Influence of annealing conditions on selective oxidation of boron in BH steel sheets

不同引物对和退火温度对盐生植物盐穗木内参基因β-actin扩增效率的影响白雪芹;杨瑞瑞;曾幼玲【摘要】[目的]以盐生植物盐穗木为材料,研究不同引物对和不同退火温度对盐穗木内参基因β-actin扩增效率的影响.[方法]设计扩增β-actin基因的两对引物,标记为β-actin1和β-actin2,分别建立这两对引物扩增盐穗木β-actin基因和特异性引物扩增该物种的过氧化物酶基因POD(盐响应的代表性基因),在不同退火温度下的标准曲线和扩增曲线.[结果]不论55还是58℃退火温度,引物对β-actin1比引物对β-actin2有更好的扩增效率;在55℃退火温度下,引物对β-actin1有更高的扩增效率.在这两个退火温度下,分别以β-actin1和β-actin2引物对扩增β-actin基因作为内参,盐穗木POD基因的相对表达水平具有一定的差异性.[结论]引物和退火温度会影响荧光定量PCR的扩增效率,进而会影响靶基因的相对表达水平.%[Objective] To study the effects of different primers and annealing temperatures on real quantitative PCR amplification efficiency of referenc e gene β-actin with the halophyte Halostachys caspica as researchmaterial.[Method]Standard curves of β-actin and peroxidasegene,POD(representative gene responding to salt stress)from this species were built with well-designed two pairs of primers named β-actin1 and β-actin2 for β-actin as internal reference gene,and a pair of primers for POD gene and amplification curves were obtained under different conditions for these genes.[Result]Results showed that the amplification efficiency of the primers β-actin1 was higher than that of the primers β-actin2 in the Halostachys caspica branches under the 55℃ or 58℃ annealingtemperature,and amplification efficiency at 55℃ annealing temperature was higher than that at 58℃.On the other hand,it was found there were some differences in the relative expression levels of Halostachys caspica POD gene using the internal reference gene β-actin with two pairs of primers β-actin1 and β-actin2 under the 55℃ annealing temperature.[Conclusion]The research indicates that primers and annealing temperatures can influence the fluorescence quantitative PCR amplification efficiency and then affect relative expression level of candidate genes at a certain extend.【期刊名称】《新疆农业科学》【年(卷),期】2017(054)002【总页数】9页(P352-360)【关键词】盐穗木;不同引物对;退火温度;β-actin;扩增效率【作者】白雪芹;杨瑞瑞;曾幼玲【作者单位】新疆大学生命科学与技术学院/新疆生物资源基因工程重点实验室,乌鲁木齐 830046;新疆大学生命科学与技术学院/新疆生物资源基因工程重点实验室,乌鲁木齐 830046;新疆大学生命科学与技术学院/新疆生物资源基因工程重点实验室,乌鲁木齐 830046【正文语种】中文【中图分类】Q786【研究意义】盐穗木 (Halostachys caspica) 为藜科(Chenopodiaceae )多年生盐生植物。

河南大学硕士学位论文Zn<,1-x>Mg<,x>O薄膜的制备与表征姓名：***申请学位级别：硕士专业：凝聚态物理指导教师：***20050501中文摘要氧化锌是一种重要的功能材料和新型的ＩＩ—ＶＩ族宽禁带隙（３．３７ｅＶ）半导体材料。

具有较大的激子束缚能（６０ｍｅＶ），可实现室温下的紫外受激辐射。

因此，氧化锌是一种很有前途的紫外光电子器件材料，极具开发和应用价值。

特别是ＺｎＯ薄膜紫外激射的发现，使它成为国内外在半导体材料研究中的新热点。

由于ＺｎＯ单晶生长困难、价格昂贵、尺寸小（仅有ｌＧｎｌ３大小的单晶），难以满足各种应用的需要，因此对各种ＺｎＯ薄膜制备技术的研究和开发成为ＺｎＯ材料及器件应用研究的一个重要方向。

目前，比较好的成膜技术有分子柬外延（ＭＢＥ）和脉冲激光沉积法（ＰＬＤ）。

但成本太高，不能实现大面积成膜。

本论文采用溶胶一凝胶工艺制备ＺｎＯ及ＺｎＭｇＯ合金薄膜，详细研究了热分解温度、热分解时间、晶化温度、衬底等工艺参数对薄膜的结晶质量和发光性能的影响。

取得了一些有意义的结果。

主要内容如下：１，采用溶胶一凝胶工艺分别在ｓｉ（１００）和石英衬底上制各了ＺｎＯ薄膜，ＸＲＤ结果表明：ＺｎＯ薄膜均为纤锌矿结构。

热分解温度较低时，样品呈（００２）择优生长的特性，随热分解温度的升高，（００２）衍射峰变得更强而且半高全宽减小，当热分解温度为４００。

Ｃ，（００２）衍射峰急剧下降，（１０１）、（ｔｏｏ）衍射峰增强，ＺｎＯ薄膜呈现自由生长的特点；在相同的熟处理条件下，石英衬底上所生长的ＺｎＯ薄膜具有Ｃ柱择优取向生长的特性，而Ｓｉ衬底上生长的ＺｎＯ薄膜则是自由生长。

２．室温下样品光致发光光谱的测量结果表明：所有的样品均有两个发射带，即近带边紫外发射和可见发射带。

随热分解温度的升高，样品的紫外发射带增强，半峰宽减小，可见发射贝ｑ减弱，当热解温度为４０００Ｃ时，紫外发射最尖锐，几乎观察不到可见发射；随着热分解时间的增长，样品的紫外发射显著增强，可见发射减弱；对于在４００。

退火消除内应力的机理引言：退火是一种通过加热和冷却材料来减轻或消除内应力的热处理方法。

它被广泛应用于金属和玻璃等材料的生产中，以提高材料的机械性能和耐腐蚀能力。

本文将探讨退火消除内应力的机理，揭示背后的科学原理。

第一部分：内应力的形成内应力是由于材料在制造、加工和使用过程中受到外力影响而产生的一种力学现象。

例如，金属在冷加工过程中，由于塑性变形和晶界滑移等原因，会形成内应力。

这些内应力会导致材料的变形、开裂和失效。

第二部分：退火的基本原理退火是一种热处理方法，通过加热和冷却材料来改变其结构和性能。

退火的基本原理是通过加热材料使其达到高温状态，然后缓慢冷却，使材料中的晶体重新排列，从而减轻或消除内应力。

第三部分：退火的工艺过程退火的工艺过程通常包括加热、保温和冷却三个阶段。

首先将材料加热到退火温度，使其达到均匀的高温状态。

然后将材料保温一段时间，使晶体结构发生改变。

最后，缓慢冷却材料，使晶体结构稳定下来。

第四部分：退火对内应力的影响退火可以通过两种方式来减轻或消除内应力：晶界扩散和塑性变形。

在退火过程中，高温状态下的材料会发生晶界扩散，使晶体结构重新排列，从而减轻内应力。

同时，退火还可以促进材料的塑性变形，使内应力得到释放。

第五部分：退火的应用和效果退火广泛应用于金属和玻璃等材料的生产中。

通过退火，可以改善材料的机械性能和耐腐蚀能力，提高材料的可加工性和使用寿命。

退火还可以减少材料的变形和开裂，提高制品的成形性能和质量。

结论：退火是一种通过加热和冷却材料来减轻或消除内应力的热处理方法。

它通过改变材料的结构和性能来提高材料的机械性能和耐腐蚀能力。

退火的机理包括晶界扩散和塑性变形，通过这些机制可以减轻或消除内应力。

退火在金属和玻璃等材料的生产中得到广泛应用，对提高制品的成形性能和质量具有重要作用。

参考文献：[1] Zhang S, Zhang L, Yang Z, et al. Effects of annealing temperature on mechanical properties and microstructure of a cold-rolled Al-Mg-Si alloy[J]. Journal of Materials Processing Technology, 2018, 257: 52-59.[2] Wang H, Chen H, Cui Y, et al. Influence of annealing temperature on microstructure and mechanical properties of cold-rolled Al-Mg-Si alloy[J]. Materials Science and Engineering: A, 2017, 708: 290-298. [3] Zhou X, Cui Y, Gao S. Effects of annealing temperature on microstructure and mechanical properties of cold-rolled Al-Mg-Si alloy[J]. Materials Science and Engineering: A, 2018, 721: 1-7.。

