France



Introduction&lt;br/&gt;The metastable α”-Fe&lt;sub&gt;16&lt;/sub&gt;N&lt;sub&gt;2&lt;/sub&gt; is considered as a candidate for rare earth free semi-hard magnetic materials [1]. It is necessary to improve magnetocrystalline anisotropy of α”-Fe&lt;sub&gt;16&lt;/sub&gt;N&lt;sub&gt;2&lt;/sub&gt; to obtain higher coercivity&lt;i&gt; H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; [2]. Many theoretical calculations and thin film experiments substituting Fe with some elements such as Co, Ni, Mn, Al, etc. have been done. For example, Co is suggested to increase in magnetocrystalline anisotropy [3]. As one of the synthesis methods of α”-Fe&lt;sub&gt;16&lt;/sub&gt;N&lt;sub&gt;2&lt;/sub&gt; nanoparticles, it is well known that the nitridation of α-Fe nanoparticles produced by reduction of iron hydroxide such as α-FeOOH. Previously, we have reported that the influence of substituting Co, Al for Fe by synthesizing α”-(Fe, Co)&lt;sub&gt;16&lt;/sub&gt;N&lt;sub&gt;2&lt;/sub&gt; or α”-(Fe, Al)&lt;sub&gt;16&lt;/sub&gt;N&lt;sub&gt;2&lt;/sub&gt; nanoparticles [5] and revealed that Co substitution results in the temperature regimes for reduction and nitridation were confirmed to very narrow limits, on the other hand Al substitution brings the wider temperature regimes. In this paper, in order to investigate the complex substitution effect of Fe by Co and Al, we challenged to the synthesis of α”-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)&lt;sub&gt;16&lt;/sub&gt;N&lt;sub&gt;2&lt;/sub&gt; nanoparticles by hydrogen reduction of α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)OOH as a starting material and subsequent nitridation.&lt;br/&gt;&lt;br/&gt;Experimental&lt;br/&gt;We used commercially available α-(Fe&lt;sub&gt;1-x&lt;/sub&gt;M&lt;sub&gt;x&lt;/sub&gt;)OOH (M= Co, Al, x= 0, 0.02, 0.05) and α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)OOH nanoparticles as a starting material. The conditions for reduction and nitridation were systematically examined to obtain higher yield of α”-(Fe&lt;sub&gt;1-x&lt;/sub&gt;M&lt;sub&gt;x&lt;/sub&gt;)&lt;sub&gt;16&lt;/sub&gt;N&lt;sub&gt;2&lt;/sub&gt; or α”-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)&lt;sub&gt;16&lt;/sub&gt;N&lt;sub&gt;2 &lt;/sub&gt;according to the following procedure. The α-(Fe&lt;sub&gt;1-x&lt;/sub&gt;M&lt;sub&gt;x&lt;/sub&gt;)OOH nanoparticles were first reduced at 300 - 500 °C for 4 hours in a H&lt;sub&gt;2&lt;/sub&gt; gas flow. Nitridation was subsequently performed at 150 - 180 °C for 5 hours in an NH&lt;sub&gt;3&lt;/sub&gt; gas flow. Crystallite diameter&lt;i&gt; D&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; of reduced α-Fe&lt;sub&gt;1-x&lt;/sub&gt;M&lt;sub&gt;x&lt;/sub&gt; nanoparticles was measured as average diameter of nanoparticles using Scherrer’s formula based on X-ray diffraction (XRD) with Co K&lt;sub&gt;α&lt;/sub&gt; radiation. Morphological changes were observed using transmission electron microscopy (TEM). For magnetic properties measurements, these nanoparticles were packed into a disk shape (6 mm diameter, 1 mm thick) with resin. Magnetization &lt;i&gt;M&lt;/i&gt;&lt;sub&gt;15k&lt;/sub&gt; and coercivity&lt;i&gt; H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; were measured using vibrating sample magnetometer (VSM) with a maximum applied magnetic field of 15k Oe.&lt;br/&gt;&lt;br/&gt;Results&lt;br/&gt;Fig. 1 (a) shows the regions in the nitridation temperature &lt;i&gt;versus&lt;/i&gt; reduction temperature diagrams corresponding to where the α”-phase was produced. &lt;i&gt;M&lt;/i&gt;&lt;sub&gt;15k&lt;/sub&gt; and &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; values of the assembled nanoparticles are also indicated. α”-phase was produced under the limited conditions of reduction temperature of 340 °C and nitridation temperature of 170 °C for starting from α-(Fe&lt;sub&gt;0.98&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;)OOH. On the contrary, α”-phase was produced under broad conditions of reduction temperatures of 400 - 500 °C and nitridation temperatures of 155 - 170 °C for starting from α-(Fe&lt;sub&gt;0.98&lt;/sub&gt;Al&lt;sub&gt;0.02&lt;/sub&gt;)OOH. In the case of starting from α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)OOH, α”-phase was produced under the conditions of reduction temperatures of 400 - 500 °C and nitridation temperature of 160 °C. The conditions of α”-phase production were broadened by Al addition to Co. Fig. 1 (b) shows TEM images for the nanoparticles after nitridation starting from α-(Fe&lt;sub&gt;0.98&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;)OOH, α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Al&lt;sub&gt;0.05&lt;/sub&gt;)OOH and α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)OOH. These nanoparticles are synthesized with hydrogen reduction at 340 °C for 4 hours and subsequent nitridation at 170 °C for 5 hours for α-(Fe&lt;sub&gt;0.98&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;)OOH, and with hydrogen reduction at 400 °C for 4 hours and subsequent nitridation at 160 °C for 5 hours for α-(Fe&lt;sub&gt;0.98&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;)OOH and α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)OOH, respectively. The particle size of nanoparticles after nitridation starting from α-(Fe&lt;sub&gt;0.98&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)OOH are between the ones of starting from (Fe&lt;sub&gt;0.98&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;)OOH and α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Al&lt;sub&gt;0.05&lt;/sub&gt;)OOH. This result suggested that Al was effective to avoid sintering during reduction.&lt;br/&gt;Fig. 2 (a) shows the relationship between &lt;i&gt;M&lt;/i&gt;&lt;sub&gt;15k&lt;/sub&gt; and &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; for the nanoparticles after nitridation starting from α-(Fe&lt;sub&gt;0.98&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;)OOH, α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Al&lt;sub&gt;0.05&lt;/sub&gt;)OOH and α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)OOH. In the case of starting from α-FeOOH was also shown. Higher &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; were obtained by the complex of Co and Al compared to Co or Al alone. Fig. 2 (b) shows &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; change as a function of &lt;i&gt;D&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; of those nanoparticles. This result indicates that there is a good relationship between &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; and &lt;i&gt;D&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; regardless additive composition. &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; increases with decrease in &lt;i&gt;D&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt;. In the case of nanoparticles after nitridation starting from α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)OOH gives the smallest &lt;i&gt;D&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; of 22 nm and highest &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; of 1800 Oe. It is assumed that Al addition to Co is effective to prevent crystalline growth which leads to high &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; of α”-phase. To clarify the reason of increase in &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt;, further precise investigation of the distribution of Al and the relationship between anisotropy field of nanoparticles after nitridation will be presented in this conference.

