Computation-guided design of Ni-Mn-Sn ferromagnetic shape memory

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Computation-guided design of Ni-Mn-Sn ferromagnetic shape memory alloy with giant magnetocaloric effect and excellent mechanical properties and high working temperature via multi-element doping Kun Zhang, Chang-long Tan, Wenbin Zhao, Erjun Guo, and Xiaohua Tian ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b08640 • Publication Date (Web): 28 Aug 2019 Downloaded from pubs.acs.org on August 30, 2019

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ACS Applied Materials & Interfaces

Computation-guided design of Ni-Mn-Sn ferromagnetic shape memory alloy with giant magnetocaloric effect and excellent mechanical properties and high working temperature via multielement doping Kun Zhang†, Changlong Tan†*, Wenbin Zhao‡, Erjun Guo‡, Xiaohua Tian† † School ‡ School

of Science, Harbin University of Science and Technology, Harbin 150080, China. of Materials Science and Engineering, Harbin University of Science and Technology, Harbin 150080, China

ABSTRACT: Ni-Mn-Sn ferromagnetic shape memory alloys (FSMAs) have promise for application in efficiency solid-state refrigeration. However, the simultaneous achievement of giant magnetocaloric effect (MCE) and excellent mechanical properties and high working temperature in these materials is always the challenge. Computation-guided materials design techniques provide an efficient way to design and identify new magnetocaloric materials. Herein, a new strategy of multidoping is presented. Firstly, we conduct a detailed and comprehensive first-principle study and predict that Ni-Mn-Sn FSMAs with co-doping 6.25 at.% Cu and 6.25 - 12.5 at.% Co can realize the multi-objective optimization of magnetocaloric material. Then it is confirmed by experiment and we report on Ni42Co6Mn37Sn9Cu6 FSMA exhibiting a large magnetic entropy change (34.8 J/kg·K), of a large value in the prevalent MCE materials at high temperature (~344 K), and whose compression stress and strain (~1072.0 MPa and ~11.9%) are both the largest among Ni-Mn-based MCE materials. Notably, the effect of Co and Cu doping is not simply stacked because that they play opposite roles in Curie temperature (TC) and martensitic transformation temperature (TM). So, achieving the balance of their effect to combine their merits in a very narrow window is the key step. This approach of multi-element doping holds promises to be extended to other magnetocaloric materials to enhance their multiple properties simultaneously. KEYWORDS: Ni-Mn-Sn alloy, high temperature, mechanical properties, magnetocaloric effect, multi-element doping

INTRODUCTION As the foundation of every country’s sustainable development strategies, efficient energy use and environmental protection have been regarded as a global focus. Nowadays, vapor compression refrigeration technology faces the challenges of replacing refrigerant and improving refrigeration efficiency. Looking for alternative refrigeration techniques is to be imminent. Solid-state refrigeration has recently been demonstrated as a promising environmentally friendly and of higher energy efficiency cooling technology1-5. Particularly, magnetic refrigeration based on the magnetocaloric effect (MCE) shows great competitiveness after the discovery that Gd5Ge2Si2 alloys show a giant MCE around room temperature 6,7. Then giant MCE based on a first-order phase transition has been observed in materials including the La-Fe-Si related compounds8-10, as well as Mn-Fe-Pbased11-13, Fe2P-based14,15 and other materials16,17. As a magnetocaloric material, in addition to the large MCE, a reasonably good mechanical property is another critical factor so that they could be processed into required shapes for improving the heat-exchange performance and could be used for a longer service lifetime. Nevertheless, these materials above show the poor mechanical properties and the intrinsic brittleness, which severely restrict their application. Moreover, so far, most of these materials just can be used below or near room temperature, but in many cases and extreme environment that the working temperature will exceed room temperature, e.g. airconditioning in the summer. Unfortunately, the existing magnetocaloric materials cannot solve all of these problems at the same time. So, developing a new magnetocaloric material with giant MCE, higher working temperature and

excellent mechanical properties is of great practical value and full of challenges. Ni-Mn-Sn ferromagnetic shape memory alloys (FSMAs) capable of a direct magnetic field-induced reverse martensitic transformation (MFIRMT) and excellent magnetocaloric effects have promise for application in efficiency solid state refrigeration18-22. The strong magnetostructural coupling under the martensitic transformation (MT) would leads to a drastic difference of magnetization (ΔM), which directly affects the MCE. Moreover, researchers have found many amazing multifunctional properties based on MFIRMT23-28. It has been reported in 2013 that the refrigerating effect could be enhanced by combing the elastocaloric effect and magnetocaloric effect in Ni43Mn40Sn10Cu7 FSMAs25. Compared with the currently mainstream materials La-FeSi, Gd-Si-Ge and so on, Ni-Mn-Sn magnetic alloys are nontoxic, green and low cost. These excellent properties make Ni-Mn-Sn alloy system has extensive application prospect in magnetocaloric refrigeration. Though the problems of brittleness and low working temperature remain, it is possible to design a satisfactory material due to the strong sensitivity of properties to the chemical composition in NiMn-Sn alloy. Here, we plan to explore a high-temperature Ni-Mn-Sn alloy with excellent mechanical properties and giant MCE by element doping. Usually, the MCE can be quantified by isothermal magnetic entropy change (ΔSM), and a large ΔSM would be expected in the alloys with a large ΔM. Up to now, the investigating on the magneto structural transition by the fourth element doping has been explored by researchers all over the world29-37. There should be emphasized that Co substitution is more interesting because very little Co concentration can strengthen ferromagnetic interactions and affect the Curie temperature (TC). Due to that Co doping

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greatly changes the magnetic configuration in Ni-Mn based Heusler alloys, Co doping shows a huge influence on the ΔM in Ni-Mn based SMAs30,38,39. Liu at al. found that the ΔM across the MT increased remarkably in Ni45Co5Mn37In13 (50 emu/g) in contrast to the Ni49Co1Mn37In13 (0.03 emu/g)40. Also, Co doping in Ni-Mn-Sn alloys has the same effect on enhancing the ΔM19, 33. Moreover, Krenke et. al made a research on Ni50-xCoxMn37Sn13 alloy and found that addition of 3 at.% Co doping could improve the TC to 335 K from 311 K21. Cong et. al observed a relatively high TC of 396 K in Ni44Co6Mn39Sn11 alloy41. It seems that Co doping is suitable for enhancing MCE and improving the Tc in Ni-Mn-Sn FSMAs. Unfortunately, a big disadvantage to the substitution Ni by Co in Ni-Mn-Sn FSMAs is shifting the TM to a lower temperature42, 43. However, TM is one of the indispensable factors in deciding working temperature. In order to improve the working temperature of Ni-Mn-Sn FSMAs, the key is to shift the TM and Curie temperature (TC) to high temperature at the same time. Wang et. al investigate the martensitic transformation and magnetocaloric effect in Cu doped Ni43Mn46-xCuxSn11 alloys. These alloys can not only obtain a giant low-field ΔSM but also shift the TM to higher temperature44. In our previous work, we also found that the substitution Mn by Cu in Ni-Mn-Sn FSMAs can remarkably improve the TM. Moreover, we also observed that the mechanical properties can be significantly enhanced by Cu doping37. Even so, it is not perfect due to that Cu doping in Ni-Mn-Sn alloys always decreases the TC. Whether Cu or Co doping, they can only enhance one or two aspect of the properties, and single element doping could not achieve the design of Ni-Mn-Sn FSMAs with excellent mechanical properties and giant MCE at high temperature. Naturally, we anticipate that Cu and Co codoping in Ni-Mn-Sn FSMAs may achieve the material design objectives. Notice, however, the effect of Co and Cu doping is not simply aggregated because that they play opposite roles in TC and TM. How to achieve the balance of their effect is the key step. Therefore, the greatest challenge is to find the effective composition range by adjusting the doping content of Co and Cu that can improve TC and TM and ΔM at the same time. Relying on a significant amount of repeated experiments with continuous trial and error, the traditional research and development mode of materials is not only costly, but also difficult to achieve success. Computational design can provide great theoretical guidance and predictability in new material design, which also has the invention has simple, fast and effective test method, low cost and short research period45-47. Up to now, the first-principle calculations based on density function theory have become one of the most powerful tools in investigating Ni-Mn based FSMAs48-52. Predict first by theories calculation and then be demonstrated by targeted experiment, which provide us an efficient model of designing a new Ni-Mn-Sn FSMAs with well improved performance. For exploring a new high-temperature Ni-Mn-Sn magnetic alloy with giant MCE and excellent mechanical properties, we first conduct a detailed and comprehensive first-principle study of the influence of Co and Cu introduction. Calculated results predict that Ni-Mn-Sn FSMAs with 6.25 at.% Cu and 6.25 - 12.5 at.% Co can realize

