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Effect of B addition on the microstructure and superelastic properties of a Ti-26Nb alloy
Materials Science & Engineering A 644 (2015) 85–89
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of B addition on the microstructure and superelastic properties of a Ti-26Nb alloy Yazan Al-Zain a,n, Hee Young Kim b,nn, Shuichi Miyazaki b,c,nnn a
Department of Industrial Engineering, The University of Jordan, Amman 11942, Jordan Division of Materials Science, University of Tsukuba, Tsukuba, Ibaraki 305-8573, Japan c Foundation for Advancement of International Science, Tsukuba, Ibaraki 305-0821, Japan b
art ic l e i nf o
a b s t r a c t
Article history: Received 26 May 2015 Received in revised form 17 July 2015 Accepted 18 July 2015 Available online 20 July 2015
The mechanical properties and shape memory behavior of Ti–26Nb–(0–1.0)B and Ti–27Nb alloys were investigated. The stress for inducing martensite and the critical stress for slip deformation were greatly affected by increasing B content up to 0.1 at%, and less affected by further increase in B content up to 0.7 at%. The martensitic transformation start temperature decreased suddenly with a slope of 350 K/at% B when B content was less than 0.1 at% and decreased with a smaller slope of 15 K/at% B with further increase in B content up to 1.0 at%. In all samples tested, only superelasticity was observed at room temperature. TEM investigations revealed that the solubility of B in Ti–26Nb alloy was around 0.1 at%. In spite of the low solubility limit, B was effective to stabilize the superelastic behavior by effectively increasing the critical stress for slip deformation. On the other hand, while the stress for inducing martensite increased effectively by the addition of 1.0 at% Nb to the Ti–26Nb alloy, the critical stress for slip deformation remained almost constant and hence superelasticity was deteriorated. & 2015 Elsevier B.V. All rights reserved.
Keywords: Ti-based shape memory alloys Superelasticity Shape memory Martensitic transformation
1. Introduction Ti–Ni shape memory alloys (SMAs) belong to the family of the so-called smart materials. These alloys are capable of restoring their original shape when deformed followed by either load removal alone or heating after load removal. In the former case the property exhibited is referred to as superelasticity (SE) while in the latter case the property exhibited is referred to as shape memory effect (SME). The excellent SE and SME exhibited by Ti–Ni alloys [1,2] render them superior to other classes of SMAs. However, human body allergy to pure Ni has been pointed out in some reports [3]. Accordingly, scientists started searching for SMAs that are free of Ni. Several studies showed that SE and SME are exhibited in some β-type Ni-free Ti-alloys. Such examples include Ti–Nb-base alloys [4–38] and Ti–Mo-base alloys [39–45]. The SE and SME exhibited by the binary Ti–Nb alloys were not as good as those exhibited by the Ti–Ni alloys due to the relatively low critical stress for slip deformation (sCSS) and small transformation strain. Several researchers showed that the SE and SME of n
Corresponding author. Corresponding author. Corresponding author at: Division of Materials Science, University of Tsukuba, Tsukuba, Ibaraki 305-8573, Japan. E-mail addresses: [email protected] (Y. Al-Zain), [email protected] (H.Y. Kim), [email protected] (S. Miyazaki). nn
nnn
http://dx.doi.org/10.1016/j.msea.2015.07.054 0921-5093/& 2015 Elsevier B.V. All rights reserved.
the binary Ti–Nb alloys could be enhanced by the addition of a third or more elements. Examples are Ti–Nb–O [34], Ti–Nb–N [35], Ti–Nb–Mo [13], Ti–Nb–Al [36], Ti–Nb–Mo–Sn [15], Ti–Nb–(Ge, Ga) [8], Ti–Nb–Al–B [37], Ti–Nb–Ta [10], and Ti–Nb–Zr [38]. According to these reports, the addition of Zr, Ge, Sn, Al, O or N has effectively increased the superelastic strain. Moreover, the addition of O and N has effectively stabilized the SE by increasing the magnitude of sCSS. In this research, the effect of B addition on the martensitic transformation start temperature (Ms), microstructure, mechanical properties and SE of the Ti–26Nb alloy was systematically investigated. In addition, the SE of the B-added Ti–26Nb alloy was compared to that of the Ti–27Nb alloy.
2. Experimental procedure Ti–26 at% Nb–(0–1.0) at% B and Ti–27 at% Nb alloys (all compositions are hereafter described in atomic per cent ) were prepared by an Ar-arc melting method using high purity titanium, niobium and boron and melted six times to avoid segregation. The ingots were homogenized at 1273 K for 7.2 ks followed by coldrolling up to a 98.5% reduction in thickness. Specimens were cut using an electric arc discharge machine and then heat-treated at 1173 K for 1.8 ks in an Ar atmosphere, followed by quenching in water at room temperature (RT). In order to remove the lightly
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oxidized surface layer, all specimens were etched using a solution of H2O, HNO3 and HF (5:4:1) in a water bath at a temperature of about 313 K. The shape memory properties were characterized by a tensile testing machine at RT. The gage length and width of the specimens for tensile tests were 20 mm and 1.5 mm, respectively. The Ms was estimated from thermal cycling tests under various stress levels. Finally, microstructure was investigated using an optical microscope and the internal structure was observed by a transmission electron microscope (TEM), using a JOEL 2010F microscope operated at 200 kV.
