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MM-Fe-B based gap magnet with excellent energy density
Intermetallics 115 (2019) 106626
Contents lists available at ScienceDirect
Intermetallics journal homepage: http://www.elsevier.com/locate/intermet
MM-Fe-B based gap magnet with excellent energy density Weiqiang Liu a, Zhipeng Zhang a, Ming Yue a, *, Zhi Li a, Dan Wu a, Zhao Zhou a, Hao Chen a, Yuqing Li a, Zaisheng Pang b, Xi Yu b a
College of Materials Science and Engineering, Key Laboratory of Advanced Functional Materials, Ministry of Education of China, Beijing University of Technology, Beijing, 100124, China b Ganzhou Fortune Electronic Co., Ltd., Ganzhou, 341000, China
A R T I C L E I N F O
A B S T R A C T
Keywords: MM-Fe-B magnet Grain boundary diffusion Magnetic properties Coercivity TbHx nanoparticles Magnetic domains
TbHx nanoparticles were applied to prepare sintered MM-Fe-B magnet by using grain boundary diffusion tech nique. Under optimal processing conditions, the coercivity and maximum energy product substantially increase by 383% and 151%, reaching to 6.81 kOe and 20.18 MGOe, respectively. Further investigation indicates that the interchange between Tb atoms and Ce atoms leads to a core/shell structure with (MM,Tb)2Fe14B phase in the surface region of MM2Fe14B matrix grains in the magnet. The modified structure successfully restrains the magnetic domain reversal process, resulting in remarkable enhancement of coercivity and maximum energy product of the magnet.
1. Introduction Misch metal (MM), also called cerium misch metal, is an alloy made from a mixture of rare-earth metals that includes approximately 55% cerium (Ce), 25% lanthanum (La), and other rare earth metals such as neodymium (Nd) and praseodymium (Pr). Herbst studied the intrinsic magnetic properties of R2Fe14B materials, including La2Fe14B (Js¼13.8kG, HA¼20kOe),Ce2Fe14B(Js¼11.7kG, HA¼26kOe),Pr2Fe14B (Js¼15.6kG, HA¼75kOe) and Nd2Fe14B(Js¼16.0kG, HA¼73kOe).Due to the low Js and HA of La2Fe14B and Ce2Fe14B, the intrinsic magnetic properties of MM2Fe14B are not as good as Nd2Fe14B, but it bears excellent potential as a kind of gap magnet between sintered Nd-Fe-B magnet and hard ferrite [1–4]. Moreover, massive applications of MM2Fe14B-based magnets are benefit for utilization of rare earth resource in equilibrium and for energy conservation [5–7]. To date, however, obtaining high energy density [for example (BH)max>20 MGOe] in MM2Fe14B-based magnets is still facing great challenge. By using melt-spinning, different researchers achieved coercivity higher than 6 kOe [7] in MM2Fe14B ribbons, but the magnetically isotropic nature undermine their energy density [8]. Recently, Shang et al. pre pared anisotropic sintered MM-Fe-B magnet, while the low coercivity of only 1.08 kOe also leads to low (BH)max of 7.6 MGOe [9]. Considering the limited resources and price hikes of Nd and Pr, it is urgent to develop a new approach to fabricate sintered MM-Fe-B magnet with high energy density.
Grain boundary diffusion (GBD) technique is an effective way to improve the coercivity of sintered Nd-Fe-B magnets [10–12]. In partic ular, the formation of (Nd, Tb)2Fe14B or (Nd, Dy)2Fe14B is achieved by introducing Terbium (Tb) or Dysprosium (Dy) into the surface region of Nd2Fe14B matrix grains, and the magnet can achieve substantially enhanced coercivity with negligible remanence loss [13–15]. In our recent work, TbHx nanoparticles were applied to improve the magnetic properties of sintered Nd-Fe-B magnet using grain boundary diffusion (GBD) technique, and excellent overall magnetic property (OMP), the sum of Hci and (BH)max, of 85.6 was obtained [16]. Here, we use similar treatment to sintered MM-Fe-B magnet to improve their poor coercivity, and thus to achieve good (BH)max in the magnet. Furthermore, the microstructure, magnetic domain evolution, and magnetic hardening mechanism of the magnet were investigated and discussed. 2. Experiments Misch metals were bought from Bayan Obo in Baotou, China, con taining about 25 wt% La, 55 wt% Ce, 5 wt% Pr and 15 wt% Nd with a purity of about 99.8 wt%. Alloy with a nominal composition of MM30.8FebalAl0.3Cu0.15Co1Ga0.35Zr0.22B0.96 was prepared by strip cast. Subsequently, hydrogen decrepitation (HD) and jet milling (JM) were used to obtain fine powders with diameters of 3–5 μm. The fine powders were arranged under the magnetic field of 1.8 T and pressed into cuboid. In order to restrain the growth of the grains, the lower sintering
* Corresponding author. E-mail address: [email protected] (M. Yue). https://doi.org/10.1016/j.intermet.2019.106626 Received 27 June 2019; Received in revised form 27 September 2019; Accepted 28 September 2019 Available online 11 October 2019 0966-9795/© 2019 Elsevier Ltd. All rights reserved.
