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Dimension changes induced by phase transformation behaviors in Ni-Mn-Cu alloys
Materials Science & Engineering A 689 (2017) 115–121
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Dimension changes induced by phase transformation behaviors in Ni-MnCu alloys
MARK
⁎
Shuiyuan Yanga, , Jialin Wua, Yuding Liua, Rongpei Shib, Cuiping Wanga, Xingjun Liua a b
Fujian Key Laboratory of Materials Genome, College of Materials, Xiamen University, Xiamen 361005, PR China Department of Materials Science and Engineering, Ohio State University, 2041 College Road, 477 Watts Hall, Columbus, OH 43210, USA
A R T I C L E I N F O
A BS T RAC T
Keywords: Ni-Mn-Cu Phase transformation Microstructure Dimension effect
The phase transformation behaviors, microstructural evolutions and dimension changes of Ni50-xCuxMn50 (x=8, 10, 12) (at%) were investigated. The results show that the phase transformation sequence of Ni50xCuxMn50 alloys obtained by quenching from 900 °C during heating and cooling is: γ(fcc) → non-martensite L10 phase→14M martensite↔B2↔γ(fcc). As these alloys are annealed under a variety of temperatures between 500 °C and 900 °C, different microstructural evolutions occur. The interesting changes that occurred during the heating and cooling of the above phase transformation are observed in Ni50-xCuxMn50 alloys. Due to the transformations of non-martensite L10 phase→14M martensite→B2, two consecutive dimension expansions are observed during heating for the alloys quenched from 900 °C. The largest dimension expansion can reach to 2.2% in Ni40Cu10Mn50 alloy. The reversible transformation behaviors of 14M martensite↔B2 and B2↔γ result in the reversible dimension changes. During cycling or as the alloys quench from 700 °C, the irreversible phase transformation behaviors and the corresponding dimension changes gradually recede and even disappear.
1. Introduction In general, some planes and directions might contract while others expand under the structural transformation from one phase to the other, such as reversible martensitic transformation. The changes of the crystal structure might be accompanied by the increase of decrease of the sample dimension. To date, this dimension change is usually too small to be utilized directly in practice. For example, only about 0.03% of dimension change was observed in Ni3Ta shape memory alloy (SMA) [1]. However, because no additional deformation is required at service, large and reversible dimension change is desirable for the applications in actuators. The two-way shape memory effect (TWSME) is an effective thermomechanical training used to obtain a large dimension change during reversible martensitic transformation without deformation. Although Ni-Ti SMAs have been widely utilized for the actuators due to their excellent TWSMEs and room temperature ductility, their developments for high temperature applications are hindered because of low martensitic transformation temperatures [2,3]. Several other SMAs also exhibit available TWSMEs, such as Ni3Ta [1], Ni54Mn25Ga20.7Gd0.3 [4] alloys and Ni-Mn-Ga single crystals [5,6]. However, the requirements of practical production cannot be satisfied by the room temperature ductility [7–11]. These deficiencies motivate us to further explore potential functional alloys with desired ductility
⁎
and dimension change, in which they can be processed by traditional technique at room temperature and operate at high temperature. Ni-Mn binary SMAs undergo a reversible martensitic transformation between high temperature B2 parent and low temperature tetragonal L10 martensite [12,13]. Unfortunately, because of its extreme brittleness, Ni-Mn binary SMAs have not received much attentions. For instance, the cracks were directly observed after quenching in Ni45Mn55 alloy [12] and the compressive fracture strain of Ni50.2Mn49.8 alloy was only about 2% [14]. Additionally, the maximum shape recovery rate was about only 30% under small compressive pre-strain of 2% [15]. In order to improve the ductility and shape recovery properties of Ni-Mn SMAs, such as Ni-Mn-Al and Ni-Mn-Ti alloys, ternary alloying has been used. The grain refinement with Ti alloying and the formation of γ phase with Al alloying may result in the ductility improvements. Therefore, in order to promote the formation of γ phase, Cu is selected as the alloying elements into Ni50Mn50 master alloy in this study. The approach is different from the previous investigations in which a wide composition region of B2 parent forms at high temperature in ternary Ni-Mn-Ti/Al systems [13,19] while alloying Al or Ti elements [12–18]. It results from the concurrence of B2 phase in Ni-Ti, Ni-Al and Ni-Mn binary systems [12,20]. However, Ni-Cu binary system has an infinite solid solution of γ phase [20], but B2 phase
Corresponding author. E-mail address: [email protected] (S. Yang).
