United States

Sm(Fe0.8Co0.2)12 thin films with the ThMn12 type crystal structure are expected as a new high-performance permanent magnet material because they possess excellent magnetic properties such as high saturation magnetization µ0Ms of 1.78 T, anisotropy field µ0HA of 12 T and Curie temperature TC of 879 K [1]. Recently, it has been clarified that a coercive force µ0Hc as high as 1.2 T can be successfully obtained in the Sm(Fe-Co) thin films with B addition. It is considered that this is mainly due to the formation of RFe12 grains with columnar structure separated by the B-rich amorphous grain boundary phase [2]. In addition, µ0Hc was further increased by the diffusion of Si into the grain boundary phase in the Sm(Fe-Co)-B thin films. Then, it is predicted that µ0Hc will be further increased by reducing the soft magnetic phase existing at the vicinity of the initial interface between V under layer and Sm(Fe-Co)-B main layer [3]. In this study, a Sm seed layer was inserted between the V under layer and the Sm(Fe-Co) layer to suppress the initial growth of soft magnetic phase such as α-Fe and α-(Fe-Co), and the crystal structure and magnetic properties were investigated in detail. The samples were prepared by using the ultra-high vacuum magnetron sputtering system. First of all, a V buffer layer of 20 nm was deposited onto the MgO (100) single crystal substrate at substrate temperature Ts of 673 K. Then, the thickness tSm of the Sm seed layer was changed from 0 to 5 nm and while that of Sm(Fe-Co)-B layer was fixed at 100 nm. Finally, a V cover layer of 10 nm was deposited. The B composition was changed up to 5 at.%. It was confirmed that the introduction of the Sm seed layer into Sm(Fe-Co) layer clearly decreases the intensity of the peaks obtained from α-(Fe-Co) and increases the intensity of the (002) and (004) peaks in the RFe12 phase. In addition, improvement of remanent magnetization ratio while maintaining µ0Hc above 1.0 T was confirmed. It was thought that the Sm seed layer plays an important role to suppress the formation of α-(Fe-Co) phase and to improve the orientation of the 1:12 phase. The detailed results will be discussed at the conference.&lt;br&gt;
**References:**&lt;br&gt;[1] Y. Hirayama, Y. K. Takahashi, S. Hirosawa and K. Hono, Scr. Mater., 138 (2017) 62-65.&lt;br&gt;
[2] Sepehri-Amin, Y. Tamazawa, M. Kambayashi, G. Saito, Y. K. Takahashi, D. Ogawa, T. Ohkubo, S. Hirosawa, M. Doi, T. Shima, K. Hono, Scr. Mater., 194 (2020) 337-342.&lt;br&gt;
[3] A. Boyachkin, H. Sepehri-Amin, M. Kambayashi, Y. Mori, T. Ohkubo, Y. K. Takahashi, T. Shima, K. Hono, Acta. Mater., 227 (2022) 117716.&lt;br&gt;




MMM 2022

Suppression of the formation of soft magnetic phase for Sm(Fe Co) thin films by introducing Sm seed layer

