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Improvement of the stability of superelasticity and elastocaloric effect of a Ni-rich Ti-Ni alloy by precipitation and grain refinement
Scripta Materialia 162 (2019) 230–234
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Improvement of the stability of superelasticity and elastocaloric effect of a Ni-rich Ti-Ni alloy by precipitation and grain refinement Hong Chen a, Fei Xiao a,d,⁎, Xiao Liang a, Zhenxing Li a, Zhu Li a, Xuejun Jin b,⁎⁎, Na Min c, Takashi Fukuda d,⁎⁎ a
State Key Lab of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dong Chuan Road, Shanghai 200240, PR China Institute of Advanced Steels and Materials, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, People's Republic of China Key Laboratory for Microstructures, School of Materials Science and Engineering, Shanghai University, Shanghai 200444, China d Department of Materials Science and Engineering, Graduate School of Engineering, Osaka University, 2-1, Yamada-oka, Suita, Osaka 565-0871, Japan b c
a r t i c l e
i n f o
Article history: Received 20 October 2018 Received in revised form 6 November 2018 Accepted 11 November 2018 Available online xxxx Keywords: Ti-Ni Shape memory alloys Superelasticity Elastocaloric effect Thermomechanical treatment
a b s t r a c t In this paper, we report that the superelasticity and the elastocaloric effect of Ti-50.8Ni (at.%) alloy are drastically improved by a combination of precipitation and grain refinement strengthening. The process is as follows: first, aging at 773 K after solution treatment to form coherent fine precipitate of Ti3Ni4; second, cold working followed by annealing at 673 K to form nanocrystalline grain containing Ti3Ni4. The decay of the superelasticity and the elastocaloric effect by 100 stress cycles is b10%. © 2018 Published by Elsevier Ltd on behalf of Acta Materialia Inc.
Superelasticity in Ti-Ni based shape memory alloys (SMAs) is now widely used in many applications, including medical devices [1–3]. It uses large recoverable strain caused by stress-induced martensitic transformations (MT). The MT in Ti-Ni alloys shows a large latent heat compared with other shape memory alloys. The value of the latent heat is reported to be approximately 0.8 kJ/mol [4,5]. The large latent heat of Ti-Ni alloys is expected to be used for refrigeration [6–8]. Because of the latent heat, the temperature of the specimen increases through adiabatic stress-induced MT, and it decreases by the reverse MT. This temperature change of the specimen is referred to as the elastocaloric effect (eCE) [9–11]. The adiabatic temperature decrease of Ti-Ni alloy expected from the above latent heat is approximately 40 K. In experiments, adiabatic temperature decreases between 17 K and 21 K are reported by [12,13]. This temperature decrease is much larger than that in magnetocaloric materials whose typical adiabatic temperature change is below 10 K [14]. Because of the large eCE, Ti-Ni SMAs have attracted much attention as a candidate for refrigerant material [15].
⁎ Corresponding author at: State Key Lab of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dong Chuan Road, Shanghai 200240, PR China. ⁎⁎ Corresponding authors. E-mail addresses: [email protected] (F. Xiao), [email protected] (X. Jin), [email protected] (T. Fukuda).
https://doi.org/10.1016/j.scriptamat.2018.11.024 1359-6462/© 2018 Published by Elsevier Ltd on behalf of Acta Materialia Inc.
Although Ti-Ni SMAs exhibit large superelasticity and eCE [16,17], these properties decay through stress cycling [18–21]. The functional decay is a critical issue to limit the application of SMAs [22]. The main reason for the decay is the accumulation of defects introduced during the stress-inducing MT [23,24]. To prevent the accumulation of defects, we need to further strengthen Ti-Ni alloys through new processes. It is well known that the Ti3Ni4 precipitate induced by heat treatment and the grain refinement induced by severe plastic deformation are two important methods to strengthen the matrix for Ti-Ni SMAs [16]. The effectiveness of these two methods has separately been confirmed in past decades [25–27]. We may expect that the combination of the two strengthening methods could be very effective to prevent the accumulation of defects. Miyazaki et al. [28] once tried to investigate such coupling effect on superelasticity by employing a Ti-50.6Ni (at.%) alloy that underwent cold rolling (CR) treatment and then aged at 673 K. However, it was found that the Ti3Ni4 hardly precipitated in the nanocrystalline grains formed by cold rolling [29]. There is still a lack of reports on the coupling effect on eCE. Now it is largely accepted that the formation of nanocrystalline grains formed by cold rolling is very effective for strengthening Ti-Ni alloys, but the fine grain prevents the formation of Ti3Ni4, which is also effective for improving the strength of these alloys. To obtain the combined effect of precipitation strengthening and grain refinement strengthening in Ti-Ni alloys, we need to find new processes. The inspiration of this study was that although Ti3Ni4 does not precipitate in nanocrystalline grain, it precipitates easily in coarse grains, and the
