TB17钛合金高温压缩变形行为
TB17钛合金等温时效析出行为
TB17钛合金等温时效析出行为王哲;王新南;祝力伟;朱知寿【摘要】Transmission Electron Microscope ( TEM) , X-Ray Diffraction( XRD) and Optical Microscope( OM) were employed to inves-tigate the aging precipitation behavior of a new type of ultra-high strength TB17 titanium alloy. The results show that during heat solu-tion treated in theβphasefield followed by aging the secondaryαphase is nucleated, precipitated and grew on theβphase matrix,and the precipitated phase is lamellar structure which has burgers relation with the matrix. Th e secondaryαphase contentis increased rap-idly and finally reach a steady-state as aging time increased and the final product of aging consists ofαphase andβphase. there is a good linearity relationship between the content of secondary α phase and the hardness of age hardening. The TB17 titanium alloy iso-thermal phase transformation kinetics can be described by JMAK equation.%采用X射线衍射( XRD)、透射电子显微镜( TEM)、光学显微镜( OM)等对新型超高强韧 TB17钛合金次生α相转变动力学进行研究。
高温高压条件下钛合金的变形行为研究
高温高压条件下钛合金的变形行为研究近几十年来，钛合金因其优异的力学性能、耐腐蚀性和耐高温性能而被广泛应用于航空航天、船舶制造、核工程和医疗等领域。
然而，在复杂的工况条件下，如高温高压、大变形速率等情况下，钛合金的力学性能会发生变化，这对于工业应用来说是一个重要的问题。
因此，对钛合金在高温高压条件下的变形行为进行研究，对于掌握其力学性能、优化材料结构和加工工艺，具有重要的实际意义。
高温高压下钛合金的变形行为高温高压是指在高温（800-1200℃）和高压（100-300MPa）的条件下，钛合金在拉伸或压缩过程中发生的变形行为。
研究表明，在高温高压下，钛合金的力学性能发生了明显的改变，包括屈服强度、塑性变形、变形硬化和断裂行为等。
这与高温高压下的材料微观结构和变形机制有着密切的关系。
在高温高压下，钛合金的变形机制主要有晶界滑移、晶内滑移、晶体旋转和相变等。
其中，晶界滑移是最主要的变形机制，这是因为钛合金晶界的热稳定性较差，容易形成滑移位错。
同时，高温高压下，晶内滑移也会逐渐占据主导地位，这是由于高温下晶内扩散速率增加，使得空位浓度增大，导致滑移位错的形成。
除此之外，高温高压下相变现象也容易发生，钛合金的相变严重影响了其力学性能。
高温高压下的钛合金变形行为研究方法要对高温高压下钛合金的变形行为进行研究，需要运用一定的实验方法。
当前，常用的方法包括微观分析和力学测试等。
微观分析是研究钛合金变形行为的一种重要方法，常用的方法包括TEM、SEM、EBSD和XRD等。
其中TEM和SEM主要用于观察钛合金在变形过程中的微观结构和变形机制，可以直观地观察到钛合金中的晶界位错、滑移带和裂纹等。
EBSD可以对钛合金晶体的取向进行测量，揭示晶体的疲劳与变形行为。
XRD常用于研究钛合金的晶体结构和相变。
力学测试是研究钛合金变形行为的另一种重要方法，常用的方法包括拉伸、压缩、扭转和剪切等测试。
这些测试可以测量钛合金在高温高压下的力学性能和变形行为，包括屈服强度、应变硬化率、断裂韧性、动态变形等。
热变形参数对TB15钛合金动态再结晶行为的影响
热变形参数对TB15钛合金动态再结晶行为的影响董显娟;张殿;王宇航;吴轩轩;涂泽立;魏科;黄龙;徐勇【期刊名称】《塑性工程学报》【年(卷),期】2024(31)1【摘要】采用热模拟压缩试验机对TB15钛合金进行变形温度为810~930℃、应变速率为0.001~10 s^(-1)、压下量为60%的等温热压缩试验,研究了热变形参数对TB15钛合金动态再结晶行为的影响。
结果表明:TB15钛合金不同应变速率下的应力-应变曲线呈现出不同特征。
随应变速率增大,流动软化曲线呈现“V”型特征。
当应变速率在0.001~0.1 s^(-1)范围时,随变形温度升高和应变速率减小,动态再结晶(DRX)体积分数和DRX晶粒尺寸增大,且DRX晶粒尺寸增幅较DRX体积分数更加显著。
在低应变速率变形时,原始β晶粒被小角度晶界分成形状规整的多边形亚晶粒,小角度晶界通过连续吸收位错和渐进晶格旋转的方式向大角度晶界转变从而实现连续动态再结晶。
随着应变速率增大,再结晶晶粒形核位置从原始β晶粒内部逐渐向晶界附近转移,在原始β晶界形成均匀细小的“项链状”再结晶晶粒,此时易于由晶界弓弯机制引发不连续动态再结晶。
【总页数】10页(P50-59)【作者】董显娟;张殿;王宇航;吴轩轩;涂泽立;魏科;黄龙;徐勇【作者单位】南昌航空大学航空制造工程学院;江西星火军工工业有限公司;南昌航空大学通航(民航)学院【正文语种】中文【中图分类】TG142.14【相关文献】1.TC11钛合金β相区热变形动态再结晶过程的研究2.钛合金热变形时的动态回复和再结晶3.钛合金热变形时的动态回复和再结晶4.应变速率对TA15钛合金β热变形动态再结晶组织的影响5.热变形参数对含铌钢奥氏体未再结晶区变形相变的影响因版权原因，仅展示原文概要，查看原文内容请购买。
tb17钛合金室温变形及时效析出行为
Volume29Number10The Chinese Journal of Nonferrous Metals October2019 DOI：10.19476/j.ysxb.1004.0609.2019.10.11TB17钛合金室温变形及时效析出行为刘洪骁1，董洪波1，王喆鑫1，蔡增1，朱知寿2(1.南昌航空大学航空制造工程学院，南昌330063；2.北京航空材料研究院先进钛合金航空科技重点研究院，北京100095)摘要：针对TB17钛合金，研究不同工艺参数下的室温变形及变形后的时效析出行为。
通过组织观察及XRD分析，在应变量为0.22以下的室温变形状态以及经室温变形后在(350℃，0.5h)的低温时效过程中均发生马氏体相变；变形过程中产生的形变诱导马氏体及低温时效过程中产生的等温马氏体都属于板条型马氏体，且均在晶界附近析出；不同点在于等温马氏体呈平行规则分布，而形变诱导马氏体呈现不规则分布；在500℃长时间时效的条件下马氏体相全部分解，细小的α相沿着变形带结构析出，形成大量交错分布的条状α相。
在分步变形实验中，等温马氏体在二次变形中起到了降低变形抗力与减小弹性模量的作用。
关键词：形变诱导马氏体；等温马氏体；低温时效；弹性模量文章编号：1004-0609(2019)-10-2306-06中图分类号：TG166.5文献标志码：ATB17是我国新一代具有自主知识产权的航空用近β型钛合金，固溶处理状态具有中高强度与高塑性，冷成型性和可焊性好，经时效处理后强度可高达1400 MPa[1−2]，是理想的高强度航空材料。
β类钛合金在高温单相区固溶处理后能获得单一的亚稳β相，经快冷均能保留至室温组织中。
由于β相所属的体心立方比密排六方的α相滑移系要多，因此β类钛合金拥有良好的室温变形性能。
在之前的生产和研究中，时效工艺往往设置在变形完成后进行，其实质是将形变强化与析出强化相结合，这是因为室温变形过程中会产生大量的位错以及晶内剪切带等亚结构，一方面，能为相变提供形核点而细化晶粒[3−4]，另一方面大幅度提高材料的强度[5−8]；析出强化主要通过在时效过程中析出大量细小的α相达到提高强度的效果[9−13]，α相大量的析出造成α/β相界面的增多，阻碍位错运动最终使得材料强度提高，利用低温时效ω相辅助α相形核也只是增强弥散强化效果，但是这样做获得的效果通常只局限在使变形后的材料强度进一步提高，却无法实现材料在变形过程中塑性与韧性等性能的改善。
纯钛在热压缩过程中的变形行为
Deformation behaviour of commercially pure titanium during simple hot compressionZhipeng Zeng *,Yanshu Zhang,Stefan JonssonDepartment of Materials Science and Engineering,Royal Institute of Technology,SE-10044Stockholm,Swedena r t i c l e i n f o Article history:Received 26September 2008Accepted 1December 2008Available online 13December 2008Keywords:Commercially pure titanium Zener–Hollomon parameter High angle grain boundarya b s t r a c tCommercially pure titanium (CP Ti),grade II,is subjected to hot compression at temperatures ranging from 673to 973K with 50K intervals and strain rates of 0.001,0.01,0.1and 1s À1up to 60%height reduc-tion.By analysing work hardening rate vs.flow stress,the deformation behaviour can be divided into three groups,viz.three-stage work hardening,two-stage work hardening and flow softening.By plottingthe data in a T vs.log _ediagram,the present and previous data fall into three distinct domains which can be separated by two distinct values of the Zener–Hollomon parameter.The microstructure after deforma-tion is characterized by optical microscopy and electron back scattered diffraction.The formation off 10 11g twins is related to the Zener–Hollomon parameter.Geometric dynamic recrystallization seems most appropriate when describing the grain refinement process of CP Ti during hot compression.Ó2008Elsevier Ltd.All rights reserved.1.IntroductionCommercially pure titanium (CP Ti)is of great importance in many industrial applications due to its highly attractive properties,such as good deformability at high temperature,low density,high biocompatibility and excellent corrosion resistance [1].It is chem-ically inert and biologically more compatible than Ti–6Al–4V,which is currently the material of choice for most medical im-plants.However,it is hard for CP Ti to strengthen to a level compa-rable to Ti–6Al–4V [2];therefore,developing higher strength CP Ti is an attractive work for medical applications.In order to increase the strength of CP Ti,grain-size refining is an effective approach,and generally,large deformation is used to accomplish this.How-ever,at room temperature,CP Ti has poor formability and limited ductility due to the intrinsic characteristics of the HCP structure with few slip systems being activated.In order to obtain large deformations,it is thus necessary to deform CP Ti at higher tem-peratures where non-basal slip systems can be activated resulting in an improved ductility.On the other hand,higher temperature leads to increased recovery reducing the hardness,if too high tem-peratures or too low strain rates are chosen.It is therefore impor-tant to gain a good insight into the deformation mechanisms and a fundamental understanding of the deformation process during hot working in order to find good processing parameters with respect to deformation temperature and deformation rate.Deformation behaviour of CP Ti during cold and hot working has been studied extensively by many researchers [3–9].Nemat-Nasser et al.