2017 Volume 58 Issue 5 Pages 825-828
This paper reports improvement of the coercivity of Nd-Fe-B powder by Nd-Cu sputtering. The total rare earth element content of the Nd-Fe-B powder increased by 0.8 mass% after Nd-Cu sputtering, and the sputtered Nd-Cu was observed at the surface of Nd-Fe-B powder by microstructural observation. The coercivity (Hc) of the Nd-Cu-sputtered Nd-Fe-B powder improved after annealing at temperatures above 600℃, and Hc of the sputtered powder reached a maximum value of 1312 kA m−1 after annealing at 900℃. This Hc value was 470 kA m−1 higher than that of Nd-Fe-B powder without sputtering. Isotropic green compacts were prepared from Nd-Fe-B powders with and without Nd-Cu sputtering, and the sputtering also improved Hc in the compacts. The compact prepared from powder with sputtering had Hc of 630 kA m−1 higher than that of the compact prepared from powder without sputtering after annealing at 900℃. Microstructure observation showed that the improvement of Hc was due to an increase in the Nd-rich phase in the compact by Nd-Cu sputtering.
Nd-Fe-B sintered magnets are used in many applications including in the motors of hybrid vehicles, an application that demands high thermal stability. Accordingly, high coercivity of about 2400 kA m−1 at room temperature is required for use at high temperatures. The coercivity of Nd-Fe-B sintered magnets is improved by the addition of Dy, but this also decreases the magnetization of the magnets. Therefore, the development of Dy-free Nd-Fe-B sintered magnets with high coercivity is needed.
One useful method for improving the coercivity of Nd-Fe-B sintered magnet is to refine the grains of the Nd2Fe14B phase.1–4) Une and Sagawa prepared fine Nd-Fe-B powders with an average particle size of approximately 1.1 µm from strip-cast (SC) Nd-Fe-B alloy by He-gas jet milling (He-JM)5) and achieved a high coercivity of 1590 kA m−1 in grain-refined Nd-Fe-B sintered magnets without Dy. Thus, the preparation of ultrafine Nd-Fe-B powder is effective for further enhancing the coercivity of sintered magnets. Nakamura et al. fabricated ultrafine Nd-Fe-B powders with an average particle size of 0.33 µm from Nd-Fe-B alloy subjected to hydrogenation-disproportionation-desorption-recombination (HDDR) treatment through hydrogen decrepitation (HD) and He-JM.6–8) They reported that the coercivity of the ultrafine powder was dependent on annealing temperature and that the maximum coercivity of the ultrafine powder was 1210 kA m−1. Une et al. prepared ultrafine-grained Nd-Fe-B sintered magnets and reported a coercivity of around 1034 kA m−1.9) Nakamura et al. attributed the low coercivity of the ultrafine powder to its lack of Nd-rich phase because of an insufficient rare earth element content. It is well known that Nd-rich phase liquefies and spreads at the surface of Nd2Fe14B grains during heat treatment, thus increasing the coercivity.10) To increase the amount of Nd-rich phase at the surface of the Nd-Fe-B ultrafine powder, our group applied a grain boundary diffusion (GBD) process between the HDDR+HD and He-JM processes.11) The GBD-treated ultrafine powder exhibited a high coercivity of 1592 kA m−1 because of the increased TRE content. However, applying GBD-treatment to our powder preparation process caused wide size distribution of the Nd-Fe-B powders.11)
Sputtering of Nd-rich alloy on the surface of Nd-Fe-B powder can be another effective method for increasing the amount of Nd-rich phase. Sputtering can deposit Nd-rich alloy having wide range of compositions, and can possible deposit thin and homogeneous Nd-rich layer onto Nd-Fe-B powders with maintaining powder size distribution. However, there have been no reports on the use of this technique. We therefore focused on applying sputtering to Nd-Fe-B powder in order to increase the amount of Nd-rich phase.
In this study, Nd-Cu alloy, which is an effective Nd-rich alloy for improving coercivity by the GBD process12,13), was sputtered onto Nd-Fe-B powder with an average particle size of 4.5 µm and the changes in the magnetic properties and microstructure were investigated.
The starting material was an SC Nd-Fe-B alloy with a composition of Nd27.4Pr0.1Fe70.4B1.05Al0.05Nb0.25 (mass%). The alloy was exposed to a hydrogen atmosphere in an HD process and was then jet-milled using N2 gas to fabricate a powder with an average particle size of 4.5 µm. Magnetron sputtering with a base pressure of <7 × 10−5 Pa was used to deposit Nd-Cu alloy on the Nd-Fe-B powder using Nd85Cu15 (mass%) alloy as the sputtering target. Nd-Cu was sputtered onto the Nd-Fe-B powder for 15 min and then the powder was stirred for 1 min. The sputtering and stirring were repeated 8 times. After sputtering, the powders were annealed at 500–1000℃ for 30 min under vacuum (<2 × 10−3 Pa). Isotropic green compacts with dimensions of 3.5 × 7.0 × 1.5 mm were prepared from Nd-Fe-B powders with or without Nd-Cu sputtering by pressing at 107 kg m−2. The green compacts were annealed at 500–1000℃ for 30 min under vacuum.
