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One step towards MnAl-based permanent magnets - Differences in magnetic, and microstructural properties from an intermediate annealing step during synthesis
Journal of Solid State Chemistry 274 (2019) 229–236
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One step towards MnAl-based permanent magnets - Differences in magnetic, and microstructural properties from an intermediate annealing step during synthesis Samrand Shafeie a, Hailiang Fang a, Daniel Hedlund b, Axel Nyberg b, Peter Svedlindh b, Klas Gunnarsson b, Martin Sahlberg a, * a b
Department of Chemistry – Ångstr€om Laboratory, Uppsala University, Sweden Department of Engineering Sciences, Uppsala University, Box 534, SE-751 21, Uppsala, Sweden
A R T I C L E I N F O
A B S T R A C T
Keywords: Powder diffraction MnAl Permanent magnets High temperature synthesis Magnetic measurements
The influence of an additional annealing step during synthesis on the preparation of MnAl based permanent magnet alloys has been investigated. Bulk samples of Mn55Al45C2 alloys were synthesized using induction heating through drop synthesis from 1400 C. Samples produced using cooling directly from 1400 C (from the melt), and from 1400 C to an intermediate annealing step at 1200 C for ~30 min before cooling were compared with respect to differences in phase purity, microstructure and magnetic properties. We found that the phase purity was significantly enhanced using the route with an intermediate annealing step at 1200 C. From XRD the relative content of the τ-phase was improved from ~91 wt% for the sample cooled directly from 1400 C to ~95.1–99.5 wt% for the sample exposed to an intermediate annealing step before cooling. Additionally, EBSD, and SEM with EDS indicates a clear difference in the phase composition and differences in the distribution of the magnetic τ-phase and the non-magnetic ε-, β-, and γ2-phases. Magnetic properties also indicate an improvement in saturation magnetization for the sample exposed to the extra annealing step during synthesis. Our results suggest that an intermediate annealing step in the production of MnAl based alloys will provide a simple way of achieving better phase purity and magnetic properties in the bulk alloy.
1. Introduction The τ-phase with the ordered tetragonal L10 type structure in the MnAl-system has been reported as a suitable rare-earth free permanent magnet material [1]. From theory, the calculated saturation magnetization Ms ¼ 597 kA/m and magnetocrystalline anisotropy K1 ¼ 1:7 106 J/m3, yield a theoretical maximum for the energy product of ðBHÞmax 110 kJ/m3 [1,2]. Together with a magnetic ordering temperature of Tc 650 K, these properties indicate that the material has the potential to replace ferrites and epoxy resin bonded Nd–Fe–B magnets for applications (e.g., permanent magnet motors and wind turbine generators). Based on the ðBHÞmax , which is usually referred to as the “figure-of-merit” for permanent magnet materials, the MnAl-based permanent magnets fill an important gap between the Nd–Fe–B based and the ferrite based permanent magnets [3]. Previous studies have reported the stabilization of the τ-phase by small additions of carbon (~2 at.%) [4], which also increases the usability of this material.
Due to its comparably low cost [1] in relation to its magnetic properties, it becomes a viable material worth optimizing for certain applications where it may replace expensive high performance Nd–Fe–B magnets. It is therefore important to find ways to optimize its magnetic properties. In our previous studies [5,6] we reported the direct formation of the τ-phase from drop synthesis and the possibility to create a homogeneous τ-phase powder through cryomilling with subsequent heat treatments [7]. Ideally a permanent magnet should have a square shaped hysteresis loop to optimize ðBHÞmax with large values for the saturation magnetization and the magnetic coercivity. Permanent magnet alloys like MnAl require a grain morphology with size comparable to the single domain limit as well as shape that promotes alignment of the grain easy axes along a common direction [8]. A first step to move forward in this direction with the MnAl system is to find a synthesis and processing approach that will optimize the purity of the magnetic phase. Here, the high magnetocrystalline anisotropy of the τ-phase is a necessary, but not
* Corresponding author. E-mail address: [email protected] (M. Sahlberg). https://doi.org/10.1016/j.jssc.2019.03.035 Received 7 December 2018; Received in revised form 11 March 2019; Accepted 17 March 2019 Available online 20 March 2019 0022-4596/© 2019 Elsevier Inc. All rights reserved.
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sufficient, prerequisite for the τ-phase to be used as a permanent magnet. Further optimization and processing is needed to render the grain sizes smaller to increase the magnetic remanence of these material. In the MnAl-system, grains with the metastable τ-phase are commonly obtained through two main routes: 1) cooling from the ε-phase at a cooling rate ~10–20 C/min [9], or 2) quenching from the ε-phase followed by annealing at 300–600 C [10]. The phase diagram of MnAl shows that the ε-phase is stable in the temperature region 900–1200 C depending on the Mn:Al ratio. Fast cooling from the melt to the temperature region where the τ-phase starts to form will likely lead to the nucleation of a thermodynamically stable phase with the highest melting point (Mn-rich) followed by the formation of the phase with a slightly lower melting point (Al-rich). At high temperatures, δ-Mn may act as the first phase that nucleates, followed by the ε-phase and the γ2-phase. In the solid state, these phases will likely not dissolve into the main ε-phase. The γ and the δ-Mn phases most likely transform into γ2 and β-Mn phases, respectively [11]. The latter two phases may also form during the ε → τ transformation where the metastable τ -phase breaks down into γ2 and β-Mn phases at temperatures < 800–900 C if sufficient undercooling is not achieved [9]. To inhibit the initial nucleation of high-melting point Mn-rich phases that leave behind an Al-rich matrix with a lower melting point, fast cooling from the melt should be avoided. Thus, to obtain a more homogeneous and phase pure τ-phase, an intermediate step with annealing in the ε-phase region to homogenize the solid should be used before cooling to room temperature to initiate the ε → τ transformation. The formation of the τ–phase is characterized by the presence of crystalline defects like twin and anti-phase boundaries [3,12]. Both defects provide nucleation sites as well as pinning potentials for domain walls, and strongly influences the magnetic properties of the material [3]. Other groups have reported improvements in the synthesis of high purity τ-phase using strip casting [13], fast cooling from the melt with subsequent annealing between 650 and 350 C [10,14], or annealing at 1050 C in air for several days after arc-melting of ingots [15]. However, few of these may be cost-efficient for large scale production of high purity τ-phase MnAl permanent magnets. Although, other preparation methods produce similar saturation magnetization values (see e.g. Refs. [16,17]), our method also shows the importance of transforming from a homogenized solid rather than from the liquid melt to avoid the formation of secondary phases with higher melting points e.g., Mn-rich phases that shifts the bulk composition of the melt. Here, we report a significant improvement in the synthesis procedure for MnAl type materials through induction melting. We present the different microstructural and magnetic properties achieved through the change in the synthesis procedure of drop synthesized MnAlC-material with the composition Mn55Al45C2. From the Mn55Al45C2 ingots produced through drop synthesis either by 1) cooling directly to room temperature from the melt at 1400 C (hereafter referred to as MAC1400) or by 2) cooling first to 1200 C and annealing for ~30 min prior to cooling to room temperature (hereafter referred to as MAC1200). The microstructure and magnetic properties are reported for the two different routes together with the change in the inter-grain ε-phase content based on the two heating procedures.
