A large barocaloric effect associated with paramagnetic martensitic transformation in Co50Fe2.5V31.5Ga16 quaternary Heusler alloy

Scripta Materialia 177 (2020) 1–5

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A large barocaloric effect associated with paramagnetic martensitic transformation in Co50 Fe2.5 V31.5 Ga16 quaternary Heusler alloy Hongwei Liu a,b, Zhe Li b,∗, Yuanlei Zhang b, Zhenting Ni a,b, Kun Xu b, Yongsheng Liu a,∗ a

Department of Physics, Shanghai University of Electric Power, Shanghai 200090, China Center for Magnetic Materials and Devices and Key Laboratory for Advanced Functional and Low Dimensional Materials of Yunnan Higher Education Institute, Qujing Normal University, Qujing 655011, China

b

a r t i c l e

i n f o

Article history: Received 5 July 2019 Revised 26 September 2019 Accepted 2 October 2019

Keywords: Heusler alloy Martensitic transformation Barocaloric effect

a b s t r a c t A martensitic transformation (MT) with paramagnetic behavior has been developed in Co50 Fe2.5 V31.5 Ga16 quaternary Heusler alloy near room temperature. Accompanying the transformation, this alloy shows a good two-way shape memory effect, which results in a considerable relative volume change between two phases. Because of this, it has been also found that the MT is strongly sensitive to the applied hydrostatic pressure. By utilizing an indirectly estimated method, the obtained maximum reversible isothermal entropy (ST ) and adiabatic temperature change (Tad ) resulted from hydrostatic pressure induced MT, respectively, achieve ∼31 J/kg K and 6 K under 5 kbar, performing a large barocaloric effect. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Barocaloric effect (BCE) generally refers to entropy or temperature changes of solid-state substance that occurs when an external hydrostatic pressure is applied at isothermal or adiabatic conditions. This effect is expected to be used for solid-state refrigeration techniques, which can greatly reduce both wasting energy and environment pollution that seem unpreventable in conventional refrigeration devices based on compression and expansion cycles of gases [1,2]. Up to date, the BCE has been developed in different solids, such as ferroic compounds [3–9], itinerant-electron system [10,11], fluorides [12,13], superionic conductor [14], hybrid perovskite [15] and plastic crystals [16,17]. The first-order structural transformation (FOST) is a commom feature of these materials. As well-known, the Ni-Mn based metamagnetic shape memory alloys (MSMAs) with Heusler structure as a class of representative barocaloric materials have been extensively studied during the past decade [3,18–22]. The mechanism of outstanding BCE presented by this system can be ascribed to the strong coupling of magnetism and structure at first-order martensitic transformation (MT), which brings about a large relative volume change between two phases, thus giving rise to a prominent transition entropy change (Str ). However, previous studies have also proposed that the Str of MSMAs mainly originates from contribution of structure variation in crystallography, whereas the contribution of magnetic degrees of freedom plays a contradictory role [23–25]. Obviously, such a ∗

Corresponding authors. E-mail addresses: [email protected], [email protected] (Z. Li), ysliu@ shiep.edu.cn (Y. Liu). https://doi.org/10.1016/j.scriptamat.2019.10.003 1359-6462/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