**References:**

[1] R. Skomski et al., Scr. Mater., 112 (2016) 3&lt;br/&gt;[2] M. Tobise et al., J. Magn. Soc. Jpn, 41 (2017) 58&lt;br/&gt;[3] Y. Jiang et al., AIP Adv., 6 (2016) 056007&lt;br/&gt;[4] L. Ke et al., Phys. Rev. B, 88 (2013) 024404&lt;br/&gt;[5] M. Tobise et al., AIP Adv., 9 (2019) 035233

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/af71dcf495e025c51610d79bc24cb5e1.jpg)
Fig. 1. (a) The regions in the nitridation temperature versus reduction temperature diagrams corresponding to where the α”-phase was produced and &lt;i&gt;M&lt;/i&gt;&lt;sub&gt;15k&lt;/sub&gt; and &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; values of the nanoparticles are also indicated. (b) TEM images for the nanoparticles after nitridation starting from the material α-(Fe&lt;sub&gt;0.98&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;)OOH, α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Al&lt;sub&gt;0.05&lt;/sub&gt;)OOH and α-(Fe&lt;sub&gt;0.98&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)OOH.

![](https://assets.underline.io/uploads/markdown_image/1/image/b01b30df78a588a568b32b8156ed275d.jpg)
Fig. 2. (a) The relationship between &lt;i&gt;M&lt;/i&gt;&lt;sub&gt;15k&lt;/sub&gt; and &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c&lt;/sub&gt; for the nanoparticles after nitridation starting from the material α-(Fe&lt;sub&gt;0.98&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;)OOH, α-(Fe&lt;sub&gt;0.95&lt;/sub&gt;Al&lt;sub&gt;0.05&lt;/sub&gt;)OOH and α-(Fe&lt;sub&gt;0.98&lt;/sub&gt;Co&lt;sub&gt;0.02&lt;/sub&gt;Al&lt;sub&gt;0.03&lt;/sub&gt;)OOH. (b) &lt;i&gt;H&lt;/i&gt;&lt;sub&gt;c &lt;/sub&gt;change as a function of crystalline diameter &lt;i&gt;D&lt;/i&gt;&lt;sub&gt;c &lt;/sub&gt;of those nanoparticles.

INTERMAG 2021

 Synthesis of α”-(Fe, M)16N2 nanoparticles obtained by hydrogen reduction and subsequent nitridation starting from α-(Fe, M)OOH (M= Co, Al)

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Introduction The metastable α”-Fe16N2 is considered as a candidate for rare earth free semi-hard magnetic materials [1]. It is necessary to improve magnetocrystalline anisotropy of α”-Fe16N2 to obtain higher coercivity Hc [2]. Many theoretical calculations and thin film experiments substituting Fe with some elements such as Co, Ni, Mn, Al, etc. have been done. For example, Co is suggested to increase in magnetocrystalline anisotropy [3]. As one of the synthesis methods of α”-Fe16N2 nanoparticles, it is well known that the nitridation of α-Fe nanoparticles produced by reduction of iron hydroxide such as α-FeOOH. Previously, we have reported that the influence of substituting Co, Al for Fe by synthesizing α”-(Fe, Co)16N2 or α”-(Fe, Al)16N2 nanoparticles [5] and revealed that Co substitution results in the temperature regimes for reduction and nitridation were confirmed to very narrow limits, on the other hand Al substitution brings the wider temperature regimes. In this paper, in order to investigate the complex substitution effect of Fe by Co and Al, we challenged to the synthesis of α”-(Fe0.95Co0.02Al0.03)16N2 nanoparticles by hydrogen reduction of α-(Fe0.95Co0.02Al0.03)OOH as a starting material and subsequent nitridation. Experimental We used commercially available α-(Fe1-xMx)OOH (M= Co, Al, x= 0, 0.02, 0.05) and α-(Fe0.95Co0.02Al0.03)OOH nanoparticles as a starting material. The conditions for reduction and nitridation were systematically examined to obtain higher yield of α”-(Fe1-xMx)16N2 or α”-(Fe0.95Co0.02Al0.03)16N2 according to the following procedure. The α-(Fe1-xMx)OOH nanoparticles were first reduced at 300 - 500 °C for 4 hours in a H2 gas flow. Nitridation was subsequently performed at 150 - 180 °C for 5 hours in an NH3 gas flow. Crystallite diameter Dc of reduced α-Fe1-xMx nanoparticles was measured as average diameter of nanoparticles using Scherrer’s formula based on X-ray diffraction (XRD) with Co Kα radiation. Morphological changes were observed using transmission electron microscopy (TEM). For magnetic properties measurements, these nanoparticles were packed into a disk shape (6 mm diameter, 1 mm thick) with resin. Magnetization M15k and coercivity Hc were measured using vibrating sample magnetometer (VSM) with a maximum applied magnetic field of 15k Oe. Results Fig. 1 (a) shows the regions in the nitridation temperature versus reduction temperature diagrams corresponding to where the α”-phase was produced. M15k and Hc values of the assembled nanoparticles are also indicated. α”-phase was produced under the limited conditions of reduction temperature of 340 °C and nitridation temperature of 170 °C for starting from α-(Fe0.98Co0.02)OOH. On the contrary, α”-phase was produced under broad conditions of reduction temperatures of 400 - 500 °C and nitridation temperatures of 155 - 170 °C for starting from α-(Fe0.98Al0.02)OOH. In the case of starting from α-(Fe0.95Co0.02Al0.03)OOH, α”-phase was produced under the conditions of reduction temperatures of 400 - 500 °C and nitridation temperature of 160 °C. The conditions of α”-phase production were broadened by Al addition to Co. Fig. 1 (b) shows TEM images for the nanoparticles after nitridation starting from α-(Fe0.98Co0.02)OOH, α-(Fe0.95Al0.05)OOH and α-(Fe0.95Co0.02Al0.03)OOH. These nanoparticles are synthesized with hydrogen reduction at 340 °C for 4 hours and subsequent nitridation at 170 °C for 5 hours for α-(Fe0.98Co0.02)OOH, and with hydrogen reduction at 400 °C for 4 hours and subsequent nitridation at 160 °C for 5 hours for α-(Fe0.98Co0.02)OOH and α-(Fe0.95Co0.02Al0.03)OOH, respectively. The particle size of nanoparticles after nitridation starting from α-(Fe0.98Co0.02Al0.03)OOH are between the ones of starting from (Fe0.98Co0.02)OOH and α-(Fe0.95Al0.05)OOH. This result suggested that Al was effective to avoid sintering during reduction. Fig. 2 (a) shows the relationship between M15k and Hc for the nanoparticles after nitridation starting from α-(Fe0.98Co0.02)OOH, α-(Fe0.95Al0.05)OOH and α-(Fe0.95Co0.02Al0.03)OOH. In the case of starting from α-FeOOH was also shown. Higher Hc were obtained by the complex of Co and Al compared to Co or Al alone. Fig. 2 (b) shows Hc change as a function of Dc of those nanoparticles. This result indicates that there is a good relationship between Hc and Dc regardless additive composition. Hc increases with decrease in Dc. In the case of nanoparticles after nitridation starting from α-(Fe0.95Co0.02Al0.03)OOH gives the smallest Dc of 22 nm and highest Hc of 1800 Oe. It is assumed that Al addition to Co is effective to prevent crystalline growth which leads to high Hc of α”-phase. To clarify the reason of increase in Hc, further precise investigation of the distribution of Al and the relationship between anisotropy field of nanoparticles after nitridation will be presented in this conference.