~30 min before cooling to room temperature (MAC1200). The top part of the MAC1200 ingot is included in the study for comparison to investigate a possible influence of the cooling rate on the material properties. 2.2. Sample preparation Sample preparation for scanning electron microscopy (SEM) was made on pieces from the synthesized ingots. The pieces were put in a conducting polymer resin, and were polished down to 0.25–1 μm. For electron backscatter diffraction (EBSD) the pieces were polished down to < 50 nm in surface roughness using an Al2O3 suspension containing Al2O3 nanoparticles (<50 nm) as the final polishing step. For the microstructural analysis, elemental mapping and surface magnetic properties, the samples were cut into cubes with the dimensions 5 5 10 mm3. For EBSD analysis, separate samples were polished and taken out of the polymer resin in order to minimize image drifting effects from charge buildup during acquisition. Samples were also demagnetized using a AC demagnetizer which operates with a frequency of 32 kHz with a field smaller than 4 Oe (320 A/m) by moving the sample in and out of the demagnetizer. This was done to ensure that there was as little excess stray field buildup on the sample as possible. 2.3. Diffraction X-ray powder diffraction (XRD) was made on crushed pieces from the ingots that were crushed in an agate mortar together with ethanol. The XRD data was collected on a Bruker D8 diffractometer equipped with a Lynx-eye position sensitive detector using Cu Kα1 radiation (λ ¼ 1.540598 Å) and a Bruker D8 Advance equipped with a Lynx-eye energy dispersive position sensitive detector using Cu Kα radiation. The obtained powder diffraction data was refined using the Rietveld method [19] (FULLPROF software [20]) to verify the structural parameters and to calculate the weight percentage of the respective phases. The background was fit using a Chebychev polynomial function with a total of 6 parameters, in addition, the cell parameters, occupancies, and peak shape were refined with a total of 13 parameters. The additional phases were included in the final refinement to estimate relative weight fractions for the different phases. 2.4. Magnetic properties Temperature dependent magnetization (M) versus temperature (T) measurements were performed using a 7400 Lake Shore VSM equipped with a high temperature furnace. The temperature dependence of the magnetization was measured while decreasing the temperature from 750 K to 300 K in a constant applied magnetic field of 80 kA/m. A physical properties measurement system (PPMS) from Quantum Design was used for room temperature high field measurements. Powder samples were enclosed in separate capsules and magnetic hysteresis loops were measured in the magnetic field range of 7162 kA/m. Atomic force/magnetic force microscopy (AFM/MFM) was used to study the topography and magnetic microstructure of the samples. The polished samples were analyzed in a Digital Instruments Dimension 3100 microscope equipped with a low-coercivity magnetic tip. Post-processing of the images were done in the software Gwyddion [21] and are available in the supplementary information.
2. Experimental 2.1. Synthesis
2.5. Microstructural properties MnAl-based ingots with the nominal composition Mn55Al45C2 were prepared by induction melting through drop synthesis of high purity Mn (99.999%), Al (99.999%) and C (99.999%) as described elsewhere [18]. For this study, however, two ingots of approximately 20 g were produced through different annealing procedures: 1) the MnAlC melt was kept at 1400 C for ~5–10 min before the furnace was shut off and cooled to room temperature, (MAC1400); 2) the MnAlC melt was kept at 1400 C for ~5–10 min before lowering the temperature to 1200 C and kept for
The microstructure and the overall elemental mapping were obtained using a high resolution FEG Zeiss LEO1550 scanning electron microscope (SEM) equipped with an Oxford instrument energy dispersive X-ray spectrometer (EDS). For SEM and EDS an acceleration voltage of 20 kV was used. Microstructural information related to grain orientation and phase analysis were obtained using a high resolution FEG ZEISS Merlin equipped with a Nordlys Max detector for electron back scatter 230
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diffraction (EBSD) analysis and the AZtec software for X-ray mapping and elemental analysis. For EBSD analysis, an acceleration voltage of 25 kV was used.
Table 1 Estimated average weight fractions of the different phases for MAC1400, MAC1200 middle, MAC1200 top. Sample
3. Results MAC1400 MAC1200 middle MAC1200 top
3.1. XRD From Fig. 1a it is found that the MAC1400 sample consists of a majority of τ-phase. However, small residues of β-Mn and γ2-Mn phases are present, as summarized in Table 1. From the Rietveld refinements ~91.0 wt% is τ-phase. From the results shown in Fig. 1b and c it is estimated that the middle and top parts of the MAC1200 ingot consist of ~95.1 wt% and ~99.5 wt% τ-phase, respectively. From Table 1, the secondary phases (β-Mn and γ2-Mn) are found to decrease significantly after the additional heat treatment in the MAC 1200 sample. The largest decrease is observed in the top part of MAC 1200 (see Table 1). Tables with detailed information from the refinements can be found in supplementary information. A slight change in peak width (full width at half maximum, FWHM) is observed for the MAC1200 samples taken from the middle and top parts of the ingot, which could indicate differences in residual strain between the samples.
Phase content (wt. %)
τ
γ2
β
91.0(4) 95.1(6) 99.5(2)
8.6(4) 4.8(2) 0.5(2)
0.4(1) 0.2(1) –
large relative amount of the τ-phase (~73%) based on the indexed diffraction pattern of different regions. Also, the main phase detected at the grain boundaries can be indexed as the ε-phase (~22%) similar to what we have observed in our previous study [7] where the Al-rich ε-phase is expelled to the grain boundaries. Small amounts of γ2 and β-Mn phases are detected (~4.5% and ~0.5%, respectively). EBSD from different regions of the MAC1400 sample, however, indicate that the size of the ε-phase rich regions vary strongly over the sample and that the ~22% obtained from indexing the EBSD data is probably overestimated (powder diffraction yields 91% τ-phase in the sample). In MAC1400 where the ε-phase is present between the grains, it appears that the average grain orientation (of the ε-phase) is the same or nearly the same for the whole ε-phase region. The ε-rich regions are, however, very small in the MAC1200 samples. For the middle region in the MAC1200 ingot (see Fig. 7), the amount of τ, ε-, γ2-, and β-phases can instead be estimated to ~89%, ~1.1%, ~2.2% and ~1.4%, respectively. The rest are nonindexed points. From the EBSD maps of the MAC1400 and the MAC1200 samples, the average grain size appears to be larger for the MAC1200 sample (≳ 50 μm2), while the MAC1400 sample appears to have smaller grains (<50 μm2). The fewer but larger grains in MAC1200 appear in comparison to MAC1400 to be crystallographically more oriented. Additionally, while the majority of the inter-grain phase is identified as the ε-phase in the MAC1400 sample, it appears to be dominated by the γ2-phase and to a smaller extent by the β-Mn and ε-phases in MAC1200. Overall, the β-phase appears to form at the grain boundaries of the larger τ-phase grains, while the ε- and γ2-phases form continuous regions between the larger τ-phase grains.
