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β Ti–47Zr–5Al–3V alloy
Journal of Alloys and Compounds 665 (2016) 1e6
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effect of heat treatment on the microstructure and tensile properties of deformed a/b Tie47Zre5Ale3V alloy Yindong Shi a, b, *, Guosheng Zhang b, Ming Li b, Defeng Guo b, Zhixiao Zhang a, Bingning Wei b, Jingtao Li b, Xiangyi Zhang b, ** a b
Equipment Manufacturing College, Hebei University of Engineering, Handan, 056038, PR China State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, 066004, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 13 September 2015 Received in revised form 16 November 2015 Accepted 2 January 2016 Available online 5 January 2016
Microstructure and tensile properties of a new a/b Tie47Zre5Ale3V (wt%) alloy are investigated in the present study. Both a/b and b thermal annealing at 650 Ce800 C for 1 h and subsequent aging at 450 C for 2 h are introduced to investigate the effect of thermal annealing and aging on the microstructure and tensile properties of the Tie47Zre5Ale3V alloy after cold rolling. After thermal annealing, a fully equiaxed microstructure, a bi-modal microstructure and a fully lamellar microstructure are obtained with increasing the annealing temperature (650e800 C), respectively. The 450 C/2 h aging leads to further precipitation of extremely fine a platelets from the annealed samples. With increasing the annealing temperature, the tensile strength (sb) of annealed samples increases gradually from ~1350 to 1420 MPa, whereas the elongation to failure (εf) first increases and then decreases and the highest εf ~10.6% is attained at 750 C. For all annealed samples, the 450 C/2 h aging leads to increased strength and decreased ductility and the 750 C/1 h annealing plus 450 C/2 h aging results in the best combination of ultrahigh strength (sb ~1510 MPa) and good ductility (εf ~9.3%). The ultrahigh strength results from the strengthening of fine secondary a platelets, while the good ductility is attributed to the formation of large primary ap grains and the fine b grain size. © 2016 Elsevier B.V. All rights reserved.
Keywords: Severe plastic deformation Heat treatment Strength Ductility Titanium alloys
1. Introduction In the last decades, titanium alloys have attained very fast development and increasing attention due to their important applications in aerospace industry and others. Amongst titanium alloys, a/b two-phase alloys, e.g., Tie6Ale4V, have the most wide applications and cover currently more than 50% the application market share of titanium alloys [1]. This is due to their attractive physical and mechanical properties, such as low density, a wide range combination of strength, ductility and toughness as well as excellent corrosion resistance [1,2]. The mechanical properties of titanium alloys are mainly derived from the allotropic modification from high temperature bcc b phase to low temperature hcp a phase, i.e., the size, volume fraction, morphology and distribution of
* Corresponding author. Equipment Manufacturing College, Hebei University of Engineering, Handan, 056038, PR China ** Corresponding author. E-mail addresses: [email protected] (Y. Shi), [email protected] (X. Zhang). http://dx.doi.org/10.1016/j.jallcom.2016.01.023 0925-8388/© 2016 Elsevier B.V. All rights reserved.
precipitated a phase [3e8]. It is well known that the microstructure and mechanical properties of a/b titanium alloys is sensitive to the thermo-mechanical processing route [1], such as deformation temperature and degree, thermal annealing temperature and cooling rate as well as aging temperature and time. Therefore, the optimization of mechanical properties of a/b titanium alloys can be expected by employing appropriate thermo-mechanical processing procedures [3e6]. Tie47Zre5Ale3V (wt%) alloy is a new a/b high strength titanium alloy recently designed for aerospace applications [9e11]. A good combination of high strength and ductility can be achieved in the Tie47Zre5Ale3V alloy by engineering a hierarchicalnanolaminated structure that consists of primary ap large grains and fine a lamellae (bimodal structure) or of lamellae with different sizes, e.g., in width with nanometer and sub-micrometer scales [10,11]. However, as a newly designed alloy, the effect of cold deformation and heat treatment on the microstructure and mechanical properties of the Tie47Zre5Ale3V alloy has not been systematically studied and its microstructure-property relationship is not fully understood. Therefore, it is of significance to understand
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Y. Shi et al. / Journal of Alloys and Compounds 665 (2016) 1e6
the microstructure-property relationship in Tie47Zre5Ale3V alloy for its property optimization. In the present study, we investigate the evolution of microstructure and tensile properties of the alloy subjected to cold deformation followed by the a/b and b thermal annealing as well as subsequent aging treatment. 2. Experiment Tie47Zre5Ale3V (wt%) with a composition belonging to an a/b alloy was prepared by melting sponge Ti (99.7 wt%), sponge Zr (Zr þ Hf > 99.5 wt%), industrially pure Al (99.5 wt%) and V (99.9 wt %) using a ZHT-001 type vacuum non-consumable electro-arc furnace. After heat forging and rolling at 900 C in the b phase field (b transus temperature Tb ~789 C determined by DSC measurements), the sheets were solution treated (ST) at 850 C for 1 h and then quenched into water. As shown in Fig. 1 (a), the as-quenched samples exhibited a coarse b microstructure with an average grain size of ~300 mm and the fine lath-shaped structures of a00 martensite which was confirmed by the XRD analysis (see Fig. 1 (b)), indicating the martensite phase transformation of b / a00 took place during water quenching. The as-quenched samples were subjected to severe cold rolling at room temperature (RT) to over 90% reduction. Following cold rolling, heat treatments, i.e., thermal annealing and aging, were conducted. The cold rolled samples were firstly subjected to thermal annealing at 650e800 C for 1 h under vacuum and then air cooled to RT. Subsequently, the thermal annealed samples were aged at 450 C for 2 h under vacuum and then air cooled to RT. Microstructure and phase compositions of the Tie47Zre5Ale3V alloy subjected to cold rolling and heat treatments were characterized and determined using an optical metallography (OM), a transmission electron microscope (TEM) and an X-ray diffraction (XRD) with Cu Ka radiation. The b grain size was determined from approximating the equivalent sphere diameters of b grains by
Fig. 1. (a) OM microstructure and (b) XRD pattern of Tie47Zre5Ale3V alloy after solution treatment and water quenching.