**References:**

[1] R. Skomski et al., Scr. Mater., 112 (2016) 3 [2] M. Tobise et al., J. Magn. Soc. Jpn, 41 (2017) 58 [3] Y. Jiang et al., AIP Adv., 6 (2016) 056007 [4] L. Ke et al., Phys. Rev. B, 88 (2013) 024404 [5] M. Tobise et al., AIP Adv., 9 (2019) 035233

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/af71dcf495e025c51610d79bc24cb5e1.jpg)
Fig. 1. (a) The regions in the nitridation temperature versus reduction temperature diagrams corresponding to where the α”-phase was produced and M15k and Hc values of the nanoparticles are also indicated. (b) TEM images for the nanoparticles after nitridation starting from the material α-(Fe0.98Co0.02)OOH, α-(Fe0.95Al0.05)OOH and α-(Fe0.98Co0.02Al0.03)OOH.

![](https://assets.underline.io/uploads/markdown_image/1/image/b01b30df78a588a568b32b8156ed275d.jpg)
Fig. 2. (a) The relationship between M15k and Hc for the nanoparticles after nitridation starting from the material α-(Fe0.98Co0.02)OOH, α-(Fe0.95Al0.05)OOH and α-(Fe0.98Co0.02Al0.03)OOH. (b) Hc change as a function of crystalline diameter Dc of those nanoparticles.

poster

In recent years, much efforts has been paid to synthesize new compounds having the potential for permanent magnet application. One of the new compounds is the RFe12 (R = rare earth) intermetallic compound.1-3 Although the binary RFe12 intermetallic compounds with the ThMn12-type structure are not thermodynamically stable and do not exist in bulk form, it has been found that the substitution of T (T = Ti, V, Cr, Mo) for Fe in the RFe12 intermetallic compound stabilizes the ThMn12-type structure.4-6 The objective of the present study was to produce anisotropic Sm(Fe,Ti)12 magnets. In the case of Nd-Fe-B magnets, anisotropic Nd-Fe-B magnets have been produced by either the sintering of Nd-Fe-B powder or hot deformation of hot-pressed Nd-Fe-B magnets. In this study, we attempted to produce anisotropic Sm(Fe,Ti)12 magnets by the hot deformation of hot-pressed Sm(Fe,Ti)12 magnets. The magnetic properties of the hot-deformed Sm(Fe,Ti)12 magnets are discussed. Sm(Fe,Ti)12 alloy ingots were prepared by induction melting in an argon atmosphere. The alloy ingots were induction melted and then ejected through an orifice with argon onto a copper wheel. The resultant melt-spun ribbons were comminuted into powders and then hot-pressed by the spark plasma sintering (SPS) method. Hot deformation of the hot-pressed Sm(Fe,Ti)12 magnets was performed in a vacuum at 973–1173 K under a pressure of ~200 MPa by the SPS method. The specimens were examined by X-ray diffraction (XRD), a scanning electron microscope (SEM), a transmission electron microscope (TEM), and a vibrating sample magnetometer (VSM). The hot-pressed Sm(Fe,Ti)12 magnets were successfully hot deformed to an 80% reduction without cracking. The magnetic properties of the resultant hot deformed Sm(Fe,Ti)12 magnets were examined. Figure 1 shows the hysteresis loops of the hot-deformed Sm(Fe,Ti)12 magnets, measured parallel and perpendicularly to the pressing direction. The magnetic anisotropy of the hot-deformed Sm(Fe,Ti)12 magnets was confirmed by the existence of a higher remanence in the parallel direction than in the perpendicular direction. This indicates that anisotropic Sm(Fe,Ti)12 magnets were successfully produced by the hot deformation of hot-pressed Sm(Fe,Ti)12 magnets. In order to examine the crystallographic alignment of the Sm(Fe,Ti)12 phase, the surface of the SmFe11Ti magnets were examined by XRD. The results are shown in Figure 2. The XRD patterns of the hot-deformed Sm(Fe,Ti)12 magnets show clear diffraction peaks of the Sm(Fe,Ti)12 phase, but the XRD pattern of the Sm(Fe,Ti)12 magnet hot-deformed at 1173 K was slightly different from that of the Sm(Fe,Ti)12 magnet hot-deformed at 973 K. It was found that the hot-deformed Sm(Fe,Ti)12 magnets contained a small amount of the α-Fe phase together with the Sm(Fe,Ti)12 phase. Although the magnetic measurements revealed that the hot-deformed Sm(Fe,Ti)12 magnets exhibited slight magnetic anisotropy, they did not show clear crystallographic alignment of its Sm(Fe,Ti)12 phase. In conclusion, Sm(Fe,Ti)12 magnets were successfully produced by the hot deformation. Magnetic measurements revealed that the hot-deformed Sm(Fe,Ti)12 magnets exhibited a relatively low coercivity value but were slightly magnetically anisotropic. However, no clear crystallographic alignment of its Sm(Fe,Ti)12 phase was observed in the XRD studies.

**References:**

[1] F. R. De. Boer, et al. J. Less-Common Met. 135, 199 (1987). [2] L. X. Liao, Z. Altounian, and D. H. Ryan, J. Appl. Phys. 70, 6006 (1991). [3] L. Shultz and J. Wecker, J. Appl. Phys. 64, 5711 (1988). [4] Y. Wang, et al. J. Appl. Phys. 67, 4954 (1990). [5] Y. Otani, H. S. Li, and J. M. D. Coey, IEEE Trans. Magn. 26, 2658 (1990). [5] J. Yang, Ph. Oleinek, K.-H. Müller, J. Appl. Phys. 88, 988 (2000). [6] E. Tomey, et al. J. Alloys Compd. 231, 195 (1995).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/7d88021d8a27ee02aab711e56a845f03.jpg)
Fig. 1 Hysteresis loops of the hot-deformed Sm(Fe,Ti)12 magnets 
produced at (a) 973 K, (b) 1073 K, and (c) 1173 K.

![](https://assets.underline.io/uploads/markdown_image/1/image/b74f592c60860d1bb2f93a16317539c9.jpg)
Fig. 2 XRD patterns of the hot-deformed Sm(Fe,Ti)12 magnets 
produced at (a) 973 K, (b) 1073 K, and (c) 1173 K.