http://dx.doi.org/10.1016/j.msea.2017.02.027 Received 6 October 2016; Received in revised form 7 February 2017; Accepted 7 February 2017 Available online 09 February 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
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Table 1 The designations of the studied alloys with their chemical compositions measured for the samples annealed at 900 °C followed by an ice-water quenching. The standard deviations of the measured compositions are ± 0.5 at%. Samples
Ni42Cu8Mn50 Ni40Cu10Mn50 Ni38Cu12Mn50
Designation
Cu8 Cu10 Cu12
Chemical compositions (at%) Ni
Mn
Cu/Co
42.49 40.91 39.05
49.24 48.65 48.71
8.27 10.44 12.24
ceases to exist. Therefore, the present studied Ni-Mn-Cu ternary alloys exhibit a single γ phase after high temperature heat treatment. Subsequently, an observation is conducted of a phase transformation behavior induced an interesting dimension change in Ni-Mn-Cu alloys. Next, an investigation of their phase transformation behaviors and corresponding microstructural evolutions were investigated. 2. Experimental procedures Polycrystalline buttons with about 40 g with the nominal compositions of Ni50-xCuxMn50 (x=8, 10, 12) (at%) alloys were prepared from pure metals. This was done by using a non-consumable tungsten electrode to melt them under a high purity argon atmosphere. In order to ensure compositional homogeneity, each button was re-melted five times. Prior to an ice-water quench, all of the buttons were further cut into small bulks, sealed into vacuum quartz ampoules, and annealed at 900 °C for 24 h. By using an average of five measurements, electron probe microanalysis (EPMA) (JEOL, JXA-8100) was used to determine chemical compositions of each sample, which were given in Table 1. Afterwards, prior to an ice-water quench, the above bulks quenched were again annealed at 800 °C, 700 °C, 600 °C and 500 °C for 72, 240, 240 and 240 h respectively, all followed an ice-water quench. Throughout the rest of the text, the studied alloys are denoted as designation-annealing temperatures for convenience reasons. For example, Cu8-900 sample represents the Ni42Cu8Mn50 alloy annealed at 900 °C for 24 h, and Cu8-800 sample represents that Cu8-900 sample was further annealed at 800 °C for another 72 h. The microstructures were characterized by optical microscopy (OM) using Leica DMI 5000 M. For both heating and cooling, the phase transformation temperatures were determined by differential scanning calorimetry (DSC) (Netzsch STA 404) at a rate of 10 °C min−1. By using a PANalytical X'pert PRO X-ray diffractometer with Cu Kα radiation and transmission electron microscopy (TEM, JEOL 2100) the crystal structure was identified. Precision Ion Polishing System using Gatan PIPS II Model 695 obtained the thin foils for TEM observation. For both heating and cooling, the in-situ high-temperature X-ray diffraction was used to investigate the phase transformation behaviors of the alloys with a rate of 10 °C min−1. Several cylindrical specimens (ϕ3 mm×5 mm) were cut from the quenched buttons. All the studied samples used for the investigations of the dimension change by TMA tests were perpendicular orientation to the cooling direction. In order to investigate the dimension changes, each sample was directly heated through phase transformation temperature without the compressive deformation in order to investigate the dimension changes. Subsequent to the removal of the height change due to thermal expansion property the dimension changes were calculated by using a Thermomechanical Analyzer (Netzsch TMA 402) with a rate of 10 °C min−1 for both heating and cooling.
Fig. 1. X-ray diffraction patterns (a–c) and optical micrographs (d–f) of the alloys quenched from 900 °C. (a) and (d) Cu8-900 alloy; (b) and (e) Cu10-900 alloy; (c) and (f) Cu12-900 alloy.
of all alloys quenched from 900 °C. Contrary to Ni-Mn-Ti/Al alloys [12–18], all ternary alloys consist of a single fcc γ phase. In this study, replacing Ni by adding 8% Cu in Ni50Mn50 alloy leads to the formation of a single γ phase when annealing at 900 °C, suggesting that the alloying of Cu at high temperatures allows the boundary of γ phase to expand. This shifts the B2 phase region to a lower temperature in NiMn-Cu systems. In contrast, because of the concurrence of B2 phase in Ni-Ti, Ni-Al and Ni-Mn binary systems [12,13,19,20], the alloying of Ti or Al is able to expand the B2 phase region. For example, the B2 phase still exists even up to 1100 °C in Ni-Mn-Al ternary system [13,19] when compared with Ni-Mn binary system [12,20]. Complex phase transformation behaviors upon heating and cooling are clearly observed and presented in Fig. 2 where are the DSC curves of all studied alloys quenched from 900 °C. During heating (Fig. 2a), four different types of transformation peaks are identified: (1) Very broad exothermal peaks around 300–550 °C are available in all the samples (peak 1), (2) subtle phase transformations are identified and highlighted by red dashed circles (peak 2), (3) three sharp endothermic peaks are present (peak 3), and (4) indicated by black dashed circles are the small endothermic peaks that were present in all of the samples at very high temperatures (peak 4). From the cooling curves in Fig. 2b, it is found that during cooling, the reverse transformation processes of aforementioned peaks (3) and (4) take place and correspond to the transformation peaks (5) and (6). Table 2 provides details about the transformation starting and finishing temperatures during heating. In order to investigate the complex phase transformation behaviors mentioned above, all of the samples were also annealed at 500 °C,
3. Results and discussion 3.1. Phase transformation behaviors and microstructural evolutions Fig. 1 shows the optical micrographs and X-ray diffraction patterns 116
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Fig. 2. DSC heating (a) and cooling (b) curves of the alloys quenched from 900 °C.
from γ phase at 350 °C. In their paper, Adachi and Wayman [12] revealed tempering results in formation of non-martensite L10 phase. Secondly, there are additional peaks due to the presence of 14M martensite at 550 °C, which occur more rapidly as the temperature increases to 650 °C. This indicates the formation of a 14M martensite. Meanwhile, resulting from the reverse martensitic transformation from 14M martensite to B2 parent, the B2 also begins to appear at 650 °C. The DSC curves and Table 2 give all of the related phase transformation behaviors of the studied alloys. In Fig. 2, the irreversible phase transformation (1) is identified to the formation of nonmartensite L10 phase during heating. During heating, the small and irreversible peak (2) is related to the transformation from the L10 phase to 14M martensite. The reversible martensitic transformation between 14M martensite and B2 parent is indicated by the reversible transformation peaks (3) and (5). The diffusion transformation between B2 phase and γ phase results in the peaks (4) and (6). Therefore, the phase transformation behaviors of the present Cu8-900, Cu10-900 and Cu12-900 alloys are: γ→L10 phase→14M martensite↔B2↔γ, with the “→” symbol representing that this transformation only occurs during heating and is irreversible, and the “↔” symbol indicating that this transformation is reversible.
