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Plastic deformation mechanism of Ti–Nb–Ta–Zr–O alloy at cryogenic temperatures
Materials Science & Engineering A 765 (2019) 138293
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Plastic deformation mechanism of Ti–Nb–Ta–Zr–O alloy at cryogenic temperatures
T
Wei-dong Zhanga, Zhenggang Wua, Yong Liub, Hongbin Beic, Bin Liub, Jingwen Qiud,∗ a
College of Materials Science and Engineering, Hunan University, Changsha, 410082, PR China State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, PR China c Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN, 37831, USA d Hunan Provincial Key Laboratory of Advanced Materials for New Energy Storage and Conversion, Hunan University of Science and Technology, Xiangtan, 411201, PR China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Gum metal Cryogenic mechanical properties Twinning β′ phase ω phase
The deformation behavior and mechanism of Ti–36Nb–2Ta–3Zr-0.35O alloy at room and cryogenic temperatures (77 and 200K) were investigated. The microstructure near the fracture was observed with electron backscattered diffraction analysis (EBSD) and transmission electron microscopy (TEM). The results show that the plastic deformation behavior of the alloy is sensitive to test temperature and cold deformation history of the alloy. For the alloy without cold deformation, both tensile strength and ductility have a significant enhancement at cryogenic temperatures, mainly due to the occurrence of mechanical twinning. Moreover, with the decrease of the temperature, the density of mechanical twinning increases. For the alloy cold swaged by 85%, an increase in strength accompanied by a decrease in uniform elongation was found at cryogenic temperatures. The unique temperature dependence of the deformation behavior was due to the low stability of β phase. At 200K, ω phase precipitated in β phase; while at 77K, an interesting transformation (β phase to β′ phase) was observed.
1. Introduction The Ti–36Nb–2Ta–3Zr-0.35O (wt.%) alloy (also known as gum metal) has obtained increased interests in recent years due to its novel properties, such as high strength, low modulus, superelasticity, superplasticity, Invar-like thermal expansion, and Elinvar-like thermal dependence of the elastic modulus at room temperature [1–10]. These novel properties were achieved by selecting a unique composition that simultaneously satisfies three electronic parameters: electron per atom ratio (e/a of about 4.24), bond order (Bo of about 2.87) and d electron orbital energy level (Md of about 2.45 eV). Moreover, a certain plastic deformation in the preparation process is another major requisite for the above novel properties of gum metal. After Saito et al. developed gum metal in 2003 [1], there are extensive researches to elucidate the deformation mechanism of gum metal. Saito et al. [1,11,12] claimed that the deformation mechanisms of gum metal is different from the typical titanium alloy. During the cold deformation, no dislocation movement, twinning and phase transformation take place, and instead, the unstable β lattice readily forms giant faults by ideal shear, thus accommodates the plastic strain. However, the deformation mechanism of gum metal is still
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controversial. Stress-induced phase transformation [13–17] and deformation induced twinning [16–19] have been reported to be active in a deformed gum metal by some other researches. Talling et al. [13,14] reported that the superelastic behavior of gum metal is associated with the reversible stress induced martensitic transformation (β → α"). Furthermore, cold work can result in the formation of large amount of stress-induced ω phase [6,15–17], and the formation of ω phase probably plays an important role in increasing the tensile strength. In the research of S. Shin et al. [19], mechanical twins were detected after thermomechanical-cycling processes applied to gum metal, resulting in a twinning-induced softening effect (about 18% decrease in microhardness). Besides, Yang et al. [20] found that the dominant deformation mechanisms are related to not only the extent of deformation but also the stability of the β phase. The research of S. Shin et al. [21] also proved that the β-phase stability dependence of deformation mode has a decisive significance effect on the mechanical properties of gum metal. Moreover, to further understand the deformation mechanism of gum metal, some other experiments under extreme conditions were performed, for example, the deformation at low temperatures (below 273K). Saito et al. [1] reported that the elastic deformability of the
Corresponding author. E-mail address: [email protected] (J. Qiu).
https://doi.org/10.1016/j.msea.2019.138293 Received 4 June 2019; Received in revised form 22 July 2019; Accepted 13 August 2019 Available online 14 August 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 765 (2019) 138293
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displacement data were presented in the insets in Fig. 1. To compute the plastic components of the strain, a line was fit to the linear elastic portion of the stress-strain curves and the amount of the elastic strain at a given point on the curve was subtracted from the total strain to obtain the plastic strain. And in general, with decreasing temperature, the flow stress increases and the stress-strain curves shift up. Besides, EP0.85 exhibits shorter necking regions (from the point of maximum stress to the point of fracture) on its 77 and 200K stress-strain curves. Fig. 2 summarizes the 0.2% offset yield strength, ultimate tensile strengths and uniform elongation to fracture. The individual data points in these figures are average of three tensile tests. he strength values for both alloys show similar temperature-dependent trend. For both alloys, the highest values of the yield strength (YS) and ultimate tensile strength (UTS) are both obtained at 77 K. An increase in temperature results in a monotonic decrease in both the yield and ultimate tensile strengths. It should be noted, however, that the strength of the EP0.85 is much larger than that of the EP0.0 at the same temperature. In addition, the uniform elongation values for both alloys show opposite temperature-dependent trend. For the EP0.0, the uniform elongation increases with the decrease of temperature, while for the EP0.85, the uniform elongation decreases with the decrease of temperature. Due to the absent of strain extensometer in the tensile tests at cryogenic temperatures, it is difficult to calculate accurately the work hardening exponent. Therefore, the difference between UTS and YS is used to qualitatively describe the work hardening capability. Temperature dependent extent work hardening of EP0.0 and EP0.85 was shown in Fig. 3. The work hardening capability (the difference between UTS and YS) of the EP0.0 also decreases with the increase of temperature. The largest difference between values of UTS and YS for the EP0.0 appears at 77 K. The work hardening capability of the EP0.85 is lower than that of EP0.0, and no clear temperature dependence was found. The unique temperature dependent strength and ductility of EP0.0 also reported in FCC metals and alloys with low stacking fault energy. At cryogenic temperatures, deformation twining occurs in these FCC materials, such as FeCoCrNi and FeCoCrNiMn equiatomic solid solution alloys [25,26], in which deformation induced twining can provide a higher work hardening capability and also contribute to the enhanced ductility. However, the stacking fault energy of EP0.0 is much higher than the critical stacking fault energy below which twining occurs (~45 mJ*m−2). For titanium alloys, the strength always increases with the decreasing temperature (below room temperature) and the elongation exhibits an opposite tendency, which is due to thermal influences on the dislocation motion [27,28]. In order to shed some light on this unique temperature dependent strength and ductility, the microstructures of EP0.0 and EP0.85 after tensile tests were characterized and analyzed. Fig. 4 shows the microstructures of tensile-fractured EP0.0 samples deformed at 77K, 200K and 273K. The microstructure at 77 K has a colony-like morphology which is full of parallel bands with various thickness and spacing (Fig. 4a). To identify details of these bands, a black line crossing the band was drawn and misorientations along the line were measured and plotted (Fig. 4d). It shows that whenever the line passes the individual band, there will be a ~53° misorientation, indicating that these bands are composed of Σ11{332}β 113β twins. The densities of the twin bands are largely reduced when the testing temperature is increased (as shown in Fig. 4b), and no twin bands are observed when the temperature is increased to 293 K (Fig. 4c). Besides, only Σ11{332}β 113β twins and no other types of twins are detected in the EP0.0 after the tensile tests at 200K and 77K. For EP 0.0, the enhanced yield strength (Fig. 1a) under lower temperatures (200K and 77 K) will make the critical stress that needs to be reached in order to activate deformation twinning easier to be achieved. Fig. 5 shows the TEM microstructure near tensile fracture of EP0.0 and EP0.85 deformed at room temperature. It can be seen that the ω phase formed in β phase (Fig. 5a). The morphology of ω plates (about
