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Excellent superelasticity and fatigue resistance of Cu-Al-Mn-W shape memory single crystal obtained only through annealing polycrystalline cast alloy
Materials Science & Engineering A 749 (2019) 249–254
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Excellent superelasticity and fatigue resistance of Cu-Al-Mn-W shape memory single crystal obtained only through annealing polycrystalline cast alloy
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Shuiyuan Yanga, , Jixun Zhangb, Xinren Chenb, Mengyuan Chib, Cuiping Wangb, Xingjun Liuc,b a
College of Materials of Xiamen University, Shenzhen Research Institute of Xiamen University, PR China Fujian Key Laboratory of Materials Genome, Xiamen University, Xiamen 361005, PR China c State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology University, (Shenzhen) Guangdong 518055, PR China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Abnormal grain growth Single crystal Phase separation Superelasticity Fatigue resistance
In this study, the cast Cu-13.96Al-9.84Mn-0.51W shape memory alloy consisted of L21-Cu2AlMn parent with completely coherent bcc A2(W) nanoparticles due to bcc phase separation. The abnormal grain growth happened only through annealing cast polycrystalline alloy. The results showed that a continuous misorientation gradient within the matrix grains with a deviation of below 2° formed after the dissolution of A2(W) nanoparticles during annealing. This continuous misorientation gradient may result in the abnormal grain growth. The single crystal specimens were directly obtained from one of large grains. The single crystal close to [100] orientation exhibited very excellent superelasticity and fatigue resistance. Here full superelasticity up to 12% and the largest superelastic strain up to 8.5% were displayed, being comparable to commercial Ni-Ti shape memory alloys. Furthermore, the single crystal almost remained full superelasticity during 5 × 103 cycles. From the points of raw material cost, preparation technology and functional properties, the Cu-Al-Mn-W single crystal shows good potential as new superelastic material.
1. Introduction Cu-Al-based shape memory alloys (SMAs) are commercially functional materials in various engineering fields as actuators or sensors. Because of easier fabrication, lower manufacturing cost, and excellent conductivity of heat and electricity in Cu-Al-based SMAs [1,2], they have been attracting the attention of many researchers all the time. The main Cu-Al-based SMAs are Cu-Al-Ni and Cu-Al-Be alloys developed from binary Cu-Al alloys through alloying [1–6]. However, the development and application of Cu-Al-based SMAs are still very slow. One of main reasons is their polycrystalline brittleness that seriously restricts the production, processing and application of alloys [1,5]. The brittleness of Cu-Al-based SMAs is closely related to coarse grains, high elastic anisotropies, and precipitates along grain boundaries. Thus the grain refinement is the most important method to deal with this shortcoming so far [7–10]. The metastable ferromagnetic Heusler (L21 structure) Cu-Al-Mn alloys are also one of Cu-Al-based SMAs. Generally, a two-stage orderdisorder transformation of A2→B2→L21 occurs during cooling for CuAl-Mn SMAs [11,12]. The high ordered degree of L21 parent is also a
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key factor for the polycrystalline brittleness of Cu-Al-Mn SMAs. Recently, Omori et al. reported that the abnormal grain growth (AGG) of Cu-17Al-11.4Mn (at%) [13,14] and Fe-34Mn-15Al-7.5Ni (at%) [15,16] alloys. The ultra-large single crystals of more than a few centimeters could be obtained both in two systems. Their method was cyclic heat treatment between single bcc phase region at high temperature and two-phase of (bcc + fcc) region at low temperature. The AGG phenomenon resulted from the formation of sub-grains (about 100 µm) induced by the coherency loss of fcc precipitates during thermal cycling. Their finding greatly promotes the development of Cu-Al-Mn SMAs. The preparation of Cu-Al-Mn single crystals through cyclic heat treatment is much simpler comparing to the directional solidification [17–20]. However, the requirements of this method are still very strict, including the anneal temperature, cooling rate and numbers of thermal cycle. Our research group recently designed a unique microstructure of cast Cu-Al-Mn-Mo SMAs based on the idea of bcc two-phase separation [21]. The cast Cu-Al-Mn-Mo alloys consisted of L21-Cu2AlMn parent with completely coherent bcc A2(Mo) nanoparticles. Interestingly, the AGG directly happened only through annealing their polycrystalline
Corresponding author. E-mail address: [email protected] (S. Yang).
https://doi.org/10.1016/j.msea.2019.02.033 Received 27 December 2018; Received in revised form 9 February 2019; Accepted 9 February 2019 Available online 11 February 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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coherent with bcc A2(W) nanoparticles due to bcc phase separation (Fig. 2c). Additionally, it was found from the composition mapping analysis in Fig. 2e - i, the L21 parent was the Cu-rich, the bcc A2(W) nanoparticles were the W-rich, and the γ1(Cu9Al4) phase is the Al-rich and Mn-poor.
cast alloys. At this time, the large grains and single crystals over cmscale were obtained. Our results further confirmed that the AGG may be closely related to a continuous misorientation gradient within the matrix grains resulted from the dissolution of A2(Mo) nanoparticles. We consider that W and Mo both have A2 structure, and little solid solubility in Cu. Thus we speculate that bcc two-phase separation may also occur in Cu-Al-Mn-W alloys, and similar microstructure of cast alloys may be obtained. Therefore in this study, the Cu-14.2Al-10Mn-0.5W (wt%) alloy was designed, and the microstructure and abnormal grain growth of the alloy, superelasticity and fatigue resistance of single crystals were investigated.