direct measurement of their areas from over 30 OM or 50 TEM images using a digital micrograph analysis software. The statistical a grain size/lamellas width distributions were obtained from TEM measurements on more than 200 grains/lamellas. Tensile tests of samples with a gage dimension of 5 2.2 0.35 mm3 were performed at RT using an Instron 5948 machine under a constant cross-head speed of an initial strain rate of 1 103 s1. The strain was measured with a non-contacting video extensometer. For each condition, four specimens were used for tests, and the standard deviations of ultimate tensile strength and elongation to failure were less than 30 MPa and 1%, respectively. The tension direction was parallel to the rolling direction of the samples.
3. Results and discussion Fig. 2 (a) shows the tensile engineering stressestrain curves of the Tie47Zre5Ale3V alloy subjected to water quenching, cold rolling and heat treatments. The as-quenched sample (see curve 1 in Fig. 2 (a)) exhibits low yield strength (sy ~280 MPa) and tensile strength (sb ~945 MPa) as well as high elongation to failure (εf ~14.1%), and a so-called double yielding or strain plateau is also observed, which is associated with the stress-induced martensite transformation [12,13]. After cold rolling (see curve 2 in Fig. 2 (a)), the strength increases dramatically to sb ~1530 MPa accompanied with considerably reduced ductility of εf ~3.1%. Following cold rolling, thermal annealing is introduced, which enhances the ductility and decreases the strength of the samples (see curves 3e5 in Fig. 2 (a)). With increasing the thermal annealing temperature (650e800 C), the sb of thermally annealed samples increases
Fig. 2. Tensile engineering stressestrain curves of Tie47Zre5Ale3V alloy subjected to water quenching, cold rolling, thermal annealing and aging treatments. Curve 1, asquenched; curve 2, as-rolled; curve 3, 650 C/1 h; curve 4, 750 C/1 h; curve 5, 800 C/1 h; curve 6, 650 C/1 h þ 450 C/2 h; curve 7, 750 C/1 h þ 450 C/2 h; curve 8, 800 C/1 h þ 450 C/2 h.
Y. Shi et al. / Journal of Alloys and Compounds 665 (2016) 1e6
gradually from ~1340 to 1420 MPa, whereas the εf first increases and then decreases (see Fig. 2 (b)). The higher εf of ~10.6% is obtained in the 750 C/1 h annealed sample, compared with those (7.0% and 6.1%, respectively) of the 650 C/1 h and 800 C/1 h annealed samples (see Fig. 2 (b)). For all thermally annealed samples, the 450 C/2 h aging leads to increased strength and decreased ductility (see curves 6e8 in Fig. 2 (a)), and the 750 C/1 h thermal annealing plus 450 C/2 h aging results in the best combination of ultrahigh strength (sb ~1510 MPa) and good ductility (εf ~9.3%) (see curve 7 in Fig. 2 (a)). For revealing the factors governing the evolution of tensile properties of the Tie47Zre5Ale3V alloy subjected to cold rolling and heat treatments, the phase compositions and microstructures are investigated by XRD analyses and OM and TEM observations. XRD patterns of the Tie47Zre5Ale3V alloy subjected to water quenching, cold rolling and heat treatments are given in Figs. 1 (b) and 3. As shown in Fig. 1 (b), b phase and a00 martensite peaks coexist in the as-quenched sample, while only peaks corresponding to the a0 martensite are detected in the cold rolled sample (see Fig. 3 (a)), demonstrating that the martensite phase transformations of b
Fig. 3. XRD patterns of Tie47Zre5Ale3V alloy subjected to cold rolling, thermal annealing and aging treatments. (a) as-rolled, (b) 650 C/1 h, (c) 750 C/1 h, (d) 750 C/ 1 h þ 450 C/2 h, and (e) 800 C/1 h.
