Giant and reversible room-temperature magnetocaloric effect in Ti-doped Ni-Co-Mn-Sn magnetic shape memory alloys

Acta Materialia 134 (2017) 236e248

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Giant and reversible room-temperature magnetocaloric effect in Ti-doped Ni-Co-Mn-Sn magnetic shape memory alloys Y.H. Qu a, D.Y. Cong a, *, X.M. Sun a, Z.H. Nie b, W.Y. Gui a, R.G. Li a, Y. Ren c, Y.D. Wang a a

State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, No. 30 Xueyuan Rd, Haidian District, Beijing, 100083, People’s Republic of China b School of Materials Science and Engineering, Beijing Institute of Technology, Beijing, 100081, People’s Republic of China c X-ray Science Division, Argonne National Laboratory, Argonne, IL, 60439, USA

a r t i c l e i n f o

a b s t r a c t

Article history: Received 11 April 2017 Accepted 6 June 2017 Available online 8 June 2017

The practical applications of magnetocaloric materials for magnetic refrigeration require not only large magnitude but also reversibility, as well as room-temperature operating temperature, of the magnetocaloric effect. However, it is a complex problem to simultaneously meet these demands in the Heuslertype Ni-(Co)-Mn-X (X ¼ In, Sn, Sb, Ga) magnetic shape memory alloys which are a promising candidate for magnetocaloric materials, owing to the interdependence of magnetostructural transformation parameters. Here, through synergic tuning of magnetostructural transformation parameters via alloying with Ti in Ni42-xTixCo9Mn39Sn10 alloys, we achieved a large reversible room-temperature magnetocaloric effect. By alloying with a proper amount of Ti, the martensitic transformation temperature was brought down to room temperature and the sensitivity of transformation temperature to magnetic field change was greatly enhanced with the transformation entropy change still remaining a large value, while the thermal hysteresis and the transformation interval only slightly increased. Thus, the field required to induce the complete and reversible transformation was significantly reduced. As a result, a large reversible room-temperature magnetic entropy change DSm of 18.7 J kg1 K1 under 5 T was achieved in the Ni41Ti1Co9Mn39Sn10 alloy. This is the first report on reversible DSm in Ni-Mn-Sn-based Heusler alloys and the DSm we achieved represents the highest reversible DSm under 5 T reported heretofore in Ni-Mnbased Heusler alloys. Furthermore, the Ni41Ti1Co9Mn39Sn10 alloy shows good compressive properties and stable martensitic transformation during thermal cycling, beneficial for potential magnetocaloric applications. This study is instructive for the development of high-performance magnetocaloric materials for room-temperature magnetic refrigeration. © 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Magnetic shape memory alloy Magnetocaloric effect Reversibility Magnetostructural transformation Martensitic transformation

1. Introduction Near-room-temperature refrigeration has become an indispensable technology in contemporary society. Conventional refrigeration technology based on vapor compression utilizes ozone depleting volatile refrigerants (e.g. chlorofluorocarbons) and is getting to its technical limits in accomplishing further improvements [1,2]. During the past decades, scientists and engineers have been endeavoring to explore environment-friendly and highly efficient refrigeration technologies to take the place of the traditional vapor-compression-based refrigeration. Magnetic

* Corresponding author. E-mail address: [email protected] (D.Y. Cong). http://dx.doi.org/10.1016/j.actamat.2017.06.010 1359-6454/© 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

refrigeration technology employing the magnetocaloric effect (MCE) has been considered as such a cooling technology and is expected to be promising for applications in future solid-state refrigeration devices for domestic and industrial use [2,3]. Development of high-performance magnetocaloric materials is of great significance to the implementation of magnetic refrigeration. To date, several classes of alloy systems exhibiting first-order magnetostructural transformation and consequently large MCE have been found, including Gd-Si-Ge [4], La-Fe-Si [5], Fe-Rh [6], Mn-FeP-As [7], Mn-Co-Ge [8], and Ni-(Co)-Mn-X (X ¼ In, Sn, Sb, Ga) Heusler alloys [9e12]. Among them, Ni-(Co)-Mn-X (X ¼ In, Sn, Sb, Ga) Heusler alloys have attracted great attention owing to their large MCE as a result of magnetic-field-induced phase transformation and less brittleness. Aside from large MCE, these Heusler alloys also exhibit other exciting multifunctional properties, such as

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magnetic-field-induced shape memory effect [13], magnetoresistance [14,15], and elastocaloric effect [16,17]. Recently, considerable efforts have been exerted to develop high-performance magnetocaloric materials in Ni-(Co)-Mn-X (X ¼ In, Sn, Sb, Ga) magnetic shape memory alloy systems [18e21]. For Ni-(Co)-Mn-X (X ¼ In, Sn, Sb, Ga) magnetic shape memory alloys, their first-order magnetostructural transformation usually involves significant magnetic and thermal hysteresis [22], which is a drawback of these alloys but is not rarely overlooked in the field of magnetocaloric materials. The large hysteresis could make the MCE irreversible or only partially reversible during the second and subsequent cycles of increasing and decreasing magnetic field [10,23], owing to the coexistence of martensite and austenite phases across the field-induced transformation. From the perspective of practical applications, high-performance magnetocaloric materials should meet at least the following important requirements for magnetic refrigeration [5,18]: (1) the MCE must be reversible or at least partially reversible during cyclic application and removal of magnetic field; (2) the magnetic-field-induced isothermal entropy change and adiabatic temperature change (the two parameters characterizing the MCE) must be significant at a magnetic field as low as possible; (3) the operating temperature of the MCE must be near room temperature; (4) the magnetocaloric materials should have good mechanical properties and mechanical stability during operation. However, to the best of our knowledge, there are few reports on room-temperature large reversible magnetic entropy change under a field below 5 T combined with good mechanical properties in Ni-(Co)-Mn-X (X ¼ In, Sn, Sb, Ga) magnetocaloric materials. It was reported that the reversibility of MCE in Ni-(Co)-Mn-X (X ¼ In, Sn, Sb, Ga) magnetic shape memory alloys relies not only on the sum of thermal hysteresis (DThys) and phase transformation interval (DTint) but also on the sensitivity of the transformation temperature to the field change (which equals to DM/DStr according to the Clausius-Clapeyron relation [13,21], where DM is the magnetization difference between austenite and martensite across transformation and DStr is the transformation entropy change) [24]. The minimum applied field D(m0H)min required to induce the complete reversible phase transformation can be expressed as D(m0H)min¼(DThysþDTint)/(DM/DStr) [25]. In order to facilitate the achievement of large reversible MCE, the magnetocaloric materials should possess low DThys, narrow DTint, large DM and opportune DStr. Atomic substitution by other elements has been proved to be an effective way to tune magnetostructural transformation parameters, such as phase transformation temperature, DThys, DTint, DM, and DStr. Therefore, besides composition change in the quaternary systems, addition of fifth element in Ni-Co-Mn-X (X ¼ In, Sn, Sb, Ga) alloys provides further opportunity to enhance their magnetocaloric performance. Several investigations were performed on the effect of alloying element such as Pd [26], Fe [27], Nb [28] on phase transformation and magnetocaloric performance of Ni-Co-Mn-X (X ¼ In, Sn, Sb, Ga) alloys. Although the magnetocaloric performance was indeed enhanced to some extent, large reversible room-temperature MCE has not yet been achieved, possibly due to the interdependence of magnetostructural transformation parameters which makes it difficult to get the optimal combination of these parameters [24,28]. Actually, synergic tuning of such parameters is essential for obtaining optimum magnetocaloric performance. Recently, it was found that addition of an appropriate amount of Ti can effectively adjust the transformation temperatures and improve the mechanical properties of Ni-Co-Mn-Sn alloys, without affecting the thermal hysteresis [29]. Thus, it could be feasible to simultaneously achieve room-temperature large reversible MCE and good mechanical properties in Ni-(Co)-Mn-X (X ¼ In, Sn, Sb,