600 °C, 700 °C and 800 °C according to the temperature ranges of each transformation in Table 2. Figs. 3 and 4 show the optical micrographs and the X-ray diffraction patterns of these alloys. From the results of the microstructural observations and X-ray diffraction patterns, it is found that: (i) for alloys annealed at 500 °C, the morphological feature of martensite is not observed, showing two-phase microstructure of non-martensite L10 phase and fcc γ phase, (ii) Cu8-600, Cu8-800, Cu10-600, Cu10-800, Cu12-600 and Cu12-700 samples have a twophase microstructure of modulated 14M martensite and γ phase, (iii) Cu8-700 and Cu10-700 samples are a single 14M martensite, and (iv) Cu12-800 sample is a single fcc γ phase. The TEM characterizations of Cu8-(500, 600, 700, 800) samples were carried out in order to further confirm the crystal structure of the alloys, and the results are shown in Fig. 5. The TEM results in Fig. 5a for Cu8-500 sample does not seem to be a martensite structure. The striations inside big grains are not straight, and the electron diffraction pattern (Fig. 5b) only exhibits a L10 structure, which does not look like twinning. The fcc γ phase is confirmed by the bright field (BF) image and the corresponding selected area diffraction pattern (SADP) in Fig. 5c and d. However, a 14M-martensite structure is found to co-exist with fcc γ phase for Cu8-600 sample as shown in Fig. 5e–h. Moreover, a single phase of 14M martensite is identified for Cu8-700 sample in Fig. 5i and j. While a two-phase microstructure consisted of 14M martensite and fcc γ phase is observed for the Cu8-800 sample in Fig. 5k - n. As shown in Fig. 6, in-situ high-temperature X-ray diffraction analyses are carried out for Cu8-900 alloy that has a single γ phase before testing in order to identify the phase behaviors of the studied alloys. It is noted that the transformation peaks of L10→14M are not clear in Fig. 2. The transformation starting and finishing temperatures can not be determined in Fig. 2, because this transformation may be occur in a wide temperature range. According to the following TMA curves, this transformation occurs between 470 °C and 610 °C for Cu8-900 alloy. Therefore, during high temperature X-ray diffraction test, the 350 °C, 550 °C and 650 °C were selected based on the phase transformation temperature ranges in Table 2. Initially, the diffraction peaks of the L10 phase are already present transformed
3.2. Dimension changes Fig. 7 shows the dimension changes characterized by TMA tests, upon heating (solid curve) and cooling (dashed curve) of all alloys quenched from 900 °C. As can be seen clearly, the complex dimension changes occur during heating and cooling. A dimension expansion and contraction upon heating is indicated by the stages AB and CD for each alloy. The dimension expansions in the stage AB finish with a two-step manner with two consecutive stages AB1 and B1B and a slope change occurring at B1. The dimension changes are about 0.9% (AB1) and 1.2% (B1B) for Cu8-900 sample, 0.8% (AB1) and 1.4% (B1B) for Cu10-900 sample, and about 0.2% (AB1) and 0.3% (B1B) for Cu12-900 sample, respectively. In contrast to AB, contractions for all alloys in CD finish in a one-step manner. The dimension contractions are about 0.5%, 1.2%
Table 2 The phase transformation starting and finishing temperatures of the alloys quenched from 900 °C during heating determined by DSC (*) and TMA (#) tests. The standard deviations of the measured values are ± 3 °C. Samples
γ→L10 (°C) Starting
Cu8-900 Cu10-900 Cu12-900
*
297 /287*/292*/-
L10→14M (°C) Finishing *
508 /-/528*/-
Starting #
-/476 -/489# -/546#
β(B2)→γ (°C)
14M→β(B2) (°C) Finishing #
-/611 -/633# -/652#
117
Starting #
-/611 -/633# -/652#
Finishing *
Starting #
676 /662 669*/656# 673*/692#
*
793 /802 -/656# -/794#
Finishing #
881*/847# -/834# -/-
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Fig. 3. Microstructural evolutions of the studied alloys. (a1) Cu8-500; (a2) Cu8-600; (a3) Cu8-700; (a4) Cu8-800; (b1) Cu10-500; (b2) Cu10-600; (b3) Cu10-700; (b4) Cu10-800; (c1) Cu12-500; (c2) Cu12-600; (c3) Cu12-700; (c4) Cu12-800.
