Magnetic and Structural Properties of MnCoGe with Minimal Fe and Sn Substitution

Abstract

The magnetic and structural properties of MnCo_0.92Fe_0.08Ge_1−xSn_x (x = 0.02, 0.05, and 0.1) were investigated. A first-order magnetic transition was observed, which was accompanied by the martensitic transformation from a hexagonal Ni₂In-type to an orthorhombic TiNiSi-type structure in the vicinity of 320 K for x = 0.02 and 170 K for x = 0.05 with thermal hysteresis of approximately 20 K. The magnetization of MnCo_0.92Fe_0.08Ge_0.95Sn_0.05 was estimated to be 102 Am² kg⁻¹ in 5 T at 10 K. High field magnetization curves up to 55 T for MnCo_0.92Fe_0.08Ge_0.95Sn_0.05 indicated the field-induced martensitic transformation. Mössbauer spectroscopy experiments revealed that Fe atoms occupied 79%–85% in the Co-site and 21%–15% in the Mn-site.

1. Introduction

Ferromagnetic (FM) functional materials, which show martensitic transformation (MT), have attracted attention as magnetic field-controlled materials, because some of them exhibit a large magnetic field-induced strain,^1–3) magnetocaloric effect⁴⁾ and, magnetoresistance effect.⁵⁾ FM MnCoGe-based compounds are one of the growing groups of magnetic functional materials. The martensite phase (M-Phase) is characterized by the stoichiometric MnCoGe ferromagnet with a Curie temperature T_C of 355 K, which crystallizes in an orthorhombic TiNiSi-type structure at room temperature (RT).⁶⁾ The compound transforms without diffusion from the M-phase to a parent phase (P-phase) with a hexagonal Ni₂In-type structure at MT temperature T_M of 398–458 K.⁶⁾ This MT exhibits a large volume change (3.9%).⁶⁾ The magnetic moment m of MnCoGe was reported to be 3.9 μ_B for the M-phase and 2.6 μ_B for the P-phase.⁶⁾

It has been reported that T_C and T_M of MnCoGe-based compounds were controlled by substituting Al,⁷⁾ V,^8,9) Cr¹⁰⁾ and Fe^11–18) for Mn or Co. Other controls include off-stoichiometric composition^19–23) and interstitial modification,²⁴⁾ etc. In particular, the Fe-substituted MnCoGe compound showed interesting magnetic and structural properties.^11–18) MnCo_1−yFe_yGe with y ∼ 0.08^12,16) or y = 0.22¹³⁾ showed a first-order magnetic transition (FOMT), which was accompanied by MT from the paramagnetic (PM) P-phase to the FM M-phase. MnCo_0.94Fe_0.06Ge exhibits FOMT and a large change in entropy (ΔS ∼ −27 J kg⁻¹ K⁻¹) at ∼315 K.¹³⁾ In our previous reports, the phase diagram of MnCo_1−yFe_yGe was shown (Fig. 1),¹⁶⁾ and MnCo_0.92Fe_0.08Ge showed FOMT from the PM to FM state, and this was accompanied by MT in the vicinity of 275 K.^17,18) The cell volume expanded by 4.1% during MT,¹⁷⁾ and ΔS was estimated to be −14 J kg⁻¹ K⁻¹.¹⁸⁾ Magnetic field-induced MT from the PM P-phase with low m to the FM M-phase with high m in MnCo_0.92Fe_0.08Ge was measured at 276 K in magnetic fields from H up to μ₀H = 5 T.¹⁷⁾

Fig. 1

Phase diagram of MnCo_1−yFe_yGe (ref. 16). T_C and T_M indicate the Curie temperature and martensitic transformation temperatures, respectively. Dotted and dashed lines are to guide the eye.

However, as seen in Fig. 1, this phenomenon was observed for a very small range of Fe-content y for MnCo_1−yFe_yGe,¹⁶⁾ and the T of FOMT and MT could not be controlled sufficiently for application. It is known that T_C and T_M are also controlled by substituting Al,²⁵⁾ Ga²⁶⁾ and Sn²⁷⁾ for Ge. To assess the potential of the Mn(CoFe)Ge-based compounds for application, it is important to study the magnetic and structural properties of Mn(CoFe)Ge-based compounds systematically.

