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Ferroelectric and piezoelectric properties of tungsten substituted SrBi 2Ta 2O 9ferroelectric ceramicsIndrani Coondoo *,S.K.Agarwal a ,A.K.Jha ba Superconductivity and Cryogenics Division,National Physical Laboratory,Dr K.S.Krishnan Road,New Delhi 110012,India bDepartment of Applied Physics,Delhi College of Engineering,Bawana Road,Delhi 110042,India1.IntroductionDefects in crystals signiﬁcantly inﬂuence physical and various other properties of materials [1].For instance,as it is well known,doping by other elements leads to signiﬁcant changes in the electrical properties of silicon.Historically,‘‘defect engineering’’has been developed in the ﬁeld of semiconducting materials such as compound semiconductors as well as in diamond,Si and Ge [2–4].Subsequently,the concept of defect engineering has been applied to other functional materials,and the signiﬁcant improve-ment in material properties have been achieved in high transition-temperature superconductors [5],amorphous SiO 2[6],photonic crystals [7]and also in the ﬁeld of ferroelectrics,such as BaTiO 3,Pb(Ti,Zr)O 3(PZT),etc.[8,9].Various structural and electrical properties of bismuth layer-structured ferroelectrics (BLSF)are also strongly affected on deviation from stoichiometric composi-tions and defects have been recognized as a crucially important factor [10–13].It has been found that in BLSF small changes in chemical composition result in signiﬁcantly altered dielectric and ferroelectric properties including dielectric constant and remanent polarization.In SrBi 2Ta 2O 9(SBT)and SrBi 2Nb 2O 9(SBN),orthor-hombic structural distortions with non-centrosymmetric spacegroup A 21am cause spontaneous ferroelectric polarization (P s )along a axis [14,15].SBT,a member of the BLSF family,has occupied an important position among the Pb-free ferroelectric memory materials [16–18].Tungsten (W 6+)has recently been investigated as a dopant for bismuth titanates and lanthanum doped bismuth titanates,in which the remanent polarization was observed to enhance when a small amount of Ti 4+was substituted by W 6+[19,20].With the objective to improve structural,dielectric and ferroelectric proper-ties,the hexavalent tungsten (W 6+)was chosen as a donor cation for partial replacement of the pentavalent tantalum (Ta 5+)SBT.In this report,the effect of tungsten substitution in SBT (SBTW),on the microstructural,ferroelectric and piezoelectric properties is reported.The results including the improvement in polarization properties have been discussed.2.ExperimentalSamples of compositions SrBi 2(W x Ta 1Àx )2O 9(SBWT),with x =0.0,0.025,0.050,0.075,0.10and 0.20were synthesized by solid-state reaction method taking SrCO 3,Bi 2O 3,Ta 2O 5and WO 3(all from Aldrich)in their stoichiometric proportions.The powder mixtures were thoroughly ground and passed through sieve of appropriate size and then calcined at 9008C in air for 2h.The calcined mixtures were ground and admixed with about 1–1.5wt%polyvinyl alcohol (Aldrich)as a binder and then pressed at $300MPa into disk shaped pellets.The pellets were sintered at 12008C for 2h in air.Materials Research Bulletin 44(2009)1288–1292A R T I C L E I N F O Article history:Received 3October 2008Received in revised form 5December 2008Accepted 6January 2009Available online 15January 2009Keywords:A.CeramicsC.X-ray diffractionD.FerroelectricityA B S T R A C TTungsten substituted samples of compositions SrBi 2(W x Ta 1Àx )2O 9(x =0.0,0.025,0.050,0.075,0.10and 0.20)were synthesized by solid-state reaction method and studied for their microstructural,electrical conductivity,ferroelectric and piezoelectric properties.The X-ray diffractograms conﬁrm the formation of single phase layered perovskite structure in the samples with x up to 0.05.The temperaturedependence of dc conductivity vis-a`-vis tungsten content shows a decrease in conductivity,which is attributed to the suppression of oxygen vacancies.The ferroelectric and piezoelectric studies of the W-substituted SBT ceramics show that the remanent polarization and d 33values increases with increasing concentration of tungsten up to x 0.05.Such compositions with low conductivity and high P r values should be excellent materials for highly stable ferroelectric memory devices.ß2009Elsevier Ltd.All rights reserved.*Corresponding author.Present address:Liquid Crystal Group,National Physical Laboratory,Dr K.S.Krishnan Road,New Delhi 110012,India.Tel.:+919810361727;fax:+911125170387.E-mail address:indrani_coondoo@ (I.Coondoo).Contents lists available at ScienceDirectMaterials Research Bulletinj o ur n a l h o m e p a g e :w w w.e l se v i e r.c om /l oc a t e /m a t r e sb u0025-5408/$–see front matter ß2009Elsevier Ltd.All rights reserved.doi:10.1016/j.materresbull.2009.01.001X-ray diffractograms of the sintered samples were recorded using a Bruker diffractometer in the range 108 2u 708with CuK a radiation.The sintered pellets were polished to a thickness of 1mm and coated with silver paste on both sides for use as electrodes and cured at 5508C for half an hour.Electrical conductivity was performed using Keithley’s 6517A Electrometer.The polarization–electric ﬁeld (P –E )hysteresis measurements were done at room temperature using an automatic P –E loop tracer based on Sawyer–Tower circuit.Piezoelectric charge co-efﬁcient d 33was measured using a Berlincourt d 33meter after poling the samples in silicone–oil bath at 2008C for half an hour under a dc electric ﬁeld of 60–70kV/cm.3.Results and discussion3.1.Structural and micro-structural studiesThe phase formation and crystal structure of the ceramics were examined by X-ray diffraction (XRD),which is shown in Fig.1.The XRD patterns of the samples show the characteristic peaks of SBT.The peaks have been indexed with the help of a computer program–POWDIN [21]and the reﬁned lattice parameters are given in Table 1.It is observed that a single phase layered perovskite structure is maintained in the range 0.0 x 0.05.Owing to the same co-ordination number i.e.6and the smallerionic radius of W (0.60A˚)in comparison to Ta (0.64A ˚),there is a high possibility of tungsten occupying the tantalum site.The observance of unidentiﬁed peak of very low intensity in the compositions with x >0.05indicates the solubility limit of W concentration in SBT.The unidentiﬁed peak is possibly due to tungsten not occupying the Ta sites in the structure as the intensity of this peak is observed to increase with tungsten content.Composition and sintering temperature inﬂuences the micro-structure such as grain growth and densiﬁcation of the specimen,which in turn control other properties of the material [11,13].The effects of W substitution on the microstructure have been examined by SEM and the obtained micrographs are shown in Fig.2.It shows the microstructure of the fractured surface of the studied samples.It is clearly observed that W substitution has pronounced effect on the average grain size and homogeneity of the grains.Randomly oriented and anisotropic plate-like grains are observed in all the samples.It is also observed that the average grain size increases gradually with increasing W content.The average grain size in the sample with x =0.0is $2–3m m while that in the sample with x =0.20the size increases to $5–7m m.3.2.Electrical studiesThe electrical conductivity of ceramic materials encompasses a wide range of values.In insulators,the defects w.r.t.the perfect crystalline structure act as charge carriers and the consideration of charge transport leads necessarily to the consideration of point defects and their migration [22].Many mechanisms were put forward to explain the conductivity mechanism in ceramics.Most of them are approximately divided into three groups:electronic conduction,oxygen vacancies ionic conduction,and ionic and p-type mixed conduction [22].Intrinsic conductivity results from the movement of the component ions,whereas conduction resulting from the impurity ions present in the lattice is known as extrinsic conductivity.At low temperature region (ferroelectric phase),the conduction is dominated by the extrinsic conduction,whereas the conduction at the high-temperature paraelectric phase ($300–7008C)is dominated by the intrinsic ionic conduction [23,25].Fig.3shows the temperature dependence of dc conductivity (s dc )for the undoped and doped SBT samples.The curves show that the conductivity increases with temperature.This is indicative of negative temperature coefﬁcient of resistance (NTCR)behavior,a characteristic of dielectrics [22].It is observed in Fig.3that throughout the temperature range,the dc conductivity of the doped samples are nearly two to three orders lower than that of the undoped sample.Two predominant conduction mechanisms indicated by slope changes in the two different temperature regions are observed in Fig.3.Such changes in the slope in the vicinity of the ferro-paraelectric transition region have been observed in other ferroelectric materials as well [23,24].In addition,it is also observed (Table 2)that the activation energy calculated using the Arrhenius equation [22]in the paraelectric phase increase from $0.80eV for the undoped sample to $2eV for the doped samples.The X-ray photoemission spectroscopic study has conﬁrmed that when Bi 2O 3evaporates during high-temperature processing,vacancy complexes are formed in the (Bi 2O 2)2+layers [26].As a result,defective (Bi 2O 2)2+layers are inherently present in SBT.The undoped SBT shows n-type conductivity,since when oxygen vacancies are created,it leaves behind two trapped electrons [27]:O o !12O 2"þV o þ2e 0(1)where O o is an oxygen ion on an oxygen site,V o is a oxygen vacant site and e 0represents electron.The conductivity in the perovskites can be described as an ordered diffusion of oxygen vacancies [28].Their motion is manifested by enhanced ionic conductivity associated with an activation energy value of $1eV [26].These oxygen vacancies can be suppressed by addition of donors,since the donor oxide contains more oxygen per cation than the host oxide it replaces [29].It has been reported that conductivity in Bi 4Ti 3O 12(BIT)can be signiﬁcantly decreased,up to three orders of magnitude with the addition of donors,such as Nb 5+and Ta 5+at the Ti 4+sites [23,30].A few other studies on layered perovskites have also reported a decrease inconductivityFig.1.XRD patterns of SrBi 2(W x Ta 1Àx )2O 9samples sintered at 12008C.Table 1Lattice parameters of SrBi 2(W x Ta 1Àx )2O 9samples.Concentration of W a (A ˚)b (A ˚)c (A ˚)0.0 5.5212 5.513924.92230.025 5.5214 5.520225.10790.05 5.5217 5.519925.05850.075 5.5191 5.504525.05670.10 5.5142 5.506125.0850.205.51335.493925.0861I.Coondoo et al./Materials Research Bulletin 44(2009)1288–12921289with addition of donors [23,24,31].In the present study,the Ta 5+-site substitution by W 6+in SBT can be formulated using a defect chemistry expression as WO 3þV o!Ta 2O 512W Ta þ3O o (2)It shows that the oxygen vacancies are reduced upon the substitution of donor W 6+ions for Ta 5+ions.Hence,it is reasonable to believe that the conductivity in SBT is suppressed by donor addition.As per the above discussion,the high s dc observed in the undoped SBT (Fig.3)can be attributed to the motion of oxygen vacancies.As already discussed,the doped samples show reduced conductivity because the transport phenomena involving oxygen vacancies are greatly reduced.The high E a value of $1.75–2eVcorresponding to the high-temperature region in the doped ceramics is consistent with the fact that in the donor-doped materials,the ionic conduction reduces [32].The activation energy E a in the low temperature ferroelectric region (Table 2)corre-sponds to extrinsic conduction.At lower temperatures the extrinsic conductivity results from the migration of impurity ions in the lattice.Some of these impurities may also be associated with lattice defects.Pure SBT has large number of Schottky defects (oxygen vacancies)in addition to impurity ions whereas in the doped samples,due to charge neutrality,there is relatively less content of oxygen vacancies.Thus,in the doped samples the conductivity in the low temperature region is largely due to the impurity ions only.This explains the high activation energy in pure SBT in the low temperature region compared to doped samples (Table 2).In the high-temperature region,the value of E a in the doped samples is observed to increase with W concentration up to x =0.05but beyond that,it decreases (Table 2).The decrease in the activation energy for samples with x >0.05suggests an increase in the concentration of mobile charge carriers [33].This observation can be ascribed to the existence of multiple valence states of tungsten.Since tungsten is a transitional metal element,the valence state of W ions in a solid solution most likely varies from W 6+to W 4+depending on the surrounding chemical environment [34].When W 4+are substituted for the Ta 5+sites,oxygen vacancies would be created,i.e.one oxygen vacancy would be created for every two tetravalent W ions entering the crystal structure,whichFig.3.Variation of dc conductivity with temperature in SrBi 2(W x Ta 1Àx )2O 9samples.Fig.2.SEM micrographs of fractured surfaces of SrBi 2(W x Ta 1Àx )2O 9samples with (a)x =0.0,(b)x =0.025,(c)x =0.050,(d)x =0.075,(e)x =0.10and (f)x =0.20Table 2Activation energy (E a )in the high-temperature paraelectric region and low temperature ferroelectric region;Curie temperature (T c )in SrBi 2(W x Ta 1Àx )2O 9samples.Concentration of W E a (high temp.)(eV)E a (low temp.)(eV)T c (8C)0.00.790.893110.025 1.920.593080.05 1.960.543250.075 1.940.543380.10 1.860.573680.201.740.54390I.Coondoo et al./Materials Research Bulletin 44(2009)1288–12921290explains the increase in the concentration of mobile charge carriers which ultimately results in an decrease in the E a beyond x>0.05. Hence it is reasonable to conclude that W ions in the SBWT exists as a varying valency state,i.e.at lower doping concentration they exist in hexavalent state(W6+)and at a higher doping concentra-tion,they tend to exist in lower valency states[8].The P–E loops of SrBi2(Ta1Àx W x)2O9are shown in Fig.4.It is observed that W-doping results in formation of well-deﬁned hysteresis loops.Fig.5shows the compositional dependence of remanent polarization(2P r)and the coerciveﬁeld(2E c)of SrBi2(Ta1Àx W x)2O9samples.Both the parameters depend on W content of the samples.It is observed that2P rﬁrst increases with x and then decreases while2E cﬁrst decreases with x and then increases(Fig.5).The optimum tungsten content for maximum2P r ($25m C/cm2)is observed to be x=0.075.It is known that ferroelectric properties are affected by compositional modiﬁcation,microstructural variation and lattice defects like oxygen vacancies[10,35,36].In hard ferroelectrics, with lower valent substituents,the associated oxide vacancies are likely to assemble in the vicinity of domain walls[37,38].These domains are locked by the defects and their polarization switching is difﬁcult,leading to an increase in E c and decrease in P r[38]. On the other hand,in soft ferroelectrics,with higher valent substituents,the defects are cation vacancies whose generation in the structure generally increases P r.Similar observations have been made in many reports[38–41].Watanabe et al.[42]reported a remarkable improvement in ferroelectric properties in the Bi4Ti3O12ceramic by adding higher valent cation,V5+at the Ti4+ site.It has also been reported that cation vacancies generated by donor doping make domain motion easier and enhance the ferroelectric properties[43].Further,it is known that domain walls are relatively free in large grains and are inhibited in their movement as the grain size decreases[44].In the larger grains, domain motion is easier which results in larger P r.Also for the SBT-based system,it is known that with increase in the grain size the remanent polarization also increases[45,46].Based on the obtained results and above discussion,it can be understood that in the undoped SBT,the oxygen vacancies assemble at sites near domain boundaries leading to a strong domain pinning.Hence,as observed,well-saturated P–E loop for pure SBT is not obtained.But in the doped samples,the suppression of the oxygen vacancies reduces the pinning effect on the domain walls,leading to enhanced remanent polarization and lower coerciveﬁeld.Also,the increase in grain size in tungsten added SBT,as observed in SEM micrographs(Fig.2)contribute to the increase in polarization values.In the present study,the grain size is observed to increase with increasing W concentration.However, the2P r values do not monotonously increase and neither the E c decreases continuously with increasing W concentration(Fig.5). The variation of P r and E c beyond x>0.05,seems possibly affected by the presence of secondary phases(observed in XRD diffracto-grams),which hampers the switching process of polarization [47–50].Also,beyond x>0.05the increase in the number of charge carriers in the form of oxygen vacancies leads to pinning of domain walls and thus a reduction in the values of P r and increase in E c is observed.Fig.6shows the variation of piezoelectric charge coefﬁcient d33 with x in the SrBi2(Ta1Àx W x)2O9.The d33values increases with increase in W content up to x=0.05.A decrease in d33values is observed in the samples with x!0.075.The piezoelectric coefﬁcient,d33,increases from13pC/N in the sample with x=0.0to23pC/N in the sample with x=0.05.It is known that the major drawback of SBT is its relatively higher conductivity,which hinders proper poling[51].High resistivity is therefore important for maintenance of poling efﬁciency at high-temperature[52,53].The W-doped SBT samples show an electrical conductivity value up to three orders of magnitude lower than that of undoped sample(Fig.3).The positional variation of2P r and2E c in SrBi2(W x Ta1Àx)2O9samples.Fig.6.Variation of d33in SrBi2(W x Ta1Àx)2O9samples.Fig. 4.P–E hysteresis loops in SrBi2(W x Ta1Àx)2O9samples recorded at roomtemperature.I.Coondoo et al./Materials Research Bulletin44(2009)1288–12921291decrease in conductivity upon donor doping improve the poling efﬁciency resulting in the observed higher d33values.Moreover, since the grain size increases with W content in SBT,it is reasonable to believe that the increase in grain size will also contribute to the increase in d33values[54].The decrease in the value of d33for samples with x!0.075is possibly due to the presence of secondary phases as observed in diffractograms[1,51,55]and the increase in oxygen vacancies for samples with x>0.05.4.ConclusionsX-ray diffractograms of the samples reveal that the single phase layered perovskite structure is maintained in the samples with tungsten content x0.05.SEM micrographs reveal that the average grain size increases with increase in W concentration. The temperature dependence of the electrical conductivity shows that tungsten doping results in the decrease of conductivity by up to three order of magnitude compared to W free SBT.All the tungsten-doped ceramics have higher2P r than that of the undoped sample.The maximum2P r($25m C/cm2)is obtained in the composition with x=0.075.The reduced conductivity allows high-temperature poling of the doped samples.Such compositions with low loss and high P r values should be excellent materials for highly stable ferroelectric memory devices.The d33value is observed to increase with increasing W content up to x0.05.The value of d33 in the composition with x=0.05is$23pC/N as compared to$13 pC/N in the undoped sample.AcknowledgmentsThe authors sincerely thank Prof.P.B.Sharma,Dean,Delhi College of Engineering,India for his generous support and providing ample research infrastructure to carry out the research work.The authors are thankful to Dr.S.K.Singhal,Scientist, National Physical 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材料科学与工程专业英语