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Ga) alloys via alloying with a proper amount of Ti. Nevertheless, this possibility has not been explored. Compared with other Ni-(Co)-Mn-X (X ¼ In, Sb, Ga) magnetic shape memory alloys, Ni-(Co)-Mn-Sn alloys have low cost of raw materials and contain no toxic elements. Furthermore, in comparison with Ni-(Co)-Mn-In alloys, the composition dependence of martensitic transformation temperatures in Ni-(Co)-Mn-Sn alloys is not that sensitive, which makes it easy to manipulate the chemical composition (precise control of the composition, which is difficult in practice, is not necessary) and to fabricate the alloys with desired working temperature. Although reversible magnetic entropy change at ~230 K (far below room temperature) was recently reported in a Ni-Co-Mn-In alloy [24], to our knowledge large reversible room-temperature MCE has never been achieved in Ni(Co)-Mn-Sn alloys which possess the attractive advantages as mentioned above. Therefore, a Ni-Co-Mn-Sn alloy was selected as the starting material in the present work. By synergic tuning of magnetostructural transformation parameters via Ti alloying in NiCo-Mn-Sn magnetic shape memory alloys, we successfully achieved a large reversible room-temperature magnetic entropy change of 18.7 J kg1 K1 under 5 T in the Ni41Ti1Co9Mn39Sn10 alloy, which is the highest value reported heretofore in Ni-Mn-based Heusler alloys. Furthermore, good compressive properties were also obtained in this alloy. This study is instructive for the development of high-performance magnetocaloric materials for roomtemperature magnetic refrigeration. 2. Experimental The Ni42Co9Mn39Sn10 alloy was chosen as the starting material and a series of Ni42-xTixCo9Mn39Sn10 (at.%) (x ¼ 0, 0.5, 1, and 1.1) alloys were prepared. The reason for choosing the specific composition of Ni42Co9Mn39Sn10 is that relatively high DM and large DStr are expected across phase transformation in this alloy based on the information from the literature [30], and further optimization of the magnetostructural transformation parameters via Ti alloying could be anticipated. The polycrystalline button ingots of Ni42-xTixCo9Mn39Sn10 were prepared by arc melting the high-purity constituent elements for five times on a water-cooled copper hearth under argon protection. To compensate for the Mn loss during melting, an additional 1 wt% of Mn was added in advance. The button ingots were then broken into small pieces which were subsequently remelted and cast into a copper mold protected under argon. To ensure homogeneity, the obtained rod ingots were sealed into a quartz tube evacuated and back filled with pure argon and then annealed at 1223 K for 9 h followed by water quenching. The characteristic temperatures of phase transformation were measured by differential scanning calorimetry (DSC) with heating and cooling rates of 10 K/min. The specific heat capacity (Cp) was measured by the modulated DSC technique, using a standard sample of sapphire with a well-known Cp. In order to study the crystal structure evolution during phase transformation, temperature-dependent in-situ synchrotron high-energy X-ray diffraction (HEXRD) experiments were performed at the 11-ID-C beamline at the Advanced Photon Source, Argonne National Laboratory, USA. A monochromatic X-ray beam with a wavelength of 0.1174 Å was used. The diffraction patterns were collected by a twodimensional (2D) large area detector. The isofield magnetization vs. temperature (i.e. M(T) curve) and isothermal magnetization vs. field (i.e. M(H) curve) measurements were performed using a physical property measurement system (PPMS) with a magnetic field up to 14 T. The M(H) measurements were carried out adopting the loop process method [24,31]. The main measuring steps are as follows: (1) before recording M(H)

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curves, the sample was initially cooled down to 225 K (well below martensitic transformation temperature) in a zero field to guarantee a pure martensite state; (2) then the sample was heated to the target temperature in a zero field and the M(H) curves were recorded during several cycles of increasing and decreasing field. Afterwards, the aforementioned steps were repeated for the next target temperature. The mechanical properties were examined by compressive tests on a sample of F3  6 mm3 at ambient temperature using a mechanical testing machine (Instron-5966). 3. Results and discussion 3.1. Phase transformation and composition optimization The DSC curves of the Ni42-xTixCo9Mn39Sn10 (x ¼ 0, 0.5, 1, and 1.1) alloys are demonstrated in Fig. 1a. The sharp exothermic and endothermic peaks arising from the forward martensitic (M) and reverse austenitic (A) transformations can be clearly seen. The corresponding start, peak and finish temperatures Ms, TM, Mf, As, TA, and Af are determined from the DSC curves and listed in Table 1. The obvious thermal hysteresis DThys (DThys ¼ TA-TM) suggests the firstorder nature of the transformation. The Curie temperature of austenite (Tc) can be detected as an inflection point at a higher temperature during both cooling and heating (with negligible hysteresis) (Fig. 1a). Fig. 1b shows the phase transformation temperatures and Tc of Ni42-xTixCo9Mn39Sn10 as a function of Ti content x. It can be seen that Ms, Mf, As, Af, TM and TA decrease with x increasing from 0 to 1.1, and the martensitic transformation of the alloy with x ¼ 1 takes place just around room temperature. It is well known that for NiMn-based Heusler alloys the martensitic transformation temperatures decrease with the valence electron concentration e/a decreasing [32,33]. Here, the substitution of Ti for Ni leads to a decrease in e/a, which in turn results in the decrease of transformation temperatures with Ti content increasing. For a clear representation, the e/a dependence of the martensitic transformation temperatures is also shown in Fig. 1b. On the other hand, with Ti content increasing the Tc of Ni42-xTixCo9Mn39Sn10 remains almost constant, as seen from Fig. 1a and b. The possible reason is that the substitution of Ti for Ni does not change the Ni-Mn and Mn-Mn atomic distances that are closely related to the value of Tc [33,34]. As inferred from the information shown above, with Ti content increasing the distance between Tc and TA (or TM) becomes larger (Fig. 1d), which is beneficial for achieving higher sensitivity of transformation temperature to the field change, as detailed later. The transformation entropy change associated with the temperature-induced martensitic transformation is closely related to the maximum attainable magnetic entropy change (an important parameter characterizing the MCE) resulting from the magnetic-field-induced transformation [30]. Therefore, we determined the transformation entropy changes for the forward martensitic transformation (DSM) and reverse austenitic transformation (DSA) from the DSC data by applying an integral calculation process [35], and the results are shown in Table 1 and Fig. 1c. Clearly, both the DSM and DSA of Ni42-xTixCo9Mn39Sn10 exhibit a monotonic decrease with increasing Ti content. It was reported that the transformation entropy change of Ni-Mn-based magnetic shape memory alloys is strongly related to the distance between Tc and the transformation temperature, due to the interaction between the magnetic and vibrational contributions to the total transformation entropy change [36,37]. Since the lattice entropy change and magnetic entropy change have opposite sign [38,39], the larger the distance between Tc and the transformation temperature, the smaller the transformation entropy change [40e44]. In the present case, the distance between Tc and TA (or TM) increases with increasing Ti content (see Fig. 1d), which accounts for the decrease