H. Chen et al. / Scripta Materialia 162 (2019) 230–234
0.1
B19’
B2
0.8 kJ/mol Cooling
Ms=241 K
HR1123
Heat Flow(mW/mg)
precipitate could remain in the specimen after cold rolling if it is introduced before cold rolling. In this study, we demonstrate that the combined strengthening can be realized by aging heat-treatment before cold rolling. The specimen produced by this process shows excellent stability in superelasticity and eCE compared with the specimens subjected to only grain refinement strengthening or precipitation strengthening. A hot-rolled Ti-50.8Ni (at.%) sheet was purchased from Fasten-PLT Materials Science Co., Ltd. The sheet was heat-treated at 1123 K for 1 h. We call this state as HR1123. Three kinds of additional thermomechanical treatments were made on the HR1123. We call them CR673 and CR773 and NPCR673. The CR673 and the CR773 samples were prepared as follows. First, the HR1123 sheet was aged at 773 K for 2 h to induce coherent precipitate of Ti3Ni4. Then, the sheet was cold rolled with thickness reduction of 50%. Subsequently, the sheet was annealed at 673 K for 10 min (CR673) or at 773 K for 10 min (CR773). The NPCR673 was prepared by the same procedure as CR673 but was not aged at 773 K before cold rolling. We expected that the NPCR673 shows the grain refinement strengthening effect, the CR773 shows the precipitation strengthening effect and the CR673 shows both effects. The transformation characteristics of the specimens were measured by a differential scanning calorimeter (DSC NETZSCH model 200F3) with a cooling/heating rate of 10 K/min. The microstructures of the specimens were observed by a JEOL 2100 transmission electron microscope (TEM) at room temperature (~293 K). Samples for TEM observation were mechanically ground to a thickness of approximately 80 μm, then double jet electro-polished by using a 10 vol% acetic acid and 90 vol% methanol solution at 283 K. Dog-bone shaped specimens for tensile tests were cut along the rolled direction by using a slow wireelectrode cutting machine. Tensile tests were made using an Instron5966 mechanical testing machine with a loading/unloading strain rate of 5 × 10−3 s−1. The eCE was detected at a fast unloading strain rate of 2 s−1 to achieve an adiabatic condition. The temperature change (ΔT) of the specimen was monitored by a T-type thermocouple welded on the surface of the specimen. Fig. 1 shows the DSC heat flow of the HR1123, CR773, CR673 and NPCR673 specimens in a temperature range of 130–390 K. The HR1123 specimen exhibits one exothermic peak in the cooling process and one endothermic peak in the subsequent heating process. The latent heat Q is ~0.84 kJ/mol and the hysteresis between the cooling and heating process (Af-Ms) is ~30 K. These results are consistent with previous reports of Ti-Ni SMAs exhibiting B2-B19′ martensitic transformation [30]. Other specimens (CR773, CR673 and NPCR673) exhibit two separated exothermic peaks in the cooling process and a merged endothermic peak in the heating process. This behavior implies that these specimens show successive B2-R-B19′ martensitic transformation. The formation of the R-phase in these three specimens will be attributed to the stress-field introduced by the cold rolling treatment or to the Ti3Ni4 precipitate. We can roughly estimate the ΔT caused by reverse MT from the latent heat Q in the heating process as Q/Cp, where Cp is the specific heat. Assuming Cp is 3R for all the specimens, where R is the gas constant, the ΔT caused by the reverse MT is estimated to be −33.7 K, −36.1 K, −24.1 K and −24.1 K for HR1123, CR773, CR673 and NPCR673 samples, respectively. The B2-R transformation start temperature Ms (R) is 294 K for the CR773 sample, 324 K for the CR673 sample, and 314 K for the NPCR673 sample. We can speculate two reasons which cause the difference in Ms (R). One is the difference in the composition of the B2-matrix. As confirmed later by TEM observations, the CR673 and CR773 specimens include the Ti3Ni4 phase, which is formed by aging treatment before cold rolling. The formation of this phase reduces the Ni content of the B2-matrix, which results in the increase in Ms (R). This will be the reason why Ms (R) of the CR673 is higher than that of the NPCR673. Another reason for the difference in Ms (R) is density of the nucleation sites
231
B19’
Af=264 K
B2
Heating
-0.8 kJ/mol
CR773
R B2 0.3 kJ/mol Ms=294 K
B19’ R 0.4 kJ/mol B19’ R B2 -0.9 kJ/mol
CR673
Af=304 K
0.3kJ/mol
0.2 kJ/mol
Ms=324 K Af=330 K
-0.6 kJ/mol 0.2 kJ/mol
NPCR673
150
0.3 kJ/mol M =314 K s Af=320 K
-0.6 kJ/mol
200
250
300
350
400
Temperature(K) Fig. 1. DSC cooling and heating curves of the Ti-50.8Ni (at.%) alloy after different thermomechanical treatments (HR1123, CR773, CR673, NPCR673). The latent heats Q and characteristic temperatures are indicated in the figure.