Sm(Fe0.8Co0.2)12 thin films with the ThMn12 type crystal structure are expected as a new high-performance permanent magnet material because they possess excellent magnetic properties such as high saturation magnetization µ0Ms of 1.78 T, anisotropy field µ0HA of 12 T and Curie temperature TC of 879 K [1]. Recently, it has been clarified that a coercive force µ0Hc as high as 1.2 T can be successfully obtained in the Sm(Fe-Co) thin films with B addition. It is considered that this is mainly due to the formation of RFe12 grains with columnar structure separated by the B-rich amorphous grain boundary phase [2]. In addition, µ0Hc was further increased by the diffusion of Si into the grain boundary phase in the Sm(Fe-Co)-B thin films. Then, it is predicted that µ0Hc will be further increased by reducing the soft magnetic phase existing at the vicinity of the initial interface between V under layer and Sm(Fe-Co)-B main layer [3]. In this study, a Sm seed layer was inserted between the V under layer and the Sm(Fe-Co) layer to suppress the initial growth of soft magnetic phase such as α-Fe and α-(Fe-Co), and the crystal structure and magnetic properties were investigated in detail. The samples were prepared by using the ultra-high vacuum magnetron sputtering system. First of all, a V buffer layer of 20 nm was deposited onto the MgO (100) single crystal substrate at substrate temperature Ts of 673 K. Then, the thickness tSm of the Sm seed layer was changed from 0 to 5 nm and while that of Sm(Fe-Co)-B layer was fixed at 100 nm. Finally, a V cover layer of 10 nm was deposited. The B composition was changed up to 5 at.%. It was confirmed that the introduction of the Sm seed layer into Sm(Fe-Co) layer clearly decreases the intensity of the peaks obtained from α-(Fe-Co) and increases the intensity of the (002) and (004) peaks in the RFe12 phase. In addition, improvement of remanent magnetization ratio while maintaining µ0Hc above 1.0 T was confirmed. It was thought that the Sm seed layer plays an important role to suppress the formation of α-(Fe-Co) phase and to improve the orientation of the 1:12 phase. The detailed results will be discussed at the conference. 
**References:** [1] Y. Hirayama, Y. K. Takahashi, S. Hirosawa and K. Hono, Scr. Mater., 138 (2017) 62-65. 
[2] Sepehri-Amin, Y. Tamazawa, M. Kambayashi, G. Saito, Y. K. Takahashi, D. Ogawa, T. Ohkubo, S. Hirosawa, M. Doi, T. Shima, K. Hono, Scr. Mater., 194 (2020) 337-342. 
[3] A. Boyachkin, H. Sepehri-Amin, M. Kambayashi, Y. Mori, T. Ohkubo, Y. K. Takahashi, T. Shima, K. Hono, Acta. Mater., 227 (2022) 117716.

poster

#### Welcome to the 67th Annual Conference on Magnetism and Magnetic Materials, MMM 2022!

The Conference was held from October 31 to November 4 and offers both in-person and on-demand content.

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Nowdays, Nd-Fe-B magnetic powder is usually prepared by the strip casting (SC), hydrogen decrepitation (HD) and jet milling (JM), so as to obtain high performance. However, due to the harsh hydrogen absorption conditions of Sm-Co alloy, the HD is rarely used for its powder preparation, and there is no research on the hydrogenation behavior of SC flake. Therefore, it is urgent to introduce a new powder preparation process for mass production. In this study, the hydrogen absorption behavior of Sm(CoFeCuZr)z SC flake was systematically investigated. 
The SC flake with nominal composition of Sm(Co0.065Fe0.28Cu0.05Zr0.02)7.6 was prepared by a single-roller SC technology. The SC flake composes of the free surface, wheel surface and cross section. The cross section consists of I-columnar grains region, III-fine grains region, and II-mixed region of fine grains and columnar grains [Fig1. (a-b)]. The weight of each sample involved in hydrogen absorption is 1.8 kg. The hydrogen absorption curves show that: when the hydrogen pressure is higher than 4.5MPa, the SC flake can break into powders by absorbing hydrogen at 20 °C. When the temperature increases to 50 °C, even if the hydrogenthe pressure reduces to 4 MPa, the SC flake can still absorb hydrogen and break. Increasing the hydrogen pressure can promote the particle refinement to a certain extent. However, increasing the temperature can obviously accelerate the HD process, but if the temperature is higher, the particle size will be coarse. The powders with small particle size, uniform distribution and full of cracks can be obtained under 50 °C-6 MPa [Fig2. (b)], which is conducive to the subsequent preparation of fine powder. The particles size distribution of the powder after JM are uniform [Fig2. (c-d)]. In conclusion, the hydrogen pressure and temperature are the key factors for HD process. The SC, HD and JM can be used as a new powder preparation process, which is expected to prepare high-performance Sm(CoFeCuZr)z magnets in the future. 
**References:** [1] A. Eldosouky, I. Škulj, J. Sustain. Metall., vol. 4, p. 516–527(2018). 
[2] A. Kianvash, I.R. Harris, J. Mater. Sci., vol. 20, p. 682–688(1985). 