3.3. Microstructural evolution during annealing In order to investigate the relationship between the AGG and the γ1(Cu9Al4) phase, the cast sample was annealed at 1023 K for 5 min and quenched, the γ1(Cu9Al4) phase had dissolved to the matrix (Fig. 3b, e and g). However, the AGG did not take place at this time. The result indicates that the existence of γ1(Cu9Al4) phase does not apply any influence to the AGG. According to the previous results, bcc two-phase separation of Cu-Al-Mn-Mo alloys happened around 1073 K [21]. Thus it is speculated that the AGG of Cu-Al-Mn-W alloy should occur only after the dissolution of A2(W) nanoparticles around 1073 K. The results of Fig. 3c, d and f confirm this speculation. When the above cast sample was annealed at 1173 K for 10 min, the AGG took place (Fig. 3c), and large grains formed after 2 h (Fig. 3d and f). Our previous investigation revealed that a continuous misorientation gradient within the matrix grains due to the dissolution of bcc precipitates was responsible for the AGG in Cu-Al-Mn-based alloys with bcc phase separation [21]. Fig. 4 shows the EBSD quasi-colored orientation mapping (OM) and the corresponding grain reference orientation deviation (GROD) of the cast alloy after annealing. When the cast alloy was annealed at 1173 K for 5 min, the γ1(Cu9Al4) phase disappeared, but the AGG still did not happen. At this time, the orientation of each grain was random (Fig. 4a), and no misorientation existed among the grains (Fig. 4b). After annealing for 10 min, few large grains formed (Fig. 4c). Here, there were no any sub-grains that reported by Omori et al. in Cu-Al-Mn ternary alloy [13,14]. Similar to those of CuAl-Mn-Mo alloys [21], a continuous misorientation gradient below 2° deviation was observed from grain core to grain boundaries (Fig. 4d).
2. Experimental Button ingot (40 g) of nominal Cu-14.2Al-10.0Mn-0.5W (wt%) alloy was prepared through arc-melting. The cast ingot was sealed into vacuum quartz tubes under argon atmosphere, then annealed at 1173 K for 2 h and followed by water quenching. The microstructure and crystal structure were observed and analyzed by optical microscopy, transmission electron microscopy (TEM) and the selected area electron diffraction (SAED). The chemical composition was measured by energy dispersive spectrometer (EDS) with an average value of five measurements. The actual composition was Cu-13.96Al-9.84Mn-0.51W. The composition mapping analysis of each element was determined by electron probe microanalysis (EPMA). The crystallographic orientation and the orientation of the single crystals were determined by electron backscatter diffraction (EBSD) using a field emission scanning electron microscope (FE-SEM). The superelastic properties were measured by compression test using cylindrical single crystal specimens (3 mm diameter, 5 mm height). The specimen was compressed to different pre-strains (εpre) and then unloading to a zero stress condition. The SE (εSE) strain was confirmed and measured directly by stress-strain curves [22]. The height of the sample was measured before loading (h0) and after unloading (h1). The residual strain (εr) was calculated as εr = (h0 - h1)/h0 × 100%. The εe implies the elastic strain when unloading to a zero stress condition. Therefore, the εSE strain was calculated using the formulas εSE = (εpre εe - εr). The fatigue resistance was carried out by compression test with 5 × 103 cycles under a constant applied deformation of 8.5%.
3.4. Superelasticity and fatigue resistance of single crystal Cylindrical single crystal specimens were directly prepared from one of large grains (Fig. 5a). The martensitic transformation temperature was measured by the differential scanning calorimetry (DSC) at a heating and cooling rate of 10 °C min−1. However, the transformation temperature of the single crystal was not detected from room temperature to −140 °C. It indicates that the transformation temperature may be below −140 °C or the transformation enthalpy is too small to detect. The orientation of single crystal along the compression direction was close to [100] (with the deviation of about 7.8°) (Fig. 5b). The fracture strain and stress were 13.2% and 706 MPa, respectively. Exhilaratingly, a very large stress plateau corresponding to the stress-induced martensite was observed. It indicates that this single crystal should possess excellent SE properties. Complete and perfect superelasticity up to 12% and the largest SE strain up to 8.5% were obtained in Fig. 6. Its SE property is compared with Ni-Ti SMAs [28–30], and CuZn-Al, Cu-Al-Ni as well as Fe-Mn-Al-Ni single crystals [3,4,31], see Table 1 [32–37]. Furthermore, the single crystal possesses excellent fatigue resistance as shown in Fig. 7. A stress-strain cycle of 5 × 103 times with a constant pre-strain of 8.5% was carried out (Fig. 7a). During cycling, this single crystal always exhibited full superelasticity behavior. Full superelasticity of 8.5% and SE strain of 6.8% were displayed during the first cycle. After 5 × 102 cycles and 5 × 103 cycles, full superelasticity decreased to 8.1% and 6.5%, the SE strain decreased to 5.9% and 4.3% (Fig. 7b and c). The superelasticity property begun to noticeably degrade after about 3000 cycles (Fig. 7d).
3. Results 3.1. Abnormal grain growth Fig. 1 shows the macroscopic morphologies of Cu-13.96Al-9.84Mn0.51W alloy before and after annealing at 1173 K for 2 h followed by water quenching. When the polycrystalline cast alloy was annealed at 1173 K, the AGG happened. At this time, only two grains exceptionally grew. As a result, large grains and single crystals over cm-scale were obtained. 3.2. Microstructure of the cast alloy Fig. 2 shows the microstructure of the cast Cu-13.96Al-9.84Mn0.51 W polycrystalline alloy. The cast alloy consisted of three-phase of L21 parent + γ1(Cu9Al4) phase + A2(W) phase. The A2(W) phase had two morphologies, visibly large particles (Fig. 2a) and fine nanoparticles (Fig. 2b). Furthermore, the L21 parent was completely
Fig. 1. Abnormal grain growth of Cu-13.96Al-9.84Mn-0.51 W alloy. 250
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Fig. 2. Microstructure of Cu-13.96Al-9.84Mn-0.51W cast alloy. (a) Optical microscopy. (b) TEM bright-field (BF) image. (c) Corresponding SAED pattern of L21 matrix and bcc A2(W) nanoparticles (solid circle in b). (d) Corresponding SAED pattern of γ1(Cu9Al4) phase (dashed circle in b). (e - i) Composition mapping analysis of each element.
