An attempt was made to fabricate the sintered La-base M-type single phase ferrite magnet, which was known to be difficult to synthesize, by Na-dopping. La-base M-type single phase (LaNa)Fe8-xCoxO19 sintered ferrite magnet was successfully fabricated by dopping and substituting with Na = 0.2 and Co = 0.1. For this sintered magnet, M-phase fraction decreased with increasing Co-substitution. (La0.8Na0.2)Fe7.9Co0.1O19 sintered magnet has good magnetic properties; M-s = 4.7 kG, B-r = 4.59 kG, which is considered to be superior to commercial ferrite magnet.

In this study, we investigated the synthesis of Na-La system M-type ferrite with different mole ratios. Some part of Fe in the M-type ferrite was substituted by Co, and variation of fraction of the M-type ferrite as function of mole ratio was investigated. Sample with/without Co-substitution composed of multiphases of M-type ferrite and other phases. Fraction of M-type ferrite in the samples was depending on mole ratio in the raw material. Curie temperatures of Co-free and Co-substituted M-type ferrite were approximately 432 degrees C and 428 degrees C, respectively. Saturation magnetisation of the Co-free Na0.5La0.5 system M-type ferrite was approximately 72.5 emu/g.

Nd 2 Fe 14 B is one of the most popular permanent magnets (PMs) possessing the best energy product (BH) max among the common PM materials.

Effective low melting-point Dy-diffusion for grain boundary diffusion treatment for enhancing coercivity in Nd-Fe-B-type magnet was found in the DyF3-LiF binary system. Efficacy of the low melting-point (DyF3-LiF) diffusion source as Dy-diffusion source for enhancing coercivity in the diffusion processed Nd-Fe-B-type magnet was investigated. Speedier and more profound coercivity enhancement in the Nd-Fe-B-type magnet was achieved by grain boundary diffusion using the low melting-point (DyF3-LiF) diffusion source with respect to solid DyF3 single salt. Since the liquid in the (DyF3 -LiF) mixture contained plenty of Dy atoms already freed from DyF3 and they were in better contact with magnet, speedier and profuse diffusion of Dy atoms through Nd-rich grain boundary was possible in the magnet coated with low melting-point (DyF3-LiF) diffusion source.

A grain boundary diffusion process (GBDP) was adopted to improve magnetic properties of Dy-free highly coercive Nd-Fe-B permanent magnet by coating thin layers of Nd and Cu in grain boundaries,For GBDP of Nd and Cu,Nd and Cu were coated by wet process,e.g.,electrochemical and electroless on Nd-Fe-B magnets,which was fabricated by hot-deformed/die-upset with meltspun specimen.Heat treatment was performed for 20 min at 600 ℃ followed by several different cooling conditions.The cooling conditions after heat treatment were varied to understand distribution and microstructural effects of Nd and Cu species in grain boundaries.The coercivity increased from 1.565 to 1.637 T in oil cooling rate but remanence decreased,while remanence jumped with little decrease in coercivity in furnace cooling.Microstructure analyses suggested that the coercivity was closely related to the cooling rate as well as distribution of Nd.The mechanism of coercivity enhancement due to the cooling rate was discussed based on the results presented here and those in the literature.

Texture development in the Ce-substituted Nd-Fe-B-type die-upset hybrid magnet, which was fabricated using starting materials of Ce-substituted (Nd0.55Ce0.45)(15) Fe72.2Co6.6Ga0.6B5.6 HDDR-treated alloy powder and meltspun Nd-13.6Fe73.6Co Ga-6.6(0).B-6(5.6) flakes (MQU-F) without Ce, was investigated. Noticeably better texture developed in die-upset magnet from the MQU-F flakes alone with respect to magnet from the HDDR-treated powder alone. Better texture developed also in the MQU-F flake regions in the die-upset hybrid magnet with respect to the HDDR particle regions, and overall texture in the hybrid magnet was dominantly controlled by the texture in the MQU-F flake regions. Overall texture in the hybrid magnet was not as good as the weighted average of texture expected from texture of the single alloy magnets from the IIDDR powder alone and the MQU-F flakes alone.