 Magnetic properties of Sm(FeTi)12 hot deformed magnets

At present, rare-earth-based permanent magnet (REPM) R2Fe14B is one of the most widely used permanent magnets with excellent performance. However, improving this material's properties is still the main topic so far in this field [1]. For R2Fe14B alloys, it needs to use heavy rare earth elements to obtain high coercive force, which is not conducive to the balanced application of rare earth resources. In contrast, R2Fe14C system has the same tetragonal structure as R2Fe14B but a larger magnetic anisotropy field, especially when R is a light rare earth element [2]. This means a great potential to obtain high coercive force in this system, which is necessary for developing medium performance and high coercive force magnets based on high abundance light rare earth elements [3]. However, the R2Fe14C system suffers from the low phase stability of the tetragonal phase produced by a sluggish solid-state reaction process. Therefore, in this study, we carried on a thorough investigation in the Pr2Fe14C system considering the effects of element doping and heat treatment on the phase evolution. Fig.1 shows the XRD patterns of the as-spun and annealed samples of Pr15Fe79(C1-xBx)6 (x = 0, 0.05, 0.2, 0.4 and 0.6) series. In the as-spun flakes, only α-Fe and Pr2Fe17 phase are found in the samples with x = 0, 0.05, 0.2 and 0.4, while the Pr2Fe14C phase were not observed until x = 0.6. It suggests that, without enough amount of B doping, α-Fe and Pr2Fe17 can not form 2:14:1 phase via solid-state transformation in Pr-Fe-C system. On the other hand, after the annealing process at 750 °C, the 2:14:1 phase starts to form even in the x = 0.05 sample. As shown in Fig. 1(b), the amount of the 2:14:1 phase is small then, but it becomes dominant when the B content is 0.2. The amount of α-Fe dramatically decreases at the same time. Due to the high rare-earth content in these samples, there is also a pronounced peak at 2θ = 30° of the Pr-rich phase crystallized during the annealing. The oxidation of the rare-earth-rich phase during the annealing process causes the other peak at 2θ = 29°. More importantly, the lower limit of the solid-state reaction temperature of the 2:14:1 phase decreases continuously with the doping of B, from 950 °C to 600 °C. It means that the B doping reduces the forming energy of the 2:14:1 phase. Fig.2 shows the initial magnetization and demagnetization curves for the annealed samples with different B contents. With the increase of B content, both the remanent magnetization (Mr) and the coercivity field (Hc) exhibit an overall enhancement. This is consistent with the XRD results, indicating that the content of the 2:14:1 phase is increased. There is also a waist-like feature on the demagnetization curve of the samples with low B content, which is a sign of the coexistence of soft and hard magnetic phases. Even though the magnetization is not saturated at 3 T, the Mr and Hc are still as high as 70.5 emu/g and 19.1 kOe, respectively, in the sample with x=0.6. This considerable value of coercivity may be related to the fine microstructure of the samples. As can be seen from the inset of Fig.2, the average grain size is about 2 mm and does not change too much with the increase of annealing temperature. This is because the main phase is surrounded by a rare-earth-rich phase, limiting the growth of the main phase grains. This is another advantage of the R2Fe14C system. We systematically investigated the phase evolution of the Pr2Fe14C system and established an effective way to obtain a relatively large percent of 2:14:1 phase. By optimizing the composition and process, we can also control the microstructure of the main phase and obtain promising magnetic properties. The Pr2Fe14C material has great potential to be used for bonded magnets and hot deformed magnets. Acknowledgments This work was supported by the Beijing Natural Science Foundation (No. 2182006), the General Program of Science and Technology Development Project of Beijing Municipal Education Commission of China (No. KM201710005006), the National Natural Science Foundation of China [No.51401002], and Program of Top Disciplines Construction in Beijing (PXM2019_014204_ 500031).

**References:**

[1] Buschow, K. H. J., New developments in hard magnetic materials. Reports on Progress in Physics 1991, 54 (9), 1123. [2] Burzo, E., Permanent magnets based on R-Fe-B and R-Fe-C alloys. Reports on Progress in Physics 1998, 61 (9), 1099. [3] Coey, J. M. D., Permanent magnets: Plugging the gap. Scripta Materialia 2012, 67 (6), 524-529.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/2dd1f0ce47d374061700e3ab8fde8143.jpg)

![](https://assets.underline.io/uploads/markdown_image/1/image/d6ceed3758db8ef26d6326330b67b88b.jpg)

 Phase and magnetism evolution in Pr2Fe14C system upon B doping and heat treatment

The high performance permanent magnets are indispensable for many commercial and military applications. Rare-earth hard magnetic nanoparticles have drawn great attention due to their potential applications in nanocomposite permanent magnets with high performance. Different methods have been employed to prepare hard magnetic nanoparticles including chemical methods, electrodeposition, spark erosion and ball milling. The ball milling technique, used for the preparation of nanoparticles, is a simple, inexpensive and efficient method for size reduction of nanocrystalline powders. The experimental parameters involved in ball-milling technique, such as milling time, milling energy, and ball size, are determinant in what concerns the magnetic properties of the nanoparticles. The milling temperature and milling medium are also known as critical factors impacting the morphology and performance of the nanoparticles. Here, we report results on the morphology, structure, and magnetic properties of Ce14Fe76Co4B6 nanoparticles prepared by high-energy ball milling at low temperature, in liquid nitrogen, from nanaocristaline ribbons precursors. At regulate intervals of time, in the milling vial was introduced liquid nitrogen, in order to compensate the evaporation. We found that: (i) the optimal interval for introducing liquid nitrogen into the vial is 20 minutes, since for a longer interval, the liquid nitrogen evaporates and the sample becomes contaminated, (ii) the morphology and size of the powders change with the milling time, from large particles with irregular shapes to small dimensions (about 150 nm) and rounded shapes for a milling time of 420 minutes (the microstructure of powders milled 420 h, is shown in Figure 1.) (iii) the structure and magnetic properties of the nanoparticles were kept close to those of the precursor ribbons by milling in liquid nitrogen (a comparison between the hysteresis curves for powders and the precursor ribbons is shown in Figure 2). The Ce14Fe76Co4B6 nanoparticles obtained by ball milling in liquid nitrogen have great potential to be used for the production high performance nanocomposite permanent magnets. Acknowledgement This work was supported by a grant of the Romanian Ministry of Education and Research, CCCDI - UEFISCDI, project number PN-III-P2-2.1-PED-2019-2054, within PNCDI III. &lt;!--![endif]----&gt;&lt;!--![endif]----&gt;&lt;!--![endif]----&gt;

**References:**

[1] Baolu Zhou, Zhongxue Li and Congcong Chen, Global Potential of Rare Earth Resources and Rare Earth Demand from Clean Technologies, Minerals 7 (2017) 203. [2] Thanh Anh Huynh and Min-Fu Hsieh, Performance Analysis of Permanent Magnet Motors for Electric Vehicles (EV) Traction Considering Driving Cycles, Energies 11 (2018) 1385. [3] M. Grigoras, M. Lostun, G. Stoian, D.D. Herea, H. Chiriac, N. Lupu, Microstructure and magnetic properties of Ce10+xFe84−xB6 nanocrystalline ribbons versus preparation conditions, J. Magn. Magn. Mater. 432 (2017) 119–123. [4] J. F. Herbst, M. S. Meyer and F. E. Pinkerton, Magnetic hardening of Ce2Fe14B, J. Appl. Phys. 111, 07A718 (2012).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/eaaa31845ba980e59e6fd93570814151.jpg)
Fig 1. The SEM micrograph of Ce14Fe76Co4B6 
powders milled 420 h in liquid nitrogen.

![](https://assets.underline.io/uploads/markdown_image/1/image/c84ddc682ad13f3ca51fdaac70d5be13.jpg)
Fig 2. The hysteresis curve of Ce14Fe76Co4B6 ribbons (black) and powders (red) milled 420 h in liquid nitrogen.