ZnO基材料的压电、铁电、介电与多铁性质研究进展作者：门保全， 郑海务， 张大蔚， 马兴平， 顾玉宗， MEN Bao-quan， ZHENG Hai-wu，ZHANG Da-wei， MA Xing-ping， GU Yu-zong作者单位：门保全,MEN Bao-quan(河南大学物理系,微系统物理研究所,光伏材料省重点实验室,开封,475004;河南农业职业学院,郑州,451450)， 郑海务,张大蔚,马兴平,顾玉宗,ZHENG Hai-wu,ZHANG Da-wei,MA Xing-ping,GU Yu-zong(河南大学物理系,微系统物理研究所,光伏材料省重点实验室,开封,475004)刊名：硅酸盐通报英文刊名：BULLETIN OF THE CHINESE CERAMIC SOCIETY年，卷(期)：2009，28(4)被引用次数：0次1.郑海务.孙利杰.张杨退火温度对6H-SiC衬底上ZnO薄膜发光性质的影响[期刊论文]-人工晶体学报 20072.Lin Y C.Chen M Z.Kuo C C Electrical and optical properties of ZnO:Al film prepared on polyethersulfone substrate by RF magnetron sputtering 20093.Kim S.Seo J.Jang H W Effects of ambient annealing in fully 002-textured ZnO:Ga thin films grown on glass substrates using RF magnetron co-sputter deposition 20094.Sanchez N.Gallego S.Muňoz Magnetic States at the Oxygen Surfaces of ZnO and Co-Doped ZnO 20085.Xin M J.Chen Y Q.Jia C Electro-codeposition synthesis and room temperature ferromagnetic anisotropy of high concentration Fe-doped ZnO nanowire arrays 20086.Kumar R.Singh A P.Thakur P Ferromagnetism and metal-semiconducting transition in Fe-doped ZnO thin films 20087.Onodera A.Tamaki N.Jin K Ferroelectric properties in piezoelectric semiconductor Zn1-xMxO(M=Li,Mg) 19978.zgürü.Alivov Y I.Liu C A comprehensive review of ZnO materials and devices 20059.Pearton S J.Norton D P.Ip K Recent progress in processing and properties of ZnO 200510.Onodera A.Tamaki N.Kawamura Y Dielectric activity and ferroelectricity in piezoelectric semiconductor Li-doped ZnO 199611.Onodera A.Yoshio K.Satoh H Li-substitution effect and ferroelectric properties in piezoelectric semiconductor ZnO 199812.Joseph M.Tabata H.Kawai T Ferroelectric behavior of Li-doped ZnO thin films on Si(100) by pulsed laser deposition 199913.Dhananjay.Nagaraju J.Krupanidhi S B Effect of Li substitution on dielectric and ferroelectric properties of ZnO thin films grown by pulsed-laser ablation 200614.Ni H Q.Lu Y F.Liu Z Y Investigation of Li-doped ferroelectric and piezoelectric ZnO films by electric force microscopy 200115.Yang Y C.Song C.Wang X H Giant piezoelectric d33 coefficient in ferroelectric vanadium doped ZnO films 200816.Zhang Y J.Wang J B.Zhong X L Influence of Li-dopants on the luminescent and ferroelectric properties of ZnO films 200817.Lin Y H.Ying M H.Li M Room-temperature ferromagnetic and ferroelectric behavior inpolycrystalline ZnO-based thin films 200718.Zhang K M.Zhao Y P.He F Q Piezoelectricity of ZnO films prepared by sol-gel method[期刊论文]-Chinese Journal of Chemical Physics 200719.Yu L G.Zhang G M.Zhao X Y Fabrication of lithium-doped zinc oxide film by anodic oxidation andits ferroelectric behavior 200920.Zou C W.Li M.Wang H J Ferroelectricity in Li-implanted ZnO thin films21.Dhananjay Singh S.Nagaraju J Dielectric anomaly in Li-doped zinc oxide thin films grown by sol-gel route 200722.Dhananjay.Nagaraju J.Choudhury P R Growth of ferroelectric Li-doped ZnO thin films for metal-ferroelectric-semiconductor FET 200623.Dhananjay.Nagaraju J.Krupanidhi S B Off-centered polarization and ferroelectric phase transition in Li-doped thin films grown by pulsed-laser ablation 200724.Yang Y C.Song C.Wang X H V5+ ionic displacement induced ferroelectric behavior in V-doped ZnO films 200725.Yang Y C.Song C.Wang X H Cr-substitution-induced ferroelectric and improved piezoelectric properties of Zn1-xCrxO 200826.Schuler L P.Valanoor ler P The effect of substrate materials and postannealing on the photoluminescence and piezo properties of DC-sputtered ZnO 200727.Wang X B.Song C.Li D M The influence of different doping elements on microstructure,piezoelectric and resistivity of sputtered ZnO film 200628.Wang X B.Li D M.Zeng F Microstructure and properties of Cu-doped ZnO films prepared by dcreactive magnetron sputtering 200529.Juang Y.Chu S Y.Weng H C Phase transition of Co-doped ZnO 200730.Ghosh C K.Malkhandi S.Mitra M K Effect of Ni doping on the dielectric constant of ZnO and its frequency dependent exchange interaction 200831.Spaldin N A Search for ferromagnetism in transition-metal-doped piezoelectric ZnO 200432.Yang Y C.Zhong C F.Wang X H Room temperature multiferroic behavior of Cr-doped ZnO films 20081.学位论文杜朝玲Sr<,m-3>Bi<,4>Ti<,m>O<,3m+3>铁电氧化物和ZnO基稀磁半导体的拉曼光谱研究2007拉曼光谱学是研究物质元激发、结构和其它等物理性质的重要手段之一。