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Intermetallics 115 (2019) 106626
Fig. 1. (a) The demagnetization curves of the original and GBD treated MM-Fe-B magnets. Inset table: magnetic properties of both magnets. (b) The distribution of magnetic properties of sintered Nd-Fe-B magnets, sintered MM-Fe-B magnets and Ferrite magnets.
temperature of 1283 K was carried out. The ethanol and TbHx nanoparticles prepared by the evaporationcondensation technology were mixed to obtain heavy rare earth sus pensions. TbHx nanoparticles suspensions were sprayed into the MM-FeB magnets with thickness of 4 mm and formed a TbHx layer with 0.5 wt % of the original magnets. The heat treatment was carried out in vacuum at 1198 K for 8 h, followed by further annealing at 773 K for 3 h. The magnetic properties of the samples were measured by permanent mag netic measuring instrument (NIM-500C). The scanning electron micro scope (SEM, Nova Nano200) with energy dispersive X-ray spectroscopy and electron-probe microanalyzer (EPMA, JEOL JXA-800) with a wavelength-dispersive X-ray detector were applied to observe the microstructure of the magnets. The Magneto-optical Kerr Optical Mi croscope (MOKE, BH-786IP-PK NEOARK CORP) was used to study the magnetic domains and their evolution in the range from 1.4 T to 0.7 T.
hardening for MM-Fe-B magnets. Note that the contents of La and Ce in the MM2Fe14B matrix grains are much higher than the content of Nd, and the magnetocrystalline anisotropy field (HA) of both La2Fe14B and Ce2Fe14B is much lower than that of Nd2Fe14B. It is therefore expected that the coercivity of MM-Fe-B magnets is low. To explore the coercivity enhancement in GBD treated magnet with TbHx nanoparticles, the microstructure of the magnets was analyzed as shown in Fig. 2(e–j). The concentration of Tb decreases gradually from the surface to the inside of the magnet, but it goes a considerable depth along with the grain boundary phase, as shown in Fig. 2(f). Moreover, Tb is mainly enriching in the surface region of MM2Fe14B matrix grains [Fig. 2(j)], forming (MM, Tb)2Fe14B phase with substantially enhanced magnetocrystalline anisotropy. Based on the above observation, it is proposed that Tb atoms first go from the surface to the inside of the magnet via liquid grain boundary phase during the GBD process, and then go into the surface region of the MM2Fe14B matrix grains by exchanging with Ce (or La and Nd) atoms. Our observation and assumption have good agreement with the reporting that Tb tends to enter the Nd2Fe14B matrix phase rather than the Nd-rich phase when the magnet is diffused with Tb [17]. On the other hand, it is found that La prefers to enrich in grain boundary phase, while Nd has a uniform distribution throughout the whole magnet, as shown in Fig. 2(g) and (i). To explore the quantitative changes of rare earth elements distri bution after Tb infiltration, the concentration distribution of La, Ce, Nd, and Tb going through the grain boundary phase and the MM2Fe14B matrix phase of the original and GBD treated MM-Fe-B magnets were examined, as shown in Fig. 3. In the original magnets (Fig. 3(a) and (b)), the counts of La, Ce, and Nd have negligible fluctuation, indicating that the distribution of La, Ce and Nd elements is uniform in the matrix grains. In the GBD treated MM-Fe-B magnets (Fig. 3(c) and (d)), the counts of La and Nd are stable; however, the counts of Ce increase first and then decrease, and the variant trend of Tb count is opposite to that of Ce counts, demonstrating that distribution of Ce element in MM2Fe14B matrix grains has been changed by Tb infiltration. Note that the counts of Tb goes from 200 in the core rapidly to 1300 in the surface region of the grains, for comparison, the counts of Ce goes from 2200 to 1700. Such exchange results in a typical core (MM2Fe14B)/shell [(MM, Tb)2Fe14B] structure. Moreover, it is worth to note that La also has a preference distribution near grain boundary, which is beneficial to the magnetic isolation of matrix grains and therefore good to coercivity of