[3]investigated the two-and three-stage work hard-ening phenomenon in CP Ti (grade II).They recorded 29stress–strain curves between 77and 1000K and between 10À3and 8000s À1.By re-examination of their data,4of the experimental points were found to fall in the two-stage regime at high temper-atures while all others fell in the three-stage regime.Ray et al.[4]also investigated the hot compression of CP Ti (grade II).They reported 6points between 1023and 1123K and between 1and 10s À1.By re-examination of their data,4of them were found to fall in the flow softening region whereas 2of them were found to fall in the two-stage work hardening region.Tanaka et al.[5]studied the creep behaviour of CP Ti below the yield stress at temperatures of 298–873K.However,their study only focused on low strain rate and the total stain is very limited.Ungár et al.[6]studied the slip-system activity and dislocation density using an X-ray line-broadening technique.They stated that the deformation behaviour of CP titanium changes from a heavily-twinned to an untwinned mode with increasing deformation tem-perature.However,the influence of strain rate was not taken into account.In Refs.[7–9],the researchers have extensively studied the equal channel angular extrusion (ECAE)technique at elevated tem-peratures in order to improve the strength of CP Ti.However,the ECAE process differs from the traditional forging process as it only imposes shear strain on the material.In addition,it is difficult to estimate the strain rate of the process.Previous studies lead to the understanding of various deforma-tion mechanisms operating in CP Ti and are very useful in charac-terizing the flow behaviour of CP Ti.Nevertheless,further investigations have to be done due to the insufficiencies of these studies as mentioned above.0261-3069/$-see front matter Ó2008Elsevier Ltd.All rights reserved.doi:10.1016/j.matdes.2008.12.002*Corresponding author.Tel.:+4687906542;fax:+468203107.E-mail address:zhipengz@kth.se (Z.Zeng).Materials and Design 30(2009)3105–3111Contents lists available at ScienceDirectMaterials and Designj o u r n a l h o m e p a g e :w w w.e l s e v i e r.c o m /l o c a t e /m a t d esThe present work is primarily concerned with the microstruc-ture evolution under various hot compression tests on CP Ti.Elec-tron backscattered diffraction (EBSD)in a scanning electron microscope (SEM)was used to monitor the evolution of the high angle grain boundaries (HAGBs,>15°)and the low angle grain boundaries (LAGBs,2–15°).By plotting the work hardening rate against the flow stress,the flow behaviour of CP Ti was divided into three groups.The deformation behaviour in each group was char-acterized by analysing the microstructure.2.Experimental proceduresThe as-received CP Ti,grade II,with chemical composition (wt.%)of C,0.01;Fe,0.01;H,0.0005;N,0.004;O,0.12;and Ti in balance was supplied by Sandvik Materials Technology,Sandviken,Sweden in form of a thick-walled extruded tube.The material was annealed at 1073K for 1h in argon atmosphere and air cooled in order to obtain new grains and remove residual stress.With this procedure,an equiaxed microstructure with an average grain size of $40l m was obtained which is shown in Fig.1.Cylindrical spec-imens,10mm in diameter and 15mm in height,were machined from the annealed material.The compression axis of the specimens was always parallel to the extrusion direction of the supplied tube.Hot compression was conducted in a Gleeble 1500thermal simula-tor.Prior to hot compression,each specimen was heated to the deformation temperature for three minutes to ensure a homoge-nous temperature distribution throughout the specimen.The deformation temperature was measured by a thermocouple spot welded to the center region of the specimen surface.A tantalum foil with a thickness of 0.05mm was put between the anvil and the specimen in order to reduce the friction and to prevent stick-ing.The deformation strain,temperature and strain rate were automatically controlled and recorded by the Gleeble 1500thermal simulator pression tests were conducted from 673to 973K with 50K intervals varying the strain rates from 0.001to 1s À1.The samples were deformed to a true compression strain of 0.92and immediately cooled in order to retain the deformed microstructures.It should be pointed out that the EBSD image in Fig.1evidently shows that the material consist of 100%a -phase CP Ti and that there is no b -phase present.As both the annealing and deformation temperatures were below the transformation temperature (1157K),the possible influence of b -phase on the deformation behaviour of CP Ti is eliminated.The deformed specimens were cut along their length axis and prepared by mechanical grinding using grit papers with different particle sizes from 320to 1200mesh.Prior to metallographic examination,the investigated surfaces were polished by diamond and silica pastes and etched for 5s in a solution containing 10%HF +5%HNO 3+85%H 2O.In order to achieve the surface quality required for EBSD exam-inations,samples polished by 1200mesh papers were electropo-lished at room