competition during MT has put forward a difficult challenge to maximize the Str for their caloric effect. Very recently, by decreasing the contribution of magnetic entropy change to the Str , a giant BCE has been further observed in (Ni50 Mn31.5 Ti18.5 )99.8 B0.2 shape memory Heusler alloy, which shows a toneless MT without long-range ferromagnetic ordering [26]. These findings predict that the thermal response could be optimized in the other Heusler alloys undergoing MT not only with the noticeable relative volume change but also with the negligible change of magnetization. In recent years, a paramagnetic MT from a high-symmetry L21 cubic structure to a low-symmetry D022 tetragonal structure has been reported in a series of non-stoichiometry Co-V-Ga ternary Heusler alloys [27]. Since then, a shape memory effect associated with metamagnetic MT has been also found in an individual composition of Co50 V34 Ga16 [28]. However, it is worth noting that the relative volume change at MT exhibited by these ternary alloys is inconspicuous [27,28] when compared with many FOST materials [2]. This hints that the MT of Co-V-Ga system is not sensitive to external hydrostatic pressure, and thereby possibly leads to a weak BCE. Based on this motivation, in this work, we have used Fe atoms to partially replace V sites in Co50 V34 Ga16 ternary parent and prepared Co50 Fe2.5 V31.5 Ga16 quaternary Heusler alloy. It has been interestingly found that this alloy displays a paramagnetic MT near room temperature with an accompanying considerable relative volume change. Associated with such behaviors, a large BCE has been obtained by applied relatively low hydrostatic pressure. The polycrystalline Co50 Fe2.5 V31.5 Ga16 alloy was fabricated by arc-melting technique, and its weight loss after melting is less than 0.3%. The obtained ingot was annealed in an evacuated quartz tube

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Fig. 1. (a) X-ray diffraction pattern at room temperature. (b) The temperature dependence of magnetization during cooling and heating measured under a magnetic field of 0.5 kOe. (c) The reciprocal susceptibility as function of temperature deduced from M(T) data of high-temperature region (open squares), the red solid line is a linear fit to the data based on Curie-Weiss law. Inset shows the isothermal magnetization curves measured at selected temperatures. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

for 24 h at 1473 K, and then quenched in ice water. The actual composition was determined by energy dispersive spectroscopy that corresponds to Co49.79 Fe2.52 V31.12 Ga16.57 . The crystal structure was identified by powder X-ray diffraction at room temperature. The magnetization data were acquired through vibrating sample magnetometer by using the stable temperature mode. Heat flow data were collected from commercial differential scanning calorimeter (μDSC7 EVO) at applied constant hydrostatic pressures (0, 0.2, 0.4, 0.6, 0.8 and 1 kbar), which were generated by high-pressure nitrogen gas. The scanning rate was 1 K/min. The transition strain were performed by standard strain-gauge technique under various hydrostatic pressures. For this measurement, a piece of rectangular specimen with dimensions 1 × 2 × 3 mm3 was dipped in a Teflon capsule filled with Daphane oil and placed inside the piston-type pressure cell. The heating and cooling rates were 1 K/min. Fig. 1(a) shows the X-ray diffraction pattern of Co50 Fe2.5 V31.5 Ga16 alloy at room temperature. From the refinement results, we can find that it possesses a L21 cubic structure at austenitic phase with the lattice parameter of ∼0.5779 nm. Fig. 1(b) illustrates the temperature dependence of magnetization M(T) of this alloy. Both a slight change of magnetization and an obvious thermal hysteresis can be clearly distinguished near room temperature, indicating that it undergoes the first-order direct and reverse MTs during cooling and heating. Here, the estimated characteristic temperatures of these two transformations, i.e., Ms , Mf , As and Af , equal to ∼276 K, ∼267 K, ∼277 K and ∼288 K, where both Ms and Af denote the direct MT starting temperature and the reverse MT finishing temperature, respectively. By using the reciprocal susceptibility as function of temperature based on Curie-Weiss law, as shown in Fig. 1(c), a paramagnetic Curie temperature of austenite (TCA ) is roughly evaluated to be ∼122 K away from the Mf . This means that the occurrence of MT for this alloy is not accompanied by rearrangement of magnetic moments, which consequences in a fact that the magnetization presents an approximatively linear variation outside the transformation as the temperature is changed [see Fig. 1(b)]. Moreover, the isothermal