Ce-containing (Nd,Ce)-Fe-B-type hybrid magnet was fabricated by die-upset technique using two different types of materials: Ce-substituted (Nd0.55Ce0.45)(15)Fe72.2Co6.6Ga0.6B5.6 HDDR powder and melt-spun Nd13.6Fe73.6CO6.6Ga0.6B5.6 flakes. Magnetic performance of the hybrid magnet was superior to that of the single alloy magnet when they had identical overall composition: H-i(c)= 6.8 kOe, M-r = 10.6 kG, and (BH)(max) = 19.8 MGOe for the Ce-containing hybrid magnet and H-i(c )= 3.9 kOe, M-r = 9.8 kG, and (BH)(max) = 11.6 MGOe for the magnet from the single alloy HDDR powder. Die-upset hybrid magnet consisting of two constituent materials showed smooth and single-material-like demagnetization behavior, and this was attributed to the exchange interaction between neighbouring grains in the magnet.

Nd-Fe-B-type die-upset magnet with high electrical resistivity was fabricated by hot-deforming the mixture of melt-spun Nd-Fe-B-type flakes (MQU-F: Nd13.6Fe73.6Co6.6Ga0.6B5.6) and Dy-containing salts: eutectic (DyF3-LiF) salt mixture and DyF3 single salt. Profound electrical resistivity enhancement was feasible in the Nd-Fe-B-type die-upset magnet by adding Dy-containing salts. More profound electrical resistivity enhancement was achieved in the magnet added with dielectric eutectic (DyF3-LiF) salt mixture with respect to the magnet added with single DyF3 salt. This was attributed to better electrical insulation between the flakes by forming more continuous coverage of the flake interface with the easily melted dielectric salt. Coercivity of the die-upset magnet was also profoundly enhanced by optimal addition of Dy-containing salts, and this was attributed to substitution of some Nd in the Nd2Fe14B-type grains near flake surface by Dy atoms from the added salt. Kerr microscopy revealed that for both the magnets with or without salt addition, formation of reverse domain initiated mostly inside the flake. Reversed domain started to form at higher reverse field for the magnet added with Dy-containing salt than for the magnet without salt addition. Practical demagnetization occurred largely by formation of new reverse domains at random places rather than enlargement of previously formed reverse domain for both the magnets with or without salt addition.

Anisotropic ceramics-bonded Nd-Fe-B-type magnet with high electrical resistivity was fabricated with the intention of suppressing induction of eddy current thus lowering operating temperature of the magnet used as rotor magnet in high-speed motor. The ceramics-bonded Nd-Fe-B-type magnet was fabricated by consolidating mixture of hydrogenation decomposition desorption recombination powder and oxide ceramics with the low melting point. Anisotropic ceramics-bonded (15 vol%) Nd-Fe-B-type magnet had remarkably enhanced electrical resistivity (similar to 730 mu Omega.cm) with respect to the magnet without ceramics binder (similar to 220 mu Omega.cm). Thanks to lowtemperature consolidation of Nd-Fe-B-type particles using oxide ceramic binder with low melting point, the detrimental reaction between the magnetic particle surface and oxide was profoundly suppressed, thus retaining the high coercivity of initial magnetic particles even in the ceramics-bonded magnets. Ceramics-bonded (15 vol%) magnet, which had good room temperature magnetic performance (H-i(c) = 12.7 kOe, M-r = 9.4 kG, and (BH)(max) = 17.0 MGOe) still had reasonably good performance [H-i(c) = 5.0 kOe, M-r = 8.7 kG, and (BH)(max) = 10.2 MGOe] at 150 degrees C.