temperature at 17V in a solution consisting of 25%H 2SO 4+15%HF +60%CH 3COOH.The prepared samples were analysed with a JSM-7000F FEG-SEM operating at 15kV and equipped with an EBSD orientation imaging system.The Kikuchi patterns were obtained with an ultra-sensitive CCD camera and processed with the Channel 5software from HKL technology.3.Results3.1.True stress–strain curvesFig.2a–d shows the stress–strain curves for CP Ti from the uni-axial compression tests at temperatures from 673to 973K with 50K intervals and strain rates of 0.001,0.01,0.1and 1s À1.As ex-pected,the flow stress is sensitive to both temperature and strain rate.It decreases with increasing temperature and also with decreasing strain rate.At high strains,the flow stress is systemat-ically decreasing with decreasing temperature from high values at 673K to low values at 973K.When deforming at 0.001s À1and above 823K,or at 0.01s À1and above 873K,or at 0.1s À1and 973K,the material exhibits a slight transitional drop in flow stress,indicating thermal softening.However at the rest of the temperatures,the stress–strain curves show a work hardening character.At 0.001and 0.01s À1and below 773K,the flow stress shows a steady state flow for true strains from 0.75to 0.92.For compression at strain rates of 0.1and 1s À1,the stress–strain behaviour as shown in Fig.2c and d is different from those deformed at 0.001and 0.01s À1.All these stress–strain curves show the characters of work hardening with the exception of that for strain rates of 0.1and 1s À1at 673K for which the stress–strain curve exhibits a thermal softening behaviour.This flow softening is attributed to the heat generation of the deformation.Although the Gleeble 1500simulator for the thermal compression is equipped with a feedback system using electrical current for tem-perature control,there is insufficient time for the adiabatic heat generated within the sample to be conducted into the cooler anvils because of the poor thermal conductivity of CP Ti.However,the analysis of the deformation behaviour presented in this study focus on the first part of the stress–strain curves,up to a true strain of about 0.5where the influence of adiabatic heating is not so strong.Fig.2a–d also reveals the strain rate sensitivity of pure titanium.At the same temperature,the flow stress increases with increasing strain rate.For instance,at temperature of 673K and strain rate of 0.001s À1,the steady flow stress is 228MPa,in contrast,it is 264MPa at strain rate of 0.01s À1for the same temperature.From the aforementioned stress–strain character,it can be con-cluded that the work hardening effect is pronounced at higher strain rate and lower temperature.For the higher temperature and lower strain rate,the stress–strain curves show transient flow softening behaviour.The data in Fig.2a–d were analysed by fitting each experimen-tal curve with a polynomial expression and taking its derivative with respect to strain in order to obtain the work hardening rate.Then,the work hardening rate was plotted against flow stress as shown in Fig.3.By inspecting these plots,it was found that the curves could be divided into three groups,viz,A:three-stageworkFig.1.EBSD map showing the microstructure of annealed CP Ti used as starting material.3106Z.Zeng et al./Materials and Design 30(2009)3105–3111hardening,B:two-stage work hardening and C:flow softening.For clarity,an example from each group is presented in Fig.4.As the shapes of the curves are quite different,it is easy to re-examine previous data in the literature dividing them into the different groups.3.2.The deformation condition mapBy plotting all our data and previous data into a map with log _eand temperature as axes,Fig.5was constructed.As seen,all data consistently fall into three domains.At low temperatures and any strain rate,the domain of three-stage work hardening is lo-cated.At higher temperature,the domain for two-stage work hard-ening is found and at even higher temperatures and lower strain rates,the domain of flow softening is found.Whenever the combined effect from changes in temperature,T ,and strain rate,_e,is considered,the Zener–Hollomon parameter,Z ,is a natural choice of parameter.It is defined as:Z ¼_e exp Q;ð1Þwhere R is the gas constant and Q is the activation energy.In the present study,the value of Q is taken from the result for CP Ti re-ported by Frost and Ashby [10],242kJ/mol.Thus,by using Eq.(1)it is possible to draw contour lines for different Z -values in Fig.5.By testing different Z -values,it was found that two border lines could be drawn which divided the map into three regions,almost perfectly consistent with the experimental information.Only one exception was found at 823K and 1s À1.The approximate values of the border lines are,log Z =16.2and 12.7,respectively,as indi-cated in Fig.5.The characteristic shapes of the stress–strain curves and resulting microstructure after deformation,will be discussed in the following sections.3.3.Three-stage work hardeningThe work hardening rate in Fig.3a shows an initial high de-crease followed by an increase up to a peak value and then a sec-ond decrease extending into negative values.By inspecting the log Z -values of each curve,it is found that they reflect the size of the peak directly.The highest peak has a log Z -value of 18.8whereas the lowest peak has a value of 15.4.The microstructures of the samples producing the curves in Fig.4were examined after full deformation i.e.after a true strain of 0.92.The optical micrograph of the sample deformed at 723K and 0.1s À1(curve A in Fig.4),is illustrated in Fig.6a.As seen,elon-gated coarse grains are mixed with refined grains.The microstruc-ture in Fig.6a is similar to that after warm rolling reported by Chun and Hwang [11].In the coarse grains,deformation twinning with common orientation is observed.From the EBSD map in Fig.6b,the deformation twinning is identified as f 10 