dimension changes that occurred during heating and cooling in Fig. 7. The second TMA tests were carried out for the same sample after the first TMA test in Fig. 7 in order to confirm the reversibility above transformation behaviors and dimension changes, the results are shown in Fig. 8. It is learned that during heating, the irreversible phase transformation behavior of L10 phase→14M recedes and even disappears during the second TMA tests. It indicates that the irreversible phase transformation behaviors must be a temporal effect induced by rapid quenching process. However, the reversible transformations, 14M (martensite)↔B2↔γ(fcc) are still reversible along with the dimension changes. Fig. 9 shows that before testing (Fig. 3), the TMA curves of Cu8-700, Cu10-700 and Cu12-700 alloys have a single 14M martensite or two-phase structure of dominant 14M martensite +γ phase. Due to the difference in the microstructure, the alloys quenched from 700 °C exhibits the reversible phase transformations of 14M (martensite)↔B2↔γ(fcc) and the corresponding reversible
and 0.2% for Cu8-900, Cu10-900 and Cu12-900 alloys, respectively. Moreover, it is found that the dimension changes at the stages CD and B1B are reversible. The corresponding reverse dimension changes happen at the stages EF (expansion) and GH (contraction) upon cooling, respectively. However, the dimension changes at the stage AB1 are irreversible. The dimension changes in Fig. 7 can be well understood according to the detailed analysis of phase transformation behaviors upon heating and cooling. In each studied alloy, the reversible dimension changes at stages CD (contraction) and EF (expansion) result from the diffusion transformation of the B2 phase↔γ phase at high temperatures. Twostep dimension expansions at the stages AB1 (irreversible) and B1B (the reverse dimension contraction occurs at the stage GH upon cooling) are caused by the transformation behaviors of L10 phase→14M martensite and reversible martensitic transformation between 14M martensite↔ B2 parent. Table 3 represents measurements taken of all of the
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Fig. 4. X-ray diffraction patterns of the alloys annealed at different temperatures (a) 500 °C, (b) 600 °C, (c) 700 °C and (d) 800 °C.
mations, namely the formation of 14M martensite and B2 parent. NiMn binary SMAs usually undergo a reversible martensitic transformation between high temperature B2 parent and low temperature tetragonal L10 martensite [12,13], without 14M martensite. The present results show that the Cu addition is beneficial for the formation of 14M martensite in Ni-Mn alloys, which is similar to the Ni-Mn-Al alloys [13]. We also investigated the effects of other elements (Co and Fe) on the dimension changes of Ni50Mn50 master alloy. It is found that such large dimension change disappears in Ni-Mn-Co/Fe alloys due to the absence of 14M martensite. Additionally, the dimension change differs when varying the Cu content. It mainly depends on the unit-cell volume and the corresponding unit-cell numbers, the volume fraction of each phase during transformation. They are also closely related to the alloy compositions, the degree of phase transformation during heating etc. The individual influence of above-mentioned factors is still difficult to quantify. It needs more further works.
dimension changes. Even though they can compare with some TWSMEs in several SMAs [1,4,8,22], the current dimension changes obtained are clearly larger than those reported [1,21]. Additionally, the mechanical properties of the alloys were also studied by compression tests. The results show that those single γ(fcc) phase alloys are still no fracture when compressed to a deformation of 50%, showing a good ductility. When these alloys were annealed during 600–800 °C, the compressive fracture strain clearly decreases with the increase of the 14M martensite content. The thermoelastic martensitic transformation usually results in a small dimension change during thermo cycle. In this study, the samples used for TMA tests are single γ(fcc) phase. During the following heating, other phase transformation behaviors (γ(fcc)→non-martensite L10 phase→14M martensite) previously happen before reverse martensitic transformation (14M martensite↔B2). Therefore, such large dimension expansion during heating results from two phase transfor-
Fig. 5. Typical microstructures of Cu8-500, Cu8-600, Cu8-700 and Cu8-800 alloys observed by TEM tests. (a–d) Cu8-500 alloy, the BF image (a) and the corresponding SADP (b) of the non-martensite L10; the BF image (c) and the corresponding SADP (d) of fcc γ phase. (e–h) Cu8-600 alloy, the BF image (e) and the corresponding SADP (f) of 14M martensite; the BF image (g) and the corresponding SADP (h) of fcc γ phase. (i and j) Cu8-700 alloy, the BF image (i) and the corresponding SADP (j) of 14M martensite; (k–n) Cu8-800 alloy, the BF image (k) and the corresponding SADP (l) of 14M martensite; the BF image (m) and the corresponding SADP (n) of fcc γ phase. The red circles mean that the SADP are taken in these areas.
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Table 3 The observed dimension changes during different phase transformations in Fig. 8. The symbols (+) and (–) respectively represent the dimension expansion and contraction during heating and cooling. The standard deviations of the measured values are ± 0.1%. Samples
Cu8-900 Cu10-900 Cu12-900
Dimension change (%) L10→14M
14M↔β(B2)
β(B2)↔γ
+0.9 +0.8 +0.2
+1.2/–0.4 +1.4/–0.3 +0.2/–0.2
–0.5/+0.2 –1.1/+0.4 −0.2/+0.3
Fig. 6. In-situ high-temperature X-ray diffraction patterns of the Cu8-900 alloy during heating.
Fig. 8. The second TMA tests of the same samples after the first TMA tests in Fig. 7.
Fig. 7. Dimension changes during heating and cooling for different alloys quenched from 900 °C characterized by TMA tests. The symbols (+) and (−) respectively represent the dimension expansion and contraction during heating and cooling.
4. Conclusions Fig. 9. TMAs curves of the Cu8-700, Cu10-700 and Cu12-700 samples.
The phase transformation behaviors, and the corresponding microstructural evolutions and dimension changes in Ni50-xCuxMn50 (x=8, 120
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10, 12) alloys were investigated in this study. The results reveal that:
References
(1) The alloys quenched from 900 °C exhibit complex phase transformation behaviors of γ(fcc)→non-martensite L10 phase→14M (martensite)↔B2↔γ(fcc). The phase transformation behaviors result in the different microstructures while these alloys quenched from 900 °C are further annealed at different temperatures (500– 800 °C). The irreversible phase transformation behaviors during heating are confirmed to be a tempering effect resulted from quenching process. (2) Interesting dimension changes induced by phase transformation behaviors are obtained in Ni-Mn-Cu alloys. Two consecutive dimension expansions are observed during heating due to the transformations of non-martensite L10 phase→14M martensite→ B2 for the alloys quenched from 900 °C. 2.2% in Ni40Cu10Mn50 alloy is the largest dimension expansion during heating. The reversible dimension changes resulted from the reversible transformation of B2↔γ. During cycling or for the alloys quenching from 700 °C, the irreversible phase transformation of non-martensite L10 phase→14M martensite and the corresponding dimension changes gradually recede and even disappear.