from where the R-phase nucleates. The higher Ms (R) in the CR673 sample compared with the CR773 sample implies that the density of nucleation sites in the CR673 specimen is higher than that of the CR773 specimen. Presumably, the defects introduced by cold-rolling which act as nucleation sites are largely reduced by the heat-treatment at 773 K, while the reduction is smaller at 673 K. The microstructure of each specimen was examined by TEM. Fig. 2 displays the bright field images and the corresponding selected area diffraction patterns (SADP) of the present specimens at room temperature. The typical grain size of the HR1123 specimen was several micrometers. Inside of the grain is plane without precipitates as seen in (a). The SADP can be indexed by the B2 parent phase. After aging the HR1123 specimen at 773 K for 2 h, we confirmed that fine precipitate of Ti3Ni4 with an average size of approximately 100 nm is formed homogeneously in the grain, which is shown in Supplementary material. Through cold rolling and subsequent heat-treatments, the grain size decreased to the nanoscale, as seen in Fig. 2(b–d). The grain size of the CR673 specimen and the NPCR673 specimen is nearly the same (~40 nm). This means that grain size is not influenced by the existence of Ti3Ni4, which is formed by the aging treatment before cold rolling. However, the grain size of the CR773 specimen (~100 nm) is significantly larger than that of the CR673 specimen. In addition, the shape of each grain is different. The grain of the CR673 specimen is elongated to the rolling direction, which is characteristic of rolled specimen. On the other hand, the grain of the CR773 specimen is equiaxed, which is characteristic of recrystallized specimens. It is speculated that the influence of cold rolling is largely retained in the CR673 and NPCR673 specimens, but not in the CR773 specimen. Due to the small grain size and large stress field, it is difficult to identify the Ti3Ni4 precipitates in the TEM bright field images of the CR773, CR673 and NPCR673 specimens. Nevertheless, reflections from Ti3Ni4 phase were detected for the CR773 and CR673 specimens while not for the NPCR673 specimen in the SADP. This implies that the Ti3Ni4 phase, which was formed by the aging treatment before cold rolling, remained after cold rolling and subsequent heat-treatment at 673 K or 773 K. On the other hand, the absence of the Ti3Ni4 phase in NPCR673 implies that precipitation of Ti3Ni4 after cold rolling is suppressed, possibly due to the small grain size. It is consistent with a previous report by
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Prokofiev et al. in which they confirmed that the precipitation of Ti3Ni4 was suppressed in grains of cross-sectional equivalent diameter below approximately 150 nm [29]. Incidentally, according to the DSC results,
(a) HR1123
110B2
101B2
000B2
500nm
(b) CR773
200nm
(c) CR673
200nm
(d) NPCR673
310B2 220B2 211B2 200B2 212P 110B2 202P 100B2
310B2 220B2 211B2 200B2 212P 110B2 202P 100B2
310B2 220B2 211B2 200B2 110B2 100B2
200nm
(e) R phase (f) 10nm
TS
the R-phase is expected to exist at room temperature in the CR673 and NP CR673 specimens. To identify the R-phase, we made a high resolution observation for the CR673. Fig. 2 (e) is a high resolution image of CR673 specimen, and (f) is the fast Fourier transformation of the red rectangular area. We notice 1/3〈110〉 reflections characteristic to the R-phase. Since the 1/3〈110〉 appears only in one direction, the red rectangular region will be a single variant of the R-phase. The solid curves in Fig. 3 are the stress-strain curves of the initial cycle of all the examined specimens. All the tests were made sufficiently above the Af temperature. They show a clear plateau corresponding to the stress-induced MT. The shape recovery is almost perfect, except for the HR1123 specimen in which a residual strain of ~3.6% can be seen. We evaluated the critical stress for the transformation by the tangent method, and the temperature dependence of the critical stress is shown in Fig. 4. It increases linearly as temperature increases for all the specimens. From the slope in Fig. 4, we can obtain the entropy change ΔS of the MT by using the ClausiusClapeyron relation dσ/dT = − ΔS/Δε, where Δε is the transformation strain. Then, we can estimate the adiabatic temperature change (ΔT) caused by the stress-induced MT as ΔT=−TΔS/CP. The value of ΔT estimated for the HR1123, CR773, CR673 and NP CR673 specimens are −31.5 K, −27.7 K, −28.5 K and −25.8 K, respectively, for the stress removing process. These values are nearly the same as those estimated from the latent heat mentioned above except for CR773 specimen. The experimental value of ΔT obtained by the stressreducing process is −8.2 K, −24.3 K, −26.4 K and −24.0 K for the HR1123, CR773, CR673 and NPCR673 specimens, respectively. The experimental value is approximately 0.92 times the estimated value for all the specimens except for HR1123 specimen. The large deviation between the estimated vale and the experiment in HR1123 specimen is due to