11g twinning.3.4.Two-stage work hardeningThe work hardening rate in Fig.3b shows an initial high de-crease followed by a slow decrease down to zero or negative val-ues.By inspecting the log Z -values of each curve,it is found thatthe flow stress rate at a given work hardening is generally reflected by the log Z -value.The right-most curve has a log Z -value of 15.8whereas the left-most one has a value of 12.7.The optical micrograph of the sample deformed at 723K and 0.01s À1(curve B in Fig.4),is illustrated in Fig.7a after a total strain of 0.92.As seen an inhomogeneous microstructure,similar to that in Fig.6a,appears.However,fewer deformation twins are ob-served.From the EBSD map in Fig.7b,it is found that elongated,high angle grains have formed and that the grain size is larger than in Fig.6b.3.5.Flow softeningThe work hardening rate in Fig.3c only shows the initial high decrease crossing zero on its way down to negative values.By inspecting the log Z -values of each curve,it is found that the flow stress at a given work hardening rate is generally reflected by the log Z -value.The right-most curve has a log Z -value of 12.4whereas the left-most one has a value of 10.0.However,there is an exception for the two right-most curves having the opposite trend.The flow stress curves having low values of log Z ,exhibit a quite broad stress peak which is distinctly different from that observed in DRX of austenitic steels [12,13].Generally,austenite exhibits a more pronounced softening than that observed in the present work.The flow softening in the present work is similar to that ob-served in a -Ti by other authors [14,15].The operatingsofteningFig.2.Stress–strain curves showing the flow stress of CP Ti in compression at various temperatures and strain rates of (a)0.001s À1,(b)0.01s À1,(c)0.1s À1and (d)1s À1.Z.Zeng et al./Materials and Design 30(2009)3105–31113107mechanism may be influenced by texture softening,as reported in aluminum [16].The optical micrograph of the sample deformed at 973K and 0.01s À1(curve C in Fig.4),is illustrated in Fig.8a after a total strain of 0.92.As seen,an inhomogeneous microstructure is found but quite different from the ones in Figs.6a and 7a,as there is no defor-mation twinning observed in the optical micrograph.The EBSD map in Fig.8b shows that the boundaries of the coarse grains are serrated and elongated vertical to the compression axis.The wave-lengths of these serrations are close to the subgrain sizes.Some ofthe high angle grain boundaries,HAGBs,are discontinuous but linked with low angle grain boundaries,LAGBs,implying that there is a generation of LAGBs during deformation and that the misfit an-gle increases gradually,finally,producing HAGBs.Whenever the process has reached a misorientation above 15°,the boundary is interpreted as a HAGB.4.Discussions4.1.The deformation condition mapAs pointed out previously,the deformation condition map is di-vided into three regions which can be separated quite precisely by two Z -values.The success of the border lines is attributed to these-Fig.3.Work hardening rate of CP Ti as functions of flow stress for (a)three-stage work hardening,(b)two-stage work hardening and (c)flowsoftening.Fig.4.Typical curves for three-stage work hardening,two-stage work hardening and flow softening in CPTi.Fig.5.Deformation condition map showing the deformation character of present and previous data on CP Ti.The two lines are drawn using Eq.(2)with the marked values of log Z .A,B and C denote the conditions for the curves shown in Fig.4.3108Z.Zeng et al./Materials and Design 30(2009)3105–3111lected activation energy,Q ,as will be shown in the following discussion.A border line can be drawn by rewriting Eq.(1)into a function of temperature vs.strain rate:T ¼Q 2:303R Á1log Z Àlog _e;ð2Þwhere log Z is used as a fitting parameter.By taking the derivative with respect to log _e,the slope of the border line at constant Z ,can be found as:@T @log _e j log Z ¼Q 2:303R Á1ðlog Z Àlog _e Þ2;ð3Þwhere 2.303was used for ln10.Then,by considering the slope at aspecific point,ðlog _e;T Þ,the following relation is found:@T _e j log Z ðlog _e ¼0Þ¼2:303RT 2Á1ð4Þrecognizing that the slope of a border line is inversely proportionalto Q .Thus,the success of reproducing the boundaries between the domains in the deformation condition map is directly reflected by a successful choice of Q .Evidently,the value of 242kJ/mol,reported by Frost and Ashby [10],is consistent with the data presented in Fig.5.4.2.Deformation twinningAs confirmed by EBSD mapping,f 10 11g twinning is active in the present material.In Fig.6a and b,f 10 11g twin boundaries,crossing the coarse grains,are observed.Similar to other HCP met-als,the deformation of CP Ti begins with slip,but since the number of active slip systems is less than for BCC and FCC metals,deforma-tion twinning begins to play a vital role,even at low strains.The reason for this is that a general deformation can only be achieved by five independent,simultaneously active,slip and twinning sys-tems [17].From the presented results in Figs.6–8,it is found that the acti-vation of f 10 11g twin exhibit a strong dependence on the defor-mation temperature and strain rate.Figs.6a and 7a indicate that the deformation twins are more profuse at higher strain rate for the same deformation temperature,viz.at higher Z -values.As known from the results reported by Salem et al.[18],a profuse deformation twinning is the reason leading to the second stage of the three-stage work hardening.As CP Ti is a typical HCP metal,the directions for easy crystallographic slip are the h 11 20i or h a i close-packed directions.In order to accommodate strain in the c-Fig.6.Microstructure of CP Ti deformed at 723K and 0.1s À1to 0.92true strain and sectioned along the compression axis.