3.2. SEM From the SEM images and the elemental mapping analysis of the MAC1400 and MAC1200 samples a clear difference is observed in the microstructure of the two samples (see Fig. 2). For the MAC1400 sample (see Fig. 2a), the backscatter image indicates brighter regions (typically indicative of higher average atomic weight) between the large grains. However, from the elemental map using EDS (see Fig. 3), these brighter regions can be connected with Al-rich phases, thus indicating possible electron scattering effects in the backscatter image (see arrows Fig. 2a). The grains, however, are rich in both Mn and Al, with C evenly distributed over the whole region. The C content is less reliable due to the use of carbon-based polymer resin that may redeposit on the surface of the sample and influence the C map. Compared to the MAC1200 sample (Fig. 2b), the difference between the Al rich regions and the Mn-rich regions are more distinct, and clearly indicate the existence of an Alrich matrix around the Mn-rich grains in the MAC1400 sample. Furthermore, Fig. 2 indicates a microstructure consisting of larger grains for the MAC1200 sample. From the EDS maps (see Figs. 4 and 5), fewer regions of Al-rich phases are observed for MAC1200 compared to MAC1400. Similar microstructure and distribution of Al rich phases are observed for both the top and middle regions of the MAC1200 ingot. However, the amount of Al rich phases is lower in the top part than in the middle, in accordance with the XRD analysis (Table 1).
3.4. Magnetometry Room temperature magnetic hysteresis loops of the MAC1400 and MAC1200 samples are shown in Fig. 8 and the results are summarized in Table 2. From the extracted data the coercivity (Hc ) follows the trend MAC 1400 > MAC 1200 top > MAC 1200 middle, while the saturation magnetization follows the opposite trend. The MS of the middle part, 622 kA/m is close to the theoretical value of 597 kA/m [1]. In addition, magnetization versus temperature results are shown in Fig. 9 for the MAC1400 and MAC1200 middle part samples, indicating that the Curie temperature defined by the peak temperature in the ∂M= dT versus temperature curve is largest (560 K) for the MAC1400 sample.
3.3. EBSD EBSD data from the MAC1400 sample (see Fig. 6) clearly indicates a
Fig. 1. XRD of ground pieces from a) MAC1400, b) MAC1200 middle part and c) MAC1200 top part (λ ¼ 1.540598 Å). 231
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Fig. 2. Electron backscatter image of a) MAC1400 sample and b) MAC1200 sample taken from the middle of the ingot.
Fig. 3. (a) SEM image of the MAC1400 sample together with EDS mapping of Al (b), Mn (c) and C (d).
Fig. 4. (a) SEM image of the MAC1200 middle part together with EDS mapping of Al (b), Mn (c) and C (d).
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Fig. 5. (a) SEM image of the MAC1200 top part together with EDS mapping of Al (b), Mn (c) and C (d).
Fig. 6. MAC1400 (cooled from melt), where the different figures correspond to a) secondary electron image of the area of interest, b) pattern quality map (band contrast) that indicates the grain boundaries and diffuseness of the diffraction, c) the grain orientation map (IPF Z) and d) a map over indexed phases (red ¼ β-Mn, blue ¼ ε - MnAl, yellow ¼ γ MnAl and green ¼ τ - MnAl). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
the melt at ~1400 C, before cooling directly to room temperature (MAC1400), and the sample where an additional hold at 1200 C for ~30 min before cooling to room temperature (MAC1200) was implemented. From the structural refinement of the XRD data (see Fig. 1), the amount of secondary equilibrium phases (e.g., β-Mn and γ2 phases) decreases significantly when the additional annealing step at 1200 C for ~30 min was implemented. Also, the XRD peak width (FWHM) increases significantly from 0.20 to 0.34 , using a Lorentzian fit of the strongest peak of the τphase ((101) at ~41 in 2θ), suggesting increased internal strain. The improved phase purity for the sample taken from the top part of the MAC1200 ingot can be attributed to the difference in cooling rate near the middle and the top parts of the ingot. More specifically, the middle of the ingot is expected to have a slightly lower cooling rate, which is expected to favor the formation of the undesired secondary
3.5. AFM/MFM From AFM/MFM analysis of MAC1400, a difference between the Alrich and the Mn-rich regions is clearly observed (see Fig. 10 (a) and (b)). By correlating the MFM to the EDS-mapping (Figs. 2 and 3), the Mn rich regions clearly show magnetic contrast unlike the Al-rich regions. The same response is observed in all samples, where the Mn-rich and the Alrich regions are observed, indicating that the same behavior is present in both samples (cf. SI). 4. Discussion Clear differences in the crystallographic and magnetic properties were observed between the sample prepared by conventional heating to
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Fig. 7. MAC1200 near middle regions (annealed at 1200 C for 30min), where the different figures corresponds to a) secondary electron image of the area of interest, b) pattern quality map (band contrast) that indicates the grain boundaries and diffuseness of the diffraction, c) the grain orientation map (IPF Z) and d) a map over indexed phases (red ¼ βMn, blue ¼ ε - MnAl, yellow ¼ γ - MnAl and green ¼ τ MnAl). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 9. Magnetization (M) and ∂M=∂T versus temperature for MAC1400 and MAC1200 middle part samples. H ¼ 80 kA/m. Fig. 8. Room temperature magnetic hysteresis loops for MAC1400, MAC1200 middle part and MAC1200 top part samples.
peak width, however, may relate to the internal strains that are created during the transformation from the ε-phase [22]. It has been reported that the transformation occurs through a massive transformation [23] that takes place through the nucleation of the τphase at the interface of the εphase, thus, requiring a large number of interfaces for the transformation to occur. The interfaces that remain after the transformation has been completed, and are likely strained, may give rise to the strong peak broadening that is observed. In addition, twin formation is expected when a phase transition goes from a high symmetry phase to a low symmetry phase, and such twin formation has been reported in the MnAl-system previously [3,12]. From the EBSD data, we observed that the small crystallite domains in MAC1400 increase in size in the MAC1200 sample, which would lead to the conclusion that the crystallite size may not be the main contributor to the peak broadening. Also, from SEM images and EDS mapping (see e.g., Figs. 2 and 3), we find that the Al-rich regions decrease significantly in size in MAC1200 compared to MAC1400. This is possibly a result of the time allowed for the equilibration to occur in the temperature region where the εphase is stable, before cooling.
Table 2 Saturation magnetization (Ms), coercivity (Hc), magnetic remanence (Mr), magnetic transition temperature (Tc) and XRD phase purity.* was measured in a magnetic field of 800 kA/m and show similar behavior as the others, however, the transition is more smeared out due to 10 times higher field.
