Microstructure and Mechanical Properties of HighStrength Graphite Particles Coated TiO2 Embedded SiC

Effects of annealing treatment on the microstructure and mechanical properties of the AlN–SiC–TiB 2ceramic composites prepared by SHS–HIPLijuan Zhou a ,⇑,Hongbo Li b ,Yongting Zheng ba School of Materials Science and Engineering,Shandong University of Technology,ZiBo 255049,PR China bCenter for Composite Materials and Structures,Harbin Institute of Technology,Harbin 150001,PR Chinaa r t i c l e i n f o Article history:Received 14July 2012Received in revised form 10November 2012Accepted 12November 2012Available online 20November 2012Keywords:AlN–SiC–TiB 2ceramic Annealing treatment MicrostructureMechanical propertiesa b s t r a c tAnnealing treatments were carried out at different temperatures (600–1200°C)for various holding times (2–10h)to evaluate the microstructural and mechanical properties changes of the AlN–SiC–TiB 2ceramic composites prepared by SHS–HIP.The experimental results show that the annealing treatment is bene-ficial to improve the mechanical properties.In the meantime,the desolution of SiC-rich solid solution and precipitation of fine TiB 2grains occurred.Due to the precipitation strengthening by the SiC-rich solid solution and the grain boundary strengthening by fine TiB 2particles,the improvement of the mechanical properties was obvious with higher annealing temperatures and longer holding times.The highest HRA,bending strength and fracture toughness were 94.1,517.5MPa and 5.79MPa m 1/2,respectively,after the annealing treatment at 1200°C for 10h.Ó2012Elsevier B.V.All rights reserved.1.IntroductionAlN–SiC solid solution has a wide range of engineering appli-cations due to their excellent mechanical properties at high tem-peratures and has been investigated considerably in the past decades years [1,2].Many preparation methods,such as hot pressing [3],reaction synthesis [4],spark plasma synthesis sinter-ing [5]and combustion synthesis [6],were adopted to fabricate AlN–SiC ceramics.TiB 2is another important material for high temperature applications because of its high melting point,mechanical properties and electrical conductivity,and relatively low coefficient of thermal expansion.So AlN–SiC–TiB 2ceramics have wide potential applications in the field of electrical conduc-tivity at high temperatures.It is known that combustion synthesis (self-propagating high-temperature synthesis,SHS)is a novel method for preparation of AlN–SiC–TiB 2ceramics with many advantages,such as high efficiency and energy saving.However,combustion synthesis is a high temperature and velocity reaction process,which often result in large amount of non-equilibrium phases and high residual stress,and this deteriorates the proper-ties of the combustion synthesized products.So annealing treat-ment at elevated temperatures (800–1200°C)has been introduced to reduce the residual thermal stress and improve the mechanical properties [7,8].In this paper,the effects of annealing treatment on properties of the AlN–SiC–TiB 2ceramics prepared by self-propagating high tem-perature synthesis and hot isostatic pressing (SHS–HIP)were stud-ied.The influences of the annealing temperature and holding time on the composition,microstructure and mechanical properties were discussed in detail.2.Experimental procedureThe AlN–SiC–TiB 2ceramics were prepared by SHS–HIP reported in reference [9].The reaction for preparing AlN–SiC–TiB 2ceramics is as follows:Al þSiC þTiB 2þN 2!AlN þSiC þTiB 2ð1ÞThe samples with Al 35wt%–SiC 35wt%–TiB 230wt%were chosen to carry out the annealing treatment tests.The annealing treatments were carried out on the box-type electric furnace (SXK-8-16,Longjiang electrical furnace works,Harbin,China)in vacuum atmosphere.The specimens were enclosed and vacuumed in quartz glass tube.During the annealing treatment test,four different temperatures (600,800,1000and 1200°C)and five different holding times (2,4,6,8and 10h)were adopted to study the effects of annealing temperature and holding time on the microstructure and mechanical properties of the AlN–SiC–TiB 2ceramics.The composition of the annealing treated samples was analyzed by X-ray dif-fraction (D/max-rB,Rigaku,Japan).The microstructures were observed by scanning electron microscope (SEM,HITACHI S-4700),energy dispersive X-ray spectrometer,and transmission electron microscopy (TEM,Philips CM12/STEM,Holland).The specimen of 3mm Â4mm Â36mm and 2mm Â4mm Â22mm in dimension for bending strength and fracture toughness tests were sliced respectively.The three-point bending strength and fracture toughness were tested on electronic uni-versal test machine with crosshead speed of 0.5mm/min (Instron5569,USA).The Rockwell hardness test was carried out with a ½00ball and 60kg load applied for 30s on the Rockwell hardness test machine (HR-150A,China).0925-8388/$-see front matter Ó2012Elsevier B.V.All rights reserved./10.1016/j.jallcom.2012.11.080Corresponding author.Tel./fax:+865332781357.E-mail address:zhoulijuan@ (L.Zhou).3.Results and discussion3.1.MicrostructureFig.1shows the XRD patterns of theafter the annealing treatment for10h atAs thefigure indicated,oxidationnot occurred and the main phases wereTiB2after the annealing treatment.temperature and nitrogen pressure duringcontent of hexagonal boron nitride(h-BN)ucts.With the increasing of the annealingtion intensity of h-BN phase decreased,TiN phase appeared at the annealing10h,and slightly increases as the annealingto1200°C.The formation of the smallfavorable to the improvement of thermalfracture toughness because of theirand micro-cracking interactionkind of structural ceramic with excellenttion of h-BN phase during preparation anding treatment will not decrease the properties of the products,onthe contrary,it is helpful to the improvement of mechanical properties.Fig.2shows the XRD patterns of AlN–SiC–TiB2ceramics after annealing treatment at1200°C for different holding times.It can be found that the composition of the annealing treated ceramics was almost unchanged at1200°C for8h.When the holding time increased to10h,a small quantity of TiN was formed in the prod-uct,which was consistent with the XRD results shown in Fig.1. Furthermore,the X-ray diffraction peaks of2H–AlN–SiC phase were becoming narrower and sharper with the increasing of the annealing holding time.Meantime,the XRD peaks of the2H–AlN–SiC phase appeared to shift to larger angles,which indicated 3.2.Strengthening and toughening mechanismFig.3shows the BSE images of AlN–SiC–TiB2samples annealed for10h at600and1200°C.AlN–SiC matrix shows gray color and the dispersed white particles are TiB2phase.Fig.3a shows the sam-ple surface annealed at600°C for10h.White columnar TiB2were dispersed inhomogenously in the matrix,and also the sample sur-face presented many pores with different sizes and shapes.When