(a)Optical microstructure,(b)EBSD orientation map showing HAGBs (>15°)as bold lines and LAGBs (2–15°)as fine ones.White arrows mark some of the twins.The scale bar indicates 20l m.Note that the investigated area is different from(a).Fig.7.Same as for Fig.6,but after deformation at 723K and 0.01s À1.The scale bar in (b)indicates 10l m.Z.Zeng et al./Materials and Design 30(2009)3105–31113109axis,the pyramidal slip or twin systems with h a þc i directions must be operating.In the present study,as only few twins are found in Fig.7a and none in Fig.8a,the pyramidal h a þc i slip should be a significant deformation mode.4.3.The formation of HAGBsThe EBSD maps in Figs.6b,7b and 8b indicate that the refined grains consist of HAGBs and LAGBs.Evidently,new HAGBs appear in all deformation conditions,and by comparing Figs.6b and 7b with Fig.1it is found that the old grains are refined and elongated vertical to the compression axis.Some segments of the HAGBs are isolated from any other boundary.They are most probably formed by the rotation of isolated subgrains owing to differences in the crystallographic rotation field cause by the formation of different configurations of slip systems in neighbouring regions.As shown in Fig.6b,the new interfaces formed by f 10 11g twins,crossing the coarse grains,are HAGBs.As a result,thef 10 11g twins partially contribute to the refinement of the grains and the formation of the HAGBs.Generally,the formation off 10 11g twins is observed in CP Ti deformed above 673K [19].They are known to form a misorientation angle of 57.26°between matrix and twin for CP Ti [20].In Fig.8b,it can be seen that the boundaries of the coarse grains consist of serrated HAGBs.These HAGBs are probably formed by the evolution of serrated LAGBs within the old grains.This kind of formation of serrated grains has been termed geometric dy-namic recrystallization (GDRX)by McQueen et al.[21].The process can cause the original grain boundaries to become serrated with the simultaneous formation of internal subgrains and is generally observed in high stacking fault energy (SFE)materials,where pro-nounced dynamic recovery takes place [22].As the SFE of pure tita-nium is very high (>300mJ/m 2[23]),dynamic recovery for the annihilation of dislocation is expected to be the dominant soften-ing process during high temperature deformation.With the strain increasing,the original grains progressively become elongated and ultimately fine grains form and penetrate the original serrated high angle boundaries.This mechanism of GDRX can easily be confused with discontin-uous dynamic recrystallization (DRX).McQueen and Bourell [24]have pointed out that the HCP alloys of Ti exhibits relatively high degrees of recoverability so that DRX can never be observed in them.For DRX,a necklace-type structure of new grains around the coarse grains is expected to be observed,as reported by Deh-ghan-Manshadi [25].However,this kind of necklace structure does not appear in the present compression tests.The observed struc-tures in the present work are quite similar to those observed in the EBSD studies of hot torsion of aluminum [26,27].Consequently,the term geometric dynamic recrystallization,GDRX,seems most appropriate when describing the grain refinement process of CP Ti during hot compression.5.ConclusionsThe deformation behaviour of CP Ti during hot compression tests is investigated.The following conclusions can be drawn from the studies:(1)The deformation behaviour of CP Ti can be divided into threetypes,viz.three-stage work hardening,two-stage work hardening and flow softening.(2)When the experimental data is mapped in a T vs.log _edia-gram,the present and previous data fall into three distinct domains which can be successfully separated by border lines at 16.2and 12.7for log Z ,where Z represents the Zener–Holl-omon parameter.The slopes of the border lines are inversely proportional to the activation energy,Q .The success of dividing the domains demonstrates that the present results are consistent with a previously reported value of Q .(3)The f 10 11g twin is found to be the predominant deforma-tion twin at high Z -values.In addition,a higher Z -value cor-responds to a more pronounced second-stage peak of the three-stage work hardening.(4)For the three-stage work hardening,the f 10 11g twin plays a vital role to accommodate strain in the c -axis of CP Ti.However,for two-stage work hardening and flow softening,the pyramidal slip with h a þc i directions is the main mech-anism for producing elongation and shortening in the c -axis direction.(5)GDRX,seems most appropriate when describing the grainrefinement process of CP Ti during hot compression.Thef 10 11g twin partially contributes to the refinement of the grains and the formation of theHAGBs.Fig.8.Same as Figs.6and 7but after deformation at 973K and 0.01s À1.The scale bar in (b)indicates 200l m.3110Z.Zeng et al./Materials and Design 30(2009)3105–3111AcknowledgementsThe authors gratefully acknowledge thefinancial support by the China Scholarship Council.AB Sandvik Materials Technology,Sand-viken,Sweden is thanked for supplying the material.References[1]Collings EW,Ho JC.Physical properties of titanium alloy in the sciencetechnology and application of titanium.In:Jaffee RI,Promisel NE,editors.Proceedings of thefirst international conference on titanium.London:Pergamon Press,Oxford;1970.p.331.[2]Stolyarov VV,Zhu YT,Alexandrov IV,Lowe TC,Valiev RZ.Grain refinement andproperties of pure Ti processed by warm ECAP and cold rolling.Mater Sci Eng A 2003;343:43–50.[3]Nemat-Nasser S,Guo WG,Cheng JY.Mechanical properties and deformationmechanisms of a commercially pure titanium.Acta Mater1999;47:3705–20.[4]Ray K,Poole WJ,Mitchell A,Hawbolt EB.In:Weiss I,Srinivasan R,Bania P,Eylon D,Semiatin SL,editors.Advances in the science and technology of titanium alloy processing,PA:TMS Warrendale;1997.p.201–8.[5]Tanaka Hisamune,Yamada Tomoyasu,Sato Eiichi,Jimbo Itaru.Distinguishingthe ambient-temperature creep region in a deformation mechanism map of annealed CP-Ti.Scripta Mater2006;54:121–4.