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Acknowledgements We acknowledge the financial supports from the Fundamental Research Funds for the Central Universities, grant number 20720160078, the National Natural Science Foundation of China, grant number 51571168.
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Dimension changes induced by phase transformation behaviors in Ni-MnCu alloys
MARK
⁎
Shuiyuan Yanga, , Jialin Wua, Yuding Liua, Rongpei Shib, Cuiping Wanga, Xingjun Liua a b
Fujian Key Laboratory of Materials Genome, College of Materials, Xiamen University, Xiamen 361005, PR China Department of Materials Science and Engineering, Ohio State University, 2041 College Road, 477 Watts Hall, Columbus, OH 43210, USA
A R T I C L E I N F O
A BS T RAC T
Keywords: Ni-Mn-Cu Phase transformation Microstructure Dimension effect
The phase transformation behaviors, microstructural evolutions and dimension changes of Ni50-xCuxMn50 (x=8, 10, 12) (at%) were investigated. The results show that the phase transformation sequence of Ni50xCuxMn50 alloys obtained by quenching from 900 °C during heating and cooling is: γ(fcc) → non-martensite L10 phase→14M martensite↔B2↔γ(fcc). As these alloys are annealed under a variety of temperatures between 500 °C and 900 °C, different microstructural evolutions occur. The interesting changes that occurred during the heating and cooling of the above phase transformation are observed in Ni50-xCuxMn50 alloys. Due to the transformations of non-martensite L10 phase→14M martensite→B2, two consecutive dimension expansions are observed during heating for the alloys quenched from 900 °C. The largest dimension expansion can reach to 2.2% in Ni40Cu10Mn50 alloy. The reversible transformation behaviors of 14M martensite↔B2 and B2↔γ result in the reversible dimension changes. During cycling or as the alloys quench from 700 °C, the irreversible phase transformation behaviors and the corresponding dimension changes gradually recede and even disappear.
1. Introduction In general, some planes and directions might contract while others expand under the structural transformation from one phase to the other, such as reversible martensitic transformation. The changes of the crystal structure might be accompanied by the increase of decrease of the sample dimension. To date, this dimension change is usually too small to be utilized directly in practice. For example, only about 0.03% of dimension change was observed in Ni3Ta shape memory alloy (SMA) [1]. However, because no additional deformation is required at service, large and reversible dimension change is desirable for the applications in actuators. The two-way shape memory effect (TWSME) is an effective thermomechanical training used to obtain a large dimension change during reversible martensitic transformation without deformation. Although Ni-Ti SMAs have been widely utilized for the actuators due to their excellent TWSMEs and room temperature ductility, their developments for high temperature applications are hindered because of low martensitic transformation temperatures [2,3]. Several other SMAs also exhibit available TWSMEs, such as Ni3Ta [1], Ni54Mn25Ga20.7Gd0.3 [4] alloys and Ni-Mn-Ga single crystals [5,6]. However, the requirements of practical production cannot be satisfied by the room temperature ductility [7–11]. These deficiencies motivate us to further explore potential functional alloys with desired ductility
⁎
and dimension change, in which they can be processed by traditional technique at room temperature and operate at high temperature. Ni-Mn binary SMAs undergo a reversible martensitic transformation between high temperature B2 parent and low temperature tetragonal L10 martensite [12,13]. Unfortunately, because of its extreme brittleness, Ni-Mn binary SMAs have not received much attentions. For instance, the cracks were directly observed after quenching in Ni45Mn55 alloy [12] and the compressive fracture strain of Ni50.2Mn49.8 alloy was only about 2% [14]. Additionally, the maximum shape recovery rate was about only 30% under small compressive pre-strain of 2% [15]. In order to improve the ductility and shape recovery properties of Ni-Mn SMAs, such as Ni-Mn-Al and Ni-Mn-Ti alloys, ternary alloying has been used. The grain refinement with Ti alloying and the formation of γ phase with Al alloying may result in the ductility improvements. Therefore, in order to promote the formation of γ phase, Cu is selected as the alloying elements into Ni50Mn50 master alloy in this study. The approach is different from the previous investigations in which a wide composition region of B2 parent forms at high temperature in ternary Ni-Mn-Ti/Al systems [13,19] while alloying Al or Ti elements [12–18]. It results from the concurrence of B2 phase in Ni-Ti, Ni-Al and Ni-Mn binary systems [12,20]. However, Ni-Cu binary system has an infinite solid solution of γ phase [20], but B2 phase
Corresponding author. E-mail address: [email protected] (S. Yang).
