Effects of Ag addition on mechanical properties and microstructures of Al-SCu-0.5Mg alloy

第27卷第4期粉末冶金材料科学与工程2022年8月V ol.27 No.4 Materials Science and Engineering of Powder Metallurgy Aug. 2022 DOI:10.19976/ki.43-1448/TF.2022030单级时效处理对2A97铝锂合金组织、力学性能和腐蚀性能的影响游任轩，马运柱，汤娅，赵心阅，刘文胜(中南大学粉末冶金研究院，长沙 410083)摘要：为了确定单级时效制度对2A97铝锂合金组织、力学性能和腐蚀性能的影响，采用室温拉伸、晶间腐蚀、剥落腐蚀、电化学腐蚀和TEM观察等方法，对不同单级时效处理后的合金组织和性能进行表征测试。

结果表明：随着时效温度升高，2A97铝锂合金达到最佳力学性能所需时间缩短。

随时效温度升高、时效时间延长，合金的抗腐蚀性能下降。

165 ℃时效60 h后，合金的抗拉强度、屈服强度和伸长率分别达到549 MPa、484 MPa和8.8%，晶间腐蚀等级为4级，剥落腐蚀评级EC。

关键词：2A97铝锂合金；单级时效；显微组织；力学性能；腐蚀性能中图分类号：TG146.21文献标志码：A 文章编号：1673-0224(2022)04-398-11Effects of single-stage aging treatment on microstructure, mechanical properties and corrosion properties of 2A97 Al-Li alloysYOU Renxuan, MA Yunzhu, TANG Ya, ZHAO Xinyue, LIU Wensheng(Powder Metallurgy Research Institute, Central South University, Changsha 410083, China)Abstract: In order to determine the effects of single-stage aging regime on the microstructures, mechanical properties and corrosion properties of 2A97 Al-Li alloys, the methods of room temperature tensile, intergranular corrosion, exfoliation corrosion, electrochemical corrosion and TEM observation were used to investigate the microstructure, mechanical properties and corrosion properties of 2A97 Al-Li alloys after different single-stage aging treatments. The alloys were tested and characterized. The results show that with increasing the aging temperature, the aging time for 2A97 Al-Li alloys to obtain the best mechanical properties decreases. With increasing the aging temperature and aging time, the corrosion resistance of the alloys decreases. After aging at 165 ℃for 60 h, the tensile strength, yield strength and elongation of the alloys are 549 MPa, 484 MPa, and 8.8%, respectively, the intergranular corrosion is grade 4, and the exfoliation corrosion grade is EC.Keywords: 2A97 Al-Li alloys; single-stage aging; microstructure; mechanical property; corrosion property铝锂合金由于低密度、高比强度、高比刚度和高弹性模量的特点，在航空航天领域有着越来越广泛的应用。

江苏大学硕士学位论文内生颗粒增强铝基复合材料的组织优化与力学性能研究姓名：许可申请学位级别：硕士专业：材料加工工程指导教师：赵玉涛20070501３．２．２Ａ３５６一Ｚｒ（ＣＯ。

）：体系原位制备复合材料的微观组织图３．２ａ为Ａ３５６－Ｚｒ（Ｃ０３）２体系原位反应生成复合材料的ＳＥＭ组织。

图３．２ｂ为其相应的商倍组织，从图中可以看到，白色颗粒弥散分布于基体中，颗粒呈球形或橄榄形，其颗粒尺寸大部分为３ｔｕｎ～４１ｕｎ，部分小于２９ｍ。

经电子探针分析（图３．３），白色颗粒为Ａ１３ｚｒ和Ａ１２０３颗粒。

（ａ）低倍组织（ｂ）高倍组织图３．２Ａ３５６－Ｚｒ（Ｃ０３ｈ体系原位反应生成复合材料的微观组织Ｆｉｇ．３．２ＭｉｃｒｏｓｔｒｕｃｔｕｒｅｏｆｔｈｅｃｏｍｐｏｓｉｔｅｓｆａｂｄｃａｔｅｄｆｒｏｍＡ３５６－Ｚｒ（Ｃ０３ｈｓｙｓｔｅｍ．目Ⅻ㈣Ｗｔ％ｍ鳍Ｊ●Ｅ∞ｌｍ喇￥％Ａｔ％ＯＫ４ｔ２８５４２４继５９７２４５７６纠１００００ｔ００∞ｉ。

（ａ）Ａ１３Ｚｒ相（ｂ）Ａ１２０３相图３．３Ａ３５６－Ｚｒ（ＣＯ，）２体系反应合成复合材料的微观组织中颗粒相分析Ｆｉｇ．３．３ＡｎａｌｙｓｉｓｏｆｐａｒｔｉｃｕｌａｔｅｐｈａｓｅｉｎｔｈｅｃｏｍｐｏｓｉｔｅｓｓｙｎｔｈｅｓｉｚｅｄｉｎｔｈｅｓｙｓｔｅｍＡ３５６－Ｚｒ（Ｃ０３ｈ图３．４ａ为Ａ３５６一Ｚｒ（Ｃ０３）２体系原位反应生成复合材料经［ＩＦ深腐蚀试样的微观组织，清楚的显示了复合材料基体Ａ３５６中的Ｓｉ相形貌：图３．４ｂ为其相应的高倍组织。