[6]Ungár T,Glavicic MG,Balogh L,Nyilas K,Salem AA,Ribárik G,et al.The use ofX-ray diffraction to determine slip and twinning activity in commercial-purity (CP)titanium.Mater Sci Eng A2008;493:79–85.[7]Nagasekhar AV,Chakkingal U,Venugopal P.Candidature of equal channelangular pressing for processing of tubular commercial purity-titanium.J Mater Process Technol2006;173:53–60.[8]Stolyarov VV,Zhu YT,Lowe TC,Islamgaliev RK,Valiev RZ.A two step SPDprocessing of ultrafine-grained titanium.Nano Mater1999;11:947–54.[9]Semiatina SL,DeLo DP.Equal channel angular extrusion of difficult-to-workalloys.Mater Des2000;21:311–22.[10]Frost HJ,Ashby MF.Deformation mechanism maps.Oxford:Pergamon Press;1982.p.43.[11]Chun YB,Hwang SK.Static recrystallization of warm-rolled pure Ti influencedby microstructural inhomogeneity.Acta Mater2008;56:369–79.[12]Jonas JJ,Sellars C,Tegart M.Strength and structure under hot workingconditions.Int Metall Rev1969;14:1–24.[13]Jonas JJ,Sakai T.Deformation,processing and microstructure.In:Krauss G,editor.Ohio:ASM,Metals Park;1984.p.185–243.[14]Senkov ON,Jonas JJ.Dynamic strain aging and hydrogen-induced softening inalpha titanium.Metall Mater Trans A1996;27A:1877–87.[15]Weiss I,Semiatin SL.Thermomechanical processing of alpha titanium alloys–an overview.Mater Sci Eng A1999;263:243–56.[16]Kassner ME,Wang MZ,Perez-Prado MT,Alhajeri rge-strain softening ofaluminum in shear at elevated temperature.Metall Mater Trans A 2002;33A:3145.[17]Von Mises R.Mechanik der plastischen Formanderung von Kristallen.Z AngewMath Mech1928;8:161–85.[18]Salem Ayman A,Kalidindi Surya R,Doherty Roger D.Strain hardening regimesand microstructure evolution during large strain compression of high purity titanium.Scripta Mater2002;46:419–23.[19]Chun YB,Yu SH,Semiatin SL,Hwang SK.Effect of deformation twinning onmicrostructure and texture evolution during cold rolling of CP-titanium.Mater Sci Eng A2005;398:209–19.[20]Christian JW,Mahajan S.Deformation twinning.Progr Mater Sci1995;39:1–157.[21]McQueen HJ,Knustad O,Ryum N,Solberg JK.Microstructural evolution in Aldeformed to strains of60at400°C.Scripta Met1985;19:73–8.[22]Kassner ME,McMahon ME.The dislocation microstructure of aluminum.Metall Trans1987;18:835–46.[23]Bacon DJ,Martin JW.The atomic structure of dislocations in h.c.p.metals:I.Potentials and unstressed crystals.Philos Mag1981;4:883–900.[24]McQueen HJ,Bourell DL.Thermomechanical processing of titanium,zirconium,magnesium,and zinc in the hcp structure.J Mater Shaping Technol1988;5:163–89.[25]Dehghan-Manshadi A,Barnett MR,Hodgson PD.Hot deformation andrecrystallization of austenitic stainless steel:Part I.Dynamic recrystallization.Metall Mater Trans A2008;39A:1359–70.[26]Barnett MR,Montheillet F.The generation of new high-angle boundaries inaluminium during hot torsion.Acta Mater2002;50:2285–96.[27]Gourdet S,Montheillet F.An experimental study of the recrystallizationmechanism during hot deformation of aluminium.Mater Sci Eng A 2000;283:274–88.Z.Zeng et al./Materials and Design30(2009)3105–31113111。
tb17钛合金β相区动态再结晶行为及转变机理
重点实验室，北京 １０００９５） ＺＨＵ Ｈｏｎｇｃｈａｎｇ１，ＬＵＯＪｕｎｍｉｎｇ１，ＺＨＵ Ｚｈｉｓｈｏｕ２ （１ＳｃｈｏｏｌｏｆＭａｔｅｒｉａｌｓＳｃｉｅｎｃｅａｎｄＥｎｇｉｎｅｅｒｉｎｇ，Ｎａｎｃｈａｎｇ Ｈａｎｇｋｏｎｇ Ｕｎｉｖｅｒｓｉｔｙ，Ｎａｎｃｈａｎｇ３３００６３，Ｃｈｉｎａ；２ Ａｖｉａｔｉｏｎ ＫｅｙＬａｂｏｒａｔｏｒｙｏｆ ＳｃｉｅｎｃｅａｎｄＴｅｃｈｎｏｌｏｇｙｏｎＡｄｖａｎｃｅｄＴｉｔａｎｉｕｍ Ａｌｌｏｙｓ，ＡＥＣＣ ＢｅｉｊｉｎｇＩｎｓｔｉｔｕｔｅｏｆＡｅｒｏｎａｕｔｉｃａｌＭａｔｅｒｉａｌｓ，Ｂｅｉｊｉｎｇ１０００９５，Ｃｈｉｎａ）
Ｖｏｌ．４８ Ｎｏ．２
Ｆｅｂ．２０２０ ｐｐ．１０８－１１３
犜犅１７钛合金β相区动态再结晶 行为及转变机理
Ｄｙｎａｍｉｃｒｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎｂｅｈａｖｉｏｒａｎｄｔｒａｎｓｆｏｒｍａｔｉｏｎ ｍｅｃｈａｎｉｓｍｉｎβｐｈａｓｅｒｅｇｉｏｎｏｆＴＢ１７ｔｉｔａｎｉｕｍａｌｌｏｙ
朱 鸿 昌１ ，罗 军 明１ ，朱 知 寿２
犃犫狊狋狉犪犮狋：Ｔｈｅｔｈｅｒｍａｌｃｏｍｐｒｅｓｓｉｏｎｅｘｐｅｒｉｍｅｎｔｏｆ ＴＢ１７ｔｉｔａｎｉｕｍ ａｌｌｏｙｉｎｔｈｅβｐｈａｓｅｒｅｇｉｏｎ ｗａｓ ｃａｒｒｉｅｄｏｕｔｂｙ Ｇｌｅｅｂｌｅ３８００ｈｏｔｃｏｍｐｒｅｓｓｉｏｎｓｉｍｕｌａｔｏｒ．Ｔｈｅｄｙｎａｍｉｃｒｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎｂｅｈａｖｉｏｒａｎｄ ｔｒａｎｓｆｏｒｍａｔｉｏｎ ｍｅｃｈａｎｉｓｍｉｎｔｈｅβｐｈａｓｅｒｅｇｉｏｎｏｆｔｈｅＴＢ１７ｔｉｔａｎｉｕｍａｌｌｏｙｗｅｒｅｓｔｕｄｉｅｄ．Ｔｈｅｒｅｓｕｌｔｓ ｓｈｏｗｔｈａｔｔｈｅｄｙｎａｍｉｃｒｅｃｏｖｅｒｙ （ＤＲＶ）ａｎｄｄｙｎａｍｉｃｒｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎ （ＤＲＸ）ｏｃｃｕｒｉｎｔｈｅβｐｈａｓｅ ｒｅｇｉｏｎｄｕｒｉｎｇｄｅｆｏｒｍａｔｉｏｎｐｒｏｃｅｓｓｏｆｔｈｅＴＢ１７ｔｉｔａｎｉｕｍａｌｌｏｙ．Ｔｈｅｒｅａｒｅｔｗｏｄｙｎａｍｉｃｒｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎ ｎｕｃｌｅａｔｉｏｎｓｉｔｅｓａｔｄｉｆｆｅｒｅｎｔｓｔｒａｉｎｒａｔｅｓ．Ａｔｌｏｗｓｔｒａｉｎｒａｔｅ，ｉｔｍａｉｎｌｙｎｕｃｌｅａｔｅｓｉｎｓｉｄｅｔｈｅｇｒａｉｎｓ，ａｎｄ ａｔｈｉｇｈｓｔｒａｉｎｒａｔｅ，ｉｔｉｓｎｅａｒｔｈｅｇｒａｉｎｂｏｕｎｄａｒｙ．ＡｃｃｏｒｄｉｎｇｔｏＥＢＳＤａｎｄＴＥＭ ａｎａｌｙｓｅｓ，ｔｈｅ ｍａｉｎ ｍｅｃｈａｎｉｓｍｈａｐｐｅｎｅｄａｔｌｏｗｓｔｒａｉｎｒａｔｅｉｓｃｏｎｔｉｎｕｏｕｓｄｙｎａｍｉｃｒｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎ （ＣＤＲＸ）ｗｈｉｃｈ ｍａｉｎｌｙ ｉｓｃｏｎｔｒｏｌｌｅｄｂｙｓｕｂｇｒａｉｎｒｏｔａｔｉｏｎ．Ｄｉｓｃｏｎｔｉｎｕｏｕｓｄｙｎａｍｉｃｒｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎ （ＤＤＲＸ）ｏｃｃｕｒｓａｔｈｉｇｈ ｓｔｒａｉｎｒａｔｅｓ，ｔｈｅｍａｉｎｆｏｒｍｏｆｄｅｆｏｒｍａｔｉｏｎｉｓｇｒａｉｎｂｏｕｎｄａｒｙｓｈｅａｒａｃｃｏｍｐａｎｉｅｄｂｙｓｕｂｇｒａｉｎｒｏｔａｔｉｏｎ． Ａｌｔｈｏｕｇｈｔｈｅｔｗｏｄｙｎａｍｉｃｒｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎｓａｒｅｔｒａｎｓｆｏｒｍｅｄｉｎ ｄｉｆｆｅｒｅｎｔ ｗａｙｓ，ｔｈｅｅｓｓｅｎｃｅｉｓｔｏ ｆｏｒｍｎｅｗ ｄｙｎａｍｉｃｒｅｃｒｙｓｔａｌｌｉｚｅｄｇｒａｉｎｓｔｈｒｏｕｇｈｔｈｅｐｒｏｐａｇａｔｉｏｎ，ｓｌｉｐａｎｄｃｅｌｌｓｔｒｕｃｔｕｒｅｅｖｏｌｕｔｉｏｎｏｆ ｄｉｓｌｏｃａｔｉｏｎｓ． 犓犲狔狑狅狉犱狊：ＴＢ１７ｔｉｔａｎｉｕｍａｌｌｏｙ；βｔｈｅｒｍａｌｄｅｆｏｒｍａｔｉｏｎ；ｄｙｎａｍｉｃｒｅｃｏｖｅｒｙ；ｃｏｎｔｉｎｕｏｕｓｄｙｎａｍｉｃｒｅｃｒｙｓ ｔａｌｌｉｚａｔｉｏｎ；ｄｉｓｃｏｎｔｉｎｕｏｕｓｄｙｎａｍｉｃｒｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎ
钛合金机壳的变形原因
钛合金机壳的变形原因
钛合金机壳的变形原因有多种，包括但不限于以下几点：
1. 加工过程：钛合金机壳在加工过程中，可能由于温度、外力、应力和其他因素的作用，导致其原子结构发生变化，从而产生变形。
此外，钛合金机壳的生产和加工工艺较为复杂，一些处理不当的情况下会导致钛合金材料在使用过程中出现变形、开裂等问题。
2. 温度变化：钛合金在不同的温度下有不同的力学性能。
当钛合金处于高温状态时，其结构整体较为松散，因此很容易发生变形。
而当钛合金处于低温状态时，其脆性较大，也容易发生变形。
此外，在高温状态下，钛合金可能会发生氧化反应，导致变形加剧。
3. 低的模量：钛合金的模量比钢低40%左右，比铝低60%左右，模量越低，变形就越大。
因此，相对于其他金属材料，钛合金容易变形。
4. 工艺问题：钛合金机壳的生产和加工工艺较为复杂，一些处理不当的情况下会导致钛合金材料在使用过程中出现变形、开裂等问题。
综上所述，钛合金机壳的变形主要与应力、温度、外力、低的模量以及工艺问题等多种因素有关。
因此，在钛合金的使用过程中，需要针对不同的因素进行相应的处理和控制，以保证钛合金材料的稳定性和可靠性。
TC17钛合金热变形过程中片状组织演变规律
球化现 象不明显 ；随着变形量 的增加 ，片状 组织被不同程度 的弯 曲、破碎 ，球化程度随着变形量 的增加逐渐变大 。变 形温度 对球 化过程也起一定 的作用 ，随着变形温度的升高 ，球化效果越来越 明显 ，这与较高 的变形温度会提高位错或
原 子 的迁 移 能 力 使 片 状 组 织 有 足 够 的能 量 通 过 界 面 迁 移 实 现 断 裂 、球 化 有 关 。 关 键 词 ：Ｔ １ 金 ；片 层 组 织 ； 化 过 程 ； 效 应 变 Ｃ ７合 球 等
Ｚ ｏ ｉ ｎ ｕ Ｗ ａ ｇ Ｘｉ ｏ ｉ ｇ ， Ｘｕ Ｂ ｎ ， Ｍａ Ｘｉｎ Ｗ ａ ｇ Ｋ ｉｕ ｎ ， Ｚ ｎ ｅｄ ｎ ｈ ｕ Ｊａ ｈ ａ ， ｎ ａ ｙｎ ｉ ｏｇ， ｎ ａｘ ａ ｅ ｇ Ｗ ｉｏ ｇ