3. Results & discussion Fig. 1(a) shows the demagnetization curves of the original and GBD treated MM-Fe-B magnets. The maximum energy product of the original magnet reaches to only 8.05 MGOe, and moreover, the intrinsic coer civity is only 1.41 kOe, which limits its practical application. After diffusion, the magnets exhibit much better properties with (BH)max of 20.18 MGOe and coercivity of 6.81 kOe, meaning remarkable increment of 150.7% and 383%, respectively. Meanwhile, the remanence of GBD treated magnets decreases slightly of 0.29 kG (3%) compared to that of original magnets. The moderate energy density of our magnet makes it a proper gap magnet between sintered Nd-Fe-B magnets and Ferrites. In detail, as shown in Fig. 1(b), the (BH)max of Ferrite magnets is less than 5.5 MGOe, and sintered Nd-Fe-B magnets are in the range of 30–52 MGOe, and our newly developed magnet exhibits a medium (BH)max of 20 MGOe. Furthermore, it is worth to note that the fabrication cost of our new magnet is much lower than that of commercial sintered Nd-Fe-B magnets, allowing its good potential in practical applications. Fig. 2(a–d) displays the BSE image and concentration distribution mapping of La, Ce, and Nd in the original MM-Fe-B magnet. It can be found that La and Ce are tending to enrich in the grain boundary phases, which consistent with our previous results [2–4]. Although such ag gregation behavior is beneficial to fabricate La/Ce containing magnets with high magnetic properties, it is difficult to achieve magnetic
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the magnet. For a deeper understanding of the GBD process on the coercivity enhancement of MM-Fe-B sintered magnets, the magnetic domains reversal status of the original and GBD treated MM-Fe-B magnets after fully positive magnetization were observed from 1.4 T to 0.7 T, as shown in Fig. 4. The original MM-Fe-B magnets show single domains status, indicating that the magnets have been fully magnetized at 1.4 T. When the magnetic field reduces to 0 T, it can be found the most grains exhibit maze domain pattern [Fig. 4(b)], which means reverse magnetic domains have nucleated before 0 T. After the magnetic field turns to 0.1 T and 0.2 T, the change of reverse domain status is not obvious, as shown in Fig. 4(c) and (d), indicating that most grains have already completely transferred into the demagnetization status. This result ex plains the low coercivity of original MM-Fe-B magnets. When the mag netic field turns to 0.7 T, almost all of the domains have been reversed magnetized completely. For the GBD treated MM-Fe-B magnets, it ex hibits single domain status at 1.4 T. Different from the original magnets, about half of grains in GBD treated magnets still keep single domain status, marked by red dotted line at 0 T in Fig. 4(g). The maze domains area enlarges with the increase of reverse magnetic field to 0.2 T. However, there still have some grains keeping their original single domain state when the reverse magnetic field reaches to 0.7 T, which means that the reverse nucleation field of these grains is obviously intensified by GBD treatment. Therefore, it can be concluded that GBD treatment plays an important role in inhibiting the nucleation of the reverse domain. On the basis of our finding, it can be pointed out that the magnetic domain reversal modes between the original and GBD treated MM-Fe-B sintered magnets are significantly different. For the MM-Fe-B sintered magnet, the nucleation of the reverse magnetic domain determines the final coercivity. As shown in Figs. 2 and 4(b), most La2Fe14B and Ce2Fe14B phases with low HA in MM2Fe14B grains result in low coer civity. More importantly, the MM2Fe14B grain surface regions usually have defects and therefore exhibit lower magnetocrystalline anisotropy than that of the inner grains. Therefore, they will prior to nucleate. After that, the expansion of the reverse domain would very fast as increasing the reverse magnetic field. For the GBD treated MM-Fe-B magnets, Tb atoms substitute Ce, La and Nd atoms in the surface region of MM2Fe14B grains and forms (MM, Tb)2Fe14B phase. It needs to point out that the HA of the Tb2Fe14B compound is 220kOe, which is much higher than that of the MM2Fe14B compound with 48kOe [4]. As a result, the HN of the surface regions with the (MM,Tb)2Fe14B phase is obviously intensified, therefore restrain the nucleation of reverse domains. Therefore, the grain boundary diffusion could achieve the magnetic hardening for MM-Fe-B magnets, and the coercivity could be obviously intensified. 