the annealed temperature was1200°C,the dispersion of TiB2 was relatively homogenous andfiner,and the porosity of the sur-face was relatively decreased,as shown in Fig.3b.It is noted that some light gray phase exists between gray ma-trix and white particles,as shown in Fig.4.Line scanning for the light gray region was conducted to determine the phase composi-tion,as shown in Fig.4.The results show that the light gray phase contained large amount of Si element and tiny content of Al,and no Ti element was detected,which can be indicated that the light gray phase was the SiC-rich phase precipitated during the annealing treatment.After the precipitation of SiC,the AlN-rich and SiC-rich solid solutions were formed correspondingly,which was consis-tent with the phase diagram analysis[10].Phase separation leads to grain size refinement,and then results in the improvement of the yield strength.According to Hall–Petch equation,the relationship between the yield strength and grain size is described as follows[11]:ry¼r0þkdÀ1=2ð2Þwhere r y is the yield strength,r0is a materials constant for the starting stress for dislocation movement,k is the strengthening coefficient,and d is the average grain diameter.As the equation indicated,the grain size refinement is in favor of the improvement of yield strength of the materials.Fig.5shows the TEM image and SAED pattern of AlN–SiC–TiB2 ceramic annealed at1200°C for10h.Light gray matrix was AlN–SiC and the black particles embedded in the matrix were TiB2 phase,as shown in Fig.5a.The SAED pattern of the TiB2particles and the calibration were shown in Fig.5b and c,respectively.The average size of these TiB2particles was much less than those in reactant,and also,these particles were completely embedded in the AlN–SiC matrix,which indicated that thefine TiB2particles were precipitated during the annealing treatment.Since combustion synthesis is a high temperature and velocity reaction process,many dislocations,grain boundaries andFig.1.XRD patterns of AlN–SiC–TiB2ceramics annealed at different temperatures for10h.Fig.2.XRD patterns of AlN–SiC–TiB2ceramics annealed at1200°C for different holding times.500L.metastructures can be formed in the products.In addition,boron and titanium could dissolve in the AlN–SiC matrix or in the grain boundaries because of high reaction temperature.What is more,the standard enthalpy of formation (D H )of TiB 2is relatively low (D H =À293kJ mol À1),and the fine TiB 2particles can be precipi-tated from the supersaturated and non-equilibrium AlN–SiC ma-trix during the high temperature annealing treatment.Precipitating of these fine TiB 2particles decreased the lattice distortion,which also released the inner stress and improved the order degree of the lattice.Meantime,composition gradient formed by the precipitation between the matrix and TiB 2phases enhanced the interface bonding strength.In summary,the improvement of the strength and toughness of the AlN–SiC–TiB 2ceramics can be attributed to the particle dis-persion strengthening effect of the precipitated SiC-rich solid solution.And also the fine TiB 2particles precipitated uniformly at the phase interface enhanced the strength of the grain boundaries.nano-particles in AlN–SiC–TiB 2ceramic annealed at 1200°C for 10h:(a)TEM image of AlN–SiC–TiB 2composite;(b)SAED of light gray phase in the AlN–SiC–TiB 2ceramic annealed at 1200°C for 10h:(a)SEM image of the surface;(b)line Fig.3.Backscattered electron images of AlN–SiC–TiB 2samples annealed at (a)600°C,10h;(b)1200°C,10Fig. 6.Typical SEM fractograph in annealing treated AlN–SiC–TiB2ceramic composite.Fig.7.Rockwell hardness of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.Fig.8.Bending strength of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.Fig.9.Fracture toughness of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.annealing temperature reached800and1200°C.The highest HRA,bending strength and fracture toughness were94.1, 517.5MPa and 5.79MPa m1/2,respectively,after the annealing treatment at1200°C for10h.4.ConclusionsThe effects of annealing treatments on the microstructure and mechanical properties of AlN–SiC–TiB2ceramics prepared by SHS–HIP were studied.The annealing treatment can promote the precipitation SiC-rich solid solution in AlN–SiC matrix andfine TiB2particles on the grain boundaries,which improve the mechan-ical properties obviously.With the annealing temperature increas-ing from600to1200°C and holding times increasing from2to 10h,the strengthening effect was more evident.The highest HRA,bending strength and fracture toughness of AlN–SiC–TiB2 ceramics were94.1,517.5MPa and5.79MPa m1/2,respectively, after the annealing treatment at1200°C for10h.AcknowledgmentsThis work has beenfinancially supported by‘‘A Project of Shandong Province Higher Educational Science and Technology Program(J09LD01)’’and‘‘Shandong Province Natural Science Foundation(ZR2010EM045)’’.References[1]L.J.Zhou,Y.T.Zheng,S.Y.Du,H.B.Li,Mater.Sci.Forum546–549(2007)1505–1508.[2]vrenko,J.Desmaison,A.D.Panasyuk,M.Desmaison-Brut,E.Fenard,J.Eur.Ceram.Soc.25(2005)1781–1787.[3]K.Strecker,M.J.Hoffmann,J.Eur.Ceram.Soc.25(2005)801–807.[4]X.M.Yue,G.J.Zhang,Y.M.Wang,J.Eur.Ceram.Soc.19(1999)293–298.[5]M.Hotta,J.Hojo,J.Eur.Ceram.Soc.30(2010)2117–2122.[6]L.Mei,J.T.Li,Acta Mater.56(2008)3543–3549.[7]A.Varma,A.S.Rogachev,A.S.Mukasyan,S.Hwang,Adv.Chem.Eng.24(1998)79–226.[8]D.Sciti,S.Guicciardi,A.Bellosi,J.Eur.Ceram.Soc.21(2001)621–632.[9]L.J.Zhou,Y.T.Zheng,S.Y.Du,Key Eng.Mater.353–358(2007)1517–1520.[10]A.Zangvil,R.Ruh,J.Am.Ceram.Soc.71(1988)884–890.[11]H.Conrad,J.Narayan,Scripta Mater.42(2000)1025–1030.L.Zhou et al./Journal of Alloys and Compounds552(2013)499–503503。

顺丁橡胶的链结构主要包括顺、反-1，4-结构和1，2-结构单元，还包括重均相对分子质量（M w）、数均相对分子质量（M n）和相对分子质量分布（M w/M n）等。

高聚物的微观链结构直接影响其聚集态的结构，进一步对宏观的物理机械性能产生重要的影响。

发现随着顺-1，4链节含量的增加，顺-1，4链节含量可高达99%，顺丁橡胶生胶的结晶速率增快，硫化胶的强度明显提升，在顺式含量高于98.5%时，对混炼胶的挤出性能和自粘性能没有产生明显的影响[1]。

顺丁橡胶的顺-1，4-结构含量影响顺丁橡胶硫化胶的物理机械性能，当顺式-1，4 结构含量由 96.4%提高至98.6%时，其硫化胶拉伸强度提高了20.7%，断裂伸长率提高了 25.8%，在65℃时tanδ值降低了11.1%，在 0℃时tanδ值升高了5.1%，T g 降低了5.5℃，因此，顺丁橡胶顺式-1，4 结构含量由 96.4%提高至 98.6%时，可以明显提高硫化胶的物理机械性能，同时降低滚动阻力及生热、提高抗湿滑性并提高耐磨性[2-4]。

分子量是判断橡胶性能和加工行为的重要依据[5]，一般来讲，橡胶的大部分物理机械性能随着分子量的增加而提高，但是当分子量达到一定数值后，由于分子链过长，分子链的体积庞大，往往容易发生缠结，导致橡胶的弹性下降，门尼粘度增大，反而对加工性能产生不利的影响。

所以为了取得更好的加工性能和综合使用性能，必须对高聚物的分子量进行一个合理的控制。

橡胶是分子量大小不一的同系物的混合体系，所以，整个体系的分子量会呈现出很大的分散性，分子量太高或者太低都会导致硫化胶的性能变差，因此，必须进一步掌握橡胶的分子量分布指标。

采用钕催化剂合成较高顺式含量、较窄分子量分布的聚丁二烯橡胶，经过研究表明，钕系聚丁二烯具有较好的加工性能、良好的回弹能力及低生热和低滞后损失[6]。

研究了顺丁橡胶随着分子量分布的变窄，即相对分子质量分布指数从4.0降至2.8，抗湿滑性能提高了8%左右，滚动阻力可以降低9%左右。