Nd-Fe-B-type magnet is exclusively used as a rotor magnet in the traction motor of hybrid electric vehicle（HEV） and electric vehicle（EV）,but its overly high operating temperature is a lingering problem attached to the magnet.The major cause of the high operating temperature is eddy current,which is readily generated in the highly conductive metallic magnet under alternating magnetic field from stator ripple.In this article,temperature rise in the Nd-Fe-B-type magnet with varying electrical resistivity under alternating magnetic field is discussed with the intention of highlighting the importance of enhancing the electrical resistivity for reducing the operating temperature of the Nd-Fe-B-type rotor magnet.Temperature rise in the Nd-Fe-B-type magnet（dielectric salt-added die-upset magnet） with high electrical resistivity is noticeably lower compared to the magnet（commercial sintered rotor magnet） with lower electrical resistivity,substantiating the theory that enhancing the electrical resistivity in the rotor magnet is fairly effective for suppressing the over-rise of its operating temperature during operation.Die-upset process is revealed to be particularly pertinent for the fabrication of highly dense salt-added magnet with high electrical resistivity.

Variation of intrinsic coercivity of the HDDR-treated Nd12.5Fe80.8B6.4Ga0.3 alloy after heating in various modes was investigated. Influence of vacuum degree and cooling after heating on the coercivity of the HDDR material was examined. The heat-treated HDDR material had consistently higher coercivity when it was quenched after heating compared to when slow-cooled. Higher coercivity in the quenched material was attributable to the grain boundary with lower Fe content. HDDR-treated material heated in high vacuum showed consistently higher coercivity than the material heated in lower vacuum, and this was attributed to less heavily oxidized surface. Reduced coercivity of the HDDR-treated material heated at moderate temperature was noticeably recovered at higher temperature, and this was attributed to lower Fe content in the grain boundary.

Texture development in hydrogenation–disproportionation–desorption–recombination (HDDR)-processed NdFeB magnets is discussed with a focus on the crystallographic alignment of the Nd2Fe14B and Fe2B phases. By careful control of the dynamic HDDR-processing parameters, shape changes in the HDDR-treated samples were minimized by suitable modifications to the sample dimensions to trace the orientation relation between the c-axes of the Nd2Fe14B and Fe2B phases. The magnetic anisotropy was significantly dependent on the hydrogen pressure during the HD process and highly aligned Fe2B phases, with the same c-axis orientation as the initial Nd2Fe14B phase, were formed under the dynamic HDDR treatment at a hydrogen pressure of 30 kPa under controlled HD conditions. During the DR reaction following the HD process, the orientation of the initial Nd2Fe14B was recovered by a highly aligned Fe2B phase. It was demonstrated that the Fe2B phase could mediate texture during the HDDR processing owing to the atomic configuration and crystallographic relation between the Nd2Fe14B and Fe2B phases.

Magnetic interaction effects on the magnetic switching volume for Cr/CoSm/Cr magnetic films deposited with different sputtering pressures (7 similar to 10 mTorr) of underlayer (Cr) was investigated. The magnetic interaction mechanism of all the Cr/CoSm/Cr magnetic films investigated in this study was dipole interaction, and the dipole interaction was stronger at the samples fabricated at low sputtering pressure. The magnetic field (H-Delta M) at which the maximum dipole interaction occurred in each sample was larger than the coercivity (H-c). The magnetic switching volume was larger for the sample with stronger dipole interaction, and it showed maximum value in the H-Delta M region and then decreased in higher field region above the H-Delta M. Therefore, it is considered that the intensity of the dipole interaction is closely related to the magnetic switching volume.

Ceramics-bonded magnet with remarkably high electrical resistivity was fabricated by hot-pressing the mixture of Nd13.6Fe73.6Co6.6Ga0.6B5.6 alloy melt-spun flakes and dielectric Bi2O3-SiO2-B2O3 ceramics powder with low melting point. Coercivity of the ceramics-bonded magnet decreased with increasing the addition of ceramics binder, and this was attributed to the increased demagnetizing factor. Thin oxidized layer on the flake surface formed by reaction between the flake and oxide binder also contributed to reducing coercivity in the ceramics-bonded magnet. Highly resistive ceramics-bonded magnet containing 30 vol% ceramics binder still had good magnetic performance and high mechanical strength at 175 oC: iHc = 5 kOe, Mr = 4.8 kG, (BH)max = 4.3 MGOe, and over 900 MPa.

