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Grinding-induced martensite stabilization in Mn50Ni33.5Sn8Co8.5 alloy
Materials Letters 107 (2013) 239–242
Contents lists available at SciVerse ScienceDirect
Materials Letters journal homepage: www.elsevier.com/locate/matlet
Grinding-induced martensite stabilization in Mn50Ni33.5Sn8Co8.5 alloy G.T. Li a, Z.H. Liu a,n, X.Q. Ma a, S.Y. Yu b, Y. Liu c a b c
Department of Physics, University of Science and Technology Beijing, Beijing 100083, China School of Physics, Shandong University, Jinan 250100, China School of Mechanical and Chemical Engineering, the University of Western Australia, Crawley, WA 6009, Australia
art ic l e i nf o
a b s t r a c t
Article history: Received 1 April 2013 Accepted 8 June 2013 Available online 14 June 2013
We found that the martensitic transformation in Mn50Ni33.5Sn8Co8.5 alloy is very sensitive to the surface stress produced during grinding. The grinding-induced martensite in a Mn50Ni33.5Sn8Co8.5 alloy is found to be more thermally stable than the thermally-induced martensite, as indicated by a large increase in the reverse transformation temperature and its latent heat. Annealing the stress-induced martensite was found effective in restoring the austenite structure. The stress-induced martensite was also found to exhibit much enlarged coercive force of 3 kOe, compared to 330 Oe for the thermally-induced undeformed martensite. This was attributed to the pinning of domain walls by internal stress. These observations were noted in contrast to that observed in most NiMn-based ferromagnetic martensitic alloys. & 2013 Elsevier B.V. All rights reserved.
Keywords: Phase transformation Shape memory materials Martensite stabilization Magnetic materials
1. Introduction Since the discovery of ferromagnetic shape memory alloy NiMnGa, intensive efforts have been made to develop new alloys with the ability to exhibit magnetic field induced strains [1,2], magnetic field induced transformations [3–5], and magnetically controlled shape memory effect [6]. Most of the alloys developed are evolved from the Heusler structure, a highly ordered intermetallic structure. A main drawback for the success of these alloys is the poor mechanical properties, more specifically their high brittleness [7–9]. As functional materials, shape memory alloys are used in applications often involving repeated transformation or deformation cycles. For successful applications and reliable designs, understanding of the transformation behavior and functional properties of the alloys subjected to various transformation/ deformation are important. For conventional shape memory alloys, e.g., NiTi, NiTiNb, and CuZnAl, it has been reported that deformation results in stabilization of the deformed martensite, manifested as the increase of the reverse transformation temperatures [10,11]. This stabilization effect has been observed after deformation via either martensite reorientation or stress-induced martensite transformation in specimens of both single crystalline and polycrystalline materials [10,12]. This effect is attributed to deformation induced internal plastic deformation and residual stresses. In this paper, we report the observation of grindingstress-induced martensite stabilization in a ferromagnetic shape
memory alloy Mn50Ni33.5Sn8Co8.5. The martensitic transformation and the magnetic behavior of the stress-induced martensite are determined. 2. Materials and methods A polycrystalline Mn50Ni33.5Sn8Co8.5 button ingot was prepared using an arc melting furnace in argon atmosphere from high purity (99.99%) elemental metals. The ingot (∼5.5 g) was heattreated at 1073 K for 24 h in vacuum followed by quenching into water to ensure composition homogeneity. Part of the ingot was crushed into small pieces and ground into powder with particle size less than 100 μm. A powder sample was sealed in the quartz tube under the vacuum condition and annealed at 600 K for 1 h and then, slowly cooled to room temperature. The structures of the powder samples were characterized by means of x-ray diffraction (XRD) using a Rigaku x-ray diffractometer with Cu Kα radiation. Differential scanning calorimetry (DSC) analysis was conducted using a NETZSCH STA449F3 with simultaneous thermal analyzer having a heating/cooling rate of 10 K/min. The ac susceptibility measurement was carried out with an ac field of 796 A/m and a frequency of 77 Hz using an ac susceptometer. Magnetization properties were measured using Quantum Design Physical Properties Measurement System (PPMS-13). 3. Results and discussions
n
Corresponding author. Tel.: +86 10 62332139. E-mail addresses: [email protected], [email protected] (Z.H. Liu).