adhesive［胶］allotropic［同素异形的amorphous［无定形的，非晶的anion［负离子］apuy［适当地］austenite［奥氏体］bainte「贝氏体binary isomorphous system［二元匀晶系统］Burgers vector［柏氏矢量］cadmium［镉］canon［正离子］carbide［碳化物Cast Iron［铸铁cementite［渗碳体］ceramic［陶瓷］chloride［氯chromium［铬］composite materials［复合材料cordinate system［坐标系统covalent bond［共价键crystal structure［晶体结构］crystallinity［结晶度ddiy［塑性deteriorative［劣化］dislocation［位错］edge dislocation［刃形位错］equilibrium［平衡］eutectic［共晶的eutectoid［共析的ferrite［铁素体fractional［分数的，部分的，相对的gain［晶粒grain boundary［晶界hase［相hexagonal dose－packed［密排六方的hexagonal［六方的］hypereutectoid［过共析hypoeutectoid［亚共析ion［硅Ionic Dond［离子键］isotherm［等温线lamellae［薄片］lattice［空间点阵，晶格］lever law［杠杆定律liquidus line［液相线martensite［马氏体］martensitic transformation［马氏体相变metallic bond［金属键］microstucture［显微组织monoclinic［单斜的nickel［镍nitride［氮化物］non－crystalline［非晶的］orthorhombic［正交的］parentheses［括孤］pearlite［珠光体periodic table［元素周期表phase diagram［相图］」phase transformation［相交］」point defect［点缺陷］polarize［极化polyethylene［聚乙烯polymerization［聚合］prism［棱镜］proeutectoid［先共析体provoke［诱发］reciprocal［倒数recrystallization［再结晶rhombohedral［菱方的screw dislocation［螺形位错］skew［歪斜］smal－（or low）angle grain boundary［小角度晶界sodium［钠solar cell ［太阳能电池］solid solution strengthening［固溶强化solidus line［固相线solute［溶质solvus line［溶解度曲线spiral［螺旋形的］stifness［刚度］strucure［组织］synthesis［合成tetragonal［四方的tetrahedron［四面体thermoplastics［热塑性塑料］thermosets［热固性塑料tie linc［连接线］tilt boundary［倾侧晶界］translucent［半透明的］triclinic［三斜的troley［石油twin boundary［李晶界unit cell［晶胞vacancy［空位valence electron［价电子］Van de Waals bond［范德华键①材料科学是研究材料的加工、组织、性能和功能之间关系的科学。

外文翻译---浇铸钛和金的显微结构和机械性能

附录一浇铸钛和金的显微结构和机械性能摘要：通过感应熔化的方法而获得的Ti21523合金，研究热处理和冷却凝固率对其显微结构和机械性能的影响和作用。

结果表明：通过增加冷却凝固率，可以使Ti21523合金的显微结构从单一化特征及大尺寸的粒状结构变成了具有优良性能的小尺寸粒状结构。

通过采用不同的方法和对不同时期的合金进行处理，合金相位逐渐在粒状晶体的内部和粒状晶体的边界上沉淀。

由于沉淀物晶相的改变，合金承受拉力的性能和伸长率同时被改良。

在σb=1. 406Gpa、δ=4. 5%时，将会获得一种具有良好性能的合金,在临界区域里使用这种合金会让我们收到满意的效果。

关键字浇铸Ti21523合金；冷却凝固率；机械性能1 介绍钛合金以其优良的机械性能，在飞机、航空航天和其它领域中，受到了人们的关注和认可，尤其是在较高特殊作用力的环境之下。

在降低航天器的质量并改进的它的运输适宜性上，该合金受到了关注。

为了满足以上两种情况，一种被叫做贝它钛的重要钛合金逐渐得到发展和优化。

由于其具有高抗力、弹性系数和伸长率等良好的综合性能，合金 Ti215V23Cr23Sn23Al(Ti21523) 已经变成了潜在的选择材料被用于在那些贝它类型合金之中。

从以上的论述中我们可以知道，Ti21523合金在室温有较好的可使用性，同时也适用于寒冷的工作环境之下。

不幸地是，由于合金的高处理成本以及诸如低可塑性和高刚度等缺点，使其在制造复杂的联合体和薄壁件时存在许多问题，成为影响其在航空航天业中广泛应用的关键所在。

为了降低其合成成本并达到其易于重新塑造的弹性，精密铸造技术被引入到了这个领域中。

但是由于铸造出来的合金其贝它晶粒较大且机械性能很低，故此Ti21523合金的使用受到了极大的限制。

由于热处理对Ti21523合金的力有影响，因此Ti21523合金还是可以改善其伸长率并提高它的机械性能的。

关于热处理对Ti21523合金的影响的研究首先在美国和前苏联开展。

TC11合金两相区退火对微观组织与微观织构的影响

TC11合金两相区退火对微观组织与微观织构的影响李赛毅;文浩;雷力明;张慧儒【摘要】采用金相、扫描电镜和电子背散射衍射(EBSD)研究两相区退火温度和时间对热压缩态 TC11(Ti-6.5Al-3.5Mo-1.5Zr-0.3Si)双相钛合金组织和微观取向的影响.结果表明：初始魏氏组织在850℃热压缩后被破坏，形成较为细小、扭折的片层组织，晶粒取向呈非均匀分布；在后续700和900℃退火过程中，α相变形组织和亚结构发生静态再结晶而转变为等轴状晶粒，晶粒尺寸较退火前的更为细小，再结晶程度和等轴组织比例随保温时间的延长而增加，片层组织的球化程度、晶粒取向和形貌的均匀性与合金的再结晶程度相关.在900℃退火时，α相的再结晶程度较700℃退火时的更为明显；经过120 min 退火后，合金发生完全再结晶，得到较为均匀、细小的等轴状组织.【期刊名称】《中国有色金属学报》【年(卷),期】2012(000)011【总页数】7页(P3022-3028)【关键词】钛合金;再结晶;退火;微观组织;微观织构【作者】李赛毅;文浩;雷力明;张慧儒【作者单位】中南大学材料科学与工程学院,长沙 410083; 中南大学有色金属材料科学与工程教育部重点实验室,长沙 410012;中南大学材料科学与工程学院,长沙410083;中航商用航空发动机有限责任公司,上海 200241; 北京航空材料研究院钛合金研究室,北京 100095;中南大学材料科学与工程学院,长沙 410083; 华南理工大学机械工程学院,广州 510640【正文语种】中文【中图分类】TG146.4TC11钛合金比强度高，具有优异的中温力学性能和成形性，是一种重要的航空和宇航结构材料。