Two types of re-substituted (FeCo)2B-type alloys with stoichiometric (Fe0.675Co0.3Re0.025)2B and 1.5 at% boron-excess compositions were prepared by suction casting and subsequent mechanical milling and annealing. Phase evolution in the (FeCo)2B-type alloys in the course of mechanical milling and annealing was investigated. Single-phase re-substituted (FeCo)2B-type material was prepared from stoichiometric (Fe0.675Co0.3Re0.025)2B by the combination of suction casting and mechanical milling. The two alloys had the same phase constitution of (Fe, Co)2B-type and amorphous phases in heavily milled state. However, the heavily milled alloys had different phase constitution after annealing; the stoichiometric alloy consisted of single (Fe, Co)2B-type phase, while the boron-excess alloy consisted of two-phase mixture of (Fe, Co)2B-type and (Fe, Co)B-type phases. The two heavily milled alloys had similar coercivity of around 1 kOe after full annealing.

The grain boundary diffusion (GBD) process with rare-earth hydride was performed to increase the coercivity of hydrogenation–disproportionation–desorption–recombination (HDDR) powder. Before the GBD process, we investigated the effect of post-annealing of the initial HDDR powder on its magnetic properties. Low-temperature annealing reduced the coercivity of the HDDR powder. However, the coercivity decline decreased with increasing annealing temperature, becoming similar to that of the initial powder at 900°C. After the GBD process at 850°C for 1 h, the coercivity increased by about 4 kOe with 4 wt.% NdH x -Cu, forming a thick and continuous grain boundary phase. In addition, the coercivity and remanence of the HDDR powder produced by the GBD process with NdH x -Cu were higher when using NdH x in spite of the same amount of diffusion as at 2 wt.%.

