VSEBINA – CONTENTS IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES Identification of the material parameters of a unidirectional fiber composite using a micromodel Identifikacija parametrov materiala enosmernega kompozita z uporabo mikromodela H. Srbová, T. Kroupa, R. Zem~ík . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 431 Microstructure and mechanical properties of carbon/carbon-silicon carbide composites prepared by sol-gel processing Mikrostruktura in mehanske lastnosti kompozitov ogljik/ogljik-silicijev karbid, pripravljenih po sol-gel metodi K. Krnel, Z. Stadler, T. Kosma~. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 435 Study of the microstructure and oxidation behavior of YSZ and YSZ/Al2O3 TBCs with HVOF bond coatings [tudij mikrostrukture in vedenja pri oksidaciji YSZ in YSZ/Al2O3 TBC z HVOF naneseno za{~itno prevleko A. C. Karaoðlanlý, G. Erdoðan, Y. Kahraman, A. Türk, F. Üstel, Ý. Özdemir. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 439 Microstructure development of the Ni-GDC anode material for IT-SOFC Razvoj mikrostrukture Ni-GDC anodnega materiala za srednjetemperaturne SOFC K. Zupan, M. Marin{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 445 Modeling of PM10 emission with genetic programming Modeliranje emisije PM10 z genetskim programiranjem M. Kova~i~, S. Sen~i~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 453 Effect of tempering on the room-temperature mechanical properties of X20CrMoV121 and P91 steels Vpliv popu{~anja na mehanske lastnosti jekel X20CrMoV121 in P91 pri sobni temperaturi F. Kafexhiu, F. Vodopivec, J. Vojvodi~ Tuma . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 459 Structure and properties of AlMgSi alloys after ECAP and POST-ECAP ageing Struktura in lastnosti zlitin AlMgSi, staranih pred ECAP in po njem M. Fujda, M. Matvija, T. Kva~kaj, O. Milkovi~, P. Zubko, K. Nagyová . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 465 Application of a Taguchi-based neural network for forecasting and optimization of the surface roughness in a wire-electrical-discharge machining process Uporaba Taguchijeve nevronske mre`e za napovedovanje in optimiranje povr{inske hrapavosti pri postopku `i~ne erozije Y. Kazancoglu, U. Esme, M. K. Kulekci, F. Kahraman, R. Samur, A. Akkurt, M. P. Ipekci . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 471 Prediction of the thermodynamic properties for liquid Al-Mg-Zn alloys Napovedovanje termodinami~nih lastnosti teko~e zlitine Al-Mg-Zn D. @ivkovi}, Y. Du, Lj. Balanovi}, D. Manasijevi}, D. Mini}, N. Talijan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 477 Friction-stir welding of aluminium alloy 5083 Varjenje s trenjem in me{anjem aluminijeve zlitine 5083 D. Klob~ar, L. Kosec, A. Pietras, A. Smolej . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 483 Influence of segregations on the fracture toughness KIc of high-strength spring steel Vpliv izcej na lomno `ilavost KIc visokotrdnostnega vzmetnega jekla B. Sen~i~, V. Leskov{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 489 Mechanical and tribological characteristics of stir-cast Al-Si10Mg and self-lubricating Al-Si10Mg/MoS2 composites Mehanske in tribolo{ke lastnosti z me{anjem ulitih kompozitov Al-Si10Mg in samomazalnih kompozitov Al-Si10Mg/MoS2 K. Somasundara Vinoth, R. Subramanian, S. Dharmalingam, B. Anandavel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 497 Computer-aided modeling of the rubber-pad forming process Ra~unalni{ko modeliranje preoblikovalnega procesa z vmesnikom iz gume M. Benisa, B. Babic, A. Grbovic, Z. Stefanovic . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 503 STROKOVNI ^LANKI – PROFESSIONAL ARTICLES Effect of fly-ash amount and cement type on the corrosion performance of the steel embedded in concrete U~inek koli~ine lete~ega pepela in vrste cementa na korozijo jekla v betonu A. R. Boða, Ý. B. Topçu, M. Öztürk . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 511 ISSN 1580-2949 UDK 669+666+678+53 MTAEC9, 46(5)429–554(2012) MATER. TEHNOL. LETNIK VOLUME 46 [TEV. NO. 5 STR. P. 429–554 LJUBLJANA SLOVENIJA SEP.–OCT. 2012 Effect of the delta-ferrite content on the tensile properties in Nitronic 60 steel at room temperature and 750 °C Vpliv vsebnosti delta ferita na natezne lastnosti jekla Nitronic 60 pri sobni temperaturi in pri 750 °C A. Gigovi}-Geki}, M. Oru~, S. Muhamedagi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 519 Physical regularities in the cracking of nanocoatings and a method for an automated determination of the crack-network parameters Fizikalne zakonitosti pokanja nanoprevlek in metoda za avtomatsko dolo~evanje parametrov mre`e razpok P. Maruschak, V. Gliha, I. Konovalenko, T. Vuherer, S. Panin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 525 Laboratory assessment of micro-encapsulated phase-change materials Laboratorijska ocena mikroenkapsuliranih materialov s fazno premeno M. Ostrý, R. Pøikryl, P. Charvát, T. Ml~och, B. Bakajová. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 531 Content of Cr and Cr (VI) in a welding fume by different Cr content in an experimental coating of a Cr-Ni rutile electrode Vsebnost Cr in Cr (VI) v varilnem dimu pri razli~ni vsebnosti Cr v pla{~u rutilne elektrode Cr-Ni R. Begi}, M. Jenko, M. Godec, ^. Donik. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 535 Use of a two-dimensional pseudo-homogeneous model for the study of temperature and conversion profiles during a polymerization reaction in a tubular chemical reactor Uporaba dvodimenzionalnega psevdohomogenega modela za {tudij temperature in profila pretvorbe med reakcijo polimerizacije v cevastem kemijskem reaktorju M. Marghsi, D. Benachour . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 539 Theoretical and experimental estimation of the working life of machine parts hard faced with austenite-manganese electrodes Teoreti~no in eksperimentalno ugotavljanje zdr`ljivosti strojnih delov, opla{~enih s trdimi avstenitno-manganskimi elektrodami V. Lazi}, A. Sedmak, D. Milosavljevi}, I. Nikoli}, S. Aleksandrovi}, R. Nikoli}, M. Mutavd`i} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 547 H. SRBOVÁ et al.: IDENTIFICATION OF THE MATERIAL PARAMETERS OF A UNIDIRECTIONAL FIBER ... IDENTIFICATION OF THE MATERIAL PARAMETERS OF A UNIDIRECTIONAL FIBER COMPOSITE USING A MICROMODEL IDENTIFIKACIJA PARAMETROV MATERIALA ENOSMERNEGA KOMPOZITA Z UPORABO MIKROMODELA Hana Srbová, Tomá{ Kroupa, Robert Zem~ík University of West Bohemia in Pilsen, Department of Mechanics, Univerzitní 22, 306 14 Plzeò, Czech Republic hsrbova@kme.zcu.cz Prejem rokopisa – received: 2011-10-20; sprejem za objavo – accepted for publication: 2012-02-14 The paper is focused on the identification of material parameters of the substituents of an unidirectional carbon-epoxy long-fiber-reinforced composite. Simple tensile tests using thin coupons with various fiber orientations were performed and force-displacement diagrams were obtained. A model of a unit cell is created in MSC.Marc. Fibers are considered to form a non-linear, elastic, transversely isotropic material and the matrix is considered to be an elasto-plastic isotropic material. The unit cell is loaded by a uniaxial stress up to the same level of loadings as the experimental samples. The sum of the squared differences of displacements between the numerically and experimentally obtained force-displacement diagrams is minimized within an identification process. The parameters of the linear relation between the Young’s modulus of fibers and strain in the fiber-axis direction, and three shape coefficients of the matrix work-hardening function are searched. The identification process is performed using the MSC.Marc, OptiSlang optimization software and Matlab. Keywords: unidirectional fiber composite, non-linear behavior, optimization, identification, matrix work-hardening function, representative volume element, unit cell, micromodel Cilj dela je bila identifikacija parametrov materiala za nadomestek pri enosmernem ogljik-epoksi dolgovlaknatem oja~enem kompozitu. Enostavni raztr`ni preizkusi z uporabo odrezkov z razli~no orientacijo vlaken so bili izvr{eni in dobljeni so bili diagrami sila – pomik. Model z enotno celico je bil ustvarjen v MSC.Marc. Vlakna so bila upo{tevana kot nelinearen elasti~en, pre~no izotropen material, matica pa je bila upo{tevana kot elastoplasti~en izotropen material. Vsota kvadratov razlik v pomiku med numeri~no in eksperimetalano dose`enimi diagrami sila – deformacija je bila minimalizirana z identifikacijskim procesom. Iskani so bili parametri linearne odvisnosti med Young modulom vlaken in deformacijo v smeri osi vlaken ter trije oblikovni koeficienti za deformacijsko utrditev matice. Proces identifikacije je bil izvr{en z uporabo MSC.Marc, OptiSlang-softvera za optimizacijo in Matlaba. Klju~ne besede: enosmerni vlaknati kompozit, nelinearno vedenje, optimizacija, identifikacija, funkcija deformacijske utrditve matice, reprezenta~ni element volumna, celica enote, mikromodel 1 INTRODUCTION Composite materials are widely used in all fields of industry such as aerospace, sport, automotive and transportation. Frequently used composites are based on a carbon-fibers and epoxy matrix for its high specific strength and stiffness. The knowledge of the material characteristics is crucial for the accuracy of the nume- rical models used in a designing process. The above type of composite shows a significant non-linear behavior. Therefore, complex non-linear material models must be used in order to achieve a good agreement with the experimental data even for the simple tensile tests. The modeling of large structures requires the use of macromodels, i.e., homogenized material models. The parameters of a macromodel can be assessed either by using a combination of a finite-element model with the mathematical optimization technique and experimental data or by using a micromodel of a unit-cell element, which is a periodically repeated volume fraction, with the knowledge of mechanical properties of all the constituents. A micromodel of the composite material can be advantageous for deeper analyses of the phenomena such as the influence of heterogeneities or microdamage mechanisms, etc. 2 EXPERIMENT Tensile tests of the thin coupons made of unidirec- tional long-fiber carbon-epoxy composite SE84LV- HSC-450-400-35 were performed on the testing machine ZWICK/ROELL Z050. The coupons were cut by a water jet from one large plate. Materiali in tehnologije / Materials and technology 46 (2012) 5, 431–434 431 UDK 66.017:519.61/.64 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)431(2012) Figure 1: Geometry of composite coupons (mm) 1 Slika 1: Geometrija odrezkov kompozita (mm) 1 The fiber direction forms the angles of 0°, 15°, 30°, 45°, 60°, 75° and 90° with the direction of the loading force (Figure 1). There were 10 specimens tested for each angle. Cracked specimens1 are shown in Figure 2. The specimens loaded along the fiber direction are fractured due to a fiber failure. All the specimens loaded at a different angle are fractured due to a matrix failure. The resulting force-displacement diagrams are shown in Figure 3. 2.1 Micromodel A finite-element model (micromodel) of a periodically repeated volume (unitcell, Figure 4) of the unidirectio- nal composite material was created in the finite-element system MSC.Marc2. A perfect honeycomb distribution of the fibers and a fiber-volume ratio of 55% were assumed (Table 1). Table 1: Geometry ratios of a unit cell Tabela 1: Geometri~ma razmerja enotne celice Fiber radius r Short side length 1.28 r Long side length 2.22 r Assuming the uniaxial stress across the whole specimen, the behavior of the material can be simulated by loading the unit cell with the normal stress  corresponding to the external force F:  = F A (1) where A is a cross-section of the specimen. The global coordinate system (xyz) is given with the force direction (x) and the direction perpendicular to the composite surface (z). The local coordinate system (123) is defined with the unit-cell edges, where the axis directions correspond to the fiber direction (1) and the directions perpendicular to it (Figure 5). The loading force is transformed to the local coordinate system using the transformation:             ⎡ ⎣ ⎢ ⎢ ⎢ ⎤ ⎦ ⎥ ⎥ ⎥ = − cos sin sin cos sin cos s 2 2 2 2 2 2 in cos sin cos sin cos cos            − ⎡ ⎣ ⎢ ⎢ ⎢ ⎤ ⎦ ⎥ ⎥ ⎥ ⎡ ⎣ ⎢ ⎢ ⎢ ⎤ ⎦ 2 ⎥ ⎥ ⎥ (2) where  is the angle of rotation between the local and the global coordinate systems3. The results from the finite-element analysis (strains) are transformed back to the global coordinate system using the transformation4: H. SRBOVÁ et al.: IDENTIFICATION OF THE MATERIAL PARAMETERS OF A UNIDIRECTIONAL FIBER ... 432 Materiali in tehnologije / Materials and technology 46 (2012) 5, 431–434 Figure 4: Three-dimensional mesh of a unit cell Slika 4: Tridimemzionalna mre`a enotne celice Figure 2: Cracked specimens with aluminum tabs1 Slika 2: Razpokani vzorci z aluminijevo podlago1 Figure 5: Rotated coordinate systems Slika 5: Rotirani koordinatni sistem Figure 3: Measured force-displacement diagrams (grey) for each fiber angle and the corresponding averaged values (black) Slika 3: Izmerjeni diagrami sila – pomik (sivo) za vsak kot vlakna in ustrezne povpre~ne vrednosti (~rno)         x y xy ⎡ ⎣ ⎢ ⎢ ⎢ ⎤ ⎦ ⎥ ⎥ ⎥ = −cos sin sin cos sin cos sin 2 2 2 2              cos sin cos sin cos cos sin2 2 2 2− ⎡ ⎣ ⎢ ⎢ ⎤ ⎦ ⎥ ⎥  ⎡ ⎣ ⎢ ⎢ ⎤ ⎦ ⎥ ⎥ (3) The unit cell must also respect the periodical boundary conditions (shown schematically in Figure 6): Δ Δ = − Δ = − u u u v v v w w w = −B A B A B A (4) where u, v and w are the translation differences of a pair of opposing nodes in directions 1, 2 and 3, respectively. These differences must remain constant for all the pairs of the corresponding nodes on the opposite sides3. In MSC.Marc, the periodical boundary conditions were implemented using a combination of links defined in the Fortran subroutine and springs. 2.2 Material models The experimental results from the tensile tests show a non-linear behavior of the composite even when loaded in the fiber direction (Figure 1). In order to capture this phenomenon a non-Hookean material model was considered for the fibers. The dependence of the longitu- dinal Young’s modulus of fibers on strain is: E E g11 11 11 0 111( ) ( ) = + (5) where g is the coefficient describing the measure of non-linearity and E011 is the initial Young’s modulus of fibers in the longitudinal direction5. The fiber is modeled as a transversely isotropic material6,7. The standard material constants given by the manufacturer are in Table 2. Table 2: Material parameters of the fiber given by the manufacturer Tabela 2: Parametri materiala vlaken, dobljeni od proizvajalca E011 (GPa) 230.00 E22 = E33 (GPa) 15.00 G12 = G23 = G31 (GPa) 50.00 12 = (–) 0.30 31 (–) 0.02 Vf (–) 0.55 The work-hardening function which respects a non-linear behavior of the matrix was proposed in the following form:      = ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ ⎡ ⎣ ⎢ ⎤ ⎦ ⎥ E E m p m p n n 0 1 (6) where p is an equivalent plastic deformation1. The matrix material was modeled to be isotropic having a Poisson’s ratio of m = 0.3 (given by the manufacturer). 2.3 Identification process The average curve for the experimentally obtained force-displacement diagrams was calculated for each angle of the fiber direction. These averaged diagrams are considered as target curves for the further analysis. Hereafter, the unit cell was loaded with the stress components corresponding with the uniaxial loading of the samples (2). The unit cell is loaded up to the range corresponding to the maximum value of the loading force in the target curve. The displacement dependence on the axial force is obtained by transforming the unit- cell strains back to the global coordinate system (3). The numerically obtained force-displacement dia- grams are subsequently compared with the target force- displacement curves. H. SRBOVÁ et al.: IDENTIFICATION OF THE MATERIAL PARAMETERS OF A UNIDIRECTIONAL FIBER ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 431–434 433 Figure 6: Equivalently deformed opposite boundaries of a hetero- geneous unit cell Slika 6: Ekvivalentno deformirane nasprotne meje heterogene enotne celice Table 3: Identified material parameters Tabela 3: Identificirani parametri materiala g (–) 23.23 E011 (GPa) 189.93 Em (GPa) 7.17 0 (kPa) 88.15 n (–) 1.56  (°) – 0.36 Figure 7: Equivalent plastic-strain contours in the matrix for  = 30° at the maximum load Slika 7: Ekvivlenten kontur plasti~ne deformacije matice za  = 30° pri najve~ji obremenitvi An optimization process was performed using Matlab, the optimization system OptiSlang and MSC.Marc. The goal was to find the best combination of all material coefficients by minimizing the sum of the squared differences of the numerical and experimental displacements l calculated as: e e l l l i i Ni N = = −⎡ ⎣⎢ ⎤ ⎦⎥ ∑ ∑∑ =    ( )Δ Δ Δ EXP FEA EXP 2 1 (7) Besides the material parameters from relations (5) and (6) an inaccuracy of the cutting of the samples was taken into account. This inaccuracy  was attributed to the angle  (Figure 5). The identified material parameters are summarized in Table 3, an example of the plastic strain in the matrix is shown in Figure 7, and the resulting force-displacement diagrams are compared in Figure 8. 3 CONCLUSION The tensile tests of the unidirectional fiber-reinforced carbon-epoxy composite coupons were performed for different angles of the fibers. A micromodel of the composite material was created. Parameters of the non-linear material models of both constituents (matrix and fibers) were identified in the optimization process. The parameters were identified by minimizing the error between numerically and experimentally obtained force- displacement diagrams. Moreover, a manufacturing inaccuracy during the specimen cutting was taken into account in the optimization. Future research will be aimed at the effects of the material imperfections, such as fiber undulation or inclusions in the matrix, and the modeling of the material-failure processes. Acknowledgement The work has been supported by the projects GA P101/11/0288 and the European project NTIS – New Technologies for Information Society No. CZ.1.05/1.1.00/02.0090. 4 REFERENCES 1 T. Kroupa, V. La{, R. Zem~ík, Improved Non-Linear Stress–Strain Relation for Carbon–Epoxy Composites and Identification of Material Parameters, Journal of Composite Materials, 45 (2011) 9, 1045–1057 2 MSC.Software Corporation: MSC.Marc User’s Guide, 2000 3 H. Srbová, Analysis of fiber composite from micromechanics point of view (in Czech), Diploma thesis, University of West Bohemia, Plzeò 4 D. Roylance, Transformation of Stresses and Strains, Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge 2001 5 I. M. Djordjevi}, D. R. Sekuli}, M. M. Stevanovi}, Non-Linear Elastic Behavior of Carbon Fibres of Different Structural and Mechanical Characteristic, Journal of the Serbian Chemical Society, 72 (2007) 5, 513–521 6 P. P. Camanho, C. G. Dávila, S. T. Pinho, J. J. C. Remmers, Mechanical Response of Composites, Springer – Verlag, 2008 7 V. La{, Mechanics of Composite Materials (in Czech), University of West Bohemia, Plzeò H. SRBOVÁ et al.: IDENTIFICATION OF THE MATERIAL PARAMETERS OF A UNIDIRECTIONAL FIBER ... 434 Materiali in tehnologije / Materials and technology 46 (2012) 5, 431–434 Figure 8: Comparison of the experimental and numerical force- displacement diagrams for all the angles  Slika 8: Primerjava eksperimentalnih in numeri~nih diagramov sila – pomik za vse kote  K. KRNEL et al.: MICROSTRUCTURE AND MECHANICAL PROPERTIES OF CARBON/CARBON-SILICON CARBIDE ... MICROSTRUCTURE AND MECHANICAL PROPERTIES OF CARBON/CARBON-SILICON CARBIDE COMPOSITES PREPARED BY SOL-GEL PROCESSING MIKROSTRUKTURA IN MEHANSKE LASTNOSTI KOMPOZITOV OGLJIK/OGLJIK-SILICIJEV KARBID, PRIPRAVLJENIH PO SOL-GEL METODI Kristoffer Krnel1, Zmago Stadler2, Toma` Kosma~1 1Engineering Ceramics Department, Jo`ef Stefan Institute, Jamova 39, 1000 Ljubljana, Slovenia 2MS Production, Pot na Lisice 17, 4260 Bled, Slovenia kristof.krnel@ijs.si Prejem rokopisa – received: 2011-10-23; sprejem za objavo – accepted for publication: 2012-05-24 In this work the preparation of a C/C composite with SiC precipitates in the matrix was studied. The formation of SiC particles in the matrix was achieved by substituting the phenolic resin used for the preparation of the composites with a phenolic resin–silica precursor prepared with the sol-gel method. The change of the matrix-phase composition resulted in improved mechanical properties of the composites which was attributed to the change in the interface between the matrix and the fibres. The silicon-carbide particles precipitating from the silicon containing a matrix are present directly at the interface increasing the bonding strength between the matrix and the fibres. Keywords: ceramic-matrix composites, carbon fibres, silicon carbide, sol-gel methods, mechanical properties V delu smo raziskovali pripravo C/C-kompozitov z izlo~ki SiC v matri~ni fazi. SiC-delce smo pripravili z zamenjavo dela fenolne smole, ki se uporablja za pripravo kompozitov, s prekurzorjem fenolna smola-SiO2, ki smo ga pripravili po sol-gel metodi. Zaradi spremembe matri~ne faze so se izbolj{ale mehanske lastnosti materiala, kar smo pripisali spremembi na fazni meji med vlaknom in matri~no fazo. Delci silicijevega karbida so se namre~ izlo~ali na povr{ini vlaken in s tem pove~ali trdnost spoja med matrico in vlaknom. Klju~ne besede: kompoziti s kerami~no matrico, ogljikova vlakna, silicijev karbid, sol-gel, mehanske lastnosti 1 INTRODUCTION Non-oxide composites with a ceramic matrix (C/C, C/SiC and SiC/SiC) have aroused great interest in the past years as a high-temperature structural material for use as a construction material in modern engines, braking systems, gas turbines and heat exchangers.1,2 On the one hand, we have C/C composites, which are, com- pared with the ceramic materials, relatively tough and elastic, but have low corrosion and wear resistance and lower thermal conductivity, and, on the other hand, there are C/SiC and SiC/SiC composites with good corrosion and wear resistance and higher thermal conductivity, but they are more rigid and brittle. Recently, new types of C/C-SiC composites were developed, which combine the good properties of all the above-mentioned systems.3 These materials preserve the structure of a C/C com- posite in the core, which gives the material the required toughness and elasticity; however, on the surface there is a layer of a C/SiC composite with good corrosion and wear resistance, which makes it ideal for use in the advanced braking systems. The weakness of these materials is a thermal-expansion mismatch between C/C in the core and C/SiC on the surface leading to an inten- sive cracking of the surface layer, which is amplified also because of the low thermal conductivity of the C/C core perpendicular to the carbon fibres leading to high-tempe- rature differences between the surface and the core.4 The aim of this work was to prepare a C/C composite with SiC precipitates in the matrix since such precipi- tates can improve the mechanical properties and possibly improve the thermal conductivity of the matrix phase. The formation of SiC particles in the matrix was studied by substituting the phenolic resin used for the prepa- ration of the composites with a phenolic-resin-silica precursor (a ceramic-forming polymer) prepared by the sol-gel method.5,6 The microstructure and its influence on the mechanical properties of such composites were also investigated. 2 EXPERIMENTAL The materials used in this study were as follows: for the preparation of all composites staple fibre fabrics (SGL Technik, Germany) were used. The samples of ceramic-matrix composites (CMCs) were prepared by impregnating the fabrics with a precursor produced via the sol-gel procedure from a phenolic resin (resolic type; Fenolit, Slovenia), TEOS (Acros Organic, USA), abso- lute ethanol (Carlo Erba Reagenti, Italy) and deionised Materiali in tehnologije / Materials and technology 46 (2012) 5, 435–438 435 UDK 666.3/.7:66.017 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)435(2012) water with an addition of the mass fraction of concen- trated HCl (37 %; Merck, Germany) used as a catalyst. 35.5 % of water was mixed with 33 % of ethanol and 1.5 % of HCl. 30 % of TEOS was added and the solution was mixed for 5 min to obtain a stable sol which was mixed with the phenolic resin for another 10 min in various ratios. The name of the samples, e.g. 50/50, denotes the mass ratio between the phenolic resin and the silica gel used for the preparation of the creamer to be 50 % phenolic resin and 50 % silica gel. Al the samples were prepared with a polymer infiltration and pyrolysis process (PIP) with 5 subsequent impregnations and pyrolysis cycles. The carbonisation conditions were as follows: 950 °C, 2 h, a heating rate of 2 °C/min in Ar gas. An additional heat treatment at 1600 °C for 2 h in the flowing argon with a heating rate of 20 °C/min was conducted to allow the matrix phase to crystallise. All the samples were characterised by XRD (D4 Endeavor, Bruker AXS, Germany), scanning (JSM-5800, JEOL, Japan) and a transmission electron microscope (JEM- 2000 FX, JEOL, Japan). The flexural strength was measured with a 3-point bending test with a span length of 80 mm (1362, Instron, UK). 3 RESULTS AND DISCUSSION Figure 1 shows the microstructures of the samples prepared from the precursor synthesised using the sol-gel procedure after the last cycle of the PIP process. In all the figures the carbon fibres are surrounded by a brighter matrix phase that is rich in silicon. The presence of the silicon was confirmed using an EDXS analysis shown in Figure 2. In all the samples the presence of some poro- sity and inhomogeneity can be seen; there is almost no brighter, silicon-rich phase present in between the fibres in the fabric. The reason for this is a relatively high viscosity of the phenolic-resin-silica ceramer used for the impregnation of the fabrics and the spaces between the fibres, which are not well filled. They are only filled later, during the re-impregnation of the composite with a pure phenolic resin. The variation of the SiO2-gel/ phenolic-resin ratio does not have any influence on the microstructures of the composites after the PIP process. The microstructures of the samples after the crystalli- sation heat treatment at 1600 °C are shown in Figure 3. At this temperature the precipitation of small nanometric particles can be observed in the matrix phase, preferen- tially around the carbon fibres. There is also a relatively large amount of free carbon present. After the thermal treatment the influence of the SiO2-gel/phenolic-resin ratio can be observed, and with an increased amount of SiO2 gel the number of SiC precipitates increases as well. The precipitates shown in Figure 4 were analysed using transmission electron microscopy. The results are presented in Figure 5. Electron diffraction patterns of the nanocrystalline material, consisting of a few broad- ened rings and many scattered spots could not be easily analyzed, thus simulated electron-diffraction patterns of K. KRNEL et al.: MICROSTRUCTURE AND MECHANICAL PROPERTIES OF CARBON/CARBON-SILICON CARBIDE ... 436 Materiali in tehnologije / Materials and technology 46 (2012) 5, 435–438 Figure 2: SEM microstructures and an EDX analysis of the CMC 50/50 sample after PIP Slika 2: Mikrostruktura in EDS-analiza vzorca CMC 50/50 po PIP Figure 1: SEM microstructures of the samples prepared from a sol-gel precursor after PIP: a) sample CMC 70/30, b) sample CMC 60/40 and c) sample CMC 50/50 Slika 1: Mikrostruktura vzorca, pripravljenega iz sol-gel prekurzorja po PIP: a) vzorec CMC 70/30, b) vzorec CMC 60/40 in c) vzorec CMC 50/50 possible candidates were calculated and compared to the experimental diffraction patterns. Based on these calculations it was concluded that nanocrystals have a structure of hexagonal (6H) silicon carbide that are embedded in an amorphous carbon matrix. The com- parison of the experimental results with the calculated data for the nanocrystalline material is presented in Figure 2b. In our first attempts the flexural strength of the samples prepared by using the phenolic-resin-silica ceramer was approximately the same as the flexural strength of C/C composites prepared by using only phenolic resin and the values were around 60 MPa.4 However, the microstructure analysis showed that the samples prepared by using the phenolic-resin-silica pre- cursor were more porous and contained larger amounts of inhomogenieties. For that reason the new set of samples was prepared with more attention paid to the preparation step. The results of the flexural-strength measurements of the new, more homogenous samples, before and after the crystallisation thermal treatment, are presented in Table 1. The flexural strength of the samples prepared from the phenolic-resin-silica precursor is somewhat higher than that of the carbon-carbon composite prepared by using just a phenolic resin (sample CMC 100/0). These results are somewhat in contrast with the results of a similar study by Chen-Chi et al.7 The presence of the silicon in the amorphous matrix changes the structure of the carbon-fibre-matrix interface and influences the strength. From the results it can also be seen that the crystallisation heat treatment further improves the strength of the composites prepared with the ceramer. The reasons for that are most probably the increased strength of the matrix phase as well as the increased strength of the interface between the fibres and the K. KRNEL et al.: MICROSTRUCTURE AND MECHANICAL PROPERTIES OF CARBON/CARBON-SILICON CARBIDE ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 435–438 437 Figure 5: a) TEM microstructure showing nano-precipitates in the carbon matrix, b) experimental and calculated electron-diffraction patterns Slika 5: a) TEM-posnetek, ki prikazuje nanoprecipitate v ogljikovi matrici, b) eksperimentalna ter izra~unana elektronska difrakcija Table 1: Flexural strength of the samples (SD) after their preparation via PIP and after additional heat treatment at 1600 °C for 2h in the flowing Ar Tabela 1: Upogibne trdnosti vzorcev (in standardna deviacija) po pripravi preko PIP in dodatni termi~ni obdelavi pri 1600 °C 2h v pretoku Ar CMC 100/0 CMC 70/30 CMC 60/40 CMC 50/50 Flexural strength after PIP (SD), MPa 59 (4) 62 (5) 67 (5) 69 (6) Flexural strength after heat treatment (SD), MPa 58 (3) 68 (5) 72 (6) 73 (4) Figure 4: SEM microstructure of a carbon matrix showing nanosized precipitates Slika 4: Mikrostruktura matri~ne faze, v kateri so vidni nanometrski izlo~ki Figure 3: SEM microstructures of the samples prepared from a sol-gel precursor after crystallisation thermal treatment: a) sample CMC 70/30, b) sample CMC 60/40 and c) sample CMC 50/50 Slika 3: SEM-posnetek vzorcev, pripravljenih s sol-gel prekurzorjem po termi~ni obdelavi: a) vzorec CMC 70/30, b) vzorec CMC 60/40 in c) vzorec CMC 50/50 matrix, resulting from the precipitation of the silicon- carbide particles around the fibres. In Figure 6 the fibre/matrix interface of the CMC 50/50 sample is shown. The interface of the sample containing silicon carbide shows the presence of the SiC particles at the matric-fibre interface that is influencing the strength of the interface resulting in a higher strength of the com- posite. The interface was analysed using also the trans- mission electron microscope to verify the presence of the SiC particles at the interface. The micrograph of the interface between the fibre and the SiC particles contain- ing the matrix is presented in Figure 7. It can be seen that the SiC particles are precipitating at the interface without any phase visible in between. Obviously, the interface between the matrix and the fibres is changed, which is affecting the mechanical properties of the com- posite. The bonding is stronger, but not too strong, so the pull-out of the fibres is still visible on the fracture surfaces of the samples. This effect is also the reason for high strength and toughness of these composite mate- rials. 4 CONCLUSIONS The preparation of the C/C-SiC composites is possible by replacing the phenolic resin with a phenolic- resin–silica ceramer. In the case of the phenolic-resin- silica ceramer the matrix phase contained nanoprecipi- tates of SiC after the crystallisation heat treatment at 1600 °C in an inert atmosphere. The changing of the matrix phase improved the mechanical properties of the composites, which was attributed to the change in the interface between the matrix and the fibres. The silicon- carbide particles precipitating from the silicon containing a matrix are present directly at the interface increasing the bonding strength between the matrix and the fibres. The presence of the SiC particles on the interface was also confirmed by the TEM microscopy. 5 REFERENCES 1 R. Naslain, Design, preparation and properties of non-oxide CMCs for application in engines and nuclear reactors: an overview, Comp Sci Tech., 64 (2004), 155–170 2 W. J. Sherwood, CMCs Come Down to Earth, Am Ceram Soc Bull., 82 (2003), 25–27 3 M. Zornik, EP 0 818 636 B1, Fahrzeugbrems- bzw. Fahrzeug- Kupplungsscheibe aus mit SiC beschichtetem C-C Werkstoff, 1997 4 K. Krnel, Z. Stadler, T. Kosma~, Preparation and properties of C/C-SiC nano-composites, J Eur Ceram Soc., 27 (2007), 1211–1216 5 J. W. Li, J. M. Tian, L. M. Dong, Synthesis of SiC precursors by a two-step sol-gel process and their conversion to SiC powders, J Eur Ceram Soc., 77 (2000), 1853–1857 6 J. M. Tian, J. W. Li, L. M. Dong, Synthesis of SiC precursors by sol-gel process, J Inorg Mat., 14 (1999), 297–301 7 C. C. M. Ma, J. M. Lin, W. C. Chang, T. H. Ko, Carbon/carbon nanocomposites derived from phenolic resin – silica hybrid ceramers: microstructure, physical and morphological properties, Carbon, 40 (2002) 7, 977–984 K. KRNEL et al.: MICROSTRUCTURE AND MECHANICAL PROPERTIES OF CARBON/CARBON-SILICON CARBIDE ... 438 Materiali in tehnologije / Materials and technology 46 (2012) 5, 435–438 Figure 7: TEM microstructure of the carbon fibre/matrix interface Slika 7: TEM-posnetek fazne meje matrica/vlakno Figure 6: SEM microstructure of the carbon fibre/matrix interface Slika 6: Mikrostruktura fazne meje matrica/vlakno A. C. KARAOÐLANLÝ et al.: STUDY OF THE MICROSTRUCTURE AND OXIDATION BEHAVIOR OF YSZ ... STUDY OF THE MICROSTRUCTURE AND OXIDATION BEHAVIOR OF YSZ AND YSZ/Al2O3 TBCs WITH HVOF BOND COATINGS [TUDIJ MIKROSTRUKTURE IN VEDENJA PRI OKSIDACIJI YSZ IN YSZ/Al2O3 TBC Z HVOF NANESENO ZA[^ITNO PREVLEKO Abdullah Cahit Karaoðlanlý1, Garip Erdoðan2, Yaºar Kahraman3, Ahmet Türk2, Fatih Üstel2, Ýsmail Özdemir1 1Department of Metallurgical and Materials Engineering, Bartin University, 74100 Bartin, Turkey 2Department of Metallurgical and Materials Engineering, Sakarya University, 54187 Sakarya, Turkey 3Department of Mechanical Engineering, Sakarya University, 54187 Sakarya, Turkey cahitkaraoglanli@gmail.com, karaoglanli@bartin.edu.tr Prejem rokopisa – received: 2011-11-11; sprejem za objavo – accepted for publication: 2012-03-27 A significant improvement in efficiency has been achieved by using thermal barrier coatings (TBCs) in gas turbines and diesel engines. A typical TBC is a multilayered coating system that comprises an oxidation-resistant metallic bond coating (BC) and a thermally insulating ceramic top coating (TC). Under service conditions an Al2O3 inter-layer, the thermally grown oxide (TGO), forms in the interface between the bond and the top coating, by a chemical reaction between the metallic aluminum from the BC material and the oxygen that comes from the environment through the pore channels of the TC. The aim of the present study is to describe the TGO formation on metallic bond coats deposited using the high-velocity oxygen fuel (HVOF) spraying technique. Therefore, TBCs that consist of a YSZ top (ZrO2 + 8 % Y2O3) and YSZ-Al2O3 double-layer systems with CoNiCrAlY bond coats were deposited on Inconel 718 super-alloy substrates. The bond coats were applied via HVOF, with the ceramic top coats being applied by atmospheric plasma spraying (APS) as well. The oxidation behaviors of the TBC systems were investigated. The oxidation tests were performed at 1000 °C in an air atmosphere for (8, 24, 50) h. The formation and growth of the TGO layers and the microstructural changes during the oxidation tests were scrutinized systematically. The results indicate that the TBC coating with the YSZ-Al2O3 double layer had a higher oxidation resistance and a lower TGO layer growth than that of the traditional TBC system. Likewise, the initial state of the porosity plays a critical role in enhancing or limiting the growth of the TGO scale in the TBC. Keywords: thermal barrier coatings (TBCs), oxidation behavior, thermally grown oxide (TGO), high-velocity oxygen fuel (HVOF), atmospheric plasma spraying (APS) Dose`eno je bilo ob~utno izbolj{anje u~inkovitosti plinskih turbin in dieselskih motorjev z uporabo toplotnih za{~itnih prevlek (TBC). Zna~ilni TBC je ve~slojni varovalni sistem, ki vklju~uje prevleko, odporno proti oksidaciji, s kovinsko vezjo (BC) in toplotno izolativno kerami~no vrhnjo plastjo (TC). Pri obratovalnih razmerah vmesni sloj Al2O3 omogo~a nastanek oksidov (TGO) na stiku med vezivnim in vrhnjim slojem, s kemijsko reakcijo kovinskega aluminija iz BC-materiala in kisika, ki iz okolice prodira skozi pore TC. Namen te {tudije je opisati nastanek TGO na kovinskem vezivu, nanesenem z nabrizgavanjem s kisikovim plamenom z veliko hitrostjo (HVOF). TBC, ki sestoji na vrhu iz YSZ (ZrO2 + 8 % Y2O3) ter z dvoslojnim sistemom YSZ-Al2O3 z vezivno plastjo CoNiCrAlY so bili naneseni na podlago iz superzlitine Inconel 718. Vezivna plast je bila nanesena z HVOF, kerami~ni vrhnji sloj pa z atmosfersko plazmo (APS). Preiskovane so bile zna~ilnosti oksidacije TBS-sistema. Preizkusi oksidacije so bili izvr{eni na zraku pri 1000 °C, za (8, 24 in 50) h. Med preizkusi oksidacije so bili sistemati~no preiskovani nastanek in rast TGO-plasti ter spremembe v mikrostrukturi. Rezultati ka`ejo, da ima TBC-prevleka YSZ-Al2O3 z dvojnim slojem bolj{o odpornost proti oksidaciji in manj{o rast TGO-plasti v primerjavi z navadnim TBC-sistemom. Videti je, da ima za~etna poroznost klju~no vlogo pri pospe{evanju ali zaviranju rasti TGO-plasti na TBC. Klju~ne besede: termi~ni varovalni sloj (TBC), vedenje pri oksidaciji, termi~na rast oksida (TGO), kisikov plamen z veliko hitrostjo (HVOF), atmosfersko plazemsko nabrizgavanje (APS) 1 INTRODUCTION Many attempts have been made to understand the role of TGO, formed at the interface between the bond coat and the top coat during elevated-temperature service conditions, which strictly governs the lifetime of the TBC. The thickness, the roughness of the TGO, the adherence quality of the bond coat to the substrate, the type and shape of the oxides present in the vicinity of the TGO during oxidation are the main issues in controlling the degradation of the TBC1–4. Likewise, the oxidation behavior of the TBC is strongly linked to the bond coat properties, which affect the durability of the TBC. The bond coat is deposited conventionally by LPPS, HVOF, plasma and also cold gas dynamic spray: a method re- cently preferred to avoid complex oxide formation as well5–8. Generally, the forming of a dense, homogeneous, -Al2O3 oxide scale is preferred as it is relatively stable, chemically and thermally, which means the degradation of the -Al2O3 is negligible and also has a low ion diffusivity, which causes a slow growth rate and prevents further oxidation9–11. Oxidation-based damage, which is the result of stresses developed at the interface of the top coat and the bond coat during TGO growth, is a common failure of TBC since these stresses result in the Materiali in tehnologije / Materials and technology 46 (2012) 5, 439–444 439 UDK 621.793:542.943:620.18 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)439(2012) spallation-induced failure of the topcoat. In order to retard or avoid such a failure in the TBC, i.e., better oxidation resistance, any methods employed should facilitate the slow growth of the TGO scale, which favors good adherence of the TGO12–14. To achieve this, apart from employing several kinds of methods to deposit the bond coat, a thermal barrier thin film deposited by EB-PVD, CVD, etc. was employed over the bond coat in order to inhibit the formation of undesirable mixed oxides as they have a fast growth rate leading to accelerated TBC failure15–17. In addition, the introduction of alumina powders to the YSZ might possibly reduce the inward diffusion of oxygen from the topcoat and thus make it difficult to have rapid growth of the TGO. It was claimed in studies that Al2O3 present in the topcoat exhibited better oxidation resistance compared to the YSZ without Al2O3 powders18,19. In this work, the YSZ topcoat with a conventional composition and with YSZ/Al2O3 double- layer systems were applied on Inconel 718 super-alloy substrates to investigate and compare their oxidation behaviors. The oxidation results for the traditional TBC that has the YSZ system was compared with the YSZ/Al2O3 double-layer TBC system. The microstruc- tural differences and the formation and growth of the TGO layers in the isothermal oxidation resistance of these TBC systems are discussed. 2 EXPERIMENTAL METHODS 2.1 Materials and coating-deposition methods The Inconel 718 Ni-based super-alloy, in disc-shaped coupons, was used as the substrate. Co38Ni32.5Cr21Al8Y0.5, ZrO2–8 % Y2O3 and Al2O3 powder were used as the starting materials. A Microtrack S3500 laser particle-size analyzer was used to determine the powder size distribution. The mean diameters were determined to be d50 = 33 μm, 38.52 μm, 33.36 μm for the CoNiCrAlY, ZrO2–8 % Y2O3, and Al2O3 powders, respectively. Half of the TBC samples consisted of a CoNiCrAlY BC and a ZrO2–8 % Y2O3 TC and the other half of the TBC sam- ples were composed of a CoNiCrAlY BC, a ZrO2–8 % Y2O3 TC and an Al2O3 top coat over the YSZ. The HVOF technique was used to produce bond coats and the ceramic top coatings were produced by the APS method using a fully automated MultiCoat System from Sulzer Metco. All the spraying parameters are shown in Table 1. 2.2 Microstructural Characterization The microstructures of the TBC systems were investigated by scanning electron microscopy (SEM, Tescan VEGA II, SBU Bruker EDX, Czech Republic). The porosity of the bond coatings was measured using optical image-analysis software (Olympus a4i). The coating microhardness (HV0.3) was determined using a microhardness tester (Shimadzu, Japan) with a load of 300 g for 15 s from the bond coats and the top coats. The oxidation tests of the TBC system produced were conducted by means of a high-temperature furnace (Nabertherm, Germany) with an air atmosphere. 3 RESULTS AND DISCUSSION 3.1 Microstructure of the powders and the coatings Two types of TBC were prepared using the HVOF method to produce bond coats that included CoNiCrAlY. The APS method was used to produce ceramic top coats, which included traditional YSZ and a double layer of YSZ and Al2O3 in which the Al2O3 was a top coat over the YSZ in a second system. The thickness of the bond and the top coats of both systems were about 100 μm and 300 μm, respectively. The YSZ top coat used in the first system was 300 μm. The YSZ and Al2O3 ceramic top coatings used in the second system were both 150 μm. The type, components, thicknesses and spray systems of the coating layers are shown in Table 2. Table 2: Type, components, thicknesses and spray systems of the coating layers Tabela 2: Vrsta, komponente, debelina in sistem za nana{anje prevlek Type of TBC system Component Thickness of layers, μm Spray system 1 CoNiCrAlY YSZ 100 300 HVOF APS 2 CoNiCrAlY YSZ Al2O3 100 150 150 HVOF APS APS Figure 1 shows the morphology of the as-received ZrO2+Y2O3, Al2O3 and CoNiCrAlY powders. As can be seen from this figure, the CoNiCrAlY powder has a spherical morphology, while the Al2O3 is angular. Fig- ures 2a and b show the cross-sectional microstructure of A. C. KARAOÐLANLÝ et al.: STUDY OF THE MICROSTRUCTURE AND OXIDATION BEHAVIOR OF YSZ ... 440 Materiali in tehnologije / Materials and technology 46 (2012) 5, 439–444 Table 1: Spraying parameters for deposition of the coatings Tabela 1: Parametri nabrizgavanja pri nana{anju prevlek HVOF CoNiCrAlY Bond Coatings APS YSZ Top Coatings APS Al2O3 Over Top Coatings Combustion medium Voltage (Plasma) Voltage (Plasma) O2 (880 L/min) and kerosene (25 L/h) 70 V 70 V Powder Carrier Gas Current (Plasma) Current (Plasma) Argon (15 L/min) 650 A 600 A Powder Feed Rate Carrier Gas Carrier Gas 50 g/min 3 nlpm 2.5 nlpm Powder feed gas flow H2 (Plasma) H2 (Plasma) 12 L/min 13 L/min 13 L/min Stand-off distance Argon (Plasma) Argon (Plasma) 330 mm 45 L/min 45 L/min Spraying Distance Spraying Distance 100 mm 120 mm Traverse Speed Traverse Speed 300 mm/s 300 mm/s traditional TBC and Al2O3-YSZ gradient double layer TBC system. The HVOF-CoNiCrAlY bond coats have relatively less porosity and cracks in the TBCs. The ceramic top coats contain porosity and some crack-like discontinuities during the spraying process. A smaller amount of porosity for the BC of both TBC systems was measured to be approximately 1.0 %. On the other hand, like for the top coat the porosity values were not significantly different from each other. The porosity level of the top coat was found to be appro- ximately 5.0 % in both TBC systems. But as is clear from Figure 2, the size and distribution of the porosity in the TC of the two-layer YSZ/Al2O3 TBC are quite different from the traditional values. 3.2 Oxidation tests The TBC specimens were subjected to oxidation tests. These oxidation tests were carried out in an air atmosphere at 1000 °C for (8, 24 and 50) h. Typical SEM microstructures of the whole TBC-systems are shown in Figures 3 and 4. As shown in these figures, the TGO was formed at the ceramic/bond-coat interface due to oxygen pene- tration through the ceramic layer. Various formations of A. C. KARAOÐLANLÝ et al.: STUDY OF THE MICROSTRUCTURE AND OXIDATION BEHAVIOR OF YSZ ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 439–444 441 Figure 3: Cross-sectional microstructures at the bond coat/ceramic layer interface for YSZ top coats with CoNiCrAlY coatings after oxidation at 1000 °C for (8, 24, 50) h Slika 3: Mikrostruktura prereza vezivne plasti/kerami~ne plasti z vrhnjo plastjo YSZ za CoNiCrAlY-prevleko po oksidaciji (8, 24 in 50) h na 1000 °C Figure 2: SEM micrographs of as-sprayed thermal barrier coatings: a) APS YSZ with HVOF bond coat and b) YSZ/Al2O3 with HVOF bond coat system Slika 2: SEM-posnetek nabrizganega sloja toplotne prevleke: a) APS YSZ z HVOF vezivno plastjo in b) YSZ/Al2O3 z HVOF vezivno plastjo Figure 1: SEM micrographs of the morphology of: a) ZrO2–8 % Y2O3, b) Al2O3 and c) CoNiCrAlY powders Slika 1: SEM-posnetek morfologije: a) ZrO2–8 % Y2O3, b) Al2O3 in c) CoNiCrAlY-prahov discontinuities between the bond layer and the ceramic top layer can be clearly seen in Figures 3 and 4. When the oxidation properties of the two different TBC systems are compared, it was clear that the oxidation in the conventional TBC system (CoNiCrAlY bond layer with YSZ top coat) developed faster than that of the YSZ/Al2O3 double-layer TBC system, and the TGO growth rate in the former system was found to be higher. As seen from the interface microstructures given in Figure 5, the TGO structure in the YSZ/Al2O3 double-layer TBC system is more uniform and is mainly composed of Al2O3, which was confirmed by an EDX analysis. In the conventional YSZ system, due to the increasing oxidation process, complex oxides developed in the TGO layer and affected the growth behaviors of the TGO. This was caused by the prevention of the oxygen penetration to the bond coat from the surface due to the Al2O3 layer and hence a slowing down of the oxygen attack in YSZ/ Al2O3 coating systems. As a result, the decrease of the Al2O3 content in the TGO layer caused by bond-coat oxidation is delayed, the degradation of the uniform structure is retarded and in this way an increase in the volume of the TGO occurs at a lower rate. Similar results showing an increase in the oxidation resistance of the coatings due to the Al2O3 layer depending on the temperature and the time exist in the literature18,20–25. Figure 6 indicates that the thickness of the TGO layer increased with increasing exposure time for both A. C. KARAOÐLANLÝ et al.: STUDY OF THE MICROSTRUCTURE AND OXIDATION BEHAVIOR OF YSZ ... 442 Materiali in tehnologije / Materials and technology 46 (2012) 5, 439–444 Figure 6: TGO thickness measurements as a function of oxidation time at 1000 °C, respectively ( YSZ top coatings with CoNiCrAlY bond coatings,  YSZ/Al2O3 top coatings with CoNiCrAlY bond coatings) Slika 6: Debelina TGO v odvisnosti od ~asa oksidacije pri 1000 °C ( vrhnja plast YSZ z vezivno plastjo CoNiCrAlY,  vrhnja plast YSZ/Al2O3 z vezivno plastjo CoNiCrAlY) Figure 5: Microstructures at the bond coat/ceramic-layer interface after oxidation at 1000 °C for 50 h; a) TGO in traditional YSZ coating; b) TGO in YSZ/Al2O3 coating Slika 5: Mikrostruktura stika vezivne plasti/kerami~ne plasti po oksidaciji 50 h na 1000 °C; a) TGO in navadna YSZ-plast; b) TGO in YSZ/Al2O3-plast Figure 4: Cross-sectional microstructures at the bond coat/ceramic layer interface for YSZ/Al2O3 top coats with CoNiCrAlY coatings after oxidation at 1000 °C for (8, 24, 50) h Slika 4: Mikrostruktura prereza vezivne plasti/kerami~ne plasti z vrhnjo plastjo YSZ/Al2O3 za CoNiCrAlY-prevleko po oksidaciji (8, 24 in 50) h na 1000 °C kinds of TBC specimens. The specimens with the YSZ top coatings with CoNiCrAlY bond coatings show a higher rate of TGO thickness growth than the samples with YSZ/Al2O3 top coatings with CoNiCrAlY bond coatings. This difference could be attributed to the initial porosity state of the as-sprayed TBC samples and/or the Al2O3 layer acting as a diffusion barrier due to its low diffusion coeficient for oxygen ions. This effect has been observed in many other studies20–25. Therefore, the greater the increase of the TGO thickness of the specimen with traditional YSZ top coatings with CoNiCrAlY bond coatings at the same exposure time compared to the specimen YSZ/Al2O3 top coatings with HVOF-BC could be attributed to this mechanism. After increasing the exposure time the TGO layer thickness increased to higher values. 3.3 Mechanical properties The microhardness value for the bond and top coats were taken from the average value of all the measure- ment points. Figures 7 and 8 present the Vickers micro- hardness measurements before and after the oxidation tests for the bond and the top layers. The bottom and top limit lines in the graph show the maximum and minimum hardness values. The microhardness of the substrate Inconel 718 super-alloy was in the range 310–340 HV. The mean values of the bond and top-coat micro- hardness changed before and after the oxidation tests with increasing time. The microhardness of the bond coats decreased with increasing time at 1000 °C for traditional and two-layered coatings. The decline in microhardness of the bond coatings was possibly linked to the decrease in the density of the bond coats. The decrease of the microhardness in the HVOF bond coats is related to the thermal relaxation of the residual stress present in the as-sprayed coating due to the high temperature. The fact that the decrease in microhardness after 8 h is higher while the decrease after 24 h is lower and no change is observed after 50 h should support this theory, i.e., thermal relaxation occurs very quickly, and since almost no change in microhardness is observed after 8 h it can be concluded that this initial change is due to thermal relaxation. The microhardness of the top coats increased with increasing time at 1000 °C for the traditional and two-layered ceramic coatings. According to the literature, the situation mentioned above is caused by the decreasing density of the porosity and the increasing density of the ceramic top coating depending on increasing time26. 4 CONCLUSIONS Traditional YSZ and YSZ/ Al2O3 double-layer TBCs were produced using the APS technique and bond coats were deposited using the HVOF spraying technique. The following results were obtained: During oxidation of the TBCs, the TGO layer was formed along the interface of the BC/TC layer. The thickness of the TGO in the traditional YSZ coating is higher in comparison with the YSZ/Al2O3 coating after oxidation at 1000 °C for different oxidation times. According to the TGO growth in both TBC systems, the TGO thickening became steady state in the YSZ/Al2O3 two-layer system and, on the other hand, the TGO A. C. KARAOÐLANLÝ et al.: STUDY OF THE MICROSTRUCTURE AND OXIDATION BEHAVIOR OF YSZ ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 439–444 443 Figure 8: Microhardness values of top coats for two kinds of TBCs, respectively ( YSZ top coats with CoNiCrAlY bond layer,  YSZ top coats with YSZ/Al2O3 top layer and CoNiCrAlY bond layer,  Al2O3 top coats with YSZ/Al2O3 top layer and CoNiCrAlY bond layer) Slika 8: Vrednosti mikrotrdote za dve vrsti TBC, ( vrhnja plast YSZ z vezivno plastjo CoNiCrAlY,  vrhnja plast YSZ z vrhnjo plastjo YSZ/Al2O3 in vezivno plastjo CoNiCrAlY, vrhnja plast Al2O3 z vrhnjo plastjo YSZ/Al2O3 in vezivno plastjo CoNiCrAlY) Figure 7: Microhardness values of bond coats for two kinds of TBCs, respectively (CoNiCrAlY bond coats with YSZ top layer,  CoNiCrAlY bond coats with YSZ/Al2O3 top layer) Slika 7: Vrednosti mikrotrdote vezivnih plasti za dve vrsti TBC, ( vezivna plast CoNiCrAlY z vrhnjo plastjo YSZ,  vezivna plast CoNiCrAlY z vrhnjo plastjo YSZ/Al2O3) thickness in the traditional TBC system was still increasing. The different formations of the discontinuities between the bond and the ceramic top layers were observed. The ceramic top coats contained porosity and some crack-like discontinuities for both kinds of TBC. After a prolonged oxidation time the number of cracks was much larger in the traditional YSZ ceramic top coat with the CoNiCrAlY bond coating system. The initial porosity state of the as-sprayed TBC samples and/or the acting of the Al2O3 layer as a diffusion barrier at high temperatures have a great influence on determining the TGO growth, since they change the penetration behavior of the oxygen from the surface. In the conventional YSZ system, due to the increasing oxidation process, complex oxides developed in the TGO layer and affected the growth behaviors of the TGO. The mean values of the bond and top-coat microhardnesses changed before and after the oxidation tests with increasing time. The microhardness of the bond coats decreased and that of the top coats increased with increasing time at 1000 °C for the traditional and two-layered coatings. 