- Email: [email protected]
Coercivity enhancement of hot-deformed Nd-Fe-B magnets by the eutectic grain boundary diffusion process using Nd62Dy20Al18 alloy
Scripta Materialia 129 (2017) 44–47
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Regular article
Coercivity enhancement of hot-deformed Nd-Fe-B magnets by the eutectic grain boundary diffusion process using Nd62Dy20Al18 alloy Lihua Liu a,b, H. Sepehri-Amin a, T. Ohkubo a, M. Yano c, A. Kato c, N. Sakuma c, T. Shoji c, K. Hono a,b,⁎ a b c
Elements Strategy Initiative Center for Magnetic Materials, National Institute of Materials Science, Tsukuba 305-0047, Japan Graduate School of Pure and Applied Science, University of Tsukuba, Tsukuba 305-8577, Japan Toyota Motor Corporation, Advanced Material Engineering Div., Susono 410-1193, Japan
a r t i c l e
i n f o
Article history: Received 24 August 2016 Received in revised form 16 October 2016 Accepted 16 October 2016 Available online xxxx Keywords: Permanent magnets Nd-Fe-B Coercivity Hot-deformed magnets
a b s t r a c t The eutectic grain boundary diffusion process was applied to a 2-mm-thick hot-deformed Nd-Fe-B magnets using Nd62Dy20Al18 alloy as a diffusion source, realizing the coercivity enhancement from 0.91 T to 2.75 T with relatively small remanence deterioration from 1.50 T to 1.30 T. In contrast, the conventional grain boundary diffusion process using Dy-vapor resulted in the degradation of coercivity as the grains are catastrophically coarsened at the processing temperature of 900 °C. Scanning transmission electron microscopy showed the formation of Dy-rich shell at the sides of the Nd2Fe14B grains and the diffusion of Al into the Nd2Fe14B grains, explaining the significant improvement in coercivity. © 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Nd-Fe-B based permanent magnets have been used for traction motors in (hybrid) electric vehicles owing to its ability to induce strong magnetic flux density. Since the application environment causes temperature increase under demagnetization fields, not only high remanence but also high coercivity is required. In order to attain the coercivity around 0.8 T at the operation temperature of 200 °C, about 10 wt% Dy is alloyed in commercial Nd-Fe-B sintered magnets with room temperature coercivity higher than 3 T. The partial substitution of Nd with Dy increases the anisotropy field of (Nd1-xDyx)2Fe14B, thereby increasing the coercivity of the magnet [1]. However, the magnetic moment of Dy is antiferromagnetically coupled with that of Fe atom in the (Nd1-xDyx)2Fe14B compound, resulting in a deterioration of magnetization as coercivity increases [2]. Since the natural abundance of Dy is quite limited compared to that of Nd, the development of Dy-free or Dy-saving high coercivity Nd-Fe-B permanent magnets is strongly needed. Grain size refinement of Nd-Fe-B sintered magnets is one of the effective methods to increase the coercivity without using Dy [3–5]. Sagawa et al. [6] have developed the press-less process (PLP) for sintering the submicron sized powders that were processed by helium-jet milling, realizing the coercivity of 2 T in the sintered magnets with 1 μm grain size. However, further refinement of grain size is difficult using the conventional powder metallurgy route considering the severe oxidation of Nd-rich phases. On the other hand, hot-deformed ⁎ Corresponding author at: Elements Strategy Initiative Center for Magnetic Materials, National Institute of Materials Science, Tsukuba 305-0047, Japan. E-mail address: [email protected] (K. Hono).
http://dx.doi.org/10.1016/j.scriptamat.2016.10.020 1359-6462/© 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
magnets produced from isotropic melt-spun ribbons followed by proper hot-press and hot-deformation exhibit strong [001] texture and the fine grain size that is comparable to the single domain size of Nd2Fe14B phase [7–9]. Unlike the conventional sintering process, powder size can be larger than 100 μm, inside which nano-sized grains are dispersed; therefore, individual grain are not exposed to oxygen even if the process is carried out without rigorous control of oxygen. The typical coercivity value, μ0Hc, of hot-deformed magnets is around 1.2 T with a remanence, μ0Mr, of 1.4 T, which are comparable to those of sintered magnets [10,11]. Strong exchange interaction between Nd2Fe14B grains through ferromagnetic intergranular phase is the main reason for this relatively low coercivity of the hot-deformed magnets [12–15]. More recently, Liu et al. reported a good correlation between the Nd content in the intergranular phase and the coercivity using atom probe tomography [11], and they suggested that the modification of the intergranular phase would enhance the coercivity further. Sepehri-Amin et al. demonstrated a dramatic enhancement of coercivity by the eutectic grain boundary diffusion (GBD) process for hydrogen-disproportionation-decomposition-recombination (HDDR) processed Nd-Fe-B powders using Nd70Cu30 eutectic alloy as a diffusion source [4]. This eutectic GBD process was later extended to Nd-Fe-B hotdeformed magnets [16], which enhanced the coercivity from 1.5 T to 2.3 T without using Dy. In order to explore the extent of coercivity enhancement by this eutectic GBD process, Liu et al. [17] applied various Nd1-xMx eutectic alloys to the 2-mm-thick Nd-Fe-B hot-deformed magnets, obtaining the highest coercivity of 2.5 T at room temperature using Nd90Al10 as a diffusion source. However, the coercivity measured at 200 °C was only 0.6 T because of poor temperature dependence of
L. Liu et al. / Scripta Materialia 129 (2017) 44–47