0167-577X/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.matlet.2013.06.027
Fig. 1 shows XRD patterns of the Mn50Ni33.5Sn8Co8.5 alloy in three conditions. The as-annealed bulk sample (curve (a)) shows a
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Fig. 1. XRD spectra of Mn50Ni33.5Sn8Co8.5 in different states: (a) as-annealed ingot; (b) powder sample in as-ground state; (c) powder sample in as-annealed state; Inset in Fig. 1 shows the temperature dependence of ac susceptibility for as-annealed ingot sample.
pure austenite structure with bcc fundamental lattice reflections of (220), (400) and (422), consistent with similar Mn2NiX (Ga,Sn,and Sb) alloys [13–15]. A superlattice diffraction at (111) is also obvious, due to the ordering of the structure. The lattice parameter is determined to be 0.601 nm. Sample (b) is a powder sample in asground state. The diffraction pattern is indexed to be body-centred tetragonal non-modulated martensite structure. The lattice parameters are determined to be a¼ b¼0.546 nm, and c¼0.697 nm. The incommensurability between the as-ground powder and the as-annealed bulk sample is apparently due to the mechanical effect of the crushing and grinding for powder making. Pattern (c) was from the annealed powder sample, which showed a nearly bcc austenite structure with a little amount of martensite. The inset in Fig. 1 shows the temperature dependence of ac susceptibility of the Mn50Ni33.5Sn8Co8.5 annealed ingot. The sample underwent the martensitic transformation at Ms ¼ 245 K during cooling and the reverse transformation at As ¼255 K during heating. It is obvious that the sample is expected to be in austenite state at room temperature. However, the as-ground powder sample (Curve (b)) showed complete martensite structure, demonstrating the mechanical grinding induced the martensitic transformation of the alloy. This behavior has not been observed in most NiMn-based ferromagnetic martensitic alloys. Fig. 2 shows the transformation behavior of the as-ground powder sample. Fig. 2(a) shows ac susceptibility measurement of the sample. For comparison the transformation behavior of an undeformed bulk specimen is also shown. The bulk sample showed a typical martensitic transformation hysteresis loop between 200 and 300 K and a Curie transition of the austenite at 473 K. The critical temperatures for martensitric transformation are estimated to be Ms ¼245 K, Mf ¼ 182 K, As ¼255 K and Af ¼294 K. The as-ground powder sample showed the reverse transformation at As ¼ 425 K on the first heating, which is 170 K above that of the reverse transformation of the thermally induced martensite of the undeformed sample. The reverse transformation is followed immediately by the Curie transition of the austenite at T AC (470 K). The occurrence of the Curie transition of the austenite prevented
the determination of the finish temperature of the reverse transformation. During the subsequent thermal cycle, the as-ground sample showed martensitic transformation at between 100 and 300 K. It is obvious that the reverse transformation of the asground sample occurred at practically the same temperature as the bulk sample but the forward transformation of the powder sample was much delayed. Consequently, the transformation thermal hysteresis between the forward and the reverse transformations of the powder sample increased significantly. Fig. 2(b) shows the DSC measurement of the transformation behavior of the as-ground powder sample upon the first heating. It is seen that the reverse transformation of the grinding-stressinduced martensite occurred and finished at 425 K and 512 K, respectively. In comparison with the undeformed bulk sample (shown in Fig. 2(c)), the temperature interval of the reverse transformation of the grinding-induced martensite is increased. The heat effect associated with the first reverse transformation is determined to be 12.45 J/g, which is much higher than that of the undeformed-bulk sample, 3.94 J/g, determined by Fig. 2(c). The transformation heat of a thermoelastic martensitic transformation measured by DSC is a combined effect of chemical enthalpy change (latent heat of the transformation), stored elastic strain enthalpy change, and irreversible energy consumed in the form of work [16]. The grinding-stress-induced martensite and the thermal induced martensite have the same crystal