![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/3a91a63ffdcd168be69eae52baafc706.png) 
Fig. 1. (a) Back-Scaterred Electron of the flake, (b) Microstructure model 
![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/eaf8e345a7e968e456221232a4d04af2.png) 
Fig. 2. (a) Hydrogen absorption curves, (b) Hydrogen absorption curve under 50 °C-6 MPa, (c-d) Powder micrograph and particle size distribution after JM

Effect of Hydrogen Pressure and Temperature on Hydrogen Absorption Behavior of Sm(CoFeCuZr)z Strip Casting Flake

The 2:17 type SmCo permanent magnets are well known for their excellent magnetic properties, good thermal stability and strong corrosion resistance, which are widely used in high temperature applications [1-2]. There are few reports on the preparation of SmCo magnet using strip-casting alloys. In this paper, the effect of casting-alloy methods on the microstructure and magnetic properties of Sm(CoFeCuZr)z magnets was investigated. 
The magnetic properties of the PM magnet prepared by mechanical crushing and ball milling are better than those of the SC magnet, especially the orientation. In detail, the magnetic properties of PM magnet are Br=11.63kG, Hcj=25.65kOe, (BH)max=30.97MGOe, while the magnetic properties of the SC magnet are Br=11.45kG, Hcj=24.00kOe, (BH)max=30.03MGOe, respectively. Both random magnets mainly consist of 1:5 and 2:17R phases [Fig. 1(a)]. After oriented, the intensity ratio of (006) and (104) peaks is used to characterize the orientation degree [3], I(006)/I(104)=24.33 in the PM magnet, while this value is 18.23 in the SC magnet. This means that PM magnet has better orientation degree. The EBSD results confirm above result. Meanwhile, there are a few grains with poor orientation in the inverse pole figure [Fig. 1(b)] of the SC magnet, and such grains will reduce the orientation. Figure 1(c) shows the (0001) pole figure (PF) and inverse pole figure (IPF) of normal direction, from which we can see the MRD values in PF and IPF of PM magnet are both higher than those of SC magnet. 
The BSE photographs of the different alloy cross-sections show that the PM alloy consists of columnar grains region, equiaxed grains region [Fig. 2a], while the SC alloy consists mainly of columnar grains region and fine grains region. Due to the existence of fine grains region in the SC[Fig. 2b], it is difficult to obtain all single grain in the subsequent ball-milling process. So the SC magnet has worse orientation than the PM magnet. 
**References:** [1] J.F. Liu, Y. Ding, G. C. Hadjipanayis, J. Appl. Phys., vol.85, p. 2800–2804(1999). 
[2] Wei Sun, Minggang Zhu, Yikun Fang,. J. Magn. Magn. Mater., vol.378, p. 214-216(2015). 
[3] Q. J. Zheng, S. Xia, A. Dozier., Mater. Sci. & Tech., vol. 22, p. 1476-1482(2006). 

![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/c339b910bf7abc99282a1b38f337d010.png) 
Fig. 1 (a)The XRD patterns of random and oriented magnets, (b) inverse pole figure of magnets, (c)the (0001) pole figure and inverse pole figure of normal direction 
![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/ca8a537d84e29b5b347c86532a6702de.png) 
Fig. 2 The Cross-sectional BSE of (a) plate ingot alloys and (b) strip-casting alloys.