Material PropertiesEffects of processing parameters on the mechanical propertiesof polypropylene random copolymerSenol Sahin,Pasa Yayla*Mechanical Engineering Department,Engineering Faculty,Kocaeli University,41040Kocaeli,TurkeyReceived9June2005;accepted19July2005AbstractThe mechanical properties of polypropylene random copolymer(PP-R)with different processing parameters were studied. Special attention is devoted to the investigation of the influence of masterbatch addition on the variation in the mechanical properties of injection moulded PP-R.Tensile,instrumented Charpy impact,Shore D hardness,differential scanning calorimeter(DSC)and Vicat softening temperature(VST)tests were conducted on the test samples containing different colour masterbatches varying from0.5to10wt%.The observed changes in the mechanical behaviour are explained by the type and level of masterbatch content.The natural UV weathering performance of the PP-R material was studied from the masterbatch type point of view.The effect of processing parameters on material performance was studied on samples which were directly obtained from extruded pipes and on injection moulded samples.Finally,the effects of storage time on the polymer properties were investigated.q2005Elsevier Ltd.All rights reserved.Keywords:Polypropylene random copolymer;Processing parameters;Masterbatch types;Masterbatch contents;Ultraviolet degradation; Storage time1.IntroductionThe use of plastic materials in pipe applications is well established because of the lightweight,high performance, and excellent corrosion performance they can offer compared with metallic materials such as iron and copper. Having achieved high level of penetration in different applications,varying from water supply to gas distribution networks,from sanitary and heating systems to waste water collection and discharge systems,the use of plastics is expected to continue growing steadily at a rate of about5–10%per year[1].In many applications it is necessary to pigment the resin to specific colour for modifying the optical appearance for design,styling and functional purposes.Because of the importance of pigment addition to the base polymer,many studies[2,3]have been undertaken to understand their effects on the performance of plastics,especially since it has been shown that incorporating additives into polymeric materials during fabrication often affects rheological[4], mechanical[5,6]and optical properties[7]in an unpredict-able,and sometimes detrimental manner.PP can be coloured by two different methods.In thefirst method, certain types of colouring pigments are added to the natural base polymer at a certain pre-defined percent by the converters during either injection moulding or extrusion of thefinal product.In the second method,the base polymer is coloured by the raw material producer by compounding during the production process.Since the second method gives more uniform colouring,both the raw material producers and,generally,the end-users prefer compounded polymer.The polymer producers claim thatthe *Corresponding author.Tel.:C902623351148;fax:C902623352812.E-mail address:pyayla@.tr(P.Yayla).development of coloured PP-R raw material has to be done by the polymer producers due to the fact that specifically designed compounding equipment is used to obtain proper pigment and additive distribution without damaging the molecular structure[8].Furthermore,it is generally claimed that coloration with masterbatches done by the converters leads in general to poor pigment dispersion,resulting in pigment agglomerates acting as defects in the polymer matrix and impairing the mechanical properties of the finished product.In fact,when a colour masterbatch is used during the converting,the distribution level achieved in the extruder is usually not acceptable,resulting in uneven distribution[8].Despite these disadvantages,colouring by the polymer converter is cheaper and gives them some additional logistic advantages.Moreover,it is fairly seldom that a specific colorant is added to the reactants during a polymerisation process,unless such polymer resin is required in a large volume.Polypropylene random copolymer(PP-R)is one of the fastest growing of plastics being used in sanitary and heating applications.The overall mechanical properties of PP are strongly influenced by testing and processing parameters of the polymer.Hence,knowledge of relations between structure and mechanical properties of polymers enables the manufacturers to produce materials with certain morphologies by altering the processing conditions[9]. The influence of testing parameters on the overall performance was discussed in Part1of this work[10]. Determining the effects of processing parameters on the mechanical properties of PP-R is the objective of this study. Special attention is devoted to the investigation of influence of masterbatch addition on the variation in the mechanical properties.2.Experiments2.1.MaterialsThe base polymer used in this study is a natural colour polypropylene random copolymer(PP-R),produced by Borealis S.A.,trade name Borealis RA130 E.The properties of this natural PP-R are given in the work by Sahin and Yayla[10].It is known that colouring in general and the method of colouring in particular,might have some influence on the mechanical properties of the polymer[11].In this study,the effect of colouring methods on the overall mechanical performance of PP-R is investigated.In addition to the method of colouring,the types of pigment used for colouring could influence the overall mechanical performance of the material.In general,inorganic pigments yield more stable and better mechanical properties than organic anic pigments, however,are known for their high colour strength, brightness and good transparency[12].In order to determine the effects of different colouring on material properties,four different colour masterbatches were used.The masterbatches used to make thefinal product in different colours were supplied by local masterbatch manufacturers.They were supplied in granule form and the carrier resin for pigments for the all types of masterbatches was PP-R.The typical properties of masterbatches used in this study are detailed in Table1.Four different coloured PP-R materials,compounded during polymerisation by the above named producer,were used to prepare test samples.From the producers’point of view,there is no significant difference in the typical properties of these different colour compounds.2.2.Specimen preparationFour different groups of samples were prepared in this investigation.For thefirst group,an un-compounded natural material was coloured during the injection process. For this group of samples,different types of masterbatches at pre-defined ratios,varied between0.5,1,2.5,5and 10%in weight,were added during the injection process of the test samples.Before introducing the resin and masterbatches into the injection moulding machine,the masterbatch and the resin were put in a mixer and mixed for about30min.For the second group,a compounded PP-R material in blue,white,grey and green was used to prepare test samples by injection moulding.For these two groups,the detailed information on the injection mould and injection parameters was given elsewhere[10].For the third and fourth groups,the test samples were directly cut and extracted from the extruded pipe manufactured from white colour compounded,and1%white colour masterbatch added materials,respectively.These third andTable1Typical properties of masterbatches used in this studyM1M2M3M4 Colour White Green Blue Black Pigment type Organic Organic/Inorganic Organic/Inorganic Inorganic Total pigment concentration(%)60503040Meltflow index(gr/10min) (2308C,21.2N)6.017.934.1!0.01S.Sahin,P.Yayla/Polymer Testing24(2005)1012–10211013fourth groups of samples are coded as W1and W2,respectively.3.Mechanical testsUnless otherwise mentioned,all tensile tests were carried out at a crosshead speed of 50mm/min and,before testing,all samples were conditioned at room temperature for a period of 30days.All the results are average of three tests.The effects of processing parameters,masterbatch types and content and natural UV weathering on the properties of material were monitored using tensile,Charpy impact,Shore D hardness,and DSC tests.The details on these three tests were outlined elsewhere [10].3.1.Microstructural analysisAs the addition of any types of additives may alter the crystallisation characteristics of PP [13],a differential scanning calorimeter (DSC)analysis was used to evaluate thermal and morphological characteristics and the degree of crystallinity in the moulded samples containing different types of masterbatches at different concentrations.Tests were carried out on a Rheometric Scientific Polymer Laboratories instrument.Samples,each having a weight of about 13mg,were extracted from the middle sections of injection moulded samples,shown in Fig.1(c)of Sahin and Yayla [10].In the DSC tests,each sample was heated from 30to 2008C at a rate of 108C/min under a nitrogen atmosphere.Both thermal and crystallisation parameters were obtained from the heating scans.The level of crystallinity was calculated with the Eq.(1).c ZD H scc100(1)where D H sc is the melting enthalpy of the semi-crystalline material to be studied,and D H c is the enthalpy of 100%crystalline material.Since,a specific value of D H c for PP-Rdoes not exist,and the enthalpy of fusion for 100%crystalline polypropylene is almost independent of the isotacticity and equals 207J/g [14],the value of D H c for PP-R is taken as 207J/g.The melting temperature was taken as the peak temperature in the curves.The peak area,calculated automatically by the DSC instrument,was taken as the melting enthalpy.Typical thermal histories of some PP-R samples are given in Fig.1.It is evident that there is an endothermic melting peak in the heating scans.From the thermograms,it could be deduced that the onset of melting temperature is around 1108C and the melting point,the maximum of endothermic of melting peak in the scan,is around 1458C.From the figure,the crystallinity could be calculated as 30%.It has been demonstrated that there is very little difference between samples containing the natural and 1%of different colour masterbatches.Another important feature to be considered in Fig.1is that the DSC scans of all samples,namely natural,1%white colour masterbatch containing PP-R,white compound,W1and W2samples differ remarkably from each other,not only in their overall history but also in their melting temperature peaks.These variations are mainly due to the method that the samples are coloured,pounding or 1%masterbatch addition,and the way the samples are prepared,i.e.injection or extrusion.The cooling history of the extruded pipes and injected samples is rather different,resulting in some morphological variations in the samples.The early peaks and other multiple peak phenomena in the DSC tests are attributed to the compositional hetero-geneity of the crystal morphology of the polymer [15,16],and to lower-molecular weight polymer,which melted very early (i.e.wax,processing aid,dispersion aid)[17].3.2.Colouring effectsThe colouring of plastic products can be achieved easily by adding a small percentage of colour masterbatches during processing.To determine the effects of masterbatch concentration on the tensile properties of PP-R,tensile tests on samples containing 0.5,1,2.5,5and 10%masterbatches,detailed in Table 1,were carried out.Fig.2depicts the variation of yield stress with different colour masterbatch content.The figure shows that the addition of masterbatch diminishes the yield stress for masterbatch contents of up to about 0.5%.Moreover,the yield stress recovers with increasing the masterbatch content,and increasing the content enhances the yield stress as well.It is worth pointing out that the yield stresses of white,green and blue colour compounds are very comparable and more or less similar to that of 0.5%masterbatch samples.The decrease in the yield stress with masterbatch content is due to the high MFI value of the masterbatch (see Table 1).On the other hand,the recovery and increase in the yield stress is attributed to the reinforcing effect of pigments in the masterbatch [11],as well as the nucleating effects of these pigments [18,19].203550658095110125140155170185–1.8–1.5–1.2–0.9–0.6–0.30.00.3 Natural 1% White White W2 W1H e a t F l o w [m c a l /s ]Temperature [°C]Fig.1.DSC thermal history of natural and coloured PP-R showing remarkable influence of the method of colouring and the way the sample is prepared.S.Sahin,P.Yayla /Polymer Testing 24(2005)1012–10211014The last point in Fig.2to be mentioned is that the yield stresses of samples extracted from the pipe,namely W1and W2,are significantly lower than that of the injection moulded test samples.The difference between W1and W2is thought to be mainly due to the method of colouring,and the differences between W1,W2and others are due to differences between the crystallisation histories of injection and extrusion of the samples.This is another illustrative example of the strong interrelationship between structure,processing,morphology and final polymer [20].Fig.3shows the variation of yield strain with masterbatch content of the samples.The figure clearly reveals that adding the masterbatch reduces the yield strain,and this decline is almost linear for all masterbatch types.The yield strain also depends on colour compound and is much lower for the samples extracted directly from the pipe.The decline in yield strain as a function of masterbatch content is attributed to the reinforcing effect of pigments in the masterbatches.Y i e l d S t r e s s [N /m m 2]Masterbatch Content [%]Fig.2.Variation of yield stress with masterbatch contents for PP-R material.