structure, so the chemical enthalpy changes of the two samples are same. Therefore the increased heat effect of the deformed sample relative to that of the undeformed sample is attributed to the increase of irreversible energy and stored elastic energy. The irreversible contribution mainly arises from the friction stress requires to move the interface, free energy changes related to defects induced by the transformation (stacking faults, twin boundaries), and internal plastic deformation [16]. For the grinding-stress-induced martensitic transformation, an external deformation is realized by accumulating internal microstrains produced by the formation of oriented martensite variant in each grain, thus an internal plastic deformation is expected to occur. This indicates a coordinating plastic deformation which is accompanied with the oriented martensite in a polycrystalline matrix, resulting in the increase of the transformation heat and the enlargement of the transformation temperature interval during the first reverse transformation. During the second thermal cycling, the small increase of the thermal hysteresis, as seen in Fig. 2(a) can be attributed to the irreversible defects and plastic deformation. The increase in the characteristic temperature for the reverse transformation in the as-ground sample indicates a stabilization effect of the grinding-stress-induced martensite. The internal plastic deformation is believed to be responsible for the stabilization effect. The plastic deformation introduces internal stress into the sample, which creates resistance to the reverse transformation of the oriented martensite [17]. Upon the reverse transformation, an internal mechanical work is performed during the process of reverse transformation of the oriented martensite, with the interior of the grains performing the work against the resistance of the grain boundaries. This increases the need for the driving force, causing the stabilization effect. Fig. 3 shows the magnetization hysteresis cycles of two samples at 5 K. Sample (a) was a thermally-induced martensite in an undeformed-bulk sample. Sample (b) was grinding-stressinduced martensite powder (the sample is the same as Fig. 1 (b)). The loops were measured in a magnetization sequence of 50 kOe-−50 kOe-50 kOe. At above | 713| kOe, the loops were highly reversible; thus only the central part of the hysteresis cycles within 7 13 kOe is shown. It is seen that the thermally-induced martensite exhibited typical soft magnetization behavior, with a coercive force (Hc) of 330 Oe. The stress-induced martensite
G.T. Li et al. / Materials Letters 107 (2013) 239–242
241
Fig. 2. (a) Temperature dependence of ac susceptibility for as-annealed ingot sample and as-ground powder sample; (b) DSC curve of as-ground powder sample during heating from room temperature; and (c) DSC curve of the as-annealed ingot sample.
for the reverse transformation was observed in the grindingstress-induced martensite due to the plastic deformation. The endothermic heat of the reverse transformation of the stressinduced martensite was increased relative to that of the undeformed bulk sample. The stress-induced martensite also exhibited magnetic hardening, with enlarged coercive force 3 kOe relative to 330 Oe of the thermally-induced martensite.
Acknowledgments
Fig. 3. Magnetization hysteresis curves measured at 5 K for thermally-induced martensite and stress-induced martensite (as-ground powder sample).
showed a much widened hysteresis, with a coercive force of 3 kOe. The increase of the coercive field is attributed to pinning effect of domain wall caused by the plastic deformation.
4. Conclusions In summary, in this paper we reported the observation of grinding-stress-induced martensite stabilization effect in Mn50Ni33.5Sn8Co8.5. The experimental evidences demonstrated that the martensitic transformation of this alloy was very sensitive to the stress. A large increase of the transformation temperature
This project is supported by the National Natural Science Foundation of China in Grant nos. 51001010 and 50901043, by Research Fund for the Doctoral Program of Higher Education of China (Grant no. 20100006120001), the Fundamental Research Funds for the Central Universities (Grant no. FRF-BR-10-008A), and by the Scientific Research Foundation for the Returned Overseas Chinese Scholars, State Education Ministry.