Fig. 1. (a) DSC curves for the Ni42-xTixCo9Mn39Sn10 alloys. Tc denotes the Curie temperature of austenite. (b) Martensitic transformation temperatures and Curie temperature Tc as a function of Ti content x and electron concentration e/a, respectively. (c) Transformation entropy changes accompanying martensitic (DSM) and reverse (DSA) transformations and (d) differences between Curie temperature and phase transformation temperature (Tc-TM for martensitic transformation and Tc-TA for reverse transformation), as a function of Ti content x.

Table 1 Characteristic temperatures (with the unit K) and transformation entropy changes (with the unit J kg1 K1) for martensitic and reverse transformations in Ni42-xTixCo9Mn39Sn10 (x ¼ 0, 0.5, 1, 1.1) alloys. The letters “M” and “A” denote martensitic and reverse transformations, respectively. x

Ms

Mf

As

Af

TM

TA

DThys Af - Mf DSM

DSA

0 0.5 1.0 1.1

365.9 337.0 291.6 274.6

356.4 327.4 279.3 254.4

371.2 341.6 297.2 277.9

378.3 349.0 305.4 291.5

362.9 332.0 289.5 266.3

375.9 345.3 303.1 286.7

13.0 13.3 13.6 20.4

33.1 30.8 22.6 16.5

21.9 21.6 25.6 37.1

32.6 28.3 20.5 12.8

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of DSA (or DSM) in the Ni42-xTixCo9Mn39Sn10 alloys. Actually, the decrease of DSA, on one hand, is advantageous for increasing the sensitivity of transformation temperature to the field change, but on the other hand, could reduce the maximum attainable magnetic entropy change. As mentioned earlier, the minimum magnetic field required to induce the complete and reversible transformation from martensite to austenite can be estimated as D(m0H)min¼(D ThysþDTint)/(DM/DStr), where DThysþDTint ¼ Af-Mf. In order to determine DM, M(T) curves under 5 T were measured for the Ni42xTixCo9Mn39Sn10 alloys, and the results are shown in Fig. 2a. The magnetostructural transformation from weak magnetic martensite to ferromagnetic austenite is observed in all alloys, and the determination of DM is illustrated in the figure. The DM/DSA, which approximately equals the sensitivity of transformation temperature to the field change according to the Clausius-Clapeyron relation [13,21], is computed and shown as a function of Ti content in Fig. 2b. Remarkably, DM/DSA considerably increases with Ti content x increasing, from 2.5 K/T for x ¼ 0 to 6.3 K/T for x ¼ 1.1. The significant increase of DM/DSA is also attributed to the enlargement of the distance between Tc and the transformation temperature, since DSA decreases (Fig. 1c) while DM tends to increase (Fig. 2a) with Tc-

Fig. 2. (a) M(T) curves measured under a magnetic field of 5 T for the Ni42-xTixCo9Mn39Sn10 alloys. The determination of phase transformation temperatures and DM is illustrated in the figure. (b) DM/DSA (left axis) and the minimum magnetic field D(m0H)min required to induce the complete reversible phase transformation (right axis) as a function of Ti content x.

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TA increasing. On the other hand, as seen from Table 1, the thermal hysteresis DThys firstly remains almost unchanged for x ¼ 0 to 1 and then increases when x reaches 1.1. The sum of DThys and DTint, namely Af-Mf, increases slightly with x increasing from 0.5 to 1 but dramatically when x increases to 1.1 (see Table 1). Clearly, the variations of DThysþDTint and 1/(DM/DSA) with x display opposite trends; the competition between them leads to first decrease and then increase of D(m0H)min with x increasing, which is shown in Fig. 2b. In consequence, the lowest value of D(m0H)min is found in the Ni42-xTixCo9Mn39Sn10 with x ¼ 1. As demonstrated above, through synergic tuning of the magnetostructural transformation parameters, we achieved a relatively low critical field to induce the complete reversible transformation in the optimal composition of Ni41Ti1Co9Mn39Sn10 which still has a large DSA. Furthermore, the martensitic transformation of this alloy occurs near room temperature. Therefore, a large reversible room-temperature MCE could be expected in this alloy. We will focus on this specific alloy in the following sections. 3.2. Crystal structure information of Ni41Ti1Co9Mn39Sn10 Since the reversibility and hysteresis of phase transformation strongly affect the magnetic-field-induced transformation during field cycling and its accompanying MCE, we carried out temperature-dependent in-situ synchrotron HEXRD experiments to provide insights into the reversibility and hysteresis of transformation from the structure point of view. Fig. 3a and b displays the experimental HEXRD patterns of Ni41Ti1Co9Mn39Sn10 measured at 320 K and 220 K, respectively. The diffraction pattern at 320 K can be well indexed according to the L21 Heusler austenite structure with lattice parameter a0 ¼ 5.960 Å. The presence of the superlattice reflections 111, 311, 331 and 511 (see Fig. 3a) indicates the second neighbor ordering of the L21 structure [10]. The diffraction pattern at 220 K (Fig. 3b) can be well indexed according to the six-layered modulated (6M) monoclinic martensite structure [30]. The lattice parameters are aM ¼ 4.393 Å, bM ¼ 5.528 Å, cM ¼ 25.830 Å, and b ¼ 93.82 ; the indexation of the pattern is shown in Fig. 3b. The evolution of the HEXRD patterns during cooling and heating is demonstrated in Fig. 3c and d, respectively. As indicated from Fig. 3c, with decreasing temperature from 300 K to 280 K, the intensity of the Bragg peaks for the L21 austenite structure decreases while the peaks for the 6M martensite structure emerge and develop, indicating the martensitic transformation process. As shown in Fig. 3d, with the increase of temperature, the peak intensity for the 6M martensite structure becomes weaker at temperatures above 290 K and the martensite peaks almost disappear at 310 K. This suggests that the reverse transformation occurs in this temperature range. The above results are in general agreement with those obtained from the DSC measurement (Fig. 1a). It should be emphasized that, during the cooling and heating processes, the martensite and austenite can reversibly transform into each other, with a low hysteresis similar to that determined from the DSC measurement. This indicates the good reversibility of the transformation induced by temperature change. Additionally, it is worth mentioning that, unlike the case of Ni43xTixCo7Mn43Sn7 (x ¼ 0.5, 1.0, 2.0 and 4.0) alloys which exhibit gphase precipitates [29], no second phase is detected from the HEXRD patterns of the present Ni41Ti1Co9Mn39Sn10 (Fig. 3), which is beneficial for obtaining good magnetocaloric performance. According to the geometric nonlinear theory of martensite, the reversibility of the martensitic transformation is to a large extent governed by the crystalline symmetry and geometric compatibility of martensite and austenite [10,45]. If the lattice vectors of lowtemperature martensite and high-temperature austenite are known, the transformation stretch tensor and its corresponding