(The bars W,white colour compound;GN,green colour compound;B,blue colour compound;GR,grey colour compound;W1,White colour compound pipe;W2,1%white colour masterbatch containing pipe).Y i e l d S t r a i n [%]Masterbatch Content [%]Fig.3.Variation of yield strain with masterbatch contents for PP-R material.S.Sahin,P.Yayla /Polymer Testing 24(2005)1012–10211015The variation of Charpy impact energy with masterbatch content is shown in Fig.4,showing that for all types of masterbatches the total Charpy impact energies increase with masterbatch content up to 1%,after which the impact energy remains constant.However,the impact energy increases almost linearly with the white masterbatch content.Regarding the compounded samples,the green and blue compounded samples showed the highest impact resistance.Fig.4indicates more clearly that the inorganic based pigments (namely green,blue,and black compounds and masterbatch added samples)give better impactproperties than organic compounds (namely white,PP-R polymer).As pointed out elsewhere [10],the Charpy impact crack initiation and propagation resistance of the material are rather sensitive to the test temperature.The lower transition temperature is around 08C,and above 858C the material becomes too ductile to break.The present investigation made it clear that neither the content nor the type of masterbatch had any effect on this brittle-ductile transition.The effect of masterbatch content on melting tempera-ture,extracted from the DSC scans,is shown in Fig.5indicating that,except for the white masterbatch,addingC V [k J /m 2]Masterbatch Content [%]Fig.4.Variationof Charpy impact energy (C v )with masterbatch contents for PP-R material.M e l t i n g P o i n t [°C ]Masterbatch Content [%]Fig.5.Variation of melting point [T m ]with masterbatch contents for PP-R material.S.Sahin,P.Yayla /Polymer Testing 24(2005)1012–10211016masterbatch does not have a significant influence on the melting point of the polymer.The melting points of coloured and compounded samples do not differ signifi-cantly from each other.Fig.6indicates the volume percentage of crystallinity of compounded and coloured PP-R samples.It could be concluded from the figure that the addition of masterbatch and compounding diminishes the crystallinity and the level of crystallinity differs from one masterbatch to another,which could be attributed to the different level of nucleating activity of each pigment [18].Therefore,different types of pigments incorporated in PP-R result in different degrees of crystallisation.These results agree with the findings of Kening et al.[21]and Krisher and Marshall [22],which showed that incorporating pigments into PP affected its mechanical properties mainly in a positive fashion.3.3.Time effectsIt is known that the properties of the polymeric material may significantly change just after conversion,i.e.by extrusion or injection processes.However,the history of this property change is not known for this material,so it is worth investigating what changes take place and how they depend on time.After injection moulding,the PP-R samples were conditioned at 228C and 50%relative humidity for a wide range of times—between 5min and 23months—in a box which excluded light.Fig.7shows the effect of storage time on yield strength and Charpy impact resistance of 1%white and 1%green masterbatch containing PP-R material.It is seen that the impact strength decreases considerably with storage time,and that it stabilizes about 30days after production.As for the yield stress,Fig.7shows that the yield stress increases gradually with storage time and that after around 30days itc [%]Masterbatch Content [ % ]Fig.6.Variation of crystallinity with masterbatch contents for PP-R material.05101520253035Conditioning Time [hour]σy [N /m m 2]C V [k J /m 2]Fig.7.Effects of storage time on yield stress (s y )and Charpy impact energy (C v )of 1%white and 1%green masterbatch containing PP-R material.Conditioning Time [hour]T m [°C ]c [%]Fig.8.Effect of storage time on melting temperature (T m )and crystallinity (c)of natural PP-R material.S.Sahin,P.Yayla /Polymer Testing 24(2005)1012–10211017remains fairly unchanged.The figure also reveals that time dependent tensile strength and impact resistance values are not sensitive to the type of masterbatch.The variation of melting temperature and crystallinity with storage time is shown given in Fig.8,indicating that both crystallinity and melting temperature increase with storage time and that both values stabilise after about 30days.Similarly,the variation of Shore D hardness with storage time is plotted in Fig.9,pointing out that the hardness increases gradually after the injection of the samples and remains relatively unchanged after more than 30days after production.Vicat softening temperature tests (VST)were utilised to determine the softening temperature of the material.A Zwick Vicat softening temperature tester at 50N force and 508C/h heating rate was utilised to determine temperature at which the indentor penetrates 1mm into the material.This value is particularly important for the conversion of the material into the product.The variation of VST with storage time is given in Fig.10,showing a minor increase with conditioning time for natural PP-R.3.4.Natural UV weatheringEvery polymeric material exposed to direct sunlight undergoes some damage.Plastic pipes made from PP-R could be used outdoors and thus may undergo ultraviolet (UV)degradation.This is especially the case for PP-R pipes used in solar heating systems.For this type of application,it is particularly important to decide what colour of PP-R pipe to use and how its mechanical properties deteriorate with the UV exposure time.Exposure of many plastics to ultraviolet radiation causes a loss in their mechanical properties.The mechanical property most severely affected is usually theV S T [°C ]Conditioning Time [hour]Fig.10.Effect of storage time on VST for natural PP-R material.H a r d n e s s (S h o r e D )Conditioning Time [hour]Fig.9.Effect of storage time on Shore D hardness for natural and 1%w white masterbatch containing PP-R material.50010001500200025003000350040004500500055000.00.51.01.52.02.53.03.54.0Januar 02December T m 5.3°C ϕm 85.9%October T m 16.7°C ϕm 70.4%September T m 21.7°C ϕm 66.8%November T m 11.2°C ϕm 74.4%August T m 24.6°C ϕm 73.3%July T m26.1°Cϕm 68.6%June T m 22.5°C ϕm 60.3%A b s o r b e d T o t a l S o l a r R a d i a t i o n E n e r g y [G J /m 2]Solar Exposure Time [hour]Fig.11.Absorbed total solar radiation energy variation as a function of UV exposure time considered in this study.T m and 4m represent the average temperature and humidity,respectively.S.Sahin,P.Yayla /Polymer Testing 24(2005)1012–10211018ductility of the material [23].Pigment systems,on the other hand,influence to what extent materials are affected by ultraviolet radiation exposure.Some colours,such as black,may make the material less susceptible to ultraviolet radiation degradation than others [24].Several series of tensile,Charpy impact,Shore D and DSC tests were carried out on test samples exposed to sunshine for different times,up to six months during the summer time.All samples were suspended so that they could freely rotate,enabling each surface to be UV affected uniformly.Fig.11shows the total absorbed energy as a function of exposure time.The variation of tensile properties with exposure time is shown in Fig.12.The figure reveals that both yield stress and yield strain decreases with UV exposure time for all samples.Most significantly,the greatest decrease is seen in the natural PP-R.The least UV affected sample is the PP-R containing 1%black masterbatch.The effect of natural UV weathering time on the Charpy impact resistance of PP-R material is shown in Fig.13,showing that for all the samples the Charpy impact energydiminishes with the UV exposure time.However,black samples are the least and natural samples the most affected.Fig.14shows the variation of the Shore hardness with UV exposure time.The figure shows that just after the beginning of natural UV exposure,the hardness of the samples increases with exposure time for all samples.It is also worth mentioning that the least affected are the natural colour samples,which have higher opacities.The hardness increase in blue and green samples is pronounced.Black samples are exceptions,mainly due to the carbon black content of the masterbatch.As stated by Turton and White [24],pigments that simply reflect or scatter UV,preventing it from penetrating far into the plastics,also limit degradation to a region close to the surface.Fundamental studies in the area of polymer morphology almost universally employ DSC.Thermal analysis has also played a significant role in the degradation studies of semicrystalline polymers [25].DSC tests on UV degraded samples could be a good indication of morphologicalWeathering Time [hour]εy [%]σy [N /m m 2]Fig.12.Natural UV weathering time effect on yield stress (s y )and yield strain (3y )of natural and 1%coloured PP-R material.Weathering Time [hour]C V [k J /m 2]Fig.13.Natural UV weathering time effect on Charpy impact strength (C v )of natural 1%coloured PP-R material.Weathering Time [Hour]H a r d n e s s [S h o r e D ]Fig.14.Natural UV weathering time effect on Shore D hardness of natural and 1%coloured PP-R material.7085100115130145160175190–1.6–1.1–0.6–0.10.40.90 hour 504 hour 2304 hour 4776 hourH e a t F l o w [m c a l /s ]Temperature [°C]Fig.15.Natural UV weathering time effect on the DSC thermal history of natural PP-R material.S.Sahin,P.Yayla /Polymer Testing 24(2005)1012–10211019changes in the material.The DSC scans obtained in this study show that the degree of UV degradation sensitivity differs from one colour to other.From the comparison of Figs.15–17,it could be deduced that the most affected sample is the natural (Fig.15)and the least effected ones are the black,blue (Fig.16),white (Fig.17)and green coloured samples,respectively,showing strong structural modifi-cations due to UV degradation.This dependence is attributed to the fact that the different pigments used in the masterbatches have different UV absorption/screening mechanisms,resulting in different protective mechanisms against UV exposure [26,27].The effect of adding different colour masterbatches on the UV performance of PP-R material was also evaluated using the melt flow index (MFI)study.MFI was measured at 2308C at 21.6N according to ISO 1133standard on UV exposed samples containing 1wt%blue,black,green and white colour masterbatch.For the sake of comparison,the variation of MFI with exposure time for natural PP-R was also measured.The results in Fig.18show that natural andgreen samples were the most affected and black,blue and white samples were the least UV affected.The results of these MFI measurements correlate with the DSC results in Figs.15–17.The increase in MFI is attributed to the chain scission and formation of new groups [25].4.ConclusionsThe results of the present investigation indicate that the way the test samples are produced,i.e.injection or extrusion,strongly influences the overall mechanical properties of polymers.Mechanical and DSC tests showed that both the type and content of masterbatches in PP-R influence not only the degree of crystallinity but also the structure and mechanical properties of the polymer.The addition of colour masterbatches increases the short term yield stress and impact strength but reduces the yield strain .The yield and Charpy impact resistance of both natural and 1%masterbatch containing PP-R material continuously vary during the room temperature conditioning period of about 30days,and then remain fairly constant.The property changes in mechanical and thermal properties,particularly in yield strength and Charpy impact resistance,after the injection moulding could not solely be explained through ‘post-crystallisation’.Apart from the slight variation in crystallinity,the influence of the polymer parameters fixed in the synthesis,namely average molar mass,molar mass distribution and chain regularity [28]as well as the reduction of free volume of the amorphous phase has to be taken into account [29].For all naturally aged PP-R samples coloured in different ways the yield strength,yield strain and Charpy impact energy decrease with increasing natural UV weathering time.Shore D hardness,on the other hand,increased with the degradation time.Tensile and impact properties were affected mostly by UV in samples having higher opacities.7085100115130145160175190–1.7–1.5–1.3–1.1–0.9–0.7–0.5–0.30 hour 504 hour 2304 hour 4776 hour H e a t F l o w [m c a l /s ]Temperature [°C]Fig.16.Natural UV weathering time effect on the DSC thermal history of 1%blue coloured PP-R material.Weathering Time (hour)M F I (g r /10 m i n )Fig.18.Effect of natural UV weathering time on the variation of MFI values for PP-R material containing 1%white,green,blue and black masterbatch and natural PP-R.7090110130150170190–2.0–1.5–1.0–0.50.0504 hour 2880 hour 4320 hour H e a t F l o w [m c a l /s ]Temperature [°C]Fig.17.Effect of Natural UV weathering time on the variation of DSC melting curves for PP-R material containing 1%white masterbatch.S.Sahin,P.Yayla /Polymer Testing 24(2005)1012–10211020Morphology,crystallization and melting behaviour is also affected by the addition of masterbatch to a degree which depends on masterbatch type.This study provides a clear view of the influence of masterbatch type and content on both mechanical proper-ties.Similarly,the UV degradation of the PP-R is very much masterbatch,thus pigment,dependent,making it clear that pigments vary in their UV stability.With this work it is evidenced that the MFI tests on UV degraded samples are quite useful to assess the extent of degradation.This has practical implications for durability assessment and stabil-ization strategies.AcknowledgementsWe are greatly indebted to colleagues at both industrial as well as university laboratories for their positive interest, and especially to Turkkablo A.O.,Emas A.S.,and Pipelife A.S.The raw materials supplied by Borealis S.A.are gratefully acknowledged.Dr P.S.Leevers of Imperial College UK,and Prof.Dr E.Cavusoglu of Arili Plastik of Turkey are also appreciated for their valuable discussion and comments on the manuscript.References[1]G.Beer,Polypropylen(PP),Kunststoffe86(10)(1996)1460–1463.[2]R.L.Gray,R.E.Lee,Influence of co-additive interactions onstabilizer performance,ANTEC’96(1996)2683–2687. 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压缩条件下AZ31镁合金第Ⅱ阶段加工硬化邹罡;张康;张诗昌【摘要】在温度为25℃～300℃、应变速率为3×10-2 s-1～ 10-4 s-1的条件下,对挤压态的AZ31镁合金沿挤压方向进行了压缩试验,试验研究了加工硬化率随应力的变化关系,以及产生第Ⅱ阶段加工硬化的原因.结果表明,在一定温度及应变速率范围内,加工硬化率随应力增加而增加,当加工硬化率达到峰值时即出现第Ⅱ阶段加工硬化.产生第Ⅱ阶段加工硬化的主要原因是压缩过程中,镁合金组织产生了{ 1012}拉伸孪晶,随着孪晶数量增加,大量孪晶界会阻碍位错运动,造成加工硬化率升高.【期刊名称】《轻合金加工技术》【年(卷),期】2017(045)005【总页数】5页(P52-56)【关键词】镁合金;第Ⅱ阶段加工硬化;孪晶;加工硬化率【作者】邹罡;张康;张诗昌【作者单位】广东省建筑科学研究院集团股份有限公司,广东广州 510500;武汉科技大学材料与冶金学院,湖北武汉 430081;武汉科技大学材料与冶金学院,湖北武汉 430081【正文语种】中文【中图分类】TG146.22根据加工硬化率与应力之间的关系，金属材料的加工硬化通常具有三个阶段特征[1]：第Ⅰ阶段的特点是加工硬化率很低，且与应力无关，一般出现在单晶材料中；第Ⅱ阶段则对应着较高的加工硬化率，且在某一给定的变形条件下，加工硬化率为一确定常数；第Ⅲ阶段加工硬化率与应力呈线性关系，即随着应力增加，加工硬化率直线降低。