incomplete reverse transformation. The small deviation (~8%) in other specimens is probably due to insufficient adiabatic condition of the experiment. The stress-strain curves and adiabatic temperature change of the initial cycle mentioned above imply that there is no significant difference in superelasticity and elastocaloric effect among CR773, CR673 and NPCR673 specimens. However, the evolution of the stress-cycle is significantly different among these three specimens, as described below. The evolution of the stress-strain curves and adiabatic temperature decrease in the stress-removing process for the CR773, CR673 and NPCR673 specimens are shown in Fig. 3(b–d) and (b′–d′). The test temperature is approximately 10 K above Af for each specimen. The transformation strain and adiabatic temperature decrease show visible decay as the number of stress cycles increases for all specimens. The decay is the most significant in the CR773 specimen, both Δε and ΔT decreases by ~45% through 100 cycles (Δε decreases from 5.6% to 2.7% and ΔT reduces from −24.3 K to −13.3 K). The decay for the CR673 specimen is ~10% and that for the NPCR673 specimen is ~20%. We briefly discuss the reason why the decay of Δε and ΔT is small in the CR673 specimen when compared with other specimens. It is reported that the decrease of the Δε during stress cycling is caused by the accumulation of defects [23,24,31]. Therefore, we may regard that the introduction of defects through stress cycling is smallest for the CR673 specimen. The nanocrystalline microstructure of the CR673 specimen caused by cold rolling and subsequent heat treatment at 673 K is probably the main reason for the strengthening. However, we notice
Fig. 2. Bright field images, selected area diffraction patterns and high resolution TEM for the Ti-50.8Ni (at.%) alloy after different thermomechanical treatments: HR1123 (a), CR773 (b), CR673 (c) and NPCR673 (d). The high resolution TEM of the CR673 specimen (e) and corresponding fast Fourier transformation of the red rectangular area (f). B2 and P on index indicate the parent phase and Ti3Ni4, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
H. Chen et al. / Scripta Materialia 162 (2019) 230–234
(a)
HR1123 T=308K
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29.1K
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Stress (MPa)
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Fig. 3. Initial cycle stress-strain curve and the adiabatic temperature change, ΔT, of the HR1123 specimen (a, a′). Stress-strain curves and the adiabatic temperature change, ΔT, as a function of time of the Ti-50.8Ni (at.%) alloy after different thermomechanical treatments, CR773 (b, b′), CR673 (c, c′), and NPCR673 (d, d′) specimens. Solid curves are the results of the first cycle, thin curves are those of every 10 cycles. The testing temperature and ΔT caused by the reverse martensite transformation are indicated in the figure.
Critical Stress(MPa)
that the decay of the CR673 specimen is obviously small when compared with the NCR673 specimen. The only difference between the two specimens is the existence of Ti3Ni4. Although the grain size of
800
600
400
300
320
340
6.1MPa/K 5.2MPa/K 6.0MPa/K 5.8MPa/K
HR1123 CR773 CR673 NPCR673
360
380
Temperature(K)
Fig. 4. The temperature dependence of the critical stress for induced martensite transformation for the Ti-50.8Ni (at.%) alloy after different thermomechanical treatments (HR1123, CR773, CR673, NPCR673).
the CR673 specimen is several tens nanometers, it is likely that Ti3Ni4 still contribute as effective precipitate which prevents the movement of dislocations. The stable eCE for the CR673 specimen indicates that it could be a good candidate for an eCE material. When the SMAs are used as refrigeration materials, the energy efficiency of the materials (η) is an important parameter for the cooling technologies. According to Ref. [32], η is defined as η = ∣ Qh/W∣, where W is the total mechanical work done to the refrigerant material in a cycle, and Qh is heat flow to the material from cold thermal bath in the cycle. We approximate | Qh | as |-Cp∙ΔT| using experimentally obtained ΔT, and |W| as the enclosed area of the stress-strain curve [6]. The value of η thus evaluated for the CR673 specimen is 7.4 for the first cycle and 12.1 for the 100th cycle. The increase of η is due to the decrease of stress hysteresis. The high value of η indicates that the CR673 specimen exhibits good refrigerating capacity [33]. In conclusion, Ti3Ni4 phase was successfully introduced in nanocrystalline Ti-50.8Ni alloy by an aging heat-treatment before cold rolling and subsequent heat-treatment. The specimen prepared by this process exhibits stable superelasticity and elastocaloric effect because of the combination of precipitation strengthening and grain-refinement strengthening. This method could be applied for the preparation of highly stable shape memory alloys, which can be used for solid-state refrigeration. Supplementary data to this article can be found online at https://doi. org/10.1016/j.scriptamat.2018.11.024.
