Low-cost Cu-based shape memory single crystals obtained by abnormal grain growth showing excellent superelasticity

Materialia 5 (2019) 100200

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Low-cost Cu-based shape memory single crystals obtained by abnormal grain growth showing excellent superelasticity Shuiyuan Yang a,∗, Jixun Zhang a, Mengyuan Chi a, Yuhua Wen b, Xinren Chen a, Cuiping Wang a, Xingjun Liu a,c,∗ a

College of Materials, Fujian Key Laboratory of Materials Genome, Xiamen University, Xiamen 361005, China School of Manufacturing Science and Engineering, Sichuan University, Chengdu 610065, China c State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology University, Shenzhen, Guangdong 518055, China b

a r t i c l e Keywords: Abnormal grain growth Single crystal Nanoparticles Phase separation Superelasticity

i n f o

a b s t r a c t Metal single crystals are generally obtained by directional solidification with sophisticated equipment and process. This method is not only very expensive but also very time-consuming. In this study, we report a new approach to produce Cu-Al-Mn-Mo shape memory alloys with single crystals through annealing cast polycrystalline alloys. The microstructure of the cast alloys has a common characteristic that the L21 -Cu2 AlMn parent is completely coherent with bcc A2(Mo) precipitates due to bcc phase separation. The dissolution of bcc A2(Mo) nanoparticles during annealing process results in abnormal grain growth because of the formation of a continuous misorientation gradient within the grains. The single crystals obtained from large grains exhibit excellent superelasticity. When the orientation of single crystals is close to [−210], the largest full superelasticity is up to 10%, and the largest superelastic strain is up to 6.7%. It is expected that such microstructural design can also be applied to the production of large grains and even single crystals of other alloy systems with phase separation.

1. Introduction Metallic alloys are usually prepared by the conventional method of solidification and annealing. They typically have polycrystalline structure due to the nucleation and growth of differently oriented grains. The grains are separated by grain boundaries. During the subsequent annealing process, some of the grain boundaries will disappear to lower total energy through the diffusion of atoms. This process is normal grain growth. When equilibrium is reached, the grain growth stops. Usually, metal single crystals are superior to their polycrystalline counterparts, especially for intermetallic compounds [1–7]. Bulk single crystals are typically prepared by directional solidification using a set of sophisticated equipment and technology with remarkably high cost and time consumption [4,5,8–12], such as Bridgman and Czochralski methods. Large grains and even single crystals can also be produced by abnormal grain growth (AGG) using special thermo-mechanical treatments, such as strain-annealing [13–16] and dynamic AGG methods [17–19]. However, the AGG phenomenon can take place only in wire, plate, or sheet samples subjected to macroscopic deformation. So far, few papers report AGG in a bulk ingot without macroscopic deformation.

∗

Recently, Omori et al. reported that large single crystals of more than a few centimeters could be obtained in Cu-Al-Mn [20,21] and FeMn-Al-Ni [22,23] shape memory alloys (SMAs) through a cyclic heat treatment between single bcc phase region at high temperature and two-phase of (bcc + fcc) region at low temperature. They suggested that the formation of sub-grains of about 100 𝜇m induced by the coherency loss of the growing fcc precipitates was responsible for the AGG phenomenon. Their findings greatly promote the development and application of Cu-Al-Mn SMAs. Although the preparation of Cu-Al-Mn single crystals through cyclic heat treatment is much simpler comparing to previous methods (directional solidification [4,5,8–12], the AGG induced special strain-annealing [13–16] and dynamic AGG methods [17–19]), the conditions of this method are still very strict, including the anneal temperature, cooling rate, and numbers of thermal cycle. In this study, a new and simpler approach to obtain bulk large grains and even single crystals over several centimeters in Cu-Al-Mn based SMAs through annealing cast polycrystalline alloys was reported. Single crystals of about 55 × 36 × 18 mm3 and 40 × 20 × 6 mm3 grew from hundreds of micrometers in the cast Cu-Al-Mn-Mo alloys, when only annealing the cast alloys at 1173 K for 24 hours followed by water quenching. Microstructural observations revealed that many coherent bcc nanoparticles precipitated inside the cast bcc matrix due to bcc phase separation.