μ0Hc partly due to the dissolution of Al in the Nd2Fe14B phase. A higher level of coercivity is required for the applications such as automotive traction motors and wind turbines. Heavy rare earth (HRE) grain boundary diffusion process was developed for enhancing the coercivity of sintered magnets with the grain size of about 5 μm using Dy-vapor, Dy2O3, DyF3, Tb3O4 and TbF4 [18– 21]. However, this method cannot be employed to ultrafine-grainsized hot-deformed Nd-Fe-B magnets since the high temperature annealing required for the HRE GBDP results in catastrophic grain growth. Recently, Sepehri-Amin reported that Dy-rich interface can be developed without any grain coarsening in Nd-Fe-B hot-deformed magnet using Nd60Dy20Cu20 powder as a diffusion source followed by annealing at low temperature of 650 °C [22]. On the other hand, along with the remarkable enhancements of coercivity by using Nd-Cu [4,16,23], Nd-Al [17], Pr-Cu [23–25] or Nd-DyCu [22] as the diffusion source, the diffusion-processed magnets always experienced a significant degradation in remanence due to the infiltration of excessive non-ferromagnetic RE-rich liquid phase into the magnet. Akiya et al. [26] applied an expansion constraint on the c-plane of the magnet during infiltration of RE-rich liquid phase and achieved coercivity enhancement while retaining a relatively high level of remanent magnetization, suggesting controlling the volume fraction of Ndrich intergranular phase and its distribution would be key factors to obtain high-performance hot-deformed Nd-Fe-B magnets. The aim of this work is to enhance the coercivity of the hot-deformed Nd-Fe-B magnet with a thickness of 2 mm by the diffusion of Nd-Dy-Al eutectic alloy and to modify the composition of the grain boundary phase by introducing Dy-rich shells in Nd2Fe14B grains. The mechanism of the coercivity enhancement is also addressed by microstructure studies using scanning electron microscopy (SEM) and aberration corrected scanning transmission electron microscopy (STEM). Hot-deformed magnet used in this study was in the size of 4 × 4 × 2 mm3 with the composition of Nd13.2Fe76Co5.6B4.7Ga0.5 (at.%). The height reduction of the hot-deformation was 75% at 780 °C. Nd62Dy20Al18 alloy ribbons were prepared by the melt-spinning technique. The hot-deformed Nd-Fe-B magnets with entire surface completely covered by the ribbons were heat treated at 700 °C, which was selected based on the melting temperature of the Nd62Dy20Al18 alloy, for 1 h. The amount of ribbons was controlled to be about 20 wt% of the starting material. For comparison, Dy-vapor diffusion process was applied to the hot-deformed magnets at 900 °C. All the samples were diffusion processed without any stress or constraint applied. Inductively coupled plasma analysis showed that about 0.27 wt% Dy was introduced in the sample after the process. 0.5-mm-thick plate-like specimens were sliced from the sample parallel to the c-planes from the surface and the center of the sample. Magnetic measurements were carried out using a superconducting quantum interface device vibrating sample magnetometer (SQUIDVSM) applying a maximum magnetic field of 7 T. The magnetization value was determined using the magnetic moment from SQUID-VSM which was calibrated with Ni standard sample, and the density value measured by Gas Displacement Pycnometry System; in this work AccyPyc II 1340 Pycnometer with helium as the inert gas was used for obtaining accurate density measurement. Overall microstructural characterization was conducted using Carl ZEISS CrossBeam 1540EsB and the surfaces of the samples were cleaned using focused ion beam (FIB) before SEM observation. Scanning transmission electron microscopy energy-dispersive spectroscopy (STEM EDS) maps were constructed using Nd-Lα, Dy-Mα, Fe-Kα, Co-Kα and Cu-Kα spectra. TEM specimens were prepared with the lift-out method by a focused ion beam on FEI Nanolab Helios 650. Fig. 1(a) shows the magnetization curves of the hot-deformed magnets diffusion-processed by Nd62Dy20Al18 eutectic alloy at 700 °C and Dy-vapor at 900 °C. The coercivity (μ0Hc) of the hot-deformed magnet is enhanced from 0.91 T to 2.75 T by the Nd62Dy20Al18 grain boundary diffusion process, while the remanence drops from 1.50 T to 1.30 T. In
45
Fig. 1. (a) Magnetization curves of the hot-deformed, Nd62Dy20Al18 eutectic diffusion processed and Dy-vapor diffusion processed hot-deformed magnets, and (b) temperature dependence of the coercivity of the hot-deformed, Nd70Cu30, Nd90Al10 and Nd62Dy20Al18 diffusion-processed samples.
contrast, the hot-deformed magnet treated by the Dy vapor at 900 °C for 4 h showed a slight change in coercivity to 1.0 T without much change in remanence (μ0Mr). The S-shaped initial magnetization curve of the hot-deformed magnet has a high-susceptibility region up to around 0.9 T, followed by a lower susceptibility part then increases again to saturation. This kind of behavior indicates the magnetization process takes place by the domain wall displacement in the multi-domain grains; i.e., the domain walls are pinned at the intergranular phase, and then get depinned to reach saturation at a higher external magnetic field [27]. The depinning field increases substantially in the sample diffusion processed with Nd62Dy20Al18. On the other hand, the Dy-vapor diffusion processed sample is magnetized in one stage, suggesting all grains are multi-domain particles that can be reversed with domain wall displacements. Fig. 1(b) shows the temperature dependence of coercivity for the hot-deformed magnets and the samples diffusion processed with Nd62Dy20Al18. For comparison, the data of the samples diffusion processed with Nd70Cu30 and Nd90Al10 are also shown [17]. The sample diffusion processed with Nd-Cu shows μ0Hc of 0.82 T at 180 °C. The Nd-Al diffusion-processed magnet exhibits only 0.76 T at 180 °C while its room temperature coercivity is higher than that of the Nd-Cu diffusion processed sample, 2.52 T [17]. This is because Al is dissolved in the Nd2Fe14B phase to form Nd2(Fe,Al)14B phase, by which the Curie temperature decreases while the anisotropy field is increased due to the reduction of Ms. The Nd62Dy20Al18 diffusion-processed sample shows the highest coercivity 2.75 T at room temperature, and the coercivity at 180 °C is also the highest, 0.91 T. Fig. 1(a), (b) and (c) are the backscattered electron SEM images of the samples (a) as hot-deformed, (b) Nd62Dy20Al18 diffusion processed, and (c) Dy-vapor diffusion-processed samples, respectively, observed with the c-axis in-plane upward direction. The as hot-deformed sample
46
L. Liu et al. / Scripta Materialia 129 (2017) 44–47
Fig. 2. BSE SEM images of (a) hot-deformed, (b) Nd62Dy20Al18 eutectic diffusion-processed and (c) Dy-vapor diffusion-processed magnets with c-axis in-plane.