cold-worked gum metal increases with the decrease of temperature, exceeding 4% at 77K. Wei et al. [22] found that the yield stress increased monotonically with increasing the deformation temperature (193 K–233 K) for the 1.2–1.8 wt%O alloys. Tane et al. [23] reported that the formation of the ω phase and the low stability of the β phase were the dominant factors determining the elastic properties below the room temperature. However, these studies only focused on the elastic deformation behavior and properties of gum metal. So far, the plastic deformation behavior and mechanism of the gum metal at cryogenic temperatures has seldom been reported. In this work, the deformation behaviors of Ti–36Nb–2Ta–3Zr-0.35O alloy with and without cold deformation at different temperatures from 77K to 293K were investigated. And the purpose of this study is to understand the effect of cold deformation on the plastic deformation behaviors and mechanisms of Ti–36Nb–2Ta–3Zr-0.35O alloy at cryogenic temperatures. 2. Materials and methods 2.1. Material preparation The Ti–36Nb–2Ta–3Zr-0.35O alloy used in this work was prepared by powder metallurgy technique (cold isostatic pressing and vacuum sintering). Details of the alloy preparation and alloy composition can be found in our previous work [24]. In this work, the round bars after solution heat treatment at 1273K for 1h in argon (the sample was named as EP0.0) and after cold swaging by 85% (reduction in cross section area, the sample was named as EP0.85) were used. 2.2. Mechanical properties tests and material characterization Flat dog-bone-shaped specimens with a gage length of 9.5 mm were cut from the alloy bars. Tensile tests were performed using an Instron screw-driven testing machine at a strain rate of 10−4 s−1and temperatures of 77, 200, and 273 K. Three nominally identical specimens were tested at each temperature. The uniform elongation was obtained by a new method which has been reported by Wu et al. [25]. Vickers microhardness indents spaced 1 mm apart were made along the specimen gage lengths using a Vickers Hardness tester with a force of 50g. The change in the distance between adjacent indents, excluding the two indents on either side of the fracture plane, was used to calculate the uniform elongations to fracture. 2.3. Material characterization FEI Quanta 650 FEG-SEM equipped with EBSD system was used to characterize the microstructures of the tensile-fractured materials. Both of the EBSD samples and the TEM samples were cut from uniform deformation region of the tensile test sample. The electron backscatter diffraction (EBSD) measurements were performed with an electron probe current of 3.0 nA at an acceleration voltage of 20 kV. The qualification and analysis of the EBSD results were performed using Channel 5 software. EBSD maps are displayed as inverse pole figure (IPF) maps in the direction of tensile axis (for the tensile-deformed microstructure). The TEM samples were prepared by twin-jet electrochemical polishing technique in an etching solution that contains 10% sulfuric acid (HSO4) and 90% methanol (CH3OH) at 20V and 238K. Then, the foils were examined using a Titan G2 60-300Transmission Electron Microscope. 3. Results and discussion Fig. 1 shows the engineering stress-strain curves of EP0.0 and EP0.85. The uniform elongations were calculated by a new way, that is averaging the change of distances between adjacent indents made in the specimen gauge region. Only the plastic components of the tensile load2
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Fig. 1. Engineering stress vs. strain of EP0.0 (a) and EP0.85 (b). The insets in (a) and (b) are the stress vs. plastic strain curves.
Fig. 2. 0.2% offset yield strengths, ultimate tensile strengths (UTS) and uniform elongations to fracture of EP0.0 (a) and EP0.85 (b).
detected in the micro giant fault. And the HRTEM image (Fig. 6c) shows obvious difference in lattice structure between β phase and ω phase, which indicates that ω phase forms in β phase during the tensile tests at 200K. From the bright-field image (Fig. 6d) and dark-field image (Fig. 6e), it can be seen that some precipitates were formed during the tensile test of EP0.85 at 77K. However, selected area electron diffraction pattern shows a single-phase-like pattern and no extra spots corresponding to the second phase was found (Fig. 6d). Besides, in Fig. 6f, there is no apparent difference in lattice structure and no distinct interface between the secondary particle and the matrix. These results indicate that the precipitate may be β′ phase, which has different composition from β phase (solute-lean β′ and solute-rich β) but the same crystal structure. This β phase decomposition phenomenon (β → β′+β) was also reported in other β-Ti alloys previously [29–32]. Fig. S2 shows the results of the research about β′ phase in Ti–25Nb–3Mo–3Zr–2Sn (wt.%) alloy [30]. The TEM microstructures and the atomic-scale HRTEM images of β′ phase in the EP0.85 and Ti–25Nb–3Mo–3Zr–2Sn alloy are similar. Besides, according to EDX map of β′ phase in Ti–25Nb–3Mo–3Zr–2Sn alloy (Fig. S2d), the β′ phase precipitate is lean in Nb element. Therefore, the precipitate in EP0.85 after tensile tests at 77K can be determined as β′ phase. This is the first time the formation of β′ phase is detected in β titanium alloy at the cryogenic temperature. Both of the above two phase transformations are reported to be the stepping stone of the transformation from β phase to α phase, furthermore, β′ phase and ω phase can be considered as precursors of α phase [31–33].
Fig. 3. Temperature dependent extent work hardening of EP0.0 and EP0.85.
100 nm in width and several hundred nanometers in length) in the EP0.0 specimen after tensile tests is shown in Fig. 5b. And there is no ω phase in the TEM microstructure of EP0.0 before tensile tests (Fig. S1b). So, the mall volume fraction of ω plates in the EP0.0 sample forms during the tensile tests at room temperature instead of already existing in EP0.0 before tensile deformation. Fig. 5c shows some micro shear bands which were also reported in other studies and were depicted as “giant faults” [1,8]. The SAED pattern in Fig. 4c indicates that the EP0.85 sample is dominated by the β-phase. In general, there are very few dislocations observed in the tensile-fractured EP0.85 sample deformed at 293K. The microstructures of the tensile-fractured EP0.85 sample deformed at 200K and 77K are shown in Fig. 6a-f. The SAED patterns with the zone axis of [110] β and morphology of the ω phase are shown together (Fig. 6a). The dark field image was obtained from the diffuse sports of ω phase marked by arrows which appeared at the 1/3 position and 2/3 position. From the bright field image (Fig. 6a), ω phase was
4. Discussion For bcc alloys, the formation of deformation twins strongly depends on the stability of β phase. As shown in Fig. 7, the Bo-Md phase stability diagrams is used to evaluate the stability of β-phase [34–37]. In Fig. 7, the β/β+ω phase boundary, represented by a solid line, indicates a 3
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Fig. 4. EBSD microstructures near the tensile fracture of EP0.0 deformed at 77 K (a), 200K (b) and293 K (c). Corresponding misorientation as a function of distance for the black line (d). Coloring is the inverse pole figure map projected on tensile direction.