Our purpose in this study was to clarify the influence of substituting Sn for Ge on the magnetic and structural properties of MnCo_0.92Fe_0.08Ge.

2. Experimental Procedure

Polycrystalline ingots of MnCo_0.92Fe_0.08Ge_1−xSn_x (x = 0.02, 0.05 and 0.1) were prepared by arc-melting a mixture of nominal amounts of pure elements (Mn, 99.9%; Co, 99.9%; Fe, 99.99%; Sn, 99.99%; Ge, 99.999%) in an argon atmosphere. The obtained ingot was turned over and re-melted several times. The ingot was annealed at 1123 K for 120 h in a quartz tube with a vacuum and then slowly cooled to RT for 10 h. X-ray powder diffraction (XRD) measurements were made using Cu-Kα radiation at RT.

Differential scanning calorimetry (DSC; NETZSCH) was carried out in an N₂ atmosphere in the temperature T range of 120–400 K. The heating and cooling rates for DSC measurements were 10 K min⁻¹. T dependence of the XRD profile was measured using Cu-Kα radiation for 10 ≤ T ≤ 290 K.²⁸⁾ The powder was fixed with vacuum grease on a copper sample holder.²⁸⁾ Magnetization M measurements were performed using a superconducting quantum interference device (SQUID) magnetometer (Quantum Design) with a T range of 10–370 K and in magnetic fields from H up to μ₀H = 5 T. High-field M in the powder sample was measured using an induction method in pulsed magnetic fields up to μ₀H = 55 T for 100 ≤ T ≤ 220 K. The pulse width was 50 ms. ⁵⁷Fe Mössbauer spectroscopy experiments with a 1.85 GBq ⁵⁷Co(Rh) source were performed using a conventional constant acceleration method at RT and 10 K. The velocity scale was calibrated with α-Fe, which had a hyperfine field B_hf of 33.1 T and offset value of the isomer shift δ of 0.1 mm s⁻¹. The values of the hyperfine parameters were refined using NORMOS.²⁹⁾

3. Results and Discussion

Figure 2 shows the XRD patterns of MnCo_0.92Fe_0.08Ge_1−xSn_x (x = 0.02, 0.05 and 0.1) at RT. Here, hkl_h and hkl_o denote the Miller indices for the hexagonal P-phase and the orthorhombic M-phase, respectively. No extra diffraction peak was detected in the sample. For x = 0.05 and 0.1, MnCo_0.92Fe_0.08Ge_1−xSn_x showed a single phase of the P-phase. The lattice parameters and unit-cell volume are listed in Table 1. Here, a_h, c_h and V_h denote the lattice parameters and unit-cell volume for the P-phase. For x = 0.02, the two-phase coexistence of the P- and M-phase was detected at RT. In Table 1, a_o, c_o and V_o denote the lattice parameters and unit-cell volume of the M-phase. The XRD pattern of x = 0.02 is similar to the previous result for MnCo_0.92Fe_0.08Ge (x = 0).¹⁶⁾ In the P-phase, c_h and V_h increase by substituting Sn for the Ge atom. The axes and volumes of the two structures are related accordingly: a_o = c_h, b_o = a_h, c_o = $\sqrt{3}$a_h and V_o = 2V_h.⁶⁾ The volume difference ΔV = (V_o − 2V_h)/(2V_h) between the P- and M-phase for x = 0.02 was evaluated to be ΔV = 7.8% at RT. This volume change is greater than that for MnCo_1−yFe_yGe (ΔV = 5.9%)¹⁶⁾ at RT. Therefore, when MT occurs in MnCo_0.92Fe_0.08Ge_1−xSn_x, a large volume change is expected.

Fig. 2

X-ray powder diffraction patterns for MnCo_0.92Fe_0.08Ge_1−xSn_x at room temperature.

Table 1 Lattice parameters and unit-cell volume of MnCo_0.92Fe_0.08Ge_1−xSn_x at room temperature.