由图可见，原位反应所得复合材料中的共晶Ｓｉ相仍以针状形态存在。

（ａ）低倍组织（ｂ）高倍组织图３．４Ａ３５６－Ｚｒ（Ｃ０３）２体系原位合成复合材料的基体Ａ３５６中Ｓｉ相形貌（ＳＥＭ，深腐蚀试样）Ｆｉｇ．３．４ＭｏｒｐｈｏｌｏｇｉｅｓｏｆＳｉｉｎｔｈｅＡ３５６ｍａｔｒｉｘｏｆｔｈｅｃｏｍｐｏｓｉｔｅｓｓｙｎｔｈｅｓｉｚｅｄｆｒｏｍＡ３５６－Ｚｒ（Ｃ０３）２ｓｙｓｔｅｍ（ＳＥＭ，ｄｅｅｐ－ｅｔｃｈｅｄｓａｍｐｌｅ）．３．３Ａ３５６－Ｋ：ＺｒＦ６体系原位制备复合材料的微结构３．３．１＾３５６－ｇ。

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。

固溶-时效对Al-Zn-Mg-Sc-Zr合金板材组织与性能的影响商宝川;尹志民;周向;何振波;林森【摘要】采用力学性能、电导率测试、金相和电子显微分析技术,研究固溶-时效处理对Al-Zn-Mg-Sc-Zr铝合金板材组织与性能的影响.结果表明:Al-Zn-Mg-Sc-Zr 合金板材的最佳热处理制度为(470 ℃,1 h,水淬)+(120 ℃,24 h);在此条件下,合金的抗拉强度、屈服强度、伸长率、硬度和电导率分别为587 MPa、564 MPa、8.95%、155HB和34.5%(IACS);固溶过程中,适当提高固溶温度或延长固溶时间,合金中过剩相逐渐减少,基体过饱和程度增加;时效过程中,固溶体析出η' (MgZn2)和η (MgZn2)相,随时效时间延长,晶内析出相η'粗化,晶界上平衡相也粗化,与此同时,晶界无析出带宽化;合金的高强度来源于微量Sc、Zr引起的亚晶强化、Al3(Sc,Zr)粒子和η'相的析出强化.【期刊名称】《中国有色金属学报》【年(卷),期】2010(020)011【总页数】7页(P2063-2069)【关键词】Al-Zn-Mg合金;固溶;时效;组织;力学性能;电导率【作者】商宝川;尹志民;周向;何振波;林森【作者单位】中南大学,材料科学与工程学院,长沙,410083;中南大学,材料科学与工程学院,长沙,410083;中南大学,材料科学与工程学院,长沙,410083;东北轻合金有限责任公司,哈尔滨,150060;东北轻合金有限责任公司,哈尔滨,150060【正文语种】中文【中图分类】TG146.2俄罗斯全俄轻合金研究院与全俄复合材料研究院合作，在中强可焊Al-Zn-Mg合金基础上，复合添加钪、锆开发了一种新型可焊Al-Zn-Mg-Sc-Zr合金，它是航空航天工业一种有发展前景的高强耐蚀可焊结构材料，还可用作舰船、舟桥、地翼船等的装甲板[1−3]。

7xxx系合金是典型的时效强化型合金，同一合金采用不同的固溶时效处理制度，可使合金的性能产生显著的差异，而无论是过时效处理，还是回归再时效处理，都是在T6处理基础上完成的，因此，确定合理的固溶时效处理制度就变得非常重要。

W机械工程材料IX)I ： 10.11973/jxgccl202101006焊前和焊后调质处理下25C r 2N i 4M o V 钢焊接接头的组织及性能张敏,仝雄伟，李洁，许帅，贾芳(西安理工大学材料科学与工程学院，西安710048)摘要：对比研究了焊前和焊后调质处理条件下25C r 2N i 4M o V 钢焊接接头的显微组织、力学 性能和耐腐蚀性能，调质处理工艺为920 °CX1 h 油淬+580 °C X 2 h 回火，焊接工艺为手工焊条电 弧焊。

结果表明：焊前调质处理的接头焊缝组织为板条马氏体+S-铁素体+ M 23C S 碳化物，焊后调 质处理使焊缝中的S-铁素体溶解，形成了板条马氏体+回火索氏体+M 23C S 碳化物；焊后调质处理 条件下，焊缝中的板条马氏体细小均匀，M 23C 6碳化物呈颗粒状分布于原奥氏体晶界和马氏体板条 晶界处，焊缝的强度、冲击初性和耐腐蚀性能均优于焊前调质处理的。