5 REFERENCES 1 A. G. Evans, D. R. Mumma, J. W. Hutchinson, G. H. Meierc, F. S. Pettit, Mechanisms controlling the durability of thermal barrier coatings, Progress in Materials Science, 46 (2001) 5, 505–553 2 Y. Li, C. J. Li, Q. Zhang, G. J. Yang, C. X. Li, Influence of TGO Composition on the Thermal Shock Lifetime of Thermal Barrier Coatings with Cold-sprayed MCrAlY Bond Coat, Journal of Thermal Spray Technology, 19 (2010) 1–2, 168-177 3 W. R. Chen, R. Archer, X. Huang, B. R. Marple, TGO Growth and Crack Propagation in a Thermal Barrier Coating, Journal of Thermal Spray Technology, 17 (2008) 5–6, 858–864 4 L. SwadŸba, G. Moskal, B. Mendala, T. Gancarczyk, Characteri- sation of air plasma sprayed TBC coating during isothermal oxidation at 1100 °C, Journal of Achievements in Materials and Manufacturing Engineering, 21 (2007) 2, 81–84 5 W. J. Brindley, Properties of Plasma-Sprayed Bond Coats, Journal of Thermal Spray Technology, 6 (1997) 1, 85–90 6 Y. Li, C. J. Li, G. J. Yang, L. K. Xing, Thermal fatigue behavior of thermal barrier coatings with the MCrAlY bond coats by cold spraying and low-pressure plasma spraying, Surface and Coatings Technology, 205 (2010), 2225–2233 7 W. O. Soboyejo, P. Mensah, R. Diwan, J. Crowe, S. Akwaboa, High temperature oxidation interfacial growth kinetics in YSZ thermal barrier coatings with bond coatings of NiCoCrAlY with 0.25 % Hf, Materials Science and Engineering A, 528 (2011), 2223–2230 8 S. Saeidi, K. T. Voisey, D. G. McCartney, The Effect of Heat Treat- ment on the Oxidation Behavior of HVOF and VPS CoNiCrAlY Coatings, Journal of Thermal Spray Technology, 18 (2009) 2, 209–216 9 N. Mu, T. Izumi, L. Zhang, B. Gleeson, Compositional Factors Affecting the Oxidation Behavior of Pt-Modified -Ni+ ’-Ni3Al- Based Alloys and Coatings, Materials Science Forum, 239 (2008) 595–598, 239–247 10 P. Richer, M. Yandouzi, L. Beauvais, B. Jodoin, Oxidation behaviour of CoNiCrAlY bond coats produced by plasma, HVOF and cold gas dynamic spraying, Surface and Coatings Technology, 204 (2010), 3962–3974 11 Q. Zhang, C. J. Li, C. X. Li, G. J. Yang, S. C. Lui, Study of oxidation behavior of nanostructured NiCrAlY bond coatings deposited by cold spraying, Surface and Coatings Technology, 202 (2008), 3378–3384 12 J. A. Thompson, T. W. Clyne, The effect of heat treatment on the stiffness of zirconia top coats in-plasma sprayed TBCs, Acta Mater., 49 (2001), 1565–1575 13 J. A. Haynes, M. K. Ferber, W. D. Porter, Thermal Cycling Behavior of Plasma-Sprayed Thermal Barrier Coatings with Various MCrAIX Bond Coats, Journal of Thermal Spray Technology, 9 (2000) 1, 38–48 14 F. H. Yuan, Z. X. Chen, Z. W. Huang, Z. G. Wang, S. J. Zhu, Oxidation behavior of thermal barrier coatings with HVOF and detonation-sprayed NiCrAlY bond coats, Corrosion Science, 50 (2008), 1608–1617 15 J. R. V. Garcia, T. Goto, Thermal barrier coatings produced by che- mical vapor deposition, Science and Technology of Advanced Materials, 4 (2003), 397–402 16 F. Pedraza, C. Tuohy, L. Whelan, A. D. Kennedy, High Quality Aluminide and Thermal Barrier Coatings Deposition for New and Service Exposed Parts by CVD Techniques, Materials Science Forum, 461–464 (2004), 305–312 17 M. H. Li, X. F. Sun, S. K. Gong, Z. Y. Zhang, H. R. Guan, Z. Q. Hu, Phase transformation and bond coat oxidation behavior of EB-PVD thermal barrier coating, Surface and Coatings Technology, 176 (2004), 209–214 18 M. Saremi, A. Afrasiabi, A. Kobayashi, Microstructural analysis of YSZ and YSZ/Al2O3 plasma sprayed thermal barrier coatings after high temperature oxidation, Surface and Coatings Technology, 202 (2008), 3233–3238 19 Q. Yu, A. Rauf, C. Zhou, Microstructure and Thermal Properties of Nanostructured 4 % Al2O3-YSZ Coatings Produced by Atmospheric Plasma Spraying, Journal of Thermal Spray Technology, 19 (2010) 6, 1294–1300 20 Kh. G. S. Thomas, U. Dietl, Thermal barrier coatings with improved oxidation resistance, Surface and Coatings Technology, 68/69 (1994), 113–115 21 Q. Yu, A. Rauf, C. Zhou, Microstructure and Thermal Properties of Nanostructured 4 % Al2O3-YSZ Coatings Produced by Atmospheric Plasma Spraying, Journal of Thermal Spray Technology, 19 (2010) 6, 1294–1300 22 C. Ren, Y. D. He, D. R. Wan, Fabrication and Characteristics of YSZ–YSZ/Al2O3 Double-Layer TBC, Oxidation of Metals, 75 (2011), 325–335 23 J. Müller, M. Schierling, E. Zimmermann, D. Neuschütz, Chemical vapor deposition of smooth a- Al2O3 films on nickel base superalloys as difusion barriers, Surface and Coatings Technology, 120–121 (1999), 16–21 24 H. Bolta, F. Koch, J. L. Rodet, D. Karpov, S. Menzel, Al2O3 coatings deposited by filtered vacuum arc-characterization of high tempe- rature properties, Surface and Coatings Technology, 116–119 (1999), 956–962 25 A. C. Karaoglanli, E. Altuncu, I. Ozdemir, A. Turk, F. Ustel, Struc- ture and durability evaluation of YSZ + Al2O3 composite TBCs with APS and HVOF bond coats under thermal cycling conditions, Surface and Coatings Technology, 205 (2011) 2, 369–S373 26 H. Guo, X. Bi, S. Gong, H. Xu, Microstructure Investigation On Gradient Porous Thermal Barrier Coating Prepared By EB-PVD, Scripta Mater., 44 (2001), 683– 687 A. C. KARAOÐLANLÝ et al.: STUDY OF THE MICROSTRUCTURE AND OXIDATION BEHAVIOR OF YSZ ... 444 Materiali in tehnologije / Materials and technology 46 (2012) 5, 439–444 K. ZUPAN, M. MARIN[EK: MICROSTRUCTURE DEVELOPMENT OF THE Ni-GDC ANODE MATERIAL FOR IT-SOFC MICROSTRUCTURE DEVELOPMENT OF THE Ni-GDC ANODE MATERIAL FOR IT-SOFC RAZVOJ MIKROSTRUKTURE Ni-GDC ANODNEGA MATERIALA ZA SREDNJETEMPERATURNE SOFC Klementina Zupan, Marjan Marin{ek Faculty of Chemistry and Chemical Technology, University of Ljubljana, A{ker~eva 5, 1000 Ljubljana, Slovenia klementina.zupan@fkkt.uni-lj.si Prejem rokopisa – received: 2011-12-02; sprejem za objavo – accepted for publication: 2012-03-02 The NiO-GDC-based material is a potential candidate for an anode material for the low-temperature SOFCs. In this work a modified combustion synthesis was used for the preparation of NiO-GDC. The main advantage of the preparation method employed was that after the synthesis both phases, NiO and GDC, in the ash product were randomly distributed on a nanometre scale. The citrate-nitrate (c/n) ratios in the combustion-reaction mixtures varied from 0.15 to 0.18. The prepared powders were isostatically pressed into pellets, sintered at 1200 °C, 1250 °C, 1300 °C, 1350 °C or 1400 °C, reduced and subsequently submitted to a microstructure analysis. The crystallite sizes of both phases in the as-prepared powders, as well as the grain sizes of nickel in the final reduced samples greatly depended on the slight variation of the c/n ratio in the starting reaction gel mixture. In the as-synthesized samples, crystallite sizes were calculated to be 4.3 nm or 40.0 nm for the GDC phase and 7.6 nm or 48.0 nm for the NiO phase for the samples with the c/n ratios of 0.15 or 0.18, respectively. After sintering under different conditions and reductions, the final average particle size of Ni varied from 71 nm to 146 nm or from 143 nm to 254 nm, while the average size of GDC grains ranged from 84 nm to 193 nm or from 96 nm to 247 nm for the samples with the c/n ratios of 0.15 or 0.18, respectively. The temperatures from 1200 °C to 1250 °C were recognized as the most appropriate temperature interval that provided good connectivity between the grains and the smallest one-phase regions in the final Ni-GDC cermets with an average Ni-particle diameter of around 70 nm. Keywords: combustion synthesis, nanocomposites, microstructure, fuel cells, Ni-GDC Materiali na osnovi NiO-GDC spadajo med potencialne kandidate za izdelavo anod v srednjetemperaturnih SOFC. NiO-GDC smo pripravili z modificirano zgorevalno sintezo. Najve~ja prednost metode je ta, da sta po sintezi obe fazi NiO in GDC naklju~no porazdeljeni na nanometrskem nivoju. Citratno-nitratno razmerje c/n v reakcijskih zmeseh je bilo 0,15 in 0,18. Pripravljeni prah smo po sintezi izostatsko stisnili v tablete, jih sintrali pri temperaturah 1200 °C, 1250 °C, 1300 °C, 1350 °C in 1400 °C, reducirali ter izvedli kvantitativno analizo mikrostruktur. Razmerje c/n v za~etni raztopini mo~no vpliva na velikost kristalitov faz (NiO in GDC) v vzorcu po sintezi, kot tudi na velikost zrn faz v sintranih in reduciranih vzorcih. Najmanj{a nikljeva zrna (povpre~na velikost okoli 70 nm) v kon~nem Ni-GDC-kompozitu keramika-kovina so nastala po sintranju in kasnej{i redukciji v temperaturnem intervalu med 1200 °C in 1250 °C. Klju~ne besede: zgorevalna sinteza, nanokompoziti, mikrostruktura, gorivne celice, Ni-GDC 1 INTRODUCTION Fuel cells are environmentally friendly energy con- verters that can transform chemical energy directly to electricity, resulting in high-energy conversion efficien- cies. Solid-oxide fuel cells (SOFCs) have several advantages over the other types of fuel cells, including flexibility of the fuels used, high-reaction kinetics due to high-temperature operation and relatively inexpensive materials formed in thin layers. However, the high- temperature operation results in a number of inherent challenges, such as mechanical stress, electrode sintering and low start-up time. First, it is difficult to obtain gas-tight seal between the chambers. Moreover, an addition of a large amount of steam to hydrocarbon fuels is required to avoid carbon deposition on the anode, resulting in a complicated water management in the SOFC systems.1,2 One approach to overcoming the above challenges is to reduce the operating temperatures of SOFCs to 800 °C or less to enable the metal materials (stainless steels) to be used as interconnect materials.3,4 Another approach towards addressing the above challenges is to design an SOFC with only one gas chamber. This type of SOFC is called a "single chamber SOFC" (SC-SOFC), wherein both anode and cathode are exposed to the same fuel-oxidant gas mixture. As a result, the gas-sealing problem can be inherently avoided since no separation between fuel and air is required, while carbon deposition is less of a problem due to the presence of a large amount of oxygen in the mixture. Materials used as the components in two-chamber SOFCs (LSM strontium- doped lanthanum manganite as a cathode, yttrium stabilized zirconia-YSZ as an electrolyte and Ni-YSZ cermet as an anode) are not appropriate for the use in an SC-SOFC due to their insufficient selectivity toward electrode reactions under operating conditions. However, there are only a few materials selective enough that can operate at the temperatures below 700 °C. Currently the most widely adopted SC-SOFC materials are gadolinium stabilized ceria (GDC) as electrolyte, Ni-GDC as anode and strontium-doped samarium cobaltite (SSC) as cathode. GDC is a better oxygen ionic conductor than Materiali in tehnologije / Materials and technology 46 (2012) 5, 445–451 445 UDK 546:66.017:669.018.95 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)445(2012) YSZ, as it has a higher ionic conductivity than YSZ in the temperature range from 300 °C to 700 °C. Anode materials are usually prepared by mechani- cally mixing the separately synthesized powders of NiO and GDC that were formed, sintered and reduced to obtain Ni-GDC cermets.5–7 In this way, an accurate chemical composition can be attained, but it is rather difficult to achieve a uniform distribution of separate phases in the anode composite that will result in a non-homogeneous structure and a poor anode perfor- mance. However, homogeneous composite powders of NiO-GDC were synthesized using a hydroxide co-preci- pitation.8,9 The co-precipitation method is still a solution-based technique with several consecutive steps and, as such, it is rather time consuming. Consequently, the preparation of complex metal oxides by using combustion synthesis has recently become an interesting area of research based on the promising results of this technique surpassing the conventional method.10,11 The main advantage of the combustion synthesis is the ability to produce complex oxide powders directly from the precursor solution. Therefore, the combustion synthesis could be, in principle, a good method for preparing the composite powder of Ni-GDC with a uniform distri- bution of fine Ni particles within the GDC framework. The aim of this work was a synthesis and a sub- sequent thermal treatment of nano-scaled highly sinterable NiO-GDC dispersions prepared with the citrate-nitrate combustion synthesis technique. High sinterability of the prepared powders is of the prime importance for the preparation of dense bodies at the relatively low temperatures (close to 1200 °C). With an examination of the densification process of such ceramic powders, we also demonstrated a microstructure evolution from the synthesized NiO-GDC nano-powder to the Ni-GDC anode cermet. 2 EXPERIMENTAL PROCEDURES The NiO-GDC powders were prepared with a modified citrate-nitrate combustion synthesis. The combustion system was based on the citrate-nitrate redox reaction. In this combustion method, the starting substances were Ni(NO3)2 · 6H2O, Ce(NO3)3 · 6H2O and Gd(NO3)3 · 6H2O, nitric acid (65 %) and citric acid (analytical reagent grade). All solid compounds were dissolved with minimum additions of water in the amounts that allowed the Ni volume content in the final composite to be 50 %. The Ce(NO3)3 · 6H2O and Gd(NO3)3 · 6H2O additions were found to allow the GDC composition to be Ce0.8Gd0.2O1.9. The reactive mixture was prepared by mixing the five reactant solutions and then kept over a water bath at 60 °C under vacuum (p = 5–7 mbar) until it transformed into a light green gel (at least 5 hours). The initial citrate/nitrate molar ratios in the starting solution were 0.15 and 0.18. The correspond- ing citrate-nitrate gel was then gently milled in an agate mortar and uni-axially pressed into pellets (F = 12 mm, h = 30 mm, p = 17 MPa). These samples were placed on a corundum plate and ignited at the top of the pellet with a hot tip to start an auto-ignition reaction. The samples were characterized by the X-ray powder diffraction technique using a PANalytical X’Pert PRO MPD appara- tus. Data were collected in the 2 range from 20° to 70° in steps of 0.033° for 1 s/°. After the synthesis, powders were milled in an agate mortar, uni-axially pressed into pellets (100 MPa) and subsequently also iso-statically pressed (750 MPa). Formed pellets were sintered at different temperatures (1200 °C, 1250 °C, 1300 °C, 1350 °C and 1400 °C) for 2 hours. Material shrinkage during the sintering was measured separately using a BÄHR DIL 802 dilatometer. For the microstructure determi- nation, the sintered tablets were polished (3 μm and 0.25 μm diamond pastes), thermally etched, reduced at 900 °C (2 h) in an H2/Ar atmosphere and subsequently analyzed with SEM FE Zeiss ULTRA plus. The quanti- tative analyses of the microstructures were performed on digital images (images were digitized into pixels with 255 different grey values) using the Zeiss KS300 3.0 image-analysis software. 3 RESULTS AND DISCUSSION The main benefit of the combustion synthesis is the exothermic effect during the fuel (citrate) and nitrate redox reaction that is accompanied by a considerable gas release preventing the formation of hard agglomerates. The combustion reaction takes just a few seconds to turn a reaction mixture into the final product, while a much longer time is required for the same process to be conducted during the calcination process. In the case of an NiO-GDC citrate-nitrate self-sustaining reaction, the maximum temperature gradient (i.e., the heating rate) K. ZUPAN, M. MARIN[EK: MICROSTRUCTURE DEVELOPMENT OF THE Ni-GDC ANODE MATERIAL FOR IT-SOFC 446 Materiali in tehnologije / Materials and technology 46 (2012) 5, 445–451 Figure 1: XRD diffraction spectra of the as-prepared powder mixtures of NiO-GDC (obtained from the reactive gels with c/n = 0.15 or c/n = 0.18) Slika 1: Rentgenska pra{kovna posnetka vzorcev NiO-GDC, priprav- ljenih iz raztopin s citratno-nitratnim razmerjem c/n = 0,15 in c/n = 0,18 inside the reaction zone (calculated on the basis of the temperature-profile measurements) was 1164 K s–1. Such a high temperature gradient and short reaction times resulted in a unique powder mixture composed of nano-sized particles. The one-phase particle size and the degree of crystallization were found to be the functions of the citrate-nitrate initial ratio. In both samples (c/n = 0.18 or 0.15) the two main phases corresponded to GDC and NiO (Figure 1). According to the Sherer X-ray broadening of the peaks, the average crystallite sizes were calculated to be 4.3 nm and 40 nm for GDC and 7.6 nm and 48 nm for NiO for the samples with the c/n ratios of 0.15 and 0.18, respectively. The ideal microstructure of the final Ni-GDC cermet is composed of very small grains of both phases (preferably sub-micrometre or nano-sized), where GDC mainly acts as a matrix to support the Ni electro-catalyst and hinder its fusion under the cell operating conditions. At the same time, GDC is also used to extend the Ni-GDC-gas triple-phase boundary (TPB) into the anode. The length of the TPB is related to the reaction rate for the electrochemical oxidation of hydrogen12,13 in an operating fuel cell. For this reason, Ni-GDC also has to exhibit continuity of both phases throughout the cermet, since it must serve both as an electronically and ionically conductive material. In principle, high conductivity is achieved when a good contact between particles is ensured; this is normally accomplished through sintering. However, sintering also means grain growth, which is in contradiction to the desire to preserve the NiO-GDC nano-distribution. However, if the sintering temperatures required for the NiO-GDC densification can be lowered, then fine oxide mixtures can also be preserved in the sintered structures. The determined sintering temperature of the prepared NiO-GDC mixture was influenced by a slight variation of the citrate-nitrate ratio (0.15 or 0.18) in the starting solutions. Being sintered at a significantly lower temperature (Tsintr. = 1120 °C for the sample with the c/n ratio of 0.15 and Tsintr. = 1280 °C for the sample with the c/n ratio of 0.18), the sample 0.15 sinters through two consecutive stages as a result of the inter- and/or intra-agglomerate sintering, while the shrinkage in sample 0.18 can be described with only one broad densification process (Figure 2). The relatively low sintering temperature determined for the sample with the c/n ratio of 0.15 may be very important from the applicability point of view, since a single cell is normally prepared in several sintering processes. First, the anode and electrolyte layers are co-sintered at higher temperatures (up to 1400 °C) and then the cathode layer is applied and co-sintered at the temperatures of up to 1200 °C. Successful sintering of the NiO-GDC anode material at the temperatures below 1200 °C may result in diverse SOFC-preparation procedures, in which all the layers are co-sintered in a single step. In order to obtain more detailed information about the microstructure development, the NiO-GDC green bodies prepared from the as-synthesized powders were sintered under various sintering conditions. Sintering temperatures were defined according to the obtained shrinkage curves. The presented microstructures (Figure 3) revealed that the nano-sized cermet mixture was no longer present in the sintered and, subsequently, reduced cermets. Instead, the one-phase dominance grew, but remained well within the sub-micrometre range if the cermets had been sintered at the relatively low temperatures. Additionally, it is evident that a fine distribution of phases (metal, ceramic and porosity) can be obtained when using the sintering temperatures no higher than 1200–1250 °C for the samples with a c/n ratio of 0.15, in which GDC serves as a continuous framework, within which the sub-micrometre Ni grains and pores are dispersed. However, in the case of the samples with a c/n ratio of 0.18 (sintered at 1200 °C and 1250 °C) an excessive grain growth of the newly formed metallic Ni particles (during the reduction) may be found. This phenomenon is a consequence of a fairly different surface energy of Ni and GDC, causing the hetero-grains at the interface to lose their chemical affinity. Consequently, the coarsening of the Ni phase proceeds appreciably due to a poor adhesion of the metal to the ceramic material at the elevated temperatures. From a practical point of view, an excessive growth of the Ni grains results in the TPB length reduction. In the samples with a c/n ratio of 0.18, the grain growth of Ni was noticeable, while in the samples with a c/n ratio of 0.15, the grain growth of both phases (in the reduced state) was controlled predominantly with the sintering temperature and no excessive enlargement of Ni-grains, after the reduction, was observed. Dense structures after sintering ensure good contact between the particles and a continuity of both phases K. ZUPAN, M. MARIN[EK: MICROSTRUCTURE DEVELOPMENT OF THE Ni-GDC ANODE MATERIAL FOR IT-SOFC Materiali in tehnologije / Materials and technology 46 (2012) 5, 445–451 447 Figure 2: Relative shrinkage versus temperature of the NiO-GDC tablets (iso-statically pressed) prepared from the reaction mixtures with the c/n ratios of 0.18 or 0.15 Slika 2: Relativni skr~ek NiO-GDC izostatsko stisnjenih vzorcev, pripravljenih iz reakcijskih zmesi s citratno-nitratnim razmerjem c/n 0,18 in 0,15 K. ZUPAN, M. MARIN[EK: MICROSTRUCTURE DEVELOPMENT OF THE Ni-GDC ANODE MATERIAL FOR IT-SOFC 448 Materiali in tehnologije / Materials and technology 46 (2012) 5, 445–451 Figure 3: Microstructures of the sintered Ni-GDC samples after the reduction Slika 3: Mikrostrukture sintranih vzorcev Ni-GDC po redukciji (NiO and GDC). The continuity of the phases is essential not only with regard to conductivity, but also for the microstructure stability. The reason for the inhomo- geneous microstructure in the reduced samples may be due to the inhomogeneous and partly porous micro- structure after sintering. One good example of an insufficiently stable NiO-GDC microstructure is shown in the case of a sample with a c/n ratio of 0.18 (sintered at 1250 °C), in which more porous regions, accompanied with the grains of both phases that are significantly larger than in an average sample, can be found (Figure 4). One possible reason for the differences in the microstructure development in the samples with a c/n ratio of 0.18 or 0.15 is in the nature of the combustion reaction itself. A slight increase in the citrate-nitrate molar ratio from 0.15 to 0.18 may lead to a carbon residue after the synthesis (0.18 sample), followed by a porosity enlargement due to carbon burning during the sintering process. Additionally, the peak combustion temperature in the reaction mixture with a c/n ratio of 0.18 is close to 1180 °C, while the peak temperature in the reaction mixture with a c/n ratio of 0.15 does not exceed 600 °C. Higher combustion temperatures and higher temperature gradients within the combusting reaction mixture may cause the formation of hard agglomerates. The microstructure parameters important for an exact cermet analysis and obtained with a detailed quantitative microstructure analysis of the sintered and reduced samples are summarized in Tables 1 and 2. For statistic- ally reliable data, several different regions were analyzed in each case. Parameters d , dx and dy, dpor and  are represented as the diameter of the area-analogue circle – DCIRCLE, the intercept lengths in the x and y directions – FERETX, Y, and the maximal intercept length – FERETMAX and FCIRCLE, respectively. Porosity  was determined as the microstructural porosity (on the basis of a microstructure analysis). According to the results of the quantitative micro- structure analysis, the porosity of the sintered samples decreased with the sintering temperature as expected. The higher sintering temperature also resulted in a pronounced grain growth. The average particle size of NiO exceeds the GDC average particle size in all the samples; however, the difference between average particle sizes is less pronounced at higher sintering temperatures. From this fact we can deduce that NiO in a NiO-GDC composite sinters first and more intensely. The average particle sizes of NiO and GDC in the samples with a c/n ratio of 0.15 range from 92 nm to 230 nm and from 84 nm to 193 nm, respectively, while in the samples with a c/n ratio of 0.18 they range from 127 nm to 249 nm and from 96 nm to 247 nm, respectively. K. ZUPAN, M. MARIN[EK: MICROSTRUCTURE DEVELOPMENT OF THE Ni-GDC ANODE MATERIAL FOR IT-SOFC Materiali in tehnologije / Materials and technology 46 (2012) 5, 445–451 449 Figure 4: Microstructures of the samples with the c/n ratios of 0.18 and 0.15 sintered at 1250 °C Slika 4: Mikrostrukture vzorcev (razmerje c/n: 0,18 in 0,15) sintranih pri 1250 °C Additionally, with higher sintering temperatures, the grains of both phases became rounder ( parameter). Comparing the sintered and reduced samples, the size of the Ni-grains is reduced relative to the NiO particle size and, consequently, the porosity is increased. When analysing the microstructure of the metal-ceramic anode layer, we find that its appropriate volume porosity is between 30 % and 40 % 14–17. Considering that about 41.1 % of the initial NiO volume is transformed into pores during the reduction of NiO to Ni, we find that the porosity of the sintered materials should be close to 10 %. For a high TPB value, an average particle size of both phases should remain as small as possible; for suitable electrical properties, the contact between the particles should be good. For these reasons, the sintering tempera- ture for a sample with a c/n ratio of 0.15 is in the range from 1200 °C to 1250 °C, while for a sample with a c/n ratio of 0.18 the sintering temperature below 1300 °C may be considered to be too low due to the clear microstructural instability of the final Ni-GDC compo- site, which was characterized for its exaggerated Ni growth. The final average particle sizes of Ni and GDC in a sample with a c/n ratio of 0.15 were 73 nm and 86 nm after the sintering at 1200 °C , and 71 nm and 97 nm after the sintering at 1250 °C, respectively. Nevertheless, the nano-sized initial oxide mixture and careful powder treatment enabled the preparation of well-sintered materials with the relative densities above 90 %, as well as subsequently reduced composites, in which the average particle sizes of both phases are still well within the sub-micrometer range. 4 CONCLUSIONS The nano-sized NiO-GDC composites were prepared using a citrate-nitrate self-sustaining reaction from the initial reactive gel. After the synthesis, the average crystallite sizes were calculated to be 4.3 nm and 40 nm for GDC and 7.6 and 48 nm for NiO for the samples with the c/n ratios of 0.15 and 0.18, respectively. Such crystallites, partially agglomerated, were the starting material for a pellet preparation and the subsequent microstructure-development investigations. Relatively dense bodies (not more than 10 % of the residual porosity) were prepared at the sintering temperatures as low as 1200 °C for the samples with a c/n ratio of 0.15. Higher sintering temperatures did not significantly decrease the porosity. For the samples with a c/n ratio of 0.18, the sintering temperatures lower than 1300 °C were recognised as insufficient due to a clear microstructure instability, where more porous regions accompanied with the grains of both phases, larger than in an average sample, can be found. The average grain size of the GDC K. ZUPAN, M. MARIN[EK: MICROSTRUCTURE DEVELOPMENT OF THE Ni-GDC ANODE MATERIAL FOR IT-SOFC 450 Materiali in tehnologije / Materials and technology 46 (2012) 5, 445–451 Table 1: Quantitative microstructure analysis of the sintered NiO-GDC samples Tabela 1: Rezultati kvantitativne analize mikrostruktur sintranih vzorcev NiO-GDC T/° C c/n /% Mikr. (nm)  dx/nm dy/nm dpor/nmNiO GDC NiO GDC NiO GDC NiO GDC 1200 0.15 12 92 84 0.70 0.70 107 98 106 97 187 1250 0.15 11 121 95 0.77 0.78 138 108 137 108 158 1300 0.15 8 143 128 0.75 0.74 161 147 162 146 167 1350 0.15 6 189 163 0.80 0.76 210 186 209 186 177 1400 0.15 3 230 193 0.78 0.71 257 223 255 227 236 1200 0.18 13 127 96 0.71 0.72 142 107 143 108 358 1250 0.18 10 155 122 0.72 0.73 174 137 176 138 274 1300 0.18 7 182 152 0.72 0.73 203 173 204 172 335 1350 0.18 6 214 197 0.74 0.73 235 218 235 216 382 1400 0.18 4 249 247 0.77 0.72 269 267 268 267 370 Table 2: Quantitative microstructure analysis of the reduced Ni-GDC samples Tabela 2: Rezultati kvantitativne analize mikrostruktur reduciranih vzorcev Ni-GDC T /°C c/n /% Mikr. (nm) ø dx/nm dy/nm dpor/nmNi GDC Ni GDC Ni GDC Ni GDC 1200 0.15 34 73 86 0.75 0.70 84 102 80 100 352 1250 0.15 29 71 97 0.75 0.78 82 108 77 108 304 1300 0.15 26 85 125 0.71 0.79 100 140 100 137 351 1350 0.15 25 109 154 0.72 0.79 133 176 125 165 407 1400 0.15 23 146 192 0.72 0.78 173 217 168 214 646 1200 0.18 40 169 95 0.70 0.71 191 106 203 107 378 1250 0.18 36 143 121 0.71 0.76 160 135 165 135 438 1300 0.18 34 168 150 0.74 0.78 194 168 183 167 385 1350 0.18 35 193 195 0.75 0.77 224 217 193 216 413 1400 0.18 21 254 244 0.76 0.78 301 264 275 260 361 phase increased with the increasing sintering tempe- rature. The NiO grains also grew with the increasing sintering temperature; however, the dimensions of the final Ni grains formed during the reduction were very much influenced by the microstructure stability of the GDC framework formed during the sintering. The smallest Ni grains, still well in the sub-micrometer range, with the average particle sizes of 73 nm and 71 nm were obtained in the sample with a c/n ratio of 0.15 sintered at 1200 °C and 1250 °C, respectively. Acknowledgements This investigation was supported by the Centre of Excellence for Low-Carbon Technologies and the Slo- venian Research Agency. 5 REFERENCES 1 J. F. Fergus, Oxide anode materials for solid oxide fuel cells, Solid State Ionics, 177 (2006), 1529–1541 2 M. Yano, A. Tomita, M. Sano, T. Hibino, Recent advances in single- chamber solid oxide fuel cells: A review, Solid State Ionics, 177 (2007), 3351–3359 3 S. Ping Jiang, A review of wet impregnation – An alternative method for the fabrication of high performance and nano-structured electrodes of solid oxide fuel cells, Materials Science and Engi- neering A, 418 (2006), 199–210 4 Y. J. Leng, S. H. Chan, S. P. Jiang, K. A. Khor, Low-temperature SOFC with thin film GDC electrolyte prepared in situ by solid-state reaction, Solid State Ionics, 170 (2004), 9–15 5 T. Ishihara, T. Shibayama, H. Nishiguchi, Y. Takita, Nickel-Gd- doped CeO2 cermet anode for intermediate temperature operating solid oxide fuel cells using LaGaO3-based perovskite electrolyte, Solid State Ionics, 132 (2000), 209–216 6 C. Xia, M. Lui, Microstructures, conductivities, and electrochemical properties of Ce0.9Gd0.1O2 and GDC-Ni anodes for low-temperature SOFCs, Solid State Ionics, 152–153 (2002), 423–430 7 S. Zha, W. Rauch, M. Liu, Ni-Ce0.9Gd0.1O1.95 anode for GDC elec- troyte-based low-temperature SOFCs, Solid State Ionics, 166 (2004), 241–250 8 C. Ding, H. Lin, K. Sato, T. Hashida, Synthesis of NiO- Ce0.9Gd0.1O1.95 nanocomposite powders for low-temperature solid oxide fuel cell anodes by co-precipitation, Scripta Materialia, 60 (2009), 254–256 9 C. Ding, H. Lin, K. Sato, T. Kawada, J. Mizusaki, T. Hashida, Impro- vement of electrochemical performance of anode-supported SOFCs by NiO-Ce0.9Gd0.1O1.95 nanocomposite powders, Solid State Ionics, 181 (2010), 1238–1243 10 A. Ringuede, J. A. Labrincha, J. R. Frade, A combustion synthesis method to obtain alternative cermet materials for SOFC anodes, Solid State Ionics, 141-142 (2001), 549–557 11 P. Duran, J. Tartaj, F. Capel, C. Moure, Processing and characteri- zation of a fine nickel oxide/zirconia/composite prepared by polymeric complex solution synthesis, J. Eur. Ceram. Soc., 23 (2003), 2125–2133 12 S. P. Jiang, Y. Y. Duan, J. G. Love, Fabrication of high-performance Ni/Y2O3-ZrO2 cermet anodes of solid oxide fuel cells by ion impregnation, J. Electrochem. Soc., 149 (2002), A1175 13 B. de Boer, M. Gonzales, H. J. M. Bouwmeester, H. Verweij, The effect of the presence of fine YSZ particles on the performance of porous nickel electrodes, Solid State Ionics, 127 (2000), 269–276 14 W. Hu, H. Guan, X. Sun, S. Li, M. Fukumoto, I. Okane, Electrical and Thermal Conductivities of Nickel-Zirconia Cermets, J. Am. Ceram. Soc., 81 (1998) 8, 2209–2212 15 D. W. Dees, T. D. Claar, T. E. Eas~er, D. C. Fee, F. C. Mrazek, Conductivity of Porous Ni/ZrO2-Y2O3 Cermets, J. Electrochem. Soc., 134 (1987) 9, 2141–2146 16 T. Kawashima, M. Hishinuma, Analysis of Electrical Conduction Paths in Ni/YSZ Participate Composites Using Percolation Theory, Mater. Trans., 37 (1996) 7, 1397–1403 17 U. Anselmi-Tamburini, G. Chiodelli, M. Arimondi, F. Maglia, G. Spinolo, Z. A. Munir, Electrical Properties of Ni/YSZ Cermets Obtained by Combustion Synthesis, Solid State Ionics, 110 (1998), 35–43 K. ZUPAN, M. MARIN[EK: MICROSTRUCTURE DEVELOPMENT OF THE Ni-GDC ANODE MATERIAL FOR IT-SOFC Materiali in tehnologije / Materials and technology 46 (2012) 5, 445–451 451 M. KOVA^I^, S. SEN^I^: MODELING OF PM10 EMISSION WITH GENETIC PROGRAMMING MODELING OF PM10 EMISSION WITH GENETIC PROGRAMMING MODELIRANJE EMISIJE PM10 Z GENETSKIM PROGRAMIRANJEM Miha Kova~i~1, Sandra Sen~i~2 1[TORE STEEL, d. o. o., @elezarska cesta 3, 3220 [tore, Slovenia 2KOVA, d. o. o., Teharska cesta 4, 3000 Celje, Slovenia miha.kovacic@store-steel.si Prejem rokopisa – received: 2012-01-18; sprejem za objavo – accepted for publication: 2012-04-05 To implement sound air-quality policies, regulatory agencies require tools to evaluate the outcomes and costs associated with various emission-reduction strategies. However, the applicability of such tools can also remain uncertain. It is furthermore known that source-receptor models cannot be implemented through deterministic modeling. The article presents an attempt of PM10 emission modeling carried close to a steel production area with the genetic programming method. The daily PM10 concentrations, daily rolling mill and steel plant production, meteorological data (wind speed and direction – hourly average, air temperature – hourly average and rainfall – daily average), weekday and month number were used for modeling during a monitoring campaign of almost half a year (23. 6. 2010 to 12. 12. 2010). The genetic programming modeling results show good agreement with measured daily PM10 concentrations. In future we will carry out genetic programming based dispersion modeling according to the calculated wind field, air temperature, humidity and rainfall in a 3D Cartesian coordinate system. The prospects for arriving at a robust and faster alternative to the well-known Lagrangian and Gaussian dispersion models are optimistic. Keywords: steel plant, PM10 concentrations, modeling, genetic programming V okviru uveljavljanja uredb o kvaliteti zraka, s ciljem zmanj{evanja emisij, nadzorne agencije zahtevajo ovrednotenje emisij in stro{kov, povezanih z njimi. Uporabnost takih orodij je v splo{nem negotova. Prav tako je znano, da pri modelih tipa vir-sprejemnik te`ko uporabimo deterministi~no modeliranje. V ~lanku je predstavljen poskus modeliranja emisije delcev PM10 na podro~ju `elezarne z metodo genetskega programiranja. Osnova za modeliranje so bili podatki, zbrani v obdobju ve~ kot pol leta (od 23. 6. 2010 do 12. 12. 2010): dnevne koncentracije PM10, produktivnost jeklarne, valjarne, meteorolo{ki podatki (hitrost in smer vetra, temperatura zraka – urno povpre~je ter padavine – dnevno povpre~je) ter dan v tednu in zaporedna {tevilka meseca. Rezultati modeliranja dnevnih koncentracij PM10 z genetskim programiranjem ka`ejo na dobro ujemanje z eksperimentalnimi podatki. V prihodnosti bomo izvedli modeliranje z genetskim programiranjem v kartezijskem 3D koordinatnem sistemu z upo{tevanjem izra~unanega vetrovnega polja, temperature zraka, vla`nosti in padavin. Mo`nosti za uporabo robustnih in hitrej{ih alternativ Lagrangovih in Gaussovih disperzijskih modelov so optimisti~ne. Klju~ne besede: `elezarna, koncentracije PM10, modeliranje, genetsko programiranje 1 INTRODUCTION Particulate matter (PM) pollution is, especially in residential areas near industrial areas, a problem of great concern. This is not only because of the adverse health effects but also because of reduced visibility; on a global scale, effects on the radiative balance are also of great importance1–3. To reduce PM levels in the air a deep knowledge of the contributing sources, background emissions, the influence of the meteorological conditions, as well as of PM10 formation and transport processes is needed. However, current state-of-the-art PM10 modeling does not allow us to quantitatively model the whole range of emissions behavior, which is why the dispersion modeling is thus increasingly connected with intelligent algorithms such as artificial neural networks4–9 and evolutionary computation9. The objective of this work was to model PM10 emissions close to a steel plant area in Slovenia by means of a genetic programming method. Genetic pro- gramming has been proven to be an effective optimi- zation tool for multicriterial and multiparametrical problems10–13. The genetic programming system for PM10 emission modeling imitates the natural evolution of living organisms, where in the struggle for natural resources the successful entities gradually become more and more dominant in adapting to the environment in which they live; the less successful ones, meanwhile, are only rarely present in subsequent generations. In the proposed concept the mathematical models for PM10 concentration prediction undergo adaptation. During the simulated evolution more and more successful organisms (PM10 emission models) emerge on the basis of given data (wind speed and direction – hourly average, air temperature – hourly average, rainfall – daily average, weekday and month number). In order to allow for a self-contained paper the basic terms and experimental setup are stated in the beginning. Afterwards the idea of the proposed concept is presented. In the conclusion the main contributions of the per- formed research are summarized, while guidelines for further research are provided. Materiali in tehnologije / Materials and technology 46 (2012) 5, 453–457 453 UDK 669:519.61/.64:351.777.6 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)453(2012) 2 EXPERIMENTAL SETUP 2.1 Sampling sites Figure 1 shows the locations of the sampling sites, rolling mill, steel plant and residential areas. Due to rolling mill and steel plant PM10 contribution also several source categories influence the PM10 concen- trations. These include combustion and non-combustion traffic sources, urban background concentrations, along with both contributions that are transported by regionally and long-range. 2.2 Sampling Samples for this study were collected between 23. 6. 2010 and 12. 12. 2010. Sampling was performed 1.5 m above the ground. PM10 samples were collected for 24 h on Mondays using low-volume samplers equipped with EPA-equivalent size-selective inlets. Particles with diameter 10 μm (PM10) were collected on cellulose esters membranes with high collection efficiencies (99 %). In total 172 PM10 samples for each sampling site were available. Before and after the samplings were made the filters were exposed for 24–48 h on open but dust-protected sieve-trays in an air-conditioned weighing room. The gravimetric determination of the mass was carried out using an analytical microbalance (precision 1 μg) located in the weighing room. In order to remove static electricity from filters the balance is equipped with a special kit in a Faraday shield. The limit value of the EU directive – i.e. a daily mean PM10 concentration – is 50 μg/m3. At the sampling site 1 and 2 the measured PM10 concentration exceeded limit value four times and five times, respectively. Figure 2 shows the measured PM10 concentrations during the study period for the sampling sites. 2.3 Meteorological data Hourly average air temperature, wind speed and direction and daily rainfall data were made available to the authors by the Slovenian Environment Agency. Figure 3 shows the hourly average temperatures during the study period. Figure 4 shows the frequency distribution of wind direction and wind speed obtained based on wind direction and speed data measured every hour during the study period. Figure 5 shows the daily rainfall during the period of the study. The hourly data based on electric arc and rolling mill production was collected during the study period. During M. KOVA^I^, S. SEN^I^: MODELING OF PM10 EMISSION WITH GENETIC PROGRAMMING 454 Materiali in tehnologije / Materials and technology 46 (2012) 5, 453–457 Figure 4: Frequency distribution of wind direction and wind speed Slika 4: Frekven~na porazdelitev smeri in hitrosti vetra Figure 2: The measured PM10 concentrations during the study period for the sampling sites Slika 2: Izmerjene koncentracije PM10 v obdobju {tudije za lokaciji vzor~enja Figure 3: The hourly average temperatures during the study period Slika 3: Urno povpre~je temperature zraka v obdobju {tudije Figure 1: Topographic view of the study area Slika 1: Topografski prikaz podro~ja {tudije the study period, the electric arc furnace was stopped for 28 465 min and the rolling mill was stopped for 8 213 min. Figure 6 shows the minutes of stopping per day for the electric arc furnace and rolling mill during the study period. 3 GENETIC PROGRAMMING MODELING Genetic programming is probably the most general evolutionary optimization method. The organisms that undergo adaptation are in fact mathematical expressions (models) for the PM10 concentrations prediction in the present work. The concentration prediction is based on the available function genes (i.e., basic arithmetical functions) and terminal genes (i.e., independent input parameters, and random floating-point constants). In the present case the models consist of the following function genes: addition (+), subtraction (–), multiplication (*) and division (/), and the following terminal genes: weekday (WEEKDAY) and month number (MONTH), wind speed [m/s] (SPEED), wind direction [°] (DIRECTION), air temperature [°C] (TEMP), rainfall [mL] (RAIN), electro arc furnace efficiency [min/h] (EAF), rolling mill efficiency [min/h] (ROLLING). In order to ascertain the influence of seasons and traffic during workday hours the weekday and month number were also added as terminal genes. One of the randomly generated mathematical models – # – is schematically represented in Figure 7 as a program tree with included function genes (*, + ,/) and terminal genes (TEMP, RAIN, EAF and a real number constants 2 and 5.1). Random computer programs of various forms and lengths are generated by means of the selected genes at the beginning of the simulated evolution. The varying of the computer programs is performed by means of the genetic operations during several iterations, known as generations. After the completion of the variation of the computer programs a new generation is obtained. Each generation is compared with the experimental data. The process of changing and evaluating organisms is repeated until the termination criterion of the process is fulfilled. The maximum number of generations is chosen as a termination criterion in the present algorithm. The following evolutionary parameters were selected for the process of simulated evolutions: 500 for the size of the population of organisms, 100 for the maximum number of generations, 0.4 for the reproduction pro- bability, 0.6 for the crossover probability, 6 for the maximum permissible depth in the creation of the population, 10 for the maximum permissible depth after the operation of crossover of two organisms, and 2 for the smallest permissible depth of organisms in gene- rating new organisms. Genetic operations of repro- duction and crossover were used. For selection of organisms the tournament method with tournament size 7 was used9–13. 100 independent civilizations of mathe- matical models for prediction of the PM10 concentration were developed. The best evolution sequence of 100 generations was computed in 8 h and 41 min on 2.39 GHz processor and 2 GB of RAM by an AutoLISP based in-house coded computer program. The model fitness f has been defined as: f P M Ni i i n = − + ⋅ = ∑ ( ) 10000 1 (1) where n is the size of sample data and, Pi is predicted PM10 concentration, Mi is measured PM10 concen- tration and N is the number of all cases when: P M M Pi i i i< ∧ > ∨ < ∧ >50 50 50 50 (2) The limit value of the EU directive, i.e. a daily mean PM10 concentration, is 50 μg/m3. The number N tells us when the prediction is above that limit value, when in M. KOVA^I^, S. SEN^I^: MODELING OF PM10 EMISSION WITH GENETIC PROGRAMMING Materiali in tehnologije / Materials and technology 46 (2012) 5, 453–457 455 Figure 6: Minutes of stopping per day for the electric arc furnace and rolling mill during the study period Slika 6: Dnevni zastoji elektrooblo~ne pe~i in valjarne v minutah v obdobju {tudije Figure 7: Randomly generated mathematical model for the PM10 concentrations prediction, represented in program tree form. Slika 7: Naklju~no ustvarjen matemati~ni model napovedovanja koncentracije PM10, predstavljen kot programsko drevo Figure 5: Daily rainfall during the study period Slika 5: Dnevne padavine v obodbju {tudije order to assure PM10 concentration exceedance predic- tion by developed predictive model it should in fact be below the limit, and also when prediction by developed predictive model should be above the limit. The simulated evolution in one run of the genetic programming system (out of 100) produced the follow- ing best model for prediction of PM10 concentration for sampling site 1: (3) with fitness of 1019.95, number N = 0 and average deviation of 5.96 μg/m3. The best evolutionary developed model (out of 100) for prediction of PM10 concentration for sampling site 1 is: (4) with fitness of 11 124.67, number N = 1 (on the 30. 6. 2010 the measured PM10 concentrations were 53.6 μg/m3 and predicted 21.41 μg/m3), and average devi- ation of 6.54 μg/m3. Figures 8 and 9 show measured and predicted PM10 concentrations for sampling sites 1 and 2, respectively. 4 CONCLUSIONS This paper presented the possibility of the PM10 concentration prediction close to a steel plant area with genetic programming. The daily PM10 concentrations, daily rolling mill and steel plant production, meteoro- logical data (wind speed and direction – hourly average, air temperature – hourly average and rainfall – daily average), weekday and month number were used for modeling during a monitoring campaign of almost half a year (23. 6. 2010 to 12. 12. 2010). The special fitness function for genetic programming system was designed in order to assure also PM10 limit value exceedance prediction. For each sampling site the best models for PM10 prediction were obtained from 100 runs of the genetic programming system. The model for sampling sites 1 and 2 predicts concentrations within an average error range of 5.96 μg/m3 and 6.54 μg/m3, respectively. All exceedances of the EU directive limit value (50 μg/m3) were administered at sampling site 1, but only 4 out of 5 of these occurred at sampling site 2. In the future we will carry out genetic programming based dispersion modeling according to the calculated wind field, air temperature, humidity and rainfall in a 3D Cartesian coordinate system. The prospects for arriving at a robust and faster alternative to the well-known La- grangian and Gaussian dispersion models are optimistic. 5 REFERENCES 1 G. M. Marcazzan, M. Ceriani, G.Valli, R.Vecchi, Source apportion- ment of PM10 and PM2.5 in Milan (Italy) using receptor modeling, The Science of the Total Environment, 317 (2003) 1–3, 137–147 2 J. G. Watson, Visibility: science and regulation, Journal of the Air and Waste Management Association, 52 (2002) 6, 628–713 3 E. Vrins, N. Schofield, Fugitive dust emission by an ironmaking site, Journal of Aerosol Science, 31 (2000), 524–525 M. KOVA^I^, S. SEN^I^: MODELING OF PM10 EMISSION WITH GENETIC PROGRAMMING 456 Materiali in tehnologije / Materials and technology 46 (2012) 5, 453–457 Figure 9: Measured and predicted PM10 concentrations [μg/m3] for sampling site 2 Slika 9: Izmerjene in napovedane koncentracije PM10 [μg/m3] za lokacijo vzor~enja 2 Figure 8: Measured and predicted PM10 concentrations [μg/m3] for sampling site 1 Slika 8: Izmerjene in napovedane koncentracije PM10 [μg/m3] za lokacijo vzor~enja 1 4 J. Kukkonena, L. Partanena, A. Karppinena, J. Ruuskanenb, H. Junninenb, M. Kolehmainenb, H. Niskab, S. Dorlingc, T. Chatter- tonc, R. Foxalld, G. Cawleyd, Extensive evaluation of neural network models for the prediction of NO2 and PM10 concentrations, compared with a deterministic modelling system and measurements in central Helsinki, Atmospheric Environment, 37 (2003), 4539–4550 5 H. Zhou, K. Cen, J. Fan, Modeling and optimization of the NOx emission, characteristics of a tangentially fired boiler with artificial, neural networks, Energy, 29 (2004), 167–183 6 J. Hooyberghsa, C. Mensinka, G. Dumontb, F. Fierensb, O. Brasseurc, A neural network forecast for daily average PM10 concentrations in Belgium, Atmospheric Environment, 39 (2005), 3279–3289 7 P. Perez, J. Reyes, An integrated neural network model for PM10 forecasting, Atmospheric Environment, 40 (2006), 2845–2851 8 G. Grivas, A. Chaloulakou, Artificial neural network models for prediction of PM10 hourly concentrations, in the Greater Area of Athens, Greece, Atmospheric Environment, 40 (2006), 1216–1229 9 M. Kova~i~, P. Uratnik, M. Brezo~nik, R. Turk, Prediction of the bending capability of rolled metal sheet by genetic programming, Materials and Manufacturing Processes, 22 (2007), 634–640 10 M. Kova~i~, B. [arler, Application of the genetic programming for increasing the soft annealing productivity in steel industry, Materials and Manufacturing Processes, 24 (2009) 3, 369–374 11 M. Kova~i~, Genetic programming and Jominy test modeling, Materials and Manufacturing Processes, 24 (2009) 7, 806–808 12 M. Kova~i~, S. Sen~i~, Critical inclusion size in spring steel and genetic programming, RMZ – Materials and Geoenvironment, 57 (2010) 1, 17–23 13 J. R. Koza, Genetic Programming III., Morgan Kaufmann, San Francisco, 1999, 3–16 M. KOVA^I^, S. SEN^I^: MODELING OF PM10 EMISSION WITH GENETIC PROGRAMMING Materiali in tehnologije / Materials and technology 46 (2012) 5, 453–457 457 F. KAFEXHIU et al.: EFFECT OF TEMPERING ON THE ROOM-TEMPERATURE MECHANICAL PROPERTIES ... EFFECT OF TEMPERING ON THE ROOM-TEMPERATURE MECHANICAL PROPERTIES OF X20CrMoV121 AND P91 STEELS VPLIV POPU[^ANJA NA MEHANSKE LASTNOSTI JEKEL X20CrMoV121 IN P91 PRI SOBNI TEMPERATURI Fevzi Kafexhiu, Franc Vodopivec, Jelena Vojvodi~ Tuma Institute of metals and technology, Lepi pot 11, 1000 Ljubljana, Slovenia fevzi.kafexhiu@imt.si Prejem rokopisa – received: 2012-01-23; sprejem za objavo – accepted for publication: 2012-04-19 The effect of tempering time and temperature on the room-temperature tensile properties and hardness of two martensitic creep-resistant steels, X20CrMoV121 and P91, was investigated. Samples cut from industrial tubes were tempered for 17520 h at 650 °C and for 8760 h at 750 °C. On the tempered samples the yield stress, tensile strength, and hardness at room temperature were determined and an SEM examination was carried out. It was found that the effect of tempering at 750 °C on the microstructural changes, room-temperature tensile properties and hardness was greater for both steels than the effect of tempering at 650 °C. The changes in the yield stress, tensile strength and hardness of both steels at a given tempering temperature were found to be very similar. Therefore, a general mathematical expression with specific coefficients for each property was deduced. These results are part of a larger investigation aimed at establishing a correlation between the particle spacing, yield stress, creep rate and hardness, which could be useful in an evaluation of the lifetime issues relating to the thermal-power-plant components. Keywords: tempering, microstructure, mechanical properties, X20CrMoV121 and P91 steels Vpliv ~asa in temperature popu{~anja na raztr`ne lastnosti in trdoto pri sobni temperaturi je bil raziskan pri martenzitnih jeklih X20CrMoV121 in P91, ki sta odporni proti lezenju. Preizku{anci so bili izrezani iz industrijskih cevi in popu{~eni do 17520 h pri 650 °C in 8760 h pri 750 °C. Na popu{~enih vzorcih so bile dolo~ene meja plasti~nosti, raztr`na trdnost in trdota pri sobni temperaturi, mikrostruktura pa preiskana v SEM. Ugotovljeno je bilo, da je vpliv popu{~anja pri 750 °C na spremembo mikrostrukture, raztr`ne lastnosti pri sobni temperaturi in trdoto ve~ji pri obeh jeklih, kot vpliv popu{~anja pri 650 °C. Spremembe meje plasti~nosti, trdnosti in trdote so bile podobne pri obeh jeklih pri dani temperaturi popu{~anja. Razvita je bila zato matemati~na odvisnost s specifi~nimi koeficienti za vsako lastnost. Rezultati so del {ir{e raziskave, katere cilj je opredeliti korelacije med razdaljo med izlo~ki, mejo plasti~nosti, hitrostjo lezenja in trdoto, ki bi bile koristne pri oceni preostale trajnostne dobe komponent termoelektrarn. Klju~ne besede: popu{~anje, mikrostruktura, mehanske lastnosti, jekli X20CrMoV121 in P91 1 INTRODUCTION In recent years there has been an increased demand to improve the efficiency of steam power plants for economical and environmental reasons.1–4 A straight- forward way to achieve this is to raise the inlet temperature and pressure of the steam that passes through the turbines. This directly saves the fuel and reduces the CO2 emissions.5 The problems with higher steam temperature and pressure are largely material related. The microstructures of the materials operating under such conditions change with time and, consequently, several degradation mecha- nisms such as creep, fatigue, thermal fatigue, creep- fatigue, progressive embrittlement, corrosion/oxidation, etc., are accelerated. Among these damage mechanisms, the most important are the damages caused by an increase in the creep deformation. The main candidate materials for building the plants with more advanced steam parameters are 9–12 % chromium steels.6, 7 The risk of a failure due to creep deformation and other damage mechanisms is always present. Therefore, periodical checking of their properties and residual lifetime after a determined period of operation of the power plants is always necessary. The checking of the creep rate and the creep strength is expensive and time consuming. For this reason, simpler methods using faster and less expensive tests that make it possible to identify the changes in the properties of the steels already employed in the vital parts of a power plant, have been developed. One of these methods is checking the room-temperature mechanical properties and the microstructure after a certain tempering time, simulating the changes in the microstructure and the properties that occur after longer operation periods (in real conditions). It has been shown recently that the time when the creep failure occurs is related to the yield stress and the tensile strength at creep temperature8,9 and that hardness is related to creep life.10 It was also shown11 that within a certain range of the room-temperature yield stress (350 MPa to 650 MPa) the accelerated creep rate at 580 °C decreases continuously from 8 · 10–7 s–1 to 5 · 10–9 s–1. Materiali in tehnologije / Materials and technology 46 (2012) 5, 459–464 459 UDK 669.14.018.44:621.785.72:620.17 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)459(2012) 2 EXPERIMENTAL WORK In the present work, the X20CrMoV121 and P91 steels were chosen for an investigation. The samples were cut from the pipelines with  = 38 mm × 8 mm and  = 82 mm × 14.5 mm. The quantometer chemical com- positions for both steels are given in Table 1. Before extracting the specimens for the room-tem- perature tests and examinations, the samples of both steels were tempered for discrete times up to 17520 h at 650 °C and for a shorter time up to 8760 h at 750 °C to simulate the changes in the microstructure that take place under real operating conditions in the power plants, and their effect on the room-temperature tensile properties and hardness. Static-tensile tests at room (ambient) temperature were performed on the specimens extracted from the previously tempered samples. The tests were initially performed on the specimens prepared from the as-delivered tubes and then on the specimens tempered for 2 h, 4320 h and 8760 h at 650 °C and 750 °C, and up to 17520 h at 650 °C. All the tensile tests were performed on a 500-kN static-dynamic testing machine in the Laboratory for Mechanical Testing at the Institute of Metals and Technology. A part of these results was published in the proceedings of IPSSC.12 With the aim to assess the changes in the microstruc- ture as a function of tempering time and temperature, the SEM specimens were prepared with the standard metallographic techniques. A Jeol – JSM6500F Field Emission Scanning Elec- tron Microscope (FE-SEM) was used to acquire images at three different magnifications, namely, 2000-, 5000- and 10000-times, with the working parameters of the 15-kV acceleration voltage, 7-nA probe current and 10-mm working distance. Images were acquired from the specimens in the initial (as-delivered) state and from those tempered for 2 h, 4320 h and 8720 h (1 year) at both 650 °C and 750 °C. In this way, the microstructural evolution of both steals as a function of tempering time and temperature could be observed. The specimens prepared for the SEM imaging were usable also for the Vickers hardness measurements. The HV5 measurements were carried out with an Instron 2100B Vickers hardness tester. The measurements were performed before the isothermal tempering, i.e., in the as-delivered state and within the used tempering times. Three measurements were performed over the whole specimen area at a suitable distance from the specimen edge to avoid any edge inaccuracy. 3 RESULTS AND DISCUSSION A comparison between a decrease in the yield stress (y) and the tensile strength (m) at both tempering temperatures indicates a similarity in the changing of these two properties for both steels. From Figures 1 and 2 it can be seen that the effect of tempering at 650 °C on the reduction of m and y is higher for the X20CrMoV121 steel, where y drops by 34 N/mm2 and m by 55 N/mm2, than for the P91 steel, where y drops by 19 N/mm2 and m by 24 N/mm2. It is F. KAFEXHIU et al.: EFFECT OF TEMPERING ON THE ROOM-TEMPERATURE MECHANICAL PROPERTIES ... 