http://dx.doi.org/10.1016/j.msea.2017.02.027 Received 6 October 2016; Received in revised form 7 February 2017; Accepted 7 February 2017 Available online 09 February 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
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Table 1 The designations of the studied alloys with their chemical compositions measured for the samples annealed at 900 °C followed by an ice-water quenching. The standard deviations of the measured compositions are ± 0.5 at%. Samples
Ni42Cu8Mn50 Ni40Cu10Mn50 Ni38Cu12Mn50
Designation
Cu8 Cu10 Cu12
Chemical compositions (at%) Ni
Mn
Cu/Co
42.49 40.91 39.05
49.24 48.65 48.71
8.27 10.44 12.24
ceases to exist. Therefore, the present studied Ni-Mn-Cu ternary alloys exhibit a single γ phase after high temperature heat treatment. Subsequently, an observation is conducted of a phase transformation behavior induced an interesting dimension change in Ni-Mn-Cu alloys. Next, an investigation of their phase transformation behaviors and corresponding microstructural evolutions were investigated. 2. Experimental procedures Polycrystalline buttons with about 40 g with the nominal compositions of Ni50-xCuxMn50 (x=8, 10, 12) (at%) alloys were prepared from pure metals. This was done by using a non-consumable tungsten electrode to melt them under a high purity argon atmosphere. In order to ensure compositional homogeneity, each button was re-melted five times. Prior to an ice-water quench, all of the buttons were further cut into small bulks, sealed into vacuum quartz ampoules, and annealed at 900 °C for 24 h. By using an average of five measurements, electron probe microanalysis (EPMA) (JEOL, JXA-8100) was used to determine chemical compositions of each sample, which were given in Table 1. Afterwards, prior to an ice-water quench, the above bulks quenched were again annealed at 800 °C, 700 °C, 600 °C and 500 °C for 72, 240, 240 and 240 h respectively, all followed an ice-water quench. Throughout the rest of the text, the studied alloys are denoted as designation-annealing temperatures for convenience reasons. For example, Cu8-900 sample represents the Ni42Cu8Mn50 alloy annealed at 900 °C for 24 h, and Cu8-800 sample represents that Cu8-900 sample was further annealed at 800 °C for another 72 h. The microstructures were characterized by optical microscopy (OM) using Leica DMI 5000 M. For both heating and cooling, the phase transformation temperatures were determined by differential scanning calorimetry (DSC) (Netzsch STA 404) at a rate of 10 °C min−1. By using a PANalytical X'pert PRO X-ray diffractometer with Cu Kα radiation and transmission electron microscopy (TEM, JEOL 2100) the crystal structure was identified. Precision Ion Polishing System using Gatan PIPS II Model 695 obtained the thin foils for TEM observation. For both heating and cooling, the in-situ high-temperature X-ray diffraction was used to investigate the phase transformation behaviors of the alloys with a rate of 10 °C min−1. Several cylindrical specimens (ϕ3 mm×5 mm) were cut from the quenched buttons. All the studied samples used for the investigations of the dimension change by TMA tests were perpendicular orientation to the cooling direction. In order to investigate the dimension changes, each sample was directly heated through phase transformation temperature without the compressive deformation in order to investigate the dimension changes. Subsequent to the removal of the height change due to thermal expansion property the dimension changes were calculated by using a Thermomechanical Analyzer (Netzsch TMA 402) with a rate of 10 °C min−1 for both heating and cooling.
Fig. 1. X-ray diffraction patterns (a–c) and optical micrographs (d–f) of the alloys quenched from 900 °C. (a) and (d) Cu8-900 alloy; (b) and (e) Cu10-900 alloy; (c) and (f) Cu12-900 alloy.
of all alloys quenched from 900 °C. Contrary to Ni-Mn-Ti/Al alloys [12–18], all ternary alloys consist of a single fcc γ phase. In this study, replacing Ni by adding 8% Cu in Ni50Mn50 alloy leads to the formation of a single γ phase when annealing at 900 °C, suggesting that the alloying of Cu at high temperatures allows the boundary of γ phase to expand. This shifts the B2 phase region to a lower temperature in NiMn-Cu systems. In contrast, because of the concurrence of B2 phase in Ni-Ti, Ni-Al and Ni-Mn binary systems [12,13,19,20], the alloying of Ti or Al is able to expand the B2 phase region. For example, the B2 phase still exists even up to 1100 °C in Ni-Mn-Al ternary system [13,19] when compared with Ni-Mn binary system [12,20]. Complex phase transformation behaviors upon heating and cooling are clearly observed and presented in Fig. 2 where are the DSC curves of all studied alloys quenched from 900 °C. During heating (Fig. 2a), four different types of transformation peaks are identified: (1) Very broad exothermal peaks around 300–550 °C are available in all the samples (peak 1), (2) subtle phase transformations are identified and highlighted by red dashed circles (peak 2), (3) three sharp endothermic peaks are present (peak 3), and (4) indicated by black dashed circles are the small endothermic peaks that were present in all of the samples at very high temperatures (peak 4). From the cooling curves in Fig. 2b, it is found that during cooling, the reverse transformation processes of aforementioned peaks (3) and (4) take place and correspond to the transformation peaks (5) and (6). Table 2 provides details about the transformation starting and finishing temperatures during heating. In order to investigate the complex phase transformation behaviors mentioned above, all of the samples were also annealed at 500 °C,
3. Results and discussion 3.1. Phase transformation behaviors and microstructural evolutions Fig. 1 shows the optical micrographs and X-ray diffraction patterns 116
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Fig. 2. DSC heating (a) and cooling (b) curves of the alloys quenched from 900 °C.
from γ phase at 350 °C. In their paper, Adachi and Wayman [12] revealed tempering results in formation of non-martensite L10 phase. Secondly, there are additional peaks due to the presence of 14M martensite at 550 °C, which occur more rapidly as the temperature increases to 650 °C. This indicates the formation of a 14M martensite. Meanwhile, resulting from the reverse martensitic transformation from 14M martensite to B2 parent, the B2 also begins to appear at 650 °C. The DSC curves and Table 2 give all of the related phase transformation behaviors of the studied alloys. In Fig. 2, the irreversible phase transformation (1) is identified to the formation of nonmartensite L10 phase during heating. During heating, the small and irreversible peak (2) is related to the transformation from the L10 phase to 14M martensite. The reversible martensitic transformation between 14M martensite and B2 parent is indicated by the reversible transformation peaks (3) and (5). The diffusion transformation between B2 phase and γ phase results in the peaks (4) and (6). Therefore, the phase transformation behaviors of the present Cu8-900, Cu10-900 and Cu12-900 alloys are: γ→L10 phase→14M martensite↔B2↔γ, with the “→” symbol representing that this transformation only occurs during heating and is irreversible, and the “↔” symbol indicating that this transformation is reversible.