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the multi-objective optimization of material design. Subsequently, according to the results of calculations, Ni48xCoxMn37Sn9Cu6 (x = 0, 6, 8, 10, 12) alloys were prepared and investigated. Ultimately, we report on Ni42Co6Mn37Sn9Cu6 FSMA exhibiting a large magnetic entropy change (34.8 J/kg·K), of a large value in the prevalent MCE materials at high temperature (~344 K), and whose compression stress and strain (~1072.0 MPa and ~11.9%) are both the largest among Ni-Mn-based MCE materials. These results are of practical importance, because it shows that Ni-Co-Mn-SnCu alloys could be excellent working materials for magnetic refrigeration at both room temperature and high temperature. Furthermore, the strategy of two elements codoping and the design model with combing calculation and experiment provide the reference and instruction for designing new magnetic shape memory alloys.

METHODS Calculation Methods All the calculations were performed with the plane-wave pseudopotential method implemented by using CASTEP software53-55, which is based on density functional theory method. We choose the generalized gradient approximation of Perdew, Burke, and Ernzerhof (GGA-PBE) as the exchange-correlation functional56. The structure is optimized with the Broyden-Fletcher-Goldfarb-Shanno (BFGS) method57, and convergence is assumed when the forces on atoms are less than 0.03 eV/Å. The cut-off of the energy plane-wave was set to 400 eV and the MonkhorstPack grid with parameters of 8 × 8 × 8 was used to sample the Brillouin zone58. The interaction between nuclei and electrons was described by ultra-soft pseudopotentials59. The convergence tolerance for the calculations was selected as 1 × 10-6 eV/atom. For the cubic L21 austenite of Ni8xCoxMn6Sn2-yCuy (x = 0, 1, 2, 3; y = 0, 1), 16 atoms unit cell was built with some Ni and Sn atoms replaced by Co and Cu atoms respectively. The tetragonal martensitic phase has been found by calculating the stability of L21 austenitic phase with respect to volume-conserving tetragonal distortions. Namely, the ground state of the martensitic phase can be found through adjusting the c/a ratio. Then, the total energy of both austenitic phase and martensitic phase, electronic structure, the density of state (DOS) and magnetic moments have been calculated.

Figure 1. The unit cell of the cubic structure for parent of the (a) Ni8Mn4Sn4, (b) Ni7CoMn6SnCu alloy. The parent phase of stoichiometric Ni8Mn4Sn4 orders with the cubic L21 structure as shown in Figure 1a, however, the austenite phase with cubic L21 structure cannot transform to the martensitic phase layered structure in the


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Ni2MnSn alloys60. Therefore, we focus on the offstoichiometrical Mn-riched Ni2Mn1.5Sn0.5 with substitution Lattice parameters (nm)

Alloys

x=0, y=0 x=0, y=1 x=1, y=0 x=2, y=0 x=3, y=0 x=1, y=1 x=2, y=1 x=3, y=1

to occupy the Sn sites in the Ni2MnSn alloys. Here, Mn atoms at Mn sites are marked as MnMn, while Mn atoms that

a Cub. Tet. Cub. Tet. Cub. Tet. Cub. Tet. Cub. Tet. Cub. Tet. Cub. Tet. Cub. Tet.

6 5.5 5.9 5.34 6 5.57 6 5.57 6 5.57 5.9 5.5 5.9 5.5 5.9 5.5

c

c/a

7.15

1.3

7.21

1.35

6.96

1.25

6.96

1.25

6.96

1.25

7.15

1.3

7.15

1.3

7.15

1.3

of Ni and Sn by Co and Cu separately. The corresponding crystallographic structure is shown in Figure 1b. As reported in previous research, the excess Mn atoms prefers

Moment (μB) Total

Ni

MnMn

MnSn

2.02 2.07 2.22 1.99 2.27 2.08 6.95 2.13 7.01 2.26 6.76 2.17 6.87 2.37 6.91 2.60

0.19 0.24 0.29 0.19 0.37 0.20 0.66 0.19 0.65 0.19 0.74 0.21 0.75 0.23 0.74 0.23

3.40 3.30 3.34 3.20 3.36 3.29 3.42 3.25 3.36 3.20 3.36 3.23 3.31 3.16 3.25 3.12

-3.54 -3.45 -3.41 -3.26 -3.55 -3.46 3.57 -3.44 3.51 -3.41 3.44 -3.26 3.39 -3.22 3.34 -3.19

Cu

Co

0.03 0.08

0.03 0.08 0.02 0.09 0.02 0.09

1.17 0.63 1.48 0.63 1.42 0.68 1.52 0.7 1.46 0.88 1.41 1.01

substitute Sn atoms in the Sn sites from stoichiometrical Ni2MnSn

Table 1. Equilibrium lattice parameters, total and partial spin moments of the cubic austenite (Cub.) and tetragonal nonmodulated martensite (Tel.) for Ni8-xCoxMnSn2-yCuy (x = 0, 1, 2, 3; y = 0, 1) alloys. alloy are marked as MnSn. Moreover, two situations of the magnetic moments are considered. Hereafter, “AP” represents the antiparallel magnetic interactions between the MnMn and MnSn and “P” represent the parallel magnetic interactions of that.

mechanical property test. The compressive stress-strain curves were tested by an Instron 5569 testing system at room temperature. And the crosshead displacement speed was set to 0.1 mm/min. To study the dominant fracture behaviour, fracture surfaces were scanned by SEM.