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Fig. 3. (a, b) High-energy X-ray diffraction (HEXRD) patterns of Ni41Ti1Co9Mn39Sn10 measured at 320 K for austenite (a) and 220 K for martensite (b). (c, d) Evolution of the HEXRD patterns of Ni41Ti1Co9Mn39Sn10 during the cooling (c) and heating (d) processes.

middle eigenvalue l2, which is intimately related to the hysteresis behavior, can be determined [45e47]. If l2 is very close to 1 (which indicates the good geometric compatibility), the stress transition layer at the interface between martensite and austenite could be greatly reduced or even eliminated, and thus the martensitic transformation occurs with a very narrow thermal hysteresis [46]. To assess the geometric compatibility in the present Ni41Ti1Co9Mn39Sn10 alloy, we calculated the l2 value with the lattice parameters of austenite and martensite determined from the HEXRD patterns. The l2 turns out to be 0.9955, which is comparable to that reported for Ni50.4Mn34.8In15.8 with l2 ¼ 0.9972 [10] and a hysteresis of 12 K. The unit cell volume change across transformation is determined to be 1.46%. The close approximation of l2 to 1 (i.e. the good geometric compatibility) in the Ni41Ti1Co9Mn39Sn10 alloy accounts for its good reversibility and low hysteresis of transformation, which are of great importance for achieving stable functional performance during field cycling. 3.3. Magnetoresponsive behaviors of Ni41Ti1Co9Mn39Sn10 3.3.1. Isofield magnetization analysis The M(T) curves measured under magnetic fields of 0.05 T and 5 T for the Ni42Ti1Co9Mn39Sn10 alloy are shown in Fig. 4a. From the M(T) curve under the low field of 0.05 T, the phase transformation temperatures As, Af, Ms, and Mf are determined by the tangent line method to be 297.6 K, 305.2 K, 290.4 K, and 279.7 K, respectively, which are close to the values determined from the DSC measurement (see Table 1). Fig. 4b demonstrates the temperature dependence of dM/dT derived from the M(T) curves shown in Fig. 4a. The temperatures at which the maxima on dM/dT vs. T curves for the heating and cooling processes under 0.05 T occur, TA and TM, are determined to be 303.6 K and 289.4 K, respectively. Actually these temperatures correspond to the peak temperatures for reverse and

martensitic transformations as determined from the DSC measurement (see Fig. 1a). The thermal hysteresis of transformation DThys, estimated as TA-TM, is 14.2 K, which is in general consistency with the DSC result (see Table 1). From the M(T) curve under 5 T, a large magnetization difference DM (~108 emu/g for heating branch and ~110 emu/g for cooling branch) between the two phases across transformation can be determined (Fig. 4a), which is the highest value among the Ni-Co-Mn-Sn alloys reported so far [21,48]. Under 5 T, the phase transformation temperatures As, TA Af, Ms, TM and Mf are determined from the M(T) curve to be 270.8 K, 278.7 K, 281.9 K, 264.8 K, 262.2 K, and 248.5 K, respectively. Clearly, the phase transformation temperatures shift to the lower temperature region under the high magnetic field of 5 T. Compared with the values under 0.05 T, TA (TM) decreases by ~24.9 K (~27.2 K) upon the application of a field of 5 T, with the rate DTA/m0DH (DTM/m0DH) being ~5.0 K/T (~5.5 K/T). As is known, the decrease of TA induced by magnetic field can be well described by the Clausius-Clapeyron relation, DTA/m0DH ¼ DM/DSA [21,49]. Using our experimental data, DM/DSA is calculated to be ~4.8 K/T, which is in agreement with the DTA/m0DH value (~5.0 K/T). Similarly, DM/DSM is calculated to be ~5.4 K/T, which is consistent with the experimental DTM/m0DH value (~5.5 K/T). For magnetocaloric materials exhibiting first-order magnetostructural transformation, substantial reversibility of magneticfield-induced transformation occurs when the shift of the whole M(T) loop is larger than the thermal hysteresis [50]. In other words, to achieve a reversible MCE with a field of m0H, there should be a gap between the M(T) loop under zero field and that under m0H. In this case, for Ni-(Co)-Mn-X (X ¼ In, Sn, Sb) magnetic shape memory alloys, a reversible field-induced transformation and accompanying reversible MCE could be expected within the temperature interval bounded by the start of the reverse transformation under an applied field m0H (As under m0H) and the start of the forward

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Fig. 4. (a) M(T) curves measured under magnetic fields of 0.05 T and 5 T for Ni41Ti1Co9Mn39Sn10. (b) Temperature dependence of dM/dT derived from the M(T) curves measured under 0.05 T and 5 T for Ni41Ti1Co9Mn39Sn10. The determination of phase transformation temperatures is illustrated in the figure.

martensitic transformation under zero field (Ms under 0 T) [51]. As seen from Fig. 4a, there is indeed a gap between the M(T) loops under 5 T and 0.05 T for the presently studied Ni42Ti1Co9Mn39Sn10, which corresponds to the temperature interval between 270.8 K (As under 5 T) and 290.4 K (Ms under 0.05 T). Therefore, it is expected that reversible MCE under 5 T could be obtained in this temperature interval. It should be pointed out that, within the above-mentioned temperature interval between As under m0H and Ms under 0 T, the magnetic-field-induced transformation is reversible but not necessarily complete; the complete reversible field-induced transformation (i.e. the reversible field-induced transformation between full martensite and full austenite) only occurs within the temperature interval between Af under m0H and Mf under 0 T with a precondition that Af under m0H must be lower than Mf under 0 T. 3.3.2. Isothermal magnetization analysis In order to examine the magnetic-field-induced transformation and its reversibility, we performed M(H) measurements during cyclic increasing and decreasing magnetic field in the temperature range from 273 K to 287 K. As mentioned before, the loop process method [24,31] is used for these measurements. The resultant M(H)