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/ a0 and a00 / a0 take place during cold rolling [13,14]. Furthermore, the a0 peaks widen evidently, indicating that high density dislocations are also introduced during cold rolling. The combined effect of deformation-induced martensite phase transformations (b / a0 and a00 / a0 ) and deformation-introduced high density dislocations leads to dramatically increased strength and reduced ductility of the cold rolled sample (see curve 2 in Fig. 2 (a)). After thermal annealing (see Fig. 3 (b), (c) and (e)), the phase compositions are characterized by the coexistence of a and b phases, indicating the reverse phase transformation of a0 / a þ b, and the volume fraction of b phase increases gradually with increasing the annealing temperature (650e800 C). For all the annealed samples, the subsequent 450 C/2 h aging treatment results in weakened b peaks and intensified a peaks (see Fig. 3 (d)), indicating the a phase precipitates from the b matrix (b / a þ b) [4,15,16], as confirmed by the TEM observation in the following section. Fig. 4 shows the optical micrographs and corresponding b grain size distributions of the Tie47Zre5Ale3V alloy subjected to cold rolling and heat treatments. After the thermal annealing at 750 C and 800 C for 1 h, both samples exhibit equiaxed b grains (see Fig. 4 (a) and (b)). However, the b grain size (~7 mm) of the 750 C/ 1 h annealed sample (see Fig. 4 (c)) is much smaller than that (~65 mm) of the 800 C/1 h annealed sample (see Fig. 4 (d)). This can be attributed to the formation of primary ap grains induced by a þ b thermal annealing (750 C/1 h) [14,15], which effectively resists the growth of b grains [1]. The refinement of the b grain microstructure from ~300 mm (as-quenched sample) to 7 mm (the 750 C/1 h annealed sample) is attributed to the processing of strain-induced martensite phase transformation during cold rolling and its reverse transformations during thermal annealing [17,18], which is an effective means of obtaining fine-grained b grain microstructure. For comparison, the b grain size distribution with an average grain size of ~200 nm of the 650 C/1 h annealed sample is shown in Fig. 4 (e). It is apparent that b grains of annealed samples coarsen dramatically from ~200 nm to ~65 mm with increasing the annealing temperature (650e800 C), which is a typical behavior for metals and alloys subjected to severe deformation and thermal annealing. Fig. 5 represents the TEM images and selected area electron diffraction (SAED) pattern of the Tie47Zre5Ale3V alloy subjected to cold rolling and heat treatments. After the thermal annealing at 650 C for 1 h (see Fig. 5 (a)), an ultrafine equiaxed microstructure is obtained with an average grain size of ~250 nm (see the bottom right in Fig. 5 (a)). The SAED pattern (see the top right in Fig. 5 (a)) confirms the coexistence of a and b phases consistent with the XRD results (see Fig. 3 (b)), comparatively smaller b grains (indicated by arrows) are formed at the a grain boundaries and detailed observations approve that a grains are uniformly mixed with b grains. The ultrafine grain size contributes to the high strength of sb ~1340 MPa, whereas limited ductility of εf ~7.0% is due to the low work hardening capacity of ultrafine-grained microstructure [19,20] (see curve 3 in Fig. 2 (a)). The formation of the ultrafine microstructure is also due to the processing of strain-induced martensite transformation and its reverse transformation [17,18]. Previous studies show that a fully equiaxed microstructure can be obtained in a/b titanium alloys when the cooling rate from the recrystallization annealing temperature is sufficiently low, only the primary ap grains will grow during the cooling process and no a lamellae are formed within the b grains, and this processing route will result in very coarse grain size up to several tens of micrometers [1]. However, through the use of a combination of severe cold rolling and appropriate thermal annealing, this study achieves an ultrafine (~200 nm) equiaxed microstructure in the Tie47Zre5Ale3V alloy, which is expected to lead to superior mechanical properties [21,22]. After the thermal annealing at 750 C
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Fig. 4. OM microstructures and b grain size distributions of Tie47Zre5Ale3V alloy subjected to cold rolling and thermal annealing at different temperatures for 1 h, (a) (c) 750 C and (b) (d) 800 C, (e) b grain size distributions obtained from TEM analysis for 650 C.
for 1 h (see Fig. 5 (b)), a typical bi-modal microstructure is formed that is composed of large microscale primary ap grains with an average grain size of ~1.4 mm (see the inset in Fig. 5 (b)) formed at prior grain boundaries and nanoscale secondary a platelets with an average lamellas width of ~75 nm (see the top section of Fig. 5 (e)). The nanoscale secondary a platelets result in the higher strength of sb ~1390 MPa than that (sb ~1340 MPa) of the 650 C/1 h annealed sample (see Fig. 2 (b)), while large primary ap grains stabilize the deformation and contribute to the high ductility of εf ~10.6% [14,15]. After the thermal annealing at 800 C for 1 h (see Fig. 5 (c)), a fully lamellar microstructure is produced with an average a lamellas width of ~70 nm (see the bottom section of Fig. 5 (e)). The fully fine a platelets lead to much higher strength of sb ~1420 MPa and lower ductility of εf ~6.1% (see curve 5 in Fig. 2 (a)), compared with those (sb ~1390 MPa and εf ~10.6%) of the sample with the bi-modal microstructure (see curve 4 in Fig. 2 (a)). Following the 750 C/ 1 h thermal annealing, the 450 C/2 h aging treatment leads to increasing precipitation of extremely fine secondary a platelets (indicated by the triangles in Fig. 5 (d)), which further enhance the strength from sb ~1420 MPa to sb ~1510 MPa with a little sacrifice of ductility (see curve 7 in Fig. 2 (a)). After tensile deformation of the
750 C/1 h thermal annealed plus the 450 C/2 h aged sample (see Fig. 5 (f)), high-density dislocations are accumulated in the primary ap grains and comparatively coarse a plates (indicated by arrows in Fig. 5 (f)), which contributes to the high work hardening capacity and the good ductility of εf ~9.3% (see curve 7 in Fig. 2 (a)). The b grain size, which is determined by the volume fraction and size of primary ap grains in a bi-modal microstructure, is an important factor governing the mechanical properties of a/b titanium alloys [1,2,5]. The a colony size is usually limited by the prior b grain size [23]. Compared with that of the 800 C/1 h annealed sample (see Fig. 4 (b) and (d)), the much finer b grain size (see Fig. 4 (a) and (c)) will lead to much smaller a colony size and slip length for the 750 C/1 h annealed sample, which is a parameter contributing to the higher ductility of the 750 C/1 h annealed sample. On another hand, fine b grain size can effectively limit the deteriorative effect of grain-boundary agb phase on mechanical properties, which is another parameter contributing to the good ductility of the 750 C/1 h annealed sample. In addition, the volume fraction of b phase is another important factor governing the mechanical properties and more residual ductile b phase contributes to better ductility [15,24]. Therefore, the less residual ductile b
Y. Shi et al. / Journal of Alloys and Compounds 665 (2016) 1e6
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Fig. 5. TEM images and primary ap grain size/secondary a lamellas width distributions of microstructures developed in Tie47Zre5Ale3V alloy subjected to cold rolling, thermal annealing and aging treatments. (a) 650 C/1 h (the insets shows the SAED pattern (top right) and grain size distribution (bottom right), respectively, and arrows indicate b grains), (b) 750 C/1 h (the inset shows the primary ap grain size distribution), (c) 800 C/1 h, (d) 750 C/1 h þ 450 C/2 h (triangles show the extremely fine a lamellas precipitated during the 450 C/2 h aging), (e) a lamellas width distributions corresponding to (b) (the top section) and (c) (the bottom section), and (f) microstructure after tensile deformation of (d) (arrows show the accumulated dislocations in coarse a plates).