Experimental

CALCULATION AND DESIGN

The polycrystalline alloys of nominal composition (at.%) Ni48-xCoxMn37Sn9Cu6 (x = 0, 6, 8, 10, 12) were arc-melted from high purity elements under inert Ar atmosphere. To make sure the alloys consist a uniform composition, the alloys were melted eight times. The cast samples were annealed under pressure of 10-3 Pa at 1073 K in seals quartz tubes for 48 h, followed by quenching in ice water to homogenize the alloys. The microstructure was observed by scanning electron microscopy (SEM) and optical microscope (OM) at room temperature. The composition determination of the phase was obtained using the energy-dispersive X-ray spectrometer (EDS), which is attached to the SEM. Here, each of these results is the average of five values. The crystal structure analyses were determined at room temperature by X-ray diffraction (XRD) using Cu-Kα radiation. The phase transformation temperatures of the alloys were measured by differential scanning calorimeter (PerkinElmer DSC 8000) on modulated mode with a cooling/heating rate of 15 K/min. The samples with the size of φ 3mm × 6 mm were cut from the sample alloys for the

In this part, we conduct a detailed and comprehensive first-principle study of the influence of Co and Cu introduction, investigating the substitution by a single element and the synchronous substitution by both elements. The effect of elements Co and Cu co-doping on TM, TC and magnetic structure is revealed. We computationally designed Ni8-xCoxMnSn2-yCuy (x = 0, 1, 2, 3; y = 0, 1) alloys and predicted an approximate composition range of Ni-Co-MnSn-Cu alloy to realize multi-objective optimization in the magnetocaloric material. Magnetic Properties Usually, a large ΔSM would be expected in the alloys with a large ΔM. To investigate the effect of Co and Cu doping on the structure and magnetic properties in Ni-Mn-Sn alloys, equilibrium lattice parameters, total and partial spin moments of the cubic austenite (Cub.) and tetragonal nonmodulated martensite (Tel.) for Ni8-xCoxMnSn2-yCuy (x = 0, 1, 2, 3; y = 0, 1) alloys are calculated as shown in Table 1.



For purpose of exhibiting the effect of Co or Cu doping visually, the variation of the total energy as a function of the lattice parameter for austenitic type alloys with two magnetic configurations (P and AP states) is presented in Figure 2. Here, we just show four feature alloys: (a) alloy without Co and Cu addition, (b) alloy with a Ni atom replaced by a Co atom, (c) alloy with a Sn atom replaced by a Cu atom, (d) alloy with a Ni atom and a Sn atom replaced by a Co atom and a Cu atom respectively. It is clearly shown in Figure 2a, without of Co/Cu addition, the AP state with lower energy of Ni8Mn6Sn2 is such stable than the P state. When 6.25 at.% Ni is substituted by Co, the curves of AP state and P state are approaching each other around the equilibrium lattice parameter. Meanwhile, a crossover between the two curves (AP and P states) appears when the lattice constant is larger than 6.05 Å. While, the curve of AP state is also lower that the curve of P state around the equilibrium lattice parameter. Similar behaviour can be observed in Figure 2c when 6.25 at.% Sn is substituted by Cu. However, compared to the results shown in Figure 2a and 2b, the difference is that the P state and AP state have nearly the same energy at the ground state lattice constant of ~5.90 Å, and the crossover appears when the lattice constant is smaller than 5.90 Å. Despite the energy of the P and AP states at the ground state lattice constant are similar, the AP state is still a little lower in energy. These phenomena above indicate that the substitution of Co or Cu can enhance the stability of P state. While, in these alloys, antiferromagnetic coupling between MnMn and MnSn remains the dominant. When 6.25 at.% Co and 6.25 at.% Cu co-doped, as shown in Figure 2d, the curves show a marked difference. The AP and P state curves swap places. The P state curve moves down and is overall below the AP state curve, which obviously indicates that Co and Cu co-doped turn the magnetic coupling between the Mn atoms. After Co and Cu co-doping, the ferromagnetic coupling is preferable. It does mean that the alloy shows a magnetic transition from antiferromagnetic to ferromagnetic through the Co and Cu co-doping in the austenite phase.

Figure 2. Calculated total energies as functions of lattice constant for austenitic type alloys (a) Ni8Mn6Sn2, (b) Ni8Mn6SnCu, (c) Ni7CoMn6Sn2, and (d) Ni7CoMn6SnCu in P and AP states. Through the above analysis, Co and Cu mono doping both enhance the magnetic properties of the austenite phase, and their modulation effect is codirectional. Most important of all, Co and Cu co-doping has a synergistic promoting effect

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on the magnetic properties of the austenite phase. What is more, from the four groups of curves, it is clear that, before and after Co or Cu doping, the equilibrium lattice constants have changed. Combing the detailed equilibrium lattice parameters shown in Table 1, Co doping makes the lattice almost no change, and Cu doping results in a slight shrinkage of the lattice. This is due to the radius of the Co atom (1.67 Å) is similar to that of Ni atom (1.62 Å), while the radius of the Sn atom (1.72 Å) is much larger than the radius of Cu atom (1.57 Å). As a result, Co and Cu co-doping results in a slight shrinkage of the unit-cell volume In order to obtain further insight into the magnetic properties of the alloys with Co/Cu doping from their electronic structure, the total densities of state (DOS) are given for alloys Ni8Mn6Sn2, Ni7CoMn6Sn2, Ni8Mn6SnCu, Ni7CoMnSnCu, and Ni6Co2MnSnCu in austenite. The Fermi level (EF) is marked by a vertical line at the energy of 0 eV, which can be seen in Figure 3. According to the energy band model, the difference between the majority spin and minority spin DOS below the EF results in a change of the total magnetic moment. The black line shows the total DOS of the Ni8Mn6Sn2 alloy without anything doping. Compared with it, the majority spin DOS below the EF of alloys with Co/Cu mono-doping (Ni7CoMn6Sn2 and Ni8Mn6SnCu) have a slight increase, whereas the minority spin DOS remains almost unchanged. Moreover, we can clearly see that the total DOS of Ni7CoMnSnCu, and Ni6Co2MnSnCu shows a marked difference from the others. After Co and Cu codoping, the majority spin DOS below the EF increases, but those of the minority spin decreases slightly. The increase of the total magnetic moments is due primarily to this. These results are consistent with calculated results of the total magnetic moments and atoms magnetic moments (listed in Table 1). From Table 1, we compare each set of data one-by-one. Though the magnetic moment of Cu is very small, the total magnetic moment of Ni8Mn6SnCu increases slightly. This is mainly due to the contraction of the lattice, which leads to enhance the interaction between the Mn and Ni atoms and finally increase the total magnetic moment. Ferromagnetic and antiferromagnetic interaction between the Mn atoms is relative to the interatomic spacing. The total magnetic moment of Ni7CoMn6Sn2 also has a slight increase. Contrast with Cu-doping, the enhancement of the magnetism is mainly due to that the partial moment of the Co atom is larger than that of atom Ni. Meanwhile, the magnetic coupling of Co and Mn is parallel. From Table 1, we can find that the doped atom of Co exhibits a relatively large magnetic moment (~1.17 μB) than that of the Ni atom (~0.37 μB). These calculated values are agreeing well with the results in Ref. [61]. When Co and Cu co-doping, the total magnetic moment show a sudden increase from 2.02 μB to 6.76 μB. As mentioned in Figure 2, Co and Cu co-doping enhance the magnetic properties of the alloys in the austenitic phase. And more notably, the magnetic moment of alloys with Co and Cu co-doping has been changed from negative to positive. Namely, Co and Cu co-doping turn the interaction between the Mn atoms from antiferromagnetic coupling to ferromagnetic. This explains the enormous increase in the total magnetic moment. As the further increase of Co-doping content, the total magnetic moments of Ni8-xCoxMnSnCu (x = 0, 1, 2, 3) alloys continue to increase slightly because of the larger magnetic of the Co atom.