Effect of casting alloy methods on magnetic properties of Sm(CoFeCuZr)z magnet

Equiatomic rare earth intermetallic compounds RNi (R = heavy rare earth such as Gd, Tb, Dy, Ho and Er) exhibit excellent magnetocaloric effect near their ferromagnetic ordering temperature and alloying the rare earth site has been found to be an efficient method to tailor the magnetic transition temperature and the associated magnetic entropy change [1]. In the present work, three rare earth elements Tb, Ho and Er in equimolar concentration occupy the rare earth site in RNi compound. Polycrystalline Tb0.33Ho0.33Er0.33Ni has been prepared by arc-melting under argon atmosphere and characterized by powder X-ray diffraction and magnetization measurements. Multicomponent Tb0.33Ho0.33Er0.33Ni compound has FeB-type orthorhombic crystal structure (space group Pnma, No.62) at room temperature. The sample orders ferromagnetically at ~59 K (TC) [Fig. 1a]. Paramagnetic susceptibility is Curie-Weiss like [Inset in Fig. 1a]. It should be recalled that the compounds with individual rare-earths viz. TbNi, HoNi and ErNi order ferromagnetically at about 67 K, 36 K and 11 K respectively. Magnetization has been measured at 5 K as a function of magnetic field up to 70 kOe. Magnetization saturates in applied magnetic field [Fig. 1b] and the saturation magnetization value is about 9.9 µB/f.u. A minor hysteresis with a coercive field of ~930 Oe is observed. By measuring isothermal magnetization data around TC, magnetocaloric effect is estimated. A moderate isothermal magnetic entropy change has been obtained near TC that is of the same order as in the end member RNi compounds. Understanding magnetic properties of such multicomponent rare earth intermetallic compounds should pave a way to comprehend the complex magnetism of rare earth high-entropy alloys. 
**References:** 1. X. Q. Zheng, B. Zhang, H. Wu, F. X. Hu, Q. Z. Huang and B. G. Shen, J. Appl. Phys. 120 (2016) 163907 

![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/a983eea10e8c486b321b35c4b26ebe3e.png) 
Fig. 1. (a) Magnetization vs temperature and (b) magnetization vs field of Tb0.33Ho0.33Er0.33Ni compound.

Understanding the magnetism of multicomponent rare earth intermetallic compound Tb0.33Ho0.33Er0.33Ni

We have fabricated isotropic Pr-Fe-B thick-film magnets with the thickness of 0.25 mm on a stainless shaft using a PLD (Pulsed Laser Deposition) method and applied them to an ultra-stepping motors with the diameter of less than 3 mm by taking advantage of Pr2Fe14B phase with a large anisotropic magnetic field at room temperature [1]. In this study, we focused on the fact that the linear expansion coefficient of Nd element (9.6×10-6 1/K ) is closer to that of stainless steel(10～11×10-6 1/K) than that of Pr element( 6.7×10-6 1/K ). Moreover, evaluation on the temperature dependence of magnetic properties and observation in the mechanical characteristic of Nd-Fe-B thick-film magnets on a stainless substrate (shaft) were carried out. 
Figure 1 shows the temperature dependence of the energy product at the operating point of (Nd or Pr)-Fe-B thick-film magnets fabricated on stainless substrates. At 125 °C, the energy product of a Nd-Fe-B thick-film exceeds those of a Pr-Fe-B one and a Nd-Fe-B bond magnet. The result is considered to be attributed to the different temperature dependence of each anisotropic magnetic field for Pr2Fe14B and Nd2Fe14B phases. We also observed the mechanical characteristic of the both films on a stainless substrate and shaft. Several as-deposited Pr-Fe-B film magnets with the Pr contents exceeding 15 at % on the substrate showed exfoliation phenomenon. On the other hand, as-deposited and annealed Nd-Fe-B films (Nd contents : 14 ～ 18 at. %) on the substrate didn’t show mechanical destruction. Although almost all the films showed good adhesion on the stainless shafts as displayed in Fig. 2, one Pr-Fe-B thick-film magnet had a slight crack at the boundary of the shaft and the film. Consequently, enhancement in coercivity of Nd-Fe-B thick-film using an additive at room temperature is a promising method to propose an alternative material of a previously reported Pr-Fe-B thick-film. 
**References:** [1] M. Nakano et.al., IEEE Trans. Magn. Vol. 56, 7516303(2020). 

![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/12a99c2023d5eae1a86663002d5f32c2.png) 
Fig 1 Energy product (Pc =1.0) of (Nd or Pr)-Fe-B thick film magnets as a function of operating temperature. 
![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/24daf7821456da0e79c5dbe705fb128b.png) 
Fig. 2 Exfoliation phenomenon of as-deposited (Nd or Pr)-Fe-B thick-films on stainless shafts.