References [1] Ullakko K, Huang JK, Kantner C, O’Handley RC, Kokorin VV. Appl Phys Lett 1996;69:1966–8. [2] Wu GH, Yu CH, Meng LQ, Chen JL, Yang FM, Qi SR, et al. Appl Phys Lett 1999;75:2990–2. [3] Kainuma R, Imano Y, Ito W, Sutou Y, Morito H, Okamoto S, et al. Nature 2006;439:957–60. [4] Kainuma R, Imano Y, Ito W, Morito H, Sutou Y, Oikawa K, et al. Appl Phys Lett 2006;88:192513. [5] Yu SY, Liu ZH, Liu GD, Chen JL, Cao ZX, Wu GH, et al. Appl Phys Lett 2006;89:162503.
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G.T. Li et al. / Materials Letters 107 (2013) 239–242
[6] Wang WH, Wu GH, Chen JL, Gao SX, Zhan WS, Wang Z, et al. Appl Phys Lett 2000;77:3245–7. [7] Chernenko VA, L’vov V, Pons J, Cesari E. J Appl. Phys 2003;93:2394–9. [8] Straka L, Heczko O. IEEE Trans Magn 2003;39:3402–4. [9] Sutou Y, Kamiya N, Omori T, Kainuma R, Ishida K, Oikawa K. Appl Phys Lett 2004;84:1275–7. [10] Liu Y, Favier D. Acta Mater 2000;48:3489–99. [11] Piao M, Otsuka K, Miyazaki S, Horikawa H. Mater Trans JIM 1993;34:919–29.
[12] Lin HC, Wu SK. Metall Trans A 1993;24:293–9. [13] Liu GD, Dai XF, Yu SY, Zhu ZY, Chen JL, Wu GH, et al. Phys Rev B 2006;74:054435. [14] Luo HZ, Liu GD, Feng ZQ, Li YX, Ma L, Wu GH, et al. J Magn Magn Mater 2009;321:4063–6. [15] Helmholdt RB, Buschow KHJ. J Less-Common Met 1987;128:167–71. [16] Ortin J, Planes A. Acta Metall 1988;36:1873–89. [17] Liu Y. Mater Sci Eng A 2004;378:459–64.
Contents lists available at SciVerse ScienceDirect
Materials Letters journal homepage: www.elsevier.com/locate/matlet
Grinding-induced martensite stabilization in Mn50Ni33.5Sn8Co8.5 alloy G.T. Li a, Z.H. Liu a,n, X.Q. Ma a, S.Y. Yu b, Y. Liu c a b c
Department of Physics, University of Science and Technology Beijing, Beijing 100083, China School of Physics, Shandong University, Jinan 250100, China School of Mechanical and Chemical Engineering, the University of Western Australia, Crawley, WA 6009, Australia
art ic l e i nf o
a b s t r a c t
Article history: Received 1 April 2013 Accepted 8 June 2013 Available online 14 June 2013
We found that the martensitic transformation in Mn50Ni33.5Sn8Co8.5 alloy is very sensitive to the surface stress produced during grinding. The grinding-induced martensite in a Mn50Ni33.5Sn8Co8.5 alloy is found to be more thermally stable than the thermally-induced martensite, as indicated by a large increase in the reverse transformation temperature and its latent heat. Annealing the stress-induced martensite was found effective in restoring the austenite structure. The stress-induced martensite was also found to exhibit much enlarged coercive force of 3 kOe, compared to 330 Oe for the thermally-induced undeformed martensite. This was attributed to the pinning of domain walls by internal stress. These observations were noted in contrast to that observed in most NiMn-based ferromagnetic martensitic alloys. & 2013 Elsevier B.V. All rights reserved.