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ACS Applied Materials & Interfaces Figure 4. Total and spin projected partial density of states of alloys Ni8Mn6Sn2, Ni7CoMn6Sn2, Ni8Mn6SnCu, and Ni7CoMn6SnCu in austenite. The insets are the MnMn and MnSn 3d partial DOS of each alloys. The arrow indicates the direction of spin polarization.

Figure 3. Total densities of states for alloys Ni8Mn6Sn2, Ni7CoMn6Sn2, Ni8Mn6SnCu, Ni7CoMn6SnCu, and Ni6Co2Mn6SnCu in austenite. The inset is the near EF total densities of states in the energy range marked by the dotdashed rectangle. The Fermi level is indicated by the vertical line at 0 eV Figure 4. shows the total and spin projected partial density of states (PDOS) of alloys Ni8Mn6Sn2, Ni7CoMn6Sn2, Ni8Mn6SnCu, and Ni7CoMnSnCu in austenite. As can be seen from Figure 4a, the Ni, MnMn and MnSn 3d states play a key role in DOS in both majority spin channel and minority spin channel below the EF while MnSn 3d and MnMn 3d states dominate the minority spin channel above the EF. Apparently, we can find that the DOS of Ni 3d states in both majority and minority spin channels are almost totally occupied. This DOS distribution of Ni gives rise to a relatively small magnetic moment of ~ 0.19 μB. Differently, the DOS of MnMn 3d states in majority channel is almost totally occupied but the minority states are empty. DOS of MnMn 3d states in both majority and minority spin channels is split at the EF, which makes a bigger magnetic moment ~ 3.40 μB due to the crystal-field effects. This DOS distribution makes Mn mostly dominate the magnetic properties. Particularly, the MnSn and MnMn 3d partial DOS of each alloys are illustrated in the insets. The arrow indicates the direction of spin polarization. Obviously, the spin direction is opposite between MnSn and MnMn. Compared the four insets, Co/Cu mono-doping cannot change the MnSn 3d state. However, for the alloy with Co and Cu co-doping, the MnSn 3d has been changed. Thus, the spin direction is parallel between MnSn and MnMn, which is the root cause of the total magnetic moment increasing.

Next, we study the influence of Co/Cu doping on the stability of the martensitic phase for both parallel and antiparallel magnetic interactions in Ni8Mn6Sn2. The ground state of the martensitic phase can be found through adjusting the c/a ratio. Here, the volume is assumed to be constant between the austenitic phase and martensitic phase. It is a routine mean used usually in this type Heusler alloys62. The detailed calculation results are shown in Figure 5. Without doping, the AP state of Ni-Mn-Sn is preferred for the whole c/a range. This result indicates that the magnetic coupling between MnMn and MnSn is antiparallel both in the austenitic phase and martensitic phase. When the total energy reaches minimum, the c/a ratio is 1.30. For Co doping at the Ni site, the magnetic interactions between MnMn and MnSn in both the austenitic phase and martensitic phase remain unchanged. When the total energy reaches minimum, the c/a ratio is changed to 1.25. Surprisingly, for Cu doping at the Sn site, the curves of AP state and P state are approaching each other at c/a = 1 (austenitic phase), though the calculated results show that magnetic state of both austenite and martensite remains unchanged. Here, when the total energy reaches minimum, the c/a ratio is changed to 1.30. Eventually, for Co and Cu co-doping, the P state becomes favored energetically around the c/a = 1 (austenitic phase). Whereas, the AP state is more stable when the martensitic transformation occurs, which indicates the strong magnetostructural coupling under the martensitic transformation appears. These results lead to a large ΔM of 4.59 μB, which is why the MFIMT in Ni-Mn-Sn alloys can be obtained with Co and Cu co-doping. Moreover, it demonstrated that giant MCE may be obtained, because a large ΔSM would be expected in the alloys with a large ΔM.

Figure 5. Calculated total energies as functions of the c/a for martensitic type alloys (a) Ni8Mn6Sn2, (b) Ni8Mn6SnCu, (c) Ni7CoMn6Sn2, and (d) Ni7CoMn6SnCu in P and AP states. The stability of P and AP states were compared in the figure. Curie temperature



Ni-Mn-Sn based FSMAs meet the requirement of the Stone condition of ferromagnetism63,64. The total energy difference between the paramagnetic and the ferromagnetic state (ΔE) can offer an estimate of TC 65. On the basis of the Heisenberg model and the molecular field theory, the relationship between ΔE and TC can be shown as 𝛥𝐸 = ― 𝑘𝑏𝑇𝑐𝜉 , (1) where 𝜉 is the ratio M / M0 of the magnetic moment M at T ≠ 0 K and the equilibrium magnetic moment M0 at T = 0 K52. In general, to predict the value, 𝜉 is treated as a constant approximatively66. Though many investigators reported that Tc could be predicted successfully52,66-,68. It should be indicated that there are deviations in the calculated values and the experimental values. Even so, we can obtain the result that the TC is directly proportional to the ΔE. To study the influence of Co/Cu doping on the TC, we calculated the total energies difference ΔE of Ni8-xCoxMn6Sn2 (x = 0, 1, 2, 3) and Ni8Mn6Sn2-yCoy (y = 0, 0.5, 1). It can be found that the total energy of ferromagnetic cubic phase is always smaller than that of the paramagnetic cubic phase in all samples, which is consistent with the fact that the paramagnetic cubic is the high-temperature phase. Figure 6 shows the influence of Co/Cu doping content on ΔE of Ni8-xCoxMn6Sn2 and Ni8Mn6Sn2-yCuy alloys in austenite. Apparently, the increasing of the Cu doping content results in a linear decrease of the ΔE, while the increasing of the Co doping content results in a linear increase of the ΔE. We can draw a conclusion that Co/Cu mono-doping shifts the TC to the opposite directions in Ni-Mn-Sn FSMAs.

between Co and Mn is much higher than that between Ni and Mn, indicating a stronger Co-Mn bounding interaction as compared to the Ni-Mn one. It is due to the introduction of the stronger exchange interaction between Co and Mn through Co doping that the whole exchange interaction of the Ni7CoMnSn system is enhanced, resulting in the increase of the Curie temperature of the alloy.

Figure 7. Charge density of parent phase of Ni7CoMnSn alloy in the (110) plane Martensitic transformation temperature The total DOS plots of (a) Ni8Mn6Sn2, (b) Ni8Mn6SnCu, (c) Ni7CoMn6Sn2, and (d) Ni7CoMnSnCu for both cubic austenitic phase and tetragonal martensitic phase are shown in Figure 8. Moreover, the enlargement of the region near the EF is shown as the inset. It can be seen that the states at the EF are lower in martensitic DOS compared with the austenite. Then the conduction electrons concentration becomes fewer after the phase transition. In the minority spin DOS of the austenite, a sharp peak at the EF can be observed, showing the bonding character between the atoms in the alloys around the EF, which this is deep related to the MT. For the alloy with Co mono-doping, the obvious difference is that the peak at the EF has nearly disappeared in austenite. The DOS of the martensite remains almost unchanged. It indicates that the stability of the austenite is enhanced, and the TM decreases. Combined with the PDOS as shown in Figure 4 mentioned above, it can be found that the DOS peak at the EF can be attributed to the Ni 3d states in the austenite.