241

curves measured during two cycles of increasing and decreasing field at different temperatures for Ni41Ti1Co9Mn39Sn10 are shown in Fig. 5a. Clearly, these M(H) curves exhibit the characteristic features of a metamagnetic transformation; a sudden change in the slope of M(H) is observed at a critical field m0Hcr followed by a fast and progressive increase in magnetization, indicative of a magneticfield-induced transformation. To reveal the nature of the fieldinduced transformation, the Arrott plots derived from the M(H) curves shown in Fig. 5a are plotted in Fig. 5b. According to the Banerjee’s criterion, a magnetic transition is regarded as first-order if the Arrott plot displays a negative slope; otherwise, it is secondorder [52]. The obvious appearance of negative slopes of the Arrott plots in Fig. 5b confirms that the field-induced transformation in the studied alloy is of first order, which could result in large MCE. As seen from the M(H) curves measured during two field cycles shown in Fig. 5a, at a given temperature (e.g. 285 K) the low-field M(H) curve recorded during the first increasing field does not coincide with that recorded during the first decreasing field (here, low field denotes m0H < 0.5 T), with the latter being slightly above the former. This suggests that the field-induced austenite during the first increasing field does not fully, although a high fraction of such austenite does, transform back to martensite during the first decreasing field. Notably, the low-field M(H) curve during the second increasing field coincides well with that recorded during the second decreasing field. The M(H) curves recorded during the third and following field cycles (not shown here) are the same as that recorded during the second field cycle. This indicates that the martensite that is transformed back during the first decreasing field could reversibly transform to austenite and back during the second and all subsequent cycles of increasing and decreasing field, which could result in a reversible MCE, while the residual field-induced austenite after removing field in the first field cycle does not participate in the transformation any more, similar to the case reported in Ref. [50]. The critical field needed for the start of the field-induced transformation, m0Hcr, is determined as illustrated in Fig. 5a. The temperature dependence of m0Hcr is shown in Fig. 5c, from which one can see that m0Hcr decreases almost linearly with increasing temperature. It is worth noting that the slope of m0Hcr vs. T curve for the second field cycle is almost equal to that for the first field cycle, both with a high slope of 0.19 T/K. Similar linear relationship was also observed in a Mn0.99Fe0.01As alloy [53]. The magnetic hysteresis of the field-induced transformation, DHhys, is also determined as illustrated in Fig. 5a. Clearly, at a given temperature, the DHhys for the second field cycle is lower than that for the first field cycle. It is well known that the magnetic hysteresis DHhys is closely related to the thermal hysteresis DThys [54]. Recently, it was well demonstrated that the thermal hysteresis could be significantly narrowed by following a minor loop of hysteresis (instead of completely transforming the material between martensite and austenite) [55,56], because with the minor loop process it is not necessary to nucleate new phase during back transformation and thus the required energy is reduced [55]. In our case, the decrease of DHhys in the second field cycle is attributed to the minor loop process of the magnetic-field-induced transformation, which is equivalent to the minor loop process of temperature-induced transformation reported in Refs. [6,55,56]. At the beginning of the second field cycle, there already exists the residual field-induced austenite formed during the first field cycle, and it is not necessary to nucleate new austenite for the field-induced transformation during the second increasing field. Therefore, the field-induced transformation in the second field cycle occurs at a lower field (see e.g. the M(H) curve at 287 K in Fig. 5a) than that in the first field cycle (at the beginning of this cycle the sample is in full martensite state), leading to the reduction of DHhys in the second field cycle.

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Fig. 5. (a) M(H) curves measured during the first (solid lines) and second (dashed lines) cycles of increasing and decreasing magnetic field at different temperatures for Ni41Ti1Co9Mn39Sn10. The dotted lines denote the extrapolation of the magnetization of pure austenite and pure martensite. The determination of magnetic hysteresis DHhys and critical field (m0Hcr) for magnetic-field-induced transformation is illustrated in the figure. (b) Arrott plots (i.e., M2 vs. m0H/M) derived from the M(H) curves at different temperatures shown in (a). (c) Temperature dependence of the critical field (m0Hcr) for field-induced transformation, with the data derived from the M(H) curves shown in (a).

3.3.3. Evaluation of magnetocaloric performance It is well known that the realistic magnetic entropy change of magnetocaloric materials with first-order magnetostructural transformation can be determined from the isofield magnetization data or isothermal magnetization data. It should be mentioned that, here the magnetic entropy change denotes the total entropy change induced by the magnetic field. To get a comprehensive and reliable evaluation of the MCE, we used different methods to estimate the magnetic entropy change in the studied Ni41Ti1Co9Mn39Sn10 alloy. Firstly, we estimated the magnetic entropy change DSm with the transformation fraction method, since the transformed fraction of the material directly reflects the yielded MCE [10,25,57]. According to this method, the DSm induced by a field change from an initial value m0Hi to a final value m0Hf can be estimated by the following relation:









DSm ¼ Df $DStr ¼ f T; Hf  f ðT; Hi Þ $DStr

(1)

where f ðT; HÞ is the fraction of austenite at a given temperature and magnetic field [57], Df reflects the transformed fraction, and DStr is the transformation entropy change at a given magnetic field which can be determined from the Clausius-Clapeyron relation (DStr ¼ DM/(DTA/m0DH)) [13,21,49]. Based on the M(T) curve under a given field (Fig. 4a), the mass fraction of austenite can be estimated with:

f ðTÞ ¼

MðTÞ  MM ðTÞ MA ðTÞ  MM ðTÞ

(2)

where M(T) is the experimentally measured magnetization shown on the M(T) curve; MM(T) and MA(T) are the magnetization of pure martensite and pure austenite, respectively, which can be roughly extracted by linearly extrapolating the heating M(T) curve from the transformation point (see Refs. [25,57] for more details). With Eq. (2), the f(T) curves under 0.05 T and 5 T, and their difference, i.e. the Df(T) curve for a field change from 0.05 T to 5 T, were determined from the heating M(T) curves (Fig. 4a) and demonstrated in Fig. 6a. The DStr (see Eq. (1)) under 5 T was estimated to be 21.6 J kg1 K1. The DSm estimated with this transformation fraction method based on M(T) data (denoted as TF_M(T)) is shown as a function of temperature in Fig. 6b. As can be seen, the maximum DSm for a field change from 0.05 T to 5 T is 21.3 J kg1 K1. Furthermore, the mass fraction of austenite can also be determined from the M(H) curve at a given temperature [25], according to the following equation:

f ðHÞ ¼

MðHÞ  MM ðHÞ MA ðHÞ  MM ðHÞ

(3)

where M(H) is the measured magnetization displayed on the M(H) curve (Fig. 5a), and MM(H) and MA(T) denote the magnetization of