region of the least stable single phase β. The boundary for the martensite starting temperature (Ms = RT) is given by a dotted line. Martensite phases such as ω, ⍺' or ⍺ depending on the alloy composition, will be able to exist predominantly below the dotted line. Besides, deformation mechanisms also could be predicted by the Bo-Md phase stability diagrams (Fig. 7). Above the solid line, the slip will be detected in the titanium alloy during the deformation. Twining and martensitic transformation are the possible deformation mechanisms below the solid line. The position of Gum metal (nominal composition: Ti–36Nb–2Ta–3Zr-0.3O) is presented by a symbol (TNTZNC) in Fig. 7. The actual composition of the alloy in this study is Ti–36Nb–2Ta–3Zr-
0.35O, in which the oxygen content is higher than 0.3%. As the oxygen has been proved as a β-phase stabilizer in Gum Metal [1,2,34] and decrease the decreases the Ms transition temperature, the position of the alloy in this study should be higher. Besides, during the tensile tests at room temperature, a small amount of ω phases were formed in the EP0.0 and the deformation mechanism was related to dislocation motion. Therefore, the position of EP0.0 at room temperature should be close to the solid line in Fig. 7 (TNTZOHT). With the temperature decreasing, the β-phase stability of EP0.0 decreases. So, the position of EP0.0 at 200K and 77K should be lower and the deformation mechanism of the EP0.0 at cryogenic temperatures prefers twinning. This speculation is further proved by the deformation twins in the
Fig. 5. TEM microstructure near the fracture of EP0.0 (a, b) and EP0.85 (c) deformed at 293K. The insets in (a) and (c) are the [110] diffraction patterns. 4
β
zone axis selected area
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Fig. 6. TEM microstructure near the fracture of EP0.85 deformed at 200K (a, b, c) and 77K (d, e, f), in which, (a) and (d) are bright-field images. (b) and (e) are darkfield images. (c) and (f) are HRTEM images. The insets in (a) and (d) are the [110] β zone axis selected area diffraction patterns.
twin boundaries can be considered as another kind of grain boundaries, the “effective grain size” will be much smaller than the real measured grain size. According to the “Hall-Petch effect” [38], the refinement of grains can significantly contribute to the enhanced strain hardening. The mechanism of twinning-induced plasticity (TWIP) in β-type Ti alloys has been reported by Min et al. [39,40]. Due to the formation of twins, there are many twin-twin and twin-grain boundaries in the EP0.0. During the tensile tests, the tensile stress distributes not only in the grain boundaries but also in the twin-twin and twin-grain boundaries. Moreover, twinning transfer was often present between two neighboring grains and secondary twinning occurred within the primary twins caused by the twin-twin intersection. So, the deformation twinning also could provide the accommodation of local stress concentrations. As a result, the tensile stress would distribute more evenly and the stress concentration would be abated. So, the formation of deformation twinning in EP0.0 not only induces the strain hardening but also postpone the necking stability and hence increase the ductility. With the decrease of temperature, the fracture surface of shrinkage decreases, while the uniform elongation of EP0.0 increases, is also attributed to the effect of twinning on the necking stability. Deformation twinning ability and activities has been shown to be strongly dependent on the stacking fault energy of metallic materials. As mentioned above, twinning tends to occur in FCC alloys with low stacking fault energy. However, the twinning boundary migration energy for BCC alloys is smaller than that for FCC metals [41]. Once the deformation twins formed in EP0.0, the mobility of the twinning boundary is higher and the hindering effect on the dislocation motions is weaker than those in FCC alloys, which will also significantly contribute to the enhanced plasticity of EP0.0. Temperature dependence of deformation behavior of the EP0.85 can also be explained by the stability of the β phase. According to Satio et al. [1,42], there are a certain amount of Zr–O atom clusters in gum metal after cold swage. The heterogeneous elemental distributions in
Fig. 7. β phase stability as a function of Bo and.Md
morphology of the tensile-fractured EP0.0 (Fig. 4). Moreover, more twins which the tensile stress activates contribute to the plastic deformation with the testing temperature decreasing. For EP0.0, both tensile strength and uniform elongation of EP0.0 could be enhanced by the deformation twins. The strong dependencies of strain hardening on temperature could also be attributed to the enhanced twinning activities at lower temperatures. The mechanical twins formed during tensile rests will be able to serve as obstacles to dislocation motions. Furthermore, interactions between dislocations and twins will cause the accumulation of a high density of sessile dislocations within the twin lamellae, leading to increased twin strength with deformation. These accumulated dislocations are potentially effective barriers to dislocation motion and can provide additional strengthening within the induced mechanical twins and increase the critical stress required to induce plastic deformation across the twins. In addition, the strain hardening in EP0.0 is also due to the dynamic microstructural refinement induced by twinning since the deformation-
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short range results in different β phase stability. The β-phase stability of some areas not containing Zr–O atom alusters in EP0.85 is lower than that of EP0.0. With the temperature decreasing, the β-phase stability of EP0.85 decreases. So, the position of EP0.85 at 200K and 77K should be closer to the boundary for the martensite start temperature than that of EP0.0 and it is very possible for EP0.85 that stress induced martensitic transformation occurred during the tensile tests. The relationship between the critical stress for martensitic transformation and temperature can be described by the Clausius-Clapeyron relationship [43]: dσ/dT = −ΔH/T·ε = −ΔS/ε
Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.msea.2019.138293. References [1] T. Saito, T. Furuta, J. Hwang, et al., Multifunctional alloys obtained via a dislocation-free plastic deformation mechanism, Science 300 (2003) 464–467. [2] M. Besse, P. Castany, T. Gloriant, Mechanisms of deformation in gum metal TNTZ-O and TNTZ titanium alloys: a comparative study on the oxygen influence, Acta Mater. 59 (2011) 5982–5988. [3] W. Zhang, Y. Liu, H. Wu, et al., Microstructural evolution during hot and cold deformation of Ti-36Nb-2Ta-3Zr-0.35O alloy, Trans. Nonferrous Metals Soc. China 26 (2016) 1310–1316. [4] V. Vorontsov, N. Jones, K. Rahman, et al., Superelastic load cycling of gum metal, Acta Mater. 88 (2015) 323–333. [5] W. Guo, M. Quadir, S. Moricca, et al., Microstructural evolution and final properties of a cold-swaged multifunctional Ti-Nb-Ta-Zr-O alloy produced by a powder metallurgy route, Mater. Sci. Eng. A 575 (2013) 206–216. [6] M. Lai, C. Tasan, D. Raabe, Deformation mechanism of ω-enriched Ti-Nb-based gum metal: dislocation channeling and deformation induced ω-β transformation, Acta Mater. 100 (2015) 290–300. [7] S. Kuramoto, T. Furuta, J. Hwang, et al., Elastic properties of gum metal, Mater. Sci. Eng. A 442 (2006) 454–457. [8] L. Wei, H. Kim, S. Miyazaki, Effects of oxygen concentration and phase stability on nano-domain structure and thermal expansion behavior of Ti-Nb-Zr-Ta-O alloys, Acta Mater. 