Corresponding authors. E-mail addresses: [email protected] (S. Yang), [email protected] (X. Liu).

https://doi.org/10.1016/j.mtla.2018.100200 Received 20 September 2018; Accepted 18 December 2018 Available online 23 December 2018 2589-1529/© 2018 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

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Materialia 5 (2019) 100200

Fig. 1. Macroscopic morphologies of Cu-14Al-10Mn-1Mo (a to d) and Cu-13Al-9Mn-1Mo alloys (e and f). (a, b, e) Cast alloys. (c, d, f) Alloys annealed at 1173 K for 24 hours and water quenching.

During the annealing process, the dissolution of the nanoparticles into the matrix resulted in a continuous misorientation gradient within the grains. Consequently, AGG occurred, and large grains resulted. Several single crystals obtained from large grains showed excellent full superelasticity (SE) of 10% and superelastic strain of 6.7%, comparing with commercially traditional SMAs. 2. Experimental procedures Bulk metal ingots (40 g or 100 g) were prepared through arc-melting under an argon atmosphere. The purity of each raw material was above 99.9 atom%. The bulk ingots were re-melted five times to ensure composition uniformity. Metal ingot (about 1000 g) was prepared through a high vacuum induction melting furnace using ten pre-melted small ingots prepared by arc-melting. The cast ingots were sealed into vacuum quartz tubes under argon atmosphere, then annealed at 1173 K for 24 hours followed by water quenching. The AGG phenomenon was directly observed by macroscopic morphology on optical camera. Optical microscopy was used to observe the microstructure. Transmission electron microscopy (TEM) was used to analyze the microstructure and crystal structure by selected area electron diffraction (SAED), the distribution of each element by composition mapping analysis. In-situ high temperature TEM was used to observe the precipitation and dissolution of nanoparticles during cooling and annealing. The distribution of each element was also analyzed through electron probe microanalysis (EPMA). Electron backscatter diffraction (EBSD) technology using a field emission scanning electron microscope (FE-SEM) was adopted to investigate the crystallographic orientation of the polycrystalline alloys and the orientation of the single crystals. Differential scanning calorimetry (DSC) with a heating and cooling rate of 10 K•min−1 was used to study the phase transformation behaviors. In-

situ high-temperature confocal microscopy was used to observe the migration of the grain boundaries during annealing. The cylindrical samples (3 mm diameter, 3 mm height) for in-situ high-temperature confocal microscope test was annealed at 1173 K for 30 minutes and quenched, and then mechanically polished without corrosion. The superelasticity was measured by compression tests using cylindrical single crystal specimens (3 mm diameter, 5 mm height) taken from large grains. The different specimens were compressed to different pre-strains (𝜀pre ) and then unloaded to a zero stress condition. The SE (𝜀SE ) strain was directly measured by stress-strain curves. 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 ). 3. Results 3.1. Abnormal grain growth Fig. 1 shows the macroscopic morphologies of Cu-14Al-10Mn-1Mo and Cu-13Al-9Mn-1Mo before and after annealing at 1173 K for 24 hours followed by water quenching (respectively correspond to 1Mo-3 and 1Mo-5 alloys in Table S1), (weight%, hereinafter). The chemical compositions of Cu-14Al-10Mn-1Mo and Cu-13Al-9Mn-1Mo alloys were Cu13.78Al-9.88Mn-0.95Mo and Cu-12.90Al-8.86Mn-1.03Mo, respectively. The cast alloys were fine grains of hundreds of micrometers (Fig. 1a, 1b and 1e). After the annealing, few grains exceptionally grew, and the large single crystals of about 55 × 36 × 18 mm3 (Fig. 1c and d) and 40 × 20 × 6 mm3 were obtained (Fig. 1f, still existing some fine grains within the single crystal).

