain at break (%) PP 217 (± 27) 22 (± 1.7) 187 (± 47) PP 1 % ZnSt/Ag 179 (± 25) 22 (± 0.5) 217 (± 23) PP 3 % ZnSt/Ag 172 (± 12) 22 (± 1.1) 68 (± 12) PP 5 % ZnSt/Ag 155 (± 29) 21 (± 1.2) 54 (± 17) PP 10 % ZnSt/Ag 143 (± 10) 18 (± 0.8) 35 (± 3) Figure 1: Scanning electron micrographs of: a) ZnSt/Ag particles and b) PP 10 % of mass fractions of ZnSt/Ag composite. The Ag nanoparticles can be recognized as white dots. Slika 1: Posnetka z vrsti~nim elektronskim mikroskopom: a) ZnSt/Ag delci in b) kompozit PP z 10 % masnega dele`a ZnSt/Ag. Nanodelci Ag se ka`ejo kot bele to~ke. ZnSt content without the effect of the Ag NPs presence on it. Changes of the mechanical properties correspond to the results observed in the case of the composites based on PP and unmodified ZnSt.5 The antimicrobial activity of the composites is summarized in Table 2. There is a slight inhibition of growth of both bacterial strains when the concentration is 5 % of mass fractions of ZnSt (corresponding to 0.08 % of mass fractions of Ag). The samples with the highest filling studied (0.16 % of mass fractions of Ag) exhibit a noteworthy activity against Escherichia coli and Staphylococcus aureus. Antimicrobial activity is in agreement with studies describing incorporation of Ag NPs into hydrophobic polymer matrices.2 4 CONCLUSIONS The hybrid systems of ZnSt and Ag nanoparticles were prepared by MW synthesis and incorporated into a PP matrix. The prepared composites showed a promising antimicrobial activity at concentration above 5 % of mass fractions. The mechanical properties were notice- ably influenced at 10 % of mass fractions of the ZnSt/Ag loading into the PP matrix. However, a noticeable impro- vement of tensile strength of the composites was ob- served already at 1 % of mass fractions of ZnSt/Ag, while the tensile strain remained unchanged in com- parison with the unmodified PP. The proposed antimicrobial modification of com- monly used additives, such as zinc stearate, represents a promising way of nanoparticle handling and applica- tions. Acknowledgments This work was supported by the grant of Technology Agency of the Czech Republic (grant No. TE02000006), the Ministry of Education, Youth and Sports of the Czech Republic – Program NPU I (grant No. LO1504) and by Internal Grant Agency of the Tomas Bata University in Zlín (grant No. IGA/CPS/2015/004). 5 REFERENCES 1 N. Iqbal, M. R. Abdul Kadir, N. A. N. Nik Malek, N. Humaimi Mahmood, M. Raman Murali, T. Kamarul, Rapid microwave assisted synthesis and characterization of nanosized silver-doped hydroxy- apatite with antibacterial properties, Material Letters, 89 (2012), 118–122, doi:10.1016/j.matlet.2012.08.057 2 P. Bazant, L. Munster, M. Machovsky, J. Sedlak, M. Pastorek, Z. Kozakova, I. Kuritka, Wood flour modified by hierarchical Ag/ZnO as potential filler for wood–plastic composites with enhanced surface antibacterial performance, Industrial Crops and Products, 62 (2014), 179–187, doi:10.1016/j.indcrop.2014.08.028 3 D. Breitwieser, M. M. Moghaddam, S. Spirk, M. Baghbanzadeh, T. Pivec, H. Fasl, V. Ribitsch, C. O. Kappe, In situ preparation of silver nanocomposites on cellulosic fibers--microwave vs. conventional heating, Carbohydrate Polymers, 94 (2013), 677–686, doi:10.1016/ j.carbpol.2013.01.077 4 X. Zhao, Y. Xia, Q. Li, X. Ma, F. Quan, C. Geng, Z. Han, Biopoly- mers regulate silver nanoparticle under microwave irradiation for effective antibacterial and antibiofilm activities, Colloids and Sur- faces A: Physicochemical and Engineering Aspects, 444 (2014), 180–188, doi:10.1016/j.colsurfa.2013.12.008 5 Ñ. V. Panin, L. A. Kornienko, T. Nguyen Suan, L. R. Ivanova, M. A. Poltaranin, The effect of adding calcium stearate on wear-resistance of ultra-high molecular weight polyethylene, Procedia Engineering, 113 (2015), 490–498, doi:10.1016/ j.proeng.2015.07.341 G. JANDIKOVA et al.: ANTIMICROBIAL MODIFICATION OF POLYPROPYLENE WITH SILVER NANOPARTICLES ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 869–871 871 Table 2: Antimicrobial activity of prepared composites determined according to ISO 22196 Tabela 2: Protimikrobna aktivnost pripravljenih kompozitov, dolo~ena po ISO 22196 Sample S. aureus E. coli CFU/cm2 R CFU/cm2 R PP 2,4E+05 – 7,2E+05 – PP + 5 % ZnSt 1,1E+05 0.4 1,9E+05 0.6 PP + 10 % ZnSt 9,6E+04 0.4 5,0E+04 1.2 PP + 5 % ZnSt/Ag 1,6E+05 0.2 5,5E+04 1.1 PP + 10 % ZnSt/Ag 9,5E+02 2.4 0,0E+00 7.0 S. S. ISLAM et al.: A NEW WIDEBAND NEGATIVE-REFRACTIVE-INDEX METAMATERIAL 873–877 A NEW WIDEBAND NEGATIVE-REFRACTIVE-INDEX METAMATERIAL NOVI [IROKOPASOVNI METAMATERIAL Z NEGATIVNIM LOMNIM KOLI^NIKOM Sikder Sunbeam Islam1, Mohammad Rashed Iqbal Faruque1, Mohammad Jakir Hossain1, Mohammad Tariqul Islam2 1Space Science Centre (ANGKASA) 2Universiti Kebangsaan Malaysia, Faculty of Engineering and Built Environment, Department of Electrical, Electronic and Systems Engineering, 43600 UKM, Bangi, Selangor, Malaysia sikder_islam@yahoo.co.uk Prejem rokopisa – received: 2015-06-30; sprejem za objavo – accepted for publication: 2015-12-15 doi:10.17222/mit.2015.144 This paper reveals the design and analysis of a new wideband negative-refractive-index (NRI) metamaterial unit cell. The proposed metamaterial unit-cell exhibits resonance in the C-band and displays negative permittivity and permeability there with a wideband NRI property. It also shows a wider negative peak of the refractive index in the major area of the C- and X-band and minor area of the S- and Ku-band, and a maximum 3-GHz negative bandwidth was achieved compared to the reference metamaterial. In the basic design, a square-shaped copper resonator was constructed with a metal strip on the FR-4 substrate material. The measured result was presented and it shows good conformity with the simulated result. Moreover, an analysis was performed with the same design by replacing the substrate material with the popular Rogers RT 6010 instead of the FR-4 material and then it shows NRI properties in the C-, X- and Ku-band. Keywords: metamaterials, negative refractive index, wideband ^lanek obravnava zgradbo in analizo osnovne celice novega, {irokopasovnega metamateriala z negativnim lomnim koli~nikom. Predlagana osnovna celica iz metamateriala ka`e resonanco v C-pasu in ka`e negativno permitivnost ter permeabilnost z lastnostmi NRI {irokega pasu. Ka`e tudi {ir{i negativni vrh lomnega koli~nika v ve~ini podro~ja C-pasu in X-pasu in v manj{em podro~ju S-pasu in Ku-pasu. Najve~ja {irina negativnega pasu (3 GHz) je bila dose`ena v primerjavi z referen~nim metamaterialom. Osnovna zgradba je bila konstruirana kot bakren rezonator {tirikotne oblike, s kovinskim trakom na podlagi iz FR-4 materiala. Predstavljeni so izmerjeni rezultati, ki ka`ejo dobro ujemanje z rezultati simulacije. Poleg tega je bila izvedena analiza z enako zgradbo in z nadome{~anjem podlage s popularnim Rogers RT 6010 namesto FR-4 materialom in prikazane so NRI lastnosti v C-, X- in Ku-pasu. Klju~ne besede: metamateriali, negativni lomni koli~nik, {irokopasovnost 1 INTRODUCTION Metamaterials are engineered (at the atomic level) materials that have unique and extraordinary properties not found in nature. A metamaterial as a composite material, usually gains these properties due to the arrangement of its constituents (in a unit cell) rather than individual properties. There are some exotic properties that are not possible with naturally available materials, but can be achieved with metamaterials, like negative permittivity (<0) or negative permeability (μ<0), nega- tive refractive index, inverted Snell’s law, etc. In 1967 the Russian physicist Victor Veselago1 predicted that it is possible to develop a material of such reverse characte- ristics that it will behave opposite to the natural material. It was also stated by him that a material could exhibit a negative refractive index if it gains negative permittivity and permeability. Around 30 years later in 2000 D. R. Smith et al.2 successfully demonstrated a composite material with such negative properties. Due to these uncommon characteristics it can used in many important applications, like antenna design, EM absorption reduc- tion, electromagnetic cloaking operation, filter deign, sensor design, etc.3–7 A metamaterial with both negative permittivity () and negative permeability (μ) is called a double-negative (DNG) metamaterial or a negative refractive index (NRI) or negative index material (NIM) or a metamaterial with either permittivity or permeability negative is called a single negative (SNG) metamaterial. However, a metamaterial with the DNG property can only exhibit the negative refractive index property properly. There are many metamaterials found in the literature, but not enough metamaterials with a double negative property are found. However, very few of them were designed to exhibit DNG characteristics in the C-band of the microwave region. H. Benosman et al.8 presented a double-negative metamaterial, but their metamaterial was applicable for the Ku-band only. O. Turkmen et al.9, showed a metamaterial for X-band operation, but their metamaterial was not double negative. A. Dhouibi et al.10, proposed a metamaterial for C-band applications, but they claimed these property for an epsilon negative (ENG) metamaterial. S.S. Islam et al.11 designed an S-band metamaterial, but their Materiali in tehnologije / Materials and technology 50 (2016) 6, 873–877 873 UDK 67.017:537.87:620.1 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)873(2016) metamaterial was showing ENG properties as well. Moreover, recently in 12, a DNG material was introduced where a maximum 1.05-GHz bandwidth of the negative refractive index region was claimed. In this study, a new double-negative metamaterial is revealed that exhibits negative refractive index property in the major region of C- and X-band of microwave spectra with a wider bandwidth. Commercially available finite-difference time-domain (FDTD) based CST-micro- wave studio software was used to retrieve the S-para- meters for the unit cell. For further investigation, the structure was then designed on a Rogers RT 6010 substrate material instead of FR-4 material and an analysis was performed. 2 DESIGN AND METHODOLOGY Figure 1a shows the geometry of the proposed square-shaped metamaterial unit cell. The proposed metamaterial unit cell consists of a simple square-shaped copper structure with a vertical copper stripe in the mid- dle of the material, all having a thickness of 0.035 mm. The copper strip in the middle was placed in such a way that it maintains an equal gap for the two opposite sides of the metal strip. The outer length and width of the unit cell were denoted by a = b = 10 mm. The width of the square-shaped structure was expressed by w = 1 mm. The tiny gap at the two ends of the metal strip was symbolized as s = 0.5 mm. The length and width (e) of the metal strip were 7 mm and 1 mm. The distance from the central metal strip to the square border was denoted by, c = d = 3.5 mm. The structure was designed on a 20 mm × 20 mm square-shaped FR-4 substrate material having a dielec- tric constant of 4.3 and a loss tangent of 0.025. The thickness of the substrate was 1.6 mm. In this study, commercially available finite-difference time-domain (FDTD) based computer simulation technology (CST) microwave studio software was adopted for the design and calculation of the reflection (S11) and transmission (S21) coefficients of the unit cell. These parameters were used to compute the effective parameters (permittivity, permeability and refractive index) for the unit cell. Figure 1 displays the simulation arrangements. For the simulation, the designed unit cell was placed between two waveguide ports of positive, negative of z-axis, and exited by transverse electromagnetic (TEM) waves. The rest of the axes were defined as perfect electric conduc- tor (PEC) and perfect magnetic conductor (PMC) boun- dary conditions. The frequency-domain solver was used for the whole simulation. The simulation was executed for the frequency range of 1–15 GHz. For the compu- tation of the effective parameters, the Nicolson-Ross- Weir method was utilized to avoid the inverse cosine- branch-index problem.13 However, as part of further investigation, the unit cell was rotated by 90° and the S parameters were estimated for gaining the effective parameters using the same method. In this study, the open-space measurement techno- logy was adopted. The open-space measurement techno- logy was chosen to observe the realistic effect. For measurement purpose, a prototype of 160 mm × 200 mm was fabricated that contains 8×10 unit cell. The fabri- cated prototype is seen in Figure 2a. The fabricated prototype was placed between two horn antennas. The antennas were acting as the transmitting and receiving end and they were connected to an Agilent E8363D vector network analyzer to calculate the S parameters. However, the distance between the prototype and the horn antenna was kept at 35 cm to avoid the near-field effect. As a part of calibration process, measurements with and without the prototype were performed as well. 3 RESULTS AND DISCUSSION Figure 2b displays the simulated magnitude of the transmission coefficient (S21) for the z-axis wave pro- pagation through the unit cell. It is evident from Figure 2b that for the z-axis wave propagation a clear resonance is seen in the range of the C-band at the frequency of 7.48 GHz of the microwave spectra. Figure 2c shows the measured magnitude of the transmission coefficient where it has been compared with the simulated one. It is apparent from Figure 2c that the measured result shows almost good conformity with the simulated result. However, a slight distortion is found in the measured magnitude of S21 than the simulated result that might have occurred due to the noise effect in the open-space measurement process and fabrication error. Figures 3a and 3b show the real magnitude of the effective permittivity and the permeability against the frequency for the z-axis wave propagation through the unit cell. It is clear from Figure 3a that the permittivity curve has two clear resonances at the frequencies of S. S. ISLAM et al.: A NEW WIDEBAND NEGATIVE-REFRACTIVE-INDEX METAMATERIAL 874 Materiali in tehnologije / Materials and technology 50 (2016) 6, 873–877 Figure 1: a) Design of the unit cell, b) simulation geometry in CST software Slika 1: a) Zgradba osnovne celice, b) simulacija geometrije s CST programsko opremo (a) (b) 2.94 GHz and 7.45 GHz. Moreover, the negative portion of this curve is found from 2.91 GHz to 5.57 GHz, which covers almost 2.66 GHz of bandwidth, more than 1 GHz bandwidth from the frequency of 7.42 GHz to 8.71 GHz, and nearly 3 GHz bandwidth from the frequency of 10.75 GHz to 13.65 GHz. Similarly, the permeability curve of Figure 3b displays a negative region from the frequency of 3.74 GHz to 7.51 GHz that covers almost 3.77 GHz bandwidth. Another negative portion is visible there from the frequency of 11.15 GHz to 14.77GHz, which also covers nearly 3.62 GHz bandwidth. For this design, for the varying magnetic field, a charge builds up in the gaps between the metal strip and the ring. At low frequency the current remains in phase with the applied field, but at higher frequency it starts lagging and produces a negative permeability at that frequency. Figure 4 reveals the real magnitude of the refractive index () against frequency for the unit cell. In this paper, it was mentioned earlier that for a material the refractive index curve would be negative if its per- mittivity and permeability curve appears negative simul- taneously. Therefore, from Figure 4 it is apparent that for the proposed material the refractive index curve exhi- bits a negative magnitude from the frequency of 3.74 GHz to 5.57GHz; 7.42 GHz to 7.51 GHz and 11.15 GHz to 13.65 GHz. It is notable that two wide bandwidths of 1.83 GHz (3.74 GHz to 5.57GHz) and 3.73 GHz (11.15 GHz to 13.65 GHz) are seen as the negative region in the refractive index curve. These bandwidths have fallen in the few portion of S- and Ku-band and major portion of C- and X-band of the microwave region. Moreover, in these negative regions of the refractive index curve, the permittivity and permeability curve are also found to be displaying a negative peak. As a result, the material can be characterized as a double-negative (DNG) metamaterial in these regions of microwave spec- tra. Another important feature is in the frequency range between 7.42 GHz and 7.51 GHz where the refractive index curve was found to be negative, both the simulated and measured transmittance (S21) are also found to be exhibiting sharp resonance clearly at the frequency of 7.48 GHz with a refractive index = -4.40. Thus, it S. S. ISLAM et al.: A NEW WIDEBAND NEGATIVE-REFRACTIVE-INDEX METAMATERIAL Materiali in tehnologije / Materials and technology 50 (2016) 6, 873–877 875 Figure 2: a) Prototypes for measurement, b) simulated transmission coefficient (S21) for z-axis wave propagation, c) comparision of measured and simulated result for S21 Slika 2: a) Prototipi za merjenje; b) simuliran koeficient prenosa (S21) pri napredovanju vala po z-osi, c) primerjava izmerjenega in simuli- ranega rezultata za S21 Figure 4: Real magnitude of refractive index () versus frequency Slika 4: Realni obseg lomnega koli~nika () v odvisnosti od frekvence Figure 3: a) Real magnitude of permittivity () against frequency, b) real magnitude of permeability (μ) versus frequency Slika 3: a) Realna magnituda permitivnosti () proti frekvenci, b) realna magnituda permeabilnosti (μ) v odvisnosti od frekvence reveals that for the z-axis wave propagation the material is practically applicable for C-band applications in the the microwave spectra. As a part of a further investigation, the Rogers 6010 substrate material was used instead of the FR-4 substrate material for the unit cell. Figures 5a and 5b depicts the transmission coefficient as well as the real magnitude of permittivity () against frequency for the unit cell on the Rogers 6010 substrate material consecutively. According to Figure 5a, the transmission coefficient shows two resonances at the frequencies of 5.14 GHz and 10.06 GHz. These frequencies are in the range of C-band and X-band. The permittivity curve in Figure 5b reveals a negative magnitude from the frequency of 3.29 GHz to 6.55 GHz, which covers more than 3 GHz band- width in the C-band. Similarly, it also shows negative peak from the frequency of 8.80 GHz to 10.30 GHz in the X-band and 12.84 GH to 15 GHz in the Ku-band. In Figures 6a and 6b the real magnitude of the per- meability and refractive index is depicted. It is apparent from Figure 6a that the permeability curve shows a negative peak from the frequency of 4.33 GHz to 8.90 GHz and 13.25 GHz to 15 GHz. Therefore, from the permittivity and permeability curve of Figures 5b and 6a it is evident that the material exhibits a double-negative property from the frequency of 4.33 GHz to 6.55 GHz, 8.80 GHz to 8.89 GHz and 13.25 GHz to 15 GHz. Usually, the properties of permittivity and per- meability are most likely affected by the polarization due to the internal architecture of the material. When electro- magnetic waves enter anisotropic materials, which have unequal lattice axes, it is affected by the polarization inside the material. As a result, the value of the per- mittivity and permeability changes due to changes in the design. In the same way, the refractive index curve is also affected by the polarization. Similarly, in the refractive index curve in Figure 6b, it is clear that the curve shows negative magnitude from the frequency of 4.33 GHz to 6.55 GHz, 8.80 GHz to 8.89 GHz and 13.32 GHz to 15 GHz and these frequency ranges cover 2.22 GHz, 9 MHz, 2.32 GHz bandwidth in the microwave ranges. These regions of negative refrac- tive index curve fall in the range of C, X and Ku-band of microwave spectra. Moreover, these frequencies obey the permittivity and permeability curves in Figures 5b and 6a as well. 4 CONCLUSIONS In this paper, a new square-shaped negative refractive index metamaterial was demonstrated that exhibits a wider negative peak in the major area of the C-and X-band and the minor area of S- and Ku-band. A more than 3-GHz wider bandwidth of the negative peak was achieved for the proposed metamaterial than the latest reference metamaterial. The measured result also agrees well with the simulated result. Moreover, the material shows a negative refractive index zone in the C-, X and Ku-band of the microwave spectra as well as when it is designed on the Rogers 6010 substrate material. How- ever, C- and X-, Ku-band are widely used for long-dist- ance and satellite communications. So, this metamaterial can be practically applied in these frequency bands, S. S. ISLAM et al.: A NEW WIDEBAND NEGATIVE-REFRACTIVE-INDEX METAMATERIAL 876 Materiali in tehnologije / Materials and technology 50 (2016) 6, 873–877 Figure 6: a) Real magnitude of permeability (μ) against frequency, b) real magnitude of refractive index () versus frequency for the unit cell on the Rogers 6010 substrate material Slika 6: a) Realen obseg magnitude permeabilnosti (μ) v odvisnosti od frekvence, b) realen obseg lomnega koli~nika () v odvisnosti od frek- Figure 5: a) Transmission coefficient (S21) for the unit cell on the Rogers 6010 substrate material, b) real magnitude of permittivity () against frequency for the unit cell on the Rogers 6010 substrate Slika 5: a) koeficient prenosa (S21) osnovne celice na podlagi iz mate- riala Rogers 6010, b) realna magnituda dielektri~ne konstante () v odvisnosti od frekvence osnovne celice na podlagi iz Rogers 6010 especially for wider bandwidth application besides the other metamaterials in the microwave range. 5 REFERENCES 1 V. G. Veselago, The electrodynamics of substances with simulta- neously negative values of  and μ, Soviet Physics Uspekhi, 10 (1968), 509–514, doi:10.1070/PU1968v010n04ABEH003699 2 D. R. Smith, W. J. Padilla, D. C. Vier, S. C. Nemat-Nasser, S. Schultz, Composite medium with simultaneously negative per- meability and permittivity, Physical Review Letters, 84 (2000), 4184–418, doi:10.1103/PhysRevLett.84.4184 3 S. S. Islam, M. R. I. Faruque, M. T. Islam, An Object-Independent ENZ Metamaterial-Based Wideband Electromagnetic Cloak, Scientific Reports, 6 (2016) 33624, 1–10, doi:10.1038/srep33624 4 J. Carver, V. Reignault, Engineering of the metamaterial-based cut-band filter, App. Phys. A, 117 (2014), 513–516, doi:10.1007/s00339-014-8694-7 5 X. Yang, D. Sun, T. Zuo, X. Chen, K. Huang, Analysis and reali- zation of improving the patch antenna gain based on metamaterials, International Journal of Applied Electromagnetics and Mechanbics, 44 (2014) 1, 17–25, doi:10.3233/JAE-131731 6 L.-W. Li, Y.-N. Li, T. S. Yeo, J. R. Mosig, O. J. F. Martin, A broad- band and high-gain metamaterial microstrip antenna, Applied Physics Letters, 96 (2010) 164101, 1–3, doi:10.1063/1.3396984 7 X. Shen, T. J. Cui, J. Zhao, H. F. W. X. Ma, Jiang, H. Li, Polari- zation-independent wide-angle triple-band metamaterial absorber, Optical Express, 19 (2011), 9401–9407, doi:10.1364/OE.19.009401 8 H. Benosman, N. B. Hacene, Design and Simulation of Double “S” Shaped Metamaterial, International Journal of Computer Science Issues, 9 (2012) 2, 534–537, (Available at: http://ijcsi.org/papers/ IJCSI-Vol-9-Issue-2-No-1.pdf) 9 O. Turkmen, E. Ekmekci, G. Turhan-Sayan, Nested U-ring reso- nators: a novel multi-band metamaterial design in microwave region, IET Microwave and Antennas Propagation, 6 (2012) 10, 1102–1108, doi:10.1049/iet-map.2012.0037 10 A. Dhouibi, S. N. Burokur, A. de Lustrac, A. Priou, Study and ana- lysis of an electric Z-shaped meta-atom, Advanced Electromagnetics, 1 (2012) 2, 64–70, doi:10.7716/aem.v1i2.82 11 S. S. Islam, M. R. I. Faruque, M. T. Islam, Design of a New ENG Metamaterial for S-Band Microwave Applications, Journal of Elec- trical and Electronics Engineering, 7 (2014) 2, 13–16, (Available at: http://electroinf.uoradea.ro/index.php/jeee.html) 12 M. I. Hossain, M. R. I. Faruque, M. T. Islam, M. H. Ullah, A New Wide-Band Double-Negative Metamaterial for C- and S-Band Appli- cations, Materials, 8 (2015) 1, 57–71, doi:10.3390/ma8010057 13 S. S. Islam, M. R. I. Faruque, M. T. Islam, A New Direct Retrieval Method of Refractive Index for Metamaterials, Current Science, 109 (2015) 2, 337–342, (Available at: http://www.currentscience.ac.in/ php/toc.php?vol=109&issue=02) S. S. ISLAM et al.: A NEW WIDEBAND NEGATIVE-REFRACTIVE-INDEX METAMATERIAL Materiali in tehnologije / Materials and technology 50 (2016) 6, 873–877 877 D. [TEFKOVÁ et al.: EVALUATION OF THE DEGREE OF DEGRADATION USING THE IMPACT-ECHO METHOD ... 879–884 EVALUATION OF THE DEGREE OF DEGRADATION USING THE IMPACT-ECHO METHOD IN CIVIL ENGINEERING OCENA STOPNJE DEGRADACIJE V GRADBENI[TVU Z UPORABO METODE ODMEVA ZVO^NIH VALOV Daniela [tefková, Kristýna Tim~aková, Libor Topoláø, Petr Cikrle Brno University of Technology, Faculty of Civil Engineering, Veveøí331/95, 602 00 Brno, Czech Republic stefkova.d@fce.vutbr.cz Prejem rokopisa – received: 2015-06-30; sprejem za objavo – accepted for publication: 2015-12-15 doi:10.17222/mit.2015.150 Non-destructive methods such as Impact-echo method are based on the acoustic properties of the material that are dependent on its condition. It allows the studied progress development of micro-defects in the structure of the material and is thus suitable for monitoring the building structure’s condition. This acoustic method allows us to identify and locate defects and is thus suitable for monitoring the building structure’s condition. The application of this method is widespread; it can be used in mechanical engineering, power engineering and in many industries as well as in the construction industry. Impact-echo uses a short-time mechanical impulse (a hammer blow) that is applied to a surface of the test sample and is detected by means of piezoelectric sensors placed on the surface of the sample. The impulse is reflected by the surface but also by micro-cracks and defects of the specimen that are under investigation. From thus obtained signal the frequency spectrum is determined and is found to be the dominant resonance frequency using fast Fourier transformations. The dominant frequencies give an account of the condition of the structure or determine the location of flaws, at which the waves are rebounded. The signal is digitized by means of a data processing system to be transferred into a computer memory. A piezoelectric MIDI sensor takes the signal response and it is brought to the input of an oscilloscope TiePie engineering Handyscope HS3 two-channel with 16-bit resolution. This paper reports the results of measurements by the Impact-echo method on three applications in civil engineering. The results are obtained in the laboratory during the hardening process in quasi-brittle materials such as alkali-activated slag mortars, the degradation of concrete samples by corrosion caused by the action of chlorides and the degradation of composite materials based on cement by high temperature. Keywords: predominant frequency, high temperature, impact, mortar, specimens Neporu{na metoda, kot je npr. odmev zvo~nih valov, temelji na akusti~nih lastnostih materiala, ki so odvisne od stanja materiala. Metoda omogo~a {tudij napredovanja mikronapak v strukturi materiala in je zato primerna za pregled stanja gradbene strukture. Ta akusti~na metoda omogo~a odkrivanje in dolo~anje polo`aja napake in je zato primerna za pregled stanja zgradbe. Metoda ima {iroko uporabnost, saj se lahko uporablja v strojni{tvu, energetiki ter v mnogih drugih industrijah kot tudi v gradbeni{tvu. Metoda odmeva zvo~nih valov uporablja kratke mehanske udarce (udarec s kladivom) po povr{ini preizku{anega vzorca, odmeve pa se dolo~i s piezoelektri~nim senzorjem, ki je tudi name{~en na povr{ini. Impulz odseva povr{ino in tudi mikrorazpoke in napake v vzorcu, ki se ga preiskuje. Iz tako dobljenega signala, se dolo~i spekter frekvenc in s pomo~jo hitre Fourierjeve transformacije se poi{~e prevladujo~a resonan~na frekvenca. Prevladujo~e frekvence poka`ejo stanje zgradbe ali pa dolo~ijo lokalne pomanjkljivosti, na katerih valovi posko~ijo. Signal se digitalizira s sistemom obdelave podatkov, da se ga lahko prenese v spomin ra~unalnika. Piezoelektri~ni sensor MIDI prevzame odbit signal, ki se ga usmeri na vhod dvokanalnega osciloskopa TiePie Engineering Handyscope HS3, z lo~ljivostjo 16 bitov. ^lanek predstavlja rezultate meritev z metodo odmeva zvo~nih valov na treh primerih v gradbeni{tvu. Rezultati so dobljeni v laboratoriju med procesom utrjevanja v kvazi krhkem materialu, kot je z alkalijami aktivirana `lindra v maltah, degradacijo betonskih vzorcev zaradi korozije, povzro~eno s kloridi in visoko temperaturno degradacijo kompozitnega materiala na osnovi betona Klju~ne besede: prevladujo~a frekvenca, visoka temperatura, udarec, malta, vzorci 1 INTRODUCTION In civil engineering, efficient non-destructive quality control plays an important role in the optimization of resources for manufacturing, maintenance and safety. The impact-echo method is a useful non-destructive technique for flaw detection in concrete. It is based on monitoring the surface motion, resulting from a short- time mechanical impulse. This method overcomes many of the barriers associated with flaw detection in concrete that occur during ultrasonic methods. One of the key features of this method is the transformation of the recorded time-domain waveform of the surface motion into the frequency domain. The impact gives rise to modes of vibration and the frequency of these modes is related to the geometry of the tested object and the presence of flaws.1 A short-time mechanical impulse, generated by tapp- ing a hammer against the surface of a concrete structure (Figure 1), produces low-frequency stress waves that propagate into the structure.2,3 Thus generated waves propagate through the specimen structure and reflect from the defects located in the volume of the specimen or in the surface. Surface displacements caused by the reflected waves are recorded by a transducer located adjacent to the impact.4 The signal is digitized by an analogue/digital data system and transmitted to a com- puter memory. This signal describes the transient local vibrations, which are caused by the mechanical wave Materiali in tehnologije / Materials and technology 50 (2016) 6, 879–884 879 UDK 669.9:620.19:67.017 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)879(2016) multiple reflections inside the structure. The dominant frequencies of these vibrations give an account of the condition of the structure that the waves pass through.5,6 The signal analysis from the impact-echo method is most frequently performed by using frequency spectra obtained from the fast Fourier transform. Fourier ana- lysis converts time to frequency and vice versa. A fast Fourier transform (FFT) is an algorithm to compute the discrete Fourier transform (DFT) and it is inverse. The results of fast Fourier transforms are widely used for many applications in engineering, science, and mathe- matics. Equation (1) is the expression of the Fourier transform for a continuous function:7 x t a a nt T b nt Tn n ( ) cos sin= + ⋅ ⎛⎝ ⎜ ⎞ ⎠ ⎟ + ⋅ ⎛⎝ ⎜ ⎞ ⎠ ⎟⎡ ⎣⎢ ⎤ ⎦⎥ 0 2 2 2π π n = ∞ ∑ 1 (1) a T x t nt T tn T = ⋅ ⎛⎝ ⎜ ⎞ ⎠ ⎟∫ 2 2 0 ( ) cos π d b T x t nt T tn T = ⋅ ⎛⎝ ⎜ ⎞ ⎠ ⎟∫ 2 2 0 ( ) sin π d 2π T =  where an and bn can be calculated from function x(t) using the following relations. 2 EXPERIMENTAL PART For the impact-echo method a short-time mechanical impulse (a hammer blow) was applied to the surface of the specimen during the test and was detected by means of a piezoelectric sensor (Figure 2). The impulse reflects from the surface but also from micro-cracks and defects present in the specimen that are under investigation. The frequency analysis can be carried out from the response signal by means of a fast Fourier transform and thus dominant resonant frequencies are found. An MIDI piezoelectric sensor was used to pick up the response and the respective impulses were directed into the input of an oscilloscope TiePie engineering Handyscope HS3 two- channel with a 16-bit resolution. 2.1 Material for the hardening process of alkali-acti- vated slag mortars The mixture consisted of 450 g of fine-grained granulated blast furnace slag [tramberk 380 (specific surface area 380 m2 kg–1), 180 g of sodium silicate (water glass) with a modulus of 1.6, 1350 g of silica sand (maximum grain size of 2.5 mm) and 95 mL of water. The amount of admixtures was 0.1 % of mass fractions with respect to the slag. CNTs were added in the form of a well-dispersed aqueous dispersion containing 1 % of mass fractions of multi-walled carbon nanotubes D. [TEFKOVÁ et al.: EVALUATION OF THE DEGREE OF DEGRADATION USING THE IMPACT-ECHO METHOD ... 880 Materiali in tehnologije / Materials and technology 50 (2016) 6, 879–884 Figure 2: Images of measurement with the Impact-Echo method Slika 2: Posnetka merjenja z metodo odmeva zvo~nih valov Figure 1: Principle of the Impact echo method Slika 1: Princip metode odmeva zvo~nih valov (Graphistrengths CW 2-45). Since CNTs are commonly not water-soluble, the dispersion also contained carboxy- methyl cellulose (68 g/L) as a dispersing agent.8 The slurry was poured into steel moulds 40 mm × 40 mm × 160 mm to set. The samples were demoulded after 24 h and one set was tested (marked 0d) and the second set was immersed in water for another 28 d before testing (marked 28 d). 2.2 Material for the degradation of concrete samples by corrosion caused by the action of chlorides The mixture for the production of concrete samples was composed of cement CEM II/B – S 32.5 and 1350 kg of sand with a fraction of aggregate 0–4 mm and 225 L of water. The high water ratio resulted in easier penetration of the degradation agents into the concrete structure. These samples of dimensions 50 mm × 50 mm × 330 mm were reinforced with one standard reinforcing bar of diameter 10 mm and a length of 400 mm passing through the centre of the beam. The samples were demoulded after 24 h and placed in the water for 27 d, then the dried samples with natural humidity were exposed to accelerated degradation by chlorides. The samples were immersed into a 5 % water solution of NaCl for 16 h and then subsequently placed into a drying oven with an air temperature of 40 °C for 8 h. 2.3 Specimens intended to be subjected in the heating process Mortars of dimensions 40 mm × 40 mm × 160 mm were produced using a type CEM I 42.5 R Portland cement (^eskomoravský Cement-Heidelberg Cement Group) and a water-to-cement ratio (w/c = 0.46) and quartz sand from Filtra~ní písky, s.r.o. for the preparation of the test mortar mixture in a ratio of 1 to 3, in com- pliance with the ^SN 721200 standard. The specimens were left in the moulds for 24 h, then cured in water for 27 d and finally air-cured for 31 d at laboratory tem- perature (25±2 °C) and a relative humidity of 53±5 %. After initial curing, the specimens were dried at a temperature of 60 °C for two d. Subsequently, the speci- mens were subjected to gradual heating in a furnace at 200 °C, 400 °C, 600 °C, 800 °C, 1000 °C and 1200 °C. The temperature increase rate was 5 °C/min. A dwell of 60 min at each temperature was provided, in order to find out the effect of temperature on the specimens. After heat treatment, the specimens were left to cool down under laboratory conditions. 3 RESULTS AND DISCUSION 3.1 Material for the hardening process of alkali- activated slag mortars The experiment was employed to determine the microstructural changes during the hardening process of the alkali-activated slag composite with different admix- tures. Changes in the density of the material due to the process of hardening as well as the creation of micro- cracks due to the time of curing are reflected in the shift of dominant frequencies. Figure 3 shows the shift of the resonance frequency during 14 d after demoulding without curing. The frequency of the reference specimen (AAS) increased by about 37 % during the first 48 h from the start of the measurement and then decreased to a steady value of around 18 % from the initial value. The dominant frequency of the AAS+CNT (AAS+HPMC) specimen started 21 % (15 %) above the initial dominant frequency of the reference specimen (AAS). The domi- nant frequency of the AAS+CNT (AAS+HPMC) increased by about 35 % (30 %) from the initial value during the first 24 h and then decreased to a steady value of around 23 % (25 %) from the initial value. The pro- cess of hardening and the formation of a hard and dense structure caused an initial increase of the dominant frequencies. At a later time, the frequencies are again slightly shifted towards lower values. This phenomenon is probably associated with the drying process, which is followed by the shrinkage of the AAS matrix and the formation of micro-cracks. Figure 4 shows the shift of the resonance frequency during 14 d after 28 d of curing. The dominant frequency decreased for all the measure- D. [TEFKOVÁ et al.: EVALUATION OF THE DEGREE OF DEGRADATION USING THE IMPACT-ECHO METHOD ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 879–884 881 Figure 4: Change of relative dominant frequencies over time – after 28 d of curing Slika 4: Spreminjanje relativne prevladujo~e frekvence s ~asom – po 28 dneh utrjevanja Figure 3: Change of relative dominant frequencies over time - without curing Slika 3: Spreminjanje relativne prevladujo~e frekvence s ~asom – brez utrjevanja ment times to a stable value. For the AAS specimen there was a stable value of about 70 % of the initial value. For AAS+CNT (AAS+HPMC) there was a de- cline of about 25 % (20 %) from the initial value. This decline is mainly associated with the drying process, which is followed by shrinkage of the AAS matrix and probably the formation of micro-cracks. Whereas that dominant frequency obtained for specimens with admix- tures was higher than for the reference sample, then both admixtures have a positive effect on the formation of a structure of alkali-activated slag composite. Both cellu- lose derivatives that were added to mixture are able to retain water. These admixtures prevent the material from rapidly drying and the subsequent formation of the micro-cracks caused by drying shrinkage, which gene- rally occurs during hardening of the samples. Carbon nanotubes employed in one set of samples can act as a micro-reinforcement, it participates on the improvement of the mechanical properties. 3.2 Material for degradation of concrete samples by corrosion caused by the action of chlorides Measuring the effect of corrosion on reinforced con- crete samples was carried out on two sets of samples. The first set of samples was a reference and the second set of samples was subjected to corrosion by using accelerated degradation by chlorides in an aqueous NaCl solution. The measurements were always carried out after 20 cycles of alternating immersion in an aqueous NaCl solution and drying in an oven. In Figure 5 we can see the relative change of the dominant frequencies, when the piezoelectric sensor was placed at one end of the concrete beam and hit by a steel hammer at the oppo- site end of the concrete beam in the longitudinal direc- tion. Figure 6 shows the relative change of dominant fre- quencies for the same samples, but measurements were carried out so that the sensor was placed at one end un- covered reinforcement and the hit was made at the opposite end of reinforcement in the longitudinal direc- tion. The signal passing through the degraded sample changes the frequency by reflections on the inhomoge- neities, which is an indicator of changes in the structure. The results are referenced to a relative value and are displayed in the form of an arithmetic average (obtained from three independent measurements for reference sam- ples and from six measurements for degraded samples) and standard deviations as error bars. In the first case, for measurements of the samples on concrete, the dominant frequencies of reference samples was a decrease of 2.4 %, which can be explained by a lack of water for the hydration of the concrete, therefore this decrease is not as significant for immersed samples, only about 1.0 %. In the case of measurements on reinforcement the pheno- menon associated with hydration is not apparent, it is only within the measurement error. Between 20 and 40 cycles of degradation occurred in the samples to the formation of observable cracks due to expansion of the corrosion products. This phenomenon is reflected in a significant decrease of the monitored frequencies. For measurements on concrete the decrease was 2.5 % and for a measurement on reinforcement it was 1.9 %. Other frequency changes are no longer significant. The results obtained correspond to the processes that occur during the degradation of the reinforced concrete and reveal damage reinforced concrete. D. [TEFKOVÁ et al.: EVALUATION OF THE DEGREE OF DEGRADATION USING THE IMPACT-ECHO METHOD ... 882 Materiali in tehnologije / Materials and technology 50 (2016) 6, 879–884 Figure 7: Shift of dominant frequency induced by degradation at elevated temperatures (Arrangement U0-S0) Slika 7: Premik prevladujo~e frekvence povzro~en z razpadanjem pri povi{anih temperaturah (Postavitev U0-S0) Figure 5: Change of the relative dominant frequencies during degra- dation cycles (impact and sensor on material) Slika 5: Spreminjanje relativne prevladujo~e frekvence med cikli pro- padanja (udarec in senzor na materialu) Figure 6: Change of relative dominant frequencies during degradation cycles (impact and sensor on reinforcement) Slika 6: Spreminjanje relativne prevladujo~e frekvence med cikli pro- padanja (udarec in senzor na betonu) 3.3 Specimens intended to be subjected in a heating process The mortar specimens of the compositions were exposed to the temperatures of 200, 400, 600, 800, 1000 and 1200 °C. The test results of impact-echo are pre- sented below. Figure 7 presents the change of dominant frequency versus temperature at which the specimens of mortar compositions were subjected (arrangement U0-S0; longitudinal waves). For this measurement, the sensor was placed at the specimen’s end at its centre line direction, while the specimen was hit at the opposite end at the centre line direction – arrangement U0-S0. Longi- tudinal waves, which propagate within the sample at a speed of about 5100 m/s, can affect the mortar element oscillations. The exposure at elevated temperatures causes a decrease of the dominant frequency, leading to the conclusion that the material’s elastic modulus for each composition also decreases (E = 4·(f·L)2). Predo- minant frequencies are shifted towards to the lower fre- quency range in the course of the degradation. The change is more rapid in the temperature range 400–600 °C, where are intense impurities changed. It is clear that the predominant frequencies shifted down to- wards the lower frequency region. For the specimen that underwent a thermal stress at a temperature of 1200 °C it is clear that the predominant frequencies shifted upwards towards the higher-frequency region. It is evident that a structural change, accompanied by the creation of new crystal phases, takes place in the specimen at tempera- tures of about 1200 °C. Figure 8 shows the change of dominant frequency versus temperature when the arrangement was the U1-S1 one. In this case, transverse waves (gradual waves) are predominantly spread throughout the specimens. The difference between the U0-S0 and U1-S1 arrangements is that in the latter the measurement took place with the sensor being placed at the mid-point and perpendicular to the specimen. The specimen was hit at the mid-point opposite to the sensor. The dependence of frequency on temperature was similar to that observed when the U0-S0 arrangement was used, however the frequency values were lower in the case of the U1-S1 arrangement. This is similar to the case of the U0-S0 arrangement (Figure 7). The comparison of Fig- ures 7 and 8 indicates that the frequency change is slower when the arrangement U1-S1 is applied. In gene- ral, acoustic methods illustrated the physical changes in the structure of all the tested materials. A reduction of the predominant frequency values was observed. More- over, it was also observed in every case of elevated tem- perature. The heating up to 110 °C resulted in a loss of capillary water and a reduction of the cohesive forces (weakening of bonds) due to moisture evaporation. At about 170 °C decomposition of gypsum occurred, result- ing in expansive spalling. Between 250 and 300 °C the hydrated cement phases were decomposed, while above 300 °C it resulted in Ca(OH)2 decomposition. Further temperature increases up to 300 °C or 400 °C intensified the cement paste’s thermal decomposition and degrada- tion. All the mentioned changes resulted in the embrittle- ment and hardening of the tested materials. Thus, the observed reduction of frequency values was assumed to be due to the formation of micro-cracks. 4 CONCLUSIONS The paper deals with the results of measurements using the Impact-echo method on three applications in civil engineering. The aim of this paper was to study the application of the impact-echo method for the detection of flaws in composite materials during different stress situations (setting and hardening in air, degradation by corrosion caused by chlorides and exposing to elevated temperature). It is known that the impact response signal of a specimen is composed of frequencies corresponding to the modes of vibration of the specimen. A shift of the dominant frequency to a lower value is a key indication of the presence of the flaw. From the results obtained in the framework of our research group and the results demonstrated in this paper it can be summarized that the frequency inspection carried out by means of the Impact-echo method makes a convenient tool to assess the quality and lifetime of these composite materials when exposed to stress situations. Acknowledgement This paper has been worked out under the project GA^R No.16-02261S supported by Czech Science Foundation and the project No. LO1408 "AdMaS UP – Advanced Materials, Structures and Technologies", supported by Ministry of Education, Youth and Sports under the "National Sustainability Programme I" and under the project No.J-16-3254 supported by Faculty of Civil Engineering of BUT. D. [TEFKOVÁ et al.: EVALUATION OF THE DEGREE OF DEGRADATION USING THE IMPACT-ECHO METHOD ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 879–884 883 Figure 8: Shift of dominant frequency induced by degradation at elevated temperatures (Arrangement U1-S1) Slika 8: Premik prevladujo~e frekvence povzro~en z razpadanjem pri povi{anih temperaturah (Postavitev U1-S1) 5 REFERENCES 1 N. J. Carino, Structures Congress and Exposition 2001, Proceedings, American Society of Civil Engineers, Washington, DC, 2001, 1–18 2 B. Kucharczyková, P. Misák, T. Vymazal, Determination and evaluation of the air permeability coefficient using Torrent Permeability Tester, Russian Journal of Nondestructive Testing, 46 (2010) 3, 226–233, doi:10.1134/ S1061830910030113 3 T. Vymazal, N. @i`ková, P. Misák, Prediction of the risks of design and development of new building materials by fuzzy inference systems, Ceramics-Silikáty, 53 (2009) 3, 216–445 4 I. Pl{ková, M. Matysík, Z. 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Schmid, The Effect of the Carbon Nanotubes on the Mechanical Fracture Properties of Alkali Activated Slag Mortars, In Dynamic of Civil Engineering and Trans- port Structures and Wind Engineering, Applied Mechanics and Materials, Donovaly, 2014, 243–246, doi:10.4028/www.scientific. net/AMM.617.243 D. [TEFKOVÁ et al.: EVALUATION OF THE DEGREE OF DEGRADATION USING THE IMPACT-ECHO METHOD ... 884 Materiali in tehnologije / Materials and technology 50 (2016) 6, 879–884 R. SOKOLAR: NON-TRADITIONAL WHITEWARE BASED ON CALCIUM ALUMINATE CEMENT 885–889 NON-TRADITIONAL WHITEWARE BASED ON CALCIUM ALUMINATE CEMENT NETRADICIONALNI PORCELAN NA OSNOVI KALCIJ ALUMINATNEGA CEMENTA Radomir Sokolar Brno University of Technology, Faculty of Civil Engineering, Veveri 95, 602 00 Brno, Czech Republic sokolar.r@fce.vutbr.cz Prejem rokopisa – received: 2015-07-01; sprejem za objavo – accepted for publication: 2015-10-27 doi:10.17222/mit.2015.182 The article introduces the differences in the properties of whiteware (porosity, strength, thermal expansion coefficient) when a non-traditional binder is used. Pure calcium aluminate cement and a mixture of kaolin and calcium aluminate cement compared with traditional plastic raw material for whiteware – kaolin – are used for the preparation of whiteware bodies with a constant content of non-plastic raw materials: K-Na feldspar and quartz sand. The results are discussed in connection with the micro- structure of the fired body of prepared whitewares (mineralogical composition). Calcium aluminate cement (CAC) in whiteware raw-material mixtures is an interesting alternative to kaolin for a higher strength of the green and fired bodies. Using calcium aluminate cement reduces the sintering temperature of the fired body and significantly changes its mineralogical composition: anorthite is the main mineralogical phase instead of mullite, which is typical for standard porcelain bodies made from raw-material mixtures based on kaolin. The coefficient of thermal expansion increases with an increasing content of CAC in the raw-materials mixture. Keywords: calcium aluminate cement, kaolin, whiteware ^lanek predstavlja razlike v lastnostih porcelana (poroznost, trdnost, koeficient toplotnega raztezka), ~e se uporabi neobi~ajno vezivo. ^isti kalcijev aluminatni cement in me{anica kaolina in kalcijevega aluminatnega cementa, v primerjavi s tradicionalno plasti~no sestavino porcelana – kaolina, so bili uporabljeni za kerami~na telesa s konstantno vsebnostjo neplasti~nih surovin; K-Na glinenec in kvar~ni pesek. Rezultati so razlo`eni v povezavi z mikrostrukturo telesa pripravljene keramike (mineralo{ka sestava) po `ganju. Kalcijev aluminatni cement (CAC), v me{anici surovin za keramiko, je zanimivo nadomestilo za kaolin, za vi{jo trdnost zelenega in `ganega telesa. Uporaba kalcijevega aluminatnega cementa zni`a temperaturo sintranja `ganega telesa in mo~no spremeni njegovo mineralo{ko sestavo; anortit je glavna mineralo{ka faza namesto mulita, ki je zna~ilen v standardnih porcelanskih telesih, izdelanih iz me{anice surovin na osnovi kaolina. Koeficient toplotnega raztezka se pove~uje z ve~anjem vsebnosti CAC v me{anici surovin. Klju~ne besede: kalcijev aluminatni cement, kaolin, porcelan 1 INTRODUCTION Whiteware is a traditional ceramic material that has been manufactured for centuries from a mixture of kao- lin, quartz and feldspar. A new type of porcelain body – the anorthite porcelain body – from feldspar, quartz and calcium aluminate cement without using any other binders and plastic ceramic raw materials (clays, kaolins) – was fabricated at a temperature of 1300 °C and the physical and mechanical properties were investigated. The addition of calcium aluminate cement as a substitute for clays exhibits a relatively high green strength and lowers the density due to the formation of anorthite in all the fired bodies.1 Increasing the strength of the green body, reducing the coefficient of linear thermal expansion and increas- ing the whiteness of the fired body can be achieved primarily by replacing the kaolin for calcium aluminate cement (70 % of Al2O3). A negative aspect of using calcium aluminate cement with a high Al2O3 content is reducing the sintering activity of the body and therefore the need for higher firing temperatures.2 The single-phase anorthite ceramic was fabricated (from a mixture of ball clay, quartz, calcite, feldspar and alumina) by slip casting and sintering at 1230 °C for 1 h. It has a high flexural strength of 103 MPa, which is higher than that of the conventional porcelain. The single-phase anorthite ceramic had a relatively low (4.9×10–6 K–1) thermal expansion coefficient, which can be easily matched with an applicable glaze and achieve an excellent thermal shock resistance.3 A new, porcelainised stoneware material based on anorthite was prepared from an undefined ratio of wollastonite, alumina, quartz, Ukrainian Ball clay and magnesia. After firing at 1220 °C we obtained a porce- lainised stoneware body containing 70 % crystalline and 30 % glassy phases with a high modulus of rupture (110 MPa) that is two times higher compared to conventional porcelainised stoneware materials based on mullite.4 Calcium aluminate cements are white (according to the Al2O3 content), high-purity hydraulic bonding agents providing controlled setting times and strength develop- ment for today’s high-performance refractory products. Materiali in tehnologije / Materials and technology 50 (2016) 6, 885–889 885 UDK 67.017:621.315.612:621.742.45 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)885(2016) There are only laboratory results in the area of re- placing kaolin with CAC for the production of whiteware without any practical industrial impact. The aim of the article is to introduce the differences in the properties (porosity, strength, thermal expansion coefficient) of whiteware when pure calcium aluminate cement or a combination of the binders – a mixture of kaolin and calcium aluminate cement – instead of kaolin is used. The results will be discussed in connection with the mineralogical composition of the fired whiteware body. 2 CHARACTERIZATION OF USED MATERIALS Calcium aluminate cement SECAR® 51 (CAC51) with 51 % of Al2O3 content (on average) was used for the experimental study. Calcium aluminate cement contains mainly CA and C12A7, C2AS and CT as minor minera- logical phases. The chemical composition is clear from Table 1. SECAR®51 is a fused hydraulic binder with a mineralogy focused on mono-calcium aluminate to give a strong hydraulic activity and impart excellent mecha- nical properties to conventional concretes. This binder is recommended when rapid hardening properties are required. It is adapted to all types of placing methods, particularly casting and gunning at a service temperature of 1400 °C and above. The surface area is about 4 m2 g–1. The equivalent mean spherical diameter (Laser PSD) of CAC51 is d(0.5) = 14 μm. Table 1: Typical chemical composition of used binders: calcium aluminate cement CAC51 and kaolin Tabela 1: Zna~ilna kemijska sestava uporabljenih veziv: kalcijev aluminatni cement CAC51 in kaolin Material CaO Al2O3 (K)Na2O SiO2 Fe2O3 MgO LOI CAC51 36.9 52.2 0.5 4.9 1.8 1.0 0.0 Kaolin 0.7 36.6 1.2 46.8 0.9 0.5 13.2 LOI – lost of ignition Kaolin Zettlitz Ia is the most renowned and the oldest product of the company Sedlecký kaolin a. s. It has been produced since 1892 when the company Zettlitzer Kaolinwerke AG was founded; the present company Sedlecký kaolin a.s. is its successor. The major features of kaolin Zettlitz Ia are: • relatively high plasticity, which enables good work- ability in the raw state, • easy deflocculation with available deflocculants (addition of 0.1 % Na2CO3), • high content of Al2O3 and a low content of alkalis, which provides high stability in fire. The original use of the kaolin Zettlitz Ia was in porcelain production. Many producers use it today as the main plastic component in bodies for plastic moulding, slip casting or isostatic pressing of tableware. Good plas- ticity in the raw state and high stability in fire are also employed in the production of ceramic rollers for fast firing kilns. A low content of harmful substances and plasticity play a significant role in pencil production. This feature is also important for use in cosmetics. The mineralogical composition of the used washed kaolin is 91 % of kaolinite, 2 % of quartz and about 7 % of mica minerals. The surface area of the kaolin is 17.5 m2 g–1. Industrially milled potassium-sodium feldspar rock and pure quartz sand were used for the experiments as non-plastic materials. The mineralogical composition of the sodium-potassium feldspar is potassium feldspar (microcline) 20.0 %, sodium feldspar (albite) 22.6 %, calcium feldspar (anorthite) 2.4 % and quartz 55.0 %. The chemical composition of the used feldspar rock (Table 2) reflects its mineralogical composition and the volume of different types of pure feldspars (potassium, sodium and calcium). The micro-milled quartz sand ST 6 from Sklopísek Støele~ Company (Czech Republic) was used. The ma- terial is produced by milling in an iron-free environment, by classification with the use of air separators. The raw material used for the production of micro-milled sands, i.e., silica flour, is treated silica sand with a SiO2 content above 99 %. The chemical purity, favourable particle size distribution, chemical inertness and the hardness of the micro-milled sands – silica flour – is appreciated in the production of glass fibres, ceramic enamels, glazes, as a filler in plastics, in the production of special mortar mixtures, tile adhesives and in the foundry industry for the production of moulds. The surface area of the used micro-milled quartz sand is 4.38 m2 g–1. Granulometries of the milled feldspar and milled quartz were determined according to the particle size distribution (laser particle size analyser Malvern Master- sizer 2000). The equivalent mean spherical diameters d(0.5) = 20.8 μm (feldspar) and 16.0 μm (quartz sand) are suitable for the production of the porcelain body.5 3 METHODOLOGY Three different raw-material mixtures on the basis of different binder bases (calcium aluminate cement R. SOKOLAR: NON-TRADITIONAL WHITEWARE BASED ON CALCIUM ALUMINATE CEMENT 886 Materiali in tehnologije / Materials and technology 50 (2016) 6, 885–889 Table 2: Chemical composition of used non-plastic materials: K-Na feldspar rock and quartz sand Tabela 2: Kemijska sestava uporabljenih neplasti~nih materialov: K-Na glinenec in kvar~ni pesek Content, in mass fractions (w/%) SiO2 Al2O3 Fe2O3 MnO TiO2 CaO MgO K2O Na2O LOI Total Feldspar 79.76 12.37 0.42 0.00 0.05 0.48 0.10 3.35 2.67 0.80 100.00 Quartz 99.60 0.20 0.03 – – 0.10 0.10 0.00 – – 100.03 LOI – lost of ignition CAC51, washed kaolin and its mixture), industrially milled K-Na feldspar rock and industrially milled glass quartz sand were prepared (Table 3). Table 3: Composition of raw-material mixtures (test samples) in mass fractions, (w/%) Tabela 3: Sestava me{anice surovin (preizkusni vzorci), v masnih odstotkih, (w/%) Sample Kaolin CAC51 Feldspar Quartz sand A 25 50 25 B 25 C 15 10 Raw-material mixtures according to Table 3 were homogenized for 24 h in a homogenizer. The dry mixture was then moistened with 9 % mass of water. The moistened mixtures were pressed through the 1-mm sieve. The pressing granulate was thus prepared and subsequently mixed for 12 h in the closed plastic vase of the homogenizer to produce a homogenous moisture. Test samples with a green-body size of 100 mm × 50 mm × 10 mm were uniaxially pressed in a steel mould at 20 MPa with 30 s of soaking time at the maximum pressure. Drying in air at a temperature of about 21 °C was followed by final drying in a laboratory drier at 110 °C to achieve a constant weight. The green bodies were fired in an electric laboratory furnace at different temperatures with a heating rate 10 °C/min and a 30-min soaking time at the maximum temperature to achieve the sintering temperature, which is defined as the temperature when the fired body has a water absorption E = 2 %. The subsequent cooling proceeded spontaneously, following the natural cooling rate of the furnace. After firing, the body properties (6 test samples) were defined according to the official standard EN ISO 10545 (vacuum water absorption Ev, modulus of rupture MOR). The mineralogical compo- sitions of the pure feldspars and fired bodies were determined by X-ray diffraction (XRD). The XRD analysis was performed with a Phillips PW 1170 diffrac- tometer using a Cu anti-cathode (1 = 0.15406 nm), accelerating voltage 40 kV, beam current 25 mA and scanning rate 1° 20’ min–1. 4 RESULTS OF EXPERIMENTS The mixtures containing CAC51 (B and C) show a significantly higher sintering activity (the ability of the body to create a compact body during the firing through the merging of the grains) according to dilatometric curves (Figure 1). The sintering activity is described by the shrinkage of the tested samples during the firing. The higher content of CAC51 in the raw-material mixture caused a lower sintering temperature of the body during the firing (Table 4). The sintering temperatures (tempe- rature when the fired body has water absorption E = 2 %) are: Sample A: approximately 1270 °C, Sample B: approximately 1150 °C Sample C: approximately 1250 °C Table 4: Water absorption of fired bodies depending on the firing tem- perature Tabela 4: Absorpcija vode v `ganih telesih v odvisnosti od tempera- ture `ganja Sample Water absorption after the firing (in mass fractions, (w/%) 1150 °C 1250 °C 1270 °C A 10.7 4.5 1.6 B 1.8 C 7.6 1.5 These are different results than when the calcium aluminate cement with a higher content of Al2O3 (70 %) was used.2 This situation reflects the different chemical composition of CA cements: CAC51 contains a higher portion of fluxing oxides (Fe2O3 and CaO especially) to CAC70. The mineralogical composition of all the tested bodies after the firing is characterized by the existence of the quartz and the glass phase. The fired body based on R. SOKOLAR: NON-TRADITIONAL WHITEWARE BASED ON CALCIUM ALUMINATE CEMENT Materiali in tehnologije / Materials and technology 50 (2016) 6, 885–889 887 Figure 2: XRD of fired bodies based on different binders: kaolin or CAC51 (M-mullite, Q-quartz, A-anorthite) Slika 2: Rentgenogram `ganih teles na osnovi razli~nih veziv- kaolin ali CAC51 (M-mulit, Q- kvarc, A-anortit Figure 1: Thermal dilatometric analysis during the firing (1200 °C, 5 °C/min without soaking time) Slika 1: Termi~na dilatometri~na analiza med `ganjem (1200 °C, 5 °C/min brez ~asa zalaganja) kaolin also contains mullite, while the fired bodies based on CAC51 contain anorthite without mullite (Figure 2). In the case of a combination of the binders, the mullite phase is missing. By comparing the modulus of rupture (MOR) of the fired samples at their sintering temperatures (with similar porosity), it is evident that anorthitic type of body (Figure 2) with a lower content of glass phase (based on calcium aluminate cement) comprises a higher modulus of rupture. This confirms the results published in 3,4. Table 5: Modulus of rupture of tested whiteware bodies after the firing at the sintering temperatures Tabela 5: Prelomni modul preizku{enih porcelanskih teles po `ganju na temperaturi sintranja Sample Modulus of rupture (MPa) A 34.4 (1270 °C) B 37.4 (1150 °C) C 33.8 (1250 °C) The thermal expansion coefficient is the main factor when considering the thermal matching between the glaze and the body, and indirectly influences the thermal shock resistance.6 The coefficient of linear thermal ex- pansion was calculated from dilatometric curves (Figure 3) in the different temperature ranges from 30 °C to 1000 °C (Table 6). The whiteware fired body based on calcium aluminate cement CAC51 (samples B and C) shows a higher coefficient of thermal expansion com- pared with the kaolin-based sample A (Figure 3). This is an unexpected result due to the formation of anorthite in the samples B and C (Figure 3). A very important tech- nical property of anorthite is its low coefficient of linear thermal expansion coefficient of 4.82×10–6 K–1 7 (com- pared with mullite 6.00×10–6 K–1).8 Different results are achieved when the CAC with 70 % of Al2O3 has been used as the replacement for kaolin: the coefficient of linear thermal expansion decreased from 9.24×10–6 to 8.08×.10–6 K–1.2 This may be connected with the higher content of Fe2O3 in CAC51 (Table 1). Table 6: The values of the coefficient of linear thermal expansion  (K–1) depending on the temperature range and raw-materials mixture Tabela 6: Vrednosti koeficienta linearnega toplotnega raztezka  (K–1) v odvisnosti od temperaturnega obmo~ja in me{anice surovin Temperature range A C B ·10–6 (K–1) ·10–6 (K–1) ·10–6 (K–1) 20-200 °C 5.73 6.43 6.35 20-300 °C 6.05 6.74 6.85 20-400 °C 6.42 7.13 7.31 20-500 °C 7.01 7.63 7.83 20-600 °C 8.25 8.71 8.73 20-700 °C 7.63 8.07 8.20 20-800 °C 7.15 7.71 8.20 20-900 °C 6.69 7.46 7.92 20-1000 °C 6.12 7.03 7.23 5 CONCLUSION The presence of calcium aluminate cements in the raw-materials mixture significantly changes the mine- ralogical composition of a fired whiteware body, i.e., anorthite is the main mineralogical phase instead of mullite, which is typical for standard porcelain bodies made from raw-material mixtures based on kaolin. The anorthite type of porcelain body is very suitable for the high strength of the fired body. Using calcium aluminate cement with a lower content of Al2O3 (51 %) reduces the sintering temperature of the body, impairs the whiteness of the body and increases the coefficient of linear ther- mal expansion. Acknowledgements This work was financially supported by the Czech Science Foundation, research project No. P104/13/ 23051S "Anorthite porcelain body on the basis of aluminous cement". 6 REFERENCES 1 W. Tai, K. Kimura, K. Jinnai, A new approach to anorthite porcelain bodies using nonplastic raw materials, Journal of the European Ceramic Society, 22 (2002) 4, 463, doi:10.1016/S0955-2219(01) 00317-X 2 R. Sokoláø, L. Vodová, Whiteware Bodies without kaolin, Inter- ceram., 63 (2014) 1–2, 19–21 3 X. Cheng, S. Ke, Q. Wang, H. Wang, A. Shui, P. Liu, Fabrication and characterization of anorthite-based ceramic using mineral raw ma- terials, Ceramics International, 38 (2012) 4, 3227–3235, doi:10.1016/j.ceramint.2011.12.028 4 M. U. Taskiran, N. Demirkol, A. Capoglu, A new porcelainised stoneware material based on anorthite, Journal of the European Cera- mic Society, 25 (2005) 4, 293–300, doi:10.1016/j.jeurceramsoc. 2004.03.017 5 A. De Noni, D. Hotza, V. C. Soler, E. Sanchez Vilchez, Effect of quartz particle size on the mechanical behaviour of porcelain tile subjected to different cooling rates, Journal of the European Ceramic R. SOKOLAR: NON-TRADITIONAL WHITEWARE BASED ON CALCIUM ALUMINATE CEMENT 888 Materiali in tehnologije / Materials and technology 50 (2016) 6, 885–889 Figure 3: Relative expansion of the fired bodies depending on the temperature for the coefficient of linear thermal expansion deter- mination Slika 3: Relativni raztezek `ganih teles v odvisnosti od temperature pri dolo~anju linearnega toplotnega raztezka Society, 29 (2009) 6, 1039–1046, doi:10.1016/j.jeurceramsoc.2008. 07.052 6 Y. Hirata, Theoretical analyses of thermal shock and thermal expan- sion coefficients of metals and ceramics, Ceramics International, 41 (2015) 1, 1145–1153, doi:10.1016/j.ceramint.2014.09.042 7 M. Potuzak, M. Solvang, D. Dingwell. Temperature independent thermal expansivities of calcium aluminosilicates melts between 1150 and 1973 K in the system anorthite–wollostanite–gehlenite (An–Wo–Geh): a density model, Geochim. Cosmochim., 70 (2006) 3059–3074, doi:10.1016/j.gca.2006.03.013 8 M. A. Camerucci, G. Urretavizcaya, M. S. Castro, A. L. Cavalieri, Electrical properties and thermal expansion of cordierite and cor- dierite-mullite materials, Journal of the European Ceramic Society, 21 (2001) 16, 2917–2923, doi:10.1016/S0955-2219(01)00219-9 R. SOKOLAR: NON-TRADITIONAL WHITEWARE BASED ON CALCIUM ALUMINATE CEMENT Materiali in tehnologije / Materials and technology 50 (2016) 6, 885–889 889 B. MA[EK et al.: BEHAVIOUR OF NEW ODS ALLOYS UNDER SINGLE AND MULTIPLE DEFORMATION 891–898 BEHAVIOUR OF NEW ODS ALLOYS UNDER SINGLE AND MULTIPLE DEFORMATION OBNA[ANJE NOVIH ODS ZLITIN PRI ENOJNI IN VE^KRATNI DEFORMACIJI Bohuslav Ma{ek1, Omid Khalaj1, Zby{ek Nový2, Tomá{ Kubina2, Hana Jirkova1, Jiøí Svoboda3, Ctibor [tádler1 1The Research Centre of Forming Technology, University of West Bohemia, Univerzitní 22, 306 14, Pilsen, Czech Republic 2COMTES FHT a.s., Prùmyslová 995, 334 41 Dobøany, Czech Republic 3Institute of Physics of Materials, Academy of Sciences Czech Republic, @i`kova 22, 616 62, Brno, Czech Republic khalaj@vctt.zcu.cz Prejem rokopisa – received: 2015-07-01; sprejem za objavo – accepted for publication: 2015-11-13 doi:10.17222/mit.2015.156 The application of innovative processing techniques to conventional raw materials can lead to new structural materials with specific mechanical and physical properties, which open up new possibilities of use in some areas of industry. The processing is enabled by powder metallurgy, which utilizes powders consisting of a metal matrix with dispersed stable particles achieved by mechanical alloying and their hot consolidation by rolling. New oxide dispersion strengthened (ODS) Fe–Al-based alloys are tested under different single and multiple thermomechanical treatments at different temperatures. The results show that new ODS alloys are significantly affected by the thermo-mechanical treatment, leading to microstructural changes. Their analysis is performed using different analytical methods such as optical microscopy, scanning electron microscopy and X-ray diffraction analysis. Keywords: ODS alloys, composite, steel, Fe-Al Uporaba inovativnih tehnik preoblikovanja na obi~ajnih materialih lahko privede do novih konstrukcijskih materialov s specifi~nimi mehanskimi in fizikalnimi lastnostmi, ki odpirajo nove mo`nosti uporabe v industriji. Metalurgija prahov omogo~a uporabo prahov s kovinsko osnovo z dispergiranimi stabilnimi delci, ki jih dobimo pri mehanskem legiranju in vro~i konsolidaciji z valjanjem. Nove zlitine Fe-Al, disperzijsko utrjene z oksidi (ODS), so bile preizku{ene pri razli~ni, eno- ali ve~stopenjski obdelavi pri razli~nih temperaturah. Rezultati ka`ejo, da ima termomehanska obdelava novih ODS zlitin mo~an vpliv, ki se vidi v spremembah mikrostrukture. Analiza je bila izvedena s pomo~jo razli~nih analitskih metod, kot so: svetlobna mikroskopija, vrsti~na elektronska mikroskopija in rentgenska difrakcijska analiza. Klju~ne besede: ODS zlitine, kompozit, jeklo, Fe-Al 1 INTRODUCTION The demand to increase the efficiency of processes in most industrial applications requires, in many cases, metallic materials that can operate at high temperatures, and often at high stresses, in corrosive environments. The presently used high-temperature Ni-, Co- and Fe- based alloys are strengthened by a combination of solid- solution and precipitation hardening, the effectiveness of which strongly decreases with increasing temperature. ODS alloys contain small amounts (0.5–1 % of weight fractions) of finely dispersed oxide phase (mostly yttrium), which is thermodynamically much more stable than other strengthening phases such as ’ or carbides, present in conventional high-temperature alloys.1 There- fore, the strengthening imparted by the oxide dispersions is retained up to very high temperatures because only li- mited coarsening or dissolution of the particles occurs.2,3 In addition, the presence of the fine dispersions com- bined with a very coarse-grained microstructure that is stable over long exposure times leads to excellent creep resistance up to higher temperatures than those that can be achieved with conventional wrought or cast alloys.4,5 The ODS alloys commercially produced at the end of the 20th century and the beginning of the 21st century are represented by MA 956 or MA 9576, PM 2000 or PM 20106, ODM alloys7 and 1DK or 1DS8 with a ferritic matrix by ODS Eurofer steels with a tempered ferritic- martensitic matrix9 and by austenitic Ni-ODS PM 1000 or Ni-ODS PM 3030.10 ODS alloys are produced by high-energy milling of powder mixtures consisting of the alloying elements, master alloys and the oxide disper- sion. The volume fraction of dispersed spherical oxides (usually Y2O3) is typically below 1 % and the oxides are typically of a mean size of 5–30 nm. The mechanically alloyed powder is then consolidated at high temperatures and pressures to produce the bulk material in the form of bar or tube stock. Afterwards, different thermomecha- nical treatments are applied to optimize its microstruc- ture and mechanical properties. In the consolidation step the processing temperatures are critical in order to retain the nanocrystalline structure generated during the me- chanical alloying and to impede particle coarsening and grain growth.11–14 The Ni- and Fe-based ODS alloys rely on the formation of slowly growing and strongly Materiali in tehnologije / Materials and technology 50 (2016) 6, 891–898 891 UDK 67.017:669.018:537.533:621.763 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)891(2016) adherent chromium and aluminium scales for their high-temperature oxidation/corrosion resistance. Be- cause of the lower diffusion coefficient, austenitic ODS alloys show a better creep resistance for the same oxide volume fraction and contain some minimum chromium/ aluminium content to guarantee sufficient oxidation resistance. However, the resistance to the coarsening of oxides is given by the product of the solubility of oxygen in the matrix and its diffusion coefficient;16 this factor is more advantageous for ferritic ODS alloys. Also, a sufficient content of Al and/or Cr in the ODS alloy is decisive for its oxidation resistance.17–19 This is probably the reason why the application of ferritic ODS steels dominates.15–23 The new ODS alloys consist of a ferritic Fe-Al matrix strengthened with about 6 to 10 % volume fractions of Al2O3 particles.24,25 In order to get a more detailed insight into these new groups of materials, an experimental pro- gramme was carried out to better understand their processing behaviour and their operational properties. 2 EXPERIMENTAL PART Mechanically alloyed (MA) powders were prepared in a low-energy ball mill, developed by the authors (Fig- ure 1), which enables evacuation and filling by oxygen. It has two steel containers (each 24 L) and each con- tainer is filled with 80 steel balls of diameter 40 mm. The revolution speed is variable between 20 min–1 to 75 min–1. The mechanically alloyed powders consisting of Fe10wt%Al matrix and 6 % to 10 % volume fractions of Al2O3 particles were deposited into a steel container of diameter 70 mm, evacuated and sealed by welding (Fig- ure 2). The steel container was heated up to a tempera- ture of 800–900 °C and rolled by a hot-rolling mill (Figure 3) to a thickness of 20–25 mm in the first rolling step and then heated up to a temperature of 1100 °C and rolled to a thickness of 9 mm in the second step. A 6-mm-thick sheet of the ODS alloy was produced in this way. Afterwards, the specimens were cut by water jet. In order to investigate the thermomechanical treat- ment of specimens, a servohydraulic MTS thermomecha- nical simulator (Figure 4) was used, which allows the running of various temperature-deformation paths necessary to find conditions leading to, e.g., the most effective grain coarsening by recrystallization. Several B. MA[EK et al.: BEHAVIOUR OF NEW ODS ALLOYS UNDER SINGLE AND MULTIPLE DEFORMATION 892 Materiali in tehnologije / Materials and technology 50 (2016) 6, 891–898 Figure 3: Rolling process: a) hot-rolling mill, b) steel container in rolling process Slika 3: Valjanje: a) ogrodje za vro~e valjanje, b) zbirnik med postop- kom valjanja Figure 1: Low-energy mill for mechanical alloying Slika 1: Nizko energijski mlin za mehansko legiranje Figure 2: Container for mechanical alloyed powder Slika 2: Zbirnik za mehansko legiran prah procedures of thermomechanical treatment were de- signed and carried out, which differed in the number of deformation steps characterized by different strains, strain rates and temperatures. The thermomechanical simulator also allows the combination of tensile and compressive deformation, thus accumulating a high plas- tic deformation (and a high dislocation density) in the specimen. 