phase is a negative factor contributing to the ductility of the 750 C/ 1 h annealed sample in comparison to the 800 C/1 h annealed sample (see Fig. 3 (c) and (e)). Apparently, the positive contributions of primary ap grains and fine b grain size overwhelm the negative effect of less ductile b phase. As a result, a much better ductility is obtained in the 750 C/1 h annealed sample, compared with that of the 800 C/1 h annealed sample. Previous study demonstrates that the b grain size has limited growth after aging
treatments at 440e560 C temperatures for 8 h [25]. Therefore, the growth of b grain size during 450 C/2 h aging treatment can be negligible in this study and the beneficial effect of fine b grain size on the good ductility is also suitable for the 750 C/1 h annealed plus 450 C/2 h aged sample (see curve 7 in Fig. 2 (a)). However, the 450 C/2 h aging treatment leads to reduced b phase (see Fig. 3 (d)), which can be the reason for the decreases in the ductility after aging treatment (see curve 7 in Fig. 2 (a)). Furthermore, it has been
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reported that the finer b grain size can result in finer a precipitates for a/b titanium alloys [23]. However, the effect of b grain size on the size distribution of secondary a platelets is not very apparent for the Tie47Zre5Ale3V alloy subjected to cold rolling and thermal annealing (see Fig. 5 (e)). This may be attributed to the fact that the cooling rate from the recrystallization temperature is the most important factor determining the width of a lamellae for a/b titanium alloys [1]. It is reasonable to infer that the cooling rates after the 750 C/1 h annealing and the 800 C/1 h annealing are nearly the same. As a result, the little difference of a lamellas width distributions of the 750 C/1 h annealed and the 800 C/1 h annealed samples is formed (see Fig. 5 (e)). 4. Conclusion In the present study, the microstructure characteristics and tensile properties of a new a/b Tie47Zre5Ale3V alloy subjected to cold rolling and various heat treatments were investigated systematically. The cold rolled samples was thermally annealed below (650e750 C) and above (800 C) the b transus temperature, and then aged at 450 C. The main results are summarized as follows: (1) The martensite phase transformations of b / a0 and a00 / a0 and the introduction of high density dislocations have been induced by severe cold rolling of the Tie47Zre5Ale3V alloy, which improves dramatically the strength (sb ~1540 MPa) and reduces the ductility (εf ~3.0%). (2) A fully equiaxed microstructure with ultrafine grains (~250 nm), a bi-modal microstructure composed of large microscale ap grains (~1.4 mm) and nanoscale secondary a platelets (~65 nm) as well as a fully lamellar microstructure with nanoscale a platelets (~70 nm) have been obtained after thermal annealing, respectively, with increasing the thermal annealing temperature (650e800 C). Further precipitation of extremely fine a platelets has taken place after the 450 C/ 2 h aging treatment for all annealed samples. (3) The tensile strength of the thermally annealed Tie47Zre5Ale3V alloy increases gradually from sb ~1340 MPa to sb ~1420 MPa with increasing the thermal annealing temperature (650e800 C), while the highest ductility (εf ~10.6%) is achieved with the thermal annealing temperature of 750 C, much better than those (εf ~7.0% and εf ~6.1%, respectively) of the samples thermally annealed at both 650 C and 800 C. The 450 C/2 h leads to enhanced strength and reduced ductility for all thermally annealed samples.
(4) The best combination of ultrahigh strength of sb ~1510 MPa and good ductility of εf ~9.3% is obtained after the 750 C/1 h thermal annealing plus the 450 C/2 h aging treatments. The ultrahigh strength results from the strengthening of fine secondary a platelets, while the good ductility is attributed to the formation of large microscale ap grains and fine b grain microstructure.
Acknowledgments The authors gratefully acknowledge the financial support of the National Natural Science Foundation of China (Nos. 51031004, 50871095 and 50821001).