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Materialia 5 (2019) 100200

Fig. 2. Microstructure of Cu-14Al-10Mn-1Mo cast alloy. (a) TEM bright-field (BF) image of L21 matrix and A2(Mo) phase with the SAED pattern of A2(Mo) phase in the inset. (b) TEM BF image of L21 matrix, 𝛾 1 (Cu9 Al4 ) phase and A2(Mo) nanoparticles. (c) SAED patterns of L21 matrix and 𝛾 1 (Cu9 Al4 ) phase corresponding to (b). (d) TEM BF image of L21 matrix and A2(Mo) nanoparticles. (e) SAED patterns of L21 matrix and A2(Mo) nanoparticles (top, dashed circle), as well as single L21 matrix (bottom, solid circle).

Fig. 3. (a–d) Composition mapping analysis measured by TEM-EDS. (e–h) Composition mapping analysis measured by EPMA. In (e–h), the color bar only represents the relative concentration of each element in EPMA test. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

3.2. Microstructure Microstructural observations in Fig. S1a and the TEM in Fig. 2 revealed that the cast Cu-14Al-10Mn-1Mo alloy consisted of the ordered Heusler L21 (Cu2 AlMn) matrix + 𝛾 1 (Cu9 Al4 ) phase + A2(Mo) phase. The A2(Mo) precipitates had two morphologies. One was visibly larger A2(Mo) particles in Fig. S1. Another was those A2(Mo) nanoparticles in Fig. 2. In addition, the A2(Mo) nanoparticles were completely coherent with the L21 matrix due to bcc phase separation (Fig. 2d and e), whereas the 𝛾 1 (Cu9 Al4 ) phase was not coherent (Fig. S2). The L21 matrix was Cu/Mn-rich and Mo-poor, whereas the A2(Mo) nanoparticles were more

enriched with Mo and poorer with Cu/Mn (Fig. 3). The Al concentration was almost the same in the L21 matrix as it was in A2(Mo) nanoparticles. The results showed that the 𝛾 1 (Cu9 Al4 ) phase precipitated when the alloy was annealed at 873 K, and then disappeared at 1023 K. At these times, the AGG did not happen (Fig. S1). The AGG only took place when the alloy was annealed around 1173 K. 3.3. Phase transformation and microstructural evolution The phase transformation behaviors during solidification and annealing processes was clarified by DSC tests, and the results are given in

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Materialia 5 (2019) 100200

Fig. 4. Phase transformation and evolution of macroscopic morphology of Cu-14Al-10Mn-1Mo alloy. (a) DSC curves of cast alloy during heating and cooling. (b) Macroscopic morphologies of the same cast alloy annealed at 873 K for 2 hours, 1023 K for 24 hours and 1123 K for 24 hours and water quenching. (c) Macroscopic morphologies of the same alloy was annealed at 1173 K for 30, 40 and 60 minutes followed by water quenching. In (a), the phase transformation peaks respectively were: (1) ↔ (9) for phase transformation of solid ↔ liquid (melting (peak 9) and solidification (peak 1)); (2) ↔ (7) for order-disorder transformation of B2 ↔ A2 (peak 2 for A2 → B2; peak 7 for B2 → A2); (3) ↔ (6) for dissolution (peak 6) and precipitation (peak 3) of 𝛾 1 (Cu9 Al4 ) phase; (4) ↔ (5) for order-disorder transformation of L21 (Cu2 AlMn) ↔ B2 (peak 4 for B2 → L21 ; peak 5 for L21 → B2); (8) peak 8 for dissolution of A2(Mo) nanoprecipitates.