has the grain size of 282 ± 94 nm in width and 72 ± 21 nm in height. The brightly imaging Nd-rich phases are mainly at triple junctions, the areal fraction of which is estimated to be around 5%. Grain boundaries are observed with faintly bright contrast, suggesting that the Nd-concentration along the grain boundaries is not so high in the as hot-deformed condition. The sample diffusion processed with Nd62Dy20Al18 has the grains with 295 ± 138 nm in width and 100 ± 38 nm in height. Platelet-shaped grains are uniformly distributed surrounded by brightly imaging intergranular phase. In contrast, the grain size in the sample diffusion processed with Dy-vapor is not uniform because of the abnormal grain growth during the Dy-vapor diffusion process at 900 °C, Fig. 2(c). The presence of the abnormally grown grains explains the steep increase in magnetization in the initial magnetization curve in Fig. 1; the magnetization process dominantly progress by the domain wall displacements within multi-domain grains. This result shows that the conventional Dy-vapor diffusion process is not applicable to the hotdeformed magnets for coercivity enhancement. Therefore, the eutectic diffusion process is the only effective method to improve the coercivity of hot-deformed magnets. Fig. 3 shows the STEM-EDS mapping results observed from the Nd62Dy20Al18 diffusion-processed sample with c-axis in-plane upward direction. The selected peaks for the EDS mapping are Nd Lα, Fe Kα, Co Kα Al Kα and Dy Mα. Fe is depleted and Nd, Dy, Co and Al are enriched in the intergranular phase. Dy is mainly partitioned in the intergranular phase. In some regions, Dy substitutes part of Nd of the Nd2Fe14B phase as seen from the superimposed Dy and Nd mappings. The Dy-rich shell does not envelope the Nd2Fe14B grains entirely but (Nd1-xDyx)2Fe14B regions appear to be formed at the side surfaces of the Nd2Fe14B grains. No Dy-substituted region can be seen on the flat surface of the platelet Nd2Fe14B grains. Note that the regions selected for TEM observation is near the center of the sample, indicating that this kind of microstructure is not limited to the surface region of the sample. Inhomogeneous distribution of Al indicates the eutectic decomposition in the intergranular phase into Al-lean regions (either metallic Nd or Nd3Co phase) and Al-rich regions with the composition close to the Nd82Al18 phase. The work has shown that the intergranular phase is modified to be Nd-rich and (Fe, Co)-lean by the Nd62Dy20Al8 diffusion process. Besides, the areal fraction and thickness of Nd-rich intergranular phase increase after the diffusion process; Dy-rich layers also form on the side surface
Fig. 3. TEM-EDS elemental mapping images of Nd-Dy-Al diffusion-processed sample with c-axis in-plane. The selected peaks in EDS spectrum are Fe Kα, Nd Lα, Co Kα, Dy Mα and Al Kα.
L. Liu et al. / Scripta Materialia 129 (2017) 44–47
Fig. 4. Benchmark of hard magnetic properties of the Nd-Dy-Al diffusion-processed hotdeformed magnet with commercial sintered and hot-deformed magnets with different Dy contents. Data of sintered magnets were taken from the catalogue of Shinetsu Chemical and the data of hot-deformed magnets were taken from ref. [29] and [30].
of the Nd2Fe14B grains. We also compared the samples diffusion processed with the Nd62Dy20Al8 alloy and with Dy-vapor. The excessively high heat treatment temperature required for the Dy-vapor diffusion process leads to catastrophic grain coarsening to form extremely inhomogeneous grain size distribution, which is detrimental to the coercivity of the hot-deformed Nd-Fe-B magnets. Watanabe et al. [28] reported high coercivity hot-deformed magnets processed by blending melt-spun Nd2Fe14B powders and Dy-Cu alloy, followed by hot-pressing and hot-deformation. Due to the segregation of ferromagnetic element in the intergranular phase, the coercivity in their work was limited to 2.0 T. Sepehri-Amin et al. [22] later introduced the eutectic grain boundary diffusion process by using Nd60Dy20Cu30 alloy as the diffusion source, succeeding in modifying the intergranular phase to be nonferromagnetic and introducing thin Dy-rich shell localized on the surface of Nd2Fe14B grains. They could achieve the coercivity value of 2.6 T after the diffusion process at the expanse of remanence drop from 1.40 T to 1.10 T. In this work, we can reach around 2.75 T for coercivity with remanence decrease from 1.50 T to 1.30 T. Fig. 4 shows the remanence and coercivity relationships and their approximated Dy content in commercial sintered magnets and hot-deformed magnets at room temperature [29,30]. The coercivity of the Nd-Dy-Al diffusion-processed magnet in this work is comparable to that of the sintered magnets containing about 7 wt% Dy and that of the hot-deformed magnets containing about 5.3 wt% Dy, although its overall Dy concentration was measured to be 1.38 wt% using inductively coupled plasma (ICP) analysis. Because of the low overall concentration of Dy, the remanence of the Nd-Dy-Al diffusion processed hot-deformed magnets is higher than the sintered and hot-deformed magnets with the comparable coercivity. This demonstrates the superiority in coercivity enhancement of hot-deformed magnets by the eutectic grain boundary diffusion process using Nd-Dy-Al alloys compared to conventional HRE grain boundary diffusion process. The coercivity enhancement in this work can be attributed to the better magnetic isolation formed by the thicker nonferromagnetic Ndrich intergranular phase and the formation of localized Dy-rich shell structure at the surface of Nd2Fe14B grains. The average Dy concentration in the shell of the Nd-Dy-Al diffusion-processed magnet was close to 3 at.%; it is higher than that was reported, for the Dy-vapor diffusion processed sintered magnets, 1.07 at.% Dy [31]. Around 4 at.% Al was detected in the Nd2Fe14B phase in the Nd-Dy-Al diffusion-processed magnet. Previous study has calculated that Al substitution with Fe in the Nd2Fe14B phase has similar effect as that of Dy
47
substitution for Nd [32], i.e., the Al substitutes for Fe causes the reduction in the saturation magnetization of the Nd2Fe14B phase with higher magnetocrystalline anisotropy, resulting in substantial increase in the anisotropy field. This is likely the reason why Nd-Dy-Al diffusion-processed magnet appears to have higher coercivity than Nd-Dy-Cu diffusion-processed magnet. By applying the eutectic grain boundary diffusion process using Nd62Dy20Al18 alloy, the coercivity of hot-deformed Nd-Fe-B magnet was enhanced to 2.75 T at room temperature. The Nd-Dy-Al diffusionprocessed magnet retains the coercivity of 0.91 T at 180 °C, which satisfy the requirement for application in traction motors. On the other hand, Dy-vapor diffusion-processed Nd-Fe-B hot-deformed sample suffers catastrophic grain growth, leading to little change in coercivity, 1.0 T. ICP analysis shows the overall Dy concentration in the Nd-Dy-Al diffusion-processed sample was only 1.38 wt%. The remanence value of the Nd-Dy-Al eutectic diffusion processed hot-deformed magnet was higher than those of Dy-alloyed sintered and hot-deformed magnets with the comparable coercivity. This work was supported by JST, CREST. One of the authors, LL, acknowledges NIMS for the provision of a NIMS Junior Research Assistantship.