4. Conclusion Misch metal based permanent magnet with nominal compositions of MM30.8FebalAl0.3Cu0.15Co1Ga0.35Zr0.22B0.96 was prepared. However, its intrinsic coercivity is only 1.41 kOe, which undermines its energy den sity and thus limits its practical application. To overcome this dilemma, grain boundary diffusion with TbHx nanoparticles was carried out to improve the coercivity of MM-Fe-B magnets. The GBD treated magnets exhibit better properties with the maximum energy product of 20.18 MGOe and intrinsic coercivity of 6.81 kOe. A typical core (MM2Fe14B)/ shell [(MM, Tb)2Fe14B] structure has a significant restraint on the reverse nucleation of domains, and the HN of the surface region is obviously intensified. Therefore, the coercivity of GBD treated MM-Fe-B magnets was intensified remarkably. Fig. 2. BSE image (a) and concentration distribution mapping of La (b), Ce (c), and Nd (d) in the original MM-Fe-B magnet; BSE image (e), concentration gradient of Tb from surface to inside in cross section (f), and concentration distribution mapping of La (g), Ce (h), Nd (i), and Tb (j) in the GBD treated MMFe-B magnet.
Declaration of competing interest None.
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Fig. 3. The concentration distribution of La, Ce, Nd, and Tb elements going through the grain boundary phase and the MM2Fe14B matrix phase of the original (a, b) and GBD treated (c, d) MM-Fe-B sintered magnets, the red arrows in (a) and (c) indicate the scanning path. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 4. MOKE images of magnetic domain evolution of the original (a–e) and GBD treated (f–j) MM-Fe-B magnets in the field from 1.4 T to plane is perpendicular to c-axis).
Acknowledgement
0.7 T. (The observing
References [1] J.F. Herbst, R2Fe14B materials: intrinsic properties and technological aspects, Rev. Mod. Phys. 63 (1991) 819–898. [2] Z. Li, W.Q. Liu, S.S. Zha, et al., Effects of Ce substitution on the microstructures and intrinsic magnetic properties of Nd-Fe-B alloy, J. Magn. Magn. Mater. 393 (2015) 551–554. [3] W.Q. Liu, Z.P. Zhang, M. Yue, et al., Effects of La substitution on the crystal structure and magnetization of MM-Fe-B alloy (MM¼La, Ce, Pr, Nd), J. Magn. Magn. Mater. 464 (2018) 61–64. [4] X.Q. Yu, M. Yue, W.Q. Liu, et al., Structure and intrinsic magnetic properties of MM2Fe14B (MM¼La, Ce, Pr, Nd) alloys, J. Rare Earths 34 (2016) 614–617.
This work was supported by the National Key Research and Devel opment Program of China under Grant 2016YFB0700902; National Natural Science Foundation of China under Grant 51371002; Interna tional S&T Cooperation Program of China under Grant 2015DFG52020; Program of Top Disciplines Construction in Beijing under Grant PXM2019_014204_500031.
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[5] J. Yamasaki, H. Soeda, M. Yanagida, et al., Misch metal-Fe-B melt spun magnets with 8 MGOe energy product, IEEE Trans. Magn. 22 (5) (1986) 763–765. [6] M. Zhang, B.G. Shen, F. Hu, et al., The effect of Si substitution on structure and magnetic properties in mischmetal-Fe-B ribbons, IEEE Trans. Magn. 51 (11) (2015) 1–4. [7] W.L. Zuo, S.L. Zuo, R. Li, et al., High performance misch-metal (MM)-Fe-B magnets prepared by melt spinning, J. Alloy. Comp. 695 (2016) 1786–1792. [8] Z.B. Li, L.C. Wang, X.P. Geng, et al., Variation of magnetic properties with mischmetal content in the resource saving magnets of MM-Fe-B ribbons, J. Magn. Magn. Mater. 426 (2017) 70–73. [9] R.X. Shang, J.F. Xiong, et al., Structure and properties of sintered MM-Fe-B magnets, AIP Adv. 7 (5) (2017), 056215. [10] K. Hirota, H. Nakamura, T. Minowa, et al., Coercivity enhancement by grain boundary diffusion process to Nd-Fe-B sintered magnets, IEEE Trans. Magn. 42 (10) (2006) 2909–2911. [11] B.P. Hu, E. Niu, Y.G. Zhao, et al., Study of sintered Nd-Fe-B magnet with high performance of Hcj (kOe) þ (BH)max (MGOe) > 75, AIP Adv. 3 (4) (2013), 042136.