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Fig. 6. (a) Fraction of austenite f(T) under 0.05 T and 5 T determined from the M(T) curves (heating branch) shown in Fig. 4a, and the Df(T) for a field change from 0.05 T to 5 T derived from the difference between the f(T) curves under 0.05 T and 5 T. (b) Magnetic entropy change DSm under 5 T, estimated with different methods, shown as a function of temperature. TF_M(T), TF_M(H), C-C_M(H), and Maxwell_M(H) denote transformation fraction method based on M(T) data, transformation fraction method based on M(H) data, Clausius-Clapeyron relation based on M(H) data, and Maxwell relation based on M(H) data, respectively. The M(H) data here refers to those recorded during the first field cycle shown in Fig. 5a. (c) M(H) curves measured at 314 K (austenite state) and 273 K (martensite state under 0 T). (d) Fraction of austenite f(H) at different constant temperatures determined from the M(H) curves recorded during the first field cycle shown in Fig. 5a. (e) Magnetic entropy change DSm as a function of applied magnetic field at the constant temperature of 285 K, determined from the f(H) curve at 285 K in (d). Solid line with solid squares (curve 1) and solid line with solid circles (curve 2) denote the curves determined using the TF_M(H) method with point A in (d) selected as the initial point. Solid line with open circles (curve 3) denotes the curve determined using the TF_M(H) method with point B in (d) selected as the initial point. Solid stars denote the data estimated with the C-C_M(H) method.

pure martensite and pure austenite, which can be deduced from the M(H) curves [25]. As can be seen from Fig. 5a, the M(H) curves recorded at 287 K and 285 K both saturate at the high field of 5 T and almost coincide with each other, indicating that the sample is in pure austenite state under 5 T at these two temperatures. Therefore, the magnetization of pure austenite can be extracted by linearly extrapolating the demagnetization curve at 287 K recorded during the first field cycle, as illustrated in Fig. 5a. Considering that the magnetization of both austenite and martensite rapidly changes at low fields (see Fig. 6c), which results in numerical instabilities below ~0.5 T [25], and stably increases above a certain critical field (e.g. ~0.5 T) (Fig. 6c) provided that no field-induced transformation occurs, the data below 0.5 T was not included when we estimated the fraction of austenite using Eq. (3). In this case, the magnetization of pure martensite was extracted by

linearly extrapolating the M(H) curve between 0.5 T and 2.5 T (above 2.5 T field-induced transformation starts to occur) recorded during the first increasing field at 273 K, as illustrated in Fig. 5a. It should be noted that, when we used this transformation fraction method based on M(H) data (denoted as TF_M(H)) to estimate DSm at different temperatures, the temperature dependence of the magnetization of pure austenite and pure martensite is neglected because the change in the magnetization of pure austenite and pure martensite within a narrow temperature interval is very small [21,48]; actually, we only considered a narrow temperature interval of 273e287 K. The f(H) curves at different constant temperatures, determined from the M(H) curves recorded during the first field cycle (Fig. 5a) using Eq. (3), are displayed in Fig. 6d. The DSm for a field change from 0.5 T to 5 T at different constant temperatures, estimated with this TF_M(H) method

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using Eq. (1), is also shown in Fig. 6b. It can be seen that the DSm estimated using this method is consistent with that estimated using the above-mentioned TF_M(T) method. The advantage of the TF_M(H) method is that it allows the determination of DSm as a function of field at a given temperature, which provides the direct view of the variation of DSm during the field increasing (or decreasing) process. With the f(H) curves shown in Fig. 6d, we determined the DSm vs. H curves for different constant temperatures. Considering the small difference between the transformation entropy change DStr under 0 T (DSA ¼ 22.6 J kg1 K1, determined from DSC measurement, see Table 1) and that under 5 T (DSA ¼ 21.6 J kg1 K1, determined with the Clausius-Clapeyron relation as mentioned above), the DStr under 5 T is taken for the determination of DSm vs. H curves using Eq. (1). The DSm vs. H curve for 285 K, as one example, is displayed in Fig. 6e. During the estimation process, we firstly selected the f(H) value at 0.5 T (point A in Fig. 6d) as the initial point and the continuous magnetization and demagnetization process (A/B/C, Fig. 6d) is treated as a whole process, and the obtained DSm vs. H curves are shown as curves 1 and 2 in Fig. 6e. It can be seen that DSm increases with increasing field during the magnetization process of the first field cycle (curve 1) and the maximum DSm is 21.6 J kg1 K1. Notably, after decreasing field to 0.5 T (point C0 in Fig. 6e), the DSm does not return to the initial value (point A0 in Fig. 6e), which is owing to the residual fieldinduced austenite upon the removal of field. We also calculated the DSm vs. H curve for the demagnetization process in the case that the f(H) value at 5 T (point B in Fig. 6d) is selected as the initial point, namely the demagnetization process (B/C, Fig. 6d) is treated separately from the magnetization one (A/B, Fig. 6d); the obtained jDSm j vs. H curve is shown as curve 3 in Fig. 6e. As can be seen, the curves 2 and 3 actually exhibit a mirror image relationship. To further verify our estimated results of DSm, we also used the Clausius-Clapeyron relation expressed as follows [53,58] to determine DSm:

dHcr : dT

DSm ¼ DM0 $

(4)

The use of Eq. (4) is based on the M(H) data and it allows the correct determination of DSm even in the case of coexistence of two phases, since in this equation DSm is directly related to the field-induced magnetization difference DM 0 at a given temperature. In this equation, dHcr/dT is the temperature dependence of critical field and can be determined from Fig. 5c; DM0 can be determined as the difference between the magnetization at the final field m0Hf and that at the initial field m0Hi. The DSm for a field change from 0.5 T to 5 T at different constant temperatures, estimated with this method (denoted as C-C_M(H)) based on the M(H) curves recorded during the first field cycle (Fig. 5a), is also demonstrated in Fig. 6b. In addition, the DSm vs. H curves for different constant temperatures were also determined with this method. The DSm vs. H curve for the magnetization process of the first field cycle at 285 K is shown in Fig. 6e as well, which is in good agreement with that determined with the above-mentioned TF_M(H) method. Since we employed the loop process method [24,31] to measure the M(H) curves and all the M(H) curves for the first field cycle were recorded with pure martensite as the initial state (the measurement temperatures were lower than As under 0 T), it is reliable to use the Maxwell relation to estimate DSm [24]. Therefore, based on the M(H) curves recorded during the first field cycle (Fig. 5a), we also computed the DSm for a field change from 0 to 5 T at different temperatures using the Maxwell relation [7,24]:

DSm ¼

Zm0 H 0

vM dðm0 HÞ: vT

(5)