600 °C, 700 °C and 800 °C according to the temperature ranges of each transformation in Table 2. Figs. 3 and 4 show the optical micrographs and the X-ray diffraction patterns of these alloys. From the results of the microstructural observations and X-ray diffraction patterns, it is found that: (i) for alloys annealed at 500 °C, the morphological feature of martensite is not observed, showing two-phase microstructure of non-martensite L10 phase and fcc γ phase, (ii) Cu8-600, Cu8-800, Cu10-600, Cu10-800, Cu12-600 and Cu12-700 samples have a twophase microstructure of modulated 14M martensite and γ phase, (iii) Cu8-700 and Cu10-700 samples are a single 14M martensite, and (iv) Cu12-800 sample is a single fcc γ phase. The TEM characterizations of Cu8-(500, 600, 700, 800) samples were carried out in order to further confirm the crystal structure of the alloys, and the results are shown in Fig. 5. The TEM results in Fig. 5a for Cu8-500 sample does not seem to be a martensite structure. The striations inside big grains are not straight, and the electron diffraction pattern (Fig. 5b) only exhibits a L10 structure, which does not look like twinning. The fcc γ phase is confirmed by the bright field (BF) image and the corresponding selected area diffraction pattern (SADP) in Fig. 5c and d. However, a 14M-martensite structure is found to co-exist with fcc γ phase for Cu8-600 sample as shown in Fig. 5e–h. Moreover, a single phase of 14M martensite is identified for Cu8-700 sample in Fig. 5i and j. While a two-phase microstructure consisted of 14M martensite and fcc γ phase is observed for the Cu8-800 sample in Fig. 5k - n. As shown in Fig. 6, in-situ high-temperature X-ray diffraction analyses are carried out for Cu8-900 alloy that has a single γ phase before testing in order to identify the phase behaviors of the studied alloys. It is noted that the transformation peaks of L10→14M are not clear in Fig. 2. The transformation starting and finishing temperatures can not be determined in Fig. 2, because this transformation may be occur in a wide temperature range. According to the following TMA curves, this transformation occurs between 470 °C and 610 °C for Cu8-900 alloy. Therefore, during high temperature X-ray diffraction test, the 350 °C, 550 °C and 650 °C were selected based on the phase transformation temperature ranges in Table 2. Initially, the diffraction peaks of the L10 phase are already present transformed
3.2. Dimension changes Fig. 7 shows the dimension changes characterized by TMA tests, upon heating (solid curve) and cooling (dashed curve) of all alloys quenched from 900 °C. As can be seen clearly, the complex dimension changes occur during heating and cooling. A dimension expansion and contraction upon heating is indicated by the stages AB and CD for each alloy. The dimension expansions in the stage AB finish with a two-step manner with two consecutive stages AB1 and B1B and a slope change occurring at B1. The dimension changes are about 0.9% (AB1) and 1.2% (B1B) for Cu8-900 sample, 0.8% (AB1) and 1.4% (B1B) for Cu10-900 sample, and about 0.2% (AB1) and 0.3% (B1B) for Cu12-900 sample, respectively. In contrast to AB, contractions for all alloys in CD finish in a one-step manner. The dimension contractions are about 0.5%, 1.2%
Table 2 The phase transformation starting and finishing temperatures of the alloys quenched from 900 °C during heating determined by DSC (*) and TMA (#) tests. The standard deviations of the measured values are ± 3 °C. Samples
γ→L10 (°C) Starting
Cu8-900 Cu10-900 Cu12-900
*
297 /287*/292*/-
L10→14M (°C) Finishing *
508 /-/528*/-
Starting #
-/476 -/489# -/546#
β(B2)→γ (°C)
14M→β(B2) (°C) Finishing #
-/611 -/633# -/652#
117
Starting #
-/611 -/633# -/652#
Finishing *
Starting #
676 /662 669*/656# 673*/692#
*
793 /802 -/656# -/794#
Finishing #
881*/847# -/834# -/-
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Fig. 3. Microstructural evolutions of the studied alloys. (a1) Cu8-500; (a2) Cu8-600; (a3) Cu8-700; (a4) Cu8-800; (b1) Cu10-500; (b2) Cu10-600; (b3) Cu10-700; (b4) Cu10-800; (c1) Cu12-500; (c2) Cu12-600; (c3) Cu12-700; (c4) Cu12-800.