2.1 Preparation of specimens One specimen (Figure 5) was selected from several examples regarding their most homogeneous temperature fields. The steel containers were removed from all the specimens that were cut by water jet in a longitudinal direction (Figure 6). The thickness of specimens was approximately 6 mm after grinding. Six types of material were used in this research, as described in Table 1. All these materials are based on a Fe10wt%Al ferritic matrix with different particle sizes and volume fractions in % of Al2O3. Al2O3 powder was added to prepare the composite, fine oxides in ODS alloys were obtained by internal oxidation during mecha- nical alloying and precipitated during hot consolidation. The microscopic SEM observations indicated several inhomogeneities due to sticking of the material during mechanical alloying on the walls of the milling con- tainer. These inhomogeneities can also influence the me- chanical and fracture properties of the material, but the mechanical alloying process is steadily optimized with respect to the homogeneity of the materials. Table 1: Material parameters Tabela 1: Parametri materiala Mate- rial No. Material type Milling time (h) Ferritic matrix (% of mass fractions) % of volume fractions of Al2O3 Typical particle size (nm) 1 Composite – Fe10%Al 10 300 2 ODS Alloy 100 Fe10%Al 6 50–200 3 ODS Alloy 150 Fe10%Al 6 50–150 4 ODS Alloy 200 Fe10%Al 6 30–150 5 ODS Alloy 245 Fe10%Al 7 20–50 6* ODS Alloy 245 Fe10%Al 7 20–50 * Different rolling force 2.2 Testing programme The test programme was divided into six different series. The tests are summarized in Table 2. Single deformation tests series were carried out to investigate the thermomechanical behaviour of the diffe- rent materials (1 to 4) at different temperatures regarding single tensile loading with a constant strain rate of 1 s–1 (Figure 7). In order to give a clearer comparison of the results, only the results at room temperature (RT), 800 °C and 1200 °C are presented. Multiple deformation-test series were carried out to investigate the thermomechanical behaviour of Materials B. MA[EK et al.: BEHAVIOUR OF NEW ODS ALLOYS UNDER SINGLE AND MULTIPLE DEFORMATION Materiali in tehnologije / Materials and technology 50 (2016) 6, 891–898 893 Figure 6: Position of specimens on rolled semi-product Slika 6: Polo`aj vzorcev na valjanem polproizvodu Figure 4: Treatment on thermomechanical simulator Slika 4: Obdelava na termomehanskem simulatorju Figure 5: Specimen dimensions Slika 5: Dimenzije vzorca Figure 7: Treatment no. 1 Slika 7: Obdelava {t. 1 5 and 6 at 1200 °C regarding multiple tensile loading with a constant strain rate of 1 s–1 (Figure 8) followed by two different holding times (10 s and 30 s). 3 RESULTS AND DISCUSSION 3.1 Single deformation-test series Single deformation-test series were carried out in order to investigate different materials under different conditions. Figure 9 shows the stress-strain curves for all the materials at different temperatures regarding the 5 % compression corresponding to treatment number 1. Material 2 exhibits a better strength at 30 °C and 800 °C, but at 1200 °C Material 1 shows a better strength. The hot-working behaviour of alloys is generally reflected by flow curves, which are a direct consequence of micro- structural changes: the nucleation and growth of new grains, dynamic recrystallization (DRX), the generation of dislocations, work hardening (WH), the rearrange- ment of dislocations and their dynamic recovery (DRV). In the deformed materials, DRX seems to be the pro- minent softening mechanism at high temperatures. DRX occurs during the straining of metals at high temperature, characterized by nucleation of low-dislocation-density grains and their posterior growth to produce a homo- geneous grain structure if a dynamic equilibrium is reached. Material 4 showed a strange curve shape at 800 °C. The test was repeated several times and similar beha- viour was observed. It could be concluded that it happens because of the inhomogeneity of the microstruc- ture of this material. B. MA[EK et al.: BEHAVIOUR OF NEW ODS ALLOYS UNDER SINGLE AND MULTIPLE DEFORMATION 894 Materiali in tehnologije / Materials and technology 50 (2016) 6, 891–898 Figure 8: Treatment no. 2 Slika 8: Obdelava {t. 2 Table 2: Parameters of test programme Tabela 2: Parametri programa preizkusa Test series Materialno. Treatment no. Treatment type Maximum temperature (°C) Number of tests Purpose of tests A 1 1 Single 1200, 1100, 1000, 900, 800, RT 6 Single deformation thermomechanical behaviour B 2 1 Single 6 C 3 1 Single 6 D 4 1 Single 6 E 5 2 Multiple 1200 2 Multiple deformation thermomechanical behaviourF 6 2 Multiple 2 Figure 9: Stress-strain curves (5 % compression) for: a) RT, b) 800 °C, c) 1200 °C Slika 9: Krivulje napetost-raztezek (5 % stiskanje) za: a) RT, b) 800 °C, c) 1200 °C Figure 10 shows the stress-strain curves for Mate- rials 1 to 4 at different temperatures corresponding to the 3 % tension of treatment number 1 (Figure 7). As can be seen in Figure 10, Material 2 shows a higher strength at 30 °C and 800 °C, but at 1200 °C, again Material 1 shows a better strength. All four materials have almost the same elastic modulus and none of them failed during 3 % deformation. The yield stress as well as the shape of the flow curves is sensitive to temperature. Comparing all these curves, it is found that decreasing the deforma- tion temperature increases the yield stress level, in other words, it prevents the occurrence of softening due to dynamic recrystallization (DRX) and dynamic recovery (DRV) and allows the deformed metals to exhibit work hardening (WH). For every curve, after a rapid increase in the stress to a peak value, the flow stress decreases monotonically towards a steady-state regime with a varying softening rate, which typically indicates the onset of DRX (Figure 9c). Figure 11 shows the stress-strain curves for Mate- rials 1 to 4 at different temperatures corresponding to the 50 % tension of treatment number 1 (Figure 7). All four B. MA[EK et al.: BEHAVIOUR OF NEW ODS ALLOYS UNDER SINGLE AND MULTIPLE DEFORMATION Materiali in tehnologije / Materials and technology 50 (2016) 6, 891–898 895 Figure 11: Stress-strain curves (50 % tension) for: a) RT, b) 800 °C, c) 1200 °C Slika 11: Krivulja napetost-raztezek (50 % natezna obremenitev) za: a) RT, b) 800 °C, c) 1200 °C Figure 10: Stress-strain curves (3 % tension) for: a) RT, b) 800 °C, c) 1200 °C Slika 10: Krivulja napetost-raztezek (3 % natezna obremenitev) za: a) RT, b) 800 °C, c) 1200 °C materials failed at RT, but only two materials failed below 50 % tension at higher temperatures. Material 2 failed at 34 % strain and Material 4 failed at 44 % strain at 800 °C. At 1200 °C, only Material 1 failed at 41 % and Material 2 failed at 45 % strain. From these curves, it can also be seen that the stress evolution with strain exhibits three distinct stages. In the first stage work hardening (WH) predominates and causes dislocations to polygonize into stable sub- grains. Flow stress exhibits a rapid increase with in- creasing strain up to a critical value. Then DRX occurs due to a large difference in dislocation density within the subgrains or grains. When the critical driving force of DRX is attained, new grains are nucleated along the grain boundaries, deformation bands and dislocations, resulting in the formation of equiaxed DRX grains. In the second stage, flow stress exhibits a smaller and smaller increase until a peak value or an inflection of the work-hardening rate is reached. This shows that the ther- mal softening due to DRX and dynamic recovery (DRV) becomes more and more important and it exceeds WH. In the third stage, three types of curves can be re- cognized: • Decreasing gradually to a steady state with DRX softening (Material 3 & 4 in Figure 11c), • Increasing continuously with significant work-hard- ening (Material 1 & 2 in Figure 11b), • Decreasing continuously with significant DRX soft- ening. 3.2 Multiple deformation-test series Multiple deformation-test series were carried out in order to investigate the material behaviour under various conditions. Figure 12 shows the stress-strain curves for both materials at different holding times following the 5 % tension during treatment number 2 (Figure 8). Both materials show approximately the same behaviour under the multiple tensile loading. However, Material 6 exhi- bits greater strength for both holding times (10 s and 30 s). It is supposed that the oxide particles prevent un- desirable cyclic softening, which is observed in ferritic- martensitic steels.22 Obviously, the oxide particles strengthen the material substantially, nevertheless, cyclic softening is observed at both holding times. The cyclic softening rate depends on the applied loading. A higher strain amplitude results in a higher softening rate. For instance, the softening rate was about 23 % during the second cycle in Material 5, while it decreased to 12 % during the last cycles. Although the softening in ODS steel is lower than in the ferritic-martensitic steel26, it indicates that oxide dispersion itself does not guarantee a stable cyclic behaviour and other microstructural aspects have to be taken into account. It is obvious that the stress amplitude decreased with an increasing number of cycles, while the amplitude of the plastic strain in- creased. The softening rate in Material 6 is lower than in Material 5, as observed for both holding times. The slight cyclic hardening is observed only during the first cycle in Material 6 with 30 s holding time (Figure 12b), while continuous softening behaviour is observed in the remaining part of the curve. 4 CONCLUSIONS This paper outlines the results of the characterization of the single and multiple deformation thermomecha- nical behaviour of a new generation of ODS alloys. Six materials differing from each other in the amount and size of the oxides embedded in the ferritic matrix were tested under different conditions. The advantages of all the materials are their low-cost and creep-corrosion and oxidation-resistance due to the Fe–Al-based ferritic ma- trix of the ODS alloy. It can be concluded that in general the oxide dispersion significantly strengthens the mate- rial. However, the typical form of the flow curve with DRX softening, including a single peak followed by a steady state flow as a plateau, is more recognizable at high temperatures than at low temperatures. This is be- cause at high temperatures the DRX softening compen- sates the WH, and both the peak stress and the onset of steady-state flow are therefore shifted to lower strain levels. The characteristics of softening flow behaviour coupled with DRX have been discussed for six materials and can be summarized as follows: B. MA[EK et al.: BEHAVIOUR OF NEW ODS ALLOYS UNDER SINGLE AND MULTIPLE DEFORMATION 896 Materiali in tehnologije / Materials and technology 50 (2016) 6, 891–898 Figure 12: Stress-strain curves (5 % tension) for: a) holding time 10 s, b) holding time 30 s Slika 12: Krivulja napetost-raztezek ( pri 5 % natezni obremenitvi) za: a) ~as zadr`anja 10 s, b) ~as zadr`anja 30 s 1. Decreasing deformation temperature causes the flow stress level to increase, in other words, it prevents the occurrence of softening due to DRX and dynamic recovery (DRV) and makes the deformed metals exhibit work hardening (WH). 2. For every curve, after a rapid increase in the stress to a peak value, the flow stress decreases monotonically towards a steady-state regime (a steady-state flow as a plateau due to DRX softening is more recognizable at higher temperatures). A varying softening rate typically indicates the onset of DRX, and the stress evolution with strain exhibits three distinct stages. 3. At higher temperatures, a higher DRX softening compensates the WH, and both the peak stress and the onset of steady-state flow are therefore shifted to lower strain levels. 4. The ODS alloy exhibits cyclic softening in most of the tests and its rate decreases with increasing strain. 5. The elastic part of the total strain amplitude is always higher than the plastic one in all the specimens tested, even for the highest total strain amplitudes of 15 %. This is further confirmation of the strong strengthening effect of oxide particles. Acknowledgements This paper includes results created within the projects 14-24252S Preparation and Optimization of Creep Resis- tant Submicron-Structured Composite with Fe-Al Matrix and Al2O3 Particles subsidised by the Czech Science Foundation, and LO1412 Development of West-Bohe- mian Centre of Materials and Metallurgy subsidised by the Ministry of Education, Youth and Sports from spe- cific resources of the state budget of the Czech Republic for research and development. 5 REFERENCES 1 M. Mohan, R. Subramanian, Z. Alam, P. C. Angelo, Evaluation of the Mechanical Properties OF A Hot Isostatically Pressed Yttria- Dispersed Nickel-Based Superalloy, Material Technology, 48 (2014) 6, 899–904 2 W. Quadakkers. Oxidation of ODS alloys. 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YENER et al.: ELECTROMAGNETIC-SHIELDING EFFECTIVENESS AND FRACTURE BEHAVIOR ... 899–902 ELECTROMAGNETIC-SHIELDING EFFECTIVENESS AND FRACTURE BEHAVIOR OF LAMINATED (Ni–NiAl3) COMPOSITES U^INKOVITOST ELEKTROMAGNETNE ZA[^ITE IN OBNA[ANJE PRI LOMU LAMINIRANEGA KOMPOZITA (Ni-NiAl3) Tuba Yener1, Suayb Cagri Yener2, Sakin Zeytýn1 1Sakarya University, Engineering Faculty, Department of Metallurgy and Materials Engineering, Serdivan, Sakarya, Turkey 2Sakarya University, Engineering Faculty, Department of Electrical and Electronic Engineering, Serdivan, Sakarya, Turkey syener@sakarya.edu.tr Prejem rokopisa – received: 2015-07-01; sprejem za objavo – accepted for publication: 2015-12-01 doi:10.17222/mit.2015.189 In this research Ni–NiAl3 multilayer composites were produced through reactive sintering in an open atmosphere using Ni and Al foils with a 250-μm initial thickness. The sintering was performed at 700 °C under 2 MPa of pressure for 6 h. The micro- structure and phase characterizations of the samples were performed. The hardness values of samples were determined using the Vickers indentation technique for the intermetallic and metallic regions as 765±60 HV and 90±10 HV, respectively. For the mechanical examinations, a perpendicular load was applied to the composite in order to observe the fracture behavior of the metallic-intermetallic laminate composites. SEM fracture surface analyses indicated that cracks initiated in the intermetallic region, and the crack propagation stopped when it reaches the ductile nickel phase. In addition, shielding-effectiveness measure- ments were performed. The MIL composite exhibits over 50 dB electromagnetic-shielding effectiveness against a very wide frequency range, from a few GHz to over 18 GHz. Keywords: intermetallics, MIL composites, fracture behavior, electromagnetic interference shielding V raziskavi so bili izdelani Ni–NiAl3 ve~plastni kompoziti z reakcijskim sintranjem na atmosferi in z uporabo Ni- in Al-folij z za~etno debelino 250 μm. Sintranje je bilo 6 h na 700 °C, pri tlaku 2 MPa. Na vzorcih je bila izvedena karakterizacija mikro- strukture in faz. Trdota vzorcev je bila dolo~ena po Vickersu, 765±60 HV, za podro~ja intermetalnih faz in 90±10 HV pri osnovi. Za mehanske preiskave je bila uporabljena navpi~na obremenitev, za opazovanje obna{anja kompozita pri lomljenju kovinskih in intermetalnih lamel. SEM-preiskave prelomov so pokazale, da je za~etek razpoke v podro~ju intermetalne faze in da se {irjenje razpoke ustavi, ko pride v duktilno fazo niklja. Izvedene so bile tudi meritve u~inkovitosti za{~ite sevanja. MIL kompozit ka`e u~inkovitost pred elektromagnetnim sevanjem, vi{jo od 50 dB v zelo {irokem obmo~ju frekvenc od nekaj GHz do preko 18 GHz. Klju~ne besede: intermetalne zlitine, MIL kompoziti, obna{anje pri lomu, elektromagnetna interferen~na za{~ita 1 INTRODUCTION Layered metallic-intermetallic laminate (MIL) com- posites are a new multifunctional materials group based on open air reactive sintering of chemically active metal foils under pressure.1,2 Laminate composites are being intensively studied for a number of potential applica- tions: electronic devices, structural components, armor, etc. Ceramic–ceramic, metal–ceramic, metal–metal, me- tal–ceramic–intermetallic and metal–intermetallic sys- tems have shown desirable properties.3–5 They are designed to optimize the desirable mechanical properties of intermetallics by incorporating layers of ductile rein- forcement.6 The combination of these types of materials makes the MIL composites candidates for the armament industry as armor materials that require improved mechanical and electromagnetic properties.1,6,7 In particular, nickel–tri-nickel aluminide (Ni–NiAl3) metal–intermetallic laminate (MIL) composite systems have a great potential for aerospace, automotive and military applications because of their combination of high strength, toughness and stiffness at a lower density than monolithic titanium or other laminate systems.7,8 Intermetallics of NiAl and NiAl3 have a high melting point, a low density, high strength, good corrosion and oxidation resistance at high temperature.9,10 The nickel- aluminum system is one of the most well known in terms of the formation of intermetallic phases. This system is also a priority among laminate composite systems.2,5,11 The aim of the present study is to synthesize nickel- nickel aluminide metallic-intermetallic composites and analyze their mechanical, fracture and electromagnetic shielding behaviors. The organization of the paper is as follows. After this introduction, in the second section the methodology and production of materials ear sum- marized. In the third section the experimental results are presented. In this section, the fracture behavior in terms of "physical-shielding" and then electromagnetic-shield- ing behavior of the composites are provided after expe- rimental processes. Finally, the paper ends with a con- clusion section. Materiali in tehnologije / Materials and technology 50 (2016) 6, 899–902 899 UDK 67.017:669.018.25:621.8.038 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)899(2016) 2 METHODOLOGY AND EXPERIMENTAL PART 2.1 Materials and method The MIL process consists of stacking commercial- purity Ni and Al foils in alternating layers. The pro- perties of the foils are listed in Table 1. Table 1: Properties of foils used in experiments Tabela 1: Lastnosti folij, uporabljenih pri preizkusih Foil Ni Al Thickness (ìm) 250 250 Purity (%) 99.5 99.5 Stack number 6 5 The nickel- and aluminum-foil dimensions were initially selected to completely consume the aluminum in forming the intermetallic compound with alternating layers of partially unreacted Ni metal. Each foil sheet was prepared as 10 mm × 10 mm and 60 mm × 40 mm rectangular pieces for mechanical and electromagnetic experiments, respectively. Contamination on the surface of the foils was cleaned using ethanol. After drying ra- pidly, they were laminated alternatively into nickel/alu- minum multilayer samples. Each stack consisted of 6 nickel and 5 aluminum foils, as indicated in Table 1. An initial pressure of 2 MPa is applied at room temperature to ensure good contact between the foils. A schematic representation of the Ni-Al stacks is shown in Figure 2a. The sintering process was applied in the open air, in an electrical resistance furnace at 700 °C for 6 h. After sintering, samples were ground and polished using stan- dard metallographic techniques. 2.2 Characterization Microstructure analyses of the composites were per- formed with a JEOL JSM-5600 model scanning electron microscope (SEM). The presence of phases formed in the sintered samples was determined by energy-disper- sive spectroscopy (EDS). The microhardness of compo- sites was determined using a Leica WMHT-Mod model Vickers hardness instrument under an applied load of 300 g for the intermetallic zone, and 100 g for the me- tallic zone. The composition of the phases was deter- mined by comparing the results of the microprobe analysis with the data in the binary Ni–Al phase diagram (Figure 1).12 3 EXPERIMENTAL PROCESSES AND RESULTS 3.1 SEM-EDS Analysis Figure 2b presents the cross-sectional micrographs of representative laminated composites. The presence of T. YENER et al.: ELECTROMAGNETIC-SHIELDING EFFECTIVENESS AND FRACTURE BEHAVIOR ... 900 Materiali in tehnologije / Materials and technology 50 (2016) 6, 899–902 Figure 2: a) Nickel-aluminum foils stack, b) SEM micrograph of la- minated composites produced at 700 °C/6h Slika 2: a) Sestav nikelj-aluminijevih folij, b) SEM-posnetek lami- niranega kompozita, izdelanega pri 700 °C/6 h Figure 1: The Ni–Al binary phase diagram12 Slika 1: Binarni fazni diagram Ni-Al12 Figure 3: SEM-EDS analyses of Ni-NiAl3 composites sintered at 700 °C/6h Slika 3: SEM-EDS analize Ni-NiAl3 kompozita po sintranju 6 h na 700 °C different regions indicates the different phases in the composites. It can be seen that the laminated composites consist of unreacted Ni layers (gray regions) and the formed intermetallic NiAl3 layers (dark regions). Moreover, the laminated composites are well-bonded and remain nearly fully dense. The nickel aluminide phase occurs due to the thermodynamics of the reaction bet- ween Ni and Al. The existence of liquid Al phase plays important roles in the nucleation and growth of NiAl particles and the eventual formation of continuous alter- native intermetallic layers. 3.2 Mechanical fracture behavior and hardness Intermetallics and ceramics, in general, have very little or no dislocation motion, and, hence, exhibit very little inherent or intrinsic crack-propagation resistance.3 By using laminate design and proper composites, it is aimed to produce intermetallic NiAl3 phase during the process to give a high hardness to the composite, while unreacted nickel provided moderate ductility. Due to the deflection of cracks along the Ni/NiAl interfaces, a non-catastrophic fracture was observed in the laminated composites. A weak delamination and de- bonding is seen at the metallic nickel and the inter- metallic layers interface. In a large number of cleavage cracks present in the brittle intermetallic layer. Despite the severe plastic deformation, the nickel layer was not torn. This clearly demonstrates the effect of crack stopping of the ductile reinforcing phases (Figure 4). When it comes to hardness, the values of samples were determined by using the Vickers indentation tech- nique for intermetallic and metallic region as 765±60 HV, 90±10 HV, respectively, whereas the hardness of metallic aluminum and nickel, respectively, is about 45 HV and 90 HV 3.3 Electromagnetic-shielding effectiveness Electromagnetic interference can lead to adverse con- sequences, such as malfunction or crashing of electronic systems and computers, unintentionally firing of electrically explosive devices, or be the cause of the loss of secret information to an enemy. In this respect, it is essential to protect devices from disruptive electromag- netic signals to guarantee their functionality in stable operating conditions. It is also obvious that the electro- magnetic shielding is vital in military applications.13–15 Shielding effectiveness is the ratio of impinging energy to the residual energy. When an electromagnetic wave passes through a shield, absorption and reflection take place. The residual energy is part of the remaining energy that is neither reflected nor absorbed by the shield, but emerges from the shield. Shielding effective- ness (SE) is the ratio of the field before and after T. YENER et al.: ELECTROMAGNETIC-SHIELDING EFFECTIVENESS AND FRACTURE BEHAVIOR ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 899–902 901 Figure 5: Electromagnetic-shielding effectiveness characteristics of laminated Ni–NiAl3 composites a) X Band (8.2–12.4 GHz), b) Ku Band (12.4–18 GHz) Slika 5: Zna~ilnost u~inkovitosti elektromagnetne za{~ite laminira- nega Ni-NiAl3 kompozita a) X-pas (8,2 GHz-12,4 GHz), b) Ku-pas (12,4 GHz – 18 GHz) Figure 4: Cross-sectional micrographs of Ni-NiAl3 composites after impact effect: a) 135× , b) 350× Slika 4: Posnetek preseka kompozita Ni-NiAl3 po udarcu: a) 135×, b) 350× attenuation of the electric and magnetic fields and can be expressed as Equation (1):15,16 [ ]SE dB E E i t = 20 lg (1) Where Ei and Et refer to the transmitted and incident waves, respectively. Shielding effectiveness is a function of frequency, and from the Equation (1) it is measured in dB. The shielding-effectiveness characteristics of lami- nated Ni–NiAl3 composites have been measured and the results obtained are shown in Figures 5a and 5b for the X band and Ku band, respectively. From the results, the produced MIL composites exhibit around or more 50 dB of electromagnetic- shielding effectiveness against a very wide frequency range from 8.2 GHz to over 18 GHz. That shielding level means even 99.999 % of the incident power is prevented by the produced composites. These shielding-effective- ness levels indicate that laminated composites can be remarkable candidates for shielding application also thanks to their improved mechanical properties. 4 CONCLUSIONS The conclusions of this research can be summarized as follows: • By controlling the duration of the reactive-foil sint- ering process, composites can be fabricated in which a tailored amount of residual aluminum remains at the intermetallic centerline. • Ni–NiAl3 metal–intermetallic laminate (MIL) com- posites have been successfully synthesized by reactive-foil sintering technique in open air at 700 °C for 6 h under 2 MPa pressure. The laminated struc- ture is well-bonded, nearly fully dense. • Microstructural characterization by SEM and EDS indicates that NiAl, NiAl3, Ni2Al3 are intermetallic phases in the composite. • The hardness of the fabricated laminated composite was dramatically changed. Whereas the hardness of metallic aluminum and nickel, respectively, is about 45 HV and 90 HV, the hardness of intermetallic zone is approximately 765±60 HV. • In this study the shielding effectiveness of laminated Ni–NiAl3 composites was examined in a two-fre- quency band at GHz levels and the results obtained are shown. Around 50-dB shielding-effectiveness le- vels were reached experimentally from the measure- ments. • Thus, experimental results obtained are promising for MIL composites to be appropriate candidate mate- rials for military applications with their electromag- netic as well as mechanical properties. 5 REFERENCES 1 K. S. Vecchio, Synthetic multifunctional metallic-intermetallic lami- nate composites, JOM, 57 (2005) 3, 25–31, doi:10.1007/s11837- 005-0229-4 2 B. Besen, M. Kalayci, T. Yener, S. Zeytin, Some Properties Of Ni-AlNi Metallic-Intermetallic Laminate Material, Journal of Inter- national Scientific Publications: Materials, Methods & Technologies, 7 (2013) 2, 390–396 3 R. R. Adharapurapu, K. S. Vecchio, F. Jiang, A. Rohatgi, Fracture of Ti-Al3Ti metal-intermetallic laminate composites: Effects of lamina- tion on resistance-curve behavior, Metallurgical and Materials Tran- sactions A, 36 (2005) 11, 3217–3236, doi:10.1007/s11661-005- 0092-5 4 L. Peng, H. Li, J. Wang, Processing and mechanical behavior of laminated titanium–titanium tri-aluminide (Ti–Al3Ti) composites, Materials Science and Engineering A, 406 (2005) 1, 309–318, doi:10.1016/j.msea.2005.06.067 5 X. Yang, X. Peng, F. Wang, Size effect of Al particles on the struc- ture and oxidation of Ni/Ni3Al composites transformed from electro- deposited Ni–Al films, Scripta Materialia, 56 (2007) 6, 509–512, doi:10.1016/j.scriptamat.2006.11.016 6 A. Rohatgi, D. J. Harach, K. S. Vecchio, K. P. Harvey, Resistance- curve and fracture behavior of Ti–Al3Ti metallic–intermetallic lami- nate (MIL) composites, Acta Materialia, 51 (2003) 10, 2933–2957, doi:10.1016/S1359-6454(03)00108-3 7 Y. Cao, C. Guo, S. Zhu, N. Wei, R. A. Javed, F. Jiang, Fracture beha- vior of Ti/Al3Ti metal-intermetallic laminate (MIL) composite under dynamic loading, Materials Science and Engineering A, 637 (2015), 235–242, doi:10.1016/j.msea.2015.04.025 8 K. H. Zuo, D. L. Jiang, Q. L. Lin, Mechanical properties of Al2O3/Ni laminated composites, Materials Letters, 60 (2006) 9–10, 1265–1268, doi:10.1016/j.matlet.2005.11.010 9 L. Z. Zhang, D. N. Wang, B. Y. Wang, R. S. Yu, L. Wei, Identifi- cation of lattice vacancies in the B2-phase region of Ni–Al system by positron annihilation, Journal of Alloys and Compounds, 457 (2008) 1–2, 47–50, doi:10.1016/j.jallcom.2007.03.065 10 F. L. Zhang, Z. F. Yang, Y. M. Zhou, C. Y. Wang, H. P. Huang, Fabri- cation of grinding tool material by the SHS of Ni–Al/diamond/dilute, International Journal of Refractory Metals and Hard Materials, 29 (2011) 3, 344–350, doi:10.1016/j.ijrmhm.2010.12.013 11 C. T. Wei, V. F. Nesterenko, T. P. Weihs, B. A. Remington, H. S. Park, M. A. Meyers, Response of Ni/Al laminates to laser-driven compression, Acta Materialia, 60 (2012) 9, 3929–3942, doi:10.1016/j.actamat.2012.03.028 12 ASM Handbook, Vol. 3: Alloy Phase Diagrams, ASM International, 2001 13 P. Saini, M. Aror, Microwave Absorption and EMI Shielding Beha- vior of Nanocomposites Based on Intrinsically Conducting Poly- mers, Graphene and Carbon Nanotubes, Chapter 3, In: A. De Souza Gomes (Ed.), New Polymers for Special Applications, InTech, 2012, doi:10.5772/48779 14 C. R. Paul, Introduction to electromagnetic compatibility, 2nd Edition, John Wiley & Sons, 2006, 184 15 S. Geetha, K. K. Satheesh Kumar, C. R. K. Rao, M. Vijayan, D. C. Trivedi, EMI shielding: Methods and materials-A review, Journal of Applied Polymer Science, 112 (2009) 4, 2073–2086, doi:10.1002/app.29812 16 O. Cerezci, S. ªeker, ª. Yener, B. Kanberoðlu, M. H. Niºancý, Ev, Ofislerde GSM Frekanslý Radyasyondan Bireysel Korunma, EMANET, Yýldýz Teknik Üniversitesi, Beºiktaº, Ýstanbul 2013, 372–376 T. YENER et al.: ELECTROMAGNETIC-SHIELDING EFFECTIVENESS AND FRACTURE BEHAVIOR ... 902 Materiali in tehnologije / Materials and technology 50 (2016) 6, 899–902 P. SLÁMA, J. NACHÁZEL: EFFECT OF THERMOMECHANICAL TREATMENT ON THE INTERGRANULAR CORROSION ... 903–910 EFFECT OF THERMOMECHANICAL TREATMENT ON THE INTERGRANULAR CORROSION OF Al-Mg-Si-Type ALLOY BARS VPLIV TERMOMEHANSKE PREDELAVE NA INTERKRISTALNO KOROZIJO PALIC IZ ZLITIN Al-Mg-Si Peter Sláma, Jan Nacházel COMTES FHT a.s., Prùmyslová 995, 33441 Dobøany, Czech Republic peter.slama@comtesfht.cz Prejem rokopisa – received: 2015-07-01; sprejem za objavo – accepted for publication: 2015-02-12 doi:10.17222/mit.2015.170 Al-Mg-Si-type alloys (6xxx-series alloys) exhibit good mechanical properties, formability, weldability and good corrosion resistance in a variety of environments. They often find use in the automotive industry and other applications. Some alloys, however, particularly those with higher copper levels, show increased susceptibility to intergranular corrosion. Intergranular corrosion (IGC) is typically related to the formation of microgalvanic cells between cathodic, more-noble phases and depleted (precipitate-free) zones along the grain boundaries. It is encountered mainly in Al-Mg-Si alloys containing Cu, where it is thought to be related to the formation Q-phase precipitates (Al4Mg8Si7Cu2) along the grain boundaries. The present paper des- cribes the effects of mechanical working (pressing, drawing and straightening) and artificial ageing on intergranular corrosion in a bar of the 6064 alloy. The resistance to intergranular corrosion was mapped using corrosion tests according to EN ISO 11846, method B. The corrosion tests showed that with continuing ageing and over-ageing, deep IGC changes into pitting corrosion with a smaller depth of attack. However, the corrosion resistance of the bars is impaired by post-quench mechanical working (drawing and straightening). Keywords: Al-Mg-Si-Cu alloy, 6064 alloy, extruded bars, thermomechanical treatment, intergranular corrosion, pitting corrosion Zlitine vrste Al-Mg-Si (6xxx-vrsta zlitin) ka`ejo dobre mehanske lastnosti: preoblikovalnost, varivost in dobro korozijsko odpornost v razli~nih okoljih. Pogosto se uporabljajo v avtomobilski industriji in tudi v druge namene. Vendar pa nekatere zlitine, posebno tiste z vi{jo vsebnostjo bakra, ka`ejo pove~ano ob~utljivost na interkristalno korozijo. Interkristalna korozija (IGC) je zna~ilno povezana z nastankom mikrogalvanskih celic med katodno, bolj plemenito fazo in osiroma{enim (brez izlo~kov) podro~jem, vzdol` meja kristalnih zrn. To se pojavlja predvsem v AlMgSi zlitinah, ki vsebujejo Cu in kjer se pred- postavlja, da je to povezano z nastankom izlo~kov Q-faze (Al4Mg8Si7Cu2), vzdol` meja med zrni. ^lanek opisuje vpliv mehanskega preoblikovanja (stiskanje, vle~enje, ravnanje) in vpliv umetnega staranja na interkristalno korozijo palic iz zlitine 6064. Odpornost na interkristalno korozijo je bila preslikana s pomo~jo korozijskih preizkusov, skladno s standardom EN ISO 11846, metoda B. Korozijski preizkusi so pokazali da se z nadaljevanjem staranja in prestaranjem globoke interkristalne korozije, spremenijo v jami~asto korozijo, z manj{o globino napada. Vseeno pa je korozijska odpornost palic poslab{ana z mehansko predelavo (vle~enje in ravnanje) po ga{enju. Klju~ne besede: zlitina Al-Mg-Si-Cu, zlitina 6064, iztiskane palice, termomehanska predelava, interkristalna korozija, jami~asta korozija 1 INTRODUCTION Al-Mg-Si-type alloys (6xxx-series alloys) exhibit good mechanical properties, formability, weldability and good corrosion resistance in a variety environments. They frequently find use in automotive, aviation and other applications.1,2 Some of these materials are alloyed with copper to improve their strength. In these alloys, particularly higher-copper alloys, increased suscepti- bility to intergranular corrosion (IGC) can be observed, most notably in the unaged condition and less often in the T6 temper condition. The effects of Cu as well as the opportunities for enhancing the resistance to inter- granular corrosion have received considerable attention in a number of studies.3–11 Intergranular corrosion (IGC) is typically related to the formation of microgalvanic cells between the cathodic more-noble phases and the depleted (precipitate-free) zones along the grain boun- daries. It is encountered mainly in AlMgSi alloys that contain Cu, where it is thought to be linked to the for- mation of cathodic Q-phase (Al4Mg8Si7Cu2) along the grain boundaries. The occurrence of phases along the grain boundaries was observed using scanning-trans- mission electron microscopy (STEM). The impact of Cu additions and heat treatment on IGC was described in several papers.3–6 The alloys contained 0.5-0.6 % Mg, 0.6-0.8 % Si, 0.2 % Fe, 0.2 % Mn and Cu at 0.02 through 0.7 % of mass fractions. The occurrence of IGC was monitored in 2.5 mm × 78 mm extruded flat bars. The effects of the cooling rate from the extrusion temperature were studied3, as were the effects of artificial ageing.4,5 Corrosion tests were carried out according to EN ISO 11846, method B. Corrosion was only monitored on the surface of the extruded parts. EN ISO 11846 specifies that the corrosion is monitored on the long side of the specimen. In an alloy with a Cu level of 0.02 %, no IGC was found. In an alloy with Materiali in tehnologije / Materials and technology 50 (2016) 6, 903–910 903 UDK 62-971:621.78:620.196.2:669.017.13 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)903(2016) 0.2 % Cu, IGC occurred depending on the artificial ageing time, and changed into pitting corrosion. These findings suggest that between the occurrence of IGC and pitting corrosion, there is a region in which no IGC occurs (Figure 1). Hence, over-ageing (by increasing either temperature or time) permits a transition from a region with IGC to a region with suppressed IGC. The AA6056 material for the aviation industry is used in overaged condition. It is supplied in the T78 state with an enhanced resistance to IGC. According to6,7, the T78 temper is achieved by two-stage ageing: 175 °C/6 h + 210 °C/5 h. Two-stage ageing was explored by the authors of the study.11 The alloy had a nominal composition of 1.0 % Mg, 1.2 % Si, 0.3 % Cu, 0.6 % Mn, 0.12 % Cr, 0.12 % Fe and a balance of Al. An ageing schedule specified as 180 °C/2 h + 160 °C/120 h led to better results than 175 °C/6 h + 210 °C/5 h. However, this work was carried out using specimens of rolled sheet with a 2-mm thick- ness, where the corrosion attack was monitored on the sheet surface and not on its cross-section. Thermomechanical treatment generally has a great influence on the corrosion in other types of aluminium alloys.12 In this research the effect of the thermomecha- nical treatment (extrusion, drawing and ageing) on the intergranular corrosion in bars from EN AW-6064A (AlMg1SiBi) machineable alloy was studied. AlMgSi-type machinable alloys are used in the automo- tive industry. Their improved machinability is imparted by alloying with Pb (6012 alloy) or with Bi+Pb (6262 and 6064 alloys). These alloys have higher alloy levels and contain more phases than the alloys studied in 3–11. These phases include Bi and Pb cathodic particles. 2 EXPERIMENTAL PART The chemical composition of the EN AW-6064A bars is shown in Table 1. The bars of 17 mm diameter were made by an industrial hot-extrusion process using a mul- tiple-hole die. The process temperature was 540–546 °C. Right after extrusion, the bars were water-wave cooled (T1 condition). The quenched bars were then drawn to the final diameter of 15 mm at 22 % reduction and straightened in a Schumag straightening machine (T2 temper). The final operation was artificial ageing to T8. Bars in conditions corresponding to each process step were gathered for testing. The samples are listed in Table 2. The bars that did not undergo ageing (HA1, HB2 and HF) were used in artificial ageing trials: single-stage and two-stage ageing to the under-aged, peak-aged and over-aged condition. The artificial ageing schedules are presented in Table 3. Table 1: Chemical composition of the alloy 6064A, in mass fractions (w/%) Tabela 1: Kemijska sestava zlitine 6064A, v masnih dele`ih (w/%) Sample Si Fe Cu Mn Mg Cr Pb Bi H 0.60 0.23 0.27 0.04 1.03 0.05 0.28 0.49 Table 2: Samples description Tabela 2: Opis vzorcev Sample Diameter Temper Description of thermomechanicalprocessing HA1 17 mm T1 Extruding, quenching HB2 15 mm T2 Extruding, quenching, drawing HF 15 mm T2 Extruding, quenching, drawing,straightening HC 15 mm T8 Extruding, quenching, drawing,straightening, ageing Table 3: Heat treatment HT (artificial ageing) for samples HA1, HB2, HF Tabela 3: Toplotna obdelava (umetno staranje) vzorcev HA1, HB2, HF HT One-stage HT-A Two-stage A HT-B Two-stage B 1 160 °C/4 h 1A 160 °C/4 h+220 °C/4 h 1B 160 °C/4 h+ 205 °C/4 h 2 160 °C/8 h 2A 160 °C/8 h+220 °C/4 h 2B 160 °C/8 h + 205 °C/4 h 3 180 °C/4 h 3A 180 °C/4 h+220 °C/4 h 3B 180 °C/4 h+ 205 °C/4 h 4 180 °C/8 h The progress of ageing was monitored by a HV5 hardness measurement using a DURASCAN 50 hardness tester. Tests of resistance to intergranular corrosion were conducted in accordance the EN ISO 11846 standard, method B.13 For these tests, specimens of 2 cm in length were made from the bars. Their cut surfaces were ground with P-1200 grinding papers. The original surface of the bar was not altered. Before testing, the specimens were degreased in acetone. In accordance with the standard requirements, they were etched with 5 % NaOH solution at 55 °C for 2 min. After a water rinse, they were placed in concentrated nitric acid for cleaning. The test itself involved submerging in a test solution for 24 h at room temperature. The solution was 30 g NaCl/L solution + 10 mL concentrated hydrochloric acid. P. SLÁMA, J. NACHÁZEL: EFFECT OF THERMOMECHANICAL TREATMENT ON THE INTERGRANULAR CORROSION ... 904 Materiali in tehnologije / Materials and technology 50 (2016) 6, 903–910 Figure 1: The dominant corrosion types in a material aged at 185 °C, according to5 Slika 1: Prevladujo~e oblike korozije v materialu, staranem na 185 °C po viru5 Following the test, the specimens were rinsed with water. Metallographic sections were prepared on longitu- dinal cross-sections through the specimens. The corrosion attacks on the bar surface as well as on the transverse cut surface were examined. The maximum corrosion depth was determined and documented using light microscopy. The surfaces of the specimens after corrosion testing were examined in a JEOL JSM 6380 scanning electron microscope. 3 RESULTS 3.1 Initial microstructures The microstructure of T8-temper HC bars upon drawing, straightening and ageing is shown in Figure 2a. A micrograph of the phases is in Figure 2b. The micro- structure is fully recrystallized. The grains in the surface layer are relatively fine, with a size of 70 μm. In the centre, the grains are coarser, of the order of several hundred μm. Different grain sizes in the surface and in the interior are a typical occurrence in extruded bars from Al alloys. Typically, the surface layer contains coarse grains and the interior remains unrecrystallized.1,2 The phases in the microstructure are banded and aligned in the extrusion/drawing direction. Large elon- gated particles consist of Bi or Bi+Pb. The small ones are alpha-Al15(Fe,Mn,Cu,Cr)3Si2 particles. Other small particles are Mg2Si particles. The Bi, Pb and alpha- Al15(Fe,Mn,Cu,Cr)3Si2 particles are more noble, cathodic. The Mg2Si particles are anodic. With cathodic particles, the matrix of the aluminium solid solution is etched away preferentially when placed in a corrosion environment. With anodic particles, it is the particles that are attacked. The microstructure may also contain cathodic Q-phase particles (Al4Mg8Si7Cu2). Figure 2b also shows minute particles along grain boundaries. EDS analysis revealed that they contain higher amounts of copper, which suggests that they are Q-phase particles. 3.2 Corrosion tests Specimens to be tested according to EN ISO 11846, method B, are to be alkaline pre-etched with 5-10 % NaOH solution. With this etch, the Al matrix and anodic phases are attacked. The etched surface of a specimen is shown in Figure 3. The large pits are the result of the Al matrix being etched away from around the Bi, Pb and alpha-Al15(Fe,Mn,Cu,Cr)3Si2 cathodic phases. The small pits are the locations of Mg2Si anodic particles that were P. SLÁMA, J. NACHÁZEL: EFFECT OF THERMOMECHANICAL TREATMENT ON THE INTERGRANULAR CORROSION ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 903–910 905 Figure 3: SEM micrographs of the surface of HC sample upon etching with NaOH: a) sample surface, b) transverse cut surface Slika 3: SEM-posnetka povr{ine vzorca HC, po jedkanju z NaOH: a) povr{ina vzorca, b) pre~ni prerez vzorca Figure 2: Micrographs of grains and phases in HC samples upon drawing and ageing: a) electrolytically etched with Barker’s reagent, polarised light, b) etched with Dix-Keller’s reagent Slika 2: Posnetka zrn in faz v HC vzorcu, po vle~enju in staranju: a) elektrolitsko jedkano z Baker jedkalom, polarizirana svetloba, b) jedkano z Dix-Keller jedkalom etched away. The grain boundaries were slightly attacked. In order to evaluate the corrosion, the specimens were cut longitudinally after the test. On the cross-section, the type and depth of the corrosion on the bar’s surface and on its transverse cut surface were examined. 3.2.1 Corrosion tests of materials in initial condition The initial condition evaluation was carried out on HF samples supplied in the T2 (non-aged) condition and on the HC samples supplied in the T8 (peak-aged) condition. The HC bars were drawn and aged during the 24 h following quenching. The surface corrosion is shown in Figure 4. Its evaluation is detailed in Table 4. Table 4: Evaluation of corrosion and hardness of initial samples in T2 and T8 condition Tabela 4: Ocena korozije in trdota za~etnih vzorcev po T2 in T8 obdelavi Sample Temper HV5 Place Corrosiondepth (μm) Corrosion type HF T2 108 Surface 420.5 IGC + pitting Transverse cut 493.2 IGC + pitting HC T8 124.3 Surface 217.8 Pitting,transgranular Transverse cut 607.7 Pitting The surface of the non-aged HF sample shows extensive intergranular corrosion (IGC) with a depth of more than 420 μm. In the artificially-aged HC sample (T8 peak-aged temper), the corrosion changed into the pitting type, which spreads perpendicularly to the surface to a depth of more than 200 μm. The corrosion type corresponds to transgranular corrosion. On the cross-section through the HF specimen, IGC with a depth of approximatelz 500 μm was found as well. The corrosion on the transverse cut surface of the HC sample is very extensive too. It is, however, pitting-type corrosion, which reached a depth of up to 600 μm. It follows the bands of coarse cathodic Bi, Pb and alpha-Al15(Fe,Mn,Cu,Cr)3Si2 particles (Figure 4d). Table 4 lists HV5 hardness values. The HF sample in the T2 state exhibits 108 HV5. Age-hardening to T8 increased the hardness to 124 HV5. 3.2.2 Corrosion tests after experimental heat treatment (artificial ageing) Using these tests, the impact of various artificial ageing schedules (under-ageing, over-ageing) on the corrosion in bars in various conditions was monitored: • Sample HA1 – after extruding and quenching; P. SLÁMA, J. NACHÁZEL: EFFECT OF THERMOMECHANICAL TREATMENT ON THE INTERGRANULAR CORROSION ... 906 Materiali in tehnologije / Materials and technology 50 (2016) 6, 903–910 Figure 4: Corrosion attack in as-received bars: a) HF surface – temper T2, non-aged, b) HC surface – temper T8, peak-aged, c) HF transverse cut surface, d) HC transverse cut surface Slika 4: Korozija na dobavljenih palicah: a) HF povr{ina - `arjenje T2, nestarano, b) HC povr{ina – `arjenje T8, starano, c) HF pre~ni presek, d) HC pre~ni presek Figure 6: Corrosion on the HA1 transverse cut surface upon ageing: a) 160 °C/8 h under-aged, b) 180 °C/8 h, c) 160 °C/4 h + 205 °C/4 h, d) 160 °C/4 h + 220 °C/4 h overaged Slika 6: Korozija na pre~nem prerezu HA1 po staranju: a) 160 °C/8 h podstarano, b) 180 °C/8 h, c) 160 °C/4 h + 205 °C/4 h, d) 160 °C/4 h + 220 °C/4 h prestarano Figure 5: Corrosion on the HA1 bar surface upon ageing: a) 160 °C/8 h under-aged, b) 180 °C/8 h, c) 160 °C/4 h + 205 °C/4 h, d) 160 °C/4 h + 220 °C/4 h overaged Slika 5: Korozija na povr{ini HA1 palice, po staranju: a) 160 °C/8 h, podstarano, b) 180 °C/8 h, c) 160 °C/4 h – 205 °C/4 h, d) 160 °C/4 h + 220 °C/4 h, prestarano • Sample HB2 – after extruding, quenching and drawing; • Sample HF – after extruding, quenching, drawing and straightening • Corrosion tests of specimens of HA1 extruded bars The surface corrosion of selected specimens in variously aged conditions is illustrated in Figure 5. The corrosion of the transverse cut surface is shown in Figure 6. Table 5 contains the results of the corrosion evaluation and the HV5 hardness levels, which indicate the progress of ageing. In specimens in the under-aged condition, the most extensive surface corrosion was found, involving continuous IGC with a maximum depth of more than 300 μm. In the peak-aged condition, the depth of attack decreased and IGC ceased to be conti- nuous. In the over-aged condition, only sporadic pitting corrosion can be observed with a depth of about 120 μm. Table 5: Evaluation of corrosion and hardness of samples HA1 Tabela 5: Ocena korozije in trdota vzorcev HA1 Sample HV5 US Place Corrosion depth (μm) Corrosion type HA1-2 92.5 160 °C/ 8h Surface 309.3 IGC + pittingsporadic Under-ageing Transversecut 421.1 IGC near-edge + pitting HA1-4 113.7 180 °C/ 8 h Surface 296.4 IGC 50 % peak ageing Transversecut 460.8 IGC near-edge + pitting HA1-1B 114.7 160 °C/4 h + 205 °C/4 h Surface 158.9 IGC + pitting sporadic peak ageing Transversecut 381.5 IGC + pitting HA1-1A 109.3 160 °C/4 h + 220 °C/4 h Surface 120.4 Pitting sporadic Over-ageing Transversecut 93.4 Pitting sporadic The same type of corrosion was found on the trans- verse cut surface. However, the corrosion depth was larger there: more than 400 μm. The only exception was the over-aged sample where the depth was less than 100 μm. 3.2.3 Corrosion tests of specimens of HB2 drawn bars The surface corrosion of selected specimens in variously aged conditions is illustrated in Figure 7. The corrosion of the transverse cut surface is shown in Figure 8. Results of the evaluation of corrosion are given in Table 6. Table 6: Evaluation of corrosion and hardness of samples HB2 Tabela 6: Ocena korozije in trdota vzorcev HB2 Sample HV5 US Place Corrosion depth (μm) Corrosion type HB2-3 120 180 °C/4 h Surface 382.5 IGC 60 % +pitting sporadic Under-ageing Transversecut 528.4 IGC, near-edge HB2-4 120.7 180 °C/8 h Surface 123.6 Pitting, sporadic peak ageing Transversecut 755.6 Pitting, near-edge HB2-3B 123.3 180 °C/4 h + 205 °C/4 h Surface 81.8 Pitting peak ageing Transversecut 742.6 Pitting HB2-3A 106.7 180 °C/4 h + 220 °C/4 h Surface 88.3 Pitting Over-ageing Transversecut 678.1 Pitting, near-edge In HB2 drawn bars, IGC was found only in the under-aged condition (Figure 7a). This intergranular P. SLÁMA, J. NACHÁZEL: EFFECT OF THERMOMECHANICAL TREATMENT ON THE INTERGRANULAR CORROSION ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 903–910 907 Figure 7: Corrosion on the HB2 transverse cut surface upon ageing: a) 180 °C/4 h under-aged, b) 180 °C/8 h, c) 180 °C/4 h + 205 °C/4 h, d) 180 °C/4 h + 220 °C/4 h overaged Slika 7: Korozija na pre~nem prerezu HB2 po staranju: a) 180 °C/4 h podstarano, b) 180 °C/8 h, c) 180 °C/4 h + 205 °C/4 h, d) 180 °C/4 h + 220 °C/4 h prestarano Figure 8: Corrosion on the HB2 transverse cut surface upon ageing: a) 180 °C/4 h under-aged, b) 180 °C/8 h, c) 180 °C/4 h + 205 °C/4 h, d) 180 °C/4 h + 220 °C/4 h overaged Slika 8: Korozija na pre~nem prerezu HB2 po staranju: a) 180 °C/4 h podstarano, b) 180 °C/8 h, c) 180 °C/4 h + 205 °C/4 h, d) 180 °C/4 h + 220 °C/4 h prestarano corrosion is not continuous. On the surface of the bar, the corrosion depth is approx. 400 μm. On the transverse cut surface, IGC is more frequent in the fine-grained surface layer. The depth of attack exceeds 500 μm (Figure 8). In the peak-aged and overaged conditions, the bar’s surface only exhibits pitting corrosion with a depth of about 100 μm. Besides that, corrosion spreads parallel to and beneath the surface, along the bands of coarse cathodic phases. The authors in10 describe this type of corrosion as ELA (Exfoliation-Like Attack). On the transverse cut surface, corrosion is of the pitting type as well. It is much deeper and, again, more frequent in the near-sur- face areas. 3.2.4 Corrosion tests of specimens of HF drawn and straightened bars Surface corrosion of selected specimens in variously aged conditions is illustrated in Figure 9. The corrosion of the transverse cut surface is shown in Figure 10. Re- sults of the evaluation of corrosion are given in Table 7. Table 7: Evaluation of corrosion and hardness of samples HF Tabela 7: Ocena korozije in trdota vzorcev HF Sample HV5 US Place Corrosiondepth (μm) Corrosion type HF-3 123 180 °C/4 h Surface 364.9 IGC Under-ageing Transversecut 545.9 IGC, near-edge HF-4 123.3 180 °C/8h Surface 307.1 Pitting peak ageing Transversecut 93.6 Pitting, sporadic HF-3B 115.4 180 °C/4 h + 205 °C/4 h Surface 348.9 Pitting, transgranula r Over-ageing Transversecut 471.2 Pitting HF-3A 110.3 180 °C/4 h + 220 °C/4 h Surface 366.1 Pitting, transgranula r Over-ageing Transversecut 122.3 Pitting, sporadic In specimens in underaged condition, there is deep IGC on the bar’s surface, as well as on the transverse cut surface. In the peak-aged and over-aged conditions, the bar surface only exhibits pitting corrosion that spreads perpendicularly to the surface to a depth of more than 300 μm. It is transgranular corrosion, as it penetrates the grains. On the transverse cut surfaces, the least extensive corrosion was found in the peak-aged condition (Fig- ure 10b). In the slightly-overaged condition, the corro- sion is extensive and deep (Figure 10c). In increasingly overaged specimens, the number and depth of corrosion attack locations decrease (Figure 10d). 4 DISCUSSION The main mechanism of IGC is reported to be the formation of micro-galvanic cells between cathodic more-noble phases and the depleted (precipitate-free) zones along the grain boundaries. In this case, the key cathodic phase is the Q-phase (Al4Mg8Si7Cu2), which precipitates along the grain boundaries. As a result, the grain-boundary areas become depleted of Cu and other elements. In addition, a thin Cu film forms along the grain boundaries and plays the key role in IGC growth and propagation.3–6 The entire precipitation process is thermally activated and depends on the diffusion of alloying elements. Its rate is described by an Arrhenius equation. With increasing ageing temperature and time, P. SLÁMA, J. NACHÁZEL: EFFECT OF THERMOMECHANICAL TREATMENT ON THE INTERGRANULAR CORROSION ... 908 Materiali in tehnologije / Materials and technology 50 (2016) 6, 903–910 Figure 10: Corrosion on the HF transverse cut surface upon ageing: a) 180 °C/4 h under-aged, b) 180 °C/8 h, c) 180 °C/4 h + 205 °C/4 h, d) 180 °C/4 h + 220 °C/4 h overaged Slika 10: Korozija na pre~nem prerezu HF, po staranju: a) 180 °C/4 h podstarano, b) 180 °C/8 h, c) 180 °C/4 h + 205 °C/4 h, d) 180 °C/4 h + 220 °C/4 h prestarano Figure 9: Corrosion on the HF bar surface upon ageing: a) 180 °C/4 h under-aged, b) 180 °C/8 h, c) 180 °C/4 h + 205 °C/4 h, d) 180 °C/4 h + 220 °C/4 h overaged Slika 9: Korozija na povr{ini HF palice, po staranju: a) 180 °C/4 h podstarano, b) 180 °C/8 h, c) 180 °C/4 h + 205 °C/4 h, d) 180 °C/4 h + 220 °C/4 h prestarano the Q-phase precipitates coarsen and the volume fraction of the Cu film along the grain boundaries decreases. Consequently, the susceptibility to IGC is reduced and the material typically exhibits only pitting corrosion. The EN AW-6064 alloy contains a number of other primary cathodic phases (Bi, Pb, alpha-Al15(Fe,Mn,Cu, Cr)3Si2). Their arrangement in bands with short distances between the phases helps the pitting corrosion to propagate to larger depths, most notably beneath the transverse cut surface (Figure 4d). In some cases there were great differences between the corrosion attack on the bar’s surface and on the transverse cut surface. In the extruded bars (HA1), it was found that with increasing over-ageing the large-depth IGC changes into shallower pitting corrosion, which is in agreement with the findings presented in 3–6. In the overaged condition, the corrosion penetrations on the transverse cut surface were smaller. Sporadic pitting corrosion with a depth of about 100 μm was found. In the drawn bars (HB2), the transition from IGC to shallower pitting corrosion was observed as well. Unlike the specimens from bars that had not been drawn, all the specimens in this group showed very deep corrosion (more than 500 μm) on their transverse cut surfaces (Fig- ure 8). In the drawn and straightened bars (HF), another type of corrosion was observed. In the under-aged bars, IGC was found on both the bar surface and the transverse cut surface. With ongoing ageing, IGC changes into pitting corrosion, which – on the bar surface – propagates per- pendicularly to the surface and by transgranular mecha- nism to a larger depth than the pitting corrosion in the drawn bars (Figures 9b to 9d). This corrosion type corresponds to transgranular stress corrosion cracking (SCC).14 The difference can be attributed to the variation between the internal stresses induced by drawing and straightening. Drawing typically induces tensile stress. Straightening, however, involves alternating bending loads and tensile and compressive stresses, which lead to non-uniform residual stress that promotes corrosion propagation, perpendicularly to the surface and to a larger depth. The transverse cut surface, unlike HB2 specimens, shows – in some cases – shallow sporadic pitting corrosion (Figures 10b and 10d). 5 CONCLUSION Extruded and drawn bars from the EN AW-6064A alloy were used for exploring the impact of thermo- mechanical treatment on intergranular corrosion (IGC). The effects of forming (drawing and straightening) and artificial ageing were mapped, along with the type of corrosion and corrosion depth on the bar surface and its transverse cross-section. The corrosion tests were carried out in accordance with EN ISO 11486 – method B. The results of the corrosion tests show that the ther- momechanical treatment affects both the type and depth of corrosion. The bar surface exhibited three types of corrosion: • IGC in under-aged specimens: typically extensive corrosion with a depth of more than 300 μm. • Pitting corrosion in more aged and over-aged extruded/drawn bars, where the corrosion depth was approximately 100 μm. • Transgranular pitting corrosion in more aged and over-aged bars that had undergone final straightening. Here, the corrosion depth was larger and exceeded 300 μm. With more intensive ageing and over-ageing (tempe- rature, time), IGC changed into pitting corrosion in extruded/drawn bars. There was an adverse impact of the post-drawing straightening operation on the resistance to surface corrosion in the bars, evidenced by deep trans- granular pitting corrosion. In most cases the transverse cross-sections exhibited very deep pitting corrosion with depths up to 800 μm, which followed the bands of coarse cathodic phases. Exceptions were found in severely over-aged bars (extruded or extruded and straightened), which showed sporadic pitting corrosion with depths of approximately 100 μm. Acknowledgements This paper was created by project Development of West-Bohemian Centre of Materials and Metallurgy No.: LO1412, financed by the MEYS of the Czech Republic. 6 REFERENCES 1 D. G. Altenpohl, Aluminum: Technology, Applications, and Environ- ment: A Profile of a Modern Metal, 6th ed., Minerals, Metals, and Materials Society, Warrendale, Pennsylvania 1998 2 J. E. Hatch (Ed.), Aluminium – Properties and Physical Metallurgy, ASM, Ohio 1984 3 G. Svenningsen, J. E. Lein, A. Bjorgum, J. H. Nordlien, Y. D. Yu, K. Nisancioglu, Effect of low copper content and heat treatment on intergranular corrosion of model AlMgSi alloys, Corrosion Science, 48 (2006) 1, 226–242, doi:10.1016/j.corsci.2004.11.025 4 G. Svenningsen, M. H. Larsen, J. H. Nordlien, K. Nisancioglu, Effect of high temperature heat treatment on intergranular corrosion of AlMgSi(Cu) model alloy, Corrosion Science, 48 (2006) 1, 258–272, doi:10.1016/j.corsci.2004.12.003 5 G. Svenningsen, M. H. Larsen, J. C. Walmsley, J. H. Nordlien, K. Nisancioglu, Effect of artificial aging on intergranular corrosion of extruded AlMgSi alloy with small Cu content, Corrosion Science, 48 (2006) 6, 1528–1543, doi:10.1016/j.corsci.2005.05.045 6 M. H. Larsen, J. C.Walmsley, O. Lunder, R. H. Mathiesen, K. Nisan- cioglu, Intergranular Corrosion of Copper-Containing AA6xxx AlMgSi Aluminum, J. Electrochem. Soc., 155 (2008) 11, C550–C556, doi:10.1149/1.2976774 7 T. Koval~ík, J. Stoulil, P. Sláma, D. Vojtìch, The Influence of Heat Treatment on Mechanical and Corrosion Properties of Wrought Alu- minium Alloys 2024 and 6064, Manufacturing Technology, 15 (2015) 1, 54–61 8 V. Guillaumin, G. Mankowski, Influence of Overaging Treatment on Localized Corrosion of Al 6056, Corrosion, 56 (2000), 12–23, doi:10.5006/1.3280517 P. SLÁMA, J. NACHÁZEL: EFFECT OF THERMOMECHANICAL TREATMENT ON THE INTERGRANULAR CORROSION ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 903–910 909 9 C. Gallais, A. Denquin, Y. Brechet, G. Lapasset, Precipitation micro- structures in an AA6056 aluminium alloy after friction stir welding: Characterisation and modelling, Mater. Sci. Eng. A, 496 (2008), 77–89, doi:10.1016/j.msea.2008.06.033 10 F. Eckermann, T. Suter, P. J. Uggowitzer, A. Afseth, P. Schmutz, Investigation of the exfoliation-like attack mechanism in relation to Al–Mg–Si alloy microstructure, Corrosion Science, 50 (2008) 7, 2085–2093, doi:10.1016/j.corsci.2008.04.003 11 Z. Wang, H. Li, F. Miao, W. Sun, B. Fang, R. Song, Z. Zheng, Im- proving the intergranular corrosion resistance of Al–Mg–Si–Cu alloys without strength loss by a two-step aging treatment, Mater. Sci. Eng. A, 590 (2014), 267–273, doi:10.1016/j.msea.2013.10.001 12 A. Halap, M. Popovi}, T. Radeti}, V. Va{~i}, E. Romhanji, Influence of the thermo-mechanical treatment on the exfoliation and pitting corrosion of an AA5083-type alloy, Mater. Tehnol., 48 (2014) 4, 479–483 13 EN ISO 11846, Corrosion of metals and alloys, Determination of resistance to intergranular corrosion of solution heat- treatable aluminium alloys 14 ASM Handbook, Vol. 13, Corrosion, ASM, Ohio 1987 P. SLÁMA, J. NACHÁZEL: EFFECT OF THERMOMECHANICAL TREATMENT ON THE INTERGRANULAR CORROSION ... 910 Materiali in tehnologije / Materials and technology 50 (2016) 6, 903–910 Y. GHERNOUTI et al.: VALORIZATION OF BRICK WASTES IN THE FABRICATION OF CONCRETE BLOCKS 911–916 VALORIZATION OF BRICK WASTES IN THE FABRICATION OF CONCRETE BLOCKS OCENA ODPADKOV IZ OPEKE PRI PROIZVODNJI BETONSKIH ZIDAKOV Youcef Ghernouti1, Bahia Rabehi1, Tayeb Bouziani2, Rabah Chaid1 1University M’Hamed Bougara of Boumerdes, Research Unit of Materials, Processes and Environment, Boumerdes, Algeria 2University Amar Telidji of Laghouat, Structures Rehabilitation and Materials Laboratory (SREML), Algeria y_ghernouti@yahoo.fr Prejem rokopisa – received: 2015-07-05; sprejem za objavo – accepted for publication: 2015-10-30 doi:10.17222/mit.2015.202 This work focuses on the reuse of recycled brick waste (RBW) as aggregates in the fabrication of concrete blocks. The experi- mental study was focused on six different concrete compositions with a w/c ratio of 0.56, a relatively constant compactness and a slump value of zero. The six compositions consist on a control concrete with natural sand and five compositions with 10 %, 20 %, 30 %, 40 % and 50 % of RBW as a partial substitute for the natural sand. The physical and mechanical properties of concrete blocks were studied, analyzed and compared. The obtained results showed that it is possible to manufacture concrete blocks based on RBW, and that the compressive strengths of these concrete blocks are comparable to that of the control concrete, but with an appreciable reduction in weight. The blocks made with 30 % of RBW showed an improvement in the compressive strength of 42 % and a reduction in weight of 11 % compared to the control concretes. Keywords: recycled brick waste, concrete block, compactness, slump, mechanical strength Delo je usmerjeno v ponovno uporabo recikliranih odpadkov opeke (RBW) kot sestavina pri izdelavi betonskih zidakov. Eksperimentalno delo je bilo usmerjeno v {est razli~nih sestav betona z razmerjem w/c je 0,56, z relativno enako kompaktnostjo in brez zmanj{anja vrednosti. [est sestav je predstavljalo kontrolni beton z naravnim peskom in pet sestav z dodatkom 10 %, 20 %, 30 %, 40 % in 50 % RBW, kot delnim nadomestkom za naravni pesek. Prou~evane, analizirane in primerjane so fizikalne in mehanske lastnosti cementnih zidakov. Dobljeni rezultati so pokazali, da je mogo~a izdelava cementnih zidakov na osnovi RBW. Tla~ne trdnosti teh betonskih zidakov so primerljive s tistimi iz kontrolnega betona, ob~utno pa je zmanj{anje te`e. Zidaki izdelani z 30 % RBW so pokazali izbolj{anje tla~ne trdnosti za 42 % in zmanj{anje te`e za 11 %, v primerjavi z zidaki iz kontrolnega betona. Klju~ne besede: reciklirani odpadki iz opeke, betonski zidak, kompaktnost, padec vrednosti, mehanska trdnost 1 INTRODUCTION In the past decade, Algeria has been experiencing rapid development in the construction sector. Indeed, several construction projects supported by the state have been launched. The concrete block occupied an im- portant place in this sector; this is mainly due to the simplicities related to its prefabrication and the handling facilities on site. Like any conventional concrete, con- crete block consists mostly of gravel, sand, cement and water. The concrete used for the precast blocks is charac- terized by a rather dry state in the fresh state (needed to confer an immediate unmolding of the block) and a delicate physico-mechanical behavior in the hardened state. The difference between this type of concrete and the conventional concrete lies mainly in the low cement and water content; that is to say a high dosage of aggre- gate. In the context of the judicious use of aggregates and the development of a strategy for the sustainable deve- lopment policy in the building and construction sector, the use of local resources and recycled waste, such as brick waste, is required. Indeed, the introduction of recycled brick waste (RBW) in the construction industry was the subject of several research works in recent years. Thus, the use of RBW as alternative aggregates has a particular interest as it can considerably reduce the prob- lem of waste storage and, on the other hand, can help in the preservation of natural aggregates.1 The use of clay brick as aggregates in concrete was proposed in the 1990s.2 Only a few researchers have studied the potential of using clay brick powder as a partial cement replacement to make mortar. G. Moriconi et. al.3 and L. Turanli et al.4 found that RBW, as a pozzolanic material, had the potential to suppress expansion due to the alkali-silica reaction. The possibi- lity of using RBW as a replacement for cement has been investigated in the study of Naceri et al.5 The authors found that the mechanical behavior (compressive and flexural strengths) at 7 d and 28 d of hardened mortar decreased with an increasing RBW content. However, at 90 d the mortars containing up to 10 % of the waste brick will reach a resistance comparable to those of mortars without RBW. S. Wild et al.6 and M. O’Farrell et al.7 reported that the presence of RBW influenced the compressive strength and the pore size distribution of mortar. Materiali in tehnologije / Materials and technology 50 (2016) 6, 911–916 911 UDK 628.4.036:628.477.2:620.28 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)911(2016) Recently, some researchers have studied the possibility of using RBW as aggregate to make high-strength concrete.8–14 P. B. Cachim15 reported that crushed bricks could be used as a partial replacement for natural coarse aggregate without a reduction in concrete properties for a 15 % replacement ratio; however, a reduction up to 20 % has been noted for a 30 % replacement ratio. A. K. Padmini et al.16 reported that for a given strength, the modulus of elasticity of concrete made with crushed brick is between one-half and two-thirds that of normal concrete. Moreover, the water absorption and sorptivity increased for the concrete containing crushed- brick aggregates. Furthermore, concrete containing coarse crushed bricks aggregate had a relatively lower strength during the early ages than normal aggregate concrete. This is due to the higher water absorption of crushed brick aggregates compared to natural aggre- gates.17 A. R. Khaloo18 found a decrease of 7 % in the con- crete’s compressive strength by using crushed clinker bricks as the coarse aggregate compared to natural aggregate. A. A. Akhtaruzzaman and A. Hasnat19 found that the tensile strength of concrete containing coarse crushed brick was higher than that of normal concrete by about 11 %. T. Kibriya and P. R. S. Speare20 reported that con- crete containing coarse, crushed brick had comparable compressive, tensile and flexural strengths to those of normal concrete, but the modulus of elasticity was drastically reduced. C. S. Poon and D. S. Chan21 found that the incorpora- tion of 20 % of fine crushed brick aggregate decreased the compressive strength and the modulus of elasticity of the concrete by 18 % and 13 %, respectively. The present experimental investigation constitutes a continuation of the work and aims to expand the use of this material in the prefabrication of concrete blocks. In this work, an optimizing of concrete block mixtures, based on natural sand and different percentages of RBW, was performed. Next, the influence of RBW on the physico-mechanical properties of the produced concrete blocks was tested. 2 EXPERIMENTAL PART 2.1 Materials The concrete block mixtures investigated in this study were prepared with Ordinary Portland Cement (OPC) CEM II/A 42.5. The mineralogical and chemical compositions of the cement are listed in Table 1. The aggregates used are natural sand (NS), with a maximum particle size of 2 mm and a siliceous mineralogical nat- ure. A crushed limestone gravel with a particle size bet- ween 3 mm and 8 mm and a recycled brick waste (RBW) aggregate resulting from crushing of the rejected bricks, composed mainly from the quartz and with a maximum particle size of 2 mm. The physical properties and gra- nular size analysis of all the aggregates used in this work are listed and presented in Table 2 and Figure 1. Table 2: Physical properties of aggregates used Tabela 2: Fizikalne lastnosti uporabljenih sestavin Natural sand (NS) Recycled brick waste (RBW) Gravel (3/8) Apparent density, Ad (g/cm3) 1.49 0.97 1.35 Specific gravity, SG (g/cm3) 2.59 1.21 2.64 Visual equivalent, VES (%) 72 67 / Finesse modulus, Fm 1.05 4.7 / Porosity (%) 26.6 34.9 33.3 Water absorption (%) 1.86 7.4 1.4 2.2 Formulation of concrete blocks In the mix design of this type of concrete, the com- pactness criterion and maneuverability have been con- sidered for the fresh state, since the mechanical strength of this type of concrete used in the manufacture of con- crete blocks is not a very important criterion. (The mechanical strengths of the blocks are relatively low compared to traditional concrete). The first step in for- mulating the blocks is the optimization of the aggregates dosage (natural Sand + Gravel) by choosing the most compact mixture. Then the second step is to search the Y. GHERNOUTI et al.: VALORIZATION OF BRICK WASTES IN THE FABRICATION OF CONCRETE BLOCKS 912 Materiali in tehnologije / Materials and technology 50 (2016) 6, 911–916 Table 1: Chemical and mineralogical compositions of the cement Tabela 1: Kemijska in mineralo{ka sestava cementa Chemical composition (%) Mineralogical composition (%) CaO SiO2 Al2O3 Fe2O3 MgO SO3 LOI Na2O K2O C3S C2S C3A C4AF 62.2 19.4 5.4 2.8 1.7 2.5 4.6 0.35 0.76 60 21 8 11 Figure 1: Granular size analysis of all aggregates Slika 1: Analiza velikosti zrn vseh sestavin dosage of water that verifies the criteria required for the concrete blocks (no slump and maximum compactness), while fixing the cement content (8 % to 9 % by weight of the aggregates).22 The third step is to replace some natural sand in the optimized mixture by different per- centages of RBW, from 10 % to 50 %, with an increment of 10 %. 2.2.1 Optimization of aggregates dosage In our work, we started by optimizing the dosage of dry aggregates, using as criteria the maximum compact- ness of the mixture. We used a Modified Proctor test for determining the compactness of the mixtures (gravel and sand). Calculating the compactness is performed after a period of vibration of 30 s using a standard vibrating table. The results of the compactness of the mixture (sand + gravel) are represented in Figure 2. According to the results, the most compact mixture is that which con- tains 60 % gravel and 40 % sand (compactness = 0.572). 2.2.2 Optimization of water content The second step is to find the dosage of water that verifies the criteria required for concrete blocks (no slump and maximum compactness), while fixing the cement content (8 % to 9 % of the weight of the aggre- gate). The water dosage ranges from 2 % to 12 % by weight of solid mixture (gravel, sand and cement). The slump is performed using the Abrams cone. In parallel the compactness of prepared fresh concrete was measured using a modified proctor mold. The results of the measurements of the Slump flow and compactness depending on the percentage of water are shown in Figure 3. Corresponding pictures for each mixture are shown in Table 3. From the results obtained, the most compact mixture that checks a zero slump flow is the mixture containing 10 % of water. 2.2.3 Incorporation of RBW in the optimized composition In this step, a portion of the natural sand from the optimized formulation was replaced by different per- centages (10 %, 20 %, 30 %, 40 % and 50 %) of RBW aggregate. Table 4: Formulations and dosages of the constituents in kg/m3 Tabela 4: Sestava in odmerek sestavin v kg/m3 Formu- lation Cement Gravel Water Natural sand (NS) Recycled brick waste (RBW) BNS 142 875 80 564 / BBW10 508 56 BBW20 452 113 BBW30 339 170 BBW40 508 226 BBW50 282 282 Y. GHERNOUTI et al.: VALORIZATION OF BRICK WASTES IN THE FABRICATION OF CONCRETE BLOCKS Materiali in tehnologije / Materials and technology 50 (2016) 6, 911–916 913 Figure 3: Compactness and slump flow values depending on the per- centage of water Slika 3: Kompaktnost in padec teko~nosti v odvisnosti od odstotka vode Figure 2: Compactness value depending on the percentage of sand and gravel Slika 2: Vrednosti kompaktnosti v odvisnosti od odstotka peska in gramoza Table 3: Examples of slump flow depending on the percentage (%) of water Tabela 3: Primer zmanj{anja teko~nosti v odvisnosti od odstotka (%) vode (%) of water Images ofslump flow test (%) of water Images of slump flow test (%/w) = 2 % Slump flow =0 mm Compactness = 0.43 (%/w) = 4 % Slump flow =0 mm Compactness = 0.49 (%/w) = 6 % Slump flow =0 mm Compactness = 0.5 (%/w) = 7% Slump flow =0 mm Compactness = 0.55 (%/w) = 8 % Slump flow = 0 mm Compactness = 0.55 (%/w) = 9 % Slump flow = 0 mm Compactness = 0.58 (%/w) = 10 % Slump flow = 0 mm Compactness = 0.61 (%/w) = 12 % Slump flow = 30 mm Compactness = 0.55 Table 4 gives the dosages of the constituents in the mixture for the optimized formulation of the block con- crete based on natural sand and block concrete based on RBW aggregate. 2.2.4 Optimization of slump flow and compactness For each composition, the conditions for obtaining concrete blocks for a slump value of 0 are respected. The obtained results are shown in Table 5. Table 5: Examples of slump flow for all formulations Tabela 5: Primeri zmanj{anja teko~nosti vseh sestav Composition Slump flow test Composition Slump flow test BNS Slump flow = 0 mm Compactness = 0.52 BBW10 Slump flow = 0 mm Compactness = 0.52 BBW20 Slump flow = 0 mm Compactness = 0.53 BBW30 Slump flow = 0 mm Compactness = 0.54 BBW40 Slump flow = 0 mm Compactness = 0.49 BBW50 Slump flow = 0 mm Compactness = 0.42 2.2.5 Preparation of concrete blocks Our study was performed on hollow blocks with dimensions of (10 × 20 × 40) cm. The mixing is performed on site with a concrete mixer. The mixed concrete is then loaded into the laying machine. Excess concrete is leveled using a striking surface so that the blocks can have a rough surface. The blocks are removed from the molds immediately and thoroughly on a concrete platform (Figure 4). The blocks are kept on site for 24 h, and then they are transported to the laboratory and are watered every day for 28 d. 2.3 Characterization of concrete blocks 2.3.1 Dimensional variation The tests of dimensional variation were conducted on specimens of (4× 4×16) cm with the same concrete made for the realized concrete blocks. For each formulation, four specimens were crafted. Two are left in the open air to measure the shrinkage and two are immersed in the water to measure the swelling. 2.3.2 Porosity and water absorption The measurements of porosity and water absorption are performed according to the NF P18 554 standards, on block samples previously realized. The porosity is the amount of water absorbed using a dry sample mass. It is determined using the following Equation (1): P= [(Ma – MS) / (Ma – Mà)] × 100 (1) where: Ma: weight of sample at a dry surface. Ms: weight of the sample after drying. Mà: weight of the sample after immersion in water for 24 h. 2.3.3 Compressive strength After surfacing of the lower and upper bearing faces of each block with sulfur, the compression test is performed by applying a continuous load without shock at a constant speed of 0.5 MPa/s. The test machine is a press for hard materials according to NFP 18-412; it is calibrated in terms of these standards (Figure 5). The compressive strength Rc is obtained using the following Equation (2): RC= [C / (Sb×10)] × (Sa / Sn) (2) where: C: breaking load of the block, Sb (Gross Section): Area obtained by multiplying both dimensions, thickness and length, measured in the same horizontal section, Sn (Net Section): Area in a horizontal section concrete, empty deducted, Y. GHERNOUTI et al.: VALORIZATION OF BRICK WASTES IN THE FABRICATION OF CONCRETE BLOCKS 914 Materiali in tehnologije / Materials and technology 50 (2016) 6, 911–916 Figure 4: Fabrication of blocks Slika 4: Izdelava zidakov Figure 5: Compressive strength test Slika 5: Preizkus tla~ne trdnosti Sa (Support Section): Common area of contact face and supporting face. 3 RESULTS AND DISCUSSION 3.1 Dimensional variation of the concrete blocks The results of shrinkage and swelling are shown in Figure 6. The obtained results show that for all compo- sitions of concrete blocks, the shrinkage increases rapidly until the age of 14 d, ranges from 9.43 mm/m to 10.54 mm/m, which may result in a loss of weight due to the phenomenon of setting and hardening of concrete in the early days of hydration, subsequently a dimensional stability is recorded until 28 d. The values of the shrink- age for all compositions based on RBW are lower than the composition based on natural sand; this may be due to the improvement of compactness by the incorporation of RBW. All the specimens show a considerable swelling until the age of 14 d, ranges from 13.54 mm/m to 15.58 mm/m, due to water absorption by the concrete blocks, subsequently a dimensional stability is recorded until 28 d; this is may be due to the saturation of the pores and capillaries. The swelling of the BBW40 and BBW50 specimens was not studied because it deteriorated immediately after the immersion in water. Finally, the obtained results show that the replacement of sand by RBW until 30 % does not have a significant effect on the evolution of the shrinkage and the swelling of the concrete blocks. The maximum difference is in the order of 11 % in the case of shrinkage and 9 % in the case of swelling. 