High coercive Nd-Fe-B magnets are indispensable materials to traction motors in hybrid and electric vehicles. However, substitution of heavy rare earth element (HRE) such as Dy or Tb for Nd has been necessary for high coercive Nd-Fe-B magnets despite the critical problem of supply and demand with HRE. Therefore, HRE-lean and/or HRE-free high coercive Nd-Fe-B magnets have drawn a great attention to solve HRE resource problem in the automotive industry [1]. To achieve high coercivity with reducing HRE, the control of microstructures, such as grain size and grain boundary, is of significant importance. Melt-spinning and hydrogenation-disproportionation–desorption-recombination (HDDR) are known as quite suitable method to decrease grain size down to the single domain size $( \sim 250$ nm). It addition, hot-deformation is known as a useful method to obtain anisotropic magnets with magnetic powders produced by these methods. On the other hand, the coercivity of hot-deformed Nd-Fe-B magnets was too low to be used for motors of hybrid and electric vehicles even though they had submicron grain size. It is because of the presence of crystallographic defects, low anisotropic energy at the grain surface, and exchange coupling between neighboring grains, which could be improved by grain boundary diffusion process (GBDP) with HRE compounds or non-magnetic materials. However, the GBDP of ultrafine grained materials produced with melt-spun powders should be done at a temperature, lower than about $700 ^{circ}\mathrm {C}$. The GBDP above $700 ^{circ}\mathrm {C}$ could induce remarkable grain growth and low coercivity. On the other hands, HDDR powder has relatively large grains about $250 \sim 400$ nm compared to melt-spun powders, so the grain growth does not occur at temperature up to about $850 ^{circ}\mathrm {C}$. Therefore, it can be expected that hot-deformed magnets produced with HDDR powder have an advantage for subsequent GBDP compared to that of the melt-spun powder. However, there were only a few studies examining the hot-deformation behavior of HDDR powders and the reported magnetic properties were relatively poor. On the other hand, numbers of research reveal that rare earth-rich phase is critical factor to the texture formation during the hot-deformation process. So, it is expected that the microstructure of initial alloy could effect on the deformation behavior during hot-deformation process. Therefore, in this study, effect of initial alloy on microstructure and magnetic properties during hot-deformation of Nd-Fe-B HDDR powder was investigated.Alloy with composition of Nd 12.5 Fe $_{bal}$ Ga 0.3 Nb 0.2 B 6.4 prepared by strip-casting process was used as a starting material. The alloy was subjected to HDDR process after pre-annealing (SC-HT) at $1100 ^{circ}\mathrm {C}$ for 12 hours or without pre-annealing (SC). Produced HDDR powders were then hot-pressed at $700 ^{circ}$ Cunder 400 MPa in a vacuum. The cylindrical compact with 7 mm in diameter and 6 mm in height were then die-upsetted at $800 ^{circ}\mathrm {C}$ with deformation degree of $\varepsilon =1.5$ and deformation rate of $\varepsilon $'$= \sim 0.01 \mathrm {s}^{-1}$. Figure 1 shows magnetic properties and microstructure of each initial HDDR powder. The remanence of the powders is similar. However, their coercivity is diffenent about 1 kOe. This could be attributed to that the SC&underscore;HT HDDR powder have non-uniform and discontinuous Nd-rich phase distribution in grain boundary which can decrease coercivity by magnetic coupling between neighbor grains although it is not clear from SEM image as shown in Fig. 1(b) and (c).Figure 2 shows demagnetization curve ((Fig. 2(a)) and microstructure of hot-deformed magnets which is deformed using SC&underscore;HDDR powder (Fig. 2(b)) and SC&underscore;HT (Fig.2(c)), respectively. The remanence of hot-deformed magnet with SC&underscore;HT HDDR powder is lower than that with SC&underscore;HDDR powder. The difference in the remenance could be attributed to the fact that the distribution of the Nd-rich phase of the grain boundary may be affected by the hot deformation behavior which could be confirmed from Fig. 2(b) and (c). This non-uniformed Nd-rich phase of grain boundary can make it difficult to grain boundary sliding during the hot-deformation process, which induce poor grain alignment and low remanence. Based upon these results, effect of initial alloy on microstructure and magnetic properties during hot-deformation of Nd-Fe-B HDDR powder will be discussed. Fig. 1. Demagnetization curve (a) of obtained HDDR powders and FESEM images of SC&underscore;HDDR powder (b) and SC&underscore;HT HDDR powder.Fig. 2. Demagnetization curve (a) of hot-deformed magnet and FESEM images of hot-deformed magnet with SC&underscore;HDDR powder (b) and hot-deformed magnet with SC-HT HDDR powder (c).

Feasibility of the electrophoresis deposition (EPD) technique for homogeneous and adhesive deposition of DyF3 particles on the Nd-Fe-B-type particles was studied, and coercivity enhancement in the diffusion-treated Nd-Fe-B-type particles deposited with DyF3 by EPD was investigated. HDDR-treated Nd12.5Fe80.6B6.4Ga0.3Nb0.2 particles were deposited with DyF3 particles by EPD. More homogeneous and adhesive deposition of DyF3 particles on the surface of Nd-Fe-B particles was made by the EPD with respect to conventional dip-coating, and this led to more active and homogeneous diffusion of Dy. More profound coercivity enhancement was achieved in the diffusion-treated Nd-Fe-B-type particles deposited with DyF3 by EPD compared to dip-coated particles.

Bulk-type body-centered-tetragonal Fe-Co alloy was synthesised by utilising a conventional alloy preparation technologies, such as melting, solidification, and homogenising treatments, and its magnetic properties were investigated. In the (Fe100-x,Co-x)(1-y)C-y alloy, the composition range, from which single phase body-centered-tetragonal alloy (martensite phase) was obtained, was severely limited: Co content x = 2.5, and C content y = 0.062. Tetragonality(c/a) of the synthesised body-centered-tetragonal (Fe97.5CO2.5)(0.938)C-0.062 alloy was 1.05. Magnetocrystalline anisotropy constant (K-1) of the body-centered-tetragonal (Fe97.5Co2.5)(0.938)C-0.062 alloy was measured to be 9.8 x 10(5) J/m(3)), which was 3.1 time as high as the pure iron (alpha-Fe).