材料在变形过程中，根据变形条件的不同往往表现出一个或几个加工硬化特征[1-2]。

镁合金的加工硬化随变形温度和应变速率的变化可以出现第Ⅱ或第Ⅲ阶段加工硬化[3-4]。

由于第Ⅱ阶段加工硬化与位错交互作用有关，因此产生第Ⅱ阶段加工硬化需要具备特定条件。

Hnorng-Yu Wu[3]等人研究了AZ31B-H24镁合金热轧薄板的加工硬化行为表明：只有当应变速率高于4×10-3 s-1且在室温下变形时，AZ31B-H24镁合金才产生第Ⅱ阶段加工硬化，而当温度升高到250℃时，第Ⅱ阶段加工硬化消失。

Trans. Nonferrous Met. Soc. China 28(2018) 251−258Effect of annealing temperature on joints of diffusion bonded Mg/Al alloysYun-long DING1, Jian-gang WANG2, Ming ZHAO3, Dong-ying JU1,31. Department of Material Science and Engineering, Saitama Institute of Technology,Fusaiji 1690, Fukaya, Saitama 369-0293, Japan;2. Hebei Key Laboratory of Material Near-Net Forming Technology,Hebei University of Science and Technology, Shijiazhuang 050018, China;3. Department of Materials Science and Engineering,University of Science and Technology Liaoning, Anshan 114051, ChinaReceived 22 December 2016; accepted 21 June 2017Abstract: To study the effect of annealing temperature on the joints between magnesium and aluminum alloys, and improve the properties of bonding layers, composite plates of magnesium alloy (AZ31B) and aluminum alloy (6061) were welded using the vacuum diffusion bonding method. The composite specimens were continuously annealed in an electrical furnace under the protection of argon gas. The microstructures were then observed using scanning electron microscopy. X-ray diffractometry was used to investigate the residual stresses in the specimens. The elemental distribution was analyzed with an electron probe micro analyzer. The tensile strength and hardness were also measured. Results show that the diffusion layers become wide as the heat treatment temperature increases, and the residual stress of the specimen is at a minimum and tensile strength is the largest when being annealed at 250 °C. Therefore, 250 °C is the most appropriate annealing temperature.Key words: annealing temperature; diffusion bonding; diffusion layer; residual stress; tensile strength1 IntroductionWith the rapid development of the transportation,aerospace, and defense industries, magnesium alloyshave received growing attention due to their lowdensities, high specific strengths, excellent casting ability,and outstanding vibrational energy absorption [1]. Whilethe use of magnesium alloys allows the weight ofcomponents to be reduced, this material is easilycorroded. Another group of materials used in a similarrole are aluminum alloys, which have attractivemechanical and metallurgical properties including highstrength and excellent corrosion resistance [2−5]. It iswell-known that both magnesium and aluminum alloysare both widely used in the aerospace, mechanical,electrical, and chemical industries [6−9]. Furthermore,with the growing emphasis on energy economy andenvironmental concerns, Mg alloys have become afavored choice in the automobile field. If aluminumalloys can be bonded to magnesium alloy and form somekind of composite material, not only would the flexibilityand availability of the material be substantially improved,but also the weight and cost would be reduced.At present, there is much research into this topic.Many welding methods have been used to join Mg alloysand Al alloys. These methods include soldering, electronbeam welding, resistance spot welding, explosivewelding [10], laser welding, and the vacuum diffusionbonding. However, no matter which technique is used,brittle and hard intermetallic compounds, such as Al3Mg2and Mg17Al12, form in the joints. This weakens thetensile strength of the joints. As was previously reportedelsewhere, the tensile strength of bonded specimens wasonly 37 MPa [11], despite the Mg/Al alloy being weldedwith a Zn interlayer [12].The process of annealing is always applied toplastically deformed metals and alloys in order to softenand restore the ductility and formability of materials [13].Annealing can transform the structure of crystals andeliminate defects in microstructures, thereby reducing thebrittleness and improving the mechanical characteristicsof a material. In this work, in order to investigate theseeffects, annealing treatments were applied at differentCorresponding author: Dong-ying JU; Tel: +81-48-5856826; Fax: +81-48-5855928; E-mail: dyju@sit.ac.jpDOI:10.1016/S1003-6326(18)64658-8Yun-long DING , et al/Trans. Nonferrous Met. Soc. China 28(2018) 251−258252 temperatures. In addition, the microstructures and their properties were investigated. This work represents one of the first attempts to study the effect of annealing temperature on the joints between magnesium and aluminum alloys. Based on the results of this work, the applications of the composite material formed by the diffusion bonding of aluminum alloy and magnesium alloy could be extensive. The composite materials formed by this process will lead to light-weight components, which will in turn lead to decreased depletion of resources and reduced energy usage, which can help mitigate environmental pollution.2 ExperimentalThe chemical compositions of AZ31B magnesium alloy and 6061 aluminum alloy are given in Table 1 and Table 2.Table 1 Composition of AZ31 Mg alloy (mass fraction, %)Al Zn Mn Ca Si Cu Ni Fe Mg 0.10.050.0050.005Bal.Table 2 Composition of Al6061 alloy (mass fraction, %)Mg Zn Mn Si Cu Fe Ti Cr Al 0.21750.14010.01270.0718Bal.The principle impurities present in the alloys are Fe and Ni [14]. These impurities have an adverse effect onthe uniformity of microstructures and the distribution of elements. Conversely, the presence of Mn elements can reduce the effect of impurities and refine the grain as well as improve the tensile strength of the diffusion layers. Si can improve the mechanical properties of the alloy at room temperature, and Cu can improve the strength at high temperature. Ca can refine the size of dendrites in Mg 17Al 12. Additionally, it will form the phase of Al 2Ca, which has a very high melting point [15]. Consequently, it can reduce the micro-hardness of the diffusion zone. The element Zn is present firstly for its own solid-solution strengthening and secondly because a small amount of Zn can increase the solubility of Al in Mg, thereby improving the solid solution strengthening effect of Al. Besides, Zn can also reduce the formation of intermetallic compounds [16]. Overall, Zn can reduce micro-hardness and increase the tensile strength of the diffusion layers.The first step in the experimental procedure was tocut AZ31B magnesium alloy sheets and 6061 aluminum alloy sheets to the dimensions shown in Fig. 1.Fig. 1 Dimensions of specimen (unit: mm)The oxide layers on the surface of the substrate were then polished with abrasive papers and the ground samples were wiped with acetone before joining. According to the Mg −Al phase diagram, the joining temperature was chosen as 440 °C. Specimens were successfully joined with a method called vacuum diffusion bonding under the protection of argon gas. After that, in order to refine microstructures and improve the properties of the bonding layers, annealing treatment was carried out. According to the Mg −Al phase graph and previous annealing experience, the samples were annealed using heat treatment temperatures of 200, 250 and 300 °C, and the holding time was 1 h. After heat treatment, samples were cooled to room temperature in an electric furnace.For the purpose of studying the effect of annealing temperatures on microstructures and the properties of the interfaces, a series of specimens annealed at different conditions were cut across the diffusion zone. The cut sections were then inlaid into resin to facilitate the investigation of microstructures. Using a grinder and abrasive papers (Grit 240, 600, 800, 1200), the samples were ground and polished with a polishing compound.The microstructures and elemental distribution of the joints were then studied using scanning electron microscopy (SEM) and an electron probe micro analyzer (EPMA). Using a tensile machine, the tensile strength was also investigated. The tensile speed used was 0.008 mm/s. After the measurement, the stress −strain graphs could be obtained. In order to investigate the distribution of residual stress of the specimens annealed at different temperatures, residual stress was measured by X-ray diffraction (XRD), based on the testing principle of residual stress and using X-ray of wavelength λ. Initially, the specimen was irradiated, and a diffraction angle 2θ was obtained which was later used to calculate the slope, M , of 2θ−sin 2ψ. Generally, ψ was set to be 0°, 15°, 30° and 45°. This allowed us to obtain the relationship between 2θ and sin 2ψ, and thus calculate the residual stress σ using the following equation:σ=K ·Δ2θ/Δsin2ψ=K ·M (1)K is the stress constant of XRD analysis, which can be expressed asYun-long DING , et al/Trans. Nonferrous Met. Soc. China 28(2018) 251−258 25301πcot 21180E K =-⋅⋅⋅+θυ (2)where E is the elastic modulus of the material, θ0 is the diffraction angle without stress, and υ is Poisson ratio [17].For the 6061 Al alloy, 2θ was set to be 140° and the stress constant was −163.32 MPa/(°). When 2θ was set as 155°, the stress constant was −79.14 MPa/(°) at the side of the AZ31 Mg alloy. The tube type was Cr, wavelength was K α, and the size of collimator was φ0.5 mm. In addition, 2θ values for Mg 17Al 12 and Al 3Mg 2 were set to be 150° and 145°, respectively, and the stress constants were −98.97 and −126.22 MPa/(°), respectively. This allowed the results of residual stress to be calculated. Finally, the micro-hardness was measured with a Vickers hardness tester with the load of 10 N.3 Results and discussion3.1 Microstructures of jointsTemperature is the most important parameter in the heat treatment process. The microstructures of the bonding interfaces annealed at 200, 250 and 300 °C were therefore observed. The microstructures of the joints are shown in Figs. 2(b −d). Figure 2(a) shows the micro- structures of joints that did not undergo annealing. The diffusion layers, including layer A, layer B, and layer C, are clearly visible in Fig. 2. Furthermore, as the annealing temperature rises, the width of the diffusion layers increases. This is because the diffusion rate increases with increasing temperatures.The elemental analysis results are shown in Fig. 3, and the width profiles of the diffusion layers are shown in Fig. 4, which also display the element analysis results from line scanning.Clearly, the diffusion layers become wide when the heat treatment temperature is increased. The magnesium and aluminum contents vary in layers A, B and C. Specifically, the untreated specimens exhibit a magnesium content of 800−300 counts, and the aluminum content ranges from 50 to 200 counts. The widths of layers A, B and C are 0.017, 0.018 and 0.11 mm, respectively. The widths of layers A, B and C of the specimens heat-treated at 200 °C are approximately 0.022, 0.025 and 0.13 mm, respectively. The magnesium content ranges from 600 to 300 counts across the diffusion layer from the Mg side to the Al side, while the amount of aluminum ranges from 50 counts to 100 counts in the same direction. However, when the sample is treated at 300 °C, the amount of magnesium ranges from approximately 650 to 800 counts, and the aluminum content varies in the range of 50−70 counts .Fig. 2 SEM micrographs of joints without annealing (a) and with annealing at 200 (b), 250 (c) and 300 °C (d) (Layer A: Al 3Mg 2; layer B: Mg 17Al 12; layer C: Mg-based solid solution)Yun-long DING , et al/Trans. Nonferrous Met. Soc. China 28(2018) 251−258254Fig. 3 Elemental distribution in joints from surface scanning: (a) Without annealing; (b) Annealed at 200 °C; (c) Annealed at 250 °C; (d) Annealed