magnetic curves at selected temperatures are presented in the inset of Fig. 1(c). As observed the curve measured at 300 K, a linear increase of magnetization with increasing magnetic field further manifests a paramagnetic behavior of austenite. When the temperature is decreased to 250 K, the magnetization of martensite also reveals a same feature, except for a feeble spin fluctuation induced by lattice distortion. Therefore, all these observations sufficiently demonstrated that such a MT manifested by present alloy has a paramagnetic behavior. Fig. 2(a) shows the temperature dependence of heat capacity dQ/dT at different hydrostatic pressures for Co50 Fe2.5 V31.5 Ga16 alloy after baseline subtraction, which were deduced from heat flow data by using universal thermodynamical equation. As shown in this figure, two well-defined peaks induced by direct and reverse MTs can be observed on the cooling and heating curves, which shifts to high temperatures with increasing hydrostatic pressure. It can be explained by the fact that application of pressure tends to stabilize martensite that possesses a relatively lower volume as compared to austenite. According to heat capacity data, the relative entropy at a given temperature, S (T , p) = S(T , p) − S(Tre f , p), can be derived from simultaneous integral equations as proposed in Ref. [2,9,26] assuming the dQ/dTP to be independent of pressure. The Tref represents the reference temperature corresponding to a value of 240 K in here. Following this way, the S (T, p) curves upon cooling and heating are plotted in Fig. 2(b) and (c). From the curves recorded at 0 kbar, the obtained values of Str (indicated by double arrow) from direct and reverse MTs are approximatively equivalent and determined to be ∼34 J/kg K. This value is comparable to those of some famous FOST magnetocaloric materials [29–31]. From these S (T, p) curves, we also summarize the evolution of Str with hydrostatic pressure, as evidenced in Fig. 2(d), which shows a monotonic decrease in the value of Str with increase of hydrostatic pressure. Similar behavior has been also reported in other barocaloric alloys [9,26]. By linear fittings, the estimated slops dStr /dp for direct and reverse MTs equal to ∼−2.5 J/kg K and ∼−1.7 J/kg K in the unit per kilobar, respectively.

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Fig. 2. (a) The heat capacity dQ/dT after baseline subtraction were recorded during cooling and heating at hydrostatic pressure range from 0 kbar to 1 kbar, which is deduced from heat flow data by using traditional thermodynamical formula. (b) and (c) The relative entropy as function of temperature upon cooling and heating under same hydrostatic pressures, which were calculated by using the simultaneous integral equations as described in Ref. [2,9,26], referred to the absolute entropy at 240 K. (d) The hydrostatic pressure dependence of transition entropy change (Str ) for direct and reverse MT. The red solid lines are the linear fits to the data. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

Fig. 3. (a): The temperature dependence of strain during cooling and heating at atmospheric pressure. (b): The values of strain were measured under applied hydrostatic pressures from 1 kbar to 5 kbar for direct MT. (c): The hydrostatic pressure dependence of transformation equilibrium temperatures (T0 ) upon cooling and heating, which were determined by the peak values from differential of strain and calorimetric data (green solid cycles). The red solid line is the best linear fit. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

This phenomenon stems from the contraction in volume of cubic phase with hydrostatic pressure, which reduces the difference in volume change between parent and martensitic states. [6,9,17,26]. Next, the effect of hydrostatic pressure on MT for Co50 Fe2.5 V31.5 Ga16 is also studied by means of strain measurement. Fig. 3(a) shows the temperature dependence of strain at

ambient pressure. During cooling, the strain of this sample rapidly decreases by 0.26% (indicated by black double arrow) in metrical direction caused by direct MT, implying its volume shrinkage. Upon heating, the shape deformation of the sample is completely recovered by an expansion with the same strain through the reverse MT, performing a good two-way shape memory effect