MAC1400 MAC1200 top part MAC1200 middle part
Ms/(kA/ m)
Hc/(kA/ m)
Mr/(kA/ m)
TC/ (K)
XRD phase purity/(wt%)
561 606
56 20
106 60
560 545*
91.0(4) 99.5(2)
622
19
60
550
95.1(6)
equilibrium phases (e.g., ε-, γ2-, and β-phases). We have previously shown that the cooling rate is crucial for the phase purity of the τ-phase, and that the cooling rate achieved by air cooling appears to be closest to the optimal (~20 K/min) [5]. The XRD
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Fig. 10. AFM/MFM topographical image MAC1400 (a) and magnetic image MAC1400 (b).
5. Conclusions
In general, this method will circumvent the possibility of the liquid melt to first crystallize in a Mn-rich phase with higher melting point, before crystallizing out the Al-rich phase in between with the residual elements from the melt, as can be deduced from the phase diagram [11]. From the phase diagram, the ε-phase is still the equilibrium phase at ~1200 C, and it can be assumed that a prolonged annealing at 1200 C will ensure a more homogeneous ε-phase with the desired composition prior to the transformation to the τ-phase. In contrast to the untreated MAC1400 sample that has been used as a baseline sample with significant amounts of secondary phases the MAC1200 samples show that the method significantly improves the phase purity through the equilibration step at 1200 C. However, the method still requires optimization since the two different parts of the MAC1200 ingot that we studied in more detail clearly shows that cooling rate differences are still part of the challenge in obtaining phase pure τ-phase across the whole ingot. From the magnetization measurements, we can conclude that TCis lowest for the MAC1200 sample (545–550 K). As seen from the XRD in Fig. 1, and the M vs. H curves for the MAC1200 samples, the part from the near top of the ingot is almost pure τ-phase, but with broadening in the peaks likely originating from internal strains. The C content of τ-MnAl has a large influence on the value of TC as exemplified by Wei et al. ([15]), who found TC of pure Mn54Al46 to be 648 K and that TC decreased with increasing C content to 530 K for Mn51Al46C3. In the same study, MS was found to decrease with increasing C content from 650 kA/m for Mn54Al46 to 625 kA/m for Mn51Al46C3. The obtained MS are close to values found in our own previous work on C doped τ-MnAl; 597 kA/m [6] and 615 kA/m [7]. It is also in agreement with results from other groups; 621 kA/m (maximum applied field 5 T) [12,24], 639 kA/m (Mn52Al46C2) [15] and 588 kA/m [22] (maximum applied field 7 T). The difference in coercivity between the samples can be understood as being due to the smaller domains present in the MAC1400. Here the domains are coupled to the different grains, which from EBSD, SEM, EDS and MFM show that the grains in the MAC1400 are smaller compared to the other samples. From MFM, a strong magnetic contrast is observed from the larger grains (see Fig. 10). In MFM the non-magnetic regions between the grains are clearly visible as light regions, which is consistent with the data obtained from EBSD that indicate non-magnetic ε and/or β-Mn and γ2 phases in between the τ-phase grains. The effect of annealing at 1200 C for 30 min before cooling down to room temperature appears to have a significant effect on the grain growth, and thus also on the magnetic properties. A similar concept with high temperature homogenization was recently reported by Bittner et al. [24] and appears to be an important step towards optimizing the magnetic properties of the MnAlC compounds.
In this study, we report the differences in τ-phase content and magnetic properties between a direct cooling from the melt in the MnAlC system, and the intermediate annealing step at 1200 C for ~30 min before cooling down to room temperature. We found that from XRD, the sample cooled directly from 1400 C contained ~91 wt% τ-phase (~73% from EBSD), compared to the ~95.5 wt% τ-phase (89% from EBSD) for the middle part of the sample. The trend is the same between XRD and EBSD, however, EBSD may underestimate the value due to the local nature of the method and the more surface specific analysis. The different Al-, and Mn-rich regions from the SEM analysis are further confirmed to be non-magnetic and magnetic regions, respectively, based on the MFM method. This further evidence the notion that the Mn-rich phase (higher melting point in the Mn-Al system) is solidified first, while the Al-rich (lower melting point in the Mn-Al system) matrix is solidified secondly, thus leading to the Al-rich matrix (ε-rich phase). The Al-rich matrix phase is shown to be minimized through an intermediate step at ~1200 C for ~30 min before cooling, thus ensuring enough time for the system to equilibrate in solid state before cooling down to transform the MnAlC from the non-magnetic ε-phase to the desired τ-phase for permanent magnets. The exact cooling rate for the top of the MAC1200 sample needs to be quantified for improved phase purity in industrial implementation of the proposed method. Acknowledgements Financial support from SweGrids, The Swedish Energy Agency is gratefully acknowledged. Additional funding from the Swedish Foundation for Strategic Research, project “Magnetic materials for green energy technology” is gratefully acknowledged. Appendix A. Supplementary data Supplementary data to this article can be found online at https://do i.org/10.1016/j.jssc.2019.03.035. References [1] J.M. Coey, New permanent magnets; manganese compounds, J. Phys. Condens. Matter Inst. Phys. J. 26 (6) (2014), 064211. [2] K. Kamino, T. Kawaguchi, M. Nagakura, Magnetic properties of MnAl system alloys, IEEE Trans. Magn. 2 (3) (1966) 506–510. [3] S. Bance, F. Bittner, T.G. Woodcock, L. Schultz, T. Schrefl, Role of twin and antiphase defects in MnAl permanent magnets, Acta Mater. 131 (2017) 48–56. [4] T. Ohtani, N. Kato, S. Kojima, K. Kojima, Y. Sakamoto, I. Konno, M. Tsukahara, T. Kubo, Magnetic properties of Mn-Al-C permanent magnet alloys, IEEE Trans. Magn. 13 (5) (1977) 1328–1330. 235
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[5] H. Fang, J. Cedervall, F.J.M. Casado, Z. Matej, J. Bednarcik, J. Ångstr€ om, P. Berastegui, M. Sahlberg, Insights into formation and stability of τ-MnAlZx (Z ¼ C and B), J. Alloys Compd. 692 (2017) 198–203. [6] H. Fang, S. Kontos, J. Ångstr€ om, J. Cedervall, P. Svedlindh, K. Gunnarsson, M. Sahlberg, Directly obtained τ-phase MnAl, a high performance magnetic material for permanent magnets, J. Solid State Chem. 237 (2016) 300–306. [7] H. Fang, J. Cedervall, D. Hedlund, S. Shafeie, S. Deledda, F. Olsson, L. von Fieandt, J. Bednarcik, P. Svedlindh, K. Gunnarsson, M. Sahlberg, Structural, microstructural and magnetic evolution in cryo milled carbon doped MnAl, Sci. Rep. 8 (1) (2018) 2525. [8] O. Gutfleisch, M.A. Willard, E. Bruck, C.H. Chen, S.G. Sankar, J.P. Liu, Magnetic materials and devices for the 21st century: stronger, lighter, and more energy efficient, Adv. Mater. 23 (7) (2011) 821–842. [9] Y.J. Kim, J.H. Perepezko, The thermodynamics and competitive kinetics of metastable τ phase development in MnAl-base alloys, Mater. Sci. Eng., A 163 (1) (1993) 127–134. [10] H. Kono, On the ferromagnetic phase in manganese-aluminum system, J. Phys. Soc. Jpn 13 (12) (1958) 1444–1451. https://doi.org/10.1143/JPSJ.13.1444. [11] Y. Du, J. Wang, J. Zhao, J.C. Schuster, F. Weitzer, R. Schmid-Fetzer, M. Ohno, H. Xu, Z.-k. Liu, S. Shang, W. Zhang, Reassessment of the Al–Mn system and a thermodynamic description of the Al–Mg–Mn system, Int. J. Mater. Res. 98 (9) (2007) 855–871. [12] F. Bittner, L. Schultz, T.G. Woodcock, Twin-like defects in L10 ordered τ-MnAl-C studied by EBSD, Acta Mater. 101 (Complete) (2015) 48–54. [13] Z. Shao, H. Zhao, J. Zeng, Y. Zhang, W. Yang, Y. Lai, S. Guo, H. Du, C. Wang, Y. Yang, J. Yang, One step preparation of pure, AIP Adv. 7 (5) (2017), 056213.