at 300 °CFig. 4 Elemental distribution in joints from line scanning: (a) Without annealing; (b) Annealed at 200 °C; (c) Annealed at 250 °C; (d) Annealed at 300 °CThe thicknesses of layers A, B and C are about 0.06, 0.07 and 0.36 mm, respectively. The results after treating the sample at 250 °C are shown in Fig. 4(c). The width of layer A is about 0.03 mm, and the widths of layers B and C are approximately 0.05 and 0.35 mm, respectively. The magnesium content ranges from 300 to 500 counts, while the aluminum content ranges from 50 to 70 counts. The reason for this variation in the width of the diffusion layers and element abundances is the diversity of annealing temperature. Elemental diffusion rates increase with increasing temperature, and diffusion layer is therefore the widest when the samples are annealed at 300 °C.3.2 Mechanical properties of jointsThe tensile strength after annealing at different temperatures is shown in the stress −strain graphs in Fig. 5.Yun-long DING , et al/Trans. Nonferrous Met. Soc. China 28(2018) 251−258 255Fig. 5 Stress −stain graphs: (a) Without annealing; (b) Annealed at 200 °C; (c) Annealed at 250 °C; (d) Annealed at 300 °CThe tensile strength is the highest when the annealing temperature is 250 °C, of about 56 MPa. When the samples are annealed at 200 and 300 °C, the values of tensile strength are approximately 41 and 49 MPa, respectively. Untreated specimens exhibit a tensile strength of 35 MPa. In this work, the application of annealing treatment results in a higher tensile strength. In addition to the stress −strain graphs shown above, the diffraction peak results are shown in Fig. 6 and the residual stresses are shown in Fig. 7.The diffraction peak of Al emerges when 2θ is near 139°, and the indices of the crystal face are (311). The indices of the crystal face at the diffraction peak of Mg near a 2θ value of 152° [18] are (104). The diffraction peaks of the intermetallic compounds Al 3Mg 2 and Mg 17Al 12 occur at 2θ=139.4° and 143°, respectively. The residual stress is a vector, and in this work, it is axially aligned with the specimens. Under a tensile stress, this value will be positive. In contrast, when the stress is compressive, the residual stress will be negative. It can be inferred from Fig. 7 that the residual stress at the interface (marked as 0 in Fig. 7) is tensile, and that of the untreated specimen exhibits a residual stress value of 65 MPa. When the samples undergo heat treatment at 300 and 200 °C, the residual stresses are respectively 51and 59 MPa. When being treated at 250 °C, the sampleexhibits a residual stress of approximately 44 MPa. Thus, the tensile strength of the specimen annealed at 250 °C is the highest, while that of the sample without annealing is the lowest. As the residual stresses are axial tensile stresses, and tensile stress will decrease the tensile strength, higher residual stresses will result in a lower tensile strength. On the contrary, if the residual stresses are compressive, then the higher the stress is, the higher the tensile strength will be. It is apparent that the results of the experiments for tensile strength are in good agreement with the predictions of the residual stress theory.The micro-hardness of the diffusion layers annealed at different temperatures were also investigated in this study, with the load set to be 10 N. The distributions of hardness in the annealed samples are shown in Fig. 8.The trends in the hardness distributions are broadly similar in that the hardness of the Mg side is higher than that of the Al side, and significantly increases in the diffusion zone. After being annealed at 200 °C, the hardness of the Mg side is HV 69, and HV 51 on the Al side, but is HV 201 in the diffusion bonding region. When the specimen is treated at 250 °C, the hardnesses of the Mg and Al sides are respectively HV 57 andYun-long DING , et al/Trans. Nonferrous Met. Soc. China 28(2018) 251−258256Fig. 6 Diffraction peaks from different parts of a bonded specimenFig. 7 Residual stress of specimens treated at different temperaturesFig. 8 Vickers hardness at interfaces annealed at different temperaturesHV 40. The interface has a hardness of HV 187. When the sample undergoes heat treatment at 300 °C, the hardness of the Mg side is HV 58 and is HV 42 in the Al side. The hardness is HV 212 in the bonded zone. When the sample does not undergo annealing, the hardnesses of the Mg and Al sides are respectively HV 55 and HV 39, and the hardness of interface is about HV 214. The hardness in the diffusion layer near the Al substrate is higher than that of the region near the Mg substrate. Thus, the variation in hardness is a result of the intermetallic compounds formed in varying locations in the diffusion zone, and the fundamental factor controlling hardness is the annealing temperature.To summarize the results above, it is apparent that 250 °C is the most suitable annealing temperature. The results supporting this conclusion can be explained as follows. From the elemental analysis results presented in Figs. 3 and 4, the content of elements after annealing at 250 °C remains steady. This indicates that the distribution of elements is relatively uniform, and that the microstructure is more uniform. Furthermore, the amount of Zn is relatively steady and more than that under other annealing conditions, as shown in Fig. 9.Zn has a significant effect on the diffusion zone of Mg alloys and Al alloys. It can reduce the formation of intermetallic compounds of Mg and Al. Even though intermetallic compounds can form, they will precipitate dispersedly, which can produce precipitation strengthening [19]. In other words, the amount of hardYun-long DING , et al/Trans. Nonferrous Met. Soc. China 28(2018) 251−258 257and brittle phases in the interface will decrease. Thisresults in more desirable properties than from other annealing conditions. Residual stress is at a minimum, and the tensile strength is the highest.Fig. 9 Elemental distributions of Zn from line scanning: (a) Without annealing; (b) Annealed at 200 °C; (c) Annealed at 250 °C; (d) Annealed at 300 °C4 Conclusions1) It is difficult to obtain high quality bonding strengths between magnesium and aluminum alloy sheets by diffusion bonding alone due to the formation of intermetallic compounds. The application of annealing improves the bonding strength.2) The width of the diffusion layers increases as the annealing temperature rises, because the diffusion rate becomes faster with increasing temperature.3) The annealing temperature has a great effect on microstructures and the mechanical behavior of samples. Various results suggest that 250 °C is the most suitable temperature for annealing of the diffusion zone of AZ31B magnesium alloy and 6061 aluminum alloy.AcknowledgmentsThis research work has been partially supported by the grant subsidy of the “Nano Project” for Private Universities: 2011−2014 from MEXT, Japan. This study was also supported by the “Advanced Science Research Laboratory” in Saitama Institute o f Technology, Japan.References[1]LAKSHMINARAYANAN A K, ANNAMALAI V E. Fabrication and performance evaluation of dissimilar magnesium−aluminium alloy multi-seam friction stir clad joints [J]. Transactions of Nonferrous Metals Society of China, 2017, 27: 25−35.[2]SHE Qing-yuan, YAN Hong-ge, CHEN Ji-hua, SU Bin, YU Zhao-hui, CHEN Chao, XIA Wei-jun. Microstructure characteristics and liquation behavior of fiber laser welded joints of Mg−5Zn−1Mn− 0.6Sn alloy sheets [J]. Transactions of Nonferrous Metals Society of China, 2017, 27: 812−819.[3]LIU L M, LIU F. Effect of Ce on micro-structures and properties of Mg/Al butt joint welded by gas tungsten arc with Zn –30Al –x Ce filler metal [J]. Science and Technology of Welding and Joining, 2013, 18: 414−420.[4]LIU Li-ming, LIU Fei, ZHU Mei-li. Study on Mg/Al weld seam based on Zn −Mg −Al ternary alloy [J]. Materials, 2014, 7: 1173−1187.[5]YAN Y B, ZHANG Z W, SHEN W, WANG J H, ZHANG L K, CHIN B A. Micro-structures and properties of magnesium AZ31B − aluminum 7075 explosively welded composite plate [J]. Materials Science and Engineering A, 2010, 527: 2241−2245.[6]LI Xiang-rong, LIANG Wei, ZHAO Xing-guo, ZHANG Yan, FU Xiao-peng, LIU Fen-cheng. Bonding of Mg and Al with Mg −Al eutectic alloy and its application in aluminum coating on magnesium [J]. Journal of Alloys and Compounds, 2009, 471: 408−411.[7]LIU Fei, REN Da-xin, LIU Li-ming. Effect of Al foils interlayer on micro-structures and mechanical properties of Mg −Al butt joints welded by gas tungsten arc welding filling with Zn filler metal [J]. Materials and Design, 2013, 46: 419−425.[8]ZHU Bo, LIANG Wei, LI Xian-rong. Interfacial micro-structures, bonding strength and fracture of magnesium-aluminum laminated composite plates fabricated by direct hot pressing [J]. Materials Science and Engineering A, 2011, 528: 6584−6588.[9]SHARMA C, UPADHYAY V , DWIVEDI D K, KUMAR P. Mechanical properties of friction stir welded armor grade Al−Zn−Mg alloy joints [J]. Transactions of Nonferrous Metals Society of China, 2017, 27: 493−506.[10]MA Li, HE Ding-yong, LI Xiao-yan, JIANG Jian-min. Micro- structures and mechanical properties of magnesium alloy AZ31B brazed joint using a Zn −Mg −Al filler metal [J]. J Mater Sci Technol, 2010, 26: 743−746.[11]SHANG Jing, WANG Ke-hong, ZHOU Qi, ZHANG De-ku, HUANG Jun, GE Jia-qi. Effect of joining temperature on micro-structures and properties of diffusion bonded Mg/Al joints [J]. Transactions of Nonferrous Metals Society of China, 2012, 22: 1961−1966.[12]ZHANG Yu, LUO Zhen, LI Yang, LIU Zu-ming, HUANG Zun-yue. Micro-structures characterization and tensile properties of Mg/Al dissimilar joints manufactured by thermo-compensated resistance spot welding with Zn interlayer [J]. Materials and Design, 2015, 75: 166−173.[13]SUN Hong-fei, CHAO Hong-ying, WANG Er-de. Microstructure stability of cold drawn AZ31 magnesium alloy during annealing process [J]. Transactions of Nonferrous Metals Society of China, 2011, 21(S): s215−s221.[14]YANG Ming-bo, PAN Fu-sheng, LI Zhong-sheng, ZHANG Jing. Alloying elements and their effects in Mg −Al based elevated temperature magnesium alloys [J]. Materials Review, 2005, 19: 46−49. (in Chinese)[15]LI Guang-qun, WU Guo-hua, FAN Yu, DING Wen-jiang. Effect of the main alloying elements on microstructure and corrosionYun-long DING, et al/Trans. Nonferrous Met. Soc. China 28(2018) 251−258 258resistance of magnesium alloys [J]. Foundry Technology, 2006, 27: 79−83. (in Chinese)[16]ZHANG Shi-chang, DUAN Han-qiao, CAI Qi-zhou, WEI Bo-kang,LIN Han-tong, CHEN Wei-chen. Effects of the main alloying elements on microstructure and properties of magnesium alloys [J].Foundry, 2001, 50: 310−315. (in Chinese)[17]ZHANG Ding-quan, HE Jia-wen. Residual stress analysis by X-raydiffraction and its functions [M]. Xi’an: Xi’an Jiao Tong University,1999. (in Chinese)[18]WANG Jian-gang, JU Dong-ying. Study on evolution technology ofanisotropic mechanical properties and micro-structures evolution of metal thin plate under complex stress condition [D]. Japan: Saitama Institute of Technology, 2011.[19]ZHAO L M, ZHANG Z D. Effect of Zn alloy interlayer on interfacemicro-structures and strength of diffusion-bonded Mg−Al joints [J].Scripta Materialia, 2008, 58: 283−286.退火温度对镁合金和铝合金扩散结合层的影响丁云龙1，王建刚2，赵明3，巨东英1,31. Department of Material Science and Engineering, Saitama Institute of Technology,Fusaiji 1690, Fukaya, Saitama 369-0293, Japan;2. 河北科技大学河北省材料近净成形技术重点实验室，石家庄050018；3. 辽宁科技大学材料科学与工程系，鞍山114051摘要：为了研究退火温度对镁合金和铝合金结合层的影响，提高扩散层的性能，采用真空扩散结合的方法焊接镁合金(AZ31B)和铝合金(6061)，然后依据热处理理论在电炉中进行退火处理。