3. Results and discussion 3.1. Effect of B content on microstructure and grain size. TEM investigations were carried out in order to investigate the internal structure of B-free and B-added Ti–26Nb alloys. Fig. 1 shows bright-field images of Ti–26Nb–(0, 0.1, 0.2 and 0.7)B alloys and selected-area diffraction patterns e and f taken inside the precipitates, corresponding to images c and d, respectively. It is noticed that the matrix is composed only of the parent β phase for samples containing less than 0.2 at% B, and there is no formation of precipitates. Adding a larger amount of B resulted in precipitation of TiB particles inside the β matrix. Moreover, the size and density of these precipitates increased with increasing B content above 0.2 at%. This indicates that the solubility limit of B in the Ti–26Nb alloy is about 0.1 at%. A specific orientation relationship was not identified between the matrix and the TiB
precipitates. Moreover, formation of the martensite phase was not observed in any of the specimens. This indicates that Ms for all the alloys is below RT. To investigate the effect of B addition on the average grain size (avg.), optical microscopy observation was carried out at RT. Fig. 2 shows the microstructure of Ti–26Nb–(0, 0.1, 0.2 and 0.7)B alloys. Generally speaking, avg. decreased by increasing B content from 0 to 0.7 at%. From avg. indicated on each image, it is clear that the grain size did not change by the addition of 0.1 at% B, and the small drop is considered to be within experimental errors. However, avg. dropped by higher additions up to 0.2 at% B. Finally, it decreased only slightly by further increase in B content above 0.2 at% B. It can be seen that most of the precipitates are found inside the grains while very few are found at the grain boundaries for the alloy containing 0.2 at% B. Similar observation was reported by Tang et al., they stated that after conducting solution treatment of βphase TiB precipitates were seen inside the grains [46]. In spite of the increase in the density of precipitates, similar observations were seen for the alloy containing 0.7 at% B, with very small number of precipitates found at the grain boundaries relative to the number of precipitates seen inside the grains. Hence, precipitation of borides does not have pronounced effect on the grain size. Moreover the reason why there was no profound change in the avg. when B content was between 0 and 0.1 at% (49 μm in the binary case and 47 μm in the 0.1 at% B-added alloy) is that there was no precipitation of borides when B content was less than 0.2 at%. This is consistent with the result of solubility limit investigated by TEM. Moreover, the current study and the result of the previous study on Ti–Nb–Al–B alloys [37] are in good
Fig. 1. Bright field images of the Ti–26Nb alloys with (a) 0 at% B, (b) 0.1 at% B, (c) 0.2 at% B and (d) 0.7 at% B heat-treated at 1173 K for 1.8 ks. (e) and (f) represent the diffraction patterns for TiB in (c) and (d), respectively.
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Fig. 2. Optical microscope images for the Ti–26Nb alloys with (a) 0 at% B, (b) 0.1 at% B, (c) 0.2 at% B and (d) 0.7 at% B heat-treated at 1173 K for 1.8 ks.
agreement in a way that B in the solid solution has no effect on the grain size. Also, both studies confirmed that the decrease in grain size was due to the pinning effect of borides at the grain boundaries. 3.2. Effect of B content on the martensitic transformation start temperature (Ms). Fig. 3 shows strain-temperature curves measured during thermal cycling between 173 K and 373 K under constant various stresses for the Ti–26Nb–0.7B alloy. The test was performed in a way such that the magnitude of the applied stress was increased with 50 MPa for every following cycle, the starting stress being 100 MPa. Same specimen was used throughout the test. The upper and lower halves of the curves represent heating and cooling cyles, respectively. The Ms at each curve was determined as illustrated on the figure.
Generally speaking, Ms increased with increasing the applied stress, since the application of the stress usually assists the transformation. The same test was conducted for all specimens with different B contents. Ms is usually taken at zero-stress using simple extrapolation technique shown in Fig. 4 (the solid line). Following the dashed lines, it can be seen that the degree of the increment in Ms by increasing the stress becomes smaller. This deviation is related to the plastic deformation (shown as (εp) in Fig. 3) introduced after a full cycle of forward transformation and reverse transformation. Accordingly, it is not possible in this case to determine Ms accurately by the extrapolation technique. To avoid such deviation, Ms is estimated from the first cycle. Recalling that the solubility limit of B in Ti–26Nb alloy does not exceed 0.1 at%, where TiB precipitates afterwards, the matrix content of Nb shall increase with about 0.45 at% when 1.0 at% B is added. In a previous report [31] it was shown that 1 at% addition of 300
Stress (MPa)
250 200 150 100 50 0 175 Fig. 3. Stress–temperature curves obtained at various constant stresses for the Ti– 26Nb–0.7B alloy heat-treated at 1173 K for 1.8 ks.
200
225 250 Temperature (K)
275
300
Fig. 4. Ms at various stress values for Ti–26Nb–0.7B alloys heat-treated at 1173 K for 1.8 ks.
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Table 1 sSIM, sCSS and εse for the alloys Ti–26 and 27 at% Nb and Ti–26 at% Nb–(0.1 and 0.7) at% B
Fig. 5. B content dependence of Ms for Ti–26Nb–(0–1.0)B alloys heat-treated at 1173 K for 1.8 ks.
Nb decreased Ms by 40 K. Hence, there must be a drop in Ms even if B becomes insoluble in the alloys. Fig. 5 shows the B content dependence of Ms for Ti–26Nb–(0– 1.0)B alloys. It is obvious that all the alloys posses Ms below RT, indicating that the phase at RT is the parent phase, this result is consistent with TEM results. Addition of B was effective to decrease Ms when its content was less than 0.1 at%: the effect of B on Ms was 350 K/at% B. This is due to the solid solution effect of B. Then Ms decreased with 15 K/at% B with further increase in B content. This is consistent with the discussion above; a 15 K decrease in Ms is very close to an increase of about 0.5 at% Nb in the matrix content. The results of Ms agree with the results shown by the previous study of Ti–Nb–Al–B [37] that B in solid solution was effective to decrease Ms. However, it disagrees with the effect of precipitation of B on Ms, as they showed that Ms remains constant irrespective of further increase in B content above the solubility limit. 3.3. Superelasticity and shape memory behavior of Ti-26Nb-(0-0.7)B and Ti–27Nb alloys. Tensile tests were carried out at RT in order to characterize the shape memory behvior of Ti–Nb–B alloys. Fig. 6 shows stress– strain curves for the Ti–26Nb, Ti–26Nb–0.1B, Ti–26Nb–0.7B and Ti–27Nb alloys heat treated at 1173 K for 1.8 ks. Load was applied until reaching 2.5% strain and then removed. A solid circle on each curve represents the apparent yield stress (In this study, the yield stress is set to be the stress at 0.2% offset strain), and the dashed
Fig. 6. Stress–strain curves for (a) Ti–26Nb (b) Ti–27Nb, (c) Ti–26Nb–0.1B and (d) Ti–26Nb–0.7B alloys heat-treated at 1173 K for 1.8 ks.