3.2 Porosity and water absorption of concrete blocks The evolution of porosity and water absorption are shown in Figure 7. These results show that the water absorption of all the blocks varies in the same manner as the porosity. The porosity decreases with the replace- ment content of sand by the RBW aggregate. However, up to a 40 % of replacement, an increase of the porosity is recorded. The decrease of the porosity and water ab- sorption at a less than 40 % replacement of sand by RBW aggregate, can be explained by the form of the RBW aggregate (angular form), the RBW makes it possible to improve the compactness of mixture, the contact is perfect and distribution of waste brick grains, is uniform. Indeed, it fills the voids among the grains of sand. The mixtures containing less than 40 % of RBW waste aggregates having a gravel 3/8 with two sands (natural and recycled) having a size more or less identical with the presence of some fine material for the recycled aggregate, which allows for a more compact granular skeleton. The composition with 30 % RBW has a porosity of 11 % less than the composition with natural sand. The increased porosity and water absorption beyond 40 % replacement can be explained by the large amount of RBW aggregate in the mixture; it is a porous material in comparison to the natural sand, which is a less porous material. Y. GHERNOUTI et al.: VALORIZATION OF BRICK WASTES IN THE FABRICATION OF CONCRETE BLOCKS Materiali in tehnologije / Materials and technology 50 (2016) 6, 911–916 915 Figure 7: Porosity and water absorption of concrete blocks Slika 7: Poroznost in absorpcija vode v betonskih blokih Figure 6: Shrinkage and swelling of concrete blocks Slika 6: Kr~enje in nabrekanje betonskih zidakov Figure 8: Compressive strength of concrete blocks in comparison with ordinary concrete blocks Slika 8: Tla~na trdnost cementnih zidakov v primerjavi z obi~ajnimi cementnimi zidaki 3.3 Compressive strength of concrete blocks The results of the compressive test on all the concrete blocks are shown in Figure 8. The blocks with RBW aggregate have a greater compressive strength than the blocks with natural sand, the concrete blocks containing 30 % of RBW have a gain of about 43 % compared to the concrete blocks with natural sand in compressive strength, which can be explained by the high compact- ness of this composition based on RBW, while the ab- sorption and porosity decrease in parallel. The com- pressive strength of the BBW40 concrete blocks decreases; this may be due to the decrease in compact- ness. The compressive strengths of all the realized blocks (BNS, BBW10, BBW20 and BBW30) are better than those of the usual blocks realized in the block pre- fabrication site (ordinary concrete blocks: OCB). 4 CONCLUSION This study presents the use of recycled brick waste (RBW) as sand in concrete blocks. On the basis of the obtained results, the following conclusions can be drawn: • It is possible to use RBW as a fine aggregate for the manufacturing of concrete blocks. The shrinkage and swelling of these blocks decreases according to the increase of compactness. • All the studied concrete blocks have the same density regardless of the replacement rate of natural sand by RBW. • The replacement of 30 % natural sand by the RBW enabled us to achieve concrete blocks with better characteristics: a maximum compactness, an accept- able shrinkage and swelling (similar to that of con- crete with natural sand), a low porosity and water absorption in comparison with other compositions, a weight reduction of 11 % and a higher compressive strength than the concrete blocks with natural sand (a gain of 43 %). Finally, we can conclude that the RBW can be used as a fine aggregate to produce concrete blocks, which allows us to reduce the waste inventory levels in brick and limit the deficit aggregates in some areas. 5 REFERENCES 1 F. Debieb, S. Kenai, The use of coarse and fine crushed bricks as aggregate in concrete, Constr. Build. Mater., 22 (2008) 5, 886–93, doi:10.1016/j.conbuildmat.2006.12.013 2 T. C. Hansen, Recycling of demolished concrete masonry, Rilem Report No. 6, E&FN Spon, London, 1992, 316 3 G. Moriconi, V. Corinaldesi, R. Antonucci, Environmentally-friendly mortars: a way to improve bond between mortar and brick, Mater Struct, 36 (2003) 10, 702–708, doi:10.1007/BF02479505 4 L. Turanli, F. Bektas, P. J. M. Monteiro, Use of ground clay brick as a pozzolanic material to reduce the alkali–silica reaction, Cem. Concr. Res., 33 (2003) 10, 1539–1542, doi:10.1016/S0008-8846(03) 00101-7 5 A. Naceri, H. M. C. Hamina, Use of waste brick as a partial replacement of cement in mortar, Waste Management, 29 (2009) 8, 2378–2384, doi:10.1016/j.wasman.2009.03.026 6 S. Wild, J. M. Khatib, S. D. Addis, The potential of fired brick clay as a partial cement replacement material, International congress-con- crete in the service of mankind, concrete for environment enhance- ment and products, University of Dundee, Eds. Dhir and Dyer, E&FN SPON, (1996), 685–696 7 M. O’Farrell, S. Wild, B. B. Sabir, Pore size distribution and com- pressive strength of waste clay brick mortar, Cem Concr Compos, 23 (2001) 1, 81–91, doi:10.1016/S0958-9465(00)00070-6 8 B. Rabehi, Y. Ghernouti, K. Boumchedda, Strength and compressive behaviour of ultra high-performance fibrereinforced concrete (UHPFRC) incorporating Algerian calcined clays as pozzolanic materials and silica fume, European Journal of Environmental and Civil Engineering, 17 (2013) 8, 599–615, doi:10.1080/19648189. 2013.802998 9 B. Safi, A. Aboutair, M. Saidi, Y. Ghernouti, C. Oubraham, Effect of the heat curing on strength development of ultra-high performance fiber reinforced concrete (UHPFRC) containing dune sand and ground brick waste, J. Build. Mater. Struct., 1 (2014) 1, 40–46 10 M. A. Mansur, T. H. Wee, S. C. Lee, Crushed bricks as coarse aggre- gate for concrete, ACI Mater. J., 96 (1999) 4, 478–84 11 H. L. Cheng, Study on recycled concrete made by fly-ash and waste clay-brick, China Concr. Cem. Prod., (2005) 5, 48–50 12 J. M. Khatib, Properties of concrete incorporating fine recycled aggregate, Cem. Concr. Res., 35 (2005) 4, 763–769, doi:10.1016/ j.cemconres.2004.06.017 13 F. M. Khalaf, Using crushed clay brick as coarse aggregate in con- crete, J. Mater. Civil. Eng., 18 (2006) 4, 518–526, doi:10.1061/ (ASCE)0899-1561(2006)18:4(518) 14 H. D. Yan, X. F. Chen, Experimental studies and analyses on pro- perties of clay brick recycled aggregate concrete, Tenth nation cement and concrete chemistry and application technology conference, 2007, China 15 P. B. Cachim, Mechanical properties of brick aggregate concrete, Constr. Build. Mater., 23 (2009) 3, 1292–1297, doi:10.1016/ j.conbuildmat.2008.07.023 16 A. K. Padmini, K. Ramamurthy, M. S. Mathews, Relative moisture movement through recycled aggregate concrete, Mag. Concr. Res., 54 (2002) 5, 77–384, doi:10.1680/macr.2002.54.5.377 17 M. Zakaria, J. G. Cabrera, Performance and durability of concrete made with demolition waste and artificial fly ash-clay aggregates, Waste Manage, 16 (1996) 1–3, 151–158, doi:10.1016/S0956- 053X(96)00038-4 18 A. R. Khaloo, Properties of concrete using crushed clinker brick as coarse aggregate, ACI Mater. J., 91 (1994) 4, 401–407 19 A. A. Akhtaruzzaman, A. Hasnat, Properties of concrete using crushed brick as aggregates, Concr. Int., 5 (1983) 2, 58–63 20 T. Kibriya, P. R. S. Speare, The use of crushed brick coarse aggregate concrete, Proceedings of international conference–concrete for environment enhancement and protection, Scotland, University of Dundee, 1996 21 C. S. Poon, D. Chan, The use of recycled aggregate in concrete in Hong Kong, Resour Conserv. Recycl., 50 (2007) 3, 293–305, doi:10.1016/j.resconrec.2006.06.005 22 T. Bouziani, A. Ferhat, M. Bederina, Optimisation des paramètres de formulation des bétons destinés à la préfabrication des blocs de béton par une approche RNA (Réseaux de Neurones Artificiels), Revue des Sciences et Sciences de l’Ingénieur, 2 (2011), 34–41 Y. GHERNOUTI et al.: VALORIZATION OF BRICK WASTES IN THE FABRICATION OF CONCRETE BLOCKS 916 Materiali in tehnologije / Materials and technology 50 (2016) 6, 911–916 P. SALVETR et al.: POROUS MAGNESIUM ALLOYS PREPARED BY POWDER METALLURGY 917–922 POROUS MAGNESIUM ALLOYS PREPARED BY POWDER METALLURGY POROZNE MAGNEZIJEVE ZLITINE, IZDELANE S POMO^JO METALURGIJE PRAHOV Pavel Salvetr, Pavel Novák, Dalibor Vojtìch University of Chemistry and Technology, Department of Metals and Corrosion Engineering, Technicka 5, 166 28 Prague 6, Czech Republic psalvetr@seznam.cz Prejem rokopisa – received: 2015-07-14; sprejem za objavo – accepted for publication: 2015-10-30 doi:10.17222/mit.2015-226 This paper deals with the development of porous magnesium alloys that can be used in medicine for bone fixations and implants. The individual components of the alloys were chosen so that the biodegradability of the material is maintained. The advantage of these magnesium materials should be an ability to decompose after some time. This should reduce the number of surgeries and consequently increase the comfort of patients. All the samples were prepared using the method of powder metallurgy. The influence of particular alloying elements – aluminium, zinc, yttrium – on the structure of the alloys was explored, with changes being seen in the area of the fraction of pores, the size and the shapes of the pores, according to the alloying elements and the prolongation of the time of sintering the powders. By altering the chemical composition and the time of sintering the demanded porosity was not achieved, and that is the reason why a pore-forming agent (ammonium carbonate) was added. It was removed by thermal decomposition before the powder’s sintering. By adding ammonium carbonate we managed to increase the porosity and at the same time we obtained more pores (in equivalent diameter 200–400 μm). The mechanical properties of the samples were tested in compression. In the samples without the pore-forming agent the values of the ultimate strength were larger than the values of natural bones. After adding the pore-forming agent the ultimate strength and modulus elasticity were reduced. Keywords: magnesium alloys, powder metallurgy, porosity, biomaterial ^lanek obravnava razvoj poroznih magnezijevih zlitin, ki bi lahko bile uporabne v medicini za utrjevanje kosti in vsadke. Posa- mezne komponente zlitin so bile izbrane tako, da je ostal material biolo{ko razgradljiv. Prednost teh magnezijevih materialov naj bi bila zmo`nost, da se s ~asom razgradi. To naj bi zmanj{alo {tevilo operacij in s tem pove~alo ugodje pacientov. Vsi vzorci so bili pripravljeni po postoku metalurgije prahov. Preiskovan je bil vpliv posameznih legirnih elementov (aluminij, cink, itrij) na mikrostrukturo zlitin. Spremembe so bile opa`ene pri dele`u por, velikosti in obliki por glede na legirni element in podalj{anje ~asa sintranja prahov. S spreminjanjem kemijske sestave in ~asa sintranja zahtevana poroznost ni bila dose`ena in to je tudi razlog, zakaj je bilo dodano sredstvo za nastajanje por (amonijev karbonat). To sredstvo je bilo odstranjeno s toplotnim razpadom pred sintranjem prahu. Z dodatkom amonijevega karbonata nam je uspelo pove~ati poroznost in isto~asno smo dobili ve~ por ekvivalentnega premera 200–400 μm. Mehanske lastnosti vzorcev so bile preizku{ene s stiskanjem. Vzorci brez sredstva za nastajanje por so dosegli vrednosti poru{ne trdnosti ve~je od vrednosti pri naravnih kosteh. Po dodajanju sredstva za nastanek por sta se poru{na trdnost in modul elasti~nosti zmanj{ala. Klju~ne besede: magnezijeve zlitine, metalurgija prahov, poroznost, biomaterial 1 INTRODUCTION Magnesium and its alloys have an import role among structural materials. They exceed other materials, espe- cially with their low density. Magnesium is used mainly in the construction of vehicles and the aerospace industry, where it is used in mechanically less-strained components. These special applications include, for example, steering wheels, dashboards, seats and gear- boxes.1 In the future, magnesium is also proposed to be a material able to store hydrogen.2 Nowadays, magnesium alloys are a subject of inten- sive research and development for applications in medicine as an osteosynthetic material. These materials have the ability to biodegrade, which means to decom- pose and be absorbed into a human body. The main advantages of these kinds of implants would be reducing the number of surgeries. The so far used biomaterials are based on bio-inert (non-reactive and corrosion-resistant alloys). The human body may have a problem with accepting these materials. This group comprises stainless steel, titanium and cobalt alloys.3 Biodegradable mate- rials must fulfil many requirements, i.e., they must not release toxic doses of metallic ions and both the products of the corrosion reactions and the original biomaterial must not cause any allergic reaction of the organism. Therefore, the appropriate corrosion rates should be reached. The implant must not decompose too early. For example, a screw fixation of broken bones should work for 12–16 weeks, as a minimum. During this time the implant must keep its mechanical properties, which should be similar to the mechanical properties of natural bones. This requirement is met by magnesium and its alloys quite well, as you can see in Table 1.4,5,7–10 Zinc alloys are investigated as competitive materials for bio- degradable implants.6 Materiali in tehnologije / Materials and technology 50 (2016) 6, 917–922 917 UDK 676.017.62:669.721.5:621.763 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)917(2016) Table 1: Mechanical properties of porous biomaterials Tabela 1: Mehanske lastnosti poroznih biomaterialov Material Porosity(x/%) Pore size (ìm) Compressive strength (MPa) Modulus (GPa) Refe- rence Porous Mg 29–31 250–500 20–70 – 7 Porous Mg 36–55 200–400 15–31 4–18 8 Porous Mg 35–55 100–400 12–17 1–2 9 Porous Ti 78 200–500 35 5 10 Porous HA 50–77 200–400 1–17 0.1–7 8 Natural bone – – 2–180 0.1–20 8 Magnesium alloys mostly have better mechanical properties than pure magnesium. Corrosion resistance is increased by aluminium, zinc and rare-earth metals, but alloying elements such as iron, nickel and copper make it worse. All alloying elements in biodegradable materials must maintain the biocompatibility of the alloy. The improvement of the mechanical properties happens especially by precipitation hardening or by strengthening of a solid solution. An element that can be used is zinc, which naturally occurs in body tissue. The results of biochemical and histological investigations show that the degradation of the Mg-Zn alloy should not harm the organism.5 The maximum solubility of zinc in magne- sium is 6.2 % of mass fractions, and it improves the corrosion resistance and mechanical properties. The influence of aluminium on magnesium alloys is similar: it increases the strength and corrosion resistance. In a magnesium solid solution there is dissolved a maximum of 12.7 % of mass fractions of aluminium. Magnesium alloys with aluminium are heat treatable to form the Mg17Al12 phase. Alloying with aluminium causes microporosity and according to some studies it can cause Alzheimer’s disease. Other alloying elements that influence the mechanical corrosion properties in a posi- tive way are rare-earth metals – the lanthanides, yttrium and scandium. To magnesium alloys they are usually added in relatively small amounts, where one or two ele- ments are dominant and the rest is added only in a small amount. The rare-earth metals can be divided into a group with better solubility (Y, Gd, Tb, Dy, Ho, Er, Tm, Yb, Lu) and another group with limited solubility. Together with magnesium they create a eutectic system with limited solubility. Intermetallic phases (Mg2Y, Mg24Y5) limit the movement of dislocations and increase the ultimate tensile strength.11 Lanthanum and cerium could be added in a limited amount according to a bio- logical test.12 Calcium and zirconium also have a positive influence on the mechanical properties. Zirconium re- fines the structure, while calcium hardens the magne- sium-based solid solution and forms the Mg2Ca phase. Calcium is the main component of bones and in combi- nation with magnesium it improves their healing.13,14 To ensure the osseointegration of the implant and the consequent substitution of the implant by the bone, the presence of pores with an average of 150–200 μm as a minimum is important.15 The porous structure can be produced with various processes. For sample preparation by powder metallurgy a pore.-forming agent (carbamide, ammonium bicarbonate) is added, which is removed by thermal decomposition.8,16 With this process, pores with a random distribution and size are formed. Regular ordering of the pores is achieved by low-pressure casting of the magnesium alloy into a NaCl template.17 This paper aims to prove the possibility of preparing porous magnesium alloys by sintering from elemental metallic powders with the addition of a pore-forming agent. 2 EXPERIMENTAL PART All the samples were prepared by the method of powder metallurgy from Mg, Zn, and Y powders. The powders were blended manually and uniaxially pressed at a pressure of 530 MPa to form cylindrical green com- pacts. The samples were sintered in a tubular furnace under an argon atmosphere. The sintering temperature was chosen according to the phase diagrams of the indi- vidual alloys (MgZn5: 575 °C, MgZn10: 405 °C, MgY5: 600 °C). The duration of the sintering was 2 h. For the MgZn5 and MgZn10 samples, the sintering was pro- longed up to 4 h or 24 h in order to observe the influence of the sintering time on the porosity and mechanical pro- perties. The pore-forming agent (NH4)2CO3, which was used in some samples, was removed before sintering by thermal decomposition. The decomposition was carried out at 230 °C for 1 h and the products of the decom- position were drained by a flow of argon. Subsequent sintering was conducted at the same temperature as for the samples without the addition of the pore-forming agent for 4 h. To determine the porosity and to observe the micro- structure, metallographic samples were prepared. The samples were mounted into methacrylate resin, ground by sandpapers P180–P4000 (abrasive elements SiC and Al2O3), polished by a water-based suspension of Al2O3 (Topol 2) or diamond paste D2. The microstructure of the samples was revealed by etching in Nital (2 mL HNO3 + 98 mL ethanol). The microstructure was observed with an optical metallographic microscope (Olympus PME3) and docu- mented by AxioVision image-processing software. A more detailed observation and chemical microanalysis were performed with a TESCAN VEGA 3 LMU scann- ing electron microscope equipped with an OXFORD Instruments INCA 350 EDS analyser. Macrographs for the measurement of porosity were acquired with a Carl Zeiss Neophot 2 optical metallo- graphic microscope. The porosity was evaluated with Lucia 4.8 image-analysis software as the area fraction of pores in the cross-section. The mechanical properties of the samples were tested in compression at room temperature. The ultimate com- pressive strength (UCS) and modulus of elasticity E were determined from the stress-strain curves. The measure- P. SALVETR et al.: POROUS MAGNESIUM ALLOYS PREPARED BY POWDER METALLURGY 918 Materiali in tehnologije / Materials and technology 50 (2016) 6, 917–922 ments of the mechanical properties were performed using a LabTest 5.250SP1-VM universal testing ma- chine. 3 RESULTS AND DISCUSSION In the first part of the work, the alloys were prepared by sintering mixtures of the corresponding elemental powders without the addition of the pore-forming agent. The purpose of this step was to find an alloy that can be produced using this simple production route. 3.1 MgZn In the case of the Mg-Zn alloy, the influence of the sintering duration and the amount of zinc on the structure and porosity of the samples were observed. A longer duration of sintering did not cause any changes to the microstructure of the MgZn5 alloy. The effect of a longer sintering time is a decreasing number of pores with an equivalent diameter of about 30 μm and, conse- quently, to a better-quality sintering of the powder. However, the decrease of the fractional area of the pores is not reached in a sample sintered for 24 h. The microstructure of the alloy MgZn5 (Figure 1) consists of a solid solution of Mg-Zn with a zinc content of 4 % of the mass fractions, whose grain boundaries are decorated by the intermetallic phase. This phase arises only in a thin layer and therefore its chemical composition was not determined by chemical microanalysis. Increasing the content of zinc in the alloy up to 10 % of mass fractions (MgZn10) led to an increasing number of pores with a equivalent diameter up to 100 μm, but the overall porosity decreased. Prolongation of the sintering duration from 2 to 4 h slightly decreased the porosity. The porosity of the alloys is written in Table 2. In the structure of the sample MgZn10 there are two phases: a Mg-Zn solid solution with a zinc content of 8.6 % of mass fractions (Mg) and the phase MgZn (or Mg7Zn3) with a zinc content of approximately 45 % of mass fractions. The microstructure of the MgZn10 alloy is shown in Figure 2. Table 2: Porosity of prepared materials Tabela 2: Poroznost pripravljenih materialov Alloy Area fractionpores (%) MgZn5-2h 8 MgZn5-4h 8 MgZn5-24h 10 MgZn10-2h 6 MgZn10-4h 6 MgAl5 2 MgAl3Zn1 5 MgY5 3 MgZn5+10 % of mass fractions of (NH4)2CO3 30 MgZn5+20 % of mass fractions of (NH4)2CO3 42 MgZn5+30 % of mass fractions of (NH4)2CO3 48 MgAl3Zn1+20 % of mass fractions of (NH4)2CO3 18 3.2. MgAl5 The lowest porosity and pore size were determined in the alloy with 5 % of mass fractions Al. All the pores had an equivalent diameter up to 200 μm, 94 % of them were smaller than 100 μm. The structure is formed by the Mg-Al solid solution with an Al content of 5 % of mass fractions. At the grain boundaries, the content in- creases, which may be caused by the presence of a eutec- tic phase consisting of a Mg-based solid solution and the Mg17Al12 phase. The microstructure of the alloy is shown in Figure 3. P. SALVETR et al.: POROUS MAGNESIUM ALLOYS PREPARED BY POWDER METALLURGY Materiali in tehnologije / Materials and technology 50 (2016) 6, 917–922 919 Figure 2: Microstructure of MgZn10 alloy Slika 2: Mikrostruktura zlitine MgZn10 Figure 1: Microstructure of MgZn5 alloy Slika 1: Mikrostruktura zlitine MgZn5 3.3 MgAl3Zn1 In the microstructure of the MgAl3Zn1 (AZ31) alloy, there are a large number of small pores, whose amount is similar to the alloy MgZn10. The sample consists of a solid solution of alloying elements in magnesium. The microstructure of the alloy is presented in Figure 4, the darker parts (Mg1 = 96.5 % of mass fractions of Mg, 2.5 % of mass fractions of Al, 1 % of mass fractions of Zn) have a lower content of Al, in contrast to the lighter parts (Mg2 = 95 % of mass fractions of Mg, 4 % of mass frac- tions of Al, 1 % of mass fractions of Zn), zinc is dis- persed homogeneously in the alloy. 3.4 MgY5 The porosity of the alloy alloyed with yttrium is 3.1 %. From all the investigated samples, this alloy con- tains the largest amount of micropores, whose equivalent diameter is not greater than 10 μm. These small pores were probably caused by the low diffusion coefficient of yttrium in magnesium. During sintering the mutual diffusion was suppressed and in the structure there were maintained the rest of the pores after pressing. In the SEM micrograph (Figure 5) it is obvious that particles of yttrium are not dissolved in magnesium and the solid solution is not formed. With the addition of a pore-forming agent, the alloys MgZn5 (10, 20 and 30 % of mass fractions of ammo- nium carbonate) and MgAl3Zn1 (20 % of mass fractions of ammonium carbonate) were formed. With an in- creasing content of the pore-forming agent, the porosity of the alloy MgZn5 increased. In the structure of the alloy MgAl3Zn1+20 % mass fractions of (NH4)2CO3 there was a lower porosity than in the alloy MgZn5 with the same content of the pore-forming agent (Table 2). In Figure 6 there are pores with the required size (equiva- P. SALVETR et al.: POROUS MAGNESIUM ALLOYS PREPARED BY POWDER METALLURGY 920 Materiali in tehnologije / Materials and technology 50 (2016) 6, 917–922 Figure 5: Microstructure of MgY5 alloy Slika 5: Mikrostruktura zlitine MgY5 Figure 3: Microstructure of MgAl5 alloy Slika 3: Mikrostruktura zlitine MgAl5 Figure 4: Microstructure of MgAl3Zn1 alloy Slika 4: Mikrostruktura zlitine MgAl3Zn1 Figure 6: Microstructure of MgZn5 alloy with addition of 20 % of mass fractions of (NH4)2CO3 Slika 6: Mikrostruktura zlitine MgZn5 z dodatkom 20 % masnega dele`a (NH4)CO3 lent diameter of more than 200–500 μm), which origi- nated from the thermal decomposition of (NH4)2CO3. In Figure 6 there are some cracks that may have been caused by the stress during the decomposition of the pore-forming agent. In the samples without the pore-forming agent it was found that 85–95 % of the pores do not reach the required equivalent diameter of at least 100 μm (in the alloy MgY5, as many as 99 % of the pores). By adding the pore-forming agent the number of pores greater than 100 μm increased. In Figure 7, the influence of the pore-forming agent on the size of the pores in the alloy MgAl3Zn1 is presented. The most circular pores (69 %) are contained in the structure of the MgAl3Zn1 alloy. In the other alloys the amount of circular pores is 45–60 %. Adding the pore-forming agent into the alloy MgAl3Zn1 decreased the quotient of the circular pores to 35 %. The ultimate compressive strengths of the alloys are higher in comparison with the properties of natural bone (Table 1). The ultimate compressive strength is between 203 MPa and 236 MPa. The modulus of elasticity is comparable with natural bone. A slightly lower modulus of elasticity was found in the case of the MgAl3Zn1 alloy, in contrast with the other alloys. Increased porosity and the presence of large pores cause a significant decrease in the mechanical properties. The results of the compression test (UCS, E) are shown in Figure 8. 4 CONCLUSION Magnesium alloys prepared by powder metallurgy without a pore-forming agent have an adequate modulus of elasticity and a higher ultimate compressive strength than natural bones. In the microstructure of magnesium alloys alloyed with zinc and aluminium, the solid solu- tion dominates. An exception is the microstructure of the alloy MgY5, where the particles of yttrium did not dissolve in magnesium. The porosity of the samples without the pore-forming agent is up to 10 % of the volume fractions. Approximately 90 % of pores have an equivalent diameter of less than 100 μm. The addition of the pore-forming agent – (NH4)2CO3 – to the alloys MgZn5 and MgAl3Zn1, increases the quantity of pores with an equivalent diameter of 200 μm and an area fraction of pores to 18–48 % of the volume fractions. The ultimate compressive strength reached 203–236 MPa. With an increasing porosity after adding the (NH4)2CO3 it decreases to 61 MPa. Acknowledgement This research was financially supported by Czech Science Foundation, project No. P108/12/G043. 5 REFERENCES 1 B. L. Mordike, T. Ebert, Magnesium: Properties – applications – potential, Materials Science and Engineering A, 302 (2001) 1, 37–45, doi:10.1016/S0921-5093(00)01351-4 2 D. Vojtìch, V. Knotek, Magnesium alloys for hydrogen storage, Mater. Tehnol., 46 (2012) 3, 247–250 3 G. Kla~nik, M. Zdovc, U. Kov{ca, B. Pra~ek, J. Kova~, J. Rozman, Osseointegration and rejection of a titanium screw, Mater. Tehnol., 44 (2010) 5, 261–264 4 D. Vojtìch, V. Knotek, J. ^apek, J. 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Kaplan, Porosity of 3D biomaterial scaffolds and osteogenesis, Biomaterials, 26 (2005) 27, 5474–5491, doi:10.1016/j.biomaterials.2005.02.002 16 J. ^apek, D. Vojtìch, Properties of porous magnesium prepared by powder metallurgy, Materials Science and Engineering C, 33 (2013) 1, 564–569, doi:10.1016/j.msec.2012.10.002 17 N. T. Kirkland, I. Kolbeinsson, T. Woodfield, G. J. Dias, M. P. Staiger, Synthesis and properties of topologically ordered porous magnesium, Materials Science and Engineering B, 176 (2011) 20, 1666–1672, doi:10.1016/j.mseb.2011.04.006 P. SALVETR et al.: POROUS MAGNESIUM ALLOYS PREPARED BY POWDER METALLURGY 922 Materiali in tehnologije / Materials and technology 50 (2016) 6, 917–922 G. HARDAL, B. Y. PRICE: INFLUENCE OF NANO-SIZED COBALT OXIDE ADDITIONS ON THE STRUCTURAL ... 923–928 INFLUENCE OF NANO-SIZED COBALT OXIDE ADDITIONS ON THE STRUCTURAL AND ELECTRICAL PROPERTIES OF NICKEL-MANGANITE-BASED NTC THERMISTORS VPLIV DODATKA NANODELCEV KOBALTOVEGA OKSIDA NA ZGRADBO IN ELEKTRI^NE LASTNOSTI NTC TERMISTORJEV NA OSNOVI NIKLJEVEGA MANGANITA Gökhan Hardal, Berat Yüksel Price Istanbul University, Engineering Faculty, Metallurgical and Materials Engineering Department, Avcýlar, Istanbul, Turkey berat@istanbul.edu.tr Prejem rokopisa – received: 2015-07-15; sprejem za objavo – accepted for publication: 2015-12-15 doi:10.17222/mit.2015.228 The structural and electrical properties of NiMn2O4 and Ni0.5CoxMn2.5-xO4 (where x = 0.5, 0.8 and 1.1) NTC thermistors have been investigated. The samples, prepared by conventional ceramic processing techniques, were calcinated at 900 °C for 2 h and then sintered at 1100 °C and 1200 °C for 5 h. The cubic spinel phase was observed by XRD analysis in the NiMn2O4 and Ni0.5Co0.8Mn1.7O4 samples sintered at 1100 °C for 5 h. The sintering at 1200 °C resulted in much denser microstructures with a larger grain size. The room-temperature electrical resistivity (25) and material constant (B) value of the NiMn2O4 sample sintered at 1100 °C were 7710  cm and 3930 K, respectively. The electrical resistivity of the samples decreased significantly with the addition of Co3O4. The B25/85 values of the Ni0.5CoxMn2.5-xO4 (where x = 0.5, 0.8 and 1.1) samples sintered at 1100 °C were found to be 3820 K, 3525 K and 3270 K, respectively. Keywords: cobalt oxide, electrical properties, microstructure, NTC thermistor Preiskovana je bila zgradba in elektri~ne lastnosti NiMn2O4 in Ni0.5CoxMn2.5-xO4 (kjer je x = 0,5, 0,8 in 1,1) NTC termistorjev. Vzorci, pripravljeni po obi~ajni tehniki priprave keramike, so bil kalcinirani 2 h na 900 °C in potem 5 h sintrani na 1100 °C in 1200 °C. Kubi~na {pinelna faza je bila opa`ena pri XRD-analizi, v vzorcih NiMn2O4 in Ni0.5Co0.8Mn1.7O4, sintranih 5 h na 1100 °C. Sintranje na 1200 °C je povzro~ilo mnogo bolj gosto mikrostrukturo z ve~jimi zrni. Vrednosti za elektri~no upornost pri sobni temperaturi (25) in materialne konstante (B) vzorca NiMn2O4, sintranega na 1100 °C, sta bili 7710  cm in 3930 K. Elektri~na upornost vzorcev se je ob~utno zmanj{ala po dodatku Co3O4. Vrednosti B25/85 pri vzorcih Ni0.5CoxMn2.5-xO4 (kjer je bil x = 0,5, 0,8 in 1,1) sintranih na 1100 °C so bile: 3820 K, 3525 K in 3270 K. Klju~ne besede: kobaltov oksid, elektri~ne lastnosti, mikrostruktura, NTC termistor 1 INTRODUCTION Sensors for monitoring and controlling temperature are very important, not only in our daily life but also in many industrial and laboratory applications such as aero- space and automotive industries, circuit compensation, cryogenic systems etc.1,2 NTC thermistors are useful for precision temperature measurements as their resistance decreases with increasing temperature.3 The most exten- sively used negative temperature coefficient (NTC) ther- mistor materials are nickel-manganite-based semicon- ducting materials which exhibit the spinel-type crystal structure with the general formula AB2O4.4 In the spinel structure, there are two sites available for the cations, i.e., the tetrahedral site, A-site, and the octahedral site, B-site. The distribution of the ions over the sites is as follows: Mn3+ will predominantly occupy the B-site, while Mn2+ will be placed on the A-site and the majority Ni2+ will go to the B-site.5 The electrical resistivity, , of NTC thermistors varies exponentially with temperature, T, by the well-known Arrhenius equation  = o exp (B/T), where o is the resistivity of the material at infinite temperature and B is a constant, which is a measure of the sensitivity of the materials over a given temperature.6 The material constant B, can be calculated using Equation (1): B T T T T 1 2 1 2 1 2 1 1 = − − ln ln  (1) where 1 and 2 are the electrical resistivity at tempe- ratures T1 and T2, respectively. The activation energy Ea can also found by the equation B = Ea/kB, where kB is the Boltzmann constant.7 The electrical properties of nickel-manganite-based NTC thermistors closely depend on the ratio of the com- positions (type and amount of additives), initial particle size of raw materials and processing conditions (selected synthesis method, calcination and sintering temperature, sintering time etc.). Attainment of high-density, con- trolled-grain-size microstructures and appropriate dimen- sional designs are important factors in good sensor design.8 Previous studies have been focused on the effect Materiali in tehnologije / Materials and technology 50 (2016) 6, 923–928 923 UDK 661.873:62-911-026.772:669.24:549.521.61 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)923(2016) of composition ratios and different production routes on the electrical properties of various metal-oxide-doped NTC thermistors. In this study, nano-sized cobalt-oxide- added, nickel-manganite-based NTC thermistors were fabricated by the solid-state reaction method, the effect of dopant concentration and sintering temperature on the structural and electrical properties of NTC materials were investigated. 2 EXPERIMENTAL PART The particle size of Co3O4 powder was less than 50 nm, purchased from Sigma-Aldrich. NiO (99 % purity, Alfa Aesar), Co3O4 (99.5 % purity, Sigma-Aldrich) and Mn2O3 (99 % purity, Sigma-Aldrich) powders were weighed according to the compositions of NiMn2O4 and Ni0.5CoxMn2.5-xO4 (where x= 0.5, 0.8 and 1.1). The molar ratios of these compositions are given in Table 1. The raw powder mixture was ball-milled using ZrO2 balls as a grinding media with ethyl alcohol in a jar for 5 h. The obtained slurries were dried and powders were calcinated at 900 °C for 2 h. The powders were pressed to form disc-shaped specimens and then sintered at 1100 and 1200 °C for 5 h employing a 360 °C/h heating rate in the air and then cooled naturally in the furnace. The bulk density (, g cm–3) of the sintered samples was calculated from their weights and dimensions. The phases in the sintered samples were determined by X-ray diffraction (XRD, Rigaku D/Max-2200/PC) analysis using Cu-K radiation at 60 kV/2 kW. Table 1: Molar ratio of Ni, Mn and Co in all compositions Tabela 1: Molarno razmerje Ni, Mn in Co v vseh spojinah Composition code Ni (moles) Mn (moles) Co (moles) A1 1 2 - A2 0.5 2 0.5 A6 0.5 1.7 0.8 A10 0.5 1.4 1.1 In order to calculate the lattice parameter of the sam- ples Equation (2) was applied: a d h k l= + +2 2 2 (2) where h, k and l are the miller indices, a (nm) is the lattice parameter of cubic structure, d is the interplanar spacing of the peaks corresponding to (311). The volume of the unit cell (V, nm3) for the cubic sys- tem is obtained from Equation (3): V = a3 (3) The average values of the crystallite size (D, nm) of the samples were calculated by means of X-ray line broadening method, using the Debye Scherrer formula: D = 0 9. cos   (4) where 0.9 is a constant related to crystallite shape, is the X-ray radiation wavelength in nanometres (nm),  is the full width at half-maximum (FWHM) of the peaks corresponding to (311) and  is Bragg’s angle.9 The value of  from the 2 axis of the diffraction profile must be in radians.10 The microstructure of the samples was observed using scanning electron microscopy (SEM, JEOL, JSM 5600) on fracture surfaces. The sin- tered samples were coated with silver paste to form electrodes. The electrical resistance was measured in a temperature programmable furnace between 25 °C and 85 °C in steps of 0.1 °C. The material constant, B, the activation energy, Ea, and the sensitivity coefficient, , values were calculated for the NTC thermistors. 3 RESULTS The XRD patterns of the NiMn2O4 and Ni0.5Co0.8Mn1.7O4 samples sintered at 1100 °C for 5 h are given in Figure 1. The calculated lattice parameter, unit-cell volume, , peak position corresponding to (311) and crystallite size of the samples are given in Table 2. The XRD analysis of these samples demonstrated only the cubic spinel phase (PDF No: 71-0852). A compa- rison of the XRD patterns of the sintered samples and the data is given in Table 2, the diffraction peaks of the Ni0.5Co0.8Mn1.7O4 sample shifted to higher 2 angles, and G. HARDAL, B. Y. PRICE: INFLUENCE OF NANO-SIZED COBALT OXIDE ADDITIONS ON THE STRUCTURAL ... 924 Materiali in tehnologije / Materials and technology 50 (2016) 6, 923–928 Figure 1: XRD patterns of NiMn2O4 (A1) and Ni0.5Co0.8Mn1.7O4 (A6) samples in the 2 range 20–65° Slika 1: Rentgenogram vzorcev NiMn2O4 (A1) in Ni0.5Co0.8Mn1.7O4 (A6) v podro~ju 2 med 20° in 65° as a result the lattice parameters and the unit-cell volume decreased. The value of  increased to 0.7692° and the value of average crystallite size decreased to 10.86 nm. Table 2: The lattice parameter, unit-cell volume, , peak position and crystallite size of samples sintered at 1100 °C Tabela 2: Parameter mre`e, prostornina enotne celice, , polo`aj vrhov in velikost kristalnih zrn vzorcev sintranih na 1100 °C Composition a(L) V (L3)  (311) (o) 2 (311) (o) D (nm) NiMn2O4 (A1) 0.8365 0.585 0.2538 35.6 32.87 Ni0.5Co0.8Mn1.7O4 (A6) 0.8273 0.566 0.7692 36 10.86 The bulk densities of the sintered NiMn2O4 and Ni0.5CoxMn2.5-xO4 samples are shown in Table 3. The bulk density of the A1 sample sintered at 1100 °C was found to be 4.23 g cm–3 and it increased to 4.78 g cm-3 when the sample was sintered at 1200 °C. The bulk density of the samples decreased first and then increased with the addition of Co3O4. Table 3: The bulk density of samples sintered at 1100 °C and 1200 °C for 5 h Tabela 3: Gostota osnove po 5 urnem sintranju na 1100 °C in 1200 °C Composition code  (g cm–3) 1100 °C 1200 °C A1 4.23 4.78 A2 4.05 4.43 A6 4.27 4.63 A10 4.30 4.72 The SEM micrographs of the A1, A2, A6, A10 samples sintered at 1100 and 1200 °C for 5 h are given in Figure 2. It can be seen in this figure that all the samples sintered at 1100 °C had a fine-grained microstructure with most of the pores at the grain boundaries. The grain size of A1 was larger relative to the A2, A6 and A10 samples sintered at 1100 °C. When the sintering temperature was increased to 1200 °C, all the samples had a much denser microstructure and larger grains with a number of small grains on their surface. In addition, G. HARDAL, B. Y. PRICE: INFLUENCE OF NANO-SIZED COBALT OXIDE ADDITIONS ON THE STRUCTURAL ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 923–928 925 Figure 2: SEM micrographs of sintered samples: A1 a) 1100 °C, b) 1200 °C, A2 c) 1100 °C, d) 1200 °C, A6 e) 1100 °C, f) 1200 °C, A10 g) 1100 °C, h) 1200 °C Slika 2: SEM-posnetki sintranih vzorcev: A1 a) 1100 °C, b) 1200 °C, A2 c) 1100 °C, d) 1200 °C, A6 e) 1100 °C, f) 1200 °C, A10 g) 1100 °C, h) 1200 °C the A10 sample had much bigger grains in comparison with the A2 and A6 samples sintered at 1200 °C. The plot of resistivity versus Co content (moles) and the plots of log  versus 1000/T are given in Figures 3a and 3b for all the sintered samples. The plots of log  versus 1000/T exhibited a linear dependence in the range 25–85 °C, indicating semiconducting NTC thermistor characteristics. The activation energy, the sensitivity coefficient and the material constant can also be calcu- lated from this plot. The room-temperature electrical resistances, R25, of the A1, A2, A6 and A10 samples sin- tered at 1100 °C were 1487, 360, 167 and 107 , res- pectively. For the same sintering temperature, the room temperature electrical resistivity of the A1, A2, A6 and A10 samples were calculated as 7710, 1870, 890 and 590  cm, respectively. The relationship between the B25/85 constant of the samples and the increase in Co3O4 content is given in Figure 4. The activation energy and sensitivity coeffi- cient value of the samples is given in Table 4. With in- creasing Co content, the B25/85 constant and activation energy of the samples sintered at 1100 °C decreased from 3930 K to 3270 K and from 0.338 to 0.282 eV, res- pectively. A similar tendency was also seen in the A1, A2 and A6 samples sintered at 1200 °C. For the A10 sample, the B25/85 constant and activation energy values were found to be 3620 K and 0.312 eV, respectively. The sensitivity coefficient value of all samples sintered at 1100 °C decreased from -4.426 to -3.683 %/K. When the sintering temperature increased to 1200 °C, the sensi- tivity coefficient value of all samples decreased from -4.311 to -4.078 %/K. Table 4: The activation energy and sensitivity coefficient of A1, A2, A6 and A10 samples sintered at 1100 °C and 1200 °C for 5 h Tabela 4: Aktivacijska energija in koeficient ob~utljivosti A1, A2, A6 in A10 vzorcev, sintranih 5 ur na 1100 °C in 1200 °C Composition code Ea (eV) 25 (%/K) 1100 °C 1200 °C 1100 °C 1200 °C A1 0.338 0.330 –4.426 –4.311 A2 0.329 0.313 –4.301 –4.097 A6 0.303 0.306 –3.970 –4.007 A10 0.282 0.312 –3.683 –4.078 4 DISCUSSION The cubic spinel phase was found by XRD analysis in NiMn2O4 and Ni0.5Co0.8Mn1.7O4 samples sintered at 1100 °C for 5 h. No secondary phase was found in these samples. As it is well known from the binary phase dia- gram of Mn-Ni-O, the spinel phase can only form when the ratio of Ni/(Ni+Mn) is less than 0.35 at a calcination temperature of 900 °C.11 The diffraction peaks of Ni0.5Co0.8Mn1.7O4 samples shift to the higher 2 angles, indicating a decrease in the lattice parameter with the addition of Co3O4 due to the differences between the ionic radii of the Mn and Co ions. Wu et al.12 reported that the peak shift toward higher 2 angles with the in- creasing of Co content indicates lattice constriction when G. HARDAL, B. Y. PRICE: INFLUENCE OF NANO-SIZED COBALT OXIDE ADDITIONS ON THE STRUCTURAL ... 926 Materiali in tehnologije / Materials and technology 50 (2016) 6, 923–928 Figure 4: Effect of cobalt content on B25/85 value of A1, A2, A6 and A10 samples sintered at 1100 and 1200 °C Slika 4: Vpliv vsebnosti kobalta na vrednost B25/85 vzorcev A1, A2, A6 in A10, sintranih na 1100 °C in na 1200 °C Figure 3: a) The change of resistivity as a function of cobalt content, b) the relationship between log  and 1000/T (K–1) for A1, A2, A6 and A10 samples Slika 3: a) Sprememba upornosti v odvisnosti od vsebnosti kobalta, b) odvisnost med log  in 1000/T (K–1) pri vzorcih A1, A2, A6 in A10 Co substitutes Mn. It was also reported that the decrease in the lattice parameter with the addition of Co should be attributed to the fact that the ionic radius of Co2+ (0.072 nm) is smaller than that of Mn2+ (0.080 nm) for occupying the tetrahedral sites and/or Co3+ (0.068 nm) is smaller than Mn3+ (0.072 nm) for occupying the octa- hedral sites.12,13 As can be seen in Figure 1, we observed a significant broadening and a decrease of the diffraction peak intensities in the XRD pattern of the Ni0.5Co0.8Mn1.7O4 sample. This could be attributed to a decrease in the average crystallite size as given in Table 2 due to the nano-size of the Co3O4 starting powder. Savic et al. reported that the increase in the diffraction peak width and the decrease in the peak intensities in the XRD patterns are associated with a decreasing of the crystallite size and an increasing of the strain.14 Since the desired NTC thermistor properties strongly depend on the densification and grain size, we also inve- stigated the microstructure properties of these samples. The bulk density and grain size of the A1, A2, A6 and A10 samples sintered at 1100 °C were less than the samples sintered at 1200 °C. Smaller grains result in a large number of grain boundaries, which act as scattering centres for the flow of electrons and therefore higher electrical resistivity values were obtained when the sam- ples were sintered at 1100 °C.15 As expected, the in- creasing of the sintering temperature gave rise to an increase in the bulk density and the grain size of these samples, thus the room-temperature resistivity of the samples decreased. In addition, the cation distribution in the octahedral and tetrahedral sites changes with an increasing sintering temperature in the spinel ceramics. The ratio of Mn3+/Mn4+ in the octahedral sites increases with the increasing sintering temperature and also results in a decrease in the resistivity.16 The electrical resistivity of the samples decreased significantly with the increas- ing of the Co3O4 content. Muralidharan et al. observed that the resistivity and B-value decreased with the increasing Co content. Their observation is expected as the Co2+ and Co3+ ions can also occupy the octahedral sites and contribute to the electrical conductivity along with Mn3+/Mn4+ ion pairs in the octahedral sites. This gives rise to a decrease of the resistivity, B-value and temperature coefficient of resistance.2 This phenomenon is prominent for all samples sintered at 1100 °C, while the Co content was increasing in the samples. Similar trends were also observed for the A1, A2 and A6 com- positions when the samples were sintered at 1200 °C. When the sintering temperature was increased from 1100 to 1200 °C for the A10 sample, similar resistivity values were found, but the B-values and activation energy of samples were nearly constant. Moreover, the lattice para- meters were found to be 0.8365 nm and 0.8273 nm for the NiMn2O4 and Ni0.5Co0.8Mn1.7O4 samples, respectively. This may be due to the fact that the hopping distance of the charge carriers becomes easier with the decreasing lattice parameter, thus the resistivity value decreases.17 The sensitivity coefficient and the activation-energy values of all the samples were found in the range –4.426 to –3.683 %/K and 0.282 eV to 0.338 eV, respectively. It is well known that the desired sensitivity coefficient, 25, and the activation energy of the NTC thermistors are in the range –2.2 %/K to –5.5 %/K and 0.1–1.5 eV, respec- tively.18,19 5 CONCLUSION The influence of nano-sized cobalt oxide additions on the structural and electrical properties of nickel-man- ganite-based NTC thermistors was investigated. Our results in this work indicate that a wide range of elec- trical properties of nickel-manganite-based NTC ther- mistors can be obtained by the addition of nano-sized cobalt oxide. The particularly interesting finding in this study demonstrated that the Ni0.5Co1.1Mn1.4O sample sintered at 1200 °C for 5 h has a low electrical resistivity and a high B-constant. Acknowledgements This study is supported by TÜBÝTAK (The Scientific and Technical Research Council of Turkey), Project number 3001-114M860. We would like to thank TÜBÝTAK for its financial support. 6 REFERENCES 1 R. N. Jadhav, S. N. Mathad, V. 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Ding, Preparation and cha- racterization of Ni0.6Mn2.4O4 NTC ceramics by solid-state coordina- tion reaction, Journal of Materials Science: Materials in Electronics, 24 (2013), 5183–5188, doi:10.1007/s10854-013-1542-2 18 A. Feteira, Negative Temperature Coefficient Resistance (NTCR) Ceramic Thermistors: An Industrial Perspective, Journal of the Ame- rican Ceramic Society, 92 (2009), 967–983, doi:10.1111/j.1551- 2916.2009.02990.x 19 K. Park, I. H. Han, Effect of Cr2O3 addition on the microstructure and electrical properties of Mn-Ni-Co oxides NTC thermistors, Jour- nal of Electroceramics, 17 (2006), 1069–1073, doi:10.1007/s10832- 006-8317-6 G. HARDAL, B. Y. PRICE: INFLUENCE OF NANO-SIZED COBALT OXIDE ADDITIONS ON THE STRUCTURAL ... 928 Materiali in tehnologije / Materials and technology 50 (2016) 6, 923–928 R. GOPALAKRISHNAN, K. CHINNARAJU: DURABILITY OF ALUMINA SILICATE CONCRETE BASED ... 929–937 DURABILITY OF ALUMINA SILICATE CONCRETE BASED ON SLAG/FLY-ASH BLENDS AGAINST ACID AND CHLORIDE ENVIRONMENTS ZDR@LJIVOST BETONA NA OSNOVI GLINICE IN SILIKATOV IZ ME[ANICE @LINDRA/LETE^I PEPEL NA KISLO IN KLORIDNO OKOLJE Rajagopalan Gopalakrishnan1, Komarasamy Chinnaraju2 1Sri Venkateswara College of Engineering, Department of Civil Engineering, 602 117 Sriperumbudur, India 2Anna University, Division of Civil and Structural Engineering, 600 025 Chennai, India gopalakrishnan@svce.ac.in, rajagopalan.gopalakrishnan0@gmail.com Prejem rokopisa – received: 2015-07-19; sprejem za objavo – accepted for publication: 2015-12-02 doi:10.17222/mit.2015.230 The durability of a concrete mainly depends on its resistance against acid and chloride environments. This article presents an investigation of the durability of geopolymer concrete with GBFS (Granulated Blast Furnace slag), Fly ash (class F) and alkaline activators when exposed to 5 % sulphuric acid and chloride solutions. GBFS was replaced by fly ash with different replacement levels from 0 % to 50 % in a constant concentration of 12-M alkaline activator solution. The main parameters of this study are the evaluation of the change in weight, strength and microstructural changes. The degradation was studied using Scanning Electron Microscopy (SEM) with EDAX. From the test results it is observed that the strength of the geopolymer concrete with GBFS in ambient curing performs compared well to geopolymer concrete with GBFS blended with fly ash. The acid resistance in terms of the rate of reduction of strength of GPC with GBFS is 85 %, while for 40 % replacement of fly ash to GBFS performs well with a reduction of only 53 %. Similar observations are also observed in a chloride environment in which 40 % replacement of fly ash to GBFS performs well when compared to GPC with GBFS. Hence, geopolymer concrete with 40 % replacement of fly ash for GBFS is the appropriate level of replacement, satisfying the above durability properties. Keywords: durability, geopolymer concrete, acid and chloride environment Zdr`ljivost betona je predvsem odvisna od odpornosti na kislo in kloridno okolje. ^lanek predstavlja preiskavo zdr`ljivosti geopolimernega betona z GBFS (granulirana `lindra iz plav`a), lete~ega pepela (razred F) in alkalnih aktivatorjev med izpostavitvijo 5 % `vepleni kislini in kloridnim raztopinam. V GBFS je bila dodana razli~na koli~ina: od 0 % do 50 % lete~ega pepela pri konstantni koncentraciji 12 M raztopine alkalnega aktivatorja. Glavni parametri v {tudiji so bili sprememba te`e, trdnost in mikrostrukturne spremembe. Degradacija je bila preu~evana z uporabo vrsti~nega elektronskega mikroskopa (SEM) z EDAX. Iz rezultatov preizkusov je opaziti, da je pri izpostavitvi trdnost geopolimernega betona z GBFS dobra, v primerjavi z geopolimernim betonom z GBFS s prime{anim lete~im pepelom. Odpornost na kislino, izra`eno s hitrostjo zmanj{evanja trdnosti GPC z GBFS je 85 %, medtem ko se pri 40 % nadomestitvi lete~ega pepela v GBFS, ta pona{a dobro, s samo 53 % zmanj{anjem trdnosti. Podobna opa`anja so bila tudi v kloridnem okolju, v katerem se 40 % nadomestilo lete~ega pepela v GBFS obna{a dobro, v primerjavi z GPC, ki vsebuje tudi GBFS. Torej je geopolimerni beton, s 40 % nadomestitvijo lete~ega pepela z GBFS, primeren za zgoraj omenjeno zdr`ljivost. Klju~ne besede: zdr`ljivost, geopolimerni beton, kislo in kloridno okolje 1 INTRODUCTION The durability of concrete structures, especially those built in corrosive environments, starts to deteriorate after 20 to 30 years, even though they have been designed for more than 50 years of service life. Although the use of Portland cement is unavoidable in the foreseeable future, many efforts are being made to reduce the use of Port- land cement in concrete.1 Inorganic polymer concretes, or geopolymers have been emerging as novel engineer- ing materials with the potential to form a alternative element for the construction industry.2–4 Geopolymers show substantially superior resistance to fire5 and acid attack6 and much less shrinkage than OPC Concrete. The tensile strength of geopolymer concrete falls within the range observed for OPC-based concrete. Also, the fle- xural strengths are generally higher than the standard model line for OPC-based concrete. This favourable behaviour can be attributed to the type of matrix forma- tion in the geopolymer concrete.7 It has been reported that the stress strain relationship of fly-ash-based geo- polymer concrete is almost similar to that of ordinary portland cement concrete.1 These advantages make the geopolymer concrete a strong alternative for replacing ordinary Portland cement concrete. Geopolymers are produced by a polymerization reac- tion of strong alkali liquids such as sodium hydroxide (NaOH), potassium hydroxide (KOH), sodium silicate and potassium silicate with a source material of geolo- gical origin or by a product material such as fly ash, GBFS, metakaolin. The mixture can be cured at room temperature or heat cured. Under a strong alkali solution, an alumina silicate material dissolves and forms SiO4 and AlO4 tetrahydral units. Materiali in tehnologije / Materials and technology 50 (2016) 6, 929–937 929 UDK 67.017:625.821.5:620.1 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)929(2016) Three common types of geopolymer are the poly- sialate Al-O-Si chain, polysialate siloxo Al-O-Si-Si chain and polysialate disiloxo Al-O-Si-Si-Si chain.8,9 The raw materials commonly used for preparing geopolymers are clay and metakaolin. Studies are under progress re- cently are using the waste and byproducts for geopoly- merization from waste materials.10–18 A number of research publications related to geopolymers have been published, with some reports on chemical composition or reaction processes, others relating to mechanical pro- perties and durability. The compressive strength depends on both the Si/Al ratio and the type of raw materials used.19–22 To improve the performance of these binders, a number of recent investigations have been published, giving attention to producing mixes based on blends with reactive precursors. The blends commonly involve a Ca-rich precursor such as granulated blast furnace slag (GBFS), and an alumino silicate source such as low cal- cium fly ash or metakaolin, to form the stable coexist- ence of calcium silicate hydrate (C-S-H) gels formed from the activation of GBFS and geopolymer gel (also expressed as N-A-S-H) produced from the activation of alumina silicate, which is a cementitious paste that improves the setting and strength properties.23–25 Fly ash contains mainly alumina and silica, along with other impurities like iron oxide, lime and magnesia. Due to increased industrial growth fly ash is generated in huge amounts and its accumulation over the years has become a threat to the environment. The utilization of fly ash for preparing geopolymers not only conserves the nature, but also reduces the ever-increasing burden of fly ash on the environment. GBFS is a glassy granular material essentially consisting of oxides like CaO, SiO2 and Al2O3. It is formed when molten blast-furnace slag, a byproduct in the extraction of iron is cooled, usually by immersion in water and then ground fine to improve its reactivity.18,26–30 There is a relatively small number of re- search reports expressing the structure and performance of alkali-activated GBFS/Fly ash blends cured at ambient environment, and have been mainly discussed where fly ash is added to GBFS to enhance the strength and micro- structure, which leads to a good durability. The durabi- lity of these binders in an acid and chloride environment was not investigated before; however, there is an opinion that geopolymer materials have excellent resistance in acid and chloride environments. The above resistance against acid and chloride is an important durability pro- perty concerned with serviceability for geopolymer materials used in the construction industry. T. Bakerev31 studied the resistance of geopolymer materials prepared from fly ash against 5 % sulphuric acid up to 5 months exposure and concluded that geopolymer materials have better resistance than ordinary cement concretes. Port- land cement and blended cement concretes show a dete- rioration when exposed to acid and chloride environ- ments. The demand of standard methods to evaluate the performance of cements in acid environments has led to research in different exposure conditions by various re- searchers making it difficult to correlate the results.32 This paper presents an investigation of acid and chloride attack on geopolymer materials prepared using GBFS blended with low calcium fly ash in different per- centages and sodium hydroxide, sodium silicate as acti- vators and cured in ambient conditions (25±5 °C). The effects of the fly ash addition to GBFS, weight change, visual appearance, microstructure and strength properties have also been studied. To study the microstructure methods such as SEM with EDAX have been employed. An attempt has been made to correlate the structure with reaction and properties. 2 EXPERIMENTAL PART 2.1 Materials The class F fly ash (as per ASTM C618-99) obtained from Ennore power plant and GBFS obtained from M/s Jindal, Karnataka were used for the study. The chemical analysis of GBFS and fly ash were made using the XR fluorescence method and the results were shown in Table 1. Coarse aggregate of maximum 20 mm with a specific gravity of 2.67 was used. Locally available river sand confirming to Zone II (as per IS 383) with a spe- cific gravity of 2.52 was used for the study. Sodium hydroxide in the form of flakes having a purity of 90 % and sodium silicate in the liquid having a chemical com- position of Na2O = 14.7 % SiO2 = 29.4 % H2O = 55.9 % by mass. To improve the workability of fresh concrete a superplasticizer Glenium supplied by BASF, a polycar- boxylic ether is added with the ingredients. 2.2 Test variables Fly ash with GBFS of various mixture proportions were subjected to geopolymerization. However the ratio of SiO2 to Al2O3 is maintained at approximately 2, which is a typical ratio for a geopolymer structure. Details of the batch compositions are given in Table 2. The ratio of sodium silicate solution to sodium hydroxide by mass was kept as 2.5, the ratio of alkaline liquid to the geo- polymer solids was kept as 0.4 and water to geopolymer R. GOPALAKRISHNAN, K. CHINNARAJU: DURABILITY OF ALUMINA SILICATE CONCRETE BASED ... 930 Materiali in tehnologije / Materials and technology 50 (2016) 6, 929–937 Table 1: Raw materials chemical properties Tabela 1: Kemijske lastnosti sestavnih materialov Materials Chemical composition, in mass fractions (w/%) SiO2 Al2O3 CaO Fe2O3 MgO SO3 Na2O K2O LOI GBFS 34.60 17.40 33.01 1.50 8.70 0.05 1.25 0.83 1.39 Fly ash 53.80 21.20 0.90 17.00 3.50 1.50 – – 0.48 solids as 0.24. The properties of various concrete mixes are shown in Table 3. 2.3 Sample preparation for physical testing Solutions of NaOH (12-M concentration) and Na2SiO3 were separately prepared 24 h before casting. Both the solutions were mixed together at the time of mixing. A weighed quantity of GBFS, fly ash, fine aggregate and coarse aggregate were dry mixed in a pan for about 3 min to 5 min. Wet mixing was done for another 3 mins and the required quantity of super plasti- cizer and water was added to obtain the required consis- tency. The samples were then cast into the steel moulds of size 100 mm × 100 mm × 100 mm. Compaction was done by manual strokes, followed by a compaction on a vibrating table for 20 s. The cubes were remoulded after 24 h and cured at a relative humidity of 25±5 °C to pre- vent drying effects. The required number of samples for each mix was prepared and cured under ambient con- ditions and were reported as the mean of the three samples. 2.4 Test procedure The resistance of the materials to acid and chloride attack was studied by immersion of the cubical speci- mens of size (100 mm × 100 mm × 100 mm) in a 5 % solution of concentrated sulphuric acid for acid attack and with a proportion of 4 % NaCl with 1 % magnesium sulphate solution for chloride attack for a period up to 180 d. The compressive strength was determined before the test and after (28, 60,120 and 180) d of exposure. The choice of acid solution and its concentration was based on the practical application of concrete as a construction material mainly in the sewage pipe and mining industries. The volume of solution was kept not less than four times the volume of the specimens immersed and maintained throughout the test period. The testing solutions were replaced with new solutions after 30 d until the completion of the test period. The samples were compared with all the grades of concrete that were ambient cured. The deterioration of samples was studied by SEM with EDAX. For this testing, the samples were taken from the surface at a 0–5-mm depth, exposed to the solutions of 120 d and compared with conventional ambient cured samples. The SEM analysis was done using a microscope having a magnification of 5× to 3,00,000× with a voltage of 0.3–30 kV. The coating of the samples for the analysis was done using an ion sputter with a gold target and the system was attached with the latest PIV. The resolution of the equipment varied from 3 nm, 4 nm to 10 nm. The change in mass before and after the immersion was observed for all the samples. 3 RESULTS AND DISCUSSION 3.1 Visual appearance Visual appearances of geopolymer specimens after immersion in a solution of concentrated sulphuric acid after 180 d are shown in the Figure 1. Its appearance R. GOPALAKRISHNAN, K. CHINNARAJU: DURABILITY OF ALUMINA SILICATE CONCRETE BASED ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 929–937 931 Table 2: Batch composition Tabela 2: Sestava posamezne serije Test variables Batch composition ratio GBFS & fly ash, in mass fractions (w/%) SiO2 / Al2O3 ratio CaO % GBFS Fly ash GPCA 100 0 1.99 33.01 GPCB 90 10 2.05 29.80 GPCC 80 20 2.12 26.59 GPCD 70 30 2.18 23.38 GPCE 60 40 2.23 20.17 GPCF 50 50 2.28 16.96 Table 3: Designation of test variables and mix proportions Tabela 3: Oznaka preizkusnih spremenljivk in razmerja me{anic Test variables ID in different environment Binder composition Ingredient contents kg/m 3 Ambient curing Sulphuric acid Sodium chloride GBFS % Fly ash % GBFS Fly ash Sodium hydroxide solution Sodium silicate solution Coarse aggregate Fine aggregate GPCA GPCAA GPCAC 100 0 400 0 46 115 1200 645 GPCB GPCBA GPCBC 90 10 360 40 46 115 1200 645 GPCC GPCCA GPCCC 80 20 320 80 46 115 1200 645 GPCD GPCDA GPCDC 70 30 280 120 46 115 1200 645 GPCE GPCEA GPCEC 60 40 240 160 46 115 1200 645 GPCF GPCFA GPCFC 50 50 200 200 46 115 1200 645 seems to be very slightly changed after 28 d, but there is a distinct change in the appearance of the deterioration after 180 d. The surface became softer as the duration of the test period prolonged, but could not be scratched with finger nails. The deterioration of the surface in- creased with the duration, but the amount of deteriora- tion could not be determined through a visual inspec- tion.31 The geopolymer specimens immersed in chloride solution did not reveal any changes in the surface after 28 d. Even after 180 d also, there is no severe deteriora- tion. 3.2 Change in weight Table 4 and Figure 2 give the comparative weight changes for the specimens exposed to sulphuric acid and chloride solution with the ambient cured samples after 180 d. In ambient cured samples there is a loss of weight with the replacement of fly ash to GBFS. The percentage of loss increases with the percentage of replacement. The loss of weight at 10 % replacement of fly ash to GBFS is 1.512 % and 1.678 % with 50 % replacement (GBFC). But at 40 % replacement there is loss of weight of only 1.369 %, which shows the optimum replacement of fly ash. This may be due to the limitations of the SiO2/Al2O3 ratio. The geopolymer samples (GPCA) immersed in sulphuric acid solution show a little loss of weight 0.05 %. The samples with the replacement of fly ash to GBFS show a loss of 0.43 % at 10 % with a gain of 0.88 % at 40 % replacement. Similar observations have been re- ported by31. The mass change was calculated according to ASTM C267. All the geopolymer concrete mixes show a very low mass loss of less than 3 %. Table 4: Weight change in % ambient curing with NaCl & sulphuric acid after 180 d Tabela 4: Sprememba te`e v % po 180 dneh izpostavitve okolju z NaCl in `vepleno kislino Test variables Id Weight change, % GPCA - GPCAC-GPCAA -1.32 0.35 -0.05 GPCB - GPCBC-GPCBA -1.512 0.49 -0.43 GPCC - GPCCC-GPCCA -1.622 1.05 -0.73 GPCD - GPCDC-GPCDA -1.715 0.16 0.69 GPCE - GPCEC-GPCEA -1.369 0.08 0.88 GPCF- GPCFC-GPCFA -1.678 0.08 -0.12 The geopolymer samples with GBFS (GPCA) immersed in chloride solution gain weight to 0.35 %. When it is replaced by fly ash to GBFS it gains weight, which varied from 0.35 % at 10 % replacement to a very low gain of 0.08 % at 50 % replacement. The gain in % increases with the increase of the replacement. But it shows a low value in 40 % and 50 % of 0.08 %. A mini- mal change in nominal weight loss has been observed with 40 % fly ash in the geopolymer composite, which indicates that 40 % fly ash composite with geopolymer performs the best, compared to all the other compo- sitions. The specimens were damaged beyond this 40 % fly ash and longer durations of immersion, which is in good agreement with33. Interaction of geopolymer in the acid solution may result in replacement of exchangeable cations such as Na+ in the polymer by hydrogen ion or hydronium ion.31 3.3 Compressive strength Figure 3 shows the variation of compressive strength of samples for different duration periods in ambient curing. The compressive strength increases with time in all the mixes. The geopolymer concrete with 100 % GBFS (GPC-A) shows a higher compressive strength of 74 MPa at 180 d. Its strength varies from 37 MPa at 3 d to 74 MPa at 180 d. The increase in percentage was 100, which is in good agreement with S. A. Bernal.34 The geopolymer concrete GBFS blended with fly ash at 10 % replacement (GPC-B) strength varies from 32.4 MPa at 3 R. GOPALAKRISHNAN, K. CHINNARAJU: DURABILITY OF ALUMINA SILICATE CONCRETE BASED ... 932 Materiali in tehnologije / Materials and technology 50 (2016) 6, 929–937 Figure 1: GPC samples after 180 d in sulphuric acid Slika 1: Vzorci GPC po 180 dneh v `vepleni kislini Figure 3: Compressive strength of GPC Slika 3: Tla~na trdnost GPC Figure 2: Weight change in % ambient curing with NaCl & H2 SO4 after 180 d Slika 2: Spreminjanje te`e v % pri 180 dnevni izpostavitvi okolju z NaCl in H2SO4 d to 69.8 MPa at 180 d. Similarly, other replacement levels 20 %, 30 % (GPC-C, GPC-D) strength varies 30.5, 29.6 MPa at 3 d to 67, 66 MPa at 180 d, respectively. In the above, replacement levels of fly ash to GBFS the percentage increase is 120 %. But at 40 % and 50 % (GPC-E, GPC-F) replacement levels its strength at 180 d was 44 MPa and 43 MPa, respectively, with a percentage increase of 73 %. The rate of development of strength at 180 d with reference to 28 d is 150 % in all the mixes, except in the mixes of GPC-E, GPC-F, which show only 100 %, as in Figure 4. 3.4 Effect of sulphuric acid and chloride Figure 5 show the evaluation of compressive strength for different duration periods for the samples immersed in a solution sodium chloride. It reveals that the reduc- tion of strength is more with an increase in the percent- age of replacement of fly ash to GBFS in the geopolymer concrete. The strength reduction rate increases with the duration period in all the mixes. The reduction rate from 28 d to 60 d is more and it is further increased at 120 d. The reduction of strength of geopolymer concrete with 100 % GBFS (GPC-A) is 42 % compared to the ambient cured samples at 180 d. There is a minimum reduction of strength from 120 d to 180 d. The GPC blended with fly ash at 10 % replacement (GPC-B) shows a reduction of strength 40 % at 180 d. Similarly, other replacement le- vels 20 %, 30 % (GPC-C, GPC-D) show a strength reduction of 47 % and 45 % respectively at 180 d. In the replacement of fly ash to GBFS, 40 % replacement performs well and shows a reduction rate of 24 %. This shows that GBFS can be replaced by 40 % of fly ash as the reduction rate is less compared to GPC with 100 % GBFS. The reactivity of fly ash in the chloride environ- ment is good. The detailed loss or gain in % is shown in Table 5 and Figure 6. Table 5: Strength change in % ambient curing with NaCl Tabela 5: Spreminjanje trdnosti v % pri utrjevanju v okolju z NaCl Days GPCA/GPCAC GPCB/ GPCBC GPCC/ GPCCC GPCD/ GPCDC GPCE/ GPCEC GPCF/ GPCFC 28 -9.2 -3.54 -5.2 2.8 -4.3 -6 60 -23.3 -31.2 -33.6 -17 -20 -50.4 120 -26.7 -33.1 -36.5 -31.2 -13.9 -29.8 180 -42 -40 -47.4 -45.7 -24 -31.86 Figure 7 shows the evaluation of compressive strength for the different duration periods for the test samples immersed in 5 % solution of sulphuric acid. The reduction rate continuously increases with the duration period. There is a strength reduction of 85 % compared to ambient cured samples for the geopolymer concrete prepared with 100 % GBFS at 180 d. The GPC blended with fly ash at 10 % replacement (GPC-B) shows an 83 % reduction of strength at 180 d. Similarly, other replacement levels 20 %, 30 % (GPC-C, GPC-D) show a strength reduction of 81 % and 77 %, respectively, at 180 d. In the replacement of fly ash to GBFS, 40 % replacement performs well and it shows a reduction rate of 53 %. This shows that the replacement of 40 % of fly ash is the optimum. The strength reduction is propor- tional to the replacement of fly ash. The detailed loss or gain in % is shown in Table 6 and in Figure 8. Com- R. GOPALAKRISHNAN, K. CHINNARAJU: DURABILITY OF ALUMINA SILICATE CONCRETE BASED ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 929–937 933 Figure 7: Compressive strength comparison of ambient curing with sulphuric acid Slika 7: Primerjava tla~ne trdnosti po utrjevanju v okolju z `vepleno kislino Figure 5: Compressive strength comparison of ambient curing with NaCl Slika 5: Primerjava tla~ne trdnosti pri utrjevanju v okolju z NaCl Figure 6: Strength change in % ambient curing with NaCl Slika 6: Spreminjanje trdnosti v % pri utrjevanju v okolju z NaCl Figure 4: Compressive strength development Slika 4: Razvoj tla~ne trdnosti pressive strength appears to increase for a set of samples with curing time for ambient curing, whereas for the set cured in the NaCl solution it appears to decrease. This improvement in compressive strength is attributed to the leaching of silica and aluminium at a higher Ca of NaOH.36 Table 6: Strength change in % ambient curing with sulphuric acid Tabela 6: Spreminjanje trdnosti v % pri utrjevanju v okolju z `veple- no kislino Days GPCA /GPCAA GPCB / GPCBA GPCC / GPCCA GPCD / GPCDA GPCE / GPCEA GPCF/ GPCFA 28 -25.1 -19.1 -27.4 -29.1 -30.5 -31.7 60 -56.4 -39.6 -44.1 -44.7 -42.3 -37.1 120 -71.2 -66.6 -66.8 -56.2 -33.5 -28.5 180 -85.8 -83.7 -81.7 -77 -53.9 -58.4 The development of the strength of specimens is suppressed due to the hydration of Ca-Al- silicates from the dissolution of the CAS phase by the hydroxyl ion (OH) contributed by water and aqueous NaOH so that the Al and Si become penta-valent due to the attachment of (OH). Due to this, Al–O–Al and Si–O–Si tend to R. GOPALAKRISHNAN, K. CHINNARAJU: DURABILITY OF ALUMINA SILICATE CONCRETE BASED ... 934 Materiali in tehnologije / Materials and technology 50 (2016) 6, 929–937 Figure 9: SEM images of: a) GBFS and b) fly ash Slika 9: SEM-posnetek: a) GBFS in b) lete~i pepel Figure 11: SEM images after immersion in solution of NaCl for 120 d: a) GPCAC, b) GPCCC, c) GPCEC Slika 11: SEM-posnetki po 120 dnevnem namakanju v raztopini NaCl: a) GPCAC, b) GPCCC, c) GPCEC Figure 8: Strength change in % ambient curing with sulphuric acid Slika 8: Spreminjanje trdnosti v % pri utrjevanju v okolju z `vepleno kislino Figure 10: SEM images before immersion: a) GPCA, b) GPCC, c) GPCE Slika 10: SEM-posnetki pred namakanjem: a) GPCA, b) GPCC, c) GPCE strengthen the bands. This can be schematically represented as: Ca – Si – (o-) o – Al – Ca +20H (H6) Ca – Si – o – Al – Ca – OH35 3.5 SEM with EDAX Figure 9 shows the SEM images of GBFS and fly ash. Figure 10 shows the SEM images of the specimens before immersion into the solutions of sulphuric acid and chloride. It is evident that the appearance of gel-like R. GOPALAKRISHNAN, K. CHINNARAJU: DURABILITY OF ALUMINA SILICATE CONCRETE BASED ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 929–937 935 Figure 12: SEM images after immersion in solution of H2SO4 for 120 d: a) GPCAA, b) GPCCA, c) GPCEA Slika 12: SEM-posnetki po 120 dnevnem namakanju v raztopini H2SO4: a) GPCAA, b) GPCCA, c) GPCEA Figure 13: EDAX spectrum before immersion: a) GPCA, b) GPCC, c) GPCE Slika 13: EDAX-spekter pred namakanjem: a) GPCA, b) GPCC, c) GPCE Figure 14: EDAX spectrum after immersion in solution of NaCl for 120 d: a) GPCA, b) GPCCC, c) GPCEC Slika 14: EDAX-spekter po 120 dnevnem namakanju v NaCl: a) GPCA, b) GPCCC, c) GPCEC phases in the microstructure (of the SEM) can be attributed to the development of the microstructure particularly in the GBFS phases.6 The immersion in NaCl for the 120 d period appears to cause a decrease in Na content of the samples, which may be due to the migration of Na+ ions from the speci- men samples to NaCl. The microstructures of the speci- men samples after 120 d of immersion in a solution of NaCl were studied using SEM and the results are shown in Figure 11. SEM images of the samples after immersion in H2SO4 for 120 d (Figure 12) indicate that there is a for- mation of light-coloured precipitates in the sample after immersion in H2SO4 solution. The formation of light- coloured precipitates is indicative of the degradation of the cured specimen.36 The appearance of lightly coloured precipitates in a low distribution may be attributed to a more amorphous, less-crystalline layer formation.6 A comparison of the EDAX patterns of the samples before immersion and also after immersion in the NaCl and H2SO4 media indicates the following: The parent samples as prepared contain Magnesium as is evident from Figure 13, but this magnesium content is retained even after immersion in NaCl, as evident from the Figure 14 after immersion. The magnesium content of the samples was found to be lost during immersion in the solution of H2SO4 after 120 d, which is shown in Figure 15. This fall is in the magnesium content of the samples in H2SO4. Immersion may be attributed to the migration of Mg2+ from the sam- ples to the H2SO4 medium forming MgSO4 (soluble). The results are good agreement with previous litera- tures.36 4 CONCLUSION From the experimental investigation the following conclusions can be drawn.There is an improvement in strength with respect to the age observed for a maximum period of 180 d in all the mixes. The geopolymer con- crete with GBFS (100 %) performed better than the other mixes with fly ash and GBFS combination during am- bient curing. A geopolymer concrete mix with 40 % replacement of fly ash to GBFS performed well in chlo- ride environment with a reduction of 24 % in comparison to the geopolymer concrete with GBFS (100 %) having a reduction rate of 42 %. The geopolymer concrete mix with 40 % replacement of fly ash to GBFS performed well in the acid environment with a reduction rate of 53 % in comparison to geopolymer concrete prepared with GBFS (100 %) has a reduction rate of 85 %. In general, the reduction rate increases with the increase in replacement levels of fly ash to GBFS in the geopolymer concrete in both the chloride and acid environment. The SEM and EDAX images after 120 d of immersion in acid and chloride environment show the deterioration, same as that of previous researchers. The geopolymer concrete prepared with GBFS can be replaced by fly ash (40 %) as it performs equally well and satisfies all the durability properties in both the chloride and acid environments. Acknowledgement The authors are very thankful to the management of Sri Venkateswara college of Engineering, Sri Perumbu- dur and College of Engineering Guindy, Anna Univer- sity, Chennai for providing facilities to carry out the work. 5 REFERENCES 1 B. V. 