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One step towards MnAl-based permanent magnets - Differences in magnetic, and microstructural properties from an intermediate annealing step during synthesis Samrand Shafeie a, Hailiang Fang a, Daniel Hedlund b, Axel Nyberg b, Peter Svedlindh b, Klas Gunnarsson b, Martin Sahlberg a, * a b
Department of Chemistry – Ångstr€om Laboratory, Uppsala University, Sweden Department of Engineering Sciences, Uppsala University, Box 534, SE-751 21, Uppsala, Sweden
A R T I C L E I N F O
A B S T R A C T
Keywords: Powder diffraction MnAl Permanent magnets High temperature synthesis Magnetic measurements
The influence of an additional annealing step during synthesis on the preparation of MnAl based permanent magnet alloys has been investigated. Bulk samples of Mn55Al45C2 alloys were synthesized using induction heating through drop synthesis from 1400 C. Samples produced using cooling directly from 1400 C (from the melt), and from 1400 C to an intermediate annealing step at 1200 C for ~30 min before cooling were compared with respect to differences in phase purity, microstructure and magnetic properties. We found that the phase purity was significantly enhanced using the route with an intermediate annealing step at 1200 C. From XRD the relative content of the τ-phase was improved from ~91 wt% for the sample cooled directly from 1400 C to ~95.1–99.5 wt% for the sample exposed to an intermediate annealing step before cooling. Additionally, EBSD, and SEM with EDS indicates a clear difference in the phase composition and differences in the distribution of the magnetic τ-phase and the non-magnetic ε-, β-, and γ2-phases. Magnetic properties also indicate an improvement in saturation magnetization for the sample exposed to the extra annealing step during synthesis. Our results suggest that an intermediate annealing step in the production of MnAl based alloys will provide a simple way of achieving better phase purity and magnetic properties in the bulk alloy.
1. Introduction The τ-phase with the ordered tetragonal L10 type structure in the MnAl-system has been reported as a suitable rare-earth free permanent magnet material [1]. From theory, the calculated saturation magnetization Ms ¼ 597 kA/m and magnetocrystalline anisotropy K1 ¼ 1:7 106 J/m3, yield a theoretical maximum for the energy product of ðBHÞmax 110 kJ/m3 [1,2]. Together with a magnetic ordering temperature of Tc 650 K, these properties indicate that the material has the potential to replace ferrites and epoxy resin bonded Nd–Fe–B magnets for applications (e.g., permanent magnet motors and wind turbine generators). Based on the ðBHÞmax , which is usually referred to as the “figure-of-merit” for permanent magnet materials, the MnAl-based permanent magnets fill an important gap between the Nd–Fe–B based and the ferrite based permanent magnets [3]. Previous studies have reported the stabilization of the τ-phase by small additions of carbon (~2 at.%) [4], which also increases the usability of this material.
Due to its comparably low cost [1] in relation to its magnetic properties, it becomes a viable material worth optimizing for certain applications where it may replace expensive high performance Nd–Fe–B magnets. It is therefore important to find ways to optimize its magnetic properties. In our previous studies [5,6] we reported the direct formation of the τ-phase from drop synthesis and the possibility to create a homogeneous τ-phase powder through cryomilling with subsequent heat treatments [7]. Ideally a permanent magnet should have a square shaped hysteresis loop to optimize ðBHÞmax with large values for the saturation magnetization and the magnetic coercivity. Permanent magnet alloys like MnAl require a grain morphology with size comparable to the single domain limit as well as shape that promotes alignment of the grain easy axes along a common direction [8]. A first step to move forward in this direction with the MnAl system is to find a synthesis and processing approach that will optimize the purity of the magnetic phase. Here, the high magnetocrystalline anisotropy of the τ-phase is a necessary, but not
* Corresponding author. E-mail address: [email protected] (M. Sahlberg). https://doi.org/10.1016/j.jssc.2019.03.035 Received 7 December 2018; Received in revised form 11 March 2019; Accepted 17 March 2019 Available online 20 March 2019 0022-4596/© 2019 Elsevier Inc. All rights reserved.
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sufficient, prerequisite for the τ-phase to be used as a permanent magnet. Further optimization and processing is needed to render the grain sizes smaller to increase the magnetic remanence of these material. In the MnAl-system, grains with the metastable τ-phase are commonly obtained through two main routes: 1) cooling from the ε-phase at a cooling rate ~10–20 C/min [9], or 2) quenching from the ε-phase followed by annealing at 300–600 C [10]. The phase diagram of MnAl shows that the ε-phase is stable in the temperature region 900–1200 C depending on the Mn:Al ratio. Fast cooling from the melt to the temperature region where the τ-phase starts to form will likely lead to the nucleation of a thermodynamically stable phase with the highest melting point (Mn-rich) followed by the formation of the phase with a slightly lower melting point (Al-rich). At high temperatures, δ-Mn may act as the first phase that nucleates, followed by the ε-phase and the γ2-phase. In the solid state, these phases will likely not dissolve into the main ε-phase. The γ and the δ-Mn phases most likely transform into γ2 and β-Mn phases, respectively [11]. The latter two phases may also form during the ε → τ transformation where the metastable τ -phase breaks down into γ2 and β-Mn phases at temperatures < 800–900 C if sufficient undercooling is not achieved [9]. To inhibit the initial nucleation of high-melting point Mn-rich phases that leave behind an Al-rich matrix with a lower melting point, fast cooling from the melt should be avoided. Thus, to obtain a more homogeneous and phase pure τ-phase, an intermediate step with annealing in the ε-phase region to homogenize the solid should be used before cooling to room temperature to initiate the ε → τ transformation. The formation of the τ–phase is characterized by the presence of crystalline defects like twin and anti-phase boundaries [3,12]. Both defects provide nucleation sites as well as pinning potentials for domain walls, and strongly influences the magnetic properties of the material [3]. Other groups have reported improvements in the synthesis of high purity τ-phase using strip casting [13], fast cooling from the melt with subsequent annealing between 650 and 350 C [10,14], or annealing at 1050 C in air for several days after arc-melting of ingots [15]. However, few of these may be cost-efficient for large scale production of high purity τ-phase MnAl permanent magnets. Although, other preparation methods produce similar saturation magnetization values (see e.g. Refs. [16,17]), our method also shows the importance of transforming from a homogenized solid rather than from the liquid melt to avoid the formation of secondary phases with higher melting points e.g., Mn-rich phases that shifts the bulk composition of the melt. Here, we report a significant improvement in the synthesis procedure for MnAl type materials through induction melting. We present the different microstructural and magnetic properties achieved through the change in the synthesis procedure of drop synthesized MnAlC-material with the composition Mn55Al45C2. From the Mn55Al45C2 ingots produced through drop synthesis either by 1) cooling directly to room temperature from the melt at 1400 C (hereafter referred to as MAC1400) or by 2) cooling first to 1200 C and annealing for ~30 min prior to cooling to room temperature (hereafter referred to as MAC1200). The microstructure and magnetic properties are reported for the two different routes together with the change in the inter-grain ε-phase content based on the two heating procedures.