与其它α+β型双相钛合金类似，TC11钛合金材料的制备通常需要先后在β相区和α+β两相区对铸锭进行锻压或镦拔以得到较为细小的晶粒组织[1-4]。

由于β相区变形所形成的魏氏组织在两相区变形后难以完全转变为具有较好综合性能的等轴状组织，往往需要进行两相区退火，使之发生静态再结晶而进一步球化[5-9]。

键合参数对AgAuPdRu键合线键合性能影响研究

键合参数对AgAuPdRu键合线键合性能影响研究范俊玲;朱丽霞;曹军;姚亚昔【摘要】利用扫描电镜研究φ0.025 mm的Ag-0.8Au-0.3Pd-0.03Ru银基键合线不同键合参数对无空气焊球(free air ball,FAB)形状及球焊点、楔焊点形貌的影响,结果表明,烧球电流或者烧球时间其中之一不变时,FAB直径随着烧球时间(烧球电流一定)和烧球电流(烧球时间一定)的增加而增大,烧球电流为30 rnA、烧球时间为700μs时,键合线FAB呈标准圆球形;键合压力一定时,键合线球焊点焊盘直径随着超声功率的增大而显著增加,键合功率为80 mW、键合压力0.30 N时,球焊点具有较好的微观形貌;键合压力为0.25N,超声功率为80 mW时,楔焊点具有规则的鱼尾形貌,并具有较好的连接强度.【期刊名称】《贵金属》【年(卷),期】2019(040)002【总页数】6页(P59-63,68)【关键词】银基键合线;无空气焊球(FAB);烧球时间;烧球电流;键合功率;键合压力【作者】范俊玲;朱丽霞;曹军;姚亚昔【作者单位】焦作大学化工与环境工程学院,河南焦作454000;焦作大学化工与环境工程学院,河南焦作454000;河南理工大学机械与动力工程学院,河南焦作454000;河南理工大学机械与动力工程学院,河南焦作454000【正文语种】中文【中图分类】TG146.3;TG457.19键合银线由于其优秀的电学性能，可降低器件高频噪声、降低大功率发光二极管(light-emitting diode，LED)发热量等，以及适当的成本因素，且在LED封装中可以有效降低光衰，提高转化率，键合银线的诸多优势，使其应用于微电子封装[1-6]，尤其是LED封装；但纯银线在应用过程中存在键合过程中参数窗口范围较小、强度较低，在低弧度引线封装中容易出现塌丝和线弧不稳定等缺陷[7]，以及高温条件下失效几率较高，无法满足大功率LED等器件的使用[8-9]等问题。

ti-fcc相中的孪晶结构英文

fcc twins: a structure favored by nature1. IntroductionFCC (face-centered cubic) structure, known for its high symmetry and close packing of atoms, has long been the focus of research in materials science. In recent years, the discovery of twin boundaries in FCC crystals has garnered significant attention from scientists and engineers. Twins, also known as "twin boundaries," are regions where the crystal structure is mirrored across a plane. In the case of FCC twins, the lattice is twinned across a {111} plane, resulting in a unique and fascinating crystal structure.2. Crystallography of FCC twinsThe twinning process in FCC crystals leads to the formation of two interpenetrating lattices, each related to the other by a 60-degree rotation about the twin boundary plane. This creates a distinctive microstructure with aplex network of dislocations and defects. The crystallographic orientation relationship between the twinned lattices has been extensively studied and is a subject of ongoing research.3. Mechanical properties of FCC twinsFCC twin boundaries have been found to significantly influence the mechanical behavior of materials. The presence of twin boundaries can impede dislocation motion and enhance the strength and ductility of FCC metals, making them highly desirable for structural applications. Additionally, the unique atomic arrangement at the twin boundary can give rise to unusual mechanical properties, such as superior resistance to deformation and high strain hardening.4. Twinning in natural and synthetic materialsNature has long favored the formation of twinned structures in minerals, where they can occur as lamellar or penetration twins. Synthetic materials, such as metallic alloys and intermetallicpounds, can also be engineered to exhibit FCC twin boundaries through controlled processing and alloy design. Understanding and harnessing the twinning phenomenon is therefore of great significance in both natural and engineered materials.5. Applications of FCC twinsThe presence of FCC twin boundaries has been linked to various functional properties in materials. For instance, twinning has been found to enhance the thermal stability and creepresistance of FCC metals, making them suitable for high-temperature applications. Furthermore, the presence of twins can lead to the emergence of unique electronic and magnetic properties, opening up new avenues for the development of advanced functional materials.6. Future prospects and challengesWhile significant progress has been made in the study of FCC twins, there are still many unanswered questions and challenges to be addressed. The precise mechanisms of twin formation and propagation, the influence of twin boundaries on thermal and electrical transport, and the impact of twins on the behavior of materials under extreme conditions are areas that require further investigation. Additionally, the development of novel techniques for the controlled manipulation of twinned structures represents an exciting avenue for future research and technological innovation.7. ConclusionIn conclusion, FCC twin boundaries represent a fascinating area of study with far-reaching implications for materials science and engineering. Their unique crystallographic, mechanical, and functional properties make them highly desirable for a widerange of applications. As our understanding of twin structures continues to advance, we can expect to see the emergence of new materials with enhanced performance and functionality, driven by the design and manipulation of FCC twin boundaries.。

工业纯钛TA2_紫铜T2异种金属_省略_MT焊接接头的微观组织和腐蚀性能_曹睿 (1)

分别对靠近钛母材的区域和靠间化合物层的组成，近焊缝金属的区域进行放大，并对各形态相进行能 Cu 二元相图［7］和 TiCuAl 三元谱点分析，结合 Ti［8 ］相图，钎焊界面 A 中相的成分分析结果如表 2
图2 Fig. 2 焊接接头的横截面
所示．图 4a 是钎焊界面 A 的微观组织，由于钛母材局部熔化，液态金属在钛板上润湿铺展形成一层约有 140 μm 宽的金属间化合物层．图 4b 和图 4c 分别是靠近钛母材 D 区和靠近焊缝金属 E 区放大的组织形态，结合表 2 可知，紧靠钛母材的 D 区主要由灰色 Ti2 Cu 相 3 和灰白色 TiCu 相 4 组成．靠近焊缝的 E 区主要包括大块状的 AlCu2 Ti 相 5 和白色的铜基固溶体相 6．根据图 5 、图 6 和能谱分析结果，钎焊界面 B， C 和界面 A 的结合机理和微观组织形态基本一致．同时显示界面 A、界面 B 、界面 C 金属间化合物 120 和 180 μm．钎焊界面金属层的厚度分别为 140 ，间化合物层厚度的不同是由于焊接过程中各部位的［9 ］受热和热量传输方式的差异造成的．
［2 ］
1
试验方法
2
2． 1
试验结果与讨论
试验所用的材料为工业纯钛（ TA2 ），紫铜（ T2 ）和直径为 1． 2 mm 的镍铝青铜（ EＲCuNiAl ）焊丝．钛
收稿日期： 2013 － 07 － 08 基金项目：国家自然科学基金资助项目（ 51265028 ）
焊接接头的表面成形和宏观组织特征图 1 为送丝速度 4． 5 m / min，焊丝偏铜量为 2 mm 时的焊缝表面成形．可以发现，焊缝成形良好，

钛合金tc11动态拉伸力学行为的实验研究

钛合金TC11动态拉伸力学行为的实验研究张军, 汪洋（中国科学技术大学近代力学系中科院材料力学行为和设计重点实验室安徽合肥 230027）摘要：利用MTS809材料试验机和旋转盘式间接杆杆型冲击拉伸实验装置，对双态组织两相钛合金TC11进行了应变率为0.001 s-1的准静态和190s-1的动态单向拉伸实验，获得了TC11等温和绝热拉伸应力-应变曲线；实施了应变率为190s-1的冲击拉伸复元实验，获得了TC11在高应变率下的等温应力-应变曲线。

试验结果表明，TC11的拉伸力学行为具有明显的应变硬化效应、应变率强化效应和绝热温升软化效应。

采用修正的Johnson-Cook模型较好地表征了TC11在试验应变率范围内的拉伸力学行为。

关键词：两相钛合金；动态拉伸；绝热温升软化；复元试验EXPERIMENTAL INVESTIGATION ON THE DYNAMIC TENSION BEHA VIOR OFTITANIUM ALLOY TC11Jun Zhang, Yang Wang(Department of Modern Mechanics, CAS Key Laboratory of Mechanical Behavior and Design of Materials, University of Science and Technology of China, Hefei, Anhui 230027, PR China)Abstract:Quasi-static and dynamic uniaxial tension tests for a titanium alloy TC11 with a duplex microstructure were performed using MTS809 testing system and rotating disk bar-bar tensile impact apparatus, respectively. The isothermal stress-strain curve at 0.001s-1 and the adiabatic stress-strain curve at 190s-1 were obtained. The dynamic tensile recovery test was carried out at the rate of 190s-1 and the isothermal stress-strain curve at the high strain rate was obtained. The experimental results indicate that there exists the strain hardening, strain-rate strengthening and adiabatic temperature-rise softening phenomenon in the tension behavior of TC11. A modified Johnson-Cook model was chosen to describe the tensile behavior of TC11 at different strain rates. The model results agree well with the experimental data.Keywords: Two phase titanium; Dynamic tension; Thermal softening; Recovery test0. 引言两相钛合金具有比强度高、高低温性能优异、耐腐蚀等优点，是航空、航天工程中广泛使用的结构材料。