100 (2015) 313–322. [9] L. Wang, W. Lu, J. Qin, et al., Microstructure and mechanical properties of coldrolled TiNbTaZr biomedical titanium alloy, Mater. Sci. Eng. A 490 (2008) 421–426. [10] M. Haghshenas, Multi-cycling instrumented nanoindentation of a Ti-23Nb-0.7Ta2Zr-1.2O alloy in annealed condition, Mater. Sci. Eng. A 697 (2017) 8–16. [11] T. Furuta, S. Kuramoto, J. Hwang, et al., Elastic deformation behavior of multifunctional Ti-Nb-Ta-Zr-O alloys, Mater. Trans. 46 (2005) 3001–3007. [12] S. Kuramoto, T. Furuta, J. Hwang, et al., Plastic deformation in a multifunctional TiNb-Ta-Zr-O alloy, Metall. Mater. Trans. A 37 (2006) 657–662. [13] R.J. Talling, R.J. Dashwood, M. Jackson, et al., Determination of (C11-C22) in Ti–36Nb–2Ta–3Zr–0.3O (wt.%) (Gum metal), Scr. Mater. 59 (2008) 669–672. [14] R.J. Talling, R.J. Dashwood, M. Jackson, et al., On the mechanism of superelasticity in Gum metal, Acta Mater. 57 (2009) 1188–1198. [15] T. Yano, Y. Murakami, D. Shindo, et al., Transmission electron microscopy studies on nanometer-sized ω phase produced in Gum Metal, Scr. Mater. 63 (2010) 536–539. [16] Y. Yang, P. Castany, E. Bertrand, et al., Stress release-induced interfacial twin boundary ω phase formation in a β type Ti-based single crystal displaying stressinduced a″ martensitic transformation, Acta Mater. 149 (2018) 97–107. [17] M.J. Lai, T. Li, D. Raabe, ω phase acts as a switch between dislocation channeling and joint twinning- and transformation-induced plasticity in a metastable β titanium alloy, Acta Mater. 151 (2018) 67–77. [18] E. Plancher, C.C. Tasan, S. Sandloebes, et al., On dislocation involvement in Ti-Nb gum metal plasticity, Scr. Mater. 68 (2013) 805–808. [19] S. Shin, C. Zhu, K.S. Vecchio, Observations on {332} < 113 > twinning-induced softening in Ti-Nb Gum metal, Mater. Sci. Eng. A 724 (2018) 189–198. [20] Y. Yang, G. Li, G. Chen, et al., Evolution of deformation mechanisms of Ti-22.4Nb0.73Ta-2Zr-1.34O alloy during straining, Acta Mater. 58 (2010) 2778–2787. [21] S. Shin, C. Zhang, K.S. Vecchioa, Phase stability dependence of deformation mode correlated mechanical properties and elastic properties in Ti-Nb gum metal, Mater. Sci. Eng. A 702 (2017) 173–183. [22] L. Wei, H. Kim, T. Koyano, et al., Effects of oxygen concentration and temperature on deformation behavior of Ti-Nb-Zr-Ta-O alloys, Scr. Mater. 123 (2016) 55–58. [23] M. Tane, T. Nakano, S. Kuramoto, et al., ω Transformation in cold-worked Ti-NbTa-Zr-O alloys with low body-centered cubic phase stability and its correlation with their elastic properties, Acta Mater. 61 (2013) 139–150. [24] W. Zhang, Y. Liu, H. Wu, et al., Room temperature creep behavior of Ti-Nb-Ta-Zr-O alloy, Mater. Char. 118 (2016) 29–36. [25] Z. Wu, H. Bei, Microstructures and mechanical properties of compositionally complex Co-free FeNiMnCr FCC solid solution alloy, Mater. Sci. Eng. A 640 (2015) 217–224. [26] B. Gludovatz, A. Hohenwarter, D. Catoor, et al., A fracture-resistant high-entropy alloy for cryogenic applications, Science 345 (2014) 1153–1158. [27] I.P. Semenova, J.M. Modina, A.V. Polyakov, et al., Fracture toughness at cryogenic temperatures of ultrafine-grained Ti-6Al-4V alloy processed by ECAP, Mater. Sci. Eng. A 716 (2018) 260–267. [28] G. Singh, G. Bajargan, R. Datta, et al., Deformation and strength of Ti-6Al-4V alloyed with B at cryogenic temperatures, Mater. Sci. Eng. A 611 (2014) 45–57. [29] T. Sakamoto, Y. Hiagaki, S. Kobayashi, et al., Precipitation of β' phase in a low cost beta titanium alloy, Mater. Sci. Forum 638–642 (2010) 461–464. [30] Y. Zhu, S. Zhu, M. Dargusch, et al., HAADF-STEM study of phase separation and the subsequent α phase precipitation in a β-Ti alloy, Scr. Mater. 112 (2016) 46–49. [31] A.M. Mebed, T. Miyazaki, Computer simulation and experimental investigation of the spinodal decomposition in the β Ti-Cr binary alloy system, Metall. Mater. Trans. A 29 (1998) 739–749. [32] S. Nag, Y. Zheng, R. Williams, et al., Non-classical homogeneous precipitation mediated by compositional fluctuations in titanium alloys, Acta Mater. 60 (2012)
(2)
where σ is a uniaxial stress, ε is a transformation strain, and ΔH and ΔS is the enthalpy and entropy of the martensitic transformation per unit volume, respectively. According to this relationship and some researches [22,44], with decreasing temperature, the critical stress for martensitic transformation decreases. Moreover, there are a certain amount of structural nano-disturbances in gum metals after cold swaging [1,11,12]. The disturbances would contribute to the nucleation of martensitic phases. Therefore, at cryogenic temperature, stress induced martensitic transformation is apt to occurring in the EP0.85. Both ω and β′ phase lead to a significant improvement on the tensile strength and an obvious decrease on the elongation of EP0.85. However, it needs further study for the detailed mechanisms of the formation of ω and β′ phase at different cryogenic temperatures.
5. Conclusions In summary, the mechanical behavior of Ti–36Nb–2Ta–3Zr-0.35O alloy with two different conditions, was investigated at different cryogenic temperatures, the following conclusions were drawn: 1. The mechanical behavior of Ti–36Nb–2Ta–3Zr-0.35O alloy is temperature-dependent. The values of yield and ultimate strengths of both EP0.0 and EP0.85 increase with the decrease of temperature. However, the uniform elongation of the EP0.0 increases with the decrease of temperature while that of the EP0.85 shows the opposite trend. 2. At cryogenic temperatures, the stability of β phase decreases, resulting in deformation twining formed in EP0.0 and different transitional phases formed in EP0.85: ω at 200K and β′ at 77K. 3. For EP0.0, the enhanced twining activities and dynamic microstructural refinement caused by deformation twining contribute to the strain hardening at cryogenic temperature. Besides, the enhanced ductility is due to the TWIP effect. 4. For the EP0.85, both ω and β′ phase lead to a significant improvement on the tensile strength and an obvious decrease on the elongation of EP0.85.
‘Data availability’ statement The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations.
Acknowledgment This work was supported by Research Funds for Central Universities of China (531118010305), National Natural Science Funds for Distinguished Young Scholar of China (51625404), National Natural Science Foundation of China (51604112) and Natural Science Foundation of Hunan Province of China (2017JJ3089). We also thank the Advanced Research Center of CSU for performing HRTEM examination and the Dr. Mingxing Huang from the University of Hong Kong for reviewing this paper. 6
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Phys. Soc. B 64 (1951) 747–753. [39] X. Min, X. Chen, S. Emura, et al., Mechanism of twinning-induced plasticity in βtype Ti-15Mo alloy, Scr. Mater. 69 (2013) 393–396. [40] X. Min, S. Emura, X. Chen, et al., Deformation microstructural evolution and strain hardening of differently oriented grains in twinning-induced plasticity β titanium alloy, Mater. Sci. Eng. A 659 (2016) 1–11. [41] S. Ogata, J. Li, S. Yip, Energy landscape of deformation twinning in bcc and fcc metals, Phys. Rev. B 71 (2005) 224102. [42] N. Nagasako, R. Asahi, D. Isheim, et al., Microscopic study of gum-metal alloys: a role of trace oxygen for dislocation-free deformation, Acta Mater. 105 (2016) 347–354. [43] P. Wollants, J.R. Roos, L. Delaey, Thermally and stress-induced thermoelastic martensitic transformations in the reference frame of equilibrium thermodynamics, Prog. Mater. Sci. 37 (3) (1993) 227–288. [44] M. Niinomi, M. Nakai, M. Hendrickson, et al., Influence of oxygen on omega phase stability in the Ti-29Nb-13Ta-4.6Zr alloy, Scr. Mater. 123 (2016) 144–148.