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Fig. 4. (a): The S (T, p) curves for the direct MT at 0 kbar and 5 kbar, which were deduced by Eq. (2) based on normalized strain curves. (b): The isothermal entropy change (ST ) as a function of temperature during cooling and heating under various p. The orange cycles refer to the ST determined by quasi-direct method based on calorimetric data under  p = 1 kbar. (c): The T(S , p) curves for the direct MT at 0 kbar and 5 kbar, which were deduced by Eq. (3). (d): The adiabatic temperature change (Tad ) as a function of temperature at different p. The orange cycles stand for the Tad determined by quasi-direct method based on calorimetric data under  p = 1 kbar. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

induced thermally. Such an effect is superior to that obtained in some typical shape memory alloys without mechanical training [32,33]. Fig. 3(b) further gives the strain data measured within the hydrostatic pressure from 1 kbar to 5 kbar. When compared with the strain curve observed at atmospheric pressure, except for movement of transition temperatures, we can also explored that the transition strain is enhanced with increase of hydrostatic pressure. This behavior similarities with that reported in Ni-Mn-Ga [34] and Ni-Co-Mn-Sb [35] polycrystals triggered by the cyclic thermo-mechanical training. It can be explained by the fact that the introduction of external physical pressure can drive the martensite variants to a preferential orientation along metrical direction, when they nucleate at the transition temperature. The relationship between transformation equilibrium temperature (T0 ) and hydrostatic pressure is illustrated by Fig. 3(c). It shows that T0 reveals a linear increment within the measured range, and the migration rate with respect to pressure (dT0 /dp) reach ∼2.62 K/kbar and 2.45 K/kbar for direct and reverse MTs. Additionally, the relative volume change (ɛ) between two phases at ambient pressure is also examined by using Clausius-Clapeyron (C–C) equation:

  |dT0 /dp| = εVmA /Str 

(1)

where the VmA refers to molar volume of austenite with the value of ∼7.26 cm3 /mol as deduced from the lattice parameters. By substituting related values into Eq. (2), the ɛ is determined to be ∼0.71% and ∼0.66% for direct and reverse MTs, which are prominently larger than that reported in some MSMAs [3,19,20,21]. This implies that the element substitution is quite beneficial to enlarge the volume change during MT for the Co-V-Ga system. Taking these discussions into account, a large BCE associated with hydrostatic pressure-induced MT would be a desirable attainment. In order to confirm this point as mentioned above, we then investigate the BCE of Co50 Fe2.5 V31.5 Ga16 by means of an indirectly

estimated method based on change of phase fraction with hydrostatic pressures. More details related to this method can refer to our previous works [20–22]. Considering that the influence of pressure on transforming slope and width for present alloy is weak [Fig 3(b)], the temperature dependence of phase fraction at various hydrostatic pressures f(T, p) in the transforming range can be depicted by the normalized strain curves. Moreover, as a result of the fact that the values of Str are strongly dependent on magnitude of hydrostatic pressure, the isothermal entropy change as a function of temperature can be gained by the following relation:

    ST (T ,  p) = f T , p f · Str p f − f (T , pi ) · Str ( pi ),

(2)

where the p is a difference between pi and pf which are defined as the initial and final hydrostatic pressures, respectively. Here, it is essential to point out that the values of Str (p) above 1 kbar are deduced by linearly extrapolating the data of Fig. 2(d). The computational procedure is described by Fig. 4(a) in detail. Accordingly, the ST (T, p) curves during cooling and heating are sketched in Fig. 4(b). For the  p = 1 kbar, we can see that the ST (T, p) curves for direct and reverse MTs reveal a asymmetrically sharp peak in a narrow temperature range due to appearance of conventional BCE of austenite aroused by its volume shrinkage, which is in nice agreement with those determined from the quasi-direct calorimetric method (see orange open cycles). With increase of hydrostatic pressure, the height and width of these ST (T, p) curves are developed simultaneously, which also strengthens the asymmetry on the opposite side of the peaks. As the p increases to 5 kbar, it is clear that a flat plateau can be observed in the ST (T, p) curves during cooling and heating, which means that the ST approaches a saturated value. By evaluating, the obtained maximum values at  p = 5 kbar equal to ∼33.7 J/kg K and 32.5 J/kg K for direct and reverse MTs, respectively. Such a small discrepancy leads to a relatively broad reversibility region from 275 K to 290 K