Alloy
sSIM (MPa)
sCSS (MPa)
εse (%)
Ti–26Nb Ti–27Nb Ti–26Nb–0.1B Ti–26Nb–0.7B
101 223 168 225
356 379 437 502
2.0 1.7 2.5 2.6
lines with arrows represent strain recovered as a result of heating. As shown earlier, Ms of all the alloys are below RT. Hence, the apparent yield stress is considered to be the stress for inducing the martensite phase (sSIM) from the parent β phase. In all specimens, most of shape recovery took place upon unloading, while a small amount of the permanent strain was recovered by heating after unloading. The magnitudes of sSIM for all alloys are summarized in Table 1. Taking sSIM of the Ti–26Nb alloy as the reference, it is seen that the magnitude of sSIM increased more severely by the addition of 1 at% Nb when compared to the increment caused by the addition of 0.1 at% B. It could be also seen that addition of more than 0.1 at% B to the binary alloy further increased sSIM. However, the effect of B addition on sSIM became less severe as B content increased above 0.1 at%. The increase in the value of the sSIM by addition of 1.0 at% Nb and 0.1 at% B to the binary Ti–26Nb alloy is a result of the solid solution effect of Nb and B on Ms. The slight increase in the magnitude of sSIM by further increasing B content above 0.1 at% is related to the precipitation effect of TiB and the corresponding enrichment of the matrix with Nb. 3.4. Effect of B addition on the critical stress for slip deformation (CSS). Fig. 7 shows stress–strain curves for the Ti–26Nb, Ti–26Nb– 0.1B, Ti–26Nb–0.7B and Ti–27Nb alloys heat-treated at 1173 K for 1.8 ks. Each stress–strain curve and strain recovery were obtained by a loading and unloading cycle. The test was repeated by increasing the maximum strain upon loading in the same specimen for each alloy. Heating was conducted after each cycle showing permanent strain after unloading. In Fig. 7b, εse and εp represent the superelastic recovery strain due to unloading and the plastic strain remained after heating, respectively. As shown earlier, all alloys exhibited SE. Residual strain was completely recovered by simple unloading up to the second cycle for Ti–26Nb alloy. The sCSS in this study is defined as the stress at which 0.5% plastic strain is introduced during cycling as shown in Fig. 7b. The values of sCSS for all alloys are shown in Table 1. It is clear that addition of B to the Ti–26Nb alloy was effective to increase the sCSS. On the other hand it remained almost unchanged with the addition of 1 at% Nb. This is an expected result, recalling that the atomic diameter of Nb is almost similar to that of Ti. Hence, solid solution hardening must be too small. In contrast, the atomic diameter of B is much smaller than that of Ti, giving a rise in the solid solution hardening effect to effectively increase the sCSS. The sCSS of Ti– 26Nb alloy was also increased by the addition of 0.7 at% B. For this alloy, only 0.1 at% B was soluble while 0.6 at% B precipitated in the form of TiB. These precipitates were also effective to increase the sCSS as they serve as obstacles hindering dislocation motion during deformation. However, it is seen that B in solid solution was more effective to increase the sCSS when compared to B in the form of TiB. Finally, as shown in Table 1, the εse decreased by the addition of 1 at% Nb to the Ti–26Nb. This can be explained by the large increase in the sSIM while having an almost constant value of sCSS. On the other hand, εse was increased by the addition of B. Similar values of εse were seen for both B-added alloys. SE was enhanced
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Fig. 7. Stress–strain curves obtained by cycling loading–unloading tensile tests for (a) Ti–26Nb (b) Ti–27Nb, (c) Ti–26Nb–0.1B and (d) Ti–26Nb–0.7B alloys heat-treated at 1173 K for 1.8 ks.
due to the larger increase in sCSS when compared to the increase in the sSIM, allowing transformation to proceed freely before slip deformation occurred. It is concluded that although the solubility limit of B in Ti–26Nb did not exceed 0.1 at%, B was good to stabilize the SE of Ti–Nb shape memory alloys.
4. Conclusions [1] Ti–27 at% Nb and Ti–26 at% Nb–(0–1.0) at% B alloys were fabricated by the Ar-arc melting method. SE was observed in the alloys. B was found to have a very small solubility in the binary Ti– 26Nb alloy with about 0.1 at%. Moreover, it was found that B was most effective to decrease the Ms when B content was less than 0.1 at%. The sSIM and sCSS were strongly affected by the addition of 0.1 at% B to the Ti–26Nb alloy due to solid solution effect; further increase in B content caused the change in these stresses to become smaller as a result of precipitation. B addition was effective to increase εse. Due to this reason and because sCSS was effectively increased, it is concluded that B addition is effective in stabilizing the SE of Ti–Nb alloys. [2] Addition of 1 at% Nb to the Ti–26Nb alloy resulted in the deterioration of the SE. This is because there was a large increase in the sSIM while sCSS remained almost unchanged.
Acknowledgment This work was partially supported by JSPS KAKENHI Grant numbers 26249104 and 25102704, and Fiscal 2014 JASSO Followup Research Fellowship.
References [1] S. Miyazaki, Y. Ohmi, K. Otsukua, Y. Suzuki, J. Phys. 43 (1982) 255–260. [2] S. Miyazaki, K. Otsuka, Trans. Iron Steel Inst. Jpn. 29 (1989) 353–377. [3] S. Shabalovskaya, J. Cunnick, J. Anderegg, B. Harmon, R. Sachdeva, in: Proceedings of the First International Conference on Shape Memory and Superelastic Technologies, Pacific Grove, California, USA, 1994, pp. 209–214. [4] E. Takahashi, T. Sakurai, S. Watanabe, N. Masahashi, S. Hanada, Mater. Trans. 43 (2002) 2978–2983. [5] H. Hosoda, Y. Fukui, T. Inamura, K. Wakashima, S. Miyazaki, K. Inoue, Mater. Sci. Forum 425–432 (2003) 3121–3126. [6] H. Hosoda, Y. Kinoshita, Y. Fukui, T. Inamura, K. Wakashima, H.Y. Kim, S. Miyazaki, Mater. Sci. Eng. A 438–440 (2006) 870–874. [7] Y.L. Hao, S.J. Li, S.Y. Sun, R. Yang, Mater. Sci. Eng. A 441 (2006) 112–118. [8] T. Inamura, Y. Fukui, H. Hosoda, K. Wakashima, S. Miyazaki, Mater. Sci. Eng. C 25 (2005) 426–432. [9] N. Sakaguchi, M. Niinomi, T. Akahori, Mater. Trans. 45 (2004) 1113–1119.