460 Materiali in tehnologije / Materials and technology 46 (2012) 5, 459–464 Table 1: Chemical compositions of the X20CrMoV121 and P91 steels in mass fractions Tabela 1: Kemi~na sestava jekel X20CrMoV121 in P91 v masnih dele`ih Chemical composition, w/% Elements C Si Mn P S Cr Ni Mo V Cu Nb Al N X20CrMoV121 0.2 0.29 0.52 0.019 0.011 11 0.64 0.94 0.31 0.059 0.024 0.032 0.017 P91 0.1 0.38 0.48 0.012 0.002 7.9 0.26 0.98 0.23 0.14 0.11 0.016 0.064 Figure 2: Actual and calculated dependences of the tensile strength of the X20CrMoV121 and P91 steels on the tempering time at 650 °C and 750 °C Slika 2: Dejanska in izra~unana odvisnost raztr`ne trdnosti jekel X20CrMoV121 in P91 od ~asa popu{~anja pri 650 °C in 750 °C Figure 1: Actual and calculated dependences of the yield stress of the X20CrMoV121 and P91 steels on the tempering time at 650 °C and 750 °C Slika 1: Dejanska in izra~unana odvisnost meje plasti~nosti jekel X20CrMoV121 in P91 od ~asa popu{~anja pri 650 °C in 750 °C also obvious that for both steels the reduction of m is higher than the reduction of y. It should be also pointed out that after a longer tempering time, i.e., 17520 h at the above temperature, the reduction of y is, surprisingly, identical for both steels, i.e., 36 N/mm2, and similar behaviour was observed in the reduction of m, i.e., 43 N/mm2 for X20CrMoV121 and 41 N/mm2 for P91. On the other hand, the effect of tempering at 750 °C on the reduction of m and y is higher and their mutual correlation is different than in the case of tempering at 650 °C. For X20CrMoV121 y drops by 163 N/mm2 and m by 195 N/mm2, whereas for P91 y drops by 216 N/mm2 and m by 183 N/mm2. The effect of the duration and the temperature of tempering on the hardness of both steels, shown in Figure 3, is similar to the effect on the yield stress and the tensile strength. In X20CrMoV121, after 8760 h of tempering at 650 °C, the hardness is reduced by 15 HV, whereas in P91, under the same tempering conditions, the hardness is reduced by 11 HV. At a higher temperature and the same tempering time, i.e., at 750 °C for 8760 h, the hardness reduction in X20CrMoV121 is 71 HV, whereas in P91 this reduction is 68 HV. The reduction of hardness indicates that the size, the amount and the distribution of precipitates have a lower effect on the hardness than on the yield stress of the steels investigated. This could be explained with the fact that the yield stress is more related to deformation hardening than hardness. There is a clear similarity between the dependences of hardness HV, tensile strength m and yield stress y on the tempering time and the temperature and, for all three properties the general mathematical expression was deduced: y t k k t x( ) = −1 2 where y(t) stands for either m, y or HV as a function of tempering time t, k1 and k2 are constants and x represents the exponent, which can take the values of 1/2, 1/3, etc., depending on the way the curve obtained from the equation (1) best fits the experimental data of each property. Experimental vs. theoretical curve fittings are given in Figures 1, 2 and 3. They are obtained by using the values for x, k1 and k2 given in Table 2 and applying them in equation (1) for all three properties, both steels and both tempering temperatures. The value of exponent x was appropriated to 1/3, k1 holds the values for each of the measured properties at the as-delivered state, whereas the k2 was calculated with the least-square method using the R-project for statistical computing.13 It should be pointed out that the fitting for the yield stress and the tensile strength is quite good, whereas for the hardness this equation does not give a good fit for a longer tempering time at 650 °C. Table 2: Parameters for calculating the yield stress, tensile strength and hardness change of the X20CrMoV121 and P91 steels for both tempering temperatures with equation (1) Tabela 2: Parametri za izra~un meje plasti~nosti, raztr`ne trdnosti in trdote za jekli X20CrMoV121 in P91 pri obeh temperaturah po- pu{~anja z ena~bo (1) Properties Parameters X20CrMoV121 P91 650 °C 750 °C 650 °C 750 °C y k1 527 527 546 546 k2 1.444 7.682 1.197 10.231 x 1/3 1/3 1/3 1/3 m k1 753 753 712 712 k2 2.096 9.4 1.456 8.539 x 1/3 1/3 1/3 1/3 HV k1 238 238 228 228 k2 0.43 3.548 0.292 3.256 x 1/3 1/3 1/3 1/3 The influence of tempering on the microstructures of both steels is greater after tempering at 750 °C than at F. KAFEXHIU et al.: EFFECT OF TEMPERING ON THE ROOM-TEMPERATURE MECHANICAL PROPERTIES ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 459–464 461 Figure 3: Actual and calculated dependences of the hardness of the X20CrMoV121 and P91 steels on the tempering time at 650 °C and 750 °C Slika 3: Dejanska in izra~unana odvisnost trdote jekel X20CrMoV121 in P91 od ~asa popu{~anja pri 650 °C in 750 °C Figure 4: Microstructure of the X20CrMoV121 steel at the initial (as-received) state Slika 4: Mikrostruktura jekla X20CrMoV121 v za~etnem (dobavlje- nem) stanju 650 °C (Figures 4 to 9), because the diffusivity of the carbide-forming elements (Cr, Mo, Fe, V, and Nb), found in the solid solution in the ferrite matrix is temperature dependent, and is higher at higher temperatures.14–17 Due to tempering, both the size and the inter-particle spacing of carbide particles increase. In addition, the carbide-particles distribution changes from stringers along the grain and sub-grain boundaries to a uniform structure. Images in Figures 4 and 7 show the as-deli- vered-state microstructures of the X20CrMoV121 and P91 steels, respectively. In both cases, the majority of particles are cementite Fe3C, containing also chromium, or they are chromium carbides Cr23C6, containing also iron and molybdenum.18,19 The carbide particles are found in the stringers distributed along the grain and sub-grain boundaries of martensite, and there is no difference between the size of the Fe3C, M23C6 and VC particles. After 8760 h of tempering at 650 °C the preci- pitates are almost evenly distributed and there is a difference between the size of VC (small white particles) and M23C6, which grow larger in both steels (Figures 5 and 8). In addition, the grain and subgrain boundaries of martensite are much less pronounced and some of them have already disappeared. From Figures 6 and 9 it is obvious that the tempering at 750 °C for 8760 h causes much greater changes in the microstructures of both steels, where the size of the M23C6 particles and the spacing between them is drasti- cally increased. On the other hand, the number density of these particles has clearly dropped, so the Ostwald ripening effect, where larger particles coarsen at the expense of smaller ones is quite obvious in this case. F. KAFEXHIU et al.: EFFECT OF TEMPERING ON THE ROOM-TEMPERATURE MECHANICAL PROPERTIES ... 462 Materiali in tehnologije / Materials and technology 46 (2012) 5, 459–464 Figure 7: Microstructure of the P91 steel at the initial (as-received) state Slika 7: Mikrostruktura jekla P91 v za~etnem (dobavljenem) stanju Figure 5: Microstructure of the X20CrMoV121 steel after 8760 h of tempering at 650 °C Slika 5: Mikrostruktura jekla X20CrMoV121 po 8760 h popu{~anja pri 650 °C Figure 8: Microstructure of the P91 steel after 8760 h of tempering at 650 °C Slika 8: Mikrostruktura jekla P91 po 8760 h popu{~anja pri 650 °C Figure 6: Microstructure of the X20CrMoV121 steel after 8760 h of tempering at 750 °C Slika 6: Mikrostruktura jekla X20CrMoV121 po 8760 h popu{~anja pri 750 °C Since carbide precipitates represent the most import- ant strengthening mechanism in 9–12% Cr steels, and having in mind the fact that mechanical properties deteriorate as both the size and the inter-particle spacing increase, it is obvious that they are in some kind of mutual correlation, investigated and confirmed by many authors.14–18 4 CONCLUSIONS The effect of the tempering time and the temperature of the creep-resistant martensitic steels, X20CrMoV121 and P91, which differ in chemical composition, on the room-temperature tensile properties and hardness was determined. According to the results obtained, the following is concluded: • The effect of tempering at both temperatures separa- tely is the same for both steels investigated. • When tempering at the temperature of 650 °C, the changes in yield stress y and tensile strength m are relatively small, whereas when tempering at a higher temperature, i.e., 750 °C, the changes are greater for both properties, and for steel X20CrMoV121 y is decreased by 163 N/mm2 and m by 195 N/mm2, whereas for steel P91 y is decreased by 216 N/mm2 and m by 183 N/mm2 after 8760 h of tempering at 750 °C. • The dependence of y and m on the tempering time and the temperature is for both steels virtually the same and can be mathematically expressed with the relation: y t k k t x( ) = −1 2 All three parameters, k1, k1 and x have different values depending on the temper- ing conditions, the steel chemical composition, etc. • Hardness evolution is similar to that of y and m and it could therefore be expressed with the same mathe- matical relation, however, with different values for all three parameters. • The effect of microstructural changes, particle size and spacing as well as particle distribution is caused by coarsening the particles. The coarsening rate depends on diffusion, which is greater at higher temperatures. Acknowledgement The authors are indebted to the company TE[ [o{tanj for supporting the investigation and to Mr. D. Kmeti~ for the mechanical tests. 5 REFERENCES 1 J. Hald, Microstructure and long-term creep properties of 9–12 % Cr steels, International Journal of Pressure Vessels and Piping, 85 (2008), 30–37 2 M. Yoshizawa, M. Igarashi, Long term creep deformation characte- ristics of advanced ferritic steels for USC power plants, International Journal for Pressure Vessels and Piping, 84 (2007), 37–43 3 R. L. Klueh, Elevated temperature ferritic and martensitic steels and their application to future nuclear reactors, International Materials Reviews, 50 (2005), 287–310 4 C. Scheu, F. Kauffmann, G. Zies, K. Maile, S. Straub, K. H. Mayer, Requirements for microstructural investigations of steels used in modern power plants, Zeitschrift für Metallkunde, 96 (2005), 653–659 5 D. V. Thornton, K. H. Mayer, European high temperature materials development for advanced steam turbines, in: Advanced heat resistant steels for power generation, ed. R. Viswanathan and J. Nutting, 708 (1999), 349–364 6 A. Shibli, F. Starr, Some aspects of plant and research experience in the use of new high strength martensitic steel P91, International Journal of Pressure Vessels and Piping, 84 (2007), 114–122 7 F. Abe, T. U. Kern, R. Viswanathan, Creep-resistant steels, Wood- head Publishing, CRC Press, Cambridge, England, 2008 8 F. Abe, Heat to heat variation in long term creep strength of some ferritic steels, in: Creep and Fracture in: Creep & Fracture in High Temperature Components, ed. I. A. Shibli, S. R. Holdsworth, DEStech Publ. Inc, 2009, 5–18 9 F. Masuyama, T. Tokumaga, N. Shimahata, T. Yamamoto, M. Hirano, Comprehensive approach to creep life assessment of martensitic heat resistant steels, in: Creep & Fracture in High Temperature Compo- nents, ed. I. A. Shibli, S. R. Holdsworth, DEStech Publ. Inc, 2009, 19–30 10 D. J. Allen, A hardness normalized model of creep rupture for P91 steel, in: Creep & Fracture in High Temperature Components, ed. I. A. Shibli, S. R. Holdsworth, DEStech Publ. Inc, 2009, 659–688 11 F. Vodopivec, J. Vojvodi~ - Tuma, M. Jenko, R. Celin, B. [u{tar{i~, Dependence of accelerated creep rate at 580 °C and room tempera- ture yield stress for two creep resistant steels, Steel research, 81 (2010) 7, 576–580 12 F. Kafexhiu, F. Vodopivec, J. V. Tuma, Tempering effects on the microstructure, mechanical properties and creep rate of X20CrMoV121 and P91 steels, Proceedings of 4th Jo`ef Stefan International Postgraduate School Students Conference, Ljubljana, 2012, 241–246 13 R-project, http://www.r-project.org/ 14 F. Vodopivec, M. Jenko, R. Celin, B. @u`ek, D. A. Skobir Balanti~, Creep resistance of microstructure of welds of creep resistant steels, Mater. Tehnol., 45 (2011) 2, 139–143 F. KAFEXHIU et al.: EFFECT OF TEMPERING ON THE ROOM-TEMPERATURE MECHANICAL PROPERTIES ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 459–464 463 Figure 9: Microstructure of the P91 steel after 8760 h of tempering at 750 °C Slika 9: Mikrostruktura jekla P91 po 8760 h popu{~anja pri 750 °C 15 D. A. Skobir Balanti~, M. Jenko, F. Vodopivec, R. Celin, Effect of change of carbide particles spacing and distribution on creep rate of martensite creep resistant steels, Mater. Tehnol., 45 (2011) 6, 555–559 16 F. Vodopivec, D. Kmeti~, J. Vojvodi~-Tuma, D. A. Skobir Balanti~, Effect of operating temperature on microstructure and creep resist- ance of 20CrMoV121 steel, Mater. Tehnol., 38 (2004) 5, 233–239 17 F. Vodopivec, M. Jenko, J. Vojvodi~-Tuma, Stability of MC carbide particles size in creep resisting steels, Metalurgija 45 (2006) 3, 147–153 18 D. A. Skobir Balanti~, F. Vodopivec, M. Jenko, S. Spai}, B. Markoli, Zeitschrift für Metallkunde, 95 (2004) 11, 1020–1024 19 D. A. Skobir Balanti~, F. Vodopivec, M. Jenko, S. Spai}, B. Markoli, The influence of tempering on the phase composition of the carbide precipitates in X20CrMoV121 steel, Mater. Tehnol., 37 (2003) 6, 353–358 F. KAFEXHIU et al.: EFFECT OF TEMPERING ON THE ROOM-TEMPERATURE MECHANICAL PROPERTIES ... 464 Materiali in tehnologije / Materials and technology 46 (2012) 5, 459–464 M. FUJDA et al.: STRUCTURE AND PROPERTIES OF AlMgSi ALLOYS AFTER ECAP AND POST-ECAP AGEING STRUCTURE AND PROPERTIES OF AlMgSi ALLOYS AFTER ECAP AND POST-ECAP AGEING STRUKTURA IN LASTNOSTI ZLITIN AlMgSi, STARANIH PRED ECAP IN PO NJEM Martin Fujda1, Milo{ Matvija1, Tibor Kva~kaj2, Ondrej Milkovi~1, Pavol Zubko1, Katarína Nagyová1 1Department of Materials Science, Faculty of Metallurgy, Technical University of Ko{ice, Slovakia 2Department of Metals Forming, Faculty of Metallurgy, Technical University of Ko{ice, Slovakia martin.fujda@tuke.sk Prejem rokopisa – received: 2012-02-10; sprejem za objavo – accepted for publication: 2012-03-22 The mechanical properties and the microstructures of the EN AW 6082 and EN AW 6063 aluminium alloys subjected to pre-ECAP solutionizing heat treatment, equal channel angular pressing (ECAP), and post-ECAP artificial ageing are compared. The quenched alloy states were severely deformed at room temperature by the ECAP technique following route BC. Repetitive ECAP caused formation of the ultra-fine subgrain microstructure with a high dislocation density and high strain hardening of the alloys, thus exhibiting an improvement in strength, but also a degradation of the analysed alloys’ ductility. The application of the optimal artificial-ageing regimes after a severe plastic deformation improved only the tensile ductility of the alloys, while their strength was slightly decreased because the softening effect caused by the solid-solution recovery and the relaxation of the internal stress dominated over the hardening effect caused by the expected metastable ''- and '-Mg2Si phase precipitation during the artificial ageing treatment. Keywords: aluminium alloys, ECAP, ageing, microstructure, mechanical properties Primerjali smo lastnosti EN AW 6082 in EN AW 6063 aluminijevih zlitin, raztopno `arjenih pred stiskanjem skozi pravokotni kanal (ECAP), po stiskanju ECAP in umetno staranih po ECAP-postopku. Zlitine v ga{enem stanju so bile mo~no deformirane pri sobni temperaturi s tehniko ECAP po poti Bc. Ponovljena ECAP je povzro~ila nastanek ultra drobnozrnate podmikrostrukture z veliko gostoto dislokacij, velikim utrjevanjem zlitin in s tem pove~anje trdnosti ter poslab{anje duktilnosti preiskovanih zlitin. Optimalni re`im umetnega staranja po veliki plasti~ni deformaciji je povzro~il samo izbolj{anje natezne duktilnosti materiala, njihova trdnost pa je bila zmanj{ana zaradi u~inka meh~anja pri popravi trdne raztopine in spro{~anju notranjih napetosti, ki prevladujejo nad u~inkom utrjevanja z izlo~anjem izlo~kov metastabilne ''- in '-Mg2Si faze med postopkom umetnega staranja. Klju~ne besede: aluminijeve zlitine, ECAP, staranje, mikrostruktura, mehanske lastnosti 1 INTRODUCTION At present, the equal-channel angular pressing (ECAP) technique is a very effective method of severe plastic deformation, and has been widely used for producing ultra-fine grain microstructures of Al-based alloys with significantly improved mechanical proper- ties, including enhanced superplasticity, high strength, etc.1–10 The severe plastic deformation (SPD) made by the ECAP process also markedly increases the density of lattice defects in the solid solution, and can thus accele- rate the precipitation process of the strengthening particles during the post-ECAP ageing treatment applied to the age-hardenable aluminium alloys.11–13 However, the ultra-fine-grained age-hardenable alloys often exhibit low tensile ductility at room temperature. Therefore, it is important to obtain ductility comparable to that of the conventional coarse-grained age-hardenable Al-based alloys. EN AW 6063 and 6082 are the most used AlMgSi alloys that are suitable for different miscellaneous structural applications in the building, automotive and aircraft industries due to their strength-modification options, low density, good corrosion properties, and good weldability.14,15 The optimal combination of heat treatment and the SPD by repetitive ECAP led to a significantly increased strength and a relatively good ductility of these alloys caused by the ultra-fine grained structure formation and the strengthening precipitation of the '-Mg2Si phase particles during the post-ECAP ageing treatment.4,16–21 The aim of the present work is to investigate the effect of the SPD by repetitive ECAP and the subsequent artificial ageing on the microstructure and the mechanical properties of the EN AW 6082 and EN AW 6063 aluminium alloys. 2 EXPERIMENTAL WORK The experiments were carried out on the EN AW 6082 and EN AW 6063 aluminium alloys, the chemical composition of which is presented in Table 1. The analyzed alloys in the form of extruded rods, subjected to artificial ageing (T5 temper), were used as the initial states. Prior to the deformation in an ECAP die, the specimens of the initial states were solution annealed for Materiali in tehnologije / Materials and technology 46 (2012) 5, 465–469 465 UDK 669.715:621.785.7:620.17 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)465(2012) 1.5 h at 550 °C (6082 alloy) or at 510 °C (6063 alloy) and subsequently cooled to room temperature by water quenching. The purpose of these heat-treatment regimes was to obtain a supersaturated solid solution. Table 1: Chemical composition of the EN AW 6082 and EN AW 6063 aluminium alloys (w/%) Tabela 1: Kemijska sestava aluminijevih zlitin EN AW 6082 in EN AW 6063 v masnih dele`ih (w/%) alloy Mg Si Mn Fe Zn Cu Al EN AW 6082 0.60 1.0 0.49 0.21 0.02 0.06 bal. EN AW 6063 0.44 0.52 0.03 0.21 0.02 0.03 bal. The quenched specimens were subjected to a deformation in an ECAP die with a channel intersection angle F = 90°, and an arc of curvature Y = 37° up to 4 (6082 alloy) or 6 passes (6063 alloy), respectively. The repetitive ECAP of the specimens of size  = 10 mm × 100 mm was realized at room temperature following route BC (the sample was rotated in the same sense by 90°). After ECAP, the specimens were subjected to artificial ageing in the temperature range of 60–180 °C as a function of time to determine the post-ECAP artificial-ageing condition for the highest peak hardness. The microstructures of the investigated alloys obtained after quenching, deformation in the ECAP die, and after the post-ECAP peak-ageing treatment were analyzed in the central zones of the specimens’ cross- sections using a transmission electron microscopy (TEM). The samples for TEM were prepared from the 1-mm-thick slices that were ground and polished to a thickness of about 0.15 mm. These slices were finally electrolytically thinned in a solution of 25 % HNO3 and 75 % CH3OH at a temperature of –30 °C. The average grain size of the solid solution for all the analysed alloy states was determined according to ASTM E112. The influence of the severe plastic deformation by the ECAP process and post-ECAP artificial ageing on the mechanical properties of the analyzed alloys was evaluated with the Vickers hardness measurement (HV 10) and a tensile test. The tensile test (the initial strain rate of 2.5 × 10–4 s–1) was carried out on short specimens (d0 = 5 mm, l0 = 10 mm) made from quenched, ECAP- processed and post-ECAP peak-aged alloy billets. Subsequently, the characteristics of the strength (Rp0.2 – yield strength and Rm – tensile strength), the uniform tensile elongation (Ag), the tensile elongation (A), and the reduction in the area (Z) were determined. 3 RESULTS AND DISCUSSION 3.1 Hardness The Vickers hardness of the analyzed alloy states is summarized in Table 2. Figures 1 and 2 show the variation of the Vickers hardness of the ECAPed analyzed alloys as a function of artificial ageing time at a temperature range of 60–180 °C. A significant increase in the alloys’ hardness value (by about 120 % for the 6082 alloy and about 202 % for the 6063 alloy) was achieved by the ECAP processing of the quenched alloy states. A post-ECAP artificial-ageing treatment was applied to increase the hardness of the ECAPed alloys. The ECAPed alloy state reached peak hardness after about 60 h of the artificial ageing at 80 °C, but the M. FUJDA et al.: STRUCTURE AND PROPERTIES OF AlMgSi ALLOYS AFTER ECAP AND POST-ECAP AGEING 466 Materiali in tehnologije / Materials and technology 46 (2012) 5, 465–469 Figure 2: Vickers hardness of the ECAPed EN AW 6063 alloy state as a function of artificial ageing time at various ageing temperatures Slika 2: Vickersova trdota zlitine EN AW 6063 po ECAP kot funkcija ~asa umetnega staranja pri razli~nih temperaturah staranja Table 2: Vickers hardness of the analyzed alloy states Tabela 2: Vickersova trdota analiziranih stanj zlitin alloy/state Q E post-ECAP peak-aged 60 °C 80 °C 100 °C 140 °C 180 °C EN AW 6082 60.2 132.5 130.7 131.1 130.5 127.3 126.6 EN AW 6063 34.6 104.4 102.3 103.0 102.6 100.4 – Q-quenched, E-ECAPed Figure 1: Vickers hardness of the ECAPed EN AW 6082 alloy state as a function of artificial ageing time at various ageing temperatures Slika 1: Vickersova trdota zlitine EN AW 6082 po ECAP kot funkcija ~asa umetnega staranja pri razli~nih temperaturah staranja maximum values were only slightly lower (by about 1.4 HV) than those of the ECAPed alloy states. The increase of the ageing temperature from 60 °C to 100 °C resulted in the peak-ageing-time shortening. However, the maximum alloys’ hardness was the highest after being aged at 80 °C. During the post-ECAP ageing at higher temperatures (140 °C and 180 °C), the hardness of the ECAPed alloy states was decreased in a very short time. 3.2 Microstructure Figure 3 shows the recrystalized microstructure of the 6082-alloy quenched state that consists of fine solid- solution grains (average size: 14.6 μm). However, very coarse grains (average size: 115.6 μm) of the recrystal- ized solid solution were observed after the quenching of the 6063 alloy (Figure 4). In the case of the 6082 alloy (containing the mass fraction 0.49 % of Mn), no considerable solid-solution grain growth during the applied solution annealing at 550 °C was prevented by the grain-boundary pinning effect of the fine, dispersive, Mn-rich phase particles (Figure 3; particle size: 30–300 nm). The role of these dispersive particles is to inhibit the recrystallization process and the matrix grain growth.22–25 Comparing Figures 3 and 4, it is obvious that the volume fraction of the fine dispersive particles in the coarse-grained matrix of the 6063 alloy (Fe-rich phase particles with the average size of 150 nm) was lower than in the 6082-alloy matrix. Both types of dispersive particles were mostly distributed homogeneously throughout the equiaxed solid-solution grains, which exhibited a relatively low dislocation density in both quenched alloy states. During the repetitive ECAP applied to the quenched alloy state, the solid-solution grains were hetero- geneously refined. A cell-dislocation substructure, equiaxed ultra-fine subgrains, and high dislocation M. FUJDA et al.: STRUCTURE AND PROPERTIES OF AlMgSi ALLOYS AFTER ECAP AND POST-ECAP AGEING Materiali in tehnologije / Materials and technology 46 (2012) 5, 465–469 467 Figure 6: Microstructure of the post-ECAP aged EN AW 6063 alloy treated for 60 h at 80 °C, TEM Slika 6: Mikrostruktura zlitine EN AW 6063 po ECAP, starane 60 h pri 80 °C, TEM Figure 4: Microstructure of the quenched EN AW 6063 alloy state (510 °C + water quenching), TEM Slika 4: Mikrostruktura ga{ene zlitine EN AW 6063 (510 °C + ga{e- nje v vodi), TEM Figure 3: Microstructure of the quenched EN AW 6082 alloy state (550 °C + water quenching), TEM Slika 3: Mikrostruktura ga{ene zlitine EN AW 6082 (550 °C + ga{e- nje v vodi), TEM Figure 5: Microstructure of the ECAPed EN AW 6082 alloy state, TEM Slika 5: Mikrostruktura zlitine EN AW 6082 po ECAP, TEM density within the subgrains were observed in the micro- structure of the ECAPed alloys (Figure 5). The average subgrain size (180 nm) measured for the ECAPed 6082 alloy was a little bit smaller than the size of the subgrains (220 nm) of the quenched 6063-alloy state formed during ECAP. This was a result of the supersaturation of the severely deformed solid solution, obtained for the quenched 6082 alloy prior to the use of repetitive ECAP that was higher than that obtained for the 6063 alloy with lower Mg and Si contents. With an increase in the content of the alloying elements in the solid solution of the Al-based alloys, the microstructure formed after ECAP becomes finer and the dislocation density inside the ultrafine grains is higher.26 After the peak ageing of the ECAPed alloy states (Figure 6), a slight dislocation recovery was observed in the subgrains, still having a very high dislocation density, and in the dislocation cells. In addition, the ultra-fine subgrains grew to a size of about 250 nm and 280 nm during the peak ageing of the 6082 and 6063 alloys, respectively. 3.3 Tensile test The values of the tensile strength (Rm), the yield strength (Rp0.2) and the tensile ductility (Ag, A, Z), which are presented in Table 3, and the stress-strain curves shown in Figures 7 and 8, provide a comparison of the mechanical properties determined for the quenched, ECAPed and post-ECAP peak-aged (for 60 h at 80 °C) states of the analyzed alloys. An elimination of the strengthening effect of the nanoparticles of the ''- and '-Mg2Si phases through their dissolution into the solid solution of the quenched alloy states is indicated by their low yield strength, ultimate tensile strength and high ductility. The ultimate tensile strength and, especially, the yield stress of the quenched alloy states increased significantly (Rp0,2 by about 293MPa and 283 MPa for the 6082 and 6063 alloys, respectively) due to the repetitive ECAP, while the tensile ductility (A, Z) deteriorated. The tensile-deformation behaviour of these deformed alloy states (Figures 7 and 8) is typical for the severely strain-hardened (ECAPed) AlMgSi alloys19,21 for which the low values of the uniform tensile elongation (Ag) and a high Rp0,2/Rm ratio have been found. Despite the fact that the hardness values of the post-ECAP peak-aged alloy states are only a little bit lower than those of the ECAPed ones, the yield strength and the ultimate tensile strength of the ECAPed alloys decreased by about 5.5 % and 2.5 %, respectively, during the applied peak-ageing treatment. It is obvious that the applied ageing of the ECAPed alloys resulted in an improved tensile ductility (mainly the uniform tensile elongation Ag). This softening of the alloys occurred because the softening effect caused by the microstructure low recovery and the relaxation of the internal stresses dominates the hardening effect caused by the expected sequence precipitation of the metastable Mg2Si-phase (GP-zones, ''- and '-phase) particles.27,28 Table 3: Mechanical properties of the analyzed alloy states Tabela 3: Mehanske lastnosti analiziranih stanj zlitin Alloy State Rp0.2/MPa Rm/ MPa Ag/ % A/ % Z/ % EN AW 6082 Q 115 243 14.1 31.9 72.5 E 408 420 1.6 19.8 39.8 pEa 384 408 4.7 22.1 42.7 EN AW 6063 Q 34 135 20.2 36.4 86.5 E 317 326 2.3 22.5 43.1 pEa 300 319 5.2 24.5 42.9 Q-quenched, E-ECAPed, pEa-post-ECAP aged 4 CONCLUSIONS The severe plastic deformation of the quenched EN AW 6082 and EN AW 6063 aluminium alloys realized M. FUJDA et al.: STRUCTURE AND PROPERTIES OF AlMgSi ALLOYS AFTER ECAP AND POST-ECAP AGEING 468 Materiali in tehnologije / Materials and technology 46 (2012) 5, 465–469 Figure 8: Stress-strain curves of the EN AW 6063 alloy states Slika 8: Krivulja napetost – raztezek za razli~na stanja zlitine EN AW 6063 Figure 7: Stress-strain curves of the EN AW 6082 alloy states Slika 7: Krivulja napetost – raztezek za razli~na stanja zlitine EN AW 6082 with the ECAP process caused a refinement and an intensive strain hardening of the alloy solid solution. The result was the increased hardness and strength of the alloys and, on the other hand, a decrease in the alloys’ tensile ductility. The increased supersaturation of the severely deformed solid solution obtained for the quenched 6082 alloy was a reaction to its higher strengthening caused by the ECAP processing and was higher than that obtained for the 6063 alloy with the lower Mg and Si contents. The artificial post-ECAP peak-ageing of the alloys was effective only in slightly improving the tensile ductility of the ECAPed alloys, which was a consequence of the low recovery of the ECAPed alloy and the relaxation of internal stresses. Acknowledgements This work was supported by the Scientific Grant Agency of the Slovak Republic within the grant project VEGA No. 1/0866/09. 5 REFERENCES 1 V. M. Segal, V. I. Reznikov, A. E. Drobyshevskiy, V. I. Kopylov, Russian Metallurgy, 1 (1981), 99–105 2 M. Furukawa, Z. Horita, M. Nemoto, T. G. Langdon, Journal of Materials Science, 36 (2001), 2835–2843 3 J. Bidulská, R. Ko~i{ko, R. Bidulský, M. Actis Grande, T. Doni~, M. Martikán, Acta Metallurgica Slovaca, 16 (2010), 4–11 4 Z. Horita, T. Fujinami, M. Nemoto, T. G. Langdon, Journal of Materials Processing Technology, 117 (2001), 288–292 5 R. K. Islamgaliev, N. F. Yunusova, R. Z. Valiev, N. K. Tsenev, V. N. Perevezentsev, T. G. Langdon, Scripta Materialia, 49 (2003), 467–472 6 S. Lee, M. Furukawa, Z. Horita, T. G. Langdon, Materials Science and Engineering, A 342 (2003), 294–301 7 R. Z. Valiev, N. A. Enikeev, T. G. Langdon, Metallic Materials, 49 (2011), 1–9 8 Z. Horita, S. Lee, S. Ota, K. Neishi, T. G. Langdon, Materials Science Forum, 357–359 (2001), 471–476 9 K. Turba, P. Malek, M. Cieslar, Metallic Materials, 45 (2007), 165–170 10 M. J. Starink, N. Gao, M. Furukawa, Z. Horita, C. Xu, T. G. Lang- don, Reviews on Advanced Materials Science, 7 (2004), 1–12 11 W. J. Kim, C. S. Chung, D. S. Ma, S. I. Hong, H. K. Kim, Scripta Materialia, 49 (2003), 333–338 12 M. Y. Murashkin, M. V. Markushev, Y. V. Ivanisenko, R. Z. Valiev, Solid State Phenomena, 114 (2001), 91–96 13 J. K. Kim, W. J. Kim, Solid State Phenomena, 124–126 (2007), 1437–1440 14 T. Kva~kaj, R. Bidulský (Eds.), Aluminium Alloys, Theory and Applications, InTech, Rijeka 2011 15 J. R. Davis (Ed.), Aluminium and Aluminium Alloys, ASM Spe- cialty Handbook, ASM International, Ohio 1993 16 J. Zrník, Z. Nový, T. Kva~kaj, V. Berná{ek, D. Ke{ner, M. Slámová, Acta Metallurgica Slovaca, 10 (2004), 277–284 17 B. Cherukuri, T. S. Nedkova, R. Srinivasan, Materials Science and Engineering, A 410-411 (2005), 394–397 18 J. C. Werenskiold, H. J. Roven, Materials Science and Engineering, A 410–411 (2005), 174–177 19 H. J. Roven, H. Nesboe, J. C. Werenskiold, T. Seibert, Materials Science and Engineering, A 410–411 (2005), 426–429 20 P. Leo, E. Cerri, P. P. De Marco, H. J. Roven, Journal of Materials Processing Technology, 182 (2007), 207–214 21 W. J. Kim, J. Y. Wang, Materials Science and Engineering, A 464 (2007), 23–27 22 N. Parson, J. Hankin, K. Hicklin, C. Jowett, Comparison of the extrusion performance and product characteristics of three structural extrusion alloys: AA6061, AA6082 and AA6005A, In.: Proceedings of the 7th International Aluminium Extrusion Technology Seminar, Vol. 1, Aluminium Association, Washington DC, 2000, 1–12 23 G. F. Vander Voort (Ed), Metallography and Microstructures, ASM Handbook, Vol. 9, ASM International, Ohio 2004 24 A. L. Dons, Journal of Light Metals, 1 (2001), 133–149 25 L. Lodgaard, N. Ryum, Materials Science and Engineering, A 283 (2000), 144–152 26 J. May, M. Dinkel, D. Amberger, H. W. Höppel, M. Göken, Metallur- gical and Materials Transactions, A 38 (2007), 1941–1945 27 M. Hockauf, L. W. Meyer, D. Nickel, G. Alisch, T. Lampke, B. Wielage, L. Krüger, Journal of Materials Science, 43 (2008), 7409–7417 28 K. Hockauf, L. W. Meyer, M. Hockauf, T. Halle, Journal of Materials Science, 45 (2010), 4754–4760 M. FUJDA et al.: STRUCTURE AND PROPERTIES OF AlMgSi ALLOYS AFTER ECAP AND POST-ECAP AGEING Materiali in tehnologije / Materials and technology 46 (2012) 5, 465–469 469 Y. KAZANCOGLU et al.: APPLICATION OF A TAGUCHI-BASED NEURAL NETWORK FOR FORECASTING ... APPLICATION OF A TAGUCHI-BASED NEURAL NETWORK FOR FORECASTING AND OPTIMIZATION OF THE SURFACE ROUGHNESS IN A WIRE-ELECTRICAL-DISCHARGE MACHINING PROCESS UPORABA TAGUCHIJEVE NEVRONSKE MRE@E ZA NAPOVEDOVANJE IN OPTIMIRANJE POVR[INSKE HRAPAVOSTI PRI POSTOPKU @I^NE EROZIJE Yigit Kazancoglu1, Ugur Esme2, Mustafa Kemal Kulekci2, Funda Kahraman2, Ramazan Samur3, Adnan Akkurt4, Melih Turan Ipekci5 1Izmir University of Economics, Department of Business Administration, 35330, Balcova-Izmir, Turkey 2Mersin University Tarsus Technical Education Faculty, Department of Mechanical Education, 33140 Tarsus-Mersin, Turkey 3Marmara University Technical Education Faculty, Department of Metal Education, 34722 Goztepe-Ýstanbul, Turkey 4Gazi University Industrial Arts Education Faculty, Department of Technological Education, 06830 Golbasi-Ankara, Turkey 5Gazi University, Department of Mechanical Engineering, 06570 Maltepe-Ankara, Turkey uguresme@gmail.com Prejem rokopisa – received: 2012-02-16; sprejem za objavo – accepted for publication: 2012-06-01 Wire-electrical-discharge machining (WEDM) is a modification of electro-discharge machining (EDM) and has been widely used for a long time for cutting punches and dies, shaped pockets and other machine parts on conductive materials. WEDM erodes workpiece materials by a series of discrete electrical sparks between the workpiece and an electrode flushed or immersed in a dielectric fluid. The WEDM process is particularly suitable for machining hard materials as well as complex shapes. In this paper, a neural network and the Taguchi design method have been implemented for minimizing the surface roughness in a WEDM process. A back-propagation neural network (BPNN) was developed to predict the surface roughness. In the development of a predictive model, machining parameters of open-circuit voltage, pulse duration, wire speed and dielectric flushing pressure were considered as the input model variables of the AISI 4340 steel. An analysis of variance (ANOVA) was used to determine the significant parameter affecting the surface roughness (Ra). Finally, the Taguchi approach was applied to determine the optimum levels of machining parameters. Keywords: WEDM, Taguchi-design method, neural network, surface roughness @i~na erozija (WEDM) je modifikacija potopne erozije (EDM) in se uporablja `e dolgo ~asa za izdelavo rezilnih in drugih orodij ter strojnih delov. WEDM erodira obdelovanec z iskrenjem med obdelovancem in elektrodo, ki se spira, oziroma je potopljena v dielektri~no teko~ino. WEDM-postopek je {e posebej primeren za obdelavo trdih materialov, kot tudi za kompleksne oblike. V tem ~lanku sta bili uporabljeni nevronska mre`a in Taguchijeva metoda na~rtovanja za zmanj{anje hrapavosti povr{ine. Pri razvoju modela za napovedovanje hrapavosti so bili upo{tevani parametri obdelave: elektri~na napetost, trajanje pulza, hitrost `ice in tlak dielektri~ne teko~ine za spiranje kot vhodne spremenljivke pri obdelavi jekla AISI 4340. Analiza variance (ANOVA) je bila uporabljena za dolo~anje vplivnih parametrov, ki vplivajo na povr{insko hrapavost (Ra). Na koncu je bil uporabljen Taguchijev pribli`ek za dolo~itev optimalnih parametrov procesa. Klju~ne besede: WEDM, Taguchi oblikovanje, nevronska mre`a, povr{inska hrapavost 1 INTRODUCTION The technologies of the wire-electrical-discharge machining (WEDM) have been emphasized significantly and have improved rapidly in recent years due to the requirements in various manufacturing fields. WEDM is used to produce complex two- and three-dimensional shapes through electrically conductive workpieces by using a wire electrode1–3. As shown in Figure 1, the sparks are generated between the conductive workpiece and a wire electrode flushed with, or immersed in, a dielectric fluid2. The degree of accuracy of the workpiece dimensions obtainable and the fine surface finishes make WEDM particularly valuable for the applications involving the manufacture of stamping dies, extrusion dies and prototype parts.2 Materiali in tehnologije / Materials and technology 46 (2012) 5, 471–476 471 UDK 621.9.025.5:620.191.35 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)471(2012) Figure 1: Working principle of a WEDM process4 Slika 1: Shemati~en prikaz WEDM-postopka4 The most important performance measures in WEDM are cutting speed, workpiece surface roughness and cutting width.2 Discharge current, discharge capaci- tance, pulse duration, pulse frequency, wire speed, wire tension, average working voltage and dielectric flushing conditions are the machining parameters that affect the performance measures.1–4 An optimization of the process parameters is the key step of the Taguchi method in achieving high quality without an increased cost. This is because an optimization of the process parameters can improve the quality, and the optimum process parameters obtained with the Taguchi method are insensitive to the variation of environmental conditions and other noise factors. Basically, the classical process-parameter design is complex and not easy to use.5,6 An advantage of the Taguchi method is that it emphasizes the mean-perfor- mance characteristic value close to the target value rather than a value within certain specification limits, thus improving the product quality. Additionally, the Taguchi method for experimental design is straightforward and easy to apply to many engineering situations, making it a powerful yet simple tool. It can be used to quickly narrow the scope of a research project or to identify problems in a manufacturing process from the data already in existence.5–8 A large number of experiments have to be carried out when the number of the process parameters increases. To solve this task, the Taguchi method uses a special design of orthogonal arrays to study the entire process-para- meter space with only a small number of experiments. Using an orthogonal array to design the experiment can help the designers to study the influence of multiple controllable factors on the average quality characteristics and variations in a fast and economic way, while using a signal-to-noise ratio (S/N) to analyze the experimental data can help the designers of a product, or a manu- facturer, to easily find out optimum parametric combi- nations.5–8 Investigations into the influences of the machining input parameters on the performance of WEDM have been widely reported. Huang et al.1 determined the effect of the WEDM process parameters (pulse-on time, pulse-off time, table feed rate, flushing pressure and workpiece surface, and machining history) on the gap width, the surface roughness and the white layer depth of a machined workpiece surface using the Taguchi method. Tosun et al.2 investigated the effect and optimization of the machining parameters on the cutting width and the material-removal rate (MRR) in WEDM based on the Taguchi method. The experimental studies were con- ducted with a varying pulse duration, open-circuit voltage, wire speed and dielectric flushing pressure. Mohammadi et al.3 applied the statistical analysis of WEDM turning on MRR using the Taguchi techniques. They considered the effects of the input parameters (power, time-off, voltage, servo, wire speed, wire tension and rotational speed) on the responses (MRR, surface roughness). Lin et al.9 investigated the effects of the machining parameters in electrical discharge machining on the machining characteristics of the SKH 57 high- speed steel. Moreover, they determined the optimum combination levels of the machining parameters based on the Taguchi method. Tosun and Cogun10 experimen- tally investigated the effects of the cutting parameters on the wire-electrode wear in WEDM. On the basis of ANOVA and F-Test, they found that the most effective parameters of the wire-wear ratio are the open-circuit voltage and pulse durations. Tarng et al.11 developed a neural-network system to determine the settings of pulse duration, pulse interval, peak current, open-circuit voltage, servo-reference voltage, electric capacitance and wire speed for an estimation of the cutting speed and the surface finish. Scott et al.12 used a factorial-design method to determine the optimum combination of control parameters in WEDM, the measures of the machining performance being the MRR and the surface finish. It was found that discharge current, pulse duration and pulse frequency are significant control factors. Spedding and Wang13 deve- loped mathematical models using the response-surface methodology (RSM) and an artificial neural network in the WEDM process. They considered the effects of the input parameters (pulse width, time between two pulses, wire mechanical tension and wire-feed speed) on the responses (cutting speed and surface roughness). Yuan et al.14 carried out a multi-objective optimization based on the Gaussian process regression to optimize the high-speed WEDM process, considering mean current, on-time and off-time as the input parameters and MRR and surface roughness as the output responses. Kung and Chiang15 developed mathematical models using RSM to investigate the influences of the machining parameters on the performance characteristics of MRR and surface roughness in a WEDM process. Esme et al.16 constructed both a mathematical model and a neural-network model to predict, reproduce and compare the types of surface roughness under different machining conditions. On the basis of a literature review some insight has been gained Y. KAZANCOGLU et al.: APPLICATION OF A TAGUCHI-BASED NEURAL NETWORK FOR FORECASTING ... 472 Materiali in tehnologije / Materials and technology 46 (2012) 5, 471–476 Table 1: Comparison between neural networks and the Taguchi design method17 Tabela 1: Primerjava med nevronskimi mre`ami in Taguchi metodo oblikovanja17 Comparison Neural networks Taguchi design Computational time Long Medium Model development Yes* No Optimization Through a model Straight Understanding Moderate Normal Software availability Available Available Optimization sensitivity High Normal Application rate (usage frequency) Frequent Rare * No factor-interaction effects into the use of neural networks and Taguchi design methods for modeling and optimizing different WEDM processes. The main advantage of the neural-network method is that it provides modeling and predictions. The main advantage of the Taguchi design method is that it provides an optimization and analyzes the effect of each process parameter on the responses. Table 1 shows a comparison between neural networks and the Taguchi modeling and optimization method17. In the present work, two of the techniques, namely, the neural network with a back-propagation network (BPN) and the Taguchi design method have been employed. Qwiknet 2.23 software and Taguchi Soft program were used for the NN modeling and the Taguchi optimization technique, respectively. Open voltage, pulse duration, wire speed and dielectric flushing pressure were selected as the input factors, whereas surface roughness (Ra) was selected as the response. A BPN model was developed for the prediction of the surface roughness. An analysis-of-variance (ANOVA) table was used to determine the significant WEDM parameter affecting the surface roughness. An approach to determine the optimum machining-parameter setting was proposed on the basis of the Taguchi design method. 2 EXPERIMENTAL SET-UP AND THE TEST PROCEDURE In this experimental study, all experiments were conducted on an Acutex WEDM machine. The WEDM machining set-up is shown in Figure 2. The work material, electrode and other machining conditions are given in Table 2. Table 2: Constant machining condition set-up Tabela 2: Podatki o nastavitvi naprave Workpiece AISI 4340 Wire material CuZn37 Workpiece dimensions (mm) 150 × 150 × 10 Table feed rate (mm/min) 8.2 Pulse-interval time (μs) 18 Wire diameter (mm) 0.25 Wire tensile strength (N/mm2) 900 Machining cut-off length (mm) 0.8 Traversing length (mm) 5 Open-circuit voltage (150–250 V), pulse duration (600, 800, 1000) ns, wire speed (6, 8, 10) mm/min and dielectric flushing pressure (10, 12, 14) kg/cm2 were selected as the input parameters and surface roughness was selected as the output parameter. Surface-roughness (Ra) measurements were made by using a Phynix TR-100 portable surface-roughness tester with  = 0.03 mm for the cut-off length. Three measurements were taken and their average was calculated as Ra value. According to the Taguchi orthogonal-design concept an L18 mixed orthogonal-array table was chosen for the experiments. The orthogonal-array table used in the Taguchi design method was applied to BPN as testing data. BPN was developed to predict the surface roughness. The optimum machining-parameter combination was obtained by using an analysis of the signal-to-noise (S/N) ratio. The signal-to-noise (S/N) ratio is a measure of the magnitude of the data set relative to the standard deviation. If the S/N is large, the magnitude of the signal is large relative to the noise, as measured with the standard deviat- ion.8,18,19 There are several S/N ratios available depending on the types of characteristics. The nominal ratio is the best, higher is better and lower is better. We would select the S/N if the system is optimized when the response is as small as possible.1–5 The S/N ratio for the LB (lower is better) characteristic is calculated by using Equations (1) and (2)5: L n yj i k n = = ∑1 2 1 (1)  j jL= −10 lg (2) where yi is the response value, Lj is the loss function, nj is the S/N ratio. 3 EXPERIMENTAL RESULTS AND DATA ANALYSIS A neural network based on back propagation is a multilayered architecture made up of one or more hidden layers placed between the input and output layers. The components of the input pattern consisted of the control variables of the machining operation (open-circuit voltage, pulse duration, wire speed and dielectric flush- ing pressure), whereas the output-pattern components represented the measured factor (surface roughness). Table 3 shows a Taguchi L18 orthogonal-array plan of the experiment and a training set for the neural-network application. Y. KAZANCOGLU et al.: APPLICATION OF A TAGUCHI-BASED NEURAL NETWORK FOR FORECASTING ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 471–476 473 Figure 2: WEDM machining set-up Slika 2: WEDM-naprava The orthogonal-array table used in the Taguchi design method was applied to BPN as testing data. The network structure was selected to be of the 4 : 5 : 1 type. The used BPN model is shown in Figure 3. The testing validity of the regression analysis and the neural-network results was achieved by using the input parameters according to the design matrix given in Table 4. The performance of each BPN was calculated with the absolute error (%) of the tested subset. The average absolute error was calculated as 1.08 %. The surface roughness for various machining conditions can be predicted in a quick and accurate manner; the BPN results showed that the predicted values were very close to the experimental values. The value of the multiple coefficient R2 is 0.99, which means that the explanatory variables explain 99 % of the variability in the response variable. The predicted values of the surface roughness were compared with the experimental values as shown in Figure 4. The effect and optimization of machining settings for the minimum surface roughness was investigated experi- mentally. The optimum machining-parameter combina- tion was obtained by analyzing the S/N ratio. ANOVA was used to consider the effects of the input factors on the response and was performed on experimental data. Y. KAZANCOGLU et al.: APPLICATION OF A TAGUCHI-BASED NEURAL NETWORK FOR FORECASTING ... 474 Materiali in tehnologije / Materials and technology 46 (2012) 5, 471–476 Figure 4: Comparison of experimental and predicted values Slika 4: Primerjava eksperimentalnih in napovedanih vrednosti Table 5: Results of ANOVA for the surface roughness Tabela 5: Rezultati analize variance (ANOVA) za hrapavost povr{ine Parame- ter code Factors DF SS F MS Contribution percentage (%) A Open-circuitvoltage 1 1.404 1378.58 1.401 49.36 B Pulse duration 2 1.440 711.42 0.720 50.03 C Wire speed 2 0.015 7.32 0.007 0.52 D Flushingpressure 2 0.003 1.68 0.001 0.10 Error 10 0.020 – 0.003 – Total 17 2.844 – 100 Figure 3: 4 : 5 : 1 (4 inputs, 1 hidden layer with 5 neurons and 1 output) type of the BPN algorithm used for modeling Slika 3: 4 : 5 : 1 (4 vhodni podatki, 1 skrit nivo s 5 nevroni in 1 iz- hodni podatek) vrsta BPN algoritma, uporabljenega pri modeliranju Table 3: L18 orthogonal array and a neural-network training set Tabela 3: L18 ortogonalna matrika za usposabljanje nevronske mre`e Exp. no. Open voltage (V) Pulse duration (ns) Wire speed (m/min) Flushing pressure (kg/cm2) Surface roughness (μm) 1 150 600 6 10 2.08 2 150 600 8 12 2.20 3 150 600 10 14 2.21 4 150 800 6 10 2.48 5 150 800 8 12 2.51 6 150 800 10 14 2.52 7 150 1000 6 12 2.79 8 150 1000 8 14 2.82 9 150 1000 10 10 2.86 10 250 600 6 14 2.62 11 250 600 8 10 2.69 12 250 600 10 12 2.72 13 250 800 6 12 3.10 14 250 800 8 14 3.06 15 250 800 10 10 3.09 16 250 1000 6 14 3.36 17 250 1000 8 10 3.40 18 250 1000 10 12 3.45 Table 4: Test set for the validity of the constructed neural network Tabela 4: Zbirka podatkov za preizku{anje postavljene nevronske mre`e Exp. no. Open voltage (V) Pulse duration (ns) Wire speed (m/min) Flushing pressure (kg/cm2) Surface roughness (μm) Experi- mental Predicted 1 150 600 8 10 2.20 2.20 2 250 600 6 12 2.62 2.58 3 150 800 8 10 2.53 2.52 4 250 800 6 14 2.95 2.97 5 150 1000 8 10 2.85 2.90 6 150 1000 6 14 2.72 2.68 7 250 1000 10 10 3.56 3.46 8 250 800 10 12 3.13 3.15 Average maximum error: 1.08 % The confidence level was selected as 95 %. The results of ANOVA for the surface roughness are shown in Table 5. After analyzing Table 4, it is observed that the open-circuit voltage and the pulse duration have a great influence on the obtained surface roughness. The wire speed and dielectric flushing pressure do not affect significantly the obtained surface roughness. The plot of the mean-factor effects is shown in Figure 5. The S/N graph for the surface roughness is shown in Figure 6. It is evident that open-circuit voltage (49.36 %) and pulse duration (50.03 %) have the most significant effect on the surface roughness, which means that by increasing these two parameters we also increase the surface roughness. Wire speed (0.52 %) has little effect on the surface roughness. The effect of dielectric flushing pressure (0.10 %) is negligible. Optimum factor levels and S/N ratios obtained at the end of the Taguchi design technique are summarized in Table 6. Based on the S/N ratio plot in Figure 6, the optimum machining parameters for the surface roughness are open-circuit voltage at level 1, pulse duration at level 1, wire speed at level 1 and dielectric flushing pressure at level 1 (A1B1C1D1). 4 CONFIRMATION TESTS A confirmation experiment is the final step in the first iteration of designing an experiment process.5,8 The purpose of the confirmation experiment is to validate the conclusions drawn during the analysis phase. The confirmation experiment is performed by conducting a test with a specific combination of the factors and levels previously evaluated.8,20 In this study, after determining the optimum conditions and predicting the response under these conditions, a new experiment was designed and conducted with the optimum levels of the welding parameters. The final step is to predict and verify the improvement of the performance characteristic. The predicted S/N ratio  using the optimum levels of the welding parameters can be calculated as:5,8,20  ( )   m i m i n + − = ∑ 1 (3) where m is the total mean of the S/N ratio, i is the mean of the S/N ratio at the optimum level, and n is the number of the main welding parameters that signifi- cantly affect the performance.8,18–20 The result of the experimental confirmation using the optimum surface roughness is shown in Table 7. Table 7: Results of the confirmation experiments for the surface roughness Tabela 7: Rezultati potrditvenih eksperimentov za hrapavost povr{ine Initial parameters Optimumparameters Parameter level A1B3C3D1 A1B1C1D1 Surface roughness (μm) 2.86 2.10 S/N ratio (dB) –9.13 –6.57 The improvement in the S/N ratio from the initial machining parameters to the optimum machining parameters is 2.56 dB. Based on the result of the Y. KAZANCOGLU et al.: APPLICATION OF A TAGUCHI-BASED NEURAL NETWORK FOR FORECASTING ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 471–476 475 Figure 6: S/N graph for the surface roughness Slika 6: S/N-diagram za hrapavost povr{ine Figure 5: The effect of machining parameters on the surface rough- ness Slika 5: U~inek parametrov obdelave na hrapavost povr{ine Table 6: Optimum factor levels and their S/N ratios Tabela 6: Nivoji optimalnih faktorjev in njihova S/N-razmerja Code Factors S/N ratio (dB) Level 1 Level 2 Level 3 Max-min Rank A Open-circuitvoltage –7.89 –9.66 – 1.77 2 B Pulseduration –7.63 –8.88 –9.83 2.20 1 C Wire speed –8.65 –8.80 –8.88 0.23 3 D Flushingpressure –8.73 –8.84 –8.75 0.11 4 Average S/N = –8.80 confirmation test, the surface roughness is decreased 1.36 times. 5 CONCLUSIONS This study presents a prediction, optimization and modeling of the surface roughness of the AISI 4340 steel in a wire-electrical-discharge machining (WEDM) process based on the Taguchi-based neural network with the back-propagation algorithm method. The following conclusions can be drawn from this study: • The main WEDM parameters that affect the surface roughness of the machined parts were determined as pulse duration and open-circuit voltage among four controllable factors influencing the surface roughness using ANOVA, • A neural network based on the back-propagation network (BPN) algorithm was constructed for pre- dicting the surface roughness. The predicted values were found to be very close to the experimental values, • The optimum parameter combination for the mini- mum surface roughness was obtained by using the Taguchi design method with an analysis of the S/N ratio, • The confirmation test supports the finding that the surface roughness is greatly decreased by using the optimum design parameters, • The obtained results indicate that the BPN model agreed well with the Taguchi analysis. 6 REFERENCES 1 J. T. Huang, Y. S. Lmiao, W. J. Hsue, Journal of Materials Processing Technology, 87 (1999), 69–81 2 N. Tosun, C. Cogun, G. Tosun, Journal of Materials Processing Technology, 152 (2004), 316–322 3 A. Mohammadi, A. F. Tehrani, E. Emanian, D. Karimi, Journal of Materials Processing Technology, 205 (2008), 283–289 4 D. Scott, S. Boyina, K. P. Rajurkar, Int. J. Prod. Res., 29 (1991), 2189–2207 5 D. C. Montgomery, Design and Analysis of Experiments, Wiley, Singapore 1991 6 S. Fraley, M. Oom, B. Terrien, J. Z. Date, The Michigan Chemical Process Dynamic and Controls Open Text Book, USA, 2006 7 U. Esme, The Arabian Journal for Science and Engineering, 34 (2009), 519–528 8 S. C. Juang, Y. S. Tarng, Journal of Materials Processing Technology, 122 (2002), 33–37 9 Y. C. Lin, C. H. Cheng, B. L. Su, L. R. Hwang, Materials and Manufacturing Process, 21 (2006), 922–929 10 N. Tosun, C. Cogun, Journal of Materials Processing Technology, 134 (2003), 273–278 11 Y. S. Tarng, S. C. Ma, L. K. Chung, International Journal of Machine Tools & Manufacture, 35 (1995), 1693–1701 12 D. Scott, S. Boyina, K. P. Rajurkar, International Journal Production Research, 11 (1991), 2189–2207 13 T. A. Spedding, Z. Q. Wang, Precision Engineering, 20 (1997), 5–15 14 J. Yuan, K. Wang, T. Yu, M. Fang, International Journal of Machine Tools & Manufacture, 48 (2008), 47–60 15 K. Y. Kung, K. T. Chiang, Materials and Manufacturing Process, 23 (2008), 241–250 16 U. Esme, A. Sagbas, F. Kahraman, Iranian Journal of Science & Technology, Transaction B, Engineering, 33 (2009), 231–240 17 K. Y. Benyounis, A. G. Olabi, Advances in Engineering Software, 39 (2008), 483–496 18 D. S. Holmes, A. E. Mergen, Signal to noise ratio–What is the right size, www.qualitymag.com/.../Manuscript%20Holmes%20&%20 Mergen.pdf, USA, 1996, 1–6 19 Y. Kazancoglu, U. Esme, M. Bayramoglu, O. Guven, S. Ozgun, Mater. Tehnol., 45 (2011) 2, 105–110 20 P. J. Ross, Taguchi Techniques for Quality Engineering, 2nd ed., McGraw-Hill, New York 1996 Y. KAZANCOGLU et al.: APPLICATION OF A TAGUCHI-BASED NEURAL NETWORK FOR FORECASTING ... 