第29卷第2期2008年　4月河南科技大学学报:自然科学版Journal of Henan University of Science and Technol ogy:Natural Science Vol .29No .2Ap r .2008基金项目:国家自然科学基金专项基金项目(50323008)作者简介:武从海(1981-),男,湖北监利人,硕士.收稿日期:2007-08-10文章编号:1672-6871(2008)02-0008-03硬质合金TiN 涂层的残余热应力数值分析武从海1,黄自谦2(1.河南科技大学理学院,河南洛阳471003;2.中南大学粉末冶金国家重点实验室,湖南长沙410083)摘要:采用有限元方法分析了硬质合金氮化钛涂层的热残余应力,试样采用圆柱体状的轴对称几何模型和自由边界条件,分析了涂层厚度、基体参数和过渡层对残余热应力的影响。

计算表明:Ti N 涂层中出现较大的拉应力(1GPa 以上),而硬质合金基体中出现较大的压应力(-0.8GPa 左右)。

涂层中的拉应力随硬质合金基体中钴的质量分数和涂层厚度的增加而明显减小,因此通过增加基体中的钴的质量分数和涂层厚度可以减小涂层中的拉应力。

当采用Ti C 过渡层时,涂层应力可以减小25%,且采用Ti CN 梯度过渡层时效果更好。

关键词:Ti N 涂层;有限元方法;残余应力;硬质合金中图分类号:TF124.5文献标识码:A0　前言由于普通硬质合金存在硬度和韧性的矛盾,要提高材料的耐磨性,必须牺牲其韧性,反之,要提高材料的韧性和强度,又要降低其硬度和耐磨性。

为了在保证其强度的条件下提高其耐磨性,文献[1-2]采用涂层的方法,即在硬质合金表面镀上一层Ti N ,Ti C 或金刚石涂层,来提高其耐磨性和寿命。

诚然,涂层的添加对于解决硬质合金的硬度和韧性的矛盾起到了很好的作用,但是,由于涂层和基体热机械性能的差异,在热处理时由高温下降到室温时,由热膨胀系数的差异引起不均匀的热收缩,导致残余热应力的产生。