[10] H.Y. Kim, S. Hashimoto, J.I. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Mater. Sci. Eng. A 417 (2006) 120–128. [11] Y.B. Wang, Y.F. Zheng, Mater. Lett. 62 (2008) 269–272. [12] B.L. Wang, Y.F. Zheng, L.C. Zhao, Mater. Sci. Eng. A 486 (2008) 146–151. [13] Y. Al-Zain, H.Y. Kim, H. Hosoda, T.H. Nam, S. Miyazaki, Acta Mater. 58 (2010) 4212–4223. [14] Y. Al-Zain, H.Y. Kim, T. Koyano, H. Hosoda, T.H. Nam, S. Miyazaki, Acta Mater. 59 (2011) 1464–1473. [15] Y. Al-Zain, Y. Sato, H.Y. Kim, H. Hosoda, T.H. Nam, S. Miyazaki, Acta Mater. 60 (2012) 2437–2447. [16] Y. Al-Zain, H.Y. Kim, T. Koyano, H. Hosoda, S. Miyazaki, Scr. Mater. 103 (2015) 37–40. [17] M. Tahara, H.Y. Kim, H. Hosoda, S. Miyazaki, Acta Mater. 57 (2009) 2461–2469. [18] M.F. Ijaz, H.Y. Kim, H. Hosoda, S. Miyazaki, Scr. Mater. 72–73 (2014) 29–32. [19] H. Tobe, H.Y. Kim, T. Inamura, H. Hosoda, T.H. Nam, S. Miyazaki, J. Alloy. Compd. 577 (2013) S435–S438. [20] H. Kim, J.I. Kim, Mater. Res. Bull. 48 (2013) 5131–5135. [21] J.I. Kim, Y.C. Park, J.M. Ock, J. Alloy. Compd. 577 (2013) S453–S458. [22] L.W. Ma, H.S. Cheng, C.Y. Chung, B. Yuan, Mater. Sci. Eng. A 561 (2013) 427–433. [23] S.J. Kim, J.I. Kim, Mater. Res. Bull. 48 (2013) 5125–5130. [24] D.C. Zhang, Y.F. Mao, Y.L. Li, J.J. Li, M. Yuan, J.G. Lin, Mater. Sci. Eng. A 559 (2013) 706–710. [25] S. Guo, Q.K. Meng, X.N. Cheng, X.Q. Zhao, J. Mech. Behav. Biomed. Mater. 38 (2014) 26–32. [26] S. Dubinskiy, S. Prokoshkin, V. Brailovski, K. Inaekyan, A. Korotitskiy, Mater. Charact. 88 (2014) 127–142. [27] H.C. Hsu, S.C. Wu, S.K. Hsu, W.H. Kao, W.F. Ho, Mater. Sci. Eng. A 579 (2013) 86–91. [28] M. Tane, T. Nakano, S. Kuramoto, M. Niinomi, N. Takesue, H. Nakajima, Acta Mater. 61 (2013) 139–150. [29] F. Sun, Y.L. Hao, S. Nowak, T. Gloriant, P. Laheurte, F. Prima, J. Mech. Behav. Biomed. Mater. 4 (2011) 1864–1872. [30] M. Tahara, H.Y. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Acta Mater. 59 (2011) 6208–6218. [31] H.Y. Kim, Y. Ikehara, J.I. Kim, H. Hosoda, S. Miyazaki, Acta Mater. 54 (2006) 2419–2429. [32] H.Y. Kim, T. Sasaki, K. Okutsu, J.I. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Acta Mater. 54 (2006) 423–433. [33] H.Y. Kim, J.I. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Mater. Sci. Eng. A 438– 440 (2006) 839–843. [34] J.I. Kim, H.Y. Kim, H. Hosoda, S. Miyazaki, Mater. Trans. 46 (2005) 852–857. [35] M. Tahara, H.Y. Kim, H. Hosoda, S. Miyazaki, Funct. Mater. Lett. 2 (2009) 79–82. [36] Y. Fukui, T. Inamura, H. Hosoda, K. Wakashima, S. Miyazaki, Mater. Trans. 45 (2004) 1077–1082. [37] Y. Horiuchi, T. Inamura, H.Y. Kim, K. Wakashima, S. Miyazaki, H. Hosoda, Mater. Trans. 48 (2007) 407–413. [38] J.I. Kim, H.Y. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Mater. Sci. Eng. A 403 (2005) 334–339. [39] W.F. Ho, C.P. Ju, J.H. Chern Lin, Biomaterials 20 (1999) 2115–2122. [40] T. Grosdidier, M.J. Philippe, Mater. Sci. Eng. A 291 (2000) 218–223. [41] H.Y. Kim, Y. Ohmatsu, J.I. Kim, H. Hosoda, S. Miyazaki, Mater. Trans. 45 (2004) 1090–1095. [42] T. Maeshima, M. Nishida, Mater. Trans. 45 (2004) 1096–1100. [43] T. Maeshima, M. Nishida, Materials Trans. 45 (2004) 1101–1105. [44] T. Zhou, M. Aindow, S.P. Alpay, M.J. Blackburn, M.H. Wu, Scr. Mater. 50 (2004) 343–348. [45] Y. Sutou, K. Yamaguchi, T. Takagi, T. Maeshima, M. Nishida, Mater. Sci. Eng. A 438–440 (2006) 1097–1100. [46] F. Tang, S. Nakazawa, M. Hagiwara, Mater. Sci. Eng. A 315 (2001) 147–152.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of B addition on the microstructure and superelastic properties of a Ti-26Nb alloy Yazan Al-Zain a,n, Hee Young Kim b,nn, Shuichi Miyazaki b,c,nnn a
Department of Industrial Engineering, The University of Jordan, Amman 11942, Jordan Division of Materials Science, University of Tsukuba, Tsukuba, Ibaraki 305-8573, Japan c Foundation for Advancement of International Science, Tsukuba, Ibaraki 305-0821, Japan b
art ic l e i nf o
a b s t r a c t
Article history: Received 26 May 2015 Received in revised form 17 July 2015 Accepted 18 July 2015 Available online 20 July 2015
The mechanical properties and shape memory behavior of Ti–26Nb–(0–1.0)B and Ti–27Nb alloys were investigated. The stress for inducing martensite and the critical stress for slip deformation were greatly affected by increasing B content up to 0.1 at%, and less affected by further increase in B content up to 0.7 at%. The martensitic transformation start temperature decreased suddenly with a slope of 350 K/at% B when B content was less than 0.1 at% and decreased with a smaller slope of 15 K/at% B with further increase in B content up to 1.0 at%. In all samples tested, only superelasticity was observed at room temperature. TEM investigations revealed that the solubility of B in Ti–26Nb alloy was around 0.1 at%. In spite of the low solubility limit, B was effective to stabilize the superelastic behavior by effectively increasing the critical stress for slip deformation. On the other hand, while the stress for inducing martensite increased effectively by the addition of 1.0 at% Nb to the Ti–26Nb alloy, the critical stress for slip deformation remained almost constant and hence superelasticity was deteriorated. & 2015 Elsevier B.V. All rights reserved.