Fig. 3. Microstructural evolution of Cu-13.96Al-9.84Mn-0.51 W alloy during annealing. (a) Cast sample. (b–d) The same cast sample (in a) was annealed at 1023 K for 5 min (b), 1173 K for 10 min (c), and 1173 K for 2 h (d). (e) Optical metallograph accords to the sample in b. (f) Optical metallograph accords to the sample in d. (g) X-ray diffraction patterns of the cast sample (black) and the sample annealed at 1023 K for 5 min (red).
4. Discussion
Cu-rich bcc#1 phase + W-rich bcc#2 phase (those A2(W) nanoparticles). Then the Cu-rich bcc#1 phase transforms to the ordered L21Cu2AlMn matrix via the order-disordered transformation of A2 → B2, and B2 → L21 during the process of cooling to room temperature [11,24–27]. In this study, the normal grain growth may be inhibited in the cast state due to the A2(W) particle pinning. When the cast alloy is annealed at high enough temperature, those fine A2(W) particles dissolve back into the matrix. At this time, the particle pinning is eliminated, which may promote a few grains to grow exceptionally. Additionally, the formation of continuous misorientation gradient within the matrix grains may provide extra driving force for the rapid migration of grain boundaries. More details about the AGG of Cu-Al-Mn-based alloy with bcc phase separation have been explained in our previous report [21].
4.1. Abnormal grain growth From the results of Fig. 2, it is confirmed that the microstructure of the cast Cu-13.96Al-9.84Mn-0.51W alloy is similar to those of Cu-AlMn-Mo alloys [21]. The bcc two-phase separation happens in Cu13.96Al-9.84Mn-0.51W alloy. Combining with the phase diagram of binary Cu-W system [23], it is possible that a liquid separation region of Cu-rich and W-rich liquids exists in the Cu-Al-Mn-W system. During the solidification, the W-rich liquid solidifies into visibly large A2(W) particles, and the Cu-rich liquid solidifies into the Cu-rich bcc phase. With further decreasing temperature, the solidified Cu-rich bcc phase undergoes bcc two-phase separation. That is the Cu-rich bcc phase → 251
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Fig. 4. Crystallographic orientations of Cu-13.96Al-9.84Mn-0.51 W alloy. (a) Quasi-colored orientation mapping (OM) of the alloy annealed at 1173 K for 5 min and water quenching. (b) Grain reference orientation deviation (GROD) mapping corresponds to (a). (c) OM of the alloy annealed at 1173 K for 10 min and water quenching. (d) GROD mapping corresponds to (c). In OM figures (a and c), the colors of crystal direction is given in the inverse pole figure. In GROD mappings (b and d), the color value of every pixel is calculated as the misorientation angle of this pixel with respect to a reference average orientation of one grain.
9.84Mn-0.51W single crystal and only 8% for Cu-12.0Al-9.0Mn-0.5Mo single crystal. Secondly, although the orientation of single crystals is both close to [100] direction, the degree of deviation is different. The deviation is 7.8° for Cu-13.96Al-9.84Mn-0.51W single crystal, whereas is larger 10.4° Cu-12.0Al-9.0Mn-0.5Mo single crystal. Thirdly, the strength of two single crystals is much different. For Cu-13.96Al9.84Mn-0.51W single crystal, the critical stress to induce martensite is about 220 MPa, and the fracture strength is about 706 MPa. However, the critical stress and fracture strength are much higher in Cu-12.0Al9.0Mn-0.5Mo single crystal. For Cu-Al-Mn ternary SMAs, the strength of both polycrystalline alloys and single crystals is low [17–19], thus the fatigue resistance of the alloys is relatedly poor [19]. The advisable increase of alloy strength is beneficial to the improvement of alloy properties. However, when the critical stress is too high, even close to
4.2. Superelasticity and fatigue resistance As we all known, the SE characteristics of shape memory single crystals are closely related to their orientations. When the orientation is different, the single crystal with the same composition may exhibit completely different SE characteristics. In the previous report, the SE properties of Cu-Al-Mn-Mo single crystals with several different orientations were studied [21]. Therefore, here only the SE properties of Cu-13.96Al-9.84Mn-0.51W and Cu-12.0Al-9.0Mn-0.5Mo single crystals with close to [100] orientation are discussed. In this study, a higher SE strain of 8.5% was obtained in Cu-13.96Al9.84Mn-0.51W single crystal than 4.8% of Cu-12.0Al-9.0Mn-0.5Mo single crystal [21]. There are three possible reasons for this issue. Firstly, the applied pre-strains are different, 12% for Cu-13.96Al-
Fig. 5. Mechanical properties of Cu-13.96Al-9.84Mn-0.51W single crystal. (a) Schematic illustrations of the tested single crystal and the analysis locations for the orientation of single crystal. (b) The orientations of single crystal respectively accord to (A) and (B) locations in (a) along the compression direction. (c) Compressive stress-strain curve. 252
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the yield strength or fracture strength of the alloy, the plastic deformation is also more likely to occur during the stress-induced martensite. At this time, the SE properties begin to decrease. With regard to the reason that resulted in the strength of single crystals is so different with different additions of Mo or W into Cu-Al-Mn alloys is not clear yet. This issue needs more investigations. H. Kato, et al. reported the cyclic stress-strain response of superelastic Cu-9.8Al-9.9Mn (wt%) shape memory single crystal with [001] orientation during a strain-cycling of 7.8% over a thousand cycles [19]. Their results revealed that the critical stress to induce the martensitic transformation gradually decreased at an almost constant rate from about 130 MPa to zero only after about 50 cycles. Because the residual martensite was accumulated during repeated cycling, the residual strain obviously increased and the SE strain decreased. The fatigue behavior is related to the dislocation formed during cycling. In the present study, several completely coherent A2(W) nanoparticles exist within L21 parent for Cu-Al-Mn-W single crystal before deformation. Thus the critical stress (above about 220 MPa) to induce martensite is higher than that of Cu-9.8Al-9.9Mn single crystal. Although the critical stress gradually decreases while increasing the cycles (Fig. 7a–c), the rate of decrease is clearly smaller. The existence of completely coherent A2(W) nanoparticles and the improvement of strength may make the alloy not easy to produce the defects or dislocation during repeated cycling, so the fatigue properties of Cu-Al-Mn-W single crystal can be improved. Because the cost of Cu-based polycrystalline SMAs is only one-tenth that of Ni-Ti-based SMAs, thus the Cu-based SMAs have wide applications. However, the polycrystalline brittleness of Cu-based alloys seriously hinder their further development so far. Additionally, the Cubased single crystals that prepared through directional solidification are not only very expensive but also time-consuming. In this study, from the points of raw material cost and preparation of single crystal, the present Cu-Al-Mn-W single crystal fabricated only through annealing cast polycrystalline alloy shows more superiority of low cost and simple technology. Furthermore, the present Cu-Al-Mn-W single crystal possesses full superelasticity up to 12% and the largest SE strain up to
Fig. 6. Superelasticity of Cu-13.96Al-9.84Mn-0.51W single crystal.