The results obtained with this method (denoted as Maxwell_M(H)) are also displayed in Fig. 6b. It is clear that the DSm data obtained from all the methods mentioned above agree well with each other, confirming the validity of the DSm values we obtained. On the other hand, this also indicates that it is reliable to use the TF_M(T) method based on M(T) data and TF_M(H), CC_M(H) and Maxwell_M(H) methods based on the M(H) data recorded during the first field cycle to estimate the DSm values in the present case. From the viewpoint of practical applications, only the reversible MCE is useful for magnetic refrigeration. Therefore, it is important to determine the reversible magnetic entropy change DSm. As mentioned before, the magnetic-field-induced transformation is reversible during the second and subsequent field cycles. In this sense, the most straightforward way to estimate the reversible DSm is to use the transformation fraction method (TF_M(H)) based on the M(H) data recorded during the second field cycle (Fig. 5a). With regard to other methods, the Maxwell_M(H) method is inapplicable because of the coexistence of residual austenite and martensite at the beginning of the M(H) measurements for the second field cycle; in this case, the Maxwell_M(H) method could lead to spurious results [59]. The TF_M(T) method is not able to determine the reversible part of the DSm based on M(T) data in the case that the field-induced transformation is only partially reversible during the first field cycle. The use of the C-C_M(H) method to determine the reversible DSm requires the determination of dHcr/dT from the consecutive M(H) measurements at different temperatures, but the M(H) measurements for the second field cycle at different temperatures were interrupted by the measurements for the first field cycle at such temperatures. Therefore, it seems that the TF_M(H) method based on the M(H) data recorded during the second field cycle is the only feasible and reliable method to determine the reversible DSm in our case where the field-induced transformation is not fully reversible during the first field cycle. With the TF_M(H) method, the fraction of austenite as a function of field, i.e. f(H) curves, at different constant temperatures is determined from the M(H) curves recorded during the second field cycle (Fig. 5a) and shown in Fig. 7a. Subsequently, using Eq. (1) we determined the DSm vs. H curves for different constant temperatures. As an example, the curve for 285 K is displayed in Fig. 7b. Similar to the estimation process for the DSm vs. H curves shown in Fig. 6e, we firstly selected point D in Fig. 7a as the initial point and obtained the curves I and II in Fig. 7b. Then we also selected point E in Fig. 7a as the initial point for estimating the DSm vs. H curve for the demagnetization process and obtained the curve III. As can be seen from the curves I and II for the magnetization and demagnetization processes, after decreasing field to 0.5 T (F0 in Fig. 7b) the DSm returns to the initial value (D0 in Fig. 7b), indicating that the DSm is reversible. Therefore, a reversible DSm with the maximum value of 18.7 J kg1 K1 is obtained at this temperature. For comparison, the DSm vs. H curve (see Fig. 6e) determined from the M(H) curve recorded during the first field cycle at 285 K is also included in Fig. 7b. Clearly, the DSm vs. H curves (determined with the f(H) value at 5 T selected as the initial point) for the second decreasing field (curve III in Fig. 7b) and first decreasing field (curve 3 in Fig. 7b) processes coincide well with each other, because the martensite that is transformed back during the first decreasing field could reversibly transform to austenite and back during the second and all subsequent field cycles. This indicates that, during the practical application process, one can firstly magnetize the material

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Fig. 7. (a) Fraction of austenite f(H) at different constant temperatures determined from the M(H) curves recorded during the second field cycle shown in Fig. 5a. (b) Magnetic entropy change DSm as a function of applied magnetic field at the constant temperature of 285 K. The red curves denote the results determined from the f(H) curve at 285 K in (a), which correspond to the second field cycle shown in Fig. 5a. Curves I and II denote the curves determined using the TF_M(H) method with point D in (a) selected as the initial point. Curve III denotes the curve determined using the TF_M(H) method with point E in (a) selected as the initial point. The dark yellow curves (corresponding to the first field cycle shown in Fig. 5a) are taken from Fig. 6e for comparison, with curves 1, 2 and 3 interpreted therein. (c) Magnetic entropy change DSm for a field change from 0.5 T to 5 T, estimated with TF_M(H) method based on the M(H) curves recorded during the second (red solid line with solid circles) and first (dark yellow line with open diamonds) field cycles, shown as a function of temperature. The data determined from the second field cycle (red solid circles) are reversible DSm, while those determined from the first field cycle (dark yellow open diamonds) are included for comparison. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

(this is like a kind of training process) and then utilize the reversible MCE during the demagnetization process and all subsequent field cycles. The reversible DSm for a field change from 0.5 T to 5 T at different constant temperatures were determined with this TF_M(H) method and shown in Fig. 7c. It can be seen that the maximum value of the

Fig. 8. The reversible DTad under 5 T, estimated with Eq. (6), shown as a function of temperature. The inset shows the zero-field specific heat capacity Cp as a function of temperature measured during heating for the Ni41Ti1Co9Mn39Sn10 alloy.

reversible DSm is 18.7 J kg1 K1. For comparison, the DSm values for the same field change determined from the M(H) curves recorded during the first field cycle are also included in Fig. 7c. Clearly, the DSm values obtained during the second field cycle are smaller than those obtained during the first cycle, because at any given temperature the residual field-induced austenite formed during the first decreasing field does not participate in the transformation any more in the following field cycles. To our knowledge, this is the first report on the reversible DSm in Ni-Mn-Sn-based Heusler alloys and, furthermore, the reversible DSm we achieved (18.7 J kg1 K1) represents the highest value under the same field reported so far in Ni-Mn-based Heusler alloys. In addition, here, for the first time, we used the transformation fraction method based only on isothermal and isofield magnetization measurements to determine the reversible magnetic entropy change. With this, we provide a simple and reliable method to evaluate the reversible MCE in magnetocaloric materials with first-order magnetostructural transformation. Moreover, the TF_M(H) method we used here allows the determination of DSm as a function of field at a given temperature, providing the direct view of the variation of DSm during the field increasing (or decreasing) process. Aside from DSm, the adiabatic temperature change DTad is the other important parameter for characterizing MCE. As the transformation mass fraction directly reflects the cooling effect, the fractional DTad is roughly estimated with the following relation [10]:

T Cp

T Cp

DTad ¼  $DStr $Df ¼  $DSm :

(6)