（ ．Ｂ ｏｈｎＩ ｎ＆ Ｓｅｌ ｏ， ｔ． Ｓ ａｇａ ２ ０ ４ ，Ｃ ｉａ １ ａｓａ ｒ ｏ ｔｅ Ｃ ． Ｌｄ， ｈｎｈｉ ０９ ０ ｈｎ ） （ ．Ｓａ ｅ ａｏａｏ ｆ ｏｉｉｃｔｎＰｏｅｓ ｇ ｏｔｗ ｓ ｒ ｏｙ ｃｎｃｌ ｎｖｒｔ， ｉａ １０ ２ Ｃ ｉａ ２ ｔｔ Ｋ ｙＬ ｂｒｔ ｙｏ Ｓｌ ｆ ａ ｏ ｒｃｓ ｎ ，Ｎ ｒ ｅｔ ｎＰｌｔ ｈ ｉ ｉ ｓｙ Ｘ ’ｎ７ ０ ７ ， ｈｎ ） ｅ ｒ ｄｉ ｉ ｉ ｈ ｅ ｅ ａＵ ｅｉ （ ．Ｗｅｔ ｎＳｐ ｒ ｎ ｕｔ ｇＴｃ ｎｌ ｉｓ ｏ，Ｌｄ， ｉａ ， １０ ８ Ｃ ｉａ ３ ｓ ｒ ｕ ｅ ｏ ｄｃｉ ｅｈ ｏ ｇｅ Ｃ ． ｔ． Ｘ ’ｎ ７０ １ ， ｈｎ ） ｅ ｃ ｎ ｏ
作 用 ，由于压 气机 盘不 同部 位 承受 的应 力状 态 不 同 ，
故 对性 能 的要 求 不 同 ，传 统 的单 一 组 织 已 不 能充 分 发 挥材 料 的潜 在 性 能 ］ 。而 采用 可 控 应 变 等 温模 锻
TC17钛合金热变形过程中片状组织演变规律
TC17钛合金热变形过程中片状组织演变规律周建华;王晓英;徐斌;马雄;王凯旋;曾卫东【期刊名称】《钛工业进展》【年(卷),期】2012(029)005【摘要】通过等温锻造试验和有限元模拟,研究了TC17钛合金在α+β两相区变形过程中热加工工艺参数对显微组织演变的影响规律.通过组织观察分析发现:随着变形程度和变形温度的增加,TC17钛合金中的片层组织逐渐向球化组织转变.变形量对片状组织的球化起决定作用,当变形量为小于20％时,仅有少数片状α相发生了弯折或扭曲,球化现象不明显；随着变形量的增加,片状组织被不同程度的弯曲、破碎,球化程度随着变形量的增加逐渐变大.变形温度对球化过程也起一定的作用,随着变形温度的升高,球化效果越来越明显,这与较高的变形温度会提高位错或原子的迁移能力使片状组织有足够的能量通过界面迁移实现断裂、球化有关.【总页数】4页(P15-18)【作者】周建华;王晓英;徐斌;马雄;王凯旋;曾卫东【作者单位】宝山钢铁股份有限公司,上海200072;宝山钢铁股份有限公司,上海200072;宝山钢铁股份有限公司,上海200072;西北工业大学凝固技术国家重点实验室,陕西西安710072;西北工业大学凝固技术国家重点实验室,陕西西安710072;西部超导材料科技有限公司,陕西西安710018;西北工业大学凝固技术国家重点实验室,陕西西安710072【正文语种】中文【相关文献】1.TC17钛合金在α+β两相区加热过程中片状组织的粗化行为 [J], 王云利;赵兴东;徐建伟;曾卫东2.片状TA15钛合金在热变形过程中的动态球化动力学 [J], 吴成宝;杨合;樊晓光;孙志超3.TC17钛合金在热变形过程中的组织演变规律 [J], 徐斌;王晓英;周建华;王凯旋;曾卫东4.TC17钛合金在热变形过程中的组织演变规律 [J], 徐斌;王晓英;周建华;王凯旋;曾卫东5.片状TA15钛合金在热变形过程中的动态球化动力学 [J], 吴成宝;杨合;樊晓光;孙志超因版权原因，仅展示原文概要，查看原文内容请购买。
TC17钛合金在α+β两相区加热过程中片状组织的粗化行为
TC17钛合金在α+β两相区加热过程中片状组织的粗化行为王云利;赵兴东;徐建伟;曾卫东【摘要】研究了TC17钛合金分别在800℃和840℃两相区加热过程中片状组织的粗化行为.结果表明,随着加热时间的延长,片状组织逐渐粗化,粗化速率逐渐降低.片状α相的粗化过程是一个元素扩散过程,元素从势能较高的片状组织末端或缺陷位置向势能较低的平滑位置迁移,导致片状α相粗化.经计算,800℃下粗化指数为0.37,840℃下粗化指数为0.42,根据LSW粗化理论,TC17钛合金片状组织的粗化行为由体扩散和界面反应共同控制.此外,基于Gibbs-Thompson理论,建立了TC17钛合金片状组织的粗化速率预测模型.%The coarsening behavior of TC17 titanium alloy at 800℃and 840℃was investigated in this research.The results show that the thickness of alpha lamellae continues to increase with heat treatment time .However, the coarsening rate decreases with heat treatment time.The coarsening of alpha lamellae is a process of element diffusion. Elements will migrate from the large curvature (the termination or defect position ) to the small curvature (the flat interface), which leads to the dissolution of the termination tip and the coarsening of the adjacent plate.The coarsening index is calculated as 0.37 and 0.42 for 800 ℃and 840 ℃.According to the LSW theory, the coarsening behavior of TC17 titanium alloy is controlled by bulk diffusion and interface reaction.The prediction model of coarsening rate of TC17 titanium alloy is successfully established based on Gibbs-Thompson theory.【期刊名称】《钛工业进展》【年(卷),期】2018(035)003【总页数】4页(P30-33)【关键词】TC17钛合金;两相区加热;粗化行为【作者】王云利;赵兴东;徐建伟;曾卫东【作者单位】中国航发沈阳黎明航空发动机有限责任公司,辽宁沈阳 110043;中国航发沈阳黎明航空发动机有限责任公司,辽宁沈阳 110043;西北工业大学凝固技术国家重点实验室,陕西西安 710072;西北工业大学凝固技术国家重点实验室,陕西西安 710072【正文语种】中文【中图分类】TG146.230 引言钛合金具有高的比强度、良好的高温性能以及损伤容限性能，已被广泛应用于航空领域[1]。
热变形参数对TC17合金的片状α球化过程的影响
-6-
合金片状组织开始球化最小真应变仅为 0.6（变形量约为 45％） ，达到 0.8 时就可以完全 球化。 3． 降低应变速率、提高变形温度不利于片状组织的球化。在应变速率低于 10-3s-1 时，即使 应变量超过 0.8 也未发生片状α球化。 因此，为了得到较为细小均匀等轴α组织，应该在较高的变形速率下适当降低变形温度 和提高变形程度。
图1
TC17 合金 920℃锻后空冷高倍组织
3. 实验结果与分析
3.1 高温变形行为
TC17 钛合金高温压缩变形时真应力真应变曲线如图 2 所示。
(a)
300 True Stress (MPa)
830 C
(MPa)
o
250 200 150 100 50
(b)
850 C 1s
-1
o
True stress
& ε =10-1s-1
(d)
& ε =1s-1
不同应变速率压缩的显微组织(850℃，ε=0.8)
在应变为 0.8, 变形速率为 0.1s-1 时变形温度对 TC17 合金显微组织的影响如图 5 所示。 图中可见，在 870℃变形时，原始的晶界α只是被切断或发生弯折，片状α组织没有发生明显 的球化；850℃变形后片状组织大部分球化，但图中还可见少量原始平直α相，片状α球化后 并非理想的球形，呈长椭圆形；在相对较低的 830℃时变形，片状状组织等轴化发生非常充 分，得到等轴α相的晶粒度也较高温下变形的组织更为细小。
TC17钛合金的热压缩变形行为研究
要化学成分（质量分数，%）如表 1 所示。铸锭经 β 相区和 α+β 两相区多火次锻造至 Φ350mm 棒材，其显微组织如 图 1 所示，主要由初生等轴 α 相和 β 转变组织组成，α 相 尺 寸 小 于 10μm。经 金 相 法 测 得 α+β/β 相 转 变 温 度 为 900℃。 2.2 试验方法
M 机械加工与制造 achining and manufacturing
TC17 钛合金的热压缩变形行为研究
董轶1，李巍1，周阳洋1，张雪敏1，刘继雄2，屠孝斌1，王少阳1
（１、宝鸡钛业股份有限公司，陕西 宝鸡 ７２１０１４ ；２． 宝钛集团有限公司，陕西 宝鸡 ７２１０１４）
摘 要 ：在 Ｇｌｅｅｂｌｅ－３８００ 热 模 拟 机 上 对 ＴＣ１７ 钛 合 金 进 行 热 压 缩 实 验，研 究 ＴＣ１７ 钛 合 金 棒 材 在 变 形 温 度 为 ８１０℃～ ９３０℃、应变速率为 １０－２ｓ－１￣１０１ｓ－１ 以及变形程度为 ２０％ ～ ６０％ 条件下的流变行为 ；利用金相显微镜分析 ＴＣ１７ 钛合金棒材在不同变形条件下的组织演化规律。结果表明 ：当应变速率与变形温度固定时，不同变形量对于 ＴＣ１７ 钛合金的流动应力曲线影响较小 ；当变形量与变形速率固定时，变形温度越高时，流变应力值越低，应力 － 应变 曲线越稳定 ；当变形温度与变形量固定时，峰值应力随应变速率的减小而降低。变形中的软化机制主要以动态回复和 动态再结晶为主。 关键字 ：ＴＣ１７ ；热压缩 ；流变应力 中图分类号 ：Ｏ６１４．４１＋２ 文献标识码 ：Ａ 文章编号 ：１００２－５０６５（２０２０）０８－００２３－４
The thermal compression deformation behavior of TC17 titanium alloyቤተ መጻሕፍቲ ባይዱ
钛合金的热变形行为研究
钛合金的热变形行为研究钛合金作为一种重要的结构材料，具有高强度、低密度和良好的耐腐蚀性能，在航空航天、汽车工业和生物医学领域得到广泛应用。
然而，钛合金的热变形行为是其应用过程中需要深入研究的重要问题之一。
因此，本文将探讨钛合金的热变形行为及其对材料性能的影响。
首先，钛合金的热变形行为可以通过压缩试验来研究。
压缩试验可以模拟材料在高温下的变形过程，对于分析材料力学性能的变化规律具有重要意义。
实验结果显示，钛合金在高温下具有较好的塑性变形能力，但随着变形温度的升高，材料的塑性变形能力逐渐减弱，甚至出现明显的脆性断裂。
这是由于高温下钛合金晶体结构发生相变，导致其力学性能发生巨大变化。
因此，在设计钛合金零件时，需要考虑材料的热变形行为对其力学性能的影响。
另一方面，钛合金的热变形行为还与合金元素的含量和微观结构有关。
例如，钛合金中添加的铝、钒等元素可以显著改善其高温下的塑性变形能力，提高材料的耐热性能。
同时，通过调整钛合金的晶粒大小和晶界分布，可以进一步优化材料的热变形行为。
实验结果表明，钛合金的细晶粒结构可以显著提高材料的塑性变形能力和抗蠕变性能。
这是由于细晶粒结构可以增加晶界的数量和长度，阻碍位错的运动，从而提高材料的塑性变形能力。
此外，钛合金的热变形行为还与变形速率和应变温度有密切关系。
在高变形速率下，钛合金的流动应力显著增加，材料的变形能力降低。
而在较低的变形速率下，材料可以充分发挥其塑性变形能力。
另外，随着应变温度的升高，钛合金的变形能力逐渐增加，但当温度超过某一临界值后，材料会出现快速蠕变现象，导致材料的强度和塑性急剧下降。
因此，在实际应用中，需要合理选择变形速率和应变温度，以避免材料的变形失效。
总之，钛合金的热变形行为是其应用过程中需要重点关注的问题。
通过研究钛合金的热变形行为，可以深入了解材料的力学性能和变形机制，为钛合金材料的设计和加工提供理论依据。
未来，随着科技的不断发展，钛合金的热变形行为研究将进一步深入，为相关领域的创新提供更多的支持。
片层组织TC17钛合金高温变形行为研究
片层组织TC17钛合金高温变形行为研究王华;冀胜利;王凯旋;曾卫东【摘要】通过热压缩试验研究了具有初始片层组织的TC17钛合金在780～860℃和应变速率0.001～10 s-1范围内的热变形行为和组织演变.分析了该合金在两相区变形的应力-应变曲线特征,其流变应力本构关系可以用双曲正弦方程和Zener-Hollomon参数描述,得到TC17合金在两相区变形的平均激活能为488.86 kJ·mol-1.显微组织分析发现:TC17合金在两相区变形时组织演变的主要特征是片层组织球化;热变形参数严重影响片层组织球化过程的进行,加大变形量、降低应变速率以及提高变形温度可以提高片状组织的动态球化程度.【期刊名称】《钛工业进展》【年(卷),期】2010(027)006【总页数】4页(P16-19)【关键词】TC17钛合金;片层组织;流变应力方程;动态球化【作者】王华;冀胜利;王凯旋;曾卫东【作者单位】海军装备部,陕西,西安,710021;陕西宏远航空锻造有限责任公司,陕西,三原,713801;西北工业大学材料学院,陕西,西安,710072;西北工业大学材料学院,陕西,西安,710072【正文语种】中文TC17(Ti-5A l-4M o-4C r-2Sn-2Zr)合金是一种富β稳定元素的α+β型两相钛合金,该合金具有强度高、断裂韧性好、淬透性高等优点,主要用于制造航空发动机风扇、压气机盘[1]。
整体叶盘结构是高推比航空发动机的重要选择。
航空发动机压气机盘各部位的服役环境差异很大,叶片部位和轮盘部位的性能要求不同。
由于叶片要求高强度和高周疲劳性能,而轮盘要求高的断裂韧性和良好的高温性能。