The particle size distributions of the powders were measured by a dry laser diffraction method. The compositions of the powders were determined by X-ray fluorescence (XRF) analysis. Microstructures were observed by field-emission scanning electron microscopy (FE-SEM) and magnetic properties were measured by a vibrating sample magnetometer after applying a pulsed magnetic field of 6360 kA m−1.
Figure 1 shows the particle size distributions of the Nd-Fe-B powders with and without Nd-Cu sputtering. The distribution profiles are almost identical. The average particle size is defined as the size at which the cumulative distribution reaches 50% (d50). In addition, the relative span (RS) is defined as a dimensionless value representing the particle size distribution width by eq. (1).
| \[RS=(d_{90} - d_{10} )/ d_{50}\] | (1) |
Particle size distribution of Nd-Fe-B powders with and without Nd-Cu sputtering.
Here, d10 and d90 are defined as the size at which the cumulative distribution reaches 10% and 90%, respectively. As the particle size distribution becomes narrower, RS approaches zero. The d10, d50, d90, and RS of the Nd-Fe-B powders with and without Nd-Cu sputtering are shown in Table 1. The Nd-Fe-B powders with and without sputtering had d50 of 4.66 and 4.54, respectively, indicating that the particle size of the Nd-Fe-B powder was slightly increased by sputtering. The RS of powders with and without sputtering were 1.35 and 1.40, respectively. This indicates that the particle size distribution was not changed by Nd-Cu sputtering.
| Sample | d10 (µm) | d50 (µm) | d90 (µm) | RS |
|---|---|---|---|---|
| Without sputtering | 1.57 | 4.54 | 7.93 | 1.40 |
| With sputtering | 1.67 | 4.66 | 7.99 | 1.35 |
The compositions of the Nd-Fe-B powders with and without Nd-Cu sputtering were analyzed by XRF, and it was found that the total rare earth element (TRE) content increased by 0.8 mass% after Nd-Cu sputtering for 120 min. Figure 2 shows secondary electron (SE) images and elemental mappings of the Nd-Fe-B powder. Figure 2(a) and (c) were taken from the powder without sputtering, and Fig. 2(b), (d), and (e) were taken from the powder with sputtering. Comparing Fig. 2(a) and (b), no obvious morphological changes were observed after sputtering. From the elemental mappings (Fig. 2(c) and (d)), Cu was detected in some particles after sputtering. Figure 2(e) shows a magnified image of the circled particle in Fig. 2(d), in which Cu was detected at the surface of the particle.
SE and elemental mapping images of Nd-Fe-B powders; (a) (c) without and (b) (d) (e) with sputtering.
Figure 3 shows highly magnified SE images of Nd-Fe-B powders with and without sputtering. A smooth surface was observed on the powder without sputtering (Fig. 3(a)), and a change in the surface morphology was observed after sputtering (Fig. 3(b)). Surface roughness of several nanometers was observed, as indicated by the white arrow. Based on the SE images (Fig. 3) and elemental mappings (Fig. 2(e)), nanometer-sized Nd-Cu particles were deposited on the surface of Nd-Fe-B powder by sputtering. These results indicate that the sputtering process was effective for increasing the TRE content of Nd-Fe-B powder while maintaining the particle size distribution. The Nd-Fe-B powders were then annealed and the changes in the magnetic properties were measured.
Magnified SE images of the Nd-Fe-B powder (a) without and (b) with sputtering.
Figure 4 shows demagnetization curves of the Nd-Fe-B powders before and after annealing at 900℃. Before annealing (Fig. 4(a)), the coercivities (Hc) of the powders with and without sputtering were 36 kA m−1 and 33 kA m−1, respectively. The remanences (σr) of both powders were almost the same, and the magnetizations of the powders with and without sputtering measured under an applied magnetic field of 1500 kA m−1 (σ1500) were 128 A m2 kg−1 and 133 A m2 kg−1, respectively. This shows that the magnetic properties were unchanged by sputtering. In contrast, the coercivity was dramatically changed by annealing. Figure 4(b) shows the demagnetization curves of powders after annealing at 900℃. In both powders with sputtering and powders without sputtering, Hc was increased by annealing at 900℃. The increase in the coercivity was larger in the powder with sputtering than in the powder without sputtering. Figure 5 shows Hc and σr of the powders with and without sputtering as a function of annealing temperature. In both powders, Hc was increased by annealing above 600℃. Hc of the powder without sputtering reached 970 kA m−1 after annealing at 700℃, and was decreased by annealing above 800℃. In contrast, Hc of the powder with sputtering was increased by annealing at 800–900℃ and reached a maximum value of 1312 kA m−1 after annealing at 900℃. This maximum value was 470 kA m−1 higher than that of the powder without sputtering after annealing at 900℃. Figure 5(b) shows remanence σr as a function of annealing temperature. In both the powder with sputtering and the powder without sputtering, the remanence was increased by annealing above 600℃, the same as Hc. For annealing above 700℃, σr remained unchanged at a value of around 77 A m2 kg−1 in both powders. These results indicate that Nd-Cu sputtering and subsequent annealing can improve the coercivity of Nd-Fe-B powder without decreasing its remanence.