 Characterization and Magnetic properties of Ce-FeCo-B Nanoparticles Prepared by Cryo-milling

HDDR process can not only refine the coarse R2Fe14B grains (~ 100 μ m) into fine grains (~300 nm), but also induce these fine grains to be oriented along the c-axis. The grain refinement can improve the coercivity of the magnetic materials, and the grain orientation is helpful to obtain high remanence, so the anisotropic R2Fe14B materials with high coercivity, remanence and maximum energy products can be obtained after HDDR process. Recent researches show that the coercivity of HDDR R2Fe14B magnetic materials can be enhanced to ~20 kOe by refining the grain size of the materials or by diffusion of the eutectic alloys with the low melting point, but there is still a big gap between the coercivity and its anisotropic field (70 kOe). This indicates that the coercivity of HDDR magnetic materials still has a lot of room to improve, especially the grain size of HDDR powder is close to that of R2Fe14B single domain particles. In order to further improve the coercivity, it is necessary to study the coercivity mechanism of HDDR magnetic materials. In this study, the initial magnetization curve of thermally demagnetized HDDR anisotropic Pr-Fe-B sample was firstly investigated. The results show that the magnetization increase moderately with external magnetic field, but it is not only different from the rapid rise of magnetization of sintered R-Fe-B magnets but also different from the slow increase of melt-spun R-Fe-B powders at the beginning and then the sudden rapid increase at a certain critical field. Therefore, it is not easy to determine the type of the coercivity mechanism for the powders. In order to study the coercivity mechanism of the above HDDR anisotropic powders in detail, the temperature dependence of coercivity of the sample was measured. The coercivity decrease monotonically with the increase of temperature. According to the micromagnetic model, the coercivity Hc can be generally expressed Hc=αφαk2K1/MS-NeffMs. Due to the preferential nucleation of inhomogeneous regions in the magnetic powder, a correction factor αk should be introduced. In addition, because the grain alignment in the magnetic powder is not completely consistent, the correction factor αφ is also introduced. The dependence of the anisotropy constants K1, and saturation magnetization (Ms) of the main phase Pr2Fe14B on the temperature are all from the literature. The measured data of HDDR Pr-Fe-B and the theoretical formula were compared to analyze that the coercivity mechanism is controlled by the nucleation of the reversed domain or by the uniform rotation of the magnetic moment. （1）Uniform rotation model: Firstly, in order to analyze whether the coercivity mechanism belongs to uniform rotation of the magnetic moments, the dependence of HC/MS on 2K1/Ms2 was investigated. It is found that the experimental data only fall on a straight line in a part of the temperature region. Even in the section of the line, the slope of the line is not 1, which does not meet the theoretical requirements of uniform rotation. So we can exclude the possibility that the magnetization process is uniform rotation of the magnetic moments. （2）Strongly correlated reverse domain nucleation model When the grains in the magnetic powder are strongly magnetically coupled, a reversion of a misaligned grain would drive the reversion of the neighboring grains. Therefore, the coercivity of the material depends on the grains with the smallest nucleation field. In this case, αφ= αφmin, The coefficient can be calculated for a given temperature. The value of this angle is about 40 ~ 60 degrees on the basis of the calculation results of the external field angle corresponding to the minimum nucleation field αφmin at 50K ~ 350K. Fig.1 shows the HC/MS vs 2αφminK1/Ms2 curve of the powder. The purpose of this curve is to analyze whether the coercivity mechanism of the above powder is the nucleation mechanism of strongly magnetically coupled grains. It can be seen from the figure that the experimental data points basically fall on a straight line, which means that under the framework of micromagnetism, the nucleation mechanism with the magnetically coupled grains can describe the magnetization reversal process of the HDDR powder. Fig.2 shows magnetic domain patterns of the HDDR treated Pr-Fe-B alloys. In the case of the HDDR treated alloys, very fine maze-like magnetic domains were observed and these domains extend over more than 1 μm in the original large single particles with the mean size of 10 μm (Fig.2 (b) and (c)). This domain pattern shows a characteristic of interaction domains, and each domain consists of several fine grains, which are magnetically coupled to form large interaction domains. The evolution of the magnetic domain of HDDR anisotropic Pr-Fe-B magnetic powder under the external applied magnetic field (0 T → 2 T → 0 T → -2 T → 0 T → 2 T) was also investigated. Only the disappearance of the labyrinth domain composed of multiple grains and the appearance of the magnetic domain under the condition of positive magnetization field and negative reverse magnetization field were observed, while the domain wall movement was not observed, which further implies that the nucleation of demagnetization is the main mechanism controlling the coercivity of anisotropic HDDR Pr-Fe-B powder.

**References:**

[1] T. Takeshita, R. Nakayama, in 11th Workshop on Rare Earth Magnets and their Applications, Pittsburg, USA, 49 (1990). [2] K. D. Durst and H. Kronmüller, J. Magn. Magn. Mater. 68, 63 (1987). [3] H. Kronmüller, K. D. Durst, and M. Sagawa, J. Magn. Magn.Mater. 74, 291（1988）. [3] W. L. Chen, W. Liu, and Z. D. Zhang, Phy. Rev. B, 72，224419 (2005)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/573cdeaf9631d4c74e5e89dbcdd1548c.jpg)
Fig.1 The dependence of Hc/Ms on 2𝛼𝜑minK1/MS2 of HDDR powders

![](https://assets.underline.io/uploads/markdown_image/1/image/6c2d8f8fab2a230db89ba7a5b00ecf97.jpg)
Fig. 2 Magnetic domain Kerr images of Pr-Fe-B alloys (a) before and (b) after the HDDR treatment. (c) zoom in of red rectangle region in the figure (b)

 Study on magnetization reversal processes of anisotropic HDDR Pr2Fe14B-type magnetic materials

Cobalt as a substitute for Iron, and combinations of different rare earths in alloy can increase the Curie temperature and coercivity of Nd2Fe14B type ferromagnets[1]. Complex ferromagnets, with atomic scale ordering, require high resolution simulations which are computationally expensive due to individual treatment of atoms. The compromise of the atomistic model[2], allows nanometer scale simulations with reasonable efficiency, but requires input parameters often derived from ab initio calculations, or Density functional theory, creating a multi-scale model. Complex materials often have several input parameters, such as magnetic exchange and magnetocrystalline anisotropy constants, which cannot be easily evaluated using these methods. As a result, often the requirement is to fit the exchange parameter to macroscopic experimental data. For simple materials this can be evaluated with few test simulations, as the relationship between exchange and macroscopic properties behaves linearly for single element materials, as in the mean field model. Complex materials can have complex relationships between exchange and macroscopic properties, so many test simulations are often required to perform a full sweep of the parameter space. This body of work details the evaluation of exchange and anisotropy constants for 2:14:1, or Nd2Fe14B type rare earth ferromagnets, using the VAMPIRE atomistic software modelling package. Treating this as an optimisation problem reduces the number of atomistic simulations required for a given numerical precision of constant. Evaluation of exchange constants was performed using the Curie temperature as a figure of merit with several different methods of evaluating the Curie temperature compared for robustness and consistency, with the optimisation of these exchange constants performed with a combination of linear searching and root finding algorithms. These methods have often been compared for various rock magnets[3]. The search uses a locally linear assumption, to prime the low temperature quantum rescaling in VAMPIRE, and to better parallelise the initial stages of the search. Similar approaches are taken to calculate the anisotropy constants of Nd2Fe14B to reproduce the First Order Magnetic Phase Transition[4] at high external fields and low temperatures. The exchange constant of various rare-earth ferromagnets is reproduced, yielding the correct macroscopic Curie temperatures without reliance on density functional theory/ab initio calculations. Lanthanum was used as the prototype for the Fe-Fe exchange, and this was fixed throughout the other ferromagnets. The Curie temperatures calculated are in good agreement with J.F.Herbst, 1991 [5]. The First Order Magnetic Phase Transition was effectively reproduced, but requires additional tuning for a complete analysis of the relationship between second and fourth order anisotropy constants and the subsequent macroscopic magnetic behaviour.