476 Materiali in tehnologije / Materials and technology 46 (2012) 5, 471–476 D. @IVKOVI] et al.: PREDICTION OF THE THERMODYNAMIC PROPERTIES FOR LIQUID Al-Mg-Zn ALLOYS PREDICTION OF THE THERMODYNAMIC PROPERTIES FOR LIQUID Al-Mg-Zn ALLOYS NAPOVEDOVANJE TERMODINAMI^NIH LASTNOSTI TEKO^E ZLITINE Al-Mg-Zn Dragana @ivkovi}1, Yong Du2, Ljubi{a Balanovi}1, Dragan Manasijevi}1, Du{ko Mini}3, Nade`da Talijan4 1University of Belgrade, Technical Faculty, Bor, Serbia 2 State Key Laboratory of Powder Metallurgy, Central South University, Changsha, Hunan, China 3University of Pri{tina, Faculty of Technical Sciences, Kosovska Mitrovica, Serbia 4University of Belgrade, Institute of Chemistry, Technology and Metallurgy, Belgrade, Serbia dzivkovic@tf.bor.ac.rs Prejem rokopisa – received: 2012-02-22; sprejem za objavo – accepted for publication: 2012-03-29 The results of a thermodynamic-property prediction for liquid Al-Mg-Zn alloys using the general solution model are presented in this paper. Calculations were done in nine sections of the system with different molar ratios of Mg:Zn, Zn:Al and Al:Mg in the temperature range of 900–1200 K. Partial and integral molar quantities – including the activities for all three components, the integral molar excess Gibbs energies and the integral molar enthalpies of mixing – were obtained. Some of the calculation results were compared with the experimental data available in the literature, showing a good agreement with it. Keywords: thermodynamics of alloys, Al-Mg-Zn system, general solution model V ~lanku so predstavljeni rezultati raziskav termodinami~nih lastnosti teko~ih zlitin Al-Mg-Zn, napovedanih z uporabo splo{nega modela raztapljanja. Izra~uni so bili izvr{eni v devetih prerezih sistema z razli~nimi molarnimi dele`i Mg:Zn, Zn:Al in Al:Mg v obmo~ju temperatur 900–1200 K. Dobljene so bile parcialne in celotne molarne koli~ine, vklju~no z aktivnostmi za vse tri komponente, skupni molarni prese`ek Gibssove energije in skupna molarna entalpija me{anja. Ugotovljeno je dobro ujemanje izra~unanih rezultatov z razpolo`ljivimi eksperimentalnimi podatki iz literature. Klju~ne besede: termodinamika zlitin, sistem Al-Mg-Zn, splo{ni model raztapljanja 1 INTRODUCTION The so-called ZA alloys – zinc-aluminum-based alloys – have a wide application in different fields of industry1,2. The ternary Al-Mg-Zn system belongs to this group of materials, which are of interest as the lead-free solders for die attach1–4. Therefore, different properties of this system were investigated in order to define it more completely5–9. The thermodynamics and the phase equilibria of the Al-Mg-Zn system have been examined widely10–20. Most of the literature data is related to the phase-diagram determination10–17. A complete reference compilation concerning the experimental data obtained for the above-mentioned ternary alloys up to 1998 can be found in the work of Liang et al.19, while the last review is given in an article by Raghavan15 from 2010. The liquidus projection of the Al-Mg-Zn system is shown in Figure 1, according to Refs. 14 and 17. Among the numerous researches, there are only a few thermodynamic studies17–19. Experimental thermodyna- mic investigations of the Al-Mg-Zn system in the liquid state were done for the chosen sections at the temperatu- res of 883 K and 933 K using vapor-pressure measure- ments17, EMF18 and mixing calorimetry19, while the thermodynamic assessments can be found in20,21. Considering the available literature and the lack of a complete thermodynamic data with respect to the wider temperature and concentration ranges, the results of the thermodynamic-property prediction for the liquid Al-Mg-Zn alloys in the temperature interval of 900–1200 K, using the general solution model, are given in this paper as a contribution to a full thermodynamic description of this ternary system. 2 THEORETICAL FUNDAMENTALS The general solution model for the calculation of the thermodynamic properties of ternary systems based on the known binary thermodynamic data has been provided by Chou22,23. It breaks down the boundary between symmetrical and asymmetrical models, and has already been proved in some practical examples24,25 as the correct and accurate model. This model was developed for multicomponent systems and its basic equations are as follows22: ΔG x x A A x x AE i i j i j m j ij ij i j ij k k i j m = + ⋅ − + = ≠ = ≠ ∑ ∑ , , ( ) 1 0 1 1 1 xk i ij k( )( ) ( )2 1 − ⎛ ⎝ ⎜ ⎜ ⎜ ⎞ ⎠ ⎟ ⎟ ⎟ ⎡ ⎣ ⎢ ⎢ ⎢ ⎤ ⎦ ⎥ ⎥ ⎥ (1) Materiali in tehnologije / Materials and technology 46 (2012) 5, 477–482 477 UDK 544.3:669.017.13:669.715'721'5 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)477(2012) where Aoij, A1ij, A2ij are the regular-solution parameters for the binary system ij independent of the composition, relying only on the temperature: GEij = XiXj (A o ij + A 1 ij (Xi – Xj) + A 2 ij (Xi – Xj) 2 + ... + Anij (Xi – Xj) 2) (2) where Xi and Xj indicate the mole fractions of compo- nents i and j in the ij binary system, which is expressed as: X x xi ij i k i ij k k k i j m ( ) ( ) ( ) , = + = ≠ ∑  1 (3) and where the coefficient entered as  i ij k ( ) ( ) in Eq.(3) presents the similarity coefficient of component k to component i in the ij system, and is defined as:    i ij k ij ik ij ik ji jk( ) ( ) ( , ) ( , ) ( , ) = + (4) where (ij,ik) is the function related to the excess Gibbs free energy of the ij and ik binaries: ( , ) ( )ij ik G G Xij E ik E i x X i i = − = = ∫ Δ Δ 2 0 1 d (5) In all the equations given, GE and GEij refer to the integral molar excess free energies for the multicom- ponent and binary systems, respectively, while x1, x2, x3 refer to the mole fraction of the components in the investigated multicomponent system. 3 RESULTS AND DISCUSSION Thermodynamic calculations in the Al-Mg-Zn ternary system were carried out in nine sections along the lines of the following constant molar ratios: Mg : Zn = 1 : 3, 1 : 1, 3 : 1 – the sections from the Al corner; Zn-Al = 1 : 3, 1 : 1, 3 : 1 – the sections from the Mg corner; and Al : Mg = 1 : 3, 1 : 1, 3 : 1 – the sections from the Zn corner. The basic data necessary for the calculation was taken from the literature20,26,27. The Redlich-Kister polynomials for the constitutional binaries in the investigated ternary Al-Mg-Zn system are presented in Table 1. Table 1: Redlich-Kister parameters for the liquid phase in the con- stitutional binaries of the Al-Mg-Zn system Tabela 1: Redlich-Kisterjevi parametri za staljeno fazo v sestavnih binarnih sistemih iz sistema Al-Mg-Zn System ij Al-Mg (20) Mg-Zn (20) Al-Zn (26) Aoij (T) –12000+8.566*T –77729.24+ 680.52266*T –95*T*ln(T)+ 40E–3*T2 10465.55– 3.39259*T A1ij (T) 1894–3*T 3674.72+0.57139*T / A2ij (T) 2000 –1588.15 / The prediction was done according to the funda- mentals of the latest version of the general-solution model22,23. Based on the starting data in Table 1, similarity coefficients were determined and further calculations were carried out for 81 alloys in all the selected cross sections of the investigated ternary Al-Mg-Zn system in the temperature interval of 900–1 200 K, as shown with Eqs.(1–5). The integral molar enthalpies of mixing were additionally calculated according to following expression: d d ( / )Δ ΔG T T H T E M = − 2 (6) The results of the thermodynamic predictions, includ- ing the values of the ternary integral molar excess Gibbs D. @IVKOVI] et al.: PREDICTION OF THE THERMODYNAMIC PROPERTIES FOR LIQUID Al-Mg-Zn ALLOYS 478 Materiali in tehnologije / Materials and technology 46 (2012) 5, 477–482 Figure 1: Al-Mg-Zn liquidus projection: a) 17 and b) 14 Slika 1: Projekcija likvidusa Al-Mg-Zn: a) 17 in b) 14 D. @IVKOVI] et al.: PREDICTION OF THE THERMODYNAMIC PROPERTIES FOR LIQUID Al-Mg-Zn ALLOYS Materiali in tehnologije / Materials and technology 46 (2012) 5, 477–482 479 Figure 3: Dependence of the integral molar enthalpies of mixing on the composition and temperature in the Al-Mg-Zn system: a) sections from the zinc corner; b) sections from the aluminum corner; c) sections from the magnesium corner Slika 3: Odvisnost skupne molarne entalpije me{anja od sestave in temperature v sistemu Al-Mg-Zn: a) prerez iz cinkovega kota; b) prerez iz aluminijevega kota; c) prerez iz magnezijevega kota Figure 2: Dependence of the integral molar excess energy on the composition and temperature in the Al-Mg-Zn system: a) sections from the zinc corner; b) sections from the aluminum corner; c) sections from the magnesium corner Slika 2: Odvisnost skupne molarne prese`ne energije od sestave in temperature v sistemu Al-Mg-Zn: a) prerez iz cinkovega kota; b) prerez iz aluminijevega kota; c) prerez iz magnezijevega kota energy, the ternary molar enthalpy of mixing and the activities of all three components in the liquid phase, were calculated for all the investigated sections at the investigated temperatures, and presented in Table 2 and Figures 2 to 4, respectively. The calculated activity values for all three components were used for the construction of the iso-activity diagrams at 1000K and shown in Figure 5. Negative values of the integral molar excess Gibbs energies were obtained for most of the concentration range at all the investigated temperatures (Figure 2). The most negative value of about –3.5 kJ/mol was present in the section from the aluminum corner with a molar ratio of Mg : Zn = 1 : 1 for the low aluminum concentrations, while the highest positive values of about 0.2 kJ/mol were noticed for the higher contents of zinc and aluminum in sections Mg : Zn = 1 : 1 and Al : Mg = 3 : 1. In the case of the integral molar enthalpies of mixing, the minimum value of -5kJ/mol was noticed for the low aluminum contents in the section Mg : Zn = 1 : 1, while the maximum value of about +3 kJ/mol was obtained for the low magnesium contents in section Al : Zn = 1 : 1. Different deviations from Raoult law were detected considering three constituent metals in the Al-Mg-Zn system. Aluminum shows a positive deviation in the whole composition range of the investigated ternary system, moving towards almost an ideal behavior in the case of the section with a molar ratio of Mg : Zn = 3 : 1. On the other hand, magnesium shows a uniform negative deviation for all the examined sections of the system, D. @IVKOVI] et al.: PREDICTION OF THE THERMODYNAMIC PROPERTIES FOR LIQUID Al-Mg-Zn ALLOYS 480 Materiali in tehnologije / Materials and technology 46 (2012) 5, 477–482 Figure 5: Iso-activity diagrams for the constitutive elements in the ternary Al-Mg-Zn system at 1000 K Slika 5: Diagram izoaktivnosti za sestavne elemente v ternarnem sistemu Al-Mg-Zn pri 1000 K Figure 4: Activity dependence on the composition and temperature in the investigated Al-Mg-Zn system: a) sections from the zinc corner; b) sections from the aluminum corner; c) sections from the magnesium corner Slika 4: Odvisnost aktivnosti od sestave in temperature v preiskovanem sistemu Al-Mg-Zn: a) prerez iz cinkovega kota; b) prerez iz aluminijevega kota; c) prerez iz magnezijevega kota while zinc behaves differently – showing a slightly positive deviation for section Al : Mg = 3 : 1 and negative deviations in the other two sections. The temperature influence on the calculated thermo- dynamic properties was not significant in the investi- gated interval 933–1200 K. The described tendencies indicate a prevalent exi- stence of the mutual mixing tendencies between the constitutive components in the Al-Mg-Zn system at the investigated temperatures, where magnesium and zinc exhibit a more significant mixing tendency than alumi- num. The calculated thermodynamic quantities were com- pared with the available literature data at the temperature of 933 K 19,20 in order to test the accuracy of the applied prediction model. These comparisons are shown in Figure 6 for different examples – the magnesium acti- vity (Figure 6a), the magnesium chemical potential (b) and the integral molar enthalpies of mixing for the three sections from the zinc corner (c). As can be seen, a good agreement was noticed between the results of this work and the reference experimental data19,20. 4 CONCLUSION The calculation of the thermodynamic properties in the ternary Al-Mg-Zn system was done by applying the general solution model. On the basis of the thermo- dynamic parameters from the constituent binary subsystems, the integral molar excess Gibbs energies and the integral molar enthalpies of mixing were calculated for the whole system, in nine sections from different corners, in the temperature range of 900–1 200 K. The obtained data showed a mostly negative deviation from Raoult law, indicating predominantly mutual mixing tendencies in the investigation system. We found that: (i) experimental investigation and thermodynamic-property determination at the selected temperatures are rather difficult to perform due to the evaporation of zinc and oxidation of magnesium in the case of the investigated Al-Mg-Zn alloys; (ii) there is a good agreement between the available experimental data and the data calculated in this paper; and (iii) due to the incomplete thermodynamic data relating to the investigated system recorded in the reference literature, the predicted results from this paper can be taken as relevant thermodynamic data relating to the examined multicomponent ZA-based system. This can be done because the accuracy of the model, used in different cases, had already been proven as cited in literature24,25 and it is important to continuously examine the Al-Mg-Zn alloys28 and other Al-based ternary alloys29,30. D. @IVKOVI] et al.: PREDICTION OF THE THERMODYNAMIC PROPERTIES FOR LIQUID Al-Mg-Zn ALLOYS Materiali in tehnologije / Materials and technology 46 (2012) 5, 477–482 481 Figure 6: Comparison of calculated and reference-literature experi- mental values19,20 Slika 6: Primerjava izra~unanih podatkov z literaturnimi eksperimen- talnimi vrednostmi19,20 Table 2: Characteristic dependencies of the integral molar excess energies and the integral molar enthalpies of mixing on the compo- sition of the ternary Al-Mg-Zn alloys expressed as GE (J/mol) = Ax2 + Bx + C and HM (J/mol) = Dx2 + Ex + F at the investigated tem- peratures Tabela 2: Zna~ilna odvisnost skupne prese`ne molarne energije in skupne molarne entalpije me{anja od sestave ternarne Al-Mg-Zn zlitine, izra`ena kot GE (J/mol) = Ax2 + Bx + C in HM (J/mol) = Dx2 + Ex + F pri preiskovanih temperaturah 933K Section A B C D E F Mg:Zn=1:3 –7444.08 10598.36 –3114.3 –10228 15172 –4917 Mg:Zn=1:1 –4584.48 8210.277 –3522.5 –5510.5 11645 –6040.4 Mg:Zn=3:1 –564.357 2931.14 –2279.32 2085.3 2368.8 –4300.3 Al:Zn=1:3 13523.87 –14194.9 1171.436 22679 –24340 1791.2 Al:Zn=1:1 11246.21 –12525.6 1697.977 20338 –22847 2566.1 Al:Zn=3:1 8269.261 –9230.4 1314.962 16806 –22847 2028.8 Al:Mg=1:3 8058.779 –8123.02 –463.269 13330 –11380 –2168.7 Al:Mg=1:1 2852.942 –2223.04 –930.149 4233.1 –1474.4 –2892.2 Al:Mg=3:1 –2062.89 2678.725 –717.506 –3290 5185.5 –1942.2 1000K Section A B C D E F Mg:Zn=1:3 –7301.1 10323.32 –2990.65 –10006 14727 –4694.9 Mg:Zn=1:1 –4608.32 8045.501 –3351 –5214.3 11053 –5744.2 Mg:Zn=3:1 –844.82 3049.588 –2143.19 2307.4 1924.6 –4078.2 Al:Zn=1:3 12810.45 –13429.3 1126.216 21791 –23452 1791.2 Al:Zn=1:1 10540.78 –11745.5 1636.327 19746 –22255 2566.1 Al:Zn=3:1 7633.816 –8521.86 1264.693 16510 –18548 2028.8 Al:Mg=1:3 7728.575 –7927.42 –334.914 12441 –10492 –2168.7 Al:Mg=1:1 2754.998 –2275.25 –785.379 3640.8 –882.13 –2892.2 Al:Mg=3:1 –1990.84 2513.915 –628.224 –3586.1 5481.6 –1942.2 1100K Section A B C D E F Mg:Zn=1:3 –7124.355 9967.950 –2827.029 –9799.6 14315 –4488.7 Mg:Zn=1:1 –4692.6 7872.956 –3122.95 –4939.3 10503 –5469.2 Mg:Zn=3:1 –1299.82 3281.329 –1960.96 2513.6 1512.1 –3872 Al:Zn=1:3 11814.41 –12358.4 1059.365 20966 –22627 1791.2 Al:Zn=1:1 9524.121 –10620.6 1544.861 19196 –21705 2566.1 Al:Zn=3:1 6700.34 –7480.99 1189.845 16235 –18273 2028.8 Al:Mg=1:3 7304.49 –7707.24 –142.743 11616 –9666.6 –2168.7 Al:Mg=1:1 2644.862 –2392.36 –568.803 3090.8 –332.13 –2892.2 Al:Mg=3:1 –1868.6 2251.542 –494.81 –3861.1 5756.6 –1942.2 1200K Section A B C D E F Mg:Zn=1:3 –6979.237 9651.942 –2674.257 –9743.3 14202 –4432.4 Mg:Zn=1:1 –4817.76 7751.966 –2909.49 –4864.3 10353 –5394.2 Mg:Zn=3:1 –1784.52 3551.037 –1789.76 2569.9 1399.6 –3815.7 Al:Zn=1:3 10843.84 –11317.1 993.472 20741 –22402 1791.2 Al:Zn=1:1 8512.729 –9505.34 1454.325 19046 –21555 2566.1 Al:Zn=3:1 5765.485 –6441.27 1115.435 16160 –18198 2028.8 Al:Mg=1:3 6905.864 –7515.85 50.01534 11391 –9441.6 –2168.7 Al:Mg=1:1 2538.665 –2517.19 –351.668 2940.8 –182.13 –2892.2 Al:Mg=3:1 –1749.7 1990.177 –361.146 –3936.1 5831.6 –1942.2 Acknowledgment The results of this paper were obtained in the frame of Project OI 172037 financed by the Ministry of Science and Technological Development, the Republic of Serbia, and a bilateral scientific and technological cooperation project between the Republic of Serbia and the People’s Republic of China (2011–2012). 5 REFERENCES 1 T. Shimizu, H. Ishikawa, I. Ohnuma, K. Ishida, Journal of Electronic Materials, 28 (1999), 1172 2 M. Rettenmayr, P. Lambracht, B. Kempf, C. Tschudin, Journal of Electronic Materials, 31 (2002) 4, 278 3 Lj. Balanovi}, D. @ivkovi}, A. Mitovski, D. Manasijevi}, @. @iv- kovi}, Journal of Thermal Analysis and Calorimetry, 103 (2011) 3, 1055 4 P. Bro`, D. @ivkovi}, J. Medved, N. Talijan, D. Manasijevi}, G. Klan~nik, Experimental and theoretical study of thermodynamic properties and phase equilibria in ternary Al-Zn-X alloys, in COST MP0602 Book (Volume 3 Chapter 6), in print 5 G. Bergman, J. L. T. Waugh, L. Pauling, Acta Crystalographica, 10 (1957), 254 6 T. Takeuchi, S. Murasaki, A. Matsumoro, U. Mizutani, Journal of Non Crystalline Solids, 156–158 (1993), 914 7 A. Niikura, A. P. Tsai, N. Nishiyama, A. Inoue, T. Matsumoto, Mate- rials Science Engineering, 181/182 A (1994), 1387 8 N. K. Mukhopaghyay, J. Bhatt, A. K. Pramanick, B. S. Murty, P. Paufler, Journal of Materials Science, 39 (2004), 5155 9 M. Zhu, G. Yang, D. Wan, Z. Wang, Y. Zhou, Rare Metals, 28 (2009) 4, 401 10 P. Donnadieu, A. Quivy, T. Tarfa, P. Ochin, A. Dezellus, M. G. Har- melin, P. Liang, H. L. Lukas, H. J. Seifert, F. Aldinger, G. Effenberg, Zeitschrift für Metallkunde, 88 (1997) 12, 911 11 D. Petrov, A. Watson, J. Grobner, P. Rogl, J. C. Tedenac, M. Bulano- va, Turkevich, Alluminium-magnesium-zinc, Ternary Alloys Sys- tems, Vol.11A3, G. Effenberg, S. Ilyenko, Ed., Springer, Germany, 2006 12 Y. P. Ren, G. W. Qin, W. L. Pei, Y. Gio, H. D. Zhao, H. X. Li, M. Jiang, S. M. Hao, Journal of Alloys and Compounds, 481 (2009), 176 13 M. Ohno, D. Mirkovi}, R. Schmid-Fetzer, Materials Science Engi- neering, 421A (2006), 328 14 V. Raghavan, Journal of Phase Equilibria and Diffusion, 28 (2007), 203 15 V. Raghavan, Journal of Phase Equilibria and Diffusion, 31 (2010), 29 16 A. L. Voskov, G. F. Voronin, Russian Journal of Physical Chemistry A, 84 (2010), 525 17 Z. Kozuka, J. Moriyama, I. Kushima, Journal of the Electrochemical Society of Japan, 28 (1960), 298 18 A. M. Pogodaev, E. E. Lukashenko, Russian Metallurgy, 6 (1974), 74 19 Y. B. Kim, F. Sommer, B. Predel, Journal of Alloys and Compounds, 247 (1997), 43 20 P. Liang, T. Tarfa, J. A. Robinson, S. Wagner, P. Ochin, M. G. Har- melin, H. J. Seifert, H. L. Lukas, F. Aldinger, Thermochimica Acta, 314 (1998), 87 21 H. Liang, S. L. Chen, Y. A. Chang, Metallurgical and Materials Tran- sactions, 28A (1997), 1725 22 K. C. Chou, K. Wei S., Metallurgical and Materials Transactions, 28B (1997), 439 23 K. C. Chou, CALPHAD, 19 (1995), 315 24 D. @ivkovi}, @. @ivkovi}, Y. H. Liu., Journal of Alloys and Com- pounds, 265 (1998), 176 25 D. @ivkovi}, I. Katayama, L. Gomid`elovi}, D. Manasijevi}, R. No- vakovi}, International Journal of Materials Research, 98 (2007) 10, 1025 26 S. Sabine an Mey, Zeitschrift für Metallkunde, 84 (1993) 7, 451 27 A. T. Dinsdale, A. Kroupa, J. Vizdal, J. Vre{tal, A. Watson, A. Ze- manova, COST 531 Database for Lead-free Solders, Ver. 3.0, 2008 28 H. R. Zaid, A. M. Hatab, A. M. A. Ibrahim, Journal of Mining and Metallurgy, Section B-Metallurgy, 47 (2011) 1, 31 29 G. Klan~nik, J. Medved, J. Min. Metall. Sect. B-Metall., 47 (2) B (2011), 179 30 X. Fang, M. Song, K. Li, Y. Du, J. Min. Metall. Sect. B-Metall., 46 (2) B (2010), 171 D. @IVKOVI] et al.: PREDICTION OF THE THERMODYNAMIC PROPERTIES FOR LIQUID Al-Mg-Zn ALLOYS 482 Materiali in tehnologije / Materials and technology 46 (2012) 5, 477–482 D. KLOB^AR et al.: FRICTION-STIR WELDING OF ALUMINIUM ALLOY 5083 FRICTION-STIR WELDING OF ALUMINIUM ALLOY 5083 VARJENJE S TRENJEM IN ME[ANJEM ALUMINIJEVE ZLITINE 5083 Damjan Klob~ar1, Ladislav Kosec2, Adam Pietras3, Anton Smolej2 1Faculty of Mechanical Engineering, University of Ljubljana, A{ker~eva 6, 1000 Ljubljana, Slovenia 2Faculty of Natural Sciences and Engineering, University of Ljubljana, A{ker~eva 12, 1000 Ljubljana, Slovenia 3Instytut Spawalnictwa, Ul. B³. Czes³awa 16/18 Gliwice, Poland damjan.klobcar@fs.uni-lj.si Prejem rokopisa – received: 2012-02-22; sprejem za objavo – accepted for publication: 2012-03-16 A study was made of the weldability of 4-mm-thick aluminium-alloy 5083 plates using friction-stir welding. A plan of experiments was prepared based on the abilities of a universal milling machine, where the tool-rotation speed varied from 200 r/min to 1250 r/min, the welding speed from 71 mm/min to 450 mm/min and the tool tilt angle was held constant at 2°. The factors feed per revolution (FPR) and revolution per feed (RPF) were introduced to get a better insight into the friction-stirring process. Samples for microstructure analyses, Vickers micro-hardness measurements and special miniature tensile-testing samples were prepared. The microstructure was prepared for observation on a light microscope under a polarised light source. A set of optimal welding parameters was determined at a FPR of 0.35 mm/r, at which quality welds can be made with a minimal increase in the weld hardness and an up to 15 % drop in the tensile strength. Keywords: friction-stir welding, EN-AW 5083, welding parameters, mechanical properties, welding defects Izdelana je {tudija varivosti 4 mm debele plo~evine iz aluminijeve zlitine 5083 pri varjenju s trenjem in me{anjem. Na~rt eksperimentov je bil pripravljen na podlagi sposobnosti univerzalnega frezalnega stroja. Spreminjali smo hitrost vrtenja orodja od 200–1250 r/min, hitrost varjenja 71–450 mm/min, kot nagiba orodja pa je bil konstanten pri 2°. Vpeljana sta bila faktorja podajanje na vrtljaj (FPR) in obratov na podajanje (RPF), s katerima bolj nazorno prika`emo vpliv parametrov procesa. Iz izdelanih varov smo pripravili vzorec za analizo mikrostrukture, vzorec za meritev trdote po Vickersu ter posebne miniaturne epruvete za natezni preizkus. Mikrostruktura je bila pripravljena za opazovanje na svetlobnem mikroskopu v polarizirani svetlobi. Optimalni parametri varjenja so bili ugotovljeni pri FPR 0,35 mm/r, pri ~emer dobimo kakovostne zvare, z minimalnim pove~anjem trdote vara in do 15-odstotnim padcem natezne trdnosti. Klju~ne besede: varjenje s trenjem in me{anjem, EN-AW 5083, varilni parametri, mehanske lastnosti, napake v varu 1 INTRODUCTION The 5083 aluminium alloy exhibits good corrosion resistance to seawater and the marine atmosphere, mode- rate mechanical properties and a high fatigue-fracture resistance. It has good formability, machinability and weldability using arc processes (metal inert gas – MIG or tungsten inert gas – TIG) or resistance welding.1,2 This alloy is used for the production of welded components for shipbuilding and railway vehicles, different panels and platforms for boats and trains, storage tanks, cryogenics, pressure vessels, piping, tubing, welded tank trailers and welded dump bodies for the automotive industry, collapsible bridges, armour plates and the bodies of military vehicles. It can be subject to inter- crystalline and stress-corrosion cracking after under- going an unsuitable thermal treatment (welding). It should not to be used above 65 °C for an extended time if later exposed to a corrosive environment.3 If the aluminium alloys are friction-stir processed (FSP) then superplastic properties are obtained, as a consequence of the grain refinement.4–7 The surface of the aluminium alloys can be modified using shot pining and laser shot pinning, with a consequent influence on the microhard- ness, residual stresses, fatigue strength and corrosion resistance.8–10 Friction-stir welding (FSW) is a solid-state joining method that is energy and environmentally friendly and versatile. FSW joints have a high fatigue strength, require less preparation, and little post-weld dressing. These welds have fewer defects than fusion welds and the process enables the welding of dissimilar metals.11 FSW has attracted significant research interest from industries like aerospace and transportation. Many studies were made on the weldability of 5083 aluminium alloy.12–14 Some researchers studied the influence of FSW parameters on fatigue life.12,14 They discovered that the rotational speed governs defect occurrence and a strong correlation between the frictional power input, the tensile strength and the low-cycle fatigue life is obtained. Han et al.15 investigated the optimal conditions for FSW in correlation with welds’ mechanical properties. These mechanical properties were similar to the base alloy at tool rotations between 500 r/min and 800 r/min at a weld-tool travel speed of 124 mm/min. Hirata et al.16 investigated the influence of the FSW parameters on the grain size and the formability. They discovered that a decrease of the frictional heat flow during FSW Materiali in tehnologije / Materials and technology 46 (2012) 5, 483–488 483 UDK 669.715:621.791.1 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)483(2012) decreases the grain size, increases the ductility and improves the formability. Sato et al.17 studied the influence of an oxide array on the formability of tailored blanks.18 They discovered that a band of collective oxide particles could act as an initiation site for cracking during the forming processes. The research of Zucchi et al.19 investigated the pitting and stress-corrosion cracking (PSCC) resistance of the 5083 aluminium alloy during FSW and MIG welding. The FSW welds showed better resistance to PSCC than the base alloy and much better resistance than MIG welds. In this study the weldability of the 5083 aluminium alloy using friction-stir welding (FSW) was investigated. Welding parameters, i.e., the tool-rotation speed, the welding speed and the tilt angle have an influence on the formation of welding defects, the weld apices appear- ance, the microstructure and the weld strength. The aim of this research was to discover the welding parameters providing a weld microstructure without defects. The FSW employed a tool-rotation speed from 200 r/min to 1250 r/min, a welding speed from 71 mm/min to 450 mm/min and the tool-tilt angle was held constant at 2°. The factors feed per revolution (FPR) and revolutions per feed (RPF) were introduced to get a better insight into the friction-stirring process. The RPF gives infor- mation about the heat input per weld length. Miniature samples for tensile testing were prepared from the welds. The welds were examined under the light of an optical microscope and the Vickers hardness was measured. 2 EXPERIMENTAL 2.1 Dimensions and composition of the workpieces The standard EN-AW 5083 aluminium alloy with chemical composition in mass fractions: 4.35 % Mg, 0.42 % Mn, 0.12 % Si, 0.087 % Cr, 0.29 % Fe, 0.019 % Zn, 0.013 % Ti and the rest Al, and temper O, was used for testing. The workpiece dimensions were 180 mm × 60 mm × 4 mm. The physical and mechanical properties of the alloy for temper O were not determined but taken according to the standard (Table 1).3 2.2 FSW tool The FSW tool was made from standard EN 42CrMo4 steel20. A basic FSW tool geometry was used with a threaded pin that was 3.9 mm long (M6 × 1.5) and the concave shoulder (= 16 mm) for producing pressure under the tool shoulder (Figure 1). 2.3 Friction-stir welding A plan of experiments was prepared regarding the capabilities of the universal milling machine used (Prvomajska ALG 100E). Different combinations of tool rotations and welding speeds were tested at a constant tilt angle of 2°. The FSW tool rotated from 200 r/min to 1250 r/min, and the welding speeds changed from 71 mm/min to 450 mm/min. The factors of feed per revolution (FPR) in μm/r and revolution per feed (RPF) in r/mm were introduced to better distinguish between the different welding parameters. The FPR varied from 56 μm/r to 2250 μm/r. RPF, which represents the "frictional heat input" per weld length, was between 17.6 r/min and 0.44 r/mm. A backing plate underneath the workpiece enabled the creation of pressure under the tool shoulder by preventing the aluminium alloy from flowing away from the seam. The two workpieces were clamped in a vice. 2.4 Preparation of samples and testing From the FSW welds the miniature tensile test samples (Figure 2) were sectioned perpendicular to the D. KLOB^AR et al.: FRICTION-STIR WELDING OF ALUMINIUM ALLOY 5083 484 Materiali in tehnologije / Materials and technology 46 (2012) 5, 483–488 Table 1: Physical and mechanical properties of aluminium alloy 50833 Property /kg m–3 Rm/MPa Rp0.2/MPa E/GPa Tsol/°C Tliq/°C AA5083-O 2.660 275–300 125–149 71 580 640 Figure 1: a) FSW tool geometry and b) experimental FSW Slika 1: a) Geometrija orodja za FSW in b) potek eksperimenta FSW welding direction, and weld cross-sections for analyses of the microstructure and the macrostructure were prepared. Before sectioning the samples with a water jet, the workpiece surfaces were milled to remove the weld underfill and the toe flash. The uniaxial tensile tests were made using a compu- ter-controlled Zwick/Roell Z050 tensile testing machine. The measurements were made using Testexpert software. The strain was measured with extensometer fixed directly on the sample. The samples for analysis of the microstructure and microstructure were sectioned, grinded and polished. The samples for the macrostructure analysis were etched using Keller reagent (1125 mL HCl, 558 mL HNO3, 200 mL HF and 1 500 mL H2O) and the microstructure was analysed using an optical microscope. The samples for the microstructure analysis were anodized with Baker’s reagent. The microstructure was examined using an optical microscope under polarised light and with a digital camera for acquiring the pictures. The Vickers micro-hardness HV1 (load equal to 9.807 N) was measured across the welds. 3 RESULTS AND DISCUSSION 3.1 Visual assessment of the FSW welds Figure 3 shows a top view of the FSW welds. The end of the weld is indicated with a hole, which is a negative of the FSW tool pin. A visual assessment of the weld apices reveals smooth weld apices for FPR between 50 μm/r and 1000 μm/r, i.e., for a RPF between 20 r/min and 1 r/mm (Figure 3). For sample 7 (Figure 3a) the frictional heat input was the highest (RPF = 17.6 r/mm). For this sample the tool moved a little too much into the workpiece, due to the higher frictional heat input, which then softened the material. For samples 5 and 6 (Figures 3b, c) the RPF was 2.86 r/mm and 1.77 r/mm and the weld apices were smooth. When the tool speed was increased to a 0.71 r/mm and 0.44 r/mm (Figures 3d, e) the heat input become too small and the weld apices become rough with traces of material tearing. The research of the influence of the width of the joint gap on the weldability showed that gaps wider than 0.5 mm could not be successfully welded, due to an inability to move the tool into the material to overcome the lack of material. 3.2 Weld microstructure Macrographs of the selected FSW welds are shown in Figure 4. Figure 4a presents a macrograph of sample 7, which was welded with the highest frictional heat input D. KLOB^AR et al.: FRICTION-STIR WELDING OF ALUMINIUM ALLOY 5083 Materiali in tehnologije / Materials and technology 46 (2012) 5, 483–488 485 Figure 4: Macrostructure of FSW welds obtained at 200 r/min: a) sample 7 (FPR= 56 μm/r), b) sample 1 (FPR = 350 μm/r), c) sample 2 (FPR = 1400 μm/r) and c) sample 3 (FPR = 2250 μm/r) Slika 4: Makrostruktura FSW varov, izdelanih pri 200 r/min: a) vzorec 7 (FPR= 56 μm/r), b) vzorec 1 (FPR = 350 μm/r), c) vzorec 2 (FPR = 1400 μm/r) in c) vzorec 3 (FPR = 2250 μm/r) Figure 2: Drawing of miniature tensile test sample Slika 2: Risba miniaturnega vzorca za natezni test Figure 3: A top view of the FSW welds Slika 3: Pogled na temena FSW varov (RPF = 17.6 r/mm). The weld is without defects, except for possible under-fill, due to the higher heat input. A trace of the oxide line is present across the weld that was welded at 200 r/min and a 71 mm/min welding speed (FPR = 0.355 mm/r and RPF = 2.81 r/mm) (Figure 4b). The presence of Al2O3 on the surface of the touching planes in the weld joint before the welding is the reason for such a defect. This is why the oxide layer should be removed from the weld joint prior to welding. Figure 4c shows the FSW weld with a "worm hole" defect or "tunnelling" defect. This weld was produced at 200 rpm and a 280 mm/min welding speed (FPR = 1.4 mm/r and RPF = 0.71 r/mm). The "worm hole" defect appears if the welding is carried out with insufficient heat input or if the welding force in the axial direction is not large enough. The weld microstructures on the top of the weld, inside the weld and at weld root are shown in Figure 5. The grain size is very small at the top of the weld (Figure 5b), which was in the vicinity of the tool shoulder. Small-sized grains are obtained across the whole weld (Figure 5c). At the weld root the material is not stirred to the bottom of the workpiece (Figure 5d). The oxide surface of contacting the workpieces is clearly seen as a line of oxides. Such an oxide line/layer could represent the initiation site for cracking during loading during forming or exploitation. The analysis of the grain size showed that with a lower RPF, i.e., frictional heat input, the grain size decreases (Figure 6). For a higher frictional heat input the weld is heated well above the temperatures of recrystallization up to the temperatures where the grain growth takes place. When the RPF was 17.6 r/mm, the grain size of the weld and the base alloy were almost identical (Figures 5 and 6a). When the RPF was approximately 2.8 r/mm, the grain size becomes approximately half the size of the base alloy (Figures 6b, c). The reason for this could be smaller heat input and the lower temperature of the workpiece. When welding with a RPF of 0.44 r/mm, the heat input was so low so that the stirring, i.e., cold deformations had a major role in grain refinement. In this case the grains were very small (Figure 6c) and the weld became harder. 3.3 Hardness The Vickers hardness HV1 was measured across the weld in the middle of the weld (2 mm below the surface) over a total distance of 26 mm. The hardness is shown for the samples 7, 1, 0 and 3 (Figure 7). The centre of the weld is shown with the "dash-dot" line and the D. KLOB^AR et al.: FRICTION-STIR WELDING OF ALUMINIUM ALLOY 5083 486 Materiali in tehnologije / Materials and technology 46 (2012) 5, 483–488 Figure 7: Hardness HV1 across the weld, HAZ and base alloy Slika 7: Trdota HV1 preko vara, TVP in osnovnega materiala Figure 5: Microstructure (polarised light microscopy images) of FSW weld produced at 200 r/min FPR = 350 μm/r (sample 1): a) weld with HAZ and base alloy, b) weld apices and HAZ, c) weld and d) weld root Slika 5: Mikrostruktura (polarizirana svetlobna mikroskopija) FSW vara, izdelanega pri 200 r/min in FPR = 350 μm/r (vzorec 1): a) var z TVP in osnovno zlitino, b) teme vara in TVP, c) var in d) koren vara Figure 6: Microstructure (polarised light microscopy images) of FSW weld produced at: a) 1250 r/min and RPF = 17.6 r/mm (sample 7), b) 200 r/min and RPF = 2.82 r/mm (sample 1), c) 1250 r/min and RPF = 2.78 r/mm (sample 0) and d) 200 r/min and RPF = 0.44 r/mm (sample 3) Slika 6: Mikrostruktura (polarizirana svetlobna mikroskopija) FSW varov, izdelanih pri: a) 1250 r/min in RPF = 17,6 r/mm (vzorec 7), b) 200 r/min in RPF = 2,82 r/mm (vzorec 1), c) 1250 r/min in RPF = 2,78 r/mm (vzorec 0) in d) 200 r/min in RPF = 0,44 r/mm (vzorec 3) advancing side of the weld is on the right-hand side of the plot of Figure 7. When welding with a higher frictional heat input of 17.6 r/mm, the whole workpiece was heated above the temperature of recrystallization, where the grain growth occurs. The hardness was the lowest among all the compared samples (74 HV1 in the base metal and HAZ, and 82 HV1 in the weld). When welding with optimal welding parameters (sample 0 and 1), the hardness across the weld was slightly higher (84 HV1), similar to the base alloy where it was 80 HV1. A higher frictional heat input for the advan- cing side of the weld resulted in a lower hardness in the HAZ (75 HV1). When the frictional heat input was very low (0.44 r/mm) the weld hardness increased up to 105 HV1. Here, a deformational hardening was the dominating process due to a lower heat input. As a result of the higher frictional heat input on the advancing side of the weld, the hardness on the advancing side is generally lower than on the retreating side. 3.4 Tensile properties The tensile strength of the base alloy used for the experimental workpiece was not measured for the experimental workpiece, but taken from the literature data (Table 1). The yield strength of the aluminium alloy 5083 is between 125 MPa and 149 MPa and the ultimate tensile strength between 275 MPa and 300 MPa (Table 1). Since non-standard test specimens were used, the results could not easily be compared with the results from the literature. The ultimate tensile strength of the tensile test specimens was generally in the range of the base aluminium alloy (Figure 8). When welding with 2.82 r/mm (sample 1), the tensile strength was even higher, i.e., 320 MPa. When welding with almost the same frictional heat input (sample 0), a strain at a tensile test of 60 % was measured, indicating the good potential for formability of the weld. When the heat input was low, i.e., the RPF was 0.44 r/mm, a strain of 5% was achieved as a consequence of the deformation-hardened microstructure. 4 CONCLUSIONS Based on the analysed results the following can be concluded: Smooth weld apices could be obtained when welding with an FPR between 50 μm/r and 1000 μm/r, i.e., for a RPF between 20 r/min and 1 r/mm. When welding with the high frictional heat input (20 r/mm): a) a risk of over-plunging and excessive flash generation is present, b) the stirred material heats well above the recrystallization temperature and the grain growth occurs, c) the hardness of the weld, the HAZ and the closer base alloy drops below the initial base-alloy hardness and d) an approximately 15 % lower tensile strength compared to the base alloy is obtained. When welding with a medium heat input in the optimal range of welding parameters (3 r/mm): a) the weld hardness increases slightly compared to the base alloy, b) the grains refine to half the size of the base alloy, due to the deformational hardening combined with the recrystallization, c) the weld has a higher tensile strength, and d) for higher tool rotations, a strain of around 60 % was measured, indicating the good forming potential of the weld. When welding with a low heat input (RPF = 1 r/mm): a) the weld apices become rough and a tearing takes place, due to a too low frictional heat input, b) a "tunnel- ing" or "elongated cavity" defect is usually present, c) the weld hardness increases since the deformation hardening becomes the dominating process, and d) an approximately 15 % lower tensile strength compared to base alloy is obtained. A lower hardness was observed for the advancing side of the weld due to the slightly higher heat input compared to the retreating side. Acknowledgement The authors would like to thank M. Hr`enjak and N. Breskvar for the preparation of the samples and the microscopy, and A. Skumavc for reviewing the paper and editing. The research was sponsored by the Slovenian Research Agency (ARRS) under the project L2-4183 entitled "Friction stir welding and processing of aluminium alloys". 5 REFERENCES 1 J. Tu{ek, IEEE Trans. Plasma Sci., 28 (2000) 5, 1688 2 P. Podr`aj, I. Polajnar, J. Diaci, Z. Kari`, Science and Technology of Welding and Joining, 13 (2008) 3, 215 3 http://aluminium. matter.org.uk/aluselect/ 4 A. Smolej, B. Skaza, B. Markoli, D. Klob~ar, V. Dragojevi}, E. Sla- ~ek, Materials Science Forum, 706–709 (2012) 706, 395 5 R. Mishra, Z. Ma, Materials Science and Engineering: R: Reports, 50 (2005) 1–2, 1 6 Z. Ma, R. Mishra, Scripta Materialia, 53 (2005) 1, 75 7 Z. Ma, S. Sharma, R. Mishra, Scripta Materialia, 54 (2006) 9, 1623 D. KLOB^AR et al.: FRICTION-STIR WELDING OF ALUMINIUM ALLOY 5083 Materiali in tehnologije / Materials and technology 46 (2012) 5, 483–488 487 Figure 8: Results of tensile tests Slika 8: Rezultati nateznega testa 8 U. Trdan, J. Grum, M. R. Hill, Materials Science Forum, 681 (2011) 480 9 U. Trdan, J. L. Ocaña, J. Grum, Strojni{ki vestnik – Journal of Mechanical Engineering, 57 (2011) 05, 385 10 S. @agar, J. Grum, Strojni{ki vestnik – Journal of Mechanical Engineering, 57 (2011) 04, 334 11 T. Debroy, H. K. D. H. Bhadeshia, Science and Technology of Welding & Joining, 15 (2010) 4, 266 12 D. G. H. M. James, G. R. Bradley, International Journal of Fatigue, 25 (2003) 12 13 D. Hattingh, C. Blignault, T. Vanniekerk, M. James, Journal of Mate- rials Processing Technology, 203 (2008) 1–3, 45 14 H. Lombard, D. Hattingh, A. Steuwer, M. James, Engineering Frac- ture Mechanics, 75 (2008) 3–4, 341 15 S. J. L. Min-Su Han, J. C. Park, S. C. Ko, Y. B. Woo, S. J. Kim, Trans. Nonferrous Met. Soc. China, 19 (2009), 17 16 T. Hirata, T. Oguri, H. Hagino, T. Tanaka, S. Chung, Y. Takigawa, K. Higashi, Materials Science and Engineering: A, 456 (2007) 1–2, 344 17 Y. S. Sato, F. Yamashita, Y. Sugiura, S. H. C. Park, H. Kokawa, Scripta Materialia, 50 (2004) 3, 365 18 J. Tu{ek, Z. Kampu{, M. Suban, J. Mater. Process. Technol., 119 (2001) 1/3, 180 19 F. Zucchi, G. Trabanelli, V. Grassi, Materials and Corrosion, 52 (2001), 853 20 Steel Selector Metal Ravne. 2012; Available from: http://www. metalravne.com/selector/steels/vcmo140.html D. KLOB^AR et al.: FRICTION-STIR WELDING OF ALUMINIUM ALLOY 5083 488 Materiali in tehnologije / Materials and technology 46 (2012) 5, 483–488 B. SEN^I^, V. LESKOV[EK: INFLUENCE OF SEGREGATIONS ON THE FRACTURE TOUGHNESS ... INFLUENCE OF SEGREGATIONS ON THE FRACTURE TOUGHNESS KIc OF HIGH-STRENGTH SPRING STEEL VPLIV IZCEJ NA LOMNO @ILAVOST KIc VISOKOTRDNOSTNEGA VZMETNEGA JEKLA Bojan Sen~i~1,2, Vojteh Leskov{ek2,3 1[TORE STEEL, d. o. o., @elezarska cesta 3, 3220 [tore, Slovenia 2Jo`ef Stefan International Postgraduate School, Jamova 39, 1000 Ljubljana, Slovenia 3Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia bojan.sencic@store-steel.si Prejem rokopisa – received: 2012-03-19; sprejem za objavo – accepted for publication: 2012-05-07 The results of this investigation showed that using the proposed method it was possible to draw, for the normally used range of working hardnesses, a tempering diagram (Rockwell-C hardness – fracture toughness KIc – tempering temperature) for the vacuum-heat-treated high-strength spring-steel grade 51CrV4. Based on measurements of the mechanical properties we have also created a classic tempering diagram, i.e., Tensile strength Rm – Yield stress Rp0.2 – Elongation A5/% – Necking Z/% – Tempering temperature, and a tempering diagram, i.e., Hardness HRc – Impact toughness Charpy-V – Tempering temperature. According to these tempering diagrams we can conclude that the investigated spring steel 51CrV4 is suitable for the production of high-strength springs when using the proper heat treatment. Fractographic and metallographic analyses of the KIc test specimens showed the presence of segregations in the steel. Therefore, we focused on examining the impact of segregations on the fracture toughness KIc. We found that the widths of the positive segregation bands and the matrix bands between the samples vary considerably. We have also discovered that the number and the width of the segregation bands influence significantly the fracture toughness KIc due to the presence of bainite in the matrix bands. Keywords: fracture toughness, segregations, high-strength spring steel, vacuum heat treatment, tempering diagrams, microstructure Pri opravljenih preiskavah smo `eleli ugotoviti, ali lahko standardizirano preizku{anje lomne `ilavosti (ASTM E399-90) nadomestimo z nestandardnim postopkom preizku{anja lomne `ilavosti s cilindri~nim nateznim preizku{ancem z zarezo po obodu in utrujenostno razpoko v dnu zareze. Inovativen na~in preiskav je pokazal, da lahko s predlagano metodo konstruiramo diagram popu{~anja (trdota Rockwell-C – lomna `ilavost KIc – temperatura popu{~anja) za vakuumsko toplotno obdelano visokotrdnostno vzmetno jeklo. Na osnovi meritev mehanskih lastnosti smo izdelali poleg klasi~nega diagrama popu{~anja: natezna trdnost Rm – meja plasti~nosti Rp0.2 – raztezek A5/% – kontrakcija Z/% – temperatura popu{~anja, {e diagram popu{~anja trdota HRc – udarna `ilavost Charpy-V – temperatura popu{~anja. Iz izdelanih diagramov popu{~anja lahko ugotovimo, da je preiskovano vzmetno jeklo 51CrV4 primerno za izdelavo visokotrdnostnih vzmeti ob ustrezno izvedeni toplotni obdelavi. Med fraktografsko in metalografsko analizo KIc preizku{ancev smo odkrili prisotnost izcej v jeklu. Zato smo se posvetili ugotavljanju vpliva izcejanja na lomno `ilavost KIc. Ugotovili smo, da se {irina trakov pozitivnih izcej in trakov matriksa pri razli~nih vzorcih precej razlikuje. Prav tako smo odkrili, da {tevilo in {irina izcej zaradi prisotnosti bainita v matriksu pomembno vplivata na lomno `ilavost KIc. Klju~ne besede: lomna `ilavost, izcejanje, visokotrdnostno vzmetno jeklo, vakuumska toplotna obdelava, diagrami popu{~anja, mikrostruktura 1 INTRODUCTION A producer of spring steel must provide a technical description of the steel, which includes the chemical composition and the basic mechanical, physical and technological properties of the steel. Among the tech- nological properties, information concerning the heat treatment is very important for the manufacturer of the springs. Charpy-V notch (CVN) impact-test values are used in toughness specifications for spring steels, even though the fracturing energy is not directly related to the spring design. It is surprising that there is no demand for the fracture toughness KIc value (the plain-strain stress- intensity factor at the onset of unstable crack growth) in the delivery conditions for spring-steel producers. To the spring designer the KIc values are more useful than the CVN values, because the design calculations for the springs from high-strength steels should also take into account the strength and the toughness of the materials in order to prevent rapid and brittle fracture. A spring’s durability is limited by the plastic deformation, the fatigue and the fracturing. Therefore, the use of spring steel with the following properties is recommended: high ductility and toughness at operating temperatures from –40 °C to +50 °C and good harden- ability, which provides the required mechanical proper- ties. As a consequence of the manufacturing route, steels with a similar chemical composition may behave differently due to the variety of mechanical properties. Assuming that the chemical composition and initial microstructure of the steel correspond to those pre- scribed for the steel grade 51CrV4 (DIN 17221 and DIN 17222), the mechanical properties for a specific application depend mainly on the appropriately selected heat-treatment parameters. Materiali in tehnologije / Materials and technology 46 (2012) 5, 489–496 489 UDK 669.14.018.252:539.42:621.785.72 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)489(2012) An investigation was conducted to determine whether standardized fracture-toughness testing (ASTM E399- 90), which is difficult to perform reliably for hard and low-ductility materials, could be replaced with a non-standard testing method using circumferentially notched and fatigue-precracked tensile specimens. The aim of this innovative investigation approach was to enable us to draw, for the normally used range of working hardness, combined tempering diagrams (Rock- well-C hardness – fracture toughness KIc – tempering temperature) for the vacuum-heat-treated high-strength spring-steel grade 51CrV4. In addition, we also wanted to create, for the spring steel 51CrV4, at a selected austenitizing temperature, a classic tempering diagram, i.e., tensile strength Rm – yield stress Rp0.2 – elongation A5 – necking Z – tempering temperature, and a tempering diagram, i.e., hardness HRc – impact toughness Charpy-V – tempering tempe- rature. Using these tempering diagrams we wanted to confirm the suitability of the investigated steel for the production of high-strength springs with the required tensile strength between 1500 MPa and 1800 MPa. In accordance with the plan of experiments the heat treatment was carried out on the basis of trial prelimi- nary research and modelling results by measurements of the mechanical properties, analysis of the fractured surfaces and the examination of the microstructure of the KIc samples. Then, based on the mechanical properties the tempering diagrams for the spring steel 51CrV4, at a selected austenitizing temperature, were created. Fractographic and metallographic analyses of the KIc test specimens showed the presence of segregations in the steel. Therefore, we also focused on examining the impact of segregations on the fracture toughness KIc. Using optical and electron (SEM + EDS) microscopy we determined the number and the width of the positive segregations bands and of the matrix bands just under the fractured surface of the KIc test specimens. We studied the influence of the number and the width of the segregation bands on the fracture toughness KIc. 2 EXPERIMENTAL 2.1 Hardness and fracture-toughness tests Circumferentially notched and fatigue-precracked tensile-test specimens1 with the dimensions indicated in Figure 1 were used for the investigation. The Rock- well-C hardness (HRc) was measured on individual groups of KIc test specimens using a Wilson 4JR hard- ness machine. The advantage of the KIc test specimens used here over standardized CT specimens (ASTM E399-90) is in the radial symmetry, which makes them particularly suitable for studying the influence of the microstructure of metallic materials on the fracture toughness. The advantage of these specimens is related to the heat transfer, which ensures a uniform microstructure. Due to the high notch sensitivity of hard and brittle metallic materials, such as continuous-cast spring-steel grade 51CrV4, it is very difficult – sometimes even almost impossible – to create a fatigue crack in the test specimen. However, with the KIc test specimens the fatigue crack can be created with rotating-bending loading before the final heat treatment2. The second advantage of such test specimens is that plain-strain conditions can be achieved using specimens with smaller dimensions than those of conventional CT test speci- mens3. For the linear elastic behaviour up to fracture of such specimens4 the following equation is applied: K P D D dIc = − +⎛⎝ ⎜ ⎞ ⎠ ⎟ 3 2 127 172/ . . (1) where P is the load at failure, D is the outside diameter, and d is the notched-section diameter of the test specimen. Equation (1) is valid as long as the condition 0.5 < d/D < 0.8 is fulfilled. The measurements of the fracture toughness were performed at room temperature using an Instron 1255 tensile-test machine. A cross-head speed of 1.0 mm/min was used for the standard tensile tests on specimens with a nominal test length of 100 mm. In the tests two specially prepared cardan fixed jaws, ensuring the axiality of the tensile load, were used. During the tests the tensile-load/displacement relationship until failure was recorded. In all cases this relationship was linear, and the validity of equation (1) for the tests was confirmed. 2.2 Impact test In order to obtain the tempering diagram, i.e., hardness HRc – Charpy-V – tempering temperature, we B. SEN^I^, V. LESKOV[EK: INFLUENCE OF SEGREGATIONS ON THE FRACTURE TOUGHNESS ... 490 Materiali in tehnologije / Materials and technology 46 (2012) 5, 489–496 Figure 1: Circumferentially notched and fatigue-precracked KIc test specimen. Dimensions in mm. Slika 1: Cilindri~ni natezni preizku{anec za merjenje lomne `ilavosti z zarezo po obodu in utrujenostno razpoko v dnu zareze. Dimenzije so v milimetrih. measured the impact toughness using the Charpy impact test, known also as the Charpy V-notch test (ISO 148). The measurement with an instrumented Charpy hammer allows us to estimate the total impact work, the work needed for crack initiation and the work necessary for crack propagation. The hardness HRc was measured on an Instron B 2000 device according to the standard SIST EN ISO 6508-1. 2.3 Tensile test The standard tensile test (SIST EN ISO 6892-1) was applied to measure the tensile strength Rm, the yield stress Rp0.2, the elongation A5 and the necking Z. 2.4 Material, sampling and vacuum heat treatment Samples from continuous-cast, high-strength, spring- steel, grade 51CrV4, delivered as hot-rolled and soft- annealed bars of dimensions 100 mm × 25 mm × 6000 mm were used. The circumferentially notched and fatigue-precracked KIc test specimens were cut from the middle of the bar in the rolling direction with a fatigue crack at the notch root in the transverse direction. They were heat treated in a horizontal vacuum furnace with uniform, high-pressure gas-quenching using nitrogen (N2) at a pressure of 5 bar. After the first preheat (650 °C) the specimens were heated at a rate of 10 °C/min to the austenitizing tem- perature of 870 °C, soaked for 10 min, gas quenched to 80 °C, and then single tempered for one hour at different temperatures between 200 °C and 575 °C. At each tem- pering temperature 16 test specimens for the tensile test (Rm specimen), for the determination of the fracture toughness (KIc specimen) and of the Charpy-V toughness (CVN specimen), as well as two metallographic samples  19 mm × 9 mm were heat treated. 2.5 Measurement of the number and the width of the segregations A fractured surface of a KIc specimen was observed on the SEM, Figure 2. The dotted line shows the orientation of the MnS sulphides. An example of the MnS inclusion inside the crack is shown in the circle. B. SEN^I^, V. LESKOV[EK: INFLUENCE OF SEGREGATIONS ON THE FRACTURE TOUGHNESS ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 489–496 491 Figure 3: Combined image of segregations just under the fractured surface Slika 3: Sestavljena slika izcej tik pod prelomno povr{ino Figure 2: SEM image of fractured surface of KIc specimen Slika 2: SEM posnetek prelomne povr{ine preizku{anca za merjenje lomne `ilavosti The presence of MnS in the crack was confirmed by EDS. Inclusions of MnS are present in positive segre- gations that are oriented in the rolling direction. To determine the number of segregations the KIc specimens were cut perpendicularly to the orientation of the MnS sulphides. The number of segregations just under the fractured surface was counted using an optical microscope, as shown in Figure 3, on a sample etched in 4 % picral to make the segregations visible. The dark bands as positive segregations were confirmed by the presence of sulphides and by measuring the microhardness. The assessment of the number and the width of the segregations is shown in Figure 4. 