Figure 3 shows the DSC curves obtained for MnCo_0.92Fe_0.08Ge_1−xSn_x with x = 0.02 (a), 0.05 (b) and 0.1 (c). For x = 0.02, exothermic and endothermic peaks with a thermal hysteresis (TH) were apparent during cooling and heating within the T range of 290–370 K, which corresponded to MT and reverse transformations (RMT),¹⁷⁾ respectively. Based on the DSC peaks for x = 0.02, the MT starting and finishing T, M_s and M_f, and the RMT starting and finishing T, A_s and A_f were determined, and are listed in Table 2. Here, the values for M_s, M_f, A_s and A_f were defined by the intersection of the baseline and the tangent line at the greatest angle between the peaks. The DSC curves for x = 0.05 indicated that the MT and RMT between the P- and M-phase occurred at 177–208 K. In addition to these large peaks, small exothermic and endothermic peaks were detected at 260 K in the P-phase, which correspond to T_C, since no TH was seen.¹⁸⁾ The DSC curves for x = 0.1 indicated that MT did not occur and T_C was 257 K in the P-phase. The determined T_C are also listed in Table 2.

Fig. 3

DSC curves for MnCo_0.92Fe_0.08Ge_1−xSn_x compounds where x = 0.02 (a), x = 0.05 (b) and x = 0.10 (c).

Table 2 Curie temperature T_C, the martensitic transformation starting and finishing temperatures M_s and M_f, and the reversed transformation starting and finishing temperatures A_s and A_f for MnCo_0.92Fe_0.08Ge_1−xSn_x. The values were determined using DSC measurements.

Figure 4 shows the typical XRD profiles of MnCo_0.92Fe_0.08Ge_0.95Sn_0.05 (x = 0.05) at several temperatures for the cooling (a) and heating (b) processes. In this figure, the XRD profiles are shown, with the peak of the copper 111 reflection (43.2° ≤ 2θ ≤ 43.8°) removed, as it corresponds to the diffraction from the copper sample holder.⁵⁾ During the cooling process, strong Bragg peaks of the P-phase were only observed above 210 K. The peaks of the M-phase started to appear and develop with decreasing T from 170 K (∼M_s). In contrast, the peak intensity of the P-phase became smaller with decreasing T. However, the XRD profiles showed that the P-phase remained at 10 K. This may have been caused by partial distortion at low T resulting from the pulverization of the original specimen into a powder sample and/or by the suppression of the large volume expansion due to the frozen grease on the sample holder.^17,20) With increasing T, the intensity of the M-phase decreased above 170 K (∼A_s) and was not detected over 210 K (∼A_f). The peaks of the P-phase were only seen at 250 K. These characteristic temperatures are consistent with M_s, M_f, A_s and A_f as determined by the DSC measurement. The T dependence of the lattice parameters and cell volume is shown in Fig. 5. The value of a_o increased by 11.6% along the c_h axis at 170 K with decreasing T, while the other lattice parameters shrank by 0.1% (c_o/√3) and 6.3% (b_o) along the a_h axes, leading to a cell volume change of 4.4%. This volume change was in good agreement with that (4.1%) reported previously for MnCo_0.92Fe_0.08Ge (4.1%) at 276 K.¹⁷⁾

Fig. 4

X-ray diffraction patterns of MnCo_0.92Fe_0.08Ge_0.95Sn_0.05 at various temperatures during cooling (a) and heating (b) processes from 290 to 10 K.

Fig. 5

Temperature dependence of the lattice parameters and cell volume of MnCo_0.92Fe_0.08Ge_0.95Sn_0.05. The circles and squares indicate the data for the hexagonal (parent phase) and orthorhombic (martensite phase) structures, respectively. The open and closed symbols indicate the data for the heating and cooling processes, respectively.

The T dependence of the M (M-T curve) for x = 0.02 (a), 0.05 (b) and 0.1 (c) in μ₀H = 1 T is shown in Fig. 6. The M-T curves for x = 0.02 showed a FOMT between the FM and PM state with a TH in the vicinity of 320 K. Considering the XRD pattern (Fig. 2) and DSC curves (Fig. 3(a)) for x = 0.02, the observed TH was due to the MT between the P- and M-phase. That is, the TH indicates the two-phase coexistence of the P-phase with low M and M-phase with high M. The M of x = 0.02 was 107 Am² kg⁻¹ at 1 T and 10 K, which corresponds to 3.6 μ_B per formula unit. This value is slightly smaller than that of the M-phase for MnCo_0.92Fe_0.08Ge.¹⁷⁾ The M-T curve for x = 0.1 was similar in appearance to Brillouin function curves with a second order magnetic transition. The M and m of x = 0.1 were 80.0 Am² kg⁻¹ and 2.7 μ_B per formula unit. These values are consistent with previous reports for the P-phase.³⁰⁾

Fig. 6

Temperature dependence of the magnetization of MnCo_0.92Fe_0.08Ge_1−xSn_x compounds where x = 0.02 (a), 0.05 (b) and 0.1 (c) at 1 T and 5 T.