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for ST (indicated by shade area), and its largest value appears at 285 that is estimated to be ∼31 J/kg K. This reversible ST surpasses those reported in a great deal of giant barocaloric materials including MSMAs [3,18–22], itinerant-electron system [10,11], as well as some ferroelectric [8,36] and negative thermal expansion compounds under applying a similar value of p [4]. The other key parameter for judging BCE of a given material is the Tad . On the basis of thermodynamic theory, the temperature dependence of entropy under different pressures can be expressed by below formula:

S (T , p) = S (T , p = 0) + ST (T ,  p).

(3)

By utilizing above formula, two selected T(S , p) curves at 0 kbar and 5 kbar are established in the inset of Fig. 4(c). From these curves, the Tad as a function of temperature can be simply deduced by following expression:

     Tad (T ,  p) = T S , p − T S , p = 0 S .

(4)

Fig. 4(d) shows the Tad as a function of temperature for change of p from 0 kbar to 5 kbar. Apparently, our plots of Tad (T, p) curves obtained during cooling and heating exhibit the similar behaviors as that observed in Fig. 4(b). When the p is increased to 5 kbar, the peak value of Tad achieves 10.7 K and 9.2 K for direct and reverse MTs, which are substantially higher than that reported in a lot of MSMAs [3,19–22] and can compete with non-ferromagnetic (Ni50 Mn31.5 Ti18.5 )99.8 B0.2 at similar condition [26]. However, it is remarkable that a flat plateau appeared in ST (T, p) data [Fig. 4(b)] cannot be found in these Tad (T, p) curves even applied our largest p. This induces a reduction of reversibility region (indicated by shade area), and the obtained maximum reversible Tad within this region amounts to 6 K at ∼283 K. In addition, because the relative volume change at MT in this studied alloy is still noticeably lower than that gained in some stateof-the-art barocaloric alloys [5,6,37], the optimization of the ratio between ɛ and Str by tailoring composition should take into consideration to further enhance the BCE for Co-V-Ga system. In summary, the paramagnetic MT associated with a considerable volume change has been developed in Co50 Fe2.5 V31.5 Ga16 quaternary shape memory alloy near room temperature. Due to lack of negative contribution from magnetic entropy change during transformation, a large BCE associated with hydrostatic pressure induced MT has been successfully realized in the studied alloy. The obtained barocaloric parameters overmatches those reported in some other barocaloric materials with FOST, which indicates that the Co-V-Ga system appears to be a potential candidate for solid-state refrigeration techniques. Acknowledgments This work was supported by the National Natural Science Foundation of China (nos. 51661029, 11674215, 51861032, 51971128). Applied Basic Research key project of Yunnan (no.2018FA031) and Project of Science and Technology Commission of Shanghai Municipality (no.190205010 0 0). Program for Innovative Research Team in Science and Technology in university of Yunnan Province (IRTSTYN). We also express our thanks to Prof. Bing Li and Dr. Zhao Zhang of Institute of Metal Research, Chinese Academy of Sciences, for testing the sample using a μDSC7 EVO and some useful discussions. Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2019.10. 003.