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G. Lütjering, J.C. Williams, Titanium, second ed., Springer, Berlin, 2007. D. Banerjee, J.C. Williams, Acta Mater. 61 (2013) 844e879. P. Schmidt, A. El-Chaikh, H.J. Christ, Metall. Trans. A 42 (2011) 2652e2667. O.M. Ivasishin, P.E. Markovsky, Yu.V. Matviychuk, S.L. Semiatin, C.H. Ward, S. Fox, J. Alloys Comp. 457 (2008) 296e309. I. Weiss, S.L. Semiatin, Mater. Sci. Eng. A 243 (1998) 46e65. A. Bhattacharjee, V.K. Varma, S.V. Kamat, A.K. Gogia, S. Bhargava, Metall. Trans. A 37 (2006) 1423e1433. R. Chait, Metall. Trans. A 6 (1975) 2301e2303. J.C. Chesnutt, F.H. Froes, Metall. Trans. A 8 (1977) 1013e1017. S.X. Liang, M.Z. Ma, R. Jing, C.L. Tan, R.P. Liu, Mater. Sci. Eng. A 541 (2012) 67e72. Z. Zhang, G. Zhang, D. Guo, M. Li, Y. Shi, X. Li, X. Zhang, Mater. Lett. 131 (2014) 240e243. M. Li, D. Guo, T. Ma, Y. Shi, G. Zhang, X. Zhang, Mater. Sci. Eng. A 606 (2014) 396e400. T.W. Duerig, J. Albrecht, D. Richter, P. Fischer, Acta Metall. 30 (1982) 2161e2172. M.K. Koul, J.F. Breedis, Acta Metall. 18 (1970) 579e588. Z. Zhang, M. Li, D. Guo, Y. Shi, X. Zhang, H.-E. Schaefer, Mater. Sci. Eng. A 594 (2014) 321e323. Y. Shi, M. Li, D. Guo, G. Zhang, Q. Zhang, X. Li, X. Zhang, Mater. Sci. Eng. A 615 (2014) 464e468. O.M. Ivasishin, P.E. Markovsky, S.L. Semiatin, C.H. Ward, Mater. Sci. Eng. A 405 (2005) 296e305. W. Xu, X. Wu, M. Calin, M. Stoica, J. Eckert, K. Xia, Scr. Mater. (2009) 1012e1015. M.-H. Cai, C.-Y. Lee, S. Kang, Y.-K. Lee, Scr. Mater. (2011) 1098e1101. Y.T. Zhu, X. Liao, Nat. Mater. 3 (2004) 351e352. D. Jia, Y.M. Wang, K.T. Ramesh, E. Ma, Y.T. Zhu, R.Z. Valiev, Appl. Phys. Lett. 79 (2001) 611e613. S. Ankem, C.A. Green, Mater. Sci. Eng. A 263 (1999) 127e131. H. Nishizawa, E. Sudekai, W. Liu, H. Hashimoto, Mater. Trans. JIM 36 (1995) 924e927. G. Lutjering, Mater. Sci. Eng. A 243 (1998) 32e45. ment, A. Lenain, P.J. Jacques, JOM 65 (2007) 50e53. N. Cle Y. Chen, Z. Du, S. Xiao, L. Xu, J. Tian, J. Alloys Comp. 586 (2014) 588e592.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effect of heat treatment on the microstructure and tensile properties of deformed a/b Tie47Zre5Ale3V alloy Yindong Shi a, b, *, Guosheng Zhang b, Ming Li b, Defeng Guo b, Zhixiao Zhang a, Bingning Wei b, Jingtao Li b, Xiangyi Zhang b, ** a b
Equipment Manufacturing College, Hebei University of Engineering, Handan, 056038, PR China State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, 066004, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 13 September 2015 Received in revised form 16 November 2015 Accepted 2 January 2016 Available online 5 January 2016
Microstructure and tensile properties of a new a/b Tie47Zre5Ale3V (wt%) alloy are investigated in the present study. Both a/b and b thermal annealing at 650 Ce800 C for 1 h and subsequent aging at 450 C for 2 h are introduced to investigate the effect of thermal annealing and aging on the microstructure and tensile properties of the Tie47Zre5Ale3V alloy after cold rolling. After thermal annealing, a fully equiaxed microstructure, a bi-modal microstructure and a fully lamellar microstructure are obtained with increasing the annealing temperature (650e800 C), respectively. The 450 C/2 h aging leads to further precipitation of extremely fine a platelets from the annealed samples. With increasing the annealing temperature, the tensile strength (sb) of annealed samples increases gradually from ~1350 to 1420 MPa, whereas the elongation to failure (εf) first increases and then decreases and the highest εf ~10.6% is attained at 750 C. For all annealed samples, the 450 C/2 h aging leads to increased strength and decreased ductility and the 750 C/1 h annealing plus 450 C/2 h aging results in the best combination of ultrahigh strength (sb ~1510 MPa) and good ductility (εf ~9.3%). The ultrahigh strength results from the strengthening of fine secondary a platelets, while the good ductility is attributed to the formation of large primary ap grains and the fine b grain size. © 2016 Elsevier B.V. All rights reserved.
Keywords: Severe plastic deformation Heat treatment Strength Ductility Titanium alloys
1. Introduction In the last decades, titanium alloys have attained very fast development and increasing attention due to their important applications in aerospace industry and others. Amongst titanium alloys, a/b two-phase alloys, e.g., Tie6Ale4V, have the most wide applications and cover currently more than 50% the application market share of titanium alloys [1]. This is due to their attractive physical and mechanical properties, such as low density, a wide range combination of strength, ductility and toughness as well as excellent corrosion resistance [1,2]. The mechanical properties of titanium alloys are mainly derived from the allotropic modification from high temperature bcc b phase to low temperature hcp a phase, i.e., the size, volume fraction, morphology and distribution of
* Corresponding author. Equipment Manufacturing College, Hebei University of Engineering, Handan, 056038, PR China ** Corresponding author. E-mail addresses: [email protected] (Y. Shi), [email protected] (X. Zhang). http://dx.doi.org/10.1016/j.jallcom.2016.01.023 0925-8388/© 2016 Elsevier B.V. All rights reserved.