关键词：25C r 2N i 4M o V 钢；焊缝；调质；显微组织；力学性能；耐腐蚀性能中图分类号：TG444文献标志码：A文章编号：1000-3738(2021)01-0034-07Microstructure and Properties of 25Cr2Ni4MoV Steel Welded Joint underPre-welding and Post-welding Quenching and Tempering TreatmentZHANG Min, TONG Xiongwei, LI Jie. XU Shuai. JIA Fang(School of Materials Science and Engineering ，Xi’an University of Technology, Xi’an 710048，China)Abstract ： M icrostructure, mechanical properties and corrosion resistance of 25Cr2Ni4M oV steel welded jointwere compared and studied under conditions of pre-welding and post-welding quenching and tem pering treatm ents.T he quenching and tem pering proceSvS was oil quenching at 920 °C for 1 h and tem pering at 580 °C for 2 h. The welding process was manual electrode arc welding. The results show that by pre-welding quenching and tem pering, the m icrostructure of the joint weld zone consisted of lath m artensite» netw ork S-ferrite and M 23C 6 carbide. A fter the post-welding quenching and tem pering，the 5-ferrite in the weld was dissolved, and the lath m artensite, tem pered sorbite and M 23C 6 carbide were formed. U nder the post-welding quenching and tem pering condition, the lath m artensite in the weld was small and uniform , and the M 23C 6 carbide distributed in granular shapes on original austenite grain boundaries and m artensite lath grain boundaries ； the strength, impact toughness and corrosion resistance were better than those by the pre-welding quenching and tem pering treatm ent.Key words ： 25Cr2Ni4M oV steel ； weld zone ； quenching and tem pering ； m icrostructure ； mechanicalproperties ； corrosion resistance25Cr2Ni4M〇V 钢中马氏体的形成，但若奥氏体化温 度过高，得到的板条马氏体较粗大[1]。

Gd含量对Al-Zn-Mg-Cu-Zr合金微观组织与力学性能的影响梅飞强;王少华;房灿峰;孟令刚;贾非;郝海;张兴国【摘要】采用铸锭冶金工艺制备6种Gd含量不同的A1-Zn-Mg-Cu-Zr-xGd合金.采用金相观察、力学性能测试、扫描电镜、电子探针及透射电镜等分析手段,研究质量分数x,分别为0%、0.10%、0.15%、0.20%、0.25%和0.30%的Gd对基体合金铸态及时效态显微组织和力学性能的影响.结果表明：Gd含量对A1-Zn-Mg-Cu-Zr合金的微观组织和力学性能的影响显著,当Gd含量低于0.25%时,随Gd含量的增加细化效果、强度以及伸长率都增加；当Gd含量为0.25%时,铸态组织中基体晶粒最小,达到32μm左右；此时T6态合金组织的强度和伸长率达到最高,抗拉强为624.54 MPa,屈服强度为595.00 MPa,伸长率为13.3%,且固溶组织具有良好的抗再结晶作用；而当Gd含量超过0.25%时,合金的微观的组织与力学性能变差.【期刊名称】《中国有色金属学报》【年(卷),期】2012(000)009【总页数】9页(P2439-2447)【关键词】A1-Zn-Mg-Cu-Zr合金;Gd;显微组织;力学性能【作者】梅飞强;王少华;房灿峰;孟令刚;贾非;郝海;张兴国【作者单位】大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024【正文语种】中文【中图分类】TG146.2+1超高强铝合金是20世纪60年代以航空航天材料为背景发展起来的一种高强度铝合金材料[1−3]。

一般将屈服强度为500 MPa以上的铝合金称为超高强铝合金[4]。

多重时效析出第二相对Al-Mg-Si合金电导率的影响袁生平;蒲雄;张国君;刘刚;王瑞红;孙军;陈康华【摘要】电导率的变化能够灵敏地反应Al-Mg-Si合金的时效析出过程,然而溶质原子及时效析出第二相对电导率的单独影响尚不清楚.Al-Mg-Si合金中含有3种成分和形貌不同的第二相.通过实验及模型化系统地研究Al-Mg-Si合金中多重析出第二相对其电导率的影响.结果表明:由于棒状β″相或针状β'相能够分别在473和523 K时有效地阻碍传导电子的移动,因此Al-Mg-Si合金的电导率主要依赖于棒状β″相(T=473 K)或针状β'相(T=523 K)的影响.模型预测结果与实验结果吻合良好,验证了模型的有效性.【期刊名称】《中国有色金属学报》【年(卷),期】2010(020)011【总页数】5页(P2070-2074)【关键词】Al-Mg-Si合金;多重析出;电导率【作者】袁生平;蒲雄;张国君;刘刚;王瑞红;孙军;陈康华【作者单位】西安交通大学,材料科学与工程学院,金属材料强度国家重点实验室,西安,710049;西安交通大学,材料科学与工程学院,金属材料强度国家重点实验室,西安,710049;西安交通大学,材料科学与工程学院,金属材料强度国家重点实验室,西安,710049;西安交通大学,材料科学与工程学院,金属材料强度国家重点实验室,西安,710049;西安交通大学,材料科学与工程学院,金属材料强度国家重点实验室,西安,710049;西安交通大学,材料科学与工程学院,金属材料强度国家重点实验室,西安,710049;中南大学,粉末冶金国家重点实验室,长沙,410083【正文语种】中文【中图分类】TG249.9铝合金具有高的比强度和比刚度等优异的力学性能，广泛地应用于航空航天及汽车领域[1−3]。