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References [1] X. Moya, S. Kar-Narayan, N.D. Mathur, Nat. Mater. 13 (2014) 439–450. [2] L. Mañosa, A. Planes, Adv. Mater. 29 (2017) 1603607. [3] L. Mañosa, D. González-Alonso, A. Planes, E. Bonnot, M. Barrio, J.L. Tamarit, S. Aksoy, M. Acet, Nat. Mater. 9 (2010) 478–481. [4] D. Matsunami, A. Fujita, K. Takenaka, M. Kano, Nat. Mater. 14 (2014) 73–78. [5] R.R. Wu, L.F. Bao, F.X. Hu, H. Wu, Q.Z. Huang, J. Wang, X.L. Dong, G.N. Li, J.R. Sun, F.R. Shen, T.Y. Zhao, X.Q. Zhen, L.C. Wang, Y. Liu, W.L. Zuo, Y.Y. Zhao, M. Zhang, X.C. Wang, C.Q. Jin, G.H. Rao, X.F. Han, B.G. Shen, Sci. Rep. 5 (2015) 18027. [6] P. Lloveras, E. Stern-Taulats, M. Barrio, J.L. Tamarit, S. Crossley, W. Li, V. Pomjakushin, A. Planes, L. Mañosa, N.D. Mathur, X. Moya, Nat. Commun. 6 (2015) 8801. [7] E. Stern-Taulats, A. Gràcia-Condal, A. Planes, P. Lloveras, M. Barrio, J.L. Tamarit, S. Pramanick, S. Majumdar, L. Mañosa, Appl. Phys. Lett. 107 (2015) 152409. [8] E. Stern-Taulats, P. Lloveras, M. Barrio, E. Defay, M. Egilmez, A. Planes, J.L. Tamarit, L. Mañosa, N.D. Mathur, X. Moya, APL Mater. 4 (2016) 091102. [9] A. Aznar, P. Lloveras, J.-Y. Kim, E. Stern-Taulats, M. Barrio, J.L. Tamarit, C.F. Sánchez-Valdé, J.L. Sánchez-Llamazares, N.D. Mathur, X. Moya, Adv. Mater. 31 (2019) 1903577. [10] L. Mañosa, D. González-Alonso, A. Planes, M. Barrio, J.L. Tamarit, I.S. Titov, M. Acet, A. Bhattacharyya, S. Majumdar, Nat. Commun. 2 (2011) 595. [11] S. Yuce, M. Barrio, B. Emre, E. Stern-Taulats, A. Planes, J.L. Tamarit, Y. Mudryk, K.A. Gschneidner Jr., V.K. Pecharsky, L. Mañosa, Appl. Phys. Lett. 101 (2012) 071906. [12] I.N. Flerov, A.V. Kartashev, M.V. Gorev, E.V. Bogdanov, S.V. Mel’nikova, M.S. Molokeev, E.I. Pogoreltsev, N.M. Laptash, J. Fluor. Chem. 183 (2016) 1–9. [13] M.V. Gorev, E.V. Bogdanov, I.N. Flerov, Scr. Mater. 139 (2017) 53–57. [14] A. Aznar, P. Lloveras, M. Romanini, M. Barrio, J.L. Tamarit, C. Cazorla, D. Errandonea, N.D. Mathur, A. Planes, X. Moya, L. Mañosa, Nat. Commun. 8 (2017) 1851. [15] J.M. Bermúdez-García, M. Sánchez-Andújar, S. Castro-García, J. López-Beceiro, R. Artiaga, M.A. Señarís-Rodríguez, Nat. Commun. 8 (2017) 15715. [16] B. Li, Y. Kawakita, S. Ohira-Kawamura, T. Sugahara, H. Wang, J.F. Wang, Y.N. Chen, S. Kawaguchi, S. Kawaguchi, K. Ohara, K. Li, D.H. Yu, R. Mole, T. Hattori, T. Kikuchi, S. Yano, Z. Zhang Z. Zhang, W.J. Ren, S.C. Lin, O. Sakata, K. Nakajima, Z.D. Zhang, Nature 567 (2019) 506–510. [17] P. Lloveras, A. Aznar, M. Barrio, Ph. Negrier, C. Popescu, A. Planes, L. Mañosa, E. Stern- Taulats, A. Avramenko, N.D. Mathur, X. Moya, J.