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KRACÍK et al.: THE SIZE EFFECT OF HEAT-TRANSFER SURFACES ON BOILING 939–944 THE SIZE EFFECT OF HEAT-TRANSFER SURFACES ON BOILING VPLIV VELIKOSTI POVR[IN, KI PRENA[AJO TOPLOTO NA VRENJE Petr Kracík, Marek Balas, Martin Lisy, Jiøí Pospí{il Brno University of Technology, Institute of Power Engineering, Faculty of Mechanical Engineering, Technická 2896/2, 616 69 Brno, Czech Republic kracik@fme.vutbr.cz Prejem rokopisa – received: 2015-07-31; sprejem za objavo – accepted for publication: 2015-12-14 doi:10.17222/mit.2015.245 A sprinkled tube bundle is frequently used in technology processes where an increase or decrease of a liquid temperature in a very low-pressure environment is required. Phase transitions of the liquid very often occur at low temperatures at pressures ranging in the thousands of pascals, which enhances the heat transfer. This paper focuses on the issue of a heat-transfer coefficient that is experimentally examined at the surface of a tube bundle. The tube is located in a low-pressure chamber where the vacuum is generated using an exhauster via an ejector. The tube consists of smooth copper tubes of 12 mm diameter placed horizontally one above another. Heating water flows in the bundle from the bottom towards the top at an average input tempe- rature of approximately 40 °C and an average flow rate of approximately 7.2 L min–1. A falling film liquid at an initial tempe- rature of approximately 15 °C at an initial tested pressure of approximately 97 kPa (atmospheric pressure) is sprinkled onto the tubes’ surface. Afterwards, the pressure in the chamber is gradually decreased. When reaching the minimum pressure of approximately 3 kPa (abs) the water partially evaporates at the lower part of the bundle. Consequently, the influence of the falling film liquid temperature increase is tested. This gradually leads to the boiling of water in a significant part of the bundle and the residual cooling liquid that drops back to the bottom of the vessel is almost not heated anymore. In this paper we present the influences of the size of the heat-transfer surfaces. Keywords: sprinkled tube bundle, water, under pressure, heat transfer Pr{enje po snopu cevi se pogosto uporablja v tehnolo{kih procesih, kjer se zahteva povi{anje ali zmanj{anje temperature teko~ine v okolju z nizkim tlakom. Fazni prehod teko~ine se pogosto pojavi pri nizkih temperaturah in tlakih v obmo~ju nekaj tiso~ paskalov, kar vpliva na prenos toplote. ^lanek se nana{a na koeficient prenosa toplote, ki je eksperimentalno dolo~en na povr{ini snopa cevi. Cev je name{~ena v nizkotla~ni komori, kjer se vakuum ustvarja s pomo~jo aspiratorja preko ejektorja. Snop cevi sestavljajo gladke bakrene cevi, premera 12 mm, ki so name{~ene horizontalno ena nad drugo. Voda za ogrevanje te~e v snop od spodaj proti vrhu s povpre~no temperaturo okrog 40 °C in povpre~no hitrostjo pretoka okrog 7,2 L min–1. Padajo~ vodni film z za~etno temperaturo okrog 15 °C in z za~etnim tlakom okrog 97 kPa (atmosferski tlak) pr{i po povr{ini cevi. Tlak se v komori postopno zni`uje. Ko dose`e minimalni tlak okrog 3 kPa (absolutni tlak) voda delno izhlapi na spodnjem delu snopa. Posledi~no je preizku{en vpliv nara{~anja temperature padajo~e vode. To postopno privede do vrenja vode na ve~jem delu povr{ine snopa in preostala hladilna teko~ina, ki kaplja na dno posode, se skoraj ne segreje ve~. ^lanek predstavlja vpliv velikosti povr{ine kjer se prena{a toplota. Klju~ne besede: pr{enje po snopu cevi, podtlak, prenos toplote 1 INTRODUCTION Due to a decreasing supply of fossil fuels and their increasing price, the minimization of energy consump- tion becomes an important priority, followed by saving the primary fuel entering energy processes that are supposed to achieve the maximum efficiency possible, and last but not least, using so-called renewable sources of energy. Among these renewable sources, a biomass combustion technology is mainly used to generate ther- mal energy and electricity in the Czech Republic. Current research aims to reflect these demands well. For instance, research is conducted in the field of optimi- zation of technology for wood mass preparation1,2 before combustion or further utilization for pellets’ production. At the Department of Power Engineering long-term re- search focuses, apart from other things, on the utilization of waste thermal energy, which is found, for example, in condensers at large energy units, for the possible gene- ration of cool in absorption units. One of the basic elements of an absorption circuit is an evaporator, inside of which the heat-carrier substance is sprayed onto a tube bundle. Due to a low pressure environment inside the container where the bundle is located the falling film liquid at the tube bundle boils. The heat necessary for boiling is derived from a cooled substance flowing inside the tube bundle. Under ideal conditions water boils at the whole surface of an exchanger, but in practice it must be con- sidered that in original spots of contact between the water and the exchanger wall the water will not boil at the tubes’ surface, but the cooling liquid will merely be heated-up. This paper deals with this very situation – heat transfer behaviour when heating a sprinkled water film that boils in a low-pressure environment for a real tube bundle. Materiali in tehnologije / Materials and technology 50 (2016) 6, 939–944 939 UDK 539.4.016:621.785:620.195 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)939(2016) Research in the area of sprinkled exchangers can be divided into two major parts. The first part is research on heat transfer and a determination of the heat-transfer coefficient with respect to sprinkled tube bundles for various liquids, whether boiling or not. For water as the falling film liquid, they were, for example3–7. The second part is testing of the sprinkle modes for various tube dia- meters, tube pitches and tube materials and a determi- nation of individual modes’ interface. This area is mainly researched in8–10. All the results published so far for water as the falling film liquid apply to one to three tubes, for which the mentioned relations studied are determined in rigid labo- ratory conditions, defined strictly in advance. The sprinkled tubes were not viewed from the operational perspective where there are more tubes and various mo- des can occur in different parts with various heat-transfer values. 2 EXPERIMENTAL PART For the purposes of examination of the heat transfer for sprinkled tube bundles a test apparatus was con- structed (Figure 1). A tube bundle at which a heat transfer from a heated water flowing inside the tubes into a falling film liquid is studied is placed in a vessel where the low pressure is created by an exhauster through an ejector. The test apparatus chamber is a cylindrical vessel with a length of 1.2 m with three apertures in which the tube bundle of an examined length of 940.0 mm is in- serted. The tube bundle is installed in two fitting metal sheets which define the sprinkled area. The bundle con- sists of eight copper tubes of diameter 12.0 mm situated horizontally, one above another, with a distribution tube above them, with apertures of the diameter from 1.0 mm to 9.2 mm. The bundle can be operated using only the first four or six tubes too. Two closed loops are connected to the chamber. A heating one and a sprinkling one. The heating liquid flowing inside the tubes is intended for an overpressure up to 1.0 MPa. The second loop contains flowing falling film liquid. There is a pump, a regulation valve, a flow meter and plate heat exchangers attached to both loops. The plate heat exchanger at the heating loop is connected to a gas boiler, which supplies heat to the heating liquid. The sprinkling loop uses two plate heat exchangers. In the first one the falling film liquid is cooled by cold drinking water from the water mains and the falling film liquid is cooled in the second exchanger by drinking water cooled in a cooler, which regulates the temperature up to 1.0 °C. In order to enable visual control, the heat- ing loop also includes a manometer and a thermometer. The thermal status in individual loops is measured by wrapped unearthed T-type thermocouples on the agents’ input and output from the vessel. All the thermocouples were calibrated in the CL1000 Series calibration furnace, which maintains a given temperature with an accuracy of ±0.15 °C. None of the thermocouples exceeded the error ±0.5 °C within the studied range from 28 °C to 75 °C. That is why the total error for the temperature measurement is set uniformly for all the thermocouples along the whole studied range ±0.65 °C. There are three vacuum gauges measuring the low pressure. The first vacuum gauge is designed for visual control and is a mercury meter, the second one is a P. KRACÍK et al.: THE SIZE EFFECT OF HEAT-TRANSFER SURFACES ON BOILING 940 Materiali in tehnologije / Materials and technology 50 (2016) 6, 939–944 Figure 1: Measurement apparatus diagram Slika 1: Shema merilnih naprav digital vacuum gauge Baumer TED6 and enables measurements within the whole desired low-pressure range, but it is less accurate with lower pressure values. To allow a precise measuring of the low spectrum, a third digital vacuum gauge in the range 2.0 kPa to 0 Pa is used. The accuracy of a vacuum gauge, the results of which have been used for the assessment, is 0.5 % from the measured range, i.e., ±0.5 kPa. Electromagnetic flow meters Flomag 3000 attached to both loops measure the flow rate. The flow meters’ range is 0.0078 to 0.9424 L s–1, where the accuracy is 0.5 % from the measured range, i.e. ±0.00467 L s–1. All the examined quantities are either directly (thermo- couples) or via transducers scanned by measuring cards DAQ 56 with a frequency of 0.703 Hz. 3 METHODOLOGY OF THE DATA ASSESSMENT The assessment of the measured data is based on the thermal balance between the operating liquid circulating inside the tubes and a sprinkling loop according to the law of conservation of energy. Heat transfer is realized by convection, conduction and radiation. At lower tem- peratures the heat transferred by radiation is negligible; therefore, it is excluded from further calculations. The calculation of the studied heat-transfer coefficient is based on Newton’s heat-transfer law and Fourier’s heat-conduction law that have been used to form the following relation in Equation (1):   0 0 0 1 2 1 1 2 1 2 = − − ⋅ ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ ⎡ ⎣⎢ ⎤ ⎦⎥ π π π r k r r ri i is s ln (1) where o [W m–2K–1] is the heat transfer coefficient at the sprinkled tubes’ surface i [W m–2 K–1] is the heat-transfer coefficient at the inner side of a tube set for a fully developed turbulent flow according to11, ro; ri [m] are the outer and inner tube radii S [W m–1 K–1] is the thermal conductivity kS [W m–1 K–1] is the heat admittance based on the above-mentioned laws governing heat transfer, which is calculated from the heat balance of the heating side of the loop, which is why the following Equation (2) must be valid:   ( )lnQ k L T M c t t t ts s p= = +⎛ ⎝ ⎜ ⎞ ⎠ ⎟ ⋅ −Δ 34 3 4 3 42 (2) where M34 [kg s–1] is the mass flow of heating water cp [J kg–1 K–1] is the specific heat capacity of water at constant pressure related to the mean temperature inside the loop L [m] is the total length of the bundle Tln [K] is a logarithmic temperature gradient where a counter-current exchanger was considered The final heat-transfer coefficient for various falling film liquid flow rates is recalculated into the Nusselt number, due to its more common application for a results comparison. For sprinkled exchangers the following form is normally used3,4 for the recalculation of the heat-transfer coefficient into the Nusselt number: [ ]Nu v g =  0 2 3 3 – (3) where v [m2 s–1] is the kinematic viscosity, g [m s–2] is the acceleration due to gravity and [W m–1 K–1] is the liquid film’s thermal conductivity. 4 RESULTS Tube bundles comprising four, six and eight copper tubes were tested and the results are given in this paper. Constant flow rates of the sprinkling liquid and the heating liquid of the required temperature were set first; the pressure in the chamber was then slowly lowered. The initial pressure in the chamber at that moment always equaled the atmospheric pressure. Three temperature gradients of sprinkling water and heating water 20/40, 15/40 and 20/50 were tested. The first number designates the temperature of the sprinkling water in the distribution tube; the other number desig- nates the temperature of the heating water entering the bundle. The thermal differences were (20, 25 and 30) °C. The flow rate of the heating water in all the experiments was kept at approximately 7.2 L min–1. The flow rates of the sprinkling liquid were carefully selected and re- mained "constant" throughout the experiments. Figure 2 shows a recording of the main quantities measured in one of the experiments. Only four tubes were heated in this experiment, which lasted 16 min and 37 s (horizontal axis of the chart). The average tempe- rature of the heating water entering the bottom tube (T3) was 40.2 °C ± 0.3 °C, and the average flow rate of the heating water (V2) was 7.16 ± 0.03 L min–1. The tem- perature of the heating water leaving the exchanger is designated as (T4) in the chart. The average temperature P. KRACÍK et al.: THE SIZE EFFECT OF HEAT-TRANSFER SURFACES ON BOILING Materiali in tehnologije / Materials and technology 50 (2016) 6, 939–944 941 Figure 2: Record of the main measured quantities Slika 2: Zapis glavnih izmerjenih koli~in of the sprinkling water entering the distribution tube was 15.4 °C ± 0.3 °C, and average flow rate (V1) was 4.03 ± 0.02 L min–1. The mass flow of the sprinkling liquid related to the length of the sprinkled area (which is 0.940 mm for all three exchangers) is more commonly used for the comparison. The said mass flow was 0.0713 ± 0.0004 kg/(s m), i.e., the average Reynolds number was 304.4 ± 2.1 [-]. Other values in the chart include the current pressure in the testing chamber (p) (in absolute values), and the temperature of the sprinkling water (T2) beyond the tube bundle where the water was heated. Figure 3 shows the quantities derived from the measurements presented in the previous figure. The quantities in the chart depend on the duration of the experiment. The quantities represented in the chart include the current heat flow extracted from the heating water (Q_34), the current heat flow absorbed by the sprinkling liquid (Q_12), the heat-transfer coefficient on the inside tube wall (alpha_i) and the analyzed heat- transfer coefficient on the surface of the tube (alpha_o). The chart also presents the current pressure in the testing chamber. The course of the analysed heat-transfer coeffi- cient clearly shows that decrease in pressure in the chamber results in a slightly increased coefficient. The coefficient depends mostly on the temperature of the heating water and the sprinkling water, and not so much on their flow rates. Altogether 21 experiments on tube bundles with four, six and eight tubes were performed, that is seven experi- ments for each bundle. Table 1 presents the measured parameters that helped identify, among others, the mass flow of the sprinkling liquid. The parameters were rela- ted to the length of the sprinkled area, the corresponding Reynolds number, and the heat flow extracted from the P. KRACÍK et al.: THE SIZE EFFECT OF HEAT-TRANSFER SURFACES ON BOILING 942 Materiali in tehnologije / Materials and technology 50 (2016) 6, 939–944 Table 1: Core values of the experiments Tabela 1: Temeljne vrednosti eksperimentov tubes T1 (°C) T2 (°C) T3 (°C) T4 (°C) V1 (L/min) V2 (L/min)  (kg/(s.m)) Re [-] q (kW/m2) 4 19.7 ± 0.1 32.1 ± 0.2 40.1 ± 0.1 32.7 ± 0.1 4 ± 0 7.2 ± 0 0.0711 ± 0.0005 326 ± 2.4 26.03 ± 0.38 20.2 ± 0.1 29.4 ± 0.3 40.3 ± 0.3 31.7 ± 0.2 6 ± 0 7.2 ± 0 0.1065 ± 0.0004 476.9 ± 2.9 30.26 ± 0.64 15.4 ± 0.2 30.2 ± 0.4 40.2 ± 0.3 30.9 ± 0.3 4 ± 0 7.2 ± 0 0.0713 ± 0.0004 304.3 ± 1.5 32.66 ± 0.29 15.7 ± 0.1 28.5 ± 0.3 39.9 ± 0.4 29.9 ± 0.2 5 ± 0 7.2 ± 0 0.0889 ± 0.0003 373.4 ± 1.4 34.97 ± 0.78 15.5 ± 0.2 27.3 ± 0.2 40.6 ± 0.1 29.8 ± 0.1 6 ± 0 7.1 ± 0 0.1065 ± 0.0006 440.3 ± 3.3 37.57 ± 0.4 19.7 ± 0.2 38.4 ± 0.4 50.2 ± 0.2 38.8 ± 0.2 4 ± 0 7.2 ± 0 0.0709 ± 0.0005 348.4 ± 2.7 39.76 ± 0.48 20.1 ± 0.2 33.6 ± 0.2 50.2 ± 0.2 37.1 ± 0.1 6 ± 0 7.2 ± 0 0.1064 ± 0.0005 498.9 ± 2.3 46.08 ± 0.46 6 20.2 ± 0.1 35.6 ± 0.2 40.6 ± 0.1 32.1 ± 0.1 4 ± 0 7.2 ± 0 0.0709 ± 0.0002 339.8 ± 1.6 19.94 ± 0.21 19.6 ± 0.2 32.0 ± 0.2 39.7 ± 0.2 29.6 ± 0.3 6 ± 0 7.2 ± 0 0.1065 ± 0.0003 487.2 ± 2.9 23.8 ± 0.33 15.4 ± 0.1 33.2 ± 0.3 40.4 ± 0.3 30.3 ± 0.2 4 ± 0.1 7.1 ± 0 0.0713 ± 0.0017 309.3 ± 7.9 23.71 ± 0.49 15.5 ± 0.2 32.3 ± 0.2 40.2 ± 0.3 28.2 ± 0.2 5 ± 0 7.2 ± 0 0.0888 ± 0.0002 389.1 ± 1.7 28.2 ± 0.36 15.4 ± 0.3 31.3 ± 0.5 40.5 ± 0.2 28.0 ± 0.3 6 ± 0 7.2 ± 0 0.1065 ± 0.0003 444.4 ± 3.5 29.03 ± 0.38 20.2 ± 0.2 40.8 ± 0.8 50.4 ± 0.3 36.0 ± 0.2 5 ± 0 7.2 ± 0 0.0885 ± 0.0005 448.8 ± 5.2 34.03 ± 0.38 19.9 ± 0.1 38.5 ± 0.5 49.5 ± 0.1 34.2 ± 0.1 6 ± 0 7.2 ± 0 0.1063 ± 0.0003 524.7 ± 3.6 35.64 ± 0.28 8 20.3 ± 0.1 37.2 ± 0.4 40.3 ± 0.2 31.1 ± 0.1 4 ± 0 7.2 ± 0 0.0709 ± 0.0005 346.5 ± 3.2 16.28 ± 0.31 20.2 ± 0.1 34.7 ± 0.2 40.3 ± 0.1 28.9 ± 0.1 6 ± 0 7.2 ± 0 0.1064 ± 0.0004 505 ± 2.3 19.94 ± 0.13 15.3 ± 0.2 36.3 ± 0.3 40.2 ± 0.1 28.7 ± 0.2 4 ± 0 7.2 ± 0 0.071 ± 0.0003 324.7 ± 2.5 20.25 ± 0.22 15.6 ± 0.3 34.1 ± 0.4 39.9 ± 0.4 27.1 ± 0.3 5 ± 0 7.2 ± 0 0.0887 ± 0.0003 397 ± 2.9 22.3 ± 0.53 15.3 ± 0.5 32.8 ± 0.6 39.7 ± 0.3 26.1 ± 0.5 6 ± 0 7.2 ± 0 0.1063 ± 0.0006 467.4 ± 6.5 23.82 ± 0.51 20.0 ± 0.2 45.8 ± 1.2 50.4 ± 0.2 36.3 ± 0.4 4 ± 0 7.2 ± 0 0.0709 ± 0.0003 378 ± 3.5 24.93 ± 0.81 20.2 ± 0.1 41.2 ± 0.7 50.0 ± 0.3 33.0 ± 0.3 6 ± 0 7.2 ± 0 0.1065 ± 0.0003 542.2 ± 4.2 29.73 ± 0.27 Figure 3: Quantities derived from the measurements Slika 3: Koli~ine pridobljene iz meritev Figure 4: Dependency of the Nusselt number on the pressure in the testing chamber and the heat flow extracted from the heating liquid – four tubes Slika 4: Odvisnost Nusseltovega {tevila od tlaka v preizkusni komori in toplotni tok pridobljen z teko~ine za ogrevanje – {tiri cevi heating water, which was related to the size of the heat- exchanging surface. The heat-transfer coefficient on the surface of the tube bundle was later converted to the Nusselt number. Values acquired at pressures ranging from 5.0 kPa(abs) to 95.0 kPa(abs) in 5.0 kPa increments, and values acquired at atmospheric pressure were selected from the data measured in each experiment. The values were then averaged for each pressure. This helped us to develop matrices of the same sizes for a particular bundle length. The matrices were interpolated using a cubic curve in Matlab software. Contour graphs, shown in Figures 4 to 6, were developed using these new matrices. The pre- sented Nusselt number in the figures is dependent on the pressure in the chamber (vertical axis) and the extracted heat flow of the heating water (horizontal axis). The range of the Nusselt number is identical for all three charts, i.e., 0.15 to 0.46 [-], so that individual phases can be compared. However, this procedure was disabled to observe slight increases in the Nusselt number as the pressure decreases in particular measurement sessions. All three analysed sizes of the tube bundle clearly show three vertical zones that do not connect; these are a result of a different temperature gradient. The tempera- ture gradient is 20/40 in the zone 1, 15/40 in the zone 2, and 20/50 in the zone 3. Temperature gradient 15/40 with the highest Nusselt number seems to be the best option for all three analysed lengths of the bundle. The Nusselt number dropped by almost 0.2 [-] when the temperature of heating water increased. The Nusselt number also drops at very low pressures. The decrease is caused by the fact that heat extracted from the heating liquid is also used for evaporation and not only for heating the heating water. However, the decrease is not very significant and the boiling occurs only at the bottom section of the tube bundle. Concerning a tube bundle with four tubes, the maxi- mum Nusselt number was attained at a specific heat flow of approximately 35 kW m–2 and 38 kW m–2, which corresponds to a temperature gradient of 15/40 °C. The maxima ranged at values around 0.44 [-]. The larger the surface area and the lower the specific loading, the higher the values. In the case of a tube bundle with six tubes, the heat flow is approximatelz 28 kW m–2. In the case of a tube bundle with eight tubes, the maximum reached almost 0.46 [-] at a thermal load of approxi- mately 24 kW m–2. 5 CONCLUSIONS Sprinkled tube bundles are most commonly used as evaporators since they quickly separate the vapor phase and the liquid phase thanks to a thin liquid film. How- ever, the practice proves that there is no boiling on the first affected tubes, only the liquid is heated. We describe in this paper how the surface area of the tubes affects the heat transfer coefficient on the tube surface. This effect was tested at three thermal differences: (15, 20 and 30) °C. After the flow rate of the heating and sprinkling liquids and their temperatures were set, only the pressure in the testing chamber was changed during the experiment. The initial pressure in the chamber was atmospheric pressure. The pressure in the chamber and the exhaustion of the air-vapour mixture from the chamber (vapour was formed at the bottom part of the tube bundle at low pressures) had no major impact on the coefficient. It was the exhaustion itself which affected the boundary layers on the tube bundle. A significant change of the coeffi- cient may come only if the boiling on the tube bundle increases. The impact of the heat-exchanging surface area on the heat-transfer coefficient is evident in zone 3 of parti- P. KRACÍK et al.: THE SIZE EFFECT OF HEAT-TRANSFER SURFACES ON BOILING Materiali in tehnologije / Materials and technology 50 (2016) 6, 939–944 943 Figure 6: Dependency of the Nusselt number on the pressure in the testing chamber and the heat flow extracted from the heating liquid – eight tubes Slika 6: Odvisnost Nusseltovega {tevila od tlaka v preizkusni komori in toplota pridobljena iz teko~ine za ogrevanje – osem cevi Figure 5: Dependency of the Nusselt number on the pressure in the testing chamber and heat flow extracted from the heating liquid – six tubes Slika 5: Odvisnost Nusseltovega {tevila od tlaka v preizkusni komori in toplota pridobljena iz teko~ine za ogrevanje – {est cevi cular bundles where the thermal gradient was 20/50 °C. The increase in the surface area results in a decrease in the specific loading of the area but the heat-transfer coefficient increases. Acknowledgement Presented results were obtained in frame of the project NETME CENTRE PLUS (LO1202), created with financial support from the Ministry of Education, Youth and Sports of the Czech Republic under the “National Sustainability Programme I". 6 REFERENCES 1 J. Beniak, J. Ondru{ka, V. ^a~ko, Design Process of Energy Effec- tive Shredding Machines for Biomass Treatment, Acta Polytechnica, 52 (2012) 5, 1210–2709 2 J. Beniak, P. Kri`an, M. Matú{, M. Ková~ová, The Operating Load of a Disintegration Machine, Acta Polytechnica, 54 (2014) 1, 1210–2709, doi:10.14311/AP.2014.54.0001 3 L. H. Chien, Ch. H. Cheng, A Predictive Model of Falling Film Eva- poration with Bubble Nucleation on Horizontal Tubes, HVAC, 12 (2006) 1, 1078–9669, doi:10.1080/10789669.2006.10391168 4 J. J. Lorenz, D. Yung, A Note on Combined Boiling and Evaporation of Liquid Films on Horizontal Tubes, Journal of Heat Transfer, 101 (1979) 1, 0022–1481, doi:10.1115/1.3450914 5 W. L. Owens, Correlation of Thin Film Evaporation Heat Transfer Coefficients for Horizontal Tubes, Proceedings of the Fifth Ocean Thermal Energy Conversion Conference, Miami Beach, Florida, 1978 6 W. H. Parken, L. S. Fletcher, V. Sernas, J. C. Han, Heat Transfer Through Falling Film Evaporation and Boiling on Horizontal Tubes, Journal of Heat Transfer, 112 (1990) 3, 0022–1481, doi:10.1115/ 1.2910449 7 V. Sernas, Heat Transfer Correlation for Subcooled Water Films on Horizontal Tubes, Journal of Heat Transfer, 101 (1979) 1, 0022–1481, doi:10.1115/1.3450913 8 R. Armbruster, J. Mitrovic, Patterns of Falling Film Flow over Horizontal Smooth Tubes, Proceedings of the 10th international heat transfer conference, Brighton, UK, 1994, 3 9 X. Hu, A. M. Jacobi, The Intertube Falling Film: Part 1 - Flow Characteristics, Mode Transitions, and Hysteresis, Journal of Heat Transfer, 118 (1996) 3, 0022–1481, doi:10.1115/1.2822678 10 J. F. Roques, V. Dupont, J. R. Thome, Falling Film Transitions on Plain and Enhanced Tubes, Journal of Heat Transfer, 124 (2002) 3, 0022–1481, doi:10.1115/1.1458017 11 M. Jícha, Heat and mass transfer, Brno, CERM, 2001, 1, 160 P. KRACÍK et al.: THE SIZE EFFECT OF HEAT-TRANSFER SURFACES ON BOILING 944 Materiali in tehnologije / Materials and technology 50 (2016) 6, 939–944 A. GRAJCAR et al.: EFFECT OF GAS ATMOSPHERE ON THE NON-METALLIC INCLUSIONS ... 945–950 EFFECT OF GAS ATMOSPHERE ON THE NON-METALLIC INCLUSIONS IN LASER-WELDED TRIP STEEL WITH Al AND Si ADDITIONS VPLIV PLINSKE ATMOSFERE NA NEKOVINSKE VKLJU^KE V LASERSKO VARJENEM TRIP JEKLU Z DODATKOM Al IN Si Adam Grajcar1, Maciej Ró¿añski2, Ma³gorzata Kamiñska3, Barbara Grzegorczyk1 1Silesian University of Technology, Institute of Engineering Materials and Biomaterials, Konarskiego Street 18a, 44100 Gliwice, Poland 2Institute of Welding, Bl. Czes³awa Street 16-18, 44100 Gliwice, Poland 3Institute of Non Ferrous Metals, Sowinskiego Street 5, 44100 Gliwice, Poland adam.grajcar@polsl.pl Prejem rokopisa – received: 2015-08-12; sprejem za objavo – accepted for publication: 2015-11-04 doi:10.17222/mit.2015.253 The present study aims to characterize the weldability of a multiphase, automotive steel containing Al and Si additions from the point of view of its tendency to form non-metallic inclusions. Laser welding tests of 2-mm-thick sheets were performed using the keyhole-welding mode and a solid-state laser. The tests were carried out in air and with the use of an argon atmosphere. The distribution, type and chemical composition of the non-metallic inclusions formed in the base metal and fusion zones were analysed. The effect of applying the protective gas on the type and amount of non-metallic inclusions was determined using light and scanning electron microscopy. The chemical composition of the identified particles was assessed using the EDS method. It was found that a protective gas has a beneficial effect on reducing the non-metallic inclusions, but only to a limited extent. The boundary between the complex oxides and the pure aluminium oxides was determined to be 2–3 μm. Keywords: TRIP steel, laser welding, non-metallic inclusions, protective gas, Ar atmosphere, oxidation Namen te {tudije je opredeliti varivost s stali{~a tvorbe nekovinskih vklju~kov v ve~faznem jeklu, ki vsebuje Al in Si, za avto- mobilsko industrijo. Izvedeni so bili varilni preizkusi z laserjem na 2 mm debelih plo~evinah z V zarezo. Preizkusi so bili izvedeni na zraku in v za{~itni atmosferi argona. Analizirana je bila razporeditev, vrsta in kemijska sestava nekovinskih vklju~kov v osnovnem materialu in v coni zlivanja. Vpliv uporabe za{~itnega plina na vrsto in koli~ino nekovinskih vklju~kov je bil dolo~en s pomo~jo svetlobne in vrsti~ne elektronske mikroskopije. Kemijska sestava najdenih delcev je bila ugotovljena z metodo EDS. Ugotovljeno je bilo, da ima za{~itni plin ugoden vpliv na zmanj{anje nekovinskih vklju~kov le v omejenem obsegu. Meja med kompleksnimi oksidi in ~istimi oksidi aluminija je bila dolo~ena kot 2-3 μm. Klju~ne besede: TRIP jeklo, lasersko varjenje, nekovinski vklju~ki, za{~itni plin, atmosfera Ar, oksidacija 1 INTRODUCTION Multiphase steels with a transformation-induced plasticity (TRIP) effect belong to the advanced high- strength steels (AHSS) used in the modern automotive industry. The beneficial balance between strength and ductility requires increased contents of Mn, Al and Si. Their total content in TRIP steels can reach 4 %.1–3 Medium-Mn and high-Mn alloys require an even higher concentration of alloying elements.4–8 Whereas a lot of efforts focus on the relationships between hot-working, heat-treatment, microstructure and mechanical pro- perties, the weldability of AHSS has not received much attention so far. M. S. Weglowski et al.9 reported that the maximum hardness in the fusion zone of 0.07C-1Mn-0.4Cr dual-phase steel (which belongs to the 1st generation of AHSS) reaches 340 HV. The hardness increases to approximately 430 HV with increasing carbon and manganese contents in the 0.13C-1.3Mn-0.2Cr-0.2Cu steel.10 Another problem in high-strength, multiphase steels is the heat-affected zone (HAZ) softening because of the tempering of the pre-existing martensite. Since the AHSS contain a high alloying content their weldability should also be affected by the presence of non-metallic inclusions. Different sulphide and oxide particles have been identified by M. Amirthalingam et al.11 and A. Graj- car et al.12 in different types of TRIP steels containing Mn, Si and Al additions. It is well known that brittle oxides and ductile manganese sulphides affect the fatigue endurance limit, the fatigue-crack propagation rate, the fracture toughness, the anisotropy of the tensile properties with respect to the rolling direction and the weld quality, too.13,14 Recently, A. Grajcar et al.12 analysed various types of non-metallic inclusions in laser-welded Fe-1.5Mn- 0.9Si-0.4Al TRIP steel. Numerous oxide-type particles have been revealed in the fusion zone formed under air conditions. Therefore, the aim of the present work is to investigate the effect of a protective gas on the quantity, type and chemical composition of non-metallic inclu- sions in Si- and Al-alloyed TRIP steel. Materiali in tehnologije / Materials and technology 50 (2016) 6, 945–950 945 UDK 621.791.754'293:544.022.344.3:621.791.725 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)945(2016) 2 EXPERIMENTAL PART The work addresses the laser welding of thermo-me- chanically processed, 2-mm-thick, TRIP steel sheets with the following chemical composition: 0.24 % C, 1.55 % Mn, 0.87 % Si, 0.4 % Al, 0.03 % Nb, 0.023 % Ti, 0.004 % S, 0.01 % P and 0.0028 % N. Rare-earth ele- ments (REE) such as mischmetal (~50 % Ce, ~20 % La, ~20 % Nd) were added to modify the chemical compo- sition and the shape of the non-metallic inclusions. The 25-kg, vacuum-melted ingot was cast using a protective atmosphere of argon. The Si+Al addition is required in TRIP steels to prevent carbide precipitation because C is needed to enrich the retained austenite.1,2,11 The Nb+Ti micro-additions are added to increase the strength level as a result of the precipitation strengthening and grain refinement. These phenomena are especially useful for thermo-mechanically processed steels.15,16 The initial hot-working included the hot forging and rough rolling of the ingot to a thickness of 5 mm. The fundamental thermo-mechanical rolling consisted of 3 passes with a finishing rolling temperature of 850 °C. The major step of the controlled cooling following the hot rolling was isothermal holding of the sheets at a tem- perature of 350 °C within 600 s. The final thickness of the sheets was 2 mm. Laser welding generally includes two major modes: keyhole welding and conductive welding.17 The conduc- tive welding utilizes the natural thermal conduction into the material. The welds obtained in this welding mode have good quality and do not contain pores and spalls. However, the keyhole welding is more efficient. Because of this the present tests were carried out using keyhole welding. In this treatment mode the power density is much higher compared to the conductive welding and the fusion depth is much higher compared to the diameter of the liquid pool. The welding tests were performed using a solid-state laser integrated with a robotized laser-treat- ment system. The welding station is equipped with the TruDisk 12002 solid-state laser type Yb:YAG characte- rized by a maximum power of 12 kW. The heat input value of 0.048 kJ/mm was applied. The welding tests were performed using an Ar atmosphere. To assess the effect of the protective atmosphere, tests under air con- ditions were performed too. The distribution, type and chemical composition of the non-metallic inclusions formed in the base metal (BM) and fusion zone (FZ) of both types of samples (using Ar gas and without any protective gas) were com- pared. Samples for light microscopy (LM) and scanning electron microscopy (SEM) were prepared. The chemical composition of the non-metallic inclusions was assessed using EDS point analyses. The distribution of the parti- cular alloying elements was revealed using mapping. The quantitative measurements of the chemical composition for the identified inclusions were carried out using a JCXA 8230 X-ray micro-analyser with an accelerating voltage of 15 kV. The microstructures of the BM, the heat-affected zone (HAZ) and the FZ were revealed using the SUPRA 25 SEM at an accelerating voltage of 20 kV after nital etching. 3 RESULTS AND DISCUSSION 3.1 Distribution of non-metallic inclusions The use of different gas atmospheres influences the quantity of non-metallic inclusions. The distributions of non-metallic inclusions formed in the fusion zones under the conditions of an air atmosphere and Ar protective gas are compared in Figure 1. The intense oxidation takes place when the laser welding was conducted without any protective gas. As a result numerous particles of different sizes can be observed in the fusion zone. The clear boun- dary near the fusion line is easily visible in Figure 2 for both the air- and Ar-treated samples. The amount of particles in the heat-affected zone is a few times lower when compared to the fusion zones. The size of the non-metallic inclusions is similar when the protective gas was applied. The quantity of particles is lower compared to the laser treatment with- out any protective atmosphere. However, the amount of non-metallic inclusions for Ar-protected samples is higher than would be expected. This means that the Ar atmosphere has a limited efficiency in the reduction of harmful, non-metallic inclusions formed in the fusion zone. The explanation is the essence of the keyhole- welding technique. The high-density power in the region of the laser beam’s exposure creates a gas-dynamic A. GRAJCAR et al.: EFFECT OF GAS ATMOSPHERE ON THE NON-METALLIC INCLUSIONS ... 946 Materiali in tehnologije / Materials and technology 50 (2016) 6, 945–950 Figure 1: Distribution of non-metallic inclusions in the fusion zones formed under conditions of air atmosphere and Ar protective gas Slika 1: Razporeditev nekovinskih vklju~kov v podro~ju zlivanja, nastalih na zraku in v za{~itni atmosferi Ar channel, i.e., a deep and narrow capillary filled with gases and metal steams.17,18 These turbulent gases partially break the protective atmosphere of the argon. That is why the amount of non-metallic inclusions for laser welding tests with the use of Ar gas is higher than expected. 3.2 Microstructure of the steel The microstructure of the base metal consists of polygonal ferrite (F) grains elongated along the hot-roll- ing direction, bainite and retained austenite (Figure 3). Retained austenite (RA) is the most favourable structural constituent of TRIP steels due to its beneficial effect on increasing the steel’s plasticity. This phase is usually found as small, blocky, granules along ferrite grain boundaries or between bainitic ferrite laths. Hence, a large fraction of the microstructure constitutes bainite- austenite (BA) constituents. The rapid cooling rate typical for laser welding influences the microstructure of the fusion zone. Typical SEM microstructures of the fusion zone are shown in Figure 4. Both microstructures are characterized by the presence of martensite laths. Due to the chemical con- trast typical for observations using back-scattered elec- trons (BSE) white interlath retained austenite (RA) can also be observed. The amount of non-metallic inclusions is slightly smaller for the samples welded using the argon gas. Moreover, a lot of sub-micron-sized particles are formed in this steel. 3.3 Non-metallic inclusions The amount of non-metallic inclusions in the base metal is very small (Figures 1 and 2). It is related to the high metallurgical cleanliness of the laboratory-melted steel and the use of rare-earth elements. A result of the mischmetal addition is a partial or total substitution of A. GRAJCAR et al.: EFFECT OF GAS ATMOSPHERE ON THE NON-METALLIC INCLUSIONS ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 945–950 947 Figure 3: SEM microstructure of the base metal consisting of ferrite (F), bainite-austenite constituents (BA) and retained austenite (RA) Slika 3: SEM-posnetek mikrostrukture osnovnega materiala, ki jo sestavljajo ferit (F), bainit-avstenit (BA) in zaostali avstenit (RA) Figure 2: Distribution of non-metallic inclusions near the fusion line in the fusion zones and heat-affected zones formed under conditions of air atmosphere and Ar protective gas Slika 2: Razporeditev nekovinskih vklju~kov blizu linije podro~ja zlivanja in toplotno vplivane cone, nastalih na zraku in v za{~itni atmosferi Ar Figure 4: Martensite laths and globular non-metallic inclusions formed in the fusion zones under conditions of: a) air atmosphere and b) Ar protective gas Slika 4: Martenzitne late in globularni nekovinski vklju~ki, nastali v coni zlivanja: a) na zraku in b) v za{~itni atmosferi Ar Mn in sulphide inclusions and Al in oxide inclusions by the REE. These elements have a higher chemical affinity for sulphur and oxygen compared to Mn and Al. The example of such a particle is shown in Figure 5. In this case Ce and La totally replaced Mn and Al, forming a globular oxysulphide. Such particles are hardly deformed during hot rolling and reduce the anisotropy of the me- chanical properties of flat products.4,15 The mapping of the globular inclusions indicates that Ce and La as well as S and O are distributed uniformly within the particle (Figure 6). Figure 7 presents typical non-metallic inclusions formed in the fusion zone of the sample laser-welded in A. GRAJCAR et al.: EFFECT OF GAS ATMOSPHERE ON THE NON-METALLIC INCLUSIONS ... 948 Materiali in tehnologije / Materials and technology 50 (2016) 6, 945–950 Figure 7: Complex oxides of various size formed in the fusion zone of the sample welded in the air atmosphere: a), b) spectrum of the particle from point 1, c) spectrum of the particle from point 2 and d) the chemical composition of the large particle (point 1) determined using EDS Slika 7: Sestavljeni oksidi razli~nih velikosti: a) nastali v coni zlivanja vzorca zvarjenega na zraku, b) spekter delca v to~ki 1, c) spekter delca v to~ki 2 in d) kemijska sestava velikega delca (to~ka 1), dolo~ena s pomo~jo EDS Figure 5: Complex oxysulphide containing La and Ce: a) formed in the base metal, b) spectrum of the inclusion and c) its chemical composition determined using EDS Slika 5: Sestavljeni oksisulfid, ki vsebuje La in Ce: a) nastal v osnovnem material, b) spekter vklju~ka in c) njegova kemijska sestava, dolo~ena s pomo~jo EDS Figure 6: Elemental mapping of the particle from Figure 5 Slika 6: Razporeditev elementov v delcu prikazanem na Sliki 5 the air atmosphere. There are a few large inclusions with a size ranging from approximately 3 μm to 7 μm and numerous small particles smaller than 1 μm. All the inclusions have a globular shape. EDS analyses of the large particles indicate that these are complex oxides containing Al, Mn and Si (Figures 7b to 7d). It indicates that there is a strong oxidation of the alloying elements in the air atmosphere. Elemental mapping indicates that there is no sulphur in the inclusions. Moreover, it should be noted that Mn and Si are located only at the largest inclusions, whereas the smaller ones are pure aluminium oxides (Figure 8). The use of a protective atmosphere of argon causes a decrease in the amount of non-metallic inclusions. It is especially true for large particles, the quantity of which is much smaller (Figure 9a). As previously, the spectral lines in Figures 9b and 9d do not contain sulphur peaks. The intense evaporation during the keyhole welding re- sults in the partial oxidation of the same alloying ele- ments, like for the samples welded in the air atmosphere. Hence, many oxides can be identified in Figure 9. The analysis of the chemical composition indicates that the concentrations of Mn and Si decrease with a decrease in the particle diameter (Figures 9c and 9e). The presence of Fe in the spectrum of the smaller particle (point 2 in Figure 9) is caused by a similar diameter of the inclu- sion and a beam diameter. The occurrence of Mn and Si only in the largest inclusions is confirmed by Figure 10, where these elements are located in the largest particle of a diameter of approximately 4 μm. Some content of titanium was also revealed. A. GRAJCAR et al.: EFFECT OF GAS ATMOSPHERE ON THE NON-METALLIC INCLUSIONS ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 945–950 949 Figure 8: Elemental mapping of the particles from Figure 7 Slika 8: Razporeditev elementov v delcih, prikazanih na Sliki 7 Figure 10: Elemental mapping of the particles from Figure 9 Slika 10: Razporeditev elementov v delcih, prikazanih na Sliki 9 Figure 9: Complex oxides of various sizes formed in the fusion zone of the sample: a) welded in the Ar atmosphere, b), d) spectra of the particles from points 1 and 2, c) the chemical composition of the particles from point 1 and e) point 2 determined using EDS Slika 9: Sestavljeni oksidi razli~nih velikosti nastalih v coni zlivanja: a) vzorca zvarjenega v atmosferi Ar, b), d) spekter delcev iz to~ke 1 in 2, c) kemijska sestava delcev iz to~ke 1 in e) to~ke 2, dolo~enih s po- mo~jo EDS Ti, Al, Si and Mn are also located in the largest inclu- sion in Figure 11. However, the amount of such large inclusions is reduced compared to the samples welded in the air atmosphere. The region around such large inclu- sions is usually free of any other particles. The limit bet- ween Mn- and Si-containing large inclusions and small pure aluminium oxides can be assessed as 2–3 μm. 4 CONCLUSIONS The effect of the use of an Ar atmosphere during laser welding in the keyhole-welding mode was assessed for the Al-Si-bearing TRIP steel. It was found that the intense evaporation of gases and metal steams partially destroys the protective atmosphere of the argon. As a result, some oxidation of the weld pool takes place. The amount of non-metallic inclusions in the fusion zone of the samples welded using Ar is smaller compared to the samples treated without a protective gas. It concerns especially the large inclusions with a diameter between 3 μm and 7 μm. The Ar atmosphere has no effect on the chemical composition of the formed particles. In both cases globular oxides of various sizes are formed in the fusion zone. The particles with a diameter ranging from 3 7 μm to 7 μm constitute the complex inclusions con- taining Al, Si and Mn (sometimes also Ti). The nume- rous particles smaller than 2–3 μm are pure aluminium oxides. Acknowledgment This work was financially supported with statutory funds of Faculty of Mechanical Engineering of Silesian University of Technology in 2015. 5 REFERENCES 1 B. Masek, C. Stadler, H. Jirkova, P. Feuser, M. Selig, Transforma- tion-induced plasticity in steel for hot stamping, Mater. Tehnol., 48 (2014), 555–557 2 A. Kokosza, J. Pacyna, Formation of medium carbon TRIP steel microstructure during annealing in the intercritical temperature range, Archives of Metallurgy and Materials, 59 (2014), 1017–1022 3 S. Wiewiorowska, Analysis of the influence of drawing process para- meters on the mechanical properties of TRIP-structure steel wires, Archives of Metallurgy and Materials, 58 (2013), 573–576 4 A. Grajcar, M. Opiela G. Fojt-Dymara, The influence of hot-working conditions on a structure of high-manganese steel, Archives of Civil and Mechanical Engineering, 9 (2009), 49–58 5 S. Lasek, E. Mazancova, Influence of thermal treatment on structure and corrosion properties of high manganese triplex steels, Meta- lurgija, 52 (2013), 441–444 6 L. A. Dobrzañski, A. Grajcar, W. Borek, Microstructure evolution of C-Mn-Si-Al-Nb high-manganese steel during the thermomechanical processing, Materials Science Forum, 638-642 (2010), 3224–3229, doi:10.4028/MSF.638-642.3224 7 M. Jab³oñska, G. Niewielski, R. Kawalla, High-manganese TWIP steel – technological plasticity and selected properties, Solid State Phenomena, 212 (2014), 87–90, doi:10.4028/SSP.212.87 8 A. Grajcar, R. 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Grzegorczyk, Study on non-metallic inclusions in laser-welded TRIP-aided Nb-microalloyed steel, Archives of Metallurgy and Materials, 59 (2014), 1163–1169, doi:10.2478/amm-2014-0203 13 J. Maciejewski, The effects of sulfide inclusions on mechanical pro- perties and failures of steel components, Journal of Failure Analysis and Prevention, 15 (2015), 169–178, doi:10.1007/s11668-015-9940-9 14 I.J. Park, S.M. Lee, M. Kang, S. Lee, Y.K. Lee, Pitting corrosion behavior in advanced high strength steels, Journal of Alloys and Compounds, 619 (2015), 205–210, doi:10.1016/j.jallcom.2014. 08.243 15 J. Górka, Weldability of thermomechanically treated steels having a high yield point, Archives of Metallurgy and Materials, 60 (2015), 469–475, doi:10.1515/amm-2015-0076 16 R. Celin, J. Bernetic, D.A. Skobir Balantic, Welding of the steel grade S890QL, Mater. Tehnol., 48 (2014), 931–935 17 A. Lisiecki, Welding of thermomechanically rolled fine-grain steel by different types of lasers, Archives of Metallurgy and Materials, 59 (2014), 1625–1631, doi:10.2478/amm-2014-0276 A. GRAJCAR et al.: EFFECT OF GAS ATMOSPHERE ON THE NON-METALLIC INCLUSIONS ... 950 Materiali in tehnologije / Materials and technology 50 (2016) 6, 945–950 Figure 11: Elemental mapping of the complex oxides formed in the fusion zone of the sample welded in the Ar atmosphere Slika11: Razporeditev elementov v sestavljenih oksidih, nastalih v coni zlivanja, vzorca zvarjenega v atmosferi Ar T. M. CHENTHIL JEGAN et al.: MACHINING PARAMETERS INFLUENCING IN ELECTRO CHEMICAL MACHINING ... 951–960 MACHINING PARAMETERS INFLUENCING IN ELECTRO CHEMICAL MACHINING ON AA6061 MMC PARAMETRI STROJNE OBDELAVE, KI VPLIVAJO NA ELEKTROKEMIJSKO STROJNO OBDELAVO AA6061 MMC Chenthil Jegan Thankaraj Mariapushpam1, Durairaj Ravindran2, Manaharan Dev Anand3 1St. Xavier's Catholic College of Engineering, Department of Mechanical Engineering, Kanyakumari District, India 2National Engineering College, Department of Mechanical Engineering, K.R.Nagar, Kovilpatti, Thuthukoodi District, India 3Noorul Islam Centre for Higher Education, Department of Mechanical Engineering, Kumaracoil, Kanyakumari District, India optrajegan@yahoo.co.in Prejem rokopisa – received: 2015-08-18; sprejem za objavo – accepted for publication: 2015-11-05 doi:10.17222/mit.2015.260 In this study the mechanical and microstructural behaviours of AA6061 reinforced with silicon carbide (SiC) and AA6061 reinforced with boron carbide (B4C), obtained from the enhanced stir-casting process, were investigated using scanning electron microscopy (SEM) and X-ray diffraction (XRD) with various weight percentages, i.e., 2.5 %, 5 % and 7.5 %. The testing shows that the tensile and microhardness properties of the AA6061 were improved in both the reinforced aluminium-matrix composites. The influence of the electrochemical machining process parameters like current, voltage, electrolyte concentration, feed rate, gap and flow rate were considered as the input parameters. The output responses are the material removal rate (MRR), the surface roughness (SR) and the radial overcut (ROC). The study shows that the dominant output parameter MRR was directly proportional to the input parameters current, voltage and feed rate. The SR was significantly influenced by the input parameters current, feed rate and gap. The ROC was considerably balanced by the input parameters current and feed rate. Keywords: metal-matrix composites, material removal rate, electrochemical machining, surface roughness, radial overcut V {tudiji so bile preiskovane mehanske lastnosti in mikrostruktura AA6061, izdelanega z naprednim postopkom ulivanja s preme{avanjem in oja~anega z razli~nim masnim dele`em silicijevega karbida (SiC) ali borovega karbida (B4C) (2,5 %, 5 % in 7,5 %). Preiskave so bile izvedene s pomo~jo vrsti~nega elektronskega mikroskopa (SEM) in z rentgensko difrakcijo (XRD). Preizku{anje ka`e, da sta natezna trdnost in mikrotrdota AA6061 narasli pri obeh vrstah kompozita na osnovi aluminija. Za vpliv vhodnih procesnih parametrov pri elektrokemijski strojni obdelavi so bili upo{tevani: tok, napetost, koncentracija elektrolita, hitrost podajanja, re`a in hitrost pretoka. Izhodni odgovori so bili: hitrost odstranjevanja materiala (MRR), hrapavost povr{ine (SR) in pove~anje radiusa (ROC). [tudija ka`e, da je prevladujo~i izhodni parameter MRR neposredno proporcionalen vhodnim parametrom; toku, napetosti in hitrosti podajanja. Na SR so mo~no vplivali vhodni parametri: tok, hitrost podajanja in re`a. ROC je bil posebej uravnote`en z vhodnima parametroma, tokom in hitrostjo podajanja. Klju~ne besede: kompoziti na kovinski osnovi, hitrost odstranjevanja materiala, elektrokemijska strojna obdelava, hrapavost povr{ine, pove~anje radiusa 1 INTRODUCTION Electrochemical machining (ECM) is a well-known process used for the manufacture of various sophisticated parts, such us turbine blades, rifle bores, hip-joint im- plants, micro-components as well as many other applica- tions. ECM provides an economical and effective method for shaping high-strength, heat-resisting mate- rials into complex shapes and producing high-quality products from composites and other hard materials.1 In ECM the machining is done at low voltage compared to other processes with a high metal removal rate. It is suit- able for mass production work and low labour require- ments. ECM is one of the most widely used advanced machining processes to make complicated shapes of varying sizes with electrically conducting, but difficult to machine, materials such as super alloys, Ti alloys, alloy steel, tool steel, stainless steel, etc.2 These materials are extensively used in aerospace, automobile, space, nuc- lear, defence, cutting tools, dies and mould making applications. The material used for ECM tools should be electrically conductive and easily machinable to the required geometry. The various materials used for this purpose include copper, brass, stainless steel, titanium, and copper-tungsten. Tool insulation controls the side electrolyzing current and hence the amount of oversize. Spraying or dipping is generally the simplest method of applying insulation. Teflon, urethane, phenol, epoxy, and powder coatings are commonly used for tool insulation.3 The material removal rate of an aluminium work- piece has been obtained by electrochemical machining using a NaCl electrolyte at different current densities and compared with the theoretical values.4 It has been observed that the resistance of the electrolyte solution decreases sharply with increasing current densities. The increase in the peak current increases the MRR, TWR and ROC significantly in a nonlinear fashion, MRR and ROC increased with the increase in the pulse on time and the gap voltage was found to have some effect on the three responses.5,6 The influence of electrochemical pro- cess parameters such as the applied voltage, electrolyte Materiali in tehnologije / Materials and technology 50 (2016) 6, 951–960 951 UDK 621.7:67.017:661.665.1:661.665.3 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)951(2016) concentration, electrolyte flow rate and tool feed rate on the metal removal rate and surface roughness to fulfil the effective utilization of electrochemical machining can be used.7 For non-passivating electrode systems, the reduction in electrolyte concentration and an increase in its tem- perature improve the quality of surfaces.8 The use of a gap voltage between 20 V and 25 V saves energy and reduces the production cost.9 The machining current increases linearly with the tool feed rate. Sparking causes damage to both the tool and the workpiece due to the critical feed rate, which is because of the rapid moment of the tool towards the workpiece. The feed rate was the main parameter affecting the material removal rate.10 So it is better to maintain the feed rate according to the anodic dissolution rate for proper machining.11 The electrolyte temperature, pressure variations in the inter electrode gap and the choice of an optimum gap voltage also avoid the occurrence of sparking and the consequent loss of the tool and workpiece.12 A metal-matrix composite (MMC) contains matrix materials where reinforcements can be made from poly- mers, ceramics or metals. MMCs are divided into composites reinforced by fibres (fibrous composites) and composites filled with fine particles that are insoluble in the base-metal-strengthened composites.13 Aluminium- matrix composites consist of a uniform distribution of strengthening ceramic particles embedded within an aluminium matrix. Many researchers discovered alumi- nium materials and found they exhibit higher strength and stiffness, in addition to isotropic behaviour at a lower density, when compared to the un-reinforced aluminium matrix.14–16 The ceramic’s ability to withstand high velocity impacts and the high toughness of the metal matrix, which helps in preventing total shattering, are one of the main reason for AMC strength. The com- posites possess good mechanical properties at high temperature and thus an AMC can be a favourite choice for cost-effective alternatives and shows potential in large-scale applications such as automotive, aerospace and airframe applications. It is proved that among the aluminium alloys, AA6061 was quite a popular choice as a matrix material.17 In our previous work the ductile and microhardness properties of silicon-carbide, boron-car- bide-based AMCs were analysed and it was reported that both AMCs are suitable for unconventional machining.18 The fabrication of a MMC with various ceramic particles such as Si3N4, TiB2, B4C and machining the fabricated MMC individually was analysed by many resear- chers.19–21 Fabrication by the pressure-less infiltration process under a nitrogen gas atmosphere and the grinding of the aluminium-based MMC reinforced with SiC particles show that the physical and chemical compatibility bet- ween the SiC particles and the Al matrix is the main concern in the preparation of SiC/Al composites.22 Due to the low coefficients of thermal expansion for maxi- mizing heat dissipation and minimizing thermal stress, high-performance thermal management materials are most commonly used in the packaging of micropro- cessors, power semiconductors, high-power laser diodes, light-emitting diodes and micro-electromechanical systems.23,24 Limited research work has been reported on AMCs reinforced with B4C due to higher raw-material costs and poor wetting. Stir casting is accepted as a general com- mercial technique for producing MMCs.25 Boron carbide was an attractive reinforcement for aluminium and its alloys, showing many of the mechanical and physical properties required of an effective reinforcement, in particular high stiffness and hardness. These factors combined with a density less than that of solid alumi- nium indicate that large specific property improvements are possible.26 Boron-carbide-particulate-reinforced alu- minium composites possessed a unique combination of high specific strength, high elastic modulus, good wear resistance and good thermal stability compared to the corresponding non-reinforced matrix alloy system.27 This study covers the fabrication and machining of an aluminium-based MMC reinforced with SiC particles. The mechanical properties of the fabricated MMC and the influence on the ECM machining process parameters were analysed with experiments. The mechanical and microstructural properties of AA6061 reinforced with silicon carbide and AA6061 reinforced with boron carbide attained from the enhanced stir casting method were discussed. The influence of the ECM process para- meters current, voltage, electrolyte concentration, feed rate, gap and flow rate on the predominant output para- meters material removal rate, surface roughness and radial overcut were also analysed. 2 SPECIMEN PREPARATION USING AA6061 The matrix material for the study was AA6061. The composite material consists of AA6061 alloy as a matrix material reinforced with three different weight per- centage of SiC and varying weight percentages of B4C (2.5 %, 5 % and 7.5 %) prepared through the enhanced stir-casting technique. The SiC has better mechanical properties such as high hardness, low density and retains its properties even at higher temperatures. The B4C is a hard reinforcement particle that has neutron absorbing characteristics in nature. The percentage of silicon content in AA6061 is high compared to other aluminium alloys and its melting point is low. The chemical com- position of AA 6061 by weight percentage is Cu 0.1 %, Mg 0.4 %, Si 10 %, Mn 0.3 %, Ni 0.1 %, Zn 0.1 %, Pb 0.05 % and Sn 0.2 %.The average particle size of the SiC is 200 mesh, with a density of 3.2 g/cm–3 and a thermal conductivity 3.2 W cm–1 K–1 and also an average particle size of the B4C is 200 mesh. The test specimens were prepared using the simplest and most commercially used technique known as T. M. CHENTHIL JEGAN et al.: MACHINING PARAMETERS INFLUENCING IN ELECTRO CHEMICAL MACHINING ... 952 Materiali in tehnologije / Materials and technology 50 (2016) 6, 951–960 enhanced stir-casting technique. In the stir-casting process the pre-heated ceramic particles were mixed with a vortex of molten alloy created by the rotating impeller. As a result of the interaction between the suspended ceramic particles and the moving solid-liquid interface during solidification, there was a possibility of inhomo- geneity in the reinforcement distribution. Generally, it is possible to incorporate up to 30 % of ceramic particles in the size range 5 μm to 100 μm in a variety of molten aluminium alloys. AA6061 was placed inside the cru- cible and the temperature was set 1000 °C. Some 1 % of the degasser Hexa Chloro Ethane was added to the melted AA6061. The molten metal was stirred and the crucible was held with forks to eliminate the gases. The same process was repeated for various weight percen- tages of SiC-reinforced AMC and B4C-reinforced AMC. 3 ELECTROCHEMICAL MACHINING For this experiment the whole work was carried out with an ECM set up having a power supply of 415 V, 3-phase AC, 50 Hz and it consists of three major sub systems: the machining cell, the control panel and the electrolyte circulation tank. The parameter that is able to change the output parameters by increasing and decreas- ing the level is known as the controlling parameters or the input parameters. In this paper the considered input parameters are the current, voltage, electrolyte concen- tration, feed rate, gap and flow rate. The output para- meters considered are the MRR, SR and ROC. The various levels of input parameters selected for the machining are listed in Table 1. Table 1: Input parameters and levels Tabela 1: Vhodni parametri in njihovi nivoji Variables Values of different levels 1 2 3 4 5 Current (A) 90 120 150 180 210 Voltage (V) 8 10 12 14 16 Electrolyte concentration (g/L) 3.34 6.67 10 13.34 16.67 Feed rate (mm/min) 0.1 0.2 0.3 0.4 0.5 Gap (mm) 0.1 0.2 0.3 0.4 0.5 3.1 Machining process The machine cell has a tool area of 300 mm2, a cross-head stroke 150 mm, a job holder 100 mm open- ing, 50 mm depth and 100 mm width. A DC servo-type tool feed motor was used for the tool movement. The control panel consists of an electrical output rating ranging from 0 A to 300 A DC for any voltage from 0 μm to 20 V, tool feed of 0.2 to 2 mm/min, while the supply given to the machining was 3-phase AC with 50 Hz. The specimen to be machined was fixed in the machine vice. The tool was brought near the job with the help of press buttons provided on the control panel and a table-lifting arrangement, maintaining a particular gap. The tool progress was maneuverered vertically by the servomotor and is governed by a microcontroller-based programmable drive. The c cathode tool is made of non-reacting copper material. The process parameters like current, voltage, electrolyte concentration, feed rate, gap and flow rate were set. The process was started in the presence of an electrolyte flow. This electrolyte flow was adjusted using a flow-control valve. After the desired time interval, a hooter gives an indication of the completion of the time and the process. The specimen prepared was a cylindrical blank of 16 mm in diameter and 32 mm in height. The electrolyte composition used was NaCl solution. 4 RESULTS AND DISCUSSION The hardness of the specimens was measured by a Rockwell hardness and a tensile test. The Rockwell hardness number measures the overall response of the material and it is relatively insensitive to localized effects. The Rockwell scale is a hardness scale based on the indentation hardness of a material. The Rockwell test determines the hardness by measuring the depth of the penetration of an indenter under a large load compared to the penetration made by a preload. The chief advantage of Rockwell hardness is its ability to display hardness values directly, thus obviating tedious calculations involved in other hardness-measurement techniques. The Rockwell Hardness number for AA 6061 and the other experimental MMC are shown in Figure 1. It is clear that the hardnesses of the MMCs are closer to one another, but the values are high compared with AA 6061. For microstructure analysis the machined samples were polished using silicon carbide paper (60, 80, 120, 220 and 400) grit and finally using a soft cloth with fine alumina powder as a slurry. Kerosene was used for cleaning and polishing to prevent the embedding of foreign particles in the sample. The samples were then etched using the modified diamond paste for 140 s. The long etching time was due to the large oxide content of T. M. CHENTHIL JEGAN et al.: MACHINING PARAMETERS INFLUENCING IN ELECTRO CHEMICAL MACHINING ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 951–960 953 Figure 1: Rockwell hardness number for experimental samples Slika 1: Rockwell trdota preizkusnih vzorcev the aluminium powder. The light microscope photo- graphs of the seven combinations of composite samples using an optical microscope of 20× lens are shown in Figure 2. Figure 2a shows various elements present in AA6061. Figures 2b to 2d represents AA6061+ rein- forced silicon carbide particles. AA6061+ reinforced boron carbide particles are shown in Figures 2d to 2f. It can be shown that the silicon carbide particles and boron carbides are homogenously distributed in the aluminium matrix. Scanning electron microscopy (SEM) was used to reveal the morphological features in the machined sam- ples. For the study of their microstructure the specimen preparation or polishing is important. The procedure for preparing the specimen involved the selection of the specimen and the mounting of the specimen in the machine, obtaining a flat specimen surface, intermediate and fine grinding, rough polishing using diamond pow- der with an oil-soluble paste, fine polishing with alumina powder along with distilled water and etching using dilute hydrofluoric acid. The morphological structure of the AMC was obtained using a ZEISS (NIIST) SEM at an accelerating voltage of 20 kV. The AA6061 speci- mens were mounted with conductive adhesives and coated with gold powder. Then, the SEM images were taken at the centre of the machined surface of the sample specimen. From the above examination both MMCs exhibited similar microstructural features in terms of homogeneous particle distributions associated with similar carbide particle sizes. The SEM analysis of the B4C-reinforced composites and SiC-reinforced composites exhibited very similar microstructures. However, it was possible to observe some large size pores that were not in the B4C-reinforced composites. The distribution of the particles was also altered by the growth of the grains. In the B4C-reinforced MMC microstructure both coarser grains and finer grains were identified, but the grain size T. M. CHENTHIL JEGAN et al.: MACHINING PARAMETERS INFLUENCING IN ELECTRO CHEMICAL MACHINING ... 954 Materiali in tehnologije / Materials and technology 50 (2016) 6, 951–960 Figure 2: Light microscope (20×) images of seven samples showing the distribution of the reinforcement in the matrix: a) AA6061, b) AA6061 + 2.5 % SiC, c) AA6061 + 5 % SiC, d) AA6061 + 7.5 % SiC, e) AA6061 + 2.5 % B4C, f) AA6061 + 5 % B4C, g) AA6061 + 7.5 % B4C Slika 2: Svetlobna mikroskopija (20×) sedmih vzorcev, ki ka`e razporeditev delcev za utrditev v osnovi: a) AA6061, b) AA6061 + 2,5 % SiC, c) AA6061 + 5 % SiC, d) AA6061 + 7,5 % SiC, e) AA6061 + 2,5 % B4C, f) AA6061 + 5 % B4C, g) AA6061 + 7,5 % B4C Figure 3: SEM microstructure of: a) SiC-reinforced AA6061 and b) B4C-reinforced AA6061 Slika 3: SEM-posnetek mikrostrukture: a) AA6061, oja~an s SiC in b) AA6061, oja~an z B4C was less compared to that of SiC. Similarly, the pore size was also less for the B4C-reinforced MMC. The SEM microstructure of the AMCs is shown in Figure 3. From the Figures 3a and 3b it is evident that the distributions of the SiC and B4C particles were homogeneous in the AA6061. Most of the ceramic particles were found within the grain boundaries. The distribution of the particles becomes intra granular. The basic atomic structure of AA6061 is body-centred cubic, so that the B4C particles were bonded together in rein- forcement. Good wettability was also seen in both MMCs, but it was especially high in the B4C. To determine the rate of penetration of each alumi- nium alloy, a series of partial infiltrations were per- formed in order to study the growth rate at short times (less than 2 h). XRD micro-diffraction of the SiC-rein- forced AMC and the B4C-reinforced AMC was shown in Figure 4. The XRD peak list of the SiC-reinforced AA6061 and B4C-reinforced AA6061 are listed in Tables 2 and 3, respectively. Table 2: Peak list for SiC reinforced AA6061 Tabela 2: Seznam vrhov pri AA6061, oja~anem s SiC Pos. 2 (°) Height (cm) FWHM 2 (°) d-spacing (nm) Rel. Int. (%) 28.4044 59.37 0.0816 0.313967 8.57 38.4426 692.59 0.1632 0.233979 100.00 44.6516 307.84 0.1632 0.202778 44.45 47.2508 38.12 0.1224 0.192212 5.50 49.0100 4.93 0.0816 0.185716 0.71 56.0546 19.32 0.2856 0.163931 2.79 57.6437 0.75 0.2448 0.159784 0.11 65.0870 156.16 0.1224 0.143194 22.55 78.1848 161.46 0.2448 0.122159 23.31 82.3912 47.66 0.2040 0.169540 6.88 87.9251 6.83 0.4896 0.110964 0.99 Table 3: Peak List for B4C reinforced AA6061 Tabela 3: Seznam vrhov pri AA6061, oja~anem z B4C Pos. 2 (°) Height (cm) FWHM 2 (°) d-spacing (nm) Rel. Int. (%) 28.5168 55.08 0.1632 0.312755 5.71 36.5506 6.86 0.1224 0.245644 0.71 38.5654 963.85 0.1632 0.233262 100.00 44.8543 372.10 0.1836 0.201909 38.61 47.3974 42.42 0.1632 0.191651 4.40 56.2284 12.82 0.3264 0.163466 1.33 65.2212 206.18 0.0816 0.142932 21.39 76.4518 7.85 0.3264 0.124490 0.81 78.3073 240.23 0.1632 0.121998 24.92 82.5530 64.93 0.2448 0.116766 6.74 83.4711 3.89 0.4080 0.115714 0.40 88.0604 10.43 0.3264 0.110828 1.08 From Figure 4 it was observed that the waves were almost the same and symmetrical. The ambiguous struc- ture was available in both AMCs. As compared to the SiC-reinforced particle, the other one possessed good hardness and an even distribution of particles. The start- ing angle was also high for the B4C-reinforced AA6061. In the SiC-reinforced AA6061, the SiC particle peaks were identified at angles of 28.4044, 44.6516, 49.0100, 57.6437, 65.0870 and 87.9251. The remaining peak values in Table 3 and 4 were related to the AA6061 alloy. Similarly, in the B4C-reinforced AA6061, the B4C particle peaks were obtained at angles of 28.5168, 36.5506, 44.8543, 56.2284, 65.2212, 76.4518 and 88.0604. The ductility of the MMC was improved by the addition of ceramic particles with AA6061. The stress vs. elongation analyses of the AA6061 samples were compared with the AA6061 + SiC MMC, and it was shown in Figure 5. From Figure 5a it was noted that the strength of the specimen increases with an increase in the addition of silicon carbide particles. The maximum strength was obtained for sample 3, with aluminium and SiC, in a ratio of 100:7.5. Similarly, for the AA6061 + B4C MMC samples the results were shown in Figure 5b. In AA6061 + B4C MMC also better tensile strength was obtained with high percentage of B4C reinforcement. From the analysis it was concluded that the ductile nature of AA6061 increased with the addition of SiC and B4C particles. T. M. CHENTHIL JEGAN et al.: MACHINING PARAMETERS INFLUENCING IN ELECTRO CHEMICAL MACHINING ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 951–960 955 Figure 4: XRD patterns of: a) SiC-reinforced AA6061 and b) B4C-reinforced AA6061 Slika 4: Rentgenograma: a) AA6061, oja~anega z SiC in b) AA6061, oja~anega z B4C The tensile strength vs. strain analysis of the samples is shown in Figure 6 from which it is noted that the ultimate strength, plasticity and elasticity increased with the percentage of added ceramics. The ultimate point and yield point of the B4C + AA6061 MMC were high compared to those of the SiC + AA6061 MMC. The ductility of both MMCs were high and it can be treated using unconventional machining process as well as conventional machining, and also the workability of the material was very good. Thus SiC-reinforced AA6061 and B4C-reinforced AA6061 MMC were suited for metal-forming processes such as hot and cold forging, rolling, drawing and spinning. 4.1 Factors affecting ECM process parameters 4.1.1 Factors influencing the material removal rate In ECM the current is the major influencing parameter on the MRR of aluminium-based alloys. The values of the MRR obtained in experimentation with different levels of current for the selected samples were shown in Figure 7a. From Figure 7a it was observed that when the current was less the MRR was also less. The AA6061 + 7.5 % B4C gave a lower MRR compared to AA6061 + 7.5 % SiC. When the current was low the material removal rate was low and it gradually increased with increases in the current. A high current will lead to more material removal and optimum material removal appeared at a current value of 150 A. For better ECM indices, higher accuracy, and a better surface finish, it is essential to choose the proper current density. Low values of current efficiency may indicate a failure to choose the optimum machining conditions that lead to high removal rates and surface roughness. The influence of voltage over MRR for different samples is shown in Figure 7b. From the graph it was clear that the material removal rate was proportional to rate of change of voltage. When the maximum voltage was applied in between the workpiece and the tool, the maximum material can be removed from the workpiece. If the potential differences in between the copper elec- trode and aluminium alloys were very low, minimum material removal occurred from the workpiece. The influence of electrolyte concentration over MRR is shown in Figure 7c. For the consideration of electro- lyte concentration, best material removal occurred at a 10 g/L concentration. A very low electrolyte concen- tration produced minimum material removal due to the lack of ionic particles present in the electrolyte. At high range values the ionic concentration was high and the reaction phase at this stage was lagging. Hence, a lower rate of material removal was caused at a high concen- tration of electrolyte than with the medium electrolyte concentration of 10 g/L. The influence of feed rate on MRR is shown in Figure 7d. The directly proportional effect was shown in discussion of the feed rate of the tool on the aluminium alloys. When the feed rate was 0.1 mm/min, then mate- rial removal was minimum. This process gradually increased with increases in the feed rate. It was a maxi- mum at the maximum feed rate of 0.5 mm/min. The influence of the gap on MRR is shown in Figure 7e. The gap plays a vital role in electrochemical ma- chining. In electrochemical machining there is no direct physical contact in between the copper electrode and the aluminium workpiece. The reaction was carried out in the presence of an electrolyte that was circulated in bet- ween the tool and the workpiece. Therefore, to set the correct gap is an important factor to produce high mate- rial removal from the workpiece. In our experiment the maximum material removal could be obtained from the gap with an optimal value at 0.3 mm. Only a minimum amount of material was removed from the workpiece at a T. M. CHENTHIL JEGAN et al.: MACHINING PARAMETERS INFLUENCING IN ELECTRO CHEMICAL MACHINING ... 956 Materiali in tehnologije / Materials and technology 50 (2016) 6, 951–960 Figure 6: Stress strain diagram of MMC samples Slika 6: Diagram odvisnosti napetost-raztezek MMC vzorcev Figure 5: Stress vs. elongation of: a) SiC-reinforced AA6061 and b) B4C-reinforced AA6061 Slika 5: Napetost v odvisnosti od raztezka: a) AA6061, oja~an z SiC in b) AA6061 oja~an z B4C gap of 0.5 mm. If the gap is a maximum then a lack of reaction will be carried out in between the tool and the minimum MRR will be obtained in the workpiece. The influence of flow rate on MRR is shown in Figure 7f. The electrolyte flow rate across the tool and the workpiece stamped the noted impressions. During machining the chips were formed on the surface of the workpiece. Electrolyte flow plays the major role in the material removal process with the removal of contami- nated chip particles presented on the surface of the machining area. If the flow rate of the electrolyte is too low to splash out the removed chips material on the ma- chined surface, then the minimum amount of material will be removed from the workpiece. From Figure 7f it is clear that the material removal rate was increased by increasing the flow of electrolyte. The maximum flow rate can be obtained from 8 L/min. Beyond this level it was slightly reduced due to the high flow of electrolyte, which will cause a lean electrolyte concentration. 4.1.2 Factors influencing surface roughness Generally the ECM process is used for the machining of hard materials with good surface finish. But material removal rate and surface roughness are inversely propor- tional to each other. Therefore, if the material removal rate is high, a poor surface finish will be obtained in the machining process. The influence of current on the SR for the selected samples is shown in Figure 8a. It was found that there was an inverse effect between the surface roughness and the current. At low current conditions a good surface finish was obtained and it was gradually decreased with an increase in the current. At high current conditions the least surface roughness was obtained. The best value of surface roughness was achieved at minimum material removal and with respect to the hardness of material. The influence of voltage on the SR for AMCs is shown in Figure 8b. If the potential difference in bet- ween the electrode and the workpiece is very high, then a poor surface roughness can be achieved from the alumi- mium workpiece. At low voltage the surface finish was high, and it was slightly increased at 10 V. Then it was smoothly reduced with an increase in the voltage and finally poor surface roughness occurred due to the maximum voltage level. The SR obtained at the different levels of electrolyte concentration for selected samples is shown in Figure 8c. In the figure the surface roughness and electrolyte concenration were directly proportional to each other. If the electrolyte concentration was less, a poor surface finish obtained. It is increased with increasing the con- centration of electrolyte. A high range of surface rough- ness can be achieved at a high electrolyte concentration. It was mainly due to the lubrication and good reaction between the electrode and the workpiece. The influence of feed rate on SR is shown in Figure 8d. For considering the surface roughness feed rate and inter electrode gap plays the same role. A good surface T. M. CHENTHIL JEGAN et al.: MACHINING PARAMETERS INFLUENCING IN ELECTRO CHEMICAL MACHINING ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 951–960 957 Figure 7: Influencing parameters with MRR: a) current, b) voltage, c) electrolyte concentrations, d) feed rate, e) gap, f) flow rate Slika 7: Parametri, ki vplivajo na MRR: a) tok, b) napetost, c) koncentracija elektrolita, d) hitrost podajanja, e) re`a, f) hitrost pretoka finish was obtained with a proper gap and a minimum feed rate. The feed rate was increased when the surface roughness of the workpiece was reduced. The value of SR for different levels of gap is plotted and shown in Figure 8e. The optimal value of the inter electrode gap was maintained for a better surface rough- ness. The minimium electrode gap produced the mini- mum material removal rate and hence provided a good surface finish. A high surface roughness was obtained for a 0.2 mm gap and the surface finish was reduced with increases in the gap between the tool and the workpiece. The electrolyte flow rate governs the surface finish in the right way. The influence of flow rate on SR was shown in Figure 8f. A very low flow rate creates a poor surface finish due to the non-flushing of the chip mate- rials on the surface of the workpiece. A good surface roughness was obtained at a 6 L/min flow of electrolyte. The flushing pressure was high for very high electrolyte flow condtions like 10 L/min. A very high flushing pressure tends to produce a vertex on the machining zone, and it will produce abrasive action on the surface of the machined zone. Hence, under high flow rate con- ditions, the surface roughness was compatively low. 4.1.3 Factors influencing radial overcut The ROC values obtained at different levels of current for the selected samples are shown in Figure 9a. The ROC was increased with an increase in the current. If the current was low then the minimum material was removed from the surface of the machined zone. Hence, the ROC was low. When the current was increased gra- dually, then the ROC was increased. When a high current was applied in between the tool and the workpiece, then the maximum amount of material was removed from the workpiece and a high ROC was produced. The influence of the applied voltage on the ROC is shown in Figure 9b. When the applied voltage between the electrodes was high, a high ROC was obtained. The optimal value of the ROC was achieved at a 12 V poten- tial difference. At very low and very high voltages the flow of electrons through the electrodes is improper and hence it will tend to produce a variable machining rate. So the extreme levels of voltages were produced maxi- mum ROC and minimum ROC was happened at mid- range values. The effects of the electrolyte concentration on radial overcut for different AMCs are shown in Figure 9c. When the electrolyte concentration was high, the ROC was low. High lubrication was provided in a high con- centration of electrolyte solutions and hence a better surface finish was obtained. The concentrated electrolyte solution has a good number of ions that will act as the catalyst to improve the reaction between the electrode and the workpiece. In the above condition an enormous amount of hydrogen was generated and an unwanted reaction in the machining zone was avoided. So, the minimum ROC was obtained with high electrolyte con- centrations. The influence of feed rate on the ROC is shown in Figure 9d. The material removal rate and the ROC were directly proportional to each other. The higher feed rate of the tools produced a poor surface roughness and a high material removal rate. The higher rate of material removal from the workpiece increased the ROC. The T. M. CHENTHIL JEGAN et al.: MACHINING PARAMETERS INFLUENCING IN ELECTRO CHEMICAL MACHINING ... 