然而，随着材料多功能化以及实际生产的需求，铝合金不仅要具备优异的力学性能，而且要具备良好得电学性能。

时效处理对2024-T3搅拌摩擦焊接头组织及性能的影响乔文广1，杨新岐1，董春林2，吴海亮11天津大学材料学院，天津 (300072)2北京赛福斯特技术有限公司，北京 (100024)E-mail：qiaowgsx@摘要：本文对AA2024-T3进行了搅拌摩擦焊，并焊后对接头进行了时效处理，其中时效℃，45/12h℃。

采用了金相、硬度、拉伸、塑性弯曲试验规范为：自然时效35天，180/10h方法对处理过的接头经行了评定。

研究结果表明，三种时效规范的接头，其各区域的组织微℃的焊缝析出的S相观上也未能发现明显的区别，但是经过X-Rad衍射分析，发现45/12h℃的分布趋势一样，只是比其它两种的多。

从接头整体硬度分布来看，自然时效与45/12h℃的接头硬度分布毫无规律，但热机影响区的硬度升高并大于后者整体比前者高；180/10h焊缝硬度。

由于接头存在缺陷，导致三种时效制度的接头抗拉强度、延伸率下降。

关键词：2024铝合金，搅拌摩擦焊，时效处理，组织及性能中图分类号：TG41.引言搅拌摩擦焊是一种新的金属焊接技术，解决了铝合金等低熔点材料的焊接，获得了无气孔、裂缝等缺陷的高质量焊缝[1-2]。

目前，搅拌摩擦焊已经被广泛的应用与航空航天、船舶、铁路、汽车等领域。

国内外评定铝合金FSW接头时，主要考察焊态条件下接头区域微观组织结构、力学性能和疲劳性能，很少考虑焊后时效热处理对FSW接头使用性能的影响。

铝合金的特点是具有时效强化效应，即使在自然放置状态下其性能也会随时间而变化。

因此，研究时效工艺对FSW接头性能的影响具有重要工程意义。

但是，目前，关于基体材料的时效行为在一些国外文献中有介绍[3-13]，但是关于搅拌摩擦焊接头的特别少，尤其是关于2024的，几乎没有。

2.试验方法试验材料为轧制状态的AA2024-T3，其化学成分如表1。

材料原始厚度为2.6mm，为了减少包铝对接头的影响，对材料进行了化学洗涤，化洗后材料的厚度为2.4mm，其机械性能如表2。

第 4 期第 164-175 页材料工程Vol.52Apr. 2024Journal of Materials EngineeringNo.4pp.164-175第 52 卷2024 年 4 月SiC p 分布对SiC p /2024Al 复合材料组织和性能及变形行为的影响Effects of SiC p distribution on microstructure ，mechanical properties and deformation behavior of SiC p /2024Al composites薛鹏鹏1，邓坤坤1*，聂凯波1，史权新1，刘力2（1 太原理工大学 材料科学与工程学院，太原 030024；2 兴县经开区铝镁新材料研发有限公司，山西 吕梁 033603）XUE Pengpeng 1，DENG Kunkun 1*，NIE Kaibo 1，SHI Quanxin 1，LIU Li 2（1 College of Materials Science and Engineering ，Taiyuan Universityof Technology ，Taiyuan 030024，China ；2 Xingxian EconomicDevelopment Zone Aluminum Magnesium New MaterialResearch and Development Company Limited ，Luliang 033603，Shanxi ，China ）摘要：采用半固态搅拌铸造法制备出SiC p /2024Al 复合材料，并通过热挤压和多向锻造（MDF ）的多步变形调控SiC p 的分布，研究SiC p 分布对SiC p /2024Al 复合材料组织和性能的影响。