脉冲电流下纳米结构形成

Nanostructure formation in the surface layer of metals under inﬂuence of high-power electric current pulseA.Vinogradov ÆA.Mozgovoi Æzarev ÆS.Gornostai-Polskii ÆR.Okumura ÆS.HashimotoReceived:9April 2009/Accepted:15June 2009/Published online:1July 2009ÓSpringer Science+Business Media,LLC 2009Abstract The possibility to tailor the microstructure of metals is explored utilising a skin-effect for surface treat-ment.The theoretical simulation of the electric and mag-netic ﬁelds in a metallic cylinder shows that melting followed by rapid quenching can occur in a skin layer of 5–10-l m thickness if the amplitude of a single electric pulse of several nanoseconds duration is of the order of hundreds kiloamperes.The experiments using the SUS304stainless steel show that besides a thin amorphous layer,a speciﬁc nano-twin structure can form at the near-surface region.The appearance of nano-twins is explained considering the stress components arising at the surface layer and in the bulk of the specimen during shock wave propagation caused by temperature gradients and the Lorentz force.It is shown that the high stress amplitudes can arise locally,furnishing the required conditions for twin nucleation and resulting in intensive plastic deformation of the sub-surface layer.IntroductionMany physical and mechanical properties of materials depend strongly on the surface state.For instance,hard-ness,fatigue life and wear resistance beneﬁt from surfacestrengthening which can be achieved in a variety of tech-niques including severe plastic deformation,plating,ion implantation,laser irradiation,etc.The ultimate properties have been achieved in materials in their nano-and amor-phous states [1].Both nano-and amorphous structures are commonly obtained in metals from melts of very speciﬁc chemical compositions by utilising a variety of rapid quenching techniques where a rather high critical cooling rate of 105–107K/s is achieved as required for amorphi-sation.At lower quenching rates,ordinary crystallisation or formation of ultraﬁne grain or nano-structure is observed with characteristic grain dimensions from few nanometres to few hundred nanometres,respectively.Numerous rapid quenching techniques allow obtaining thin ﬁlms,ribbons,ﬂakes and powders of micrometre dimensions [2].In con-trast,severe plastic deformation is capable of producing bulk articles with the ultraﬁne grain and nano-structure [3].In this study,we explore the possibility to create novel microstructures by combining both rapid quenching and severe plastic deformation.We make use of a skin-effect in conductors,arising from the short electric current pulse ﬂowing through the conductor,for their surface heating up to melting followed by rapid cooling due to heat transfer the bulk.It was supposed that various ultraﬁne grain,nano-and even amorphous structures can be produced in the surface layer of metals in this way [4].The ﬁrst theoretical attempt has been made in ref.[4]to simulate the behaviour of the electric and magnetic ﬁelds in a metallic cylinder subject to a rapid discharge of a bank of capacitors as shown schematically in Fig.1.It was demonstrated that melting followed by rapid quenching can occur in a skin layer of 5–10-l m thickness if the amplitude of the passing electric pulse of several nano-seconds duration is high enough,i.e.of the order of 102kA.Furthermore,it appeared that since solidiﬁcation of the molten layer occurs within a microsecond or sub-microsecond timeA.Vinogradov (&)ÁS.HashimotoOsaka City University,Osaka 558-8585,Japan e-mail:alexei@imat.eng.osaka-cu.ac.jpA.Mozgovoi Ázarev ÁS.Gornostai-Polskii ÁR.Okumura Institute of Experimental Physics,Sarov 607190,Russia Present Address:R.OkumuraMaterials Engineering Department,Denso Corp.,Kariya 448-8661,JapanJ Mater Sci (2009)44:4546–4552DOI 10.1007/s10853-009-3689-zscale,very high quenching rates of108–109K/s are achiev-able,providing an appealing novel opportunity for amorphi-sation or signiﬁcant modiﬁcation and nano-structurisation of the surface layer in conductors.The possibility to obtain thickness of5–10l m amorphous layer in the commercial SUS304stainless steel was then successfully demonstrated by the same authors[5],which agrees with common estimates of the skin layer thickness[6].Theoretical calculations in this study show that the extremely high stresses may arise from the temperature gradients and Lorenz force action resulting in intensive plastic deformation of the sub-surface layer.The experiments under this study utilising an ultra-high power generator producing a short electric current pulseﬂowing through copper or SUS304stainless steel cylindrical speci-mens reveal that a speciﬁc nano-twin structure can be formed in the surface layer and in the bulk of the specimens depending on the electric pulse parameters.The appearance of nano-twins is explained from the theoretical consideration of the stress components arising at the surface vicinity and in the bulk of the specimen.Use of skin-effect for surface heating,melting and rapid quenchingLet us consider an endless conducting cylinder with the symmetry axis Z in a circuit shown in Fig.1.The proce-dure,which allows evaluating the parameters of the external circuit such as the resistance R,capacitance C, inductance L and the initial voltage U0,as well as the dimensions of the metallic cylinder l and r0,is given in[4] to ensure melting and rapid quenching in the surface layer.The initial system of differential equations describing a given problem can be written in a classic Maxwell’s form for the electric E and magnetic Hﬁelds together with a heat transfer equation:rÂH¼1 q E;rÂE¼Àl a o H o t;q m c V o To t¼rÁðk r TÞþ1qj E j2;ð1Þwhere the right-hand side of the last equation includes theJoule heating term proportional to j E j2as a heat source.Here,q m is the materials density,l a is the magneticpermeability,c V is the speciﬁc heat capacity and k is theheat conductivity of the material.Apparently,the system ofequations(1)should be completed with the ordinaryelectric equations for the external circuit given asLd Id tþRIÀUðtÞþlE zðt;r0Þ¼0;d Ud t¼ÀIðtÞC;ð2Þwith initial conditions setting the initial discharge voltageU0Uð0Þ¼U0;Ið0Þ¼0:ð3ÞIn order to model the dynamics of various processes inthe actual experimental installation utilising the electricalenergy of the bank of capacitors discharging on a low-inductance load,the system of equations(2)and(3)wassolved together.The details of the problem formulation andsolution have been reported in[4,5].For instance,for acopper cylinder having the length l=10mm and radiusr0=0.1–0.5mm at T0=300K,C=0.1l F,R=0.3Ohm,L=5nHn and U0=100kV,the peak currentmagnitude reaches180–250kA while the pulse duration isas short as40–50ns.The maximum thickness of the moltenlayer R m was found to be10l m,and the melting time was230ns.Cooling until complete solidiﬁcation occurredwithin approximately1l s.The maximum temperature onthe surface reached almost2,000K,which is between themelting point T m=1,356K and vaporization temperatureT v=2,868K.The maximum quenching rate,q T/q t,was of29108K s-1.This quenching rate is still not enoughfor amorphisation of pure metals;it is,however.within therange of typical quenching rates attained in melt-spinningtechniques used for production of a broad variety of amor-phous alloys[2].Using the model constituted by Eqs.1and3,one can show that the depth of the molten layer varies in arange of0.1–40l m,depending on the specimen dimensionsand initial voltage provided that the parameters of theexperiment have been the realistically chosen[4,5].Hence,utilising a high-power generator described brieﬂy in[7],which was modiﬁed toﬁt the cylindrical specimens of1-mmdiameter,the proposed scheme vas validated using Cu andSUS304stainless steel.It was demonstrated that formationof the amorphous(glassy-like)layer in the intimate prox-imity(from zero to about30l m depending on the initialvoltage U0)to a free surface of the SUS304steel(Fe–18Cr–8Ni)is possible after melting followed by rapid quenchingcaused by the short high power electric pulse.For this study,the most interestingﬁnding reported in ref.[4]is that aspeciﬁc twin structure is formed in the subsurfacelayer,extending up to 100l m inside the material from the surface as illustrated in Fig.2(the details of experimental procedure,specimen preparation and observation can be found in [5]).Transmission electron microscopy (TEM)with the selected area electron diffraction pattern (SAEDP)analysis reveals that this structure differs drastically from that is commonly found in the bulk of the as-received specimens.After the ‘skin’treatment,the structure near the surface consists primarily of regions which are multiply twinned along {111}planes with the nanometre scale twin spacing.The very ﬁne twin lamellae,parallel to each other can be readily found here and there in TEM foils.Apparently,the density of twin population reduces with depth from the surface,c.f.Fig.2a–d.Figure 2b shows that the twins in the bulk of the specimen are limited to the initial grain size which is of the order of 2–3l m.The appearance of such a very ﬁne nano-scaled twin structure,induced by the ‘skin’treatment,is difﬁcult to explain by annealing-type effects caused by the temperature rise at the specimen surface due to the skin electric current.Since the cooling rate has been proven very high [4,5],the time interval over which the surface layer is exposed to the high temperature inﬂuence is too short to facilitate diffusion processes controlling the annealing effects.Besides,the twin structure appearance does not resemble typical annealing twin patterns,for example [8].On the other hand,it is well known that the ﬁne twin lamellae with the nano-scale spacing can be produced mechanically at high stresses due to plastic deformation [9–11].Appearanceof mechanical twins assumes two factors of primary importance:(i)high stresses approaching a sizable fraction of the theoretical cohesive strength s /l =10with l —shear modulus of the material,(ii)impeded alternative accom-modative mechanisms such as dislocation motion.In the course of conventional plastic deformation,the dislocation motion can be blocked by high internal stresses arising from dislocation accumulation during strain hardening.In the experiments under this study,the dislocation motion is limited due to a very short inﬂuence from the electric pulse.The mechanical twins form very fast,with the velocity of sound [10]and can be highly possible candidates for alternative sources of plastic deformation.Therefore,the internal stresses arising from the inﬂuence of electric and magnetic ﬁelds during the ‘‘skin’’treatment of the metal surface should be evaluated to shed light on the feasibility of twin nucleation.This will be accomplished in the next section.Stress calculation Problem statementThe pressure caused by the electromagnetic ﬁeld acting on the specimen in the geometry shown in Fig.1is pro-portional to the Pointing vector,i.e.to the vectorproductFig.2Nano-scaled twinstructure below the surface of a ‘‘skin-’’treated SUS304stainless steel a 25l m below the surface;SEADP is indicative of the twinnedstructure;b 100l m below the surface;c and d 100l m below the surface.The initial discharge voltage is 70kV (a ,b )and 90kV (c ,d ),respectivelyE 9H .The displacement vector U ,will have both axial and radial components U ={U r ,0,U z }.In order to calculate the stresses which arise in the metallic specimen when the high power electric pulse is passing through,the equation for the vector U should be added to the system of Maxwell equations for electric and magnetic vectors.The equation describing the mechanical behaviour of an isotropic deforming media is written as q m o 2Uo t¼r Á½r þF ;ð4Þwhere the Lorentz force term and the linear isotropic expansion term are incorporated in the right hand side;[r ]is the stress tensor,F is the net bulk force vector.In the axially symmetrical case the stress tensor takes a form½r ¼r r 0r rz0r u 0r zr 0r z 0@1A ;ð5Þwith the