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Acknowledgement This work was funded by the National Youth Science Foundation (No. 51871151, No. 51501113), the National Key R&D Program of China (No. 2017YFB0703003, No. 2017YFB0406000), National Natural Science Foundation of China (U1564203) and Iketani Science and Technology Foundation (No. 0301012-A). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]
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Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Improvement of the stability of superelasticity and elastocaloric effect of a Ni-rich Ti-Ni alloy by precipitation and grain refinement Hong Chen a, Fei Xiao a,d,⁎, Xiao Liang a, Zhenxing Li a, Zhu Li a, Xuejun Jin b,⁎⁎, Na Min c, Takashi Fukuda d,⁎⁎ a
State Key Lab of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dong Chuan Road, Shanghai 200240, PR China Institute of Advanced Steels and Materials, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, People's Republic of China Key Laboratory for Microstructures, School of Materials Science and Engineering, Shanghai University, Shanghai 200444, China d Department of Materials Science and Engineering, Graduate School of Engineering, Osaka University, 2-1, Yamada-oka, Suita, Osaka 565-0871, Japan b c
a r t i c l e
i n f o
Article history: Received 20 October 2018 Received in revised form 6 November 2018 Accepted 11 November 2018 Available online xxxx Keywords: Ti-Ni Shape memory alloys Superelasticity Elastocaloric effect Thermomechanical treatment
a b s t r a c t In this paper, we report that the superelasticity and the elastocaloric effect of Ti-50.8Ni (at.%) alloy are drastically improved by a combination of precipitation and grain refinement strengthening. The process is as follows: first, aging at 773 K after solution treatment to form coherent fine precipitate of Ti3Ni4; second, cold working followed by annealing at 673 K to form nanocrystalline grain containing Ti3Ni4. The decay of the superelasticity and the elastocaloric effect by 100 stress cycles is b10%. © 2018 Published by Elsevier Ltd on behalf of Acta Materialia Inc.
Superelasticity in Ti-Ni based shape memory alloys (SMAs) is now widely used in many applications, including medical devices [1–3]. It uses large recoverable strain caused by stress-induced martensitic transformations (MT). The MT in Ti-Ni alloys shows a large latent heat compared with other shape memory alloys. The value of the latent heat is reported to be approximately 0.8 kJ/mol [4,5]. The large latent heat of Ti-Ni alloys is expected to be used for refrigeration [6–8]. Because of the latent heat, the temperature of the specimen increases through adiabatic stress-induced MT, and it decreases by the reverse MT. This temperature change of the specimen is referred to as the elastocaloric effect (eCE) [9–11]. The adiabatic temperature decrease of Ti-Ni alloy expected from the above latent heat is approximately 40 K. In experiments, adiabatic temperature decreases between 17 K and 21 K are reported by [12,13]. This temperature decrease is much larger than that in magnetocaloric materials whose typical adiabatic temperature change is below 10 K [14]. Because of the large eCE, Ti-Ni SMAs have attracted much attention as a candidate for refrigerant material [15].
⁎ Corresponding author at: State Key Lab of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dong Chuan Road, Shanghai 200240, PR China. ⁎⁎ Corresponding authors. E-mail addresses: [email protected] (F. Xiao), [email protected] (X. Jin), [email protected] (T. Fukuda).
https://doi.org/10.1016/j.scriptamat.2018.11.024 1359-6462/© 2018 Published by Elsevier Ltd on behalf of Acta Materialia Inc.
Although Ti-Ni SMAs exhibit large superelasticity and eCE [16,17], these properties decay through stress cycling [18–21]. The functional decay is a critical issue to limit the application of SMAs [22]. The main reason for the decay is the accumulation of defects introduced during the stress-inducing MT [23,24]. To prevent the accumulation of defects, we need to further strengthen Ti-Ni alloys through new processes. It is well known that the Ti3Ni4 precipitate induced by heat treatment and the grain refinement induced by severe plastic deformation are two important methods to strengthen the matrix for Ti-Ni SMAs [16]. The effectiveness of these two methods has separately been confirmed in past decades [25–27]. We may expect that the combination of the two strengthening methods could be very effective to prevent the accumulation of defects. Miyazaki et al. [28] once tried to investigate such coupling effect on superelasticity by employing a Ti-50.6Ni (at.%) alloy that underwent cold rolling (CR) treatment and then aged at 673 K. However, it was found that the Ti3Ni4 hardly precipitated in the nanocrystalline grains formed by cold rolling [29]. There is still a lack of reports on the coupling effect on eCE. Now it is largely accepted that the formation of nanocrystalline grains formed by cold rolling is very effective for strengthening Ti-Ni alloys, but the fine grain prevents the formation of Ti3Ni4, which is also effective for improving the strength of these alloys. To obtain the combined effect of precipitation strengthening and grain refinement strengthening in Ti-Ni alloys, we need to find new processes. The inspiration of this study was that although Ti3Ni4 does not precipitate in nanocrystalline grain, it precipitates easily in coarse grains, and the