Table 1 Superelastic strain in various shape memory alloys. The symbol (†) indicates that the properties are obtained in single crystal. Alloy composition Cu-13.96Al-9.84Mn-0.51 W Fe-34Mn-15Al-7.5Ni Ti-42.4Ni-8Cu Ti-51.3Ni Ti-24Zr-10Nb-2Sn Cu68.4Al27.8Ni3.8 Cu67.9Zn16.1Al16 Cu71.9Al16.6Mn9.3Ni2B0.2 Cu-Al16.915-Mn10.4475Co0.5 Cu-24.25Zn-8.57Al-0.07Zr Cu-17.0Al-11.4Mn Zu45-Au30-Cu25 Co-32Ni-28Al Ni51-Fe22-Ga27
SE (%) †
8.5 9.5†/5.2 7.0 6.0 6.0 8.6† 8.5† 7.5 7.1 4.5 6.5†/5† 8.0 6.3 6.2†
References This work [30] [28] [29] [31] [3] [4] [32] [33] [34] [17,18] [35] [36] [37]
Fig. 7. Fatigue resistance of Cu-13.96Al-9.84Mn-0.51 W single crystal. (a) Stress-strain curves of 5 × 103 cycles with a constant pre-strain of 8.5%. (b and c) Stressstrain curves of 5 × 102 cycles (b) and 5 × 103 cycles (c). (d) Relationship between full superelasticity, SE strain and number of cycle. 253
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8.5%, and excellent fatigue resistance, being compared to Ni-Ti-based SMAs. Therefore, the Cu-Al-Mn-W single crystal shows good potential as new superelastic material.
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5. Conclusions In summary, a novel microstructure of cast Cu-Al-Mn-W SMA was designed. The cast alloy consisted of L21-Cu2AlMn parent with completely coherent bcc A2(W) nanoparticles. This microstructure induced a continuous misorientation gradient within the matrix grains from grain core to grain boundary during annealing, which might result in the AGG. Furthermore, the present Cu-Al-Mn-W shape memory single crystal exhibited excellent SE properties and fatigue resistance. So that our findings may enable wide applications of low-cost Cu-based SE materials through various engineering fields as advanced functional materials. Acknowledgements This work was supported by the financial supports of the Shenzhen Science and Technology Project, Grant no. JCYJ20170306142550151, the Natural Science Foundation of Fujian Province, China, Grant no. 2018J01078. References [1] K. Otsuka, C.M. Wayman (Eds.), Shape Memory Materials, Cambridge Univ. Press, Cambridge, 1998, pp. 1–48. [2] R. Dasgupta, A look into Cu-based shape memory alloys: present scenario and future prospects, J. Mater. Res. 29 (2014) 1681–1698. [3] K. Otsuka, C.M. Wayman, K. Nakai, H. Sakamoto, K. Shimizu, Superelasticity effects and stress-induced martensitic transformations in CuAlNi alloys, Acta Metall. 24 (1976) 207–226. [4] T. Saburi, Y. Inada, S. Nenno, N. Hori, Stress-induced martensitic transformations in Cu-Zn-Al and Cu-Zn-Ga alloys, J. De. Phys. C4 43 (1982) 633–638. [5] V.H.C. de Albuquerque, T.A. de, A. Melo, R.M. Gomes, S.J.G. de Limaa, J.M.R.S. Tavares, Grain size and temperature influence on the toughness of a CuAlBe shape memory alloy, Mater. Sci. Eng. A 528 (2010) 459–466. [6] X.J. Liu, I. Ohnuma, R. Kainuma, K. Ishida, Phase equilibria in the Cu-rich portion of the Cu–Al binary system, J. Alloy. Compd. 264 (1998) 201–208. [7] V. Sampath, Studies on the effect of grain refinement and thermal processing on shape memory characteristics of Cu–Al–Ni alloys, Smart Mater. Structs. 14 (2005) S253–S260. [8] V. Sampath, Improvement of shape-memory characteristics and mechanical properties of copper–zinc–aluminum shape-memory alloy with low aluminum content by grain refinement, Mater. Man. Proc. 21 (2006) 789–795. [9] P. Zhang, A.B. Ma, S. Lu, G.G. Liu, P.H. Lin, J.H. Jiang, C.L. Chu, Effect of grain refinement on the mechanical properties of Cu–Al–Be–B shape memory alloy, Mater. Des. 32 (2011) 348–352. [10] J. Yang, Q.Z. Wang, F.X. Yin, C.X. Cui, P.G. Ji, B. Li, Effects of grain refinement on the structure and properties of a CuAlMn shape memory alloy, Mater. Sci. Eng. A 664 (2016) 215–220. [11] R. Kainuma, N. Satoh, X.J. Liu, I. Ohnuma, K. Ishida, Phase equilibria and Heusler phase stability in the Cu-rich portion of the Cu–Al–Mn system, J. Alloy. Compd. 266 (1998) 191–200. [12] M.O. Prado, P.M. Decorte, F. Lovey, Martensitic transformation in Cu-Mn-Al alloys, Scr. Mater. 33 (1995) 877–883.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Excellent superelasticity and fatigue resistance of Cu-Al-Mn-W shape memory single crystal obtained only through annealing polycrystalline cast alloy
T
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Shuiyuan Yanga, , Jixun Zhangb, Xinren Chenb, Mengyuan Chib, Cuiping Wangb, Xingjun Liuc,b a
College of Materials of Xiamen University, Shenzhen Research Institute of Xiamen University, PR China Fujian Key Laboratory of Materials Genome, Xiamen University, Xiamen 361005, PR China c State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology University, (Shenzhen) Guangdong 518055, PR China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Abnormal grain growth Single crystal Phase separation Superelasticity Fatigue resistance
In this study, the cast Cu-13.96Al-9.84Mn-0.51W shape memory alloy consisted of L21-Cu2AlMn parent with completely coherent bcc A2(W) nanoparticles due to bcc phase separation. The abnormal grain growth happened only through annealing cast polycrystalline alloy. The results showed that a continuous misorientation gradient within the matrix grains with a deviation of below 2° formed after the dissolution of A2(W) nanoparticles during annealing. This continuous misorientation gradient may result in the abnormal grain growth. The single crystal specimens were directly obtained from one of large grains. The single crystal close to [100] orientation exhibited very excellent superelasticity and fatigue resistance. Here full superelasticity up to 12% and the largest superelastic strain up to 8.5% were displayed, being comparable to commercial Ni-Ti shape memory alloys. Furthermore, the single crystal almost remained full superelasticity during 5 × 103 cycles. From the points of raw material cost, preparation technology and functional properties, the Cu-Al-Mn-W single crystal shows good potential as new superelastic material.