components given as r rr u r z rrz0B B B @1C C C A ¼k 1Àm m 11011Àm m 10111Àm m00001À2mm0B B @1C C A e r Àa T ðT ÀT 0Þe u Àa T ðT ÀT 0Þe z Àa T ðT ÀT 0Þe rz 0B B @1C C A ð6ÞComponents of the strain tensor [e ]are given as:e r ¼o U r o r ;e u ¼U r r ;e z ¼o U z o z ;e rz ¼12o U r o z þo U z o r:ð7ÞHence,the initial system of equations takes a form:r ÂH ¼1q E ;r ÂE ¼Àl a o H o t;q m co T o t ¼r Áðk r T Þþ1qj E j 2;q m o 2Uo t2¼ðk þ2l Þrðr ÁU ÞÀl r Âðr ÂU ÞÀa T K r T þl a 1ÁE ÂH ;l ¼E u 2ð1þm Þ;k ¼E u m ð1þm Þð1À2m Þ;K ¼E u3ð1À2m Þ:ð8ÞHere,E u ,l and K are the Young’s,shear and bulk moduluses,respectively,m is the Poisson ratio,e ij are the strain tensor components,U r ,U z are the displacement vector components,a T is the linear heat expansion coefﬁcient and k is the Lame constant.The right-hand side of the third equation includes the Joule heating termproportional to j E j 2as a heat source.The system is said tobe thermally isolated because the energy losses due to irradiation have been proven negligible [4].Using Raleigh damping,which is commonly used to provide a source of energy dissipation in analyses of structures responding to dynamic loads and having H ={0,H ,0},E ={E r ,0,E z }and U ={U r ,0,U z },the equations (8)in cylindrical coordinates {r,z }take a ﬁnal form:l a 1q o H u o t ¼1r r o 2H u o r 2þo H u o r ÀH u rþo 2H uo z 2;E z ¼q r o o r rH u ÀÁ;E r ¼Àq o H u o z;q m c o T o t ¼k o 2T o r þo 2T o zþk r o T o r þ1q E 2rþE 2z ÀÁ;q m o 2U r ¼ðk þ2l Þo 2V r þo V rþl o 2V r o z 2þðk þl Þo 2V zo r o z Àl a q E z H u Àa T K o T o r þa R q mo U ro t ;q m o 2U z o t 2¼ðk þ2l Þo 2V z o z 2þl o 2V z o r 2þ1r o V z o rþðk þl Þo 2V r o r o z þ1r o V ro zþl a q E r H u Àa T K o T o z þa R q mo U zo t;V r ¼U r þb R o U r o t ;V z ¼U z þb R o U zo t:ð9Þwhere,U r and U z are the radial and axial components of thedisplacement vector U ,respectively,and a R and b R are the coefﬁcients of Raleigh damping,which are commonly expressed as a R ¼4pnf min f maxmin max;b R ¼nmin max ;ð10Þwhere n is the damping efﬁciency ranging between 0and 1,f min and f max stand for the lower and upper frequency in the working frequency range,respectively.For numeric cal-culations,f min and f max are chosen as 0.5an 90MHz,respectively.Initial and boundary conditions are set as H u ð0;r ;z Þ¼0;U z ð0;r ;z Þ¼0;U r ð0;r ;z Þ¼0;H u ðt ;r 0Þ¼I ðt Þ2p r 0;T ð0;r ;z Þ¼T 0;o T o r r ¼0¼0;o T o r r ¼r 0¼0;o E z o rr ¼0¼0ð11ÞThe free surface state determines boundary conditions for U z and U r in the Lame equation on the surface of the specimen:r r j y ¼r 0¼ð2l þk Þo U r o r þk o U z o z þU rr¼0;r u y ¼r 0¼ð2l þk ÞU r r þko U z o z þo U ro r¼0;r z j y ¼r 0¼ð2l þk Þo U z o z þk o U r o r þU rr¼0:ð12ÞThe boundary conditions on the Z axis in the centre ofthe specimen (r =0)take a form (due to the axial symmetry):U r ðt ;0;z Þ¼0;o U z o rt ;r ¼0;z¼0:ð13ÞThus,together with Eqs.2and 3,the problem is fully deﬁned and can be solved numerically.Numerical estimates and discussionEquations 6,7,9and10were solved together with electric equations (2)and initial and boundary conditions (3,11–13).Numeric solutions were obtained using ﬁnite element method within a mesh constructed in r -z space having 1440elements and 794nodes.The Lagrange polynomial func-tions of the second order were taken for approximations.For illustrative purposes,the calculations of internal stresses caused by high-power short electric pulse passing through a conductive material have been performed for Cu cylinder of r 0=0.5-mm initial radius and 5-mm length (regretfully,for the SUS304steel,only few required material constants are known,which makes it impossible to perform quantitative modelling).All the other conditions were set the same as in [4],i.e.T 0=300K,C =0.1l F,R =0.3ohm,L =5nH and U 0=100kV.Figures 3and 4show the cumulative von Mises stresses arising from the linear thermal expansion and Lorentz force componentsr Von ÀMises ¼ﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃr 2r þr 2h þr 2z Àr r r h Àr r r z Àr h r z þ3r 2rzq ð14Þand its components calculated in cylindrical coordinates(r ,h ,z)for r =0(specimen axis)and r =0.5mm (specimen surface)in the middle of the specimen z =2.5mm.The shock-wave character of the stress tem-poral and spatial distribution is evident,particularly in Fig.5,where the snapshots of the axial component of von Mises stresses distribution over the specimen are shown in subsequent instants of time.The shock wave propagates from the surface to the bulk and back in an oscillating manner with attenuation caused by Raleigh damping introduced in Eqs.9.The duration of wave propagation depends on materials’properties and may vary from sev-eral microseconds to several milliseconds which is quite enough to activate subsonic or even trasonic dislocation motion and mechanical twinning,which can be viewed as a cooperative motion of several partial dislocations.While the r rz component is negligible both at the surface and in the bulk,the radial r r and axial r z and r u components may attain very high peak magnitudes of several gigapascals,which is not unexpected for shock wave loading.The average magnitude of stress components near the surface is notably higher than that in the bulk which accounts for the higher possibility to ﬁnd the twinned domains in the sur-face proximity,i.e.decreasing twin density with the depth from the surface.Gerland et al.[12]used a surface treatment with shock waves caused by explosion in a thin layer applied to a 316L stainless steel.Similar to the results in this study,the near surface structure was found to be composed ofmechanicalFig.3Cumulative von Mises stress on the axis r =0and on the surface r =0.5mm of Cu cylinder subjected to theinﬂuence of the ultrahigh power short electric current pulse.The electric current magnitude I and surface temperature T areplotted for comparison and time scaling.T m —melting temperaturetwins,the density of which decreased with increasing depth.Firrao et al.[13]have also shown that mechanical twins are readily formed in the 304stainless steel subjected to stress impact caused by explosions of small charges in the surface vicinity.The beneﬁts for the mechanical properties of the nanostructured steels produced by explo-sion inﬂuence is argued in [12,13].Ample evidence exists in the literature showing twin formation due to severe plastic deformation caused by intense ultrasonic shot peening [14],mechanical attrition [15]or shock pulse loading [16].The minute mechanisms of twin formation under shock wave loading have yet to be explored,possibly with the help of atomistic molecular dynamic simulations in a way similar to those performed in [17].The results of this study may serve as a starting point for further quan-titative modelling.They suggest that the high local stress conditions required for twin nucleation in the surface layer are fulﬁlled.Indeed,assuming the shear stress required for twin nucleation given as [18]s %c SFE b 1;ð15Þwhere c SFE is the stacking fault energy and b 1is the Burgers vector of the twinning dislocation b 1¼a 6121h i for an f.c.c.crystal with a —lattice parameter,and taking c SFE for the 304Fe–Cr–Ni steel of 17mJ/m 2[19],one estimates s &0.12GPa,which is notably smaller than the possible peak values of stress components in elastic waves arising in the body due to the ‘skin’treatment by a high-power electric pulse.Furthermore,twinning can be plausibly expected in the materials with much higher stacking fault energy under the same treatment conditions.Signiﬁcant changes in the microstructure of skin-treated specimens and the formation of twin colonies in the sub-surface layer suggest that macrocharacteristics of treated metals may alter considerably.The Vickers microhardness HV was measured using a Shmiadzu dynamic Ultra-hard-ness tester DUH-1at 20-g load and 20-s exposure time.Ten to twenty measurements along the specimen gauge part were averaged.The depth of indentation did not exceed 2–3l m.The HV values in the specimens subjected to ‘skin’treatment with the capacitor discharge from 70,75and 80kV are found to be of 200±23HV,161±25HV and 140±27,respectively.All these values are notably higher than that for the reference of the ‘‘as-received’’SUS304sample—124±6HV,which is consistent with formation of the hardened twin nano-structure in thesurfaceFig.4Stress–tensor components (normal and shear)on the axis (a )and on the surface (b )of Cu cylinder subjected to the inﬂuence of the ultrahigh power short electric currentpulseFig.5Snapshots showing the von Mises stress distribution over the cylindrical specimen of 0.5-mm radius and 5-mm length in subsequent instants of timelayer.Interestingly,that discharging capacitors from smal-ler voltage (70kV)results in the highest hardness of the surface layer which agrees with microstructural observa-tions:populations of twins are easier to see in the 70-kV-treated specimens while the density of twins reduces with increasing voltage (this has been noticed earlier in ref.[5]).Furthermore,this is also consistent with numeric evaluation of the temperature proﬁle in the subsurface layer,Fig.6:the thickness of the heated layer is found to be typically of the order of 100l m almost independently of the discharge voltage,which is considerably wider than the skin layer (50–10l m).However,the maximum temperature reached in the proximate vicinity of surface reduces sharply with voltage reduction.Even though this temperature reduces very quickly as we have discussed above (see [5]for details),this can be another argument why the nano-twin structure is more clearly seen after discharge from 70kV.One can notice in Fig.5that maximum von Mises stresses are obtained in the shock-wave within the 100-l m sub-surface layer,being in accordance with TEM obser-vations showing that twin colonies are readily observed in the subsurface region up to about a 100-l m thickness,whereas they are rarely or not at all observed in the bulk of the specimen.Although,the only qualitative comparison between the results of calculations and structural observa-tions are still possible,a fair agreement is seen between the numeric estimates,microstructural observations and microhardness assessment.References1.Gleiter H (1989)Progr Mater Sci 33:2232.Beck H,Guntherodt H-J (eds)(1981)Glassy metals (topics in applied physics),vol 46.Springer,Germany3.Valiev RZ,Islamgaliev RK,Alexandrov IV (2000)Progr Mater Sci 45:1034.Vinogradov A,Lazarev SG,Mozgovoi AL,Hashimoto S (2005)Philos Mag Lett 85:5755.Vinogradov A,Lazarev SG,Mozgovoi AL,Gornostai-Polskii SA,Okumura R,Hashimoto S (2007)J Appl Phys 101:033510–033510-7ndau LD,Lifshitz EM (1984)Electrodynamics of continuous media,vol 8.Pergamon,Oxford,UK7.Voronin VV,Tananakin VA,Pavlov SS,Tsiberev VP,Voronov SL (1977)In:Proceedings 11th IEEE international pulsed power conference,Baltimore,Maryland,USA,p 15668.Jones AR (1981)J Mater Sci 16:1374.doi:10.1007/BF010338549.Nartia N,Takamura J (1992)In:Nabarro FRN (ed)Dislocations in solids,vol 9.Elsevier,Amsterdam10.Rigsbee JM,Benson RB (1977)J Mater Sci 12:406–409.doi:10.1007/BF0056628411.Friedel J (1964)Dislocations.Pergamon Press,Oxford,UK 12.Gerland M,Presles HN,Mendez J,Dufour JP (1993)J Mater Sci28:1551.doi:10.1007/BF0036334813.Firrao D,Matteis P,Scavino G,Ubertalli G,Ienco MG,Pellati G,Piccardo P,Pinasco MR,Stagno E,Montanari R,Tata ME,Brandimarte G,Petralia S (2006)Mat Sci Eng A 424:2314.Liu G,Wang SC,Lou XF,Lu J,Lu K (2001)Scripta Mater44:179115.Zhang HW,Hei ZK,Liu G,Lu J,Lu K (2003)Acta Mater51:187116.Bakalinskaya ND,Zubov VI,Nadezhdin GN,Petrov YuN,Svechnikov VI,Stepanov GV (1988)J Strength Mater 20:120517.Gumbsch P,Gao H (1999)Science 283:96518.Hirth JP,Lotte J (1982)Theory of dislocations.Wiley,New York 19.Murr LE (1975)In:Herndan VA (ed)Interfacial phenomena inmetals and alloys.Techbooks,HerndanFig.6Solutions of Eqs.8and 2showing temperature proﬁles in the subsurface layer of Cu cylinder of 0.5-mm radius and 10-mm length after 0.4l s since the capacitor discharge from different initial voltages.T m —melting temperature。