研究结果表明：热挤压变形促使SiC p 沿挤压方向（ED ）排布；而经过多向锻，SiC p 分布得到显著改善，由定向排布转变为均匀分布。

Trans. Nonferrous Met. Soc. China 22(2012) 262−267Structure stability and mechanical properties of high-pressure die-cast Mg −Al −Ce −Y-based alloyZHANG Jing-huai 1, LIU Shu-juan 2, LENG Zhe 1, ZHANG Mi-lin 1, MENG Jian 3, WU Rui-zhi 11. Key Laboratory of Superlight Materials & Surface Technology of Ministry of Education,Harbin Engineering University, Harbin 150001, China;2. School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, China;3. State Key Laboratory of Rare Earth Resources Utilization, Changchun Institute of Applied Chemistry,Chinese Academy of Sciences, Changchun 130022, ChinaReceived 28 February 2011; accepted 23 June 2011Abstract: With the aim to further improve the mechanical properties of Mg −Al −RE-based alloy, Mg −3.0Al −1.8Ce −0.3Y −0.2Mn alloy was prepared by high-pressure die-casting technique. The microstructure, thermal stability of intermetallic phases and mechanical properties were investigated. The results show that the alloy is composed of fine primary α-Mg dendrites and eutectic in the interdendritic regions. The intermetallic phases in eutectic are Al 11(Ce,Y)3 and Al 2(Ce,Y) with the former being the dominant one. The thermal stability of Al 11(Ce,Y)3 is conditioned. It is basically stable at temperature up to 200 °C within 800 h, while most of the Al 11(Ce,Y)3 intermetallics transform to Al 2(Ce,Y) at higher temperature of 450 °C for 800 h. The alloy exhibits remarkably improved strength both at room temperature and 200 °C, which is mainly attributed to the reinforcement of dendrite boundaries with Al 11(Ce,Y)3 intermetallics, small dendritic arm spacing effect as well as the solid solution strengthening with Y element. Key words: magnesium alloy; Mg −Al −Ce −Y alloy; Al 11RE 3; structure stability; mechanical properties1 IntroductionIn recent years, magnesium alloy in the form of high-pressure die-cast (HPDC) components has attracted much attention in automotive applications as advanced light material. Mg −Al-based alloys such as AZ91D and AM60B are used extensively in some non-critical parts [1]. However, these alloys are unsuitable for more critical components such as transmission and engine parts because of their poor mechanical properties at temperature above 125 °C [2, 3]. Further studies show that it is ascribed to the coarsening of β-Mg 17Al 12 phase in the eutectic region [4] and the discontinuous precipitate of lamellar Mg 17Al 12 from α-Mg matrix [5]. Moreover, a widely accepted view is that both diffusion controlled dislocation climbing and grain boundary sliding are the creep mechanisms in Mg −Al-based alloys [6]. Therefore, the design of alloys with good heatresistance should be based on the reinforcement of both the grain/dendrite boundaries and the α-Mg matrix.The Mg −Al −RE series alloys such as AE42 are the typical heat-resistant alloys designed for lightweight applications in the automotive industry [7]. However, there still appear inconsistencies in the previous reports as to the stability of the main strengthening phase Al 11RE 3 in Mg −Al −RE alloys. PETTERSON et al [8] (AE42 heat treated at 150, 200 and 250 °C for 100 h), HUANG et al [9] (AE42, 450 °C for 15 h), ZHU et al [5] (AE42, 200 °C for 2 weeks), DARGUSCH et al [10] (AE42+1%Sr, 200 °C for 2 weeks) and RZYCHO Ń et al [11] (AE44, 175 °C for 3000 h) reported that the intermetallic phase Al 11RE 3 has high thermal stability, and no decomposition is observed. In contrast, according to the work by POWELL et al [12] (AE42, 175 °C for 1000 h with and without stress) and our recent work [13] (AE44, 200 °C for 100 h under 70 MPa stress), Al 11RE 3 is unstable and partially decomposes to Al 2RE.Foundation item: Project (HEUCFR1128) supported by the Fundamental Research Funds for the Central Universities, China; Project (2010AA4BE031)supported by the Key Project of Science and Technology of Harbin City, China; Projects (20100471015, 20100471046) supported by the China Postdoctoral Science Foundation; Project (LBH-Z09217) supported by the Heilongjiang Postdoctorial Fund, ChinaCorresponding author: ZHANG Jing-huai; Tel: +86-451-82533026; E-mail: jinghuaizhang@ DOI: 10.1016/S1003-6326(11)61169-2ZHANG Jing-huai, et al/Trans. Nonferrous Met. Soc. China 22(2012) 262−267 263In this work, a new heat-resistant HPDC Mg alloy was designed. Relatively low Al content was used to improve the die castability, form high melting point intermetallics and try to avoid the formation of Mg17Al12. Ce with low solid solubility in Mg (0.74%, mass fraction) was added to form large amount of Al−Ce intermetallic particles to reinforce the grain/dendrite boundaries [14]; Y with high solubility in