Fig. 5. Grain boundary migration during annealing observed through in-situ high-temperature confocal microscope test, in which the tested Cu-14Al-10Mn1Mo cast alloy was pre-annealed at 1173 K for 30 minutes and water quenching. (a) 580.1 K. (b) 1038.4 K. (c) 1059.4 K. (d) 1173.4 K.

Fig. 4a. Combined with the results of the later microstructural observations, the phase transformation behaviors were analyzed in Fig. 4a. In Fig. 4b, the macroscopic photographs revealed that the grain size of the cast alloy did not change as increasing the annealing temperature up to 1023 K. The grains exceptionally grew when annealing at 1123 K for 24 hours. Furthermore, when annealing the cast alloy at 1173 K, the AGG occurred only after 30 minutes (Fig. 4c). The in-situ observation using a high-temperature confocal microscope also confirmed that the grain boundaries started to rapidly migrate at about 1059 K (Fig. 5, Supplementary Movie S1). In-situ high temperature TEM observation revealed that the microstructure did not change as increasing the temperature up to 873 K, but the 𝛾 1 phase disappeared at 973 K (Fig. 6, Supplementary Movie S2). The completely coherent A2(Mo) nanoparticles started to dissolve into the matrix at 1073 K. The results confirmed that the AGG did not occur when the 𝛾 1 phase dissolved back to the matrix. It implies

that the existence of 𝛾 1 phase does not apply any influence to AGG and the formation of large grains. Interestingly, after the A2(Mo) nanoparticles dissolved back into the matrix for a period of time when annealing at 1173 K, some regions lower than hundreds of nanometer within the grains showed obvious contrast differences (Fig. 6). In TEM, the bright and dark areas are closely related to the composition (atomic number) and the sample thickness. In the present study, when the A2(Mo) precipitates dissolve into the matrix, the sample is single phase structure. Thus it is suggested that the compositions are not responsible for those bright areas in Fig. 6 and Movie S2. According to the later analysis (4.1 section), during the dissolution of the A2(Mo) nanoparticles, the coherent strain energy between A2(Mo) nanoparticles and bcc matrix due to bcc phase separation should release, and a certain amount of Mo atoms inside the region of A2(Mo) nanoparticles must move outward. Thus a stress field and lattice distortion within the grains must be introduced. As a result, the tested sample is subjected to internal stress, and some areas become thinner, according to those bright areas in Movie S2 and Fig. 6. The phase evolution corresponding to solidification and annealing processes is schematically illustrated in (Fig. 7). Combining with the phase diagram of binary Cu-Mo [24], it is found that a liquid separation region of Cu-rich and Mo-rich liquids exists in the Cu-14Al-10Mn-1Mo alloy (>1262 K). During solidification, the Mo-rich liquid solidifies into the visibly large A2(Mo) 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 of Cu-rich bcc phase → Cu-rich bcc#1 phase + Mo-rich bcc#2 phase (The DSC test did not determine the temperature of bcc phase separation during solidification, Fig. 4a). Then the Cu-rich bcc#1 phase transforms to the ordered L21 (Cu2 AlMn) matrix via the order-disordered transformation of A2 → B2 (∼ 993 K) and B2 → L21 (∼ 835 K) during the process of cooling to room temperature. Similar ordered transformation of Cu-Al-Mn alloys has been reported in the previous investigation [25,26]. Accordingly, the A2(Mo) nanoparticles (Mo-rich bcc#2 phase) precipitates due to bcc two-phase separation. Additionally, the 𝛾 1 (Cu9 Al4 ) phase also precipitates at 873 K, and again dissolves into the matrix around 1023 K (Fig. S1b and 1c). The reversible phase transformations that occur during annealing are illustrated in Fig. 7b. 3.4. Mechanism of abnormal grain growth The electron backscatter diffraction (EBSD) method was used to further investigate the reason for the AGG phenomenon in the present alloys. Fig. 8 shows the EBSD quasi-colored orientation mapping (OM)

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Fig. 6. In-situ high temperature TEM observations of the dissolution of A2(Mo) nanoparticles and the formation of a continuous misorientation gradient during annealing for Cu-14Al-10Mn-1Mo alloy.