IN792合金

Microstructural Evolution in Directionally Solidiﬁed Ni-Base Superalloy IN792+HfS.M.Seo1)†,I.S.Kim1),J.H.Lee2),C.Y.Jo1),H.Miyahara3)and K.Ogi4)1)High Temperature Materials Research Group,KIMS,Changwon,641-010,Korea2)Dept.of Metall.&Mater.Eng.,Changwon Univ.,Changwon,641-773,Korea3)Dept.of Mater.Sci.&Eng.,Kyushu Univ.,Fukuoka,819-0395,Japan4)Oita National College of Technology,Oita,870-0152,Japan[Manuscript received September3,2007]Microstructural evolution during directional solidiﬁcation(DS)of Ni-base superalloy IN792+Hf has been in-vestigated with an emphasis on theγ precipitates and M C-type carbides.The quantitative image analyses revealed that the increase in the solidiﬁcation rate up to100µm/s at constant thermal gradient of178K/cm resulted in aﬁne and uniform distribution ofγ precipitates.The relationship between the as-castγ size and cooling rate was also determined for DS IN792+Hf.In the mean time,the M C carbide size was found to be dependent both on the solidiﬁcation rate and the S/L interface morphology while the area fraction of M C carbide was signiﬁcantly inﬂuenced by the S/L interface morphology.KEY WORDS:IN792;M C corbide;Gamma prime;Solidiﬁcation rate1.IntroductionNi-base superalloys are extensively used for tur-bine blades and vanes in aero-and industrial gas tur-bine engines.The mechanical properties of these al-loys depend on grain structure,dendrite arm spacing,γ precipitates,γ/γ eutectic and various types of sec-ondary phase such as carbide and boride[1].IN792+Hf is a Ni-base superalloy that contains about12wt pct Cr with a high Ti/Al ratio in its al-loy chemistry.A high Ti/Al compositional ratio in this alloy,however,promotes the formation of low strength eta(Ni3Ti)phase in preference to the de-sirableγ (Ni3Al)phase due to the strong segrega-tion propensity of Ti during solidiﬁcation.Recently, Seo et al.[2]have shown that microsegregation also re-sulted in the formation of Cr-rich boride phases in the vicinity of eta phase in the as-cast directional solid-iﬁcation(DS)IN792+Hf.In addition,the formation of these phases was found to be deactivated with de-creasing the solidiﬁcation rate because the solid-state diﬀusion at lower solidiﬁcation rates decreased the mi-crosegregation.Even though the eﬀect of solidiﬁca-tion rate on the eta and boride formation in IN792+Hf has been well established,limited information is still available regarding the eﬀect of solidiﬁcation rate on the microstructural features such as M C carbide and γ precipitates that formed during the solidiﬁcation and the subsequent cooling.In the present study,a series of directional solidi-ﬁcation experiments were carried out over a range of solidiﬁcation rates and the inﬂuence of solidiﬁcation rate on the microstructural evolution,especially onγ precipitates and M C carbides,was investigated.2.ExperimentalThe material used in the present study is the Ni-base superalloy IN792+Hf whose chemical comp-†Senior researcher,to whom correspondence should be ad-dressed,E-mail:castme@kims.re.kr.osition is listed in Table1.Specimens of5.0mm in diameter and80mm in length were directionally solid-iﬁed under Ar atmosphere with various solidiﬁcation rates(R),R=0.5–100µm/s and constant thermal gra-dient(G)at the S/L interface,178K/cm.The S/L interface was preserved by quenching the specimens after a desired volume fraction of original liquid was solidiﬁed.The DS specimens for microstructural observation were prepared by the standard metallographic proce-dures and examined by using an optical microscope and a scanning electron microscope(SEM).Comput-erized image analysis was also performed to quantita-tively analyze the size distribution ofγ precipitates and M C carbides.3.Results and Discussion3.1As-cast microstructureDS experiments were carried out with the solidiﬁ-cation rates of R=0.5–100µm/s under constant ther-mal gradient G=178K/cm.For these DS conditions, the S/L interface morphology of the alloy developed from planar(R=0.5µm/s),to cellular(R=1.0µm/s), and to coarse andﬁne dendritic(R≥5.0µm/s)with gradually increasing R.Figure1(a)shows the typical as-cast microstructure of DS IN792+Hf solidiﬁed at R=50µm/s.The as-cast microstructure was charac-terized by the dendrite core and interdendritic region, which was composed of rosette shapedγ/γ eutectic, M C carbide,eta and Cr-rich boride phases(Fig.1(b)). Theﬁneγ particles also precipitated in the entireγmatrix during subsequent cooling after solidiﬁcation.Most of the M C carbides existed near interden-dritic area.Since the major M C forming elements Ti, Ta and Hf exhibited a partitioning tendency to liquid, these elements would be rejected into interdendritic liquid during solidiﬁcation[2].Therefore,as the solid-iﬁcation proceeds,M C carbide forming elements are enriched in interdendritic liquid,which results in the facilitation of nucleation and growth of M C carbide in interdendritic area.Table 1Chemical composition of Ni-base superalloy IN792+Hf (wt pct)Al Co Cr Hf Mo Ta Ti W C B Zr Ni 3.478.712.10.891.84.23.984.30.0720.0160.03Bal.Fig.1Typical as-cast microstructure of DS IN792+Hf solidiﬁed at R =50m/s:(a)optical micrograph and (b)de-tailed SEM micrograph near interdendriticregionFig.2Solidiﬁcation paths of IN792+Hf predicted byThermo-Calc equilibrium and Scheil modelThe eta and boride phases always appeared in front of the coarse γ/γ eutectic as shown in Fig.1(b).Considering that the solidiﬁcation of γ/γ eutectic proceeds toward the coarse γ [3],eta and boride phases expected to be developed from the residual liquid just after the completion of γ/γ eutectic reac-tion.3.2Solidiﬁcation pathIn order to examine the solidiﬁcation sequence of IN792+Hf during solidiﬁcation,thermodynamic cal-culations were performed using Thermo-Calc soft-ware with Ni-Data developed by Thermo Tech Ltd.(UK).Figure 2presents the solidiﬁcation paths of IN792+Hf alloy calculated by Thermo-Calc equilib-rium and Scheil model.In the equilibrium model,solidiﬁcation products were γphase,M C and small amount of M 3B 2.However,the Scheil calculation accounting for the non-equilibrium solidiﬁcation fea-tures predicted the solidiﬁcation of γ/γ eutectic and eta phase in addition to primary γ,M C and M 3B 2.The predicted solidiﬁcation path by Scheil model is asfollows:liquid (L)→primary γ(1608K)→M C car-bide (1596K)→γ/γ eutectic (1467K)→M 3B 2boride (1448K)→eta phase (1261K).Both equilibrium and Scheil model predicted that M C carbide formed at 1596K,which is about 12K lower than the crystallization temperature of primary γphase.This result is comparable to the microstruc-tural observation of Sun et al.[4]who reported that M C carbide forms just below the liquidus tempera-ture of IN792+Hf.In addition,the solidiﬁcation se-quence of eta and boride phases predicted by Scheil calculation corresponds to the microstructural obser-vation result.However,the Scheil model predicted a very low crystallization temperature of eta phase (1261K)compared with the reported value of about 1402K [2].This discrepancy might be caused by the Hf solubility in eta phase.The Thermo-Calc Scheil model predicted little solubility of Hf in eta phase while the experimental result reported by Seo et al.[2]clearly showed that more than 12wt pct of Hf is dis-solved in eta phase.Therefore,considering that Hf is one of the major elements comprising the eta phase,the Scheil calculation with little Hf solubility in eta phase might delay the eta phase formation to lower temperature.3.3γ precipitatesThe γ precipitates are the primary strengthen-ing phase for Ni-base superalloys.A ﬁne and uni-formly distributed γ size results in desirable mechan-ical properties.The relationship between the mor-phology/size of γ precipitates and the solidiﬁcation rate is presented in Fig.3,where the microstructure was observed at the similar position of the DS sam-ple,such as the solidiﬁcation fraction (f s )is about 0.15.The γ precipitates were very large,and showed an irregular and split shape when the solidiﬁcation rate is very low (Fig.3(a)and (b)).However,the γ particles became obviously ﬁne and their morphology developed from irregular to cuboidal with increasing the solidiﬁcation rate.From the SEM micrographs shown in Fig.3,the size distribution of γ precipitates and their averageFig.3Eﬀect of solidiﬁcation rate on the morphology and size ofγ particles in the dendrite core region:(a)R=0.5µm/s,(b)1.0µm/s,(c)5.0µm/s,(d)10µm/s,(e)25µm/s and(f)50µm/sFig.4Size distribution of particles in the dendrite core(a)–(f)and their average size as a function of cooling rate (G·R)Fig.5M C carbide morphology developed from various solidiﬁcation rates:(a)R=1.0µm/s,(b)5.0µm/s,(c)10µm/s,(d)25µm/s,(e)50µm/s and(f)100µm/sFig.6Eﬀect of the solidiﬁcation rate (and the S/L in-terface morphology)on the average size and area fraction of M C carbidesize were determined and were summarized in Fig.4.In the case of lower solidiﬁcation rates,the γ particle size distributed over a wide range (from about 0.5µm to over 1.5µm for R =0.5µm/s)while it showed uni-form size distribution as the solidiﬁcation rate gradu-ally increased.Figure 4(g)shows the eﬀect of cooling rate (G ·R )on the average γ sizes in comparison with the reported values [5,6].The average γ particle size decreased with increasing the cooling rate in the double logarithmic plot.The linear regression on the basis of the data was derived as follows:d γ =0.33(G ·R )−0.334(1)where d γ is the average size of the γ precipitates.Although the as-cast γ size linear-logarithmically de-creased with increasing the cooling rate in the present study,signiﬁcant diﬀerences still remained compared with the previous studies (Fig.4(g)).This result in-dicates that the as-cast γ size might be an alloy de-pendent,i.e .chemical composition,total amount of γ forming elements and segregation behavior of al-loying elements may inﬂuence on the as-cast γ size in Ni-base superalloys.In addition,the volume fraction of γ precipitates appears obviously low at R =0.5,1.0and 5.0µm/s,where the interface morphologies are planar,cellu-lar and coarse dendritic,respectively (Fig.3(a)–(c)).This result is expected to be related with the interface morphology.Macro-segregation occurs in the planar and cellular interface formed at relatively low solidi-ﬁcation rates of 0.5and 1.0µm/s.This appears to occur some in the coarse dendritic interface morphol-ogy at 5.0µm/s in the presence of convection which forms inevitably in the Bridgman type directional solidiﬁcation [9].The γ forming elements (Al,Ti,Ta,Hf)must be lack at the low solidiﬁcation fraction of DS samples (f s =0.15)due to macro-segregation.Macro-segregation due to the interface morphologies,such as the planar,cellular,and coarse dendritic in-terfaces,is expected to change the volume fraction of γ in the γmatrix.3.4MC carbideM C-type carbide that formed during directional solidiﬁcation of Ni-base superalloy strengthens longi-tudinal grain boundaries at elevated temperatures.Italso has a signiﬁcant eﬀect on the solidiﬁcation behav-ior of Ni-base superalloys [7,8].Figure 5shows the mor-phology evolution of M C carbide during directional solidiﬁcation of IN792+Hf under various solidiﬁcation rates.In the lower solidiﬁcation rates of R =1.0and 5.0µm/s,the morphology of MC carbide exhibited a faceted blocky shape (Fig.5(a)–(b)).As the solidiﬁ-cation rate increased,small script type M C carbides started to form together with large blocky shaped M C carbides (Fig.5(c)),and ﬁnally most of the M C car-bide morphologies changed to dendritic script type when the solidiﬁcation rate is higher than 25µm/s (Fig.5(d)–(f)).Figure 6shows the variation of average M C car-bide size and area fraction as a function of solidiﬁca-tion rate.The average M C carbide size was found to be dependent on the S/L interface morphology as well as the solidiﬁcation rate.As the S/L inter-face morphology changes from cellular to dendritic (R =1.0µm/s to 5.0µm/s),the average M C carbide size slightly increased.However,in dendritic solid-iﬁcation conditions (R ≥5.0µm/s),the M C carbide size rapidly decreased with increasing the solidiﬁca-tion rate at slower rates of R =5.0–25µm/s,and this tendency became sluggish at relatively high solidiﬁca-tion rate range (R ≥25µm/s).The steep decrease in M C carbide size in the solidiﬁcation range of R =5.0–25µm/s,appears to be due to the evolution of script type M C carbide and the coarse inter-dendritic spac-ing may provide the larger growth of M C carbide.The area fraction of M C carbide appears to be rather dependent on the S/L interface morphology than the solidiﬁcation rate.The area fraction of M C carbide increased when the S/L interface changes from cellular to dendritic morphology.However,the increase in the solidiﬁcation rate did not have a signiﬁcant eﬀect on the fraction of M C carbide,about 0.76%,where the S/L interface was dendritic morphology.The area fraction of M C carbide at 1.0µm/s,showing the cellular interface,is expected to be low due to the macro-segregation of carbide form-ing elements (Ti,Ta and Hf)at the low solidiﬁcation fraction of f s =0.15.The lower area fraction of γ ,which contains Ti,Ta and Al shown in Fig.3(a),may also reveal the reducing of these elements due to the micro-and macro-segregation.4.Conclusions(1)The solidiﬁcation path of IN792+Hf alloy pre-dicted by Thermo-Calc Scheil model is as follows:L →primary γ→M C →γ/γ eutectic →M 3B 2→eta phase.(2)The increase in the solidiﬁcation rate up to 100µm/s at a constant thermal gradient of 178K/cm resulted in a ﬁne and uniform distribution of γ pre-cipitates within supersaturated γmatrix.The as-cast γ size appeared to be alloy dependent and a following relationship between the γ size and cooling rate was established for DS IN792+Hf:d γ =0.33(G ·R )−0.334(3)The M C carbide size was found to be depen-dent on the S/L interface morphology as well as thesolidiﬁcation rate while the area faction of M C car-bide was strongly related to the S/L interface mor-phology during directional solidiﬁcation.The average M C size decreased with increasing the solidiﬁcation rate,but its area fraction was nearly constant where the S/L interface exhibits dendritic morphology. AcknowledgementThis work was supported by the National Research Laboratory Project of Korean Ministry of Science and Technology.The authors also acknowledge theﬁnancial support of Japan Society for the Promotion of Science (JSPS)through the RONPAKU fellowship.REFERENCES[1]C.T.Sims,N.S.Stoloﬀand W.C.Hagel:Superalloys II,Wiley,New York,NY,1986,97.[2]S.M.Seo,I.S.Kim,J.H.Lee,C.Y.Jo,H.Miyahara andK.Ogi:Metall.Mater.Trans.A,2007,38A,883. [3]Y.Zhu,S.Zhang,L.Xu,J.Bi,Z.Q.Hu and C.X.Shi:Su-peralloys1988,eds.D.N.Duhl,TMS,Warrendale,PA, 1988,703.[4]W.R.Sun,J.H.Lee,S.M.Seo,S.J.Choe and Z.Q.Hu:Mater.Sci.Eng.A,1999,271,143.[5]K.O.Yu,J.J.Nichols and M.Robinson:J.Metals,1992,21.[6]X.Guo,H.Fu and J.Sun:Metall.Mater.Trans.A,1997,28A,997.[7]S.Tin and T.M.Pollock:Meter.Sci.Eng.A,2003,A348,111.[8]S.Tin,T.M.Pollock and W.Murphy:Metall.Mater.Trans.A,2001,32A,1743.[9]J.H.Lee,S.Liu and R.Trivedi:Metall.Mater.Trans.A,2005,36A,3111.。