dimension changes that occurred during heating and cooling in Fig. 7. The second TMA tests were carried out for the same sample after the first TMA test in Fig. 7 in order to confirm the reversibility above transformation behaviors and dimension changes, the results are shown in Fig. 8. It is learned that during heating, the irreversible phase transformation behavior of L10 phase→14M recedes and even disappears during the second TMA tests. It indicates that the irreversible phase transformation behaviors must be a temporal effect induced by rapid quenching process. However, the reversible transformations, 14M (martensite)↔B2↔γ(fcc) are still reversible along with the dimension changes. Fig. 9 shows that before testing (Fig. 3), the TMA curves of Cu8-700, Cu10-700 and Cu12-700 alloys have a single 14M martensite or two-phase structure of dominant 14M martensite +γ phase. Due to the difference in the microstructure, the alloys quenched from 700 °C exhibits the reversible phase transformations of 14M (martensite)↔B2↔γ(fcc) and the corresponding reversible
and 0.2% for Cu8-900, Cu10-900 and Cu12-900 alloys, respectively. Moreover, it is found that the dimension changes at the stages CD and B1B are reversible. The corresponding reverse dimension changes happen at the stages EF (expansion) and GH (contraction) upon cooling, respectively. However, the dimension changes at the stage AB1 are irreversible. The dimension changes in Fig. 7 can be well understood according to the detailed analysis of phase transformation behaviors upon heating and cooling. In each studied alloy, the reversible dimension changes at stages CD (contraction) and EF (expansion) result from the diffusion transformation of the B2 phase↔γ phase at high temperatures. Twostep dimension expansions at the stages AB1 (irreversible) and B1B (the reverse dimension contraction occurs at the stage GH upon cooling) are caused by the transformation behaviors of L10 phase→14M martensite and reversible martensitic transformation between 14M martensite↔ B2 parent. Table 3 represents measurements taken of all of the
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Fig. 4. X-ray diffraction patterns of the alloys annealed at different temperatures (a) 500 °C, (b) 600 °C, (c) 700 °C and (d) 800 °C.
mations, namely the formation of 14M martensite and B2 parent. NiMn binary SMAs usually undergo a reversible martensitic transformation between high temperature B2 parent and low temperature tetragonal L10 martensite [12,13], without 14M martensite. The present results show that the Cu addition is beneficial for the formation of 14M martensite in Ni-Mn alloys, which is similar to the Ni-Mn-Al alloys [13]. We also investigated the effects of other elements (Co and Fe) on the dimension changes of Ni50Mn50 master alloy. It is found that such large dimension change disappears in Ni-Mn-Co/Fe alloys due to the absence of 14M martensite. Additionally, the dimension change differs when varying the Cu content. It mainly depends on the unit-cell volume and the corresponding unit-cell numbers, the volume fraction of each phase during transformation. They are also closely related to the alloy compositions, the degree of phase transformation during heating etc. The individual influence of above-mentioned factors is still difficult to quantify. It needs more further works.
dimension changes. Even though they can compare with some TWSMEs in several SMAs [1,4,8,22], the current dimension changes obtained are clearly larger than those reported [1,21]. Additionally, the mechanical properties of the alloys were also studied by compression tests. The results show that those single γ(fcc) phase alloys are still no fracture when compressed to a deformation of 50%, showing a good ductility. When these alloys were annealed during 600–800 °C, the compressive fracture strain clearly decreases with the increase of the 14M martensite content. The thermoelastic martensitic transformation usually results in a small dimension change during thermo cycle. In this study, the samples used for TMA tests are single γ(fcc) phase. During the following heating, other phase transformation behaviors (γ(fcc)→non-martensite L10 phase→14M martensite) previously happen before reverse martensitic transformation (14M martensite↔B2). Therefore, such large dimension expansion during heating results from two phase transfor-
Fig. 5. Typical microstructures of Cu8-500, Cu8-600, Cu8-700 and Cu8-800 alloys observed by TEM tests. (a–d) Cu8-500 alloy, the BF image (a) and the corresponding SADP (b) of the non-martensite L10; the BF image (c) and the corresponding SADP (d) of fcc γ phase. (e–h) Cu8-600 alloy, the BF image (e) and the corresponding SADP (f) of 14M martensite; the BF image (g) and the corresponding SADP (h) of fcc γ phase. (i and j) Cu8-700 alloy, the BF image (i) and the corresponding SADP (j) of 14M martensite; (k–n) Cu8-800 alloy, the BF image (k) and the corresponding SADP (l) of 14M martensite; the BF image (m) and the corresponding SADP (n) of fcc γ phase. The red circles mean that the SADP are taken in these areas.
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Table 3 The observed dimension changes during different phase transformations in Fig. 8. The symbols (+) and (–) respectively represent the dimension expansion and contraction during heating and cooling. The standard deviations of the measured values are ± 0.1%. Samples
Cu8-900 Cu10-900 Cu12-900
Dimension change (%) L10→14M
14M↔β(B2)
β(B2)↔γ
+0.9 +0.8 +0.2
+1.2/–0.4 +1.4/–0.3 +0.2/–0.2
–0.5/+0.2 –1.1/+0.4 −0.2/+0.3
Fig. 6. In-situ high-temperature X-ray diffraction patterns of the Cu8-900 alloy during heating.
Fig. 8. The second TMA tests of the same samples after the first TMA tests in Fig. 7.
Fig. 7. Dimension changes during heating and cooling for different alloys quenched from 900 °C characterized by TMA tests. The symbols (+) and (−) respectively represent the dimension expansion and contraction during heating and cooling.
4. Conclusions Fig. 9. TMAs curves of the Cu8-700, Cu10-700 and Cu12-700 samples.
The phase transformation behaviors, and the corresponding microstructural evolutions and dimension changes in Ni50-xCuxMn50 (x=8, 120
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10, 12) alloys were investigated in this study. The results reveal that:
References
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Acknowledgements We acknowledge the financial supports from the Fundamental Research Funds for the Central Universities, grant number 20720160078, the National Natural Science Foundation of China, grant number 51571168.
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