References [1] S. Hirosawa, K. Tokuhara, Y. Matsuura, H. Yamamoto, S. Fujimura, M. Sagawa, J. Magn. Magn. Mater. 61 (1986) 363. [2] E.B. Boltich, E. Oswald, M.Q. Huang, S. Hirosawa, W.E. Wallace, E. Burzo, J. Appl. Phys. 57 (1985) 4106. [3] H. Sepehri-Amin, Y. Une, T. Ohkubo, K. Hono, M. Sagawa, Scr. Mater. 65 (2011) 396. [4] H. Sepehri-Amin, T. Ohkubo, T. Nishiuchi, S. Hirosawa, K. Hono, Scr. Mater. 63 (2010) 1124. [5] W.B. Cui, Y.K. Takahashi, K. Hono, Acta Mater. 59 (2011) 7768. [6] R. Goto, M. Matsuura, S. Sugimoto, N. Tezuka, Y. Une, M. Sagawa, Proc. REPM Slovenia (2010) 193. [7] R.W. Lee, E.G. Brewer, N.A. Schaffel, IEEE Trans. Magn. 21 (1985) 1958. [8] R.K. Mishra, T.Y. Chu, L.K. Rabenberg, J. Magn. Magn. Mater. 84 (1990) 88. [9] P. Tenaud, A. Chamberod, F. Vanoni, Solid State Commun. 63 (1987) 303. [10] W.F. Li, T. Ohkubo, K. Hono, M. Sagawa, J. Magn. Magn. Mater. 321 (2009) 1100. [11] J. Liu, H. Sepehri-Amin, T. Ohkubo, K. Hioki, A. Hattori, T. Schrefl, K. Hono, Acta Mater. 61 (2013) 5387. [12] K. Khlopkov, O. Gutfleisch, D. Hinz, K.-H. Müller, L. Schultz, J. Appl. Phys. 102 (2007) 023912. [13] L.H. Lewis, Y. Zhu, D.O. Welch, J. Appl. Phys. 76 (1994) 6235. [14] L.H. Lewis, Y.M. Zhu, D.O. Welch, Scr. Metall. Mater. 33 (1995) 1775. [15] A. Kirchner, J. Thomas, O. Gutfleisch, D. Hinz, K.-H. Müller, L. Schultz, J. Alloys Compd. 365 (2004) 286. [16] H. Sepehri-Amin, T. Ohkubo, S. Nagashima, M. Yano, T. Shoji, A. Kato, T. Schrefl, K. Hono, Acta Mater. 61 (2013) 6622. [17] L. Liu, H. Sepehri-Amin, T. Ohkubo, M. Yano, A. Kato, T. Shoji, K. Hono, J. Alloys Compd. 666 (2016) 432. [18] H. Nakamura, K. Hirota, M. Shimao, T. Minowa, M. Honshima, IEEE Trans. Magn. 41 (2005) 3844. [19] K. Hirota, H. Nakamura, T. Minowa, M. Honshima, IEEE Trans. Magn. 42 (2006) 2909. [20] H. Nagata, BM NEWS (The Japan Association of Bonded Magnetic Materials (JABM)) 40 (2008) 47 (in Japanese). [21] H. Suzuki, Y. Satsu, M. Komuro, J. Appl. Phys. 105 (2009) 07A734. [22] H. Sepehri-Amin, J. Liu, T. Ohkubo, K. Hioki, A. Hattori, K. Hono, Scr. Mater. 69 (2013) 647. [23] T. Akiya, J. Liu, S. Sepehri-Amin, T. Ohkubo, K. Hioki, A. Hattori, K. Hono, J. Appl. Phys. 115 (2014) 17A766. [24] H. Sepehri-Amin, L. Liu, T. Ohkubo, M. Yano, T. Shoji, A. Kato, T. Schrefl, K. Hono, Acta Mater. 99 (2015) 297. [25] Z. Lin, J. Han, M. Xing, S. Liu, R. Wu, C. Wang, Y. Zhang, Y. Yang, J. Yang, Appl. Phys. Lett. 100 (2012) 052409. [26] T. Akiya, J. Liu, H. Sepehri-Amin, T. Ohkubo, K. Hioki, A. Hattori, K. Hono, Scr. Mater. 81 (2014) 48. [27] D.I. Paul, J. Appl. Phys. 53 (1982) 1649. [28] N. Watanabe, M. Itakura, M. Nishida, J. Alloys Compd. 557 (2013) 1. [29] H. Haiduka, Spec. Steel 63 (2014) 45. [30] K. Hioki, A. Hattori, T. Iriyama, IEEJ Trans. Fundam. Mater. 136 (2016) 484. [31] U.M.R. Seelam, T. Ohkubo, T. Abe, S. Hirosawa, K. Hono, J. Alloys Compd. 617 (2014) 884. [32] W. Rodewald, W. Fernengel, IEEE Trans. Magn. 24 (1988) 1638.