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Microstructure and mechanical properties of twinned Al 0.5CrFeNiCo 0.3C 0.2high entropy alloy processed by mechanical alloying and spark plasmasinteringSicong Fang,Weiping Chen,Zhiqiang Fu ⇑School of Mechanical and Automotive Engineering,South China University of Technology,Guangzhou,Guangdong 510640,Chinaa r t i c l e i n f o Article history:Received 9July 2013Accepted 29August 2013Available online 7September 2013Keywords:High entropy alloys Mechanical alloying Spark plasma sintering Nanoscale twins Microstructurea b s t r a c tMost of multi-component high entropy alloys (HEAs)only consist of metallic elements.In the present paper,by introducing nonmetallic carbon element,non-equiatomic Al 0.5CrFeNiCo 0.3C 0.2HEA has been successfully prepared by mechanical alloying (MA)and spark plasma sintering (SPS)process.Alloying behavior,microstructure,phase evolution and mechanical properties of the alloy were investigated systematically.During the MA process,a supersaturated solid solution with both face-center cubic (FCC)and body-center cubic (BCC)structures was formed within 38h of milling.However,a major FCC phase,a BCC phase,Cr 23C 6carbide and an ordered BCC phase were observed after SPS.The FCC phase is enriched in Fe–Ni,the BCC phase is enriched in Ni–Al and the ordered BCC phase is especially enriched in Al,respectively.In addition,nanoscale deformation twins obviously presented only in partial FCC phase after SPS.The compressive strength and Vickers hardness of Al 0.5CrFeNiCo 0.3C 0.2high entropy alloy are 2131MPa and 617±25HV,respectively.Ó2013Elsevier Ltd.All rights reserved.1.IntroductionTraditional alloys are typically composed of one principal ele-ment that occupies at least 50at.%in the composition,with minor additions of other elements to obtain deﬁnite microstructure and properties,such as Fe-,Al-,Cu-and Mg-based alloys [1].However,high entropy alloys (HEAs)that proposed by Yeh et al.in 2004have broken this conventional paradigm [2,3].This kind of alloys are de-ﬁned as alloys which consist of at least ﬁve principal elements,and the concentration of each constituent element ranges from 5to 35at.%.Previous researches show that HEAs can be processed to form simple solid solution structures instead of intermetallics and other complicated compounds [4–6].This phenomenon is commonly attributed to the high conﬁgurational entropy in the solid solution state of HEAs [2,7].Furthermore,HEAs have also exhibited interesting properties such as high hardness and high strength [8,9],good thermal stability [10],outstanding wear and oxidation resistance [3,11],which offer great potential for engi-neering applications.The HEA systems explored in the past decade show that metal-lic elements are the most commonly used,e.g.Al,Cr,Fe,Co,Ni,Cu,Ti,etc.[12–14].It is known that the proper addition of nonmetallic elements like C in some traditional alloys is favorable to theirstructures and properties [15,16].However,to the best of our knowledge,HEAs with addition of C element have been rarely investigated and reported.In contrast,AlCrFeNiCo HEA system pre-pared by arc-melting has been extensively studied in existing liter-atures [5,17,18].Wang et al.[19]have investigated microstructure and mechanical properties of Al x CrFeNiCo (06x 62)HEAs,ﬁnding that the as-cast Al x CrFeNiCo alloys can possess FCC and/or BCC structure(s)depending on the aluminum content.Increasing con-centration of Al in this alloy system can lead to the formation of BCC phase,which possess high strength and high hardness while inferior plasticity.Apparently,the concentration of Al element should be in an appropriate range to achieve optimum properties.Hence,considering all the factors discussed above,Al 0.5CrFeNiC-o 0.3C 0.2HEA was designed.Moreover,the most widely studied processing route for HEAs is arc-melting (casting),and only a few reports deal with mechanical alloying (MA)[20–23].As a widely used solid state processing route,MA can easily fabricate nanocrystalline materials with good homogeneity from elemental powders,thus MA can be an ideal way to prepare HEAs [20,24,25].In addition,SPS can rapidly con-solidate alloy powders to high density by applying pressure and passing electric pulse current within short soaking time [26].HEAs synthesized by MA and SPS have been reported to possess good densiﬁcation characteristics,as well as high strength and high hardness [27–29].Hence,MA and SPS were combined to prepare Al 0.5CrFeNiCo 0.3C 0.2HEA.Alloying behavior and phase evolution0261-3069/$-see front matter Ó2013Elsevier Ltd.All rights reserved./10.1016/j.matdes.2013.08.099Corresponding author.Tel.:+862087113832;fax:+862087112111.E-mail addresses:kopyhit@ ,fu.zhiqiang@ (Z.Fu).during milling and consolidation processesied.Microstructure and mechanical properties ofwere investigated.2.Experiment proceduresHEA with a nominal composition of Al0.5pressed in molar ratio)was prepared by drythen by wet milling for4h in ethanol.Al,Cr,Fe,elemental powders with purity higher than99.9size of645l m(325mesh)were mechanicallywas carried out in a high energy planetaryPlanetary Ball Mill)at300rpm with a ball toof10:1under argon atmosphere.Highvials and tungsten carbide balls(10mm inthe milling media.In order to conﬁrm the alloymilling,powder samples were taken out after27,38,42h respectively.The42h ball milledsubsequently sintered by Dr.Sinter Model SPS-825Spark PlasmaSintering System(Sumitomo Coal Mining Co.Ltd.,Japan)at 1273K for8min at a pressure of30MPa under vacuum(residual cell pressure<8Pa).The samples were heated to873K within 4min,while from873K to1173K and from1173K to1273K, heating rates of75K minÀ1and50K minÀ1were used, respectively.The milled powders and the bulk alloy after sintering were ana-lyzed by a Bruker D8ADVANCE X-ray diffractometer(XRD)with a Cu K a radiation.The microstructure of the alloy was revealed by etching in aqua regia and observed using scanning electron micros-copy(SEM,Zeiss Supra40,Carl Zeiss NTS GmbH,Germany).Thin-foil specimens were prepared by mechanical thinning followed by ion milling at room temperature and were analyzed by a transmis-sion electron microscopy(200kV TEM,JEM-2100,JEOL,Japan) with selected area electron diffraction(SAED)analysis.According to GB/T7314-2005[30],the room-temperature compressive prop-erties of the cylindrical samples(U3mmÂ4.5mm in size)were measured with an Instron5500testing system at a strain rate of 1Â10À3sÀ1.Three compression tests were performed to obtain average value.Hardness measurement was conducted using a Dig-ital Micro Hardness Tester HVS-1000Vickers hardness instrument under a load of300gf.The reported hardness value is an average of at least10measurements.3.Results and discussion3.1.XRD analysis and microstructure during MAFig.1shows the XRD patterns of Al0.5CrFeNiCo0.3C0.2HEA pow-ders prepared under different milling durations.It can be seen that diffraction patterns of all alloy elements are included in the initial blending powder.As the milling time increases,drastic decrement of diffraction intensity is observed after6h of milling.The peaks of Al,Co and C elements dissolve most rapidly.According to Chen et al.[31],the alloying sequence for the Cu0.5NiAlCoCrFeTiMo alloy system correlates best with the melting point of the component elements.Thus rapid dissolution of both Al and Co could be asso-ciated with their lower melting point than other elements.In con-trast,the early disappearance of diffraction peak of C may result from its smallest atomic fraction in the alloy.Many diffraction peaks can hardly be seen when the milling time reaches up to 27h.The complete disappearance of all the elemental peaks and the formation of solid solution are founded within38h of milling. Predominant peaks corresponding to an face-centered cubic(FCC) and a body-centered cubic(BCC)crystal structures are respectively observed.Subsequently,the powder was subjected to wet ball milling in ethanol for4h with the aim to obtainﬁne metallic pow-der for being conductive to sintering.As the milling time extends to42h,the diffraction peaks exhibit no change except for a minor broadening and a signiﬁcant increment of diffraction intensity.The crystal size and lattice strain of the BCC and FCC phases with different milling time have been calculated from the X-ray peak broadening using Scherrer’s formula after deducting the instrumental contribution.The calculated results are listed in Table1.The crystalline sizes of both BCC and FCC phases after 42h of milling are slightly reﬁned compared with38h of milling, while the lattice strain of these two phases increase as the milling time prolongs.The crystallite reﬁnement can be attributed to the circulation of crushment and agglomeration during the MA process.Reasons for the increment of lattice strain include size mismatch effect between the constituent elements,increasing grain boundary fraction and high dislocation density imparted by MA[21].The crystallite reﬁnement and high lattice strain may account for the above-mentioned intensity increment and peak broadening in the diffraction.The SEM images of the Al0.5CrFeNiCo0.3C0.2HEA powders of dif-ferent milling time are shown in Fig.2.The primitive powder shows a granular size of less than40l m.In the early period of MA(as shown in Fig.2(b)),the particles cold weld together to form even larger particles than that of primitive powder.Subsequently, when the milling time reaches15h(Fig.2(c)),most of the cold welded agglomerations are crushed down to smaller particles. The27h milled alloy powder reveals an average particle size of approximately5l m as shown in Fig.2(d),and the particles cold weld again when the milling time reaches38h(Fig.2(e)).This cir-culation of crushing and cold welding induced by the ball mill gradually reduces the crystalline size and facilitates the diffusion and alloying among different elements.The38h milled alloy pow-der is subsequently wet ball milled for4h with alcohol as the mill-ing media.It can be seen from Fig.2(f)that the elliptoid particles are fractured,exhibiting a lamellar structure.The particle size for theﬁnal powder is much smaller than that of38h milled powder.1.XRD patterns of Al0.5CrFeNiCo0.3C0.2HEA powders with different milling time.Table1The crystalline size and lattice strain of Al0.5CrFeNiCo0.3C0.2HEA during MA.Milling time(h)Crystalline size(nm)Lattice strain(%)BCC FCC BCC FCC 3812130.710.734211120.760.77974S.3.2.Phase evolution and microstructure after SPSFig.3illustrates the XRD pattern of Al0.5CrFeNiCo 0.3C 0.2HEA detailed analysis of the XRD pattern suggests a BCC phase,a Cr 23C 6carbide phase and an visible.Except the ordered BCC phase,the other three phases are calculated to be (BCC),10.652Å(Cr 23C 6),respectively.After the FCC phase reduces from 0.77%after MA phase from 0.76%to 0.27%,which conﬁrmsthe nearly annihilation of defects introduced during MA after sin-tering.As hereinbefore mentioned,the main phases of Al 0.5CrFe-NiCo 0.3C 0.2HEA powder after 42h of milling are the BCC and FCC solid-solution phases,demonstrating that densiﬁcation at 1273K for 8min has resulted in phase evolution.The MA process could in-duce large strain and defects which might lead to the extension of solubility,thus the milled powders are generally in a non-equilib-rium state.It is certain that a reordering process happened during SPS,leading to the metastable supersaturated solid solutions of MA to more stable phases.This reordering can be correlated with the above-mentioned annihilation of defects introduced by severe plastic deformation during MA.Furthermore,the huge pulsed elec-tric current during the SPS process can also facilitate the phase evolution.Fig.4shows the SEM micrographs of bulk Al 0.5CrFeNiCo 0.3C 0.2HEA consolidated by SPS.Few porosities can be seen in low magni-ﬁed image (Fig.4(a)).Actually,relative density of the sintered Al 0.5-CrFeNiCo 0.3C 0.2HEA sample,which is calculated with respect to theoretical density,reaches up to 99.6%.Two distinctive areas are visible in Fig.4(a),viz.,irregular bulk areas and irregular pot hole areas.The irregular bulk areas are most likely to be the main phase because they account for a higher volume fraction.While the pot hole areas can be consist of phases which are removed by corrosive agent.High magniﬁed image (Fig.4(b))shows that some nanoscale black spots are dispersed in irregular pot hole areas.Most of these spots might be numerous nanoscale phases,and a small minority of spots might be ultraﬁne porosities.In summary,bulk Al 0.5CrFe-NiCo 0.3C 0.2HEA might consist of at least three types of phases.The TEM bright ﬁeld image and corresponding selected area electron diffraction (SAED)patterns of Al 0.5CrFeNiCo 0.3C 0.2HEAAl 0.5CrFeNiCo 0.3C 0.2HEA powders with different milling time.