Development of rare earth thick

For high performance and high efficiency, rare earth magnets have been widely used in recent motors. Rare earth is very unstable as a future industry due to limited resources and monopoly of some countries. For this reason, as the price of rare earth magnets and overall materials has increased recently, the issues of cost reduction and weight reduction are increasing. In this paper, iron usage and weight were reduced by changing the existing EPS motor to the STM model [1]. However, the STM model showed low performance compared to the amount of magnet used because large leakage occurred in the lips. To compensate for these shortcomings and to reduce the amount of magnet used, consequent poles were applied. When Consequent pole is applied, the usage of rare earth magnet is greatly reduced. Therefore, the weight and price are greatly reduced. In general, the consequent pole reduces the use of magnets by half, so the performance is also reduced by half. However, in the STM model, the magnetic flux leakage from the lips is greatly reduced by removing the poles. Therefore, the performance is reduced by about 35%, not by half. Therefore, even though the use of magnets has been reduced by half, by applying STM as a consequent pole rather than a general consequent pole with the same performance, many advantages such as weight reduction and reduced use of magnets can be brought. The application of this motor is EPS. For EPS, cogging torque is the most important performance factor. However, when the Consequent pole is applied, only the N pole or the S pole exists and the rest is made of iron, so a lot of cogging torque and torque ripple occurs. In this paper, after applying the consequent pole to the STM model, various methods are applied to reduce the cogging torque. Finally, it can be seen that the weight and magnet usage are reduced with the same performance, and the cogging torque is similarly generated. 
**References:** [1] S. -W. Song, I. -J. Yang, S. -H. Lee, D. -H. Kim, K. S. Kim and W. -H. Kim, "A Study on New High-Speed Motor That Has Stator Decreased in Weight and Coreloss," in IEEE Transactions on Magnetics, vol. 57, no. 2, pp. 1-5, Feb. 2021, Art no. 8201105, doi: 10.1109/TMAG.2020.3012482. 

![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/0abfd8dfb908f1e8bb0f63c25f68ac61.png) 
Fig. 1 Consequent pole applied to STM structure 
![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/fed13d62fa9eecb87a4406d56e74f4dd.png) 
Fig. 2 STM structure magnetic flux line

A study on the application of consequent poles to the STM model to reduce weight and use of magnets

The use of Nd-Fe-B magnets with Nd2Fe14B, which are used in high performance motors, is increasing. However, it is predicted that inexpensive Nd will become unavailable in near future, and there is a need to reduce the amount of Nd used and develop new magnets [1]. Exchange-coupled nanocomposite magnets consisting of magnetic hard and soft nanoparticles, NPs have been proposed as candidate [2]. The Sm(Fe0.8Co0.2)12 with ThMn12 structure has higher saturation magnetization, Ms, and uniaxial magnetocrystalline anisotropy, Ku than Nd2Fe14B and is expected to be a magnetic hard phase in the Fe-based system without Nd [3]. We investigate the effect of the direction of the α-Fe coating surface and the volume fraction of α-Fe, VFe on the demagnetization process and the maximum energy product, (BH)max of exchange-coupled Sm(Fe0.8Co0.2)12 /α-Fe nanocomposite NPs using LLG simulations [4]. Sm(Fe0.8Co0.2)12 NPs is as reference model #1. The nanocomposite NPs of Sm(Fe0.8Co0.2)12 coated with α-Fe on top/bottom, and both sides are model #2 and #3. For both # 2 and #3, α-Fe layer thickness is 1 nm ≤ tFe ≤ 7 nm. The size of NPs are x = y = 50 nm and z = 100 nm. The magnetic properties of Sm(Fe0.8Co0.2)12 and α-Fe were taken from references [3, 5-7]. The cell volume is 1 nm3. External field, Hex is applied parallel to c-axis of Sm(Fe0.8Co0.2)12 along to z-axis of NPs. Fig. 1 shows VFe dependence of (BH)max for #2 and #3. The (BH)max of #1 was 630 kJ/m3 , while the (BH)max of #2 and #3 with optimized tFe were larger, 636 kJ/m3 for #2 with VFe = 4 % (tFe = 2 nm) and 657 kJ/m3 for #3 with VFe = 12 % (tFe = 3 nm). Fig. 2 shows distribution of demagnetizing field, Hd at Hex = 0 T for #2 and #3. The Hd acting on top/bottom surfaces of NPs was larger for #2, which was coated on top/bottom surfaces with α-Fe, which is larger than Ms of Sm(Fe0.8Co0.2)12, than for #3, which was coated on sides with α-Fe. Therefore, remanence and (BH)max of #3 are larger than those of #2. From the above discussion, it is clear that side coating of magnetic soft phase is effective in increasing (BH)max of Sm(FeCo)12/α-Fe NPs. 
**References:** [1] S. Hirosawa, M. Nishino and S. Miyashita, Adv. Nat. Sci: Nanosci. Nanotechnol., 8, (2017) 013002. 
[2] R. Skomski and J. M. D. Coey, Phys. Rev. B, 48, (1993) 15812. 
[3] Y. Hirayama et al., Scripta Mater., 138, (2017) 62–65. 
[4] M. J. Donahue and D. G. Porter, “OOMMF User’s Guide, Version 1.0,” NISTIR 6376, National Institute of Standards and Technology, Gaithersburg, MD (Sept 1999). ; M. J. Donahue, D. G. Porter (2021), "OOMMF: Object Oriented MicroMagnetic Framework," https://nanohub.org/resources/oommf. (DOI: 10.21981/8RRA-5656). 
[5] Felix Jimenez-Villacorta and Laura H. Lewis, Chapter 7 "Advanced Permanent Magnetic Materials" Nanomagnetism, edited by J. Gonzalez (OCP Publishing Group) (2014). 
[6] D. Ogawa, T. Yoshioka et al., J. Magn. Magn. Mater., 497, (2020) 165965. 
[7] T. Fukazawa, H. Akai et al., J. Magn. Magn. Mater., 469, (2019) 296-301. 