Figure 6. Effect of Co/Cu doping content on total energy difference (ΔE) between paramagnetic and ferromagnetic states of Ni8-xCoxMn6Sn2 and Ni8Mn6Sn2-yCuy in austenite. To further reveal the mechanism of the increasing of TC, the electron exchange integral between atoms is discussed. For the electron exchange integral between any two atoms, the more the charge accumulation in the middle region of the two neighboring atoms, the greater the positive contribution in the exchange integral will be, which leads to a large exchange integral (the exchange energy). Charge density can intuitively reflect the situation of charge accumulation. In general, high charge accumulation between neighboring atoms indicates a strong hybridization between the orbitals of the atoms and thus a strong bounding interaction. Figure 7 shows the charge density of Ni7CoMnSn ferromagnetic austenite in the (110) plane. It demonstrates clearly that the charge accumulation

Page 6 of 16

Figure 8. The total DOS plots for the alloys (a) Ni8Mn6Sn2, (b) Ni8Mn6SnCu, (c) Ni7CoMn6Sn2, and (d) Ni7CoMn6SnCu for cubic austenite and tetragonal


Page 7 of 16

martensite phase. The vertical dashed line is chosen as Fermi level EF. The dot-dashed rectangle range near the Fermi level is shown on an expanded scale in the inset. To investigate the influence of Co and Cu doping on MT, the differences in total energy ΔE1 (ΔE1 = EAP(Cub.) - EAP(Tetra.)) between the austenite phase and martensite phase were calculated. Lots of investigators ever reported that total energy can be used to assess the phase stability. Then the conclusion can be obtained that the ΔE1 is critical to the MT behaviour and the larger ΔE1 corresponds to a higher TM 69,70. As shown in Figure 9, it can be seen that ΔE shows a 1 significant decreasing with the increase of Co content. As the relationship mentioned before, a smaller ΔE1 implies a lower TM. This result indicates that substitution of Co for Ni in Ni-Mn-Sn alloy decreases the TM. Moreover, value of ΔE1 becomes negative when x is more than 2. It means that the martensitic transition will not occur when the Co content is more than 12.5 at%, which is due to the TM is too low. However, for Cu doping, it shows the opposite effect on the martensitic transition. The ΔE1 shows a significant increase with increase of Cu content. The substitution of Cu for Sn in Ni-Mn-Sn alloy increases the TM. y=1

0.6

0.4 x=0

y=0 x=1

0.2

0.0

-0.2

achieve simultaneous improvement of TM and TC and ΔM. These researches can help in understanding the distinct properties between Co and Cu during co-doping and provide an insight in to tuning multiple properties effectively.

Figure 10. Effect of Co/Cu doping content on energy difference between austenite L21 structure and the tetragonal martensite structure (ΔE1) and total energy difference between paramagnetic and ferromagnetic states (ΔE) per formula unit for Ni8-xCoxMn6Sn2 and Ni8Mn6Sn2yCuy alloys.

REALIZATION BY EXPERIMENT

y = 0.5

ΔE1（ eV/f.u.（

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


x=2

Ni8-xCoxMn6Sn2 Ni8Mn6Sn2-yCuy x=3

-0.4

Co/Cu content Figure 9. Effect of Co/Cu doping content on energy difference between austenite L21 structure and the tetragonal martensite structure (ΔE1) per formula unit for Ni8-xCoxMn6Sn2 and Ni8Mn6Sn2-yCuy alloys. Calculation prediction From the single-element doping results presented above, what can be obtained is that Co-doped alloys deficient in Ni, TM decreases and TC increases, whereas Cu-doped alloys deficient in Sn, TM increases and TC decreases. As mentioned before, it is critical for improving both TM and TC to high temperatures. Then, we focus on the effect of simultaneous introduction of Co and Cu. Figure 10 exhibits the influence of Co/Cu doping content on the ΔE and the ΔE1 per formula unit for Ni8-xCoxMn6Sn2 and Ni8Mn6Sn2-yCuy alloys. 6.25 at% Cu is doped firstly, then the content of Co doping increases gradually. From the two curves, we can observe the obvious complementary relationship. Co-doping and Cu-doping complement each other perfectly in the effect on TM and TC. Specifically, the results offer an approximate composition range (Ni-Mn-Sn alloys with 6.25 at.% Cu and 6.25 - 12.5 at.% Co) to improve the working temperature Ni-Mn-Sn FSMAs. Next, through a slight composition adjustment may

For now, the influence of Co and Cu co-doping on TM, TC and the magnetic structure was revealed by using firstprinciple calculations. Calculated results offer an approximate composition range of Ni-Co-Mn-Sn-Cu alloy to realize multi-objective optimization. In addition, it may be the appropriate composition range for magnetocaloric refrigeration in these alloys. According to the results of calculations, sample alloys were prepared and investigated. In the next parts, we will mainly concentrate on the microstructure, martensitic transition behaviour, magnetic and mechanical properties of Ni48-xCoxMn37Sn9Cu6 (x = 0, 6, 8, 10 and 12) alloys. Microstructure Figure 11 shows the XRD patterns of Ni48-xMn38Sn9Cu6 (x = 0, 6, 8, 10, 12) alloys at room temperature. In the figure, the black and red dotted lines are used to indicate the diffraction angles corresponding to the diffraction peaks. For Ni48Mn37Sn9Cu6 alloy, the typical tetragonal nonmodulated martensite phase (donated as M) peaks of (222), (400) and (622) can be detected. It indicates that the alloy is a martensitic phase at room temperature. While, in addition to the peaks from martensite, some other diffraction peaks appear at around 49.6°, 73.1° and 87.7°. The new phase can be confirmed as a face-centered cubic (fcc) γ phase (donated as γ) with the main diffraction peaks of (200), (202), and (311) 71,72. Then, we find that the peaks of γ phase can be detected in all samples and the position of γ phase diffraction peaks almost remained unchanged. For Ni42Co6Mn37Sn9Cu6 alloy, the XRD patterns are similar to that of Ni48Mn37Sn9Cu6 alloy. However, when the Co-doping content is 8 at.%, a satellite peak exists in the M(222) peak, which shows the appearance of some austenitic feature in that alloy. Further increase of Co content, the peaks of the martensitic phase disappear and the typical diffraction



peaks of L21 structure appear. Accordingly, we can draw that by substituting Co in Ni site in Ni48Mn37Sn9Cu6 alloy, the room temperature phase changes to austenite, thus indicating the stabilization of the austenite phase with Co doping.

Page 8 of 16

Figure 11. XRD patterns of Ni48-xCoxMn38Sn9Cu6 (x = 0, 6, 8, 10 and 12) alloys taken at room temperature. The dotted line represents the diffraction Angle corresponding to the diffraction peak.