H. Chen et al. / Scripta Materialia 162 (2019) 230–234
0.1
B19’
B2
0.8 kJ/mol Cooling
Ms=241 K
HR1123
Heat Flow(mW/mg)
precipitate could remain in the specimen after cold rolling if it is introduced before cold rolling. In this study, we demonstrate that the combined strengthening can be realized by aging heat-treatment before cold rolling. The specimen produced by this process shows excellent stability in superelasticity and eCE compared with the specimens subjected to only grain refinement strengthening or precipitation strengthening. A hot-rolled Ti-50.8Ni (at.%) sheet was purchased from Fasten-PLT Materials Science Co., Ltd. The sheet was heat-treated at 1123 K for 1 h. We call this state as HR1123. Three kinds of additional thermomechanical treatments were made on the HR1123. We call them CR673 and CR773 and NPCR673. The CR673 and the CR773 samples were prepared as follows. First, the HR1123 sheet was aged at 773 K for 2 h to induce coherent precipitate of Ti3Ni4. Then, the sheet was cold rolled with thickness reduction of 50%. Subsequently, the sheet was annealed at 673 K for 10 min (CR673) or at 773 K for 10 min (CR773). The NPCR673 was prepared by the same procedure as CR673 but was not aged at 773 K before cold rolling. We expected that the NPCR673 shows the grain refinement strengthening effect, the CR773 shows the precipitation strengthening effect and the CR673 shows both effects. The transformation characteristics of the specimens were measured by a differential scanning calorimeter (DSC NETZSCH model 200F3) with a cooling/heating rate of 10 K/min. The microstructures of the specimens were observed by a JEOL 2100 transmission electron microscope (TEM) at room temperature (~293 K). Samples for TEM observation were mechanically ground to a thickness of approximately 80 μm, then double jet electro-polished by using a 10 vol% acetic acid and 90 vol% methanol solution at 283 K. Dog-bone shaped specimens for tensile tests were cut along the rolled direction by using a slow wireelectrode cutting machine. Tensile tests were made using an Instron5966 mechanical testing machine with a loading/unloading strain rate of 5 × 10−3 s−1. The eCE was detected at a fast unloading strain rate of 2 s−1 to achieve an adiabatic condition. The temperature change (ΔT) of the specimen was monitored by a T-type thermocouple welded on the surface of the specimen. Fig. 1 shows the DSC heat flow of the HR1123, CR773, CR673 and NPCR673 specimens in a temperature range of 130–390 K. The HR1123 specimen exhibits one exothermic peak in the cooling process and one endothermic peak in the subsequent heating process. The latent heat Q is ~0.84 kJ/mol and the hysteresis between the cooling and heating process (Af-Ms) is ~30 K. These results are consistent with previous reports of Ti-Ni SMAs exhibiting B2-B19′ martensitic transformation [30]. Other specimens (CR773, CR673 and NPCR673) exhibit two separated exothermic peaks in the cooling process and a merged endothermic peak in the heating process. This behavior implies that these specimens show successive B2-R-B19′ martensitic transformation. The formation of the R-phase in these three specimens will be attributed to the stress-field introduced by the cold rolling treatment or to the Ti3Ni4 precipitate. We can roughly estimate the ΔT caused by reverse MT from the latent heat Q in the heating process as Q/Cp, where Cp is the specific heat. Assuming Cp is 3R for all the specimens, where R is the gas constant, the ΔT caused by the reverse MT is estimated to be −33.7 K, −36.1 K, −24.1 K and −24.1 K for HR1123, CR773, CR673 and NPCR673 samples, respectively. The B2-R transformation start temperature Ms (R) is 294 K for the CR773 sample, 324 K for the CR673 sample, and 314 K for the NPCR673 sample. We can speculate two reasons which cause the difference in Ms (R). One is the difference in the composition of the B2-matrix. As confirmed later by TEM observations, the CR673 and CR773 specimens include the Ti3Ni4 phase, which is formed by aging treatment before cold rolling. The formation of this phase reduces the Ni content of the B2-matrix, which results in the increase in Ms (R). This will be the reason why Ms (R) of the CR673 is higher than that of the NPCR673. Another reason for the difference in Ms (R) is density of the nucleation sites
231
B19’
Af=264 K
B2
Heating
-0.8 kJ/mol
CR773
R B2 0.3 kJ/mol Ms=294 K
B19’ R 0.4 kJ/mol B19’ R B2 -0.9 kJ/mol
CR673
Af=304 K
0.3kJ/mol
0.2 kJ/mol
Ms=324 K Af=330 K
-0.6 kJ/mol 0.2 kJ/mol
NPCR673
150
0.3 kJ/mol M =314 K s Af=320 K
-0.6 kJ/mol
200
250
300
350
400
Temperature(K) Fig. 1. DSC cooling and heating curves of the Ti-50.8Ni (at.%) alloy after different thermomechanical treatments (HR1123, CR773, CR673, NPCR673). The latent heats Q and characteristic temperatures are indicated in the figure.