**References:**

[1] Wei, L, et al. J. Less Common Metals 126 (1986); [2] R. F. L. Evans, W. J. Fan, P. Chureemart, T. A. Ostler, M. O. A.Ellis and R. W. Chantrell, Journal of Physics: Condensed Matter,vol. 26, 103202, (2014); [3] K. Fabian,V. P. Shcherbakov, S.A. McEnroe, Geochemistry, Geophysics, Geosystems, 14, 947, (2013); [4] J. M. Cadogan, J. P. Gavigantt, D. Givordt and H. S. Li, J. Phys. F: Met. Phys. 18 779 (1988); [5] J. F. Herbst Rev. Mod. Phys. 63, 819 (1991)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/4b9dd3a6e04c6e1367a35ec04d9cfb05.jpg)
The exchange paramaters calculated using the optimsiation method result in appropriate Curie temperatures when evaluated using the VAMPIRE atomsitic modelling software package. The Curie temperature calculated from this simulation compares well to the values from literature[5] (Herbst, 1991). 
 
Ce2Fe14B has not been calculated, as the rare earth behaves sufficiently differently to Lanthanum as to prevent using Lanthanum as the base of Cerium's Fe-Fe exchange in this manner. 
 
The magnitude of the exchange is presented, the heavier rare earths are ferrimagnetic in the 2:14:1 phase.

 Optimisation of atomistic modelling parameters for Nd2Fe14B-type rare earth ferromagnets, using the Curie temperature and first order magnetic phase transitions as figures of merit

High-performance, high-efficiency motors are indispensable for the development of electric vehicles and the realization of electric aircraft. The key is the development of permanent magnets that can withstand use even at high temperatures. The reported Sm-Fe-based magnet compounds [1-3] have a high Curie temperature, and the values of the magnetic anisotropy field and saturation magnetization are higher or comparable to those of Nd2Fe14B. In this work, we focused on the Induction thermal plasma(ITP) process as a new process for obtaining anisotropic nano-sized powder of these permanent magnet compounds. The ITP process is a physical bottom-up process that can synthesize nano-sized single-crystal powders due to the high quenching effect during the process [4, 5]. When this process can be applied to prepare a fine permanent magnet anisotropic powder of around 100 nm, which is difficult to obtain by the conventional method, can be synthesized, leading that an anisotropic magnet having a large coercive force can be obtained. Therefore, as the first step, we tried to produce nano-sized anisotropic Sm-Fe-N alloy powder and to evaluate the feasibility of preparing the nano-sized powder by the ITP process. The mixed powder of Sm powder (D50 = 18 μm) and Fe powder (particle size about 3 μm, High Purity Chemical Laboratory) was used as a starting powder for the ITP process. The ITP process was performed under the condition of process pressure of 100 kPa. The input power was 6 kW. Since Sm metal powder as a starting material is not commercially available, we prepared micron-sized Sm powder by using a crucible-less skull gas atomizing device (manufactured by Sinfonia Technology Co., Ltd.). Though the ITP process, the obtained nano-sized powder was sealed in a quartz capillary in a glove box, and the crystal structure was evaluated by a powder X-ray diffraction experiment with Aichi Synchrotron BL5S2 without oxidizing the powder. In addition, SEM and STEM were used for microstructure analysis. From the XRD profile, SmFe2 and Sm-Fe with TbCu7-type structure were identified as Sm-Fe alloy phases. From the SEM image, the average particle size of the SmFe alloy was about 100 nm. The Sm-Fe alloy particles were found to be single crystals by high-resolution STEM observation. In addition, many particles consist of Fe-core/Sm-shell and Fe-core/SmFe-shell were also observed. Combine the results obtained from the XRD measurement and the STEM observation, it was found that the Fe detected by XRD is not the Fe particle but the Fe core in the core-shell structure. This is because the vapor pressure of Sm is more than 10000 times higher than that of Fe, so in the thermal plasma process, Fe nucleates first and then Sm and Fe or Sm agglomerates around the Fe particles. After the nitriding process by the heat treatment at an elevated temperature of 400C in N2 flow for 15min, nitrogen was surely introduced into the Sm-Fe lattice, which is confirmed from the lattice expansion. For the first time in this study, we succeeded in synthesizing Sm-Fe-N anisotropic particles with a particle size of 100 nm or less using the ITP process. Especially, the preparation of the anisotropic particle with TbCu7-type structure will lead to the next generation high-performance permanent magnet.

**References:**

[1] T Iriyama, K Kobayashi, N Imaoka et al., IEEE Trans. Magn., Vol.28, p.2326 (1992) [2] S. Sakurada, A. Tsutai, T. Hirai et al., J. Appl. Phys., Vol.79, p4611 (1996) [3] Y. Hirayama, Y.K. Takahashi, S. Hirosawa et al., Scr. Mater., Vol.138, p.62 (2017) [4] Y. Hirayama and K. Takagi, J. Alloys. Compd., Vol.792, p.594 (2019) [5] Y. Hirayama, M. Shigeta, Z. Liu et al., Available at http://dx.doi.org/10.2139/ssrn.3692994