where Cp is the specific heat capacity. The Cp of our studied

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Ni41Ti1Co9Mn39Sn10 alloy is determined with the modulated DSC technique and shown as a function of temperature in the inset of Fig. 8. With the reversible DSm and corresponding Cp values, we estimated the reversible DTad using Eq. (6). The results are demonstrated as a function of temperature in Fig. 8. As can be seen the maximum value of the reversible DTad under 5 T is 10.5 K, which is actually a large DTad. The reversible DSm values under 5 T for the typical well-known magnetocaloric materials [4,24,60e67] are schematically displayed in Fig. 9. As can be seen, the present Ni41Ti1Co9Mn39Sn10 alloy shows a large reversible DSm, which is comparable to that obtained in the most promising Gd-Si-Ge-based and La-Fe-Si-based magnetocaloric materials under the same field. Nevertheless, although through synergic tuning of magnetostructural transformation parameters the critical field to induce the complete reversible transformation in our alloy has been greatly reduced, the field (i.e. 5 T) needed to achieve the above-mentioned reversible DSm (18.7 J kg1 K1) is still high as compared with the Gd-Si-Ge-based and La-Fe-Si-based magnetocaloric materials which only need a relatively low field to induce the complete reversible magnetostructural transformation. Therefore, future efforts should be devoted to further reducing the field needed to induce the large reversible DSm, possibly through further narrowing the thermal hysteresis by optimizing the geometric compatibility between austenite and martensite or enhancing the magnetization difference between the two phases by tuning their exchange interactions. A recent study suggests that the critical filed to induce the complete reversible transformation can also be reduced by coupled magneto-mechanical loading [68]. As mentioned before, the high-performance magnetocaloric materials should possess reasonably good mechanical properties so that they could be processed into required shapes for improving the heat-exchange performance [69] and could be used for a longer service lifetime during field cycling [70]. It should be noted that the well-known Gd-Si-Ge and La-Fe-Si-based magnetocaloric materials are quite brittle [71], even lack of any machinability. In contrast, our Ni41Ti1Co9Mn39Sn10 alloy possesses good mechanical properties. The compressive stress-strain curve measured at room temperature for this alloy is illustrated in Fig. 10, from which one can see that the compressive strength is as high as ~730 MPa and the compressive strain reaches ~4.5%. This Ni41Ti1Co9Mn39Sn10 alloy can be processed into

Fig. 10. Compressive stress-strain curve measured at room temperature for the Ni41Ti1Co9Mn39Sn10 alloy.

Fig. 11. DSC curves recorded during thermal cycling across martensitic transformation (with 110 cycles of cooling and heating) for the Ni41Ti1Co9Mn39Sn10 alloy.

different shapes by machining, which is beneficial for potential applications. The functional fatigue behavior of the MCE based on magneticfield-induced transformation is of great importance for practical applications, and a reproducible martensitic transformation is a prerequisite for the functional stability. To examine the cyclic stability of martensitic transformation in the Ni41Ti1Co9Mn39Sn10 alloy, we performed thermal cycling experiments that are commonly used to measure the transformation reproducibility [45]. Fig. 11 shows the DSC curves recorded during 110 cycles of cooling and heating across martensitic transformation. As can be seen, the peak positions and peak heights of martensitic transformation remain almost unchanged during the thermal cycling process, clearly demonstrating the superior cyclic stability of the martensitic transformation in our Ni41Ti1Co9Mn39Sn10 alloy.

4. Conclusions Fig. 9. Schematic illustration of the reversible magnetic entropy change DSm under 5 T shown as a function of temperature for some of the well-known magnetocaloric materials. The numbers in the square brackets are the reference numbers.

Through synergic tuning of magnetostructural transformation parameters in Ni42-xTixCo9Mn39Sn10 magnetic shape memory alloys

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via alloying with Ti, we achieved a large reversible roomtemperature magnetocaloric effect. When the Ti content reaches 1% (i.e. x ¼ 1), the martensitic transformation temperature is brought down to room temperature, and the distance between the Curie temperature and martensitic transformation temperature is enlarged. This results in the increase of DM/DStr which is indicative of the sensitivity of the transformation temperature to the field change, while DStr still remains a large value. On the other hand, the sum of thermal hysteresis (DThys) and transformation interval (DTint), i.e. Af - Mf, only slightly increases. Thus, the field required to induce the complete reversible phase transformation is greatly reduced. Consequently, a large reversible room-temperature magnetic entropy change DSm of 18.7 J kg1 K1 under 5 T, which is corroborated by the transformation fraction method based on the M(H) data measured during the second field cycle, is achieved in the Ni41Ti1Co9Mn39Sn10 alloy. This is the first report on the reversible DSm in Ni-Mn-Sn-based Heusler alloys and the reversible DSm we achieved represents the highest reversible DSm under 5 T reported up to now in Ni-Mn-based Heusler alloys. Furthermore, the Ni41Ti1Co9Mn39Sn10 alloy shows good compressive properties and exhibits a stable martensitic transformation during thermal cycling. In addition, it contains neither expensive nor toxic elements. All these merits make this alloy attractive for potential room-temperature magnetic refrigeration applications. The implications of the present work are important for designing highperformance magnetocaloric materials for environment-friendly solid-state cooling. Acknowledgments This work is supported by the National Natural Science Foundation of China (Nos. 51471030, 11305008, and 51527801), the National High Technology Research and Development Program of China (863 Program) (No. 2015AA034101), the Fundamental Research Funds for the Central Universities (Nos. 06111023 and 06111020), and also supported by State Key Laboratory for Advanced Metals and Materials (Grant Nos. 2016-T01 and 2015ZD01). Use of the Advanced Photon Source was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Science, under Contract No. DE-AC02-06CH11357. References [1] K.A. Gschneidner Jr., V.K. Pecharsky, A.O. Tsokol, Recent developments in magnetocaloric materials, Rep. Prog. Phys. 68 (2005) 1479e1539. [2] A.M. Tishin, Y.I. Spichkin, Recent progress in magnetocaloric effect: mechanisms and potential applications, Int. J. Refrig 37 (2014) 223e229. [3] T. Samanta, D.L. Lepkowski, A.U. Saleheen, A. Shankar, J. Prestigiacomo, I. Dubenko, A. Quetz, I.W.H. Oswald, G.T. McCandless, J.Y. Chan, P.W. Adams, D.P. Young, N. Ali, S. Stadler, Hydrostatic pressure-induced modifications of structural transitions lead to large enhancements of magnetocaloric effects in MnNiSi-based systems, Phys. Rev. B 91 (2015), 020401. [4] V.K. Pecharsky, K.A. Gschneidner Jr., Giant magnetocaloric effect in Gd5(Si2Ge2), Phys. Rev. Lett. 78 (1997) 4494e4497. [5] Y.Y. Shao, J. Liu, M.X. Zhang, A. Yan, K.P. Skokov, D.Y. Karpenkov, O. Gutfleisch, High-performance solid-state cooling materials: balancing magnetocaloric and non-magnetic properties in dual phase La-Fe-Si, Acta Mater. 125 (2017) 506e512. [6] A. Chirkova, K.P. Skokov, L. Schultz, N.V. Baranov, O. Gutfleisch, T.G. Woodcock, Giant adiabatic temperature change in FeRh alloys evidenced by direct measurements under cyclic conditions, Acta Mater. 106 (2015) 15e21. [7] O. Tegus, E. Brück, K.H.J. Buschow, F.R. de Boer, Transition-metal-based magnetic refrigerants for room-temperature applications, Nature 415 (2002) 150e152. [8] J. Liu, K. Skokov, O. Gutfleisch, Magnetostructural transition and adiabatic temperature change in Mn-Co-Ge magnetic refrigerants, Scr. Mater. 66 (2012) 642e645. [9] S. Fabbrici, J. Kamarad, Z. Arnold, F. Casoli, A. Paoluzi, F. Bolzoni, R. Cabassi, M. Solzi, G. Porcari, C. Pernechele, F. Albertini, From direct to inverse giant magnetocaloric effect in Co-doped NiMnGa multifunctional alloys, Acta Mater.

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