Figure 12. Optical micrographs of solution treated Ni48-xMn38Sn9Cu6 alloys at room temperature (a)x=0; (b)x=6; (c)x=8; (d)x=10; (e)x=12. Optical micrographs of Ni48-xCoxMn37Sn9Cu6 (x = 0, 6, 8, 10, 12) FSMAs at room temperature are shown in Figure 12. The partial magnifications are shown at the upper right. In Figure 12(a - c), the grain boundaries can be observed obviously, which are marked by white dashed lines approximately. Apparently, it can be found that the grain size of samples decreases after introducing element Co. From Figure 12a, it can be seen that the Ni48Mn37Sn9Cu6 alloy consists of coarser equiaxed grains, and the grain size is about 400 μm. A single martensitic phase is observed, which is in plate shape boarded with the inter-plate interfaces. The martensite plates are found to be composed of thin lamellae with different crystallographic orientation (areas with a different colour). Moreover, a small number of granular second phases distribute sporadically, which is

mainly as a result of introduction of Cu in Ni-Mn-Sn alloy. This second phase can be confirmed to be an fcc γ phase according to the above XRD patterns. With the addition of Co, the grain size of the Ni48-xCoxMn37Sn9Cu6 is reduced significantly. When 6 at.% Co doping as shown in Figure 12b, the grain size decreases to ~200 μm. When the content of Co doping is 8 at.%, the grain size is decreased to ~100 μm. With the further doping of Co, the microstructure of the samples shows a significant difference compared to previous sample alloys. It shows the distinctive texture of the γ phase with straight and elongated γ phase grains. Excessive γ phases are brittle, the γ phase particles dispersed in the matrix act as a deterrent to its mechanical properties. From the partial magnifications, it is found that the γ phase and their volume fraction increase gradually


Page 9 of 16

with increasing doping content, then tends to distribute along the grain boundaries. As Co content is more than 8 at.%, the γ phases particles are dispersed in the austenite phase uniformly. As shown in Figure 12e, the γ phase (dark phase) exhibits the largest grain size. Figure 13 shows the BSE images of Ni48Mn37Sn9Cu6 and Ni36Co12Mn37Sn9Cu6 FSMAs. The grey phase was the matrix and the dark phase was precipitate γ phase. Without Co doping, just a small amount γ phase is observed. With 12 at.% Co doping, the amount of γ phase significantly increases. To study the composition of the phases, the corresponding compositional maps of elements Sn, Mn, Cu, and Co contents, respectively. For Ni48Mn37Sn9Cu6 alloy, we can see that the content of Sn and Cu in the γ phase is much higher than that in the matrix. For Ni36Co12Mn37Sn9Cu6 alloy, a remarkable difference between the matrix and γ phase can be seen. The content of Co and Cu in the γ phase is higher, whereas the Sn content is very low. Mn content of γ phase is also slightly higher. The content change in the matrix is the opposite. Table 2 lists the detailed EDS results, and it shows the same results about the elements content.

Elevating the operating temperature is one of the most important objects of our magnetocaloric Ni-Mn-Sn alloy design. in order to exhibit optimum properties, element Ni is substituted by Co in Ni48Mn37Sn9Cu6 FSMA to tune phase transition parameters. Figure 14 shows the DSC curves for the Ni48−xCoxMn39Sn9Cu6 (x = 0, 6, 8, and 10) alloys. Here, we just show the DSC curves of 3 samples which can show the martensitic transition behaviour in the temperature range of 273.5 K – 573.5 K The characteristic temperatures (austenite start (As), austenite finish (Af), martensite start (Ms), martensite finish (Mf) and TC) of the samples are taken as the intersection pint of two base lines, as the arrow shown in Figure 14. While, the peaks indicating TC cannot be exhibited. The characteristic phase transition parameters obtained from the DSC curves are listed in Table 3. Without Co doping, Cu mono-doping shift the TM of Ni-Mn-Sn alloy to relatively high temperature. Then, the phase transition temperatures decrease rapidly with increasing Co doping. When x reaches 8, the phase transition temperature (~ 354.1 K) is brought down to slightly below 100 ℃, which can still meet basically the requirements of high temperature. Co and Cu co-doping can achieve our design goal of tuning TM. Ni48Co0Mn37Sn9Cu6 Ni42Co6Mn37Sn9Cu6 Ni40Co8Mn37Sn9Cu6

100

Heat flow (W/Kg)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Heating 50

0

AS

Cooling

-50 300

Figure 13. Backscattered electron images of Ni48Mn37Sn9Cu6 (a) and Ni36Co12Mn37Sn9Cu6 (b) and their corresponding compositional mapping analysis of elements Sn, Mn, Cu, Co Table 2 The EDS results of Ni48-xCoxMn37Sn9Cu6 (x = 0, 6, 8, 10 and 12) (at.%). phase

Matrix

Second phase

x

Ni

Mn

Sn

Cu

Co

0 6

46.95 41.71

37.94 37.02

9.45 10.55

5.93 5.64

0 5.08

8 10

39.88 37.78

37.00 36.27

10.57 11.34

5.71 5.70

6.83 8.91

12

35.77

37.23

11.57

5.84

9.59

0 6 8

50.04 41.97 37.28

37.05 37.63 38.01

2.98 2.70 3.04

9.93 7.91 9.21

0 9.79 12.47

10

38.09

37.48

2.43

8.93

13.07

12

36.45

37.33

2.37

9.12

14.73

Optimizing of phase transition parameters

350

400

450

500

550

Temperature (K)

Figure 14. The DSC curves of Ni48-xMn38Sn9Cu6 (x = 0, 6 and 8) alloys. The thermomagnetic curves of Ni48-xCoxMn37Sn9Cu6 (x = 0, 6, 8, 10) alloys measured in an external field of 100 Oe by the same extraction method are plotted in Figure 15. Moreover, some characteristic temperatures are also obtained and are listed in Table 3. This shows that TM determined by DSC is consistent with the magnetization measurements. In Figure 15a, the reverse martensitic transition cannot be observed in the whole temperature range, which is mainly due to the Cu mono-doping makes the TM too high above the test temperature range. Also, the changes of magnetization cannot be observed. Namely, for Ni48Mn37Sn9Cu6 alloy, it is martensite phase in the test temperature range and the TM and the TC are not in this range. It is consistent with the DSC results. When 6 at.% Co doping shown in Figure 15b, the alloy undergoes an incomplete reverse martensitic transition from weakmagnetic martensite to the ferromagnetic austenite, accompanied by a jump like variation in magnetization. Owning to the limitation of the test temperature range, the full transition process cannot be shown. At the temperature


ACS Applied Materials & Interfaces of 350 K, a small bump on the magnetization curve, which may correspond to the magnetic transition of the martensitic phase. In Figure 15c, the Ni40Co8Mn37Sn9Cu6 alloy undergoes a complete reverse martensitic transition. With further increase of temperature, the alloy undergoes a transformation to the paramagnetic state at the Curie temperature of austenite (𝑇𝐴𝐶). The 𝑇𝐴𝐶 of 375.3 K is indicated as the white arrow. Apparently, TM shifts to the lower temperature. On current trends, with the further doping of Co, the Ni38Co10Mn37Sn9Cu6 alloy shows the austenitic phase in the whole temperature range due to the too low martensitic transition temperature. In this alloy, the 𝑇𝐴𝐶 Of 392.1 K is also identified as the white arrow shown. The entire analysis results indicate that at the foundation of Cu doping, Co doping decreases the TM and increases the 𝑇𝐴𝐶. Finally, through the Co and Cu co-doping and composition adjustments, the increase of the operating temperature of the Ni-Mn-Sn FSMAs by improving TM and 𝑇𝐴𝐶 to high temperatures simultaneously. All these test results are in line with the results of first-principle calculation.