Fig. 7. Schematic illustrations for the microstructure and phase evolution during the processes of solidification (a) and annealing (b).

and the corresponding grain reference orientation deviation (GROD) of the cast Cu-14Al-10Mn-1Mo alloy after annealing at 1173 K for 10 minutes (Fig. 8a and b) and 40 minutes (Fig. 8c and d). After annealing for 10 minutes, the 𝛾 1 phase has completely dissolved back to the matrix. At this time, the misorientation exists among the grains, and the orientation of each grain was random (Fig. 8a). But in each grain, no obvious misorientation gradient was observed (Fig. 8b). After annealing for 40 minutes, a large grain formed (Fig. 8c). Meanwhile, the misorientation still exists among the grains, and the orientation of each grain was random (Fig. 8c). However, an obvious continuous misorientation gradient with 2° deviation was produced from the grain core to grain boundaries within the formed large grain (Fig. 8d). Combined the results of TEM observation in Fig. 6, the continuous misorientation gradient within the matrix grains should be induced by the stress field and lattice distortion introduced by the dissolution of the A2(Mo) nanoparticles. We also confirmed that the Cu-13Al-9Mn-1Mo cast alloy possessed the same mi-

crostructure as that of the Cu-14Al-10Mn-1Mo cast alloy (Fig. S3). Furthermore, the same continuous misorientation gradient with 2° deviation was produced from the grain core to grain boundaries when the large grain formed during annealing for Cu-13Al-9Mn-1Mo alloys (Fig. S4). In order to further identify the relationship between the AGG and the dissolution of A2(Mo) nanoparticles, we annealed the cast Cu-13Al10Mn alloy without the addition of Mo. The result showed that Cu-13Al10Mn alloy was single L21 parent structure before and after annealing (Fig. 9a and b). After annealing at 1173 K for 24 h, the grains can normally grow up to millimeter level (< ∼ 2 mm), but the AGG did not occur (Fig. 9c and d). In DSC curves (Fig. 9e), the phase transformation of the dissolution of A2(Mo) nanoparticles did not be observed due to the absence of A2(Mo) phase. Furthermore, the results of the EBSD measurement also revealed that a continuous misorientation gradient within the matrix grains could not be produced when annealing the cast Cu-13Al-

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Fig. 8. Crystallographic orientations of Cu-14Al-10Mn-1Mo alloy. (a) Quasi-colored orientation mapping (OM) of the sample annealed at 1173 K for 10 minutes and water quenching. (b) Grain reference orientation deviation (GROD) mapping corresponding to (a). (c) OM of the sample annealed at 1173 K for 40 minutes and water quenching. (d) GROD mapping corresponding 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. The red arrows imply the increase direction of the misorientation gradient within large grain. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

10Mn alloy without the A2(Mo) nanoparticles (Fig. 10). Therefore, it is suggested that the exceptional grain growth of Cu-Al-Mn-Mo alloy must be related to the dissolution of the A2(Mo) nanoparticles and their resulting microstructural change. 3.5. Composition range of abnormal grain growth We also investigated the composition range for effecting AGG in Cu-Al-Mn-Mo alloys. The corresponding results are provided in Fig. S5 and S6, Table S1. The microstructure of the cast Cu-Al-Mn-Mo alloys clearly differed when the alloy compositions were different, including three situations of L21 + A2(Mo) + 𝛾 1 (Cu9 Al4 ), L21 + A2(Mo), and richCu A2 + 𝛼(Cu) + A2(Mo). However, one common feature existed in all studied alloys: the cast alloys consisted of bcc L21 parent and completely coherent bcc A2(Mo) nanoparticles. Therefore, it is also suggested that