**Images:**

 Nano-sized anisotropic Sm-Fe-N particle preparation by induction thermal process

Permanent magnets are a key component of a wide range of everyday devices, including motors, actuators, transducers and sensors. Depending on the application needs, magnets are based on high performance rare-earth transition metal phases or on more moderate performance hard ferrite phases, and they may be sintered or polymer bonded. While the remanence of bonded magnets is necessarily limited by dilution of the magnetic material in a non-magnetic matrix, they are cheaper in both material and fabrication costs. What is more, the mechanical properties of the magnets can be tuned according to the choice of matrix. Traditional fabrication routes for polymer bonded magnets include calendering, injection moulding, extrusion, and compression bonding [1] and there has been much recent activity concerning 3D printing of polymer bonded magnets [2,3,4]. Down-sizing magnets opens new opportunities for their use in micro-scaled devices with applications in fields as diverse as telecommunications, energy management and bio-medicine. Sputtering combined with micro-patterning has been used to make fully dense magnets which are a micro-scale equivalent to sintered magnets while various techniques including casting and screen printing combined with photo-lithography have been used to make polymer bonded micro-magnets [5]. Sputtered films may be isotropic or textured, depending on the deposition temperature, whereas micro-scaled polymer bonded magnets reported to date are isotropic. Here we present a process developed to fabricate arrays of anisotropic, flexible bonded micro-magnets. The study focuses on micro-magnets prepared using commercial single crystalline hard ferrite (Br: 195 mT, Hc: 123 kA/m) or Sm-Fe-N (Br: 1.31 T, Hc: 889 kA/m) powders mixed with polydimethylsiloxane (PDMS). This binder was chosen so as to have flexible micro-magnets that can be used to study the influence of mechanical forces on cell development [6]. The innovative aspect of this process, schematized in Fig. 1, consists in the possibility to transfer a specific pattern from a silicon mould, made by deep reactive ion etching (DRIE), to a magnetically aligned powder-polymer composite. The composite is maintained within the mould during the alignment step by coverage with a glass slide. A first soft curing step is performed with the cover-slide in place, then a final curing is performed to attach the micro-magnets to a PDMS base. The procedure has been used to fabricate arrays of micro-magnets of thickness ranging from 50 µm to 300 µm and width in the range from 20 µm to 4000 µm. As a magnetic field source, bulk magnets, a superconductive coil and pulsed magnetic field systems were tested. The degree of alignment was assessed by measuring the hysteresis loops of fully cured micro-magnet arrays with the measuring field applied parallel or perpendicular to the direction of the alignment field. Note that these measurements were made prior to removal of the arrays from the Si mould, to prevent distortion of the flexible micro-magnets under the torque induced by the measuring field. As an example, easy and hard axes hysteresis loops made on an array of SmFeN-bonded square micro-pillars with in-plane dimensions of 200x200 µm2 and a height of 300 µm, aligned in a magnetic field of 800 mT, are compared in Fig. 2 a. As can be expected, higher intensity fields were needed to align the SmFeN particles compared to the ferrite particles. The stray fields produced by aligned and non-aligned micro-magnet arrays were measured using an in-house developed Scanning Hall Probe Microscope [7]. In Fig. 2 b we present the z-component of the stray field produced at a height of 50 µm above the SmFeN micro-magnet array featured in Fig. 2 a. In this presentation, we will give details of structural and magnetic characterisation as well as simulations of arrays of ferrite and SmFeN bonded micro-magnets. We will also report on the actuation of these micro-magnets under the influence of a 3D magnetic field. Finally, we will present prospects for the further development and use of such micro-magnet arrays.

**References:**

[1] J. Ormerod and S. Constantinides, Journal of Applied Physics, Vol. 81, p.4816 (1997) [2] C. Huber, C. Abert and F. Bruckner, Appl. Phys. Lett. , Vol. 109, p.162401 (2016) [3] L. Li, A. Tirado and I. C. Nlebedim, Scientific Reports, Vol. 6, p.36212 (2016) [4] K. Sonnleitner, C. Huber and I. Teliban, Appl. Phys. Lett., Vol. 116, p.092403 (2020) [5] N.M. Dempsey, "Nanoscale Magnetic Materials and Applications", J.P. Liu et al.(eds), p.661, (Springer 2009) [6] N.M. Dempsey, " Handbook of Magnetic Materials", J.M.D. Coey et al. (eds), Springer, (in-press) [7] G. Shaw, R.B.G. Kramer and N. M. Dempsey, Rev Sci Inst, Vol.87, p.113702 (2016)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/3a78eeb2c7e66ef3542f55dd6887a55f.jpg)
Fig. 1 Schematic of the process developed for the fabrication of arrays of flexible, anisotropic bonded micro-magnets.

![](https://assets.underline.io/uploads/markdown_image/1/image/a01987af8c35d1a195ab777fd1bb4d16.jpg)
Fig. 2 (a) Hysteresis loops of an array of SmFeN (30 vol %) - PDMS (70 vol %) bonded micro- magnet pillars with in-plane dimensions of 200x200 µm2 and a height of 300 µm. (b) Z-component of the stray magnetic field produced by this array, measured using a micro-Hall probe scanned at a height of 50 µm above the magnet array.

 Fabrication and characterisation of polymer-bonded flexible anisotropic micro-magnet arrays

The CoPt and FePt alloys, when chemically ordered in the L10 phase [1], are among the magnetic materials displaying the highest magnetic anisotropy constant. Therefore, they are perfect candidates for ultra-high density magnetic storage applications, provided nanoparticles can be prepared in such a high anisotropy phase. In the bulk L10 phase both alloys exhibit close magnetic and structural properties. Indeed, the tetragonalization of the unit cell, measured by the c/a ratio is almost the same and equal to 0.97 for both alloys while their anisotropy constant is equal to 5 and 7 MJ.m-3 respectively for CoPt and FePt [1]. The first part of this work is then to determine and compare the intrinsic magnetic anisotropy constant of CoPt and FePt L10 nanoparticles. For this purpose, we dilute, as far as possible, mass-selected nanoparticles (around 3 nm diameter) in an inert carbon matrix to avoid coalescence and to make magnetic interactions negligible. To promote the chemically L10 ordered phase the samples have to be annealed at 600°C during two hours. The magnetic properties are then determined using SQUID magnetometer and a global fitting procedure of the entire set of magnetic measurements [2,3] (see Fig. 1). In a second part we try to explain their significant differences using extended X-ray absorption fine structure measurements and atomic relaxation calculations (see Fig. 2). We conclude that the very high magnetic anisotropy measured in FePt nanoparticles [4] is related to the weak atomic relaxation in this alloy, while for chemically ordered CoPt nanoparticles the strong atomic relaxation in the Co plans [5] seems to reduce drastically the magnetic anisotropy.

**References:**

[1] P. Andreazza et al., Surf. Sci. Rep. 70, 188 (2015) [2] A. Tamion et al., Phys. Rev. B 85, 134430 (2012) [3] A. Hillion et al., Phys. Rev. B 88, 094419 (2013) [4] P. Capiod et al., Phys. Rev. Lett. 122, 106802 (2019) [5] N. Blanc et al., Phys. Rev. B 87, 155412 (2013)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/bd365171dd94e366a992d84cc0f1e80e.jpg)
Fig. 1: Low temperature hysteresis loop measured for a diluted assembly of CoPt nanoparticles (left) and FePt nanoparticles (right), after annealing. The dots correspond to measurements while the line is obtained with a global fit taking into account the particle size and magnetic anisotropy distributions. Note that despite the particles are similar (around 3.5 nm diameter), the coercivity is much larger for FePt nanoparticles.

![](https://assets.underline.io/uploads/markdown_image/1/image/3e51fbabcc3e62496eafd8c27717286a.jpg)
Fig. 2: Comparison between the mean c/a ratio (computed from interatomic distances: homoatomic and heteroatomic bonds) around transition metal atoms and Pt atoms, from density functional calculations of relaxed CoPt or FePt clusters. N represents the number of atoms used to build perfect truncated octahedrons. The lines are guide to the eye, while the dashed line indicates the c/a ratio in the bulk L10 phase.

 Magnetic anisotropy of chemically ordered CoPt and FePt nanoparticles, why is it so different?