![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/3a61bed85d0536499bdd51d5c215ec21.png) 
Fig. 1 VFe dependence of (BH)max for #2 and #3. 
![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/a2b7eb05cc1ea092f5b586fe81fd43cf.png) 
Fig. 2 Hd distribution on the top surface and the cross-section of NPs for #2 and #3 at Hex = 0 T.

Maximum Energy Product of Exchange coupled Sm(FeCo)12/α

We will present our results on the possibility of deriving high performance magnet powder for bonded magnets production via hydrogen decrepitation (HD) and heat treatment. Applications of permanent magnets (PMs) in the energy conversion industry and in electric vehicles have attracted enormous recent interest. The flexibility of bonded magnets manufacturing allows for high degree of magnet shape customization. They are manufactured by blending rare earth magnet powder with thermoset or thermo-plastic polymers. The magnetic properties of bonded magnets, such as maximum energy product, depend on the loading of the starting powders. By developing an approach for producing bonded magnets via HD, we will enable future opportunity to recycle end-of-life sintered magnets into high performance bonded magnets. 
In this work, HD of N52 sintered magnets were performed at 1bar, 1.5bar, 2bar and 4bar pressure. The resultant powder due to the absorption of the hydrogen gas and pulverization of the material shows soft magnetic characteristics. The HD powder was characterized by DSC to find the degassing temperature. Fig.1 shows the DSC curve of 4 bar pressure produced HD powder. The peak at T1=151°C attribute to removal of H2 from Nd2Fe14B phase, T2=175°C conversion of Nd trihydride to Nd dihydride and the Peak at T3=655°C correspond to formation of Nd rich phase at grain boundary after degassing from NdH2 1. Hence, we carried out degassing in vacuum furnace at 800°C for 2h. The degassed samples show hard magnetic properties (Fig.2) however, final powder product magnetic property does not vary with pressure in HD process. Therefore, we selected the powder treated at 4bar, which shows 30 MGOe maximum energy product for bonded magnets development. The bonded magnets were developed by compression molding of a mixture of 60 vol% of magnetic powder and 30 vol% Nylon 12 at 180 °C for 5 min, followed by ambient cooling. Throughout the process, 3kPa pressure was applied to achieve maximum density. The bonded magnets had a density of 4.1 g/cc. Magnetic field alignment was performed on the bonded magnets at 180 °C under 0.8 T field to obtain 13MGOe maximum energy product. 
**References:** M.Zakotnik, E.Devlin, I.R.Harris, A.J.Williams, Journal of Alloys and Compounds,
450, 525-531, 2008 