3 RESULTS AND DISCUSSION In a classic tempering diagram for an austenitizing temperature of 870 °C the average measured values of the mechanical properties (tensile strength Rm/MPa, yield strength Rp0.2/MPa, elongation A5/% and necking Z/%) are shown as a function of the tempering temperature in the range between 300 °C and 700 °C in Figure 5. Given that we had for each selected tempering temperature a statistically relevant number of Rm speci- mens, for each group we performed a statistical analysis. As can be seen from the diagram, the minimum disper- sion of results is within ±2 across the whole range of selected tempering temperatures in tensile strength and elongation, while it is higher only for the necking. The tempering diagram indicates that the requirement from the standard DIN EN 10089:2003-4 for the elonga- tion (minimum 8 %) and necking (minimum 30 %) can be achieved after the tempering of the investigated steel in the temperature range between 300 °C and 700 °C, while the required tensile strength (1350–1650 MPa) and yield-strength (minimum 1200 MPa) can be achieved with a tempering temperature below 530 °C. The required tensile strength for high-strength spring steel (1500–1800 MPa) can be achieved if the tempering temperature is below 475 °C. The tempering diagram, hardness HRc – Charpy-V – tempering temperature, for the selected austenitizing temperature of 870 °C and the selected tempering tem- perature in the range of 200 °C to 625 °C for the high-strength steel 51CrV4 is shown in Figure 6. The diagram shows that the curves for the hardness and impact toughness Charpy-V over the entire range of tempering temperatures are similar to the curves of the hardness and the fracture toughness KIc in the tempering diagram shown in Figure 7. Similar to the fracture toughness, the impact toughness Charpy-V also increases to a temperature of 525 °C, then it decreases to 550 °C, and then the toughness increases again. This trend can be attributed to the kinetics of precipitation B. SEN^I^, V. LESKOV[EK: INFLUENCE OF SEGREGATIONS ON THE FRACTURE TOUGHNESS ... 492 Materiali in tehnologije / Materials and technology 46 (2012) 5, 489–496 Figure 5: Classic tempering diagram for continuous-cast, hot-rolled, flat, spring steel 51CrV4, for an austenitizing temperature of 870 °C Slika 5: Klasi~ni diagram popu{~anja za kontinuirno lito, vro~e valjano vzmetno jeklo 51CrV4, temperatura avstenitizacije 870 °CFigure 4: Microstructure of test specimen D99: positive segregations – tempered martensite (dark bands), 4% picral Slika 4: Mikrostruktura preizku{anca D99: pozitivne izceje – po- pu{~eni martenzit (temni pasovi), 4 % pikral Figure 6: Effect of tempering temperature on the hardness HRc and the impact toughness Charpy-V of a continuous-cast, hot-rolled, flat, spring steel 51CrV4 Slika 6: Vpliv temperature popu{~anja na trdoto HRc in udarno `ilavost Charpy-V za kontinuirno lito, vro~e valjano vzmetno jeklo 51CrV4 during the tempering. In the diagram, the dispersion of results within ±2 in the quenched and tempered condition using a temperature of 475 °C is of the same magnitude. The dispersion of results is increased above the tempering temperature of 500 °C. The results of the measurements of the Charpy-V toughness in the range of tempering temperature 525–575 °C show that the tough- ness is even reduced, which confirms the observations made when measuring the fracture toughness, so we assume that this is an area of irreversible temper embrittlement5,6. In the case of the CVN-specimens, which were quenched and tempered at the same temperature, the reason for the scattering of the results is the hetero- geneity of the investigated steel and also, but less of a factor, the geometry and surface roughness of the notch. The requirement from the standard DIN EN 10089: 2003-04 for the impact energy (minimum 8 %) can be achieved if the tempering temperature is above 200 °C. High-strength spring steels are very notch sensitive, for this occasion, it is also important to measure fracture toughness KIc, which can be described as the ability of a material to resist, under tensile loading, the progress of existing cracks. We determined the fracture toughness KIc by the use of circumferentially notched and fatigue- precracked KIc test specimens which were linear elastic loaded to fracture. The tempering diagram, hardness HRc – fracture toughness KIc – tempering temperature, for the selected austenitizing temperature of 870 °C and the selected tempering temperature in the range of 200 °C to 625 °C for the steel 51CrV4 is shown in Figure 7. Given that we had for each selected tempering temperature a statistically relevant number of KIc specimens, for each group we performed a statistical analysis. As can be seen from the diagram, the minimum dispersion of results within ±2 s is in the quenched state and in KIc specimens that were quenched and tempered at 200 °C. At higher tempering temperatures between 300 °C and 525 °C, the scattering of the results slightly increased. This trend can be attributed to the kinetics of extracting precipitates during tempering. The reason for the dispersion of the results within each group of KIc specimens that were quenched and tempered at the same temperature is the heterogeneity of the investigated steel. The highest hardness of 58 HRc is achieved in the as-quenched condition after vacuum quenching from the austenitizing temperature of 870 °C. In examining the evolution of the properties by an increase of the tempering temperatures, it is observed that the minimum scatter of results within ± 2 = 3.2 MPa m1/2 is found in the as-quenched state and after single tempering at 200 °C. At higher tempering temperatures between 300 °C and 575 °C, the scatter of the KIc results slightly increased up to ±2 = 9.4 MPa m1/2, while the scatter of the Rockwell-C hardness is up to ±2 = 1.2 HRc across the whole range of used tempering temperatures. Within each group of KIc specimens that were quenched and tempered at the same temperature, this can be attributed to the kinetics of the carbides’ precipitation during tempering at selected temperatures as well as to the B. SEN^I^, V. LESKOV[EK: INFLUENCE OF SEGREGATIONS ON THE FRACTURE TOUGHNESS ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 489–496 493 Figure 8: The microstructure of the KIc test specimen in the as-quen- ched condition (transverse direction); a) specimen with the lowest fracture toughness (KIc= 16.2 MPa m1/2, 58.1 HRc), b) specimen with the highest fracture toughness (KIc= 22.3 MPa m1/2, 58.1 HRc) Slika 8: Mikrostruktura KIc preizku{anca v kaljenem stanju (pre~na smer); a) z najmanj{o lomno `ilavostjo (KIc= 16,2 MPa m1/2, 58,1 HRc), b) z najve~jo lomno `ilavostjo (KIc= 22,3 MPa m1/2, 58,1 HRc) Figure 7: Effect of tempering temperature on the hardness HRc and the fracture toughness KIc of the continuous-cast, hot-rolled, flat, spring steel 51CrV4 Slika 7: Vpliv temperature popu{~anja na trdoto HRc in lomno `ilavost KIc za kontinuirno lito, vro~e valjano vzmetno jeklo 51CrV4 heterogeneity of the investigated steel. Since the austenitizing temperature is the same, it is clear that the fracture toughness, KIc, is a very selective mechanical property with regard to the tempering temperature. It should be noted that the KIc test specimens were taken from the middle of the bar, and therefore the micro- structures of the KIc test specimens with lowest and highest fracture toughnesses are comparable. In the as-quenched condition, the microstructure consists of untempered martensite and lower bainite, Figure 8. Strong positive (bright) and negative (dark) segregations are visible due to the lower etching intensity of the untempered martensite. The fractured surface of the KIc test specimens was examined in SEM at magnification of 12-times, Figure 9. The presence of inclusions (sulphides in cracks) in positive segregations is evident. The density of positive segregations is higher on the fractured surface of the KIc test specimen with the highest fracture toughness in comparison to the fractured surface of the KIc test specimen with the lowest fracture toughness with the B. SEN^I^, V. LESKOV[EK: INFLUENCE OF SEGREGATIONS ON THE FRACTURE TOUGHNESS ... 494 Materiali in tehnologije / Materials and technology 46 (2012) 5, 489–496 Figure 9: Fractured surfaces of KIc test specimen, tempered at 475 °C with an equal hardness; a, b) specimen with the lowest fracture toughness (KIc = 67.0 MPa m1/2, 44.8 HRc), c, d) specimen with the highest fracture toughness (KIc = 76.1 MPa m1/2 in 44.8 HRc) Slika 9: Prelomne povr{ine KIc preizku{ancev popu{~enih pri 475 °C z enako trdoto; a, b) z najmanj{o lomno `ilavostjo (KIc = 67,0 MPa m1/2, 44,8 HRc), c, d) z najve~jo lomno `ilavostjo (KIc = 76,1 MPa m1/2 in 44,8 HRc) Figure 10: Typical microstructure of KIc test specimens (KIc = 75.7 MPa m1/2, 43.8 HRc) in the longitudinal direction. Sulphide inclusions are located in the positive segregations. Slika 10: Zna~ilna mikrostruktura KIc preizku{anca (KIc = 75,7 MPa m1/2, 43,8 HRc) v vzdol`ni smeri. Sulfidi se nahajajo v pozitivnih izcejah. Figure 11: SEM microstructure of the KIc test specimens with the highest fracture toughness: a) positive segregations and b) negative segregations (matrix), longitudinal direction Slika 11: SEM-mikrostrukture KIc preizku{anca z najve~jo lomno `ilavostjo: a) pozitivna izceja, b) negativna izceja, vzdol`na smer same hardness. Also, the distribution is more even and the size of the segregations is smaller. The microstructure of the same KIc specimens just below the fractured surface was examined in an optical microscope, Figure 10. The fractured KIc specimen was cut perpendicularly to the direction of segregations determined by the position of the sulphide inclusions. Then the number of segregations in the transverse direction was counted and for the specimen with the lowest fracture toughness (KIc = 67.0 MPa m1/2, 44.8 HRc) 98 segregations and specimen with the highest fracture toughness (KIc = 76.1 MPa m1/2, 44.8 HRc) 147 segregations were found. We also measured the widths of the segregation bands. The specimen with the lowest fracture toughness has an average width of the segregation bands equal to 29 μm and that with the highest fracture toughness has an average width of the segregation bands equal to 33 μm. The microstructure in the quenched and tempered condition consists of tempered martensite and bainite (≈20 %, volume fraction). In the microstructure, non- metallic inclusions of the sulphide type located in positive segregations and oriented in the rolling direction are observed. The microstructure of the KIc test specimen was examined in the SEM and it was found that the microstructure in the positive segregations consisted of tempered martensite and in the negative (matrix) of tempered martensite with islands of bainite, Figures 11 and 12. From a comparison of the microstructures of the KIc test specimens with the lowest and highest fracture toughnesses from the same heat and with the same hardness, it can be concluded that the density, the size and the distribution of segregations have a significant influence on the fracture toughness. The increase of fracture toughness can probably be ascribed to the pre- sence of bainite in the matrix of negative segregations. To find out the influence of segregation on the fracture toughness the number of segregations and their width just under the fractured surface were assessed, Figures 13 and 14. It can be seen that there is a trend for the fracture toughness KIc to increase with a larger number of segregations. It is also evident that the average width of the segregations affects the fracture toughness KIc. The broad bands of the positive segregations increase the fracture toughness KIc, while it is lowered by narrow bands of the matrix. This can also be associated with the number of segregations. If there are more segregations, the average width of the matrix bands decreases. 4 CONCLUSIONS The investigated high-strength, spring-steel grade 51CrV4 was successfully quenched in a horizontal vacuum furnace with uniform high-pressure gas-quench- B. SEN^I^, V. LESKOV[EK: INFLUENCE OF SEGREGATIONS ON THE FRACTURE TOUGHNESS ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 489–496 495 Figure 12: Bainite grain from Figure 11 Slika 12: Kristalno zrno bainita s slike 11 Figure 14: Influence of the average width of the matrix bands and the segregation bands on the fracture toughness KIc Slika 14: Vpliv povpre~ne {irine pasov negativnih in pozitivnih izcej na lomno `ilavost KIc Figure 13: Influence of the number of segregations on the fracture toughness KIc Slika 13: Vpliv {tevila izcej na lomno `ilavost KIc ing using nitrogen (N2) at a pressure of 5 bar. The Rockwell-C hardness was (58.4 ± 0.8) HRc in the quenched state, which is high enough to obtain the required hardnesses from 35 HRc to 50 HRc after a single tempering. The obtained microstructure consisting of tempered martensite and bainite (20 % volume fraction) meets the requirements of TB1402 (Scania Standard STD512090 and STD4153), thus we can conclude that the investigated high-strength spring steel 51CrV4 with a thickness up to 20 mm is suitable for heat treatment in a vacuum furnace with uniform, high-pressure gas- quenching using nitrogen (N2) at a pressure of 5 bar or higher. Our investigation showed that standardized fracture- toughness testing (ASTM E399-90) could be replaced with a non-standard testing method using circum- ferentially notched and fatigue-precracked tensile specimens (KIc test specimen). The results of this innovative approach to the investigation have shown that using the proposed method it was possible to draw, for the normally used range of working hardness, combined tempering diagrams (Rockwell-C hardness – Fracture toughness KIc – Tempering temperature) for vacuum- heat-treated, high-strength, spring-steel grade 51CrV4. According to the tempering diagrams the tensile strength Rm – Yield stress Rp0.2 – Elongation A5 – Neck- ing Z – Tempering temperature and the tempering diagram, Hardness HRc – Impact toughness Charpy-V – Tempering temperature, we can conclude that the investigated spring steel 51CrV4 is suitable for the production of high-strength springs when the proper heat treatment is performed. The presence of segregations was confirmed by fractographic and metallographic analyses of the KIc test specimens. We found that the width of the positive segregation bands and matrix bands between the samples varied considerably. We have also discovered that the number and the width of the segregation bands influenced significantly the fracture toughness KIc due to the presence of bainite in the matrix bands. Our investi- gation shows that there is a trend of fracture toughness KIc increase with a larger number of segregations. These findings will be checked in future investiga- tions. Acknowledgement [tore Steel, d. o. o., @elezarska 3, SI-3220 [tore, Slovenia, is thanked for the supply of test material as well for the financial support. Thanks also to Prof. dr. Franc Vodopivec for helpful discussions. 5 REFERENCES 1 B. Sen~i~, V. Leskov{ek, Mater. Tehnol., 45 (2011) 1, 67–73 2 A. Sandberg, M. Nzotta, High performance matrix tool steels pro- duced via a clean steel concept, 7th Tooling Conference, Torino, 2006 3 V. Leskov{ek, Optimization of the vacuum heat treatment of high- speed steels, University of Zagreb, 1999 (Ph. D. thesis) 4 H. F. Bueckner, ASTM STP 381, 1965, 82 5 A. Shekhter, S. Kim, D. G. Carr, A. B. L. Croker, S. P. Ringer, International Journal of Pressure Vessels and Piping, 79 (2002), 611–615 6 B. Ule, F. Vodopivec, M. Pristavec, F. Gre{ovnik, @elezarski zbornik, 24 (1990), 35–40 B. SEN^I^, V. LESKOV[EK: INFLUENCE OF SEGREGATIONS ON THE FRACTURE TOUGHNESS ... 496 Materiali in tehnologije / Materials and technology 46 (2012) 5, 489–496 K. SOMASUNDARA VINOTH et al.: MECHANICAL AND TRIBOLOGICAL CHARACTERISTICS OF STIR-CAST Al-Si10Mg ... MECHANICAL AND TRIBOLOGICAL CHARACTERISTICS OF STIR-CAST Al-Si10Mg AND SELF-LUBRICATING Al-Si10Mg/MoS2 COMPOSITES MEHANSKE IN TRIBOLO[KE LASTNOSTI Z ME[ANJEM ULITIH KOMPOZITOV Al-Si10Mg IN SAMOMAZALNIH KOMPOZITOV Al-Si10Mg/MoS2 Kannappan Somasundara Vinoth1, Ramanathan Subramanian2, Somasundaram Dharmalingam3, Balu Anandavel2 1Department of Production Engineering, PSG College of Technology, 641 004 Coimbatore, India 2Department of Metallurgical Engineering, PSG College of Technology, 641 004 Coimbatore, India 3Department of Mechanical Engineering, PSG Polytechnic College, 641 004 Coimbatore, India vinothks@yahoo.com Prejem rokopisa – received: 2012-03-20; sprejem za objavo – accepted for publication: 2012-05-24 The mechanical and tribological characteristics of aluminium-molybdenum-disulphide self-lubricating composites have been investigated and compared to the Al-Si10Mg alloy. Al-Si10Mg/4MoS2 display the finest microstructures due to a higher fraction of MoS2 added. The densities of Al-Si10Mg/2MoS2 and Al-Si10Mg/4MoS2 were marginally higher than in the case of the aluminium alloy by 1 % and 2 % mass fractions, respectively. The ultimate tensile strength decreases considerably due to the additions of 2 % and 4 % MoS2 by 15 % and 22 %, respectively, compared to the Al-Si10Mg alloy. It was seen that while the Al-Si10Mg alloy shows a predominantly ductile fracture (fibrous regions), the composite specimens (with an MoS2 addition) show an increase in the mixed mode (ductile and brittle regions). Al-Si10Mg/2MoS2 and Al-Si10Mg/4MoS2 show an enormous decrease in the wear rate by 55 % and 65 %, respectively, compared with the Al-Si10Mg alloy. The decrease in the wear occurs due to the presence of an MoS2 layer, which forms a film on the wear surface. Keywords: aluminium-molybdenum-disulphide self-lubricating composites Preiskovane so bile mehanske in tribolo{ke zna~ilnosti aluminij-molibden disulfidnih samomazalnih kompozitov in primerjane z zlitino Al-Si10Mg. Al-Si10Mg/4MoS2 ima najdrobnej{o mikrostrukturo zaradi ve~jega dele`a dodanega MoS2. Gostota Al-Si10Mg/2MoS2 in Al-Si10Mg/4MoS2 je bila navidezno ve~ja za masni dele` 1 % oziroma 2 %. Kon~na natezna trdnost se je v primerjavi z zlitino Al-Si10Mg ob~utno zmanj{ala za 15 % oziroma 22 % pri dodatku 2 % oziroma 4 % masnega dele`a MoS2. Izkazalo se je, da pri zlitini Al-Si10Mg prevladuje `ilav prelom (vlaknata podro~ja), pri kompozitnih vzorcih (z dodatkom MoS2) pa me{an prelom (duktilna in krhka podro~ja). Al-Si10Mg/2MoS2 in Al-Si10Mg/4MoS2 izkazujeta ob~utno pove~anje odpornosti proti obrabi, in sicer 55 % oziroma 65 % v primerjavi z zlitino Al-Si10Mg. Zmanj{anje obrabe je zaradi sloja MoS2, ki tvori tanko plast na obrabni povr{ini. Klju~ne besede: aluminij-molibden disulfidni samomazalni kompoziti 1 INTRODUCTION Aluminium-silicon alloys and composites are being used in automotive applications like pistons, brake rotors and engine-block cylinder liners1,2. Tribological beha- viour is an important aspect in the use of aluminium metal-matrix composites in automotive applications. The wear behaviour of Al-Si alloys can be further enhanced by adding ceramic particles. Abrasive particles like silicon carbide, alumina, and diamond are added to improve the tribological behaviour by increasing the hardness of a composite3–5. Nevertheless, lubricating particles like graphite and MoS2 have also been added to improve the tribological behaviour of different materials by providing a solid lubricating layer6,7. The additions of these particles considerably affect the mechanical behaviour of the composites. There are various methods of producing composites like blending and consolidation, vapour deposition and consolidation, stir casting, infiltration process, spray deposition and consolidation, as well as in-situ reacting process8. Of all these processes, stir casting is the simplest and the most economical method. Stir-cast self-lubricating composites have been successfully developed by adding graphite particles9. It has also been suggested that these composite materials have the capacity to achieve low friction and wear of the contact surfaces without any external supply of lubrication during the sliding. However, graphite films fail in lower loads and shorter lifetimes compared with MoS210. Self-lubricating Al-MoS2 composites have been prepared by using the powder-metallurgy route11. However, neither the preparation of MoS2- based composites by stir casting nor the characterisation of Al-Si10Mg/MoS2 composites have been reported in literature. In this investigation, two self-lubricating composites of molybdenum disulphide, namely, Al-Si10Mg/2MoS2 and Al-Si10Mg/4MoS2 have been produced with the stir-casting route. The changes in the mechanical and Materiali in tehnologije / Materials and technology 46 (2012) 5, 497–501 497 UDK 66.017:620.168:539.92 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)497(2012) tribological properties caused by the addition of MoS2 are studied and compared with the Al-Si10Mg alloy. 2 MATERIALS AND METHODS 2.1 Preparation of the composite The Al-Si10Mg aluminium alloy (Table 1) with a density of 2640 kg/m3 was used in this investigation as the matrix material. The Al-Si10Mg alloy has excellent resistance to corrosion in both normal atmospheric and marine environments collectively exhibiting high strength and hardness. Table 1: Chemical composition of the aluminium alloy used (mass fractions, w/%) Tabela 1: Kemijska sestava uporabljene aluminijeve zlitine (mas. dele`i, w/%) Mg Si Fe Mn Others* Al 0.2 to 0.6 10.0 to 13.0 0.6 max 0.3 to 0.7 1.5 max balance * (Cu, Ni, Zn, Pb, Sn and Ti) The Al-Si10Mg alloy was charged into an electrical resistance-heated furnace modified for this investigation. The melting of the Al-Si10Mg alloy was carried out under argon atmosphere at 1073 K. Molybdenum- disulphide (MoS2) solid lubricant with an average particle size of 1.5 μm and a density of 4600 kg/mm3 (Figure 1) was used as the reinforcement in this investigation. The MoS2 particulates were incorporated into the molten metal and stirred continuously for ten minutes. The molten mixture was solidified in a cast-iron die in the form of a cylindrical pin with a diameter of 14 mm and a length of 70 mm. 2.2 Testing of the materials The density of composites was determined using a top-loading electronic balance (Mettler Toledo make) according to the Archimedean principle. The micro- structure of the composite specimens was identified using a Carl Zeiss Goettingen optical microscope. The specimens were metallographically polished to obtain an average roughness value of 0.8 μm. The tensile testing was carried out using a Hounsefield tensometer. The ultimate tensile strength of the specimens was calculated from the load at which a fracture occurred. The morpho- logy of worn surfaces of the composite specimens was examined by using a JEOL T100 Scanning Electron Microscope (SEM). The hardness was measured by using a Zwick hardness tester at a load of 100 g. The dry-sliding wear behaviour of the composites was studied using a pin-on-disc apparatus. The disc material was made of the EN-32 steel with a hardness of 65 HRC. The difference in weights before and after the test was taken as weight loss. The wear rate was calculated on the basis of the difference in the weights of a specimen using the following formula: Wear rate [ ]WR W D = 981.  mm3/km (1) where W/kg = mass loss, /(kg/mm3) = density of the material, D = sliding distance K. SOMASUNDARA VINOTH et al.: MECHANICAL AND TRIBOLOGICAL CHARACTERISTICS OF STIR-CAST Al-Si10Mg ... 498 Materiali in tehnologije / Materials and technology 46 (2012) 5, 497–501 Figure 1: Microstructures of the materials: a) Al-Si10Mg, b) Al-Si10 Mg/2MoS2, c) Al-Si10Mg/4MoS2 Slika 1: Mikrostruktura materiala: a) Al-Si10Mg, b) Al-Si10Mg/ 2MoS2, c) Al-Si10Mg/4MoS2 3 RESULTS AND DISCUSSION 3.1 Microstructures Optical micrographs of the Al-Si10Mg alloy and of the composites (Figures 1a to c) show as-cast (dendritic) structures consisting of silicon particles in a eutectic matrix. The microstructures of the composites (Al-Si10Mg/2MoS2 and Al-Si10Mg/4MoS2) are signifi- cantly finer, affected probably by the heterogeneous nucleation caused by MoS2 particles. Al-Si10Mg/4MoS2 exhibits the finest microstructure due to the higher fraction of MoS2 added. Figures 1b and 1c show that MoS2 particles were uniformly distributed in the matrix. 3.2 Mechanical properties The mechanical properties of the composites (den- sity, hardness, and tensile strength), given in Table 2, show the average properties of various test specimens at different positions. The density of MoS2 is higher than that of the aluminium alloy and hence an increase in the MoS2 content will raise the density of the composite. The densities of Al-Si10Mg/2MoS2 and Al-Si10Mg/4MoS2 were marginally higher than the density of the aluminium alloy by 1 % and 2 %, respectively. A similar increase in the density of the composites was achieved by adding SiC12 and Al20313 by various authors. The ultimate tensile strength [UTS] of Al-Si10Mg was approximately 218 MPa. It was reported in previous researches that an addition of alumina to AA6061 and AA7005 causes an increase in the tensile strength14. Similar results were reported for SiCp/aluminium-alloy composites15 and aluminium-alumina, aluminium-illite and aluminium-silicon carbide particle composites5. In contrast to this, the studies on an addition of alumina to the 2024 Al alloy have shown a decrease in UTS16. Similar results were obtained for an addition of graphite to aluminium17. The tensile strength decreases consi- derably due to the additions of 2 % and 4 % by mass MoS2 by 15 % and 22 %, respectively. The observed decrease in UTS may be due to various mechanisms like the particle pull-out and crack propagation, which are initiated by the presence of MoS2. The elongation of the composites decreases slightly less than in the case of the Al-Si10Mg alloy indicating that an addition of MoS2 lowers the ductility of a composite. A similar result was observed in the SiC reinforcement of the 2124, 7075 alloys and monolithic aluminium18,19. 3.3 Fracture surface Figures 2a to c show the SEM fractographs of Al-Si10Mg, Al-Si10Mg/2MoS2 and Al-Si10Mg/4MoS2, K. SOMASUNDARA VINOTH et al.: MECHANICAL AND TRIBOLOGICAL CHARACTERISTICS OF STIR-CAST Al-Si10Mg ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 497–501 499 Figure 2: Fracture analysis of the materials: a) Al-Si10Mg, b) Al-Si10Mg/2MoS2, c) Al-Si10Mg/4MoS2 Slika 2: Analiza prelomov materiala: a) Al-Si10Mg, b) Al-Si10Mg/ 2MoS2, c) Al-Si10Mg/4MoS2 Table 2: Mechanical properties of the Al-Si10Mg alloy and the composites Tabela 2: Mehanske lastnosti Al-Si10Mg zlitine in kompozitov Material Densitykg/m3 UTS MPa Hardness HV Elongation % Decrease in UTS, % Increase in hardness, % Al-Si10Mg 2640 218.45 102 1.66 - - Al-Si10Mg/2MoS2 2670 185.31 145 1.22 15 42 Al-Si10Mg/4MoS2 2697 170.28 148 1.10 22 45 respectively. From the fractographs of the tensile-test specimens (Figure 2) it can be seen that, in the aluminium-matrix alloy, the fracture was primarily transgranular with a microscopic void formation; later the progressive growth and the final coalescence around the reinforcement particles can be observed. It can also be seen that while the Al-Si10Mg alloy shows a predominantly ductile fracture (fibrous regions), the composite specimens show an increase in the mixed mode [ductile and brittle regions]. In addition, the composite samples also show the features like particle pullout, crack growth, and propagation typical of a brittle fracture. 3.4 Wear behaviour Dry-sliding wear tests were conducted according to ASTM G-99 using a pin on a disc apparatus under an applied load of 50 N for a sliding speed of 5 m/s. The wear rate plotted against the sliding distance is shown in Figure 3. The Al-Si10Mg alloy experiences the maximum wear rate. The wear mechanism was studied using a SEM micrograph of the worn surface of the Al-Si10Mg alloy (Figure 4a), revealing severe delamination that is an indication of an adhesive wear. Al-Si10Mg/2MoS2 and Al-Si10Mg/4MoS2 show an enormous decrease in the wear rate by 55 % and 65 %, respectively, compared with the Al-Si10Mg alloy. The decrease in the wear is due to the presence of the MoS2 layer, which forms a film on the wear surface. This in evident in the SEM micrograph of Al-Si10Mg/2MoS2 (Figure 4b) where the MoS2 particles form a film in certain regions, partially reducing the ploughing and delamination. 4 CONCLUSION In this research work, Al-Si10Mg/MoS2 composites were fabricated using the stir-casting technique and the mechanical and tribological characteristics were studied. The following important observations can be noted: 1. UTS, elongation percentage and hardness decrease with an addition of MoS2 particles to Al-Si10Mg. However, the densities of the composites are higher than the density of the Al-Si10Mg alloy. 2. A uniform distribution of MoS2 is observed on the optical micrographs. 3. The improved wear resistance of Al-Si10Mg/MoS2 composites is better than the wear resistance of the Al-Si10Mg alloy. 5 REFERENCES 1 P. Rohatgi, Cast Metal Matrix Composites: Past, Present and Future, Transactions of the American Foundry Society, 109 (2001), 1–25 2 M. M. Haque, A. Sharif, Study on wear properties of aluminium- silicon piston alloy, Journal of Materials Processing Technology, 118 (2001), 69 K. SOMASUNDARA VINOTH et al.: MECHANICAL AND TRIBOLOGICAL CHARACTERISTICS OF STIR-CAST Al-Si10Mg ... 500 Materiali in tehnologije / Materials and technology 46 (2012) 5, 497–501 Figure 3: Wear behaviour of the Al-Si10Mg alloy and the composites Slika 3: Obraba Al-Si10Mg-zlitine in kompozitov Figure 4: Wear-surface SEM micrographs at a load of 50 N and a speed of 5 m/s of: a) Al-Si10Mg, b) Al-Si10Mg/2MoS2 Slika 4: SEM-posnetek obrabljene povr{ine pri obremenitvi 50 N in hitrosti 5 m/s za: a) zlitine Al-Si10Mg, b) Al-Si10Mg/2MoS2 3 A. T. Alpas, J. Zhang, Effect of SiC particulate reinforcement on the dry sliding wear of aluminium-silicon alloys (A356), Wear, 155 (1992), 83 4 P. W. Ruch, O. Beffort, S. Kleiner, L. Weber, P. J. Uggowitzer, Selective interfacial bonding in Al(Si)–diamond composites and its effect on thermal conductivity, Composites Science and Technology, 66 (2006), 2677–2685 5 M. K. Surappa, P. K. Rohatgi, Preparation and properties of cast aluminium-ceramic particle composites, Journal of Materials Science, 16 (1981), 983–993 6 B. N. P. Bai, E. S. Dwarakadasa, S. K. Biswas, Scanning electron microscopy studies of wear in LM13 and LM13-graphite particulate composite, Wear, 76 (1982), 211 7 B. [u{tar{i~, L. Kosec, M. Kosec, B. Podgornik, S. Dolin{ek, The influence of MoS2 additions on the densification of water-atomized HSS powders, Journal of Materials Processing Technology, 173 (2006) 3, 291–300 8 M. K. Surappa, Aluminium matrix composites: Challenges and opportunities, Sadhana, 28 (2003), 319–334 9 P. L. Menezes, P. K. Rohatgi, M. R. Lovell, Self-Lubricating Behavior of Graphite Reinforced Metal Matrix Composites, in: M. Nosonovsky, B. Bhushan (Eds.), Green Tribology, Springer, Berlin Heidelberg 2012, 445–480 10 A. J. Haltner, C. S. Oliver, Frictional Properties of Some Solid Lubricant Films under High Load, J. Chem. Eng. Data, 6 (1961), 128–130 11 H. Kato, M. Takama, Y. Iwai, K. Washida, Y. Sasaki, Wear and mechanical properties of sintered copper–tin composites containing graphite or molybdenum disulfide, Wear, 255 (2003), 573–578 12 Y. Sahin, Preparation and some properties of SiC particle reinforced aluminium alloy composites, Materials & Design, 24 (2003), 671 13 M. Kok, Production and mechanical properties of Al2O3 particle- reinforced 2024 aluminium alloy composites, Journal of Materials Processing Technology, 161 (2005), 381–387 14 L. Ceschini, G. Minak, A. Morri, Tensile and fatigue properties of the AA6061/20 vol. % Al2O3p and AA7005/10 vol. % Al2O3p com- posites, Composites Science and Technology, 66 (2006), 333–342 15 Ü. Cöcen, K. Önel, Ductility and strength of extruded SiCp/alumi- nium-alloy composites, Composites Science and Technology, 62 (2002), 275–282 16 A. N. Abdel-Azim, Y. Shash, S. F. Mostafa, A. Younan, Casting of 2024-Al alloy reinforced with Al2O3 particles, Journal of Materials Processing Technology, 55 (1995), 199–205 17 C. B. Lin, R. J. Chang, W. P. Weng, A study on process and tribological behavior of Al alloy/Gr (p) composite, Wear, 217 (1998), 167–174 18 T. J. A. Doel, P. Bowen, Tensile properties of particulate-reinforced metal matrix composites, Composites Part A: Applied Science and Manufacturing, 27 (1996), 655–665 19 J. N. Hall, J. Wayne Jones, A. K. Sachdev, Particle size, volume fraction and matrix strength effects on fatigue behavior and particle fracture in 2124 aluminum-SiCp composites, Materials Science and Engineering A, 183 (1994), 69–80 K. SOMASUNDARA VINOTH et al.: MECHANICAL AND TRIBOLOGICAL CHARACTERISTICS OF STIR-CAST Al-Si10Mg ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 497–501 501 M. BENISA et al.: COMPUTER-AIDED MODELING OF THE RUBBER-PAD FORMING PROCESS COMPUTER-AIDED MODELING OF THE RUBBER-PAD FORMING PROCESS RA^UNALNI[KO MODELIRANJE PREOBLIKOVALNEGA PROCESA Z VMESNIKOM IZ GUME Muamar Benisa, Bojan Babic, Aleksandar Grbovic, Zoran Stefanovic University of Belgrade, Faculty of Mechanical Engineering, Kraljice Marije 16, 11120 Belgrade, Serbia bbabic@mas.bg.ac.rs Prejem rokopisa – received: 2012-03-22; sprejem za objavo – accepted for publication: 2012-06-01 The conventional way to develop press-formed metallic components requires a burdensome trial-and-error process for setting-up the technology, whose success depends largely on the operator’s skill and experience. The finite element (FE) simulations of a sheet-metal-forming process help a manufacturing engineer to design a forming process by shifting the costly press-shop try-outs to the computer-aided design environment. Numerical simulations of a manufacturing process, such as rubber-pad forming, have been introduced in order to avoid the trial-and-error procedure and shorten the development phases when tight times-to-market are demanded. The main aim of the investigation presented in this paper was to develop a numerical model that would be able to successfully simulate a rubber-pad forming process. The finite-element method was used for blank- and rubber-behavior predictions during the process. The study was concerned with a simulation and investigation of significant parameters (such as forming force and stress, and strain distribution in a blank) associated with the rubber-pad forming process and the capabilities of this process regarding the manufacturing of aircraft wing ribs. The simulation and investigation carried out identified the stress and strain distribution in a blank as well as the forming force. Experimental analyses of a rib with a lightening hole showed a good correlation between FE simulations and experimental results. Keywords: rubber-pad forming, sheet-metal bending, finite-element simulation, aircraft manufacturing Obi~ajna pot pri razvoju stiskanih kovinskih komponent zahteva za postavitev tehnologije zamuden postopek z analiziranjem napak pri preizkusih, pri ~emer je uspe{nost odvisna od izvajal~eve spretnosti in izku{enj. Simulacija postopka preoblikovanja plo~evine s stiskanjem po metodi kon~nih elementov (FEM) pomaga in`enirjem pri postavljanju preoblikovalnega postopka z nadome{~anjem dragih preizkusov v ra~unalni{kem okolju. Da bi se izognili postopkom preizku{anja z analizo napak in ko je odlo~ujo~ kratek rok za tr`enje, so bile vpeljane numeri~ne simulacije preoblikovalnega procesa, kot je preoblikovanje z vmesnikom iz gume. Glavni namen predstavljene preiskave je bil razvoj numeri~nega modela, ki bi bil sposoben uspe{ne simulacije preoblikovalnega procesa z vmesnikom iz gume. Za napovedovanje dogajanj v stiskancu in v vmesniku iz gume je bila uporabljena metoda kon~nih elementov. [tudija je obsegala simulacije in preiskave pomembnih parametrov (kot so preoblikovalna sila, napetosti ter razporeditev deformacije v preoblikovancu), povezanih s preoblikovanjem z vmesnikom iz gume in z zmogljivostmi tega procesa pri izdelavi reber letalskega krila. Izvr{ena simulacija in preiskava je odkrila napetosti in razporeditev deformacije v preoblikovancu, kot tudi sile pri preoblikovanju. Analiza reber z luknjo za zmanj{anje mase je pokazala dobro ujemanje med FE-simulacijo in eksperimentalnimi rezultati. Klju~ne besede: preoblikovanje z vmesnikom iz gume, krivljenje plo~evine, simulacija kon~nih elementov, izdelovanje letal 1 INTRODUCTION Stamping is a metal-forming process, with which the sheet metal is punched using a press tool that is mounted on a machine or a stamping press forming the sheet into the desired shape. The conventional stamping process is performed through a punch, which, together with a blank holder, forces the sheet metal to slide into a die and comply with the shape of the die itself. Computers allow us to obtain and process data to improve and accelerate an analysis of the information required to optimize the processes of metal forming and to minimize the production costs.1 Rubber-pad forming is a metalworking process where sheet metal is pressed between a die and a rubber block. In general, an elastic upper die, usually made of rubber, is connected to a hydraulic press. A rigid lower die, often called a form block, provides the mold for the sheet metal to be formed. Because the upper (male) die can be used with different lower (female) dies, the process is relatively cheap and flexible. However, rubber pads exert less pressure in the same circumstances than the non-elastic parts, which may lead to a lower accuracy of the forming process. The form-block height is usually less than 100 mm.2 In the rubber-pad forming process an aluminum blank is placed between a die and a rubber pad (flexible punch), which is held in a container to enclose the flexible punch (Figure 1a). At this stage, the flexible punch (rubber-pad) is fixed on the arm of a pressing machine and the punch is on a machine table. As the rubber-pad moves down the rubber deforms elastically and offers a counter pressure. Due to this pressure the rubber-pad and the blank flow into the cavity of the die. This process can be divided into three steps: the first, self-forming of the rubber, the second, when the blank moves to the bottom of the die and produces the outer bending, and the third when the rubber pad pushes the blank into the cavity of the die. In the aircraft industry most of the sheet parts, such as ribs, frames, doors and windows, are fabricated using Materiali in tehnologije / Materials and technology 46 (2012) 5, 503–510 503 UDK 621.7.01:519.61/.64 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 46(5)503(2012) the rubber-pad forming processes (flexible tools). The advantages of using the rubber-pad forming process instead of the conventional metallic tools are: (i) the same flexible pad can be used to form several different work-piece shapes, because a rubber pad has the ability to return to its original shape; (ii) the tool costs are lower than the costs for the conventional forming processes; (iii) the thinning of the work metal, which occurs during the conventional deep drawing, is reduced considerably; (iv) the set-up time can be reduced considerably in this process, because no die clearance or alignment checks need to be made; (v) lubrication is not necessary and good surface finish can be achieved, because no tool marks are created. However, the rubber-pad forming processes have several disadvantages, such as: (i) the lifetime of a flexible pad is limited (this depends on the severity of forming combined with the pressure level); (ii) a lack of a sufficient forming pressure results in the parts with a lower sharpness or with wrinkles, which requires the reworking of the parts to their correct shape and dimensions; (iii) a low production rate, so that they are suitable mostly for small series (typical of the aircraft industry).2–5 Several studies have been carried out to analyze the rubber-pad forming. Browne and Battikha6 presented an experimental study of the rubber-pad forming process to investigate the capability of the process and optimize the process parameters. Sala3 optimized the rubber-pad forming of an aluminum-alloy fuselage frame belonging to an AerMacchi MB-339 trainer aircraft using his own finite-element code. Several effects have been investi- gated depending on stamping velocity, geometry, heat treatment of the sheet metal and rubber-pad parameters. Dirikolu and Akdemir7 investigated the influence of rubber hardness and blank-material type on stress distribution using a 3D finite-element-simulation study of a flexible forming process. The investigation showed that the variation of pad thickness does not cause a big change in the forming stress in a blank. Madoliat and Narimani8 presented sheet forming by using a rubber pad and also investigated, experimentally and numerically, the design for the tooling set. Thiruvarudchelvan,9 presented in this overview, highlighted the role of urethanes that are considered to be the best materials for flexible tools because of their good oil and solvent resistance, good wear resistance, high thermal stability and load-bearing capacity. Ramezani and Ahmed5 carried out a numerical simulation of a rubber-pad forming process, along with an experimental validation, using a die of a flexible punch. They studied some forming parameters such as rubber type, stamping velocity, etc., and found that silicone rubber has a shorter lifetime than polyurethane and natural rubber. As a result, it cannot be used to form blanks with sharp edges. No significant change in the blank thinning was discovered for 5 different stamping velocities. Lee et al.10 have investigated the deformation characteristic using the rubber-pad bending of a structural aluminum tube. A 3D finite-element analysis was used to examine the effect of process parameters on deformation characte- ristics of an extruded aluminum tube, and the influence of the formable radius of a tube curvature on bending resistance. The relation between the bent profile of a material and the roller stroke was defined. Fabrizio and Loreddana11 studied flexible forming of thin sheets from aluminum alloys using different geometries and materials for the flexible die. They have investigated the forming force during a forming process for different dies. These investigations showed that a numerical model can help us better understand the forming procedure, and the correlations with experimental results were good. M.W. Fu and H. Li12 have presented 3D-FE simulations and investigated the deformation behavior of the flexible-die-forming process. The comparison between the conventional deep drawing and a viscoplastic carrying medium based on flexible-die forming was conducted in terms of wall-thickness reduction, hydro- M. BENISA et al.: COMPUTER-AIDED MODELING OF THE RUBBER-PAD FORMING PROCESS 504 Materiali in tehnologije / Materials and technology 46 (2012) 5, 503–510 Figure 1: a) Schematic representation of a rubber-pad forming process, b) Experimental tool set-up Slika 1: a) Shematski prikaz preoblikovalnega procesa z vmesnikom iz gume, b) eksperimentalno orodje static pressure, principle-stress distribution and damage factor. The concave and convex rubber-pad forming processes were investigated by Liu, et al.,13 using FE simulations and experimental methods. The investi- gations of the forming load, thickness variation of the formed plate and variations in the channel-width- to-rib-width ratio were also performed. A fabrication of a metallic bipolar plate for a proton membrane in fuel cells is presented in14. The FE analyses were used to describe the rubber-pad forming process and to investi- gate the main parameters (such as rubber hardness and dimensions of the rigid die). It was found that the smaller the internal radius, the harder it is to fill the cavity of a rigid die. The authors examined whether the blank filled the cavity of the rigid die by using a 3D-laser-scanning measurement system. In this paper, numerical simulations of rubber-pad forming processes are used to analyze the blank and the rubber behaviors during a production of supporting ribs, as well as to analyze different punch geometries. A non-linear FE analysis was conducted to predict stress and strain distributions, and forming forces during the rubber-pad forming process. The main goal was to develop a computer model that would be able to simulate the process and, therefore, to develop the right design for a tooling set. 2 FINITE-ELEMENT MODELING AND EXPERIMENTAL VERIFICATION Numerical simulations of the rubber-pad forming processes are complicated mainly because of a large deformation of a rubber pad. As a consequence, a mesh distortion may occur in a simulation, which can lead to inaccurate and incomplete results. This is why FE analyses must be carried out carefully and with an understanding of the physical phenomena of the rubber-pad forming process. The commercial finite-element software Ansys® was used to make an FE simulation in this study. In order to reduce the processing time and improve the precision of the calculations, 2D FE models were created for three different sheet-metal elements (straight rib, stringer and rib with a lightening hole) and analyses were carried out for each model. Figure 2 illustrates these three geometrical models. The models in FE analyses included three elements only: a rigid die, a blank and a rubber pad (flexible punch). In order to simplify the numerical model, the container of the rubber pad was not modeled. To eliminate the influence of the container, the frictionless support constraints were applied on the opposite sides of the rubber, while a displacement constraint was applied on the upper edge of the 2D rubber model (Figure 3). The die was modeled as a rigid body because the stress and strain of the die were not analyzed and the die material (steel) is much less deformable than the material of the blank (aluminum). So, the material properties attached to the die were not important, and the mesh was not generated either. This eliminated unnecessary calculations causing a decrease in both the run time and the errors in the numerical solution. Because the blank undergoes a large plastic-strain deformation during the forming process, the stress-strain test data up to a failure was required to define the blank material in the simulation (Figure 4). The blank was considered as a multilinear isotropic hardening material. In the FE simulations of the blank behavior, the von Mises yield criterion coupled with an isotropic work- hardening assumption15 was used. The rubber pad undergoes a nonlinear hyper-elastic deformation. The behavior of the nonlinear hyper-elastic and incompressible rubber-like material is usually described with the Mooney-Rivlin model that uses a strain-energy function W. The derivative of W, with M. BENISA et al.: COMPUTER-AIDED MODELING OF THE RUBBER-PAD FORMING PROCESS Materiali in tehnologije / Materials and technology 46 (2012) 5, 503–510 505 Figure 3: Constraints used in the FE simulation Slika 3: Omejitve, uporabljene pri FE-simulaciji Figure 2: Geometrical models used in the investigation: a) straight rib, b) stringer and c) rib with a lightening hole Slika 2: Geometrijski modeli, uporabljeni pri preizku{anju: a) ravno rebro, b) rebrasto rebro in c) rebro z luknjo za zmanj{anje mase respect to the strain component, determines the corresponding stress component:  ij ij W = ∂ ∂ (1) W C I I k I= − + − + − + + ∑ km k m k m 1 n ( ) ( ) ( )1 2 3 23 3 1 2 1 (2) where I1, I2 and I3 (I3 = 1) are the strain invariants, k is the bulk modulus and Ckm is the constant of the Mooney-Riviin material model for the incompressible material. Usually, two Mooney-Rivilin parameters, C10 and C01, are used to describe the hyper-elastic rubber deformation.7 In the FE models, the polyurethane rubber with the Shore A hardness of 70 (HD70) was used for the rubber pad. The values of C10 and C01 were 0.736 MPa and 0.184 MPa, respectively.5,7,13,14 An aluminum plate with a thickness of 0.6 mm was used as the blank. The aluminum properties were determined via the stress-stain curve obtained from the tensile tests,7,16 as shown in Figure 5. For this alloy, the elastic module (E) is 71G Pa and the Poisson’s ratio ( ) is 0.334. During the rubber-pad forming process, the materials exhibit large deformations and rotations. There is a friction contact at the blank interfaces, too. At the same time, geometric nonlinearities arise from a nonlinear strain-displacement relationship, as well as the nonlinearities associated with the material properties. According to that, the geometric nonlinearity option was activated in the nonlinear solution procedure. The friction behavior between the two different pairs of contact (rubber pad–blank and blank–die) was assumed to follow the Coulomb’s model.5,7 The friction coefficients for the former and latter contact pairs were considered to be 0.2 and 0.1, respectively.5,7,13 Table 1 shows the specifications of the contact region. The interface contacts of the blank–rubber pad, blank–die and die–rubber were modeled as deformable and the software used solves these tasks on the basis of the contact-target-surface approach with an adjustable impenetrability constraint that assures contact compati- bility. The CONTA 175 (a node-to-surface contact) finite element was used on the blank’s surfaces at the interface between the blank and the die and on the rubber surface at the interface between the rubber and the die. The CONTA171 (a surface-to-surface contact) was used on the surface of the rubber pad at the interface between the blank and the rubber pad. The other surfaces at each interface were modeled with the TARGE169 element. It can be summarized that, in all the interface contacts, the upper surfaces of the die and the blank were considered as a target, while the lower surface of the blank and the upper surface of the rubber pad were considered as a contact.17 As mentioned above, the container was not modeled, so – in order to fix the rubber pad correctly – frictionless supports had to be applied on the side edges of the rubber. A remote displacement was applied on the lower edge of the die. A displacement was applied on the top edge of the rubber in order to simulate a forming load on the blank (Figure 3). All deformable materials were modeled with a Plane-183 finite element (a 2-D element with 8 or 6 nodes). Plane 183 has quadratic displace- ment, plasticity, hyper-elasticity, creep, stress stiffening, large deflection and large strain-simulation capabilities. The number of nodes and elements used for the blank and the rubber pad are presented in Table 2. Table 1: Interface contacts Tabela 1: Vmesni stiki Parts in the contact Contact type Die & Blank Frictional contact (0.1) node to surface Blank & Rubber Frictional contact (0.2) surface to surface Die & Rubber Frictional contact (0.1) node to surface Table 2: Numbers of the nodes and elements for three models Tabela 2: [tevilo vozlov in elementov za tri modele Modes Blank Rubber pad Rigid die Nodes Element Nodes Element Nodes Element Straight rib 2161 1680 4253 3958 205 204 Stringer 768 157 7254 2299 304 152 Rib with a lightening hole 863 615 5490 5181 208 207 In order to validate the FE simulations results, the rubber-pad forming experiments were carried out for the stamping of an aluminum blank. An experimental set-up (shown in Figure 1b) was used, together with the assembly of a die set shown in Figure 1a. The die and the rubber-pad container were made of steel, while polyurethane rubber with a Shore A hardness of 70 (HD70) was used as a rubber pad. A hydraulic press machine (produced by REXROTH), with the maximum M. BENISA et al.: COMPUTER-AIDED MODELING OF THE RUBBER-PAD FORMING PROCESS 506 Materiali in tehnologije / Materials and technology 46 (2012) 5, 503–510 Figure 4: Experimental tensile-stress-strain curve for the aluminum blank sheet Slika 4: Eksperimentalna krivulja napetost – raztezek za aluminijevo plo~evino capacity of 160 t was used in the rubber-pad forming process. The process begins with the die placed on the base of the hydraulic press machine. The aluminum blank is then introduced between the die and the flexible punch. After this the flexible punch moves down to stamp the blank. One of the formed parts fabricated during the rubber-pad forming process and the die used during the fabrication are shown in Figure 5. 3 RESULTS AND DISCUSSION In this study, as mentioned above, the forming force was presented as a displacement applied on the upper edge of the rubber pad. Figure 6 illustrates the step-by-step forming process using the rubber pad. It is clear that the process can be divided into three stages (or steps). The first stage is a self-deformation of the flexible die (the rubber pad); the second stage includes a blank deformation (under the pressure of the rubber pad when it reaches the bottom of the rigid die); and, finally, during the third stage the blank fills the die cavities until they are completely filled. The convergences of the forming forces for each model, obtained through the FE simulations, are shown in Figure 7. This figure shows that the highest value of a forming force is present in the rib with a lightening hole (6735 N), while the lowest value is achieved in the straight rib (867 N). It can be seen that the magnitude of a forming force increases as the geometry of a rib becomes more complex, i.e., as more bending regions have to be obtained (Table 3). The FE simulation of the forming process for the stringer and the rib with a lightening hole goes through three stages/steps (corresponding to the forming process), while the straight-rib forming can be performed in the first two steps (because there is no cavity to fill). During the first step – the self-forming of the rubber – the rubber deforms elastically and offers a counter pressure, so the forming load is very small (Figure 6). The time needed for this step is short (between 0.2 s and 0.35 s in the simulation – Figure 7). After 0.2 second (the straight rib) and 0.35 second (the other models), the second step starts and the forming load increases slightly to produce the outer bending (Figure 5). During the last step, after approximately 0.65 second, the blank starts to fill the cavity of the rigid die and the forming force increases sharply (Figures 6 and 7). According to the results of the FE simulations, it is obvious that the highest value of the forming force will be obtained in the M. BENISA et al.: COMPUTER-AIDED MODELING OF THE RUBBER-PAD FORMING PROCESS Materiali in tehnologije / Materials and technology 46 (2012) 5, 503–510 507 Figure 6: Forming steps during the rubber-pad forming including the first, the second and the third steps: a) straight rib, b) stringer and c) rib with a lightening hole Slika 6: Stopnje med preoblikovanjem z gumijastim vmesnikom, vklju~no s prvo, drugo in tretjo stopnjo: a) ravno rebro, b) rebrasto rebro, c) rebro z luknjo za zmanj{anje mase Figure 5: Rib with a lightening hole fabricated by the rubber-pad forming process and the die Slika 5: Rebro z luknjo za zmanj{anje mase, izdelano z vmesnikom iz gume in orodje case of the most complex sheet-metal geometry (the rib with a lightening hole), where several bends with different radii must be produced. This is in correlation with the empirical data3,7,13,14, which means that the used FE models have been well defined. Along with the calculation of the forming force, stress and strain analyses were performed. As was expected, the stress and strain concentrations in a blank accumulate in the last two stages. Figure 8 shows the equivalent stresses (in MPa, left column) and plastic strains (in mm/mm, right column) in the blanks at the end of the forming processes for all FE models. The summarization of the maximum and minimum stress and strain values, presented in Figure 8, is given in Table 3. Table 3 shows that the maximum equivalent stress and plastic strain appear in the rib with a lightening hole (241.45 MPa and 0.206 mm/mm, respectively), while the minimum values of the equivalent stress and plastic strain are in the straight rib (224.74 MPa and 0.115 mm/mm, respectively). As can be seen in Table 3, the stress and plastic strain increase with an increased complexity of the rib geometry, as well as with an increased number of bend radii. The reason for that, according to Sala3, is that the blank is exposed not only to the tensile and tangential stresses, but also to the stress arising from the bending pressure imposed by the tool. Consequently, the thinning phenomenon occurs homo- geneously and, finally, the necking appears. The necking can induce a crack, which is not unusual in this forming process5. The crack starts when the blank undergoes stretching forces and when the ultimate stress is reached during the second or third stage of the forming process. According to Sala3 and Takuda15, the maximum plastic strain that can be considered for the forming of this type of aluminum alloy is approximately 0.186 mm/mm. This value of plastic strain was used as a reference for the crack-appearance predictions in the FE simulations presented in this paper. The value of the plastic strain obtained in the forming simulation of the rib with a lightening hole (0.206 mm/mm) indicated the possibility of a crack appearance in the outer radius of the lightening hole, while the plastic-strain values for the other two models were less than 0.186 mm/mm. M. BENISA et al.: COMPUTER-AIDED MODELING OF THE RUBBER-PAD FORMING PROCESS 508 Materiali in tehnologije / Materials and technology 46 (2012) 5, 503–510 Figure 7: Forming force in three models Slika 7: Preoblikovalna sila pri treh modelih Figure 8: Equivalent stress (left column) and plastic strain (right column) in the straight rib, the stringer and the rib with a lightening hole (from the top) Slika 8: Ekvivalentna napetost (leva kolona) in plasti~na deformacija (desna kolona) v ravnem rebru, rebrastem rebru in rebru z luknjo za zmanj{anje mase (od zgoraj navzdol) The experiments with the rubber pad showed that the FE predictions were good. Figure 9 shows the crack that appeared in the region around the rib hole during the third stage of the forming process, as predicted by the FE simulation. This was the proof of the quality of the developed finite-element model, as well as of the criterion proposed for this alloy (with the maximum plastic strain not exceeding 0.186 mm/mm). On the basis of these findings, more FE models of a rib with a lightening hole (with different values of the fillet radii) were developed and analyzed in order to find the connections between the values of a fillet radius and a plastic strain. These simulations showed that the values of stress and strain strongly depended on the rib geometry (i.e., the values of the fillet radii), but more investigations need to be performed in order to clearly define these dependencies. 