For x = 0.05, FOMT with a TH of approximately 20 K was observed at 150 ≤ T ≤ 200 K in the FM state. The FOMT of x = 0.05 originated from the MT between the P-phase with low M and M-phase with high M. M was about 102 Am² kg⁻¹ at 1 T and 10 K, which corresponds to 3.4 μ_B per formula unit. This m is smaller than the results found in neutron diffraction experiments (2.9 μ_B/Mn and 1.0 μ_B/Co)³¹⁾ and electronic structure calculations (2.98 μ_B/Mn and 0.78 μ_B/Co)³⁰⁾ for the M-phase of MnCoGe. When a μ₀H of 5 T was applied, FOMT slightly shifted toward a higher T. This result suggests that MnCo_0.92Fe_0.08Ge_0.95Sn_0.05 shows magnetic field-induced MT.

The high-field magnetization (M-H) curves of MnCo_0.92Fe_0.08Ge_0.95Sn_0.05 in μ₀H up to 55 T at several T for 100 ≤ T ≤ 220 K are shown in Fig. 7. Here, the M-H curves were measured after zero-field cooling from 270 K (P-phase). The M process with hysteresis was measured at 140 ≤ T ≤ 200 K. This T range is consistent with the results of the M-T measurement. The M-H curves at T = 100 K (< M_f) and 220 K did not show hysteresis. The M-H curves for 100 ≤ T ≤ 220 K suggested that the magnetic field-induced MT from the P-phase with low M to M-phase with high M occurred due to high magnetic fields. M and m were estimated to be 120 Am² kg⁻¹ and 4.0 μ_B/f.u. at 55 T and 170 K, respectively. This m was in good agreement with the results of neutron diffraction experiments (2.9 μ_B/Mn and 1.0 μ_B/Co)³¹⁾ and electronic structure calculations (2.98 μ_B/Mn and 0.78 μ_B/Co)³⁰⁾ for the M-phase. Therefore, the P-phase at 170 K almost transformed to the M-phase with a high μ₀H of 55 T. It is noted that M at 55 T and 100 K was slightly smaller than at 140 K and 170 K. This may be due to the increase in magnetocrystalline anisotropy in the M-phase.

Fig. 7

High-field magnetization curves of MnCo_0.92Fe_0.08Ge_0.95Sn_0.05 in magnetic fields up to 55 T at several temperatures from 100–220 K.

Our results showed that the m of MnCo_0.92Fe_0.08Ge_0.95Sn_0.05 increased, and was accompanied by an MT with a large volume expansion (ΔV = 4.4%). This is probably due to the suppression of the 3d-3d overlap after MT. According to the result of the calculation,³⁰⁾ the band structure in orthorhombic MnCoGe showed narrower 3d-band widths, a large exchange splitting between the majority and minority bands, and a decrease in the density of states at the Fermi level, which causes an enhancement of m and a decrease in magnetic energy compared with that of hexagonal MnCoGe. The observed MT with a large volume change around M_s probably occurred to minimize the magnetic and elastic energies in FM MnCo_0.92Fe_0.08Ge_0.95Sn_0.05. When H is applied to MnCo_0.92Fe_0.08Ge_0.95Sn_0.05, the decrease in free energy during the M-phase with high m is greater than for the P-phase with low m because of the addition of Zeeman energy. Therefore, the compound exhibited an increase in MT (FOMT) temperature by applying μ₀H of 5 T and magnetic field-induced MT from the P-phase with the low m to the M-phase with the high m. In MnCo_0.92Fe_0.08Ge_0.95Sn_0.05, the difference in M (ΔM) between the P- and M-phase around M_s-M_f was not large, so high H was required to induce the MT.