Demagnetization curves of powders with and without Nd-Cu sputtering (a) before and (b) after annealing at 900℃.
(a) Coercivity (Hc) and (b) remanence (σr) of the powders with and without sputtering versus annealing temperature.
To investigate the effect of Nd-Cu sputtering on compacts, isotropic green compacts were prepared from Nd-Fe-B powders with and without sputtering, and the compacts were annealed. Figure 6 shows the magnetic properties of the compacts as a function of annealing temperature. They show the same tendency as those of powders (Fig. 5). Hc of the compact prepared from powder without sputtering reached a maximum of 840 kA m−1 after annealing at 800℃, and it was decreased by annealing above 900℃. In contrast, Hc of the compact prepared from powder with sputtering increased by annealing at 800–900℃, and it reached a maximum value of 1260 kA m−1 after annealing at 900℃. This value was 630 kA m−1 higher than that of the compact prepared from powder without sputtering. The σr remained unchanged at around 75 A m2 kg−1 in both compacts after annealing above 700℃. Thus, Nd-Cu sputtering was also effective for improving the coercivity of compacts without decreasing remanence.
(a) Coercivity (Hc) and (b) remanence (σr) of isotropic compacts prepared from powders with and without sputtering versus annealing temperature.
Then, the microstructure of the compacts was observed. Figure 7 shows backscattered electron (BSE) images of the polished surface of the compacts prepared from powders with and without sputtering after annealing at 900℃. Dark and bright regions were observed in the BSE images which correspond to Nd2Fe14B and Nd-rich phases, respectively. Comparing Fig. 7(a) and (b), the amount of Nd-rich phase in the compact prepared from powder with sputtering was larger than that in the compact prepared from powder without sputtering. The area fraction of Nd-rich phase was calculated from these images by the point counting method, and the area fraction of Nd-rich phase of the compacts with and without sputtering were 5.2% and 2.3%, respectively. As described above, the TRE content of Nd-Fe-B powder was 27.5 mass%, which was 0.8 mass% higher than the stoichiometric composition of Nd26.7Fe72.3B1.0(mass%), and the TRE content of Nd-Fe-B powder was increased by 0.8 mass% by Nd-Cu sputtering. Both area fraction and TRE content approximately doubled as a result of sputtering, and therefore the BSE images in Fig. 7(a) and (b) are in agreement with the XRF analysis results. Figure 7(c) and (d) show magnified BSE images of the areas enclosed by white squares in Fig. 7(a) and (b), respectively. Nd-rich phase was observed at some triple junctions with only a small amount observed at grain boundaries in the compact prepared from powder without sputtering (Fig. 7(c)). In contrast, Nd-rich phase was observed at triple junctions and grain boundaries between Nd2Fe14B phases in the compact prepared from powder with sputtering (Fig. 7(d)). It is thought that the Nd-Cu phase liquefied and infiltrated the grain boundaries during annealing which led to increased coercivity of the compact. Consequently, the amount of Nd-rich phase in Nd-Fe-B powder was increased by Nd-Cu sputtering, and the coercivity was also improved after annealing for sputtered samples.
BSE images of the polished surface of compacts (a) (c) without and (b) (d) with sputtering after annealing at 900℃.
These results showed that sputtering of Nd-rich alloy is effective for increasing the TRE content while maintaining the particle size distribution of Nd-Fe-B powder, as well as for increasing the coercivity of the powder after annealing.
In this study, the effects of Nd-Cu sputtering on the magnetic properties and microstructure of Nd-Fe-B powder were investigated. The particle size distribution did not change due to Nd-Cu sputtering. Sputtered nanometer-sized Nd-Cu particles were found at the surface of the powder, and the TRE content was increased by 0.8 mass% by sputtering for 120 min. The coercivity of Nd-Cu sputtered Nd-Fe-B powder improved after annealing above 600℃, and the coercivity of the powder reached a maximum of 1312 kA m−1 by annealing at 900℃. This value was 470 kA m−1 higher than that of the powder without sputtering. Isotropic compacts were prepared from these powders, and the coercivities of the compacts were improved by Nd-Cu sputtering. After annealing at 900℃, the coercivity of the compact prepared from powder with sputtering was 630 kA m−1 higher than that of the compact prepared from powder without sputtering. FE-SEM observation showed an increase in Nd-rich phase and the presence of Nd-rich phase at grain boundaries in the compact prepared from the sputtered powder. Thus, sputtering is an effective process for increasing the TRE content while maintaining the particle size distribution for Nd-Fe-B powder, thus leading to improved coercivity after annealing.
This study was supported by “Future Pioneering Projects/Development of magnetic material technology for high-efficiency motors” from NEDO, JAPAN.