[12] H. Sepehri-Amin, J. Liu, T. Ohkubo, et al., Enhancement of coercivity of hotdeformed Nd–Fe–B anisotropic magnet by low-temperature grain boundary diffusion of Nd60Dy20Cu20 eutectic alloy, Scr. Mater. 69 (9) (2013) 647–650. [13] G. Bai, R.W. Gao, Y. Sun, et al., Study of high-coercivity sintered NdFeB magnets, J. Magn. Magn. Mater. 308 (1) (2007) 20–23. [14] S. Pandian, V. Chandrasekaran, G. Markandeyulu, et al., Effect of Co, Dy and Ga on the magnetic properties and the microstructure of powder metallurgically processed Nd-Fe-B magnets, J. Alloy. Comp. 364 (1–2) (2004) 295–303. [15] S. Hirosawa, Y. Mastuura, H. Yamamoto, et al., Magnetization and magnetic anisotropy of R2Fe14B measured on single crystals, J. Appl. Phys. 59 (3) (1986) 873–879. [16] D. Wu, M. Yue, W.Q. Liu, et al., Magnetic domain switching in Nd-Fe-B sintered magnets with superior magnetic properties, Mater. Res. Lett. 6 (2018) 255–260. [17] X.B. Liu, Z. Altounian, The partitioning of Dy and Tb in NdFeB magnets: a firstprinciples study, J. Appl. Phys. 111 (2012), 07A701.
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Contents lists available at ScienceDirect
Intermetallics journal homepage: http://www.elsevier.com/locate/intermet
MM-Fe-B based gap magnet with excellent energy density Weiqiang Liu a, Zhipeng Zhang a, Ming Yue a, *, Zhi Li a, Dan Wu a, Zhao Zhou a, Hao Chen a, Yuqing Li a, Zaisheng Pang b, Xi Yu b a
College of Materials Science and Engineering, Key Laboratory of Advanced Functional Materials, Ministry of Education of China, Beijing University of Technology, Beijing, 100124, China b Ganzhou Fortune Electronic Co., Ltd., Ganzhou, 341000, China
A R T I C L E I N F O
A B S T R A C T
Keywords: MM-Fe-B magnet Grain boundary diffusion Magnetic properties Coercivity TbHx nanoparticles Magnetic domains
TbHx nanoparticles were applied to prepare sintered MM-Fe-B magnet by using grain boundary diffusion tech nique. Under optimal processing conditions, the coercivity and maximum energy product substantially increase by 383% and 151%, reaching to 6.81 kOe and 20.18 MGOe, respectively. Further investigation indicates that the interchange between Tb atoms and Ce atoms leads to a core/shell structure with (MM,Tb)2Fe14B phase in the surface region of MM2Fe14B matrix grains in the magnet. The modified structure successfully restrains the magnetic domain reversal process, resulting in remarkable enhancement of coercivity and maximum energy product of the magnet.
1. Introduction Misch metal (MM), also called cerium misch metal, is an alloy made from a mixture of rare-earth metals that includes approximately 55% cerium (Ce), 25% lanthanum (La), and other rare earth metals such as neodymium (Nd) and praseodymium (Pr). Herbst studied the intrinsic magnetic properties of R2Fe14B materials, including La2Fe14B (Js¼13.8kG, HA¼20kOe),Ce2Fe14B(Js¼11.7kG, HA¼26kOe),Pr2Fe14B (Js¼15.6kG, HA¼75kOe) and Nd2Fe14B(Js¼16.0kG, HA¼73kOe).Due to the low Js and HA of La2Fe14B and Ce2Fe14B, the intrinsic magnetic properties of MM2Fe14B are not as good as Nd2Fe14B, but it bears excellent potential as a kind of gap magnet between sintered Nd-Fe-B magnet and hard ferrite [1–4]. Moreover, massive applications of MM2Fe14B-based magnets are benefit for utilization of rare earth resource in equilibrium and for energy conservation [5–7]. To date, however, obtaining high energy density [for example (BH)max>20 MGOe] in MM2Fe14B-based magnets is still facing great challenge. By using melt-spinning, different researchers achieved coercivity higher than 6 kOe [7] in MM2Fe14B ribbons, but the magnetically isotropic nature undermine their energy density [8]. Recently, Shang et al. pre pared anisotropic sintered MM-Fe-B magnet, while the low coercivity of only 1.08 kOe also leads to low (BH)max of 7.6 MGOe [9]. Considering the limited resources and price hikes of Nd and Pr, it is urgent to develop a new approach to fabricate sintered MM-Fe-B magnet with high energy density.