Keywords: Phase transformation Shape memory materials Martensite stabilization Magnetic materials
1. Introduction Since the discovery of ferromagnetic shape memory alloy NiMnGa, intensive efforts have been made to develop new alloys with the ability to exhibit magnetic field induced strains [1,2], magnetic field induced transformations [3–5], and magnetically controlled shape memory effect [6]. Most of the alloys developed are evolved from the Heusler structure, a highly ordered intermetallic structure. A main drawback for the success of these alloys is the poor mechanical properties, more specifically their high brittleness [7–9]. As functional materials, shape memory alloys are used in applications often involving repeated transformation or deformation cycles. For successful applications and reliable designs, understanding of the transformation behavior and functional properties of the alloys subjected to various transformation/ deformation are important. For conventional shape memory alloys, e.g., NiTi, NiTiNb, and CuZnAl, it has been reported that deformation results in stabilization of the deformed martensite, manifested as the increase of the reverse transformation temperatures [10,11]. This stabilization effect has been observed after deformation via either martensite reorientation or stress-induced martensite transformation in specimens of both single crystalline and polycrystalline materials [10,12]. This effect is attributed to deformation induced internal plastic deformation and residual stresses. In this paper, we report the observation of grindingstress-induced martensite stabilization in a ferromagnetic shape
memory alloy Mn50Ni33.5Sn8Co8.5. The martensitic transformation and the magnetic behavior of the stress-induced martensite are determined. 2. Materials and methods A polycrystalline Mn50Ni33.5Sn8Co8.5 button ingot was prepared using an arc melting furnace in argon atmosphere from high purity (99.99%) elemental metals. The ingot (∼5.5 g) was heattreated at 1073 K for 24 h in vacuum followed by quenching into water to ensure composition homogeneity. Part of the ingot was crushed into small pieces and ground into powder with particle size less than 100 μm. A powder sample was sealed in the quartz tube under the vacuum condition and annealed at 600 K for 1 h and then, slowly cooled to room temperature. The structures of the powder samples were characterized by means of x-ray diffraction (XRD) using a Rigaku x-ray diffractometer with Cu Kα radiation. Differential scanning calorimetry (DSC) analysis was conducted using a NETZSCH STA449F3 with simultaneous thermal analyzer having a heating/cooling rate of 10 K/min. The ac susceptibility measurement was carried out with an ac field of 796 A/m and a frequency of 77 Hz using an ac susceptometer. Magnetization properties were measured using Quantum Design Physical Properties Measurement System (PPMS-13). 3. Results and discussions
n
Corresponding author. Tel.: +86 10 62332139. E-mail addresses: [email protected], [email protected] (Z.H. Liu).
0167-577X/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.matlet.2013.06.027
Fig. 1 shows XRD patterns of the Mn50Ni33.5Sn8Co8.5 alloy in three conditions. The as-annealed bulk sample (curve (a)) shows a
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Fig. 1. XRD spectra of Mn50Ni33.5Sn8Co8.5 in different states: (a) as-annealed ingot; (b) powder sample in as-ground state; (c) powder sample in as-annealed state; Inset in Fig. 1 shows the temperature dependence of ac susceptibility for as-annealed ingot sample.
pure austenite structure with bcc fundamental lattice reflections of (220), (400) and (422), consistent with similar Mn2NiX (Ga,Sn,and Sb) alloys [13–15]. A superlattice diffraction at (111) is also obvious, due to the ordering of the structure. The lattice parameter is determined to be 0.601 nm. Sample (b) is a powder sample in asground state. The diffraction pattern is indexed to be body-centred tetragonal non-modulated martensite structure. The lattice parameters are determined to be a¼ b¼0.546 nm, and c¼0.697 nm. The incommensurability between the as-ground powder and the as-annealed bulk sample is apparently due to the mechanical effect of the crushing and grinding for powder making. Pattern (c) was from the annealed powder sample, which showed a nearly bcc austenite structure with a little amount of martensite. The inset in Fig. 1 shows the temperature dependence of ac susceptibility of the Mn50Ni33.5Sn8Co8.5 annealed ingot. The sample underwent the martensitic transformation at Ms ¼ 245 K during cooling and the reverse transformation at As ¼255 K during heating. It is obvious that the sample is expected to be in austenite state at room temperature. However, the as-ground powder sample (Curve (b)) showed complete martensite structure, demonstrating the mechanical grinding induced the martensitic transformation of the alloy. This behavior has not been observed in most NiMn-based ferromagnetic martensitic alloys. Fig. 2 shows the transformation behavior of the as-ground powder sample. Fig. 2(a) shows ac susceptibility measurement of the sample. For comparison the transformation behavior of an undeformed bulk specimen is also shown. The bulk sample showed a typical martensitic transformation hysteresis loop between 200 and 300 K and a Curie transition of the austenite at 473 K. The critical temperatures for martensitric transformation are estimated to be Ms ¼245 K, Mf ¼ 182 K, As ¼255 K and Af ¼294 K. The as-ground powder sample showed the reverse transformation at As ¼ 425 K on the first heating, which is 170 K above that of the reverse transformation of the thermally induced martensite of the undeformed sample. The reverse transformation is followed immediately by the Curie transition of the austenite at T AC (470 K). The occurrence of the Curie transition of the austenite prevented