precipitated a phase [3e8]. It is well known that the microstructure and mechanical properties of a/b titanium alloys is sensitive to the thermo-mechanical processing route [1], such as deformation temperature and degree, thermal annealing temperature and cooling rate as well as aging temperature and time. Therefore, the optimization of mechanical properties of a/b titanium alloys can be expected by employing appropriate thermo-mechanical processing procedures [3e6]. Tie47Zre5Ale3V (wt%) alloy is a new a/b high strength titanium alloy recently designed for aerospace applications [9e11]. A good combination of high strength and ductility can be achieved in the Tie47Zre5Ale3V alloy by engineering a hierarchicalnanolaminated structure that consists of primary ap large grains and fine a lamellae (bimodal structure) or of lamellae with different sizes, e.g., in width with nanometer and sub-micrometer scales [10,11]. However, as a newly designed alloy, the effect of cold deformation and heat treatment on the microstructure and mechanical properties of the Tie47Zre5Ale3V alloy has not been systematically studied and its microstructure-property relationship is not fully understood. Therefore, it is of significance to understand
2
Y. Shi et al. / Journal of Alloys and Compounds 665 (2016) 1e6
the microstructure-property relationship in Tie47Zre5Ale3V alloy for its property optimization. In the present study, we investigate the evolution of microstructure and tensile properties of the alloy subjected to cold deformation followed by the a/b and b thermal annealing as well as subsequent aging treatment. 2. Experiment Tie47Zre5Ale3V (wt%) with a composition belonging to an a/b alloy was prepared by melting sponge Ti (99.7 wt%), sponge Zr (Zr þ Hf > 99.5 wt%), industrially pure Al (99.5 wt%) and V (99.9 wt %) using a ZHT-001 type vacuum non-consumable electro-arc furnace. After heat forging and rolling at 900 C in the b phase field (b transus temperature Tb ~789 C determined by DSC measurements), the sheets were solution treated (ST) at 850 C for 1 h and then quenched into water. As shown in Fig. 1 (a), the as-quenched samples exhibited a coarse b microstructure with an average grain size of ~300 mm and the fine lath-shaped structures of a00 martensite which was confirmed by the XRD analysis (see Fig. 1 (b)), indicating the martensite phase transformation of b / a00 took place during water quenching. The as-quenched samples were subjected to severe cold rolling at room temperature (RT) to over 90% reduction. Following cold rolling, heat treatments, i.e., thermal annealing and aging, were conducted. The cold rolled samples were firstly subjected to thermal annealing at 650e800 C for 1 h under vacuum and then air cooled to RT. Subsequently, the thermal annealed samples were aged at 450 C for 2 h under vacuum and then air cooled to RT. Microstructure and phase compositions of the Tie47Zre5Ale3V alloy subjected to cold rolling and heat treatments were characterized and determined using an optical metallography (OM), a transmission electron microscope (TEM) and an X-ray diffraction (XRD) with Cu Ka radiation. The b grain size was determined from approximating the equivalent sphere diameters of b grains by
Fig. 1. (a) OM microstructure and (b) XRD pattern of Tie47Zre5Ale3V alloy after solution treatment and water quenching.
direct measurement of their areas from over 30 OM or 50 TEM images using a digital micrograph analysis software. The statistical a grain size/lamellas width distributions were obtained from TEM measurements on more than 200 grains/lamellas. Tensile tests of samples with a gage dimension of 5 2.2 0.35 mm3 were performed at RT using an Instron 5948 machine under a constant cross-head speed of an initial strain rate of 1 103 s1. The strain was measured with a non-contacting video extensometer. For each condition, four specimens were used for tests, and the standard deviations of ultimate tensile strength and elongation to failure were less than 30 MPa and 1%, respectively. The tension direction was parallel to the rolling direction of the samples.
3. Results and discussion Fig. 2 (a) shows the tensile engineering stressestrain curves of the Tie47Zre5Ale3V alloy subjected to water quenching, cold rolling and heat treatments. The as-quenched sample (see curve 1 in Fig. 2 (a)) exhibits low yield strength (sy ~280 MPa) and tensile strength (sb ~945 MPa) as well as high elongation to failure (εf ~14.1%), and a so-called double yielding or strain plateau is also observed, which is associated with the stress-induced martensite transformation [12,13]. After cold rolling (see curve 2 in Fig. 2 (a)), the strength increases dramatically to sb ~1530 MPa accompanied with considerably reduced ductility of εf ~3.1%. Following cold rolling, thermal annealing is introduced, which enhances the ductility and decreases the strength of the samples (see curves 3e5 in Fig. 2 (a)). With increasing the thermal annealing temperature (650e800 C), the sb of thermally annealed samples increases
Fig. 2. Tensile engineering stressestrain curves of Tie47Zre5Ale3V alloy subjected to water quenching, cold rolling, thermal annealing and aging treatments. Curve 1, asquenched; curve 2, as-rolled; curve 3, 650 C/1 h; curve 4, 750 C/1 h; curve 5, 800 C/1 h; curve 6, 650 C/1 h þ 450 C/2 h; curve 7, 750 C/1 h þ 450 C/2 h; curve 8, 800 C/1 h þ 450 C/2 h.
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gradually from ~1340 to 1420 MPa, whereas the εf first increases and then decreases (see Fig. 2 (b)). The higher εf of ~10.6% is obtained in the 750 C/1 h annealed sample, compared with those (7.0% and 6.1%, respectively) of the 650 C/1 h and 800 C/1 h annealed samples (see Fig. 2 (b)). For all thermally annealed samples, the 450 C/2 h aging leads to increased strength and decreased ductility (see curves 6e8 in Fig. 2 (a)), and the 750 C/1 h thermal annealing plus 450 C/2 h aging results in the best combination of ultrahigh strength (sb ~1510 MPa) and good ductility (εf ~9.3%) (see curve 7 in Fig. 2 (a)). For revealing the factors governing the evolution of tensile properties of the Tie47Zre5Ale3V alloy subjected to cold rolling and heat treatments, the phase compositions and microstructures are investigated by XRD analyses and OM and TEM observations. XRD patterns of the Tie47Zre5Ale3V alloy subjected to water quenching, cold rolling and heat treatments are given in Figs. 1 (b) and 3. As shown in Fig. 1 (b), b phase and a00 martensite peaks coexist in the as-quenched sample, while only peaks corresponding to the a0 martensite are detected in the cold rolled sample (see Fig. 3 (a)), demonstrating that the martensite phase transformations of b
Fig. 3. XRD patterns of Tie47Zre5Ale3V alloy subjected to cold rolling, thermal annealing and aging treatments. (a) as-rolled, (b) 650 C/1 h, (c) 750 C/1 h, (d) 750 C/ 1 h þ 450 C/2 h, and (e) 800 C/1 h.