Mg (12.4%, mass fraction) was expected to strengthen the α-Mg matrix [14]. The microstructure, thermal stability of intermetallic and mechanical properties at room temperature and 200 °C of HPDC Mg−Al−Ce−Y-based alloy were investigated.2 ExperimentalThe chemical composition of the experimental alloy was Mg−3.0Al−1.8Ce−0.3Y−0.2Mn (ACY320). The reference alloy prepared under the same condition was Mg−3.3Al−0.2Mn (AM30). Commercial pure Mg and Al were used. Ce, Y and Mn were added in the form of Mg−20%Ce, Mg−20%Y and Al−10%Mn master alloys (mass fraction). Specimens were die cast using a 280 t clamping force cold chamber die-cast machine. About 20 kg raw materials were melted in a mild steel crucible. Pure argon was used as protective gas and refined gas. The molten metal was hand-ladled into the casting machine and the melt temperature prior to casting was about 700 °C. The die was equipped with an oil heating/cooling system and the temperature of the oil heater was set to 220 °C. The chemical compositions of the castings were determined by inductively coupled plasma atomic emission spectrometer.The tensile samples were 70 mm in gauge length and 6 mm in gauge diameter, as shown in Fig. 1. Tensile tests were performed using Instron 5869 tensile testing machine at a strain rate of 1.1×10−3 s−1. The experimental result in the study was the average value of at least four measured specimens. Metallographic sample was cut from the middle segment of the tensile bar. The microstructures and intermetallic phases were characterized by scanning electron microscope (SEM)Fig. 1 Photo of HPDC tensile test bars obtained from ACY320 alloy equipped with an energy dispersive X-ray spectrometer (EDS) and X-ray diffractometer (XRD).3 Results and discussion3.1 MicrostructureIn order to understand the role of Ce and Y additions in the Mg−3Al-based alloy, the microstructure of HPDC AM30 alloy is observed, as shown in Fig. 2(a). It reveals that the AM30 alloy is composed of primary α-Mg dendrites surrounded by interdendritic eutectic. A spot of β-Mg17Al12 intermetallics disperses in the interdendritic regions. The average dendritic arm spacing (DAS) is about 17 μm. The remarkable changes in the microstructure can be observed with the addition of Ce and Y, as shown in Fig. 2(b). The primary α-Mg dendrites in HPDC ACY320 alloy tend to be equiaxed and the average DAS is reduced to 9 μm. More important, the amount of intermetallic compounds in the interdendritic regions is much higher than that in AM30 alloy with the similar Al content.3.2 Intermetallics and their thermal stabilityThe XRD result of the HPDC ACY320 alloy is illustrated in Fig. 3(a). It is known that Mg−Al−Zn (AZ) and Mg−Al−Mn (AM) alloys are mainly composed of α-Mg and β-Mg17Al12 phases [15], the diffraction peaks of β-Mg17Al12 are not emerged due to the addition of Ce and Y. The main secondary phases in ACY320 alloy are Al11(Ce,Y)3 and Al2(Ce,Y), while the diffraction peak of Al11(Ce,Y)3 is much more intense than that of Al2(Ce,Y), which suggests that the former is the dominant one.The magnified SEM images characterizing the secondary phases are shown in Figs. 2(c) and (d). The acicular intermetallics with length of 1−3 μm arrange in a row then the rows form layers roughly, which exhibits outstanding morphology and distribution of the intermetallics. Besides, a few polyhedral intermetallics with size of 1−5 μm can be observed by SEM observations. EDS analyses were used to identify these intermetallics. Figures 2(e) and (f) show the EDS spectra detected from acicular intermetallics and polyhedral intermetallic, respectively. Combined with the XRD result, the acicular intermetallic is identified as Al11(Ce,Y)3 and the polyhedral one corresponds to Al2(Ce,Y). It is worthwhile to note that Al11(Ce,Y)3 has higher content of Ce than Y, while Y content is much higher than Ce in Al2(Ce,Y).To study the thermal stability of Al11(Ce,Y)3 intermetallic, some samples were aged at 200 and 450 °C for 800 h, respectively. The microstructure of ACY320 alloy after heat treatment is shown in Fig. 4. There is no obvious evidence of decomposition of the intermetallics in 200 °C annealed sample, except that a few particlesZHANG Jing-huai, et al/Trans. Nonferrous Met. Soc. China 22(2012) 262−267264Fig. 2 SEM images and EDS analyses of HPDC alloys: (a) SEM images of AM30 alloy; (b, c, d) SEM images of HPDC ACY320 alloy; (e) EDS analyses of acicular intermetallics at point A ; (f) EDS analyses of polyhedral intermetallic at point BFig. 3 XRD patterns of HPDC ACY320 alloys before aging (a), after aging at 200 °C for 800 h (b) and after aging at 450 °C for 800 h (c)show certain coarsening (see Figs. 4(a)−(c)). The XRD patterns also