The shift in transformation temperatures with magnetic field provides a unique opportunity to induce reversible phase transformation at a certain temperature interval. Figure 16 presents magnetization isotherms M-H curves measured at different temperatures for Ni42Co6Mn37Sn9Cu6 alloy. From the M-T curves mentioned above, we can know that the sample alloy is martensite phase when the temperature is 200 K, 300 K and 350 K. But the difference is that the martensite phase under high temperatures exhibits paramagnetic or antiferromagnetic behaviour. While, at the low temperature of 200 K (below the 𝑇𝑀 𝐶 ), the curve shows some characteristics of ferromagnetic behaviour, which is related to the secondary magnetic deformation of martensite. At 390K, which is 1.5K below the As temperature, the sample alloy magnetized to ~5.4 emu/g at ~3.7 T, corresponding to the magnetization of the martensite. Upon increasing the external field to above 3.7 T, the transformation from the martensite to austenite was induced. However, with further increase of the magnetic field, the magnetization can hardly be saturated at 7T. 16

（ a（ Ni48Co0Mn37Sn9Cu6

0.01 0.00 0.06

H = 100 Oe

M

12

（ b（ Ni42Co6Mn37Sn9Cu6

TM C

Transition

0.03 0.00 4

（ c（ Ni40Co8Mn37Sn9Cu6

Transition

TA C

2

2

10 8 6 4

0 6 4

390K 350K 300K 200K

14

M (emu/g)

0.02

M (emu/g)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 10 of 16

（ d（ Ni38Co10Mn37Sn9Cu6

0 200

TA C

A

2 0

300

400

0

Figure 15. Thermomagnetic curves of Ni48xCoxMn37Sn9Cu6 (x = 0, 6, 8, 10) alloys measured in an external field of 100 Oe by the same extraction method. (a), Ni48Co0Mn37Sn9Cu6; (b), Ni42Co6Mn37Sn9Cu6; (c), Ni40Co8Mn37Sn9Cu6; (d), Ni38Co10Mn37Sn9Cu6. M represents the martensite, and A represents the austenite. Table 3 The characteristic temperatures of Ni48xCoxMn37Sn9Cu6 (x = 0, 6, 8, and 10) denoted as Ms, Mf, As, Af, and TC respectively. The values marked by “ * ” are obtained in the thermomagnetic curves. Alloys

As (K)

Af (K)

Ms (K)

Mf (K)

x=0 x=6

492.2 391.5 349.5 341.7*

511.5 403.9 354.1 353.7*

480.1 379.6 343.6 342.7*

456.1 363.9 333.8 328.9*

x=8 x=10

Enhancing the Magnetic Properties

TC (K)

375.3* 392.1*

20000

40000

60000

H (Oe)

T (K)

Figure 16. M-H curves of Ni42Co6Mn37Sn9Cu6 alloy at different temperatures. Figure 17 presents magnetization isotherms M-H curves measured at different temperatures for Ni40Co8Mn37Sn9Cu6 alloy. At 300 K, the magnetization shows nearly linear dependence on the applied magnetic field, corresponding to the magnetization of the paramagnetic or antiferromagnetic martensite. M-H curves measured in 335 K ≤ 347 K show the ferromagnetic behaviour and the magnetization is found to increase with increase in temperature. Though the magnetization occurs a rapid increase within the low magnetic field, the magnetization can hardly be saturated with further increase in the magnetic field. Such behaviour is attributed to antiferromagnetically coupled Mn magnetic moments in the martensitic state when Mn atoms occupy the Sn sites in offstoichiometric Ni-Mn-Sn Heusler alloys73, 74. The content of Co and Cu addition just weakens but not completely suppress the antiferromagnetic interaction between MnMn atoms, which results in a short-range ferromagnetic order in the austenite phase of the sample alloy. Moreover, metamagnetic transition and hysteresis loss can be seen in this temperature range. It should be noted that these behaviours add a considerable contribution to ΔSM, and lead to giant MCE75.


Page 11 of 16

which agrees well with the calculation results mentioned above. Figure 19. The maximum magnetic entropy changes ΔSpeak as a function of temperature and magnetic field for Ni40Co8Mn37Sn9Cu6 alloy (a) and the values of ΔSpeak as a function of the corresponding temperature in the magnetic field changes of 1 T, 2 T, 5 T and 7 T for the prevailing MCE materials such as Ni-Mn based alloy, LaFeSi, Gd-Si-Ge and so on (b) 6-15,19, 21, 24-28,33,34,76,77. The data are taken from the present work and literature (see Supplemental Material, Note 1 for details).

Figure 17. M-H curves of Ni40Co8Mn37Sn9Cu6 alloy at different temperatures.

In order to investigate the MCE, we calculated the temperature of magnetic entropy change (ΔSM) from the MH curves at different temperatures using the Maxwell relationship:

40

ΔSpeak (J/kgK)

100

M (emu/g)

80

60

40 390K 350K 300K 200K

20

0

30

Field change 1.0 T 2.0 T 5.0 T 7.0 T

20000

40000

10

0

（ b（

With further increase in measured temperature, the magnetization of the austenite phase is found to decrease, as shown in the Figure 18. This is primarily because the atomic magnetic moment orientation is gradually destroyed by thermal motion. In the whole temperature range, the sample alloy is in the austenitic phase and it shows a typical response of ferromagnetic material where the magnetization increases quickly and saturates at the low fields. The reverse demagnetization path overlapped with the forward magnetization. Though the maximum saturation magnetization is determined to be 94 emu/g, there is no sign of a metamagnetic transition up to 7 T. Comparing all the sample alloys, the saturation magnetization of the austenite phase shows a monotonic increase with increase in Co content. This indicates that Co and Cu addition enhances the magnetic interactions and increases the overall magnetization in the austenite phase,

30

10

0


（ a（ 345

350

355

360

320

360

𝐻

∂𝑀(𝐻,𝑇)

0

∂𝑇

)

400

440

𝑑𝐻 𝐻

Improving the Mechanical Properties

1.0 T 3.0 T 5.0 T 7.0 T

20

280

Figure 19a shows the temperature dependence of calculated ΔSM in the vicinity of martensitic transformation. And a large value of ΔSM (34.8 J/kg K) was obtained at 344K under the magnetic field of 7 T. In order to compare the MCE of present samples with that of other currently prevailing magnetocaloric materials, the maximum entropy change Δ 𝑆𝑝𝑒𝑎𝑘 and the corresponding temperatures are 𝑀 schematically illustrated in Figure 19b. As is shown, most of the magnetocaloric materials distribute on the left half, whose optimum temperature of use is at and below room temperature. At higher temperatures, the values of Δ𝑆𝑝𝑒𝑎𝑘 𝑀 almost less than 20 J/kg K. Here, in the present work, Ni40Co8Mn37Sn9Cu6 alloy shows the largest Δ𝑆𝑝𝑒𝑎𝑘 under a 𝑀 magnetic field of 7 T at the high-temperature range. Even though under a magnetic field of 5 T, the Δ𝑆𝑝𝑒𝑎𝑘 is still a 𝑀 relatively high value.

Ni40Co8Mn37Sn9Cu6

Temperature (K)

240

∫(

40

340

200

Temperature (K)

H (Oe)

335

(the present work)

60000

Figure 18. M-H curves of Ni38Co10Mn37Sn9Cu6 alloy at different temperatures.