Grain boundary diffusion (GBD) technique is an effective way to improve the coercivity of sintered Nd-Fe-B magnets [10–12]. In partic ular, the formation of (Nd, Tb)2Fe14B or (Nd, Dy)2Fe14B is achieved by introducing Terbium (Tb) or Dysprosium (Dy) into the surface region of Nd2Fe14B matrix grains, and the magnet can achieve substantially enhanced coercivity with negligible remanence loss [13–15]. In our recent work, TbHx nanoparticles were applied to improve the magnetic properties of sintered Nd-Fe-B magnet using grain boundary diffusion (GBD) technique, and excellent overall magnetic property (OMP), the sum of Hci and (BH)max, of 85.6 was obtained [16]. Here, we use similar treatment to sintered MM-Fe-B magnet to improve their poor coercivity, and thus to achieve good (BH)max in the magnet. Furthermore, the microstructure, magnetic domain evolution, and magnetic hardening mechanism of the magnet were investigated and discussed. 2. Experiments Misch metals were bought from Bayan Obo in Baotou, China, con taining about 25 wt% La, 55 wt% Ce, 5 wt% Pr and 15 wt% Nd with a purity of about 99.8 wt%. Alloy with a nominal composition of MM30.8FebalAl0.3Cu0.15Co1Ga0.35Zr0.22B0.96 was prepared by strip cast. Subsequently, hydrogen decrepitation (HD) and jet milling (JM) were used to obtain fine powders with diameters of 3–5 μm. The fine powders were arranged under the magnetic field of 1.8 T and pressed into cuboid. In order to restrain the growth of the grains, the lower sintering
* Corresponding author. E-mail address: [email protected] (M. Yue). https://doi.org/10.1016/j.intermet.2019.106626 Received 27 June 2019; Received in revised form 27 September 2019; Accepted 28 September 2019 Available online 11 October 2019 0966-9795/© 2019 Elsevier Ltd. All rights reserved.
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Intermetallics 115 (2019) 106626
Fig. 1. (a) The demagnetization curves of the original and GBD treated MM-Fe-B magnets. Inset table: magnetic properties of both magnets. (b) The distribution of magnetic properties of sintered Nd-Fe-B magnets, sintered MM-Fe-B magnets and Ferrite magnets.
temperature of 1283 K was carried out. The ethanol and TbHx nanoparticles prepared by the evaporationcondensation technology were mixed to obtain heavy rare earth sus pensions. TbHx nanoparticles suspensions were sprayed into the MM-FeB magnets with thickness of 4 mm and formed a TbHx layer with 0.5 wt % of the original magnets. The heat treatment was carried out in vacuum at 1198 K for 8 h, followed by further annealing at 773 K for 3 h. The magnetic properties of the samples were measured by permanent mag netic measuring instrument (NIM-500C). The scanning electron micro scope (SEM, Nova Nano200) with energy dispersive X-ray spectroscopy and electron-probe microanalyzer (EPMA, JEOL JXA-800) with a wavelength-dispersive X-ray detector were applied to observe the microstructure of the magnets. The Magneto-optical Kerr Optical Mi croscope (MOKE, BH-786IP-PK NEOARK CORP) was used to study the magnetic domains and their evolution in the range from 1.4 T to 0.7 T.