Keywords: Ti-based shape memory alloys Superelasticity Shape memory Martensitic transformation
1. Introduction Ti–Ni shape memory alloys (SMAs) belong to the family of the so-called smart materials. These alloys are capable of restoring their original shape when deformed followed by either load removal alone or heating after load removal. In the former case the property exhibited is referred to as superelasticity (SE) while in the latter case the property exhibited is referred to as shape memory effect (SME). The excellent SE and SME exhibited by Ti–Ni alloys [1,2] render them superior to other classes of SMAs. However, human body allergy to pure Ni has been pointed out in some reports [3]. Accordingly, scientists started searching for SMAs that are free of Ni. Several studies showed that SE and SME are exhibited in some β-type Ni-free Ti-alloys. Such examples include Ti–Nb-base alloys [4–38] and Ti–Mo-base alloys [39–45]. The SE and SME exhibited by the binary Ti–Nb alloys were not as good as those exhibited by the Ti–Ni alloys due to the relatively low critical stress for slip deformation (sCSS) and small transformation strain. Several researchers showed that the SE and SME of n
Corresponding author. Corresponding author. Corresponding author at: Division of Materials Science, University of Tsukuba, Tsukuba, Ibaraki 305-8573, Japan. E-mail addresses: [email protected] (Y. Al-Zain), [email protected] (H.Y. Kim), [email protected] (S. Miyazaki). nn
nnn
http://dx.doi.org/10.1016/j.msea.2015.07.054 0921-5093/& 2015 Elsevier B.V. All rights reserved.
the binary Ti–Nb alloys could be enhanced by the addition of a third or more elements. Examples are Ti–Nb–O [34], Ti–Nb–N [35], Ti–Nb–Mo [13], Ti–Nb–Al [36], Ti–Nb–Mo–Sn [15], Ti–Nb–(Ge, Ga) [8], Ti–Nb–Al–B [37], Ti–Nb–Ta [10], and Ti–Nb–Zr [38]. According to these reports, the addition of Zr, Ge, Sn, Al, O or N has effectively increased the superelastic strain. Moreover, the addition of O and N has effectively stabilized the SE by increasing the magnitude of sCSS. In this research, the effect of B addition on the martensitic transformation start temperature (Ms), microstructure, mechanical properties and SE of the Ti–26Nb alloy was systematically investigated. In addition, the SE of the B-added Ti–26Nb alloy was compared to that of the Ti–27Nb alloy.
2. Experimental procedure Ti–26 at% Nb–(0–1.0) at% B and Ti–27 at% Nb alloys (all compositions are hereafter described in atomic per cent ) were prepared by an Ar-arc melting method using high purity titanium, niobium and boron and melted six times to avoid segregation. The ingots were homogenized at 1273 K for 7.2 ks followed by coldrolling up to a 98.5% reduction in thickness. Specimens were cut using an electric arc discharge machine and then heat-treated at 1173 K for 1.8 ks in an Ar atmosphere, followed by quenching in water at room temperature (RT). In order to remove the lightly
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oxidized surface layer, all specimens were etched using a solution of H2O, HNO3 and HF (5:4:1) in a water bath at a temperature of about 313 K. The shape memory properties were characterized by a tensile testing machine at RT. The gage length and width of the specimens for tensile tests were 20 mm and 1.5 mm, respectively. The Ms was estimated from thermal cycling tests under various stress levels. Finally, microstructure was investigated using an optical microscope and the internal structure was observed by a transmission electron microscope (TEM), using a JOEL 2010F microscope operated at 200 kV.