~30 min before cooling to room temperature (MAC1200). The top part of the MAC1200 ingot is included in the study for comparison to investigate a possible influence of the cooling rate on the material properties. 2.2. Sample preparation Sample preparation for scanning electron microscopy (SEM) was made on pieces from the synthesized ingots. The pieces were put in a conducting polymer resin, and were polished down to 0.25–1 μm. For electron backscatter diffraction (EBSD) the pieces were polished down to < 50 nm in surface roughness using an Al2O3 suspension containing Al2O3 nanoparticles (<50 nm) as the final polishing step. For the microstructural analysis, elemental mapping and surface magnetic properties, the samples were cut into cubes with the dimensions 5 5 10 mm3. For EBSD analysis, separate samples were polished and taken out of the polymer resin in order to minimize image drifting effects from charge buildup during acquisition. Samples were also demagnetized using a AC demagnetizer which operates with a frequency of 32 kHz with a field smaller than 4 Oe (320 A/m) by moving the sample in and out of the demagnetizer. This was done to ensure that there was as little excess stray field buildup on the sample as possible. 2.3. Diffraction X-ray powder diffraction (XRD) was made on crushed pieces from the ingots that were crushed in an agate mortar together with ethanol. The XRD data was collected on a Bruker D8 diffractometer equipped with a Lynx-eye position sensitive detector using Cu Kα1 radiation (λ ¼ 1.540598 Å) and a Bruker D8 Advance equipped with a Lynx-eye energy dispersive position sensitive detector using Cu Kα radiation. The obtained powder diffraction data was refined using the Rietveld method [19] (FULLPROF software [20]) to verify the structural parameters and to calculate the weight percentage of the respective phases. The background was fit using a Chebychev polynomial function with a total of 6 parameters, in addition, the cell parameters, occupancies, and peak shape were refined with a total of 13 parameters. The additional phases were included in the final refinement to estimate relative weight fractions for the different phases. 2.4. Magnetic properties Temperature dependent magnetization (M) versus temperature (T) measurements were performed using a 7400 Lake Shore VSM equipped with a high temperature furnace. The temperature dependence of the magnetization was measured while decreasing the temperature from 750 K to 300 K in a constant applied magnetic field of 80 kA/m. A physical properties measurement system (PPMS) from Quantum Design was used for room temperature high field measurements. Powder samples were enclosed in separate capsules and magnetic hysteresis loops were measured in the magnetic field range of 7162 kA/m. Atomic force/magnetic force microscopy (AFM/MFM) was used to study the topography and magnetic microstructure of the samples. The polished samples were analyzed in a Digital Instruments Dimension 3100 microscope equipped with a low-coercivity magnetic tip. Post-processing of the images were done in the software Gwyddion [21] and are available in the supplementary information.
2. Experimental 2.1. Synthesis
2.5. Microstructural properties MnAl-based ingots with the nominal composition Mn55Al45C2 were prepared by induction melting through drop synthesis of high purity Mn (99.999%), Al (99.999%) and C (99.999%) as described elsewhere [18]. For this study, however, two ingots of approximately 20 g were produced through different annealing procedures: 1) the MnAlC melt was kept at 1400 C for ~5–10 min before the furnace was shut off and cooled to room temperature, (MAC1400); 2) the MnAlC melt was kept at 1400 C for ~5–10 min before lowering the temperature to 1200 C and kept for
The microstructure and the overall elemental mapping were obtained using a high resolution FEG Zeiss LEO1550 scanning electron microscope (SEM) equipped with an Oxford instrument energy dispersive X-ray spectrometer (EDS). For SEM and EDS an acceleration voltage of 20 kV was used. Microstructural information related to grain orientation and phase analysis were obtained using a high resolution FEG ZEISS Merlin equipped with a Nordlys Max detector for electron back scatter 230
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diffraction (EBSD) analysis and the AZtec software for X-ray mapping and elemental analysis. For EBSD analysis, an acceleration voltage of 25 kV was used.
Table 1 Estimated average weight fractions of the different phases for MAC1400, MAC1200 middle, MAC1200 top. Sample
3. Results MAC1400 MAC1200 middle MAC1200 top
3.1. XRD From Fig. 1a it is found that the MAC1400 sample consists of a majority of τ-phase. However, small residues of β-Mn and γ2-Mn phases are present, as summarized in Table 1. From the Rietveld refinements ~91.0 wt% is τ-phase. From the results shown in Fig. 1b and c it is estimated that the middle and top parts of the MAC1200 ingot consist of ~95.1 wt% and ~99.5 wt% τ-phase, respectively. From Table 1, the secondary phases (β-Mn and γ2-Mn) are found to decrease significantly after the additional heat treatment in the MAC 1200 sample. The largest decrease is observed in the top part of MAC 1200 (see Table 1). Tables with detailed information from the refinements can be found in supplementary information. A slight change in peak width (full width at half maximum, FWHM) is observed for the MAC1200 samples taken from the middle and top parts of the ingot, which could indicate differences in residual strain between the samples.