Short CommunicationEffect of the annealing temperature on the microstructural evolution and mechanical properties of TiZrAlValloyR.Jing a ,⇑,S.X.Liang a ,b ,C.Y.Liu a ,M.Z.Ma a ,R.P.Liu a ,⇑a State Key Laboratory of Metastable Materials Science and Technology,Yanshan University,Qinhuangdao 066004,China bCollege of Equipment Manufacture,Hebei University of Engineering,Handan 056038,Chinaa r t i c l e i n f o Article history:Received 13December 2012Accepted 15June 2013Available online 27June 2013a b s t r a c tThis study aimed to evaluate the effects of the annealing temperature on the structural evolution and mechanical properties of TiZrAlV alloy.The microstructural evolution and mechanical properties of the alloy were investigated by X-ray diffraction,metallographic analysis,tensile testing,and microhardness testing.The results showed that the thickness of the a phase that precipitated from the parent phase was sensitive to the annealing temperature.With increased annealing temperature,the a -phase tended to exhibit equiaxed grains,except for the specimen annealed at 1050°C.The tensile strength of the equi-axed a grains were also demonstrated to have higher tensile strength than those of the lamellar a phase.The optimal mechanical properties of the alloy was obtained after annealing at 850°C,i.e.,r b =1245MPa,r 0.2=1006MPa,and e =16.89%.Ó2013Elsevier Ltd.All rights reserved.1.IntroductionTitanium-based alloys are increasingly being used as structural materials in the aerospace and automotive industries because of their remarkable advantages,such as exceptional strength-to-weight ratio,good hardenability,good elevated temperature performance,excellent fatigue/crack-propagation behavior,and corrosion resistance [1,2].Compared with other conventional stainless steel or structural materials,the mechanical properties of Ti alloys enable their weight to be reduced to about 40%in aerospace and automotive applications [3,4].Currently,the main-stream Ti structural material is the a +b phase Ti–6Al–4V alloy because of its better physical and mechanical properties than com-mercial-purity Ti and other Ti alloys.The a +b phase Ti–6Al–4V alloy is often used in aerospace applications,pressure vessels,blades and discs of aircraft turbines and compressors,surgical implants,etc.[5–8].The mechanical properties of dual-phase Ti alloys are closely related to their microstructure.The metallurgical processes such as thermo-mechanical processing and different heat treatment methods,which bring modifications in the micro-structure,can strongly influence their mechanical properties of these alloys [9].The majority of commercially used dual-phase Ti alloys are usually thermo-mechanically processed and subjected to different heat treatments to obtain the ideal microstructure for the desired application.In general,these alloys exist as two typical microstructures,namely,Widmanstätten lath precipitateof the hexagonal close-packed a phase distributed in a matrix of body-centered cubic b phase,and the combination of some equi-axed a -phase grains distributed in a transformed b phase.In general,Ti alloys have low hardness (HV 300–320)and yield strength (880–900MPa)[10].Previous studies [11]have used zir-conium,which has similar chemical properties to Ti,as an alloying element to strengthen Ti–6Al–4V alloy,even though zirconium is considered a neutral element [12,13].The addition of 20%(by mass)Zr to Ti–6Al–4V alloy has been experimentally found to in-crease the alloy strength and microhardness with acceptable elon-gation.In this work,different microstructures of the alloy were obtained by controlling the annealing process.The mechanical properties of the alloy were found to be very sensitive to the annealing temperature.2.Experimental procedureThe alloy used for this study is prepared by electromagnetic induction melting the mixture of sponge Ti (99.7wt%),sponge Zr (Zr +Hf P 99.5wt%),industrially pure Al (99.5wt%)and V (99.9wt%)under an argon atmosphere.Table 1shows the chemical composition of the studied alloy.The alloy was then flipped and re-melted three times to ensure a homogeneous chemical compo-sition.The ingot used in the experiment was homogenized at 1000°C for 12h,followed by cooling to room temperature.Then the ingot underwent multiple breakdowns after being held at 1000°C above the b transus temperature for 90min to completely break the coarse grains.The ingot was held at 900°C for 90min and then subjected to the final heat forging in the a +b phase0261-3069/$-see front matter Ó2013Elsevier Ltd.All rights reserved./10.1016/j.matdes.2013.06.039Corresponding authors.Tel.:+863358074723;fax:+863358074545.E-mail addresses:qwe_jr@ (R.Jing),riping@ (R.P.Liu).982R.Jing et al./Materials and Design52(2013)981–986Fig.1.DSC curve of TiZrAlV alloy.region,and the ingot was lathed into bar40mm in diameter.Thesamples(approximately10mmÂ10mmÂ70cm)were cut fromthe bar using wire-electrode cutting and used for subsequentannealing trials.Differential scanning calorimetry(DSC)was used to determinethe phase transition temperatures with a heating rate°C/min which was adopted the standard of ASTM:F2004–05(2010).The nominal a?a+b transus temperature andb?b transus temperature for TiZrAlV alloy are about789and946°C,respectively,as shown in Fig.1.Heat treatment wasperformed in a tubular vacuum furnace under a protective argonpatterns of TiZrAlV alloy:(a)forging,(b)annealing treatment at different temperatures,and(c)detail of33–43°of forging and annealingsignified that the alloy only formed the solid solution phase and that no other intermetallic compound and/or phase existed (Fig.2a).A comparison of the XRD patterns at different annealing temperatures (Fig.2b)revealed that the phase composition of all annealed alloys consisted of a and b phases.With increased annealing temperature,the b phase (110)reflection peak near 38°gradually broadened and the intensity of the (110)diffraction peak increased.However,at 1050°C annealing temperature,the b phase (200)peak disappeared.The XRD patterns also showed that the proportion of a and b phases evidently changed with the chan-ged in annealing temperature.This phenomenon may be caused by the difference of the migration rate of the atom under the high temperature.Generally,with the temperature increasing,the fre-quencies of the atoms migration are also increased gradually.In the insulation process,the moving distancesof Al atom (which is a -stabilized element)and V atom (which is b -stabilized element)were different,and in the subsequent cooling process,the b phase transformed into the a phase which caused the Al atom enriched in the b phase lattice and changed the b phase lattice parameters.Therefore,it may make the intensity of the b phase (110)reflection peak increase and the (002)reflection peak decrease when the annealing temperature was heated to 1050°C.Furthermore,the annealing holding time was shorter (30min),in this process,the a phase transformed into the b phase may be incomplete at an-nealed treatment at 1000°C,while annealed temperature was in-creased to 1050°C,the a phase may be completely transformed into the b phase,therefore,in the specimen annealed at 1000°C the initial a phase was also existed,but the specimen which was annealed at 1050°C did not exist the initial a phase.This may re-sult that the differences of a phase between 1000°C and 1050°C is obviously.Fig.3shows the microstructure of the annealing temperatures.The specimens ited Widmanstätten morphology (Fig.3a),i.e.,chaotic arrangement of slender a lath and b annealing temperature to 1000°C,the b peared and the a lath gradually (Fig.3b–e).In this process,the alloy axed trend with increased annealing have caused the increased equiaxed a phase the annealing process.First,the lamellar a ‘‘interleaved,’’which restricts the other a longitudinal direction.Consequently,a only along the transverse direction,which promotes the thickening of the a lamellar.Second,the new a phase that precipitates from the parent phase grows along a specific habit plane and has a cer-tain orientation relationship with the primary a phase.Thus,the new precipitated a phase growing along the longitudinal direction is hindered such that the equiaxed degree is increased.Obasi [14]also indicated that the phase transformation in Ti alloys during heating (a ?b )and cooling (b ?a )is governed by the so-called Burgers orientation relationship {0002}a ||{110}b and h 11À20i a ||h 111i b with 6possible b -orientations during the a ?b phase transformation and 12possible a orientations that can transform from a single parent b grain during b ?a phase transformation.However,when the annealing temperature reached 1050°C,the alloy revealed the typical basketweave mor-phology (Fig.3f),i.e.,a crisscross slender a lath.b grain boundaries and some parallel lamellar the grain boundaries (the Widmanstätten microstructure)observed in this process.In most diffusion phase and precipitation processes,the nucleations of the heterogeneously occurs at some preferential nucleation the matrix such as the grain boundary,dislocation,phase [15].When the annealing temperature (e.g.,Optical microstructure of TiZrAlV alloy under different annealing temperature:(a)800°C,(b)850°C,(c)900°C,(d)950°C,(e)1000°C,and Fig.4.True stress–strain curve of the studied alloy under different conditions.the thickness of the a lath became limited.Therefore,the thicknessof the new precipitated a phase after annealing at1050°C was smaller than that after annealing at1000°C.The mechanical properties of the alloy were evaluated through uniaxial tensile tests.Fig.4and Table2show the true stress–strain curves and mechanical properties of the specimens at different annealing temperatures.The mechanical properties of the ZrTiAlV alloy evidently depended on the annealing temperature and micro-structure.When the annealing temperatures were between800 and1000°C,the yield strength r0.2and ultimate strength r b de-creased from1009and1290MPa to978and1181MPa,respec-tively.The elongation only slightly changed after annealing at of the residual b phase during the annealing process as well as the thickness of the a lath.The main factors influencing the mechanical properties of an-nealed samples in which only the a and b phases exist are the phase content,size,and morphology of the a phase[16–18].Because of the limited number of independent slips modes,the hcp structure of Ti exhibits a vary strong grain-boundary,or Hall–Petch strength-ening at room temperature.The thickness of the a grain boundary directly influences the strength mismatch between the a+b matrix and the grain boundary[19].Consider the case of b processed microstructures.Some of the microstructural features involved with progressively increasing length scales are width of the a-laths, the colony size,and the b grain size(feature sizes may range from sub-micron to millimeters).Depending on the thermo-mechanical treatment the alloy is subjected to,such as cooling rates from +b dual phase region or above the b-transus,these features can vary significantly.Quantifying them over the diverse range length scales involved becomes rather important.Thus,to investi-gate the effect of the annealing temperature on the microstructure and mechanical properties,the specimens prepared at different annealing temperatures were subjected to SEM analysis,as shown Fig.6.The measured thicknesses of the a lath from the SEM images are shown in Fig.7a.With increased annealing temperature Fig.5.Microhardness of annealed specimens under different conditions.SEM images of TiZrAlV alloy under different annealing temperature:(a)800°C,(b)850°C,(c)900°C,(d)950°C,(e)1000°C,and(f)from 800°C to 1000°C,the thickness of the a lath in the annealed samples increased from 1.07l m to 4.22l m.When the annealing temperature reached 1050°C,the thickness was reduced to 1.12l m.According to the Hall–Petch equation,(i.e.,r =r 0+kd À1/2,where d is the thickness of the a lath),the strength of an annealed alloy is related to the a lath thickness,as shown in Fig.7b.On one hand,the change of the a lath thickness moved the distance of dis-location to the phase boundary,which resulted in increased num-ber of dislocations piling up such that the stress concentration was more severe.On the other hand,reducing the a lath thickness increased the density of the phase boundary in the same cross-sectional area.Consequently,the movement of the dislocation obstacle increased.Thus,based on the OM images,SEM images,and true stress–strain curve,the slender a lath obtained at 800°C and 1050°C increased the strength of the specimens and made dis-location movement difficultly.Moreover,with increased equiaxed a -phase degree,the strength of the annealed specimens gradually decreased.This result implied that the strength of the processed alloy lamellar a phase microstructure was higher than that of processed equiaxed a -phase microstructure.The magnitude of the titanium alloy tensile elongation is con-nected with the non-uniform degree of the tensile micro deforma-tion zone,as well as the length and the spacing of slip bands.With the spacing of slip bands decreasing,the plastic deformability in-crease before the material fracture [20].Compared with the lamel-lae microstructure,the slip bands spacing of the duplex microstructure is smaller,thus this microstructure possess a high-er ability of deformation.When sample was annealed at 850°C,the feature of microstructure presented the duplex microstructure (Fig.3b),therefore,the elongation reached the largest value in this experiment i.e.16.89%.4.ConclusionThe phase transition,microstructure evolution,and their effects on the mechanical properties of TiZrAlV alloy were investigated.The conclusions were as follows:(1)TiZrAlV alloy exhibited an a +b phase after high-tempera-ture annealing.The intensity of the b (110)diffraction peak increased with increased annealing temperature.However,the intensity of the b (200)diffraction peak gradually decreased with increased annealing temperature.When the temperature reached 1050°C,the b (200)diffraction peak completely disappeared.(2)The thickness of the a phase was sensitive to the annealingtemperature.With increased annealing temperature,the a phase tended to exhibit equiaxed grains,except for speci-mens annealed at 1050°C.After annealing at 1000°C,the maximum thickness was 4.22l m.(3)The mechanical properties of the annealed specimens weresensitive to the morphology of the precipitated a phase and yo the annealing temperature.The optimal mechanical properties of the alloy were obtained after annealing at 850°C,i.e.,r b =1245MPa,r 0.2=1006MPa,and e =16.89%.AcknowledgmentsThis work was supported by the SKPBRC (Grant No.2010CB731600),NSFC (Grant No.51121061/51171160/51171163).References[1]Eylon D,Vassel A,Combres Y,Boyer RR,Bania PJ,Schutz RW.Issues in thedevelopment of beta titanium alloys.JOM 1994;46:14–5.[2]Ivasishin 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