3. Results and discussion 3.1. Effect of B content on microstructure and grain size. TEM investigations were carried out in order to investigate the internal structure of B-free and B-added Ti–26Nb alloys. Fig. 1 shows bright-field images of Ti–26Nb–(0, 0.1, 0.2 and 0.7)B alloys and selected-area diffraction patterns e and f taken inside the precipitates, corresponding to images c and d, respectively. It is noticed that the matrix is composed only of the parent β phase for samples containing less than 0.2 at% B, and there is no formation of precipitates. Adding a larger amount of B resulted in precipitation of TiB particles inside the β matrix. Moreover, the size and density of these precipitates increased with increasing B content above 0.2 at%. This indicates that the solubility limit of B in the Ti–26Nb alloy is about 0.1 at%. A specific orientation relationship was not identified between the matrix and the TiB
precipitates. Moreover, formation of the martensite phase was not observed in any of the specimens. This indicates that Ms for all the alloys is below RT. To investigate the effect of B addition on the average grain size (avg.), optical microscopy observation was carried out at RT. Fig. 2 shows the microstructure of Ti–26Nb–(0, 0.1, 0.2 and 0.7)B alloys. Generally speaking, avg. decreased by increasing B content from 0 to 0.7 at%. From avg. indicated on each image, it is clear that the grain size did not change by the addition of 0.1 at% B, and the small drop is considered to be within experimental errors. However, avg. dropped by higher additions up to 0.2 at% B. Finally, it decreased only slightly by further increase in B content above 0.2 at% B. It can be seen that most of the precipitates are found inside the grains while very few are found at the grain boundaries for the alloy containing 0.2 at% B. Similar observation was reported by Tang et al., they stated that after conducting solution treatment of βphase TiB precipitates were seen inside the grains [46]. In spite of the increase in the density of precipitates, similar observations were seen for the alloy containing 0.7 at% B, with very small number of precipitates found at the grain boundaries relative to the number of precipitates seen inside the grains. Hence, precipitation of borides does not have pronounced effect on the grain size. Moreover the reason why there was no profound change in the avg. when B content was between 0 and 0.1 at% (49 μm in the binary case and 47 μm in the 0.1 at% B-added alloy) is that there was no precipitation of borides when B content was less than 0.2 at%. This is consistent with the result of solubility limit investigated by TEM. Moreover, the current study and the result of the previous study on Ti–Nb–Al–B alloys [37] are in good
Fig. 1. Bright field images of the Ti–26Nb alloys with (a) 0 at% B, (b) 0.1 at% B, (c) 0.2 at% B and (d) 0.7 at% B heat-treated at 1173 K for 1.8 ks. (e) and (f) represent the diffraction patterns for TiB in (c) and (d), respectively.
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Fig. 2. Optical microscope images for the Ti–26Nb alloys with (a) 0 at% B, (b) 0.1 at% B, (c) 0.2 at% B and (d) 0.7 at% B heat-treated at 1173 K for 1.8 ks.
agreement in a way that B in the solid solution has no effect on the grain size. Also, both studies confirmed that the decrease in grain size was due to the pinning effect of borides at the grain boundaries. 3.2. Effect of B content on the martensitic transformation start temperature (Ms). Fig. 3 shows strain-temperature curves measured during thermal cycling between 173 K and 373 K under constant various stresses for the Ti–26Nb–0.7B alloy. The test was performed in a way such that the magnitude of the applied stress was increased with 50 MPa for every following cycle, the starting stress being 100 MPa. Same specimen was used throughout the test. The upper and lower halves of the curves represent heating and cooling cyles, respectively. The Ms at each curve was determined as illustrated on the figure.
Generally speaking, Ms increased with increasing the applied stress, since the application of the stress usually assists the transformation. The same test was conducted for all specimens with different B contents. Ms is usually taken at zero-stress using simple extrapolation technique shown in Fig. 4 (the solid line). Following the dashed lines, it can be seen that the degree of the increment in Ms by increasing the stress becomes smaller. This deviation is related to the plastic deformation (shown as (εp) in Fig. 3) introduced after a full cycle of forward transformation and reverse transformation. Accordingly, it is not possible in this case to determine Ms accurately by the extrapolation technique. To avoid such deviation, Ms is estimated from the first cycle. Recalling that the solubility limit of B in Ti–26Nb alloy does not exceed 0.1 at%, where TiB precipitates afterwards, the matrix content of Nb shall increase with about 0.45 at% when 1.0 at% B is added. In a previous report [31] it was shown that 1 at% addition of 300
Stress (MPa)
250 200 150 100 50 0 175 Fig. 3. Stress–temperature curves obtained at various constant stresses for the Ti– 26Nb–0.7B alloy heat-treated at 1173 K for 1.8 ks.
200
225 250 Temperature (K)
275
300
Fig. 4. Ms at various stress values for Ti–26Nb–0.7B alloys heat-treated at 1173 K for 1.8 ks.
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Table 1 sSIM, sCSS and εse for the alloys Ti–26 and 27 at% Nb and Ti–26 at% Nb–(0.1 and 0.7) at% B
Fig. 5. B content dependence of Ms for Ti–26Nb–(0–1.0)B alloys heat-treated at 1173 K for 1.8 ks.
Nb decreased Ms by 40 K. Hence, there must be a drop in Ms even if B becomes insoluble in the alloys. Fig. 5 shows the B content dependence of Ms for Ti–26Nb–(0– 1.0)B alloys. It is obvious that all the alloys posses Ms below RT, indicating that the phase at RT is the parent phase, this result is consistent with TEM results. Addition of B was effective to decrease Ms when its content was less than 0.1 at%: the effect of B on Ms was 350 K/at% B. This is due to the solid solution effect of B. Then Ms decreased with 15 K/at% B with further increase in B content. This is consistent with the discussion above; a 15 K decrease in Ms is very close to an increase of about 0.5 at% Nb in the matrix content. The results of Ms agree with the results shown by the previous study of Ti–Nb–Al–B [37] that B in solid solution was effective to decrease Ms. However, it disagrees with the effect of precipitation of B on Ms, as they showed that Ms remains constant irrespective of further increase in B content above the solubility limit. 3.3. Superelasticity and shape memory behavior of Ti-26Nb-(0-0.7)B and Ti–27Nb alloys. Tensile tests were carried out at RT in order to characterize the shape memory behvior of Ti–Nb–B alloys. Fig. 6 shows stress– strain curves for the Ti–26Nb, Ti–26Nb–0.1B, Ti–26Nb–0.7B and Ti–27Nb alloys heat treated at 1173 K for 1.8 ks. Load was applied until reaching 2.5% strain and then removed. A solid circle on each curve represents the apparent yield stress (In this study, the yield stress is set to be the stress at 0.2% offset strain), and the dashed
Fig. 6. Stress–strain curves for (a) Ti–26Nb (b) Ti–27Nb, (c) Ti–26Nb–0.1B and (d) Ti–26Nb–0.7B alloys heat-treated at 1173 K for 1.8 ks.