![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/808cb55f27f0aedb207995c564480198.png) 
Fig. 1 DSC of HD powder 
![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/613d0e1318ab6239d6cca5892fae0e61.png) 
Fig. 2 MH loop of annealed powder

Bonded Magnets Development using Hydrogen Decrepitation

While tetragonal MnAl with the L10-type structure (P4/mmm #123) is metastable, it displays good intrinsic hard magnetic properties such as large magnetization (Ms=110 emu/g), high magnetocrystalline anisotropy field (HA=40 kOe), and moderate Curie temperature (Tc=600 K),1 which make it highly attractive as a sustainable rare-earth free permanent magnet. However, the metastable behavior of the MnAl L10-phase promotes the formation of structural defects that degrade its intrinsic magnetic properties such as the local magnetocrystalline anisotropy (MCA). In this work, we study the effects of a partial replacement of Al by Ga on the stability of the L10-phase, and both the magnetization and the MCA using a first-principles density functional calculation. We find that increasing the Ga content, enhances the phase stability while the total magnetic moment per formula unit (f.u.) remains almost unchanged. The site-, atomic- and spin-resolved MCA has been estimated by evaluating the spin-orbit coupling energy using a second-order perturbation approximation. Structural distortion, chemical effects and different spin components in Mn(Al1-xGax) all contribute to the MCA. The site- and atomic-resolved MCA calculations show that the MCA energy (MAE) mainly comes from the Mn atoms. With increasing Ga content, the MAE at both Mn and Al(Ga) sites increases and the total MAE increases from 0.2 meV/f.u. for x=0 to 0.34 meV/f.u for x=1 (Fig. 1). However, the structural distortion induced by the partial replacement of Al by Ga reduces the MCA. The spin-resolved MCA and band structure calculations indicate that the microscopic origin of high MCA is mainly associated with the spin flipping behavior near the Fermi level, induced by the spin-orbit coupling. The derived effective magnetic anisotropy field increases from 37 kOe (x=0) to 46 kOe (x=1), in agreement with experiments. Doping with Ga improves the stability of the L10 structure and enhances the magnetic anisotropy field, which facilitates the development of high coercivity Mn-Al based permanent magnets. 
**References:** 1 L Pareti, F. Bolzoni, F. Leccabue, A.E. Ermakov, J. Appl. Phys. 59, 3824 (1986) 

![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/2686cbc73adcf807336315c2b6ddbdee.png) 
Fig. 1 Dependence of MAE on Ga content, x, in Mn(Al1-xGax).

Constituent contribution to the magnetocrystalline anisotropy in Mn(Al1 xGax)

Effects of La3+ and Li1+ co-substitution on structural and magnetic properties of M-type strontium hexaferrite have been studied by using crystallographic and magnetic measurements. Stoichiometric Sr1-xLaxFe12-x/2Lix/2O19 (x=0, 0.25, 0.4, 0.5, and 0.6) samples have been prepared by the conventional ceramic method using La2O3, SrCO3, Fe2O3, and Li2CO3 powders with repeated sinterings at 1250 °C for 24 hours.
XRD patterns showed that when x > 0.6, LaFeO3, SrFeO3-δ, and other spinel ferrite coexisted along with the M-type hexaferrite. Lattice parameters deduced from the GSAS refinement shows that both a and c decrease with increasing x, but the change in c is more outstanding (Table 1). Variations of the saturation magnetization Ms and the coercivity Hc with x was obtained from room temperature VSM measurement as in Table 1. Whereas the Ms decreases slightly with x, the Hc increases with the La3+ - Li1+ co-substitution throughout the doping range examined. Effective magnetic anisotropy constants K are calculated as well from the hysteresis curves using the law of approach to saturation method. K shows slight decrease with increasing x as listed in Table 1. 
XPS scan shows the Li 1s peak at ~56 eV (Fig. 1(a)). Apparently, monovalent Li1+ is non-magnetic. In view of the facts that Li1+ in the lithium ferrite (LiFe5O8) preferentially occupies the octahedral site and that the M-type hexaferrite has the spinel S-block in its structure, the decreasing trend of Ms with x strongly suggests that Li+1 ions preferentially replace Fe3+ ions either at 12k or 2a sites (rather than the 4f2 site in R-block). Non-magnetic substitution at those sites is known to give rise to negative contribution to Ms. Finally, Mössbauer spectra show gradual decrease in 2a intensity, in support of the octahedral 2a site preference of Li1+ in M-type hexaferrite (Fig. 1(b)). 
**References:** 