6247–6256. [33] S. Nag, R. Banerjee, R. Srinivasan, et al., ω-Assisted nucleation and growth of α precipitates in the Ti-5Al-5Mo-5V-3Cr-0.5Fe β titanium alloy, Acta Mater. 57 (2009) 2136–2147. [34] M. Abdelhady, K. Hinoshita, M. Morinaga, General approach to phase stability and elastic properties of β-type Ti-alloys using electronic parameters, Scr. Mater. 55 (2006) 477–480. [35] M. Abdel-Hady, H. Fuwa, K. Hinoshita, et al., Phase stability change with Zr content in β-type Ti–Nb alloys, Scr. Mater. 57 (2007) 1000–1003. [36] M. Marteleur, F. Sun, T. Gloriant, et al., On the design of new β-metastable titanium alloys with improved work hardening rate thanks to simultaneous TRIP and TWIP effects, Scr. Mater. 66 (2012) 749–752. [37] S.X. Liang, X.J. Feng, L.X. Yin, et al., Development of a new β Ti alloy with low modulus and favorable plasticity for implant material [J], Mater. Sci. Eng. C 61 (2016) 338–343. [38] E.O. Hall, The deformation and ageing of mild steel: III discussion of results, Proc.
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Plastic deformation mechanism of Ti–Nb–Ta–Zr–O alloy at cryogenic temperatures
T
Wei-dong Zhanga, Zhenggang Wua, Yong Liub, Hongbin Beic, Bin Liub, Jingwen Qiud,∗ a
College of Materials Science and Engineering, Hunan University, Changsha, 410082, PR China State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, PR China c Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN, 37831, USA d Hunan Provincial Key Laboratory of Advanced Materials for New Energy Storage and Conversion, Hunan University of Science and Technology, Xiangtan, 411201, PR China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Gum metal Cryogenic mechanical properties Twinning β′ phase ω phase
The deformation behavior and mechanism of Ti–36Nb–2Ta–3Zr-0.35O alloy at room and cryogenic temperatures (77 and 200K) were investigated. The microstructure near the fracture was observed with electron backscattered diffraction analysis (EBSD) and transmission electron microscopy (TEM). The results show that the plastic deformation behavior of the alloy is sensitive to test temperature and cold deformation history of the alloy. For the alloy without cold deformation, both tensile strength and ductility have a significant enhancement at cryogenic temperatures, mainly due to the occurrence of mechanical twinning. Moreover, with the decrease of the temperature, the density of mechanical twinning increases. For the alloy cold swaged by 85%, an increase in strength accompanied by a decrease in uniform elongation was found at cryogenic temperatures. The unique temperature dependence of the deformation behavior was due to the low stability of β phase. At 200K, ω phase precipitated in β phase; while at 77K, an interesting transformation (β phase to β′ phase) was observed.
1. Introduction The Ti–36Nb–2Ta–3Zr-0.35O (wt.%) alloy (also known as gum metal) has obtained increased interests in recent years due to its novel properties, such as high strength, low modulus, superelasticity, superplasticity, Invar-like thermal expansion, and Elinvar-like thermal dependence of the elastic modulus at room temperature [1–10]. These novel properties were achieved by selecting a unique composition that simultaneously satisfies three electronic parameters: electron per atom ratio (e/a of about 4.24), bond order (Bo of about 2.87) and d electron orbital energy level (Md of about 2.45 eV). Moreover, a certain plastic deformation in the preparation process is another major requisite for the above novel properties of gum metal. After Saito et al. developed gum metal in 2003 [1], there are extensive researches to elucidate the deformation mechanism of gum metal. Saito et al. [1,11,12] claimed that the deformation mechanisms of gum metal is different from the typical titanium alloy. During the cold deformation, no dislocation movement, twinning and phase transformation take place, and instead, the unstable β lattice readily forms giant faults by ideal shear, thus accommodates the plastic strain. However, the deformation mechanism of gum metal is still
∗
controversial. Stress-induced phase transformation [13–17] and deformation induced twinning [16–19] have been reported to be active in a deformed gum metal by some other researches. Talling et al. [13,14] reported that the superelastic behavior of gum metal is associated with the reversible stress induced martensitic transformation (β → α"). Furthermore, cold work can result in the formation of large amount of stress-induced ω phase [6,15–17], and the formation of ω phase probably plays an important role in increasing the tensile strength. In the research of S. Shin et al. [19], mechanical twins were detected after thermomechanical-cycling processes applied to gum metal, resulting in a twinning-induced softening effect (about 18% decrease in microhardness). Besides, Yang et al. [20] found that the dominant deformation mechanisms are related to not only the extent of deformation but also the stability of the β phase. The research of S. Shin et al. [21] also proved that the β-phase stability dependence of deformation mode has a decisive significance effect on the mechanical properties of gum metal. Moreover, to further understand the deformation mechanism of gum metal, some other experiments under extreme conditions were performed, for example, the deformation at low temperatures (below 273K). Saito et al. [1] reported that the elastic deformability of the
Corresponding author. E-mail address: [email protected] (J. Qiu).
https://doi.org/10.1016/j.msea.2019.138293 Received 4 June 2019; Received in revised form 22 July 2019; Accepted 13 August 2019 Available online 14 August 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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displacement data were presented in the insets in Fig. 1. To compute the plastic components of the strain, a line was fit to the linear elastic portion of the stress-strain curves and the amount of the elastic strain at a given point on the curve was subtracted from the total strain to obtain the plastic strain. And in general, with decreasing temperature, the flow stress increases and the stress-strain curves shift up. Besides, EP0.85 exhibits shorter necking regions (from the point of maximum stress to the point of fracture) on its 77 and 200K stress-strain curves. Fig. 2 summarizes the 0.2% offset yield strength, ultimate tensile strengths and uniform elongation to fracture. The individual data points in these figures are average of three tensile tests. he strength values for both alloys show similar temperature-dependent trend. For both alloys, the highest values of the yield strength (YS) and ultimate tensile strength (UTS) are both obtained at 77 K. An increase in temperature results in a monotonic decrease in both the yield and ultimate tensile strengths. It should be noted, however, that the strength of the EP0.85 is much larger than that of the EP0.0 at the same temperature. In addition, the uniform elongation values for both alloys show opposite temperature-dependent trend. For the EP0.0, the uniform elongation increases with the decrease of temperature, while for the EP0.85, the uniform elongation decreases with the decrease of temperature. Due to the absent of strain extensometer in the tensile tests at cryogenic temperatures, it is difficult to calculate accurately the work hardening exponent. Therefore, the difference between UTS and YS is used to qualitatively describe the work hardening capability. Temperature dependent extent work hardening of EP0.0 and EP0.85 was shown in Fig. 3. The work hardening capability (the difference between UTS and YS) of the EP0.0 also decreases with the increase of temperature. The largest difference between values of UTS and YS for the EP0.0 appears at 77 K. The work hardening capability of the EP0.85 is lower than that of EP0.0, and no clear temperature