hardening for MM-Fe-B magnets. Note that the contents of La and Ce in the MM2Fe14B matrix grains are much higher than the content of Nd, and the magnetocrystalline anisotropy field (HA) of both La2Fe14B and Ce2Fe14B is much lower than that of Nd2Fe14B. It is therefore expected that the coercivity of MM-Fe-B magnets is low. To explore the coercivity enhancement in GBD treated magnet with TbHx nanoparticles, the microstructure of the magnets was analyzed as shown in Fig. 2(e–j). The concentration of Tb decreases gradually from the surface to the inside of the magnet, but it goes a considerable depth along with the grain boundary phase, as shown in Fig. 2(f). Moreover, Tb is mainly enriching in the surface region of MM2Fe14B matrix grains [Fig. 2(j)], forming (MM, Tb)2Fe14B phase with substantially enhanced magnetocrystalline anisotropy. Based on the above observation, it is proposed that Tb atoms first go from the surface to the inside of the magnet via liquid grain boundary phase during the GBD process, and then go into the surface region of the MM2Fe14B matrix grains by exchanging with Ce (or La and Nd) atoms. Our observation and assumption have good agreement with the reporting that Tb tends to enter the Nd2Fe14B matrix phase rather than the Nd-rich phase when the magnet is diffused with Tb [17]. On the other hand, it is found that La prefers to enrich in grain boundary phase, while Nd has a uniform distribution throughout the whole magnet, as shown in Fig. 2(g) and (i). To explore the quantitative changes of rare earth elements distri bution after Tb infiltration, the concentration distribution of La, Ce, Nd, and Tb going through the grain boundary phase and the MM2Fe14B matrix phase of the original and GBD treated MM-Fe-B magnets were examined, as shown in Fig. 3. In the original magnets (Fig. 3(a) and (b)), the counts of La, Ce, and Nd have negligible fluctuation, indicating that the distribution of La, Ce and Nd elements is uniform in the matrix grains. In the GBD treated MM-Fe-B magnets (Fig. 3(c) and (d)), the counts of La and Nd are stable; however, the counts of Ce increase first and then decrease, and the variant trend of Tb count is opposite to that of Ce counts, demonstrating that distribution of Ce element in MM2Fe14B matrix grains has been changed by Tb infiltration. Note that the counts of Tb goes from 200 in the core rapidly to 1300 in the surface region of the grains, for comparison, the counts of Ce goes from 2200 to 1700. Such exchange results in a typical core (MM2Fe14B)/shell [(MM, Tb)2Fe14B] structure. Moreover, it is worth to note that La also has a preference distribution near grain boundary, which is beneficial to the magnetic isolation of matrix grains and therefore good to coercivity of
3. Results & discussion Fig. 1(a) shows the demagnetization curves of the original and GBD treated MM-Fe-B magnets. The maximum energy product of the original magnet reaches to only 8.05 MGOe, and moreover, the intrinsic coer civity is only 1.41 kOe, which limits its practical application. After diffusion, the magnets exhibit much better properties with (BH)max of 20.18 MGOe and coercivity of 6.81 kOe, meaning remarkable increment of 150.7% and 383%, respectively. Meanwhile, the remanence of GBD treated magnets decreases slightly of 0.29 kG (3%) compared to that of original magnets. The moderate energy density of our magnet makes it a proper gap magnet between sintered Nd-Fe-B magnets and Ferrites. In detail, as shown in Fig. 1(b), the (BH)max of Ferrite magnets is less than 5.5 MGOe, and sintered Nd-Fe-B magnets are in the range of 30–52 MGOe, and our newly developed magnet exhibits a medium (BH)max of 20 MGOe. Furthermore, it is worth to note that the fabrication cost of our new magnet is much lower than that of commercial sintered Nd-Fe-B magnets, allowing its good potential in practical applications. Fig. 2(a–d) displays the BSE image and concentration distribution mapping of La, Ce, and Nd in the original MM-Fe-B magnet. It can be found that La and Ce are tending to enrich in the grain boundary phases, which consistent with our previous results [2–4]. Although such ag gregation behavior is beneficial to fabricate La/Ce containing magnets with high magnetic properties, it is difficult to achieve magnetic