958 Materiali in tehnologije / Materials and technology 50 (2016) 6, 951–960 Figure 8: Influencing parameters with SR: a) current, b) voltage, c) electrolyte concentrations, d) feed rate, e) gap, f) flow rate Slika 8: Parametri, ki vplivajo na SR: a) tok, b) napetost, c) koncentracija elektrolita, d) hitrost podajanja, e) re`a, f) hitrost pretoka ROC was high at a very high range of the tool feed on the workpiece. The ROC values obtained in the different gaps for the selected samples were plotted in Figure 9e. The inter electrode gap was maintained with an optimal range of values. If the gap between the tool and the workpiece was too low, then the contact took place and the very minimum material was removed from the work- piece. Hence, the ROC in the low gap was a minimum. The spark in between the tool and the workpiece material might be produced because of an improper gap and irregular machining on the workpiece. Due to the irregular machining the ROC was very high. The influence of flow rate on ROC is shown in Figure 9f. When the flow of the electrolyte was very low then the sludge was stagnated at the machining zone. Further machining was on the chip with a new machin- ing surface and a poor surface finish leads to high ROC. Similarly, the high velocity of the stream of electrolyte created the flow of the abrasive action with sludge chips and hence the ROC was increased. The optimal flow of electrolyte was needed for the proper removal of sludge formation and good machining accuracy. 5 CONCLUSION In this paper the microstructure and mechanical pro- perties of AA6061 reinforced with silicon carbide and AA6061 reinforced with boron carbide obtained from an enhanced stir-casting method was investigated. Three different weight percentages, i.e., 2.5 %, 5 % and 7.5 %, of silicon carbide and boron carbide were used for the reinforcement. The presence of silicon- and boron-based particles employs a key role in the microstructure deve- lopment of the composites. The ductility of the AA6061 was increased in both the reinforced AMCs with signi- ficantly good tensile properties. The microhardness was markedly influenced by both the alloying elements. From the study it was observed that the selected AMCs were suitable for unconventional machining process like electrochemical machining. In ECM the input parame- ters influence the output variables Material Removal Rate (MRR), Surface Roughness (SR) and Radial Over- cut (ROC). The higher input parameters increased the Material Removal Rate. The higher value of Material Removal Rate during the machining produced poor-sur- face-quality materials. The RO was also in close agree- ment with the SR. Acknowledgement The author gratefully acknowledges the contributions of National Institute for Interdisciplinary Science and Technology Trivandrum, National Institute of Techno- logy Trichy and Annamalai University Chithambaram. 6 REFERENCES 1 K. P. Rajurkar Sekar. T. Marappan, R. Experimental investigations into the influencing parameters of electrochemical machining of AISI 202, Journal of Advanced Manufacturing Systems, 7 (2008) 2, 337–43, doi:10.1142/S0219686708001486 2 A. DeBarr, D. A. Eand Oliver, Electro-chemical Machining, Mac- donald & Co. Ltd, 1968 3 Metals Handbook, Vol. 16, Machining, ASM International, Materials Park, OH 1989 T. M. CHENTHIL JEGAN et al.: MACHINING PARAMETERS INFLUENCING IN ELECTRO CHEMICAL MACHINING ... 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CHENTHIL JEGAN et al.: MACHINING PARAMETERS INFLUENCING IN ELECTRO CHEMICAL MACHINING ... 960 Materiali in tehnologije / Materials and technology 50 (2016) 6, 951–960 B. SENER, H. KURTARAN: MODELING THE DEEP DRAWING OF AN AISI 304 STAINLESS-STEEL ... 961–965 MODELING THE DEEP DRAWING OF AN AISI 304 STAINLESS-STEEL RECTANGULAR CUP USING THE FINITE-ELEMENT METHOD AND AN EXPERIMENTAL VALIDATION MODELIRANJE GLOBOKEGA VLEKA PRAVOKOTNE ^A[E IZ AISI 304 NERJAVNEGA JEKLA Z METODO KON^NIH ELEMENTOV IN Z EKSPERIMENTALNIM PREVERJANJEM Bora Sener1, Hasan Kurtaran2 1Yildiz Technical University, Department of Mechanical Engineering, Istanbul, Turkey 2Gebze Technical University, Department of Mechanical Engineering, Kocaeli, Turkey borasener84@gmail.com, borasen@yildiz.edu.tr Prejem rokopisa – received: 2015-09-06; sprejem za objavo – accepted for publication: 2015-11-16 doi:10.17222/mit.2015.278 In this paper the deep drawing of a rectangular cup from AISI 304 stainless steel sheet was investigated numerically and experi- mentally. The finite-element method was used for computer modeling of the deep-drawing process. The thickness distribution predicted from the finite-element analysis was compared with experimental measurements. It was observed that the numerical results agree well with the experimental values. The minimum thickness was observed at the punch radius in both the simulation and experiment. Keywords: deep drawing, rectangular cup, finite element, stainless steel V ~lanku je bil numeri~no in eksperimentalno preiskovan globoki vlek pravokotne ~a{e iz plo~evine iz nerjavnega jekla AISI 304. Za ra~unalni{ko modeliranje procesa globokega vleka je bila uporabljena metoda kon~nih elementov. Iz analize kon~nih elementov napovedana razporeditev debeline je bila primerjana z eksperimentalnimi meritvami. Preiskava je pokazala, da se numeri~ni rezultati dobro ujemajo z eksperimentalnimi vrednostmi. Najmanj{a debelina je bila opa`ena na radiusu pesti~a, tako pri simulaciji kot tudi pri eksperimentu. Klju~ne besede: globoki vlek, pravokotna ~a{a, kon~ni element, nerjavno jeklo 1 INTRODUCTION The rectangular/square-cup deep-drawing process has specific forming characteristic. Non-uniform material flow and quite complicated deformation mechanism are seen in this process. Therefore, the deep drawing of square and rectangular cups is more difficult than that of some other shapes, such as circular cups. Many resear- chers investigated the rectangular/square-cup deep-draw- ing process experimentally and numerically. A. G. Ma- malis et al.1 investigated the effect of material and forming characteristics on the simulation of the deep drawing of square cups by using the explicit non-linear finite-element code DYNA-3D. They considered the effect of material density, punch velocity and friction coefficient. L. F. Menezes and C. Teodosiu2 studied the square-cup deep-drawing process numerically and expe- rimentally. They modeled the process by using solid elements and compared the numerical results with the experiment. E. Bayraktar and S. Altintas3 investigated the square-cup deep-drawing process and 2D-draw bending process of Hadfield steel experimentally. They evaluated the draw-in values of the flange, the principal strains in the square-cup deep drawing and compared the experimental results with that those of mild steel. Y. Ma- rumo and H. Saiki4 studied differential lubrication methods in the square-cup deep-drawing process in order to prevent any deformation concentration on the corners. Y. Harada and M. Ueyama5 investigated the drawability of pure titanium sheets in the square-cup deep-drawing process. Titanium sheets were coated by heat oxide and formed into a square with a punch. L. M. A. Hezam et al.6 developed a new technique for the deep drawing of square cups made from brass and pure aluminum. They improved the material flow by using a conical die with a square aperture at its end without a blank holder. M. Gavas and M. Izciler7 designed a blank holder with a spiral spring to reduce the friction area between the blank and the blank holder during the deep drawing of square cups. A higher drawing height, a homogenous thickness distribution and minimum earing cups were obtained in their study. M. A. Hassan et al.8 have deve- loped a new divided blank holder with a tapered base and eight tapered segments to increase the deep drawability of the square cups. They improved the drawability of thin sheets and foils and increased the limiting drawing ratio with this technique over the conventional techni- ques. L. P. Lei et al.9 studied the square-cup deep-draw- ing process for 304 stainless-steel sheet numerically and experimentally. They evaluated the effect of the blank Materiali in tehnologije / Materials and technology 50 (2016) 6, 961–965 961 UDK 621.77:004.946:669.14.018.8:620.1 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)961(2016) shape on the material flow. J. H. Lee and B. S. Chun10 investigated the effect of temperature, blank shape and holding pressure on the deep drawability of a square cup from 304 stainless-steel sheet experimentally and nume- rically. F. K. Chen and S. Y. Lin11 examined the effects of process parameters such as punch radius, die radius, die corner radius, die gap and the length-to-width ratio by both the finite-element method and the experimental approach. The authors formulated a formability index that serves as a design rule for the rectangular cup draw- ing from 304 stainless-steel sheet. Although many signi- ficant studies are carried out about the rectangular/ square-cup deep-drawing process, they are generally limited to the forming of Al and Al alloys. Very little of the literature is devoted to the rectangular-cup deep drawing of 304 stainless-steel sheets. The available literature on the rectangular-cup deep drawing of 304 stainless steel sheets is limited to warm forming. Therefore, the cold forming of a rectangular cup from 304 stainless steel sheet is investigated numerically and experimentally in this study. 2 EXPERIMENTAL PROCEDURE A die with a rectangular aperture, a conical punch that has flat surface with a rectangular shape and a circular blank holder that has a rectangular cavity are used in this study. The dimensions of the rectangular-cup tooling are given in Table 1. Deep-drawing experiments were carried out on a 160-ton-capacity double-action hydraulic press. The punch is mounted to the lower shoe. The blank holder slides around the punch in the lower shoe. The upper shoe consists of the die. In the press, the lower shoe is mounted on the press bed and the upper shoe is attached to the press ram. The blank holder is supported by the cushion pins that apply the blank hold- ing force to the blank holder during the forming process. During the deep-drawing process, the blank holder is raised to the top-most level. The blank is positioned on the blank holder. The clamped blank with the die and blank holder moves further down and forms the blank against the stationary punch under the action of the blank holder force through the cushion pins. The experimental set-up is shown in Figure 1. Table 1: Tool dimensions Tabela 1: Dimenzije orodja Die Radius, mm 35 Edge length of rectangular cavity, mm 129 × 143 Corner radius of the rectangular cavity, mm 47 Depth of the rectangular portion, mm 27 Punch Radius, mm 20 Rectangular side length, mm 120 × 134 Cone angle, deg 2 Blank holder Diameter, mm 351 Edge length of rectangular cavity, mm 127.5 × 142 Corner radius of the rectangular cavity, mm 45.5 An austenitic grade AISI 304 stainless steel was used in this study. The thickness of the material was 0.8 mm. The mechanical properties of the material are explained in Section 3. The initial blank of diameter 335 mm was drawn to a rectangular cup of height 80 mm. A 340-kN blank holder force was applied in the experiments. The operating speed was 20 mm/s. A rectangular cup from 304 stainless-steel sheet was successfully drawn using the this blank holder force as shown in Figure 2. 3 FINITE ELEMENT MODEL In the present work, the explicit non-linear finite-ele- ment (FE) code DYNAFORM 5.9.2 software is used for B. SENER, H. KURTARAN: MODELING THE DEEP DRAWING OF AN AISI 304 STAINLESS-STEEL ... 962 Materiali in tehnologije / Materials and technology 50 (2016) 6, 961–965 Figure 2: Rectangular cup Slika 2: Pravokotna ~a{a Figure 1: Experimental set-up Slika 1: Eksperimentalni sestav simulating the rectangular-cup deep-drawing process. The blank is meshed with 3333 quadrilateral elements and 3434 nodes. A Belytschko-Tsay shell element with five integration points across the thickness is used for the shell mesh of the blank. Because of the symmetry, only a quarter model is employed in the numerical simulation, as shown in Figure 3. The punch, die and the blank holder were modeled as rigid objects because of their high stiffness, while the blank was modeled as a deform- able body. A forming-one-way-surface-to-surface con- tact algorithm is used in the analysis. The friction coeffi- cients are assumed to be 0.11 for the contact between the tools (die, punch and blank holder) and the blank. This value was recommended by a previous investigation.12 The die speed employed is 1000 mm/s, which is extrem- ely slow compared to the typical wave speeds in the ma- terials to be formed (the wave speed in steel is approxi- mately 5000 m/s). In general, inertia forces will not play a dominant role for forming rates that are considerably higher than the nominal 1000 mm/s rates in the physical problem.13 The displacement of the die was taken as 80 mm, which is decided by the height of the cup. A 85-kN constant blank holder force is applied in the model (quarter of the experimental value). AISI 304 stainless-steel sheet is used in the simulation work. It is assumed that the material is isotropic and homogeneous. The strain-hardening model used is isotropic hardening. The mechanical properties of the material were determined by tensile testing. Different constitutive equations, such as Holloman, Swift, and Ludwick, were evaluated in order to represent the plastic behavior of the AISI 304 stainless steel. The nonlinear least-squares method and a trust-region algorithm were used in determining the material parameters. It was found that the Ludwick equation was the best fit to the experimental data for the AISI 304 stainless steel. This result agrees with the literature14. The flow curve was extrapolated to higher strains using this equation and was used in the simulation. A comparison of these different hardening models with experimental data is shown in Figure 4. The mechanical properties of the materials are reported in Table 2. Table 2: Mechanical properties of the material Tabela 2: Mehanske lastnosti materiala Parameters Units Value Young’s modulus (E) GPa 256.86 Poisson’s ratio (v) 0.28 Yield strength ( y) MPa 308.94 Strain hardening (n) 0.3 Coefficient of strength (K) 1528 4 RESULTS AND DISCUSSION The thickness distribution is one of the major quality characteristics in the sheet-metal formed part. Therefore, the thickness distribution of the AISI 304 stainless-steel sheet in the deep-drawing process was investigated theoretically and experimentally. The drawn component was cut along the diagonal direction and the thickness of the part along this direction was measured using a micro- meter, as shown in Figure 5. For the verification of the FEM results, the thickness variations predicted by the numerical simulation were compared with the experimental results. Figure 6 shows the comparison of the FE predictions with the experi- ment for the formed part along section YO (diagonal). It could be observed from Figure 6 that the thickness B. SENER, H. KURTARAN: MODELING THE DEEP DRAWING OF AN AISI 304 STAINLESS-STEEL ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 961–965 963 Figure 5: Location of the diagonal section in the formed part for evaluation of the thickness distribution Slika 5: Prikaz diagonalnega preseka preoblikovanega dela za oceno razporeditve debeline Figure 3: Finite-element model Slika 3: Model kon~nega elementa Figure 4: Comparison of the hardening models with experimental data Slika 4: Primerjava modelov utrjevanja z eksperimentalnimi podatki distribution predicted by the FE simulation along section YO agrees with the experiment. A minimum thickness was observed at the punch radius for both the simulation and the experiment. This is due to the appearance of the biaxial stretching state in this region. Hence, it is obvious that the potential failure site in the deep drawing of the part is located in the vicinity of the punch radius where the thinning is a maxi- mum. This phenomenon is also observed by S. Gallée12 during the deep drawing of stainless-steel sheet. Although the trend of thickness distribution predicted by the simulation matches with the experiment, more thinning was observed in the simulation (32 %) than in the experiments (26 %). The difference in thinning could be attributed to the insufficiency of the tensile test. The flow stress available from the tensile test was limited to small strains at low strain rates. The flow curve of the material was obtained up to the 0.21 plastic strain value in the tensile test. Beyond this value, the hardening curve was extrapolated to high strains using the Ludwick equa- tion. These values were probably overestimated. Thus, more restraining force to draw the material from the flange resulted in a localized thinning of 32 % in the simulation. 5 CONCLUSIONS Finite-element simulations and deep-drawing experi- ments of the AISI 304 stainless-steel sheet were carried out. The conclusions can be summarized as follow: • Different constitutive equations were evaluated to represent the plastic behavior of the AISI 304 stainless steel sheet. The Ludwick equation was the best fit to the experimental data for this steel. • Good agreement was obtained between the experi- mental and finite-element results. Maximum thinning was observed at the punch radius in both the simula- tion and experiment. This is because of the appear- ance of the biaxial stretch state at the punch radius. • Predicted thinning values were larger than the experi- mental data at the punch radius. More thinning observed in the simulation could be due to the extra- polation of the flow stress to higher strains in the software and overestimated. Acknowledgements The authors would like to thank Mr. Kani Yilmaz for his help with the experimental set-up. 6 REFERENCES 1 A. G. Mamalis, D. E. Manolakos, A. K. Baldoukas, Simulation of sheet metal forming using explicit finite element techniques: effect of material and forming characteristics, Journal of Materials Processing Technology, 72 (1997), 110–116, doi:10.1016/S0924-0136(97) 00137-4 2 L. F. Menezes, C. Teodosiu, Three-dimensional numerical simulation of the deep-drawing process using solid finite elements, Journal of Materials Processing Technology, 97 (2000), 100–106, doi:10.1016/ S0924-0136(99)00345-3 3 E. Bayraktar, S. Altintas, Square cup deep drawing and 2D-draw bending analysis of Hadfield steel, Journal of Materials Processing Technology, 60 (1996), 183–190, doi:10.1016/0924-0136(96) 02326-6 4 Y. Marumo, H. Saiki, Evaluation of the forming limit of aluminum square cups, Journal of Materials Processing Technology, 80-81 (1998), 427–432, doi:10.1016/S0924-0136(98)00196-4 5 Y. Harada, M. Ueyama, Formability of Pure Titanium Sheet in Square Cup Deep Drawing, Procedia Engineering, 81 (2014), 881–886, doi:10.1016/j.proeng.2014.10.092 6 L. M. A. Hezam, M. A. Hassan, I. M. Hassab-Allah, M. G. El-Sebaie, Development of a new process for producing deep square cups through conical dies, International Journal of Machine Tools & Manufacture, 49 (2009), 773–780, doi:10.1016/j.ijmachtools.2009. 04.001 B. SENER, H. KURTARAN: MODELING THE DEEP DRAWING OF AN AISI 304 STAINLESS-STEEL ... 964 Materiali in tehnologije / Materials and technology 50 (2016) 6, 961–965 Figure 6: Comparison of the thickness distribution along section YO (diagonal) from FE simulation with experiments Slika 6: Primerjava razporeditve debeline vzdol` preseka YO (diagonalno) iz FE simulacije z eksperimentalnimi podatki 7 M. Gavas, M. Izciler, Design and application of blank holder system with spiral spring in deep drawing of square cups, Journal of Materials Processing Technology, 171 (2006), 274–282, doi:10.1016/j.jmatprotec.2005.06.082 8 M. A. Hassan, K. I. E. Ahmed, N. A. Takakura, A developed process for deep drawing of metal foil square cups, Journal of Materials Processing Technology, 212 (2012), 295–307, doi:10.1016/ j.jmatprotec.2011.09.015 9 L. P. Lei, S. M. Hwang, B. S. Kang, Finite element analysis and design in stainless steel sheet forming and its experimental compa- rison, Journal of Materials Processing Technology, 110 (2001), 70–77, doi:10.1016/S0924-0136(00)00735-4 10 J. H. Lee, B. S. Chun, Investigation on the variation of deep drawability of STS304 using FEM simulations, Journal of Materials Processing Technology, 159 (2005), 389-396, doi:10.1016/ j.jmatprotec.2004.05.029 11 F. K. Chen, S. Y. Lin, A formability index for the deep drawing of stainless steel rectangular cups, International Journal of Advanced Manufacturing Technology, 34 (2007), 878–888, doi:10.1007/ s00170-006-0659-3 12 S. Gallée, P. Pilvin, Deep drawing simulation of a metastable austenitic stainless steel using a two-phase model, Journal of Mate- rials Processing Technology, (2010), 835-843, doi:10.1016/ j.jmatprotec.2010.01.008 13 F. Ayari, E. Bayraktar, Journal of Achievements in Materials and Manufacturing Engineering, 48 (2011) 1, 64–86 14 N. C. Da Silva, S. A. G. De Oliveria, E. H. Guimaraes, A Compa- rative Study of the Constitutive Equations to Predict Work Hardening Characteristics of Stainless Steels 304 and ACE P439A, 20th International Congress of Mechanical Engineering, Brazil, 2009, 1–7 B. SENER, H. KURTARAN: MODELING THE DEEP DRAWING OF AN AISI 304 STAINLESS-STEEL ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 961–965 965 M. CONRADI, A. KOCIJAN: SURFACE AND ANTICORROSION PROPERTIES OF HYDROPHOBIC ... 967–970 SURFACE AND ANTICORROSION PROPERTIES OF HYDROPHOBIC AND HYDROPHILIC TiO2 COATINGS ON A STAINLESS-STEEL SUBSTRATE POVR[INSKE IN PROTIKOROZIJSKE LASTNOSTI HIDROFOBNIH IN HIDROFILNIH TiO2 PREVLEK NA JEKLENI PODLAGI Marjetka Conradi, Aleksandra Kocijan Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia marjetka.conradi@imt.si Prejem rokopisa – received: 2016-04-21; sprejem za objavo – accepted for publication: 2016-05-03 doi:10.17222/mit.2016.068 We compare the wetting, morphology and anticorrosion properties of fluorosilane-modified TiO2 (FAS-TiO2/epoxy) and as-received TiO2/epoxy coatings. An array of double-layer TiO2 nanoparticles of two sizes (30 nm and 300 nm) were spin coated onto a steel substrate AISI 316L. The static water contact angles were measured to evaluate the wetting properties of the FAS-TiO2/epoxy (hydrophobic) and the as-received TiO2/epoxy (hydrophilic) coatings. The morphology of the coatings was analyzed with average surface roughness (Sa) measurements and SEM imaging. We show that the order of the deposition in a double layer composed of dual-size nanoparticles plays an important role in the surface roughness and hence the wettability. SEM images reveal a typical morphology and Sa difference between the FAS-TiO2/epoxy and the as-received TiO2/epoxy coatings, reflected in the discrepancy of the average size of the agglomerates that are coating the substrate. Potentiodynamic measurements show an enhanced corrosion resistance for the FAS-TiO2/epoxy-coated AISI 316L stainless steel compared to the as-received TiO2/epoxy-coated AISI 316L. Keywords: TiO2, epoxy, coatings, wetting, corrosion V ~lanku primerjamo omo~itvene lastnosti, morfologijo in antikorozijske lastnosti s fluorosilanom oble~enih TiO2 (FAS-TiO2/epoksi) in ~istih TiO2/epoksi prevlek. TiO2 nanodelce dveh velikosti (30 nm in 300 nm) smo na jekleno podlago tipa AISI 316L nanesli s "spin coaterjem". Omo~itvene lastnosti prevlek smo dolo~ili z meritvami stati~nih kontaktnih kotov. Le-te so pokazale hidrofobno naravo FAS-TiO2/epoksi prevlek in hidorfilno naravo ~istih TiO2/epoksi prevlek. Morfolo{ke lastnosti prevlek smo analizirali z meritvami povpre~ne hrapavosti povr{ine (Sa) ter SEM-mikroskopijo. Pokazali smo pomen vrstnega reda nalaganja nanodelcev dveh velikosti na hrapavost povr{ine in njeno omo~ljivost. SEM-posnetki prikazujejo razliko v morfologiji in hrapavosti povr{in FAS-TiO2/epoksi in ~istih TiO2/epoksi prevlek, ki se odra`a v tvorbi aglomeratov razli~nih velikosti na eni in drugi povr{ini. Potenciodinamske meritve ka`ejo izbolj{ano odpornost proti koroziji FAS-TiO2/epoksi prevlek v primerjavi s ~istimi TiO2/epoksi prevlekami na jekleni podlagi tipa AISI 316L. Klju~ne besede: TiO2, epoksi, prevleke, omo~itvene lastnosti, korozija 1 INTRODUCTION Austenitic (AISI) stainless steel is an important engi- neering material because of its generally high corrosion resistance combined with favourable mechanical proper- ties, such as its high tensile strength.1,2 Its high corrosion resistance is attributed to the presence of a passive film, which is stable, invisible, thin, durable and extremely adherent and self-repairing.3 However, in many aggres- sive environments, such as a chloride-ion-rich environ- ment, AISI 316L stainless steel is still observed to suffer from pitting corrosion.4 Therefore, the modification of metallic surfaces using various coatings is an important subject in the field of enhancing particular surface pro- perties, mechanical as well as anticorrosion properties. Epoxy coatings have been widely used for metallic- surface protection because of their good mechanical and electrical-insulating properties, chemical resistance and strong adhesion to heterogeneous substrates. However, the highly cross-linked structure of an epoxy resin often makes epoxy coatings susceptible to the propagation of cracks and damage due to surface abrasion and wear.5 Therefore, a lot of research has been done to improve the performance of epoxy coatings by adding various nano- particles, like SiO2, TiO2, ZnO, CuO etc.6 In addition, nanoparticles also enhance the corrosion-protection pro- perties of the epoxy coatings by decreasing the porosities due to the small size and high specific area. TiO2 nano- particles are well-known anticorrosion additives used in several applications, such as aerospace, marine, bio- medicine, etc. because of their unique physiochemical properties and good chemical stability.7–9 Here we report on a comparison of the surface and anticorrosion properties of double-layer, dual-size (30 nm and 300 nm) FAS-TiO2/epoxy and as-received TiO2/epoxy coatings. We show that the order of the nanoparticle deposition plays an important role in the wetting and the morphological properties of the coatings. Potentiodynamic measurements reveal that the hydro- phobic coating has better anticorrosion properties than the hydrophilic coating. Materiali in tehnologije / Materials and technology 50 (2016) 6, 967–970 967 UDK 67.017:620.193:669.148:669.295 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)967(2016) 2 EXPERIMENTAL PART Materials – Epoxy resin (Epikote 816, Momentive Specialty Chemicals B.V.) was mixed with a hardener Epikure F205 (Momentive Specialty Chemicals B.V.) in the ratio of mass fractions of 100:53 %. TiO2 nano- particles with mean diameters of 30 nm were provided by Cinkarna Celje, whereas the 300-nm particles were provided by US Research Nanomaterials, Inc. Austenitic stainless steel AISI 316L (17 % Cr, 10 % Ni, 2.1 % Mo, 1.4 % Mn, 0.38 % Si, 0.041 % P, 0.021 % C, <0.005 % S in mass fraction) was used as a substrate. Surface functionalization – For the hydrophobic effect, TiO2 particles were functionalized in 1 % of volume fractions of ethanolic fluoroalkylsilane or FAS17 (C16H19F17O3Si) solution. Steel substrate preparation – Prior to the application of the coating, the steel discs of 25 mm diameter and with a thickness of 1.5 mm were diamond polished following a standard mechanical procedure and then cleaned with ethanol in an ultrasonic bath. Coating preparation – To improve the TiO2 nano- particles’ adhesion, the diamond-polished AISI 316L substrate was spin-coated with a 300-nm layer of epoxy (as determined by ellipsometry)10 and then cured for 1 h at 70 °C and post-cured at 150 °C for another hour. The nanoparticles were then coated onto the AISI 316L + epoxy (AISI + E) surface by spin-coating 20 μL of 3 % of mass fractions of TiO2 nanoparticle ethanolic solution. We prepared dual-size, double-layer coatings consisting of 30 nm and 300 nm FAS-TiO2 nanoparticles. Both possibilities of the order of TiO2 nanoparticles were analyzed for the FAS-TiO2/epoxy coatings’ preparation: AISI+E+30+300 and AISI+E+300+30. Finally, the coatings were dried in an oven for approximately 20 min at 100 °C. The same procedure was repeated with the as- received, non-functionalized, TiO2 nanoparticles to pre- pare the TiO2/epoxy coatings. Scanning electron microscopy (SEM) – SEM analysis using a FE-SEM Zeiss SUPRA 35VP was employed to investigate the morphology of the TiO2 coatings’ surfa- ces, which were sputtered with gold prior to imaging. Contact-angle measurements – The static contact- angle measurements of water (W) on the TiO2/epoxy- coated AISI 316L substrates and on the FAS-TiO2/ epoxy-coated AISI 316L substrates were performed using a surface-energy evaluation system (Advex Instru- ments s.r.o.). Liquid drops of 5 μL were deposited on different spots of the substrates to avoid the influence of roughness and gravity on the shape of the drop. The drop contour was analysed from an image of the deposited liquid drop on the surface and the contact angle was determined by using Young-Laplace fitting. To minimize the errors due to roughness and heterogeneity, the average values of the contact angles of the drop were calculated approximately 30 s after the deposition from at least five measurements on the studied coated steel. All the contact-angle measurements were carried out at 20 °C and ambient humidity. Surface roughness – Optical 3D metrology system, model Alicona Infinite Focus (Alicona Imaging GmbH) was employed for the surface-roughness analysis. At least three measurements per sample were performed at a magnification of 20× with a lateral resolution of 0.9 μm and a vertical resolution of about 50 nm. IF-Measure- Suite (Version 5.1) software was used for the roughness analysis. The software offers the possibility to calculate the average surface roughness, Sa, for each sample, based on the general surface roughness equation (Equation 1): Sa L L z x y x y x y LL yx = ∫∫ 1 1 00 ( , ) d d (1) where Lx and Ly are the acquisition lengths of the sur- face in the x and y directions and z(x,y) is the height. The size of the analyzed area was (714×542) μm. To level the profile, corrections were made to exclude the general geometrical shape and possible measurement- induced misfits. Electrochemical measurements – Electrochemical measurements were performed on the TiO2/epoxy-coated AISI 316L stainless steel and on the FAS-TiO2/epoxy- coated AISI 316L stainless steel. The experiments were carried out in a simulated physiological Hank’s solution, containing 8 g/L NaCl, 0.40 g/L KCl, 0.35 g/L NaHCO3, 0.25 g/L NaH2PO4×2H2O, 0.06 g/L Na2HPO4×2H2O, 0.19 g/L CaCl2×2H2O, 0.41 g/L MgCl2×6H2O, 0.06 g/L MgSO4×7H2O and 1 g/L glucose, at pH = 7.8 and 37 °C. All the chemicals were from Merck, Darmstadt, Germany. The measurements were performed by using BioLogic Modular Research Grade Potentiostat/Galva- nostat/FRA Model SP-300 with EC-Lab Software and a three-electrode flat corrosion cell, where the working electrode (WE) was the investigated specimen, the refe- rence electrode (RE) was a saturated calomel electrode (SCE, 0.242 V vs. SHE) and the counter electrode (CE) was a platinum net. The specimens were immersed in the solution 1 h prior to the measurement in order to stabi- lize the surface at the open-circuit potential (OCP). The potentiodynamic curves were recorded, starting the measurement at 250 mV vs. SCE more negative than the open-circuit potential (OCP). The potential was then increased, using a scan rate of 1 mV s–1, until the trans- passive region was reached. 3 RESULTS AND DISCUSSION 3.1 Wetting properties To analyze the surface wettability, we performed five static contact-angle measurements with water (W) on different spots all over the sample and used them to de- termine the average contact-angle values of the coating with an estimated error in the reading of ±1.0°. To fabricate a surface that is as hydrophobic as possi- ble we followed the trend of increasing hydrophobicity based on dual-scale roughness.11,12 For this purpose, the surface roughness was adjusted via spin-coating 30-nm M. CONRADI, A. KOCIJAN: SURFACE AND ANTICORROSION PROPERTIES OF HYDROPHOBIC ... 968 Materiali in tehnologije / Materials and technology 50 (2016) 6, 967–970 and 300-nm FAS-TiO2 nanoparticles onto the flat AISI+E surface. The substrate was consequently modi- fied by self-assembled FAS-TiO2 nanoparticles resulting in micro- to nanoparticle-textured surfaces with a refined roughness structure. The static water contact angles, W, and average surface roughness, Sa, for both possibilities of the dual-size, double-layer, FAS-TiO2/epoxy coatings, (30 + 300) nm and reversed, (300 + 30) nm, are listed in Table 1. The measured contact angles indicate that the surface is more hydrophobic when the bottom layer is composed of 30-nm and the top layer of 300-nm FAS- TiO2. The difference in W between the two coatings is approximately 11° and this behavior can be attributed to the increased roughness implemented by the larger nano- particles on the top, which is reflected in the average surface-roughness measurements, Sa (Table 1). Table 1: Comparison of static water contact angles (W) and average surface roughness (Sa) of dual-size, double-layer FAS-TiO2/epoxy and as-received TiO2/epoxy coatings Tabela 1: Primerjava stati~nih kontaktnih kotov (W) in povpre~ne povr{inske hrapavosti (Sa) dvoplastnih FAS-TiO2/epoksi in ~istih TiO2/epoksi prevlek Contact angle Rough- ness Contact angle Rough- ness FAS-TiO2 TiO2 Substrate W/° Sa/nm W/° Sa/nm AISI+E+30+300 126.0 250.9 80.3 89.8 AISI+E+300+30 115.2 160.2 79.2 95.1 We prepared, in the same manner, a double-layer of (30 + 300) nm and (300 + 30) nm with as-received TiO2, nanoparticles. These coatings are hydrophilic due to the hydroxyl groups on the surface of the as-received TiO2 nanoparticles. The static water contact angles of both possibilities were comparable, around 80°. In addition, the average surface roughness, Sa, was much lower com- pared to the FAS-TiO2/epoxy coatings (Table 1). This result indicates that FAS functionalization significantly changes not only the wetting properties of the coating but also its morphology, as will be shown in the follow- ing section. 3.2 Surface morphology Figure 1 compares the morphology of the double- layer FAS-TiO2/epoxy (a, b) and the as-received TiO2/epoxy (c, d) coatings. SEM images reveal an obvi- ous difference in the morphology between layers of FAS-TiO2 and as-received TiO2 nanoparticles, which is reflected mostly in the different length scale of the average size of the nanoparticle agglomerates and conse- quently in a discrepancy of the average surface rough- ness, Sa, as reported in Table 1. FAS functionalization apparently does not homogenize the particle distribution as the formation of large agglomerates up to a few tenths of microns is observed (Figure 1a and 1b). In contrast, for the as-received TiO2 nanoparticle coatings, the nano- particles are more finely dispersed and agglomerates of the order of few microns are observed (Figure 1c and 1d). SEM images also reveal that TiO2 nanoparticles were not able to cover completely the underlying substrate This might additionally influence the contact-angle va- lues and the wetting properties of the FAS-TiO2 and the as-received TiO2 layers as the epoxy substrate is hydro- philic with a static water contact angle of 74.3°. This effect is, however, probably more pronounced in coatings prepared with hydrophobic FAS-TiO2 nanoparticles, as the uncovered fractions allow the water to impregnate between the nanoparticles and the agglomerates to come into contact with the exposed hydrophilic epoxy and, consequently, reduce the static water contact angle. On the other hand, this effect does not play an important role in the as-received TiO2/epoxy coatings as both the as- received TiO2 nanoparticles and the epoxy are hydro- philic. The role of the order of nanoparticle deposition seems to be more pronounced in the FAS-TiO2/epoxy coatings (Figure 1a and 1b), which is also reflected in the discrepancy in static water contact angles and the average surface roughness between AISI+E+30+300 and AISI+E+300+30, as reported in Table 1. Larger particles on the top seem to create larger agglomerates and con- sequently a rougher surface. The morphology of the as-received TiO2/epoxy coat- ings, AISI+E+30+300 and AISI+E+300+30 (Figure 1c and 1d) is, however, comparable, as are the static water contact angles and the average surface roughness (Table 1). 3.3 Potentiodynamic measurements For an analysis of the anticorrosion properties we chose the more hydrophobic coating, FAS-TiO2/epoxy coating, AISI+E+30+300. The comparison was made M. CONRADI, A. KOCIJAN: SURFACE AND ANTICORROSION PROPERTIES OF HYDROPHOBIC ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 967–970 969 Figure 1: Comparison of surface morphology of double-layer, FAS-TiO2/epoxy (a, b) and as-received, TiO2/epoxy (c, d) coatings Slika 1: Primerjava morfologije dvoplastnih FAS-TiO2/epoksi (a, b) in ~istih TiO2/epoksi (c, d) prevlek with the as-received TiO2/epoxy coating using the same order of particle deposition (30+300). Figure 2 shows the potentiodynamic behaviour of the as-received TiO2/ epoxy-coated AISI 316L and FAS-TiO2/epoxy-coated AISI 316L stainless steel in a simulated physiological Hank’s solution. We studied the polarization and the passivation behaviour of the tested material after the surface modification. After 1 h of stabilization at the OCP, the corrosion potential (Ecorr) for the as-received TiO2/epoxy-coated AISI 316L in Hank’s solution was approximately –0.13 V vs. SCE. Following the Tafel region, the alloy exhibited a broad range of passivity. The breakdown potential (Eb) was approximately 0.25 V vs. SCE. In the case of the FAS-TiO2/epoxy-coated AISI 316L stainless steel, the Ecorr in Hank’s solution was approximately –0.27 V vs. SCE. The passivation range was significantly broader, i.e., 0.4 V vs. SCE, and at lower corrosion-current densities compared to TiO2/ epoxy-coated AISI 316L specimen. The results show an enhanced corrosion resistance for the FAS-TiO2/epoxy- coated AISI 316L stainless steel compared to as-received TiO2/epoxy coated AISI 316L. 4 CONCLUSIONS We analyzed the wettability behavior of double-sized, double-layer, FAS-functionalized TiO2 and as-received TiO2 nanostructured surfaces. We showed that the order of the TiO2 nanoparticle deposition determines the sur- face roughness and hence the wettability, as confirmed by the average surface-roughness measurements and the SEM imaging. This effect was more pronounced in coat- ings with FAS-TiO2 nanoparticles. The morphology analysis also revealed a typical morphology and Sa diffe- rence between the FAS-TiO2/epoxy and the as-received TiO2/epoxy coatings reflected in a discrepancy in the average size of the agglomerates that are coating the sub- strate. The corrosion stability of double-sized, double- layer, FAS-functionalized TiO2 and the as-received TiO2 nanostructured coatings on the surface of the AISI 316L stainless steel was studied in a simulated physiological Hank’s solution. The results showed the superior corro- sion stability of the FAS-TiO2/epoxy-coated AISI 316L stainless steel compared to the as-received TiO2/epoxy- coated AISI 316L. Acknowledgement This work was carried out within the research project J2-7196: "Antibakterijske nanostrukturirane za{~itne pla- sti za biolo{ke aplikacije" of the Slovenian Research Agency (ARRS). 5 REFERENCES 1 M. A. M. Ibrahim, S. S. A. El Rehim, M. M. 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KOCIJAN: SURFACE AND ANTICORROSION PROPERTIES OF HYDROPHOBIC ... 970 Materiali in tehnologije / Materials and technology 50 (2016) 6, 967–970 Figure 2: Potentiodynamic curves for as-received TiO2/epoxy- and FAS-TiO2/epoxy-coated AISI 316L substrate in a simulated physiological Hank’s solution Slika 2: Potenciodinamske krivulje ~istih TiO2/epoksi in FAS-TiO2/ epoksi prevlek na AISI 316L podlagi, izmerjene v simulitani fiziolo{ki Hankovi raztopini B. ARH et al.: ELECTROSLAG REMELTING: A PROCESS OVERVIEW 971–979 ELECTROSLAG REMELTING: A PROCESS OVERVIEW ELEKTROPRETALJEVANJE POD @LINDRO – PREGLED PROCESA Bo{tjan Arh, Bojan Podgornik, Jaka Burja Institute of Metals and Technology, Lepi pot 11, Ljubljana, Slovenia bostjan.arh@imt.si Prejem rokopisa – received: 2016-06-15; sprejem za objavo – accepted for publication: 2016-09-02 doi:10.17222/mit.2016.108 The electroslag remelting process (ESR) is important because it provides better control of the solidification microstructure and chemical homogeneity; it also enables greater cleanliness and better mechanical properties. The manufactured high-alloyed steels and other alloys with a controlled chemical composition are used in aerospace, in thermal- and nuclear-power plants, in chemical engineering, for military equipment, special tools, etc. An overview and the basics of the ESR process are presented in this paper. Keywords: electroslag remelting, solidified microstructure, chemical homogeneity, clean steel Postopek elektropretaljevanja kvalitetnih jekel pod `lindro (EP@) ima velik pomen zaradi sposobnosti nadzora strjevalne strukture in kemijske homogenosti saj postopek omogo~a doseganje ve~je ~istosti jekla in bolj{e mehanske lastnosti. Tako izdelana visokolegirana jekla, in zlitine z garantirano kemijsko sestavo in lastnostmi, se namensko uporabljajo v letalstvu, termoelektrarnah in jedrskih elektrarnah, v kemi~ni industriji, v medicini, za voja{ko opremo, za specialna orodja, itd. V pri- spevku so opisane osnove tehnologije pretaljevanja zlitin pod `lindro. Klju~ne besede: elektropretaljevanje pod `lindro, strjevalna struktura, kemijska homogenost, ~isto jeklo 1 INTRODUCTION Nowadays, steelmaking technology enables the pro- duction of high-purity steel melts. However, during ingot casting the reoxidation of the melt occurs, thus increas- ing the inclusion content. Segregations on the macro and micro scales are also characteristic for ingot casting. These cause anisotropy in the mechanical properties of the steel. The ESR process almost completely removes the macro-segregation phenomenon in heavy steel ingots, thus ensuring a more homogeneous chemical composition and a finer microstructure with fewer and more evenly distributed non-metallic inclusions than in cast ingots.1 The influence of ESR on remelted steel is shown in Figure 1.2 This is why the ESR process is essential for heavy steel ingots that are used for the manufacturing of large generator and turbine shafts.3 The ESR process is suitable for high-quality ma- terials such as: • steel ball bearings, steel rollers, tool steels, wear- resistant steels for low and high working temperatu- res, high-speed steels for high performance, • highly alloyed stainless steels, corrosion- and acid- resistant steels and steels for high-temperature applications, • steels for aviation and aerospace technology, for medical, pharmaceutical and chemical industries, • Ni superalloys, Ti and Zr alloys for aerospace, medi- cal and chemical industry, components, • off-shore, power and aerospace engineering, reactor comp. ESR is a continuous process, where during the remelting of the consumable electrode, refining and solidification of the steel occur simultaneously. Cast, rolled or forged ingots can be used as a consumable electrode. The ESR process is based on an electrical current running via an electrode through the molten slag and ingot. Due to the high electrical resistance of the slag, the slag heats up and melts. The consumable electrode is immersed in the liquid slag where the slag heat gradually melts the tip of the electrode. Liquefied steel is dripping Materiali in tehnologije / Materials and technology 50 (2016) 6, 971–979 971 UDK 669.187.56-046:621.745:66.02 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)971(2016) Figure 1: Effect of ESR on the properties of remelted steels2 Slika 1: Vpliv ESR na lastnosti pretaljenih jekel2 from the electrode tip and is refined when passing through the liquid slag, with oxides and sulphur being bound in the slag. After passing through the slag, the steel cools down and solidifies again into a remelted ingot.3,4 The whole remelting process takes place in a water-cooled copper mould, which allows the remelted ingot to solidify quickly and very uniformly. The mould with the slag pool is moving upwards as the new ingot is formed. The design of the mould can be in the form of fixed long moulds or collar-type moulds. The use of collar-type moulds with movable moulds or a movable base plate, gives the possibility of producing ingots of any required length (Figure 2).5,6 Furthermore, the ESR enables the production of ingots with the desired shape, i.e., round, square, rectangular. In Slovenia, among other steelmaking processes, ESR technology is used in the Metal Ravne steelworks. 2 ATMOSPHERE CONTROL Due to the ever-increasing demands for material properties, different variations of the ESR process were developed to ensure these demands.6–8 IESR (electroslag remelting under a protective atmosphere of inert gas at atmospheric pressure) is a variation of the ESR where an inert gas (argon) protects the slag and metal from oxidation and the absorption of nitrogen and hydrogen from the air. The oxidation of the electrode is almost entirely avoided, thus providing better cleanliness of the ingot. However, due to the ab- sence of oxygen in the furnace atmosphere, desulphuri- zation is not optimal. PESR (electroslag remelting under increased pressure) is a variation of the ESR with an increased pressure of nitrogen and the melt solidifies under pressure. In this way a large amount of nitrogen can be introduced into the steel melt. The pressure depends on the alloy composition and the desired nitrogen content in the remelted ingot. VAC-ESR (electroslag remelting under vacuum) is a variation of the ESR that provides vacuum degassing of the melt. Dissolved gases such as hydrogen and nitrogen are removed, and the remelted metal is protected from oxidation. The process is suitable for the remelting of superalloys and titanium alloys. 3 PROCESS PARAMETERS The heat required to run the ESR process is generated by the Joule effect in the slag bath. The electrical characteristics, heat balance and electrode/ingot diameter, influence the quality of the remelted ingot.7,9 An energy input of between 1000 kwh/t and 1500 kwh/t of steel is usually required for ESR. The slag bath is considered to be a variable resistor, whose resistance is determined by the electrode distance, the effective slag resistivity and by the electrical current path. The usual slag depth is of the order of 100 mm. The shape of the liquid pool is influenced by the heat input in the process.10,11 The greater the distance between the consumable electrode and the remelted ingot, the smoother the heat distribution in the slag. When determining the electrode distance, it is important to take into consideration that a shorter current path means a higher current with concentrated heat generation under the electrode tip and an undesirable deepening of the metal pool. On the other hand, a longer current path demands a high voltage, which causes more even heat generation and a flatter, more favourable pool profile. ESR operating voltages are usually in the region of 40 V or less. A choice of AC versus DC circuit ESR9 is possible. For ingots of 20 cm in diameter or more, single phase AC gives optimum refinement and melt rate. The DC-ESR requires a lower melt rate for metal refinement. However, when the metal refinement is not the main criterion, the DC-ESR provides the highest melting rates per unit of power consumption. The present practice is to utilize a single-phase AC power supply and low electrode/ingot diameter ratio (0.4 to 0.7).11 Typically, a frequency of 50 Hz or 60 Hz is used in AC operation. However, for the largest ingots, where reactivity be- comes more important, it is better to use low-frequency power (5-10 Hz) for improved efficiency. Optimum melting rates and energy inputs depend on the ingot diameter. A. S. Ballantyne and A. Mitchell12 considered the optimum conditions for the maximum permissible melt rate at the lowest possible power with the Equation (1): Melt rate = constant × power × fill ratio (area) × × mould area / electrode distance (1) Many operators considered melt rate as proportional to the ingot diameter, which is obtained at a melt rate of the order of 0.004 kg/min/mm.7 The relationships between the melt rate, current and voltage for a 240-mm diameter ingot are shown in Figure 3. For a given B. ARH et al.: ELECTROSLAG REMELTING: A PROCESS OVERVIEW 972 Materiali in tehnologije / Materials and technology 50 (2016) 6, 971–979 Figure 2: Schematic representation of the ESR unit: a) retracting ingot, b) rising mould5,6 Slika 2: Shematski prikaz ESR naprave: a) navzdol vle~en ingot, b) dvigajo~a kokila5,6 current and ingot size there is an optional voltage that corresponds to a maximum melt rate.7 The ESR process can be controlled by a computer: from melt initiation, through power build-up, steady melt rate period, reduced melt rate period to maintain pool profile, hot-tapping sequences and melting termination.13 4 SLAG The slag plays an important role in the ESR process; it generates Joule heat for the melting of the electrode, refines the liquid metal through the absorption of non-metallic inclusions, desulphurization, protects the metal from contamination, provides lubrication for the copper mould/solidifying steel shell interface, and controls the horizontal heat transfer between solidifying metal and mould. Slags for ESR are usually based on calcium fluoride (CaF2), lime (CaO) and alumina (Al2O3). Silica (SiO2), magnesia (MgO) and titania (TiO2) may be present, depending on the alloy to be remelted and refined. The CaF2 content increases the solubility of basic components in the slag (CaO and MgO) and thus increases the effective sulphide capacity of the slag.7 In order to per- form its intended functions, the slag must have some well-defined properties, which are: • its melting point must be lower than that of the metal to be remelted, • it must be electrically efficient, • its composition must ensure the desired chemical reactions, • it must have suitable viscosity at the remelting temperature. As presented in Table 1 the concentrations of cal- cium fluoride may vary from 0 % to 100 % of mass frac- tions.7 The remaining slag constituents are mostly used for decreasing the basicity. The slag chemical composition is changed during the ESR process, due to the formation of volatile fluoride, the precipitation of high-melting-point phases and the reaction in the ESR process.14 The changes in composi- tion affect the slag’s metallurgical properties and even- tually affect the quality of the final product. The quantity of consumed slag steel depends on the remelted ingot diameter.6 Table 1: Composition of some ESR slags, in mass fractions (w/%)7 Tabela 1: Sestava nekaterih ESR `linder, v masnih odstotkih (w/%)7 CaF2 CaO MgO Al2O3 SiO2 Comments 100 Electrically inefficient, use where oxides are not permissible 70 30 Difficult starting, high conductivity, use where Al not allowed, risk of H2 pick-up 70 20 10 Good all-round slags, medium resistivity70 15 0 15 50 20 30 Good all-round slag,higher resistivity 70 30 Some risk of Al pick-up, good for avoidance of H2 pick-up, higher resisti- vity 40 30 30 Good general-purpose slags60 20 20 80 10 10 Moderate resistivity,relatively inert 60 10 10 10 10 Low melting point,"long" slag 50 50 Difficult starting,efficient electrically Many of the slags used in ESR can be described with the ternary fluorspar-lime-alumina system.7 The phase diagram shown in Figure 4 has been extensively investi- gated and defined by K. Mills.15 The main feature is an eutectic corresponding to compositions with roughly B. ARH et al.: ELECTROSLAG REMELTING: A PROCESS OVERVIEW Materiali in tehnologije / Materials and technology 50 (2016) 6, 971–979 973 Figure 3: Effect of current and voltage on melt rates7 Slika 3: Vpliv toka in napetosti na hitrost taljenja7 Figure 4: Phase diagram of the CaF2-Al2O3-CaO system according to K. Mills15 Slika 4: Fazni diagram sistema CaF2-Al2O3-CaO po K. Millsu15 equal proportions of lime and alumina. This identifies the slags with liquidus temperatures in the range 1350–1500 °C, which make them suitable for melting of a wide range of alloys, including steels and super alloys. In the case of slag with 70 % fluoride and 30 % alumina the lime is excluded as much as possible in order to pre- vent hydrogen pick-up, while there are no problems with the presence of the two liquids. The binary lime-alumina system on the other hand, has only a limited range of slags with suitable melting characteristics, while the binary calcium fluoride-lime system is used in cases where a high degree of desulphurization is required. However, its disadvantage is having a low resistivity. High lime contents also increase the risk of moisture retention or hydrogen pick-up. According to the tendency to reduce CaF2 in slags, the investigations of the slag S 2015 (Wacker Cheime) with 30 % CaF2 were performed. The results also showed that tested slags with 4.7 % CaF2 can be satis- factorily applied in the ESR-process of UTOP Mo6 steel.16 A certain amount of SiO2 addition into the ESR slag in the case of the drawing-ingot-type ESR process is important for improving the lubrication performance, controlling silicon and aluminium content in the liquid steel and modifying oxide-type inclusions.17 Further- more, the addition of SiO2 suppresses the crystallization temperature of CaF2-Al2O3-CaO slags. Furthermore, the MgO and SiO2 in fluoride-containing slags affect the slag’s surface tension.18 Although CaF2 is a crucial component in any ESR slag and it greatly decreases the melting temperature of the slag systems, it is insoluble in oxide phases. An example of an ESR slag microstructure is shown in Figure 5, where the lighter dendrite-like phase is a stable CaO-Al2O3 phase, the darker phase is fluorspar, and the small white dots the undissolved magnesia particles. The fluorspar phase contains only calcium and fluoride and is not dissolved in other microstructural constituents. Slag properties, such as electrical conductivity, ther- mal conductivity, density, viscosity and surface tension play an important role in effective melting and metal refining. K. Mills19 has produced a table of the physical properties of slags for practical purposes (Table 2). Slag resistivity affects the operating characteristics and eco- nomics of ESR. Alumina increases the resistivity of the slag and promotes good heat generation, thus enabling a reduction of the slag bulk content, which also reduces the heat loss due to the reduced area of contact between the slag and the mould wall. L. A. Kamenski et al.20 refer to "long" and "short" slags when discussing slag viscosity. Long slags remain fluid over a wide range of temperatures and are likely to give thin slag skins and therefore good ingot surfaces. Short slags rapidly become viscous on cooling and are likely to give thick slag skins and poor ingot surfaces. High calcium fluoride contents promote short slags, whereas silica and magnesia favour long slags. The slag plays an important role in ESR, from the control-of-inclusions point of view.21 The chemical and physical properties of slag also have a great effect on the removal of inclusions. 5 THERMODYNAMICS In the case of ESR of steel in an air atmosphere, chemical reactions take place and change the chemical composition of the as-cast ingot.22 The levels of some B. ARH et al.: ELECTROSLAG REMELTING: A PROCESS OVERVIEW 974 Materiali in tehnologije / Materials and technology 50 (2016) 6, 971–979 Figure 5: Microstructure of ESR slag Slika 5: Mikrostruktura ESR `lindre Table 2: Phase diagram of the CaF2-Al2O3-CaO system and the physical properties of slag at 1600 °C, according to K. Mills19 Tabela 2: Fazni diagram sistema CaF2-Al2O3-CaO in fizikalne lastnosti `lindre pri 1600 °C, po K. Millsu19 Electrical conduc- tivity Viscosity Density Surfacetension Total normal emissivity Contour  ( –1 cm–1)  (10–1 Ns m–1)  (g cm–3) s (mN m–1) TN A 6 0.15 2.47 285 0.96B 5 0.2 2.48 300 C 4 0.25 2.49 310 D 3.5 0.3 2.5 320 0.9 E 3 0.4 2.55 335 F 2.5 0.6 2.6 350 0.85 G 2 0.8 2.7 400 H 1 1.0 2.8 450 0.81 elements, such as Co, Ni, Cr, Mo, W, C remain un- changed after remelting. However, the content of Si, O, and S can be changed from 10 % to 80 %, while the content of Al and Ti can vary depending on the melting conditions (decrease or increase). Therefore, some measures need to be taken to prevent the losses of elements. This can be achieved by using special ESR variations as discussed before. Another way is control of the slag composition and regular additions to the slag, which is desirable due to steady melting conditions. The oxidation of the elements can be prevented by deoxida- tion of the slag during the melting process by additions of aluminium. The oxygen potential of the slag determines the chemistry of the ESR process.7 It affects the non-metallic inclusions and sulphur removal. Oxygen reacts with some elements in the metal and suppresses hydrogen pick up. In the slag oxygen is mostly bound as FeO, MnO and SiO2. To estimate the oxygen content in the steel it is necessary to find the relationship between FeO in the slag and oxygen in the remelted ingot.22,23 How- ever, due to the very low solubility of FeO in CaF2 slags, its activity is extremely high. The oxygen content can be estimated by thermodynamic analyses of the reactions between active components and oxygen. Silicon and manganese are elements that can react with the oxygen present in the steel and from the slag.23,24 When silicon is the strongest deoxidizer, the oxygen content of the steel is determined by the Si content.23 At constant tempe- rature and Si content in the steel, the oxygen content of the metal is higher at higher activity of the SiO2 in the slag, or by lowering the basicity of the slag. Aluminium losses in the remelted ingot are small, especially at high alumina content in the slag. On the other hand, the presence of Al2O3 in the slag reduces the oxidation of silicon. The reaction between the silicon in the electrode and Al2O3 in the slag also controls the oxidation of aluminium in the remelted ingot.25 Thus, Al content in the remelted ingot depends on the content of Al2O3 in the slag and the content of silicon in the electrode, temperature and chemical composition of the steel.7,25 The content of Al in the remelted ingot de- creases when CaF2-Al2O3-CaO slags with increased SiO2 content are used. When aluminium is used for deoxidation, up to 15 % of added Al is transferred to the molten steel. The content of titanium in the remelted steel will depend on the content of Al and Ti in the consumable electrode, the content of Al2O3 and TiO2 in the slag and the oxygen potential in the gas phase above the slag (Figure 6).26 The equilibrium between the Al and Ti content in the electrode at different TiO2 contents is presented in Figure 6. For the content of Al in electrode, the titanium loss can be minimized by the addition of TiO2 to the slag. At high contents of Al, aluminium reduces TiO2 in the slag and also regulates the ratio of Ti:TiO2. In the early stages of ESR development, the removal of sulphur was considered as one of the major objectives. The rate of desulphurization increases with the basicity of the slag. Sulphur transfer takes place mainly at two interfaces, according to the following two reactions:23 Slag/metal reaction: [S] + (O2–) = (S2-) + [O] (2) Gas/slag reaction: (S2–) + 3/2 {O2} = {SO2} + (O 2) (3) A thermodynamic analysis of the reactions shows that the desulphurisation is related to the concentration of O2– ions in the slag, the partial pressure of oxygen in the gas phase and the chemical composition of the steel.27 The transfer of sulphur from the metal to the slag is promoted by the high slag basicity and low concen- tration of oxygen in the metal. On the other hand, the sulphur transfer from slag to gas is promoted by a high partial pressure of oxygen in the atmosphere and the low basicity of the slag. The ability of the slag to take sul- phur is defined in terms of its sulphur capacity. The sul- phur capacity for the fluorspar-lime-alumina system increases as the CaF2 content is increased and by in- creasing the amount of lime to the saturation point.28 In the case of ESR under inert gas, the sulphur remains in the slag and builds up there as the process continues. In such cases the sulphur capacity is the ruling factor, and the slag bulk must be adjusted in order to continue its desulphurising action to the end of the pro- cess, i.e., the slag/metal ratio assumes greater import- ance.7 6 SOLIDIFICATION AND STRUCTURE The solidification structure of an ESR ingot is a func- tion of the local solidification time and the temperature gradient at the liquid/solid interface.4,7 To achieve a directed dendrite primary structure, a relatively high temperature gradient at the solidification front must be maintained during the entire remelting period. B. ARH et al.: ELECTROSLAG REMELTING: A PROCESS OVERVIEW Materiali in tehnologije / Materials and technology 50 (2016) 6, 971–979 975 Figure 6: Effect of Al:Ti ratio in the electrode on TiO2 content in the slag26 Slika 6: Vpliv odvisnosti Al:Ti v elektrodi na vsebnost TiO2 v `lindri26 Macrostructure of the ESR ingots is different from the macrostructure of conventionally cast ingots due to the different heat transfer and heat removal. The growth direction of the dendrites is a function of the metal pool during solidification. Thus, the gradient of dendrites with respect to the ingot axis increases with melting rate. In extreme cases the growth of directed dendrites can come to a stop. The ingot core then solidifies non-directionally in equiaxed grains, which leads to segregation and micro shrinkage. Even in the case of directional dendritic soli- dification, the micro segregation increases with the den- drite arm spacing. A solidification structure with dendrites parallel to the ingot axis yields optimal results. However, this is not always possible. A good ingot sur- face requires a minimum energy input and accordingly a minimum melting rate. Increasing the melting rate increases the difference between the gradient of the solidus and liquidus iso- therms and leads to increased pool depth.29 Hence, grains grow in radial direction instead of vertical direction. Figure 7 shows the direction of the grain growth depen- dent on the melting rate, which affects the pool depth. Figure 8 presents the predicted grain structures for different melting rates up to 600 kg/h.30 Increasing the melting rate causes a finer grain structure and changes the growth direction of the columnar structure from the axial to radial growth and deeper liquid pool at very high melting rates. Increasing the molten slag temperature also results in a coarser columnar grain structure and a reduced thickness of the refined equiaxed grain layer, both at the surface and the bottom of the ESR ingot. In spite of directional dendritic solidification, defects such as tree-ring patterns, freckles and white spots can occur in a remolten ingot.4,7,31 Macro-segregation and porosity structures in the middle of the ingot are very un- common for ESR ingots. A major attribute of ESR is its ability to produce ma- terial with reduced micro-segregation. This is linked with the local solidification time and dendrite-arm spacing.7 ESR material normally freezes in a columnar manner, which gives less micro-segregation than equiaxed struc- tures. The greater the temperature gradient, the smaller is the distance between the dendritic arm spacing and the lower is the chemical heterogeneity in the micro areas. In ESR, temperature gradients are greater than for conven- tional casting. Therefore, the secondary dendrite-arm spacing will be smaller in ESR than in conventional B. ARH et al.: ELECTROSLAG REMELTING: A PROCESS OVERVIEW 976 Materiali in tehnologije / Materials and technology 50 (2016) 6, 971–979 Figure 7: The ingot grain growth at a melting rate of: a) 0.9 kg/min and b) 1.2 kg/min32 Slika 7: Rast zrna v ingotu pri hitrostih taljenja: a) 0,9 kg/min in b) 1,2 kg/min34 Figure 9: Segregations in the hot-work tool steel: a) consumable elec- trode, b) remelted ingot Slika 9: Izceje v orodnem jeklu za delo v vro~em: a) elektroda pred pretaljevanjem, b) pretaljen ingot Figure 8: The predicted grain structure of ESR ingot30 for different melting rates; a) 136 kg/h, b) 271 kg/h, c) 407 kg/h and d) 542 kg/h Slika 8: Napovedana zrnatost ESR-ingota30 pri razli~nih hitrostih taljenja: a) 136 kg/h, b) 271 kg/h, c) 407 kg/h in d) 542 kg/h ingots. The effect of decreasing the segregation effect is shown in Figure 9, where a comparison of microstruc- tures before (Figure 9a) and after (Figure 9b) ESR pro- cessing was made for a hot-work tool steel. The micro- structure in both cases is tempered martensite. The difference in segregation bands is obvious. While the segregations are evident in the consumable electrode (Figure 9a) they are almost completely eliminated in the remelted ingot (Figure 9b). Figure 10 shows the effect of local solidification time on the dendrite spacing.32 The dendrite-arm spacing is decreased as the cooling rate is increased. Besides a more homogeneous composition and com- pact solidification structure, the removal of non-metallic inclusions is an important characteristic of ESR.33 Normally, inclusions easily initiate micro-voids and cracks at the inclusion/steel interface, which can be the origin of fatigue fracture or other defects. Also, ESR steel is not an exception.34 Many factors influence the formation of non-metallic inclusions in ESR steel, including furnace atmosphere, content of inclusions in the consumable electrode, slag amount and its compo- sition, power input, melting rate, filling ratio, etc. Most non-metallic inclusions occur due to the reactions bet- ween oxygen and elements such as manganese, silicon and aluminium.35 Deoxidization of the slag during electroslag remelting36 has an important influence on the non-metallic inclusions formation in the ESR ingot. The results of the experimental work show that the lowest number of inclusions is attained in ESR with the lowest viscosity and the highest interfacial tension.37 However, the absence of large inclusions is typical for ESR, as shown in Figure 11.38 The removal of non-metallic inclusions during ESR takes place at the tip of the electrode, where mainly absorption and dissolution of non-metallic inclusions in the slag take place.6,39 As the electrode tip is heated towards its melting point, the inclusions in the electrode are re-dissolved before the metal melts. Any other in- clusions, such as larger exogenous inclusions in the electrode, are not dissolved in the solid metal and will be exposed to the slag when the electrode tip becomes molten. If the slag composition is suitable, the tempe- rature is high enough and the dwell time is long enough the non-metallic inclusions will dissolve in the slag.40 However, at this point there may be further reactions due to the difference in equilibrium constants, as well as the possibility of the flotation of large inclusions. The metal at this point is free from non-metallic inclusions, but may have in solution elements that produce inclusions by reaction during the freezing time (sulphur removal reac- tion).41,42 The removal efficiency of inclusions increases with the reduced melting speed.43 The research work of Y.-W. Dong et al.44 was focused on the impact of fluoride-containing slag and interac- tions at the slag-metal interface on the non-metallic inclusions in steel. Results indicate that a multi-compo- nent slag (CaF2, CaO, Al2O3, SiO2, MgO) has a better capacity for controlling the amount of inclusions. Most non-metallic inclusions for multi-component slag are MgO-Al2O3 inclusions, while mainly Al2O3 inclusions exist when using conventional 70 % CaF2 – 30 % Al2O3 slag. Furthermore, the maximum inclusion size for multi-component slags was found to be smaller than for conventional binary slag. 7 CONCLUSION The main purpose of the remelting process is to control the non-metallic inclusions in the steel, remove segregations and shrinkage, and produce more homoge- nous ingots. The low speed of remelting, combined with the water-cooled mould, ensures a particularly homoge- B. ARH et al.: ELECTROSLAG REMELTING: A PROCESS OVERVIEW Materiali in tehnologije / Materials and technology 50 (2016) 6, 971–979 977 Figure 11: Inclusion size and frequency for inclusions >2.5 μm in air-melted and ESR with 3 % Cr-Mo-V steels38 Slika 11: Velikost in {tevilo vklju~kov > 2,5 μm po ESR-pretaljevanju in pretaljevanju na zraku jekla 3 % Cr-Mo-V38 Figure 10: Microstructures of ingots of ESR nickel alloy 718, with dendrite-arm spacing at different cooling rates: a) slower cooling, b) faster cooling and c) fastest cooling32 Slika 10: Mikrostruktura EP@ ingota nikljeve zlitine 718 in razdalj med dendritnimi vejami: a) po~asnej{e ohlajanje, b) hitrej{e ohlajanje in c) najhitrej{e ohlajanje 32 neous and balanced, stable solidification. The segrega- tions within a remolten ingot are thus much lower (or even eliminated) compared to open cast continuous cast billets or conventional cast ingots. For this reason most segregation-sensitive steels are ESR processed for homogenisation. The slag plays an important role in the ESR process, as it absorbs non-metallic inclusions, removes sulphur, influences the ingot surface and the melting rate as well as the overall economics of the process. That is why the chemical composition and physical properties of the chosen ESR slag is of paramount importance for a high-quality ESR ingot. In order to further improve non-metallic inclusion reduction, specialised versions of the ESR process like IESR, PESR and VAC-ESR have been developed. The application of the ESR process is appropriate for tool and high-speed steels as well as special stainless steels and special alloys intended for the most demand- ing applications. 8 REFERENCES 1 R. C. Reed, The Superalloys: Fundamentals and Applications, Cambridge University Press, 2006 2 J. Rodi~, Razvoj elektri~nega pretaljevanja jekel pod `lindro v @elezarni Ravne, @EZB, 18 (1984) 4, 105 3 R. H. 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Inter- national Conference on Metallurgy and Materials, may 2014, Brno, Czeh Republic B. ARH et al.: ELECTROSLAG REMELTING: A PROCESS OVERVIEW 978 Materiali in tehnologije / Materials and technology 50 (2016) 6, 971–979 38 E. M. Lowe, A. Hogg, Aplication of ESR to alloy steel forgings, Proccedings of conference on Electroslag refining, Iron and Steel Institute, Sheffield, 1973, 68 39 B.-H. Yoon, K.-H. Heo, J.-S. Kim, H.