1. Introduction Cu-Al-based shape memory alloys (SMAs) are commercially functional materials in various engineering fields as actuators or sensors. Because of easier fabrication, lower manufacturing cost, and excellent conductivity of heat and electricity in Cu-Al-based SMAs [1,2], they have been attracting the attention of many researchers all the time. The main Cu-Al-based SMAs are Cu-Al-Ni and Cu-Al-Be alloys developed from binary Cu-Al alloys through alloying [1–6]. However, the development and application of Cu-Al-based SMAs are still very slow. One of main reasons is their polycrystalline brittleness that seriously restricts the production, processing and application of alloys [1,5]. The brittleness of Cu-Al-based SMAs is closely related to coarse grains, high elastic anisotropies, and precipitates along grain boundaries. Thus the grain refinement is the most important method to deal with this shortcoming so far [7–10]. The metastable ferromagnetic Heusler (L21 structure) Cu-Al-Mn alloys are also one of Cu-Al-based SMAs. Generally, a two-stage orderdisorder transformation of A2→B2→L21 occurs during cooling for CuAl-Mn SMAs [11,12]. The high ordered degree of L21 parent is also a
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key factor for the polycrystalline brittleness of Cu-Al-Mn SMAs. Recently, Omori et al. reported that the abnormal grain growth (AGG) of Cu-17Al-11.4Mn (at%) [13,14] and Fe-34Mn-15Al-7.5Ni (at%) [15,16] alloys. The ultra-large single crystals of more than a few centimeters could be obtained both in two systems. Their method was cyclic heat treatment between single bcc phase region at high temperature and two-phase of (bcc + fcc) region at low temperature. The AGG phenomenon resulted from the formation of sub-grains (about 100 µm) induced by the coherency loss of fcc precipitates during thermal cycling. Their finding greatly promotes the development of Cu-Al-Mn SMAs. The preparation of Cu-Al-Mn single crystals through cyclic heat treatment is much simpler comparing to the directional solidification [17–20]. However, the requirements of this method are still very strict, including the anneal temperature, cooling rate and numbers of thermal cycle. Our research group recently designed a unique microstructure of cast Cu-Al-Mn-Mo SMAs based on the idea of bcc two-phase separation [21]. The cast Cu-Al-Mn-Mo alloys consisted of L21-Cu2AlMn parent with completely coherent bcc A2(Mo) nanoparticles. Interestingly, the AGG directly happened only through annealing their polycrystalline
Corresponding author. E-mail address: [email protected] (S. Yang).
https://doi.org/10.1016/j.msea.2019.02.033 Received 27 December 2018; Received in revised form 9 February 2019; Accepted 9 February 2019 Available online 11 February 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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coherent with bcc A2(W) nanoparticles due to bcc phase separation (Fig. 2c). Additionally, it was found from the composition mapping analysis in Fig. 2e - i, the L21 parent was the Cu-rich, the bcc A2(W) nanoparticles were the W-rich, and the γ1(Cu9Al4) phase is the Al-rich and Mn-poor.
cast alloys. At this time, the large grains and single crystals over cmscale were obtained. Our results further confirmed that the AGG may be closely related to a continuous misorientation gradient within the matrix grains resulted from the dissolution of A2(Mo) nanoparticles. We consider that W and Mo both have A2 structure, and little solid solubility in Cu. Thus we speculate that bcc two-phase separation may also occur in Cu-Al-Mn-W alloys, and similar microstructure of cast alloys may be obtained. Therefore in this study, the Cu-14.2Al-10Mn-0.5W (wt%) alloy was designed, and the microstructure and abnormal grain growth of the alloy, superelasticity and fatigue resistance of single crystals were investigated.