the determination of the finish temperature of the reverse transformation. During the subsequent thermal cycle, the as-ground sample showed martensitic transformation at between 100 and 300 K. It is obvious that the reverse transformation of the asground sample occurred at practically the same temperature as the bulk sample but the forward transformation of the powder sample was much delayed. Consequently, the transformation thermal hysteresis between the forward and the reverse transformations of the powder sample increased significantly. Fig. 2(b) shows the DSC measurement of the transformation behavior of the as-ground powder sample upon the first heating. It is seen that the reverse transformation of the grinding-stressinduced martensite occurred and finished at 425 K and 512 K, respectively. In comparison with the undeformed bulk sample (shown in Fig. 2(c)), the temperature interval of the reverse transformation of the grinding-induced martensite is increased. The heat effect associated with the first reverse transformation is determined to be 12.45 J/g, which is much higher than that of the undeformed-bulk sample, 3.94 J/g, determined by Fig. 2(c). The transformation heat of a thermoelastic martensitic transformation measured by DSC is a combined effect of chemical enthalpy change (latent heat of the transformation), stored elastic strain enthalpy change, and irreversible energy consumed in the form of work [16]. The grinding-stress-induced martensite and the thermal induced martensite have the same crystal structure, so the chemical enthalpy changes of the two samples are same. Therefore the increased heat effect of the deformed sample relative to that of the undeformed sample is attributed to the increase of irreversible energy and stored elastic energy. The irreversible contribution mainly arises from the friction stress requires to move the interface, free energy changes related to defects induced by the transformation (stacking faults, twin boundaries), and internal plastic deformation [16]. For the grinding-stress-induced martensitic transformation, an external deformation is realized by accumulating internal microstrains produced by the formation of oriented martensite variant in each grain, thus an internal plastic deformation is expected to occur. This indicates a coordinating plastic deformation which is accompanied with the oriented martensite in a polycrystalline matrix, resulting in the increase of the transformation heat and the enlargement of the transformation temperature interval during the first reverse transformation. During the second thermal cycling, the small increase of the thermal hysteresis, as seen in Fig. 2(a) can be attributed to the irreversible defects and plastic deformation. The increase in the characteristic temperature for the reverse transformation in the as-ground sample indicates a stabilization effect of the grinding-stress-induced martensite. The internal plastic deformation is believed to be responsible for the stabilization effect. The plastic deformation introduces internal stress into the sample, which creates resistance to the reverse transformation of the oriented martensite [17]. Upon the reverse transformation, an internal mechanical work is performed during the process of reverse transformation of the oriented martensite, with the interior of the grains performing the work against the resistance of the grain boundaries. This increases the need for the driving force, causing the stabilization effect. Fig. 3 shows the magnetization hysteresis cycles of two samples at 5 K. Sample (a) was a thermally-induced martensite in an undeformed-bulk sample. Sample (b) was grinding-stressinduced martensite powder (the sample is the same as Fig. 1 (b)). The loops were measured in a magnetization sequence of 50 kOe-−50 kOe-50 kOe. At above | 713| kOe, the loops were highly reversible; thus only the central part of the hysteresis cycles within 7 13 kOe is shown. It is seen that the thermally-induced martensite exhibited typical soft magnetization behavior, with a coercive force (Hc) of 330 Oe. The stress-induced martensite
G.T. Li et al. / Materials Letters 107 (2013) 239–242
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Fig. 2. (a) Temperature dependence of ac susceptibility for as-annealed ingot sample and as-ground powder sample; (b) DSC curve of as-ground powder sample during heating from room temperature; and (c) DSC curve of the as-annealed ingot sample.