-S. Sohn, Improvement of steel cleanliness by controlling slag composition, Ironnmaking Steelmak., 29 (2002) 214–217, doi:10.1179/030192302225004160 40 A. Mitchell, Oxide inclusion behavior during consumable electrode remelting, Ironmaking and Steelmaking, 1 (1974) 3, 172 41 M. E. Fraser, A. Mitchell, Mass transfer in the electroslag process, Part 1, Mass transfer model, Ironmaking and Steelmaking, 3 (1976) 5, 279 42 M. E. Fraser, A. Mitchell, Mass transfer in the electroslag process, Part 2, Mass transfer coefficients, Ironmaking and Steelmaking, 3 (1976) 5, 288 43 C. Chen, J. Wang, D. Shu, B. Sun, Removal of Iron Impurity from Aluminum by Electroslag Refining, Mater. Trans., 52 (2011) 1320–1323, doi:10.2320/matertrans.M2010435 44 Y. - W. Dong, Z. – H. Jiang, Y. - L. Cao, A. Yu, D. Hou, Effect of slag on inclusions during electroslag remelting process of die steel, Metallurgical and Materials Transactions B, 45B (2014) 8, 1315–1324, http://dx.doi.org/10.1007/s11663-014-0070-7 B. ARH et al.: ELECTROSLAG REMELTING: A PROCESS OVERVIEW Materiali in tehnologije / Materials and technology 50 (2016) 6, 971–979 979 A. STAMBOLI] et al.: CONTINUOUS VERTICAL CASTING OF A NiTi ALLOY 981–988 CONTINUOUS VERTICAL CASTING OF A NiTi ALLOY VERTIKALNO KONTINUIRNO LITJE NiTi ZLITINE Ale{ Stamboli}1,2, Ivan An`el3, Gorazd Lojen3, Aleksandra Kocijan1, Monika Jenko1,2, Rebeka Rudolf3,4 1Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia 2Jo`ef Stefan International Postgraduate School, Jamova 39, 1000 Ljubljana, Slovenia 3University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, 2000 Maribor, Slovenia 4Zlatarna Celje d.d., Kersnikova 19, 3000 Celje, Slovenia ales.stambolic@imt.si Prejem rokopisa – received: 2016-06-17; sprejem za objavo – accepted for publication: 2016-06-27 doi:10.17222/mit.2016.111 In this paper we present research that is connected to the performance of a series of experiments combined with the vacuum-induction melting and continuous vertical casting of a NiTi alloy in order to produce the strand. The theoretical chosen parameters made it possible to obtain a continuously cast strand with a diameter of 11 mm. The strand microstructures were investigated with a light and scanning electron microscope, while the chemical composition of the single phase was identified with the semi-quantitative micro-analysis energy-dispersive X-ray spectroscopy and inductively coupled plasma – optical emission spectrometry. The research showed that the microstructure is dendritic, where in the inter-dendritic region the eutectic is composed of a dark NiTi phase and a bright TiNi3–x phase. In some areas we found Ti carbides and phases rich in Fe. The micro-chemical analysis of the NiTi strand showed that the composition changed over the cross and longitudinal sections, which is proof that the as-cast alloys are inhomogeneous. In the final part, the electrochemical behaviours of NiTi strand samples were compared to a commercially available NiTi cast alloy with the same composition. Keywords: NiTi alloy, continuous vertical casting, microstructure, potentiodynamic and impedance test V tem prispevku predstavljamo raziskavo, ki je povezana z izvedbo niza preizkusov vakuumskega pretaljevanja in so~asnega kontinuirnega vertikalnega litja NiTi zlitine s ciljem odliti palico. Teoreti~no izbrani parametri so omogo~ili, da smo uspeli kontinuirno odliti NiTi palico s premerom 11 mm. Dobljeno mikrostrukturo palice smo raziskali s svetlobnim in vrsti~nim elektronskim mikroskopom, kemijsko sestavo posameznih faz pa smo identificirali s semi-kvantitativno mikro-kemi~no analizo Energijsko disperzijsko spektrometrijo in z opti~nim emisijskim spektrometrom z induktivno sklopljeno plazmo. Preiskave so pokazale, da je mikrostruktura dendritska, medtem ko s v meddendritskem prostoru nahaja evtektik, sestavljen iz temne NiTi faze in svetle TiNi3–x faze. Mestoma smo identificirali tudi Ti karbide in fazo bogato s Fe. Mikro-kemi~na analiza NiTi palice je odkrila, da se sestava spreminja po prerezu in po dol`ini, kar nakazuje, da je zlitina po strjevanju nehomogena. V zaklju~nem delu smo primerjali elektrokemijsko obna{anje vzorcev NiTi palice s komercialno dostopno valjano NiTi zlitino enake sestave. Klju~ne besede: NiTi zlitina, vertikalno kontinuirno litje, mikrostruktura, potenciodinami~ni in impedan~ni test 1 INTRODUCTION NiTi alloys are an attractive group that also include nitinol. Nitinol is a group of nearly equiatomic alloys of nickel and titanium which is located in the central region of the NiTi phase diagram and bounded by the Ti2Ni and TiNi3 phases.1 It exhibits a unique combination of good functional properties and a high mechanical strength, such as super-elasticity and a shape-memory effect, good corrosion resistance, an unusual combination of strength and ductility and excellent biomechanical compati- bility.2,3 This alloy was developed in the 1970s and its properties have enabled its use especially for biomedical purposes, first in orthodontic treatments, and later on in cardiovascular surgery for stents, guide wires, filters, etc., in orthopaedic surgery for various staples and rods, and in maxillofacial and reconstructive surgery.4 In addition to bio-engineering, nitinol has been used in aerospace, automotive, civil and structural engineering.5 Super-elastic NiTi is capable of recovering large inelastic strains spontaneously upon unloading. On the other hand, shape memory is exhibited when NiTi recovers large strain deformation upon heating. Both the super- elasticity and shape-memory effect are induced in nitinol by reversible, displacive, diffusionless, solid–solid phase transformations from a high-temperature parent phase (austenite) with a highly ordered crystal structure to a low temperature, stress-free martensite that has a less or- dered structure. Nitinol is hysteretic, and there are seve- ral transformation temperatures, including the austenite start temperature (As), the austenite finish temperature (Af) during heating and the martensite start temperature (Ms) and the martensite finish temperature (Mf) during cooling. Super-elastic behaviour will only occur if the material is loaded above its Af temperature.6–8 The common production route for a NiTi alloy with a shape-memory effect is known and has been experi- mented on laboratory equipment with the technological aspects of vacuum induction melting, hot and cold work- ing operations. The process is still being optimized with a particular focus on obtaining a small dimension in the cross-section and with stabilisation of its functional properties over its lifetime.9 Vacuum induction melting Materiali in tehnologije / Materials and technology 50 (2016) 6, 981–988 981 UDK 621.74.047:669.24:669.295 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)981(2016) (VIM) is often used as the first technique in the prepa- ration of a melt. Basically, it is a typical melting tech- nique for the production of different NiTi-based alloys. This is appreciated particularly for NiTi alloy due to the strong influence of the chemical composition on the reactivity with oxygen and other elements, leading to oxidation of the NiTi melt. In the second step, such a prepared melt is cast, which enables pouring the molten metal into a mould of the desired shape, and allowing it to solidify. When the molten metal is poured into the mould, chill crystals nucleate on the cold walls of the mould and grow inwards. Conventional casting is a batch process that produces large ingots requiring significant subsequent processing. Large mechanical equipment with high construction and operational costs is necessary to break down most ingots. These problems can be solved by using continuous vertical casting (CVC).10–13 With CVC the raw material is placed into a VIM furnace, in which the material melts. After melting, the melt is, based on gravity force, moved against the nozzle, which adjusts the rate and direction of the melt flow. The melt flows through the nozzle into a water-chilled mould, where the melt is solidified, and obtains the final strand shape. Nitinol is often subjected to deformations or stresses that result in some kinds of mechanical failures. Two very important factors must be considered when using various materials in medicine, i.e., the toxicity of the material and the failure of material. The main problem of NiTi alloys is the high Ni content. Ni releasing can in- duce toxic, allergic and hypersensitive reactions or tissue necrosis after long-term implantation. To prevent failure and Ni release, a coating of appropriate thickness must be formed on the NiTi surface. Titanium oxide coatings effectively suppress the nickel ions outleaching. The niti- nol surface is spontaneously covered by Ti dioxide because of the gain in free energy of formation for this oxide compared to the Ni oxides. However, the oxides formed on the nitinol surface always contain a certain fraction of Ni.14–19 The main goal of this work was the performance of the series of experiments combined with vacuum-induc- tion melting and continuous vertical casting of NiTi alloy in order to produce the strand. This was followed by the characterization of the obtained microstructure and finally we compared the electrochemical behaviour bet- ween a NiTi strand and commercially available nitinol. 2 EXPERIMENTAL PART 2.1 Continuous vertical casting of NiTi alloy The NiTi alloy composed of 50 % of amount frac- tions of Ni and 50 % of amount fractions of Ni was prepared with the combination of techniques: VIM and CVC. A clay-graphite crucible was filled up to 2/3 of its volume due to the high metallostatic pressure (pressure that occurs within a molten metal) with Ti pellets (99.99 % purity) and Ni tablets (99.99 % purity). By remelting the NiTi alloy with VIM at a temperature of about 1450 °C a pressure lower than 10–2 mbar was achieved in the system. The induction power during heating was for first 10 min 10 kW, then next 10 min 20 kW and in final 5 min 30 kW, while during casting it was between 25 and 30 kW. Continuous casting was operating in the mid range frequency (4 kHz). In the experiments a Cu-mould (Figure 1), a ZrO2 nozzle stabilized with Y2O3 and an Fe starter bar were applied. 2.2 Preparing of the samples for further investigation The samples for characterization were cut longitu- dinally (according to the direction of casting) and across the cross-section. For this purpose, an Accutom 50 elec- tronic saw was used for precision cutting. The grinding was performed with 320 grit SiC abrasive paper, mechanical polishing with MD-Largo discs with 9-μm diamond suspension and with peroxide grains in a chemically aggressive suspension – OP-S (colloidal silica). The sample was then etched with Kroll’s reagent (3 mL HF, 6 mL HNO3 and 100 mL of distilled water). 2.3 Analytical techniques The microstructure was investigated with a light microscope – Microphot FXA, Nikon 3CCD-Hitachi Camcorder HV-C20A and Thermal Field Emission SEM JEOL JSM-6500F equipped with energy-dispersive X-ray spectroscopy (EDS) analytical technique. Chemi- cal analyses were performed by inductively coupled plasma – optical emission spectrometry ICP-OES (Agi- lent 720). Potentiodynamic polarisation measurements and electrochemical impedance spectrometry (EIS) have been used to study the electrochemical behaviour of A. STAMBOLI] et al.: CONTINUOUS VERTICAL CASTING OF A NiTi ALLOY 982 Materiali in tehnologije / Materials and technology 50 (2016) 6, 981–988 Figure 1: Schematic presentation of copper mould with cooling sys- tem at the Faculty of Mechanical Engineering, Maribor, Slovenia Slika 1: Shema bakrene kokile s hladilnim sistemom na Strojni fakul- teti v Mariboru, Slovenija samples. All the measurements were recorded by BioLogic Modular Research Grade Potentiostat/Galva- nostat/FRA Model SP-300 with an EC-Lab Software and a three-electrode cell. In this cell, the sample was the working electrode, saturated calomel electrode (SCE, 0,242 V vs. SHE) was used as reference electrode and the counter electrode (CE) was a platinum net. The expe- riment was held in simulated physiological Hank’s solution, containing 8 g/L NaCl, 0.40 g/L KCl, 0.35 g/L NaHCO3, 0.25 g/L NaH2PO4×2H2O, 0.06 g/L Na2HPO4×2H2O, 0.19 g/L CaCl2×2H2O, 0.41 g/L MgCl2×6H2O, 0.06 g/L MgSO4×7H2O and 1 g/L gluco- se, at pH = 7.8 and 37 °C. All the chemicals were from Merck, Darmstadt, Germany. The potentiodynamic curves were recorded after 1 h of sample stabilisation at the open-circuit potential (OCP), starting the measure- ment at 250 mV vs. SCE more negative than the OCP. The potential was then increased, using a scan rate of 1 mV s–1, until the transpassive region was reached. Long-term open circuit potentiostatic electrochemical impedance spectra were obtained for the investigated samples. The impedance was measured at the OCP, with sinus amplitude of 5 mV peak to peak and a frequency range of 65 kHz to 1 mHz, in the sequence of directly after immersion after 1h, 2 h, 6 h, 12 h, 24 h, 48 h, 72 h, 96 h, 120 h, 144 h, 168 h and 192 h. The impedance data are presented in terms of Nyquist plots. For the fitting process Zview v3.4d Scribner Associates software was used. 3 RESULTS AND DISCUSSION 3.1 Continuous vertical casting of a NiTi alloy The CVC of a NiTi alloy is a complex process that requires precise process parameters. Accurate measure- ment and regulation of temperature was very difficult because the thermocouple was not in constant contact with the melt due to the potential contamination of the melt and the temperature at the crucible wall is quite different from the actual temperature of the melt. The frequency of induction is also very important for the casting, as a high frequency enables the temperature to rise and low frequency means more intensive stirring. In this case the casting was operated at a mid-range fre- quency of induction that does not provide adequate mixing power, causing an undesirable chemical compo- sition in some places of the strand. The drawing of the strand was carried out in the sequence of pull – pause, as this reduces the possibility of a reaction between the alloy and the mould, as well as the porosity of the material or the occurrence of cracks in the material. The drawing stroke had a length of between 0 and 10 mm and the pause lasted between 0 and 1 s. The drawing rate is also an important factor. When the drawing is too slow, the temperature decreases, which leads to solidification of the alloy in the nozzle and retraction of further drawing. This leads to fracture of the strand and the process ends without the desired result. The strand also breaks when the drawing rate is too fast due to the adhesion to the mould and the weakness of the thin solidified skin. With CVC a strand with diameter of 11 mm was ob- tained (Figure 2). ICP analysis for the first attempt of CVC NiTi strand showed a constant material compo- sition of 59.8 % of amount fractions of Ni, 38.9 % of amount fractions of Ti and 0.3 % of amount fractions of C, EDS analysis showed approximately 1 % of amount fractions of Fe. Deviation from the desired value (50 % of amount fractions of Ni) is probably caused by complications with stirring of the melt (better mixing takes place at a lower frequency induction, Ti is very difficult to mix). The source of Fe could be attributed to the Fe screw that was used as a starter bar. During fur- ther attempts the chemical composition of the strand varied during casting. At the beginning of drawing XRF analysis showed that the strand was rich in nickel (70.6 % of amount fractions of Ni; 27.1 % of amount fractions of Ti) and with the increasing length of the strand the nickel content decreased. Chemical compo- sition during the fracture of the strand was 52 % of amount fractions of Ni and 47 % of amount fractions of Ti. 3.2 Microstructure 3.2.1 Continuously vertical cast NiTi alloy The light microscopy of the strand cross-section reveals the dendritic microstructure (Figure 3), where inside the primary phase NiTi is located. This is accord- ing to the Ni-Ti phase diagram where the first solidified phase is NiTi. Dendrites grew in the direction from the coldest location (from the walls of the nozzle) to the middle of the strand. The orientation of dendrites is random. These dendrites are arranged in the matrix of eutectic (composed with NiTi eut + TiNi3–x). In the microstructure there are no visible defects such as cracks and porosity. A. STAMBOLI] et al.: CONTINUOUS VERTICAL CASTING OF A NiTi ALLOY Materiali in tehnologije / Materials and technology 50 (2016) 6, 981–988 983 Figure 2: a) NiTi strand, produced at Faculty of Mechanical Engineer- ing, Maribor, Slovenia and b) light microscope image of cross-section of the strand Slika 2: a) NiTi palica, lita na Strojni fakulteti v Mariboru, Slovenija in b) posnetek pre~nega prereza palice, narejen s svetlobnim mikro- skopom The NiTi strand contains between 50 % and 60 % of amount fractions of Ni. From the phase diagram (Figure 4) it is clear that this is a hypo-eutectic alloy (according to the eutectic reaction at 1118 °C: L  NiTi + TiNi3). With an ideal cooling the melt would begin to solidify in the temperature range between 1310 °C and 1118 °C. From the melt firstly the primary NiTi phase solidifies that would be continuously generated and grew until the eutectic temperature (1118 °C) is reached. At this tem- perature, the remaining melt solidifies into a eutectic structure composed of a NiTi phase and TiNi3 phase in the form of lamellas. In the real case, the cooling is non-equilibrium. Solidifying rates are large, but the diffusion rates in the solid state are too small to make it possible to achieve a homogeneous solid phase. A backscattered electrons image (Figure 5) shows a typical dendritic structure (tree-like form) that are solidified primarily (NiTi phase). At the eutectic temperature (1118 °C) solidifies typical lamellar eutectic structure (NiTi + TiNi3–x) from the residue of the melt. EDS analysis at 5 keV showed that both the dendritic phase and the dark lamellas of eutectic, have a composition of approximately 50 % of amount fractions of Ni and 50 % of amount fractions of A. STAMBOLI] et al.: CONTINUOUS VERTICAL CASTING OF A NiTi ALLOY 984 Materiali in tehnologije / Materials and technology 50 (2016) 6, 981–988 Figure 5: Backscattered-electron image of NiTi strand at 1000× mag- nification Slika 5: Posnetek povratno-sipanih elektronov NiTi palice pri 1000× pove~avi Figure 3: Light microscope image of NiTi strand at 100× magnifica- tion Slika 3: Posnetek NiTi palice na svetlobnem mikroskopu pri 100× pove~avi Figure 6: a) SE image of NiTi strand at 5000× magnification of area where the TiC inclusions are present, b), c), d) and e) elemental mapp- ing at the microstructural level by scanning electron microscopy (SEM) with energy dispersive X-ray spectrometry (EDS) in the area with TiC inclusions Slika 6: a) SE-posnetek NiTi palice pri 5000× pove~avi v obmo~ju, kjer so prisotni TiC vklju~ki, b), c), d) in e) elementna analiza na mikrostrukturni ravni z vrsti~nim elektronskim mikroskopom (SEM) z energijo disperzijsko rentgensko spektrometrijo (EDS) v obmo~ju s TiC vklju~ki Figure 4: Ni-Ti phase diagram Slika 4: Fazni diagram Ni-Ti Ti, while bright lamellas of eutectic have a composition of 33 % of amount fractions of Ti and 67 % of amount fractions of Ni. The secondary electron (SE) image (Figure 6) reveals in addition to the dendritic structure also the presence of the individual inclusions. The EDS analysis showed that the inclusions are titanium carbide (TiC). Carbon originates from the clay-graphite crucible and diffuses into the melt during the melting and reacts there with the Ti. The Gibbs free energy for the formation of TiC is very low, so the conditions for the formation of TiC are very favourable. From the results of the EDS analysis it appears that the carbon is located only in the form of carbides, and there is none in the other phases. Ni and Fe are located in the dendrites and the matrix, but not in the carbides, while titanium is present in all the phases. Another important fact is that, during CVC, there was no contamination with oxygen because no dissolved oxygen or oxides were observed in the strand. In this manner it could be concluded that the vacuum was appropriate. So far several VIM + CVC experiments for the pro- duction of NiTi strand were made. In the first attempt the chemical composition of the strand was constant, but incorrect. During further attempts it varied during draw- ing in the direction of reducing the nickel content. It was concluded that the mixing of the melt was inappropriate. Insufficient stirring was attributed to the 4-kHz inductor. To achieve better stirring, a low-frequency generator should be modulated. Costs for something like that are too high and therefore the remelting method will be fur- ther used. CVC will be held with an in advance prepared NiTi alloy. Instead of Fe starter bar, that probably intro- duced Fe impurities in the alloy, a starter bar with a Ti-tip will be applied. The vacuum by VIM was appro- priate, because no oxygen or oxides were found in the strand, but the crucible will also need to be modified due to some concentration of TiC phase in the strand. 3.2.2 Commercially available NiTi alloy The light microscope image (Figure 7a) reveals rela- tively large grains (> 20 μm); the grain boundaries are clearly noticeable and the grains have different shapes and sizes. The secondary-electron image made with SEM (Fig- ure 7b) reveals that the commercially available NiTi A. STAMBOLI] et al.: CONTINUOUS VERTICAL CASTING OF A NiTi ALLOY Materiali in tehnologije / Materials and technology 50 (2016) 6, 981–988 985 Figure 8: a), b) and c) Elemental mapping at the microstructural level by scanning electron microscopy (SEM) with energy-dispersive X-ray spectrometry (EDS) of commercial NiTi alloy Slika 8: a), b) in c) Elementna analiza komercialne NiTi zlitine na mikrostrukturni ravni z vrsti~nim elektronskim mikroskopom (SEM) z energijsko disperzijsko rentgensko spektrometrijo (EDS) Figure 7: a) Light microscope image of commercially available NiTi alloy at 100× magnification and b) SE image of commercially available NiTi alloy at 5000× magnification Slika 7: a) Posnetek komercialno dostopne NiTi zlitine na svetlobnem mikroskopu pri 100× pove~avi in b) SE slika komercialno dostopne NiTi zlitine pri 5000× pove~avi alloy consist of two phases. EDS analysis showed that the prevailing phase is NiTi, containing 50 % of amount fractions of Ni and 50 % of amount fractions of Ti. The second phase is carbon rich phase (33.3 % of amount fractions of C, 40.12 % of amount fractions of Ti, 26.59 % of amount fractions of Ni). Figure 8 shows the distribution of elements in the individual phases. 3.3 Potentiodynamic test Figure 9 shows the potentiodynamic curves for NiTi strand and commercially available NiTi alloy, while Table 1 contains the quantitative results of the measure- ments. Corrosion potential and current, and breakdown potential and current values were obtained by graphic extrapolation. The corrosion potential of the NiTi strand is 38 mV higher than for the commercially available NiTi alloy, which means that the passive layers spontaneously developed on the NiTi strand are less affected by envi- ronmental factors. The presence of a wider passivation range was observed for the commercially available NiTi alloy, while for the NiTi strand the passivity occurs in a narrower range of potentials, indicating a higher ten- dency for localized corrosion. On the surface of the nitinol a double layer is formed. The outer layer is TiO2 and the inner layer is TiNi3. When the thickness of the TiO2 layer increases, two phenomena play a competing role. First, since Ni atoms are diffusing further away from the surface, they accumulate in the region with the lowest oxidation state (close to the oxide–metal inter- face). Second, as TiNi3 appears as a line phase in the Ni–Ti phase diagram, the amount of Ni in the inter- metallic TiNi3 layer becomes saturated upon formation of this layer. As a result it will be more energetically favourable to form metallic particles within the TiO2 layer than increase the thickness of the intermetallic layer.20 Breakdown of the passive film occurs as a result of thickening of the oxide layer, leading to an increase in the size of the nickel particles in the outer oxide layer. These particles cause local stress, so the layer cracks, which facilitates the progress of corrosion. The commer- cially available NiTi alloy has higher breakdown po- tential, meaning it will form thicker oxide layer before the collapse. The corrosion rate of the commercially available NiTi alloy is lower, so it is more corrosion resistant. 3.4 Impedance test Electrochemical impedance spectroscopy (EIS) measurements were performed at open circuit potential conditions in a simulated physiological fluid for 8 days. Figure 10 shows the Nyquist impedance diagrams for the NiTi strand and the commercially available NiTi alloy. The analysed spectra proposed an equivalent circuit, considering an outer titanium oxide layer with A. STAMBOLI] et al.: CONTINUOUS VERTICAL CASTING OF A NiTi ALLOY 986 Materiali in tehnologije / Materials and technology 50 (2016) 6, 981–988 Figure 10: Nyquist diagrams for the NiTi strand and the commercially available NiTi alloy with corresponding fit after a) 12 h, b) 96 h, and c) 168 h of immersion Slika 10: Nyquistovi diagrami NiTi palice in komercialno dostopne NiTi zlitine, z ustreznimi prilegajo~imi krivuljami po ~asu izpostavljenosti: a) 12 h, b) 96 ur, in c) 168 h Figure 9: Potentiodynamic curves for NiTi strand and commercially available NiTi alloy Slika 9: Potenciodinamske krivulje NiTi palice in komercialno do- stopne NiTi zlitine Table 1: Electrochemical parameters determined from the poten- tiodynamic curves measured for the NiTi strand and the commercially available NiTi alloy Tabela 1: Elektrokemijski parametri, dolo~eni iz potenciodinamskih krivulj, izmerjenih za NiTi palico in komercialno dostopno NiTi zlitino Ecorr (mV) Icorr (μA) Ebd (mV) Ibd (μA) vcorr (mmpy) NiTi strand -287.1 0.343 348.5 6.767 3.201·10–3 com. NiTi alloy -324.9 0.328 625.8 5.948 2.828·10–3 Ecorr – corrosion potential determined from potentiodynamic curves; Icorr – corrosion current; Ebd – breakdown potential; Ibd – breakdown current; and vcorr – corrosion rate corrosion resistance R1 and an inner TiNi3 layer with resistance R2 (Figure 11), where Rs is the resistance of the solution. The use of a constant phase element (CPE) was required to account for the non-ideal capacitive response observed as a depressed semicircle when the spectra were plotted in the corresponding Nyquist diagrams. The CPE originates from the surface rough- ness and inhomogeneities present in the titanium oxide layers at the microscopic level.21 Table 2: Corrosion resistance of NiTi strand and commercially available NiTi alloy in outer (R1) and inner (R2) oxide layer, and total corrosion resistance Rp at certain time of immersion Tabela 2: Korozijska odpornost NiTi palice in komercialno dostopne NiTi zlitine v zunanji (R1) in notranji (R2) plasti oksida ter skupna odpornost proti koroziji Rp pri dolo~enem ~asu izpostavljenosti t/h R1com/ R2com / R1strand / R2strand / Rp, com / Rp, strand / 1 180400 328850 10110 457060 509250 467170 2 170510 640510 10812 638590 811020 649402 12 265080 603270 11854 736850 868350 748704 24 333940 695490 12652 764400 1029430 777052 48 387740 1048800 12903 939500 1436540 952403 72 544540 1919700 28966 1206800 2464240 1235766 96 702330 3989100 9377 1470000 4691430 1479377 120 798180 5642300 5493 1500600 6440480 1506093 144 843170 6984000 25020 1511900 7827170 1536920 168 917140 8901100 96538 1526000 9818240 1622538 192 1018900 8101300 35321 1447100 9120200 1482421 As shown in Table 2, the resistances of the outer and inner oxide layers in the commercial NiTi alloy are very similar and very high, while the difference in resistance between the outer and the inner layer by the NiTi strand is very high. This means that the outer layer of the NiTi strand has Ni particles, which are the weakest link in the corrosion resistance of the NiTi alloy. Resistance values in the outer layer of NiTi strand are so low (< 10000 ), that they present no obstacle in the progress of corrosion that can occur hazardous nickel ions outleaching from this layer into the surrounding area. Corrosion has slower progress in the inner TiNi3 layer. Figure 12 represents the polarization or a totally corrosion resistance Rp as a function of time. Rp can be calculated according to Equation (1): Rp = R1 + R2 (1) as a function of time. The slope of the commercial NiTi alloy increases rapidly with time, while the slope of NiTi strand increases slightly with time. It is clear that the corrosion resistance of the commercial NiTi alloy is much greater than that of the NiTi strand at any time. The main reasons for the poorer corrosion resistance of the NiTi strand are a lower homogeneity and a lower titanium content. 4 CONCLUSIONS From this study the following conclusions can be drawn: • a dendritic microstructure of the NiTi strand was formed while VIM+CVC, • the chemical composition of the NiTi strand varied through the cross and longitudinal sections, so the drawing process by CVC is not optimal, • TiC and Fe phases were identified in the NiTi strand, • the commercially available NiTi alloy has a higher breakdown potential than the NiTi strand, meaning it will have thicker, more stable oxide layer before the collapse, • the corrosion resistance of the commercial NiTi alloy is much greater than that of NiTi strand at any time, • 10% deficit of titanium in NiTi strand is reflected in poorer corrosion resistance properties, • despite the fact that the corrosion resistance of the NiTi strand is not sufficient, we have successfully cast NiTi strand by VIM + CVC processes, so it is evident that it is possible to produce such an alloy in this way. A. STAMBOLI] et al.: CONTINUOUS VERTICAL CASTING OF A NiTi ALLOY Materiali in tehnologije / Materials and technology 50 (2016) 6, 981–988 987 Figure 12: Rp vs time diagram for NiTi strand and commercially available NiTi alloy Slika 12: Diagram Rp proti ~asu NiTi palice in komercialno dostopne NiTi zlitine Figure 11: Equivalent circuit of two-layer model used for the interpretation of the measured impedance spectra of NiTi alloy Slika 11: Ekvivalentno vezje uporabljeno za razlago izmerjenih impedan~nih spektrov NiTi zlitine na osnovi dvoslojnega modela 5 REFERENCES 1 A. Tuissi, P. Bassani, A. Mangioni , L. Toia, F. Butera, Fabrication process and characterization of NiTi wires for actuators, SMST- 2004: Proceedings of the International Conference on Shape Memory and Superelastic Technologies, Baden-Baden, 2004, 501–508 2 P. R. Halani, I. Kaya, Y. C. Shin, H. E. 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J. Santana, R. M. Souto, Multi- scale electrochemical analysis of the corrosion of titanium and nitinol for implant applications, Electrochimica Acta, 203 (2016), 366–378, doi:10.1016/j.electacta.2016.01.146 A. STAMBOLI] et al.: CONTINUOUS VERTICAL CASTING OF A NiTi ALLOY 988 Materiali in tehnologije / Materials and technology 50 (2016) 6, 981–988 F. TEHOVNIK et al.: HOT TENSILE TESTING OF SAF 2205 DUPLEX STAINLESS STEEL 989–993 HOT TENSILE TESTING OF SAF 2205 DUPLEX STAINLESS STEEL VRO^I NATEZNI PRESKUSI DUPLEKS NERJAVNEGA JEKLA SAF 2205 Franc Tehovnik, Borut @u`ek, Jaka Burja Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia franc.tehovnik@imt.si Prejem rokopisa – received: 2016-08-04; sprejem za objavo – accepted for publication: 2016-09-06 doi:10.17222/mit.2016.242 The changes to the microstructure of the duplex stainless steel SAF 2205 during hot tensile tests were investigated. The dominant restoration mechanisms for ferrite and austenite were dynamic recovery (DRV) and dynamic recrystallization (DRX), respectively. Also, the effect of temperature on the deleterious phase precipitation was investigated. The specimens were tested with hot tensile deformation tests in the temperature range from 800 °C to 1100 °C. The tensile strength of the investigated steel decreased rapidly in the temperature range from 800 °C to 950 °C and slowly at the testing temperatures from 1000 °C to 1100 °C. It was found that SAF 2205 has excellent hot-working properties at temperatures between 950 °C and 1100 °C. The differences in the mechanical properties of austenite and ferrite along with the precipitation of the -phase represent the most important restrictions during hot working. Optical microscopy was used for the microstructure-evolution analysis in the duplex stainless steel during the hot tensile tests. Keywords: duplex stainless steel, hot tensile test, microstructural evolution, intermetallic phases, -phase Preiskovali smo spremembe mikrostrukture dupleks nerjavnega jekla SAF 2205 med vro~imi nateznimi preskusi. Prevladujo~a mehanizma meh~anja v feritu in avstenitu sta bila dinami~na poprava in/ali dinami~na rekristalizacija. V dupleks nerjavnem jeklu je bil raziskan vpliv temperature na izlo~anje {kodljivih faz. Vzorce smo vro~e natezno presku{ali v temperaturnem obmo~ju med 800 °C in 1100 °C. Natezna trdnost preiskanega jekla se je hitro zni`ala v temperaturnem obmo~ju med 800 °C in 950 °C ter po~asneje pri temperaturah preskusov med 1000 °C in 1100 °C. Ugotovili smo, da ima SAF 2205 odli~ne vro~e preoblikovalne lastnosti v temperaturnem obmo~ju med 950 °C in 1100 °C. Najve~je omejitve pri preoblikovanju predstavljajo razlike v mehanskih lastnostih med avstenitom in feritom ter izlo~anje -faze. V raziskavi je bila uporabljena opti~na mikroskopija za analizo razvoja mikrostrukture v dupleks nerjavnem jeklu med vro~imi nateznimi preskusi. Klju~ne besede: dupleks nerjavno jeklo, vro~i natezni preskusi, razvoj mikrostrukture, intermetalne faze, -faza 1 INTRODUCTION Duplex stainless steels (DSS) have the advantage of low price, pitting-corrosion resistance, stress-corrosion- cracking resistance, resistance to intergranular corrosion, high mechanical strength, corrosion-fatigue resistance, wear resistance, super plastic behaviour and good weld- ability.1–8 But they have a severe disadvantage in being difficult to hot work.2,9,10 The DSS microstructure generally consists of around 50 % -ferrite and 50 % -austenite, which gives them excellent properties, but also provides for a narrow window in the hot-working process. During hot working, ferrite and ferritic stainless steels undergo dynamic recovery (DRV),11 which means that subgrains develop in the microstructure, as is the case for the hot working of duplex stainless steels,12 while austenite and austenitic stainless steels undergo a high degree of strain hardening and then dynamic recrystallization (DRX).8 While the austenite is dominated by high-angle grain boundaries (HAB), ferrite shows a substantial amount of subgrain boundaries, i.e., low-angle grain boundaries (LAB). Such a contrasting behaviour of austenite and ferrite phases is due to the different stacking-fault ener- gies of these phases, which imposes different deforma- tion mechanisms.13 Therefore, when hot working duplex stainless steels, DRV is the initial softening mechanism. But austenite undergoes DRV only at high temperatures and low strain rates. At low strain rates and high temperatures DSS exhibit superplasticity; therefore, allowing elongations that exceed 1000 %.6 Superplasticity can be linked to grain-boundary sliding, but grain-boundary sliding also causes cavity formation, which in turn causes stress concentration, and if these stresses cannot be released at sufficient rates, the cavities nucleate.14 The cavities, provided with appropriate conditions, then undergo the processes of growth and coalescence and form larger cavities that can lead to failure.14 Another factor that negatively influences the mechanical properties is the precipitation of secondary phases,15 especially the -phase, which occurs in the temperature range from 700 °C to 1000 °C and is promoted by Cr, Mo and Si.12,16,17 The -phase reduces the ductility and toughness of the steel.12 Increased steel hardness is an indicator of -phase precipitation.12 The precipitation phenomena in duplex stainless steels occur mainly in the -ferrite, Materiali in tehnologije / Materials and technology 50 (2016) 6, 989–993 989 UDK 621.78:620.172:669.14.018.8 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 50(6)989(2016) because the diffusion rates are much faster than in the austenite. The genesis of the precipitates can be attri- buted to the large amount of alloying element and is promoted by the instability of the ferritic matrix at the high temperatures. The morphology and location of the secondary phases are well known and described,18 they precipitate both at triple points and grain boundaries, and their growth occurs towards the unstable ferrite, also due the diffusion behaviour of the involved elements. 2 MATERIALS AND EXPERIMENTAL PART Duplex stainless-steel grade SAF 2205, with the che- mical composition given in Table 1, was used in the experimental work. The initial microstructure is presented in Figure 1. Both phases have an almost fully recrystallized micro- structure produced by the annealing treatment at 1050 °C before deformation testing. The tensile test specimens of SAF 2205 were made according to DIN 50125. The specimens were cut out from the hot-rolled plates longitudinal to the rolling direction. Tensile tests were performed at elevated tem- peratures of (800, 850, 900, 950, 1000, 1050 and 1100) °C on a 500-kN INSTRON tensile test machine with a spe- cial furnace mounted to ensure high-temperature condi- tions. The specimens were heated at a rate of 700 °C/h, and were held at the deformation temperature for 15 min. The deformation speed was 5 mm/min, which gives a logarithmic strain rate of 0.0014 s–1. During the hot tensile tests the force and extension were recorded simul- taneously. At the end of the tensile tests, the specimens were air cooled. The samples for metallographic analysis of the tensile specimens were taken 10 mm from the fracture surface in the contraction area, longitudinal to the testing direction. The samples were etched with a solution of 30 g KOH, 30 g K3Fe(CN)6 and 100 mL distilled water. The etching causes the phases in the microstructure to colour differently: the -phase is dark, the -ferrite is lighter and the -austenite is the brightest in colour. The me- tallographic samples were observed with an optical microscope (Microphot FXA, Nikon). The content of the magnetic -ferrite was determined by a ferritoscope instrument FISCHER MP30. 3 RESULTS AND DISCUSSION The initial microstructure of the steel is composed from austenitic grains (light) that are oriented in the rolling direction and recrystallized ferrite grains (dark). Individual "islands" of austenite grains are located in the ferrite areas. Specimens that were tested at (800, 850 and 900) °C contain -phase, while those that were tested at higher temperatures contain no -phase. The specimen’s microstructures after the tensile tests are shown in Figures 2a to 2g. Cavitations or cracks that formed at high deformations are visible in Figures 2d to 2f at temperatures from 950 °C to 1050 °C. The Figures 2d to 2f show some cracks that nucleated at the auste- nite/ferrite interfaces and propagated along the interface towards the softer ferrite phase. The same phenomenon was also found by M. Faccoli and R. Roberti.19 Due to a notable difference in strength between the ferrite and austenite, voids are formed at the austenite/ferrite inter- faces. Deformation at high temperatures, however, allows dynamic restoration (DRV and DRX) to take affect and reduce the work hardening, thus reducing the probability of void formation. It is clear that the hot tensile strength of the steel at deformations up to  = 0.5 gradually increases at tem- peratures around 800 °C, and this can be partially said for the test at 850 °C. But at the temperature of 900 °C the tensile strength does not change significantly due to the established equilibria between the strain hardening and the softening effects. Figure 2a shows the micro- structure of the sample tested at the lowest temperature. The austenite phase appears in the form of elongated islands. In the sample tested at the lowest temperature (Figure 2a, 2b) their distribution is not homogeneous and the austenite/ferrite interfaces are affected by -pha- se precipitation. F. TEHOVNIK et al.: HOT TENSILE TESTING OF SAF 2205 DUPLEX STAINLESS STEEL 990 Materiali in tehnologije / Materials and technology 50 (2016) 6, 989–993 Table 1: Chemical composition SAF 2205 in mass fractions, (w/%) Tabela 1: Kemijska sestava jekla SAF 2205 v masnih odstotkih, (w/%) C Si Mn P S Cr Ni Cu Mo V Ti Nb Al N 0.021 0.32 1.58 0.026 0.002 22.95 5.3 0.26 2.742 0.15 0.005 0.008 0.012 0.141 Figure 1: Initial microstructure of SAF 2205 duplex stainless steel, austenite (light) and ferrite (dark), light microscopy Slika 1: Za~etna mikrostruktura dupleks nerjavnega jekla SAF 2205, avstenit (svetel) in ferit (temen), svetlobna mikroskopija At higher temperatures from 950 °C to 1100 °C there is a significant drop in tensile strength due to the soften- ing effects, and the initial strengths are not reached. Figures 2d to 2g show more equiaxed austenitic grains, a sign of the recrystallization process. As the temperature increases, the austenite volume fraction decreases and the austenite phase islands are progressively reduced in length and appear more dis- continuous; -islands which are almost totally dissolved assume a globular shape. According to A. Iza-Mendia et al.,20 austenite produces non-deformed regions at this temperature. The strain difference between the two phases is, therefore, responsible for the severe shear strains at the phase boundaries, sliding or even cracks. The stress–strain curves for the tests at different tem- peratures (from 800 °C to 1100 °C) at the strain rate of 0.0014 s–1 are shown in Figure 3. According to Figure 3, the optimum hot-working temperature range at the strain rate of 0.0014 s–1 should be between 950 °C and 1100 °C. The worst hot-working properties are at 800 °C. Generally, the flow stress rises to a maximum at the commencement of the straining before dropping to a steady-state level. The shape of the curves also changes as the material and deformation conditions are altered. At high temperatures, above 950 °C, the flow curves have the characteristic shape expected for materials that show dynamic recovery. As the deformation temperature is decreased the shape changes with rapid work harden- ing, followed by extensive flow softening. These changes become more dramatic when the austenite volume fraction is further increased. As the volume fraction of austenite in the ferrite matrix is increased, the ductility also decreases The results of the tensile strength and the reduction of area that were obtained from hot tensile tests are summed up in Figure 4. The best hot-working properties are achieved at higher temperatures, as the material has the lowest tensile strengths and the highest contractions and elongations (Figure 4). F. TEHOVNIK et al.: HOT TENSILE TESTING OF SAF 2205 DUPLEX STAINLESS STEEL Materiali in tehnologije / Materials and technology 50 (2016) 6, 989–993 991 Figure 2: Microstructures of specimens after tensile tests at different temperatures, light microscopy Slika 2: Mikrostuktura presku{ancev po vro~ih nateznih testih pri raz- li~nih temperaturah, svetlobna mikroskopija Figure 4: Tensile strength and contraction of SAF 2205, depending on the testing temperature in the range from 800 °C to 1100 °C Slika 4: Natezna trdnost in kontrakcija SAF 2205 v odvisnosti od tem- perature preizku{anja v temperaturnem obmo~ju od 800 °C do 1100 °C Figure 3: Stress–strain (logarithmic) diagram for hot tensile tests at different temperatures Slika 3: Graf napetost – (logaritemska) deformacija vro~ih nateznih preskusov pri razli~nih temperaturah There is a rapid increase in the reduction area as the temperature rises from 800 °C to 850 °C, from 850 °C on it gradually increases with temperature, while the fall in tensile strength is more continuous. The high-temperature mechanical properties and hot-working properties depend mostly on the microstruc- tural changes. DRV can be observed as a polygonisation of the subgrains; it is more pronounced at high tem- peratures and lower deformation speeds and lower loads. DRX can be observed as the formation of new, equiaxed grains. Another process that can influence the hot-work- ing properties is the precipitation of the intermetallic phases from the matrix, namely the -phase. The equilibrium phase composition for SAF 2205 was calculated with Thermo-Calc, the results are pre- sented in Figure 5. As shown in Figure 5, -phase formation does not occur at temperatures above 930 °C, while the sample at 950 °C still contains some -phase. With a decrease of temperature, the austenite weight fraction increases significantly from 28 % at 1200 °C up to 67 % at 860 °C. The change in austenite volume fraction influences the mechanical behaviour of the duplex stainless steel because of the large difference in strength between the ferrite and austenite. The content of ferrite in the samples was measured with the ferritoscope. The results of the measurements are presented in Figure 6. Figure 6 shows the evolution of the -ferrite volume fraction at different test temperatures; the amount of -ferrite increases between 850 °C and 1000 °C. According to Figure 5, the amount of -phase formed is at least a factor of 5 greater than the amount of chi phase. Predictions by Thermo-Calc are relatively good for the amount of ferrite at 1000 °C; the predicted value is about 48 % compared to 46 % measured. Above 850 °C, the amount of austenite decreases considerably with the deformation temperature, which is attributed to the    transformation. The content of ferrite rises at temperatures from 950 °C to 1100 °C; it is about 45 % at 1100 °C, the rest is austenite. The hot-working properties greatly depend on the austenite and ferrite content in the steel microstructure; their ratio depends on the temperature, as shown on Fig- ure 5. The softening mechanism in ferrite is DRV; it can be observed as the subgrain formation, while the soften- ing mechanism in austenite is the DRX, which is discontinuous, and it mostly occurs on ferrite–austenite phase grain boundaries. The lower driving force for ferrite strain softening is compensated with a higher diffusion rate and a higher mobility of dislocations in a cubic body-centred lattice, that contribute to a faster strain softening process in ferrite. Another important factor that contributes to the difficult hot-working properties of DSS is the occurrence of the -phase. The -phase is formed by the eutectoid transformation:   → + The presence of the -phase is the most prominent at 900 °C. Long exposure times at temperatures up to 900 °C and deformation are sufficient conditions for -phase formation during hot tensile tests. Duplex stainless steels are more susceptible to the precipitation of intermetallic phases than austenitic steels due to the high Cr and Mo contents and higher diffusion rates in the ferrite phase. The precipitation reaction of -phase in austenite is sluggish due to the slow diffusivity of the solute atoms. As the precipitation continues, Cr and Mo diffuse to the -phase, leading to the depletion of these elements in the ferrite, especially the Mo content. There- fore, Mo from the inner region of the ferrite diffuses to the -phase. The -phase nucleates preferentially at the ferrite/ferrite and ferrite/austenite boundaries, and then grows into the adjacent ferritic grains. Mo is the main element controlling secondary-phase precipitation. The -phase is a hard, brittle, non-magnetic intermetallic phase, with high Cr and Mo contents. F. TEHOVNIK et al.: HOT TENSILE TESTING OF SAF 2205 DUPLEX STAINLESS STEEL 992 Materiali in tehnologije / Materials and technology 50 (2016) 6, 989–993 Figure 6: Measured -ferrite weight fraction and mean values at diffe- rent temperatures in the range from 800 °C to 1100 °C Slika 6: Izmerjen masni dele` -ferita in srednje vrednosti pri raz- li~nih temperaturah v obmo~ju od 800 °C do 1100 °C Figure 5: Calculated weight fractions of equilibrium phases in SAF 2205 in the temperature range from 500 °C to 1250 °C Slika 5: Izra~unani masni dele`i ravnote`nih faz v SAF 2205 v temperaturnem obmo~ju od 500 °C do 1250 °C The partitioning of the elements between the ferrite and austenite takes place during DSS and contributes to the difference in hot strength between the ferrite and austenite. This means that some alloying elements can dissolve preferentially in one phase compared to the other, depending on the nature of the considered che- mical element: austenite or ferrite stabilizer. At both extremes, Mo is the element that segregates mostly to ferrite, whereas C and N tend to leave the ferrite. The high N content in austenite is important, as the solute- strengthening effect tends to increase the hot strength of the austenite. The partition of elements changes with temperature. The content of ferrite-forming elements in the austenite phase decreases with temperature, whereas the austenite stabilizer content increases. As a conse- quence the hot-deformation behaviour of duplex steels can be different at the beginning and at the end of the hot-deformation process. The higher the temperature, the more uniform the element partitioning is between ferrite and austenite. 4 CONCLUSIONS The following conclusions can be drawn from the experimental work: The SAF 2205 duplex stainless steel has excellent hot-working properties at temperatures between 950 °C and 1100 °C, while hot working is not recommended at 800 °C. Difficulties in hot working due to -phase precipitation can be accurately predicted by Thermo-Calc software. The -phase is mostly observed below 950 °C. The microstructure of the steel is ferrite + austenite at room temperature, ferrite + austenite + from 800 °C to 950 °C and ferrite + austenite at higher temperatures 1000 °C to 1100 °C. The specimens that were tensile tested at lower temperatures (800–900 °C) broke at lower strains due to severe -phase precipitation and diminished softening mechanisms. These specimens did not show cavity occurrence 10 mm from the fracture surface, as the fracture was more localized, while the higher temperature tests from 950 °C to 1050 °C had cavities occurring even 10 mm from the fracture surface. The highest testing temperature, however, resulted in a microstructure that showed no apparent damage. This suggests that the dominant reason for failures at temperatures from 950 °C to 1050 °C is the difference in mechanical properties between the ferrite and austenite that are further increased by the differences in the softening mechanisms. 5 REFERENCES 1 Z. Y. Liu, C. F. Dong, X. G. Li, Q. Zhi, Y. F. Cheng, Stress corrosion cracking of 2205 duplex stainless steel in H2S–CO2 environment, J. Mater. Sci., 44 (2009) 4228–4234, doi: 10.1007/s10853-009-3520-x 2 A. Momeni, K. Dehghani, Hot working behavior of 2205 austenite- ferrite duplex stainless steel characterized by constitutive equations and processing maps, Mater. Sci. Eng. A, 528 (2011) 1448–1454, doi:10.1016/j.msea.2010.11.020 3 G. W. Fan, J. Liu, P. D. Han, G. J. Qiao, Hot ductility and micro- structure in casted 2205 duplex stainless steels, Mater. Sci. Eng. A, 515 (2009), 108–112, doi:10.1016/j.msea.2009.02.022 4 N. Ortiz, F. F. Curiel, V. H. López, A. 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Magrini, Effect of isothermal heat treatments on Duplex Stainless Steels impact toughness. in Con- vegno Nazionale IGF XXII, Rome, 2013, 56–65 19 M. Faccoli, R. Roberti, Study of hot deformation behaviour of 2205 duplex stainless steel through hot tension tests, J. Mater. Sci., 48 (2013), 5196–5203, doi:10.1007/s10853-013-7307-8 20 A. Iza-Mendia, A. Pinol-Juez, J. J. Urcola, I. Gutiérrez, Micro- structural and Mechanical Behavior of a Duplex Stainless Steel under Hot Working Conditions, Metall. Mater. Trans. A, 29 (1998), 2975–2986 F. TEHOVNIK et al.: HOT TENSILE TESTING OF SAF 2205 DUPLEX STAINLESS STEEL Materiali in tehnologije / Materials and technology 50 (2016) 6, 989–993 993 C.-C. KUO, C.-M. HUANG: A HIGH-EFFICIENCY AUTOMATIC DE-BUBBLING SYSTEM FOR LIQUID SILICONE RUBBER 995–1000 A HIGH-EFFICIENCY AUTOMATIC DE-BUBBLING SYSTEM FOR LIQUID SILICONE RUBBER VISOKOZMOGLJIV SISTEM ZA ODPRAVLJANJE MEHUR^KOV V TEKO^I SILIKONSKI GUMI Chil-Chyuan Kuo, Chuan-Ming Huang Ming Chi University of Technology, Department of Mechanical Engineering, No. 84, Gungjuan Road, Taishan Dist. New, Taipei City 24301, Taiwan jacksonk@mail.mcut.edu.tw Prejem rokopisa – received: 2014-06-13; sprejem za objavo – accepted for publication: 2016-10-27 doi:10.17222/mit.2014.089 The silicone-rubber mold is regarded as an important method for reducing the cost and time to market a new product by shortening the development phase. A commercial, automatic, vacuum machine is widely used to degas in the manufacturing of a silicone-rubber mold, but the hardware is costly. A low-cost, high-efficiency degassing system was designed and implemented from a regular vacuum machine. The control style was based on a human-machine interface. It was found that the whole degassing-process sequences consist of the explosive phase, the balanced phase and the convergence phase. The time saving in the degassing process can be at least 42 %. Six predicted equations for both the balanced phase and the convergence phase are investigated and the maximum relative error of these equations can be controlled to within 6.34 %. The advantages of the developed de-bubbling system include saving labor, reducing the human error of the operator and a higher degassing efficiency. Keywords: air bubbles, de-bubbling, silicone rubber mold, rapid tooling Forma iz silikonske gume je pomembna pri zmanj{evanju stro{kov in skraj{anju ~asa razvojne faze pri uvajanja novega izdelka na trg. Pri izdelavi forme iz silikonske gume se za razplinjevanje uporabljajo drage komercialne vakuumske avtomatske naprave. Iz obi~ajne vakuumske naprave je bil izdelan poceni in u~inkovit razplinjevalni sistem. Kontrola stroja temelji na povezavi ~lovek-naprava. Ugotovljeno je, da razplinjevanje sestoji iz eksplozivne faze, faze uravnote`enja in iz faze konvergence, Postopek razplinjevanja omogo~a 42 % prihranek ~asa. Preiskanih je bilo {est predvidenih ena~b pri obeh uravnote`enih fazah in pri konvergen~ni fazi in maksimalna napaka zna{a 6,34 %. Prednosti razvitega sistema razplinjevanja so: prihranek dela, zmanj{anje ~love{kih napak operaterja in ve~ja u~inkovitost razplinjevanja. Klju~ne besede: zra~ni mehur~ki, odprava mehur~kov, forma iz silikonske gume, hitra izdelava orodij 1 INTRODUCTION To reduce the time and the cost of product develop- ment, rapid prototyping (RP) was developed.1 This offers the potential to completely revolutionize the process of manufacture. However, the features of the prototype do not usually meet the needs of the end product with the required material. Rapid tooling (RT) technologies are then developed because it is the technology that uses RP technologies and applies them to the manufacturing of mold inserts.2 Since the importance of RT goes far beyond component performance testing, RT is regarded as an important method of reducing the costs and the time to market in a new-product development process. Several RT technologies are commonly available in industry now. RT is divided into direct tooling and indi- rect tooling.3 Direct tooling means fabricating mold inserts directly from an RP machine, such as selective laser sintering.4,5 Indirect tooling means fabricating the mold insert by a master pattern fabricated using various RP technologies. Soft tooling is used for low-volume production. The materials used for soft tooling have low hardness levels, such as silicone-rubber6 and epoxy-resin composites.7 Conversely, hard tooling is associated with higher volumes of production. Materials used for hard tooling often have high hardness levels. Soft tooling is easier to work with than tooling steels because these tools are created from materials such as epoxy-based composites with aluminum particles, silicone rubber or low-melting-point alloys. It is well known that RT is capable of replacing conventional steel tooling, saving costs and time in the manufacturing process.8 Indirect soft tooling is used more frequently in the development of new products than direct tooling, because it is fast, simple and cost-effective. It is a well-known fact that a silicone rubber mold is employed frequently because it has flexible and elastic characteristics, so that parts with sophisticated geometries can be fabricated.9 A silicone- rubber mold can be used for producing low-melting- point metal parts, wax patterns and plastic parts. Air bubbles in the silicone-rubber mold are one of the most common types of defects, especially in the vicinity of the master pattern. A silicone-rubber mold with air bubbles appearing in the vicinity of the master pattern will change the appearance and dimensional accuracy of the part duplicated from this silicone-rubber mold using vacuum casting. Conventionally, de-bubbling the air bubbles with a purely manual operation mode depends Materiali in tehnologije / Materials and technology 50 (2016) 6, 995–1000 995 UDK 677.017.7:678.84:678.032 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 50(6)995(2016) significantly on the experiences of the operator. The drawbacks of this method include human error and noise pollution derived from the vacuum pump. A com- mercially available automatic vacuum machine was widely employed to degas in the manufacturing of sili- cone-rubber molds, but the hardware is costly.10 In addition, the programmable logic controller mode lacks the flexibility to modify the program. Hence, developing a low-cost and easy-to-operate automated de-bubbling system is a major concern. To meet this requirement, the objective of this work is to develop a low-cost and high-efficiency automatic de-bubbling system with a human-machine interface (HMI).11 Three vessels were used for filling the liquid silicone rubber for de-bubbling. The de-bubbling process sequences were investigated in detail. The effect of the pressure-relief process in the explosive phase on the de-bubbling efficiency for manual and automatic modes was also analyzed. Trend equations for predicting the balance phase duration and conver- gence phase duration for three vessels were investigated. The performance of the automatic de-bubbling system developed was evaluated. Comparisons of the de-bubbl- ing efficiencies for automatic and manual modes were compared. 2 EXPERIMENTAL PART Figure 1 shows a low-cost, automated de-bubbling system with an HMI. This system consists of a photo- electric sensor (EX-11EB; SUNX, Inc.), a programmable logic controller (PLC) (FX 2N-32MR, Mitsubishi), an HMI (GP37W2-BG41-24V; Pro-face, Inc.), an electro- magnetic valve (SUG 15-24VDC; Chelic, Inc.) and an electronic buzzer (TS2BCL; Tend, Inc.). A photoelectric sensor (response time 0.5 ms) was used for detecting the air bubbles during the de-bubbling process. A PLC was used as a key component of the electric control module. An HMI was used for operating the automatic de- bubbling system. The HMI consists of all the aspects of the interaction and communication between the operators and machines by using a graphical HMI unit. The electromagnetic valve (on-off reaction time < 15 ms) was used to break the vacuum automatically. An electronic buzzer was used to alert the operator when the de- bubbling process is completed. Considering the practical requirements for the different operators, the de-bubbling method of this system includes an automation mode and a manual control mode. This system can be equipped with manual and automatic modes. Three different volu- mes (250 mL, 500 mL and 1000 mL) of vessels were used for filling the liquid silicone rubber in this study. The silicone rubber (KE-1310ST; ShinEtsu, Inc.) in the liquid state was mixed with a hardener (CAT-1310S; ShinEtsu, Inc). Generally, the curing agent and silicone C.-C. KUO, C.-M. HUANG: A HIGH-EFFICIENCY AUTOMATIC DE-BUBBLING SYSTEM FOR LIQUID SILICONE RUBBER 996 Materiali in tehnologije / Materials and technology 50 (2016) 6, 995–1000 Figure 2: Brief flowchart of the automatic vacuum degassing process using an HMI Slika 2: Shema poteka postopka avtomatskega vakuumskega razpli- njevanja s pomo~jo HMI Figure 1: A low-cost, automated, de-bubbling system with an HMI Slika 1: Stro{kovno ugoden in avtomatiziran sistem za odpravo me- hur~kov s HMI rubber in a weight ratio of 10:1 were mixed thoroughly with a stirrer. After the de-bubbling process, the pressure inside the vacuum machine was changed by breaking the vacuum atmosphere. Thus, a silicone-rubber mold can be fabricated with defects caused by the air bubbles derived from the mixing process. To reduce the difference caused by the different operators in the amount of air bubbles, while mixing the liquid silicone rubber, an agitation blade for mixing the liquid silicone rubber was designed and fabricated. To confirm the center of the vessel is aligned with the center of the agitation blade, a positioning fixture was designed and fabricated. Figure 2 shows a brief flowchart of the automatic vacuum degassing process using an HMI. Due to the experimental limitations, the solenoid valve does not work when the break vacuum duration is set less than 0.03 s. Thus, the break vacuum duration was set to be 0.03 s. 3 RESULTS AND DISCUSSION The center of the vessel can be aligned with the center of the agitation blade using the positioning fix- ture. Figure 3 shows the three phases of the de-bubbling process sequences. In general, the liquid silicone rubber has a large number of air bubbles because of the mixing reaction of the materials. As can be seen from this figure, the first phase is the explosive phase, where the volume of liquid silicone rubber is drastically increased because the pressure in the vacuum chamber is lower than the atmospheric pressure, as shown in Figure 4a. The second phase is the balance phase, where the air bubbles in the liquid silicone rubber are eliminated gradually, as shown in Figure 4b. The final phase is the convergence phase, where the number and the size of the air bubbles in the liquid silicone rubber are eliminated significantly, as shown in Figure 4c. Schematic illustrations of the three phases of the de-bubbling process are shown in Figure 5. A number of air bubbles less than 15 is defined as the end of balance phase. In addition, no more air bubbles inside the liquid silicone rubber is defined as the end of convergence. A stopping vacuum of 120 s bet- ween convergence phase and convergence phase is required for reducing the convergence phase’s duration. This is because air bubbles inside the liquid silicone rubber can reach the top of the liquid silicone rubber during 120 s. The air bubbles can be eliminated comple- tely by an automatic de-bubbling system with an HMI based on the criteria discussed above. Table 1: Time saving of the explosive phase for different volumes of silicone rubber Table 1: Prihranek ~asa pri eksplozivni fazi, pri razli~nih prostorninah silikonske gume Volume of silicone rubber (mL) Percentage of volume (%) Mode Explosive phase du- ration (s) Time saving (%) 350 70 Manual 5335 61.46 Automatic 2056 300 60 Manual 3952 81.12 Automatic 746 250 50 Manual 2883 82.59 Automatic 502 200 40 Manual 706 70.00 Automatic 233 150 30 Manual 283 61.48 Automatic 109 100 20 Manual 0 0 Automatic 0 Table 1 shows the time saving of the explosive phase for different volumes of silicone rubber. As can be seen, C.-C. KUO, C.-M. HUANG: A HIGH-EFFICIENCY AUTOMATIC DE-BUBBLING SYSTEM FOR LIQUID SILICONE RUBBER Materiali in tehnologije / Materials and technology 50 (2016) 6, 995–1000 997 Figure 4: Schematic illustrations of the three phases of the de-bubbl- ing process: a) explosive phase, b) balance phase and c) convergence phase Slika 4: Shematski prikaz treh faz postopka razplinjevanja: a) eksplo- zivna faza, b) ravnote`na faza in c) konvergen~na faza Figure 3: Three phases of the de-bubbling process: a) explosive phase, b) balance phase and c) convergence phase Slika 3: Tri faze postopka razplinjevanja: a) eksplozivna faza, b) ravnote`na faza in c) konvergen~na faza the maximum time saving of the explosive phase is about 81.12 %. The time savings of the explosive phase in- crease with the increasing percentage of volume, ranging from 30 % to 60 %. The time saving of the explosive phase decreases when the percentage of volume exceeds 60 %. These results show that the vessel with a percent- age of volume of 60% of silicone rubber is the optimum value. Figure 5 shows the sequential procedures of the automatic de-bubbling process. Figure 6 shows the sequential procedures of the de-bubbling process with manual operation. It is obvious that the results for the two operation modes are the same, showing no air bubbles inside the liquid silicone rubber. This result shows that the air bubbles inside the liquid silicone rubber can be eliminated completely with the automatic operation mode. This means the system can be used for the production of a high-quality, bubble-free, silicone- rubber mold.12 Precise determination the balance phase’s duration and the convergence duration is an absolute requirement for a high-efficiency, automatic, de-bubbling process. Figure 7 shows the trend equations for the balance duration and convergence duration for vessel volumes of 250 mL, 500 mL and 1000 mL. For a vessel volume of 250 mL, the balance phase’s duration (y) can be pre- dicted from the trend equation of y = 3.316x – 30.8 by the volume of silicone rubber (x). Note that the R2 repre- sents the correlation coefficient. Generally, a higher R2 value (maximum value =1) means a better accuracy of the trend equation.16 Six predicted equations for both the balanced phase and the convergence phase are investi- gated, and the maximum relative error of these equations can be controlled within 6.34 %. This means that both the balanced phase duration and the convergence phase duration can be calculated from these equations. To evaluate the performance of the automatic degass- ing system developed, each test was carried out three times with the mean and the deviation determined. Table 2 shows the time saving of the total degassing time for three different volumes of silicone rubber. Figure 8 shows the total degassing time as a function of silicone rubber for the manual and automatic operation modes. As can be seen, a reduction in the degassing time of at least 42 % can be observed using the automatic degass- ing system with a human-machine interface. It is obvious that there is an increase in time saving of degassing with an increase in the volume of silicone rubber. Three phases are important for the degassing process, but the explosive phase is the most critical one. This is because the pressure-relief process in the automatic operation mode is significantly different from that in the manual operation mode, as shown in Figure 9. The recovery C.-C. KUO, C.-M. HUANG: A HIGH-EFFICIENCY AUTOMATIC DE-BUBBLING SYSTEM FOR LIQUID SILICONE RUBBER 998 Materiali in tehnologije / Materials and technology 50 (2016) 6, 995–1000 Figure 6: Sequential procedures of the de-bubbling process with the manual de-bubbling mode: a) liquid silicone rubber before de-bubbl- ing process, b) explosive phase, c) balance phase, d) pause 120 s, e) convergence phase and f) de-bubbling process was completed Slika 6: Sekven~ni postopki postopka razplinjevanja pri ro~nem vo- denju razplinjevanja: a) teko~a silikonska guma pred razplinjevanjem, b) eksplozivna faza, c) ravnote`na faza, d) pavza 120 s, e) konver- gen~na faza in f) zaklju~en postopek razplinjevanja Figure 5: Sequential procedures of the automatic operation mode: a) liquid silicone rubber before de-bubbling process, b) explosive phase, c) balance phase, d) pause 120 s, e) convergence phase and f) de-bubbling process was completed Slika 5: Sekven~ni postopki pri avtomatskem na~inu dela: a) teko~a silikonska guma pred razplinjevanjem, b) eksplozivna faza, c) ravno- te`na faza, d) pavza 120 s, e) konvergen~na faza in f) zaklju~en postopek razplinjevanja time needed for the pressure of the chamber reaching the degassing pressure with the automatic operation mode is shorter than that with the manual operation mode. Based on the experimental results, the advantages of this sys- tem include saving labor, reducing the human error of the operator and a higher degassing efficiency. This system has broad application prospects in the develop- ment of new products using a silicone-rubber mold. 4 CONCLUSIONS A low-cost, high-efficiency degassing system has been designed, implemented and tested in this study. The entire degassing-process sequences consist of the explosive phase, the balance phase and the convergence phase. The automatic degassing method provides three decisive advantages compared with the manual degass- ing method. The explosive phase has been proved to be a key process for a high-efficiency degassing process. A reduction of the degassing time by at least 42 % can be gained. This system has broad application prospects in the development stage for new products using rapid-tool- ing technology. Acknowledgement This work was financially supported by the Ministry of Science and Technology of Taiwan under contract nos. NSC 102-2221-E-131-012 and NSC 101-2221- E-131-007. C.-C. KUO, C.-M. HUANG: A HIGH-EFFICIENCY AUTOMATIC DE-BUBBLING SYSTEM FOR LIQUID SILICONE RUBBER Materiali in tehnologije / Materials and technology 50 (2016) 6, 995–1000 999 Table 2: Time saving of the total degassing time for three different volumes of silicone rubber Tabela 2: Prihranek ~asa od vsega ~asa razplinjevanja, pri treh razli~nih prostorninah silikonske gume Volume of silicone rubber (mL) Mode Explosive phase duration (s) Balance phase duration (s) Pause (s) Convergence phase duration (s) Total (s) Time saving (%) 110 Manual 1642 338 120 237 2337 45.64 Automatic 583 334 120 233 1270 240 Manual 1415 509 120 427 2471 42.24 Automatic 416 496 120 396 1428 430 Manual 1409 530 120 465 2525 42.07 Automatic 406 508 120 428 1462 Figure 8: Total degassing time as a function of silicone-rubber vol- ume for manual and automatic operation modes Slika 8: Celotni ~as razplinjevanja (v odvisnosti od volumna) silikon- ske gume pri avtomatskem in ro~nem na~inu upravljanja Figure 9: Schematic illustrations of the pressure relief process in the explosive phase: a) manual operation mode and b) automatic operation mode Slika 9: Shematski prikaz postopka spro{~anja tlaka v eksplozivni fazi: a) ro~ni na~in vodenja in b) avtomatski na~in vodenja Figure 7: Trend equations of: a) balance phase duration and b) con- vergence phase duration for vessel volumes of 250 mL, 500 mL and 1000 mL Slika 7: Trend ena~b za: a) trajanje ravnote`ne faze in b) trajanje kon- vergen~ne faze pri prostornini posode 250 mL, 500 mL in 1000 mL 5 REFERENCES 1 M. 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Tan, J. Y. H. Fuh, H. T. Loh, Y. S. Wong, S. C. H. Thian, L. Lu, Micro-mould fabrication for a micro-gear via vacuum casting, Journal of Materials Processing Technology, 192-193 (2007), 334–339, doi:10.1016/j.jmatprotec.2007.04.098 11 A. Jakli~, F. Vode, R. Robi~, F. Perko, B. Strmole, J. Novak, J. Trip- lat, The implantation of an online mathematical model of slab reheating in a posher-type furnace, Mater. Tehnol., 39 (2005) 6, 215–220 12 C. C. Kuo, Y. J. Wang, Development of a micro-hot embossing mold with high replication fidelity using surface modification, Materials and Manufacturing Processes, 29 (2014) 9, 1101–1110, doi:10.1080/ 10426914.2014.912312 C.-C. KUO, C.-M. HUANG: A HIGH-EFFICIENCY AUTOMATIC DE-BUBBLING SYSTEM FOR LIQUID SILICONE RUBBER 1000 Materiali in tehnologije / Materials and technology 50 (2016) 6, 995–1000 T. WEGRZYN et al.: IMPACT TOUGHNESS OF WMD AFTER MAG WELDING WITH MICRO-JET COOLING 1001–1004 IMPACT TOUGHNESS OF WMD AFTER MAG WELDING WITH MICRO-JET COOLING UDARNA @ILAVOST WMD PO MAG VARJENJU Z MIKRO-JET HLAJENJEM Tomasz Wegrzyn1, Jan Piwnik2, Aleksander Borek3, Agnieszka Kurc-Lisiecka4 1Silesian University of Technology, Faculty of Transport, Krasiñkiego 8, 40-019 Katowice, Poland 2Bialystok University of Technology, Mechanical Faculty, Wiejska 45c, 16-351 Bialystok, Poland 3Plasma-system, Towarowa 14, 41-103 Siemianowice Œl¹skie, Poland 4University of D¹browa Górnicza, Rail Transport Department, Cieplaka 1c, 41-300 D¹browa Górnicza, Poland a.kurc@wp.pl Prejem rokopisa – received: 2015-06-30; sprejem za objavo – accepted for publication: 2015-11-05 doi:10.17222/mit.2015.159 The MAG welding process with micro-jet cooling of the weld during the cooling stage was investigated. For micro-jet gases the mixtures of argon with carbon dioxide, oxygen, and nitrogen were tested. This paper presents a piece of information about a new proposal for gas mixtures during micro-jet cooling after welding. Presented is the main information about the influence of various micro-jet gas mixtures on the metallographic structure of the weld metal. The mechanical properties of the welds were presented in terms of various gas mixtures selection for micro-jet cooling. The influence of argon gas mixtures with oxygen and nitrogen for micro-jet cooling after welding are reported for the first time in the technical literature. Keywords: welding, micro-jet cooling, weld, metallographic structure, gas mixtures, GMA welding Preiskovana je bila za~etna faza postopka MAG varjenja z mikro-jet hlajenjem zvara. Za mikro-jet so bile preizku{ene me{anice argona z ogljikovim dioksidom, kisikom in du{ikom. ^lanek predstavlja del informacije o predlogu novih me{anic plinov za mikro-jet hlajenje po varjenju. Dane so informacije o vplivu razli~nih plinskih me{anic za mikro-jet na metalografske strukture zvarjenega materiala. Mehanske lastnosti zvarov so prikazane v smislu izbranih razli~nih vrst me{anic plinov za mikro-jet hlajenje. Vpliv me{anice argona s kisikom in du{ikom za mikro-jet hlajenje po varjenju je prvi~ prikazan tudi v tehni{ki literaturi. Klju~ne besede: varjenje, mikro-jet hlajenje, zvar, metalografska struktura, me{anice plinov, GMA varjenje 1 INTRODUCTION MAG is an important industrial welding process, preferred for its versatility, speed and the relative ease of adapting the process to robotic automation. Develop- ments in arc welding processes are strongly related with the need to increase productivity without losing the quality of the weld.1–5 The reduction of costs and com- petitive pricing are each day more strongly related with technological innovations.6–11 The properties of steel welded structures depend on many factors such as weld- ing technology, filler materials, state of stress. The main role of these conditions is also connected with the ma- terials, the chemical composition of steel and the weld metal deposit (WMD).12–16 The chemical composition of metal weld deposit could be regarded as a very important factor influencing the properties of the weld metal deposit (WMD). In particular, the oxygen, titanium, manganese and aluminium are regarded as the main elements that positively effect the mechanical properties and the metallographic structure of low-alloy welds. This is because of the non-metallic inclusions in weld (Figure 1) that have similar lattice parameter as the ferrite (TiO, TiN, MnO, Al2O3). Materiali in tehnologije / Materials and technology 50 (2016) 6, 1001–1004 1001 UDK 620.178.2:621.791:621.78.08 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 50(6)1001(2016) Figure 1: a) SEM micrographs showing the oxide inclusions in low-alloy WMD after welding with basic electrodes and b) EDS anal- ysis of the WMD1 Slika 1: a) SEM-posnetek prikazuje oksidne vklju~ke v malo legi- ranem WDM po varjenju z bazi~nimi elektrodami in b) EDS analiza WMD1 The welding parameters, metallographic structure and chemical composition of the weld metal deposit are regarded as important factors that influence the impact toughness properties of the deposits.9–12 In a typical low- alloy steel weld structure the best mechanical properties correspond with the maximum amount of acicular ferrite (AF) in the weld metal deposit (WMD) and the mini- mum amount of MAC phases (self-tempered martensite, retained austenite, carbide). This article focuses on mild-steel welding and covers the new possibilities of that method. Since 2011, innovative welding technology based on micro-jet cooling just after welding has been investigated. The weld metal deposit (WMD) was carried out for the standard MAG process and for the innovative welding method with micro-jet cooling. A very high per- centage of acicular ferrite (AF) in WMD was obtainable (55–73 %) for low-alloy steel welding only for micro-jet cooling after the MIG process with argon or helium.13–18 Argon and helium, as micro-jet gases, could provide a better impact toughness of the WMD (0.08 % C, 0.8 % Mn) than in the case of the classic MAG process (Table 1). Table 1 shows that argon is a more beneficial micro- jet cooling gas than helium. Also, it is shown that micro-jet cooling improves the amount of acicular ferrite in the weld. Helium is not such a beneficial micro-jet gas as argon and its mixtures in the MAG process (because of the high percentage of MAC phases). In that paper gas mixtures of argon with a small amount of oxygen and nitrogen were mainly tested because of the positive influ- ence of some oxide and nitride inclusions of acicular ferrite forming and thus the very good impact toughness of the welds. Table 1: Metallographic structure of MAG welds1 Tabela 1: Metalografske strukture MAG zvarov1 Micro-jet gases Ferrite AF MAC phases without micro-jet 43% 4% He 59% 6% Ar 63% 2% 2 EXPERIMENTAL PART The weld metal deposit was prepared by welding with micro-jet cooling with gas mixtures both for the standard MAG process and the MAG welding with micro-jet cooling. The MAG welding process was based on a shielded gas mixture of 79 % Ar and 21 % CO2. To obtain various amounts of acicular ferrite in the WMD the installed micro-jet injector was close to the MAG welding head. The main parameters of the micro-jet cooling were slightly varied: • cooling steam diameter (40 μm and 50 μm), • gas pressure (0.4 MPa and 0.5 MPa), • gas mixtures of argon (82 % Ar/18 % CO2 and 98 % Ar/2 % O2 and 98 % Ar + 2 % N2) were chosen as the micro-jet gases. A montage of the welding head and the micro-jet injector is illustrated in Figure 2. The main data about the parameters of the welding are shown in Table 2. The weld metal deposit was prepared by welding with micro-jet cooling using a larger number of parameters (Table 3). Table 2: Parameters of the welding process Tabela 2: Parametri procesa varjenja No. Parameter Value 1. Diameter of wire 1.2 mm 2. Standard current 220 A 3. Voltage 24 V 4. Shielding welding gases 82% Ar/18% CO2 5. Kind of tested micro-jetcooling gases Ar, 82% Ar/18% CO2; 98% Ar/2% O2; 98% Ar/2% N2 6. Gas pressure 0.4 MPa; 0.5 MPa 7. Number of micro-jets 1 8. Cooling stream diameter 40 μm; 50 μm Table 3: Chemical composition of WMD Tabela 3: Kemijska sestava WDM Comment Element Amount in all tested cases C 0.08% in all tested cases Mn 0.79% in all tested cases Si 0.39% in all tested cases P 0.017% in all tested cases S 0.018% 3 RESULTS AND DISCUSSION We tested and compared various welds of the stan- dard MAG process connected with those of the innova- tive micro-jet cooling. A typical weld metal deposit had a similar chemical composition in all the tested cases. The micro-jet gas could only have an influence on more or less intensive cooling conditions, but it does not have any influence on the chemical WMD composition (Table 3), except for the oxygen and nitrogen amounts in the WMD (Table 4). It is easy to deduce that the amount of oxygen and nitrogen was slightly increased in terms of the chemical composition of the micro-jet gas mixtures. After the che- mical analyses the metallographic structure of the WMD T. WEGRZYN et al.: IMPACT TOUGHNESS OF WMD AFTER MAG WELDING WITH MICRO-JET COOLING 1002 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1001–1004 Figure 2: Montage of welding head and micro-jet injector (on the right) Slika 2: Namestitev varilne glave in mikro-jet injector (na desni) (with and without micro-jet cooling) was carried out. An example of this structure was shown in Table 5. Table 4: Content of oxygen and nitrogen in WMD Tabela 4: Vsebnost kisika in du{ika v WMD Micro-jet gases Element Amount (%) Ar O 0.0350 82% Ar / 18% CO2 O 0.0380 98% Ar / 2% O2 O 0.0380 98% Ar / 2% N2 O 0.0350 Ar N 0.0055 82% Ar / 18% CO2 N 0.0055 98% Ar / 2% O2 N 0.0055 98% Ar / 2% N2 N 0.0060 Table 5: Metallographic structure of (MAG method 82% Ar/18% CO2) welds Tabela 5: Metalografska struktura zvarov (metoda MAG z 82 % Ar/18 % CO2) Micro-jet gas Gas pressure, MPa Cooling steam diameter, μm Ferrite AF MAC phases without micro-jet - - 55% 3% Ar 0.4 40 60% 2% Ar 0.4 50 63% 2% Ar 0.5 40 63% 2% Ar 0.5 50 61% 2% 98% Ar/2% O2 0.4 40 64% 2% 98% Ar/2% O2 0.4 50 66% 2% 98% Ar/2% O2 0.5 40 67% 2% 98% Ar/2% O2 0.5 50 65% 2% 82% Ar/18% CO2 0.4 40 58% 2% 82% Ar/18% CO2 0.4 50 60% 2% 82% Ar/18% CO2 0.5 40 61% 2% 82% Ar/18% CO2 0.5 50 59% 3% 98% Ar/2% N2 0.4 40 58% 2% 98% Ar/2% N2 0.4 50 59% 2% 98% Ar/2% N2 0.5 40 59% 2% 98% Ar/2% N2 0.5 50 57% 3% Table 5 shows that in all cases a gas mixture of argon with oxygen is the most beneficial choice. We also ob- served MAC (self-tempered martensite, retained auste- nite, carbide) phases on various levels. In the standard MAG welding process (without micro-jet cooling) there are usually gettable larger amounts of grain-boundary ferrite (GBF) and site-plate ferrite (SPF) fraction, mean- while in micro-jet cooling WMD both of the GBF and SPF structures were not so dominant. Ferrite with a percentage above 60 % was gettable only in one case after MAG welding with micro-jet gas mixtures: argon/oxygen or argon/carbon dioxide (Figure 3). The larger amount of MAC phases was especially gettable for the more intensive micro-jet cooling with a gas mixture of argon-oxygen (Tables 5 and 6). The heat-transfer coefficient of the various micro-jet gas mixtures influences the cooling conditions of the welds (and consequently the rise in the content of the MAC phases). This is due to the conductivity coefficients ( ·105), as shown in Table 6. Table 6: Heat-transfer coefficient of various gases used in micro-jet cooling Tabela 6: Koeficient prenosa toplote razli~nih plinov, uporabljenih pri mikro-jet hlajenju Gas Conductivity coefficients,mW/mK Ar 17.9 CO2 16.8 O2 26.3 N2 26 He 156.7 Table 7: Metallographic structure of MAG (82 % Ar/18 % CO2) welds Tabela 7: Metalografska struktura MAG zvarov (82 % Ar/18 % CO2) Welding method Micro-jetgas Test temperature, °C Impact toughness KCV, J MAG without -40 below 40 MAG with micro-jet cooling Ar -40 55 MAG with micro-jet cooling 82% Ar/ 18% CO2 -40 53 MAG with micro-jet cooling 98% Ar/ 2% O2 -40 57 MAG with micro-jet cooling 98% Ar/ 2% N2 -40 below 40 MAG without +20 177 MAG with micro-jet cooling Ar +20 191 MAG with micro-jet cooling 82% Ar/ 18% CO2 +20 189 MAG with micro-jet cooling 98% Ar/ 2% O2 +20 194 MAG with micro-jet cooling 98% Ar/ 2% N2 +20 183 Analysing Table 6, it is possible to deduce that helium could give the strongest cooling conditions, but helium was not tested in this investigation. The cooling T. WEGRZYN et al.: IMPACT TOUGHNESS OF WMD AFTER MAG WELDING WITH MICRO-JET COOLING Materiali in tehnologije / Materials and technology 50 (2016) 6, 1001–1004 1003 Figure 3: Microstructure weld metal with large amount of acicular ferrite in weld (67 %) after Ar/CO2 micro-jet cooling Slika 3: Mikrostruktura zvara z velikim dele`em igli~astega ferita v zvaru (67 %), po mikro-jet hlajenju z Ar/CO2 conditions after welding using other micro-jet gases are on a similar level. After the microscope tests, the Charpy V impact toughness of the deposited metal was assessed with 5 specimens. The Charpy tests were carried out at temperatures of –40 °C and +20 °C only. The impact toughness results are given in Table 7. It is easy to deduce that the impact toughness, espe- cially at the negative temperature of the weld metal deposit, is apparently affected by the kind of micro-jet cooling gas mixtures. Micro-jet technology always strongly improves the impact toughness of the WMD. Argon with oxygen and argon with carbon dioxide must be treated as good choices. Argon as the main element of the gas mixture with a small amount of oxygen gives better results than gas mixtures of argon with carbon dioxide and argon with nitrogen. Nevertheless, micro-jet cooling with gas mixture of argon with 2 % of nitrogen gives better results than the simple MAG welding without micro-jet cooling. This can be explained by the presence of nitride inclusions in the weld (for instance TiN) that facilitate the nucleation of ferrite AF. 4 CONCLUSIONS In low-alloy steel welding there are two general types of tests performed: impact toughness and structure. Aci- cular ferrite and MAC phases (self-tempered martensite, upper and lower bainite, retained austenite, carbides) were analysed and counted for each weld metal deposit. These two tests (microstructure and impact toughness) proved that micro-jet technology gives a beneficial modi- fication to the mechanical properties of the welds. The innovative micro-jet technology was firstly recognized with great success for MIG welding only with argon as a micro-jet gas. In this paper micro-jet cooling technology was for the first time described and tested for MAG welding process with various micro-jet gas mixtures of argon. Final conclusions: • micro-jet cooling could be treated as an important element of MAG welding process, • micro-jet cooling after welding can improve the amount of ferrite AF, the most beneficial phase in low-alloy steel WMD, • gas mixture of argon with carbon dioxide and gas mixture of argon with oxygen could be treated as better micro-jet cooling media than gas mixture of argon with nitrogen, • micro-jet cooling after welding can seriously improve the impact toughness of low-alloy steel WMD, • micro-jet cooling after welding practically does not have an influence on the MAC amount in low-alloy steel WMD. 5 REFERENCES 1 T. Wêgrzyn, Gas mixtures for welding with micro-jet cooling, Arch. Metall. 