3.3. Microstructural evolution during annealing In order to investigate the relationship between the AGG and the γ1(Cu9Al4) phase, the cast sample was annealed at 1023 K for 5 min and quenched, the γ1(Cu9Al4) phase had dissolved to the matrix (Fig. 3b, e and g). However, the AGG did not take place at this time. The result indicates that the existence of γ1(Cu9Al4) phase does not apply any influence to the AGG. According to the previous results, bcc two-phase separation of Cu-Al-Mn-Mo alloys happened around 1073 K [21]. Thus it is speculated that the AGG of Cu-Al-Mn-W alloy should occur only after the dissolution of A2(W) nanoparticles around 1073 K. The results of Fig. 3c, d and f confirm this speculation. When the above cast sample was annealed at 1173 K for 10 min, the AGG took place (Fig. 3c), and large grains formed after 2 h (Fig. 3d and f). Our previous investigation revealed that a continuous misorientation gradient within the matrix grains due to the dissolution of bcc precipitates was responsible for the AGG in Cu-Al-Mn-based alloys with bcc phase separation [21]. Fig. 4 shows the EBSD quasi-colored orientation mapping (OM) and the corresponding grain reference orientation deviation (GROD) of the cast alloy after annealing. When the cast alloy was annealed at 1173 K for 5 min, the γ1(Cu9Al4) phase disappeared, but the AGG still did not happen. At this time, the orientation of each grain was random (Fig. 4a), and no misorientation existed among the grains (Fig. 4b). After annealing for 10 min, few large grains formed (Fig. 4c). Here, there were no any sub-grains that reported by Omori et al. in Cu-Al-Mn ternary alloy [13,14]. Similar to those of CuAl-Mn-Mo alloys [21], a continuous misorientation gradient below 2° deviation was observed from grain core to grain boundaries (Fig. 4d).
2. Experimental Button ingot (40 g) of nominal Cu-14.2Al-10.0Mn-0.5W (wt%) alloy was prepared through arc-melting. The cast ingot was sealed into vacuum quartz tubes under argon atmosphere, then annealed at 1173 K for 2 h and followed by water quenching. The microstructure and crystal structure were observed and analyzed by optical microscopy, transmission electron microscopy (TEM) and the selected area electron diffraction (SAED). The chemical composition was measured by energy dispersive spectrometer (EDS) with an average value of five measurements. The actual composition was Cu-13.96Al-9.84Mn-0.51W. The composition mapping analysis of each element was determined by electron probe microanalysis (EPMA). The crystallographic orientation and the orientation of the single crystals were determined by electron backscatter diffraction (EBSD) using a field emission scanning electron microscope (FE-SEM). The superelastic properties were measured by compression test using cylindrical single crystal specimens (3 mm diameter, 5 mm height). The specimen was compressed to different pre-strains (εpre) and then unloading to a zero stress condition. The SE (εSE) strain was confirmed and measured directly by stress-strain curves [22]. The height of the sample was measured before loading (h0) and after unloading (h1). The residual strain (εr) was calculated as εr = (h0 - h1)/h0 × 100%. The εe implies the elastic strain when unloading to a zero stress condition. Therefore, the εSE strain was calculated using the formulas εSE = (εpre εe - εr). The fatigue resistance was carried out by compression test with 5 × 103 cycles under a constant applied deformation of 8.5%.
3.4. Superelasticity and fatigue resistance of single crystal Cylindrical single crystal specimens were directly prepared from one of large grains (Fig. 5a). The martensitic transformation temperature was measured by the differential scanning calorimetry (DSC) at a heating and cooling rate of 10 °C min−1. However, the transformation temperature of the single crystal was not detected from room temperature to −140 °C. It indicates that the transformation temperature may be below −140 °C or the transformation enthalpy is too small to detect. The orientation of single crystal along the compression direction was close to [100] (with the deviation of about 7.8°) (Fig. 5b). The fracture strain and stress were 13.2% and 706 MPa, respectively. Exhilaratingly, a very large stress plateau corresponding to the stress-induced martensite was observed. It indicates that this single crystal should possess excellent SE properties. Complete and perfect superelasticity up to 12% and the largest SE strain up to 8.5% were obtained in Fig. 6. Its SE property is compared with Ni-Ti SMAs [28–30], and CuZn-Al, Cu-Al-Ni as well as Fe-Mn-Al-Ni single crystals [3,4,31], see Table 1 [32–37]. Furthermore, the single crystal possesses excellent fatigue resistance as shown in Fig. 7. A stress-strain cycle of 5 × 103 times with a constant pre-strain of 8.5% was carried out (Fig. 7a). During cycling, this single crystal always exhibited full superelasticity behavior. Full superelasticity of 8.5% and SE strain of 6.8% were displayed during the first cycle. After 5 × 102 cycles and 5 × 103 cycles, full superelasticity decreased to 8.1% and 6.5%, the SE strain decreased to 5.9% and 4.3% (Fig. 7b and c). The superelasticity property begun to noticeably degrade after about 3000 cycles (Fig. 7d).
3. Results 3.1. Abnormal grain growth Fig. 1 shows the macroscopic morphologies of Cu-13.96Al-9.84Mn0.51W alloy before and after annealing at 1173 K for 2 h followed by water quenching. When the polycrystalline cast alloy was annealed at 1173 K, the AGG happened. At this time, only two grains exceptionally grew. As a result, large grains and single crystals over cm-scale were obtained. 3.2. Microstructure of the cast alloy Fig. 2 shows the microstructure of the cast Cu-13.96Al-9.84Mn0.51 W polycrystalline alloy. The cast alloy consisted of three-phase of L21 parent + γ1(Cu9Al4) phase + A2(W) phase. The A2(W) phase had two morphologies, visibly large particles (Fig. 2a) and fine nanoparticles (Fig. 2b). Furthermore, the L21 parent was completely
Fig. 1. Abnormal grain growth of Cu-13.96Al-9.84Mn-0.51 W alloy. 250
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Fig. 2. Microstructure of Cu-13.96Al-9.84Mn-0.51W cast alloy. (a) Optical microscopy. (b) TEM bright-field (BF) image. (c) Corresponding SAED pattern of L21 matrix and bcc A2(W) nanoparticles (solid circle in b). (d) Corresponding SAED pattern of γ1(Cu9Al4) phase (dashed circle in b). (e - i) Composition mapping analysis of each element.