from where the R-phase nucleates. The higher Ms (R) in the CR673 sample compared with the CR773 sample implies that the density of nucleation sites in the CR673 specimen is higher than that of the CR773 specimen. Presumably, the defects introduced by cold-rolling which act as nucleation sites are largely reduced by the heat-treatment at 773 K, while the reduction is smaller at 673 K. The microstructure of each specimen was examined by TEM. Fig. 2 displays the bright field images and the corresponding selected area diffraction patterns (SADP) of the present specimens at room temperature. The typical grain size of the HR1123 specimen was several micrometers. Inside of the grain is plane without precipitates as seen in (a). The SADP can be indexed by the B2 parent phase. After aging the HR1123 specimen at 773 K for 2 h, we confirmed that fine precipitate of Ti3Ni4 with an average size of approximately 100 nm is formed homogeneously in the grain, which is shown in Supplementary material. Through cold rolling and subsequent heat-treatments, the grain size decreased to the nanoscale, as seen in Fig. 2(b–d). The grain size of the CR673 specimen and the NPCR673 specimen is nearly the same (~40 nm). This means that grain size is not influenced by the existence of Ti3Ni4, which is formed by the aging treatment before cold rolling. However, the grain size of the CR773 specimen (~100 nm) is significantly larger than that of the CR673 specimen. In addition, the shape of each grain is different. The grain of the CR673 specimen is elongated to the rolling direction, which is characteristic of rolled specimen. On the other hand, the grain of the CR773 specimen is equiaxed, which is characteristic of recrystallized specimens. It is speculated that the influence of cold rolling is largely retained in the CR673 and NPCR673 specimens, but not in the CR773 specimen. Due to the small grain size and large stress field, it is difficult to identify the Ti3Ni4 precipitates in the TEM bright field images of the CR773, CR673 and NPCR673 specimens. Nevertheless, reflections from Ti3Ni4 phase were detected for the CR773 and CR673 specimens while not for the NPCR673 specimen in the SADP. This implies that the Ti3Ni4 phase, which was formed by the aging treatment before cold rolling, remained after cold rolling and subsequent heat-treatment at 673 K or 773 K. On the other hand, the absence of the Ti3Ni4 phase in NPCR673 implies that precipitation of Ti3Ni4 after cold rolling is suppressed, possibly due to the small grain size. It is consistent with a previous report by
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Prokofiev et al. in which they confirmed that the precipitation of Ti3Ni4 was suppressed in grains of cross-sectional equivalent diameter below approximately 150 nm [29]. Incidentally, according to the DSC results,
(a) HR1123
110B2
101B2
000B2
500nm
(b) CR773
200nm
(c) CR673
200nm
(d) NPCR673
310B2 220B2 211B2 200B2 212P 110B2 202P 100B2
310B2 220B2 211B2 200B2 212P 110B2 202P 100B2
310B2 220B2 211B2 200B2 110B2 100B2
200nm
(e) R phase (f) 10nm
TS
the R-phase is expected to exist at room temperature in the CR673 and NP CR673 specimens. To identify the R-phase, we made a high resolution observation for the CR673. Fig. 2 (e) is a high resolution image of CR673 specimen, and (f) is the fast Fourier transformation of the red rectangular area. We notice 1/3〈110〉 reflections characteristic to the R-phase. Since the 1/3〈110〉 appears only in one direction, the red rectangular region will be a single variant of the R-phase. The solid curves in Fig. 3 are the stress-strain curves of the initial cycle of all the examined specimens. All the tests were made sufficiently above the Af temperature. They show a clear plateau corresponding to the stress-induced MT. The shape recovery is almost perfect, except for the HR1123 specimen in which a residual strain of ~3.6% can be seen. We evaluated the critical stress for the transformation by the tangent method, and the temperature dependence of the critical stress is shown in Fig. 4. It increases linearly as temperature increases for all the specimens. From the slope in Fig. 4, we can obtain the entropy change ΔS of the MT by using the ClausiusClapeyron relation dσ/dT = − ΔS/Δε, where Δε is the transformation strain. Then, we can estimate the adiabatic temperature change (ΔT) caused by the stress-induced MT as ΔT=−TΔS/CP. The value of ΔT estimated for the HR1123, CR773, CR673 and NP CR673 specimens are −31.5 K, −27.7 K, −28.5 K and −25.8 K, respectively, for the stress removing process. These values are nearly the same as those estimated from the latent heat mentioned above except for CR773 specimen. The experimental value of ΔT obtained by the stressreducing process is −8.2 K, −24.3 K, −26.4 K and −24.0 K for the HR1123, CR773, CR673 and NPCR673 specimens, respectively. The experimental value is approximately 0.92 times the estimated value for all the specimens except for HR1123 specimen. The large deviation between the estimated vale and the experiment in HR1123 specimen is due to incomplete reverse transformation. The small deviation (~8%) in other specimens is probably due to insufficient adiabatic condition of the experiment. The stress-strain curves and adiabatic temperature change of the initial cycle mentioned above imply that there is no significant difference in superelasticity and elastocaloric effect