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/ a0 and a00 / a0 take place during cold rolling [13,14]. Furthermore, the a0 peaks widen evidently, indicating that high density dislocations are also introduced during cold rolling. The combined effect of deformation-induced martensite phase transformations (b / a0 and a00 / a0 ) and deformation-introduced high density dislocations leads to dramatically increased strength and reduced ductility of the cold rolled sample (see curve 2 in Fig. 2 (a)). After thermal annealing (see Fig. 3 (b), (c) and (e)), the phase compositions are characterized by the coexistence of a and b phases, indicating the reverse phase transformation of a0 / a þ b, and the volume fraction of b phase increases gradually with increasing the annealing temperature (650e800 C). For all the annealed samples, the subsequent 450 C/2 h aging treatment results in weakened b peaks and intensified a peaks (see Fig. 3 (d)), indicating the a phase precipitates from the b matrix (b / a þ b) [4,15,16], as confirmed by the TEM observation in the following section. Fig. 4 shows the optical micrographs and corresponding b grain size distributions of the Tie47Zre5Ale3V alloy subjected to cold rolling and heat treatments. After the thermal annealing at 750 C and 800 C for 1 h, both samples exhibit equiaxed b grains (see Fig. 4 (a) and (b)). However, the b grain size (~7 mm) of the 750 C/ 1 h annealed sample (see Fig. 4 (c)) is much smaller than that (~65 mm) of the 800 C/1 h annealed sample (see Fig. 4 (d)). This can be attributed to the formation of primary ap grains induced by a þ b thermal annealing (750 C/1 h) [14,15], which effectively resists the growth of b grains [1]. The refinement of the b grain microstructure from ~300 mm (as-quenched sample) to 7 mm (the 750 C/1 h annealed sample) is attributed to the processing of strain-induced martensite phase transformation during cold rolling and its reverse transformations during thermal annealing [17,18], which is an effective means of obtaining fine-grained b grain microstructure. For comparison, the b grain size distribution with an average grain size of ~200 nm of the 650 C/1 h annealed sample is shown in Fig. 4 (e). It is apparent that b grains of annealed samples coarsen dramatically from ~200 nm to ~65 mm with increasing the annealing temperature (650e800 C), which is a typical behavior for metals and alloys subjected to severe deformation and thermal annealing. Fig. 5 represents the TEM images and selected area electron diffraction (SAED) pattern of the Tie47Zre5Ale3V alloy subjected to cold rolling and heat treatments. After the thermal annealing at 650 C for 1 h (see Fig. 5 (a)), an ultrafine equiaxed microstructure is obtained with an average grain size of ~250 nm (see the bottom right in Fig. 5 (a)). The SAED pattern (see the top right in Fig. 5 (a)) confirms the coexistence of a and b phases consistent with the XRD results (see Fig. 3 (b)), comparatively smaller b grains (indicated by arrows) are formed at the a grain boundaries and detailed observations approve that a grains are uniformly mixed with b grains. The ultrafine grain size contributes to the high strength of sb ~1340 MPa, whereas limited ductility of εf ~7.0% is due to the low work hardening capacity of ultrafine-grained microstructure [19,20] (see curve 3 in Fig. 2 (a)). The formation of the ultrafine microstructure is also due to the processing of strain-induced martensite transformation and its reverse transformation [17,18]. Previous studies show that a fully equiaxed microstructure can be obtained in a/b titanium alloys when the cooling rate from the recrystallization annealing temperature is sufficiently low, only the primary ap grains will grow during the cooling process and no a lamellae are formed within the b grains, and this processing route will result in very coarse grain size up to several tens of micrometers [1]. However, through the use of a combination of severe cold rolling and appropriate thermal annealing, this study achieves an ultrafine (~200 nm) equiaxed microstructure in the Tie47Zre5Ale3V alloy, which is expected to lead to superior mechanical properties [21,22]. After the thermal annealing at 750 C
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Fig. 4. OM microstructures and b grain size distributions of Tie47Zre5Ale3V alloy subjected to cold rolling and thermal annealing at different temperatures for 1 h, (a) (c) 750 C and (b) (d) 800 C, (e) b grain size distributions obtained from TEM analysis for 650 C.