show incognizable changes before and after 200 °C aging by comparing Figs. 3(a) and (b). However, the conspicuous change of microstructure could be observed after age treatment at higher temperature of 450 °C (Figs. 4(d)−(f)). It reveals that most of the acicular intermetallics disappear and give place to fine quadrate-like particles with size of 100−400 nm in interdendritic regions. The XRD pattern illustrated in Fig. 3(c) also changes obviously after 450 °C aging, namely, the intensity of Al 11(Ce,Y)3 peaks decreases and that of Al 2(Ce,Y) peaks increases obviously. In addition, EDS in SEM mode also indicates the Al: (Ce,Y) atom ratio of these small particles is close to 2:1, not 11:3. Therefore, the present study shows a clear evidence ofZHANG Jing-huai, et al/Trans. Nonferrous Met. Soc. China 22(2012) 262−267 265Fig. 4 SEM images of HPDC ACY320 after ageing at 200 °C (a, b, c) and 450 °C for 800 h (d, e, f)the phase transition according to the reaction Al 11(Ce,Y)3 → 3Al 2(Ce, Y) + 5Al [12].The observations reported here suggest that there is a limit to the thermal stability of Al 11RE 3. Up to now, the thermal stability of the main strengthening phase Al 11RE 3 in Mg −Al −RE alloys is still a matter of debate [6, 9−14]. POWELL et al [12] investigated the microstructural stability of die-cast AE42 (Mg −4Al −2RE) alloys which were individually creep tested for 1000 h at 22, 100, 125, 150 and 175 °C. It was reported that the lamellar/acicular phase Al 11RE 3, which dominates the interdendritic microstructure of the alloy, partly decomposes into Al 2RE and Al (forming Mg 17Al 12) at temperature above 150 °C (i.e. 175 °C), and the sharp decrease in the creep resistance is attributed to the reduced presence of the lamellar/acicular Al 11RE 3 and the appearance of Mg 17Al 12. In addition, it was also pointed out that the stress is not a contributing factor to Al 11RE 3 stability. However, ZHU et al [5] reported that intermetallic phaseAl 11RE 3 in die-cast AE42 has high thermal stability with no decomposition observed at temperature up to 200 °C for 336 h, meanwhile, the continuous precipitation of Mg 17Al 12 due to the supersaturation of Al solute was observed in the Mg matrix, which was considered to be responsible for the deterioration in creep resistance at temperatures above 150 °C. In the present work, as for the alloy Mg −3.0Al −1.8Ce −0.3Y −0.2Mn after severe aging treatment at 200 °C for 800 h, neither obvious decomposition of Al 11RE 3 (i.e. Al 11(Ce,Y)3) nor the formation of Mg 17Al 12 is observed. It is considered that following aspects are related to the stability of Al 11RE 3 and the formation of Mg 17Al 12. First, the harsh degree of aging treatment containing aging temperature and time is an important factor for the composition of Al 11RE 3. This may be responsible for the different observation results by POWELL et al (175 °C for 1000 h) and ZHU et al (200 °C for 336 h) in AE42 alloy as well as ACY320 alloy in this study (200 °C for 800 h and 450 °C forZHANG Jing-huai, et al/Trans. Nonferrous Met. Soc. China 22(2012) 262−267 266800 h) at different aging temperatures. Second, the difference of RE components also affects the stability of Al11RE3 [13], which is a factor to explain the difference between Al11RE3 in AE42 and Al11(Ce, Y)3 in ACY320 alloy. Last, the lower Al content in ACY320 (3%) compared with that in AE42 (4%) causes little Al sequestered as soluted in the Mg matrix, which is responsible for no discernible precipitation of Mg17Al12 after aging in ACY320 alloy. The absence of DSC analysis to study phase transition challenges further work.3.3 Mechanical propertiesThe representative tensile stress—strain curves of HPDC ACY320 and AM30 alloys at room temperature and 200 °C are shown in Fig. 5. The tensile properties including ultimate tensile strength, tensile yield strength and elongation to failure are listed in Table 1. The strength increases dramatically by 45−65 MPa both at room temperature and 200 °C, while the elongation still keeps a high level with the addition of Ce and Y to AM30 alloy. The following aspects are considered to be related to the improved strength. The first and also the most important aspect is the formation of large amountFig. 