Alloy
sSIM (MPa)
sCSS (MPa)
εse (%)
Ti–26Nb Ti–27Nb Ti–26Nb–0.1B Ti–26Nb–0.7B
101 223 168 225
356 379 437 502
2.0 1.7 2.5 2.6
lines with arrows represent strain recovered as a result of heating. As shown earlier, Ms of all the alloys are below RT. Hence, the apparent yield stress is considered to be the stress for inducing the martensite phase (sSIM) from the parent β phase. In all specimens, most of shape recovery took place upon unloading, while a small amount of the permanent strain was recovered by heating after unloading. The magnitudes of sSIM for all alloys are summarized in Table 1. Taking sSIM of the Ti–26Nb alloy as the reference, it is seen that the magnitude of sSIM increased more severely by the addition of 1 at% Nb when compared to the increment caused by the addition of 0.1 at% B. It could be also seen that addition of more than 0.1 at% B to the binary alloy further increased sSIM. However, the effect of B addition on sSIM became less severe as B content increased above 0.1 at%. The increase in the value of the sSIM by addition of 1.0 at% Nb and 0.1 at% B to the binary Ti–26Nb alloy is a result of the solid solution effect of Nb and B on Ms. The slight increase in the magnitude of sSIM by further increasing B content above 0.1 at% is related to the precipitation effect of TiB and the corresponding enrichment of the matrix with Nb. 3.4. Effect of B addition on the critical stress for slip deformation (CSS). Fig. 7 shows stress–strain curves for the Ti–26Nb, Ti–26Nb– 0.1B, Ti–26Nb–0.7B and Ti–27Nb alloys heat-treated at 1173 K for 1.8 ks. Each stress–strain curve and strain recovery were obtained by a loading and unloading cycle. The test was repeated by increasing the maximum strain upon loading in the same specimen for each alloy. Heating was conducted after each cycle showing permanent strain after unloading. In Fig. 7b, εse and εp represent the superelastic recovery strain due to unloading and the plastic strain remained after heating, respectively. As shown earlier, all alloys exhibited SE. Residual strain was completely recovered by simple unloading up to the second cycle for Ti–26Nb alloy. The sCSS in this study is defined as the stress at which 0.5% plastic strain is introduced during cycling as shown in Fig. 7b. The values of sCSS for all alloys are shown in Table 1. It is clear that addition of B to the Ti–26Nb alloy was effective to increase the sCSS. On the other hand it remained almost unchanged with the addition of 1 at% Nb. This is an expected result, recalling that the atomic diameter of Nb is almost similar to that of Ti. Hence, solid solution hardening must be too small. In contrast, the atomic diameter of B is much smaller than that of Ti, giving a rise in the solid solution hardening effect to effectively increase the sCSS. The sCSS of Ti– 26Nb alloy was also increased by the addition of 0.7 at% B. For this alloy, only 0.1 at% B was soluble while 0.6 at% B precipitated in the form of TiB. These precipitates were also effective to increase the sCSS as they serve as obstacles hindering dislocation motion during deformation. However, it is seen that B in solid solution was more effective to increase the sCSS when compared to B in the form of TiB. Finally, as shown in Table 1, the εse decreased by the addition of 1 at% Nb to the Ti–26Nb. This can be explained by the large increase in the sSIM while having an almost constant value of sCSS. On the other hand, εse was increased by the addition of B. Similar values of εse were seen for both B-added alloys. SE was enhanced
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Fig. 7. Stress–strain curves obtained by cycling loading–unloading tensile tests for (a) Ti–26Nb (b) Ti–27Nb, (c) Ti–26Nb–0.1B and (d) Ti–26Nb–0.7B alloys heat-treated at 1173 K for 1.8 ks.
due to the larger increase in sCSS when compared to the increase in the sSIM, allowing transformation to proceed freely before slip deformation occurred. It is concluded that although the solubility limit of B in Ti–26Nb did not exceed 0.1 at%, B was good to stabilize the SE of Ti–Nb shape memory alloys.
4. Conclusions [1] Ti–27 at% Nb and Ti–26 at% Nb–(0–1.0) at% B alloys were fabricated by the Ar-arc melting method. SE was observed in the alloys. B was found to have a very small solubility in the binary Ti– 26Nb alloy with about 0.1 at%. Moreover, it was found that B was most effective to decrease the Ms when B content was less than 0.1 at%. The sSIM and sCSS were strongly affected by the addition of 0.1 at% B to the Ti–26Nb alloy due to solid solution effect; further increase in B content caused the change in these stresses to become smaller as a result of precipitation. B addition was effective to increase εse. Due to this reason and because sCSS was effectively increased, it is concluded that B addition is effective in stabilizing the SE of Ti–Nb alloys. [2] Addition of 1 at% Nb to the Ti–26Nb alloy resulted in the deterioration of the SE. This is because there was a large increase in the sSIM while sCSS remained almost unchanged.
Acknowledgment This work was partially supported by JSPS KAKENHI Grant numbers 26249104 and 25102704, and Fiscal 2014 JASSO Followup Research Fellowship.
References [1] S. Miyazaki, Y. Ohmi, K. Otsukua, Y. Suzuki, J. Phys. 43 (1982) 255–260. [2] S. Miyazaki, K. Otsuka, Trans. Iron Steel Inst. Jpn. 29 (1989) 353–377. [3] S. Shabalovskaya, J. Cunnick, J. Anderegg, B. Harmon, R. Sachdeva, in: Proceedings of the First International Conference on Shape Memory and Superelastic Technologies, Pacific Grove, California, USA, 1994, pp. 209–214. [4] E. Takahashi, T. Sakurai, S. Watanabe, N. Masahashi, S. Hanada, Mater. Trans. 43 (2002) 2978–2983. [5] H. Hosoda, Y. Fukui, T. Inamura, K. Wakashima, S. Miyazaki, K. Inoue, Mater. Sci. Forum 425–432 (2003) 3121–3126. [6] H. Hosoda, Y. Kinoshita, Y. Fukui, T. Inamura, K. Wakashima, H.Y. Kim, S. Miyazaki, Mater. Sci. Eng. A 438–440 (2006) 870–874. [7] Y.L. Hao, S.J. Li, S.Y. Sun, R. Yang, Mater. Sci. Eng. A 441 (2006) 112–118. [8] T. Inamura, Y. Fukui, H. Hosoda, K. Wakashima, S. Miyazaki, Mater. Sci. Eng. C 25 (2005) 426–432. [9] N. Sakaguchi, M. Niinomi, T. Akahori, Mater. Trans. 45 (2004) 1113–1119.
