2021 Volume 62 Issue 1 Pages 130-134
Ferromagnetic Mn–Al–C (τ-phase) can be synthesized by a single-route conventional reactive sintering method. The maximum magnetization and coercivity were 75.8 Am2/kg and 57 mT, respectively. The τ-phase fraction was evaluated to be 81 mass% for Mn55Al45C2 annealed at 1273 K. The τ-phase of Mn55Al45C2 can be synthesized at an annealing temperature from 873 to 1273 K, whereas that of Mn55Al45 cannot be synthesized. It was indicated that the phase stability of the hcp-phase (ε-phase) was improved by adding carbon, resulting in an ε–τ transformation.
Fig. 4 Ta dependence of M1.5T of Mn55Al45 and Mn55Al45C2.
L10-Mn–Al (τ-phase) is a Mn-based ferromagnetic alloy, which has received much attention as a candidate of rare-earth-free hard magnetic materials. The τ-phase has a large saturation magnetization Ms of approximately 2.4 μB/f.u. at 0 K, uniaxial magnetic crystalline anisotropy energy Ku of 1.5 MJ/m3, and Curie temperature TC of 655 K.1–3) The ferromagnetic τ-phase is reported to be metastable and obtained by massive transformation from ε-phase with a disordered hcp structure.4) Because the τ-phase is metastable, the Ms of the entire sample reduces because of the dissolution of τ-phase to a non-ferromagnetic stable γ2- or β-phase during the aging treatment. Therefore, there has been much reporting on improving the magnetic properties5–9) since the discovery of τ-phase in 1958.5) Among the studies in recent years, it has been suggested that a slight change in the composition of the ε-phase significantly influenced the fraction of the τ-phase.8)
The hard magnetic properties of Mn–Al–C have also been evaluated because the addition of carbon stabilized τ-phase.7,10–19) Anisotropic Mn–Al–C magnets were prepared by the hot extrusion technique, with residual magnetization Mr = 93 Am2/kg, and the coercivity μ0Hc = 0.3 T.10) In addition, gas atomized alloyed powder with ε-phase as the main phase were used at extrusion technique, for improving the maximum energy product (BH)max.11,12) Most of these studies showed that the synthesis of the τ-phase was due to the ε–τ transformation.
Quite recently, the synthesis of the τ-phase without an ε–τ phase transformation by high-pressure synthesis20) and electrodeposition21) were reported. In this study, we focused on reactive sintering method. Herein, a sintering method was performed to prepare a bulk Mn–Al–C from alloyed powder,15,19) but this was not reactive sintering for synthesizing the τ-phase. Reactive sintering is a method that the reaction of the powder of raw materials and the sintering proceed simultaneously. At first, the reacted phase is synthesized by diffusion at the boundary of the particles. The reaction proceeds inside the grain by using the heat generated by the exothermic reaction.22) Disadvantage and advantage of the reactive sintering are as follows. When no loading or no constraint is applied at reactive sintering, reduction of the density may occur. The reduction of density leads to the reduction of the volume magnetization and (BH)max. Meanwhile, the advantage of the reactive sintering method is single-route synthesis from a raw material. If the τ-phase can be synthesized without melting process, shortening of the processes of the synthesis and lowering the synthesis temperature are expected. In this study, the conventional single-route reactive sintering method was applied to synthesize the τ-phase for Mn–Al and Mn–Al–C.
Nominal composition of Mn55Al45Cx (x = 0.0–6.0) pellets of approximately 3 g were prepared by pressing under a pressure of 15 MPa after mixing high-purity Mn (99.9%, grain size d < 75 µm), Al (99.99%, d < 75 µm) and C (99.99%, 2 < d < 10 µm) powders. When x was changed from 0 to 6, the chemical formula was expressed from Mn55Al45 to Mn51.89Al42.45C5.66. The obtained pellets were sealed in a quartz tube with Ar gas. Afterwards, reactive sintering was carried out for 48 h, followed by cooling to room temperature for 12 h. For investigating the mechanism of phase formation of τ-phase, we confirmed whether ε-phase was synthesized or not at annealing temperature Ta. Therefore, the quenched sample was prepared by the annealing at 1073 K for 48 h, followed by quenching into the iced water. It is noted that neither loading nor constraint on a pellet was applied during the reactive sintering process. Ta was 873 ≤ Ta ≤ 1273 K. Because the melting point of Al is 933 K, the sintering at Ta ≥ 973 K was a liquid-phase reactive sintering, and the others were solid-phase sintering. According to the Al–Mn phase diagram,21,22) the ε-phase is not an equilibrium phase at Ta ≤ 1173 K.