Effect of differential speed rolling strain on microstructure and mechanical properties of nanostructured 5052AlalloyLoorentz,Young Gun Ko ⇑School of Materials Science and Engineering,Yeungnam University,Gyeongsan 712-749,South Koreaa r t i c l e i n f o Article history:Available online 5November 2012Keywords:Al alloyDifferential speed rolling MicrostructureMechanical propertiesa b s t r a c tThe present work reported the influence of differential speed rolling (DSR)strain on microstructure and mechanical properties of the nanostructured 5052Al alloy.As the amount of DSR strain increased,the deformed microstructure developed from the band-like structure of the elongated grains after one-pass DSR (%0.4)into the nanostructure of the equiaxed grains whose mean size of %700nm after four-pass DSR (%1.6).This was attributed to the fact that,by a sample rotation of 180°along the longitudinal axis,the macro shear deformation formed by one-pass DSR was intersected with that by two-pass DSR.From the microhardness contour maps of the DSR-deformed samples,the microhardness values and their uni-formity were improved with increasing amount of DSR strain.Tensile test results showed that,as the amount of DSR strain increased,the tensile strength increased significantly while sacrificing tensile duc-tility and strain hardenability.Such mechanical response of the nanostructured 5052Al alloy was dis-cussed in relation to microstructure evolution during DSR.Ó2012Elsevier B.V.All rights reserved.1.IntroductionThe processing of bulk metallic metals by means of severe plas-tic deformation (SPD)techniques has been generating great inter-est in recent years because the nanostructured materials fabricated via SPD methods,such as high pressure torsion (HPT)and equal channel angular pressing (ECAP)possessed superior mechanical properties to their coarse grained counterparts [1–6].Asymmetri-cal rolling was one of the continuous SPD techniques suitable for achieving severe grain reduction below the micrometer level,to-gether with a deep industrial potential.Among asymmetrical roll-ing methods,a differential speed rolling (DSR)was known to be desirable for enhancing the mechanical properties of the workpiec-es.DSR was one of the rolling methods utilizing two identical rolls in size where each was driven by its own motor,generating the dif-ferent rotation speeds of upper and lower rolls,so that the shear strain could be imposed uniformly through the sheet [7,8].In this regard,active research endeavors have been made re-cently,and successful applications have been reported for various materials such as Fe [9,10],Al [11,12],Ti [13,14],etc.For instance,Jiang et al.[11]demonstrated the use of DSR method resulted in severely refined grains of pure Al.Kim et al.[13]reported that the excellent combination of ultrafine grained structure and high tensile properties of commercially-pure Ti was attained by control-ling the speed ratio and deformation temperature during DSR.De-spite these previous investigations,however,a systematic study on how DSR strain influences microstructure evolution and mechani-cal properties of Al alloy will be needed.Therefore,the main pur-pose of the present work is to study the effect of amount of strain on microstructural development of Al alloy fabricated via DSR.The mechanical properties of the DSR-deformed Al alloy sam-ples are also investigated.2.Experimental proceduresThe material used in this study was a 5052Al alloy sheet with a chemical composition of 2.2Mg,0.2Cr,0.4Fe,0.25Si,0.028Ti and the balance Al in wt.%.The as-received microstructure was homogenized at 823K for 30min followed by air cooling,resulting in a coarse grained microstructure whose grain size was %95l m as shown in Fig.1(a).Prior to DSR,the sample was machined into the plate type with a dimension of 70Â30Â4mm.The principle and direction of DSR oper-ation were depicted in Fig.1(b).The diameters of the two rolls in DSR equipment were identical as 220mm.The DSR processing was performed at a roll speed ratio of 1:4for the lower and upper rolls,respectively,while the velocity of the lower roll was fixed at %3.4m/min.The sample was subjected to four-pass DSR operations with a height reduction of 30%for each pass,corresponding to the total strain of %1.6.Each sample was rotated 180°around its longitudinal axis between passes.Poulton’s reagent was used to etch the sample for optical observation.For trans-mission electron microscope (TEM)observations,the thin foils were cut from the normal direction (ND)-rolling direction (RD)plane of the deformed samples where the effect of shear deformation on microstructure evolution was clearly shown as reported earlier [15].TEM micrograph and corresponding selected area electron dif-fraction (SAED)pattern were taken by using TEM (Hitachi H-7600)operating at 120kV.Vickers microhardness tests were conducted on the ND-RD plane of the DSR-deformed samples with a load of 100g and a dwelling time of 10s.A series of individual results obtained from the polished sections with a gap of %0.2mm were recorded.These values were then plotted in the form of the contours depicting0925-8388/$-see front matter Ó2012Elsevier B.V.All rights reserved./10.1016/j.jallcom.2012.10.128Corresponding author.Tel.:+82538102537;fax:+82538104628.E-mail address:younggun@ynu.ac.kr (Y.G.Ko).the distribution of the microhardness over the ND-RD plane of samples.Tensile test was performed at room temperature on the dog-bone sample with a gauge length of 25mm and a width of6mm at a constant rate of crosshead displacement with an initial strain rate of10À3/s.3.Results and discussion3.1.MicrostructureFig.2shows the optical micrographs taken from the ND-RD plane of the DSR-deformed samples as a function of DSR strain. In spite of the high roll speed ratio of1:4used in this study,no obvious plastic failure of the samples such as surface crack and wrinkle was detected with increasing DSR operations,which was responsible for the excellent cold-workability of5052Al alloy.As apparent from Fig.2(a),the microstructure developed into the coarse elongated grains parallel to the DSR deformation direction after one-pass DSR,leading to the band-like structure with a thick-ness of%40l m.As the amount of DSR strain increased,the micro-structural observation shown in Fig.2(b)–(d)revealed that the thickness of the band structures became slender and the contour of the band boundaries was likely to be indistinct due to high amount of DSR strain.A similar trend was also found in the previ-ous study[7].Tofigure out the details of microstructural features,the bright-field TEM and SAED pattern images of the deformed samples are shown in Fig.3.The deformed microstructures tended to vary with respects to observing area and DSR strain.After one-and two-pass DSR operations,the microstructure evolution was observed to be gradual from top to bottom regions.As the amount of DSR strain increased,however,the microstructure tended to be reasonably uniform.Thus,TEM images which were obtained from the middle region of the sample were displayed in Fig.3.since the middle region represented the whole deformed microstructure.After one-pass