众所周知,等轴组织具有强度高、塑性好和良好的抗裂纹萌生能力等优点,但高温性能和断裂韧性较差;反之,片层组织具有良好的断裂韧性、抗裂纹扩展能力、高温持久强度、蠕变抗力等优点,但塑性和热稳定性较差。
锻造整体式风扇盘和压气机盘件是通过在两相区热变形促使具有原始片层组织的坯料的不同部位组织发生不同程度的球化,从而得到理想组织分布以满足各部位性能的最佳匹配。
TC17钛合金在高温与高应变率下的动态压缩力学行为研究
TC17钛合金在高温与高应变率下的动态压缩力学行为研究牛秋林;陈明;明伟伟【摘要】采用微型分离式霍普金森压杆实验系统对TC17钛合金在高温、高应变率条件下的动态力学行为进行研究,测试材料的应力应变行为,分析实验温度、应变率和应变对其动态力学性能的影响规律.实验结果表明:当应变率为3000 s-1时,TC17钛合金表现出明显的应变硬化效应,但在高温、高应变率条件下其应变硬化效应明显减弱;TC17钛合金具有应变率强化效应,但在温度升高过程中其应变率敏感性随着实验温度的升高而先减小后增大;实验温度对TC17钛合金的动态压缩力学行为的影响非常明显,温度敏感性因子随温度的升高大幅度增大.%The dynamic mechanics behaviors of TC17 titanium alloy were studied under the condi-tions of high temperatures and high strain rates by using SHPB experimental system.The stress and strain behaviors of the materials were tested,and the effects of the experimental temperatures,strain rates and strains on the dynamic mechanics properties were analyzed.The experimental results show that when the strain rate is as 3000 s-1 ,TC17 titanium alloy exhibits obvious strain hardening effects, but the strain hardening effects are weakened at high temperatures and high strain rates.TC17 titani-um alloy has strain rate strengthening effect,but during the temperature increase processes,the strain rate sensitivity decreases firstly and then increases with the temperature increase.The effects of temperature on the dynamic compressive mechanics behaviors of TC17 titanium alloy are very obvi-ous,and the temperature sensitivity factor increases with the temperature increase.【期刊名称】《中国机械工程》【年(卷),期】2017(028)023【总页数】6页(P2888-2892,2897)【关键词】材料力学;动态压缩力学行为;分离式霍普金森压杆;TC17钛合金;高应变率【作者】牛秋林;陈明;明伟伟【作者单位】湖南科技大学机电工程学院,湘潭,411201;湖南科技大学难加工材料高效精密加工湖南省重点实验室,湘潭,411201;上海交通大学机械与动力工程学院,上海,200240;上海交通大学机械与动力工程学院,上海,200240【正文语种】中文【中图分类】TG146.2TC17钛合金是一种典型的近β相的α+β两相钛合金，具有高强度、高韧性和高淬透性，可被用于制造航空发动机压气机盘、风扇盘、离心叶轮及各种零部件[1-2]。
钛合金高温变形实验报告
一、实验目的1. 研究钛合金在高温下的变形行为。
2. 探究不同温度、不同变形速度对钛合金变形性能的影响。
3. 分析钛合金高温变形过程中的组织演变规律。
二、实验材料及设备1. 实验材料：某型号钛合金板材。
2. 实验设备：高温炉、万能试验机、金相显微镜、扫描电镜等。
三、实验方法1. 实验步骤：（1）将钛合金板材切割成所需尺寸。
（2）将钛合金板材放入高温炉中，按照预定的温度和时间进行加热。
（3）将加热后的钛合金板材取出，迅速放入万能试验机中进行压缩变形实验。
（4）观察钛合金板材的变形行为，记录变形量。
（5）对变形后的钛合金板材进行金相显微镜和扫描电镜观察，分析组织演变规律。
2. 实验参数：（1）实验温度：900℃、1000℃、1100℃。
（2）变形速度：1mm/min、2mm/min、3mm/min。
四、实验结果与分析1. 钛合金在高温下的变形行为（1）随着温度的升高，钛合金的变形抗力逐渐降低，变形量逐渐增大。
（2）在900℃时，钛合金的变形抗力较高，变形量较小；在1100℃时，钛合金的变形抗力较低，变形量较大。
（3）在相同温度下，随着变形速度的增加，钛合金的变形抗力逐渐增大，变形量逐渐减小。
2. 钛合金高温变形过程中的组织演变规律（1）在900℃时，钛合金板材经过压缩变形后，组织以等轴晶为主，晶粒尺寸较小。
（2）在1000℃时，钛合金板材经过压缩变形后，组织以等轴晶和细长晶为主，晶粒尺寸有所增大。
（3）在1100℃时，钛合金板材经过压缩变形后，组织以细长晶为主，晶粒尺寸较大。
五、结论1. 钛合金在高温下具有良好的变形性能，随着温度的升高，变形抗力逐渐降低，变形量逐渐增大。
2. 钛合金高温变形过程中的组织演变规律：在900℃时，以等轴晶为主；在1000℃时，以等轴晶和细长晶为主；在1100℃时，以细长晶为主。
3. 实验结果表明，高温变形对钛合金的组织和性能具有重要影响，为钛合金高温成形工艺的优化提供了理论依据。
TB17钛合金β相区晶粒长大行为
（ Ｂ ｅ ｉ ｊ ｉ ｎ ｇ Ｉ ｎ ｓ ｔ ｉ ｔ ｕ ｔ ｅ ｏ ｆ Ａ ｅ ｒ ｏ ｎ ａ ｕ ｔ ｉ ｃ ａ ｌ Ｍａ ｔ ｅ ｉ ｒ ａ ｌ ｓ ，Ｂ ｅ ｉ ｊ ｉ ｎ ｇ １ ０ ０ ０ ９ ５ ，Ｃ ｈ ｉ ｎ ａ ）
Ａｂ ｓ ｔ ｒ ａ ｃ ｔ ：Ｔ ｈ ｅ／ ３ ｇ ｒ ａ ｉ ｎ ｇ ｒ ｏ ｗ ｔ ｈ ｂ ｅ ｈ ａ ｖ ｉ ｏ ｒ ｗ ａ ｓ ｉ ｎ ｖ ｅ ｓ ｔ ｉ ｇ ａ ｔ ｅ ｄ ｏ ｎ ａ ｎ ｅ ｗ ｈ ｉ ｇ ｈ ｓ ｔ ｒ ｅ ｎ ｇ ｔ ｈ ａ ｎ ｄ ｈ ｉ ｇ ｈ ｔ ｏ ｕ ｇ ｈ ｎ ｅ ｓ ｓ Ｔ Ｂ１ ７ ｔ ｉ ｔ ａ ｎ ｉ ｕ ｍ
０ ． ２ ３ ．Ｈｅ ａ ｔ ｔ ｒ ｅ ａ ｔ ｅ ｄ ｆ ｏ ｒ ｓ ｈ ｏ ｒ ｔ ｔ ｉ ｍｅ ，ｔ ｈ ｅ ｋ ｉ ｎ ｅ ｔ ｉ ｃ ｓ ｆ ａ ｃ ｔ ｏ ｒ ｓ ａ ｒ ｅ ｍａ ｉ ｎ ｉ ｎ ｌ ｆ ｕ ｅ ｎ ｃ ｅ ｆ ０ ｒ ｔ ｈ ｅ／ ３ ｇ ｒ ａ ｉ ｎ ｇ ｒ ｏ ｗ ｔ ｈ，ｗｈ ｉ ｌ ｅ ｌ ｏ ｎ ｇ ｔ ｉ ｍｅ ｈ ｅ ａ ｔ
要 ：以新型超高强韧 Ｔ Ｂ １ ７钛合金棒材为研究对象 ，研究 了 Ｔ Ｂ １ ７钛合金 相 区晶粒 的长大行为 。考察 了 Ｔ Ｂ １ ７钛
合金在不 同温度和保温时问的条件下 ， 晶粒 尺寸的变化 ，通过 Ｂ ｅ ｃ ｋ 公 式计 算了晶粒长大参数 ，采Ａ ｒ ｒ ｈ ｅ ｎ ｉ ｕ ｓ 公式计算 了晶粒长大激活能。结果 表明 ，加热温度及保温 时间对 Ｔ Ｂ １ ７钛合金 ／ ３ 晶粒长 大行 为具 有重要影响 。在 ８ ６ ０～１ ０ ４ ５℃ 进行等温加热 ，Ｔ Ｂ １ ７钛合金 晶粒等温长大曲线近似符合指 数关 系，』 Ｂ晶粒 长大指数 ｎ在 ０ ． １ ２～０ ． ２ ３之 间 。短 时保 温 ，卢晶粒长大过程 中动力学影 响因素 占主导 作用 ，延 长保温 时间 ，动 力学影 响因素作 用降 低。Ｔ Ｂ １ ７钛合金 晶粒 长大激活能为 ４ ８ ． ２ ６ ｋ Ｊ ／ ｍ ｏ ｌ 。 关键词 ：Ｔ Ｂ １ ７钛合金 ；口晶粒尺寸 ；晶粒长大激活能 中图分类号 ：Ｔ Ｇ１ ４ ６ ． ２ ３ 文献标识码 ：Ａ 文章编号 ：１ ０ ０ ９ — ９ ９ ６ ４ （ ２ ０ １ ６ ） ０ ６ － ０ ０ １ １ ０ － ５
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钛合金具有密度小，比强度高，耐蚀性好等优 异特性，广泛应用在航空、航天、航海及化工等领 域。为满足新一代飞机和高性能航空发动机的长 寿命与高减重设计需求，对轻质高强材料提出了 更高的要求[1-4]。TB17 钛合金是我国自主研发设 计的新型亚稳 β 型超高强韧钛合金，通过合适的 固溶强化处理，强度可达 1350 MPa 以上，并具有 较好的强度-塑性-韧性匹配，可用于制造大型结构 锻件等。
收稿日期：2018-09-17；修订日期：2018-11-29 基金项目：国家自然科学基金项目（5176401） 通讯作者：朱知寿（1966—），男，博士，研究员，主要从事航 空钛合金及应用技术研究，（E-mail）zhuzzs@。
了温度-应变速率-应变量之间的本构关系，误差分 析结果表明该方法建立的本构模型具有较高的精 确度。陈海生等[12] 基于 BP 网络对 Ti-6 Al-3 Nb-2 Zr-1 Mo 合金等温压缩流变应力进行预测，结果与 实验结果十分接近，具有非常高的准确度。
目前，国内对 TB17 钛合金热变形特征的研究 报道较少，本工作主要研究 TB17 钛合金热压缩过 程中的高温变形行为，建立 Arrhenious 本构方程， 分析热变形过程中动态再结晶行为。
1 实验材料与方法
所用材料为中国航发北京航空材料研究院研 制的 TB17 钛合金，名义成分为 Ti-6.5 Mo-2.5 Cr-2 Nb-1 Sn-1 Zr-4 Al，锻件的原始组织如图 1 所示，可 以看出显微组织为典型的双相组织，初生 α 相呈细 小短棒状，采用金相法测得其相变点约为 845 ℃。
2019 年 第 39 卷 第 3 期 第 44 – 52 页
航 空 材 料 学 报
JOURNAL OF AERONAUTICAL MATERIALS
2019，Vol. 39 No.3 pp.44 – 52
TB17 钛合金高温压缩变形行为
朱鸿昌1, 罗军明1, 朱知寿2
（1．南昌航空大学 材料科学与工程学院，南昌 330063；2．中国航发北京航空材料研究院 先进钛合金航空科技重点实验室， 北京 100095）
摘要：通过 Gleeble 3800 热模拟试验机对 TB17 钛合金在变形温度 860～980 ℃、应变速率为 0.001～1 s–1、最大变 形量为 70% 下高温变形行为进行研究。通过材料参数与真应变之间的关系，利用 Arrhenious 本构方程关系式和 Z 参数建立流变应力和变形温度、应变速率和真应变三者之间的本构关系，并对组织进行分析。结果表明：TB17 钛 合金在应变速率为 0.001～0.01 s–1、变形温度为 890～980 ℃ 下更容易发生连续动态再结晶，而在应变速率为 0.1～1 s–1 下主要发生不连续动态再结晶；误差分析结果显示计算值与实测值平均相对误差为 6%，说明建立的本构 关系模型具有较高的准确度。 关键词：TB17 钛合金；热变形；连续动态再结晶；不连续动态再结晶；Arrhenious 本构方程 doi：10.11868/j.issn.1005-5053.2018.000103 中图分类号：TG146 文献标识码：A 文章编号：1005-5053(2019)03-0044-09
2 实验结果与分析
2.1 应力-应变曲线 图 2 为 TB17 钛合金在变形温度为 860～980 ℃、
应变速率为 0.001～1 s–1 的应力-应变曲线。曲线在 开始阶段为弹性形变，即应变量很小的情况下，流 变应力随应变的增加急剧上升，此时加工硬化在变 形初始阶段占主导地位。当应变不断增加，开始出 现一个不连续屈服点，不连续屈服现象在很多钛合 金 中 都 出 现 过 。 如 TB6[13]、 Ti5553[6]、 Ti55511[14] 等。不连续屈服现象主要是由于可动位错在晶界 处快速增殖所引起的[15]。随着应变的增加，流变应 力开始下降并基本保持水平，此时加工硬化和动态 软化达到平衡。值得注意的是，应变速率较高时， 应力-应变曲线均呈现一个较宽的峰，这可能与动 态再结晶的发生和局部温升效应有关[16]。 2.2 热变形参数对显微组织的影响 2.2.1 变形温度对显微组织的影响
第3期
TB17 钛合金高温腐 蚀 ， 腐 蚀 剂 体 积 比 为 HF∶HNO3∶H2O = 1∶2∶7，用 Leica DMI 3000 M 型卧式金相显微镜观 察金相组织。
50 μm
图 1 TB17 钛合金原始组织 Fig. 1 Original microstructure of TB17 titanium alloy
实验在 Gleeble-3800 热模拟试验机上进行，试 样尺寸为 ϕ8 mm × 12 mm 的圆柱体，表面光亮且无 氧化层。试样两端垫上石墨，以减少压头与试样之 间的摩擦力；表面焊接热电偶，反馈实验过程中温 度的变化。采用电阻加热，升温速率 10 ℃/s，保温 时间 10 min，压缩结束后及时水淬，以保留高温变 形后的组织。变形温度为 860 ℃、890 ℃、920 ℃、 950 ℃、980 ℃，应变速率为 0.001 s–1、0.01 s–1、0.1 s–1 和 1 s–1，最大变形量为 70%。压缩后的试样用线切 割沿轴向切开后进行金相制样，试样用 Kroll 腐蚀