Fig. 3. Microstructural evolution of Cu-13.96Al-9.84Mn-0.51 W alloy during annealing. (a) Cast sample. (b–d) The same cast sample (in a) was annealed at 1023 K for 5 min (b), 1173 K for 10 min (c), and 1173 K for 2 h (d). (e) Optical metallograph accords to the sample in b. (f) Optical metallograph accords to the sample in d. (g) X-ray diffraction patterns of the cast sample (black) and the sample annealed at 1023 K for 5 min (red).
4. Discussion
Cu-rich bcc#1 phase + W-rich bcc#2 phase (those A2(W) nanoparticles). Then the Cu-rich bcc#1 phase transforms to the ordered L21Cu2AlMn matrix via the order-disordered transformation of A2 → B2, and B2 → L21 during the process of cooling to room temperature [11,24–27]. In this study, the normal grain growth may be inhibited in the cast state due to the A2(W) particle pinning. When the cast alloy is annealed at high enough temperature, those fine A2(W) particles dissolve back into the matrix. At this time, the particle pinning is eliminated, which may promote a few grains to grow exceptionally. Additionally, the formation of continuous misorientation gradient within the matrix grains may provide extra driving force for the rapid migration of grain boundaries. More details about the AGG of Cu-Al-Mn-based alloy with bcc phase separation have been explained in our previous report [21].
4.1. Abnormal grain growth From the results of Fig. 2, it is confirmed that the microstructure of the cast Cu-13.96Al-9.84Mn-0.51W alloy is similar to those of Cu-AlMn-Mo alloys [21]. The bcc two-phase separation happens in Cu13.96Al-9.84Mn-0.51W alloy. Combining with the phase diagram of binary Cu-W system [23], it is possible that a liquid separation region of Cu-rich and W-rich liquids exists in the Cu-Al-Mn-W system. During the solidification, the W-rich liquid solidifies into visibly large A2(W) particles, and the Cu-rich liquid solidifies into the Cu-rich bcc phase. With further decreasing temperature, the solidified Cu-rich bcc phase undergoes bcc two-phase separation. That is the Cu-rich bcc phase → 251
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Fig. 4. Crystallographic orientations of Cu-13.96Al-9.84Mn-0.51 W alloy. (a) Quasi-colored orientation mapping (OM) of the alloy annealed at 1173 K for 5 min and water quenching. (b) Grain reference orientation deviation (GROD) mapping corresponds to (a). (c) OM of the alloy annealed at 1173 K for 10 min and water quenching. (d) GROD mapping corresponds to (c). In OM figures (a and c), the colors of crystal direction is given in the inverse pole figure. In GROD mappings (b and d), the color value of every pixel is calculated as the misorientation angle of this pixel with respect to a reference average orientation of one grain.
9.84Mn-0.51W single crystal and only 8% for Cu-12.0Al-9.0Mn-0.5Mo single crystal. Secondly, although the orientation of single crystals is both close to [100] direction, the degree of deviation is different. The deviation is 7.8° for Cu-13.96Al-9.84Mn-0.51W single crystal, whereas is larger 10.4° Cu-12.0Al-9.0Mn-0.5Mo single crystal. Thirdly, the strength of two single crystals is much different. For Cu-13.96Al9.84Mn-0.51W single crystal, the critical stress to induce martensite is about 220 MPa, and the fracture strength is about 706 MPa. However, the critical stress and fracture strength are much higher in Cu-12.0Al9.0Mn-0.5Mo single crystal. For Cu-Al-Mn ternary SMAs, the strength of both polycrystalline alloys and single crystals is low [17–19], thus the fatigue resistance of the alloys is relatedly poor [19]. The advisable increase of alloy strength is beneficial to the improvement of alloy properties. However, when the critical stress is too high, even close to
4.2. Superelasticity and fatigue resistance As we all known, the SE characteristics of shape memory single crystals are closely related to their orientations. When the orientation is different, the single crystal with the same composition may exhibit completely different SE characteristics. In the previous report, the SE properties of Cu-Al-Mn-Mo single crystals with several different orientations were studied [21]. Therefore, here only the SE properties of Cu-13.96Al-9.84Mn-0.51W and Cu-12.0Al-9.0Mn-0.5Mo single crystals with close to [100] orientation are discussed. In this study, a higher SE strain of 8.5% was obtained in Cu-13.96Al9.84Mn-0.51W single crystal than 4.8% of Cu-12.0Al-9.0Mn-0.5Mo single crystal [21]. There are three possible reasons for this issue. Firstly, the applied pre-strains are different, 12% for Cu-13.96Al-
Fig. 5. Mechanical properties of Cu-13.96Al-9.84Mn-0.51W single crystal. (a) Schematic illustrations of the tested single crystal and the analysis locations for the orientation of single crystal. (b) The orientations of single crystal respectively accord to (A) and (B) locations in (a) along the compression direction. (c) Compressive stress-strain curve. 252
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the yield strength or fracture strength of the alloy, the plastic deformation is also more likely to occur during the stress-induced martensite. At this time, the SE properties begin to decrease. With regard to the reason that resulted in the strength of single crystals is so different with different additions of Mo or W into Cu-Al-Mn alloys is not clear yet. This issue needs more investigations. H. Kato, et al. reported the cyclic stress-strain response of superelastic Cu-9.8Al-9.9Mn (wt%) shape memory single crystal with [001] orientation during a strain-cycling of 7.8% over a thousand cycles [19]. Their results revealed that the critical stress to induce the martensitic transformation gradually decreased at an almost constant rate from about 130 MPa to zero only after about 50 cycles. Because the residual martensite was accumulated during repeated cycling, the residual strain obviously increased and the SE strain decreased. The fatigue behavior is related to the dislocation formed during cycling. In the present study, several completely coherent A2(W) nanoparticles exist within L21 parent for Cu-Al-Mn-W single crystal before deformation. Thus the critical stress (above about 220 MPa) to induce martensite is higher than that of Cu-9.8Al-9.9Mn single crystal. Although the critical stress gradually decreases while increasing the cycles (Fig. 7a–c), the rate of decrease is clearly smaller. The existence of completely coherent A2(W) nanoparticles and the improvement of strength may make the alloy not easy to produce the defects or dislocation during repeated cycling, so the fatigue properties of Cu-Al-Mn-W single crystal can be improved. Because the cost of Cu-based polycrystalline SMAs is only one-tenth that of Ni-Ti-based SMAs, thus the Cu-based SMAs have wide applications. However, the polycrystalline brittleness of Cu-based alloys seriously hinder their further development so far. Additionally, the Cubased single crystals that prepared through directional solidification are not only very expensive but also time-consuming. In this study, from the points of raw material cost and preparation of single crystal, the present Cu-Al-Mn-W single crystal fabricated only through annealing cast polycrystalline alloy shows more superiority of low cost and simple technology. Furthermore, the present Cu-Al-Mn-W single crystal possesses full superelasticity up to 12% and the largest SE strain up to
Fig. 6. Superelasticity of Cu-13.96Al-9.84Mn-0.51W single crystal.