-Ll Tamarit, Nat. Commun. 10 (2019) 1803. [18] L. Mañosa, E. Stern-Taulats, A. Planes, P. Lloveras, M. Barrio, J.L. Tamarit, B. Emre, S. Yüce, S. Fabbrici, F. Albertini, Phys. Status Solidi B 251 (2014) 2114–2119. [19] E. Stern-Taulats, A. Planes, P. Lloveras, M. Barrio, J.L. Tamarit, S. Pramanick, S. Majumdar, S. Yüce, B. Emre, C. Frontera, L. Mañosa, Acta Mater. 96 (2015) 324–332. [20] X.J. He, K. Xu, S.X. Wei, Y.L. Zhang, Z. Li, C. Jing, J. Mater. Sci. 52 (2017) 2915–2923. [21] X.J. He, S.X. Wei, Y.R. Kang, Y.L. Zhang, Y.M. Cao, K. Xu, Z. Li, C. Jing, Scr. Mater. 145 (2018) 58–61. [22] X.J. He, Y.R. Kang, S.X. Wei, Y.L. Zhang, Y.M. Cao, K. Xu, Z. Li, C. Jing, Z.B. Li, J. Alloy. Compd. 741 (2018) 821–825. [23] J. Liu, T. Gottschall, K.P. Skokov, J.D. Moore, O. Gutfleisch, Nat. Mater. 11 (2012) 620–626. [24] T. Gottschall, K.P. Skokov, D. Benke, M.E. Gruner, O. Gutfleisch, Phys. Rev. B 93 (2016) 184431. [25] Y.L. Zhang, X.J. He, Z. Li, K. Xu, C.Q. Liu, Y.S. Huang, C. Jing, Appl. Phys. Lett. 112 (2018) 182402. [26] A. Aznar, A. Gràcia-Condal, A. Planes, P. Lloveras, M. Barrio, J.L. Tamarit, W.X. Xiong, D.Y. Cong, C. Popescu, L. Mañosa, Phys. Rev. Mater. 3 (2019) 044406. [27] X. Xu, A. Nagashima, M. Nagasako, T. Omori, T. Kanomata, R. Kainuma, Appl. Phys. Lett. 110 (2017) 121906. [28] C.Q. Liu, Z. Li, Y.L. Zhang, Y.S. Huang, M.F. Ye, X.D. Sun, G.J. Zhang, Y.M. Cao, K. Xu, C. Jing, Appl. Phys. Lett. 112 (2018) 211903. [29] V.K. Pecharsky, K.A. Gschneidner, Phys. Rev. Lett. 78 (1997) 4494–4497. [30] H. Wade, Y. Tanabe, Appl. Phys. Lett. 79 (2001) 3302–3304. [31] O. Tegus, E. Brück, K.H.J. Buschow, F.R. de Boer, Nature 415 (2002) 150–152. [32] D.A. Miller, D.C. Lagoudas, Smart Mater. Struct. 9 (20 0 0) 640–652. [33] R. Lahoz, J.A. Puértolas, J. Alloy. Compd. 381 (2004) 130–136. [34] Z.B. Li, N.F. Zou, B. Yang, W.M. Gan, L. Hou, X. Li, Y.D. Zhang, C. Esling, M. Hofmann, X. Zhao, L. Zuo, J. Alloy. Compd. 666 (2016) 1–9. [35] C.S. Mejía, R. Küchler, A.K. Nayak, C. Felser, M. Nicklas, Appl. Phys. Lett. 110 (2017) 071901. [36] E.A. Mikhaleva, I.N. Flerov, M.V. Gorev, M.S. Molokeev, A.V. Cherepakhin, A.V. Kartashev, N.V. Mikhashenok, K.A. Sablina, Phys. Solid State 54 (2012) 1719. [37] T. Samanta, P. Lloveras, A.U. Saleheen, D.L. Lepkowshi, E. Kramer, I. Dubenko, P.W. Adams, D.P. Young, M. Barrio, J.L. Tamarit, N. Ali, S. Stadler, Appl. Phys. Lett. 112 (2018) 021907.