4 CONCLUSION A finite-element simulation of the rubber-pad forming process could be a very useful tool for understanding and improving the forming operations because it provides important data for determining the forming parameters and the operation time. The developed FE models and the method proposed in this paper have proved to be sufficiently effective in pre- dicting the final shape of the component and the regions of a possible crack appearance. The FE simulations showed that the maximum stresses and strains in all the cases were at the flanges and the corners. The minimum stress and plastic strain were achieved in the straight rib (the rib with the simplest geometry), while the maximum stress and plastic strain were found in the rib with a lightening hole (the most complex geometry). These results have been validated with the experiments, as well as with the fracture criterion used for the crack predictions. The FE simulations proved that simpler tools would reduce the lead times and enable a rapid production of small parts without a possibility of a crack appearance during the forming. On the other hand, the geometry of more complicated, but necessary, tools must be defined very carefully, with a determination of the fillet radii that will minimize the chance of a fracture. FEM can help us with this determination, too, while additional potential applications – such as 3D model simulations and tool optimization – are also possible. However, it must be noticed that the optimization procedure of the press-forming processes – owing to the presence of the hardly reproducible phenomena like friction and lubrication – should never be limited to simple numerical simulations because the above pheno- mena can contribute a lot towards saving the costs and reducing the time-to-market, currently held up by empirical trial-and-error processes. Sheet-metal-forming-simulation results, today, are reliable and accurate enough so that even the try-out tools and the time-consuming try-out processes may be eliminated, or at least reduced significantly. 5 REFERENCES 1 A. Shramko, I. Mamuzic, V. Danchenko, The application of the program QFORM 2D in the stamping of wheels for railway vehicles, Mater. Tehnol., 43 (2009) 4, 207–211 2 ASM Handbook Vol. 14B Metal Working: Sheet Forming, 2006 3 G. Sala, A numerical and experimental approach to optimize sheet stamping technologies: part II – aluminium alloys rubber-forming, Material and Design, 22 (2001) 4, 299–315 M. BENISA et al.: COMPUTER-AIDED MODELING OF THE RUBBER-PAD FORMING PROCESS Materiali in tehnologije / Materials and technology 46 (2012) 5, 503–510 509 Figure 9: Rib with a crack around the lightening hole Slika 9: Rebro z razpoko okrog luknje za zmanj{anje mase Table 3: Summarization of the engineering requirements for the rubber-pad sheet-metal forming process Tabela 3: Povzetek in`enirskih zahtev pri preoblikovalnem procesu z vmesnikom iz gume MODEL Model 1 (straight rib) Equivalent stress (MPa) Equivalent plastic strain(mm/mm) Reaction force N Model 1 Straight rib Maximum 224.74(occurs on the blank) 0.115 (occurs on the blank) 866.43 Minimum 1.2e-4 0.0 Model 2 Stringer Maximum 221.29(occurs on the blank) 0.1268 (occurs on the blank) 2 734.7 Minimum 7.032 0.0 Model 3 Rib with a lightening hole Maximum 241.45(occurs on the blank) 0.206 (occurs on the blank) 6 553.8 Minimum 0.101 0.0 4 E. L. Deladi, Static friction rubber metal contact with application to rubber pad forming process, PhD Thesis, University of Twente, 2006 5 M. Ramezani, Z. M. Ripin, R. Ahmad, Sheet metal forming with the aid flexible punch, numerical approach and experimental validation, CIRP Journal of Manufacturing Science and Technology, 3 (2010) 3, 196–203 6 D. J. Browne, E. Battikha, Optimization of aluminium sheet forming using a flexible die, Journal of Materials Processing Technology, 55 (1995) 3/4, 218–223 7 M. H. Dirikolu, E. Akdemir, Computer aided modelling of flexible forming process, Journal of Materials Processing Technology, 148 (2004) 3, 376–381 8 R. Madoliat, R. Narimani, H. Rahrovan, Investigation of sheet metal forming using rubber pad forming, The SMEIR 2005 International Conference in Manufacturing Engineering, 2005, 1–9 9 S. Thiruvarudchelvan, The potential role of flexible tools in metal forming, Journal of Materials Processing Technology, 122 (2002) 2/3, 293–300 10 J. W. Lee, H. C. Kwon, M. H. Rhee, Y. T. Im, Determination of forming limit of a structural aluminum tube in rubber pad bending, Journal of Materials Processing Technology, 140 (2003) 1–3, 487–493 11 F. Quadrini, L. Santo, E. A. Squeo, Flexible forming of thin alumi- num alloy sheets, International Journal of Modern Manufacturing Technologies, 2 (2010) 1, 79–84 12 M. W. Fu, H. Li, J. Lu, S. Q. Lu, Numerical study on the deformation behaviors of the flexible die forming by using viscoplastic pressure- carrying medium, Computational Materials Science, 46 (2009) 4, 1058–1068 13 Y. Liu, L. Hua, J. Lanm, X. Wei, Studies of the deformation styles of the rubber-pad forming process used for manufacturing metallic bipolar plates, Journal of Power Sources, 195 (2010) 24, 8177–8184 14 Y. Liu, L. Hua, Fabrication of metallic bipolar plate for proton exchange membrane fuel cells by rubber pad forming, Journal of Power Sources, 195 (2010) 11, 3529–3535 15 H. Takuda, N. Hatta, Numerical Analysis of formability of an Aluminum 2024 alloy sheet and Its Laminates with Steel Sheets, Metallurgical and Material Transactions A, 29 (1998) 11, 2829–2834 16 W. Eichlseder, Enhanced fatigue analysis – incorporating down- stream manufacturing processes, Mater. Tehnol., 44 (2010) 4, 185–192 17 H. H. Lee, Finite Element Simulation with ANSYS Workbench12, Schroff Development Corporation, Taiwan, 2010 M. BENISA et al.: COMPUTER-AIDED MODELING OF THE RUBBER-PAD FORMING PROCESS 510 Materiali in tehnologije / Materials and technology 46 (2012) 5, 503–510 A. R. BO  GA et al.: EFFECT OF FLY-ASH AMOUNT AND CEMENT TYPE ON THE CORROSION PERFORMANCE ... EFFECT OF FLY-ASH AMOUNT AND CEMENT TYPE ON THE CORROSION PERFORMANCE OF THE STEEL EMBEDDED IN CONCRETE U^INEK KOLI^INE LETE^EGA PEPELA IN VRSTE CEMENTA NA KOROZIJO JEKLA V BETONU Ahmet Raif Boða1, Ýlker Bekir Topçu2, Murat Öztürk3 1Afyon Kocatepe University, Faculty of Engineering, Civil Engineering Department, Afyonkarahisar, Turkey 2Eskiºehir Osmangazi University, Faculty of Engineering and Architecture, Civil Engineering Department, 26480 Eskiºehir, Turkey 3Selçuk University, Faculty of Engineering and Architecture, Civil Engineering Department, Konya, Turkey muratozturk@selcuk.edu.tr Prejem rokopisa – received: 2011-10-10; sprejem za objavo – accepted for publication: 2012-03-01 In this study the corrosion performance of the steel embedded in the concrete produced by using three different types of cement (CEM II/B-M (P-L) 32.5 R, CEM I 42.5 R and CEM I 52.5 R) was investigated. 300 kg/m3 and 375 kg/m3 dosages of the cement with (0, 10 and 20) % of fly-ash (FA) replacements of cement were used to produce the concretes. These concretes were cured for 28 d and 180 d. The mechanical properties of the concretes were determined and the corrosion performances of the reinforced-concrete specimens were determined using the impressed voltage test. After the impressed voltage test weight losses occurred because of the corrosion that was determined. The results of this study show that using composite cement and an FA replacement of the cement are useful in combating corrosion. Keywords: concrete, corrosion, mechanical properties, fly ash V tej {tudiji je bilo preiskovano korozijsko vedenje jekla, vgrajenega v beton, izdelan iz treh vrst cementa (CEM II/B-M (P-L) 32,5 R, CEM I 42,5 R in CEM I 52,5 R). Odmerki 300 kg/m3 in 375 kg/m3 cementa z (0, 10 in 20) % lete~ega pepela (FA) kot nadomestila za cement, so bili uporabljeni za izdelavo betona. Beton je bil presku{en po 28 d in po 180 d. Dolo~ene so bile mehanske lastnosti betona. Korozijske lastnosti armiranega betona so bile dolo~ene s pospe{enim korozijskim napetostnim preizkusom. Pri pospe{enem napetostnem korozijskem preizkusu se je zaradi korozije pojavilo zmanj{anje mase. Rezultati te {tudije ka`ejo, da je za zmanj{anje korozije ugodna uporaba kompozita cementa z lete~im pepelom, ki nadomesti cement. Klju~ne besede: beton, korozija, mehanske lastnosti, lete~i pepel 1 INTRODUCTION Corrosion of the steel embedded in concrete plays a vital role in the determination of the life and durability of the concrete structures.1 The durability of reinforced concrete is largely controlled by the capability of the concrete cover to protect the steel reinforcement from corrosion. Chemical protection is provided by the concrete’s high alkalinity and physical protection is afforded by the concrete cover acting as a barrier to the access of aggressive species.2 The corrosion of the steel in concrete is retarded by the passivating ferric-oxide film ( -FeOOH) formed in the concrete medium (that is highly alkaline with a pH of around 13).1 Corrosion of the reinforcing-steel bars is initiated to form an inactive thin layer that can be broken when immersed in carbonate, chloride or sulphate solutions.3 Corrosion of steel produces rust products that have a volume three to eight times greater than that of the original metal. This generates stress and causes cracking and spalling of the concrete cover, which further accelerates corrosion.4 Various methods have been applied to protect reinforced steel against corrosion; these methods include variation of the concrete formulation, cathodic protection, surface treatment of the rebar and addition of inhibitors and mineral admixtures.5 It is generally recognized that the incorporation of fly ash (FA) in blended cements helps to protect concrete against the chloride-induced corrosion of steel reinforcement by reducing its permeability, particularly for chloride-ion transportation, and in- creasing the resistivity of the concrete.6,7 In this study, the corrosion performance of the steel embedded in the concrete produced by using three different types of cement was investigated. Cements in 300 kg/m3 and 375 kg/m3 dosages and also (0, 10 and 20) % FA replacements of cement were used to produce the concretes. The mechanical properties of the concretes were determined. The impressed voltage test was performed on the reinforced concrete specimens. In conclusion, the best cement type, FA ratios and curing times for the corrosion performance of the steel embedded in concrete were determined. 2 GENERAL OBSERVATIONS ABOUT THE BUILDING STOCK IN TURKEY According to the studies in various regions of Turkey, the average concrete compressive strength is approxi- mately 10 MPa.8,9 Several identified improper practices for fabricating an in-situ concrete caused this signifi- Materiali in tehnologije / Materials and technology 46 (2012) 5, 511–518 511 UDK 691.328.1:620.19:620.17 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 46(5)511(2012) cantly low concrete quality. Important factors include the ignorance of the aggregate gradation, the use of unwashed sea sand and/or river sand, and the use of the aggregate sizes that are too large. One of the most negative factors contributing to a low concrete quality is the use of the sea sand or river sand in a concrete mixture. According to the obtained results, the significantly high chloride content in the sea-sand concrete is an indication of a high tendency for the corrosion of the reinforcement in the reinforced concrete structures. The observations of the building rubbles showed that most rebars had corroded due to the use of the river/sea sand, thereby reducing the efficiency of the longitudinal rebar area and the anchorage (Figure 1). 3 EXPERIMENTAL STUDIES 3.1 Materials used 3.1.1 Cement The experimental studies used the CEM I 42.5 R Portland cement and the CEM II/ B-M (P-L) 32.5 R Portland composite cement, produced by the Eskisehir cement factory according to TS EN 197-1 : 2 000 standards. The CEM II/B-M (P-L) cement contains natural pozzolans and limestone in the ratio of 21–35 % by weight. In addition, the CEM I 52.5 R Portland cement was used. The results of the chemical and physical analyses of these cements, which were provided by the factories, are given in Table 1. 3.1.2 FA In the experiment, the Tunçbilek thermal power plant’s FA was used. The chemical composition and the physical properties are given in Table 1. Because the SiO2+Al2O3+Fe2O3 content exceeds 70 % and the CaO value is under 10 %, the FA of Tunçbilek is of class F (low lime) according to the ASTM C 618 standard. 3.1.3 Aggregates The Osmaneli sand and the Sö  güt Zemzemiye crushed-stone aggregates were used. The maximum particle size of the aggregates was 31.5 mm. According to the results of the experiment, the specific gravities of the sand and the crushed stones I and II are 2620, 2710 and 2710 and the unit weights are (1550, 1720 and 1770) kg/m3, respectively. (35, 30 and 35) %, of the sand and the crushed stones I and II, respectively, were used in the grain mixture. 3.1.4 Steel reinforcements and the NaCl solution 14-mm-diameter deformed steel reinforcement was used for the preparation of the reinforced concrete specimens to attempt the corrosion tests. According to TS 708 the minimum yield strength of this steel is 420 MPa and the minimum tensile strength is 500 MPa. In the experimental setup for the corrosion testing, an industrial type of the NaCl salt was used for obtaining the solution. 3.2 Mix proportions of the concrete Three different types of cement, CEM II/B-M(P-L) 32.5R (CII-3), CEM I 42.5R (CI-4) and CEM I 52.5R (CI-5), were used in the concrete mixtures. By using 300 A. R. BO  GA et al.: EFFECT OF FLY-ASH AMOUNT AND CEMENT TYPE ON THE CORROSION PERFORMANCE ... 512 Materiali in tehnologije / Materials and technology 46 (2012) 5, 511–518 Table 1: Chemical and physical properties of the cements and fly ash Tabela 1: Kemijske in fizikalne lastnosti cementov in lete~ega pepela Chemical Composition, % Cement Type FACEM I 52.5 R CEM I 42.5 R CEM II/B-M (P-L) 32.5 R SiO2 20.47 20.74 30.88 58.25 Al2O3 5.68 5.68 8.01 16.66 Fe2O3 3.08 4.12 3.57 12.91 CaO 62.66 63.70 47.78 1.95 MgO 1.1 1.22 1.30 5.08 Na2O 0.20 0.17 0.12 0.33 K2O 0.75 0.53 1.33 1.37 SO3 2.5 2.29 1.67 0.41 Cl 0.010 0.019 0.011 0.002 Loss of ignition 1.9 1.34 6.20 2.09 Insoluble residue 0.7 0.57 0.27 – Free lime 1.0 1.29 1.31 0.16 Physical Properties Specific gravity 3.17 3.14 2.85 2.34 Specific surface, cm2/g 3700 3450 3580 Compressive Strength, MPa 2 days 27 26 13 – 7 days 41 38 27 – 28 days 59 59 43 – Figure 1: Corrosion of the reinforcement in a damaged column observed after the Van earthquake of 23 October 2011 (Murat Öztürk's archive) Slika 1: Korozija armature v po{kodovanem stebru med potresom 23. 10. 2011 (arhiv Murat Öztürk) kg/m3 and 375 kg/m3 dosages of each cement, the concrete specimens were also produced as reference specimens containing 10 % and 20 % of FA. These specimens were cured for two different periods: 28 d and 180 d. Thus, 36 series of the concrete mixture were produced. The mix proportions of the concrete are given in Table 2. 3.3 Specimen preparation and testing 3.3.1 Compressive test The compressive strength test was carried out on the specimens with the cube dimensions of 150 mm × 150 mm × 150 mm. The specimens were demoulded after 24 h and immersed in water at (20 ± 2) °C. The compressive strength tests were made after 28 d and 180 d of curing. 3.3.2 Splitting tensile test The splitting tensile test was carried out for the cylinder specimens with the dimensions of 150 mm × 300 mm. The specimens were demoulded after 24 hours and immersed in water at (20 ± 2) °C. The splitting tensile tests were carried out after 28 d of curing. 3.3.3 Tests for the ultrasonic pulse velocity and the modulus of dynamic elasticity The tests of measuring the ultrasonic pulse velocity and the modulus of dynamic elasticity were carried out for the specimens that were prepared for the splitting tensile tests. 3.3.4 Impressed voltage test The setup for this test included a DC power source, a test specimen and a plastic dish containing a 4 %-NaCl solution, two steel plates and a data logger. The impressed-voltage test setup is shown in Figure 2. The reinforced concrete specimens for the accelerated- corrosion tests were the cylinder specimens with the dimensions of 150 mm × 300 mm, in which a 14-mm-diameter steel reinforcement was centrally embedded. A 250-mm part of the steel reinforcement was embedded into each concrete cylinder. The specimens were demoulded after 24 h and immersed in water at (20 ± 2) °C. The impressed voltage tests were carried out after 28 d and 180 d of curing. The steel reinforcement (the working electrode) of the reinforced concrete specimen was connected to the positive terminal and the steel plates (the counter electrodes) were connected to the negative terminal of the DC power source applying 30 V of fixed stress to the system. A similar impressed-voltage test setup has been reported by other researchers.7,10 Every five minutes the corrosion current of every reservoir was saved using a data logger and the corrosion current-time figures were drawn using the impressed voltage test system. A. R. BO  GA et al.: EFFECT OF FLY-ASH AMOUNT AND CEMENT TYPE ON THE CORROSION PERFORMANCE ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 511–518 513 Table 2: Mix proportions of the concrete for 1 m3 Tabela 2: Dele` sestavin v 1 m3 betona Cement Type Dosage FA % Cement Sand Crush- ed Stone I Crush- ed Stone II Water FA CII-3 375 0 375 611 541 632 188 – 10 337.5 608 539 629 188 37.5 20 300 605 537 626 188 75 300 0 300 669 593 692 150 – 10 270 667 591 690 150 30 20 240 665 590 688 150 60 CI-4 375 0 375 622 553 645 188 – 10 337.5 618 550 642 188 37.5 20 300 614 547 638 188 75 300 0 300 678 603 704 150 – 10 270 675 601 701 150 30 20 240 672 598 698 150 60 CI-5 375 0 375 623 554 646 188 – 10 337.5 619 551 643 188 37.5 20 300 615 548 639 188 75 300 0 300 679 604 705 150 – 10 270 676 602 702 150 30 20 240 673 599 699 150 60 *Superplasticiser (SP) was used as 0.4 % by mass of binder (cement and FA) Figure 2: Setup for the impressed voltage test Slika 2: Skica pospe{enega napetostnega preizkusa 3.3.5 Weight loss method The method proposed in the ASTM G1-03 standard was used for the determination of the weight loss of a steel bar.11 After the impressed voltage tests, steel bars were demounted from the reinforced concrete specimens and the weight losses were found by cleaning these steel bars with a Clarke solution, which contained 1000 mL HCL, 24 g Sb2O3 and 71.3 g SnCl2 ⋅ 2H2O. 4 RESULTS AND DISCUSSION 4.1 Compressive strength Changes in the compressive strengths due to the amount of FA are shown in Figures 3 and 4. According to the figures for the 28 d and 180 d compressive strengths, the strength results of the concrete produced with the CI-5 cement are higher than those of the others. With the increase in the FA amount in the specimens, the compressive strengths of the specimens cured for 28 d decreased because the mineral admixtures like FA cannot fully complete the pozzolanic reactions.12 If Figure 4 is examined, it can be seen that the strength values increase as the dosage increases. Looking at Figure 4, by using 10 % of FA instead of the CI-5 and CI-4 cements in the 375 dosed specimens, the compressive strength of the specimens increases respectively by 4.5 % and 15.66 % according to the control specimens and it decreases at the ratio of 6.42 % in the specimens produced with CII-3. If 20 % of FA is used in the specimens, the compressive strength decreases at the ratio of 6.06 % and 12.17 % for the specimen series produced with the CI-5 and CII-3 cements, but increases at the ratio of 5.43 % for the specimen series produced with the CI-4 cement. If Figures 3 and 4 are compared to each other, it is seen that the compressive strength increases as the curing time increases. 4.2 Splitting tensile strength According to a general assessment from Figure 5, the splitting tensile strengths of the specimens produced with the CI-4 and CI-5 cements are very similar to each other. The minimum strengths are found with the speci- mens produced with the CII-3 cement. As the amount of FA increases in the specimens produced with CII-3, the splitting tensile strength decreases but it increases in the specimens produced with the other cements. According to Figure 5, the splitting tensile strength is seen to increase as the dosage of the cement is increased. 4.3 Tests for the ultrasonic pulse velocity and modulus of dynamic elasticity The ultrasonic pulse velocity and modulus of dyna- mic elasticity test results are shown in Table 3. Con- sidering the ultrasonic pulse velocity of the 375 kg/m3 and 300 kg/m3 dosage specimens, it can be seen that the ultrasonic pulse velocity of all the specimens produced with CII-3 are fine. The ultrasonic pulse velocity results of the specimens produced with CI-4 and CI-5 are seen to be better than those of CII-3. Depending on the A. R. BO  GA et al.: EFFECT OF FLY-ASH AMOUNT AND CEMENT TYPE ON THE CORROSION PERFORMANCE ... 514 Materiali in tehnologije / Materials and technology 46 (2012) 5, 511–518 Figure 5: Relationship between the splitting tensile strength and the FA amount Slika 5: Odvisnost med poru{no natezno trdnostjo in dele`em lete~ega pepela (FA) Figure 3: Relationship between the compressive strength and the FA amount in 28-day specimens Slika 3: Odvisnost med tla~no trdnostjo in dele`em lete~ega pepela (FA) v vzorcu po 28 dneh Figure 4: Relationship between the compressive strength and the FA amount in 180-day specimens Slika 4: Odvisnost med tla~no trdnostjo in dele`em lete~ega pepela (FA) v vzorcu po 180 dneh ultrasonic pulse velocity values, the best results are seen to be in the specimens produced with the CI-5 cement and without any FA. According to the assessment in Table 3, it can be seen that the modulus of dynamic elasticity decreases as the amount of FA increases. The minimum Edin values are for the specimens produced using the CII-3 cement with the dosages of 300 kg/m3 and 375 kg/m3. As seen in Table 3, for the other series, the Edin values are similar. Table 3: Test results for ultrasonic pulse velocity and modulus of dynamic elasticity Tabela 3: Rezultati preizkusov za hitrost ultrazvo~nega impulza in modul dinami~ne elasti~nosti Cement Type FA/% Ultrasonic pulse velocity /(km/s) Edin/GPa 375 Dosage 300 Dosage 375 Dosage 300 Dosage CII-3 0 4.35 4.52 45.7 49.4 10 4.46 4.46 47.9 48.0 20 4.29 4.40 43.6 46.3 CI-4 0 4.62 4.62 52.3 53.1 10 4.63 4.50 52.3 49.9 20 4.62 4.60 51.3 51.5 CI-5 0 4.66 4.63 53.5 52.8 10 4.56 4.56 50.5 50.5 20 4.65 4.55 52.1 50.5 4.4 Impressed voltage test After the impressed voltage tests, steel in the reinforced concrete corroded and the specimens were damaged. The damaged specimens are shown in Figure 6. Since the volume of the corrosion products (rust) are 2.5–6 times greater than the volume of the steel used in the concrete, these corrosion products lead to higher internal tensile stresses in the hardened concrete. Being exposed to these stresses, the hardened concrete cracks and splits off.12 The compressive strength results and damage occurrence times (DOTs) of the accelerated-corrosion specimens are shown in Table 4. As seen in Table 4, the DOT ranges between 251 h and 394 h. According to Table 4, as is the case with the 300 kg/m3 dosage specimens, an increase in the FA amount leads to a longer DOT. With respect to DOTs, it is observed that the best results are obtained for the specimens produced with CII-3 and 20 % FA cured for 28 d, and the specimens produced with CI-5 and 20 % FA cured for 180 d. A comparison between the 300 kg/m3 and 375 kg/m3 dosage specimen results shows that the DOT extends as the cement dosage increases. Generally, if the cement dosage increases, the compressive strength of the con- crete increases as well. There is a general relationship between the compressive strength of the concrete and the permeability.12 According to Table 4, the DOT is extend- ed even if the compressive strengths of the concretes produced with the CII-3 cement are low. Actually, the concretes produced with the CII-3 cement that have low compressive strengths are damaged in a shorter time than the concretes produced with the CI-4 and CI-5 cements, because the concretes produced with CII-3 are more porous than the other concretes, as can be understood from the compressive strengths. Table 4: Compressive strengths and damage occurrence times (DOTs) of the 300 – (375) dosage specimens Tabela 4: Tla~ne trdnosti in ~asi do pojava napak (DOT) vzorcev z odmerkom 300 – (375) Cement Type FA/ % Compressive Strength, MPa 300 (375) Damage Occurrence Time Hour – 300 (375) 28 days 180 days 28 days 180 days CII-3 0 36.7 (37.2) 38.9 (45.2) 291 (325) 368 (308) 10 33.8 (34.7) 41.6 (42.3) 311 (323) 381 (324) 20 31.2 (32.0) 38.3 (39.7) 325 (366) 394 (373) CI-4 0 46.5 (47.5) 48.2 (47.9) 279 (287) 303 (306) 10 43.7 (44.4) 49.2 (55.4) 310 (315) 324 (326) 20 42.6 (42.2) 48.1 (50.5) 319 (326) 327 (329) CI-5 0 53.6 (55.2) 64.7 (64.4) 251 (300) 304 (315) 10 50.7 (54.0) 63.4 (67.3) 268 (316) 343 (368) 20 49.8 (52.8) 53.6 (60.5) 287 (292) 376 (389) However, it is known from the previous studies that fly ash, ground granulated blast furnace slag, silica fume and pozzolans bind chloride ions; thus, the chloride permeability decreases.13–16 The concretes produced with the CII-3 cement are highly porous and permeable due to a high percentage of FA and natural pozzolans that do not pozzolanicly react with water and lime. Therefore, these FA and pozzolan particles react with chloride ions and bind them. In this way, the chloride permeabilities of these concretes are lower than those of the control ones. Figure 7 shows that the corrosion-time curves of the 28 d cured concretes were produced with a 300 kg/m3 dosage of the CII-3, CI-4 and CI-5 cements, and 10 % and 20 % FA replacements of the cement. According to the assessment in Figure 7, for the concretes produced with the CII-3, CI-4 and CI-5 cements, the curves 2, 3, 5, 6, 8 and 9 show that an increase in the FA amount in the concrete, used as a replacement for the cement, reduces the corrosion currents. According to these results the use A. R. BO  GA et al.: EFFECT OF FLY-ASH AMOUNT AND CEMENT TYPE ON THE CORROSION PERFORMANCE ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 511–518 515 Figure 6: Damaged specimens after the impressed voltage test Slika 6: Po{kodovani vzorci po pospe{enem napetostnem preizkusu of FA is beneficial. It has been observed that FA binds chloride ions.15,16 As a result, the corrosion currents decrease as seen in Figure 7. Figure 8 shows the current-time relations for the corrosion of the concretes that are cured for 28 d and produced by using FA as the 10 % and 20 % replacements of the cement weight with a 375 kg/m3 dosage of the CII-3, CI-4 and CI-5 cements. The series produced by using the CII-3 cement gives better test results. However, the best test results are achieved with the concretes produced with a 20 % FA replacement of the CI-4 cement. It is seen that the corrosion current of the specimens is reduced by using the FA replacement for a 375 kg/m3 dosage of the CI-5 cement with the given ratios. By using FA, the corrosion currents are reduced, but better results are achieved than with the other concrete specimens. According to Figure 8, the corrosion currents increase with time and it is seen that some series have sudden increases. The corrosion current-time relations are shown in Figure 9 for the concretes that are cured for 180 d and produced by using FA as the 10 % and 20 % replacements of the cement weight for a 300 kg/m3 dosage of the CII-3, CI-4 and CI-5 cements. Looking at Figure 9, with the concretes produced with the CI-4 and CI-5 cements, and without any FA, they are seen to have a corrosion current above 0.1 A. All the other specimens have a corrosion current under 0.1 A at the beginning. The concretes cured for 180 days and produced with a 300-kg/m3 dosage of the CI-4 cement depending on their FA ratio (0 %, 10 % and 20 %) have (0.12, 0.08 and 0.04) A corrosion currents in turn. However, the same series that were cured for 28 d have corrosion currents of (0.18, 0.12 and 0.09) A in turn. According to these results an increase in the curing time reduces the corrosion currents. The corrosion current-time relations are shown in Figure 10 for the concretes cured for 180 d and produced by using FA as the 10 % and 20 % replacements of the cement weight for the 375 kg/m3 dosage of the CII-3, CI-4 and CI-5 cements. With respect to Figure 10, the best results are achieved for the concretes produced with a 20 % FA replacement of the weight of the CII-3 cement. 4.5 Weight loss method The steel reinforcements cleaned with a Clarke solution are shown in Figure 11. The weight losses of the steel reinforcements in various concrete mixtures that were exposed to corrosion are shown in Figures 12 and 13. According to a general assessment of these figures, A. R. BO  GA et al.: EFFECT OF FLY-ASH AMOUNT AND CEMENT TYPE ON THE CORROSION PERFORMANCE ... 516 Materiali in tehnologije / Materials and technology 46 (2012) 5, 511–518 Figure 10: Variation of the corrosion current according to time (a 375-kg/m3 dosage and 180-day specimens) Slika 10: Spreminjanje korozijskega toka glede na ~as (odmerek 375 kg/m3 in vzorec po 180 dneh) Figure 8: Variation of the corrosion current according to time (a 375-kg/m3 dosage and 28-day specimens) Slika 8: Spreminjanje korozijskega toka glede na ~as (odmerek 375 kg/m3 in vzorec po 28 dneh) Figure 7: Variation of the corrosion current according to time (a 300-kg/m3 dosage and 28-day specimens) Slika 7: Spreminjanje korozijskega toka glede na ~as (odmerek 300 kg/m3 in vzorec po 28 dneh) Figure 9: Variation of the corrosion current according to time (a 300-kg/m3 dosage and 180-day specimens) Slika 9: Spreminjanje korozijskega toka glede na ~as (odmerek 300 kg/m3 in vzorec po 180 dneh) the weight loss is seen to decrease as the amount of FA, being the concrete replacement of the cement, is increased. Increasing the dosage reduces the weight loss in some series, but increases it in others. As it is seen in Figure 12, the maximum weight loss is observed with the concretes that did not have any FA and were produced with CI-4. According to Figure 13, the weight losses decrease as the amount of FA increases in the concretes. According to the increase in the curing time, by comparing Figures 12 and 13, in some series the weight loss of the reinforcements is seen to increase and in some series it decreases. 5 CONCLUSIONS After the tests it was observed that the mechanical properties of the concretes and the corrosion perfor- mances of the steel embedded in the concrete changed with different cement types, dosages, FA amounts used as the replacements for the cement and the curing time. • The compressive strength of the specimens increased after an increase in the curing time and the cement dosage. The maximum compressive strengths were observed with the concretes that were produced with the CI-5 cement. • The splitting tensile test results of the specimens that were produced with the CI-4 and CI-5 cements were very similar. The minimum strengths were observed with the specimens that were produced with CII-3. • According to a general assessment of the results, an increase in the dosage, the curing time and the FA amount used to replace the cement caused an increase in the DOT of the reinforced concrete specimens. • With respect to the assessment of the DOTs, the most positive results were observed with the specimens that were produced with the CII-3 cement. Com- patible results were observed with the specimens that were produced with the other cements and the 20 % FA amount. • Consequently, the corrosion of the steel embedded in the concrete depends significantly on the cement type, the cement dosage and the curing time. To combat the corrosion of the steel reinforcement in the concrete, the permeability of concrete by water and hazardous ions has to be prevented. This is possible by producing an impermeable concrete. For this reason, impermeable and high-quality concretes have to be produced by using various mineral admixtures such as FA and SF. 6 REFERENCES 1 T. Parthiban, R. Ravi, G. T. Parthiban, S. Srinivasan, K. R. Rama- krishnan, M. Raghavan, Neural network analysis for corrosion of steel in concrete, Corr. Sci., 47 (2005), 1625–1642 2 E. Güneyisi, T. Özturan, M. Geso  glu, Effect of initial curing on chloride ingress and corrosion resistance characteristics of concretes made with plain and blended cements, Build. and Environ., 42 (2007), 2676–2685 3 K. Sakr, Effect of cement type on the corrosion of reinforcing steel bars exposed to acidic media using electrochemical techniques, Cem. Concr. Res., 35 (2005), 1820–1826 4 K. Y. Ann, H. S. Jung, H. S. Kim, S. S. Kim, H. Y. Moon, Effect of calcium nitrite-based corrosion inhibitor in preventing corrosion of embedded steel in concrete, Cement and Concrete Research, 36 (2006), 530–535 A. R. BO  GA et al.: EFFECT OF FLY-ASH AMOUNT AND CEMENT TYPE ON THE CORROSION PERFORMANCE ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 511–518 517 Figure 13: Variation of weight loss by mixture type for 180-day spe- cimens Slika 13: Razlike v izgubi mase glede na vrsto me{anice v vzorcih po 180 dneh Figure 12: Variation of weight loss by mixture type for 28-day specimens Slika 12: Razlike v izgubi mase glede na vrsto me{anice v vzorcih po 28 dneh Figure 11: Steel reinforcements cleaned with a Clarke solution Slika 11: Z raztopino Clarke o~i{~ena jeklena armatura 5 A. A. Gürten, K. Kayakýrýlmaz, M. Erbil, The effect of thiosemi- carbazide on corrosion resistance of steel reinforcement in concrete, Const. and Buil. Mater., 21 (2007), 669–676 6 P. Chindaprasirt, C. Chotithanorm, H. T. Cao, V. Sirivivatnanon, Influence of fly ash fineness on the chloride penetration of concrete, Const. Build. Mater., 21 (2007), 356–361 7 E. Güneyisi, T. Özturan, M. Gesoðlu, A study on reinforcement corrosion and related properties of plain and blended cement concretes under different curing conditions, Cem. Concr. Compos., 27 (2005), 449–461 8 M. Y. Kaltakci, M. Kamanli, M. Ozturk, M. H. Arslan, H. H. Kork- maz, Sudden Complete Collapse of Zumrut Apartment Building and It’s Causes, Journal of Performance of Constructed Facilities, doi:http://dx.doi.org/10.1061/(ASCE)CF.1943-5509.0000337 9 M. H. Arslan, H. H. Korkmaz, What is to be Learned from Damage and Failure of Reinforced Concrete Structures During Recent Earthquakes in Turkey, Engineering Failure Analysis, 14 (2007), 1–22 10 E. Güneyisi, M. Gesoðlu, A study on durability properties of high-performance concretes incorporating high replacement levels of slag, Materials and Structures, 41 (2008), 479–493 11 ASTM G1–03, Standard Practice for Preparing, Cleaning, and Evaluating Corrosion Test Specimens, ASTM Book of Standards Volume: 3. 2. 2003 12 T. Y. Erdogan, Concrete, METU Press, Ankara, Turkey, 2003 13 K. Y. Yeau, E. K. Kim, An experimental study on corrosion resistan- ce of concrete with ground granulate blast-furnace slag, Cem. and Concr. Res., 35 (2005), 1391–1399 14 K. M. A. Hossain, M. Lachemi, Corrosion resistance and chloride diffusivity of volcanic ash blended cement mortar, Cem. and Con- crete Res., 34 (2004), 695–702 15 S. P. Shah, S. H. Ahmad, High performance concretes and applica- tions, Hodder Headline Group, 1994 16 R. B. Polder, W. H. A. Peelen, Characterisation of chloride transport and reinforcement corrosion in concrete under cyclic wetting and drying by electrical resistivity, Cem. Concr. Compos., 24 (2002), 427–435 A. R. BO  GA et al.: EFFECT OF FLY-ASH AMOUNT AND CEMENT TYPE ON THE CORROSION PERFORMANCE ... 518 Materiali in tehnologije / Materials and technology 46 (2012) 5, 511–518 A. GIGOVI]-GEKI] et al.: EFFECT OF THE DELTA-FERRITE CONTENT ON THE TENSILE PROPERTIES ... EFFECT OF THE DELTA-FERRITE CONTENT ON THE TENSILE PROPERTIES IN NITRONIC 60 STEEL AT ROOM TEMPERATURE AND 750 °C VPLIV VSEBNOSTI DELTA FERITA NA NATEZNE LASTNOSTI JEKLA NITRONIC 60 PRI SOBNI TEMPERATURI IN PRI 750 °C Almaida Gigovi}-Geki}1, Mirsada Oru~2, Sulejman Muhamedagi}1 1University of Zenica, Faculty of Metallurgy and Materials Science, Travni~ka cesta 1, Zenica, Bosnia and Herzegovina 2Metallurgical Institute "Kemal Kapetanovi}", Zenica, Bosnia and Herzegovina almaida.gigovic@famm.unze.ba Prejem rokopisa – received: 2011-10-19; sprejem za objavo – accepted for publication: 2012-03-15 This paper presents the results of the tensile testing of the austenitic stainless steel Nitronic 60 at room temperature and 750 °C in the solution-annealed condition. It also presents the results of the optical and SEM analyses of the tested samples. The microstructural analysis showed the presence of a delta-ferrite phase in the austenite matrix at room temperature. The content of the delta ferrite was calculated using a Feritscope MP30. The content of the delta ferrite depends on the chemical composition. An increase in the Si and Cr contents causes an increase in the delta-ferrite content. The results also show that an increase in the delta-ferrite content leads to an increase in the strength and a decrease in the ductility at room temperature. After testing the samples at 750 °C, the presence of a sigma phase was noticed. Precipitation of the sigma phase causes a slight increase in the strength and a decrease in the ductility of the tested material. An analysis of the fracture surface shows the presence of ductile fracture in the samples tested at room temperature and a combination of ductile and brittle fractures occurring at 750 °C. Keywords: austenitic steel, Nitronic 60, tensile properties, SEM analysis, delta ferrite, sigma phase ^lanek predstavlja rezultate nateznih preizkusov avstenitnega nerjavnega jekla Nitronic 60 v raztopno `arjenem stanju pri sobni temperaturi in pri 750 °C. Predstavljeni so tudi rezultati preiskav vzorcev s svetlobno in SEM analizo. Analiza mikrostrukture pri sobni temperaturi je pokazala prisotnost delta ferita v avstenitni osnovi. Vsebnost dela ferita je bila dolo~ena z uporabo naprave Feritscope MP30.Vsebnost delta ferita je odvisna od kemijske sestave. Pove~anje vsebnosti Si in Cr pove~a tudi vsebnost delta ferita. Iz rezultatov izhaja, da pove~anje vsebnosti delta ferita pove~a trdnost in zmanj{a preoblikovalnost pri sobni temperaturi. Po preizku{anju pri 750 °C je bila opa`ena prisotnost sigma faze. Izlo~anje te faze rahlo pove~a trdnost in poslab{a preoblikovalnost preizku{enega materiala. Analiza prelomov ka`e `ilav prelom pri vzorcih, preizku{enih pri sobni temperaturi in kombinacijo `ilavega in krhkega preloma pri 750 °C. Klju~ne besede: avstenitno jeklo, Nitronic 60, natezne lastnosti, SEM-analiza, delta ferit, sigma faza 1 INTRODUCTION Microstructure stability is the most important requirement needed to obtain proper mechanical properties for an austenitic stainless steel (ASS).1 To achieve a stable microstructure, the samples are usually solution heat treated at a temperature between 1000 °C and 1120 °C 2 and then water quenched. The micro- structure of Nitronic 60 is primarily monophasic, i.e., austenitic. However, precipitation of the delta ferrite (-ferrite) in an austenite matrix is possible, too. A higher volume fraction of the -ferrite can be achieved in the microstructure of the samples by changing the chemical composition of steel in terms of increasing the content of the -ferrite-stabilising elements1 such as Cr, Si, Ti, Mo, etc. The presence of the -ferrite, which has a BCC crystalline structure, slows the grain growth and increases the strength properties of the steel because the interphase boundaries act as strong barriers to the dislocation motion3. During the annealing at 600–900 °C the intermetallic phases and carbides precipitate from the austenite and/or the -ferrite.2,4 One of the most common phases in ASS is the sigma phase (-phase).2,5 As a result of the heat-treatment temperature, the -ferrite can transform in the austenite and the -phase.4–10 The -phase is an intermetallic compound with a complex, tetragonal, crystal structure. The chemical composition of this phase varies considerably and it is therefore difficult to define this phase in the form of unique formulas. At room temperature this phase is hard, brittle and nonmagnetic,10 therefore having a negative effect on the mechanical properties especially on the toughness and ductility. The presence of the -ferrite reduces the incubation period of the precipitation of the -phase. The rate of the -phase precipitation from the -ferrite is about 100 times more rapid than the rate of the -phase precipitation directly from austenite4. The temperature interval, in which this phase occurs with most of the commercial steels, is between 590 °C and 870 °C, but the phase decomposes at temperatures above 1000 °C. To make the samples free of the -phase, the heat-treat- ment temperature has to be higher than 1000 °C, followed by rapid cooling. The aim of the research is an investigation of the influence of -ferrite on the tensile properties of the tested materials at room and elevated temperatures Materiali in tehnologije / Materials and technology 46 (2012) 5, 519–523 519 UDK 669.14.018.8:669.112:620.17 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 46(5)519(2012) because the Nitronic 60 parts are intended to operate at elevated temperatures. For the research project we tested six melts with different contents of the alloying elements of Si, Cr, Mn and Ni in order to get the samples with different contents of the -ferrite. 2 EXPERIMENTAL WORK The melts were prepared by using a vacuum induc- tion furnace with an argon protective atmosphere with the chemical composition corresponding to the standard ASTM A276,11 as shown in Table 1. The requirements relating to the chemical composition and the tensile properties of the tested material correspond to the requirements for the steel UNS S21800.11 The average -ferrite content was determined with a Feritscope MP30 (Fisher, Germany) using broken specimens after the tensile testing at room temperature and 750 °C, as seen in Table 1. The content of the -ferrite was determined with the method that takes advantage of the fact that the -ferrite is magnetic, while austenite and the -phase are not. The average value was calculated on the basis of ten measurements. A microstructural analysis was carried out with an Olympus optical microscope and a scanning electron microscope (Jeol JSM 5610 operating at 20kV, a take-off angle of 35°, and an elapsed livetime of 60 %) using broken specimens after the tensile testing. The samples for the microstructure analysis were prepared with the standard grinding and polishing techniques and etched with an aqua regia. In addition, a fractographic analysis was carried out on broken specimens after the tensile testing with the scanning electron microscope. The tensile testing was performed on the samples obtained from a rod with a  of 15 mm. Before the testing, the samples were solution annealed at 1020 °C for 60 min followed by water quenching. The samples were tested at room temperature and 750 °C. The tensile test was carried out on a universal hydraulic machine for static testing (200 kN). The sampling and testing pro- cedures were realized in accordance with the standards BAS EN 10002-1/02 and BAS EN 10002-5/01. 3 RESULTS AND DISCUSSION 3.1 Microstructure The microstructure analysis of the samples tested at room temperature shows the presence of a two-phase microstructure. The microstructure consists of -ferrite islands in an austenite matrix.12 The -ferrite islands are elongated in the rolling direction. The precipitation of the -ferrite occurs mainly at the grain boundaries as shown in Figure 1a. Figure 1b shows the microstructure A. GIGOVI]-GEKI] et al.: EFFECT OF THE DELTA-FERRITE CONTENT ON THE TENSILE PROPERTIES ... 520 Materiali in tehnologije / Materials and technology 46 (2012) 5, 519–523 Table 1: Chemical composition and the average -ferrite content of steel Nitronic 60 Tabela 1: Kemijska sestava in povpre~na vsebnost -ferita v jeklu Nitronic 60 Melt Chemical composition, w/% -ferrite /% (Room temp.) -ferrite /% (750 °C) C Si Mn Cr Ni P S N Prescribed ASTM A 276 max. 0.10 3.5–4.5 7.0–9.0 16.0–18.0 8.0–9.0 max. 0.06 max. 0.03 0.08–0.18 1 0.04 4.41 7.4 18.0 8.1 0.007 0.005 0.18 10.43 0.52 2 0.04 3.74 8.6 18.0 8.0 0.007 0.005 0.16 6.30 0.31 3 0.04 4.25 8.4 16.0 8.8 0.006 0.010 0.14 3.37 0.77 4 0.05 3.5 7.9 16.9 8.6 0.005 0.005 0.12 1.30 0.58 5 0.04 3.5 7.2 16.6 8.0 0.006 0.012 0.15 1.14 0.59 6 0.05 3.8 8.9 17.0 9.0 0.007 0.005 0.16 0.82 0.46 Figure 1: Microstructure of steel Nitronic 60 at: a) room temperature and b) 750 °C Slika 1: Mikrostruktura jekla Nitronic 60 pri: a) sobni temperaturi in b) 750 °C of a sample tested at 750 °C exhibiting an austenite matrix and a transformed -ferrite phase. The SEM analysis of the sample tested at 750 °C with a higher magnification is shown in Figure 2, confirming the transformation of the -ferrite. The nucleation of the sigma phases predominantly occurred at the austenite/-ferrite grain boundaries, because these grain boundaries and the interfaces are the high-energy regions. The average compositions of the constituent phases were determined with an EDS analysis and presented in Table 2. The results show that the -phase has the highest content of Cr and that the content of Ni is quite high, too. The austenite is rich in Ni and Mn, but depleted in Cr and Si. In the case of the -ferrite, it is depleted in Ni and rich in Cr. Table 2: Chemical compositions of the constituent phases in mass fractions (w/%) obtained with an EDS analysis Tabela 2: Kemijska sestava faz v masnih dele`ih (w/%), dolo~ena z EDS-analizo Elements Constituent phases -ferrite -phase Austenite Si 3.50 4.96 3.12 Cr 19.15 36.30 17.50 Mn 6.28 5.94 8.52 Ni 5.06 8.02 7.22 The -ferrite is decomposed into the -phase and the austenite with a eutectoid transformation.9 The rate of the -phase precipitation from the -ferrite is more rapid than its precipitation directly from the austenite4, and the -ferrite is not in the state of equilibrium at this tem- perature. The high susceptibility of the -ferrite phase to -phase formation is associated with the chemical composition of the -ferrite phase. This phase is rich in Cr and Si, which stimulate the formation of the -phase.9,13 In the case of the alloys with the contents of Cr below 25 %, an addition of Mn and Ni can also stimulate the formation of the -phase.4,14 3.2 Fracture surface The analysis of the samples’ fracture surfaces shows that the samples tested at room temperature have ductile fracture but the samples tested at 750 °C have intergranular brittle fracture as a consequence of the presence of the -phase with a small portion of a ductile fracture.15 The ductile fracture is a result of the presence of austenite and the -ferrite, shown in Figures 3a and 3b. 3.3 Mechanical properties The results of the tensile testing at room temperature are in accordance with the standard ASTM A 276.11 However, there is no data about the tensile testing at elevated temperature in the ASTM A 276 11 standard. The results at room temperature indicate that an increase in the -ferrite content from 0.82 % to 10.43 % causes an increase in the tensile strength (TS) from 707 MPa to 821 MPa and an increase in the yield strength (YS) from 356 MPa to 467 MPa, while the reduction (Red.) and elongation (Elo.) are slightly reduced from 73 % to 63 % and from 55.3 % to 42.9 %, respectively, as shown in Figure 4. A. GIGOVI]-GEKI] et al.: EFFECT OF THE DELTA-FERRITE CONTENT ON THE TENSILE PROPERTIES ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 519–523 521 Figure 2: SEM microstructure samples tested at 750 °C Slika 2: SEM-posnetek vzorca, preizku{enega pri 750 °C Figure 3: Fracture surface – SEM: a) room temperature and b) 750 °C Slika 3: Povr{ina preloma – SEM: a) sobna temperatura in b) 750 °C The highest value of the strength and the lowest value of ductile properties were found for the melt 1 with 10.43 % of the -ferrite. The values of the strength and ductility are almost constant being up to 3 % of the -ferrite. With an increase in the -ferrite content to over 3 %, the strength is increased and the ductility is decreased. The heating of the tested material at 750 °C causes a remarkable decrease in the strength and ductility of the samples in comparison with the test results obtained at room temperature. One of the reasons for this was the transformation of the -ferrite in the -phase. The analysis of the -ferrite content after the tensile testing at 750 °C shows that the -phase content increases with an increase in the -ferrite content. The content of the transformed -ferrite was about 60 % for the melts with the -ferrite content of up to 2 % and about 95 % for the melts with the -ferrite content of over 6 %. The average content of the transformed -ferrite for the melts from 2 % to 6 % was about 70 %. Figure 5 shows that with an increase in the content of the transformed -ferrite from 77.15 % to 95.01 %, the tensile and yield strengths slightly increase from 245 MPa to 299MPa and from 180 MPa to 211 MPa, respectively. The yield strength is almost constant, being up to 77 %, while the tensile strength is decreasing. The ductility increases with an increase in the content of the transformed -ferrite from 43.9 % to 77.15 %. In the range between 77.15 % and 95.01 % the ductility slightly decreases. The value of the reduction is between 40 and 55 %, while the value of the elongation is bet- ween 25 and 45 %. The reason for the increase in the tensile strength is the precipitation of the -phase because the strength generally grows together with the growth of the precipitated intermetallic phases.16,17 The intermetallic phases cause an increase in the strength by obstructing, or stopping, the movement of dislocations. It is known from the literature sources13 that by controlling the distribution and morphology of the -phase, it is possible to improve the strength and ductility in the temperature range where the -phase precipitates. Generally, the tensile properties of steel Nitronic 60 depend on the chemical composition of the alloy. By controlling the chemical composition of the alloy we control the -ferrite content in the steel. The content of the -ferrite increases with an increased concentration of the ferrite-stabilizing elements, such as Cr and Si, and with a decreased concentration of the austenite-stabi- lizing elements, such as Mn, Ni and N. The -ferrite content can become lower than 1 % if the contents of Ni and Mn are increased to the upper allowed limit and the contents of Cr and Si are decreased to the middle of the allowed limit,18 as seen in Table 1. A reduction of the -ferrite content leads to a decrease in the strength and an increase in the ductility of steel at room temperature. At the elevated temperature (750 °C), the -ferrite transforms to the -phase, having a negative effect on the tensile properties because it is a hard and brittle phase. 4 CONCLUSION In this study, the influence of the -ferrite on the tensile properties at room temperature and 750 °C was investigated. From the presented results it could be concluded that: • The microstructure of the solution-annealed steel Nitronic 60 consisted of -ferrite islands in an auste- nite matrix at room temperature. • An increase in the -ferrite content leads to an increase in the strength and a decrease in the ductility at room temperature, especially for the steel with over 3 % -ferrite. • An increase in the content of Si and Cr causes an increase in the strength and a decrease in the ductility of steel due to the increased content of the -ferrite. • The EDS analysis shows the presence of the -phase at the ferrite/austenite boundaries, which is a result of the -ferrite transformation at 750 °C. The percen- tage of the -ferrite decomposition in the samples was up to 90 %. • The strength slightly increased and the ductility decreased after the content of the transformed -ferrite increased above 70 %. A. GIGOVI]-GEKI] et al.: EFFECT OF THE DELTA-FERRITE CONTENT ON THE TENSILE PROPERTIES ... 522 Materiali in tehnologije / Materials and technology 46 (2012) 5, 519–523 Figure 4: Effect of the -ferrite content on the tensile properties at room temperature Slika 4: U~inek vsebnosti -ferita na natezne lastnosti pri sobni temperaturi Figure 5: Effect of the content of transformed -ferrite on the tensile properties at 750 °C Slika 5: U~inek vsebnosti pretvorjenega -ferita na natezne lastnosti pri 750 °C • The SEM analysis also shows the presence of a mixture of the intergranular brittle fracture and ductile fracture at 750 °C, produced as a consequence of the presence of the sigma phase. The fracture surface of the samples tested at room temperature is ductile because of the presence of austenite and the -ferrite. Acknowledgements A part of the research described in this paper was conducted at the University of Ljubljana (Faculty of Natural Sciences and Engineering) as part of bilateral agreements between the Republic of Slovenia, and Bosnia and Herzegovina within the project "Application of new materials in the automotive industry" No: SLO-BA10-11-011. 5 REFERENCES 1 J. Janovec, B. [u{tar{i~, J. Medved, M. Jenko, Mater. Tehnol., 37 (2003) 6, 307–312 2 R. L. Plaut, C. Herrera, D. M. Escriba, P. R. Rios, A. F. Padilha, Materials Research, 10 (2007) 4 3 A. A. Astafev, L. I. Lepekhina, N. M. Batieva, https://springerlink. metapress.com/content/106486/?p=3803ac420c9f43dd80ca8a84319a 1436&pi=0, https://springerlink.metapress.com/content/ g625036m 8347/?p=3803ac420c9f43dd80ca8a84319a1436&pi=0 4 A. F. Padilha, P. R. Rios, ISIJ International, 42 (2002) 4, 325–337 5 H. S. Khatak, B. Raj, Corrosion of Austenitic Stainless Steels Mechanism, Mitigation and Monitoring, 2002 6 F. Tehovnik, F. Vodopivec, L. Kosec, M. Godec, Mater. Tehnol., 40 (2006) 4, 129–137 7 J. C. Tverberg, The Role of Alloying Elements on the Fabricability of Austenitic Stainless Steel, [cited:1. 