[10] H.Y. Kim, S. Hashimoto, J.I. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Mater. Sci. Eng. A 417 (2006) 120–128. [11] Y.B. Wang, Y.F. Zheng, Mater. Lett. 62 (2008) 269–272. [12] B.L. Wang, Y.F. Zheng, L.C. Zhao, Mater. Sci. Eng. A 486 (2008) 146–151. [13] Y. Al-Zain, H.Y. Kim, H. Hosoda, T.H. Nam, S. Miyazaki, Acta Mater. 58 (2010) 4212–4223. [14] Y. Al-Zain, H.Y. Kim, T. Koyano, H. Hosoda, T.H. Nam, S. Miyazaki, Acta Mater. 59 (2011) 1464–1473. [15] Y. Al-Zain, Y. Sato, H.Y. Kim, H. Hosoda, T.H. Nam, S. Miyazaki, Acta Mater. 60 (2012) 2437–2447. [16] Y. Al-Zain, H.Y. Kim, T. Koyano, H. Hosoda, S. Miyazaki, Scr. Mater. 103 (2015) 37–40. [17] M. Tahara, H.Y. Kim, H. Hosoda, S. Miyazaki, Acta Mater. 57 (2009) 2461–2469. [18] M.F. Ijaz, H.Y. Kim, H. Hosoda, S. Miyazaki, Scr. Mater. 72–73 (2014) 29–32. [19] H. Tobe, H.Y. Kim, T. Inamura, H. Hosoda, T.H. Nam, S. Miyazaki, J. Alloy. Compd. 577 (2013) S435–S438. [20] H. Kim, J.I. Kim, Mater. Res. Bull. 48 (2013) 5131–5135. [21] J.I. Kim, Y.C. Park, J.M. Ock, J. Alloy. Compd. 577 (2013) S453–S458. [22] L.W. Ma, H.S. Cheng, C.Y. Chung, B. Yuan, Mater. Sci. Eng. A 561 (2013) 427–433. [23] S.J. Kim, J.I. Kim, Mater. Res. Bull. 48 (2013) 5125–5130. [24] D.C. Zhang, Y.F. Mao, Y.L. Li, J.J. Li, M. Yuan, J.G. Lin, Mater. Sci. Eng. A 559 (2013) 706–710. [25] S. Guo, Q.K. Meng, X.N. Cheng, X.Q. Zhao, J. Mech. Behav. Biomed. Mater. 38 (2014) 26–32. [26] S. Dubinskiy, S. Prokoshkin, V. Brailovski, K. Inaekyan, A. Korotitskiy, Mater. Charact. 88 (2014) 127–142. [27] H.C. Hsu, S.C. Wu, S.K. Hsu, W.H. Kao, W.F. Ho, Mater. Sci. Eng. A 579 (2013) 86–91. [28] M. Tane, T. Nakano, S. Kuramoto, M. Niinomi, N. Takesue, H. Nakajima, Acta Mater. 61 (2013) 139–150. [29] F. Sun, Y.L. Hao, S. Nowak, T. Gloriant, P. Laheurte, F. Prima, J. Mech. Behav. Biomed. Mater. 4 (2011) 1864–1872. [30] M. Tahara, H.Y. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Acta Mater. 59 (2011) 6208–6218. [31] H.Y. Kim, Y. Ikehara, J.I. Kim, H. Hosoda, S. Miyazaki, Acta Mater. 54 (2006) 2419–2429. [32] H.Y. Kim, T. Sasaki, K. Okutsu, J.I. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Acta Mater. 54 (2006) 423–433. [33] H.Y. Kim, J.I. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Mater. Sci. Eng. A 438– 440 (2006) 839–843. [34] J.I. Kim, H.Y. Kim, H. Hosoda, S. Miyazaki, Mater. Trans. 46 (2005) 852–857. [35] M. Tahara, H.Y. Kim, H. Hosoda, S. Miyazaki, Funct. Mater. Lett. 2 (2009) 79–82. [36] Y. Fukui, T. Inamura, H. Hosoda, K. Wakashima, S. Miyazaki, Mater. Trans. 45 (2004) 1077–1082. [37] Y. Horiuchi, T. Inamura, H.Y. Kim, K. Wakashima, S. Miyazaki, H. Hosoda, Mater. Trans. 48 (2007) 407–413. [38] J.I. Kim, H.Y. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Mater. Sci. Eng. A 403 (2005) 334–339. [39] W.F. Ho, C.P. Ju, J.H. Chern Lin, Biomaterials 20 (1999) 2115–2122. [40] T. Grosdidier, M.J. Philippe, Mater. Sci. Eng. A 291 (2000) 218–223. [41] H.Y. Kim, Y. Ohmatsu, J.I. Kim, H. Hosoda, S. Miyazaki, Mater. Trans. 45 (2004) 1090–1095. [42] T. Maeshima, M. Nishida, Mater. Trans. 45 (2004) 1096–1100. [43] T. Maeshima, M. Nishida, Materials Trans. 45 (2004) 1101–1105. [44] T. Zhou, M. Aindow, S.P. Alpay, M.J. Blackburn, M.H. Wu, Scr. Mater. 50 (2004) 343–348. [45] Y. Sutou, K. Yamaguchi, T. Takagi, T. Maeshima, M. Nishida, Mater. Sci. Eng. A 438–440 (2006) 1097–1100. [46] F. Tang, S. Nakazawa, M. Hagiwara, Mater. Sci. Eng. A 315 (2001) 147–152.