DSR,the microstructure was mainly comprised offine lamellar bands of elongated subgrains with a width of%1l m. Due to the low-angle misorientation of the band boundaries in nat-ure which was confirmed by the individual regular spots in the SAED pattern,they seemed to be invisible through optical observa-tion.Thus,the amount of strain imposed by a single DSR was insuf-ficient to induce the formation of nanostructure having the high misorientation.Numerous dislocations were mainly detected in the vicinity of subgrain boundaries while the dislocation density was comparatively low in the matrix.As shown in Fig.3(b),the microstructure after two-pass DSR showed the equiaxed subgrains whose size was comparable to the width of lamellar bands fabricated by one-pass(%1l m).The SADP spots of the deformed sample were diffused,suggesting the fact that a misorientation dif-ference between subgrains begun to increase without a significant further reduction in grain size in order to accommodate the intense plastic strain.In Fig.3(c),it was observed that the elongated grains appeared after three-pass DSR,which was similar to that after a single pass in terms of grain morphologies,but both the width and length of the elongated grains became smaller.By four-pass DSR(Fig.3(d)),the deformed microstructure was consisted of nearly equiaxed nanostructured grains of%0.7l m,whichwere Initial microstructure of5052Al alloy and(b)schematic illustration of DSR machine and sample rotationsmaller than those by two-pass DSR.The appearance of the addi-tional rings and extra spots in SAED pattern implied the formation of high-angle boundaries.The resulting grain sizes in this study were quite comparable to the grain sizes fabricated by other SPD techniques[16,17].The development of nearly equiaxed nanostructured grains might be addressed by the fact that the macro shear bands formed by one-pass DSR crossed those by two-pass DSR,as illustrated in Fig.4. This was associated with a sample rotation of180°along the lon-gitudinal axis,allowing the elongated subgrains by odd-numbered pass to restore their original equiaxed segments after even-numbered pass in order to accommodate the intense plastic strain.Optical images of the deformed samples after(a)one-pass,(b)two-pass,(c)three-pass,and(d)four-pass and SAED pattern images of the deformed samples after(a)one-pass,(b)two-pass,(c)three-pass,and(d)Thus,the equiaxed grains would be achieved after each even-addition to the morphological change the started to be significantly diffused as the increased,which indicated a gradual incre-high-angle boundaries.The formation of boundaries would be presumably attributed to during multi-pass DSR operations.The num-formed by initial DSR deformation would with low misorientation and,thereby,they absorbed by the subgrain boundaries,resulting 3.2.Mechanical propertiesThe microhardness contour maps depicting distribution (or microstructural uniformity the ND-RD plane of the DSR-deformed Fig.5.The average microhardness value to DSR was %60Hv.As shown in Fig.5(a),of the deformed sample after one-pass DSR high rate and the high microhardness value detected in the upper side of the deformed Fig.4.Schematic illustration of shearing during multi-pass DSR operations.Microhardness contour maps of the deformed samples after (a)one-pass,(b)two-pass,(c)three-pass,and (d)achievement of the microhardness homogeneity,the effect of sam-ple rotation during multi-pass DSR should be taken into account.Since the sample was rotated around 180°along their longitudinal axis between each pass,the upper side of the sample,which was in contact with the upper roll during odd-numbered pass,was altered to the lower side during even-numbered pass.Therefore,as the amount of DSR strain was evenly distributed,the microhardness distribution was anticipated to be more homogeneous throughout the deformed sample.The engineering stress–strain curves of the DSR-deformed sam-ples and corresponding tensile data are presented in Fig.6and Ta-ble 1,respectively.The yield strength (YS),ultimate tensile strength (UTS),and total elongation of the initial sample were 65MPa,137MPa and 32%,respectively.As the amount of DSR strain increased,YS and UTS increased in a manner similar to the microhardness properties,approaching the maximum values of 380and 390MPa,respectively,whilst losing both tensile ductility and strain hardenability.In case of Al alloys,several strengthening mechanisms associated with grain,dislocation,precipitate and,so-lid solution could contribute to the mechanical strength.According to the earlier works by Straumal et al.[21]and Mazilkin et al.[22],intense plastic strain would lead to the decomposition of supersat-urated solid solution in Al–Mg and Al–Zn alloy samples subjected to HPT,causing the mechanical softening.In contrast,the tensile strength of the present sample processed by DSR was seemed to in-crease with increasing amount of strain.This was attributed to the significant difference in the amounts of Mg and Zn elements be-tween the present and previous studies.In addition,the strain-induced nanoprecipitates were not detected due to lower amount of strain imposed by DSR (%1.6)as compared to that by HPT (%6)[21].Consequently,Hall–Petch and dislocation strengthenings ap-peared to dominate the overall hardening in this study.Tensile strength results exhibited similar behavior to those found in the SPD-deformed materials.Cherukuri et al.[23]reported that the tensile strength of the nanostructured Al–Mg–Si alloy deformed by a multi-axial forging (strain;%6.5)was %350MPa.Indeed,Tsai et al.[24]demonstrated that the use of ECAP (strain;%8)for Al–Mg alloy resulted in a maximum value of %390MPa.In spite of the different strain levels,no significant difference in tensile properties was found between the reported and present results.It is deduced that multi-pass DSR (strain;%1.6)with the sample rotation of 180°around its longitudinal axis was beneficial for attaining the equiaxed nanostructured grains with a fairly uniform distribution as aforementioned,giving rise to high tensile strength.The present study investigated microstructure evolution and mechanical properties of the nanostructured 5052Al alloy pro-duced by DSR with respect to the amount of strain imposed.Since the nanostructured sample still exhibited low ductility and strain hardening,a further investigation on the post-DSR annealing behavior of the nanostructured sample would be necessary to re-store the tensile elongation of the nanostructured Al alloy de-formed by DSR.4.ConclusionsThe effect of DSR strain on microstructure evolution and mechanical properties of the nanostructured 5052Al alloy was investigated.After initial-pass DSR,the band-like structure consist-ing of the elongated grains parallel to the rolling direction ap-peared to form.As the amount of DSR strain increased,the thickness of the band-like structure tended to decrease consider-ably,achieving the nanostructured grains of %700nm in size after four-pass DSR.Hence,the yield strength of the nanostructured sample was approximately five times as high as that of the initial coarse counterpart,approaching a maximum value of 375MPa.References[1]R.Z.Valiev,ngdon,Prog.Mater.Sci.51(2006)881–981.