among CR773, CR673 and NPCR673 specimens. However, the evolution of the stress-cycle is significantly different among these three specimens, as described below. The evolution of the stress-strain curves and adiabatic temperature decrease in the stress-removing process for the CR773, CR673 and NPCR673 specimens are shown in Fig. 3(b–d) and (b′–d′). The test temperature is approximately 10 K above Af for each specimen. The transformation strain and adiabatic temperature decrease show visible decay as the number of stress cycles increases for all specimens. The decay is the most significant in the CR773 specimen, both Δε and ΔT decreases by ~45% through 100 cycles (Δε decreases from 5.6% to 2.7% and ΔT reduces from −24.3 K to −13.3 K). The decay for the CR673 specimen is ~10% and that for the NPCR673 specimen is ~20%. We briefly discuss the reason why the decay of Δε and ΔT is small in the CR673 specimen when compared with other specimens. It is reported that the decrease of the Δε during stress cycling is caused by the accumulation of defects [23,24,31]. Therefore, we may regard that the introduction of defects through stress cycling is smallest for the CR673 specimen. The nanocrystalline microstructure of the CR673 specimen caused by cold rolling and subsequent heat treatment at 673 K is probably the main reason for the strengthening. However, we notice
Fig. 2. Bright field images, selected area diffraction patterns and high resolution TEM for the Ti-50.8Ni (at.%) alloy after different thermomechanical treatments: HR1123 (a), CR773 (b), CR673 (c) and NPCR673 (d). The high resolution TEM of the CR673 specimen (e) and corresponding fast Fourier transformation of the red rectangular area (f). B2 and P on index indicate the parent phase and Ti3Ni4, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
H. Chen et al. / Scripta Materialia 162 (2019) 230–234
(a)
HR1123 T=308K
(a’)
330
29.1K
600
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315
300 -8.2K
0
-24.3K
300
100th
-13.3K
(b)
(b’)
300 330 315 300
0
1st 100th
1st
(d) NPCR673 T=328K
300 0 0
(d’)
1st 100th
2
4
6
Strain (%)
345 330 315
100th
-18.9K
0 600
100th
-23.5K
300
-26.4K
CR673 T=338K
(c’)
-24.0K
600
1st
(c)
Temperature (K)
Stress (MPa)
600
1st
CR773 T=318K
3s
330 315
20s
Time (s)
Fig. 3. Initial cycle stress-strain curve and the adiabatic temperature change, ΔT, of the HR1123 specimen (a, a′). Stress-strain curves and the adiabatic temperature change, ΔT, as a function of time of the Ti-50.8Ni (at.%) alloy after different thermomechanical treatments, CR773 (b, b′), CR673 (c, c′), and NPCR673 (d, d′) specimens. Solid curves are the results of the first cycle, thin curves are those of every 10 cycles. The testing temperature and ΔT caused by the reverse martensite transformation are indicated in the figure.
Critical Stress(MPa)
that the decay of the CR673 specimen is obviously small when compared with the NCR673 specimen. The only difference between the two specimens is the existence of Ti3Ni4. Although the grain size of
800
600
400
300
320
340
6.1MPa/K 5.2MPa/K 6.0MPa/K 5.8MPa/K
HR1123 CR773 CR673 NPCR673
360
380
Temperature(K)
Fig. 4. The temperature dependence of the critical stress for induced martensite transformation for the Ti-50.8Ni (at.%) alloy after different thermomechanical treatments (HR1123, CR773, CR673, NPCR673).
the CR673 specimen is several tens nanometers, it is likely that Ti3Ni4 still contribute as effective precipitate which prevents the movement of dislocations. The stable eCE for the CR673 specimen indicates that it could be a good candidate for an eCE material. When the SMAs are used as refrigeration materials, the energy efficiency of the materials (η) is an important parameter for the cooling technologies. According to Ref. [32], η is defined as η = ∣ Qh/W∣, where W is the total mechanical work done to the refrigerant material in a cycle, and Qh is heat flow to the material from cold thermal bath in the cycle. We approximate | Qh | as |-Cp∙ΔT| using experimentally obtained ΔT, and |W| as the enclosed area of the stress-strain curve [6]. The value of η thus evaluated for the CR673 specimen is 7.4 for the first cycle and 12.1 for the 100th cycle. The increase of η is due to the decrease of stress hysteresis. The high value of η indicates that the CR673 specimen exhibits good refrigerating capacity [33]. In conclusion, Ti3Ni4 phase was successfully introduced in nanocrystalline Ti-50.8Ni alloy by an aging heat-treatment before cold rolling and subsequent heat-treatment. The specimen prepared by this process exhibits stable superelasticity and elastocaloric effect because of the combination of precipitation strengthening and grain-refinement strengthening. This method could be applied for the preparation of highly stable shape memory alloys, which can be used for solid-state refrigeration. Supplementary data to this article can be found online at https://doi. org/10.1016/j.scriptamat.2018.11.024.
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Acknowledgement This work was funded by the National Youth Science Foundation (No. 51871151, No. 51501113), the National Key R&D Program of China (No. 2017YFB0703003, No. 2017YFB0406000), National Natural Science Foundation of China (U1564203) and Iketani Science and Technology Foundation (No. 0301012-A). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]
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