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the magnet. For a deeper understanding of the GBD process on the coercivity enhancement of MM-Fe-B sintered magnets, the magnetic domains reversal status of the original and GBD treated MM-Fe-B magnets after fully positive magnetization were observed from 1.4 T to 0.7 T, as shown in Fig. 4. The original MM-Fe-B magnets show single domains status, indicating that the magnets have been fully magnetized at 1.4 T. When the magnetic field reduces to 0 T, it can be found the most grains exhibit maze domain pattern [Fig. 4(b)], which means reverse magnetic domains have nucleated before 0 T. After the magnetic field turns to 0.1 T and 0.2 T, the change of reverse domain status is not obvious, as shown in Fig. 4(c) and (d), indicating that most grains have already completely transferred into the demagnetization status. This result ex plains the low coercivity of original MM-Fe-B magnets. When the mag netic field turns to 0.7 T, almost all of the domains have been reversed magnetized completely. For the GBD treated MM-Fe-B magnets, it ex hibits single domain status at 1.4 T. Different from the original magnets, about half of grains in GBD treated magnets still keep single domain status, marked by red dotted line at 0 T in Fig. 4(g). The maze domains area enlarges with the increase of reverse magnetic field to 0.2 T. However, there still have some grains keeping their original single domain state when the reverse magnetic field reaches to 0.7 T, which means that the reverse nucleation field of these grains is obviously intensified by GBD treatment. Therefore, it can be concluded that GBD treatment plays an important role in inhibiting the nucleation of the reverse domain. On the basis of our finding, it can be pointed out that the magnetic domain reversal modes between the original and GBD treated MM-Fe-B sintered magnets are significantly different. For the MM-Fe-B sintered magnet, the nucleation of the reverse magnetic domain determines the final coercivity. As shown in Figs. 2 and 4(b), most La2Fe14B and Ce2Fe14B phases with low HA in MM2Fe14B grains result in low coer civity. More importantly, the MM2Fe14B grain surface regions usually have defects and therefore exhibit lower magnetocrystalline anisotropy than that of the inner grains. Therefore, they will prior to nucleate. After that, the expansion of the reverse domain would very fast as increasing the reverse magnetic field. For the GBD treated MM-Fe-B magnets, Tb atoms substitute Ce, La and Nd atoms in the surface region of MM2Fe14B grains and forms (MM, Tb)2Fe14B phase. It needs to point out that the HA of the Tb2Fe14B compound is 220kOe, which is much higher than that of the MM2Fe14B compound with 48kOe [4]. As a result, the HN of the surface regions with the (MM,Tb)2Fe14B phase is obviously intensified, therefore restrain the nucleation of reverse domains. Therefore, the grain boundary diffusion could achieve the magnetic hardening for MM-Fe-B magnets, and the coercivity could be obviously intensified. 4. Conclusion Misch metal based permanent magnet with nominal compositions of MM30.8FebalAl0.3Cu0.15Co1Ga0.35Zr0.22B0.96 was prepared. However, its intrinsic coercivity is only 1.41 kOe, which undermines its energy den sity and thus limits its practical application. To overcome this dilemma, grain boundary diffusion with TbHx nanoparticles was carried out to improve the coercivity of MM-Fe-B magnets. The GBD treated magnets exhibit better properties with the maximum energy product of 20.18 MGOe and intrinsic coercivity of 6.81 kOe. A typical core (MM2Fe14B)/ shell [(MM, Tb)2Fe14B] structure has a significant restraint on the reverse nucleation of domains, and the HN of the surface region is obviously intensified. Therefore, the coercivity of GBD treated MM-Fe-B magnets was intensified remarkably. Fig. 2. BSE image (a) and concentration distribution mapping of La (b), Ce (c), and Nd (d) in the original MM-Fe-B magnet; BSE image (e), concentration gradient of Tb from surface to inside in cross section (f), and concentration distribution mapping of La (g), Ce (h), Nd (i), and Tb (j) in the GBD treated MMFe-B magnet.
Declaration of competing interest None.
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Fig. 3. The concentration distribution of La, Ce, Nd, and Tb elements going through the grain boundary phase and the MM2Fe14B matrix phase of the original (a, b) and GBD treated (c, d) MM-Fe-B sintered magnets, the red arrows in (a) and (c) indicate the scanning path. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Fig. 4. MOKE images of magnetic domain evolution of the original (a–e) and GBD treated (f–j) MM-Fe-B magnets in the field from 1.4 T to plane is perpendicular to c-axis).
Acknowledgement
0.7 T. (The observing
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This work was supported by the National Key Research and Devel opment Program of China under Grant 2016YFB0700902; National Natural Science Foundation of China under Grant 51371002; Interna tional S&T Cooperation Program of China under Grant 2015DFG52020; Program of Top Disciplines Construction in Beijing under Grant PXM2019_014204_500031.
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