[2]A.P.Zhilyaev,ngdon,Prog.Mater.Sci.53(2008)893–979.[3]X.Huang,N.Kamikawa,N.Hansen,Mater.Sci.Eng.A 493(2008)184–189.[4]Y.G.Ko,C.S.Lee,D.H.Shin,S.L.Semiatin,Metall.Mater.Trans.A 37(2006)381–391.[5]Z.Horita,ngdon,Mater.Sci.Eng.A 410–411(2005)422–425.[6]K.J.Cho,S.I.Hong,Met.Mater.Int.18(2012)355–360.[7]Loorentz,Y.G.Ko,J.Alloys Comp.536S (2012)S122–S125.[8]B.H.Cheon,J.H.Han,H.W.Kim,J.C.Lee,Korean J.Met.Mater.49(2011)243–249.[9]A.Wauthier,H.Regle,J.Formigoni,G.Herman,Mater.Charact.60(2009)90–95.[10]S.H.Lee,D.N.Lee,Int.J.Mech.Sci.43(2001)1997–2015.[11]J.Jiang,Y.Ding,F.Zuo,A.Shan,Scr.Mater.60(2009)905–908.[12]H.Jin,D.J.Lloyd,Mater.Sci.Eng.A 465(2007)267–273.[13]W.J.Kim,S.J.Yoo,H.T.Jeong,D.M.Kim,B.H.Choe,J.B.Lee,Scr.Mater.64(2011)49–52.[14]X.Huang,K.Suzuki,Y.Chino,Scr.Mater.63(2010)473–476.[15]N.Kamikawa,T.Sakai,N.Tsuji,Acta Mater.55(2007)5873–5888.[16]K.T.Park,H.J.Kwon,W.J.Kim,Y.S.Kim,Mater.Sci.Eng.A 316(2001)145–152.[17]C.P.Chang,P.L.Sun,P.W.Kao,Acta Mater.48(2000)3377–3385.[18]L.M.Dougherty,I.M.Robertson,J.S.Vetrano,Acta Mater.51(2003)4367–4378.[19]Y.G.Ko,C.S.Lee,D.H.Shin,Scr.Mater.58(2008)1094–1097.[20]D.H.Shin,I.Kim,J.Kim,K.T.Park,Acta Mater.49(2001)1285–1292.[21]B.B.Straumal,B.Baretzky,A.A.Mazilkin,F.Phillipp,O.A.Kogtenkova,M.N.Volkov,R.Z.Valiev,Acta Mater.52(2004)4469–4478.[22]A.A.Mazilkin,B.B.Straumal,E.Rabkin,B.Baretzky,S.Enders,S.G.Protasova,O.A.Kogtenkova,R.Z.Valiev,Acta Mater.54(2006)3933–3939.[23]B.Cherukuri,T.S.Nedkova,R.Srinivasan,Mater.Sci.Eng.A 410–411(2005)394–397.[24]T.L.Tsai,P.L.Sun,P.W.Kao,C.P.Chang,Mater.Sci.Eng.A 342(2003)144–151.Fig.6.Room-temperature tensile curves of the deformed samples with respect to DSR strain.Table 1Room-temperature tensile properties of the deformed 5052Al alloy samples with respect to DSR strain.Condition Yield strength (MPa)Ultimate tensile strength (MPa)Elongation (%)Initial 65±5137±1032±2One-pass 317±30360±309.4±1Two-pass 345±20381±207.4±1Three-pass 363±15386±15 5.7±0.5Four-pass380±10390±104.2±0.5and Compounds 586(2014)S205–S209S209。

International Journal of Minerals, Metallurgy and Materials Volume 25, Number 11, November 2018, Page 1294https:///10.1007/s12613-018-1682-8Corresponding author: Hamed Jamshidi Aval E-mail:h.jamshidi@nit.ac.ir© University of Science and Technology Beijing and Springer-Verlag GmbH Germany, part of Springer Nature 2018Microstructural evolution and mechanical properties of friction stir-weldedC71000 copper–nickel alloy and 304 austenitic stainless steelHamed Jamshidi AvalDepartment of Materials and Industrial Engineering, Babol Noshirvani University of Technology, Shariati Avenue, Babol, 47148-71167, Iran(Received: 20 February 2018; revised: 29 May 2018; accepted: 11 June 2018)Abstract: Dissimilar joints comprised of copper–nickel and steel alloys are a challenge for manufacturers in modern industries, as these met-als are not thermomechanically or chemically well matched. The present study investigated the effects of tool rotational speed and linear speed on the microstructure and mechanical properties of friction stir-welded C71000 copper–nickel and 340 stainless steel alloys using a tungsten carbide tool with a cylindrical pin. The results indicated that a rotational-to-linear speed ratio of 12.5 r/mm did not cause any macro defects, whereas some tunneling defects and longitudinal cracks were found at other ratios that were lower and higher. Furthermore, chro-mium carbide was formed on the grain boundaries of the 304 stainless steel near the shoulder zone and inside the joint zone, directing carbon and chromium penetration toward the grain boundaries. Tensile strength and elongation percentages were 84% and 65% of the corresponding values in the copper–nickel base metal, respectively.Keywords: dissimilar friction stir welding; copper–nickel alloy; austenitic stainless steel; microstructure; mechanical properties1. IntroductionCopper–nickel alloys exhibit substantial corrosion resis-tance and anti-algae properties against biological sediments. Pure copper is not stable in oxygenated electrolytes, espe-cially in marine and chlorine ion environments where cop-per–nickel alloys are widely used, with copper as the main component [1]. The addition of nickel to copper improves the mechanical strength, durability, and resistance to corro-sion, abrasion, and cavitation in sea and polluted water. This alloy also exhibits significant stress corrosion cracking and corrosion fatigue resistance. Corrosion resistance can be in-creased by adding more nickel to copper–nickel alloys [2]. Since these alloys can be easily assembled and welded, they are prime candidates for plumbing systems, ship bodies, and other marine structures.Generally, stainless steel plays a major role in the modern world. Welding of austenite stainless steel is known for two important properties: maintenance of corrosion resistance and prevention of crack formation. Dissimilar joints of coatings on offshore platform insulators, achieved by different tech-niques, are among copper–nickel plate applications for corro-sion prevention. Other applications include the joining of copper–nickel pipes with steel flanges and/or direct joining of these pipes with steel pipes in marine industries [3].Nevertheless, welding of dissimilar metals is always challenging because of numerous factors. These include different melting points, thermal conductivity, and thermal expansion coefficients; galvanic corrosion; the high solidi-fication rate of molten copper; entry of molten copper into steel grain boundaries (especially in the heat-affected zone (HAZ)); formation of hot cracks; high copper oxidation at high temperatures; and type of filler metal [4−9]. It is essen-tial to select the appropriate filler metal and welding para-meters for dissimilar-metal fusion welding of copper–nickel and stainless steel alloys in order to reduce probable defects (e.g., cavitation and gas cavities).Recent developments in solid-state welding have made it an alternative to fusion welding. In comparison to other welding techniques, friction-stir welding is a solid-state technique with an outstanding combination of high speed, precision, and variety. Among different welding methods, friction-stir welding of dissimilar alloys is important due to the ability to join alloys with different properties. In addition,H. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1295)different welding configurations in this method (e.g., lap and butt joints) make it applicable in different situations.Few studies have been conducted on friction-stir welding of dissimilar copper and stainless steel alloys. In this regard, Imani et al. [10] investigated a pure copper and stainless steel joint with a thickness of 3 mm using friction-stir weld-ing. It was found that the tool offset toward the copper side played a significant role in eliminating defects in the joints. In addition, Ramirez et al. [11] examined the effects of tool offset on the microstructure and mechanical properties of joints in friction-stir welding of pure copper and 316 stain-less steel with a thickness of 2 mm. They studied 0, 0.6, and 1.6 mm offsets relative to the joint interface. When a major part of the tool was on the steel side, the joint efficiency was 55% of copper base metal. Maximum joint efficiency, i.e., 87% of copper base metal, was reported in the 0.6-mm off-set relative to the joint interface.Furthermore, Najafkhani et al. [12] studied the joint of pure copper and 316 stainless steel with a thickness of 5 mm using friction-stir welding. In their study, all joints cracked from the heat-affected zone of the copper base metal. The highest tensile strength and elongation percentage were 220 MPa and 7%, respectively. In addition, Shamsujjoha et al. [13] studied the lap joint of pure copper with 1018 carbon steel using friction-stir welding. They found that the joining process at the interface was both mechanical and metallur-gical. Jafari et al. [14] also studied the friction-stir welding of pure copper and 304 stainless steel with a thickness of 3 mm. The heat input from the welding increased the grain size in the heat-affected zone and decreased joint ductility by increasing the number of welding passes.According to the literature, there are no studies on the friction-stir welding of copper−nickel and austenite stainless steel alloys. Accordingly, the present study investigated the effects of process parameters on the microstructure and me-chanical properties of friction stir-welded C71000 cop-per−nickel and 304 stainless steel alloys using a tungsten carbide tool with a cylindrical pin. Optical microscopy and scanning electron microscopy (SEM) were used to study the microstructure and detect the created phases in different zones. The mechanical properties of joints were also eva-luated by tensile and microhardness tests.2. ExperimentalIn the present study, C71000 copper−nickel and 304 aus-tenite stainless steel plates with thicknesses of 2 mm were used. Both plates were cut perpendicular to the rolled metal direction and had a dimension of 50 mm × 100 mm. The chemical compositions and mechanical properties of alloysare listed in Tables 1 and 2. The plates were welded in a buttjoint configuration. The copper−nickel alloy was on the re-treating side, while the stainless steel alloy was on the ad-vancing side. According to the literatures [10−11], 0.75 mmof the tool axis was offset to the copper−nickel alloy relativeto the joint interface. Fig. 1 shows the schematic of the tool offsetting procedure. A tungsten carbide tool with a cylin-drical pin with a height of 1.8 mm was used for welding. Fig.2 demonstrates the dimensions and geometry of the appliedtool in welding and Table 3 indicates the welding parameters.The present study selected two rotational speeds of 800 and1000 r/min and three linear speeds of 40, 60, and 80mm/min.Table 1. Chemical composition of alloy wt%ZnMnFeCrCuCNiAlloy0.90.010.05―Base0.0519.12C71000―1.20Base18.500.440.058.10SS304Table 2. Mechanical properties of alloysAlloyUltimate tensilestrength / MPaYieldstrength / MPaMicro-hardness,HV0.1Elonga-tion / %C71000338 110 9032 SS304585 210 15242Fig. 1.Schematic illustration of friction stir butt welding.Fig. 2. Tool geometrical characteristics.1296Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Table 3. Friction stir welding process parametersRotational-to-linear speed ratio / (r ⋅mm –1)Linear speed / (mm ⋅min –1)Rotational speed / (r ⋅min –1)Sample No.20.00 40 800113.33 60 800 2 10.00 80 800 3 25.00 40 1000 4 16.65 60 1000 5 12.508010006The samples were transversely cut for metallographic studies. A marble solution was used for etching the micro-structure on the stainless steel side after sanding and polish-ing, whereas a nitric acid and distilled water solution was used for the copper–nickel alloy. Scanning electron micro-scopy (SEM) and X-ray diffraction (XRD) were used to evaluate the joint interface and examine the distribution and type of intermetallic compounds in the joint cross section. The mechanical properties of the joint were investigated us-ing a tensile test according to the ASTM E8-M03 standard. The tensile test was carried out at a crosshead speed of 1 mm/min. A Vickers microhardness testing machine with a load of 3 N and test time of 15 s was used to evaluate the hardness distribution of a joint cross section.3. Results and discussion3.1. Weld appearanceThe qualitative test of the welded samples indicated that samples No. 1–5 had defects. Longitudinal cracks on the copper–nickel side or tunneling defects on the stainless steel side were observed in all defected samples. Fig. 3 shows the effects of rotational and linear speed on the appearance of welding samples No. 1, 3, 4, and 6, as representative sam-ples containing cracks, tunneling defects, and defect-free welds. It is generally difficult to explain the causes of de-fects in the samples; however, the heat input may be an in-fluential factor. Many researchers have introduced various analytical, numerical, and empirical models in order to evaluate the relationship between rotational and linear tool speed and heat input and to examine their effects on the temperature distribution in the friction-stir welding proce-dure. With a proper estimation, the rotational-to-linear speed ratio can be considered a measure of welding heat input.In this study, samples No. 3 and 4 received the least and most heat input, respectively. The lower temperature of sample No. 3 caused insufficient material flow into the stir zone. After the tool was moved forward, the flow of material stopped before arriving at the advancing side. Therefore, there was inadequate material to fill the hole on the advanc-ing side (stainless steel). The tunnel hole led to the loss of joint strength in this sample, and the two parts were easily separated.Fig. 3. Surface appearance of welded samples: (a) No. 4; (b) No. 1; (c) No. 6: (d) No. 3.Fig. 3 presents the longitudinal cracks because of a tunneling defect in sample No. 3. Fig. 4 shows the effect of welding heat input on longitudinal crack length. It can be seen that by increasing the rotational-to-linear speed ratio (increasing welding heat input), the maximum temperature in the joint increased, which led to the higher temperature gradient in the welded samples. The significant difference in thermal conductivity of copper–nickel and stainless steel al-loys (thermal conductivity of copper–nickel is 2.8 times higher than that of stainless steel) [15–16] at a high temper-ature gradient produced longitudinal cracks as a result of thermal stress in the joint. According to the visual inspection of welded samples, a tunneling defect developed in the joint at a rotational-to-linear speed ratio of less than 10.00 r/mm.H. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1297)On the other hand, at the rotational-to-linear speed ratio of 13.33 r/mm or higher, longitudinal cracks were formed at the joint interface.3.2. Macrostructure and microstructureThe evaluation of mechanical and metallurgical proper-ties was only carried out for sample No. 6 because it had no defects. The macrostructure of the joint and microstructure of different zones are shown in Figs. 5 and 6. The micro-structure of stainless steel included austenite and δ-ferrite with a grain size of (40 ± 5) µm (Fig. 5(b)). Although the quantity of ferrite phase was not significant, the presence of δ-ferrite could improve the formation of the sigma phase inalloys during friction-stir welding [17].Fig. 4. Effect of the rotational-to-linear speed ratio on cracklength.Fig. 5. Optical images of different zones of sample No. 4: (a) macrostructure of welded sample No. 4; (b) base metal of AISI 304; (c) base metal of C71000; (d) TMAZ in AISI 304 side; (e) SZ in AISI304 side; (f) TMAZ in C71000 side as marked by zone I in (a).1298 Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Fig. 6. Microstructure of stir zone of sample No. 4: (a) microstructure of zone II in Fig. 5(a); (b) SZ in C71000 side; (c) SEM image of zone I in (a); (d) SEM image of zone II in (a).The copper–nickel microstructure had a grain size of (50 ± 4) µm and an average particle size of (10 ± 3) µm in the grain boundaries. The results of energy dispersive X-ray spectroscopy (EDS) indicated that these particles were nick-el-rich oxides with iron and zinc (Fig. 7). The stir zone mostly consisted of copper–nickel alloy, which was likely due to the lower flow stress of copper–nickel alloy [18] and location of the main part of the tool on the copper–nickel side. Different behaviors of the two alloys in the etchant so-lution confirmed this finding.Fig. 7. Element mapping result of base metal C71000 alloy.As shown in Fig. 5(a), a steel layer was drawn from the advancing zone to the retreating zone (zone I). The joint cross section as a result of friction-stir procedure consisted of the stir zone (SZ), thermomechanically affected zoneH. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1299)(TMAZ), and heat-affected zone (HAZ). The mechanical behaviors of the welding materials, especially the welding zone hardness, were affected by high plastic deformation and high temperature during the friction-stir welding. The stir zone microstructure in the friction-stir weld had smaller and equiaxed grains in comparison with the base metal due to high plastic deformation and stir resulting from the tool pin.As presented in Fig. 5(f), grains in the thermomechani-cally affected zone on the copper–nickel side were elon-gated, which is exclusive to this zone [19]. The steel layer on the copper–nickel side contained recrystallized cop-per–nickel grains (Fig. 5(f)). On the other hand, the stir zone microstructure on the copper–nickel side contained equiaxed grains with a size of (15 ± 4) µm as a result of dynamic re-crystallization in this zone (Fig. 6(b)). The oxide particles observed in the copper–nickel base metal are shown in this figure. These particles were mainly at grain boundaries with a size of (5 ± 2) µm and prevented the growth of stir-zone grains.The EDS results showed that zinc and iron concentra-tions in the oxide particles increased (Figs. 7 and 8). The stir zone on the steel alloy side contained small recrystallized grains with a size of (5 ± 1) µm (Fig. 5(e)). Clearly, the grain size in the copper–nickel stir zone was greater than that of the steel-stir zone. The temperature and deformation rate in the friction-stir procedure had inverse effects on the grain size of the stir zone. In fact, an increase in the defor-mation rate led to a reduced grain size, and a rise in temper-ature increased the grain size in the stir zone [20].Fig. 8. Element mapping result of stir zone of C71000 side.The advancing side showed the highest temperature and deformation [21]. According to the stir zone microstructure results, the deformation effect was dominant on the steel side, and the grain size of stir zone reduced relative to the copper–nickel alloy side. On the contrary, elongated grains did not exist in the thermomechanically affected zone on the steel side (Fig. 5(d)). However, annealing twins were found across the base metal, whereas there were fewer twins in the thermomechanically affected zone of the steel. There were no twins in the stir zone on the steel side. An interesting point in the microstructure study was the occurrence of a specific layer-by-layer structure at the interface between copper–nickel and steel alloys near the tool shoulder (Fig. 6(a)). The SEM images of different zones in Fig. 6(a) are presented in Figs. 6(c) and 6(d).According to the line scan analysis presented in Fig. 9, the layer-by-layer structures consisted of copper-rich layers adjacent to iron-rich layers. Based on the comparison of the chemical composition of the copper-rich layer and cop-per–nickel base metal, this zone belonged to the cop-per–nickel base metal. However, the iron-rich layer did not match the chemical composition of steel base metal. The highest mass percentages of copper and chromium in the iron-rich layer were 9% and 30%, respectively. The iron-rich layer had a higher copper percentage, which in-creased to 31wt% in some layers.The high percentages of nickel and copper as austenite stabilizers could promote the formation of austenite phase. Generally, welding of austenite stainless steel can cause de-fects, including formation of the brittle phase, hot cracks, and carbide–chrome in grain boundaries. Copper, as an auste-nite-forming element, eliminates the δ-ferrite and sigma phases. Furthermore, the copper–nickel alloy limits the sigma phase by increasing the cooling rate from 600 to 800°C [22].1300Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Fig. 9. SEM image and line scans of chemical elements at the layer structure: (a, b) zone I in Fig. 6(a): (c, d) Fig. 5(f).The sigma phase is very hard and brittle. Its value in-creases by increasing the percentage of chromium, molyb-denum, and silicon, but decreases by increasing the nitrogen, nickel, and carbon contents. Prevention of sigma phase for-mation in stainless steel is difficult when the chrome per-centage is about 20wt%. When the chrome percentage is less than 20wt%, the sigma phase is not observable in auste-nite stainless steels. Due to the very low amount of chrome (up to 9wt%) in the layered structure, formation of sigma phase is not expected.Fig. 10 shows the XRD analysis of the iron-rich zone in the layered structure (point A in Fig. 6(c)); the austenite phase is the only existing phase in this zone. The high per-centage of nickel and copper prevented the formation of sigma phase as expected. The line scan analysis (Fig. 9) in-dicated that nickel concentrations reduced in layer bounda-ries but increased in the iron-rich layers due to nickel migra-tion from the interface to iron-rich layers.According to the EDS results (Fig. 11(a)) regarding point A in Fig. 6(c), the nickel and copper percentages were 24wt% and 21wt%, respectively, indicating the diffusion ofFig. 