4.2009.] Available from: www. csidesigns.com/tech/fabtech 8 J. Pilhagen, A Literature Review of the Stainless Steel 21-6-9 and its Potential for Sandwich Nozzles, Master Thesis, Lulea, 2007 9 F. Tehovnik, B. Arzen{ek, B. Arh, D. Skobir, B. Pirnar, B. @u`ek, Mater. Tehnol., 45 (2011) 4, 339–345 10 S. Ko`uh, M. Goji}, L. Kosec, RMZ-Materials and Geoenvironment, 54 (2007) 3, 331–344 11 ASTM Standard A276-96, 1997 Annual Book of ASTM Standards, section 1. Iron and steel products, vol.01.03 Steel-Plate, Sheet, Strip, Wire, Stainless Steel Bar, ASTM, 1997 12 A. Gigovi}-Geki}, M. Oru~, I. Vitez, Metallurgy, 50 (2011) 1, 21–24 13 K. H. Lo, C. H. Shek, J. K. L. Lai, Materials Science and Engineer- ing, R 65 (2009), 39–104 14 Metals Handbook: Properties and Selection: Iron, Steels and High-Performance 10th. Edition, Alloys, vol.1, ASM American Society for Metals, 1990 15 A. G. Geki}, M. Oru~, A. Nagode, H. Avdu{inovi}, RMZ-Materials and Geoenvironment, 58 (2011) 2, 121–128 16 M. Pohl, O. Storz, T. Glogowski, Materials Characterization, 58 (2007), 65–71 17 S. C. Kim, Z. Zhang, Y. Furuya, C. Y. Kang, J. H. Sung, Q. Q. Ni, Y. Watanabe, I. S. Kim, Materials Transactions, 46 (2005) 7, 1656–1662 18 M. Oru~, M. Rimac, O. Beganovi}, S. Muhamedagi}, Mater. Tehnol., 45 (2011) 5, 483–487 A. GIGOVI]-GEKI] et al.: EFFECT OF THE DELTA-FERRITE CONTENT ON THE TENSILE PROPERTIES ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 519–523 523 P. MARUSCHAK et al.: PHYSICAL REGULARITIES IN THE CRACKING OF NANOCOATINGS ... PHYSICAL REGULARITIES IN THE CRACKING OF NANOCOATINGS AND A METHOD FOR AN AUTOMATED DETERMINATION OF THE CRACK-NETWORK PARAMETERS FIZIKALNE ZAKONITOSTI POKANJA NANOPREVLEK IN METODA ZA AVTOMATSKO DOLO^EVANJE PARAMETROV MRE@E RAZPOK Pavlo Maruschak1, Vladimir Gliha2, Igor Konovalenko1, Toma` Vuherer2, Sergey Panin3 1Ternopil Ivan Pul’uj National Technical University, 46001 Ternopil, Ukraine 2University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, 2000 Maribor, Slovenia 3Institute of Strength Physics and Materials Science SB RAS, 634021 Tomsk, Russia maruschak.tu.edu@gmail.com Prejem rokopisa – received: 2011-10-20; sprejem za objavo – accepted for publication: 2012-04-05 The regularities and spatial distribution of multiple cracking of a nanocoating are investigated. It is found that in the cracking zones the relaxation of the stresses accumulated in the coating takes place; moreover, the intensity of its failure is determined by the structural level of defect accumulation. A new algorithm for a digital identification of the elements of a crack network in a nanocoating is proposed, and its adequacy is checked. Keywords: multiple cracks, nanocoating, damage, diagnostics, surface, strain Raziskovali smo zakonitosti in prostorsko porazdelitev ve~ razpok na nanoprevleki. Ugotovili smo, da pride na podro~jih z razpokami do relaksacije napetosti, nakopi~ene v prevleki. Intenzivnost njenega propadanja je dolo~ena s strukturnim nivojem nakopi~enja defektov. Predstavljen je nov algoritem za digitalno identifikacijo elementov mre`e razpok na nanoprevleki. Preverili smo ustreznost modela. Klju~ne besede: ve~ razpok, nanoprevleka, po{kodba, diagnoza, povr{ina, raztezek 1 INTRODUCTION The tensile, compressive and shear deformations play key roles in many processes of failure in a nanocoating because they are the results of a mutual influence of multiple defects and blocks of a coating1,2. The activation of deformation processes is observed in the sections of maximum coating cracking. Therefore, an important aspect is the quantitative evaluation of multiple cracking needed for the development of optimum algorithms for the identification and calculation of the parameters of a network of detected cracks2. However, there are a number of problems, in particular, a significant disper- sion of the sizes of the crack-like defects, the shape of the cracks, and the need for determining a degree of their coalescence. In the previous papers3,4, the authors offered a number of algorithms for evaluating the technical con- dition of the surface of a metallurgical equipment affected by multiple defects. However, the use of these methods for an investigation of a nanocoating requires a combination of material-science approaches, physical and mechanical approaches that will allow an analysis of the hierarchical structure of multiple cracking zones, bringing together the entire volume of information on the processes of initiation, growth and coalescence of defects into a single system, and identifying the main regularities in this process. In this paper, the structure of multiple cracking of a zirconium nanocoating, with a network of fatigue cracks, is investigated. 2 RESEARCH TECHNIQUE The main objective of the paper is to find the basic regularities in the cracking of a nanocoating and to analyse the possibility of using a digital processing method for identifying defects. The requirement to determine the main regularities in the formation of the geometrical structure of a multiple cracking of a zirconium nanocoating with a thickness of 100 μm predetermined the need to use the approaches of digital identification. Ion nanostructuring of the surface layer of the specimens made of steel 25Kh1M1F was performed using the high-current, vacuum-arc source of metallic ions on the UVN-0.2 "Quant" setup5. After reaching a vacuum of at least 3 · 10–3 Pa in the chamber, the specimens were treated with a zirconium-ion flow with the energy of 0.9–2.8 keV and the ion-current density of 0.1–0.3 mA/m2. The duration of the treatment was from 5 min to 20 min. The substrate holder with the speci- mens fixed on the specimen stage is connected directly Materiali in tehnologije / Materials and technology 46 (2012) 5, 525–529 525 UDK 620.3:621.793:620.191 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 46(5)525(2012) to the ion-acceleration schema instead of the traditional extraction of the radial-ion beam from the implanter. In this case, the acceleration of ions occurs in the dynamic self-organising boundary layer featuring a double electrical layer, which forms around the specimen surface under the negative potential5. Specimens from the ferrite-pearlite steel 25Kh1M1F in mass fractions6: w(C) = 0.23…0.29; w(Mn) = 0.40…0.70; w(Si) = 0.17…0.37; w(Mo) = 0.60…0.80; w(Ni) = 0.30; w(Cr) = 1.50…1.80; w(S)  0.025) with applied coatings were loaded, under cyclic tensioning, on the STM-100 test setup at the frequency of f = 1 Hz, max = 500 MPa, min = 0.1 max. After the failure of the speci- mens the coating condition was investigated at different values of the real transverse necking ~  of the specimens destroyed by tensioning7: ~ ln / ) F F 0 k (1) where F0 and Fk are the initial and the current areas of the cross-section, respectively. Based on the analysis of the photo images, the structural and morphological data about the cracks and the mechanisms of their formation in the zones of loca- lised tensioning and shear were analysed8. The measure- ment of the residual crack opening was determined by using the standard measurement of the Kappa software for TEM. The spatial orientation of the crack-network elements and their relation to the formation of the meso- and macroscale failure zones were determined. 3 STAGE-LIKE NATURE OF A NANOCOATING FAILURE It is known that multiple defects have a multilevel structural organisation. So, any section of a specimen is an aggregate of structural elements separated by cracks of various sizes, which are stress concentrators9,10. Failure of the coating material takes place in the vicinity of a number of defects simultaneously, which prevents the peeling of the coating. Fatigue failure of a nanocoating is a step-by-step process, taking place in the following successive steps: • formation of multiple defects – elastic-plastic defor- mation of a coating, nucleation of a network of indi- vidual cracks, their growth and partial coalescence followed by a gradual accumulation of local plastic strains in the most inhomogeneous sections of the material, Figure 1a. • formation of a longitudinal deformation relief – this phenomenon is connected with the macrolocalisation of the plastic strain in individual sections of a speci- men (Figure 1b), and shears of the grain conglo- merates in the substrate material (steel 25Kh1M1F). • activation of the transverse shears of the boundaries, formation of local microfragments of a coating – this process is preconditioned by a relative shear of the blocks of a coating due to the attainment of limiting values at the boundaries of the mesostructural ele- ments (blocks), Figures 1c, d. Further plastic deformation of the substrate material causes an intensified accumulation of plastic strains at the boundaries between the newly created blocks and the generation of damage6. At this stage, the opening of the crack-like defects is determined by the mutual effect of the elements of the block structure and the adhesive strength of the coating. Upon attainment of the ultimate state failure takes place, as well as the shear with the for- mation of a rupture, and the shear of the coating fragment. In Figure 2 the dependence of the residual crack opening () on the real necking of a specimen is pre- sented. With respect to this dependence, three sections are noticeable and they are reflected in the change of the slope angle of the curve. Within the first section, the P. MARUSCHAK et al.: PHYSICAL REGULARITIES IN THE CRACKING OF NANOCOATINGS ... 526 Materiali in tehnologije / Materials and technology 46 (2012) 5, 525–529 Figure 2: Stage-like nature of an increase in the mean value of a crack opening () with an increase in the real necking of the specimen ~  Slika 2: Stopni~asta narava srednje vrednosti odpiranja razpoke () z nara{~anjem resni~ne kontrakcije vzorca ~  Figure 1: Stage-like nature of failure of a zirconium nanocoating at various plastic strains of the coating steel ( ~  = 0; 0.23; 0.32; 0.35, respectively) Slika 1: Stopni~asta narava po{kodb cirkonijeve nanoprevleke pri raz- li~nih plasti~nih deformacijah jekla ( ~  = 0; 0,23; 0,32; 0,35 ) newly formed cracks of the fatigue origin are located perpendicularly to the direction of loading; they have an insignificant residual opening ( = 1.0…2.0 μm) and are separated from each other. So, the processes of the mutual effect of multiple defects are insignificant. The second section corresponds to the opening and growth of the already formed network of defects. The coating acquires the properties of a breakup-block medium by dividing itself into separate "islands" of the material surrounded by multiple cracks. The residual crack opening in this section varies within the range of  = 2.0…4.0 μm. The third section corresponds to the critical opening of the defects of  > 4.0 μm and a fragmentation of the coating. In this case, individual fragments of the coating are slipping relative to one another. In our opinion, local strains can be determined by using these fragments as the markers of displacements. 4 DISCUSSION OF IDENTIFICATION OF MULTIPLE CRACKS IN NANOCOATING The algorithm for identifying the crack positions consists of the following main steps: binarisation of the original grayscale image, its filtering, repeated binari- sation of the obtained image and its skeletonisation3,4. After the completion of the above operations we obtain an array of points that describe the geometrical para- meters of the network of cracks in the image. While setting the algorithm, several parameters are used, which are crucial to the correct and precise identification of a crack. The original image is a colour image obtained with an electron microscope (Figure 3). In order to simplify the analysis it is transformed into a black-and-white one by means of an adaptive binarisation, which allows us not only to identify potential objects (cracks), but also to eliminate the effect of inhomogeneous illumination during the image acquisition. The most important parameter of this stage of the algorithm, which has a significant effect on the final result of the identification of the geometrical crack parameters, is the background L edge. A wrong choice of the background edge leads to the situation when a part of an object to be identified is mistaken for the back- ground, or fictitious objects are found that are actually a part of the background. The "correctly chosen para- meter" allows a value of the background border, at which the number of the identified crack objects approximates, as much as possible, the number of the cracks available in the image that can be identified by the operator. The range of the optimal values for the said parameter was determined by an experimental method11. The crack identification algorithm and the background border effect on the final result are considered in greater detail in our previous papers8. A change in the background edge causes a displace- ment of the edges of the objects found in the image, due to which the geometrical characteristics of the cracks calculated as a result of the algorithm operation may vary a little. The result of the transformation is a monochrome image of the surface studied. It contains basic infor- mation about the location of potential objects sought; however, it also contains a large amount of noise ele- ments, which complicate the process of identification and, to a certain extent, distort the general view of the cracked surface. In order to reduce the effect of the noise elements (which are present after the binarisation of the original grayscale image) and join the close-set indivi- dual fragments of a crack into a solid object, the filtering was performed followed by a repeated binary transfor- mation of the filtered image. Filtering is carried out in two stages: first, a discrete gauss filter is used, next the obtained image is enhanced by applying the filter to increase the image contrast12. The gauss filter is effective in suppressing the noises and smearing the edges of the image objects, as a result of P. MARUSCHAK et al.: PHYSICAL REGULARITIES IN THE CRACKING OF NANOCOATINGS ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 525–529 527 Figure 4: Final image with detected cracks and a network of reference points Slika 4: Kon~na slika z odkritimi razpokami in mre`o referen~nih to~k Figure 3: Original grayscale image Slika 3: Originalen sivinski posnetek which the objects are formed that correspond to the crack positions. The contrast filter allows an increase in the contrast of the image elements. For the final distinguishing of a crack a repeated binarisation of the filtered image is used. Following all of the above transformations, the final image is obtained, describing the picture of a cracking on the surface analysed (Figure 4). In this picture, white pixels corre- spond to the background and black ones to the cracks. The most important parameter of the algorithm at this stage is the filter kernel size hF. It has an effect on the processes of "screening" the background pixels and coalescence of the separated fragments of the crack, therefore its change may significantly affect the final result of the identification. The image obtained as a result of the above transformations contains information about the shape and the area of the cracks; however, it cannot be used directly for determining the quantitative parameters, such as the number, the length, the slope of the cracks, etc. In order to approximate the aggregate of pixels, which form the detected objects/cracks, a skeletonisation is performed using the simplest array of points3,4. The skeletonisation allows us to distinguish the wireframe lines of the cracks by detecting medial lines with the thickness of one pixel in the image. Reference points of the medial lines are the final set of data, on the basis of which a conclusion is made about the crack location, direction of the crack propagation and its length12. Both of the algorithm parameters considered have an immediate effect on the position of the obtained wire- frame (medial) lines; therefore, an important issue is an assessment of their influence on the accuracy of deter- mining the identified cracks13. The L parameter is the background border of the binary transformation. The hF parameter is the filter kernel size. The corresponding corrections are made to the paper results. The back- ground border of the binary transformation is the gene- rally used parameter, which does not require any special clarification. The effect of these parameters was assessed by observing the algorithm operation with the variation of their value within a certain range. To this end, the arrays of the P point coordinates, which form the wireframe line, were fixed for each group of the parameters (hF and L). The obtained arrays of P points were used to deter- mine the displacements of the wireframe line pixels depending on the values of the algorithm parameters studied. It is clear that the other algorithm parameters have insignificant effects on the identification results. Therefore, in general, the total error of determining the position of the crack wireframe line , caused when setting the algorithm parameters, does not exceed the following value:   ≤ +L h (2) where L is the error in setting the background edge; h is the error caused by a displacement of the coordinate centre in setting the filter kernel size. With a view to assessing the effect of the variation of the above algorithm parameters on the displacement of the wireframe line, the relative background edge L was varied within 10…30 %, and the filter kernel size hF varied from 5 to 10 pixels. In this case, the array of Pμ (L, hF) points was fixed for every set of parameter values individually, where μ  (1...M), and M is the number of observations. In order to calculate the displacement of the wireframe line, the closest point of the Pμ2 (L, hF) array was found for every Pμ1 (L, hF) point, and the distance between them was calculated. As a result of the above calculations, the aggregate of sets was obtained, which characterises the displacement of the points of wireframe line μ (L, hF). Each μi element equals the dis- tance between point Pμ1i and point Pμ2j. To make the sample homogeneous by removing errors from it, the sample was censored using Wright’s criterion. In this case, the values, for which the   i S− ≥ 3 condition was fulfilled (where  is the mean value and S  is the standard deviation of the sample), were removed from the μ (L, hF) aggregate of values. A typical view of the μ (L, hF) displacement-distri- bution histogram for the image (Figure 3) at the variation of the filter kernel size from 5 to 7 pixels and the constant background edge of 20 % is shown in Figure 5. It is clear that the majority of the points shift very insignificantly (only by a few pixels). The investigation of a series of images showed that, at the constant background edge of L = 20 % and the variation of the filter kernel hF from 5 to 10 pixels, the standard deviation of the sample varies within the range of 1.7–2.2 pixels. When recalculated in the units of the real object length, the displacement varies within the range of 0.20…0.26 μm. The maximum displacement fixed for various measurements varied from 0.18 μm to 0.29 μm. Moreover, an increase in the kernel size proportionally influences the displacement of the crack wireframe lines. P. MARUSCHAK et al.: PHYSICAL REGULARITIES IN THE CRACKING OF NANOCOATINGS ... 528 Materiali in tehnologije / Materials and technology 46 (2012) 5, 525–529 Figure 5: Displacement-distribution diagram for the points of the wireframe line Slika 5: Diagram premika to~k v liniji mre`e At the constant kernel size of hF = 5 pixels and the variation of L from 5 % to 20 %, the standard deviation varies within the range of 1.14…1.95 pixels (which corresponds to 0.13…0.23 μm on the test specimen). In order to reveal the effect of the changes of any of the algorithm parameters studied on the final result – the geometrical parameters of the crack network – a number of investigations were carried out, during which the background edge varied within the range of 5…20 %, and the filter kernel size varied from 5 to 10 pixels. Figure 6 shows the crack-length-distribution graphs for the image (Figure 3) at the limiting values of the background edge and the filter kernel size. The results obtained show that, for different images within the range of the algorithm parameters investigated, the deviation of the length from the mean value is within 0–0.30 μm. With the confidence probability of 95 % this deviation does not exceed 0.26 μm. 5 CONCLUSIONS On the basis of the results of the investigations into the cracking processes in the surface of the heat-resistant steel with a nanocoating, it is established that multiple defects are caused by a plastic deformation of the substrate material. Multiple defects are nucleated in the strain-localisation areas and in the zones of microstruc- tural inhomogeneity. The new algorithm of digital defectometry has been developed, which is intended for the identification of the crack-network elements in a nanocoating. The effect of the main algorithm parameters on the measurement results relating to the geometrical characteristics – the coordinates, the lengths and the slope angles of the cracks – is investigated. The results obtained allow a compilation of such a combination of the algorithm parameters, with which the general measurement error of the cracking parameters will be minimal. It was found that the optimum parameters of the algorithm for identifying the cracks have the following limits: hF = 5…7 pixels and L = 8–14 %. With the confidence probability of 95 % the deviation of length is less than 0.30 μm. 6 REFERENCES 1 V. E. Panin (Ed.), Physical mesomechanics of heterogeneous media and computer-aided design of materials, Cambridge Interscience Publishing, Cambridge 1998, 339 2 P. Z. Iordache, R. M. Lungu, G. Epure, et al., J. of Optoelectron. Adv. Mater., 28 (2011), 550 3 P. Yasniy, P. Maruschak, I. Konovalenko, V. Gliha, T. Vuherer, R. Bishchak, Multiple cracks on continuous caster rolls surface: A three-dimensional view, Proc. of the 4th Int. conf. Processing and Structure of Materials (May 27–29), Pali}, Serbia, 2010, 7–12 4 P. Yasnii, P. Maruschak, I. Konovalenko, R. Bishchak, Materials Science, 46 (2008), 833 5 O. V. Sergeev, M. V. Fedorischeva, V. P. Sergeev, N. A. Popova, E. V. Kozlov, Increase of plasticity of maraging steels by means of ion beam nanostructuring of surface layer, Proc. of the 10th Int. Conf. on Modification of Materials with Particle Beams and Plasma Flows, Tomsk, (September 19–24), 2010, 342 6 P. Yasniy, P. Maruschak, R. Bishchak, V. Hlado, A. Pylypenko, Theoretical and Applied Fracture Mechanics, 52 (2009), 22 7 P. Yasnii, P. Maruschak, V. Hlado, D. Baran, Materials Science, 46 (2008), 144 8 I. V. Konovalenko, P. O. Maruschak, Optoelectronics, Instrumenta- tion and Data Processing, 47 (2011), 49 9 P. O. Maruschak, S. V. Panin, S. R. Ignatovich et al., Theoretical and Applied Fracture Mechanics, 57 (2012), 43 10 P. Yasniy, I. Konovalenko, P. Maruschak, Investigation into the geometrical parameters of a thermal fatigue crack pattern, WSEAS Int. Conference New aspects of engineering mechanics, structures and engineering geology, Heraklion, Crete Island, Greece, 2008, 61–66 11 I. Konovalenko, P. Maruschak, Computer analysis of digital images with quasiperiodical structure, Proc. of the Int. conf. TCSET 2012 – Modern Problems of Radio Engineering, Telecommunications and Computer Science, Lviv-Slavske, (February 21–24), 2012, 419 12 P. Yasniy, P. Maruschak, I. Konovalenko, R. Bishchak, Mechanika, 17 (2011) 3, 251 13 W. K. Pratt, Digital image processing (4th Ed.), Wiley, 2007, 807 P. MARUSCHAK et al.: PHYSICAL REGULARITIES IN THE CRACKING OF NANOCOATINGS ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 525–529 529 Figure 6: Graphs of detected crack lengths  for the limiting values of the algorithm parameters Slika 6: Graf dol`ine odkritih razpok  za omejene vrednosti para- metrov algoritma M. OSTRÝ et al.: LABORATORY ASSESSMENT OF MICRO-ENCAPSULATED PHASE-CHANGE MATERIALS LABORATORY ASSESSMENT OF MICRO-ENCAPSULATED PHASE-CHANGE MATERIALS LABORATORIJSKA OCENA MIKROENKAPSULIRANIH MATERIALOV S FAZNO PREMENO Milan Ostrý1, Radek Pøikryl2, Pavel Charvát3, Tomá{ Ml~och2, Barbora Bakajová2 1Brno University of Technology, Faculty of Civil Engineering, Veveøí 95, 602 00 Brno, Czech Republic 2Brno University of Technology, Faculty of Chemistry, Purkyòova 464/188, 612 00, Brno, Czech Republic 3Brno University of Technology, Faculty of Mechanical Engineering, Technická 2, 616 69, Brno, Czech Republic ostry.m@fce.vutbr.cz Prejem rokopisa – received: 2011-10-20; sprejem za objavo – accepted for publication: 2012-05-07 The operation of low-energy-consumption and passive houses can be based on the passive or active utilization of renewable energy sources. Thermal energy storage plays a key role in the application of renewable energy sources and it thus contributes to the reduction of global CO2 emissions. Thermal energy storage is commonly based on the sensible- or latent-heat-storage techniques. Latent-heat thermal storage is based on the absorption or release of heat when a storage material is changing phase. The thermal-storage materials suitable for latent-heat storage are called Phase-Change Materials (PCMs). PCMs have considerably higher thermal-energy-storage densities than the sensible-heat-storage materials and they are able to absorb large quantities of energy in a small range of temperatures during the phase change. Nowadays, micro-encapsulation of phase-change materials is one of the promising approaches in the integration of latent-heat storage in various applications. The most important properties of a latent-heat-storage medium are the heat of the fusion and the temperature range of the phase change. The correct determination of the physical and chemical properties is essential for a practical use of phase-change materials. The paper deals with the results of a laboratory assessment of the selected micro-encapsulated PCMs and shows a practical example of a possible integration in the building structures. Keywords: phase-change materials, latent heat, differential scanning calorimetry, micro-encapsulation Delovanje nizkoenergijskih pasivnih hi{ lahko temelji na uporabi aktivne ali pasivne uporabe obnovljivih energijskih virov. Shranjevanje toplotne energije igra klju~no vlogo pri uporabi obnovljivih virov energije in zato vpliva na zmanj{anje globalne emisije CO2. Shranjevanje toplotne energije navadno temelji na smiselnih tehnikah ali tehnikah shranjevanja latentne toplote. Shranjevanje latentne toplote temelji na absorpciji ali spro{~anju toplote, ko material za shranjevanje spreminja fazo. Materiale, ki so primerni za shranjevanje latentne toplote, imenujemo Materiali s fazno premeno (PCM). PCM imajo ob~utno vi{jo gostoto shranjenje toplotne energije v primerjavi z ob~utljivimi materiali za shranjevanje toplote in so sposobni absorbirati med fazno premeno ve~jo koli~ino energije v manj{em temperaturnem intervalu. Dandanes je vgradnja mikroenkapsulacijskih materialov s fazno premeno obetajo~a za {tevilne mo`nosti uporabe. Najpomembnej{i lastnosti medija za shranjevanje latentne toplote sta talilna toplota in temperaturno podro~je fazne premene. Pravilno dolo~anje fizikalnih in kemijskih lastnosti je bistveno za prakti~no uporabo materialov s fazno premeno. Ta ~lanek obravnava laboratorijsko oceno izbranih PCM in ka`e prakti~en primer mo`nosti njihove uporabe v gradbeni{tvu. Klju~ne besede: materiali s fazno premeno, latentna toplota, diferen~na dinami~na kalorimetrija, mikroenkapsulacija 1 INTRODUCTION Sensible-heat storage utilizes the heat capacity and the change in the temperature of a thermal-storage material during the process of charging and discharging the heat1. The amount of stored heat depends on the specific heat of the storage material, the temperature difference and the amount (mass) of the material. Any building material can generally be used for sensible-heat thermal storage but the materials with high specific heat and high density usually perform the best. The typical representatives of sensible-heat-storage materials are common building materials such as ceramic bricks or blocs, concrete, lime-cement bricks and stone. The indoor environments with the envelopes made of such materials exhibit a much higher degree of thermal stability than the light-weight envelopes (e.g., timber- frame walls). The thermal storage in common building structures has its limits. The first limit is the use of heavy-weight structures. This is a very important con- straint, especially in modern buildings. For example, glass-building envelopes would need to be supplemented with heavy-weight indoor structures and that is not always possible or desirable. This is where latent-heat storage can be employed. 1.1 Latent-heat storage Latent-heat storage is based on the absorption or release of heat when a storage material undergoes a phase change from solid to liquid1. Such thermal-storage materials are called Phase-Change Materials (PCMs). They use chemical bonds to store and release heat2. PCMs have a high ability to store thermal energy. PCMs are able to absorb large quantities of heat in a small range of temperatures during a phase change. Latent-heat storage is one of the most efficient ways of storing thermal energy3. The selection of a PCM is mainly based on its melting temperature. A PCM’s melting tempera- ture should be within the operating temperature range of Materiali in tehnologije / Materials and technology 46 (2012) 5, 531–534 531 UDK 536.65:536.42 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 46(5)531(2012) the thermal system. With respect to building use, it means within the thermal-comfort temperature range of an occupied space. 1.2 Selection of phase-change materials Phase-change materials can be chosen from both organic and inorganic materials. The organic phase- change materials melt and freeze repeatedly without a phase-change segregation and crystallize with little or no supercooling. The organic phase-change materials, e.g., the paraffins, are compatible with metals without any risk of corrosion. The paraffins have a rather poor ther- mal conductivity and they are flammable. The melting point of the alkanes increases with an increased number of carbon atoms1. The inorganic PCMs are compatible with plastics and their storage capacity is higher than the capacity of the organic PCMs due to their higher density. The inorganic PCMs, e.g., salt hydrates, are incompatible with uncoated metals. The salt hydrates are important PCMs because of the high heat of fusion and a small volume change during the process of melting and solidification. The main disadvantages of salt hydrates are their poor nucleating properties that result in supercooling. Suitable PCMs from both organic and inorganic groups are available for applications in the latent-heat- storage technology. Many phase-change materials cannot be used as latent-heat-storage mediums because of the problems with their chemical stability, toxicity, corrosion, volume change and price. The phase-change materials should meet the following thermodynamic, kinetic, chemical and economic criteria4. Thermodynamic criteria: • high heat of fusion; • melting range in the desired operating-temperature range; • high specific heat; • high thermal conductivity; • high density and low volume change; • congruent melting. Kinetic criteria: • little or no supercooling during the solidification process; • sufficient crystallization rate. Chemical criteria: • compatibility with the container; • long-term chemical stability; • no toxicity; • no flammability. Economic criteria: • availability in the required quantities; • low cost. 2 MATERIALS AND METHODS The research and development at the Brno University of Technology is focused on the utilization of the latent-heat storage in passive and active solar-heating and cooling technologies. The development of the advanced latent-heat-storage technologies is strongly dependent on the possibility to find suitable phase- change materials that fulfill the above-mentioned requirements. The second problem lies in finding a suitable technology for an integration of latent-heat- storage media in building structures. Micro-encapsulation is one of the possible approaches to the PCM integration in building struc- tures5. 2.1 Micro-encapsulated PCMs Micro-encapsulation is based on enclosing a PCM in a very small capsule. The micro-capsules can be included in the common building materials and structures. Special attention has to be paid to the choice of the material of the capsule to avoid a chemical reaction between the capsules and the building material5. Micro-capsules can be added to the composition of lime or gypsum plaster, concrete, fibrous wooden slabs and gypsum wall boards. A special composition of gypsum plaster and micro-encapsulated PCMs was developed for the application in building structures. The gypsum plaster contains 30 % of micro-capsules Micronal DS 5008 X. The plaster is the final layer of the walls and ceilings in the buildings with a low thermal mass. 2.2 Differential scanning calorimetry Differential Scanning Calorimetry (DSC) is a thermo-analytical technique where a temperature range is scanned. The difference between the amount of the heat required for changing the temperature of a sample and its reference is measured as a function of temperature. Both the sample and the reference are maintained at nearly the same temperature throughout the experiment. The reference is used to determine the heat stored in the sample by considering the difference between the signal of the sample and the reference6. The DSC heat flux has got a Siamese structure5. The sample and the reference are connected to the same metal disc. The behavior difference between the sample and the reference submitted to the same temperature excitation leads to a voltage difference between the sample and the reference. The absorbed heat in the PCM sample is deduced from the voltage5. The weight of the sample is only a few grams. A calorimeter PYRIS1 Perkin Elmer was used in the tests. M. OSTRÝ et al.: LABORATORY ASSESSMENT OF MICRO-ENCAPSULATED PHASE-CHANGE MATERIALS 532 Materiali in tehnologije / Materials and technology 46 (2012) 5, 531–534 3 RESULTS AND DISCUSSION The experiments focused on determining the thermal properties of a micro-encapsulated PCM and the plaster containing 30 % of a PCM. Figure 1 shows the results after two heating/cooling cycles for the micro-capsules with a PCM. The results were obtained with the continuous scanning mode at the rate of 1 °C/min. This rate is rather quick compared to common conditions in rooms. The black and blue curves represent heating, while the red and green curves show the results of the cooling mode. There is some super- cooling in the cooling phase that is not really proble- matic for a practical application in building structures. The presence of supercooling plays a role in discharging the heat stored in the PCM. As can be seen, the PCM displays two heat-flow peaks. The first peak is well below the comfortable indoor air temperature, thus, it has no implications for the practical use. Table 1 shows the peak temperatures for both cycles. The difference between the peaks of the cooling and heating phases is about 2 °C. Table 2 shows the heat of fusions for the heating and cooling phases. Table 1: Peak temperatures for the PCM Tabela 1: Maksimalne temperature za PCM Cycle Peak temperature T/°C cooling heating I 22.32 24.32 II 22.40 24.20 Table 2: Heat of fusions for the PCM Tabela 2: Talilna toplota za PCM Cycle Heat of fusion in J/g cooling heating I –84.16 86.83 II –81.29 86.76 Thermal stability of PCMs was determined with a thermogravimetric apparatus (TGA) Q500 TA Instru- ments. The results from TGA are shown in Figure 2. Significant weight losses start at 140 °C, which is above the commonly used temperature range. The temperature range of the indoor climate between 15 °C and 35 °C can be assumed for building applications. Figure 3 shows DSC results for a gypsum plaster with 30 % of Micronal D5008 X. The chart shows a significant reduction of the latent heat during the heating and cooling processes. Three temperature rates of (1, 10 and 20) °C/min were tested. As can be seen, the onset and the peak tempe- ratures during the heating and cooling strongly depend on the temperature ramp. Melting temperatures rise with an increased heating rate. A shift of the solidifi- cation-temperature range follows an increased rate of cooling. The risk of supercooling increases with a faster cooling rate. 4 CONCLUSION The results obtained with DSC confirm suitability of the tested PCMs and the plaster for their integration in building structures. The peak temperature during the M. OSTRÝ et al.: LABORATORY ASSESSMENT OF MICRO-ENCAPSULATED PHASE-CHANGE MATERIALS Materiali in tehnologije / Materials and technology 46 (2012) 5, 531–534 533 Figure 3: DSC results for gypsum plaster Slika 3: Rezultati DSC za mav~ni omet Figure 1: DSC results for the PCM Slika 1: Rezultati DSC za PCM Figure 2: Thermal stability of the PCM Slika 2: Toplotna stabilnost PCM heating is about 24 °C. The micro-encapsulated PCMs have a required melting range for the thermal-energy storage during the summer season. The difference between the peak-melting and solidification tempera- tures of a PCM is about 2 °C allowing a proper, natural or driven, regeneration of the heat-storage medium at night. Acknowledgement This work was supported by the Czech Science Foundation under contract No. P101/11/1047 and by the Czech Science Foundation under contract No. P104/12/1838 "Utilization of latent heat storage in phase change materials to reduce primary energy consumption in buildings". 5 REFERENCES 1 A. Sharma, V. V. Tyagi, C. R. Chen, D. Buddhi, Review on thermal energy storage with phase change materials and applications, Renewable and Sustainable Energy Reviews, 13 (2009), 318–345 2 V. V. Tyagi, D. Buddhi, PCM thermal storage in buildings: A state of art, Renewable and Sustainable Energy Reviews, 11 (2007), 1146–1166 3 M. M. Farid, A. M. Khudhair, S. A. K. Razack, S. Al-Hallaj, A review on phase change energy storage: materials and applications, Energy Conversion and Management, 45 (2004), 1597–1615 4 H. P. Garg, S. C. Mullick, A. K. Bhargava, Solar Thermal Energy Storage, D. Reidel Publishing Company, 1985, 642 5 F. Kuznik, D. David, K. Johannes, J. J. Roux, A review on phase change materials integrated in building walls, Renewable and Sustainable Energy Reviews, 15 (2011) 1, 379–391 6 H. Mehling, L. F. Cabeza, Heat and cold storage with PCM. An up to date introduction into basics and applications, Springer-Verlag, Berlin Heidelberg 2008, 308 M. OSTRÝ et al.: LABORATORY ASSESSMENT OF MICRO-ENCAPSULATED PHASE-CHANGE MATERIALS 534 Materiali in tehnologije / Materials and technology 46 (2012) 5, 531–534 R. BEGI] et al.: CONTENT OF Cr AND Cr (VI) IN A WELDING FUME ... CONTENT OF Cr AND Cr (VI) IN A WELDING FUME BY DIFFERENT Cr CONTENT IN AN EXPERIMENTAL COATING OF A Cr-Ni RUTILE ELECTRODE VSEBNOST Cr IN Cr (VI) V VARILNEM DIMU PRI RAZLI^NI VSEBNOSTI Cr V PLA[^U RUTILNE ELEKTRODE Cr-Ni Razija Begi}1, Monika Jenko2, Matja` Godec2, ^rtomir Donik2 1Faculty of Engineering, University of Biha}, Irfana Ljubijanki}a bb., 77000 Biha}, Bosnia and Herzegovina 2Institute of Metals and Technology, Lepi pot 11, Ljubljana, Slovenia razijabegic@yahoo.co.uk Prejem rokopisa – received: 2011-11-17; sprejem za objavo – accepted for publication: 2012-06-01 In the SMAW welding process welding fumes are generated, harmful to human health and the environment. A welding fume is a mixture of gaseous and solid phases, which are generated during most of the electric-arc-welding processes. This article presents the researches of the particles that constitute the solid-phase-welding fumes. The change in the chemical composition of an electrode and its components (the coating and the core) can affect the chemical composition of the particles in welding fumes. The largest amount of fumes, about 80 %, is generated from electrodes and, accordingly, the focus of research was the influence of the chemical composition of an electrode on the chemical composition of the welding-fume particles. This paper presents the research results obtained for the content of Cr and Cr (VI) oxide particles in welding fumes. The experimental work on six variants of commercial electrodes E 23 12 2 LR 12, a welding chamber collecting fume particles and a chemical analysis of the particles were applied according to the standard EN15011. The aim was to determine an experimental welding electrode that should generate the welding-fume particles with the lowest content of Cr and Cr (VI) oxide. Keywords: health, welding fumes, particle, coated electrodes, Cr (VI) Med varilnim procesom SMAW nastaja dim, ki je {kodljiv za zdravje in okolje. Dim, ki nastaja pri ve~ini varilnih procesov, je me{anica plinov in trdnih delcev. Ta ~lanek opisuje preiskavo delcev, ki so trdni del dima, ki nastane pri varjenju. Spreminjanje sestave elektrode (str`ena in obloge) vpliva na kemijsko sestavo trdnih delcev v dimu. Najve~ji dele` dima, okrog 80 %, izvira iz elektrode, zato je bila raziskava osredinjena na u~inek kemijske sestave elektrode na kemijsko sestavo delcev v dimu. Ta ~lanek navaja rezultate raziskav vsebnosti oksidnih delcev Cr in Cr (VI) v dimu pri varjenju. Eksperimentalno delo je bilo izvr{eno s {estimi razli~nimi komercialnimi elektrodami E 23 12 2 LR 12. V varilni komori zbrani delci iz dima so bili analizirani skladno s standardom EN15011. Namen je bil ugotoviti eksperimentalno elektrodo, ki proizvaja med varjenjem delce z najmanj{o vsebnostjo Cr v Cr (VI)-oksidu. Klju~ne besede: zdravje, varilni dim, delci, opla{~ene elektrode, Cr (VI) 1 INTRODUCTION No material of any source can be directly compared with the composition and structure of a welding vapour. Chromium is generated in flue gases of welding with coated high-alloyed Cr electrodes and it appears in several phases, of which the six-valent oxide of chro- mium, Cr (VI), is the most damaging. Epidemiological studies prove the Cr (VI) compounds to be occupational carcinogens. During the MAG stainless-steel welding much less Cr (VI) is generated than during SMAW. Cr (III) compounds are biologically inert because they do not enter the cell, while Cr (VI) causes cell mutation. Chromium has a low threshold limit value ( TLV), which is 0.5 mg/m3.1 2 EXPERIMENT Experimental electrodes were made according to the experimental plan shown in Table 1, based on the changes in the chromium content in an electrode and its components (a wire2 and an electrode coating), marked with labels A, B, C, D, E and F, representing six varieties of commercial electrodes E 23 12 2 LR 12. Table 1: Change in the contents of Cr in an electrode and its compo- nents4 Tabela 1: Spreminjanje vsebnosti Cr v elektrodi in v pla{~u elektrode4 No Electrode Cr content in an electrode wire2 Cr content in an electrode coating Mean Cr content in an electrode 1. A 18.2 % 20.8 % 19.3 % 2. E 22.8 % 20.1 % 3. C 29.4 % 22.8 % 4. B 19.6 % 18.1 % 19.0 % 5. F 20.0 % 19.8 % 6. D 27.4 % 22.7 % For the tests related to emissions and their qualitative and quantitative chemical analyses it is necessary to have the appropriate equipment, which primarily consists of a collecting chamber, made for the purpose of this research according to the model in the standard EN150113. The chemical composition of the welding-fume particles was obtained with the tests for six experimental electrodes, a Materiali in tehnologije / Materials and technology 46 (2012) 5, 535–537 535 UDK 621.791:621.791.04 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 46(5)535(2012) total of 18 probes. The current level was constant, I = 95 A, and the basic material was a low-carbon, unalloyed structural steel S235JRG2. 3 RESULTS The content of chromium in the welding-fume particles is shown in Table 2, determined with the AAS method. Table 2: Results of a chemical analysis of the Cr content in the fume particles4 Tabela 2: Rezultati kemijske analize vsebnosti Cr v delcih dima4 No Electrode Probe Cr content in the welding-fume particles, % Cr content Cr mean value 1 A A1 5.62 5.912 A2 6.00 3 A3 6.10 4 B B1 5.33 5.095 B2 5.20 6 B3 4.73 7 C C1 6.00 5.588 C2 5.84 9 C3 4.90 10 D D1 5.05 5.1811 D2 5.40 12 D3 5.05 13 E E1 5.10 5.0614 E2 4.95 15 E3 5.10 16 F F1 5.30 5.0517 F2 5.00 18 F3 4.85 SEM-EDS and XPS chemical analyses of fume particles were carried out at the Institute of Metal Materials and Technology, Ljubljana. As an addition to the analysis of Cr and Cr (VI) in welding fumes and particles, the performed chemical analyses also included the contents of Mo, Mn and Ni as the most influential alloying elements and other elements and compounds. The change in the content of Cr particles in welding fumes shown in Figure 1 depends on the increase in the Cr content in the lining of an experimental electrode. The functional dependence of the Cr content in welding fumes and the Cr content in the electrode coating in Figure 1 corresponds to the exponential equation: CrZD Crcoating = + ⋅ −5 05 174 10 8 1 7. ( . ) . (1) Reliability of the calculated functional dependence is relatively high with R2 = 0.98. A graphical representation of functional dependen- cies of the contents of Cr (VI) particles in welding fumes on the Cr content in electrodes and in electrode coatings is shown in Figure 2. Figure 2 corresponds to the exponential equation: Cr(VI) ZD Crcoating = + ⋅ −4 74 6 2 10 7 2 16. ( . ) . (2) The reliability of functional dependence is relatively high being R2 = 0.97. 4 DISCUSSION The test conditions for electrode A were different than for the other electrodes and for this reason electrode A was excluded from the further analysis. The shape of the curve in Figure 1 shows that a 23–24 % addition of Cr to the electrode coating causes no significant increase in the Cr content in welding fumes. However, if this amount is increased the content of Cr particles in weld- ing fumes has a much stronger growth trend. Similarly, from Figure 2 it can be concluded that no significant increase in the content of Cr (VI) particles in welding R. BEGI] et al.: CONTENT OF Cr AND Cr (VI) IN A WELDING FUME ... 536 Materiali in tehnologije / Materials and technology 46 (2012) 5, 535–537 Figure 2: Content of Cr (VI) in the fume particles depending on the Cr content in the electrode coating4 Slika 2: Vsebnost Cr (VI) v delcih dima v odvisnosti od vsebnosti Cr v pla{~u elektrode4 Figure 1: Content of Cr particles depending on the Cr content in the electrode coating4 Slika 1: Vsebnost Cr v delcih v odvisnosti od vsebnosti Cr v pla{~u elektrode4 fumes occurs due to a concentration of Cr in the electrodes above 24–25 % Cr. 5 CONCLUSION The problems of welding fumes are becoming associated with harmful emissions that increasingly affect the protection of people and environment. Throughout the world we have seen an increased use of welding needed for joining the structures that are more and more being made of alloyed steel. Therefore, the production of the welding smoke is larger and the increasing use of the high-alloyed electrodes and the resulting chemical compositions of welding fumes are becoming more harmful. Any reduction of the harmful emissions to the atmosphere increases the protection of the people and environment. This paper explores this issue and the results obtained can be applied to the development of the electrodes for SMAW with a lower level of harmful emissions allowing a satisfactory quality of a weld. Two basic requirements needed for a SMAW process and for the coated electrodes are examined. After qualitative and quantitative analyses of the particles in the welding fumes a feedback loop can be introduced, based on the formation of solid particles, helping us to make decisions on the introduction of a new alloy coating on the electrodes that can lower the amount of harmful components in the welding-fume particles.4 6 REFERENCES 1 V. E. Spiegel-Ciobanu, Von "Schweißrauche" zu "Schweißtechnische Arbeiten", Die neuen technischen Regeln für Gefahrstoffe, Hannover, TRGS 528 TÜ Bd.50, 2009, Nr. 9 2 Rodacciai, Certificato di collaudo, Italia, May 2008 3 EN ISO 15011-1, Health and safety in welding and allied processes-Laboratory method for sampling fume and gases generated by arc welding-Part 1: Determination of emission rate and sampling for analysis of particulate fume, April 2002 4 R. Begi}, Doctoral dissertation, Exploring optimal technological composition of electrode coatings in terms of minimizing welding fumes, September 2011, Faculty of Engineering, University of Biha} R. BEGI] et al.: CONTENT OF Cr AND Cr (VI) IN A WELDING FUME ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 535–537 537 M. MARGHSI, D. BENACHOUR: USE OF A TWO-DIMENSIONAL PSEUDO-HOMOGENEOUS MODEL ... USE OF A TWO-DIMENSIONAL PSEUDO-HOMOGENEOUS MODEL FOR THE STUDY OF TEMPERATURE AND CONVERSION PROFILES DURING A POLYMERIZATION REACTION IN A TUBULAR CHEMICAL REACTOR UPORABA DVODIMENZIONALNEGA PSEVDOHOMOGENEGA MODELA ZA [TUDIJ TEMPERATURE IN PROFILA PRETVORBE MED REAKCIJO POLIMERIZACIJE V CEVASTEM KEMIJSKEM REAKTORJU Mohamed Marghsi, Djafer Benachour Laboratory of Preparations, Modifications and Applications of Multiphase Polymeric Materials (LMPMP), Ferhat Abbas University, 19 000 Setif, Algeria mmarghsi@yahoo.fr Prejem rokopisa – received: 2012-01-18; sprejem za objavo – accepted for publication: 2012-03-06 A two-dimensional pseudo-homogeneous model is used to study temperature and conversion profiles during the polymerization reaction of low-density polyethylene (LDPE) in a tubular chemical reactor. This model is integrated with the Runge-Kutta 4th-order semi-implicit method, using orthogonal collocation to transform a system of complex equations into the ordinary differential ones, with respect to the heat and mass transfers involved. Ethylene polymerization has been simulated over a range of temperatures and pressures and according to the mechanisms of radical polymerization. The results of several tests, carried out under the conditions similar to those of an industrial-scale polymerization, are presented. The influences of the initial temperature T0, the total pressure Pt and the ratio L/D (the main dimensions of the reactor) on the profiles of the temperature and conversion rates are tested and analyzed to predict the behavior and performance of the tubular chemical reactor considered. The focus was on the effect of an increase in the initial temperature T0 since such a rise results in a decrease in Tc (hot spot) appearing at the entrance of the reactor on the one hand, and in an improved conversion on the other hand. An opposite effect is observed for Pt since a pressure increase will result in a rapid rise in Tc and a decrease in the conversion. The ranges of pressures and temperatures are thus limited by the system performance: excessive pressures must be avoided and working temperatures must be chosen in the range where the polymerization reaction is very fast; such conditions allow not only a good conversion, but also a resulting polymer with a low crystallinity and, thus, a low density. In the present work the effect of the L/D ratio was also studied in order to find the most suitable ratio that permits the best evacuation of the heat released during the polymerization. Keywords: modeling, tubular reactor, simulation, low-density polyethylene, pseudo-homogeneous two-dimensional model Dvodimenzijski psevdohomogeni model je bil uporabljen za {tudij temperature in profila pretvorbe med reakcijo polimerizacije polietilena z nizko gostoto (LDPE) v cevastem kemijskem reaktorju. V model je bila vklju~ena Runge-Kuttova semiimplicitna metoda 4. reda z uporabo ortogonalne kolokacije za pretvorbo sistema kompleksnih ena~b v navadne diferencialne ena~be glede na vklju~en prenos toplote in mase. Simulirana je bila polimerizacija etilena v {ir{em podro~ju temperature in tlaka skladno z mehanizmom radikalne polimerizacije. Predstavljenih je ve~ preizkusov polimerizacije, izvedenih v razmerah, podobnih industrijskim. Preizku{en in analiziran je bil vpliv za~etne temperature T0, celotnega tlaka Pt in razmerja L/D (glavne dimenzije reaktorja) na profil temperature in hitrost pretvorbe, da bi bilo mogo~e napovedati pona{anje in zmogljivost uporabljenega cevastega reaktorja. Pozornost je bila usmerjena na u~inek povi{anja za~etne temperature T0, ker to po eni strani vpliva na zni`anje Tc (vro~a to~ka) na vstopu v reaktor, po drugi pa na izbolj{anje pretvorbe. Nasproten u~inek je bil opa`en za Pt, ker se narastek tlaka izra`a v hitrem povi{anju Tc in zmanj{anju konverzije. Obmo~je tlaka in temperature je torej omejeno z zmogljivostmi sistema: treba se je izogibati prekomernemu tlaku, delovne temperature pa je treba izbrati v obmo~ju, kjer je reakcija polimerizacije zelo hitra; take razmere omogo~ajo dobro konverzijo, in nastali polimer ima majhno kristalini~nost in s tem nizko gostoto. V tem delu je bilo preu~evano tudi razmerje L/D, da bi dobili najbolj primerno razmerje, ki omogo~a najbolj{i odvod toplote, ki se spro{~a med polimerizacijo. Klju~ne besede: modeliranje, cevast reaktor, simulacija, polietilen nizke gostote, psevdohomogen dvodimenzionalni model Materiali in tehnologije / Materials and technology 46 (2012) 5, 539–546 539 UDK 66.095.26:678.742.2 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 46(5)539(2012) 1 INTRODUCTION For many years the production in the chemical industry has been based solely on experience. However, for economic reasons and to avoid extreme conditions of temperature and pressure, the use of simulation methods has become more and more necessary. Indeed, through mathematical models, it is possible to predict relation- ships between the variations of experimental or produc- tion parameters and the practice results. The polymerization of some unsaturated hydro- carbons, called olefins, is extremely important. The types of polymers obtained consist mainly of low- (LDPE), and high-density polyethylene (HDPE) and polypropy- lene (PP). LDPE, one of the most frequently produced engineering polymers in the world, is a versatile polymer with a very wide range of applications1. Therefore, its polymerization process has been the subject of numerous research works and many studies are still in progress in order to obtain improved LDPE-based materials. In this regard, a better understanding of the polymerization process requires the use of a simulation of the reaction mechanism in order to establish new procedures to reach a better conversion (which is currently 15 % to 35 % in most industrial cases) and to improve the reactor performance based on a model proposed for a tubular chemical reactor. The reactor operating conditions are often difficult to fix. They should ideally be based on the simulation calculations. In other words, an important task is to determine the optimal operating conditions relating to temperature, pressure and conversion for each point of the reactor and this must be done in a rational manner. Several mathematical models are available for che- mical reactors2–6. Some researchers and chemists have tried to model and simulate the ethylene polymerization in autoclave reactors7–9 and in tubular reactors10,11. The choice between these models is mostly dictated by the computing resources available and by the knowledge of the values of the parameters required for the simulation. Tubular chemical reactors have a key role in the chemical industry and they are always part of a larger production system12. A tubular reactor for the production of LDPE is usually very long (>1000 m). Despite this reactor length, the conversion is very low (about 15–35 %) and this is due to a high exothermicity of the reaction. Unreacted monomer is separated from the polymer and recycled in the reactor1,13. For this reason it is necessary to know the behavior of the reactor, which is contained within the proposal of a set of mathematical models that characterize it2,5,6,14. Depending on the precision required, the models can be refined to take into account, in their calculations, the phenomena which are more or less secondary. This will allow presentations that are very close to the real situation. For this purpose, our present work deals with the use of a two-dimensional pseudo-homogeneous model, based on mass and heat balances15 in order to study the behavior of a tubular chemical reactor where a poly- merization reaction proceeds. The proposed model, together with its resolution method, allows a better understanding of the reaction kinetics, particularly for the LDPE synthesis reaction, and for obtaining the temperature and conversion profiles. All the steps of this process (via the polymer chemistry as well as polyme- rization engineering) were based on modeling and simulation. Through the use of a simulation program taking into account all the parameters, several series of calculations are performed over a wide range of temperatures (60–300 °C) and pressures (800–3000 bar). It was found that it is possible to work at the temperatures that allow a moderate conversion (35 %), while avoiding an excessive pressure and producing a resulting polymer that has a low degree of crystallinity. Considering the exothermicity of this reaction, special attention was paid to the operating conditions to ensure the performances, the functioning and the stability control of the reactor. The idea is to have a better control of the heat exchanges by studying the effect of the L/D ratio (length or height/diameter of the cylindrical reactor). For such a purpose, it is better to choose a long reactor with a small diameter (i.e., a high L/D value). 