Powder X-ray diffraction (XRD) at room temperature (RT) was carried out to characterize the reacted phases. The phase fraction was evaluated by Rietveld analysis using PDXL software (Rigaku). The magnetic properties were evaluated by vibrating sample magnetometer (VSM) in magnetic fields μ0H ≤ 1.5 T at RT. Because the magnetization of the τ-phase did not saturate even at 1.5 T, magnetization at 1.5 T M1.5T was used for evaluating the saturation magnetization.
Figure 1 and Table 1 show photos of the obtained Mn–Al–C pellet at Ta = 1073 K and the size of the pellet, respectively. The volume of the reacted pellet was three times larger than that before the annealing. The annealed Mn–Al–C pellets were brittle and of low density. The density of Mn–Al–C prepared by hot deformation or microwave sintering was reported to be 4.9–5.2 g/cc,19) which is about 1.3 times larger than what was attained in the present study. The low density of the pellet is due to the weak connection of the particles compared with the melting method. In addition, unreacted Mn powder at liquid-phase reactive sintering led to the brittleness of the pellets. During the cooling, a volume change of ∼1% at the ε–τ transformation was probably the reason for the expansion of the pellet. Transformation from an hcp to L10 structure resulted in many voids in the samples and expansion of the pellets.
Photos of Mn55Al45C2 pellet. (a), (b) Bottom and lateral views of the sample before annealing. (c), (d) Bottom and lateral views of the sample after annealing.
Figure 2 shows the XRD patterns at RT for Mn55Al45 (a) and Mn55Al45C2 (b) at various Ta. At Ta below 973 K, unreacted Mn, γ2-phase and β-phase were observed in both Mn55Al45 and Mn55Al45C2 samples. When Ta increased, the ratio of the β-phase increased and the γ2-phase decreased in Mn55Al45 and Mn55Al45C2. According to the Al–Mn phase diagram,23,24) the phase boundary between the β-phase and (β + γ2) two-phase region shifts to Al-rich composition for T > 873 K, whereas that between the γ2-phase and (β + γ2) two-phase region is almost unchanged. Therefore, the change of the phase fractions was consistent with the phase equilibrium based on the lever rule. It should be noted that the τ-phase was also observed for Mn55Al45C2 at Ta > 1073 K. The fraction of the τ-phase increased with higher Ta. The maximum fraction of τ-phase was 81 mass% for Mn55Al45C2 at Ta = 1273 K. This result suggests that the τ-phase can be obtained by one-step reactive sintering. The lattice constants of the τ-phase were evaluated to be a = 2.77 Å, and c = 3.57 Å, regardless of heat treatment temperature, which is comparable to Refs. 17 and 25).
XRD patterns of Mn55Al45 (a) and Mn55Al45C2 (b) annealed at 873 ≤ Ta ≤ 1273 K.
Figure 3 shows the magnetization curves (M-H curves) of Mn55Al45 and Mn55Al45C2 at Ta = 873 K (a), 973 K (b), 1073 K (c) and 1273 K (d) at RT. The inset in Fig. 3(d) is a magnified view of the M-H curve of Mn55Al45 at Ta = 1273 K. Although the magnetization was found to be very small, but it can be surely seen from the inset of Fig. 3(d) that the Mn55Al45 annealed at Ta = 1273 K was ferromagnetic. On the other hand, the samples annealed at Ta ≤ 1173 K did not exhibit ferromagnetic properties. According to the Al–Mn phase diagram,23,24) the ε-phase is an equilibrium phase at Ta = 1273 K. This result implied that a very small amount of the ε-phase was synthesized by the reaction at 1273 K, and then a similarly small amount of τ-phase was formed during slow cooling. However, the rest of the part of the sample is occupied by the precipitated non-ferromagnetic β- and γ2-phases during the slow cooling, resulting in the quite small magnetization. Considering the M1.5T, the phase fraction of the τ-phase in the sample was less than 1 mass%, which cannot be detected in an XRD pattern. Meanwhile, the M-H curves show ferromagnetic properties at Ta ≥ 873 K in the Mn55Al45C2 sample. With increasing Ta, M steeply improved and M1.5T = 75.8 Am2/kg was obtained at Ta = 1273 K. Figure 4 shows the Ta dependence of M1.5T for Mn55Al45 and Mn55Al45C2. In Mn55Al45C2 at Ta = 1073 K, the magnetization was drastically enhanced. The obtained M1.5T at Ta ≥ 1173 K showed high magnetization over 70 Am2/kg and became saturated. However, the M1.5T in Mn55Al45 showed a constant and almost non-ferromagnetic value. It is noted that the ε-phase was not an equilibrium phase in the Al–Mn phase diagram at 1073 K. Therefore, Mn55Al45 did not show ferromagnetic properties, whereas the phase stability of the ε-phase changed by carbon addition.