dependence was found. The unique temperature dependent strength and ductility of EP0.0 also reported in FCC metals and alloys with low stacking fault energy. At cryogenic temperatures, deformation twining occurs in these FCC materials, such as FeCoCrNi and FeCoCrNiMn equiatomic solid solution alloys [25,26], in which deformation induced twining can provide a higher work hardening capability and also contribute to the enhanced ductility. However, the stacking fault energy of EP0.0 is much higher than the critical stacking fault energy below which twining occurs (~45 mJ*m−2). For titanium alloys, the strength always increases with the decreasing temperature (below room temperature) and the elongation exhibits an opposite tendency, which is due to thermal influences on the dislocation motion [27,28]. In order to shed some light on this unique temperature dependent strength and ductility, the microstructures of EP0.0 and EP0.85 after tensile tests were characterized and analyzed. Fig. 4 shows the microstructures of tensile-fractured EP0.0 samples deformed at 77K, 200K and 273K. The microstructure at 77 K has a colony-like morphology which is full of parallel bands with various thickness and spacing (Fig. 4a). To identify details of these bands, a black line crossing the band was drawn and misorientations along the line were measured and plotted (Fig. 4d). It shows that whenever the line passes the individual band, there will be a ~53° misorientation, indicating that these bands are composed of Σ11{332}β 113β twins. The densities of the twin bands are largely reduced when the testing temperature is increased (as shown in Fig. 4b), and no twin bands are observed when the temperature is increased to 293 K (Fig. 4c). Besides, only Σ11{332}β 113β twins and no other types of twins are detected in the EP0.0 after the tensile tests at 200K and 77K. For EP 0.0, the enhanced yield strength (Fig. 1a) under lower temperatures (200K and 77 K) will make the critical stress that needs to be reached in order to activate deformation twinning easier to be achieved. Fig. 5 shows the TEM microstructure near tensile fracture of EP0.0 and EP0.85 deformed at room temperature. It can be seen that the ω phase formed in β phase (Fig. 5a). The morphology of ω plates (about
cold-worked gum metal increases with the decrease of temperature, exceeding 4% at 77K. Wei et al. [22] found that the yield stress increased monotonically with increasing the deformation temperature (193 K–233 K) for the 1.2–1.8 wt%O alloys. Tane et al. [23] reported that the formation of the ω phase and the low stability of the β phase were the dominant factors determining the elastic properties below the room temperature. However, these studies only focused on the elastic deformation behavior and properties of gum metal. So far, the plastic deformation behavior and mechanism of the gum metal at cryogenic temperatures has seldom been reported. In this work, the deformation behaviors of Ti–36Nb–2Ta–3Zr-0.35O alloy with and without cold deformation at different temperatures from 77K to 293K were investigated. And the purpose of this study is to understand the effect of cold deformation on the plastic deformation behaviors and mechanisms of Ti–36Nb–2Ta–3Zr-0.35O alloy at cryogenic temperatures. 2. Materials and methods 2.1. Material preparation The Ti–36Nb–2Ta–3Zr-0.35O alloy used in this work was prepared by powder metallurgy technique (cold isostatic pressing and vacuum sintering). Details of the alloy preparation and alloy composition can be found in our previous work [24]. In this work, the round bars after solution heat treatment at 1273K for 1h in argon (the sample was named as EP0.0) and after cold swaging by 85% (reduction in cross section area, the sample was named as EP0.85) were used. 2.2. Mechanical properties tests and material characterization Flat dog-bone-shaped specimens with a gage length of 9.5 mm were cut from the alloy bars. Tensile tests were performed using an Instron screw-driven testing machine at a strain rate of 10−4 s−1and temperatures of 77, 200, and 273 K. Three nominally identical specimens were tested at each temperature. The uniform elongation was obtained by a new method which has been reported by Wu et al. [25]. Vickers microhardness indents spaced 1 mm apart were made along the specimen gage lengths using a Vickers Hardness tester with a force of 50g. The change in the distance between adjacent indents, excluding the two indents on either side of the fracture plane, was used to calculate the uniform elongations to fracture. 2.3. Material characterization FEI Quanta 650 FEG-SEM equipped with EBSD system was used to characterize the microstructures of the tensile-fractured materials. Both of the EBSD samples and the TEM samples were cut from uniform deformation region of the tensile test sample. The electron backscatter diffraction (EBSD) measurements were performed with an electron probe current of 3.0 nA at an acceleration voltage of 20 kV. The qualification and analysis of the EBSD results were performed using Channel 5 software. EBSD maps are displayed as inverse pole figure (IPF) maps in the direction of tensile axis (for the tensile-deformed microstructure). The TEM samples were prepared by twin-jet electrochemical polishing technique in an etching solution that contains 10% sulfuric acid (HSO4) and 90% methanol (CH3OH) at 20V and 238K. Then, the foils were examined using a Titan G2 60-300Transmission Electron Microscope. 3. Results and discussion Fig. 1 shows the engineering stress-strain curves of EP0.0 and EP0.85. The uniform elongations were calculated by a new way, that is averaging the change of distances between adjacent indents made in the specimen gauge region. Only the plastic components of the tensile load2
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Fig. 1. Engineering stress vs. strain of EP0.0 (a) and EP0.85 (b). The insets in (a) and (b) are the stress vs. plastic strain curves.
Fig. 2. 0.2% offset yield strengths, ultimate tensile strengths (UTS) and uniform elongations to fracture of EP0.0 (a) and EP0.85 (b).
detected in the micro giant fault. And the HRTEM image (Fig. 6c) shows obvious difference in lattice structure between β phase and ω phase, which indicates that ω phase forms in β phase during the tensile tests at 200K. From the bright-field image (Fig. 6d) and dark-field image (Fig. 6e), it can be seen that some precipitates were formed during the tensile test of EP0.85 at 77K. However, selected area electron diffraction pattern shows a single-phase-like pattern and no extra spots corresponding to the second phase was found (Fig. 6d). Besides, in Fig. 6f, there is no apparent difference in lattice structure and no distinct interface between the secondary particle and the matrix. These results indicate that the precipitate may be β′ phase, which has different composition from β phase (solute-lean β′ and solute-rich β) but the same crystal structure. This β phase decomposition phenomenon (β → β′+β) was also reported in other β-Ti alloys previously [29–32]. Fig. S2 shows the results of the research about β′ phase in Ti–25Nb–3Mo–3Zr–2Sn (wt.%) alloy [30]. The TEM microstructures and the atomic-scale HRTEM images of β′ phase in the EP0.85 and Ti–25Nb–3Mo–3Zr–2Sn alloy are similar. Besides, according to EDX map of β′ phase in Ti–25Nb–3Mo–3Zr–2Sn alloy (Fig. S2d), the β′ phase precipitate is lean in Nb element. Therefore, the precipitate in EP0.85 after tensile tests at 77K can be determined as β′ phase. This is the first time the formation of β′ phase is detected in β titanium alloy at the cryogenic temperature. Both of the above two phase transformations are reported to be the stepping stone of the transformation from β phase to α phase, furthermore, β′ phase and ω phase can be considered as precursors of α phase [31–33].
Fig. 3. Temperature dependent extent work hardening of EP0.0 and EP0.85.