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Kowalska, Analysis of deformation texture in AISI 304 steel sheets, Sol. St. Phenom., 203-204 (2013), 105–110, doi: 10.4028/www.scientific.net/SSP. 203-204.105 17 G. Golañski, J. S³ania, Effect of different heat treatments on micro- structure and mechanical properties of the martensitic GX12CrMoVNbN91 cast steel, Arch. Metall. Mater., 58 (2012) 1, 25–30, doi: 10.2478/v10172-012-0145-x 18 T. Wêgrzyn, J. Piwnik, D. Hadryœ, R. Wiesza³a, Car body welding with micro-jet cooling, J.Arch. Mater. Sci. Eng., 49 (2011), 90–94 T. WEGRZYN et al.: IMPACT TOUGHNESS OF WMD AFTER MAG WELDING WITH MICRO-JET COOLING 1004 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1001–1004 O. ÇAVUªOÐLU et al.: FORMING-LIMIT DIAGRAMS AND STRAIN-RATE-DEPENDENT MECHANICAL PROPERTIES ... 1005–1010 FORMING-LIMIT DIAGRAMS AND STRAIN-RATE-DEPENDENT MECHANICAL PROPERTIES OF AA6019-T4 AND AA6061-T4 ALUMINIUM SHEET MATERIALS MEJNI DIAGRAMI PREOBLIKOVANJA IN ODVISNOST MEHANSKIH LASTNOSTI OD HITROSTI PREOBLIKOVANJA ALUMINIJEVIH PLO^EVIN IZ AA6019-T4 IN AA6061-T4 Onur Çavuºoðlu1,2, Alan Gordon Leacock2, Hakan Gürün1 1Gazi University, Faculty of Technology, Department of Manufacturing Engineering, Ankara, Turkey 2University of Ulster, Advanced Metal Forming Research Group (AMFOR), Newtonabbey, Northern Ireland onur.cavusoglu@gazi.edu.tr Prejem rokopisa – received: 2015-08-17; sprejem za objavo – accepted for publication: 2015-27-10 doi:10.17222/mit.2015.259 The mechanical properties and formability behaviour of sheet materials depend on the deformation conditions. In this study, the variance of AA6019-T4 and AA6061-T6 aluminium sheet materials related to the strain rate of their mechanical properties was studied by applying uniaxial tensile tests on these materials at different semi-static strain rates (0.3 s–1, 0.03 s–1, 0.003 s–1, 0.0003 s–1, 0.00003 s–1). In addition, forming-limit diagrams of these materials were determined by applying Nakajima tests. When the results were analysed, it was found that the strain rate developed some mechanical properties in the AA6019-T4 and AA6061-T4 sheet materials and that the AA6061-T4 sheet material has a higher formability capability in comparison with the AA6019-T4 sheet material. Keywords: FLD, strain rate, aluminium alloy 6061, aluminium alloy 6019 Mehanske lastnosti in obna{anje pri preoblikovanju plo~evine sta odvisna od pogojev deformacije. V {tudiji je bilo prou~evano spreminjanje AA6019-T4 in AA6061-T6 aluminijeve plo~evine glede na hitrost preoblikovanja in prou~evane so bile njihove mehanske lastnosti z uporabo enoosnih nateznih preizkusov teh materialov pri razli~nih semi-stati~nih hitrostih preoblikovanja (0.3 s–1, 0.03 s–1, 0.003 s–1, 0.0003 s–1, 0.00003 s–1). Dodatno so bili dolo~eni {e diagrami mejnega preoblikovanja teh materialov z uporabo Nakajima preizkusov. Analiza dobljenih rezultatov je pokazala, da hitrost preoblikovanja vpliva na mehanske lastnosti AA6019-T4 in AA6061-T4 plo~evin in da ima AA6061-T4 plo~evina ve~jo sposobnost preoblikovanja v primerjavi s plo~evino AA6019-T4. Klju~ne besede: FLD, hitrost deformacije, aluminijeva zlitina 6061, aluminijeva zlitina 6019 1 INTRODUCTION Great importance is attached to weight-reduction technology at the present time. For this reason, studies on the reduction of the amount of vehicle fuel consump- tion and CO2 emissions have been analysed. It has been observed in the studies conducted for this purpose that improved high-strength steels and light-metal materials such as aluminium and magnesium alloys have an im- portant place.1,2 When the usage applications of alumi- nium alloys are analysed, it is clear that they have been commonly preferred in the automotive and aerospace industries on account of their superior characteristics, such as their low density, high strength, formability, corrosion resistance and high availability as a source.2,3 Tensile testing is a method commonly used for deter- mining a plurality of mechanical properties and the deformation behaviour of the materials. Many parame- ters obtained from the tensile test may vary, based on the material’s deformation conditions, such as friction, deformation speed and temperature.4 It is known that the strain rate changes many metallic materials’ properties by affecting the relation between the tensile and defor- mation.3 Studies about the effects of the strain rate on the mechanical properties of aluminium alloys are available in the literature. J. Q. Tan et al.5 analysed the tensile behaviour at high strain rates by using 7050-T7451 aluminium alloy, and they determined that an increasing deformation rate increases the amount of tensile. M. Vural and J. Caro6 observed, by analysing the tensile behaviour of 2139-T8 aluminium alloy at different tem- peratures and deformation rates, that a significant change in the behaviour of materials at low strain rates does not occur and that the yield stress at high deformation rates increases. By applying tensile tests on the AA5754 and AA5182 aluminium alloys, Smerd and others have found that when high deformation rates subtract from the semi-static strain rates of the AA5754 aluminium alloy, there is no change among the high strain rates, while a significant increase in the yield stress occurs, and they also found that the AA5182 alloy is not completely affected by changes in the strain rates. However, the increasing strain rate increases the amount of stretching in both alloys.7 In their study O. G. Lademo et al.8 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1005–1010 1005 UDK 67.017:620.172.2:669.715 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 50(6)1005(2016) determined that an increase in the tensile strength and elongation occurs by analysing the strain rate sensitivity of the AA1200 and AA3103 aluminium alloys, although the yield strength is not affected, depending on the increase in the deformation rate. In their study, D. Li and A. Ghosh9 determined the mechanical properties of aluminium alloys at different temperatures and strain rates. Forming-limit diagrams (FLDs) must be determined in order to assess the behaviour of the sheet’s formability during the sheet material’s characterization. The Naka- jima test is widely preferred for the determination of the forming-limit diagram. Deforming the sheet with a hemi- spherical punch until the sheet material, for different geometries, starts to neck in the Nakajima test and measuring changes to the shape occurring in a predeter- mined grid on the sheet material can be obtained using forming-limit diagrams. In this process, deformations of uniaxial plane strain and biaxial stretching deformations occur in the sheet material.10,11 The parameters of the deformation, such as the tem- perature and strain rate, affecting the mechanical proper- ties of materials affect the formability of sheet material and therefore also affect the forming-limit diagrams. In their studies, C. Zang et al.12,13 determined, by analysing the FLDs of AA5086 and AA5083 aluminium alloys at elevated temperatures and at different strain rates, that the ability of formability decreases due to the increase in the strain rate. T. Naka et al.14 in their study looked at the 5083 magnesium-aluminium alloy at different tempe- ratures and strain rates and reported that the strain rates decrease the formability at high temperatures, but a significant effect is not observed at room temperature. Considering the studies in the literature; it has been seen in the majority of aluminium alloys that an increase in the strain rate improves some of the mechanical pro- perties, whereas analysing its effects on the forming- limit diagrams and the strain rate at room temperature is understood not to have a significant effect.4–8,12–14 In this study, it is intended to bring in the literature AA6019-T4 and AA6061-T4 aluminium alloy sheet materials whose mechanical properties have been deter- mined, and whose forming-limit diagram has been ob- tained. 2 MATERIAL AND EXPERIMENTAL METHODS 2.1 Material AA6061-T4 and AA6019-T4 sheet metal materials that are Al-Mg-Si-based are used in the study conducted. The chemical composition of these materials is given in Table 1. Also, microstructure photographs taken during the rolling, transverse and diagonal directions of the sheet materials are shown in Figures 1 and 2. The spe- cimens for the optical investigations were etched using a mixture of acetic acid (7 mL), picric acid (25 g), ethanol (140 mL), and purified water (40 mL) for 15 s to reveal the microstructure. Tensile test samples were prepared according to ASTM E517 standards in order to determine the mecha- nical properties of the material depending on the strain rate. The standard size is shown for the tensile test in Figure 3. It is carried out according to the Nakajima test method for forming-boundary limits. Thus, test samples that have been prepared according to the ASTM E 2218-02 standard are used. The Nakajima test-sample dimensions are shown in Table 2. The cutting process was performed in a water-jet machine in order to mini- mize the thermal impacts that may occur on the sheet material during the preparation of the test samples. Also, the notch effect, which may occur during deformation, O. ÇAVUªOÐLU et al.: FORMING-LIMIT DIAGRAMS AND STRAIN-RATE-DEPENDENT MECHANICAL PROPERTIES ... 1006 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1005–1010 Figure 1: Microstructure photos of AA6061-T4 Slika 1: Posnetki mikrostrukture AA6061-T4 Table 1: Chemical composition of AA6019-T4 and AA6061-T4 sheet metal material, in mass fractions (w/%) Tabela 1: Kemijska sestava AA6019-T4 in AA6061-T4 plo~evin, v masnih odstotkih (w/%) Material Mg Si Fe Cu Cr Mn Ti Zn Al AA 6061-T4 0.8 0.4 0.5 0.15 0.15 <0.15 <0.15 <0.25 balance AA 6019-T4 1.2 0.1 <0.5 0.6 0.35 <0.1 <0.15 0.1 balance was eliminated by polishing the side surface of the test sample. Table 2: Nakajima test sample sizes Tabela 2: Velikost vzorcev za Nakajima preizkus Nakajima specimen Sample No. R (mm) 1 0 2 20 3 40 4 50 5 57.5 6 65 7 72.5 8 80 2.2 Experimental methods 2.2.1 Tensile test Tensile test samples of the AA6019-T4 and AA6061-T4 metal material sheet were prepared according to ASTM E517 standards. The tensile tests were carried out at room temperature using a mechanical deformation meter in an Instron 5500 tensile test device. Tensile tests were carried out for three different directions (RD, TD, DD) and five different strain rates (0.3 s–1, 0.03 s–1, 0.003 s–1, 0.0003 s–1, 0.00003 s–1). The average values were taken by repeating each test three times for each sheet material and each different defor- mation rate in order to reduce the margin of error. As a result of these tests, the yield strength, tensile strength, strain-hardening coefficient, total elongation and aniso- tropy values were obtained, depending on the strain rates from the the true stress vs. true strain curves of the sheet material. 2.2.2 Forming-limit diagrams (FLDs) In the second part of the study, the forming-limit diagrams were obtained using the Nakajima test method in order to determine the formability capabilities of AA6019-T4 and AA6061-T4 sheet metal. A schematic view of the Nakajima test method is presented in Figure 4. A 2.4-mm-diameter circular gridding process using an electrochemical etching method was performed on the upper surfaces of the samples for these tests. Then, the hydraulic bulging process was applied on the test samples with a hemispherical punch until the fracture on O. ÇAVUªOÐLU et al.: FORMING-LIMIT DIAGRAMS AND STRAIN-RATE-DEPENDENT MECHANICAL PROPERTIES ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 1005–1010 1007 Figure 2: Microstructure photos of AA6019-T4 Slika 2: Posnetki mikrostrukture AA6019-T4 Figure 3: ASTM E517 tensile test specimen dimensions Slika 3: Dimenzije nateznega preizku{anca po ASTM E517 Figure 5: The measurement of the grid Slika 5: Merjenje mre`e Figure 4: Nakajima test method’s schematic display of forming limit diagram Slika 4: Shematski prikaz Nakayima preizkusa za mejni diagram preoblikovanja the sheet material occurred. The measurement process of the samples proceeded in order to determine the plastic deformation occurring in the grid. The measurement processes was carried out by means of the measuring system at the University of Ulster, Northern Ireland. A photograph of the performed measurement process is presented in Figure 5. The grids in which the deforma- tion occurs on the sheet material were computerized by measuring with the help of a camera and a 3.0 GPA program. After this, the forming-limit diagrams were determined by processing the data. 3 RESULTS AND DISCUSSION 3.1 Tensile test results The variance of the sheet materials’ mechanical properties depending on the strain rate is determined by applying the tensile tests at (0.3 s–1, 0.03 s–1, 0.003 s–1, 0.0003 s–1, 0.00003 s–1) strain rates. The data obtained from the tensile-test results are given in Tables 3 and 4. Table 3: AA6019-T4 tensile test results Tabela 3: Rezultati nateznega preizkusa AA6019-T4 Strain rate Yield strength Tensile strength Hardening coefficient Elon- gation R value 0.3 224 399 0.23 0.15 0.291 0.03 221 397 0.24 0.17 0.316 0.003 219 396 0.23 0.17 0.384 0.0003 220 397 0.24 0.17 0.446 0.00003 219 390 0.22 0.18 0.469 Table 4: AA6061-T4 tensile test results Tabela 4: Rezultati nateznega preizkusa AA6061-T4 Strain rate Yield strength Tensile strength Hardening coefficient Elon- gation R value 0.3 165 325 0.25 0.18 0,46 0.03 165 322 0.24 0.18 0.468 0.003 162 324 0.25 0.19 0.479 0.0003 164 325 0.24 0.20 0.555 0.00003 155 310 0.25 0.21 0.649 The relationship between the strain rate of the AA6019 -T4 and AA6061-T4 sheet materials and the yield strength is given in Figure 6. An increase in the yield strength of the sheet materials occurs with an increased strain rate.5–7 Upon analysing Figure 6, it is seen that the yield strength of both sheet materials shows an upward tendency with the strain rate. By considering the amounts of increase in the yield strength, the AA6019-T4 sheet material’s yield strength was deter- mined to be between 220 MPa and 225 MPa at the lowest (0.00003 s–1) and the highest (0.3 s–1) strain rates. It is determined that the AA6061-T4 sheet material has the lowest yield strength at the lowest (0.00003 s–1) strain-rate value. At the other strain rates it is clear that the detected values of the yield strength were between 163 MPa and 166 MPa. According to these results, it is observed that the yield strength of AA6019-T4 and AA6061-T4 sheet materials was not significantly affected by the increase in the strain rate. Tensile strength is one of the most important parameters in the classification of sheet materials. The strain rate vs. tensile strength relationship is shown in Figure 7. Analysing the effects of the strain rate on the tensile strength, it is determined that both sheet materials have the lowest tensile strength at the lowest deformation rate. The tensile strength also increased with an increase in the strain rate. However, very small amounts of change in the tensile strength were seen at the other strain rates. In studies in the literature, it is determined that an increase occurred in the tensile strength with an increasing strain rate, but this increase was very small at room temperature.5,8 Also in this study, when the tensile strength obtained at the lowest deformation rate for the AA6019-T4 and AA6061-T4 sheet materials used in the experiments was compared with the other strain rates in the tensile strength values, it is clear that the tensile strength of both sheet materials increased by a small amount. O. ÇAVUªOÐLU et al.: FORMING-LIMIT DIAGRAMS AND STRAIN-RATE-DEPENDENT MECHANICAL PROPERTIES ... 1008 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1005–1010 Figure 7: Strain rate vs. tensile strength Slika 7: Odvisnost med hitrostjo preoblikovanja in natezno trdnostjo Figure 6: Strain rate vs. yield strength Slika 6: Odvisnost med hitrostjo preoblikovanja in mejo plasti~nosti As a result of the performed tensile tests, the strain rate vs. total elongation graphs of the AA6019-T4 and AA6061-T4 sheet materials are given in Figure 8. When the graphics are analysed, the maximum elongation of the AA6019-T4 sheet material occurred at the lowest strain rate of the total elongation. For the other defor- mation rates, it is seen to be unaffected by the changes at the strain rates. When the AA6061-T4 sheet material’s elongation behaviour was analysed, it is clear that the increasing strain rate and total elongation have a down- ward tendency. Upon comparing the elongation beha- viour of the two materials, it is seen that the AA6019-T4 sheet material has elongated less than the AA6061-T4 sheet material. Upon analysing sheet materials in the same series, it is seen that sheet materials having a high tensile strength elongated less. The materials gain strength through strain hardening during the plastic deformation of the sheet-metal ma- terial. Upon analysing graphs of the strain rate vs. strain hardening coefficient in Figure 9, it is clear that strain- hardening coefficient was in the range 0.22–0.24 for the AA6019-T4 sheet material, and 0.24-0.25 for the AA6061-T4 sheet material. Also, it is clear that the strain-hardening coefficient of both sheet materials was not significantly affected by the variance in the deforma- tion rate. The strain-hardening coefficient results in the study of D. Li and A. Ghosh9 shows that it has changed over a very small range. Therefore, the results obtained were found to be consistent with the literature. When analysing the average planer anisotropy values in Figure 10, it is seen that an amount of decrease has occurred in the plane anisotropy and increasing deformation speed. However, when evaluating the plane anisotropy coefficients, the values were found to be close to each other. It was concluded that the plane anisotropy coefficient in both sheet materials shows a slight down- ward tendency. 3.2 Forming-limit diagram Upon analysing the results obtained from tensile tests, both experimental sheet materials used in this study were determined not to have been significantly affected by the variance in the strain rate. Therefore, the forming-limit diagrams of the AA6019-T4 and AA6061-T4 sheet materials were determined by Nakajima Tests by considering the strain rate 0.003 s–1 (10 mm/s). Forming-limit diagrams are shown in Fig- ure 11. The area under the forming-limit diagrams represents a safety forming area. The upper parts of the curve show the areas occurring for the fracture in the sheet material.11–13 Upon evaluating the forming-limit diagram in Figure 11, it is clear that the AA6061-T4 sheet ma- terial has higher deformation limits than the AA6019-T4 sheet material, by showing higher elongation behaviour than the latter. Thus, it provides that the forming-limit curve obtained for the AA6061-T4 sheet material was above the AA6019-T4 sheet metal’s forming-limit curve. This case shows that the AA6061-T4 sheet material had better formability. And this case constitutes solid-solu- O. ÇAVUªOÐLU et al.: FORMING-LIMIT DIAGRAMS AND STRAIN-RATE-DEPENDENT MECHANICAL PROPERTIES ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 1005–1010 1009 Figure 10: Strain rate vs. anisotropy Slika 10: Odvisnost med hitrostjo preoblikovanja in anizotropijo Figure 8: Strain rate vs. elongation Slika 8: Odvisnost med hitrostjo deformacije in raztezkom Figure 9: Strain rate vs. hardening coefficient Slika 9: Odvisnost med hitrostjo preoblikovanja in koeficientom utrjevanja tion strengthening, owing to the fact that the rate of alloying elements in the chemical composition of the AA6019-T4 sheet material was higher than the AA6061-T4 sheet material. This also ensures that it has a higher yield and tensile strength and causes the ability of formability to decrease. Furthermore, the AA6061-T4 sheet material had greater elongation and anisotropy (R) values than the AA6019-T4 sheet material, which is another sign that the former has a higher formability than the latter. 4 CONCLUSIONS The results obtained from this study are summarized below. The lowest values for both the yield and tensile strength of the sheet material are obtained at the lowest deformation rate (0.00003 s–1). It is found that as long as the strain rate increases, the yield and tensile strengths show a slight increase. A very small range of change in the strain-hardening coefficient is observed with an increase in the strain rate. A downward tendency is seen in both sheet materials’ elongation, the plane anisotropy upon the increase of the strain rate. However, the exchange range of the elongation and the plane anisotropy are found to be too narrow. Therefore, it can be said that the AA6061-T4 and AA6019-T4 sheet materials are insensitive to the strain rate. AA6061-T4 aluminium alloy is found to have higher formability limits than the AA6019-T4 aluminium alloy. Acknowledgement The authors would like to thank Turkish Council of Higher Education for the research grant and University of Ulster, Advanced Metal Forming Research Group (AMFOR) for the necessary equipment and material support. 5 REFERENCES 1 O. Çavuºoðlu, H. Gürün, Investigation of the effects of deformation speed on the mechanical properties and deep drawing process of DP600 and DP780 sheet metal, Journal of the Faculty of Engineering and Architecture of Gazi University, 29 (2014) 4, 777–784, doi:10.17341/gummfd.76140 2 N. Wang, Z. Zhou, G. Lu, Microstructural evolution of 6061 alloy during isothermal heat treatment, J. Mater. Sci. Technol., 27 (2011) 1, 8–14, doi:10.1016/S1005-0302(11)60018-2 3 B. M. Dariani, H. G. Liaghat, M. Gerdooei, Experimental investiga- tion of sheet metal formability under various strain rates, Proc. IMechE Part B: J. Engineering Manufacture, 223 (2009), 703–712, doi:10.1243/09544054JEM1430 4 C. Kubat, A. Kiraz, The modeling of tensile test in virtual laboratory design using artificial intelligence Journal of the Faculty of Engi- neering and Architecture of Gazi University, 27 (2012) 1, 205–209, doi:10.17341/gummfd.04550 5 J. Q. Tan, M. Zhan, S. Liu, T. Huang, J. Guo, H. Yang, A modified Johnson–Cook model for tensile flow behaviors of 7050-T7451 aluminum alloy at high strain rates, Materials Science & Engineering A, 631 (2015), 214–219, doi:10.1016/j.msea.2015.02.010 6 M. Vural, J. Caro, Experimental analysis and constitutive modeling for the newly developed 2139-T8 alloy, Materials Science and Engineering A, 520 (2009), 56–65, doi:10.1016/j.msea.2009.05.026 7 R. Smerd, S. Winklera, C. Salisburya, M. Worswicka, D. Lloydb, M. Finn, High strain rate tensile testing of automotive aluminum alloy sheet, International Journal of Impact Engineering, 32 (2005), 541–560, doi:0.1016/j.ijimpeng.2005.04.013 8 O. G. Lademo, O. Engler, J. Aegerter, T. Berstad, A. Benallal, O. S. Hopperstad, Strain-rate sensitivity of aluminum alloys AA1200 and AA3103, J. Eng. Mater. Technol., 132 (2010) 4, 041007–8, doi:10.1115/1.4002160 9 D. Li, A. Ghosh, Tensile deformation behavior of aluminum alloys at warm forming temperatures, Materials Science and Engineering A, 352 (2003), 279–286, doi:10.1016/S0921-5093(02)00915-2 10 Material Properties: Determination of Process Limitations in Sheet Metal Forming - Forming Limit Diagram, gom Optical Measuring Techniques, http://www.gom.com/fileadmin/user_upload/industries/ flc_fld_EN.pdf, 18.06.2015 11 O. Anket, T. Koruvatan, Ý. Ay, The use of forming limit diagrams in forming sheet metal materials, Journal of Polytechnic, 14 (2011) 1, 39–47 12 C. Zhang, X. Chuc, D. Guinesd, L. Leotoing, J. Dinga, G. Zhao, Effects of temperature and strain rate on the forming limit curves of AA5086 sheet, Procedia Engineering, 81 (2014), 772–778, doi:10.1016/j.proeng.2014.10.075 13 C. Zhang, L. Leotoing, D. Guines, E. Ragneau, Theoretical and numerical study of strain rate influence on AA5083 formability, Journal of Materials Processing Technology, 209 (2009), 3849–3858, doi:10.1016/j.jmatprotec.2008.09.003 14 T. Naka, G. Torikai, R. Hino, F. Yoshida, The effects of temperature and forming speed on the forming limit diagram for type 5083 alu- minum–magnesium alloy sheet, Journal of Materials Processing Technology, 113 (2001), 648–653, doi:10.1016/S0924-0136(01) 00650-1 O. ÇAVUªOÐLU et al.: FORMING-LIMIT DIAGRAMS AND STRAIN-RATE-DEPENDENT MECHANICAL PROPERTIES ... 1010 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1005–1010 Figure 11: Forming-limit diagrams Slika 11: Diagrami meje preoblikovanja Y. ALTUNPAK et al.: EFFECT OF ALTERNATIVE HEAT-TREATMENT PARAMETERS ON THE AGING BEHAVIOR ... 1011–1016 EFFECT OF ALTERNATIVE HEAT-TREATMENT PARAMETERS ON THE AGING BEHAVIOR OF SHORT-FIBER-REINFORCED 2124 Al COMPOSITES VPLIV ALTERNATIVNIH PARAMETROV TOPLOTNE OBDELAVE NA STARANJE 2124 Al KOMPOZITA, OJA^ANEGA S KRATKIMI VLAKNI Yahya Altunpak1, Serdar Aslan2, Mehmet Oðuz Güler2, Hatem Akbulut2 1Abant Izzet Baysal University, Faculty of Engineering and Architecture, Mechanical Engineering, 14280 Bolu, Turkey 2Sakarya University, Faculty of Engineering, Department of Metallurgy and Materials Engineering, Esentepe Campus, 54187 Adapazari, Turkey altunpak_y@ibu.edu.tr Prejem rokopisa – received: 2015-09-10; sprejem za objavo – accepted for publication: 2015-11-24 doi:10.17222/mit.2015.287 The 2124 Al alloy and a composite of the 2124 Al alloy reinforced with 20 % of volume fractions of -Al2O3 short fibers made by squeeze casting were subjected to controlled and systematic aging treatments. The materials were solution treated at (495, 525 and 555) °C. After quenching, the matrix alloy and the composite were artificially aged at (160, 170, 180 and 190) °C up to 36 h. The aging was monitored with hardness measurements and differential scanning calorimetry. The time required to reach the peak hardness of the composite matrix during a precipitation treatment was shorter than that for the unreinforced 2124 Al. An increase in the solution-treatment temperature resulted in an increase of the composite-matrix hardness. The -Al2O3-reinforced composite exhibits no grain-boundary melting, but appears to show incipient melting around short alumina fiber interfaces at temperatures above 525 °C. The highest HV value was obtained after solutionizing at 495 °C for 6 h, followed by water quenching and aging at 190 °C for 10 h for the unreinforced matrix alloy. In the case of the reinforced alloy the highest HV value was found after solutionizing at 555 °C for 6 h, quenching and aging at 170 °C for 12 h. Keywords: aluminum matrix composite, alumina, solution temperature, aging kinetics Al zlitina 2124 in kompozit Al zlitine 2124, oja~ane z 20 % volumenskega dele`a kratkih vlaken -Al2O3 , ulitih z iztiskanjem, so bile kontrolirano in sistemati~no starane. Materiali so bili raztopno `arjeni na (495, 525 in 555) °C. Po hitrem ohlajanju sta bili osnovna zlitina in kompozit umetno starani 36 h na (160, 170, 180 in 190) °C. Staranje je bilo kontrolirano z merjenjem trdote in z diferen~no vrsti~no kalorimetrijo. ^as za doseganje najvi{je trdote kompozitnega materiala med postopkom izlo~anja je bil kraj{i kot pri osnovnem materialu Al 2124. Povi{anje temperature raztopnega `arjenja se je odrazilo na pove~anju trdote kompozita. Kompozit, oja~an z -Al2O3, ne ka`e nataljevanja po mejah zrn, vendar pa ka`e zametke taljenja na stiku s kratkimi vlakni, pri temperaturah nad 525 °C. Najvi{ja vrednost HV neoja~ane zlitine je bila dobljena po 6 urnem raztopnem `arjenju na 495 °C, ki mu je sledilo hlajenje v vodi in 10 urno staranje na 190 °C. V primeru kompozitne zlitine je bila zabele`ena najvi{ja HV vrednost po raztopnem `arjenju 6 h na 555 °C, ohlajanju v vodi in 12 urnem staranju na 170 °C. Klju~ne besede: kompozit na osnovi aluminija, aluminijev oksid, temperatura raztapljanja, kinetika staranja 1 INTRODUCTION Research efforts on aluminum alloys are focused on precipitation phenomena in which the precipitates formed from a supersaturated solution are responsible for the hardening by natural or artificial ageing in an alloy. The response of an aluminum-matrix composite to aging can be completely different from that of the unreinforced alloy.1–8 Hence, the age-hardening behavior of particulate-reinforced aluminum composites has been the subject of great interest from the scientific and technological viewpoints. The nature of the change in the hardening kinetics during the aging of composites depends on the matrix material, the type of reinforce- ment including their size, morphology and volume fraction, composite processing route, solution and aging temperatures.9–12 The pressure applied during solidification in the squeze-casting technique results in excellent feeding during solidification shrinkage. The commercialization of squeeze casting has only been used to fabricate high- integrity engineering components with reinforcement very recently.13–15 Different types of intermetallics were reported in the solidified 2xxx Al alloys. K. C. Chen and C. G. Chao3 found Cu2Mn3Al20 intermetallic particles in the 2024 matrix alloy and its -Al2O3 short-fiber-reinforced com- posites. The same intermetallic phase was also reported by T. Christman and S. Suresh16 in a 2014 Al alloy. On the other hand, C. Badini et al.17 showed the possibility of the formation of Cu2FeAl7 and (CuFeMn)Al6 in the 2618 Al SiC particle-reinforced composite. The same authors detected (CuFeMn)Al6 in the 2024 Al alloy and its composite derived from MnAl6. They pointed out that this precipitate did not grow during aging and was not affected by the solution treatment because of its large size. Aluminum alloys with a copper:magnesium weight ratio of 2:1 and higher are used for manufacturing a Materiali in tehnologije / Materials and technology 50 (2016) 6, 1011–1016 1011 UDK 622'18:621.785.7:669.715 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 50(6)1011(2016) variety of age-hardenable structural alloys. The structural changes that occur during the aging of these alloys have been extensively studied.16–19 According to S. K. Varma et al.18 and A. P. Sanino and H. J. Rack19, the precipi- tation sequence in the pseudo-binary Al-Al2CuMg alloy (Al-3 % of mass fractions of Cu-1.5 % of mass fractions of Mg) can be represented as follows in Equation (1): SSS (supersaturated solid solution)  GPB zones  S’’  S’  S (1) The solutionizing at 495 °C for 2-6 h, subsequent quenching and precipitation at 190 °C for 8–12 h, is defined as the standard heat treatment for the 2124 Al alloy.20 It is an age-hardenable alloy, whose mechanical properties are mainly controlled with the hardening precipitates contained in the material. Accordingly, the present work was undertaken to study the effect of the -Al2O3 short fibers on the aging response of the com- posite matrix. Particular emphasis was given to examine the effect of the solution treatment and aging tempera- ture on the age-hardening kinetics. 2 EXPERIMENTAL PART The 2124 aluminum matrix alloy and the Saffil fiber/2124 aluminum (4.2 Cu, 1.5 Mg, 0.6 Mn, 0.3 Fe, 0.25 Zn, in mass fractions) composite were produced by squeeze casting using 20 % volume fractions of -Al2O3 preforms supplied by I.C.I. The preform cohesion was ensured by the addition of 3–4 % silica binder. The preform was supplied in the form of discs with 100 mm in diameter and 10 mm in thickness. The liquid alumi- num alloy was squeezed into the preform at 800 °C with a 60 MPa hydraulic press to produce the composites. The pressure holding time was 75 s, to eliminate shrinkage during the solidification. Specimens were solution treated at three different temperatures of (495, 525 and 555) °C for 6 h. Thereafter, all the specimens were quenched into the water ice brine (–15 °C). The aging treatment was carried out in an electrical furnace at (160, 170, 180 and 190) °C up to 36 h. Microhardness measurements of the matrix and the composites (between the fibers) were performed using a diamond pyramid indenter and a 25-g mass. At least 6 hardness measurements were carried out for each aging condition to ensure accurate results. A group of specimens from the alloy solution treated 2124 Al matrix and the 2124 + 20 % of volume fractions of -Al2O3 composite were immediately stored in a refri- gerator at –15 °C. Discs (5 mm diameter and 0.3 mm thickness) for DSC measurements were prepared. The differential scanning calorimetry (DSC) analyses of these samples were performed using a Perkin-Elmer DSC 1700 thermal analyzer. All the samples were loaded in a DSC cell at room temperature and equilibrated for a few minutes. The heating rate was 10 K/min from 25 °C to 550 °C. Dry pure nitrogen was purged through the cell at a rate of 55 cm3/min to avoid oxidation. The data for all the DSC runs were recorded in the instrument me- mory. At least two samples of each heat treatment were analyzed. Microstructural observations were performed on me- chanically polished and etched samples of the unrein- forced alloy and composite. Surface-characterization studies were carried out using a Hitachi HHS-2R scann- ing electron microscope (SEM) with energy-dispersive spectroscopy (EDS). 3 RESULTS AND DISCUSSION 3.1 Microstructural aspects At 525 °C and 555 °C the solution treatment both alloys showed surface blistering. However, the amount of blistering was more pronounced on the solution treated and quenched unreinforced alloy. These blisters were caused by a high internal gas pressure. Typical represen- tative microstructures of the unreinforced alloy solution treated at (495, 525 and 555) °C and subsequently quen- ched are shown in Figures 1a to 1c and Figures 2a to 2c shows the light-microscope microstructures of the rein- forced composite. All the polished samples were slightly etched with Kellers’ agent to reveal grain boundaries and the dissolved intermetallics. Light microscope investiga- tions of the samples of unreinforced alloy have revealed that significant incipient melting along the grain boun- daries occurred when the solution heat treatment was at 525 °C and 555 °C (Figures 1b and 1c). The -Al2O3-reinforced composite exhibits no such grain- Y. ALTUNPAK et al.: EFFECT OF ALTERNATIVE HEAT-TREATMENT PARAMETERS ON THE AGING BEHAVIOR ... 1012 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1011–1016 Figure 1: Light micrographs of the unreinforced 2124 Al matrix solution treated at: a) 495 °C, b) 525 °C, and c) 555 °C Slika 1: Mikrostrukture neoja~ane Al zlitine 2124, raztopno `arjene na: 495 °C, b) 525 °C in c) 555 °C boundary melting, but appears to show incipient melting around the short alumina-fiber interfaces (Figures 2b and 2c). There are low-melting-point constituents, which melt first at temperatures above the initial melting point of an alloy producing incipient melting and embrittle- ment. Thus, it appears that whilst these low-melting- point constituents segregate to the grain boundaries in the unreinforced alloy they may segregate to -Al2O3 short fiber interfaces. Figures 1 and 2 show that as the solution treatment temperature is increased the amount of intermetallic appearing on the polished and etched surface decreases. The unreinforced alloy does not exhibit a visible intermetallic at any of the solution tem- peratures studied. The reduction in the intermetallics appearing on the polished surface as the solution temperature increase (Figure 2) is due to the dissolution of intermetallic particles. In order to clarify the un- dissolved phase, an SEM-EDS study was undertaken. Figure 3 shows a typical EDS spectrum performed on the undissolved phase in the composite matrix. From the EDS analysis copper, magnesium, aluminum, manga- nese, and iron peaks were observed on the undissolved phase in the heat-treated, unreinforced and composite matrix. For the EDS analysis of the undissolved phase the average elemental concentrations of these un- dissolved phases and the spectral analyses of the matrix material (2124 Al) were given in Table 1. It is clear that the contents of copper and manganese of the undissolved phase are much higher than those of the matrix material. The undissolved phase in the literature was identified as Cu2Mn3Al20, which was also found in 2024 Al by K. C. Chen and C. G. Chao3 and in 2014 Al by T. Christman and S. Suresh.16 Table 1: Spectral analysis and EDS analysis results, in mass fractions (w/%) Tabela 1: Rezultati spektralne in EDS-analiz, v masnih dele`ih (w/%) Cu Mg Mn Fe Zn Al 2124 Al alloy spectral analyses results 4.2 1.5 0.6 0.3 0.25 Bal- ance EDS analysis results of the undissolved phase in unreinforced matrix 7.55 0.15 9.05 0.04 83.2 EDS analysis results of the undissolved phase in the composite matrix 7.52 0.13 9.20 0.05 83.1 3.2 DSC analysis The DSC scans of the unreinforced 2124 alloy and 2124 Al + 20 % volume fractions of -Al2O3 composite are shown in Figure 4. For comparison purposes, results from the DSC scans of the unreinforced matrix and the composite quenched into ice brine after the solution treatment at 495 °C and 555 °C are presented in Figures 4a and 4b, respectively. The DSC traces of the compo- sites were different from each other. However, the DSC Y. ALTUNPAK et al.: EFFECT OF ALTERNATIVE HEAT-TREATMENT PARAMETERS ON THE AGING BEHAVIOR ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 1011–1016 1013 Figure 3: a) SEM micrograph of the 2124 Al + 20 % volume fractions of -Al2O3 composite and b) an EDS spectrum taken from the undissolved intermetallic phase Slika 3: a) SEM-posnetek Al kompozita 2124 + 20 % volumenskega dele`a -Al2O3 in b) EDS- spekter neraztopljene intermetalne faze Figure 2: Light micrographs of the 2124 Al + 20 % volume fractions of -A12O3 composite solution treated at: a) 495 °C, b) 525 °C, and c) 555 °C Slika 2: Mikrostrukture Al kompozita 2124 +20 % volumenskega dele`a -A12O3 , raztopno `arjenega na: a) 495 °C, b) 525 °C in c) 555 °C traces of the matrix alloy do not exhibit significant differences depending on the solutionizing temperatures. The curve of the unreinforced matrix alloy shows four zones: an exothermic reaction between 40 °C and 135 °C due to the formation of Guiner-Preston zones followed by an endothermic between 135 °C and 245 °C due to the dissolution of the Guiner-Preston zones; an exother- mic reaction between 245 °C and 335 °C due to forma- tion of S’ precipitates and finally the endothermic reac- tion between 335 °C and 485 °C due to the dissolution of these S’ precipitates (Figures 4a and 4b). However, the DSC curves for the composite do not have obvious Guiner-Preston zone formation and dissolution peaks. Increasing the solution temperature to 555 °C resulted in an increase in the amount of S’ precipitate for the com- posite, as shown in Figures 4a and 4b. Additionally, the figure also shows an acceleration of S’ precipitate for- mation in the case of the reinforced alloy. For example, the S’ precipitates are formed at 245 °C in the unrein- forced 2124 Al matrix (Figure 4a), but this temperature is approximately 230 °C for the composite material (Fig- ure 4b). This shows that the short -Al2O3 ceramic phase shifted the S’ precipitate-formation temperature. 3.3 Matrix microhardness Figure 5 shows the microhardness (HV) as a func- tion of aging time at 190 °C for both alloys. For all con- ditions, the unreinforced 2124 Al matrix alloy and the composite reached a peak hardness after aging for 10–12 h and 7–9 h, respectively, at 190 °C. A reduced hardness of the matrix was observed in the case of the reinforced alloy. The times required to attain the peak hardness for both alloys artificially aged between 160 °C and 190 °C are summarized in Table 2. First, the peak of the matrix microhardness values of the composite samples increase with increasing solutionizing temperature, while these values show a slight decrease in the unreinforced 2124 Al matrix. Second, the presence of short alumina fibers in the matrix decreases the time needed to achieve the peak microhardness. This suggests that the addition of 20 % volume fractions of -A12O3 short fibers for the matrix causes considerable acceleration in the aging kinetics of the matrix alloy. Third, a decrease in the aging temperature from 190 °C to 170 °C leads to an increase of the peak microhardness of the composite matrix. From Table 2, solutionizing at 495 °C for 6 h, followed by water quenching and aging at 190 °C for 10 h, seems to be the best aging procedure for the un- reinforced matrix alloy. In the case of the reinforced alloy the highest HV value was formed after solutioniz- ing at 555 °C for 6 h, quenching and aging at 170 °C for 12 h. Y. ALTUNPAK et al.: EFFECT OF ALTERNATIVE HEAT-TREATMENT PARAMETERS ON THE AGING BEHAVIOR ... 1014 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1011–1016 Figure 5: Microhardness variation of 2124 Al alloy and 2124 Al alloy + 20 % volume fractions of -A12O3 composite as a function of aging time at 190 °C after solution treatment at: a) 495 °C and b) 555 °C Slika 5: Spreminjanje mikrotrdote Al zlitine 2124 in Al kompozita 2124 + 20 % volumenskega dele`a -Al2O3, v odvisnosti od ~asa staranja na 190 °C, po raztopnem `arjenju na: a) 495 °C in b) 555 °C Figure 4: Results of DSC for 2124 Al alloy and 2124 Al alloy + 20 % volume fractions of -Al2O3 composite solution treated at: a) 495 °C and b) 555 °C Slika 4: Rezultati DSC Al zlitine 2124 in Al kompozita 2124 + 20 % volumenskega dele`a -Al2O3, po raztopnem `arjenju na: a) 495 °C in b) 555 °C Table 2: Times (h) required to attain peak hardness for 2124 Al and 2124 Al + 20 % volume fractions of -Al2O3 composite matrices at different aging temperatures Tabela 2: ^as (h) potreben za doseganje najvi{je trdote Al zlitine 2124 in kompozita Al zlitine 2124+ 20 % volumenskega dele`a -Al2O3 pri razli~nih temperaturah staranja Solution temp. (°C) Time (h) / Microhardness (HV0.025) Material 160 °C 170 °C 180 °C 190 °C 495 2124 Al 20/132.0 17/130.4 12/133.0 10/133.3 2124 Al composite 12/118.8 12/119.6 10/117.9 8/116.7 525 2124 Al 20/130.6 16/129.6 12/130.7 8/132.5 2124 Al composite 12/124.5 12/125.4 10/122.9 8/121.3 555 2124 Al 20/127.5 16/129.1 10/129.6 8/130.5 2124 Al composite 12/128.4 12/130.4 8/125.8 8/124.5 3.4 Discussion The use of an Al-Cu-Mg alloy for fabricating an alu- minum matrix composite leads to the formation of inter- metallic particles of Cu2Mn3Al20 type. Increasing the solution treatment temperature caused greater dissolution of the Cu-Mn-Al-based intermetallic particles, and this is related to a decrease in the number of particles on the metallographic polished and etched sample surfaces. At 525 °C and 555 °C for the solution-treated unreinforced alloy and the composites, surface blistering was observed after quenching. However, the amount of blistering was observed to be more in the unreinforced alloy when com- pared with the composite. These blisters were suggested to form due to a high internal gas pressure. Above the solution treatment of 495 °C for the unreinforced alloy, grain-boundary melting was detected. For 2124 Al + 20 % volume fractions of -Al2O3 the composite speci- mens intermetallic dissolution was almost complete at 555 °C. Regarding the heat-treated samples, the extensive dissolution of the Cu2Mn3Al20-type intermetallics and the formation of coherent precipitates after aging, indicate an apparently good solution and precipitation treatment that improves the hardness values of the 2124 Al matrix. At solution-treatment temperatures above 495 °C the peak hardness values of the matrix alloy decrease, prob- ably as a result of increasing the amount of gas porosity due to hydrogen absorption from a moist atmosphere. The increment in the peak hardness values of the com- posite was observed by increasing the solution-treatment temperature. The thermal analysis technique was employed to study the change in the enthalpy, which is associated with the formation and dissolution of the precipitates. The area of the peak in the DSC curve gives the reaction enthalpy, which is directly related to the molar heat of reaction and the volume fraction of the precipitating or dissolving phase. The corresponding temperature is related to the stability of the precipitates and to the reac- tion kinetics. The DSC curves show that the GP zone’s formation peak is suppressed in the composite matrix. This inhibition of the GP zone’s formation and its effect on the age hardening are similar to the observation of sintered aluminum powder (SAP) alloys.3 In SAP alloys, these phenomena have been attributed to a lack of quenched-in vacancies that were soaked up by the grain boundaries in the fine-grain matrix and by the Al/A12O3 particle interfaces. The observations in the present work suggest that a similar mechanism is responsible for the inhibition of the GP zone’s formation in aluminum ma- trix composite matrices. From the DSC scans and micro- hardness measurements, it was observed that the S’ formation temperature is lower in the composite than in the unreinforced matrix alloy. In general, the addition of -A12O3 fibers decreases the time required to attain peak hardness. These features encourage the nucleation of precipitates by reinforcing the ceramic phase. The addi- tion of -Al2O3 fibers to the 2124 alloy was also resulted to obtain lower age-hardening temperatures to attain time to require peak aging. It is well known that -A12O3 fibers lead to acceleration in the aging kinetics and some possible reasons for the acceleration of the aging kinetics are given in 21,22. Composite matrix has finer grains than that of the unreinforced alloy. As also evidenced from the transmission electron microscopy studies by C. M. Friend and S. D. Luxton23, the dislocation densities are almost similar in the aluminum matrix alloy and its -Al2O3 short-fiber-reinforced composites. The disloca- tions in the matrix are in the form of long lines and short lines in the composite. However, the short dislocations are agglomerated close to the fiber matrix interfaces. In this study, it is suggested that these dislocations can also accelerate the formation of the S’ phase. In addition, fiber matrix interfaces are well-known heterogeneous nucleation sites for precipitation. In the as-quenched condition, the microhardness of the composite matrix is lower than that of the unreinforced matrix. The result is similar to the 2024 Al--A12O3 composites and 2124 Al-SiC composites that were reported by K. C. Chen and C. G. Chao3 and T. S. Christman and Suresh16, respectively. When the composite is cooled from the elevated temperature of the aging process, misfit strains occur due to differential thermal contraction at the reinforcement/matrix inter- face, which are sufficient to generate dislocations. The dislocation density is able to influence the microhard- ness. B. Dutta and M. K. Surappa11 suggested an en- hanced dislocation density in the 6061 Al-A12O3 composite matrix that contributes to higher microhard- ness values compared to the unreinforced matrix alloy. K. C. Chen and C. G. Chao3 reported that a very low dislocation density was observed in a 2024 Al alloy that was reinforced with -A12O3 short fibers when compared with Al-SiC systems. They pointed out that the coeffi- cient of thermal expansion difference between Al and -A12O3 is 3:1, but this ratio is 10:1 for Al and SiC. Thus, a lower dislocation density is established in the matrix of Al--A12O3 short-fiber composites. Moreover, the porosity level is expected to be higher in the com- posite matrix due to poor wetting of the short fiber by the Y. ALTUNPAK et al.: EFFECT OF ALTERNATIVE HEAT-TREATMENT PARAMETERS ON THE AGING BEHAVIOR ... Materiali in tehnologije / Materials and technology 50 (2016) 6, 1011–1016 1015 matrix during the squeeze-casting process route since the fibers are agglomerated in some regions. Since the artifi- cial age hardening of 2124 Al is attributed to the GP zone and the S’ phase, the microhardness of the compo- sites is smaller due to the suppression of the GP zones. Consequently, the number of S’ precipitates in the com- posite matrix decreases. The increasing number of precipitates are expected to result in high strength in the age-hardenable alloys and their composites. The res- ponse of the -Al2O3 short-fiber-reinforced 2124 Al alloy matrix to aging is significantly different to that of the un- reinforced alloy in the 495 °C solution-treated condition. The grain boundaries and the addition of fiber-matrix interfaces were thought to be the preferred site for pore nucleation for the following reasons; i) the incipient melting that occurs there (incipient melting is particul- arly prone to gas porosity because of the higher solubi- lity in the liquid phase9 and, ii) the observation that when an aluminum matrix composite was charged with hydrogen, the gas migrates to the clusters in addition to the -A12O3 fiber interfaces.24 4 CONCLUSIONS Increasing the solution treatment temperature led to an increase in the dissolution of the intermetallic particles in the composite matrix. At temperatures above 525 °C, samples of the unreinforced alloy experienced significant incipient melting along the grain boundaries. The -Al2O3 reinforced composite exhibits no grain- boundary melting, but appears to show incipient melting around the short alumina fiber interfaces. 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Dollar, A. W. Thompson, Microstructure property rela- tionships and hydrogen effects in a particulate-reinforced aluminum composite, Metall. Trans. A., 22 (1991), 2445–2450, doi:10.1007/ BF02665010 Y. ALTUNPAK et al.: EFFECT OF ALTERNATIVE HEAT-TREATMENT PARAMETERS ON THE AGING BEHAVIOR ... 1016 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1011–1016 LETNO KAZALO – INDEX Letnik / Volume 50 2016 ISSN 1580-2949 © Materiali in tehnologije IMT Ljubljana, Lepi pot 11, 1000 Ljubljana, Slovenija MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY VSEBINA / CONTENTS LETNIK / VOLUME 50, 2016/1, 2, 3, 4, 5, 6 2016/1 Improvement of the casting of special steel with a wide solid-liquid interface Izbolj{anje ulivanja posebnega jekla s {irokim intervalom trdno-teko~e T. Mauder, J. Stetina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 Non-traditional non-destructive testing of the alkali-activated slag mortar during the hardening Netradicionalno neporu{no preizku{anje z alkalijami aktivirane malte med strjevanjem L. Topoláø, P. Rypák, K. Tim~aková-[amárková, L. Pazdera, P. Rovnaník . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 Multi-walled carbon nanotubes effect in polypropylene nanocomposites Vpliv ve~stenskih ogljikovih nanocevk v nanokompozitih iz polipropilena C. E. Ban, A. Stefan, I. Dinca, G. Pelin, A. Ficai, E. Andronescu, O. Oprea, G. Voicu . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11 Experimental and numerical study of hot-steel-plate flatness Eksperimentalni in numeri~ni {tudij ravnosti vro~ih plo{~ iz jekla J. Hrabovský, M. Pohanka, P. J. Lee, J. H. Kang . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17 Investigation of hole effects on the critical buckling load of laminated composite plates Preiskava vpliva luknje na kriti~no upogibno obremenitev laminiranih kompozitnih plo{~ A. Kurºun, E. Topal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23 Corrosion of the refractory zirconia metering nozzle due to molten steel and slag Korozija ognjeodporne cirkonske dozirne {obe s staljenim jeklom in `lindro K. Wiœniewska, D. Madej, J. Szczerba. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29 Effects of an epoxy-resin-fiber substrate for a -shaped microstrip antenna Vpliv z vlakni oja~ane epoksi podlage pri -obliki mikrotrakaste antene Md. M. Islam, M. R. I. Faruque, M. Tariqul Islam, H. Arshad . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33 X-ray radiography of AISI 4340-2205 steels welded by friction welding Rentgenski pregled jekel AISI 4340-2205, varjenih s trenjem U. Caligulu, M. Yalcinoz, M. Turkmen, S. Mercan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39 Thermodynamic properties and microstructures of different shape-memory alloys Termodinami~ne lastnosti in mikrostruktura razli~nih zlitin z oblikovnim spominom L. Gomid`elovi}, E. Po`ega, A. Kostov, N. Vukovi}, D. @ivkovi}, D. Manasijevi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47 The relationship between thermal treatment of serpentine and its reactivity Odvisnost med toplotno obdelavo serpentina in njegovo aktivnostjo G. Su~ik, A. Szabóová, L’. Popovi~, D. Hr{ak . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55 Deformations and velocities during the cold rolling of aluminium alloys Deformacija in hitrosti pri hladnem valjanju aluminijevih zlitin M. Mi{ovi}, N. Tadi}, M. Ja}imovi}, M. Janji} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59 Prediction of the chemical non-homogeneity of 30MnVS6 billets with genetic programming Napovedovanje nehomogenosti kemijske sestave pri gredicah 30MnVS6 s pomo~jo genetskega programiranja M. Kova~i~, D. Novak. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69 Effect of the TiBN coating on a HSS drill when drilling the MA8M Mg alloy Vpliv TiBN prevleke na HSS svedru pri vrtanju MA8M Mg zlitine F. Karaca, B. Aksakal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75 Application of the Taguchi method to select the optimum cutting parameters for tangential cylindrical grinding of AISI D3 tool steel Uporaba Taguchi metode za izbiro optimalnih parametrov odrezavanja pri tangencialnem cilindri~nem bru{enju orodnega jekla AISI D3 C. Ozay, H. Ballikaya, V. Savas. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 81 Effects of friction-welding parameters on the morphological properties of an Al/Cu bimetallic joint Vpliv parametrov tornega varjenja na morfolo{ke lastnosti Al/Cu bimetalnega spoja V. D. Mila{inovi}, R. V. Radovanovi}, M. D. Mila{inovi}, B. R. Gligorijevi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 89 Characterisation of the mechanical and corrosive properties of newly developed glass-steel composites Karakterizacija mehanskih in korozijskih lastnosti novo razvitih kompozitov steklo-jeklo O. Lyubimova, E. Gridasova, A. Gridasov, G. Frieling, M. Klein, F. Walther . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 95 1018 Materiali in tehnologije / Materials and technology 49 (2016) 6, 1017–1040 LETNO KAZALO – INDEX Phase analysis of the slag after submerged-arc welding Analiza faz v `lindri pri oblo~nem varjenju pod pra{kom M. Prijanovi~ Tonkovi~, J. Lamut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 101 Optimizing the parameters for friction welding stainless steel to copper parts Optimiranje parametrov pri tornem varjenju nerjavnega jekla na bakrene dele M. Sahin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 109 WEDM cutting of Inconel 718 nickel-based superalloy: effects of cutting parameters on the cutting quality WEDM rezanje nikljeve superzlitine Inconel 718: vpliv parametrov rezanja na kvaliteto rezanja U. Çaydaº, M. Ay . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 117 Influence of dredged sediment on the shrinkage behavior of self-compacting concrete Vpliv izkopanih sedimentov na kr~enje samozgo{~evalnega betona N. E. Bouhamou, F. Mostefa, A. Mebrouki, K. Bendani, N. Belas . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 127 Study of the properties and hygrothermal behaviour of alternative insulation materials based on natural fibres [tudij lastnosti in higrotermalno obna{anje alternativnih izolacijskih materialov na osnovi naravnih vlaken J. Zach, M. Reif, J. Hroudová . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 137 Prediction of the elastic moduli of chicken-feather-reinforced PLA and a comparison with experimental results Napovedovanje modulov elasti~nosti PLA, oja~anega s pi{~an~jim perjem in primerjava z eksperimentalnimi rezultati U. Özmen, B. Okutan Baba . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 141 Composites based on inorganic matrices for extreme exposure conditions Kompoziti z anorgansko osnovo za izpostavitev ekstremnim razmeram A. Dufka, T. Melichar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 147 The effect of EO and steam sterilization on the mechanical and electrochemical properties of titanium Grade 4 Vpliv EO in sterilizacije s paro na mehanske in elektrokemijske lastnosti titana Grade 4 M. Basiaga, W. Walke, Z. Paszenda, A. Kajzer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153 Influence of the carbide-particle spheroidisation process on the microstructure after the quenching and annealing of 100CrMnSi6-4 bearing steel Vpliv procesa sferoidizacije karbidnih delcev na mikrostrukturo jekla 100CrMnSi6-4 za le`aje po kaljenju in popu{~anju J. Dlouhy, D. Hauserova, Z. Novy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 159 2016/2 Corrosion behavior and the weak-magnetic-field effect of aluminum packaging paper Vpliv {ibkega magnetnega polja na korozijo aluminijeve embala`ne folije N. Zazi, J.-P. Chopart, A. Bilek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 165 Characteristics of the AlTiCrN+DLC coating deposited with a cathodic arc and the PACVD process Zna~ilnosti AlTiCrN+DLC prevleke, nane{ene s katodnim oblokom in PACVD postopkom K. Lukaszkowicz, E. Jonda, J. Sondor, K. Balin, J. Kubacki. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 175 Implicit numerical multidimensional heat-conduction algorithm parallelization and acceleration on a graphics card Paralelizacija in pospe{itev implicitnega numeri~nega ve~dimenzijskega algoritma prevajanja toplote na grafi~ni kartici M. Pohanka, J. Ondrou{ková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 183 Magnetic properties and microstructure of a bulk amorphous Fe61Co10Ti3Y6B20 alloy, fabricated as rods and tubes Magnetne lastnosti in mikrostruktura masivne amorfne zlitine Fe61Co10Ti3Y6B20 v obliki palic in cevi M. Nabia³ek, K. Bloch, K. Szl¹zak, M. Szota . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 189 Effect of the skin-core morphology on the mechanical properties of injection-moulded parts Vpliv morfologije skorja-jedro na mehanske lastnosti vbrizganih delov E. Hnatkova, Z. Dvorak. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 195 Recrystallization behaviour of a nickel-based superalloy Obna{anje superzlitine na osnovi niklja pri rekristalizaciji P. Podany, Z. Novy, J. Dlouhy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 199 Estimation of the number of forward time steps for the sequential Beck approach used for solving inverse heat-conduction problems Ugotavljanje {tevila vnaprej{njih ~asovnih korakov za sekven~ni Beckov pribli`ek pri re{evanju problemov inverzne toplotne prevodnosti J. Komínek, M. Pohanka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 207 Enhanced stability and electrochemical performance of a BaTiO3/PbO2 electrode via a layer obtained with layer electrodeposition Izbolj{ana stabilnost in elektrokemijska zmogljivost elektrode BaTiO3/PbO2, izdelane z elektrodepozicijo plast na plast G. Muthuraman, K. Karunakaran, I. S. Moon. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 211 Deformation behaviour of amorphous Fe-Ni-W/Ni bilayer-confined bulk metallic glasses Obna{anje deformiranega, amorfnega, na dve plasti omejenega kovinskega stekla Fe-Ni-W/Ni H. K. Lau, N. Yip, S. H. Chen, W. Chen, K. C. Chan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 217 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 1019 LETNO KAZALO – INDEX Synergistic effect of organic- and ceramic-based ingredients on the tribological characteristics of brake friction materials Sinergisti~en vpliv sestavin z organsko in kerami~no osnovo na tribolo{ke zna~ilnosti materialov za torne zavore R. Ertan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 223 Optimization of the parameters for the surfactant-added EDM of a Ti–6Al–4V alloy using the GRA-Taguchi method Optimizacija povr{insko aktivnih me{anih EDM parametrov na Ti-6Al-4V zlitini z uporabo GRA-Taguchi metode M. Kolli, K. Adepu . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 229 Determination of the cutting-tool performance of high-alloyed white cast iron (Ni-Hard 4) using the Taguchi method Dolo~anje zmogljivosti rezalnih orodij na mo~no legiranem belem litem `elezu (Ni-Hard 4) z uporabo Taguchi metode D. Kir, H. Öktem, M. Çöl, F. Gül Koç, F. Erzincanli . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 239 Use of the ABI technique to measure the mechanical properties of aluminium alloys: effect of chemical composition on the mechanical properties of the alloys Uporaba tehnike ABI za merjenje mehanskih lastnosti aluminijevih zlitin: vpliv kemijske sestave na mehanske lastnosti zlitin M. Puchnin, O. Trudonoshyn, O. Prach . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 247 Fe-Zn intermetallic phases prepared by diffusion annealing and spark-plasma sintering Fe-Zn intermetalne faze, pripravljene z difuzijskim `arjenjem in s sintranjem v iskre~i plazmi P. Pokorný, J. Cinert, Z. Pala . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 253 High-temperature oxidation of silicide-aluminide layer on the TiAl6V4 alloy prepared by liquid-phase siliconizing Visokotemperaturna oksidacija plasti silicid-aluminid, pripravljene s silikoniziranjem s teko~o fazo zlitine TiAl6V4 T. F. Kubatík . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 257 Characterization and kinetics of plasma-paste-borided AISI 316 steel Karakterizacija in kinetika plazma boriranja s pasto jekla AISI 316 R. Chegroune, M. Keddam, Z. Nait Abdellah, S. Ulker, S. Taktak, I. Gunes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 263 Investigation of the adhesion and wear properties of borided AISI H10 steel Preiskava adhezije in obrabnih lastnosti boriranega jekla AISI H10 I. Gunes, M. Ozcatal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 269 The effects of cutting conditions on the cutting torque and tool life in the tapping process for AISI 304 stainless steel Vpliv pogojev rezanja na moment pri rezanju in zdr`ljivost navojnega vreznika pri vrezovanju notranjih navojev v nerjavno jeklo AISI 304 G. Uzun, I. Korkut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 275 Chemical synthesis and densification behavior of Ag/ZnO metal-matrix composites Obna{anje Ag/ZnO kompozita s kovinsko osnovo pri kemijski sintezi in zgo{~evanju M. Ardestani . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 281 2016/3 Experimental verifications and numerical thermal simulations of automobile lamps Eksperimentalna preverjanja in numeri~ne toplotne simulacije avtomobilskih `arometov M. Guzej, J. Horsky . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 289 Tensile and compressive tests of textile composites and results analysis Natezni in tla~ni preizkusi tekstilnih kompozitov in analiza rezultatov K. Kunc, T. Kroupa, R. Zem~ík, J. Krystek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 295 Deformation behaviour of a natural-shaped bone scaffold Obna{anje naravno oblikovanega ogrodja kosti pri deformaciji D. Kytýø, T. Doktor, O. Jirou{ek, T. Fíla, P. Koudelka, P. Zlámal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 301 Printed microstrip line-fed patch antenna on a high-dielectric material for C-band applications Tiskana mikrotrakasta linijsko napajana krpasta antena na visoko dielektri~nem materialu za uporabo v C-pasu Md. M. Islam, M. R. I. Faruque, M. F. Mansor, M. T. Islam. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 307 Compressive properties of auxetic structures produced with direct 3D printing Stiskanje struktur materialov z negativnim Poissonovim razmerjem, proizvedenih z neposrednim tridimenzionalnim tiskanjem P. Koudelka, O. Jirou{ek, T. Fíla, T. Doktor . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 311 Model of progressive failure for composite materials using the 3D Puck failure criterion Model postopnega popu{~anja kompozitnega materiala z uporabo Puckovega tridimenzionalnega kriterija poru{itve L. Bek, R. Zem~ík . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 319 Physicochemical properties of a Ti67 alloy after EO and steam sterilization Fizikalno kemijske lastnosti zlitine Ti67 po EO in parni sterilizaciji W. Walke, M. Basiaga, Z. Paszenda, J. Marciniak, P. Karasinski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 323 Surface properties of a laser-treated biopolymer Lastnosti povr{ine biopolimera, obdelanega z laserjem I. Michaljani~ova, P. Slepi~ka, S. Rimpelova, P. Sajdl, V. [vor~ík . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 331 Analyzing the heat-treatment effect on the mechanical properties of free-cutting steels Analiza vpliva toplotne obdelave na mehanske lastnosti avtomatnih jekel M. K. Kulekci, U. Esme, F. Kahraman, R. Ozgun, A. Akkurt . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 337 LETNO KAZALO – INDEX 1020 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 LETNO KAZALO – INDEX Analysis of the cutting temperature and surface roughness during the orthogonal machining of AISI 4140 alloy steel via the Taguchi method Analiza temperature rezanja in hrapavosti povr{ine s Taguchi metodo pri ortogonalni strojni obdelavi legiranega jekla AISI 4140 A. R. Motorcu, Y. Isik, A. Kus, M. C. Cakir . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 343 Weldability of Ti6Al4V to AISI 2205 with a nickel interlayer using friction welding Preizku{anje varivosti pri varjenju s trenjem Ti6Al4V in AISI 2205 z vmesno plastjo niklja I. Kirik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 353 Effect of activated flux and nitrogen addition on the bead geometry of borated stainless-steel GTA welds Vpliv aktiviranega topila in dodatka du{ika na geometrijo kopeli pri GTA zvarih boriranega nerjavnega jekla G. R. Kumar, G. D. J. Ram, S. R. Koteswara Rao . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 357 Microstructural evolution during the transient liquid-phase bonding of dissimilar nickel-based superalloys of IN738LC and NIMONIC 75 Razvoj mikrostrukture med spajanjem s prehodno teko~o fazo neenakih superzlitin na osnovi niklja IN738LC in NIMONIC 75 M. G. Khakian, S. Nategh, S. Mirdamadi. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 365 Workability behaviour of Cu-TiB2 powder-metallurgy preforms during cold upsetting Preoblikovalnost Cu-TiB2 predoblik izdelanih z metalurgijo prahov med hladnim kovalnim preizkusom S. Gadakary, A. Kumar Khanra, M. J. Davidson . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 373 Effects of extrusion shear on the microstructures and a fracture analysis of a magnesium alloy in the homogenized state Vplivi stri`enja med iztiskanjem homogenizirane magnezijeve zlitine na mikrostrukturo in na analizo preloma H. J. Hu, Z. Sun, D. F. Zhang . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 381 FSW welding of Al-Mg alloy plates with increased edge roughness using square pin tools of various shoulder geometries FSW varjenje plo{~ iz Al-Mg zlitine s pove~ano hrapavostjo robov z orodjem s kvadratno konico in razli~no geometrijo bokov S. Balos, L. Sidjanin, M. Dramicanin, D. Labus Zlatanovic, A. Antic. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 387 Improvement of selective copper extraction from a heat-treated chalcopyrite concentrate with atmospheric sulphuric-acid leaching Izbolj{anje selektivne ekstrakcije bakra iz toplotno obdelanega koncentrata halkopirita z lu`enjem z `vepleno kislino na zraku E. Uzun, M. Zengin, Ý. Atýlgan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 395 Homogenization of an Al-Mg alloy and alligatoring failure: alloy ductility and fracture Homogenizacija Al-Mg zlitine in krokodiljenje: duktilnost zlitine in prelom E. Romhanji, T. Radeti}, M. Popovi}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 403 Assessment of tubular light guides with respect to building physics Ocena cevastih vodnikov svetlobe glede na gradbeno fiziko F. Vajkay, D. Be~kovský, V. Tichomirov . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 409 Creep behaviour of a short-fibre C/PPS composite Vedenje kratkih vlaken C/PPS kompozitov pri lezenju T. Fíla, P. Koudelka, D. Kytýø, J. Hos, J. [leichrt. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 413 Increasing micro-purity and determining the effects of the production with and without vacuum refining on the qualitative parameters of forged-steel pieces with a high aluminium content Pove~anje mikro~isto~e in dolo~itev u~inka proizvodnje, z vakuumskim rafiniranjem ali brez, na kvalitativne parametre kovanega jekla z visoko vsebnostjo aluminija V. Kurka, J. Pindor, J. Kosòovská, Z. Adolf . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 419 Use of the ABI technique to measure the mechanical properties of aluminium alloys: effect of heat-treatment conditions on the mechanical properties of alloys Uporaba ABI tehnike za merjenje mehanskih lastnosti aluminijevih zlitin: vpliv pogojev toplotne obdelave na mehanske lastnosti zlitin O. Trudonoshyn, M. Puchnin, O. Prach . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 427 Investigation of the effect of holding time and melt stirring on the grain refinement of an A206 alloy Preiskava vpliva ~asa zadr`evanja in me{anja taline na zmanj{anje velikosti zrn zlitine A206 N. Akar, Z. Tanyel, K. Kocatepe, R. Kayikci . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 433 Investigating the influence of cutting speed on the tool life of a cutting insert while cutting DIN 1.4301 steel Preiskava vpliva hitrosti rezanja na zdr`ljivost vlo`ka za rezanje pri rezanju jekla DIN 1.4301 R. Dubovská, J. Majerík, R. ^ep, K. Kouøil . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 439 NiAl intermetallic prepared with reactive sintering and subsequent powder-metallurgical plasma-sintering compaction Reakcijsko sintranje in zgo{~evanje s plazemskim sintranjem NiAl intermetalne zlitine A. Michalcová, D. Vojtìch, T. F. Kubatík, P. Novák, P. Dvoøák, P. Svobodová, I. Marek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 447 Microscopic characterization and particle distribution in a cast steel matrix composite Mikroskopska karakterizacija in razporeditev delcev v kompozitu z matrico litega jekla A. Kra~un, M. Torkar, J. Burja, B. Podgornik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 451 A comparison of as-welded and simulated heat affected zone (HAZ) microstructures Primerjava mikrostrukture toplotno vplivanega podro~ja varjenega in simuliranih vzorcev R. Celin, J. Burja, G. Kosec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 455 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 1021 LETNO KAZALO – INDEX Degradation of an AISI 304 stainless-steel tank Degradacija rezervoarja iz AISI 304 nerjavnega jekla M. Torkar, I. Paulin, B. Podgornik. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 461 2016/4 Predgovor urednika/Editor’s preface Matja` Torkar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 469 Organosoluble xanthone-based polyimides: synthesis, characterization, antioxidant activity and heavy-metal sorption Organsko topni poliamidi na osnovi ksantona: sinteza, karakterizacija, antioksidativna aktivnost in sorpcija te`kih kovin M. M. Lakouraj, G. Rahpaima, R. Azimi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 471 Correlation of the heat-transfer coefficient at sprinkled tube bundle Korelacija koeficienta prenosa toplote pri potresenem snopu cevi P. Kracík, L. [najdárek, M. Lisý, M. Balá{, J. Pospí{il. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 479 Mathematical modeling of a cement raw-material blending process using a neural network Matemati~no modeliranje postopka me{anja sestavin cementa s pomo~jo nevronske mre`e A. Egrisogut Tiryaki, R. Kozan, N. Gokhan Adar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 485 Possibilities of determining the air-pore content in cement composites using computed tomography and other methods Mo`nosti dolo~anja vsebnosti zra~nih por v cementnih kompozitih z uporabo ra~unalni{ke tomografije in drugih metod B. Moravcová, P. Põssl, P. Misák, M. Bla`ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 491 Material and technological modelling of closed-die forging Materialno-tehnolo{ko modeliranje kovanja v zaprtem utopu I. Vorel, [. Jení~ek, H. Jirková, B. Ma{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 499 Investigation of wear behavior of borided AISI D6 steel Preiskava obrabe boriranega jekla AISI D6 I. Gunes, S. Kanat . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 505 Investigation of Portevin-Le Chatelier effect of hot-rolled Fe-13Mn-0.2C-1Al-1Si TWIP steel Preiskava Portevin-Le Chatelier u~inka pri vro~em valjanju Fe-13Mn-0.2C-1Al-1Si TWIP jekla B. Aydemir, H. Kazdal Zeytin, G. Guven. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 511 The influence of surface coatings on the tooth tip deflection of polymer gears Vpliv povr{inskih prevlek na poves vrha zoba polimernih zobnikov B. Trobentar, S. Glode`, J. Fla{ker, B. Zafo{nik. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 517 Vapour-phase condensed composite materials based on copper and carbon Kompoziti na osnovi bakra in ogljika, kondenzirani iz plinske faze V. Bukhanovsky, M. Rudnytsky, M. Grechanyuk, R. Minakova, C. Zhang . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 523 Homogenization of an Al-Mg alloy and alligatoring failure: influence of the microstructure Homogenizacija Al-Mg zlitine in krokodiljenje: vpliv mikrostrukture E. Romhanji, T. Radeti}, M. Popovi}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 531 Metal particles size influence on graded structure in composite Al2O3-Ni Vpliv velikosti kovinskih delcev na gradientno strukturo kompozita Al2O3-Ni J. Zygmuntowicz, A. Miazga, K. Konopka, W. Kaszuwara . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 537 Static and dynamic tensile characteristics of S420 and IF steel sheets Stati~ne in dinami~ne natezne lastnosti plo~evine iz S420 in IF jekla M. Mihaliková, V. Girman, A. Li{ková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 543 Acoustic and electromagnetic emission of lightweight concrete with polypropylene fibers Akusti~na in elektromagnetna emisija lahkega betona s polipropilenskimi vlakni R. [toudek, T. Tr~ka, M. Matysík, T. Vymazal, I. Pl{ková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 547 Multi-criteria analysis of synthesis methods for Ni-based catalysts Ve~kriterijska analiza sinteznih metod na osnovi Ni katalizatorja V. Nikoli}, B. Agarski, @. Kamberovi}, Z. An|i}, I. Budak, B. Kosec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 553 Influence of structural defects on the magnetic properties of massive amorphous Fe60Co10Mo2WxY8B20-x (x = 1, 2) alloys produced with the injection casting method Vpliv strukturnih napak na magnetne lastnosti masivne amorfne zlitine Fe60Co10Mo2WxY8B20-x (x = 1, 2), izdelane z metodo litja z vbrizgavanjem J. Gondro, K. B³och, M. Nabia³ek, S. Garus . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 559 Possibilities of NUS and Impact-Echo methods for monitoring steel corrosion in concrete Mo`nosti metod NUS in Udarec-odmev za kontrolo korozije jekla v betonu K. Tim~aková-[amárková, M Matysík, Z. Chobola . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 565 Characterization of heterogeneous arc welds through miniature tensile testing and Vickers-hardness mapping Karakterizacija heterogenih zvarov s pomo~jo miniaturnih nateznih preizkusov in matri~nimi meritvami trdote po Vickersu S. Hertelé, J. Bally, N. Gubeljak, P. [tefane, P. Verleysen, W. De Waele. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 571 1022 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 Overcooling in overlap areas during hydraulic descaling Podhladitev in prekrivanje podro~ij med hidravli~nim raz{kajanjem M. Pohanka, H. Votavová . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 575 Investigation of the mechanical properties of a cork/rubber composite Raziskava mehanskih lastnosti kompozita pluta/guma R. Kottner, J. Kocáb, J. Heczko, J. Krystek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 579 Fabrication and properties of SiC reinforced copper-matrix-composite contact material Izdelava in lastnosti s SiC utrjenega kompozitnega materiala na osnovi bakra G. F. Celebi Efe, M. Ýpek, S. Zeytin, C. Bindal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 585 Investigation of the cutting forces and surface roughness in milling carbon-fiber-reinforced polymer composite material Preiskava sil rezanja in hrapavosti povr{ine pri rezkanju kompozitnega polimernega materiala, oja~anega z ogljikovimi vlakni S. Bayraktar, Y. Turgut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 591 Development of aluminium alloys for aerosol cans Razvoj aluminijevih zlitin za aerosol doze S. Kores, J. Turk, J. Medved, M. Von~ina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 601 Computer tools to determine physical parameters in wooden houses Dolo~anje fizikalnih parametrov z ra~unalni{kimi orodji v lesenih hi{ah D. Be~kovský, F. Vajkay, V. Tichomirov . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 607 Impression relaxation and creep behavior of Al/SiC nanocomposite Sprostitev vtisa in obna{anje Al/SiC nanokompozita pri lezenju Y. S. Kakhki, S. Nategh, T. S. Mirdamadi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 611 Microstructure and properties of the high-temperature (HAZ) of thermo-mechanically treated S700MC high-yield-strength steel Mikrostruktura in lastnosti visoko temperaturnega obmo~ja zvara (HAZ) termo-mehansko obdelanega jekla S700MC z visoko mejo plasti~nosti J. Górka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 617 New concept for manufacturing closed die forgings of high strength steels Nov koncept izdelave odkovkov iz visokotrdnostnih jekel v zaprtih utopih K. Ibrahim, I. Vorel, D. Bublíková, B. Ma{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 623 Helium atom scattering – a versatile technique in studying nanostructures Sipanje atomov helija – vsestranska tehnika za {tudij nanostruktur G. Bavdek, D. Cvetko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 627 2016/5 Predgovor urednika/Editor’s preface P. J. McGuiness. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 639 Effect of the addition of niobium and aluminium on the microstructures and mechanical properties of micro-alloyed PM steels Vpliv dodatka niobija in aluminija na mikrostrukturo in mehanske lastnosti mikrolegiranih PM jekel S. Gündüz, M. A. Erden, H. Karabulut, M. Türkmen . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 641 Characteristics of dye-sensitized solar cells with carbon nanomaterials Zna~ilnosti na fiksirano barvo ob~utljivih solarnih celic z ogljikovimi nanomateriali L. A. Dobrzañski, A. Mucha, M. Prokopiuk vel Prokopowicz, M. Szindler, A. Dryga³a, K. Lukaszkowicz . . . . . . . . . . . . . . . . . . . . . . . 649 The effect of the welding parameters and the coupling agent on the welding of composites Vpliv parametrov varjenja in sredstva za spajanje na varjenje kompozitov S. E. Erdogan, U. Huner . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 655 Chemical cross-linking of chitosan/polyvinyl alcohol electrospun nanofibers Kemijsko zamre`enje elektro spredenih nanovlaken iz hitosan/polivinil alkohola S. Pouranvari, F. Ebrahimi, G. Javadi, B. Maddah . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 663 Investigation of hole profiles in deep micro-hole drilling of AISI 420 stainless steel using powder-mixed dielectric fluids Preiskava profilov luknje pri globokem vrtanju mikroluknje v AISI 420 nerjavnem jeklu s pomo~jo dielektri~ne teko~ine s prime{anim prahom V. Yýlmaz . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 667 The phenomenon of reduced plasticity in low-alloyed copper Pojav zmanj{anja plasti~nosti malo legiranega bakra W. Ozgowicz, E. Kalinowska-Ozgowicz, B. Grzegorczyk, K. Lenik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 677 The effect of high-speed grinding technology on the properties of fly ash Vpliv tehnologije hitrega mletja na lastnosti lete~ega pepela K. Dvoøák, I. Hájková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 683 Investigation of the mechanical properties of electrochemically deposited Au-In alloy films using nano-indentation Preiskava mehanskih lastnosti elektrokemijsko nane{enega filma zlitine Au-In z nanovtiskovanjem S. Cherneva, R. Iankov, M. Georgiev, T. Dobrovolska, D. Stoychev . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 689 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 1023 LETNO KAZALO – INDEX Growth of K2CO3-doped KDP crystal from an aqueous solution and an investigation of its physical properties Rast KDP kristalov z dodatkom K2CO3 iz vodne raztopine in preiskava njihovih fizikalnih lastnosti A. Rousta, H. R. Dizaji . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 695 Surface treatment of heat-treated cast magnesium and aluminium alloys Obdelava povr{ine toplotno obdelanih magnezijevih in aluminijevih livnih zlitin T. Tañski, M. Wiœniowski, W. Matysiak, M. Staszuk, R. Szklarek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 699 Analysis of the structural-defect influence on the magnetization process in and above the Rayleigh region Analiza vpliva strukturnih defektov na proces magnetizacije v in nad Rayleigh podro~jem K. Gruszka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 707 Effect of sulphide inclusions on the pitting-corrosion behaviour of high-Mn steels in chloride and alkaline solutions Vpliv sulfidnih vklju~kov na jami~asto korozijo jekel z visoko vsebnostjo Mn v raztopinah kloridov in alkalij A. Grajcar, A. P³achciñska . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 713 Influence of Na2SiF6 on the surface morphology and corrosion resistance of an AM60 magnesium alloy coated by micro arc oxidation Vpliv Na2SiF6 na morfologijo povr{ine in korozijsko odpornost magnezijeve zlitine AM60, prekrite z mikrooblo~no oksidacijo A. Ayday . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 719 Mechanical properties of polyamide/carbon-fiber-fabric composites Mehanske lastnosti kompozitne tkanine iz poliamid/ogljikovih vlaken C.-E. Pelin, G. Pelin, A. ªtefan, E. Andronescu, I. Dincã, A. Ficai, R. Truºcã . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 723 Evaluation of the grindability of recycled glass in the production of blended cements Ocena sposobnosti drobljenja recikliranega stekla pri proizvodnji me{anih cementov K. Dvoøák, D. Dolák, D. V{ianský, P. Dobrovolný . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 729 Rheological properties of alumina ceramic slurries for ceramic shell-mould fabrication Reolo{ke lastnosti go{~e iz glinice za izdelavo kerami~nih kalupov J. Szymañska, P. Wiœniewski, M. Ma³ek, J. Mizera . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 735 Effect of mechanical activation on the synthesis of a magnesium aluminate spinel Vpliv mehanske aktivacije na sintezo magnezij-aluminatnega {pinela D. Kýrsever, N. K. Karabulut, N. Canikoðlu, H. Ö. Toplan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 739 Phase and microstructure development of LSCM perovskite materials for SOFC anodes prepared by the carbonate-coprecipitation method Razvoj kristalnih faz in mikrostrukture LSCM perovskitnih materialov za SOFC anode, pripravljenih s karbonatno metodo koprecipitacije K. Zupan, M. Marin{ek, T. Skalar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 743 Artificial aggregate from sintered coal ash Umetni agregat iz sintranega pepela premoga V. Cerny, R. Drochytka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 749 Investigation studies involving wear-resistant ALD/PVD hybrid coatings on sintered tool substrates Preiskave obrabne odpornosti hibridnega nanosa ALD/PVD na sintranem orodju M. Staszuk, D. Paku³a, T. Tañski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 755 Dissimilar spot welding of DQSK/DP600 steels: the weld-nugget growth To~kasto varjenje jekel DQSK/DP600: rast jedra zvara S. P. Hoveida Marashi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 761 Armour plates from Kozlov rob – analyses of two unusual finds Oklepni plo{~i s Kozlovega roba – analize dveh nenavadnih najdb T. Lazar, P. Mrvar, M. Lamut, P. Fajfar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 767 Numerical and experimental investigation of the effect of hydrostatic pressure on the residual stress in boiler-tube welds Numeri~na in eksperimentalna preiskava vpliva hidrostati~nega tlaka na zaostale napetosti v zvaru na kotlovski cevi D. Danyali, E. Ranjbarnodeh . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 775 Effect of direct cooling conditions on the microstructure and properties of hot-forged HSLA steels for mining applications Vpliv pogojev ohlajanja na mikrostrukturo in lastnosti vro~e kovanih HSLA jekel za uporabo v rudarstvu P. Skubisz, £. Lisiecki, T. Skowronek, A. ¯ak, W. Zalecki . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 783 Influence of the tool rotational speed on the microstructure and joint strength of friction-stir spot-welded pure copper Vpliv hitrosti vrtenja orodja na mikrostrukturo in trdnost torno vrtilno to~kasto zvarjenega spoja ~istega bakra I. Dinaharan, E. T. Akinlabi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 791 Measurement of bio-impedance on an isolated rat sciatic nerve obtained with specific current stimulating pulses Meritev bioimpedance na izoliranem `ivcu Ischiadicus pri podgani, vzbujenem s posebnimi tokovnimi stimulacijskimi impulzi J. Rozman, M. C. @u`ek, R. Frange`, S. Ribari~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 797 Influence of different production processes on the biodegradability of an FeMn17 alloy Vpliv razli~nih procesov izdelave na biorazgradljivost zlitine FeMn17 A. Kocijan, I. Paulin, ^. Donik, M. Ho~evar, K. Zeli~, M. Godec. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 805 1024 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 LETNO KAZALO – INDEX LETNO KAZALO – INDEX Effect of a combination of fly ash and shrinkage-reducing additives on the properties of alkali-activated slag-based mortars Vpliv kombinacije lete~ega pepela in dodatka za zmanj{anje kr~enja na lastnosti malte iz z alkalijami aktivirane `lindre V. Bílek, L. Kalina, J. Koplík, M. Mon~eková, R. Novotný . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 813 Cutting-tool performance in the end milling of carbon-fiber-reinforced plastics Zmogljivost rezilnega orodja pri rezkanju plastike, oja~ane z ogljikovimi vlakni O. Bílek, S. Rusnáková, M. @aludek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 819 Influence of solidification speed on the structure and magnetic properties of Nd10Fe81Zr1B6 in the as-cast state Vpliv hitrosti strjevanja na strukturo in magnetne lastnosti zlitine Nd10Fe81Zr1B6 v litem stanju M. Doœpia³, M. Nabia³ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 823 Metalografska preiskava in korozijska odpornost zvarov feritnega nerjavnega jekla Metallographic investigation and corrosion resistance of welds of ferritic stainless steels M. Torkar, A. Kocijan, R. Celin, J. Burja, B. Podgornik. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 829 2016/6 The microstructure of metastable austenite in X5CrNi18-10 steel after its strain-induced martensitic transformation Mikrostruktura metastabilnega avstenita po pretvorbi v napetostno inducirani martenzit v jeklu X5CrNi18-10 A. Kurc-Lisiecka, W. Ozgowicz, E. Kalinowska-Ozgowicz, W. Maziarz . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 837 The structure and morphology of the surface of duplex layers after saturation of the base layer with carbon Struktura in morfologija povr{ine dupleks plasti po nasi~enju osnovne plasti z ogljikom W. Skoneczny . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 845 Modeling of shot-peening effects on the surface properties of a (TiB + TiC)/Ti–6Al–4V composite employing artificial neural networks Modeliranje vpliva hladnega povr{inskega kovanja na lastnosti povr{ine (TiB + TiC)/Ti-6Al-4V kompozita s pomo~jo umetnih nevronskih mre` E. Maleki, A. Zabihollah . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 851 Analysis of twin-roll casting AA8079 alloy 6.35-μm foil rolling process Analiza procesa valjanja 6,35 μm folije iz zlitine AA8079 ulite med dvema valjema A. Can, H. Arikan, K. Çýnar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 861 Antimicrobial modification of polypropylene with silver nanoparticles immobilized on zinc stearate Protimikrobno spreminjanje polipropilena z nanodelci srebra, imobiliziranih na cinkovem stearatu G. Jandikova, P. Holcapkova, M. Hrabalikova, M. Machovsky, V. Sedlarik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 869 A new wideband negative-refractive-index metamaterial Novi {irokopasovni metamaterial z negativnim lomnim koli~nikom S. S. Islam, M. R. Iqbal Faruque, M. J. Hossain, M. T. Islam. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 873 Evaluation of the degree of degradation using the impact-echo method in civil engineering Ocena stopnje degradacije v gradbeni{tvu z uporabo metode odmeva zvo~nih valov D. [tefková, K. Tim~aková, L. Topoláø, P. Cikrle. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 879 Non-traditional whiteware based on calcium aluminate cement Netradicionalni porcelan na osnovi kalcij aluminatnega cementa R. Sokolar. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 885 Behaviour of new ODS alloys under single and multiple deformation Obna{anje novih ODS zlitin pri enojni in ve~kratni deformaciji B. Ma{ek, O. Khalaj, Z. Nový, T. Kubina, H. Jirkova, J. Svoboda, C. [tádler . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 891 Electromagnetic-shielding effectiveness and fracture behavior of laminated (Ni–NiAl3) composites U~inkovitost elektromagnetne za{~ite in obna{anje pri lomu laminiranega kompozita (Ni-NiAl3) T. Yener, S. C. Yener, S. Zeytýn . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 899 Effect of thermomechanical treatment on the intergranular corrosion of Al-Mg-Si-Type alloy bars Vpliv termomehanske predelave na interkristalno korozijo palic iz zlitin Al-Mg-Si P. Sláma, J. Nacházel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 903 Valorization of brick wastes in the fabrication of concrete blocks Ocena odpadkov iz opeke pri proizvodnji betonskih zidakov Y. Ghernouti, B. Rabehi, T. Bouziani, R. Chaid . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 911 Porous magnesium alloys prepared by powder metallurgy Porozne magnezijeve zlitine, izdelane s pomo~jo metalurgije prahov P. Salvetr, P. Novák, D. Vojtìch . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 917 Influence of nano-sized cobalt oxide additions on the structural and electrical properties of nickel-manganite-based NTC thermistors Vpliv dodatka nanodelcev kobaltovega oksida na zgradbo in elektri~ne lastnosti NTC termistorjev na osnovi nikljevega manganita G. Hardal, B. Y. Price . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 923 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 1025 LETNO KAZALO – INDEX Durability of alumina silicate concrete based on slag/fly-ash blends against acid and chloride environments Zdr`ljivost betona na osnovi glinice in silikatov iz me{anice `lindra/lete~i pepel na kislo in kloridno okolje R. Gopalakrishnan, K. Chinnaraju . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 929 The size effect of heat-transfer surfaces on boiling Vpliv velikosti povr{in, ki prena{ajo toploto na vrenje P. Kracík, M. Balas, M. Lisy, J. Pospí{il . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 939 Effect of gas atmosphere on the non-metallic inclusions in laser-welded trip steel with Al and Si additions Vpliv plinske atmosfere na nekovinske vklju~ke v lasersko varjenem trip jeklu z dodatkom Al in Si A. Grajcar, M. Ró¿añski, M. Kamiñska, B. Grzegorczyk . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 945 Machining parameters influencing in electro chemical machining on AA6061 MMC Parametri strojne obdelave, ki vplivajo na elektrokemijsko strojno obdelavo AA6061 MMC C. J. Thankaraj Mariapushpam, D. Ravindran, M. D. Anand . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 951 Modeling the deep drawing of an AISI 304 stainless-steel rectangular cup using the finite-element method and an experimental validation Modeliranje globokega vleka pravokotne ~a{e iz AISI 304 nerjavnega jekla z metodo kon~nih elementov in z eksperimentalnim preverjanjem B. Sener, H. Kurtaran . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 961 Surface and anticorrosion properties of hydrophobic and hydrophilic TiO2 coatings on a stainless-steel substrate Povr{inske in protikorozijske lastnosti hidrofobnih in hidrofilnih TiO2 prevlek na jekleni podlagi M. Conradi, A. Kocijan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 967 Electroslag remelting: A process overview Elektropretaljevanje pod `lindro – pregled procesa B. Arh, B. Podgornik, J. Burja. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 971 Continuous vertical casting of a NiTi alloy Vertikalno kontinuirno litje NiTi zlitine A. Stamboli}, I. An`el, G. Lojen, A. Kocijan, M. Jenko, R. Rudolf . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 981 Hot tensile testing of SAF 2205 duplex stainless steel Vro~i natezni preskusi dupleks nerjavnega jekla SAF 2205 F. Tehovnik, B. @u`ek, J. Burja . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 989 A high-efficiency automatic de-bubbling system for liquid silicone rubber Visokozmogljiv sistem za odpravljanje mehur~kov v teko~i silikonski gumi C.-C. Kuo, C.-M. Huang . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 995 Impact toughness of WMD after MAG welding with micro-jet cooling Udarna `ilavost WMD po MAG varjenju z mikro-jet hlajenjem T. Wegrzyn, J. Piwnik, A. Borek, A. Kurc-Lisiecka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1001 Forming-limit diagrams and strain-rate-dependent mechanical properties of AA6019-T4 and AA6061-T4 aluminium sheet materials Mejni diagrami preoblikovanja in odvisnost mehanskih lastnosti od hitrosti preoblikovanja aluminijevih plo~evin iz AA6019-T4 in AA6061-T4 O. Çavuºoðlu, A. G. Leacock, H. Gürün . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1005 Effect of alternative heat-treatment parameters on the aging behavior of short-fiber-reinforced 2124 Al composites Vpliv alternativnih parametrov toplotne obdelave na staranje 2124 Al kompozita, oja~anega s kratkimi vlakni Y. Altunpak, S. Aslan, M. Oðuz Güler, H. Akbulut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1011 Letnik 50 (2016), 1–6 – Volume 50 (2016), 1–6 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1017 1026 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 LETNO KAZALO – INDEX MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY AVTORSKO KAZALO / AUTHOR INDEX LETNIK / VOLUME 50, 2016, 1–6, A–@ A Adepu K. 229 Adolf Z. 419 Agarski B. 553 Akar N. 433 Akbulut H. 1011 Akinlabi E. T. 791 Akkurt A. 337 Aksakal B. 75 Altunpak Y. 1011 Anand M. D. 951 Andronescu E. 11 Andronescu E. 723 An|i} Z. 553 Antic A. 387 An`el I. 981 Ardestani M. 281 Arh B. 971 Arikan H. 861 Arshad H. 33 Aslan S. 1011 Atýlgan Ý. 395 Ay M. 117 Ayday A. 719 Aydemir B. 511 Azimi R. 471 B Balá{ M. 479, 939 Balin K. 175 Ballikaya H. 81 Bally J. 571 Balos S. 387 Ban C. E. 11 Basiaga M. 153, 323 Bavdek G. 627 Bayraktar S. 591 Be~kovský D. 409, 607 Bek L. 319 Belas N. 127 Bendani K. 127 Bilek A. 165 Bílek O. 819 Bílek V. 813 Bindal C. 585 Bla`ek M. 491 B³och K. 189, 559 Borek A. 1001 Bouhamou N. E. 127 Bouziani T. 911 Bublíková D. 623 Budak I. 553 Bukhanovsky V. 523 Burja J. 451, 455, 829, 971, 989 C Cakir M. C. 343 Caligulu U. 39 Can A. 861 Canikoðlu N. 739 Çavuºoðlu O. 1005 Çaydaº U. 117 Celebi Efe G. F. 585 Celin R. 455, 829 Cerny V. 749 Chaid R. 911 Chan K. C. 217 Chegroune R. 263 Chen S. H. 217 Chen W. 217 Cherneva S. 689 Chinnaraju K. 929 Chobola Z. 565 Chopart J.-P. 165 Cikrle P. 879 Cinert J. 253 Conradi M. 967 Cvetko D. 627 Çýnar K. 861 Çöl M. 239 ^ ^ep R. 439 D Danyali D. 775 Davidson M. J. 373 Dinaharan I. 791 Dincã I. 11, 723 Dizaji H. R. 695 Dlouhy J. 159, 199 Dobrovolný P. 729 Dobrovolska T. 689 Dobrzañski L. A. 649 Doktor T. 301, 311 Dolák D. 729 Donik ^. 805 Doœpia³ M. 823 Dramicanin M. 387 Drochytka R. 749 Dryga³a A. 649 Dubovská R. 439 Dufka A. 147 Dvoøák K. 683, 729 Dvoøák P. 447 Dvorak Z. 195 E Ebrahimi F. 663 Egrisogut Tiryaki A. 485 Erden M. A. 641 Erdogan S. E. 655 Ertan R. 223 Erzincanli F. 239 Esme U. 337 F Fajfar P. 767 Faruque M. R. I. 33, 307 Ficai A. 11, 723 Fíla T. 301, 311, 413 Fla{ker J. 517 Frange` R. 797 Frieling G. 95 G Gadakary S. 373 Garus S. 559 Georgiev M. 689 Ghernouti Y. 911 Girman V. 543 Gligorijevi} B. R. 89 Glode` S. 517 Godec M. 805 Gokhan Adar N. 485 Gomid`elovi} L. 47 Gondro J. 559 Gopalakrishnan R. 929 Górka J. 617 Grajcar A. 713, 945 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 1027 LETNO KAZALO – INDEX Grechanyuk M. 523 Gridasov A. 95 Gridasova E. 95 Gruszka K. 707 Grzegorczyk B. 677, 945 Gubeljak N. 571 Gunes I. 263, 269, 505 Guven G. 511 Guzej M. 289 Gül Koç F. 239 Gündüz S. 641 Gürün H. 1005 H Hájková I. 683 Hardal G. 923 Hauserova D. 159 Heczko J. 579 Hertelé S. 571 Hnatkova E. 195 Ho~evar M. 805 Holcapkova P. 869 Horsky J. 289 Hos J. 413 Hossain M. J. 873 Hoveida Marashi S. P. 761 Hr{ak D. 55 Hrabalikova M. 869 Hrabovský J. 17 Hroudová J. 137 Hu H. J. 381 Huang C.-M. 995 Huner U. 655 I Iankov R. 689 Ibrahim K. 623 Ýpek M. 585 Iqbal Faruque M. R. 873 Isik Y. 343 Islam M. T. 307, 873 Islam Md. M. 33, 307 Islam S. S. 873 J Ja}imovi} M. 59 Jandikova G. 869 Janji} M. 59 Javadi G. 663 Jení~ek [. 499 Jenko M. 981 Jirková H. 499, 891 Jirou{ek O. 301, 311 Jonda E. 175 K Kahraman F. 337 Kajzer A. 153 Kakhki Y. S. 611 Kalina L. 813 Kalinowska-Ozgowicz E. 677, 837 Kamberovi} @. 553 Kamiñska M. 945 Kanat S. 505 Kang J. H. 17 Karabulut H. 641 Karabulut N. K. 739 Karaca F. 75 Karasinski P. 323 Karunakaran K. 211 Kaszuwara W. 537 Kayikci R. 433 Kazdal Zeytin H. 511 Keddam M. 263 Khakian M. G. 365 Khalaj O. 891 Kir D. 239 Kirik I. 353 Kýrsever D. 739 Klein M. 95 Kocáb J. 579 Kocatepe K. 433 Kocijan A. 805, 829, 967, 981 Kolli M. 229 Komínek J. 207 Konopka K. 537 Koplík J. 813 Kores S. 601 Korkut I. 275 Kosec B. 553 Kosec G. 455 Kosòovská J. 419 Kostov A. 47 Koteswara Rao S. R. 357 Kottner R. 579 Koudelka P. 301, 311 Koudelka P. 413 Kouøil K. 439 Kova~i~ M. 69 Kozan R. 485 Kra~un A. 451 Kracík P. 479, 939 Kroupa T. 295 Krystek J. 295, 579 Kubacki J. 175 Kubatík T. F. 257, 447 Kubina T. 891 Kulekci M. K. 337 Kumar G. R. 357 Kumar Khanra A. 373 Kunc K. 295 Kuo C.-C. 995 Kurc-Lisiecka A. 837, 1001 Kurka V. 419 Kurºun A. 23 Kurtaran H. 961 Kus A. 343 Kytýø D. 301, 413 L Labus Zlatanovic D. 387 Lakouraj M. M. 471 Lamut J. 101 Lamut M. 767 Lau H. K. 217 Lazar T. 767 Leacock A. G. 1005 Lee P. J. 17 Lenik K. 677 Li{ková A. 543 Lisiecki £. 783 Lisý M. 479, 939 Lojen G. 981 Lukaszkowicz K. 175, 649 Lyubimova O. 95 M Ma{ek B. 499, 623, 891 Machovsky M. 869 Maddah B. 663 Madej D. 29 Majerík J. 439 Ma³ek M. 735 Maleki E. 851 Manasijevi} D. 47 Mansor M. F. 307 Marciniak J. 323 Marek I. 447 Marin{ek M. 743 Matysiak W. 699 Matysík M 565, 547 Mauder T. 3 Maziarz W. 837 Mebrouki A. 127 Medved J. 601 Melichar T. 147 Mercan S. 39 Mi{ovi} M. 59 Miazga A. 537 Michalcová A. 447 Michaljani~ova I. 331 Mihaliková M. 543 1028 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 LETNO KAZALO – INDEX Mila{inovi} M. D. 89 Mila{inovi} V. D. 89 Minakova R. 523 Mirdamadi S. 365 Mirdamadi T. S. 611 Misák P. 491 Mizera J. 735 Mon~eková M. 813 Moon I. S. 211 Moravcová B. 491 Mostefa F. 127 Motorcu A. R. 343 Mrvar P. 767 Mucha A. 649 Muthuraman G. 211 N Nabia³ek M. 189, 559, 823 Nacházel J. 903 Nait Abdellah Z. 263 Nategh S. 365, 611 Nikoli} V. 553 Novak D. 69 Novák P. 447, 917 Novotný R. 813 Nový Z. 159, 199, 891 O Oðuz Güler M. 1011 Okutan Baba B. 141 Ondrou{ková J. 183 Oprea O. 11 Ozay C. 81 Ozcatal M. 269 Ozgowicz W. 677, 837 Ozgun R. 337 Öktem H. 239 Özmen U. 141 P Paku³a D. 755 Pala Z. 253 Paszenda Z. 153, 323 Paulin I. 461, 805 Pazdera L. 7 Pelin C.-E. 723 Pelin G. 11, 723 Pindor J. 419 Piwnik J. 1001 Pl{ková I. 547 P³achciñska A. 713 Po`ega E. 47 Podany P. 199 Podgornik B. 451, 461, 829, 971 Pohanka M. 17, 183, 207, 575 Pokorný P. 253 Popovi} M. 403, 531 Popovi~ L’. 55 Pospí{il J. 479, 939 Pouranvari S. 663 Prach O. 247, 427 Price B. Y. 923 Prijanovi~ Tonkovi~ M. 101 Prokopiuk vel Prokopowicz M. 649 Puchnin M. 247, 427 Põssl P. 491 R Rabehi B. 911 Radeti} T. 403, 531 Radovanovi} R. V. 89 Rahpaima G. 471 Ram G. D. J. 357 Ranjbarnodeh E. 775 Ravindran D. 951 Reif M. 137 Ribari~ S. 797 Rimpelova S. 331 Romhanji E. 403, 531 Rousta A. 695 Rovnaník P. 7 Ró¿añski M. 945 Rozman J. 797 Rudnytsky M. 523 Rudolf R. 981 Rusnáková S. 819 Rypák P. 7 S Sahin M. 109 Sajdl P. 331 Salvetr P. 917 Savas V. 81 Sedlarik V. 869 Sener B. 961 Sidjanin L. 387 Skalar T. 743 Skoneczny W. 845 Skowronek T. 783 Skubisz P. 783 Sláma P. 903 Slepi~ka P. 331 Sokolar R. 885 Sondor J. 175 Stamboli} A. 981 Staszuk M. 699, 755 ªtefan A. 11, 723 Stetina J. 3 Stoychev D. 689 Su~ik G. 55 Sun Z. 381 Svoboda J. 891 Svobodová P. 447 Szabóová A. 55 Szczerba J. 29 Szindler M. 649 Szklarek R. 699 Szl¹zak K. 189 Szota M. 189 Szymañska J. 735 [ [leichrt J. 413 [najdárek L. 479 [tádler C. 891 [tefane P. 571 [tefková D. 879 [toudek R. 547 [vor~ík V. 331 T Tadi} N. 59 Taktak S. 263 Tañski T. 699, 755 Tanyel Z. 433 Tariqul Islam M. 33 Tehovnik F. 989 Thankaraj Mariapushpam C. J. 951 Tichomirov V. 409, 607 Tim~aková K. 879 Tim~aková-[amárková K. 7, 565 Topal E. 23 Toplan H. Ö. 739 Topoláø L. 7, 879 Torkar M. 451, 461, 469, 829 Tr~ka T. 547 Trobentar B. 517 Trudonoshyn O. 247, 427 Truºcã R. 723 Turgut Y. 591 Turk J. 601 Turkmen M. 39, 641 Ulker S. 263 Uzun E. 395 Uzun G. 275 U V{ianský D. 729 Waele De W. 571 Vajkay F. 409, 607 W Walke W. 153, 323 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 1029 Walther F. 95 Wegrzyn T. 1001 Wiœniewska K. 29 Wiœniewski P. 735 Wiœniowski M. 699 V Verleysen P. 571 Voicu G. 11 Vojtìch D. 447, 917 Von~ina M. 601 Vorel I. 499, 623 Votavová H. 575 Vukovi} N. 47 Vymazal T. 547 Y Yalcinoz M. 39 Yener S. C. 899 Yener T. 899 Yip N. 217 Yýlmaz V. 667 Z Zabihollah A. 851 Zach J. 137 Zafo{nik B. 517 ¯ak A. 783 Zalecki W. 783 Zazi N. 165 Zeli~ K. 805 Zem~ík R. 295, 319 Zengin M. 395 Zeytýn S. 585, 899 Zhang C. 523 Zhang D. F. 381 Zlámal P. 301 Zupan K. 743 Zygmuntowicz J. 537 @ @aludek M. 819 @ivkovi} D. 47 @u`ek B. 989 @u`ek M. C. 797 LETNO KAZALO – INDEX 1030 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY VSEBINSKO KAZALO / SUBJECT INDEX LETNIK / VOLUME 50, 2016, 1–6 Kovinski materiali – Metallic materials Improvement of the casting of special steel with a wide solid-liquid interface Izbolj{anje ulivanja posebnega jekla s {irokim intervalom trdno-teko~e T. Mauder, J. Stetina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 Corrosion of the refractory zirconia metering nozzle due to molten steel and slag Korozija ognjeodporne cirkonske dozirne {obe s staljenim jeklom in `lindro K. Wiœniewska, D. Madej, J. Szczerba. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29 X-ray radiography of AISI 4340-2205 steels welded by friction welding Rentgenski pregled jekel AISI 4340-2205, varjenih s trenjem U. Caligulu, M. Yalcinoz, M. Turkmen, S. Mercan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39 Thermodynamic properties and microstructures of different shape-memory alloys Termodinami~ne lastnosti in mikrostruktura razli~nih zlitin z oblikovnim spominom L. Gomid`elovi}, E. Po`ega, A. Kostov, N. Vukovi}, D. @ivkovi}, D. Manasijevi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47 The relationship between thermal treatment of serpentine and its reactivity Odvisnost med toplotno obdelavo serpentina in njegovo aktivnostjo G. Su~ik, A. Szabóová, L’. Popovi~, D. Hr{ak . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55 Deformations and velocities during the cold rolling of aluminium alloys Deformacija in hitrosti pri hladnem valjanju aluminijevih zlitin M. Mi{ovi}, N. Tadi}, M. Ja}imovi}, M. Janji} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59 Prediction of the chemical non-homogeneity of 30MnVS6 billets with genetic programming Napovedovanje nehomogenosti kemijske sestave pri gredicah 30MnVS6 s pomo~jo genetskega programiranja M. Kova~i~, D. Novak. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69 Effect of the TiBN coating on a HSS drill when drilling the MA8M Mg alloy Vpliv TiBN prevleke na HSS svedru pri vrtanju MA8M Mg zlitine F. Karaca, B. Aksakal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75 Effects of friction-welding parameters on the morphological properties of an Al/Cu bimetallic joint Vpliv parametrov tornega varjenja na morfolo{ke lastnosti Al/Cu bimetalnega spoja V. D. Mila{inovi}, R. V. Radovanovi}, M. D. Mila{inovi}, B. R. Gligorijevi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 89 Optimizing the parameters for friction welding stainless steel to copper parts Optimiranje parametrov pri tornem varjenju nerjavnega jekla na bakrene dele M. Sahin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 109 WEDM cutting of Inconel 718 nickel-based superalloy: effects of cutting parameters on the cutting quality WEDM rezanje nikljeve superzlitine Inconel 718: vpliv parametrov rezanja na kvaliteto rezanja U. Çaydaº, M. Ay . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 117 The effect of EO and steam sterilization on the mechanical and electrochemical properties of titanium Grade 4 Vpliv EO in sterilizacije s paro na mehanske in elektrokemijske lastnosti titana Grade 4 M. Basiaga, W. Walke, Z. Paszenda, A. Kajzer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153 Influence of the carbide-particle spheroidisation process on the microstructure after the quenching and annealing of 100CrMnSi6-4 bearing steel Vpliv procesa sferoidizacije karbidnih delcev na mikrostrukturo jekla 100CrMnSi6-4 za le`aje po kaljenju in popu{~anju J. Dlouhy, D. Hauserova, Z. Novy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 159 Corrosion behavior and the weak-magnetic-field effect of aluminum packaging paper Vpliv {ibkega magnetnega polja na korozijo aluminijeve embala`ne folije N. Zazi, J.-P. Chopart, A. Bilek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 165 Characteristics of the AlTiCrN+DLC coating deposited with a cathodic arc and the PACVD process Zna~ilnosti AlTiCrN+DLC prevleke, nane{ene s katodnim oblokom in PACVD postopkom K. Lukaszkowicz, E. Jonda, J. Sondor, K. Balin, J. Kubacki. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 175 LETNO KAZALO – INDEX Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 1031 Magnetic properties and microstructure of a bulk amorphous Fe61Co10Ti3Y6B20 alloy, fabricated as rods and tubes Magnetne lastnosti in mikrostruktura masivne amorfne zlitine Fe61Co10Ti3Y6B20 v obliki palic in cevi M. Nabia³ek, K. Bloch, K. Szl¹zak, M. Szota . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 189 Recrystallization behaviour of a nickel-based superalloy Obna{anje superzlitine na osnovi niklja pri rekristalizaciji P. Podany, Z. Novy, J. Dlouhy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 199 Use of the ABI technique to measure the mechanical properties of aluminium alloys: effect of chemical composition on the mechanical properties of the alloys Uporaba tehnike ABI za merjenje mehanskih lastnosti aluminijevih zlitin: vpliv kemijske sestave na mehanske lastnosti zlitin M. Puchnin, O. Trudonoshyn, O. Prach . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 247 Fe-Zn intermetallic phases prepared by diffusion annealing and spark-plasma sintering Fe-Zn intermetalne faze, pripravljene z difuzijskim `arjenjem in s sintranjem v iskre~i plazmi P. Pokorný, J. Cinert, Z. Pala . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 253 High-temperature oxidation of silicide-aluminide layer on the TiAl6V4 alloy prepared by liquid-phase siliconizing Visokotemperaturna oksidacija plasti silicid-aluminid, pripravljene s silikoniziranjem s teko~o fazo zlitine TiAl6V4 T. F. Kubatík . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 257 Characterization and kinetics of plasma-paste-borided AISI 316 steel Karakterizacija in kinetika plazma boriranja s pasto jekla AISI 316 R. Chegroune, M. Keddam, Z. Nait Abdellah, S. Ulker, S. Taktak, I. Gunes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 263 The effects of cutting conditions on the cutting torque and tool life in the tapping process for AISI 304 stainless steel Vpliv pogojev rezanja na moment pri rezanju in zdr`ljivost navojnega vreznika pri vrezovanju notranjih navojev v nerjavno jeklo AISI 304 G. Uzun, I. Korkut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 275 Physicochemical properties of a Ti67 alloy after EO and steam sterilization Fizikalno kemijske lastnosti zlitine Ti67 po EO in parni sterilizaciji W. Walke, M. Basiaga, Z. Paszenda, J. Marciniak, P. Karasinski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 323 Analyzing the heat-treatment effect on the mechanical properties of free-cutting steels Analiza vpliva toplotne obdelave na mehanske lastnosti avtomatnih jekel M. K. Kulekci, U. Esme, F. Kahraman, R. Ozgun, A. Akkurt . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 337 Analysis of the cutting temperature and surface roughness during the orthogonal machining of AISI 4140 alloy steel via the Taguchi method Analiza temperature rezanja in hrapavosti povr{ine s Taguchi metodo pri ortogonalni strojni obdelavi legiranega jekla AISI 4140 A. R. Motorcu, Y. Isik, A. Kus, M. C. Cakir . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 343 Weldability of Ti6Al4V to AISI 2205 with a nickel interlayer using friction welding Preizku{anje varivosti pri varjenju s trenjem Ti6Al4V in AISI 2205 z vmesno plastjo niklja I. Kirik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 353 Effect of activated flux and nitrogen addition on the bead geometry of borated stainless-steel GTA welds Vpliv aktiviranega topila in dodatka du{ika na geometrijo kopeli pri GTA zvarih boriranega nerjavnega jekla G. R. Kumar, G. D. J. Ram, S. R. Koteswara Rao . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 357 Microstructural evolution during the transient liquid-phase bonding of dissimilar nickel-based superalloys of IN738LC and NIMONIC 75 Razvoj mikrostrukture med spajanjem s prehodno teko~o fazo neenakih superzlitin na osnovi niklja IN738LC in NIMONIC 75 M. G. Khakian, S. Nategh, S. Mirdamadi. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 365 Workability behaviour of Cu-TiB2 powder-metallurgy preforms during cold upsetting Preoblikovalnost Cu-TiB2 predoblik izdelanih z metalurgijo prahov med hladnim kovalnim preizkusom S. Gadakary, A. Kumar Khanra, M. J. Davidson . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 373 Effects of extrusion shear on the microstructures and a fracture analysis of a magnesium alloy in the homogenized state Vplivi stri`enja med iztiskanjem homogenizirane magnezijeve zlitine na mikrostrukturo in na analizo preloma H. J. Hu, Z. Sun, D. F. Zhang . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 381 FSW welding of Al-Mg alloy plates with increased edge roughness using square pin tools of various shoulder geometries FSW varjenje plo{~ iz Al-Mg zlitine s pove~ano hrapavostjo robov z orodjem s kvadratno konico in razli~no geometrijo bokov S. Balos, L. Sidjanin, M. Dramicanin, D. Labus Zlatanovic, A. Antic. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 387 Homogenization of an Al-Mg alloy and alligatoring failure: alloy ductility and fracture Homogenizacija Al-Mg zlitine in krokodiljenje: duktilnost zlitine in prelom E. Romhanji, T. Radeti}, M. Popovi}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 403 Increasing micro-purity and determining the effects of the production with and without vacuum refining on the qualitative parameters of forged-steel pieces with a high aluminium content Pove~anje mikro~isto~e in dolo~itev u~inka proizvodnje, z vakuumskim rafiniranjem ali brez, na kvalitativne parametre kovanega jekla z visoko vsebnostjo aluminija V. Kurka, J. Pindor, J. Kosòovská, Z. Adolf . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 419 LETNO KAZALO – INDEX 1032 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 Use of the ABI technique to measure the mechanical properties of aluminium alloys: effect of heat-treatment conditions on the mechanical properties of alloys Uporaba ABI tehnike za merjenje mehanskih lastnosti aluminijevih zlitin: vpliv pogojev toplotne obdelave na mehanske lastnosti zlitin O. Trudonoshyn, M. Puchnin, O. Prach . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 427 Investigation of the effect of holding time and melt stirring on the grain refinement of an A206 alloy Preiskava vpliva ~asa zadr`evanja in me{anja taline na zmanj{anje velikosti zrn zlitine A206 N. Akar, Z. Tanyel, K. Kocatepe, R. Kayikci . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 433 Investigating the influence of cutting speed on the tool life of a cutting insert while cutting DIN 1.4301 steel Preiskava vpliva hitrosti rezanja na zdr`ljivost vlo`ka za rezanje pri rezanju jekla DIN 1.4301 R. Dubovská, J. Majerík, R. ^ep, K. Kouøil . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 439 NiAl intermetallic prepared with reactive sintering and subsequent powder-metallurgical plasma-sintering compaction Reakcijsko sintranje in zgo{~evanje s plazemskim sintranjem NiAl intermetalne zlitine A. Michalcová, D. Vojtìch, T. F. Kubatík, P. Novák, P. Dvoøák, P. Svobodová, I. Marek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 447 Microscopic characterization and particle distribution in a cast steel matrix composite Mikroskopska karakterizacija in razporeditev delcev v kompozitu z matrico litega jekla A. Kra~un, M. Torkar, J. Burja, B. Podgornik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 451 A comparison of as-welded and simulated heat affected zone (HAZ) microstructures Primerjava mikrostrukture toplotno vplivanega podro~ja varjenega in simuliranih vzorcev R. Celin, J. Burja, G. Kosec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 455 Degradation of an AISI 304 stainless-steel tank Degradacija rezervoarja iz AISI 304 nerjavnega jekla M. Torkar, I. Paulin, B. Podgornik. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 461 Correlation of the heat-transfer coefficient at sprinkled tube bundle Korelacija koeficienta prenosa toplote pri potresenem snopu cevi P. Kracík, L. [najdárek, M. Lisý, M. Balá{, J. Pospí{il. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 479 Investigation of wear behavior of borided AISI D6 steel Preiskava obrabe boriranega jekla AISI D6 I. Gunes, S. Kanat . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 505 Investigation of Portevin-Le Chatelier effect of hot-rolled Fe-13Mn-0.2C-1Al-1Si TWIP steel Preiskava Portevin-Le Chatelier u~inka pri vro~em valjanju Fe-13Mn-0.2C-1Al-1Si TWIP jekla B. Aydemir, H. Kazdal Zeytin, G. Guven. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 511 Homogenization of an Al-Mg alloy and alligatoring failure: influence of the microstructure Homogenizacija Al-Mg zlitine in krokodiljenje: vpliv mikrostrukture E. Romhanji, T. Radeti}, M. Popovi}. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 531 Static and dynamic tensile characteristics of S420 and IF steel sheets Stati~ne in dinami~ne natezne lastnosti plo~evine iz S420 in IF jekla M. Mihaliková, V. Girman, A. Li{ková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 543 Multi-criteria analysis of synthesis methods for Ni-based catalysts Ve~kriterijska analiza sinteznih metod na osnovi Ni katalizatorja V. Nikoli}, B. Agarski, @. Kamberovi}, Z. An|i}, I. Budak, B. Kosec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 553 Influence of structural defects on the magnetic properties of massive amorphous Fe60Co10Mo2WxY8B20-x (x = 1, 2) alloys produced with the injection casting method Vpliv strukturnih napak na magnetne lastnosti masivne amorfne zlitine Fe60Co10Mo2WxY8B20-x (x = 1, 2), izdelane z metodo litja z vbrizgavanjem J. Gondro, K. B³och, M. Nabia³ek, S. Garus . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 559 Possibilities of NUS and Impact-Echo methods for monitoring steel corrosion in concrete Mo`nosti metod NUS in Udarec-odmev za kontrolo korozije jekla v betonu K. Tim~aková-[amárková, M. Matysík, Z. Chobola . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 565 Characterization of heterogeneous arc welds through miniature tensile testing and Vickers-hardness mapping Karakterizacija heterogenih zvarov s pomo~jo miniaturnih nateznih preizkusov in matri~nimi meritvami trdote po Vickersu S. Hertelé, J. Bally, N. Gubeljak, P. [tefane, P. Verleysen, W. De Waele. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 571 Overcooling in overlap areas during hydraulic descaling Podhladitev in prekrivanje podro~ij med hidravli~nim raz{kajanjem M. Pohanka, H. Votavová . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 575 Development of aluminium alloys for aerosol cans Razvoj aluminijevih zlitin za aerosol doze S. Kores, J. Turk, J. Medved, M. Von~ina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 601 LETNO KAZALO – INDEX Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 1033 Microstructure and properties of the high-temperature (HAZ) of thermo-mechanically treated S700MC high-yield-strength steel Mikrostruktura in lastnosti visoko temperaturnega obmo~ja zvara (HAZ) termo-mehansko obdelanega jekla S700MC z visoko mejo plasti~nosti J. Górka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 617 New concept for manufacturing closed die forgings of high strength steels Nov koncept izdelave odkovkov iz visokotrdnostnih jekel v zaprtih utopih K. Ibrahim, I. Vorel, D. Bublíková, B. Ma{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 623 Effect of the addition of niobium and aluminium on the microstructures and mechanical properties of micro-alloyed PM steels Vpliv dodatka niobija in aluminija na mikrostrukturo in mehanske lastnosti mikrolegiranih PM jekel S. Gündüz, M. A. Erden, H. Karabulut, M. Türkmen . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 641 The effect of the welding parameters and the coupling agent on the welding of composites Vpliv parametrov varjenja in sredstva za spajanje na varjenje kompozitov S. E. Erdogan, U. Huner . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 655 Investigation of hole profiles in deep micro-hole drilling of AISI 420 stainless steel using powder-mixed dielectric fluids Preiskava profilov luknje pri globokem vrtanju mikroluknje v AISI 420 nerjavnem jeklu s pomo~jo dielektri~ne teko~ine s prime{anim prahom V. Yýlmaz . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 667 The phenomenon of reduced plasticity in low-alloyed copper Pojav zmanj{anja plasti~nosti malo legiranega bakra W. Ozgowicz, E. Kalinowska-Ozgowicz, B. Grzegorczyk, K. Lenik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 677 Investigation of the mechanical properties of electrochemically deposited Au-In alloy films using nano-indentation Preiskava mehanskih lastnosti elektrokemijsko nane{enega filma zlitine Au-In z nanovtiskovanjem S. Cherneva, R. Iankov, M. Georgiev, T. Dobrovolska, D. Stoychev . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 689 Surface treatment of heat-treated cast magnesium and aluminium alloys Obdelava povr{ine toplotno obdelanih magnezijevih in aluminijevih livnih zlitin T. Tañski, M. Wiœniowski, W. Matysiak, M. Staszuk, R. Szklarek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 699 Analysis of the structural-defect influence on the magnetization process in and above the Rayleigh region Analiza vpliva strukturnih defektov na proces magnetizacije v in nad Rayleigh podro~jem K. Gruszka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 707 Effect of sulphide inclusions on the pitting-corrosion behaviour of high-Mn steels in chloride and alkaline solutions Vpliv sulfidnih vklju~kov na jami~asto korozijo jekel z visoko vsebnostjo Mn v raztopinah kloridov in alkalij A. Grajcar, A. P³achciñska . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 713 Influence of Na2SiF6 on the surface morphology and corrosion resistance of an AM60 magnesium alloy coated by micro arc oxidation Vpliv Na2SiF6 na morfologijo povr{ine in korozijsko odpornost magnezijeve zlitine AM60, prekrite z mikrooblo~no oksidacijo A. Ayday . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 719 Investigation studies involving wear-resistant ALD/PVD hybrid coatings on sintered tool substrates Preiskave obrabne odpornosti hibridnega nanosa ALD/PVD na sintranem orodju M. Staszuk, D. Paku³a, T. Tañski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 755 Dissimilar spot welding of DQSK/DP600 steels: the weld-nugget growth To~kasto varjenje jekel DQSK/DP600: rast jedra zvara S. P. Hoveida Marashi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 761 Armour plates from Kozlov rob – analyses of two unusual finds Oklepni plo{~i s Kozlovega roba – analize dveh nenavadnih najdb T. Lazar, P. Mrvar, M. Lamut, P. Fajfar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 767 Numerical and experimental investigation of the effect of hydrostatic pressure on the residual stress in boiler-tube welds Numeri~na in eksperimentalna preiskava vpliva hidrostati~nega tlaka na zaostale napetosti v zvaru na kotlovski cevi D. Danyali, E. Ranjbarnodeh . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 775 Effect of direct cooling conditions on the microstructure and properties of hot-forged HSLA steels for mining applications Vpliv pogojev ohlajanja na mikrostrukturo in lastnosti vro~e kovanih HSLA jekel za uporabo v rudarstvu P. Skubisz, £. Lisiecki, T. Skowronek, A. ¯ak, W. Zalecki . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 783 Influence of the tool rotational speed on the microstructure and joint strength of friction-stir spot-welded pure copper Vpliv hitrosti vrtenja orodja na mikrostrukturo in trdnost torno vrtilno to~kasto zvarjenega spoja ~istega bakra I. Dinaharan, E. T. Akinlabi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 791 Influence of different production processes on the biodegradability of an FeMn17 alloy Vpliv razli~nih procesov izdelave na biorazgradljivost zlitine FeMn17 A. Kocijan, I. Paulin, ^. Donik, M. Ho~evar, K. Zeli~, M. Godec. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 805 LETNO KAZALO – INDEX 1034 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 Influence of solidification speed on the structure and magnetic properties of Nd10Fe81Zr1B6 in the as-cast state Vpliv hitrosti strjevanja na strukturo in magnetne lastnosti zlitine Nd10Fe81Zr1B6 v litem stanju M. Doœpia³, M. Nabia³ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 823 Metalografska preiskava in korozijska odpornost zvarov feritnega nerjavnega jekla Metallographic investigation and corrosion resistance of welds of ferritic stainless steels M. Torkar, A. Kocijan, R. Celin, J. Burja, B. Podgornik. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 829 The microstructure of metastable austenite in X5CrNi18-10 steel after its strain-induced martensitic transformation Mikrostruktura metastabilnega avstenita po pretvorbi v napetostno inducirani martenzit v jeklu X5CrNi18-10 A. Kurc-Lisiecka, W. Ozgowicz, E. Kalinowska-Ozgowicz, W. Maziarz . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 837 Analysis of twin-roll casting AA8079 alloy 6.35-μm foil rolling process Analiza procesa valjanja 6,35 μm folije iz zlitine AA8079 ulite med dvema valjema A. Can, H. Arikan, K. Çýnar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 861 Behaviour of new ods alloys under single and multiple deformation Obna{anje novih ods zlitin pri enojni in ve~kratni deformaciji B. Ma{ek, O. Khalaj, Z. Nový, T. Kubina, H. Jirkova, J. Svoboda, C. [tádler . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 891 Electromagnetic-shielding effectiveness and fracture behavior of laminated (Ni–NiAl3) composites U~inkovitost elektromagnetne za{~ite in obna{anje pri lomu laminiranega kompozita (Ni-NiAl3) T. Yener, S. C. Yener, S. Zeytýn . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 899 Effect of thermomechanical treatment on the intergranular corrosion of Al-Mg-Si-Type alloy bars Vpliv termomehanske predelave na interkristalno korozijo palic iz zlitin Al-Mg-Si P. Sláma, J. Nacházel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 903 Porous magnesium alloys prepared by powder metallurgy Porozne magnezijeve zlitine, izdelane s pomo~jo metalurgije prahov P. Salvetr, P. Novák, D. Vojtìch . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 917 The size effect of heat-transfer surfaces on boiling Vpliv velikosti povr{in, ki prena{ajo toploto na vrenje P. Kracík, M. Balas, M. Lisy, J. Pospí{il . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 939 Effect of gas atmosphere on the non-metallic inclusions in laser-welded trip steel with Al and Si additions Vpliv plinske atmosfere na nekovinske vklju~ke v lasersko varjenem trip jeklu z dodatkom Al in Si A. Grajcar, M. Ró¿añski, M. Kamiñska, B. Grzegorczyk . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 945 Machining parameters influencing in electro chemical machining on AA6061 MMC Parametri strojne obdelave, ki vplivajo na elektrokemijsko strojno obdelavo AA6061 MMC C. J. Thankaraj Mariapushpam, D. Ravindran, M. D. Anand . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 951 Modeling the deep drawing of an AISI 304 stainless-steel rectangular cup using the finite-element method and an experimental validation Modeliranje globokega vleka pravokotne ~a{e iz AISI 304 nerjavnega jekla z metodo kon~nih elementov in z eksperimentalnim preverjanjem B. Sener, H. Kurtaran . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 961 Electroslag remelting: A process overview Elektropretaljevanje pod `lindro – pregled procesa B. Arh, B. Podgornik, J. Burja. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 971 Continuous vertical casting of a NiTi alloy Vertikalno kontinuirno litje NiTi zlitine A. Stamboli}, I. An`el, G. Lojen, A. Kocijan, M. Jenko, R. Rudolf . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 981 Hot tensile testing of SAF 2205 duplex stainless steel Vro~i natezni preskusi dupleks nerjavnega jekla SAF 2205 F. Tehovnik, B. @u`ek, J. Burja . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 989 Impact toughness of WMD after MAG welding with micro-jet cooling Udarna `ilavost WMD po MAG varjenju z mikro-jet hlajenjem T. Wegrzyn, J. Piwnik, A. Borek, A. Kurc-Lisiecka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1001 Forming-limit diagrams and strain-rate-dependent mechanical properties of AA6019-T4 and AA6061-T4 aluminium sheet materials Mejni diagrami preoblikovanja in odvisnost mehanskih lastnosti od hitrosti preoblikovanja aluminijevih plo~evin iz AA6019-T4 in AA6061-T4 O. Çavuºoðlu, A. G. Leacock, H. Gürün . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1005 Effect of alternative heat-treatment parameters on the aging behavior of short-fiber-reinforced 2124 Al composites Vpliv alternativnih parametrov toplotne obdelave na staranje 2124 Al kompozita, oja~anega s kratkimi vlakni Y. Altunpak, S. Aslan, M. Oðuz Güler, H. Akbulut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1011 LETNO KAZALO – INDEX Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 1035 Anorganski materiali – Inorganic materials Non-traditional non-destructive testing of the alkali-activated slag mortar during the hardening Netradicionalno neporu{no preizku{anje z alkalijami aktivirane malte med strjevanjem L. Topoláø, P. Rypák, K. Tim~aková-[amárková, L. Pazdera, P. Rovnaník . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 Investigation of hole effects on the critical buckling load of laminated composite plates Preiskava vpliva luknje na kriti~no upogibno obremenitev laminiranih kompozitnih plo{~ A. Kurºun, E. Topal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23 Characterisation of the mechanical and corrosive properties of newly developed glass-steel composites Karakterizacija mehanskih in korozijskih lastnosti novo razvitih kompozitov steklo-jeklo O. Lyubimova, E. Gridasova, A. Gridasov, G. Frieling, M. Klein, F. Walther . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 95 Phase analysis of the slag after submerged-arc welding Analiza faz v `lindri pri oblo~nem varjenju pod pra{kom M. Prijanovi~ Tonkovi~, J. Lamut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 101 Composites based on inorganic matrices for extreme exposure conditions Kompoziti z anorgansko osnovo za izpostavitev ekstremnim razmeram A. Dufka, T. Melichar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 147 Enhanced stability and electrochemical performance of a BaTiO3/PbO2 electrode via a layer obtained with layer electrodeposition Izbolj{ana stabilnost in elektrokemijska zmogljivost elektrode BaTiO3/PbO2, izdelane z elektrodepozicijo plast na plast G. Muthuraman, K. Karunakaran, I. S. Moon. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 211 Deformation behaviour of amorphous Fe-Ni-W/Ni bilayer-confined bulk metallic glasses Obna{anje deformiranega, amorfnega, na dve plasti omejenega kovinskega stekla Fe-Ni-W/Ni H. K. Lau, N. Yip, S. H. Chen, W. Chen, K. C. Chan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 217 Synergistic effect of organic- and ceramic-based ingredients on the tribological characteristics of brake friction materials Sinergisti~en vpliv sestavin z organsko in kerami~no osnovo na tribolo{ke zna~ilnosti materialov za torne zavore R. Ertan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 223 Chemical synthesis and densification behavior of Ag/ZnO metal-matrix composites Obna{anje Ag/ZnO kompozita s kovinsko osnovo pri kemijski sintezi in zgo{~evanju M. Ardestani . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 281 Tensile and compressive tests of textile composites and results analysis Natezni in tla~ni preizkusi tekstilnih kompozitov in analiza rezultatov K. Kunc, T. Kroupa, R. Zem~ík, J. Krystek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 295 Printed microstrip line-fed patch antenna on a high-dielectric material for C-band applications Tiskana mikrotrakasta linijsko napajana krpasta antena na visoko dielektri~nem materialu za uporabo v C-pasu Md. M. Islam, M. R. I. Faruque, M. F. Mansor, M. T. Islam. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 307 Compressive properties of auxetic structures produced with direct 3D printing Stiskanje struktur materialov z negativnim Poissonovim razmerjem, proizvedenih z neposrednim tridimenzionalnim tiskanjem P. Koudelka, O. Jirou{ek, T. Fíla, T. Doktor . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 311 Improvement of selective copper extraction from a heat-treated chalcopyrite concentrate with atmospheric sulphuric-acid leaching Izbolj{anje selektivne ekstrakcije bakra iz toplotno obdelanega koncentrata halkopirita z lu`enjem z `vepleno kislino na zraku E. Uzun, M. Zengin, Ý. Atýlgan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 395 Assessment of tubular light guides with respect to building physics Ocena cevastih vodnikov svetlobe glede na gradbeno fiziko F. Vajkay, D. Be~kovský, V. Tichomirov . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 409 Creep behaviour of a short-fibre C/PPS composite Vedenje kratkih vlaken C/PPS kompozitov pri lezenju T. Fíla, P. Koudelka, D. Kytýø, J. Hos, J. [leichrt. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 413 Vapour-phase condensed composite materials based on copper and carbon Kompoziti na osnovi bakra in ogljika, kondenzirani iz plinske faze V. Bukhanovsky, M. Rudnytsky, M. Grechanyuk, R. Minakova, C. Zhang . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 523 Metal particles size influence on graded structure in composite Al2O3-Ni Vpliv velikosti kovinskih delcev na gradientno strukturo kompozita Al2O3-Ni J. Zygmuntowicz, A. Miazga, K. Konopka, W. Kaszuwara . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 537 Fabrication and properties of SiC reinforced copper-matrix-composite contact material Izdelava in lastnosti s SiC utrjenega kompozitnega materiala na osnovi bakra G. F. Celebi Efe, M. Ýpek, S. Zeytin, C. Bindal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 585 LETNO KAZALO – INDEX 1036 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 Impression relaxation and creep behavior of Al/SiC nanocomposite Sprostitev vtisa in obna{anje Al/SiC nanokompozita pri lezenju Y. S. Kakhki, S. Nategh, T. S. Mirdamadi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 611 Helium atom scattering – a versatile technique in studying nanostructures Sipanje atomov helija – vsestranska tehnika za {tudij nanostruktur G. Bavdek, D. Cvetko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 627 Characteristics of dye-sensitized solar cells with carbon nanomaterials Zna~ilnosti na fiksirano barvo ob~utljivih solarnih celic z ogljikovimi nanomateriali L. A. Dobrzañski, A. Mucha, M. Prokopiuk vel Prokopowicz, M. Szindler, A. Dryga³a, K. Lukaszkowicz . . . . . . . . . . . . . . . . . . . . . . . 649 The effect of high-speed grinding technology on the properties of fly ash Vpliv tehnologije hitrega mletja na lastnosti lete~ega pepela K. Dvoøák, I. Hájková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 683 Growth of K2CO3-doped KDP crystal from an aqueous solution and an investigation of its physical properties Rast KDP kristalov z dodatkom K2CO3 iz vodne raztopine in preiskava njihovih fizikalnih lastnosti A. Rousta, H. R. Dizaji . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 695 Rheological properties of alumina ceramic slurries for ceramic shell-mould fabrication Reolo{ke lastnosti go{~e iz glinice za izdelavo kerami~nih kalupov J. Szymañska, P. Wiœniewski, M. Ma³ek, J. Mizera . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 735 Effect of mechanical activation on the synthesis of a magnesium aluminate spinel Vpliv mehanske aktivacije na sintezo magnezij-aluminatnega {pinela D. Kýrsever, N. K. Karabulut, N. Canikoðlu, H. Ö. Toplan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 739 Phase and microstructure development of LSCM perovskite materials for SOFC anodes prepared by the carbonate-coprecipitation method Razvoj kristalnih faz in mikrostrukture LSCM perovskitnih materialov za SOFC anode, pripravljenih s karbonatno metodo koprecipitacije K. Zupan, M. Marin{ek, T. Skalar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 743 The structure and morphology of the surface of duplex layers after saturation of the base layer with carbon Struktura in morfologija povr{ine dupleks plasti po nasi~enju osnovne plasti z ogljikom W. Skoneczny . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 845 A new wideband negative-refractive-index metamaterial Novi {irokopasovni metamaterial z negativnim lomnim koli~nikom S. S. Islam, M. R. Iqbal Faruque, M. J. Hossain, M. T. Islam. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 873 Non-traditional whiteware based on calcium aluminate cement Netradicionalni porcelan na osnovi kalcij aluminatnega cementa R. Sokolar. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 885 Surface and anticorrosion properties of hydrophobic and hydrophilic TiO2 coatings on a stainless-steel substrate Povr{inske in protikorozijske lastnosti hidrofobnih in hidrofilnih TiO2 prevlek na jekleni podlagi M. Conradi, A. Kocijan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 967 A high-efficiency automatic de-bubbling system for liquid silicone rubber Visokozmogljiv sistem za odpravljanje mehur~kov v teko~i silikonski gumi C.-C. Kuo, C.-M. Huang . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 995 Organski materiali – Organic materials Deformation behaviour of a natural-shaped bone scaffold Obna{anje naravno oblikovanega ogrodja kosti pri deformaciji D. Kytýø, T. Doktor, O. Jirou{ek, T. Fíla, P. Koudelka, P. Zlámal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 301 Investigation of the mechanical properties of a cork/rubber composite Raziskava mehanskih lastnosti kompozita pluta/guma R. Kottner, J. Kocáb, J. Heczko, J. Krystek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 579 Measurement of bio-impedance on an isolated rat sciatic nerve obtained with specific current stimulating pulses Meritev bioimpedance na izoliranem `ivcu Ischiadicus pri podgani, vzbujenem s posebnimi tokovnimi stimulacijskimi impulzi J. Rozman, M. C. @u`ek, R. Frange`, S. Ribari~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 797 Polimeri – Polymers Effects of an epoxy-resin-fiber substrate for a -shaped microstrip antenna Vpliv z vlakni oja~ane epoksi podlage pri -obliki mikrotrakaste antene Md. M. Islam, M. R. I. Faruque, M. Tariqul Islam, H. Arshad . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33 LETNO KAZALO – INDEX Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 1037 Effect of the skin-core morphology on the mechanical properties of injection-moulded parts Vpliv morfologije skorja-jedro na mehanske lastnosti vbrizganih delov E. Hnatkova, Z. Dvorak. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 195 Surface properties of a laser-treated biopolymer Lastnosti povr{ine biopolimera, obdelanega z laserjem I. Michaljani~ova, P. Slepi~ka, S. Rimpelova, P. Sajdl, V. [vor~ík . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 331 Organosoluble xanthone-based polyimides: synthesis, characterization, antioxidant activity and heavy-metal sorption Organsko topni poliamidi na osnovi ksantona: sinteza, karakterizacija, antioksidativna aktivnost in sorpcija te`kih kovin M. M. Lakouraj, G. Rahpaima, R. Azimi . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 471 The influence of surface coatings on the tooth tip deflection of polymer gears Vpliv povr{inskih prevlek na poves vrha zoba polimernih zobnikov B. Trobentar, S. Glode`, J. Fla{ker, B. Zafo{nik. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 517 Investigation of the cutting forces and surface roughness in milling carbon-fiber-reinforced polymer composite material Preiskava sil rezanja in hrapavosti povr{ine pri rezkanju kompozitnega polimernega materiala, oja~anega z ogljikovimi vlakni S. Bayraktar, Y. Turgut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 591 Chemical cross-linking of chitosan/polyvinyl alcohol electrospun nanofibers Kemijsko zamre`enje elektro spredenih nanovlaken iz hitosan/polivinil alkohola S. Pouranvari, F. Ebrahimi, G. Javadi, B. Maddah . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 663 Mechanical properties of polyamide/carbon-fiber-fabric composites Mehanske lastnosti kompozitne tkanine iz poliamid/ogljikovih vlaken C.-E. Pelin, G. Pelin, A. ªtefan, E. Andronescu, I. Dincã, A. Ficai, R. Truºcã . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 723 Cutting-tool performance in the end milling of carbon-fiber-reinforced plastics Zmogljivost rezilnega orodja pri rezkanju plastike, oja~ane z ogljikovimi vlakni O. Bílek, S. Rusnáková, M. @aludek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 819 Antimicrobial modification of polypropylene with silver nanoparticles immobilized on zinc stearate Protimikrobno spreminjanje polipropilena z nanodelci srebra, imobiliziranih na cinkovem stearatu G. Jandikova, P. Holcapkova, M. Hrabalikova, M. Machovsky, V. Sedlarik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 869 Nanomateriali in nanotehnologije – Nanomaterials and nanotechnology Multi-walled carbon nanotubes effect in polypropylene nanocomposites Vpliv ve~stenskih ogljikovih nanocevk v nanokompozitih iz polipropilena C. E. Ban, A. Stefan, I. Dinca, G. Pelin, A. Ficai, E. Andronescu, O. Oprea, G. Voicu . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11 Investigation of the adhesion and wear properties of borided AISI H10 steel Preiskava adhezije in obrabnih lastnosti boriranega jekla AISI H10 I. Gunes, M. Ozcatal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 269 Influence of nano-sized cobalt oxide additions on the structural and electrical properties of nickel-manganite-based NTC thermistors Vpliv dodatka nanodelcev kobaltovega oksida na zgradbo in elektri~ne lastnosti ntc termistorjev na osnovi nikljevega manganita G. Hardal, B. Y. Price . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 923 Gradbeni materiali – Materials in civil engineering Influence of dredged sediment on the shrinkage behavior of self-compacting concrete Vpliv izkopanih sedimentov na kr~enje samozgo{~evalnega betona N. E. Bouhamou, F. Mostefa, A. Mebrouki, K. Bendani, N. Belas . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 127 Study of the properties and hygrothermal behaviour of alternative insulation materials based on natural fibres [tudij lastnosti in higrotermalno obna{anje alternativnih izolacijskih materialov na osnovi naravnih vlaken J. Zach, M. Reif, J. Hroudová . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 137 Acoustic and electromagnetic emission of lightweight concrete with polypropylene fibers Akusti~na in elektromagnetna emisija lahkega betona s polipropilenskimi vlakni R. [toudek, T. Tr~ka, M. Matysík, T. Vymazal, I. Pl{ková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 547 Evaluation of the grindability of recycled glass in the production of blended cements Ocena sposobnosti drobljenja recikliranega stekla pri proizvodnji me{anih cementov K. Dvoøák, D. Dolák, D. V{ianský, P. Dobrovolný . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 729 Artificial aggregate from sintered coal ash Umetni agregat iz sintranega pepela premoga V. Cerny, R. Drochytka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 749 LETNO KAZALO – INDEX 1038 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 Effect of a combination of fly ash and shrinkage-reducing additives on the properties of alkali-activated slag-based mortars Vpliv kombinacije lete~ega pepela in dodatka za zmanj{anje kr~enja na lastnosti malte iz z alkalijami aktivirane `lindre V. Bílek, L. Kalina, J. Koplík, M. Mon~eková, R. Novotný . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 813 Evaluation of the degree of degradation using the impact-echo method in civil engineering Ocena stopnje degradacije v gradbeni{tvu z uporabo metode odmeva zvo~nih valov D. [tefková, K. Tim~aková, L. Topoláø, P. Cikrle. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 879 Valorization of brick wastes in the fabrication of concrete blocks Ocena odpadkov iz opeke pri proizvodnji betonskih zidakov Y. Ghernouti, B. Rabehi, T. Bouziani, R. Chaid . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 911 Durability of alumina silicate concrete based on slag/fly-ash blends against acid and chloride environments Zdr`ljivost betona na osnovi glinice in silikatov iz me{anice `lindra/lete~i pepel na kislo in kloridno okolje R. Gopalakrishnan, K. Chinnaraju . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 929 Numeri~ne metode – Numerical methods Experimental and numerical study of hot-steel-plate flatness Eksperimentalni in numeri~ni {tudij ravnosti vro~ih plo{~ iz jekla J. Hrabovský, M. Pohanka, P. J. Lee, J. H. Kang . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17 Application of the Taguchi method to select the optimum cutting parameters for tangential cylindrical grinding of AISI D3 tool steel Uporaba Taguchi metode za izbiro optimalnih parametrov odrezavanja pri tangencialnem cilindri~nem bru{enju orodnega jekla AISI D3 C. Ozay, H. Ballikaya, V. Savas. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 81 Prediction of the elastic moduli of chicken-feather-reinforced PLA and a comparison with experimental results Napovedovanje modulov elasti~nosti PLA, oja~anega s pi{~an~jim perjem in primerjava z eksperimentalnimi rezultati U. Özmen, B. Okutan Baba . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 141 Implicit numerical multidimensional heat-conduction algorithm parallelization and acceleration on a graphics card Paralelizacija in pospe{itev implicitnega numeri~nega ve~dimenzijskega algoritma prevajanja toplote na grafi~ni kartici M. Pohanka, J. Ondrou{ková . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 183 Estimation of the number of forward time steps for the sequential Beck approach used for solving inverse heat-conduction problems Ugotavljanje {tevila vnaprej{njih ~asovnih korakov za sekven~ni Beckov pribli`ek pri re{evanju problemov inverzne toplotne prevodnosti J. Komínek, M. Pohanka . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 207 Optimization of the parameters for the surfactant-added EDM of a Ti–6Al–4V alloy using the GRA-Taguchi method Optimizacija povr{insko aktivnih me{anih EDM parametrov na Ti-6Al-4V zlitini z uporabo GRA-Taguchi metode M. Kolli, K. Adepu . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 229 Determination of the cutting-tool performance of high-alloyed white cast iron (Ni-Hard 4) using the Taguchi method Dolo~anje zmogljivosti rezalnih orodij na mo~no legiranem belem litem `elezu (Ni-Hard 4) z uporabo Taguchi metode D. Kir, H. Öktem, M. Çöl, F. Gül Koç, F. Erzincanli . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 239 Experimental verifications and numerical thermal simulations of automobile lamps Eksperimentalna preverjanja in numeri~ne toplotne simulacije avtomobilskih `arometov M. Guzej, J. Horsky . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 289 Model of progressive failure for composite materials using the 3D Puck failure criterion Model postopnega popu{~anja kompozitnega materiala z uporabo Puckovega tridimenzionalnega kriterija poru{itve L. Bek, R. Zem~ík . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 319 Mathematical modeling of a cement raw-material blending process using a neural network Matemati~no modeliranje postopka me{anja sestavin cementa s pomo~jo nevronske mre`e A. Egrisogut Tiryaki, R. Kozan, N. Gokhan Adar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 485 Possibilities of determining the air-pore content in cement composites using computed tomography and other methods Mo`nosti dolo~anja vsebnosti zra~nih por v cementnih kompozitih z uporabo ra~unalni{ke tomografije in drugih metod B. Moravcová, P. Põssl, P. Misák, M. Bla`ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 491 Material and technological modelling of closed-die forging Materialno-tehnolo{ko modeliranje kovanja v zaprtem utopu I. Vorel, [. Jení~ek, H. Jirková, B. Ma{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 499 Computer tools to determine physical parameters in wooden houses Dolo~anje fizikalnih parametrov z ra~unalni{kimi orodji v lesenih hi{ah D. Be~kovský, F. Vajkay, V. Tichomirov . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 607 LETNO KAZALO – INDEX Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040 1039 Modeling of shot-peening effects on the surface properties of a (TiB + TiC)/Ti–6Al–4V composite employing artificial neural networks Modeliranje vpliva hladnega povr{inskega kovanja na lastnosti povr{ine (TiB + TiC)/Ti-6Al-4V kompozita s pomo~jo umetnih nevronskih mre` E. Maleki, A. Zabihollah . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 851 Editor's Preface – Predgovor urednika Predgovor urednika/Editor’s preface M. Torkar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 469 Predgovor urednika/Editor’s preface P. J. McGuiness. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 639 Letno kazalo – Index Letnik 50 (2016), 1–6 – Volume 50 (2016), 1–6 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 995 LETNO KAZALO – INDEX 1040 Materiali in tehnologije / Materials and technology 50 (2016) 6, 1017–1040