![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/df4dab457afcb7e5118c6d0fedb9652a.png) 
Table 1. Various structural and magnetic parameters for Sr1-xLaxFe12-x/2Lix/2O19 (x=0, 0.25, 0.4, 0.5, and 0.6) 
![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/05367f421b43d801d81539c55a4d55eb.png) 
Figure 1. (a) A comparison of Li 1s XPS scan for Sr1-xLaxFe12-x/2Lix/2O19 (x=0.25 and 0.5) and (b) 57Fe Mössbauer spectrum for Sr0.4La0.6Fe11.7Li0.3O19 taken at room temperature.

Structural and magnetic characterizations of La3+ and Li1+ co substituted M

The magnetization of Mn4C with face centered cubic structure increases with the increase of temperature within a certain temperature range, which means that Mn4C may be applied to the field requiring high working temperature of magnetic materials through further research and development. Unfortunately, Mn4C is a metastable phase [1-3], so it is difficult to synthesize high-purity Mn4C. 
High-purity Mn4C powders was prepared by melt spinning, grinding and magnetic separation in this paper. It is found that high cooling rate is beneficial to obtain high-purity Mn4C material. Moreover, compared with argon gas as the grinding medium, grinding Mn4C melt-spun strips in alcohol solution is not only conducive to refining the size of the ground particles, but also conducive to improving the dispersion of the particles. Therefore, Mn4C powders with higher purity can be obtained after magnetic separation (see Fig. 1).
The magnetization of Mn4C powders remained unchanged with the increase of temperature in the range of 2 K -50 K, but increased with the increase of temperature in the range of 50 K to 300 K (see Fig. 2). This may be because there is a spin glass state in Mn4C powders. At low temperature, the spins are frozen, but the frozen spins are gradually thawed with the increase of temperature. Further investigation is in progress. 
Crystal and magnetic structure of Mn4C powders was investigated by variable temperature neutron diffraction technique. The magnetization change of Mn4C with temperature should be mainly due to its own magnetic phase transition, and is closely related to the antiparallel arrangement of the magnetic moments of Mn atoms at the symmetrical crystal positions 1a and 3c in the Mn4C lattice. The total magnetic moment of crystal cells increases with the increase of temperature. 
**References:** 1. Si, PZ, Qian, HD, Ge, HL,et al. Applied physics letters, 2018, 112(19):192407. 
2.Park ,J, Qian, HD, Si, PZ, et al. Journal of magnetism and magnetic materials, 2021, 527:167765. 
3. Tagawa, Y and Motizuki, K. Journal of physics-condensed matter, 1991, 3(12): 1753-1761 

![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/65df04ba198f154c12183b9b3f6c217e.png) 
Fig.1 XRD patterns of Mn4C obtained by (a) arc melting, (b) Suction casting, (c) melt spinning technology (50m/s),and (d) grinding and magnetic separation of the same strips as (c) in alcohol solution. (a) (b) and (c) are both heat treated and ground and magnetically separated in argon. 
![](https://s3.eu-west-1.amazonaws.com/underline.prod/uploads/markdown_image/1/image/e3091542db36f8e069365c557c6f6f8e.png) 
Fig.2 (a) Variation of magnetization of Mn4C powder with temperature under different magnetic field; (b) Hysteresis loops of Mn4C powders at 2K-300K

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