100 nm in width and several hundred nanometers in length) in the EP0.0 specimen after tensile tests is shown in Fig. 5b. And there is no ω phase in the TEM microstructure of EP0.0 before tensile tests (Fig. S1b). So, the mall volume fraction of ω plates in the EP0.0 sample forms during the tensile tests at room temperature instead of already existing in EP0.0 before tensile deformation. Fig. 5c shows some micro shear bands which were also reported in other studies and were depicted as “giant faults” [1,8]. The SAED pattern in Fig. 4c indicates that the EP0.85 sample is dominated by the β-phase. In general, there are very few dislocations observed in the tensile-fractured EP0.85 sample deformed at 293K. The microstructures of the tensile-fractured EP0.85 sample deformed at 200K and 77K are shown in Fig. 6a-f. The SAED patterns with the zone axis of [110] β and morphology of the ω phase are shown together (Fig. 6a). The dark field image was obtained from the diffuse sports of ω phase marked by arrows which appeared at the 1/3 position and 2/3 position. From the bright field image (Fig. 6a), ω phase was
4. Discussion For bcc alloys, the formation of deformation twins strongly depends on the stability of β phase. As shown in Fig. 7, the Bo-Md phase stability diagrams is used to evaluate the stability of β-phase [34–37]. In Fig. 7, the β/β+ω phase boundary, represented by a solid line, indicates a 3
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Fig. 4. EBSD microstructures near the tensile fracture of EP0.0 deformed at 77 K (a), 200K (b) and293 K (c). Corresponding misorientation as a function of distance for the black line (d). Coloring is the inverse pole figure map projected on tensile direction.
region of the least stable single phase β. The boundary for the martensite starting temperature (Ms = RT) is given by a dotted line. Martensite phases such as ω, ⍺' or ⍺ depending on the alloy composition, will be able to exist predominantly below the dotted line. Besides, deformation mechanisms also could be predicted by the Bo-Md phase stability diagrams (Fig. 7). Above the solid line, the slip will be detected in the titanium alloy during the deformation. Twining and martensitic transformation are the possible deformation mechanisms below the solid line. The position of Gum metal (nominal composition: Ti–36Nb–2Ta–3Zr-0.3O) is presented by a symbol (TNTZNC) in Fig. 7. The actual composition of the alloy in this study is Ti–36Nb–2Ta–3Zr-
0.35O, in which the oxygen content is higher than 0.3%. As the oxygen has been proved as a β-phase stabilizer in Gum Metal [1,2,34] and decrease the decreases the Ms transition temperature, the position of the alloy in this study should be higher. Besides, during the tensile tests at room temperature, a small amount of ω phases were formed in the EP0.0 and the deformation mechanism was related to dislocation motion. Therefore, the position of EP0.0 at room temperature should be close to the solid line in Fig. 7 (TNTZOHT). With the temperature decreasing, the β-phase stability of EP0.0 decreases. So, the position of EP0.0 at 200K and 77K should be lower and the deformation mechanism of the EP0.0 at cryogenic temperatures prefers twinning. This speculation is further proved by the deformation twins in the
Fig. 5. TEM microstructure near the fracture of EP0.0 (a, b) and EP0.85 (c) deformed at 293K. The insets in (a) and (c) are the [110] diffraction patterns. 4
β
zone axis selected area
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Fig. 6. TEM microstructure near the fracture of EP0.85 deformed at 200K (a, b, c) and 77K (d, e, f), in which, (a) and (d) are bright-field images. (b) and (e) are darkfield images. (c) and (f) are HRTEM images. The insets in (a) and (d) are the [110] β zone axis selected area diffraction patterns.
twin boundaries can be considered as another kind of grain boundaries, the “effective grain size” will be much smaller than the real measured grain size. According to the “Hall-Petch effect” [38], the refinement of grains can significantly contribute to the enhanced strain hardening. The mechanism of twinning-induced plasticity (TWIP) in β-type Ti alloys has been reported by Min et al. [39,40]. Due to the formation of twins, there are many twin-twin and twin-grain boundaries in the EP0.0. During the tensile tests, the tensile stress distributes not only in the grain boundaries but also in the twin-twin and twin-grain boundaries. Moreover, twinning transfer was often present between two neighboring grains and secondary twinning occurred within the primary twins caused by the twin-twin intersection. So, the deformation twinning also could provide the accommodation of local stress concentrations. As a result, the tensile stress would distribute more evenly and the stress concentration would be abated. So, the formation of deformation twinning in EP0.0 not only induces the strain hardening but also postpone the necking stability and hence increase the ductility. With the decrease of temperature, the fracture surface of shrinkage decreases, while the uniform elongation of EP0.0 increases, is also attributed to the effect of twinning on the necking stability. Deformation twinning ability and activities has been shown to be strongly dependent on the stacking fault energy of metallic materials. As mentioned above, twinning tends to occur in FCC alloys with low stacking fault energy. However, the twinning boundary migration energy for BCC alloys is smaller than that for FCC metals [41]. Once the deformation twins formed in EP0.0, the mobility of the twinning boundary is higher and the hindering effect on the dislocation motions is weaker than those in FCC alloys, which will also significantly contribute to the enhanced plasticity of EP0.0. Temperature dependence of deformation behavior of the EP0.85 can also be explained by the stability of the β phase. According to Satio et al. [1,42], there are a certain amount of Zr–O atom clusters in gum metal after cold swage. The heterogeneous elemental distributions in
Fig. 7. β phase stability as a function of Bo and.Md
morphology of the tensile-fractured EP0.0 (Fig. 4). Moreover, more twins which the tensile stress activates contribute to the plastic deformation with the testing temperature decreasing. For EP0.0, both tensile strength and uniform elongation of EP0.0 could be enhanced by the deformation twins. The strong dependencies of strain hardening on temperature could also be attributed to the enhanced twinning activities at lower temperatures. The mechanical twins formed during tensile rests will be able to serve as obstacles to dislocation motions. Furthermore, interactions between dislocations and twins will cause the accumulation of a high density of sessile dislocations within the twin lamellae, leading to increased twin strength with deformation. These accumulated dislocations are potentially effective barriers to dislocation motion and can provide additional strengthening within the induced mechanical twins and increase the critical stress required to induce plastic deformation across the twins. In addition, the strain hardening in EP0.0 is also due to the dynamic microstructural refinement induced by twinning since the deformation-
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short range results in different β phase stability. The β-phase stability of some areas not containing Zr–O atom alusters in EP0.85 is lower than that of EP0.0. With the temperature decreasing, the β-phase stability of EP0.85 decreases. So, the position of EP0.85 at 200K and 77K should be closer to the boundary for the martensite start temperature than that of EP0.0 and it is very possible for EP0.85 that stress induced martensitic transformation occurred during the tensile tests. The relationship between the critical stress for martensitic transformation and temperature can be described by the Clausius-Clapeyron relationship [43]: dσ/dT = −ΔH/T·ε = −ΔS/ε
Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.msea.2019.138293. References [1] T. Saito, T. Furuta, J. Hwang, et al., Multifunctional alloys obtained via a dislocation-free plastic deformation mechanism, Science 300 (2003) 464–467. [2] M. Besse, P. Castany, T. Gloriant, Mechanisms of deformation in gum metal TNTZ-O and TNTZ titanium alloys: a comparative study on the oxygen influence, Acta Mater. 59 (2011) 5982–5988. [3] W. Zhang, Y. Liu, H. Wu, et al., Microstructural evolution during hot and cold deformation of Ti-36Nb-2Ta-3Zr-0.35O alloy, Trans. Nonferrous Metals Soc. China 26 (2016) 1310–1316. [4] V. Vorontsov, N. Jones, K. Rahman, et al., Superelastic load cycling of gum metal, Acta Mater. 88 (2015) 323–333. [5] W. Guo, M. Quadir, S. Moricca, et al., Microstructural evolution and final properties of a cold-swaged multifunctional Ti-Nb-Ta-Zr-O alloy produced by a powder metallurgy route, Mater. Sci. Eng. A 575 (2013) 206–216. [6] M. Lai, C. Tasan, D. Raabe, Deformation mechanism of ω-enriched Ti-Nb-based gum metal: dislocation channeling and deformation induced ω-β transformation, Acta Mater. 100 (2015) 290–300. [7] S. Kuramoto, T. Furuta, J. Hwang, et al., Elastic properties of gum metal, Mater. Sci. Eng. A 442 (2006) 454–457. [8] L. Wei, H. Kim, S. Miyazaki, Effects of oxygen concentration and phase stability on nano-domain structure and thermal expansion behavior of Ti-Nb-Zr-Ta-O alloys, Acta Mater. 100 (2015) 313–322. [9] L. Wang, W. Lu, J. Qin, et al., Microstructure and mechanical properties of coldrolled TiNbTaZr biomedical titanium alloy, Mater. Sci. Eng. A 490 (2008) 421–426. [10] M. Haghshenas, Multi-cycling instrumented nanoindentation of a Ti-23Nb-0.7Ta2Zr-1.2O alloy in annealed condition, Mater. Sci. Eng. A 697 (2017) 8–16. [11] T. Furuta, S. Kuramoto, J. Hwang, et al., Elastic deformation behavior of multifunctional Ti-Nb-Ta-Zr-O alloys, Mater. Trans. 46 (2005) 3001–3007. [12] S. Kuramoto, T. Furuta, J. Hwang, et al., Plastic deformation in a multifunctional TiNb-Ta-Zr-O alloy, Metall. Mater. Trans. A 37 (2006) 657–662. [13] R.J. Talling, R.J. Dashwood, M. Jackson, et al., Determination of (C11-C22) in Ti–36Nb–2Ta–3Zr–0.3O (wt.%) (Gum metal), Scr. Mater. 59 (2008) 669–672. [14] R.J. Talling, R.J. Dashwood, M. Jackson, et al., On the mechanism of superelasticity in Gum metal, Acta Mater. 57 (2009) 1188–1198. [15] T. Yano, Y. Murakami, D. Shindo, et al., Transmission electron microscopy studies on nanometer-sized ω phase produced in Gum Metal, Scr. Mater. 63 (2010) 536–539. [16] Y. Yang, P. Castany, E. Bertrand, et al., Stress release-induced interfacial twin boundary ω phase formation in a β type Ti-based single crystal displaying stressinduced a″ martensitic transformation, Acta Mater. 149 (2018) 97–107. [17] M.J. Lai, T. Li, D. Raabe, ω phase acts as a switch between dislocation channeling and joint twinning- and transformation-induced plasticity in a metastable β titanium alloy, Acta Mater. 151 (2018) 67–77. [18] E. Plancher, C.C. Tasan, S. Sandloebes, et al., On dislocation involvement in Ti-Nb gum metal plasticity, Scr. Mater. 68 (2013) 805–808. [19] S. Shin, C. Zhu, K.S. Vecchio, Observations on {332} < 113 > twinning-induced softening in Ti-Nb Gum metal, Mater. Sci. Eng. A 724 (2018) 189–198. [20] Y. Yang, G. Li, G. Chen, et al., Evolution of deformation mechanisms of Ti-22.4Nb0.73Ta-2Zr-1.34O alloy during straining, Acta Mater. 58 (2010) 2778–2787. [21] S. Shin, C. Zhang, K.S. Vecchioa, Phase stability dependence of deformation mode correlated mechanical properties and elastic properties in Ti-Nb gum metal, Mater. Sci. Eng. A 702 (2017) 173–183. [22] L. Wei, H. Kim, T. Koyano, et al., Effects of oxygen concentration and temperature on deformation behavior of Ti-Nb-Zr-Ta-O alloys, Scr. Mater. 123 (2016) 55–58. [23] M. Tane, T. Nakano, S. Kuramoto, et al., ω Transformation in cold-worked Ti-NbTa-Zr-O alloys with low body-centered cubic phase stability and its correlation with their elastic properties, Acta Mater. 61 (2013) 139–150. [24] W. Zhang, Y. Liu, H. Wu, et al., Room temperature creep behavior of Ti-Nb-Ta-Zr-O alloy, Mater. Char. 118 (2016) 29–36. [25] Z. Wu, H. Bei, Microstructures and mechanical properties of compositionally complex Co-free FeNiMnCr FCC solid solution alloy, Mater. Sci. Eng. A 640 (2015) 217–224. [26] B. Gludovatz, A. Hohenwarter, D. Catoor, et al., A fracture-resistant high-entropy alloy for cryogenic applications, Science 345 (2014) 1153–1158. [27] I.P. Semenova, J.M. Modina, A.V. Polyakov, et al., Fracture toughness at cryogenic temperatures of ultrafine-grained Ti-6Al-4V alloy processed by ECAP, Mater. Sci. Eng. A 716 (2018) 260–267. [28] G. Singh, G. Bajargan, R. Datta, et al., Deformation and strength of Ti-6Al-4V alloyed with B at cryogenic temperatures, Mater. Sci. Eng. A 611 (2014) 45–57. [29] T. Sakamoto, Y. Hiagaki, S. Kobayashi, et al., Precipitation of β' phase in a low cost beta titanium alloy, Mater. Sci. Forum 638–642 (2010) 461–464. [30] Y. Zhu, S. Zhu, M. Dargusch, et al., HAADF-STEM study of phase separation and the subsequent α phase precipitation in a β-Ti alloy, Scr. Mater. 112 (2016) 46–49. [31] A.M. Mebed, T. Miyazaki, Computer simulation and experimental investigation of the spinodal decomposition in the β Ti-Cr binary alloy system, Metall. Mater. Trans. A 29 (1998) 739–749. [32] S. Nag, Y. Zheng, R. Williams, et al., Non-classical homogeneous precipitation mediated by compositional fluctuations in titanium alloys, Acta Mater. 60 (2012)
(2)
where σ is a uniaxial stress, ε is a transformation strain, and ΔH and ΔS is the enthalpy and entropy of the martensitic transformation per unit volume, respectively. According to this relationship and some researches [22,44], with decreasing temperature, the critical stress for martensitic transformation decreases. Moreover, there are a certain amount of structural nano-disturbances in gum metals after cold swaging [1,11,12]. The disturbances would contribute to the nucleation of martensitic phases. Therefore, at cryogenic temperature, stress induced martensitic transformation is apt to occurring in the EP0.85. Both ω and β′ phase lead to a significant improvement on the tensile strength and an obvious decrease on the elongation of EP0.85. However, it needs further study for the detailed mechanisms of the formation of ω and β′ phase at different cryogenic temperatures.
5. Conclusions In summary, the mechanical behavior of Ti–36Nb–2Ta–3Zr-0.35O alloy with two different conditions, was investigated at different cryogenic temperatures, the following conclusions were drawn: 1. The mechanical behavior of Ti–36Nb–2Ta–3Zr-0.35O alloy is temperature-dependent. The values of yield and ultimate strengths of both EP0.0 and EP0.85 increase with the decrease of temperature. However, the uniform elongation of the EP0.0 increases with the decrease of temperature while that of the EP0.85 shows the opposite trend. 2. At cryogenic temperatures, the stability of β phase decreases, resulting in deformation twining formed in EP0.0 and different transitional phases formed in EP0.85: ω at 200K and β′ at 77K. 3. For EP0.0, the enhanced twining activities and dynamic microstructural refinement caused by deformation twining contribute to the strain hardening at cryogenic temperature. Besides, the enhanced ductility is due to the TWIP effect. 4. For the EP0.85, both ω and β′ phase lead to a significant improvement on the tensile strength and an obvious decrease on the elongation of EP0.85.
‘Data availability’ statement The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations.
Acknowledgment This work was supported by Research Funds for Central Universities of China (531118010305), National Natural Science Funds for Distinguished Young Scholar of China (51625404), National Natural Science Foundation of China (51604112) and Natural Science Foundation of Hunan Province of China (2017JJ3089). We also thank the Advanced Research Center of CSU for performing HRTEM examination and the Dr. Mingxing Huang from the University of Hong Kong for reviewing this paper. 6
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