Phase content (wt. %)
τ
γ2
β
91.0(4) 95.1(6) 99.5(2)
8.6(4) 4.8(2) 0.5(2)
0.4(1) 0.2(1) –
large relative amount of the τ-phase (~73%) based on the indexed diffraction pattern of different regions. Also, the main phase detected at the grain boundaries can be indexed as the ε-phase (~22%) similar to what we have observed in our previous study [7] where the Al-rich ε-phase is expelled to the grain boundaries. Small amounts of γ2 and β-Mn phases are detected (~4.5% and ~0.5%, respectively). EBSD from different regions of the MAC1400 sample, however, indicate that the size of the ε-phase rich regions vary strongly over the sample and that the ~22% obtained from indexing the EBSD data is probably overestimated (powder diffraction yields 91% τ-phase in the sample). In MAC1400 where the ε-phase is present between the grains, it appears that the average grain orientation (of the ε-phase) is the same or nearly the same for the whole ε-phase region. The ε-rich regions are, however, very small in the MAC1200 samples. For the middle region in the MAC1200 ingot (see Fig. 7), the amount of τ, ε-, γ2-, and β-phases can instead be estimated to ~89%, ~1.1%, ~2.2% and ~1.4%, respectively. The rest are nonindexed points. From the EBSD maps of the MAC1400 and the MAC1200 samples, the average grain size appears to be larger for the MAC1200 sample (≳ 50 μm2), while the MAC1400 sample appears to have smaller grains (<50 μm2). The fewer but larger grains in MAC1200 appear in comparison to MAC1400 to be crystallographically more oriented. Additionally, while the majority of the inter-grain phase is identified as the ε-phase in the MAC1400 sample, it appears to be dominated by the γ2-phase and to a smaller extent by the β-Mn and ε-phases in MAC1200. Overall, the β-phase appears to form at the grain boundaries of the larger τ-phase grains, while the ε- and γ2-phases form continuous regions between the larger τ-phase grains.
3.2. SEM From the SEM images and the elemental mapping analysis of the MAC1400 and MAC1200 samples a clear difference is observed in the microstructure of the two samples (see Fig. 2). For the MAC1400 sample (see Fig. 2a), the backscatter image indicates brighter regions (typically indicative of higher average atomic weight) between the large grains. However, from the elemental map using EDS (see Fig. 3), these brighter regions can be connected with Al-rich phases, thus indicating possible electron scattering effects in the backscatter image (see arrows Fig. 2a). The grains, however, are rich in both Mn and Al, with C evenly distributed over the whole region. The C content is less reliable due to the use of carbon-based polymer resin that may redeposit on the surface of the sample and influence the C map. Compared to the MAC1200 sample (Fig. 2b), the difference between the Al rich regions and the Mn-rich regions are more distinct, and clearly indicate the existence of an Alrich matrix around the Mn-rich grains in the MAC1400 sample. Furthermore, Fig. 2 indicates a microstructure consisting of larger grains for the MAC1200 sample. From the EDS maps (see Figs. 4 and 5), fewer regions of Al-rich phases are observed for MAC1200 compared to MAC1400. Similar microstructure and distribution of Al rich phases are observed for both the top and middle regions of the MAC1200 ingot. However, the amount of Al rich phases is lower in the top part than in the middle, in accordance with the XRD analysis (Table 1).
3.4. Magnetometry Room temperature magnetic hysteresis loops of the MAC1400 and MAC1200 samples are shown in Fig. 8 and the results are summarized in Table 2. From the extracted data the coercivity (Hc ) follows the trend MAC 1400 > MAC 1200 top > MAC 1200 middle, while the saturation magnetization follows the opposite trend. The MS of the middle part, 622 kA/m is close to the theoretical value of 597 kA/m [1]. In addition, magnetization versus temperature results are shown in Fig. 9 for the MAC1400 and MAC1200 middle part samples, indicating that the Curie temperature defined by the peak temperature in the ∂M= dT versus temperature curve is largest (560 K) for the MAC1400 sample.
3.3. EBSD EBSD data from the MAC1400 sample (see Fig. 6) clearly indicates a
Fig. 1. XRD of ground pieces from a) MAC1400, b) MAC1200 middle part and c) MAC1200 top part (λ ¼ 1.540598 Å). 231
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Fig. 2. Electron backscatter image of a) MAC1400 sample and b) MAC1200 sample taken from the middle of the ingot.
Fig. 3. (a) SEM image of the MAC1400 sample together with EDS mapping of Al (b), Mn (c) and C (d).
Fig. 4. (a) SEM image of the MAC1200 middle part together with EDS mapping of Al (b), Mn (c) and C (d).
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Fig. 5. (a) SEM image of the MAC1200 top part together with EDS mapping of Al (b), Mn (c) and C (d).
Fig. 6. MAC1400 (cooled from melt), where the different figures correspond to a) secondary electron image of the area of interest, b) pattern quality map (band contrast) that indicates the grain boundaries and diffuseness of the diffraction, c) the grain orientation map (IPF Z) and d) a map over indexed phases (red ¼ β-Mn, blue ¼ ε - MnAl, yellow ¼ γ MnAl and green ¼ τ - MnAl). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
the melt at ~1400 C, before cooling directly to room temperature (MAC1400), and the sample where an additional hold at 1200 C for ~30 min before cooling to room temperature (MAC1200) was implemented. From the structural refinement of the XRD data (see Fig. 1), the amount of secondary equilibrium phases (e.g., β-Mn and γ2 phases) decreases significantly when the additional annealing step at 1200 C for ~30 min was implemented. Also, the XRD peak width (FWHM) increases significantly from 0.20 to 0.34 , using a Lorentzian fit of the strongest peak of the τphase ((101) at ~41 in 2θ), suggesting increased internal strain. The improved phase purity for the sample taken from the top part of the MAC1200 ingot can be attributed to the difference in cooling rate near the middle and the top parts of the ingot. More specifically, the middle of the ingot is expected to have a slightly lower cooling rate, which is expected to favor the formation of the undesired secondary
3.5. AFM/MFM From AFM/MFM analysis of MAC1400, a difference between the Alrich and the Mn-rich regions is clearly observed (see Fig. 10 (a) and (b)). By correlating the MFM to the EDS-mapping (Figs. 2 and 3), the Mn rich regions clearly show magnetic contrast unlike the Al-rich regions. The same response is observed in all samples, where the Mn-rich and the Alrich regions are observed, indicating that the same behavior is present in both samples (cf. SI). 4. Discussion Clear differences in the crystallographic and magnetic properties were observed between the sample prepared by conventional heating to
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Fig. 7. MAC1200 near middle regions (annealed at 1200 C for 30min), where the different figures corresponds to a) secondary electron image of the area of interest, b) pattern quality map (band contrast) that indicates the grain boundaries and diffuseness of the diffraction, c) the grain orientation map (IPF Z) and d) a map over indexed phases (red ¼ βMn, blue ¼ ε - MnAl, yellow ¼ γ - MnAl and green ¼ τ MnAl). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 9. Magnetization (M) and ∂M=∂T versus temperature for MAC1400 and MAC1200 middle part samples. H ¼ 80 kA/m. Fig. 8. Room temperature magnetic hysteresis loops for MAC1400, MAC1200 middle part and MAC1200 top part samples.
peak width, however, may relate to the internal strains that are created during the transformation from the ε-phase [22]. It has been reported that the transformation occurs through a massive transformation [23] that takes place through the nucleation of the τphase at the interface of the εphase, thus, requiring a large number of interfaces for the transformation to occur. The interfaces that remain after the transformation has been completed, and are likely strained, may give rise to the strong peak broadening that is observed. In addition, twin formation is expected when a phase transition goes from a high symmetry phase to a low symmetry phase, and such twin formation has been reported in the MnAl-system previously [3,12]. From the EBSD data, we observed that the small crystallite domains in MAC1400 increase in size in the MAC1200 sample, which would lead to the conclusion that the crystallite size may not be the main contributor to the peak broadening. Also, from SEM images and EDS mapping (see e.g., Figs. 2 and 3), we find that the Al-rich regions decrease significantly in size in MAC1200 compared to MAC1400. This is possibly a result of the time allowed for the equilibration to occur in the temperature region where the εphase is stable, before cooling.