5 Typical tensile stress—strain curves of HPDC alloys at room temperature (a) and 200 °C (b) of Al11(Ce,Y)3 intermetallics which provide considerable reinforcement of dendrite boundaries. Second, since the DAS is finer for the Mg−3Al-based alloys containing Ce and Y, the fine DAS effect can contribute to the observed increase in strength. Another factor is the strengtheningof the α-Mg matrix by solid solution with rare earth elements especially Y. Generally, the ductility is low for the alloy containing large amount of intermetallic particles [16]. However, the elongation of HPDC ACY320 alloy with high volume fraction of intermetallics still keeps a high level. It may be related tothe fine morphology and arrangement of Al11(Ce,Y)3 intermetallics.Table 1 Tensile and compressive properties of HPDC alloys at room temperature and 200 °CRoom temperatureAlloy Tensile yieldstrength/MPaUltimate tensilestrength/MPaElongation/%ACY320 158 255 10 AM30 116 191 9200 °CAlloy Tensile yieldstrength/MPaUltimate tensilestrength/MPaElongation/%ACY320 103 120 22 AM30 60 72 18 4 Conclusions1) With the addition of 1.8% Ce and 0.3% Y to Mg−3%Al-based alloy, the primary Mg dendrites are refined, the intermetallic phase Mg17Al12 is completely suppressed and substituted by Al11(Ce,Y)3 and Al2(Ce,Y) with the former being the dominant one at the interdendritic regions.2) Al11(Ce,Y)3 intermetallics are basically stable at temperature up to 200 °C within 800 h, while most of theAl11(Ce,Y)3 transform to Al2(Ce,Y) at higher temperatureof 450 °C for 800 h. This suggests that the thermal stability of Al11RE3 is conditioned.3) The high-pressure die-cast Mg−3.0Al−1.8Ce−0.3Y−0.2Mn alloy exhibits significantly improved strength at room temperature and 200 °C, which is the results of the reinforcement of dendrite boundaries withAl11(Ce,Y)3 intermetallics, fine dendritic arm spacing effect as well as the solid solution strengthening with Y element.References[1]KULEKEI M K. Magnesium and its alloys applications inautomotive industry [J]. Int J Adv Manuf Tech, 2008, 39(9−10):851−865.ZHANG Jing-huai, et al/Trans. Nonferrous Met. Soc. China 22(2012) 262−267 267[2]WANG Jian-li, PENG Qiu-ming, WU Yao-min, WANG Li-min.Microstructure and mechanical properties of Mg−6Al−4RE−0.4Mnalloy [J]. Transactions of Nonferrous Metals Society of China, 2006,16: s1703−s1707.[3]TONG Guo-dong, LIU Hai-feng, LIU Yao-hui. Effect of rare earthadditions on microstructure and mechanical properties of AZ91magnesium alloys [J]. Transactions of Nonferrous Metals Society ofChina, 2010, 20: s336−s340.[4]BAKKE P, WESTENGEN H. Die casting for highperformance-focus on alloy development [J]. Adv Eng Mater, 2003,5(12): 879−885.[5]ZHU S M, GIBSON M A, NIE J F, EASTON M A, ABBOTT T B.Microstructure analysis of the creep resistance of die-castMg−4Al−2RE alloy [J]. Scripta Mater, 2008, 58(6): 477−480.[6]LUO A A. Recent magnesium alloy development for elevatedtemperature application [J]. Int Mater Rev, 2004, 49(1): 13−30.[7]WANG J, LIAO R, WANG L, WU Y, CAO Z, WANG L.Investigations of the properties of Mg−5Al−0.3Mn−x Ce (x=0−3,wt.%) alloys [J]. J Alloys Compd, 2009, 477(1−2): 341−345.[8]PETTERSEN G, WESTENGEN H, HØIER R, LOHNE O.Microstructure of a pressure die cast magnesium-4wt.% aluminiumalloy modified with rare earth additions [J]. Mater Sci Eng A, 1996,207(1): 115−120.[9]HUANG Y D, DIERINGA H, HORT N, MAIER P, KAINER K U,LIU Y L. Evolution of microstructure and hardness of AE42 alloyafter heat treatments [J]. J Alloys Compd, 2008, 463(1−2): 238−245. [10]DARGUSCH M S, ZHU S M, NIE J F, DUNLOP G L.Microstructural analysis of the improved creep resistance of adie-cast magnesium-aluminium-rare earth alloy by strontium additions [J]. Scripta Mater, 2009, 60(2): 116−119.[11]RZYCHON T, KIELBUS A, CWAJNA J, MIZERA J.Microstructural stability and creep properties of die casting Mg−4Al−4RE magnesium alloy [J]. Mater Charact, 2009, 60(10):1107−1113.[12]POWELL B R, REZHETS V, BALOGH M P, WALDO R A.Microstructure and creep behavior in AE42 magnesium die-castingalloy [J]. JOM, 2002, 54(8): 34−38.[13]ZHANG Jing-huai, YU Peng, LIU Ke, FANG Da-qing, TANGDing-xiang, MENG Jian. Effect of substituting cerium-rich mischmetal with lanthanum on microstructure and mechanicalproperties of die-cast Mg−Al−RE alloys [J]. Mater Design, 2009,30(7): 2372−2378.[14]ZHANG Jing-huai, LIU H F, SUN W, LU H Y, TANG D X, MENG J.Influence of structure and ionic radius on solubility limit in theMg−RE systems [J]. Mater Sci Forum, 2007, 561−565: 143−146. [15]SU Gui-hua, ZHANG Liang, CHENG Li-ren, LIU Yong-bing, CAOZhan-yi. Microstructure and mechanical properties of Mg−6Al−0.3Mn−x Y alloys prepared by casting and hot rolling [J].Transactions of Nonferrous Metals Society of China, 2010, 20(3):383−389.[16]BAE D H, KIM S H, KIM D H, KIM W T. Deformation behavior ofMg−Zn−Y alloys reinforced by icosahedral quasicrystalline particles[J]. Acta Mater, 2002, 50(9): 2343−2356.Mg−Al−Ce−Y基压铸合金的微观结构稳定性和力学性能 张景怀1, 刘淑娟2, 冷哲1, 张密林1, 孟健3, 巫瑞智11. 哈尔滨工程大学超轻材料与表面技术教育部重点实验室，哈尔滨 150001；2. 哈尔滨工业大学材料科学与工程学院，哈尔滨 150001；3. 中国科学院长春应用化学研究所稀土资源利用国家重点实验室，长春 130022摘 要：为进一步提高Mg−Al−RE基合金的力学性能，采用高压压铸技术制备Mg−3.0Al−1.8Ce−0.3Y−0.2Mn合金，并研究其微观组织、金属间相的热稳定性和合金的力学性能。