for 1 h (see Fig. 5 (b)), a typical bi-modal microstructure is formed that is composed of large microscale primary ap grains with an average grain size of ~1.4 mm (see the inset in Fig. 5 (b)) formed at prior grain boundaries and nanoscale secondary a platelets with an average lamellas width of ~75 nm (see the top section of Fig. 5 (e)). The nanoscale secondary a platelets result in the higher strength of sb ~1390 MPa than that (sb ~1340 MPa) of the 650 C/1 h annealed sample (see Fig. 2 (b)), while large primary ap grains stabilize the deformation and contribute to the high ductility of εf ~10.6% [14,15]. After the thermal annealing at 800 C for 1 h (see Fig. 5 (c)), a fully lamellar microstructure is produced with an average a lamellas width of ~70 nm (see the bottom section of Fig. 5 (e)). The fully fine a platelets lead to much higher strength of sb ~1420 MPa and lower ductility of εf ~6.1% (see curve 5 in Fig. 2 (a)), compared with those (sb ~1390 MPa and εf ~10.6%) of the sample with the bi-modal microstructure (see curve 4 in Fig. 2 (a)). Following the 750 C/ 1 h thermal annealing, the 450 C/2 h aging treatment leads to increasing precipitation of extremely fine secondary a platelets (indicated by the triangles in Fig. 5 (d)), which further enhance the strength from sb ~1420 MPa to sb ~1510 MPa with a little sacrifice of ductility (see curve 7 in Fig. 2 (a)). After tensile deformation of the
750 C/1 h thermal annealed plus the 450 C/2 h aged sample (see Fig. 5 (f)), high-density dislocations are accumulated in the primary ap grains and comparatively coarse a plates (indicated by arrows in Fig. 5 (f)), which contributes to the high work hardening capacity and the good ductility of εf ~9.3% (see curve 7 in Fig. 2 (a)). The b grain size, which is determined by the volume fraction and size of primary ap grains in a bi-modal microstructure, is an important factor governing the mechanical properties of a/b titanium alloys [1,2,5]. The a colony size is usually limited by the prior b grain size [23]. Compared with that of the 800 C/1 h annealed sample (see Fig. 4 (b) and (d)), the much finer b grain size (see Fig. 4 (a) and (c)) will lead to much smaller a colony size and slip length for the 750 C/1 h annealed sample, which is a parameter contributing to the higher ductility of the 750 C/1 h annealed sample. On another hand, fine b grain size can effectively limit the deteriorative effect of grain-boundary agb phase on mechanical properties, which is another parameter contributing to the good ductility of the 750 C/1 h annealed sample. In addition, the volume fraction of b phase is another important factor governing the mechanical properties and more residual ductile b phase contributes to better ductility [15,24]. Therefore, the less residual ductile b
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Fig. 5. TEM images and primary ap grain size/secondary a lamellas width distributions of microstructures developed in Tie47Zre5Ale3V alloy subjected to cold rolling, thermal annealing and aging treatments. (a) 650 C/1 h (the insets shows the SAED pattern (top right) and grain size distribution (bottom right), respectively, and arrows indicate b grains), (b) 750 C/1 h (the inset shows the primary ap grain size distribution), (c) 800 C/1 h, (d) 750 C/1 h þ 450 C/2 h (triangles show the extremely fine a lamellas precipitated during the 450 C/2 h aging), (e) a lamellas width distributions corresponding to (b) (the top section) and (c) (the bottom section), and (f) microstructure after tensile deformation of (d) (arrows show the accumulated dislocations in coarse a plates).
phase is a negative factor contributing to the ductility of the 750 C/ 1 h annealed sample in comparison to the 800 C/1 h annealed sample (see Fig. 3 (c) and (e)). Apparently, the positive contributions of primary ap grains and fine b grain size overwhelm the negative effect of less ductile b phase. As a result, a much better ductility is obtained in the 750 C/1 h annealed sample, compared with that of the 800 C/1 h annealed sample. Previous study demonstrates that the b grain size has limited growth after aging
treatments at 440e560 C temperatures for 8 h [25]. Therefore, the growth of b grain size during 450 C/2 h aging treatment can be negligible in this study and the beneficial effect of fine b grain size on the good ductility is also suitable for the 750 C/1 h annealed plus 450 C/2 h aged sample (see curve 7 in Fig. 2 (a)). However, the 450 C/2 h aging treatment leads to reduced b phase (see Fig. 3 (d)), which can be the reason for the decreases in the ductility after aging treatment (see curve 7 in Fig. 2 (a)). Furthermore, it has been
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reported that the finer b grain size can result in finer a precipitates for a/b titanium alloys [23]. However, the effect of b grain size on the size distribution of secondary a platelets is not very apparent for the Tie47Zre5Ale3V alloy subjected to cold rolling and thermal annealing (see Fig. 5 (e)). This may be attributed to the fact that the cooling rate from the recrystallization temperature is the most important factor determining the width of a lamellae for a/b titanium alloys [1]. It is reasonable to infer that the cooling rates after the 750 C/1 h annealing and the 800 C/1 h annealing are nearly the same. As a result, the little difference of a lamellas width distributions of the 750 C/1 h annealed and the 800 C/1 h annealed samples is formed (see Fig. 5 (e)). 4. Conclusion In the present study, the microstructure characteristics and tensile properties of a new a/b Tie47Zre5Ale3V alloy subjected to cold rolling and various heat treatments were investigated systematically. The cold rolled samples was thermally annealed below (650e750 C) and above (800 C) the b transus temperature, and then aged at 450 C. The main results are summarized as follows: (1) The martensite phase transformations of b / a0 and a00 / a0 and the introduction of high density dislocations have been induced by severe cold rolling of the Tie47Zre5Ale3V alloy, which improves dramatically the strength (sb ~1540 MPa) and reduces the ductility (εf ~3.0%). (2) A fully equiaxed microstructure with ultrafine grains (~250 nm), a bi-modal microstructure composed of large microscale ap grains (~1.4 mm) and nanoscale secondary a platelets (~65 nm) as well as a fully lamellar microstructure with nanoscale a platelets (~70 nm) have been obtained after thermal annealing, respectively, with increasing the thermal annealing temperature (650e800 C). Further precipitation of extremely fine a platelets has taken place after the 450 C/ 2 h aging treatment for all annealed samples. (3) The tensile strength of the thermally annealed Tie47Zre5Ale3V alloy increases gradually from sb ~1340 MPa to sb ~1420 MPa with increasing the thermal annealing temperature (650e800 C), while the highest ductility (εf ~10.6%) is achieved with the thermal annealing temperature of 750 C, much better than those (εf ~7.0% and εf ~6.1%, respectively) of the samples thermally annealed at both 650 C and 800 C. The 450 C/2 h leads to enhanced strength and reduced ductility for all thermally annealed samples.
(4) The best combination of ultrahigh strength of sb ~1510 MPa and good ductility of εf ~9.3% is obtained after the 750 C/1 h thermal annealing plus the 450 C/2 h aging treatments. The ultrahigh strength results from the strengthening of fine secondary a platelets, while the good ductility is attributed to the formation of large microscale ap grains and fine b grain microstructure.
Acknowledgments The authors gratefully acknowledge the financial support of the National Natural Science Foundation of China (Nos. 51031004, 50871095 and 50821001).
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