10. XRD pattern of iron rich layer structure.nickel and copper from the copper–nickel alloy at the inter-face of steel alloy due to the close proximity of this region to the tool shoulder and high temperature of the zone. Ac-cording to the EDS results (Fig. 11(b)), regardless of the in-creased percentage of copper and nickel in the grain boun-daries of the recrystallized zone on the steel side, the high percentage of chrome indicates the increased effect of this element by moving toward the stir zone of the stainless steel. Carbon present in the grain boundaries indicates chrome carbide formation at the joint interface near the tool shoulder.H. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1301)The chemical compositions of these spots indicate that chrome and carbon move toward high-energy zones and form chrome carbide. Formation of carbide and a chrome-free zone around the grain boundary can severely degrade corro-sion resistance of the joint. Analysis of point C (Fig. 11(c)) indicates that this zone belongs to 304 stainless steel. The transient zone in the joint interface can affect the mechanical properties of the joint. Partial diffusion and formation of iron- and copper-rich layers, as shown in Fig. 9, are alsoobserved in zone I of Fig. 5(a).3.3. Hardness evaluationThe joint microhardness profile at the mid-thickness of the weld cross section is presented in Fig. 12. Hardness of the stir zone increases by moving from the steel base metal. According to the Hall-Petch equation, smaller grains have greater hardness; accordingly, hardness increases by de-creasing the grain size and increasing the particle boundary density. Hardness near the interface fluctuates considering the layer-by-layer structure. This structure produces impor-tant features, such as non-uniform hardness profiles and stress concentration zones. The stir zone on the cop-per–nickel side had a more uniform hardness profile and lower quantity. Hardness gradually decreased to the level of copper–nickel base metal by moving toward the cop-per–nickel base metal.3.4. Tensile properties and fractographyThe stress–strain curves for the base metals and joint are shown in Fig. 13. The yield strength and tensile strength of the joint are 103 MPa and 285 MPa, respectively, while elongation is 21%; these values are significantly lower than the corresponding values in the base metals. Tensile strength and elongation of joint were 84% and 65% of the corres-ponding values, respectively in the copper–nickel base metal. It should be noted that fracture occurred in the weld nugget and at the interface of steel and copper–nickel. The hardness profile shows sudden fluctuations, which cause stress con-centrations and joint strength degradation.Fig. 12. Microhardness profiles of cross-section of joint No. 6.Fig. 11. EDS analysis of points A (a), B (b), and C (c) in Fig. 6(c).1302 Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Fig. 13. Stress–strain curve of base metals and welded sample No. 6.The fractured cross section was investigated by SEM af-ter the tensile test. Fig. 14 shows the fractured section and SEM image. The SEM image of the fracture zone shows a brittle cleavage fracture, along with plastic deformation and small uniform holes on the surface. In the brittle cleavage fracture, the crack propagation corresponds to the successive and repeated breaking of atomic bonds along specific crys-tallographic planes. The fracture surface has a faceted tex-ture because of different orientations of the cleavage planes in the grains. In this type of fracture, no substantial plastic deformation occurs and the crack propagates very fast, nearly perpendicular to the direction of the applied stress. In the ductile fracture mode, spherical dimples correspond to microvoids initiating crack formation. Each dimple is half the size of the microvoid, which is formed and then sepa-rated during the fracture process. In the welded sample, brit-tle and ductile failures simultaneously occurred, which could be attributed to the transient zone (Fig. 12) and sud-den fluctuations in the hardness of the sample.Fig. 14. SEM image of fracture surface of the joint No. 6.4. ConclusionsThe present study investigated the friction-stir welding of C71000 and AISI304 stainless steel with a cylindrical pin tool and the following results were obtained.(1) Lack of proper material flow occurred as a result of low temperature at a rotational-to-linear speed ratio of 10 r/mm; therefore, there was not adequate material to fill the hole as the tool traveled forward on the advancing side (stainless steel). In case of rotational-to-linear speed ratio of greater than 20 r/mm, the high heat input produced a higher temperature gradient and resulted in the formation of longi-tudinal cracks as a result of thermal stress in the joint sec-tion.(2) The grain size on the copper–nickel side was larger than that of the stainless steel side. The stirring phenomena during friction-stir welding eliminated annealing twins in the stainless steel base metal and a uniform microstructure with small equiaxed grains formed in the stir zone. Tensile strength and elongation of joint were 84% and 65% of the corresponding values, respectively in the copper–nickel base metal. The fracture surface indicated brittle cleavage and plastic deformation behaviors.(3) Heat and plastic deformation caused element diffu-sion at copper- and iron-rich layers in the stir zone. Nickel and copper, as austenite stabilizers, led to the formation of austenite phase in the iron-rich layers. Chrome and carbon were transferred to grain boundaries, which were high-energy zones, and formed chrome carbide. The layer-by-layer structure and precipitation at the interface made the hardness profile non-uniform and formed possible stress concentra-tion zones.AcknowledgementThe author acknowledges the funding support of Babol Noshirvani University of Technology (No. BNUT/370167/97).References[1] M. Metikoš-Hukovi ć, R. Babi ć, I. Škugor, and Z. 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