-10

Ni40Co8Mn37Sn9Cu6

20

160 0

ΔS (J/kgK)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


ACS Applied Materials & Interfaces The excellent mechanical properties of these Ni-Mn-Sn alloys are indeed quite beneficial for their applications. To determine the effect of Co and Cu co-doping on mechanical properties in Ni48-xCoxMn37Sn9Cu6 (x = 0, 6, 8, 10 and 12) FSMAs, compression tests were carried out at room temperature. The samples were loaded until fracture in compression as shown in Figure 20a. Ni48Co0Mn37Sn9Cu6 Ni42Co6Mn37Sn9Cu6

1200

Stress (MPa)

3.Co8Cu6

Ni40Co8Mn37Sn9Cu6 Ni38Co10Mn37Sn9Cu6

1000

Ni36Co12Mn37Sn9Cu6

800

4.Co10Cu6

600

0

1.Co0Cu6

0

2

4

6

Figure 20. The compressive stress-strain curves of Ni48xMn38Sn9Cu6 (x = 0, 6, 8, 10 and 12) alloys at room temperature (a) and the maximum values of the maximum compressive stress and strain for the prevailing MCE materials (b). 26, 36,77-81 It is found that the compressive strength and strain are all enhanced obviously by doping Co in Ni-Mn-Sn-Cu FSMAs. Without Co doping, the Ni48Mn37Sn9Cu6 alloy has the worse mechanical properties among the tested samples. The fracture stress value is only 203.0 MPa and the strain value is only about 5.3%. Apparently, the compressive ductility of Ni48Mn37Sn9Cu6 FSMA is improved through doping of Co. Strikingly, when the doping content of Co is 8 at.%, the compressive strength reaches 1072.0 MPa and the

5.Co12Cu6

2.Co6Cu6

400

200

（ a（ 8

10

12

14

Strain (%) Ni-Mn based

1400

LaFeSi

GdAlCoZr

MnCoGe

Ni40Co8Mn37Sn9Cu6

1200

Stress (MPa)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 12 of 16

1000

(the present work)

800 600 400 200 0

（ b（ 0

2

4

6

8

Strain (%)

10

12

14

Figure 21. SEM fractography of Ni48-xMn38Sn9Cu6 (x = 0, 6, 8, 10 and 12) alloys at room temperature(a)x=0; (b)x=6; (c)x=8; (d)x=10; (e)x=12.


Page 13 of 16 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


compressive fracture strain even reaches 11.9%. The mechanical properties of this alloy with Co and Cu codoping are greatly enhanced than that of the alloy with Cu mono-doping. When the content of Co doping further increases, the compressive strength and strain decreased significantly. It is also worth noting that the alloys with 10 at.% and 12 at.% are austenitic phases, but other alloys are martensitic phases initially. This is one reason for causing the decrease of the compressive strength. In conclusion, all of these results show that the appropriate amount of Co doping improves the compressive stress and the strain of Ni48Mn37Sn9Cu6 FSMA significantly. In order to compare the mechanical properties of present samples with that of other currently prevailing magnetocaloric materials, the maximum values of their compressive stress and strain are schematically illustrated in Figure 20b. As is shown, the Gd-Al-Co-Zr show a largest stress but its strain is worst. The La-Fe-Si and Mn-Co-Ge system show poor mechanical properties. Relatively, NiMn-based alloys are less brittle. Even more importantly, the compressive stress and strain obtained in the present work are the largest for the Ni-Mn-based magnetocaloric materials reported so far. Figure 21 shows the compressive fracture surface of Ni48xCoxMn37Sn9Cu6 (x = 0, 6, 8, 10 and 12) FSMAs. The alloy without Co doping, as shown in Figure 21a, the typical brittle fracture characteristic can be confirmed through the intergranular fracture pattern. And the martensite phase with lamellar twin and the granular γ phase can be seen. We can find that the cracks propagated along the grain boundaries. At the same time, cracks will bypass the γ phases, when they ran into small and circular γ phases. As shown in Figure 21b and c, it can be seen visually that the density of alloys is enhanced with further increasing Co doping content. This result is in line with the aforementioned analysis from OM. As we all know the HallPetch relation: 𝛿𝑦 = 𝛿𝑖 +𝐾𝑑 ―1/2, the smaller the average grain size, the better the yield strength. As a consequence, grain size can affect the mechanical properties of the alloys. Here, grain size of Ni48Mn37Sn9Cu6 alloy decreases with the increase of Co content. Hence, the above relation is further demonstrated that the refinement of grain size plays a positive role in the mechanical properties. Moreover, when the content of Co increases to 6 at.% and 8 at.%, in Figure 21b-c), the fracture is of characteristics cleavage fracture with dimples and tearing ridges appearing. Particularly, for Ni40Co8Mn37Sn9Cu6 alloy, the amount of γ phase increases and the shape becomes more and more irregular. Thus, the cracks will through γ phase directly. Therefore, crack growth requires more energy, which will result in the formation and growth of the cracks would be damped. For all these reasons and more, the Ni40Co8Mn37Sn9Cu6 alloy exhibits the best mechanical properties. With the further doping of Co, it can be seen that the fracture surface of Ni38Co10Mn37Sn9Cu6 and Ni36Co12Mn37Sn9Cu6 alloy show the different characteristics. On one hand, more amount of γ phase appeared, and it is brittle phase. On the other hand, too many doping of Co shifts the TM too low, and Ni38Co10Mn37Sn9Cu6 and Ni36Co12Mn37Sn9Cu6 alloy are austenite phases. These may lead to the decreased mechanical properties of alloys when Co content is more than 10 at.%.

CONCLUSIONS In conclusion, to realize the multi-objective optimization of Ni-Mn-Sn magnetocaloric material, a new approach of Cu and Co co-doping is proposed. However, the effect of these two elements is interactive and restrictive. To balance them and find the appropriate content of co-doping in a narrow window is always difficult, but very essential. Therefore, we first make a deep investigation on the influence of Cu and Co co-doping by using first-principle calculation. The mechanism of co-doping is revealed, and the results offer a composition range of Ni-Co-Mn-Sn-Cu alloy. Ultimately, according to the results of calculations, a giant ΔSM of 34.8 J/kg·K and excellent mechanical properties (1072.0 MPa stress and 11.9 % strain) are simultaneously achieved at the high temperature of 344 K. Particularly, the temperature of use can be further improved through this method we presented. This work proves the great potential of Ni-MnSn FSMAs for the application in magnetic refrigeration at high temperature with giant MCE and outstanding mechanical performance. We believe that this strategy of multi-doping and the design model with combing calculation and experiment provide the reference and instruction for designing high-performance magnetocaloric materials and extend their application range.

AUTHOR INFORMATION Corresponding Author * (Changlong Tan) E-mail: [email protected].

Author Contributions The manuscript was written through contributions of all authors. / All authors have given approval to the final version of the manuscript. / ‡These authors contributed equally.

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT This work was supported by the National Natural Science Foundation of China (No.51871083 and No. 51471064). This work was carried out at LvLiang Cloud Computing Center of China, and the calculations were performed on TianHe-2.

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Computation-guided design of Ni-Mn-Sn ferromagnetic shape memory

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