M-H curves of Mn55Al45 and Mn55Al45C2 for Ta = 873 K (a), 973 K (b), 1073 K (c) and 1273 K (d).
Ta dependence of M1.5T of Mn55Al45 and Mn55Al45C2.
Figure 5 shows the x dependence of M1.5T at Ta = 1073 K. M1.5T = 32.5 Am2/kg at x = 4, which is the highest value among the samples annealed this temperature. At the present sintering condition, the effect of carbon content x on the synthesis of the τ-phase was saturated at x = 4. According to Pareti et al., the effect of x on the magnetic properties of the τ-phase by carbon addition was saturated at x = 2.1) Therefore, it is suggested that the magnetization of the τ-phase did not increase, but the phase fraction of the τ-phase increased up to x = 4. As described in the following section, synthesis of the τ-phase in this study is due to the ε–τ transformation. According to reports on the phase stability of the ε-phase,26–28) carbon addition enhanced the phase stability of the ε-phase. In the Al–Mn–C phase diagram, the area of the ε-phase expanded for x = 2–4 at maximum for both Mn-rich and Al-rich regions.27,28) Therefore, it is suggested that the carbon-addition-induced stabilization of the ε-phase is x ∼ 4 at maximum at 1073 K, resulting in the enhancement of the ε–τ transformation. For x > 4, because of the instability of the ε-phase, the fraction of the τ-phase reduced.
x dependence of magnetization at Ta = 1073 K.
Figure 6 shows the XRD patterns of Mn55Al45C2 annealed at 1073 K, followed by slow cooling (a) and quenching (b). When the sample was slow-cooled, the τ-, β-, and γ2-phases were observed. Meanwhile, the γ2-, β-, and ε-phases were obtained for the quenched sample. Therefore, τ-phase was not directly synthesized from raw materials but was transformed from ε-phase.
XRD pattern of Mn55Al45C2 prepared by slow-cooling (a) and quenching (b) from 1073 K.
Herein, the addition of carbon to the phase stability of ε- and τ-phase was discussed. We recently evaluated phase stability of ε-phase by carbon-addition.26) In Mn55Al45, ε-phase was not synthesized at 1073 K. With increasing carbon content, the fraction of the precipitated ε-phase increased and 66 mass% of ε-phase was obtained by adding 2 at%C. However, ε-phase was not precipitated at Ta = 973 K even adding 2 at%C. These are consistent with the quite small magnetization of the sample annealed at Ta = 973 K and rapid increase of magnetization at Ta = 1073 K, which was shown in Fig. 4. Meanwhile, according to the ab-initio calculation, when carbon atoms occupied interstitial sites of L10-type structure, τ-phase stabilized.29) Although both τ- and ε-phase was stabilized by carbon addition, it is suggested by thermal analysis for quenched ε-phase that ε-phase becomes stabilized against τ-phase by carbon addition.30,31) Consequently, carbon addition on the reactive sintering can be explained as below. Whether ε-phase was formed or not played a key role for the formation of τ-phase. The stabilization of τ-phase by carbon addition contributed to the suppression of the dissolution from τ-phase to β- and γ2-phase.
The ferromagnetic Mn–Al–C phase was synthesized by reactive sintering with a single route annealing process. The maximum magnetization of the sample is 75.8 Am2/kg. The τ-phase can be obtained for a carbon-doped sample by annealing at 1073 K. However, the τ-phase was not obtainable in Mn–Al in this condition. These results indicated that the phase stability of the ε-phase was improved by addition of carbon, leading to the synthesis of the τ-phase by the ε–τ transformation.
This work was partly supported by KAKENHI (grant no. 16H04547). Magnetization measurements were performed at the Institute for Materials Research, Tohoku University. R.K. is grateful to the JSPS Research Fellowships for Young Scientists.