2 FORMULATION OF THE MATHEMATICAL MODEL In an elementary volume of the reactor it will be assumed that the system is treated as a homogeneous one and the proposed model is subjected to the following conditions: 1. the system is stationary ∂ ∂ ∂ ∂  i t t = = 0 2. the reactor has a cylindrical shape 3. the effect of the volume variation due to the reaction, or the temperature, is negligible 4. the axial, mass and heat dispersions are assumed negligible ∂ ∂ ∂ ∂2 2 2 2 0  i z z = = 5. the diffusion and heat-exchange coefficients in the reactor remain constant 6. the pressure is constant along the reactor 7. the temperature of the reactor wall is constant. The mass and energy balances in the dimensionless form are presented as follows: ∂ ∂ ∂ ∂ ∂ ∂2      i i i i iz a y y y b R= + ⎡ ⎣⎢ ⎤ ⎦⎥ +12 2 12 1 ( , ) (1) ∂ ∂ ∂ ∂ ∂ ∂2      z a y y y b R i i= + ⎡ ⎣⎢ ⎤ ⎦⎥ +22 2 22 1 ( , ) (2) M. MARGHSI, D. BENACHOUR: USE OF A TWO-DIMENSIONAL PSEUDO-HOMOGENEOUS MODEL ... 540 Materiali in tehnologije / Materials and technology 46 (2012) 5, 539–546 where: a L Pe r12 2 0 2= ⋅mr Pe L D z mr eff, r = ⋅ a L Pe r22 2 0 2= ⋅hr Pe L C D z G PG mr eff, r = ⋅ ⋅ ⋅  b L C i i z 12 0 1 = ⋅ ⋅ ⋅ −⎛ ⎝ ⎜ ⎞ ⎠ ⎟     s b L T C H i z G PG 22 0 1= ⋅ ⋅ ⋅ ⋅ ⋅ − ⋅ −      s r( ) with the following boundary conditions: z = 0 ⇒   ,   ∀ =( , )z y 0 ⇒ ∂ ∂ ∂ ∂  i t t = = 0 y r r = = 0 1 ⇒ ∂ ∂ ∂ ∂ ∂ ∂ C r C r y y i i i i= ⋅ = → = 0 0 0 0   − ⋅ = − ⋅ −   eff, r w w T r y h T 0 0 0 ∂ ∂ ( ) ⇒ − = − − = − − ∂ ∂       y h T r T Bi w eff, r w w 0 0 0 0 ( ) ( ) with w w= T T0 Equations (1) and (2) form a system of parabolic partial derivatives, which are very difficult to resolve. Therefore, the use of the orthogonal collocation method, which assumes a simple form for the radial profiles is advantageous. For each collocation point yj we obtain: ∂ ∂     i j jk i k k N j jk i k k Nz y z a B y A ( , ) = ⋅ + ⋅ ⎡ ⎣ ⎢ ⎤ ⎦= = ∑ ∑12 1 1 1 ⎥+ b R j j12 ( , )  (3) ∂ ∂     j jk k k N j jk k k N z a B y A b R= ⋅ + ⋅ ⎡ ⎣ ⎢ ⎢ ⎤ ⎦ ⎥ ⎥ + = = ∑ ∑22 1 1 22 1 ( j j, ) (4) (j = 2, 3, … n–1) where  i k i kz y, ( , )= It is clear that equations (3) and (4) are in the form of two first-order differential systems that can be solved by the Runge-Kutta semi-implicit 4th-order method. By substituting the conversion rate in equations (3) and (4) of the model, we obtain: ∂ ∂ X z a B x y A x j jk k j jk k k N k N = − ⋅ − + ⋅ − ⎡ ⎣ ⎢ ⎤ ⎦ ⎥ == ∑∑12 11 1 1 1( ) ( ) + = − ⋅ + ⋅ ⎡ ⎣ == ∑∑ b R z a B y A j j j jk k j jk k k N k N 12 22 11 1 ( , )     ∂ ∂ ⎢ ⎤ ⎦ ⎥+ b R j j22 ( , )  (5) (j = 2,3, … N–1) After the development of these two equations (5) as a form of two systems, we obtain: F j a A x b R F j Ncol a jk k k N j j( ) ( ) ( , ) ( ) = − ⋅ − + + = − = ∑12 1 12 22 1   A b Rjk k k N j j⋅ + = ∑    1 22 ( , ) (6) where A B y Ajk k N jk j jk k N = = ∑ ∑= + ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ 1 1 1 The validation of this model, which takes into account differential terms characteristic of the system behavior, has been particularly considered. In this frame- work a simulation program was developed to predict the behavior of a tubular reactor, where a polymerization reaction takes place under high pressure. The model can be used for a reaction, whose rate is a function of the temperature and concentrations of the reacting chemi- cals. 3 REACTION KINETICS OF THE ETHYLENE POLYMERIZATION The reaction system chosen in our work is the radical synthesis of low-density polyethylene (LDPE) in a tubular chemical reactor. LDPE is produced with the radical polymerization of ethylene under high pressure (800–3000 bar) and at the temperatures ranging from 60 °C to 300 °C in the presence of the traces of oxygen and a free radical generator, the azo-bis-isobutyronitrile (AIBN). The reaction is highly exothermic, and one of the first challenges in this process is the removal of the excess heat generated. The general mechanism of this polymerization involves three main steps, i.e., initiation, propagation, and termination, as shown below: 3.1 Initiation: the step, during which a limited number of active species is created, I R R M M K K a a ⎯ →⎯ + ⎯ →⎯ 2 1 * * * where: I = initiator, R* = initial free radical, M1 * = pro- pagating free radical. 3.2 Propagation: successive reactions of monomer molecules to one active or activated end, leading to the growth of the macromolecular chain, M M M M M M K K p p 1 2 2 3 1 2 * * * * + ⎯ →⎯⎯ + ⎯ →⎯⎯ Or in general: M M Mn K n pn* *+ ⎯ →⎯⎯ +1 where: Kp1, Kp2, …, Kpn are propagation constants. 3.3 Termination: deactivation of the species or the active end, and the cessation of the chain growth. M M Pn m K n m tc* *+ ⎯ →⎯ + (by combination or coupling) M. MARGHSI, D. BENACHOUR: USE OF A TWO-DIMENSIONAL PSEUDO-HOMOGENEOUS MODEL ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 539–546 541 M M P Pn m K n m td* *+ ⎯ →⎯ + (by disproportionation) The main step of the polymerization is the chain propagation (including several elementary reactions), when the macromolecule is built. The successive addition of monomer M during the propagation can be generally described as follows: M M Mn K n p* *+ ⎯ →⎯ +1 The expression of the propagation speed is then: V K M f K K Ip p d t= ⋅ ⋅ ⋅( ) / ( ) Vp can also be expressed generally as follows: V K M V Kp p i t= ⋅( ) / 2 It is to be noted that the values of various parameters (technology, kinetics, thermodynamics, etc.) are taken from the literature16–20. Thus, we have all the data required to calculate the theoretical profiles of the temperature and the conversion in the reactor (Tables 1 and 2). Table 1: Operating conditions used for simulation Tabela 1: Delovni pogoji, uporabljeni za simulacijo L = 1390 m D = 0.05 m Vms = 0.20 m s–1 c = 1900 kg m–3 g = 530 kg m–3 Cp = 3.135 J g–1 °C–1 CA0 = 2.47 · 10–4 mol L–1 CM0 = 16.75 mol L–1 R = 8.3 J mol–1 K–1 Tw = 400 K Pt = 2250 bar Deff = 6 · 10–5 cm2 s–1 Hw = 81.677 · 10–3 J cm–2 °C–1 s–1 eff = 0.02 W m–1 K–1 Hr = –89.87 kJ mol–1  = 0.5 Table 2: Kinetic parameters of the radical polymerization of ethylene (AIBN at 60 °C)20 Tabela 2: Kineti~ni parametri radikalne polimerizacije etilena (AIBN pri 60 °C)20 Kd = 0.845 · 10–5 s–1 Kp = 0.243 · 103 L mol–1 s–1 Kt = 54 · 107 L mol–1 s–1 Ed = 123 kJ mol–1 Ep = 18.4 kJ mol–1 Et = 1.3 kJ mol–1 4 RESULTS AND DISCUSSION Knowing the great importance of some techno- economic parameters in the industry using chemical reactors, we propose to study the influence of three parameters – the initial temperature T0, the total pressure Pt and the ratio of the reactor’s main dimensions (L/D) – on the profiles of the temperature and conversion rates, to establish the optimal working conditions. It is often difficult to control, at the same time, these three parameters and to obtain reproducible results. Therefore, it is necessary to vary one parameter only while main- taining the other two constant. Some chemical reactors can certainly work at a high temperature or a high pressure, but because of their technology implementation and geometry, it is more difficult to have, simultaneously, a high operating pressure and an elevated reaction temperature. For this reason, we are mainly interested in obtaining information about the physical state of the reaction mixture. The purpose is to clarify the influence of physical conditions – mainly temperature T0 and pressure Pt as well as the L/D ratio – on the reactor functioning in order to know its behavior and also to avoid the phenomenon of thermal runaway due to excessive temperature resulting from a bad heat exchange that leads to a thermal instability. After the simulation treatments of our models, the following results are obtained: 4.1 Influence of the initial temperature (T0) In all industrial installations, the temperature measurement is particularly important to ensure the performance and to monitor the smooth running of the operations. The T0 factor is a basic parameter, from which we deduce most of the other reaction parameters such as pressure, mixture composition and geometry of the reactor (the optimal ratio L/D)21. Changing the initial temperature T0, while maintain- ing the other two parameters, Pt and L/D, constant, affects the dimensionless temperature profiles  and the rate of conversion x. M. MARGHSI, D. BENACHOUR: USE OF A TWO-DIMENSIONAL PSEUDO-HOMOGENEOUS MODEL ... 542 Materiali in tehnologije / Materials and technology 46 (2012) 5, 539–546 Figure 1: Variations of dimensionless temperature  and conversion x along the z-axis of the reactor for various values of T0 Slika 1: Spreminjanje brezdimenzijske temperature  in konverzije x vzdol` z-osi reaktorja za razli~ne vrednosti T0 Figure 1 shows a clear hotspot Tc, right at the entrance of the reactor (z = 0.10), for each value of T0 considered, where we may have to cope with the problems of thermal instability that most often occur after a failure of the cooling system (placed against the outer wall of the reactor). The value of TC decreases with an increase in T0, which means that the dimensionless temperature  decreases with an increase in the initial temperature T0 at any position along the reactor z-axis. This increase of T0 has the advantage of increasing the conversion, that is to say, the polymer molecular weight and viscosity. In general, the conversion increases slightly with increasing T0 for any value of z. The range of temperature T0 was limited with the performance of the system that allows both relatively fast reaction rates and relatively good conversions, leading to a polymer having a low crystalline content and a low density. It is noticeable that the temperature, after reaching the maximum value of TC for each T0, decreases along the z-axis, which allows us to conclude that: • The reaction is greatly accelerated by a rise in the temperature at the entrance of the reactor, where the polymerization rate is maximum. The heat generated is removed through a cooling liquid (usually water) to reduce the occurrence of hot spots, TC, and to obtain a uniform distribution of the temperature inside the reactor. • The initiation reaction has a strong thermal energy resulting in the "hot spot" observed during this step. We can say that it is likely that the activation energy of the initiation (Ed) is much superior to the other activation energies (propagation (Ep) and termination (Et)); in our case, the initiation is the result of a thermal decomposition (Table 2). Because the velocity profiles are considered flat, i.e., the flow in our reactor is assumed to be a plug flow, the transit time is the same for each species. For this reason, the radial temperature  and the conversion (Figure 2) remain constant. 4.2 Influence of the total pressure (Pt) High-pressure polymerization of ethylene in tubular reactors is an important commercial process22. The pressure is a factor involved directly in the reaction kinetics through the reaction enthalpy as given by the following expression19: H = 115.3 · [718.6 + (0.05 · T0) + (0.025 · Pt)] /(J · g–1 · mol–1) To better highlight the effect of the total pressure on the temperature and conversion profiles, the simulations are made assuming that this pressure remains constant along the reactor (an assumption of no charges losses). The results shown in Figure 3 (the profiles of dimen- sionless temperature  and the rate of conversion x as a M. MARGHSI, D. BENACHOUR: USE OF A TWO-DIMENSIONAL PSEUDO-HOMOGENEOUS MODEL ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 539–546 543 Figure 3: Variations of dimensionless temperature  and conversion x along the z-axis of the reactor for different values of Pt Slika 3: Spreminjanje brezdimenzijske temperature  in konverzije x vzdol` z-osi reaktorja za razli~ne vrednosti Pt Figure 2: Variations of dimensionless temperature  and conversion x as a function of radial direction y of the reactor for various positions along the z-axis Slika 2: Spreminjanje brezdimenzijske temperature  in konverzije x kot funkcije radialne smeri y v reaktorju za razli~ne polo`aje vzdol` z-osi function of the pressure along the z-axis of the reactor) confirm the influence of the pressure on these two parameters. Such results are in agreement with the expectations: when pressure Pt increases the conveying speeds are highly accelerated and the temperatures are higher, while the conversion rate decreases regardless of the position along the z-axis of the reactor. This con- version decrease can be explained as follows: a pressure rise increases the conveying speed that, in turn, limits the cross-linking reactions and deposit formation (crust) on the wall (these reactions cause the formation of deposits that prevent adequate heat transfer). The range of the total pressure Pt has been limited by the performance of the system in order to avoid excessive pressure ( 3000 bar) and prevent an over- heating of the system. It may be noticed that, when working at high pressures, hot spots become more consistent, thus potentially leading to a runaway of the reactor. For this reason, the heat must be controlled precisely to prevent such a runaway and to better control the molecular weight distribution of the polymer. Here we also notice that the radial temperature  and the rate of conversion x (Figure 4) remain practically constant along the radial direction y. This uniformity of the temperature and the conversion rate is due to the fact that the velocity profiles are considered flat, i.e., the flow in our reactor is assumed to be a plug flow, therefore, the transit time is the same for each species. 4.3 Influence of the ratio (L/D) Equipment geometry and dimensions are very important because, in the industrial production, all the operations are performed in a reactor having certain geometric dimensions. In practice these parameters (characteristics) are designated as "the main dimen- sions", and can be represented by the length (or height) L and the diameter D of the reactor. It is common to use a dimensionless number, characteristic of the device, designated by the ratio (L/D). Knowing the value of the L/D ratio is very important, because, from its value, one can deduce, for example, the type of a flow (i.e., the Reynolds Number) and also the reactor similarity that should then be considered. If the ratio (L/D) increases, the particle motion becomes increasingly ordered, i.e., the axial diffusion is less important. The transition (similarity) between a tubular reactor and a plug-flow reactor can then be characterized with the axial diffusion, i.e., with an increase in the ratio (L/D). Indeed, the tubular reactor is used only if the residual heat is moderate. In the opposite case, significant radial temperature differences will appear. These would cause radial gradients of the polymerization rate and the viscosity that would impair the quality of the polymer. Figures 5 and 6 show that an increase in the ratio (L/D) leads to the reduction of the temperature, which, in turn, causes an increase in the conversion rate, and this can be observed for each z position along the reactor axis. This allows us to say that a large value of the ratio (L/D) greatly influences the homogenization of the reaction medium with molecular diffusion, and that this M. MARGHSI, D. BENACHOUR: USE OF A TWO-DIMENSIONAL PSEUDO-HOMOGENEOUS MODEL ... 544 Materiali in tehnologije / Materials and technology 46 (2012) 5, 539–546 Figure 5: Variations of dimensionless temperature versus ratio (L/D) for different positions along the z-axis Slika 5: Spreminjanje brezdimenzijske temperature z razmerjem L/D za razli~ne pozicije vzdol` z-osi Figure 4: Variations of dimensionless temperature  and conversion x as a function of radial direction y of the reactor for various positions along the z-axis Slika 4: Spreminjanje brezdimenzijske temperature  in konverzije x kot funkcije radialne smeri y reaktorja za razli~ne pozicije vzdol` z-osi property makes the tubular reactors more adapted to the study and implementation of highly exothermic reac- tions. 5 CONCLUSION Taking into account all the results obtained, we observe that: • Measurements of temperatures, pressures and main dimensions of the reactor are particularly important to ensure the performance and monitoring of the functioning of the ongoing operations. • The two-dimensional model formulation and the development of the appropriate calculation programs led us to a better understanding of the evolution of the temperature and conversion rate during the synthesis reaction of low-density polyethylene in a tubular chemical reactor. • The proposed model (two dimensional) and the reso- lution method (Runge-Kutta semi-implicit 4th-order method), applied to the process of polymerization of LDPE, allow the obtention and the prediction of the influence of the initial temperature T0, the total pressure Pt and the reactor dimensions (ratio L/D) on the temperature and conversion-rate profiles in a tubular chemical reactor in a steady state. The proposed model as well as the solving method are general enough to be applied to many industrial chemical reactions, with respect to the materials (pro- duction of polymers, for instance) and the materials engineering (reactor dimensions and operating condi- tions). They allow a study and a comparison of the profiles (temperature and conversion) for different operating conditions of the reactor. Thus, they appear to be able to predict, with a reasonable accuracy, the behavior of the reactor in question. Nomenclature Bi Biot number at the reactor wall Ci0 initial concentration of the i component [mol L–1] CP G, CP Sspecific heat at a constant pressure of gas, solid [J kg–1 K–1] Deff effective diffusivity [m2 s–1] DA diffusion coefficient of the component A [m2 s–1] E activation energy [J mol–1] F efficiency factor of the initiator hW coefficient of the overall heat transfer to the wall [W m–2 K–1] (I) initiator concentration [mol L–1] Kp propagation rate constant [L mol1 s–1] Kd rate constant for the initiator dissociation [s–1] Ka rate constant for the monomer addition [s–1] Kt termination rate constant [L mol1 s–1] Ktc termination rate constant by combination [L mol1 s–1] Ktd termination rate constant by disproportionation [L mol1 s–1] L reactor length [m] M monomer (M) monomer concentration M [mol L–1] M* monomer radical Pt total pressure [bar] Pe Peclet number Pemr Peclet number on the radial matter Pehr Peclet number of the radial heat r radial distance [m] Ri reaction rate [mol kg–1 s–1] t time [s] T temperature [K] T0 initial temperature [K] TW temperature of the wall [K] Vp speed of the propagation (polymerization) [mol L–1 s–1] Vi initiation rate z average axial velocity z dimensionless axial distance Y dimensionless radial distance Z axial distance [m] eff effective thermal conductivity [W m–1 K–1]  density [kg m–3] s volume density of the catalyst (s = m/Vs) [kg m–3] G volume density of gas a bulk density (a = m/V) [kg m–3]  porosity i dimensionless concentration (i = Ci/Ci0)  dimensionless temperature ( = T/T0 ) W dimensionless temperature of the wall (W = TW/T0) i stoichiometric coefficient x conversion rate [%] Hr0 heat released during the reaction [J mol–1] M. MARGHSI, D. BENACHOUR: USE OF A TWO-DIMENSIONAL PSEUDO-HOMOGENEOUS MODEL ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 539–546 545 Figure 6: Variations of conversion rate x versus ratio (L/D) for different positions along the z-axis Slika 6: Spreminjanje hitrosti konverzije x z razmerjem L/D za razli~ne pozicije vzdol` z-osi 6 REFERENCES 1 M. Hafele, A. Kienle, M. Boll, C. U. Schmidt, Modeling and analysis of a plant for the production of low density polyethylene, Computers and Chem. Eng., 31 (2006), 51 2 L. Le Letty, A. Le Pourhiet, J. B. Gros, M. Enjalbert, Modèle Bidimensionnel de Réacteur Catalytique à Lit Fixe, Chem. Eng. J., 8 (1974), 179 3 G. Djelveh, J. B. Gros, R Bugarel, Simulation d’un réacteur cataly- tique à lit fixe (Oxydation du propène), Can. J. of Chem. Eng., 60 (1982), 146 4 J. B. Gros, R. Bugarel, Etude Comparative de Modèles de Réacteurs Catalytiques à Lit Fixe, Chem. Eng. J., 13 (1977), 165 5 G. F. Froment, K. B. Bischoff, Chemical Reactor Analysis and Design, John Wiley and Sons, New York 1979 6 N. Thérien, P. Tessier, Modélisation et Simulation de la décom- position catalytique du méthanol dans un réacteur à lit fixe, Can. J. of Chem. Eng., 65 (1987), 950 7 P. Feucht, B. Tilger, G. Luft, Prediction of molar mass distribution, number and Weight overage degree of polymerization and branching of low density polyethylene, Chem. Eng. Sic., 40 (1985) 10, 1935 8 P. Lorenzini, M. Pons, J. Villermaux, Free-radical polymerization engineering, IV: Modelling homogeneous polymerization of ethy- lene: determination of model parameters and final adjustment of kinetic coefficients, Chem. Eng. Sic., 47 (1992) 15–16, 3981 9 R. Dhib, N. Al-Nidawy, Modeling of free radical polymerization of ethylene using difunctionnal initiators, Chem. Eng. Sci., 57 (2002), 2735 10 A. Brandolin, N. J. Capiati, J. N. Farber, E. M. Valles, Mathematical model for high pressure tubular reactor for ethylene polymerization, Ind. Eng. Chem. Res., 27 (1988), 784 11 A. Baltsas, E. Papadopoulos, C. Kiparissides, Application and validation of the pseudo-kinetic rate constant method to high pressure LDPE tubular reactor, Comput. Chem. Eng., 22(Suppl.1) (1998), S95–S102 12 J. R. H. Ross, Catalyst preparation from Art to Science, Technisch Hogeschool Twente, 1984 13 M. Hafele, A. Kienle, M. Boll, C. U. Schmidt, M. Schwibach, Dyna- mic simulation of a tubular reactor for the production of low-density polyethylene using adaptive method of lines, J. of Computation and Applied Mathematics, 183 (2005), 288 14 C. R. Barkelew, Chem. Eng. Prog. Symp., Series 55 (1959) 25, 37 15 M. Marghsi, Modélisation et simulation d’un réacteur catalytique à lit fixe: Application à la synthèse du SO3, Magister’s thesis, Uni- versité Ferhat-Abbas, Algérie, 1996 16 M. Asteasuain, S. M. Tonelli, A. Brandolin, J. A. Bandoni, Dynamic simulation and optimisation of tubular polymerisation reactors in gPROMS, Comp. and Chem. Eng., 25 (2001), 509 17 D. M. Kim, P. D. Iedema, Molecular weight distribution in low-den- sity polyethylene polymerization: impact of scission mechanisms in the case of a tubular reactor, Chem. Eng. Sci., 59 (2004), 2039 18 D. M. Kim, M. Busch, H. C. J Hoefsloot, P. D. Iedema, Molecular weight distribution modeling in low-density polyethylene polymeri- zation: impact of scission mechanisms in the case CSTR, Chem. Eng. Sci., 59 (2004), 699 19 S. Agrawal, C. D. Han, Analysis of the high pressure polyethylene tubular reactor with axial mixing, AIChE J., 21 (1975), 449 20 G. Odian, Principles of polymerization, 4th ed., Wiley, New York 2004, 206 21 J. Horak, J. Pasek, Conception des réacteurs chimiques industriels sur la base des données de laboratoire, Eyrolles, Paris 1981 22 H. D. Anspon, Polyethylene. In Manufacture of Plastic, Vol. 1, W. M. Smith, ed., Reinhold, New York 1964 M. MARGHSI, D. BENACHOUR: USE OF A TWO-DIMENSIONAL PSEUDO-HOMOGENEOUS MODEL ... 546 Materiali in tehnologije / Materials and technology 46 (2012) 5, 539–546 V. LAZI] et al.: THEORETICAL AND EXPERIMENTAL ESTIMATION OF THE WORKING LIFE ... THEORETICAL AND EXPERIMENTAL ESTIMATION OF THE WORKING LIFE OF MACHINE PARTS HARD FACED WITH AUSTENITE-MANGANESE ELECTRODES TEORETI^NO IN EKSPERIMENTALNO UGOTAVLJANJE ZDR@LJIVOSTI STROJNIH DELOV, OPLA[^ENIH S TRDIMI AVSTENITNO-MANGANSKIMI ELEKTRODAMI Vuki} Lazi}1, Aleksandar Sedmak2, Dragan Milosavljevi}1, Ilija Nikoli}1, Srbislav Aleksandrovi}1, Ru`ica Nikoli}1, Milan Mutavd`i}3 1Faculty of Engineering, S. Janji} 6, 34000 Kragujevac, Serbia 2Faculty of Mechanical Engineering, Kraljice Marije 16, 11000 Beograd, Serbia 3PD „Kragujevac“, Kragujevac, Tanaska Raji}a 16, 34000 Kragujevac, Serbia vlazic@kg.ac.rs Prejem rokopisa – received: 2012-02-06; sprejem za objavo – accepted for publication: 2012-02-16 We have investigated the possibility of repairing damaged machine parts by hard facing with austenite-manganese steel electrodes. The subject is a Fe-C-Mn alloy with a microstructure of soft austenite which, after cold deformation, transforms by a shearing mechanism into a hard martensite microstructure. These steels are used mainly for parts exposed to high impact loads and intensive abrasive wear. Depending on the degree of wear, these parts can be replaced by new ones or repaired by hard facing. The selection of the optimal reparation technology for the rotational crusher’s impact beams is the subject of this study. Investigations of model samples were conducted first, followed by layers hard faced onto samples with austenite manganese and special electrodes. After this the microstructure and hardness of the welds’ characteristic zones were investigated. After reparatory hard facing the impact beams were mounted in the crusher and their behaviour was monitored periodically. Both the new and hard-faced beams’ behaviours were monitored and compared under the same working conditions. In this way, the optimal technology for hard facing was established, taking into account not only the technical indicators, but also the economic effects. Keywords: austenite manganese – hadfield steel, mining engineering equipment, hard facing, reparation Ta ~lanek predstavlja {tudij mo`nosti obnove po{kodovanih delov z nana{anjem trdih plasti iz avstenitno-manganskih jekel. Predmet raziskave je zlitina Fe-C-Mn z mehko avstenitno mikrostrukturo, ki se med hladno deformacijo s stri`nimi mehanizmi pretvori v trdo martenzitno mikrostrukturo. Ta jekla se uporabljajo predvsem za dele, izpostavljene velikim udarnim obremenitvam in mo~ni obrabi. Odvisno od stopnje obrabe se ti deli nadome{~ajo z novimi ali pa se obnovijo s trdimi nanosi. V tem delu je preu~evana izbira optimalne tehnologije obnove rotacijskih udarnih drobilnikov. Najprej je bila izvr{ena preiskava na modelnih vzorcih z nanosom trde plasti iz avstenitno-manganske posebne elektrode, nato pa preiskana {e mikrostruktura in trdota zna~ilnih varjenih podro~ij. Po obnovi z nanosom trde plasti so bile udarne plo{~e name{~ene v drobilnik in periodi~no je bilo spremljano njihovo vedenje. Primerjane so bile lastnosti novih in obnovljenih plo{~ v enakih obratovalnih razmerah. Tako je bila dolo~ena optimalna tehnologija obnavljanja delov z upo{tevanjem ne samo tehni~nih lastnosti, temve~ tudi iz ekonomskega stali{~a. Klju~ne besede: avstenitno-mangansko – Hadfield jeklo, rudarska strojna opreme, nana{anje trdih prevlek, obnavljanje 1 INTRODUCTION In an investigation of the damage caused to various parts of machines and devices it was established that in more than 50 % of cases the damage occurs due to tribological processes involving more-or-less regular working conditions1–3. Accordingly, for the design of the reparation technology for damaged parts, we must first study the possible mechanisms of wear for coupled parts. Here, it should be kept in mind that, besides repairing the parts damaged in normal conditions, hard facing is also used for parts damaged due to failures, as well as for new, flawed cast pieces. Besides, new parts are also hard faced by depositing hard alloys, which can replace the traditional procedures of carburizing and nitriding. All these facts indicate that hard facing is an important advanced technologies. The key parts of machines, assemblies and devices are frequently produced from very expensive alloys. Thus, by repairing them we are not only shortening the down times due to repairs, but also saving on expensive base materials as well as for the machining of parts. In the majority of cases the economic criterion for applying reparation is that the price of the repair cannot exceed the price of the new part. This is especially important for large-sized parts and batch production, while the reparation of unique machines and devices sometimes has to be performed regardless of the price4–10. Materiali in tehnologije / Materials and technology 46 (2012) 5, 547–554 547 UDK 621.793 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 46(5)547(2012) 2 SELECTION AND PROPERTIES OF MATERIALS RESISTANT TO ABRASIVE WEAR 2.1 Base metals The materials used most frequently for manufac- turing the working parts of civil-engineering machines exposed to abrasive wear are cast pieces made from manganese steel. When selecting a material for manu- facturing the working parts of such machines, i.e., by selecting the filler metals for their reparation or hard-facing manufacture, we should pay particular attention to the mechanism of abrasive action, since two very different basic cases are met. In the first case, at the contact surface of the working parts of civil-engineering machines, abrasive and very high specific pressures and local plastic deformation of those parts appear when the external load forces have an impact character. A typical part exposed to abrasive wear of this type is the impact beam of a stone crusher, where the abrasive element is the compact stone. In the second case, the abrasive is in the dispersed state (e.g., a small aggregate). Thus, on the contact surfaces only smaller specific pressures appear and large plastic deformations do not occur. The characteristic parts exposed to this type of wear are the teeth of loading dredger scoops, the storage crates of the transporters, etc. Manganese austenite steels, also Hadfield steels11–15, exhibit good resistance to the first type of abrasive action. They are usually supplied in the cast, hot- or cold-deformed states. Besides the basic Hadfield steel (^L3160-JUS, G-X120Mn12-DIN), which contains 1.20 % C, 12 % Mn, 0.50 % Si, 0.35 % P and 0.10 % S, the following multi-alloyed manganese steels are applied steels: (^L3161-JUS), (^L3460-JUS), (^L3462-JUS) and (^L3463-JUS), which possess good resistance to impact wear because, besides the high contents of manganese (w(Mn) = 12–17 %) and carbon (w(C) = 1–1.4 %), they are also alloyed with chromium (w(Cr) = 1–1.8 %). The hard facing reparation of these steels is usually performed by the application of basic manganese electrodes. The good resistance to abrasive wear by dispersed materials is exhibited by heat-treated, low alloyed steels, rapidly cooled cast steels and ledeburite tool steels12,15,16. The chemical composition and instructions for the application of the base metals (^L3160 and ^L3460) are given in Table 1, while the comparative marks by the JUS and DIN Standards, as well as the mechanical properties and microstructure of those materials, are given in Table 2 4–6,17. 2.2 Filler metals In the experimental investigation, various filler metals were applied (E Mn14, E Mn17Cr13, E DUR 600, ABRADUR 58 and INOX B 18/8/6) 17. In Table 3 the chemical composition and the comparative Stan- dards’ marks are presented, and in Table 4 are the hardness and applications of tested electrodes. The most frequently recommended filler metals for the hard facing of the parts for civil-engineering machines subjected to impact abrasive wear and the V. LAZI] et al.: THEORETICAL AND EXPERIMENTAL ESTIMATION OF THE WORKING LIFE ... 548 Materiali in tehnologije / Materials and technology 46 (2012) 5, 547–554 Table 1: Chemical composition and application of ^L3160 and ^L3460 Tabela 1: Kemijska sestava in uporaba ^L3160 in ^L3460 Base metal Chemical composition, mass fraction, w/% Application C Si Mn Cr P S Suitable for manufacturing the parts exposed to abrasive wear and high impact loads, such as the working parts of mills, crushers, civil engineering machines, for work with raw materials of high hardness, etc. ^L3160 Prescribed 1.20 0.50 12.00 – 0.035 0.10 Analyzed 1.20 0.48 12.35 – 0.025 0.10 ^L3460 Prescribed 1.20 0.50 13.00 1.00 0.040 0.10 Analyzed 1.20 0.55 13.14 1.12 0.035 0.15 Table 2: Comparative Standard marks, some mechanical properties and microstructure of ^L3160 and ^L3460 Tabela 2: Primerjalne oznake po standardih, nekatere mehanske lastnosti in mikrostruktura ^L3160 in ^L3460 Comparative marks Mechanical properties Microstructure JUS DIN Tensile strength, Rm/MPa Hardness, HB ^L3160 G-X120Mn12 200  200 after fast quenching Austenite ^L3460 G-X120Mn12 210  210 after fast quenching Austenite Table 3: Comparative Standard marks and chemical composition of the tested electrodes Tabela 3: Primerjava oznak po standardih in kemijska sestava preizku{enih elektrod Comparative marks Chemical composition, mass fraction, w/% S@ Fiprom Jesenice DIN8555 C Mn Cr Ni W Mo V E Mn14 E7-UM-200-KP 1.20 12.50 – – – 0.70 – E Mn17Cr13 – 0.60 16.50 13.50 – – – – E DUR 600 E 6-UM-60 0.50 – 7.50 – – – – ABRADUR 58 E 10-UM-60-GR 3.60 – 32.0 – – – – action of impact loads are austenite manganese electrodes with a large content of manganese and carbon, and the addition of other alloying elements, usually chromium, and then Ni, Mo, V and W11–13,15,17. 3 ESTIMATION OF THE WELDABILITY OF MANGANESE STEELS Manganese steels with an austenite structure are prone to overheating and, with the deposition of multi- layer welds, the appearance of cracks is possible. Thus, it is necessary to know the behaviour of those materials when a significant amount of heat is being brought in, which would enable measures preventing the growth of austenite grains and the formation of brittle phases. Depending on the content of manganese and carbon, during a slow enough cooling of a Fe-C-Mn alloy we can obtain the following: perlite, perlite-martensite, marten- site-austenite and austenite microstructures. Perlite and austenite manganese steels have practical applications; however, they have a low plasticity and, due to rapid hardening, they are difficult to machine mechanically. An increase of the plasticity of manganese austenite steels is achieved by the thermal treatment of rapid quenching, which consists of heating up to a temperature of 1093 °C 15, then heating through that temperature, followed by rapid water cooling. In this way, a pure austenite microstructure and high toughness are obtained. After the cooling, from casting or from the hot-deformation temperature, we obtain an austenite microstructure with particles of complex iron manganese carbides precipitated at the boundaries of the austenite grains. In contrast to this, after slow cooling some martensite will also appear, i.e., a predominantly marten- site or austenite-martensite microstructure. To achieve a satisfactory weldability of these man- ganese austenite steels it is necessary to prevent the abrupt growth of austenite grains due to the excessive quantity of heat that is introduced. This is why it is recommended that the hard facing is performed with short welds, the careful selection of hard-facing parameters and forced rapid cooling. In the opposite case, the abrupt growth of austenite grains can occur, which would worsen its weldability and mechanical properties. The tendency of welds that are applied with manganese electrodes to crack is significantly reduced by the addition of a certain quantity of nickel and chromium, usually up to 4 % 13,15. 4 MODEL INVESTIGATIONS 4.1 Selection of the hard-facing technology The tests were conducted on samples made of low-carbon steel (^0361) with a thickness s = 10 mm for the purpose of selecting the technological parameters of the hard-facing procedure. The samples were welded using the REWL procedure. Depending on the type and the diameter of the used electrode, the hard-facing parameters were within the limits given in Table 5 4–6,16. The method of depositing the weld layers, the order and the number of deposited layers and the appearance of the model are presented in Figure 1. It is necessary to emphasize that, for further metallo- graphic and other investigations, the samples were chosen to be welded using an electrode of diameter de = 3.25 mm. Single-pass welds had a width b = 6–12 mm and a height h = 3.2–4.6 mm. The two-layered samples were only welded with the electrodes E Mn14 and E Mn17Cr13 (Figure 1b), while the three-layered samples were welded with electrodes E DUR 600 and ABRADUR 58 (Figure 1c) with the prior deposition of the inter-layer with the electrode INOX B 18/8/6. For model investigations, neither a prior nor a posterior thermal treatment was applied. After the hard facing the samples were rapidly cooled and then with grinding a V. LAZI] et al.: THEORETICAL AND EXPERIMENTAL ESTIMATION OF THE WORKING LIFE ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 547–554 549 Table 4: Mechanical properties and application of the tested electrodes Tabela 4: Mehanske lastnosti in podro~je uporabe preizku{enih elektrod Electrode mark S@ Fiprom-Jesenice Hardness Application E Mn14 220 HB – after hard-facing 48 HRC – after forging For hard-facing of manganese steels of thickness up to 10 mm which are applied for the manufacturing of parts in railroad engineering and parts of the stone crushers E Mn17Cr13 220 HB – after hard-facing 48 HRC – after forging For hard-facing of hydraulic presses’ mallets, parts of the loading scoops of the civil engineering mechanization, parts of crushers, rails and railway-crossings E DUR 600 57–62 HRC For depositing of the hard welds from which the high wear resistance in the hot and cold states is expected, as well as good toughness and impact resistance ABRADUR 58 57–62 HRC For hard-facing of tools subjected to intensive abrasive wear in contact with minerals in a cold state INOX B 18/8/6 - For welding of the Cr- and Cr-Ni steels, for welding of two different types of steels, for depositing of welds resistant to corrosion and for depositing of the plastic inter-layer. portion of the material was removed from the back layer4–6. 4.2 Metallographic investigations and hardness measu- rements on models Manganese austenite steels have a relatively small hardness in the range from 180 HB to 250 HB (usually 200–229 HB). They are highly resistant to abrasive wear only when their working surface layers are intensively plastically cold deformed with effect of strong impact loads or of slow pressure as a result of the pressing load. However, in these steels the hardening of the surface layers is not due to the strain hardening of the austenite, but the plastic deformation initiates the phase transfor- mation of austenite to martensite. After local transfor- mation of the austenite to martensite, a hardness of the surface layers as high as 500–520 HK can be achieved (the usual hardness range is between 330 HK and 480 HK). For this reason, these steels are hard to machine by cutting and are usually treated by hot or cold plastic deformation or by casting. In some references4–6,8 the limiting depth was established at which we can still observe an increase in the surface layers’ hardness due to the austenite-to- martensite transformation initiated by the plastic cold deformation. The maximum measured hardness was 460 HK, and the width of the transformed zone was 0.50 V. LAZI] et al.: THEORETICAL AND EXPERIMENTAL ESTIMATION OF THE WORKING LIFE ... 550 Materiali in tehnologije / Materials and technology 46 (2012) 5, 547–554 Figure 2: Microstructure of the Hadfield steel (^L3160): a) before the plastic deformation and b) after the plastic deformation Slika 2: Mikrostruktura Hadfield jekla (^L3160): a) pred plasti~no deformacijo, b) po plasti~ni deformaciji Figure 1: Order of weld layers’ depositing: a) 1 layer, b) 2 layers (E Mn14, E Mn17Cr13), c) 3 layers (INOX B 18/8/6-E DUR 600, INOX B 18/8/6-ABRADUR 58), d) metallographic ground piece (block) Slika 1: Zaporedje navarjenih slojev: a) 1 sloj, b) 2 sloja (E Mn14, E Mn17Cr13) c) 3 sloji (INOX B 18/8/6-E DUR 600, INOX B 18/8/6-ABRADUR 58), d) kos za metalografske preiskave Table 5: Technological parameters of hard-facing Tabela 5: Tehnolo{ki parametri nana{anja trdih plasti Base metal thickness s, mm Electrode mark Electrode core diameter de/mm Current intensity I/A Working voltage U/V Welding speed vz/(cm/s) Welding driving energy, J/cm 10 E Mn14 3.25 120 25  0.148 16216 E Mn14 5.00 180 27  0.162 24000 E Mn17Cr13 3.25 130 25  0152 17105 E Mn17Cr13 5.00 200 28  0.168 26667 INOX B 18/8/6 3.25 100 24  0.136 14118 INOX B 18/8/6 5.00 160 26  0.178 18697 E DUR 600 3.25 120 25  0.119 20168 ABRADUR 58 3.25 130 25  0.124 20968 mm. In Figure 2 the microstructure of the Hadfield steel (^L3160), before and after plastic deformation, is presented4–6. In Figure 2 we can see that the microstructure of the Hadfield steel before the plastic deformation was purely austenite, while after the plastic deformation the needles of martensite can be seen in the austenite matrix. In some investigations1–3 it was shown that the austenite-carbide microstructure has the highest wear resistance, rather than the martensite-carbide, as would be expected from their hardness values. The reason lies in the stronger austenite-carbide, grain-boundary bonds due to the smaller difference of the lattices parameters, rather than in the martensite-carbide combination. In other words, abrasive particles are pulling out the carbide particles from the martensite matrix more easily than from the austenite. Despite the evident difficulties associated with hardness measurements and in dterming the hard-faced layers’ microstructure, we were able to determine the width of the austenite-to-martensite transformation zone and were also able to record the microstructure of that zone. These data can be of special importance in the hard facing of various working parts of technical systems that operate in such or similar working conditions. In Figures 3 and 4 are the distributions of the welds before and after the plastic deformation. The following filler metals were used: E Mn14 and E Mn17Cr13 4–6,8. From Figures 3 and 4 it is clear that there is an increase in the hardness after the plastic deformation in both tested filler metals. By comparing the results (Figure 4) related to the filler metals E Mn14 and E Mn17Cr13, it is clear that a somewhat higher hardness was obtained for the second filler metal. Also, it can be concluded that the width of the transformed austenite- to-martensite zone is larger for the welds deposited by the E Mn17Cr13 electrode. The maximum measured hardness after the cold hardening in the welds deposited by the E Mn14 electrode was about 520 HK, while the width of the transformed zone was 0.60 mm. In the same conditions for welds deposited by the E Mn17Cr13 electrode the maximum obtained hardness was 560 HK and the width of the transformed zone was 1.20 mm 4–6,8. The microstructures of the surface layer of the back weld before and after the cold plastic deformation are shown in Figures 5 and 6 4–6,8. The hardness distributions and the appearance of the formed structures in the characteristic zones of the filler metals E DUR 600 and ABRADUR 58 are presented in4–6,8. V. LAZI] et al.: THEORETICAL AND EXPERIMENTAL ESTIMATION OF THE WORKING LIFE ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 547–554 551 Figure 3: Micro hardness distribution along the welds’ cross-section before the plastic deformation: a) E Mn14 and b) E Mn17Cr13 Slika 3: Razporeditev mikrotrdote v pre~ni smeri vzdol` zvara pred plasti~no deformacijo: a) E Mn14 in b) E Mn17Cr13 Figure 4: Micro hardness distribution along the welds’ layer cross- section after the plastic deformation: a) E Mn14 and b) E Mn17Cr13 Slika 4: Razporeditev mikrotrdote v pre~ni smeri vzdol` zvara po plasti~ni deformaciji: a) E Mn14 in b) E Mn17Cr13 5 EXPERIMENTAL HARD FACING OF IMPACT BEAMS AND A DETERMINATION OF THEIR WEAR RESISTANCE 5.1 Description of the crusher plants’ operation The impact beams of crushers are simultaneously exposed to large impact loads and intensive abrasive wear since they are in working contact with rocks. In fact, the rocks are being crushed, minced and ground in order to obtain various fractions of the granules for direct insertion into roads or buildings. Besides being used for grinding rocks, similar crushing plants are also being used for the mincing of ores and coal. Impact-beam wear occurs according to the mechanism of abrasive wear of the so-called closed type. The rock is brought into the working space between the beams and the stationary crusher housing where the kinetic energy of the rotational beams is transferred to the work, which is being spent on breaking the cohesive and adhesive bonds of the rock material. The rock materials for building roads are usually of high hardness, and thus the crushers’ working parts must have high toughness and wear resistance4–7. The experimental investigations presented in this paper were conducted on impact beams and segments of the housing forming the crusher’s jaw, Figure 7. The wear of the beams and the housing was monitored during the crushing of lime stone in a crusher with a capacity of 350 t/h. This is a rotational crusher with four impact beams, which are mounted onto the rotor and are driven by an electric or IC engine, with a pulley and v-belts. Such transmissions enable the belts to slide if over- loading occurs, which protects the crusher’s vital parts against fracture. Impact beams have two alternative working surfaces, which allows the beam to be turned over and the second surface used after the first one is worn out. The crusher housing is made of flat spherical segments that are also being worn during the working process. The crusher impact beams, on which the experimen- tal hard-facing was performed, are made of manganese steel. Their dimensions were 300 mm × 120 mm × 1000 mm and their mass was around 300 kg. The same cast steel sheets were used for manufacturing the housing elements of thickness 30 mm and mass of around 20 kg. V. LAZI] et al.: THEORETICAL AND EXPERIMENTAL ESTIMATION OF THE WORKING LIFE ... 552 Materiali in tehnologije / Materials and technology 46 (2012) 5, 547–554 Figure 6: Microstructure of the weld’s back layer: a) before the plastic deformation and b) after the plastic deformation; Filler metal: E Mn17Cr13 Slika 6: Mikrostruktura zvara: a) pred plasti~no deformacijo in b) po plasti~ni deformaciji; material elektrode: E Mn17Cr13 Figure 5: Microstructure of the weld’s back layer: a) before the plastic deformation and b) after the plastic deformation; Filler metal: E Mn14 Slika 5: Mikrostruktura zvara: a) pred plasti~no deformacijo in b) po plasti~ni deformaciji; material elektrode: E Mn14 5.2 Experimental hard facing The hard facing of impact beams was performed by technologies that are similar to the model investigations, but with changed working conditions. The reason for this is the large mass of the impact beams, as the hard facing was done in real working conditions, i.e., in a quarry. This means that the welds were deposited in the most unfavourable conditions, which gives special importance to the obtained results. In the reparatory hard facing of the damaged impact beams, four beams were hard-faced, each one with a different filler metal (E Mn14, E Mn17Cr13, ABRA- DUR 58 and E DUR 600). The welds were deposited onto one of the two working surfaces of the impact beams, while the other working surface was brand new and unworn, and in this way a comparison of the working life of new and repaired beams was possible. The welds were deposited longitudinally on the working surface and the necessary weld thickness was achieved by multilayer hard facing with the thickness of each layer ranging from 10 mm at the ends to 35 mm in the middle of the impact beams where the wear was the greatest. The hard facing of the impact beams with electrodes E Mn14 and E Mn17Cr13 was done without the deposition of a plastic inter-layer. On the contrary, the weld layers realized by the ABRADUR 58 and E DUR 600 electrodes were deposited over the pre- viously deposited plastic inter-layer of INOX B 18/8/6. After the hard facing, no faults of the crack type were spotted during a visual control of the large, hard-faced working surfaces (1200 mm × 100 mm). In the reparatory hard facing of impact beams, two new beams were hard faced by depositing lateral multi-layers consisting of single-pass welds with the electrodes ABRADUR 58 (one impact beam) and E DUR 60 (the other impact beam). The distance between the deposited layers was about 100 mm, and partial welds were deposited all over the working surface. Two other impact beams were hard faced by depositing partial cross-like (honeycomb) welds over the whole working surface in such way that one impact beam was hard-faced with the E Mn14 electrode while the E Mn17Cr13 electrode was used for the other beam. The maximum weld thickness allowed is up to 10 mm, for construction reasons, since thicker welds would touch the housing during rotation. As in the first case, the hard facing with the ABRADUR 58 and E DUR 600 welds was deposited over the previously deposited plastic layer of INOX B 18/8/6, while the hard-facing welds of the electrodes E Mn14 and E Mn17Cr13 were directly deposited over the base metal. In this way, the second set of impact beams was prepared. With monitoring of the impact beams’ wear it was established that, after continuous work for 72 h, the damage to the working surfaces of the new impact beams was very severe and it was not possible to continue the crushing and adjustment of the crusher and that beams must be turned over to use their second working surface. The material losses of the hard-faced impact beams were obtained by monitoring the manufacturing process during 60 effective working hours so the two could be compared. The results of these examinations are pre- sented in Table 6 and in Figure 8 4,6,16. Based on obtained results of the wear-resistance investigations in real working conditions we can conclude that the highest resistance was achieved with the hard faced layer of the E Mn17Cr13 (4.12 %) electrode, followed by the layer deposited with the E DUR 600 (5.73 %) electrode, then the layer deposited with the E Mn14 (7.0 %) electrode, while the layer deposited with the ABRADUR 58 (8.87 %) electrode exhibited wear resistance even lower than that of the base metal – ^L3460 (7.87–8.12 %). Clearly, ABRA- DUR 58 is not suitable to be used for this type of wear. V. LAZI] et al.: THEORETICAL AND EXPERIMENTAL ESTIMATION OF THE WORKING LIFE ... Materiali in tehnologije / Materials and technology 46 (2012) 5, 547–554 553 Table 6: Material losses of crushers’ impact beams in real working conditions after 60 h of operation Tabela 6: Izguba materiala na udarnih plo{~ah drobilnika po 60-urnem delovanju v realnih razmerah Tested samples mass (impact beams) Non hard-faced beams Reparatory hard-faced beams Production hard-faced beams 1 2 3 4 1* 2* 3* 4* 1* 2* 3* 4* At the test beginning, kg 300 300 300 300 300 300 300 300 305 305 305 305 At the test end, kg 275.6 276.0 276.4 276.2 279.0 287.6 282.8 273.4 282.2 291.0 284.8 279.5 Mass loss, kg 24.4 24.0 23.6 23.8 21.0 12.4 17.2 26.6 22.5 14.0 20.2 25.5 Mass loss, % 8.12 8.00 7.87 7.93 7.00 4.12 5.73 8.87 7.38 4.59 6.62 8.36 *Note: 1- E Mn14; 2- E Mn17Cr13, 3- E DUR 600, 4- ABRADUR 58. Figure 7: Appearance of the impact beams and crusher housing seg- ments Slika 7: Videz udarnih plo{~ in deli ohi{ja drobilnika The whole hard-facing process could have been performed automatically by robots. However, for the crushers’ impact beams this is still not possible, espe- cially because the main economic effect of introducing robots into the hard-facing and welding process, increasing the economic efficiency, i.e., lowering the process costs, cannot be achieved on small batch products like beams. For smaller parts and in batch manufacturing, however, the introduction of robots for the hard-facing process would be justified18–20. 6 CONCLUSION The extended experimental investigations have shown that the working life of properly hard-faced impact beams significantly exceeds the working life of the original beams. In this way large savings in material costs are realized, the crusher’s down-time is reduced, and the range and quantity of the necessary spare parts is smaller. In terms of days, the working life of the new impact beams was about 15 d, while the working life of the repaired beams was reaching about 30 d, on average, depending on the applied filler metal. By analysing the costs of the filler metals and the price of the welders’ labour, data were obtained which have shown that the costs of the reparatory hard facing of one impact beam are 25 % lower than the costs of a set of the new beams. This means that by applying the reparatory hard facing, four damaged impact beams (one set) can be renewed for the price of one new beam. By application of reparatory hard facing (primarily with the E Mn17Cr13 electrode) the working life of parts is extended and the down time of the manufacturing process is reduced. All these positive effects came as a result of a complex theoretical model and the investigations of this paper’s authors, realized in collaboration with the user company. One of the tasks that remains for future work is how to exploit the advantages of the robotization of the welding and hard-facing process for large parts such as the crusher’s impact beams. 7 REFERENCES 1 P. Bla{kovi}, J. Balla, M. Dzimko, Tribology, Vydavatelstvo, ALFA, Bratislava, 1990 (In Slovak) 2 K. M. Mashloosh, T. S. Eyer, Abrasive wear and its application to digger teeth, Tribology International, 18 (1985) 5, 257–312 3 N. Jost, I. Schmidt, Friction-induced martensite in austenitic Fe-C steels, in: Ludema K. C. (ed): Wear of materials, ASME, N. Y., 1985, 2005–211 4 M. Mutavd`i}, Reparatory hard-facing of the machine parts and devices in the civil engineering industry mechanization, Master’s thesis, Faculty of Mechanical Engineering, Kragujevac, Serbia, 2007 (In Serbian) 5 V. Lazi}, M. Jovanovi}, N. Ratkovi}, D. Adamovi}, R. Vulovi}, Estimate of the wear resistance of the hard-faced layers deposited by the manganese electrode, Tribology in Industry, 22 (2002) 3/4, 10–17 (In Serbian) 6 M. Mutavd`i}, V. Lazi}, M. Jovanovi}, D. Josifovi}, B. Krsti}, Selection of the optimum technology of reparatory hard facing of the impact beams of the rotational crushing mills, Welding & welded structures, 2 (2007), 55–67 7 G. I. Sil’man, Alloys of the Fe-C-Mn system. Part 4. Special features of structure formation in manganese and high-manganese steels, Metal Science and Heat Treatment, 48 (2006) 1/2, 3–8 8 V. Lazi}, Optimization of the hard facing procedures from the aspect of tribological characteristics of the hard faced layers and residual stresses, Doctoral Dissertation, The Faculty of Mechanical Engi- neering, Kragujevac, 2001 (In Serbian) 9 V. Lazi}, A. Sedmak, S. Aleksandrovi}, D. Milosavljevi}, R. ^uki}, V. Grabulov, Reparation of damaged mallet for hammer forging by hard facing and weld cladding, Technical Gazette, 16 (2009) 4, 107–113 10 B. Nedeljkovi}, M. Babi}, M. Mutavd`i}, N. Ratkovi}, S. Aleksan- drovi}, R. Nikoli}, V. Lazi}, Reparatory hard facing of the rotational device knives for terrain leveling, Journal of the Balkan Tribological Association, 16 (2010) 1, 46–75 11 G. S. Zhang, J. D. Xing, Y. M. Gao, Impact Wear Resistance of WC/Hadfield Steel Composite and its interfacial characteristics, Wear, 728 (2006), 260 12 Engineering handbook – Welding II, WNT, Warszawa, 1983 (In Polish) 13 Metals Handbook, Desk Edition, Wear resistant austenitic manga- nese steel, Edited by J. R. Davis, & Associates, American Society for Metals, Metals Park, Ohio 1998 14 Y. N. Dastur, W. C. Leslie, Mechanism of work hardening in Had- field manganese steel, Metall. Trans. A, 749 (1981), 12A 15 R. W. Smith, A. DeMonte, W. B. F. Mackay, Development of high- manganese steels for heavy duty cast-to-shape applications, Journal of Materials Processing Technology, 153–154 (2004), 589–595 16 M. Mutavd`i}, R. ^uki}, M. Jovanovi}, D. Milosavljevi}, V. Lazi}, Model investigations of the filler materials for regeneration of the damaged parts of the construction mechanization, Tribology in Industry, Journal of Serbian Tribology Society, (2008) 3/4, 3–9 17 Catalogues Thyssen Marathon Edelstahl-Vosendorf, FEP-Plu`ine, Elvaco-Bijeljina, @elezarna Jesenice-Fiprom, Böhler-Kapfenberg, Messer Griesheim-Frankfurt am Main, Esab-Göteborg, Lincoln Electric, USA, Atlas zur Wärmebehandlung der Stähle 18 K. Kelleghan, Welding Report: Sorting through industry trends, February 12, 2001, http://www.robot-welding.com. 19 R. Miller, Automation and Robots, Postel Newsletter, 3 (2010) 6 http://www.postle.com 20 J. Berge, Robotic Weld Process Control, Mr. Roboto: Welding funda- mentals for managers, February 28, 2002, http://www.thefabrica- tor.com V. LAZI] et al.: THEORETICAL AND EXPERIMENTAL ESTIMATION OF THE WORKING LIFE ... 554 Materiali in tehnologije / Materials and technology 46 (2012) 5, 547–554 1, 2, 3, 4 – New beams; 1N, 2N, 3N, 4N – Reparatory hard-faced beams; 1PN, 2PN, 3PN, 4PN – Production hard-faced beams Figure 8: Graphical representation of the crushers’ impact beams material losses after 60 h of operation Slika 8: Grafi~na predstavitev izgube materiala udarnih plo{~ drobilnika po 60-urnem delovanju