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Regular article
Coercivity enhancement of hot-deformed Nd-Fe-B magnets by the eutectic grain boundary diffusion process using Nd62Dy20Al18 alloy Lihua Liu a,b, H. Sepehri-Amin a, T. Ohkubo a, M. Yano c, A. Kato c, N. Sakuma c, T. Shoji c, K. Hono a,b,⁎ a b c
Elements Strategy Initiative Center for Magnetic Materials, National Institute of Materials Science, Tsukuba 305-0047, Japan Graduate School of Pure and Applied Science, University of Tsukuba, Tsukuba 305-8577, Japan Toyota Motor Corporation, Advanced Material Engineering Div., Susono 410-1193, Japan
a r t i c l e
i n f o
Article history: Received 24 August 2016 Received in revised form 16 October 2016 Accepted 16 October 2016 Available online xxxx Keywords: Permanent magnets Nd-Fe-B Coercivity Hot-deformed magnets
a b s t r a c t The eutectic grain boundary diffusion process was applied to a 2-mm-thick hot-deformed Nd-Fe-B magnets using Nd62Dy20Al18 alloy as a diffusion source, realizing the coercivity enhancement from 0.91 T to 2.75 T with relatively small remanence deterioration from 1.50 T to 1.30 T. In contrast, the conventional grain boundary diffusion process using Dy-vapor resulted in the degradation of coercivity as the grains are catastrophically coarsened at the processing temperature of 900 °C. Scanning transmission electron microscopy showed the formation of Dy-rich shell at the sides of the Nd2Fe14B grains and the diffusion of Al into the Nd2Fe14B grains, explaining the significant improvement in coercivity. © 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Nd-Fe-B based permanent magnets have been used for traction motors in (hybrid) electric vehicles owing to its ability to induce strong magnetic flux density. Since the application environment causes temperature increase under demagnetization fields, not only high remanence but also high coercivity is required. In order to attain the coercivity around 0.8 T at the operation temperature of 200 °C, about 10 wt% Dy is alloyed in commercial Nd-Fe-B sintered magnets with room temperature coercivity higher than 3 T. The partial substitution of Nd with Dy increases the anisotropy field of (Nd1-xDyx)2Fe14B, thereby increasing the coercivity of the magnet [1]. However, the magnetic moment of Dy is antiferromagnetically coupled with that of Fe atom in the (Nd1-xDyx)2Fe14B compound, resulting in a deterioration of magnetization as coercivity increases [2]. Since the natural abundance of Dy is quite limited compared to that of Nd, the development of Dy-free or Dy-saving high coercivity Nd-Fe-B permanent magnets is strongly needed. Grain size refinement of Nd-Fe-B sintered magnets is one of the effective methods to increase the coercivity without using Dy [3–5]. Sagawa et al. [6] have developed the press-less process (PLP) for sintering the submicron sized powders that were processed by helium-jet milling, realizing the coercivity of 2 T in the sintered magnets with 1 μm grain size. However, further refinement of grain size is difficult using the conventional powder metallurgy route considering the severe oxidation of Nd-rich phases. On the other hand, hot-deformed ⁎ Corresponding author at: Elements Strategy Initiative Center for Magnetic Materials, National Institute of Materials Science, Tsukuba 305-0047, Japan. E-mail address: [email protected] (K. Hono).
http://dx.doi.org/10.1016/j.scriptamat.2016.10.020 1359-6462/© 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
magnets produced from isotropic melt-spun ribbons followed by proper hot-press and hot-deformation exhibit strong [001] texture and the fine grain size that is comparable to the single domain size of Nd2Fe14B phase [7–9]. Unlike the conventional sintering process, powder size can be larger than 100 μm, inside which nano-sized grains are dispersed; therefore, individual grain are not exposed to oxygen even if the process is carried out without rigorous control of oxygen. The typical coercivity value, μ0Hc, of hot-deformed magnets is around 1.2 T with a remanence, μ0Mr, of 1.4 T, which are comparable to those of sintered magnets [10,11]. Strong exchange interaction between Nd2Fe14B grains through ferromagnetic intergranular phase is the main reason for this relatively low coercivity of the hot-deformed magnets [12–15]. More recently, Liu et al. reported a good correlation between the Nd content in the intergranular phase and the coercivity using atom probe tomography [11], and they suggested that the modification of the intergranular phase would enhance the coercivity further. Sepehri-Amin et al. demonstrated a dramatic enhancement of coercivity by the eutectic grain boundary diffusion (GBD) process for hydrogen-disproportionation-decomposition-recombination (HDDR) processed Nd-Fe-B powders using Nd70Cu30 eutectic alloy as a diffusion source [4]. This eutectic GBD process was later extended to Nd-Fe-B hotdeformed magnets [16], which enhanced the coercivity from 1.5 T to 2.3 T without using Dy. In order to explore the extent of coercivity enhancement by this eutectic GBD process, Liu et al. [17] applied various Nd1-xMx eutectic alloys to the 2-mm-thick Nd-Fe-B hot-deformed magnets, obtaining the highest coercivity of 2.5 T at room temperature using Nd90Al10 as a diffusion source. However, the coercivity measured at 200 °C was only 0.6 T because of poor temperature dependence of
L. Liu et al. / Scripta Materialia 129 (2017) 44–47