Table 2 Saturation magnetization (Ms), coercivity (Hc), magnetic remanence (Mr), magnetic transition temperature (Tc) and XRD phase purity.* was measured in a magnetic field of 800 kA/m and show similar behavior as the others, however, the transition is more smeared out due to 10 times higher field.
MAC1400 MAC1200 top part MAC1200 middle part
Ms/(kA/ m)
Hc/(kA/ m)
Mr/(kA/ m)
TC/ (K)
XRD phase purity/(wt%)
561 606
56 20
106 60
560 545*
91.0(4) 99.5(2)
622
19
60
550
95.1(6)
equilibrium phases (e.g., ε-, γ2-, and β-phases). We have previously shown that the cooling rate is crucial for the phase purity of the τ-phase, and that the cooling rate achieved by air cooling appears to be closest to the optimal (~20 K/min) [5]. The XRD
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Fig. 10. AFM/MFM topographical image MAC1400 (a) and magnetic image MAC1400 (b).
5. Conclusions
In general, this method will circumvent the possibility of the liquid melt to first crystallize in a Mn-rich phase with higher melting point, before crystallizing out the Al-rich phase in between with the residual elements from the melt, as can be deduced from the phase diagram [11]. From the phase diagram, the ε-phase is still the equilibrium phase at ~1200 C, and it can be assumed that a prolonged annealing at 1200 C will ensure a more homogeneous ε-phase with the desired composition prior to the transformation to the τ-phase. In contrast to the untreated MAC1400 sample that has been used as a baseline sample with significant amounts of secondary phases the MAC1200 samples show that the method significantly improves the phase purity through the equilibration step at 1200 C. However, the method still requires optimization since the two different parts of the MAC1200 ingot that we studied in more detail clearly shows that cooling rate differences are still part of the challenge in obtaining phase pure τ-phase across the whole ingot. From the magnetization measurements, we can conclude that TCis lowest for the MAC1200 sample (545–550 K). As seen from the XRD in Fig. 1, and the M vs. H curves for the MAC1200 samples, the part from the near top of the ingot is almost pure τ-phase, but with broadening in the peaks likely originating from internal strains. The C content of τ-MnAl has a large influence on the value of TC as exemplified by Wei et al. ([15]), who found TC of pure Mn54Al46 to be 648 K and that TC decreased with increasing C content to 530 K for Mn51Al46C3. In the same study, MS was found to decrease with increasing C content from 650 kA/m for Mn54Al46 to 625 kA/m for Mn51Al46C3. The obtained MS are close to values found in our own previous work on C doped τ-MnAl; 597 kA/m [6] and 615 kA/m [7]. It is also in agreement with results from other groups; 621 kA/m (maximum applied field 5 T) [12,24], 639 kA/m (Mn52Al46C2) [15] and 588 kA/m [22] (maximum applied field 7 T). The difference in coercivity between the samples can be understood as being due to the smaller domains present in the MAC1400. Here the domains are coupled to the different grains, which from EBSD, SEM, EDS and MFM show that the grains in the MAC1400 are smaller compared to the other samples. From MFM, a strong magnetic contrast is observed from the larger grains (see Fig. 10). In MFM the non-magnetic regions between the grains are clearly visible as light regions, which is consistent with the data obtained from EBSD that indicate non-magnetic ε and/or β-Mn and γ2 phases in between the τ-phase grains. The effect of annealing at 1200 C for 30 min before cooling down to room temperature appears to have a significant effect on the grain growth, and thus also on the magnetic properties. A similar concept with high temperature homogenization was recently reported by Bittner et al. [24] and appears to be an important step towards optimizing the magnetic properties of the MnAlC compounds.
In this study, we report the differences in τ-phase content and magnetic properties between a direct cooling from the melt in the MnAlC system, and the intermediate annealing step at 1200 C for ~30 min before cooling down to room temperature. We found that from XRD, the sample cooled directly from 1400 C contained ~91 wt% τ-phase (~73% from EBSD), compared to the ~95.5 wt% τ-phase (89% from EBSD) for the middle part of the sample. The trend is the same between XRD and EBSD, however, EBSD may underestimate the value due to the local nature of the method and the more surface specific analysis. The different Al-, and Mn-rich regions from the SEM analysis are further confirmed to be non-magnetic and magnetic regions, respectively, based on the MFM method. This further evidence the notion that the Mn-rich phase (higher melting point in the Mn-Al system) is solidified first, while the Al-rich (lower melting point in the Mn-Al system) matrix is solidified secondly, thus leading to the Al-rich matrix (ε-rich phase). The Al-rich matrix phase is shown to be minimized through an intermediate step at ~1200 C for ~30 min before cooling, thus ensuring enough time for the system to equilibrate in solid state before cooling down to transform the MnAlC from the non-magnetic ε-phase to the desired τ-phase for permanent magnets. The exact cooling rate for the top of the MAC1200 sample needs to be quantified for improved phase purity in industrial implementation of the proposed method. Acknowledgements Financial support from SweGrids, The Swedish Energy Agency is gratefully acknowledged. Additional funding from the Swedish Foundation for Strategic Research, project “Magnetic materials for green energy technology” is gratefully acknowledged. Appendix A. Supplementary data Supplementary data to this article can be found online at https://do i.org/10.1016/j.jssc.2019.03.035. References [1] J.M. Coey, New permanent magnets; manganese compounds, J. Phys. Condens. Matter Inst. Phys. J. 26 (6) (2014), 064211. [2] K. Kamino, T. Kawaguchi, M. Nagakura, Magnetic properties of MnAl system alloys, IEEE Trans. Magn. 2 (3) (1966) 506–510. [3] S. Bance, F. Bittner, T.G. Woodcock, L. Schultz, T. Schrefl, Role of twin and antiphase defects in MnAl permanent magnets, Acta Mater. 131 (2017) 48–56. [4] T. Ohtani, N. Kato, S. Kojima, K. Kojima, Y. Sakamoto, I. Konno, M. Tsukahara, T. Kubo, Magnetic properties of Mn-Al-C permanent magnet alloys, IEEE Trans. Magn. 13 (5) (1977) 1328–1330. 235
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