MnCo_1−yFe_yGe with y ∼ 0.08^12,16) or y = 0.22¹³⁾ showed FOMT between the PM (P-phase) and the FM state (M-phase). However, Sn-substituted for Ge in MnCo_1−yFe_yGe with y = 0.08 did not exhibit this phenomenon. It is hard to control the magnetic properties of this system through low H. This is a disadvantage for the application of magnetic field-controlled materials.

Figure 8 shows the Mössbauer spectra of MnCo_0.92Fe_0.08Ge_0.95Sn_0.05 for the P-phase at RT (a) and the M-phase at 10 K (b). The spectra were fitted with two sub-spectra (Co- and Mn-site) using hyperfine interaction parameters: site-occupation, δ, quadruple splitting Q and hyperfine fields B_hf. The determined parameters are summarized in Table 3. Fe atoms occupied 79%–85% of the Co-site and 21%–15% of the Mn-site in the P- and M-phase. This result is in agreement with previous results for the P-phase of MnCo_0.88Fe_0.12Ge¹⁸⁾ and for M-phase Mn(Co_0.96Fe_0.04)Ge with ⁵⁷Fe.¹⁴⁾ According to the Mössbauer experiments by Ren et al. for (Mn_0.96Fe_0.04)CoGe, Fe atoms also occupied 66% of the Co-site and 34% of the Mn-site. This, combined with our results, indicates that Fe atoms are distributed on both sites but prefer to occupy the Co-site in the MnCoGe-based compound.¹⁴⁾ That is, it may be difficult to control the magnetic and structural properties precisely through Fe substitution for MnCoGe.

Fig. 8

Mössbauer spectra and sub-spectral fits for MnCo_0.92Fe_0.08Ge_0.95Sn_0.05 at room temperature (parent phase) (a) and 10 K (martensite phase) (b).

Table 3 Hyperfine interaction parameters of MnCo_0.92Fe_0.08Ge_0.95Sn_0.05.

The B_hf were estimated to be B_hf-Co = 19 T for the Co-site and B_hf-Mn = 7 T for the Mn-site in the M-phase at 10 K. B_hf-Co was larger than B_hf-Mn, which was consistent with the results of Ren et al.¹⁴⁾ but was different from the results deduced by the neutron diffraction experiments (2.9 μ_B/Mn and 1.0 μ_B/Co).³¹⁾ According to Ren et al., the Fe atom on the Co-site has 6 Mn nearest-neighbors, so that the transferred B_hf is larger than that on the Mn-site (4 Mn nearest-neighbors).¹⁴⁾ The Q values of both sites changed from positive values in the P-phase to negative values in the M-phase, indicating that the charge distribution around the Mn- and Co-sites was modified drastically according to the MT. Considering the lower symmetry in the M-phase, this modification of the charge distribution probably affects the magnetic anisotropy.

4. Conclusions

The structural and magnetic properties of MnCo_0.92Fe_0.08Ge_1−xSn_x (x = 0.02, 0.05 and 0.1) were investigated using XRD, DSC, M measurements and Mössbauer spectroscopy in the T range of 10–400 K. The results for x = 0.02 indicated that FOMT between the PM P-phase (hexagonal) and FM M-phase (orthorhombic) occurred in the vicinity of 320 K for x = 0.02. The compound with x = 0.05 was FM with the T_C = 260 K and showed FOMT between the FM P-phase (low M) and FM M-phase (high M) in the vicinity of 170 K. The compound with x = 0.1 had P-phase for 10 ≤ T ≤ 400 K, and its T_C was 257 K. High field magnetization curves up to 55 T for x = 0.05 indicated that the field-induced MT from the P-phase (low M) and FM M-phase (high M) occurred at 140 ≤ T ≤ 200 K. The results of Mössbauer spectroscopy experiments showed that Fe atoms occupied 79%–85% of the Co-site and 21%–15% of the Mn-site.

Acknowledgments

Temperature dependence of the XRD patterns was measured at the High Field Laboratory for Superconducting Materials, Institute for Materials Research, Tohoku University. Magnetization measurements using a SQUID magnetometer and high field magnetization measurements using a pulse-magnet were performed at the Materials Design and Characterization Laboratory and the International MegaGauss Science Laboratory, the Institute for Solid State Physics, and The University of Tokyo, respectively. Mössbauer spectroscopy experiments were carried out at the Division of Isotope Science, Research Support Center, Institute for Research Promotion, Kagoshima University.

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