Table 1 Superelastic strain in various shape memory alloys. The symbol (†) indicates that the properties are obtained in single crystal. Alloy composition Cu-13.96Al-9.84Mn-0.51 W Fe-34Mn-15Al-7.5Ni Ti-42.4Ni-8Cu Ti-51.3Ni Ti-24Zr-10Nb-2Sn Cu68.4Al27.8Ni3.8 Cu67.9Zn16.1Al16 Cu71.9Al16.6Mn9.3Ni2B0.2 Cu-Al16.915-Mn10.4475Co0.5 Cu-24.25Zn-8.57Al-0.07Zr Cu-17.0Al-11.4Mn Zu45-Au30-Cu25 Co-32Ni-28Al Ni51-Fe22-Ga27
SE (%) †
8.5 9.5†/5.2 7.0 6.0 6.0 8.6† 8.5† 7.5 7.1 4.5 6.5†/5† 8.0 6.3 6.2†
References This work [30] [28] [29] [31] [3] [4] [32] [33] [34] [17,18] [35] [36] [37]
Fig. 7. Fatigue resistance of Cu-13.96Al-9.84Mn-0.51 W single crystal. (a) Stress-strain curves of 5 × 103 cycles with a constant pre-strain of 8.5%. (b and c) Stressstrain curves of 5 × 102 cycles (b) and 5 × 103 cycles (c). (d) Relationship between full superelasticity, SE strain and number of cycle. 253
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8.5%, and excellent fatigue resistance, being compared to Ni-Ti-based SMAs. Therefore, the Cu-Al-Mn-W single crystal shows good potential as new superelastic material.
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5. Conclusions In summary, a novel microstructure of cast Cu-Al-Mn-W SMA was designed. The cast alloy consisted of L21-Cu2AlMn parent with completely coherent bcc A2(W) nanoparticles. This microstructure induced a continuous misorientation gradient within the matrix grains from grain core to grain boundary during annealing, which might result in the AGG. Furthermore, the present Cu-Al-Mn-W shape memory single crystal exhibited excellent SE properties and fatigue resistance. So that our findings may enable wide applications of low-cost Cu-based SE materials through various engineering fields as advanced functional materials. Acknowledgements This work was supported by the financial supports of the Shenzhen Science and Technology Project, Grant no. JCYJ20170306142550151, the Natural Science Foundation of Fujian Province, China, Grant no. 2018J01078. References [1] K. Otsuka, C.M. Wayman (Eds.), Shape Memory Materials, Cambridge Univ. Press, Cambridge, 1998, pp. 1–48. [2] R. Dasgupta, A look into Cu-based shape memory alloys: present scenario and future prospects, J. Mater. Res. 29 (2014) 1681–1698. [3] K. Otsuka, C.M. Wayman, K. Nakai, H. Sakamoto, K. Shimizu, Superelasticity effects and stress-induced martensitic transformations in CuAlNi alloys, Acta Metall. 24 (1976) 207–226. [4] T. Saburi, Y. Inada, S. Nenno, N. Hori, Stress-induced martensitic transformations in Cu-Zn-Al and Cu-Zn-Ga alloys, J. De. Phys. C4 43 (1982) 633–638. [5] V.H.C. de Albuquerque, T.A. de, A. Melo, R.M. Gomes, S.J.G. de Limaa, J.M.R.S. Tavares, Grain size and temperature influence on the toughness of a CuAlBe shape memory alloy, Mater. Sci. Eng. A 528 (2010) 459–466. [6] X.J. Liu, I. Ohnuma, R. Kainuma, K. Ishida, Phase equilibria in the Cu-rich portion of the Cu–Al binary system, J. Alloy. Compd. 264 (1998) 201–208. [7] V. Sampath, Studies on the effect of grain refinement and thermal processing on shape memory characteristics of Cu–Al–Ni alloys, Smart Mater. Structs. 14 (2005) S253–S260. [8] V. Sampath, Improvement of shape-memory characteristics and mechanical properties of copper–zinc–aluminum shape-memory alloy with low aluminum content by grain refinement, Mater. Man. Proc. 21 (2006) 789–795. [9] P. Zhang, A.B. Ma, S. Lu, G.G. Liu, P.H. Lin, J.H. Jiang, C.L. Chu, Effect of grain refinement on the mechanical properties of Cu–Al–Be–B shape memory alloy, Mater. Des. 32 (2011) 348–352. [10] J. Yang, Q.Z. Wang, F.X. Yin, C.X. Cui, P.G. Ji, B. Li, Effects of grain refinement on the structure and properties of a CuAlMn shape memory alloy, Mater. Sci. Eng. A 664 (2016) 215–220. [11] R. Kainuma, N. Satoh, X.J. Liu, I. Ohnuma, K. Ishida, Phase equilibria and Heusler phase stability in the Cu-rich portion of the Cu–Al–Mn system, J. Alloy. Compd. 266 (1998) 191–200. [12] M.O. Prado, P.M. Decorte, F. Lovey, Martensitic transformation in Cu-Mn-Al alloys, Scr. Mater. 33 (1995) 877–883.
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