μ0Hc partly due to the dissolution of Al in the Nd2Fe14B phase. A higher level of coercivity is required for the applications such as automotive traction motors and wind turbines. Heavy rare earth (HRE) grain boundary diffusion process was developed for enhancing the coercivity of sintered magnets with the grain size of about 5 μm using Dy-vapor, Dy2O3, DyF3, Tb3O4 and TbF4 [18– 21]. However, this method cannot be employed to ultrafine-grainsized hot-deformed Nd-Fe-B magnets since the high temperature annealing required for the HRE GBDP results in catastrophic grain growth. Recently, Sepehri-Amin reported that Dy-rich interface can be developed without any grain coarsening in Nd-Fe-B hot-deformed magnet using Nd60Dy20Cu20 powder as a diffusion source followed by annealing at low temperature of 650 °C [22]. On the other hand, along with the remarkable enhancements of coercivity by using Nd-Cu [4,16,23], Nd-Al [17], Pr-Cu [23–25] or Nd-DyCu [22] as the diffusion source, the diffusion-processed magnets always experienced a significant degradation in remanence due to the infiltration of excessive non-ferromagnetic RE-rich liquid phase into the magnet. Akiya et al. [26] applied an expansion constraint on the c-plane of the magnet during infiltration of RE-rich liquid phase and achieved coercivity enhancement while retaining a relatively high level of remanent magnetization, suggesting controlling the volume fraction of Ndrich intergranular phase and its distribution would be key factors to obtain high-performance hot-deformed Nd-Fe-B magnets. The aim of this work is to enhance the coercivity of the hot-deformed Nd-Fe-B magnet with a thickness of 2 mm by the diffusion of Nd-Dy-Al eutectic alloy and to modify the composition of the grain boundary phase by introducing Dy-rich shells in Nd2Fe14B grains. The mechanism of the coercivity enhancement is also addressed by microstructure studies using scanning electron microscopy (SEM) and aberration corrected scanning transmission electron microscopy (STEM). Hot-deformed magnet used in this study was in the size of 4 × 4 × 2 mm3 with the composition of Nd13.2Fe76Co5.6B4.7Ga0.5 (at.%). The height reduction of the hot-deformation was 75% at 780 °C. Nd62Dy20Al18 alloy ribbons were prepared by the melt-spinning technique. The hot-deformed Nd-Fe-B magnets with entire surface completely covered by the ribbons were heat treated at 700 °C, which was selected based on the melting temperature of the Nd62Dy20Al18 alloy, for 1 h. The amount of ribbons was controlled to be about 20 wt% of the starting material. For comparison, Dy-vapor diffusion process was applied to the hot-deformed magnets at 900 °C. All the samples were diffusion processed without any stress or constraint applied. Inductively coupled plasma analysis showed that about 0.27 wt% Dy was introduced in the sample after the process. 0.5-mm-thick plate-like specimens were sliced from the sample parallel to the c-planes from the surface and the center of the sample. Magnetic measurements were carried out using a superconducting quantum interface device vibrating sample magnetometer (SQUIDVSM) applying a maximum magnetic field of 7 T. The magnetization value was determined using the magnetic moment from SQUID-VSM which was calibrated with Ni standard sample, and the density value measured by Gas Displacement Pycnometry System; in this work AccyPyc II 1340 Pycnometer with helium as the inert gas was used for obtaining accurate density measurement. Overall microstructural characterization was conducted using Carl ZEISS CrossBeam 1540EsB and the surfaces of the samples were cleaned using focused ion beam (FIB) before SEM observation. Scanning transmission electron microscopy energy-dispersive spectroscopy (STEM EDS) maps were constructed using Nd-Lα, Dy-Mα, Fe-Kα, Co-Kα and Cu-Kα spectra. TEM specimens were prepared with the lift-out method by a focused ion beam on FEI Nanolab Helios 650. Fig. 1(a) shows the magnetization curves of the hot-deformed magnets diffusion-processed by Nd62Dy20Al18 eutectic alloy at 700 °C and Dy-vapor at 900 °C. The coercivity (μ0Hc) of the hot-deformed magnet is enhanced from 0.91 T to 2.75 T by the Nd62Dy20Al18 grain boundary diffusion process, while the remanence drops from 1.50 T to 1.30 T. In
45
Fig. 1. (a) Magnetization curves of the hot-deformed, Nd62Dy20Al18 eutectic diffusion processed and Dy-vapor diffusion processed hot-deformed magnets, and (b) temperature dependence of the coercivity of the hot-deformed, Nd70Cu30, Nd90Al10 and Nd62Dy20Al18 diffusion-processed samples.
contrast, the hot-deformed magnet treated by the Dy vapor at 900 °C for 4 h showed a slight change in coercivity to 1.0 T without much change in remanence (μ0Mr). The S-shaped initial magnetization curve of the hot-deformed magnet has a high-susceptibility region up to around 0.9 T, followed by a lower susceptibility part then increases again to saturation. This kind of behavior indicates the magnetization process takes place by the domain wall displacement in the multi-domain grains; i.e., the domain walls are pinned at the intergranular phase, and then get depinned to reach saturation at a higher external magnetic field [27]. The depinning field increases substantially in the sample diffusion processed with Nd62Dy20Al18. On the other hand, the Dy-vapor diffusion processed sample is magnetized in one stage, suggesting all grains are multi-domain particles that can be reversed with domain wall displacements. Fig. 1(b) shows the temperature dependence of coercivity for the hot-deformed magnets and the samples diffusion processed with Nd62Dy20Al18. For comparison, the data of the samples diffusion processed with Nd70Cu30 and Nd90Al10 are also shown [17]. The sample diffusion processed with Nd-Cu shows μ0Hc of 0.82 T at 180 °C. The Nd-Al diffusion-processed magnet exhibits only 0.76 T at 180 °C while its room temperature coercivity is higher than that of the Nd-Cu diffusion processed sample, 2.52 T [17]. This is because Al is dissolved in the Nd2Fe14B phase to form Nd2(Fe,Al)14B phase, by which the Curie temperature decreases while the anisotropy field is increased due to the reduction of Ms. The Nd62Dy20Al18 diffusion-processed sample shows the highest coercivity 2.75 T at room temperature, and the coercivity at 180 °C is also the highest, 0.91 T. Fig. 1(a), (b) and (c) are the backscattered electron SEM images of the samples (a) as hot-deformed, (b) Nd62Dy20Al18 diffusion processed, and (c) Dy-vapor diffusion-processed samples, respectively, observed with the c-axis in-plane upward direction. The as hot-deformed sample
46
L. Liu et al. / Scripta Materialia 129 (2017) 44–47
Fig. 2. BSE SEM images of (a) hot-deformed, (b) Nd62Dy20Al18 eutectic diffusion-processed and (c) Dy-vapor diffusion-processed magnets with c-axis in-plane.