the unique microstructure of the cast alloy containing nanoparticles completely coherent with the matrix as a result of bcc phase separation is the key condition for the formation of continuous misorientation gradient within the matrix grains and the occurrence of the AGG. Additionally, from the present obtained results (Fig. S5 and S6, Table S1), it was also confirmed that the existence of 𝛾 1 (Cu9 Al4 ) and 𝛼(Cu) does not apply any influence to the AGG and the formation of large grains. The relationship between the alloy composition and the AGG in Cu-Al-MnMo alloys are summarized in Fig. 11. 3.6. Superelasticity The cylindrical single crystal specimens were directly obtained from large grains, and their superelastic properties were investigated as shown in Figs. 12 and 13. The superelastic properties of all the

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Materialia 5 (2019) 100200

Fig. 9. Normal grain growth of Cu-13Al-10Mn alloy. (a) Optical metallograph and BSE image (inset) of the cast alloy. (b) Optical metallograph and BSE image (inset) of the alloy annealed at 1173 K for 24 h followed by quenching. (c) Macroscopic morphology of the cast alloy. (d) Macroscopic morphology of the cast alloy after annealed at 1173 K for 24 h followed by quenching. The result showed that Cu-13Al-10Mn alloy was single L21 parent structure before and after annealing. After annealing at 1173 K for 24 h, the grains can normally grow up to millimeter level (< ∼2 mm), but the AGG did not occur. (e) DSC curves of the cast Cu-13Al-10Mn alloy. The results confirmed that only order-disorder transformations of L21 (Cu2 AlMn) ↔ B2 ↔ A2 existed in ternary Cu-13Al-10Mn alloy.

Fig. 10. Crystallographic orientation of Cu-13Al-10Mn alloy annealed at 1173 K for 40 min and then quenched. (a) Quasi-colored orientation mapping (OM). (b) Grain reference orientation deviation (GROD) mapping corresponding to OM figure. In OM figure, the colors of crystal direction is given in the inverse pole figure. In GROD mapping, 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. It was confirmed that ternary Cu-13Al-10Mn alloy did not exhibit a continuous misorientation gradient within the matrix grains when the cast alloy was annealed at 1173 K for 40 min and then quenched. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

studied Cu-Al-Mn-Mo single crystals and polycrystalline alloys were summarized in Table S2. The Cu-13Al-10Mn alloy with coarse grains only exhibited about 3.1% of SE strain (Fig. 13). However, when the orientations of single crystals along the compression direction were close to [100], [-210], [-120], [-310] and [1-20], these single crystals should

exhibit excellent superelasticity (Fig. 12). The SE properties of the single crystals decreased when the orientations of single crystals were close to [-4-32] and [-5-42] in 0.5Mo-1 and 2Mo-3 alloys, and even disappeared with the orientations of [2-2-1], [-3-32] and [-2-32] in 2Mo-1, 0.2Mo and 1.5Mo-1 alloys (Fig. 13). The 1Mo-7 alloy with few large grains also

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Fig. 11. Relationship between the abnormal grain growth and the compositions of Cu-Al-Mn-Mo alloys.

exhibited about 3.8% SE strain. When the alloy was fine polycrystalline structure, the SE decreased in 1.5Mo-1 and 1Mo-8 alloys. 4. Discussion 4.1. Abnormal grain growth of Cu-Al-Mn-Mo alloys There are other methods that can lead to the AGG, such as the Zener pinning effect [27,28] and diffusion-induced grain boundary migration [29]. These methods heavily involve particles and their dissolution around grain boundaries. Therefore, the presence of the secondphase particles plays an important role in AGG after recrystallization. For example, AGG happened in Fe-Si steels with precipitated particles smaller than about 450 nm [30–32]; similar events folded in Cu-Al alloy with trace amounts of particles [33]. In this study, the normal grain growth may be inhibited in the cast state due to the A2(Mo) particle pinning. When the cast alloy is annealed at high enough temperature, the second-phase particles dissolves back into the matrix. At this time, the particle pinning is eliminated, which may promote a few grains to grow