(a)0h,(b)6h,(c)15h,(d)27h,(e)38h and (f)Fig.3.XRD pattern of Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.after SPS are shown in Fig.5.In order to conﬁrm the phase compo-sition of Al 0.5CrFeNiCo 0.3C 0.2HEA,EDS/TEM method was used to measure chemical compositions (in at.%)of the phases in Fig.5(a)(regions marked A–Y).Results of the chemical composi-tion analysis are listed in Table 2.It is noticed that all the regions can be classiﬁed as four different phases.Crystal structures of these phases are measured by their corresponding SAED patterns as shown from Fig.5(b–e).As a result,these four phases are measured to be an FCC phase,a BCC phase,Cr 23C 6carbide and an ordered BCC phase,respectively.It is worth mentioning that diffraction pattern of region C exhibits a FCC twin structure of the FCC phase.Clear microstructures and corresponding SAED patterns of twins will be illustrated hereinafter.As shown in Table 2,the FCC phase isdepleted in Al,Cr compared with the nominal CrFeNiCo 0.3C 0.2HEA.The BCC phase is Ni-rich,Fe-depleted and C-depleted,while the or-greatly rich in Al.The lattice parameters of measured by SAED are 3.700Å(FCC),Å(Cr 23C 6)and 4.320Å(Ordered BCC),Cr 23C 6phase also presents Fe,Ni,Co and Al elements,especially Fe shows a high concentration (12.7±0.4at.%).In conclusion,the bulk Al 0.5CrFeNiCo 0.3C 0.2HEA exhibits a Fe–Ni-rich FCC phase,a Ni ÀAl-rich BCC phase,a Cr 23C 6carbide phase and a Al-rich ordered BCC phase,which is in accordance with the XRD result (Fig.3).It is noticeable that the grain size of these phases shows a wide distri-bution ranging from several hundred nanometers to 1l m.How-ever,the Cr 23C 6and the ordered BCC phases are much ﬁner than the FCC and the BCC phases.It is a reasonable explanation that the Cr 23C 6and the ordered BCC phases are formed in a phase evo-lution and reordering process during SPS.The formation of these four phases is complicated.According to the Gibbs phase rules,the number of equilibrium phases (p )is p =n +1for the alloy that contains n elements.Since phase formation not in equilibrium conditions,the number p >n +1.However,the Al 0.5CrFeNiCo 0.3C 0.2four types of phases,showing much fewer Yeh et al.[2]has proposed that it is possible solutions when HEAs contain ﬁve or more 4.SEM micrographs of Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.(a)low-magniﬁed image,and (b)high-magniﬁed image.corresponding SAED patterns of bulk Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.(a)bright ﬁeld image,(b)SAED pattern of FCC [011]1]zone axis (Region D),(d)SAED pattern of Cr 23C 6[112]zone axis (Region B),and (e)SAED pattern of ordered BCC [012]Table 2Chemical composition (in at.%)analysis results of the phases by EDS/TEM.RegionsPhases Cr Fe Ni Al Co C A,C,F,G,I,J,K,M,N,O,P,T,V,W,X,Y FCC 21.8±0.931.8±1.229±1.1 4.4±0.68.5±0.8 4.5±0.5D,E,Q,R BCC 19.7±1.218.6±2.032.8±1.217.5±1.87.2±0.5 4.2±0.4B,L Cr 23C 663.5±2.112.7±0.4 2.9±0.4 1.8±0.3 2.6±0.316.5±1.6H,S,UOrdered BCC 1.6±0.5 2.4±0.42±0.486.8±1.7 1.2±0.46±1.0Nominal composition–25252512.57.55are randomly distributed in the crystal lattice,though the alloy exhibits high mixing entropy.It reveals that high mixing entropy is insufﬁcient to dominate the formation of phases in HEAs sys-tems.Zhang et al.[32]related the formation of simple solid solu-tion to the mixing enthalpy (D H mix )and atomic size difference (d ).Subsequently,Yang and Zhang [33]proposed a solid-solution formation rule for multi-component HEAs based on the calculation of most of reported HEAs.According to Yang et al.,two parameters can be used to estimate the phase formation behavior of HEAs:O is deﬁned as a parameter of the entropy of mixing timing the average melting temperature of the elements over the enthalpy of mixing,d is deﬁned as the mean square deviation of the atomic size of ele-ments.It is proposed that HEAs form simple crystal structures when O P 1.1and d 66.6%.These two parameters are deﬁned by Eqs.(1)and (2),respectively.X ¼T m D S mix j D H mix jð1Þd ¼ﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃX n i ¼1c i ð1Àr i = r Þ2q ð2Þwhere c i is the atomic percentage of the i th component, r ¼Pn i ¼1c i r i is the average atomic radius and r i is the atomic radius of the i th component.T m ,D S mix and D H mix are calculated as follows:T m ¼Xn i ¼1c i ðT m Þið3ÞHere,(T m )i is the melting point of the i th component of alloy.D S mix¼ÀR Xn i ¼1ðc i ln c i Þð4Þwhere c i is the mole percent of component,P ni ¼1c i¼1,and R(=8.314JK À1mol À1)is gas constant.D H mix ¼X n i ¼1;i –jX ij c i c j ð5Þwhere X ij ¼4D H mix ijis the regular solution interaction parameter between the i th and j th elements,c i or c j is the atomic percentage of the i th or j th component,and D H mix ijis the enthalpy of mixing of binary liquid alloys.Table 3presents the value of mixing enthal-py of atom-pairs of Al 0.5CrFeNiCo 0.3C 0.2HEA.The calculated results of O and d for the alloy are 1.41and 10.01%,respectively.It is obvi-ous that O well matches while d breaks the solid-solution formation rules for the multi-component high entropy alloys proposed by Yang et al.This phenomenon indicates that the formation of Cr 23C 6carbide and ordered BCC phase is reasonable in this alloy sys-tem and it well conforms to Yang’s research.Firstly,it is obvious that the value of O shows the relative predominance of T m D S mix and D H mix .Since O >1is obtained,T m D S mix should be the predomi-nant part of the free energy.It is known that mixing entropy indi-cates the tendency for the formation of random solution while enthalpy of mixing in any system indicates the tendency for order-ing or clustering,thus the formation of solid-solution phases ought to be much easier than the formation of intermetallic compoundsand other ordered phases in the Al 0.5CrFeNiCo 0.3C 0.2HEA.Obviously,the formation of FCC solid-solution phase as the main phase of Al 0.5-CrFeNiCo 0.3C 0.2HEA is mainly correlated with the alloy’s high mix-ing entropy.Secondly,the relatively large value of d indicates that the atomic size difference between components is too large for this alloy system to form entire solid-solution phases.Because the large atomic size mismatch could lead to serious lattice distortion and subsequently increase the corresponding strain energy which could lower the stability of solid-solution.Actually,the interstitial solubil-ity of carbon element in alloy is quite limited,so that carbon ele-ment has a strong tendency to exist in the form of carbonization or graphite in alloy.Moreover,the diffusion of atoms in the matrix could be suppressed due to large atomic size difference.Thus it facilitates the atomic segregation,even results in the formation of amorphous structures [35].As listed in Table 3,atomic radiuses of C (0.77Å)and Al (1.43Å)are signiﬁcantly different from the other elements.This could be a reasonable explanation for the formation of the Cr 23C 6carbide and the Al-rich ordered BCC phase.In addition,it can also be found that the mixing enthalpy of most of the atomic pairs is highly negative (shown in Table 3).Al and C are the main contributor to the negative enthalpy of mixing in this alloy,indicat-ing their strong tendency for ordering or clustering.It is interesting to note that the mixing enthalpy of C and Cr (À61kJ/mol)is the most negative.This could be another factor which contributes to the formation of the Cr 23C 6carbide in the alloy.In addition,as above-mentioned,Fe has a high concentration (12.7±0.4at.%)in Cr 23C 6phase,revealing that Fe has high solubility in this type Cr 23C 6phase.According to Table 3,the mixing enthalpy of C and Fe (À50kJ/mol)stays a highly negative level,leading to Fe atoms preferring C sites.Mechanically alloyed powders of Al 0.5CrFeNiCo 0.3-C 0.2HEA exhibit simple solid solution structure within 38h of mill-ing,which can be attributed to the formation of supersaturated solid solution.The solid solubility extension can be ascribed to the high mixing entropy effect as well as the non-equilibrium state of the MA process [36].Interestingly,twinned FCC phase is also ob-served.Fig.6shows TEM images and corresponding SAED patterns of the FCC phase with nanoscale twins.It is worth pointing out that nanoscale twins are found only in the FCC phase,which has been conﬁrmed by EDS/TEM and the corresponding diffraction patterns.The lamella thickness of the nanoscale twins shown in Fig.6(a)is less than 60nm.The corresponding SAED pattern (the matrix axisis [011]M and the twin axis is ½0 1 1 T)which is shown in the inset of the upper right of Fig.6(a)indicates the nanoscale twins belong to the FCC phase.Actually,existing researches in deformation twin-ning of nanocrystalline materials are mainly focus on FCC crystal structure metals,which is probably attributed to their favorable capability of deformation twinning [37].Fig.6(b)illustrates a noticeable twin structure of the FCC phase surrounding by grey phases.These grey phases are found to be the Cr 23C 6carbide through EDS/TEM analysis,revealing that the formation of Cr 23C 6might have effect on deformation twinning.Since Cr 23C 6shows a hard brittle texture [38],it could be presumed that partial FCC phase between the Cr 23C 6carbide is not readily deformed and con-solidated.Under the isostatic pressure of 30MPa during the SPS process,certain parallel crystal faces of partial crystals moving opposite to each other along a direction with a certain value of dis-placement distance.Thus twinning in the FCC phase may occur dur-ing the phase evolution and densiﬁcation with the aim to be more stable and attaining complete densiﬁcation [29].Moreover,a cer-tain twin system could be activated by a critical resolved shear stress which can be achieved during crystal deformation and phase evolution process.It is worth pointing out that similar nanoscale twins in CoNiFeAl 0.6Ti 0.4and CrCoNiFeAl 0.6Ti 0.4HEAs prepared by MA and SPS have been observed in previous studies of our research group [28,29].Table 3The chemical mixing enthalpy D H mix ij ;kJ =mol of binary equiatomic alloys calculated by Miedema’s approach [34].Element (atomic sizes,Å)C Ni Cr Co Fe Al C (0.77)–À39À61À42À50À36Ni (1.24)––À70À2À22Cr (1.25)–––À4À1À10Co (1.26)––––À1À19Fe (1.26)–––––À11Al (1.43)––––––S.Fang et al./Materials and Design 54(2014)973–9799773.3.Mechanical propertiesFig.7shows the room-temperature compressive stress–strain curve of bulk Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.The compressive strength and compression ratio of the alloy are 2131MPa and 3.0%,respectively.The average Vickers hardness of Al 0.5CrFeNiC-o 0.3C 0.2HEA after SPS has been measured to be 617±25HV.How-ever,plastic deformation behavior is not characterized in theindicating a limited compression plasticity of mechanical properties of Al 0.5CrFeNiCo 0.3C 0.2HEA HEAs of AlCrFeNiCo HEA system are listed The listed HEAs are prepared by casting 0.5CrFeNiCo 0.3C 0.2HEA studied in this paper.hardness of these as-cast HEAs has beenreported.Obviously,Al 0.5CrFeNiCo 0.3C 0.2alloy exhibits the highest hardness of the HEAs listed in Table 4.The high strength and high hardness of Al 0.5CrFeNiCo 0.3C 0.2HEA is possibly due to the formation of the ordered BCC phase and the Cr 23C 6carbide,the nanocrystalline structure,as well as solid solu-tion strengthening mechanism of Al atoms.It is worth pointing out that nanoscale twins might also have a considerable effect on strengthening the alloy.The strengthening mechanism of nano-scale twins is associated with the complicated interaction between dislocations and twin boundaries.It is believed that twin bound-aries are effective in blocking dislocation motion,especially when the thickness of twin/matrix lamellae decreases down to the nano-scale,and a Hall–Petch-type relationship exists for twin boundary strengthening [39,40].The limited compression ratio may owing to the brittle Cr 23C 6carbide,the ordered BCC phase,solid solution strengthening of Al atoms and ultraﬁne porosities.4.ConclusionsAl 0.5CrFeNiCo 0.3C 0.2high entropy alloy with nanocrystalline has been successfully synthesized by MA and SPS.A supersaturated so-lid solution with both FCC and BCC structures is evidently observed after MA.After SPS,Cr 23C 6carbide and an ordered BCC phase are newly formed.The TEM analysis results conﬁrm that the alloy con-sists of one FCC phase,one BCC phase,Cr 23C 6carbide and one or-dered BCC pared with the nominal composition of Al 0.5CrFeNiCo 0.3C 0.2high entropy alloy,the FCC phase is enriched in Fe–Ni,the BCC phase is enriched in Ni–Al,while the ordered BCC phase is especially enriched in Al.The grain size of these phases shows a wide distribution ranging from several hundred nanometers to 1l m.Nanoscale deformation twins present only 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