has the grain size of 282 ± 94 nm in width and 72 ± 21 nm in height. The brightly imaging Nd-rich phases are mainly at triple junctions, the areal fraction of which is estimated to be around 5%. Grain boundaries are observed with faintly bright contrast, suggesting that the Nd-concentration along the grain boundaries is not so high in the as hot-deformed condition. The sample diffusion processed with Nd62Dy20Al18 has the grains with 295 ± 138 nm in width and 100 ± 38 nm in height. Platelet-shaped grains are uniformly distributed surrounded by brightly imaging intergranular phase. In contrast, the grain size in the sample diffusion processed with Dy-vapor is not uniform because of the abnormal grain growth during the Dy-vapor diffusion process at 900 °C, Fig. 2(c). The presence of the abnormally grown grains explains the steep increase in magnetization in the initial magnetization curve in Fig. 1; the magnetization process dominantly progress by the domain wall displacements within multi-domain grains. This result shows that the conventional Dy-vapor diffusion process is not applicable to the hotdeformed magnets for coercivity enhancement. Therefore, the eutectic diffusion process is the only effective method to improve the coercivity of hot-deformed magnets. Fig. 3 shows the STEM-EDS mapping results observed from the Nd62Dy20Al18 diffusion-processed sample with c-axis in-plane upward direction. The selected peaks for the EDS mapping are Nd Lα, Fe Kα, Co Kα Al Kα and Dy Mα. Fe is depleted and Nd, Dy, Co and Al are enriched in the intergranular phase. Dy is mainly partitioned in the intergranular phase. In some regions, Dy substitutes part of Nd of the Nd2Fe14B phase as seen from the superimposed Dy and Nd mappings. The Dy-rich shell does not envelope the Nd2Fe14B grains entirely but (Nd1-xDyx)2Fe14B regions appear to be formed at the side surfaces of the Nd2Fe14B grains. No Dy-substituted region can be seen on the flat surface of the platelet Nd2Fe14B grains. Note that the regions selected for TEM observation is near the center of the sample, indicating that this kind of microstructure is not limited to the surface region of the sample. Inhomogeneous distribution of Al indicates the eutectic decomposition in the intergranular phase into Al-lean regions (either metallic Nd or Nd3Co phase) and Al-rich regions with the composition close to the Nd82Al18 phase. The work has shown that the intergranular phase is modified to be Nd-rich and (Fe, Co)-lean by the Nd62Dy20Al8 diffusion process. Besides, the areal fraction and thickness of Nd-rich intergranular phase increase after the diffusion process; Dy-rich layers also form on the side surface
Fig. 3. TEM-EDS elemental mapping images of Nd-Dy-Al diffusion-processed sample with c-axis in-plane. The selected peaks in EDS spectrum are Fe Kα, Nd Lα, Co Kα, Dy Mα and Al Kα.
L. Liu et al. / Scripta Materialia 129 (2017) 44–47
Fig. 4. Benchmark of hard magnetic properties of the Nd-Dy-Al diffusion-processed hotdeformed magnet with commercial sintered and hot-deformed magnets with different Dy contents. Data of sintered magnets were taken from the catalogue of Shinetsu Chemical and the data of hot-deformed magnets were taken from ref. [29] and [30].
of the Nd2Fe14B grains. We also compared the samples diffusion processed with the Nd62Dy20Al8 alloy and with Dy-vapor. The excessively high heat treatment temperature required for the Dy-vapor diffusion process leads to catastrophic grain coarsening to form extremely inhomogeneous grain size distribution, which is detrimental to the coercivity of the hot-deformed Nd-Fe-B magnets. Watanabe et al. [28] reported high coercivity hot-deformed magnets processed by blending melt-spun Nd2Fe14B powders and Dy-Cu alloy, followed by hot-pressing and hot-deformation. Due to the segregation of ferromagnetic element in the intergranular phase, the coercivity in their work was limited to 2.0 T. Sepehri-Amin et al. [22] later introduced the eutectic grain boundary diffusion process by using Nd60Dy20Cu30 alloy as the diffusion source, succeeding in modifying the intergranular phase to be nonferromagnetic and introducing thin Dy-rich shell localized on the surface of Nd2Fe14B grains. They could achieve the coercivity value of 2.6 T after the diffusion process at the expanse of remanence drop from 1.40 T to 1.10 T. In this work, we can reach around 2.75 T for coercivity with remanence decrease from 1.50 T to 1.30 T. Fig. 4 shows the remanence and coercivity relationships and their approximated Dy content in commercial sintered magnets and hot-deformed magnets at room temperature [29,30]. The coercivity of the Nd-Dy-Al diffusion-processed magnet in this work is comparable to that of the sintered magnets containing about 7 wt% Dy and that of the hot-deformed magnets containing about 5.3 wt% Dy, although its overall Dy concentration was measured to be 1.38 wt% using inductively coupled plasma (ICP) analysis. Because of the low overall concentration of Dy, the remanence of the Nd-Dy-Al diffusion processed hot-deformed magnets is higher than the sintered and hot-deformed magnets with the comparable coercivity. This demonstrates the superiority in coercivity enhancement of hot-deformed magnets by the eutectic grain boundary diffusion process using Nd-Dy-Al alloys compared to conventional HRE grain boundary diffusion process. The coercivity enhancement in this work can be attributed to the better magnetic isolation formed by the thicker nonferromagnetic Ndrich intergranular phase and the formation of localized Dy-rich shell structure at the surface of Nd2Fe14B grains. The average Dy concentration in the shell of the Nd-Dy-Al diffusion-processed magnet was close to 3 at.%; it is higher than that was reported, for the Dy-vapor diffusion processed sintered magnets, 1.07 at.% Dy [31]. Around 4 at.% Al was detected in the Nd2Fe14B phase in the Nd-Dy-Al diffusion-processed magnet. Previous study has calculated that Al substitution with Fe in the Nd2Fe14B phase has similar effect as that of Dy
47
substitution for Nd [32], i.e., the Al substitutes for Fe causes the reduction in the saturation magnetization of the Nd2Fe14B phase with higher magnetocrystalline anisotropy, resulting in substantial increase in the anisotropy field. This is likely the reason why Nd-Dy-Al diffusion-processed magnet appears to have higher coercivity than Nd-Dy-Cu diffusion-processed magnet. By applying the eutectic grain boundary diffusion process using Nd62Dy20Al18 alloy, the coercivity of hot-deformed Nd-Fe-B magnet was enhanced to 2.75 T at room temperature. The Nd-Dy-Al diffusionprocessed magnet retains the coercivity of 0.91 T at 180 °C, which satisfy the requirement for application in traction motors. On the other hand, Dy-vapor diffusion-processed Nd-Fe-B hot-deformed sample suffers catastrophic grain growth, leading to little change in coercivity, 1.0 T. ICP analysis shows the overall Dy concentration in the Nd-Dy-Al diffusion-processed sample was only 1.38 wt%. The remanence value of the Nd-Dy-Al eutectic diffusion processed hot-deformed magnet was higher than those of Dy-alloyed sintered and hot-deformed magnets with the comparable coercivity. This work was supported by JST, CREST. One of the authors, LL, acknowledges NIMS for the provision of a NIMS Junior Research Assistantship.