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exceptionally. This may be one reason why few grains grow abnormally in the present study. On the other hand, comparing to the traditional Cu-based SMAs [3– 5,20,21,34–37], the present Cu-Al-Mn-Mo cast alloys exhibit unique microstructure containing nanoparticles completely coherent with the matrix as a result of bcc phase separation. It may be another important condition for the AGG of Cu-Al-Mn-Mo alloy. A schematic illustration of microstructural evolution and corresponding mechanism for the AGG in the present alloys with bcc phase separation was given in Fig. 14. Firstly, the coherent strain energy between A2(Mo) nanoparticles and bcc matrix is usually high due to bcc phase separation [38] (Fig. 14a). When the A2(Mo) nanoparticles dissolve into the matrix, the energy within the matrix grains around A2(Mo) nanoparticles will increase due to the release of coherent strain energy (Fig. 14b). Secondly, the compositions of A2(Mo) nanoparticles and bcc matrix are completely different due to bcc phase separation. And the solid solubility of Mo in the matrix is very small. Accordingly, during the dissolution of A2(Mo) nanoparticles, a certain amount of Mo atoms inside the region of A2(Mo) nanoparticles must move outward. As the result of size difference among Mo, Cu, Mn and Al atoms, a stress field and lattice distortion within the grains must be introduced [39] (Fig. 14b). These two reasons may result in the contrast differences in Fig. 6 and the formation of a continuous misorientation gradient within the matrix grains in Fig. 8d. Subsequently, a number of Mo atoms must outward to release the internal stress and reduce the system energy during further annealing (Fig. 14c). The movement of atoms may be the driving force for the rapid migration of the grain boundaries. The mechanism for the AGG in the present Cu-Al-MnMo alloy was also illustrated in Fig. 7b. The present AGG phenomenon is different from the reports of Omori et al. [20,21], embodying in the following four differences. (1) Alloy composition. Omori et al. reported Cu-8.1Al-11.1Mn ternary alloy (wt%) with obviously lower Al content. Our finding is Cu-Al-Mn-Mo quaternary alloy, in which the Mo is a very crucial element. In the present study, the 1Mo-8 and 1Mo-9 alloys possessed the similar compositions and microstructure as that of Cu-8.1Al-11.1Mn alloy with a large amount of fcc 𝛼(Cu) phase. However, the AGG did not happen when directly annealing 1Mo-8 and 1Mo-9 alloys at 1173 K and quenching. (2) Microstructure of the alloys. In this study, bcc phase separation happened due to alloying Mo, which leads to those completely coherent A2(Mo) nanoparticles. (3) Mechanism of AGG. The fcc 𝛼(Cu) phase and sub-grains resulted in the AGG of Cu-8.1Al-11.1Mn alloy were not observed in the present alloys. A continuous misorientation gradient within the matrix grains due to the dissolution of the nanoparticles is responsible for the present AGG. (4) Preparation technology. The subgrains induced by the precipitation of fcc phase at low temperature is a necessary condition for the AGG occurring at high temperature in Cu8.1Al-11.1Mn alloy. However, the amount of sub-grains formed during each thermal cycle may be limited, thus the corresponding AGG is also inadequate. This requires several thermal cycles to induce AGG repeatedly. As compared to the thermal cycling process with the strict conditions for the anneal temperature, cooling rate and numbers of thermal cycle, our present approach to prepare large grains and even single crystal by the dissolution of fine nanoparticles is more easily and simpler. From the present results, it is found that the AGG is closely related to the amounts of A2(Mo) precipitates mainly depending on bcc phase separation. Although the A2(Mo) precipitates have been observed and confirmed by TEM tests in this study, its volume fraction is too small to measure. It is estimated that the volume fraction of A2(Mo) precipitates is less than 1% from Fig. 2. It is suggested when the amount of A2(Mo) particles is too much or too few, obvious continuous misorientation gradient within the grains may not have formed, thus the AGG does not happen (Fig. 11 and Fig. S5). Therefore, only when the amount A2(Mo) particles is suitable, and its distribution is remarkably non uniform, obvious continuous misorientation gradient within few grains forms. Then one or two grains abnormally grow at the expense of others and the AGG happens. The deeper mechanism of thermodynamics and kinet-