for the reverse transformation was observed in the grindingstress-induced martensite due to the plastic deformation. The endothermic heat of the reverse transformation of the stressinduced martensite was increased relative to that of the undeformed bulk sample. The stress-induced martensite also exhibited magnetic hardening, with enlarged coercive force 3 kOe relative to 330 Oe of the thermally-induced martensite.
Acknowledgments
Fig. 3. Magnetization hysteresis curves measured at 5 K for thermally-induced martensite and stress-induced martensite (as-ground powder sample).
showed a much widened hysteresis, with a coercive force of 3 kOe. The increase of the coercive field is attributed to pinning effect of domain wall caused by the plastic deformation.
4. Conclusions In summary, in this paper we reported the observation of grinding-stress-induced martensite stabilization effect in Mn50Ni33.5Sn8Co8.5. The experimental evidences demonstrated that the martensitic transformation of this alloy was very sensitive to the stress. A large increase of the transformation temperature
This project is supported by the National Natural Science Foundation of China in Grant nos. 51001010 and 50901043, by Research Fund for the Doctoral Program of Higher Education of China (Grant no. 20100006120001), the Fundamental Research Funds for the Central Universities (Grant no. FRF-BR-10-008A), and by the Scientific Research Foundation for the Returned Overseas Chinese Scholars, State Education Ministry.
References [1] Ullakko K, Huang JK, Kantner C, O’Handley RC, Kokorin VV. Appl Phys Lett 1996;69:1966–8. [2] Wu GH, Yu CH, Meng LQ, Chen JL, Yang FM, Qi SR, et al. Appl Phys Lett 1999;75:2990–2. [3] Kainuma R, Imano Y, Ito W, Sutou Y, Morito H, Okamoto S, et al. Nature 2006;439:957–60. [4] Kainuma R, Imano Y, Ito W, Morito H, Sutou Y, Oikawa K, et al. Appl Phys Lett 2006;88:192513. [5] Yu SY, Liu ZH, Liu GD, Chen JL, Cao ZX, Wu GH, et al. Appl Phys Lett 2006;89:162503.
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G.T. Li et al. / Materials Letters 107 (2013) 239–242
[6] Wang WH, Wu GH, Chen JL, Gao SX, Zhan WS, Wang Z, et al. Appl Phys Lett 2000;77:3245–7. [7] Chernenko VA, L’vov V, Pons J, Cesari E. J Appl. Phys 2003;93:2394–9. [8] Straka L, Heczko O. IEEE Trans Magn 2003;39:3402–4. [9] Sutou Y, Kamiya N, Omori T, Kainuma R, Ishida K, Oikawa K. Appl Phys Lett 2004;84:1275–7. [10] Liu Y, Favier D. Acta Mater 2000;48:3489–99. [11] Piao M, Otsuka K, Miyazaki S, Horikawa H. Mater Trans JIM 1993;34:919–29.
[12] Lin HC, Wu SK. Metall Trans A 1993;24:293–9. [13] Liu GD, Dai XF, Yu SY, Zhu ZY, Chen JL, Wu GH, et al. Phys Rev B 2006;74:054435. [14] Luo HZ, Liu GD, Feng ZQ, Li YX, Ma L, Wu GH, et al. J Magn Magn Mater 2009;321:4063–6. [15] Helmholdt RB, Buschow KHJ. J Less-Common Met 1987;128:167–71. [16] Ortin J, Planes A. Acta Metall 1988;36:1873–89. [17] Liu Y. Mater Sci Eng A 2004;378:459–64.