References [1] S. Hirosawa, K. Tokuhara, Y. Matsuura, H. Yamamoto, S. Fujimura, M. Sagawa, J. Magn. Magn. Mater. 61 (1986) 363. [2] E.B. Boltich, E. Oswald, M.Q. Huang, S. Hirosawa, W.E. Wallace, E. Burzo, J. Appl. Phys. 57 (1985) 4106. [3] H. Sepehri-Amin, Y. Une, T. Ohkubo, K. Hono, M. Sagawa, Scr. Mater. 65 (2011) 396. [4] H. Sepehri-Amin, T. Ohkubo, T. Nishiuchi, S. Hirosawa, K. Hono, Scr. Mater. 63 (2010) 1124. [5] W.B. Cui, Y.K. Takahashi, K. Hono, Acta Mater. 59 (2011) 7768. [6] R. Goto, M. Matsuura, S. Sugimoto, N. Tezuka, Y. Une, M. Sagawa, Proc. REPM Slovenia (2010) 193. [7] R.W. Lee, E.G. Brewer, N.A. Schaffel, IEEE Trans. Magn. 21 (1985) 1958. [8] R.K. Mishra, T.Y. Chu, L.K. Rabenberg, J. Magn. Magn. Mater. 84 (1990) 88. [9] P. Tenaud, A. Chamberod, F. Vanoni, Solid State Commun. 63 (1987) 303. [10] W.F. Li, T. Ohkubo, K. Hono, M. Sagawa, J. Magn. Magn. Mater. 321 (2009) 1100. [11] J. Liu, H. Sepehri-Amin, T. Ohkubo, K. Hioki, A. Hattori, T. Schrefl, K. Hono, Acta Mater. 61 (2013) 5387. [12] K. Khlopkov, O. Gutfleisch, D. Hinz, K.-H. Müller, L. Schultz, J. Appl. Phys. 102 (2007) 023912. [13] L.H. Lewis, Y. Zhu, D.O. Welch, J. Appl. Phys. 76 (1994) 6235. [14] L.H. Lewis, Y.M. Zhu, D.O. Welch, Scr. Metall. Mater. 33 (1995) 1775. [15] A. Kirchner, J. Thomas, O. Gutfleisch, D. Hinz, K.-H. Müller, L. Schultz, J. Alloys Compd. 365 (2004) 286. [16] H. Sepehri-Amin, T. Ohkubo, S. Nagashima, M. Yano, T. Shoji, A. Kato, T. Schrefl, K. Hono, Acta Mater. 61 (2013) 6622. [17] L. Liu, H. Sepehri-Amin, T. Ohkubo, M. Yano, A. Kato, T. Shoji, K. Hono, J. Alloys Compd. 666 (2016) 432. [18] H. Nakamura, K. Hirota, M. Shimao, T. Minowa, M. Honshima, IEEE Trans. Magn. 41 (2005) 3844. [19] K. Hirota, H. Nakamura, T. Minowa, M. Honshima, IEEE Trans. Magn. 42 (2006) 2909. [20] H. Nagata, BM NEWS (The Japan Association of Bonded Magnetic Materials (JABM)) 40 (2008) 47 (in Japanese). [21] H. Suzuki, Y. Satsu, M. Komuro, J. Appl. Phys. 105 (2009) 07A734. [22] H. Sepehri-Amin, J. Liu, T. Ohkubo, K. Hioki, A. Hattori, K. Hono, Scr. Mater. 69 (2013) 647. [23] T. Akiya, J. Liu, S. Sepehri-Amin, T. Ohkubo, K. Hioki, A. Hattori, K. Hono, J. Appl. Phys. 115 (2014) 17A766. [24] H. Sepehri-Amin, L. Liu, T. Ohkubo, M. Yano, T. Shoji, A. Kato, T. Schrefl, K. Hono, Acta Mater. 99 (2015) 297. [25] Z. Lin, J. Han, M. Xing, S. Liu, R. Wu, C. Wang, Y. Zhang, Y. Yang, J. Yang, Appl. Phys. Lett. 100 (2012) 052409. [26] T. Akiya, J. Liu, H. Sepehri-Amin, T. Ohkubo, K. Hioki, A. Hattori, K. Hono, Scr. Mater. 81 (2014) 48. [27] D.I. Paul, J. Appl. Phys. 53 (1982) 1649. [28] N. Watanabe, M. Itakura, M. Nishida, J. Alloys Compd. 557 (2013) 1. [29] H. Haiduka, Spec. Steel 63 (2014) 45. [30] K. Hioki, A. Hattori, T. Iriyama, IEEJ Trans. Fundam. Mater. 136 (2016) 484. [31] U.M.R. Seelam, T. Ohkubo, T. Abe, S. Hirosawa, K. Hono, J. Alloys Compd. 617 (2014) 884. [32] W. Rodewald, W. Fernengel, IEEE Trans. Magn. 24 (1988) 1638.