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Fig. 12. Superelasticity of the present studied Cu-Al-Mn-Mo single crystals.

Fig. 13. Superelasticity of the partial studied Cu-Al-Mn-Mo single crystals and polycrystalline alloys.

ics, as well as the effect of A2(Mo) precipitates, including the amount, distribution and grain size, on the AGG of Cu-Al-Mn-Mo alloys may be complex, and they need much further investigations in the next works. 4.2. Superelasticity of Cu-Al-Mn-Mo alloys The orientation of single crystal along the deformation direction has important effects the superelastic properties of the single crystals. When the orientations of the single crystals are close to [−210], [−120] and [1–20], their superealstic behaviors are similar in 1Mo-3, 1Mo-4 and 2Mo-2 single crystals. The largest full superelasticity is up to 10%, the largest SE strain is 6.7% in 1Mo-3 single crystal. The stress hysteresis

loop is the largest with the lowest critical stress to induce martensitic transformation when 1.5Mo-2 single crystal has an orientation close to [−310]. A large stress hysteresis loop implies that the single crystal can absorb a large amount of energy during superelastic cycle. The single crystal of this orientation shows a potential as a high-damping material. A small stress hysteresis loop presents in 0.5Mo-2 single crystal with the orientation of [100], which indicates lower energy loss during repeated loading and unloading. When the orientations of single crystals are close to [-3-32], [-2-32] and [2-2-1], there is no superelasticity. Since Cu-Al-based SMAs have easier fabrication and lower manufacturing cost than those of Ni-Ti-based alloys (A tenth of the cost of Ni-Tibased alloys), they have also a number of researchers’ attentions all the

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Fig. 14. Schematic illustrations for the formation mechanism of a continuous misorientation gradient with the matrix grains in the present study. (a) Cast alloy. (b) Dissolution of A2(Mo) nanoparticles at high temperature. (c) Diffusion of atoms during further annealing process at high temperature.

time. However, the polycrystalline brittleness of traditional Cu-Al-based SMAs, and expensive cost for the production of single crystals, both seriously limit their wide development and application. In this study, the method of directly annealing cast polycrystalline alloys to produce CuAl-Mn-Mo single crystals is much more versatile and cost-efficient manufacturing technique compared with traditional method (directional solidification) or the cyclic heat treatment. Furthermore, the present CuAl-Mn-Mo shape memory single crystals exhibit excellent SE properties.

5. Conclusions In summary, bulk large grains were obtained only through simply annealing cast Cu-Al-Mn-Mo polycrystalline alloys possessing bcc phase separation and completely coherent nanoparticles. Several single crystals were prepared from the large grains. Comparing to the reported methods for preparing single crystala by the special thermo-mechanical treatments and the thermal cycling, our approach is the easiest to implement with the simplest requirement and the lowest cost. Furthermore, the single crystals possess excellent superelasticity, comparing to the conventional Ni-Ti and Cu-based SMAs. From the points of raw material cost, preparation and functional properties of single crystal, this Cu-Al-Mn-Mo alloy may be suitable as new superelastic materials.

Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgments We acknowledge the financial supports of the National Key R&D Program of China (grant number 2017YFB0702901), the Natural Science Foundation of Fujian Province, China, grant number 2018J01078, the National Natural Science Foundation of China (grant numbers 51571168 and 51671138).

Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.mtla.2018.100200.

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