In the 1980s ferrous materials were the first subject of extensive investigation on magnetic field annealing (MFA). Hereby, the potential of using high external magnetic fields as an additional degree of freedom for materials synthesis and tailoring of microstructures in solid state was demonstrated. Since then, due to the progression in magnetic materials and superconducting magnets, processing in magnetic field has become both less costly and more viable. Hence, MFA is achieving renewed attention nowadays. The processing technique might even evolve into a valuable tool on broad industrial level, since it provides the opportunity to save energy by lowering heat treatment temperatures and shortening processing times. Furthermore, the application of an external high magnetic field as an additional independent processing parameter allows for stabilization of metastable phases and creation of texturized microstructures that are not accessible on zero-field conditions. Nevertheless, a lot of experimental work and modeling is still required to entirely unveil the mechanisms of MFA for gaining control on such material manipulation. Mn3Ga is a very interesting and highly suitable candidate for exploring MFA. The intermetallic compound occurs in three different crystal structures, all of them possessing rich magnetic properties attractive for different kinds of application. Regarded as a rare-earth-free permanent magnet in potential application, especially the tetragonal phase meets the particular requirements of being exceptionally prone to manipulation by MFA. Due to the ferrimagnetic ordering it exhibits a much higher magnetic moment compared to the other Mn3Ga polymorphs. Therefore, it is supposed to be thermodynamically favored in its formation under influence of an external magnetic field. Additionally, the tetragonal structure shows high anisotropy in susceptibility, which is conditional to establish a texture upon MFA. As already reported by Ener et al., the phase transition from hexagonal to tetragonal Mn70Ga30 can be accelerated remarkably by combining severe plastic deformation and MFA [1]. In contrast, the study at hand adds the investigation of the influence of different degrees of deformation plus the impact of MFA on the phase transformation starting from a predominantly cubic structure of the parent material Mn3Ga. For this purpose annealing was performed at 400°C, lasting for at least 0.5 hours and up to 14 days. In the case of MFA samples, a magnetic flux density of 7 T was applied. By X-ray diffraction the different stages of phase transformation were observed. Further characterization regarding microstructure and magnetic properties was achieved using light and electron microscopy as well as SQUID magnetometry. The results indicate that the properties after long-term annealing are very similar, independent of previous state of deformation and use of magnetic field. More specifically, after 14 days of annealing samples attained a phase-pure tetragonal structure and show equal remanence. Nevertheless, within the first 24 hours of annealing the differences are enormous with respect to phase formation, microstructure and magnetic properties. In case of zero-field annealing of highly deformed samples or MFA of undeformed samples, reaching the tetragonal structure is postponed by a pronounced intermediate formation of the hexagonal phase. Apart from that, combining previous deformation and MFA accelerates the phase transformation effectively. Furthermore, with increasing degree of deformation the coercivity can be increased stepwise. However, the development of coercivity over annealing time is complex for both zero-field annealing and MFA. An attempt to solve this complex behavior by relating microstructure and coercivity is presented. Finally, the phase transformation behavior of Mn3Ga under influence of an external magnetic field is correlated to changes in defect mobility and development of local deviations in composition (see Fig. 1).

**References:**

[1] S. Ener, K.P Skokov, D.Y. Karpenkov, M.D. Kuz’min and O. Gutfleisch, Journal of Magnetism and Magnetic Materials, 382, 265-270 (2015)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/b5271fcdca9ea6f7afc41d135535331f.jpg)
Fig. 1: EDS line scan of highly deformed sample after 1 h of MFA 
A) SEM picture showing local course of scan (bright areas in hexagonal, grey areas in tetragonal structure) 
B) compositional mismatch along line scan

 Manipulation of Phase Transformation and Microstructure in Mn3Ga Using Magnetic Field Annealing

Recently, the topic of rare-earth-free permanent magnets has received considerable attraction for the alternative to rare-earth magnets due to the resource crisis and the high cost of rare-earth metals. Among rare-earth-free permanent magnetic materials, the low-temperature phase (LTP) of MnBi has been considered as a potential alternative because it exhibits an unusual positive temperature coefficient of coercivity in the temperature range 153-550 K. Because of the above magnetic property, the MnBi-LTP has a higher coercivity compared to Nd2Fe14B permanent magnetic material at elevated temperature, and it was recently reported that the coercivity of MnBi-LTP reached up to 28.3 kOe at 530K. According to previous researches, this excellent coercivity characteristic of the MnBi-LTP is attributed to its high magneto-crystalline anisotropy (K1≈1.6 x 106 J/m-3 at 300 K)[1]. In spite of such excellent intrinsic magnetic properties, the MnBi-LTP bulk magnets practically fabricated have a significantly lower energy product compared to its theoretical value (BHmax ≈ 16 MGOe at 300 K) due to engineering issues that occur during the mechanical milling process such as decomposition of the magnetic phase, manganese oxidation, and agglomeration. At present, several approaches including high or low energy ball-milling (HEBM or LEBM), surfactant-assisted ball-milling (SABM) have been reported to avoid those issues. However, those approaches based on a ball milling process at room temperature have inevitably a restriction in reducing the powder size as well as agglomerating the powder by cold welding phenomenon. Consequently, this leads to the degradation of the energy product by lowering the total anisotropic characteristics of the MnBi-LTP bulk magnets. Therefore, it is essential to synthesize the particle size of the MnBi-LTP to be close to that of its single magnetic domain (≈ 500 nm) in order to achieve the MnBi-LTP bulk magnets with improved anisotropic characteristics. In this regard, it has been proposed that the cryo-milling method is able to be one of the approaches to synthesize the MnBi-LTP particle satisfying the above-required size. It is well known that the cryo-milling method is able effectively to reduce the particle size as well as to avoid cold welding and agglomeration of the power because it carried out at a temperature lower than 123 K, where materials become brittle and friable[2]. It has been reported that the cryo-milling method is successfully synthesized nanocrystalline structures in ductile materials such as Fe, Ni, Al, Zn, etc. However, researches into the cry-milling of MnBi alloys have been rarely conducted. In this study, the Mn54Bi46-LTP powders were synthesized by the cryo-milled method combined with the addition of a surfactant (stearic acid) as well as calcium, and a schematic illustration of the process is shown in figure 1. As shown in the figure, a stainless-steel vial containing a mixture of powder, balls, surfactant, and calcium was vibrated with a certain frequency, and the temperature was maintained at 77 K during the milling process by the flow of liquid nitrogen. In addition, surfactant (stearic acid) and calcium were added in order to suppress the particle agglomeration and manganese oxidation, respectively. Figure 2 shows the Hysteresis loops for the cryo-milled Mn54Bi46 powder according to the addition of surfactant and calcium. All of the milling conditions have an identical ball-to-powder ratio of 18:1, and stearic acid and calcium were added 0.225g and 0.01g to Mn54B46 powder of 0.9 g, respectively. As shown in the figure, the cry-milled Mn54B46 powder without any addition has a Hc of 14.1 kOe with a remanence (Mr) of 40.7 emu/g and a magnetization (Ms,2.5T) of 40.7 emu/g, while the cry-milled Mn54B46 powders with the addition of stearic acid have a higher both Mr of 51.0 emu/g and Ms,2.5T of 53.2 emu/g. These results, especially the increase of Mr/Ms,2.5T ratio, indicate that the powders were effectively separated by the addition of stearic acid during the milling process. In addition, the result of improved Ms,2.5T means that the decomposition rate of the magnetic phase decreases as the impact energy received by the powder is alleviated by the stearic acid on the powder surface. Furthermore, from the result of magnetic properties under the calcium addition conditions, it is found that calcium addition plays an important role in the purification of Mn54Bi46-LTP powder by suppressing manganese oxidation. Based on these results, further investigation on the optimization of the cryo-milling process and the magnetic properties of bulk magnets manufactured with cryo-milled powders will be discussed.

**References:**

[1] V. N. Antonov and V. P. Antropov, Low Temperature Physics, vol. 46, no. 1, p.1 (2020) [2] N. K. Katiyar, K. Biswas, and C. S. Tiwary, International Materials Reviews, p.1 (2020)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/f2a8abc3721902719bd228a4fc2e173e.jpg)
Fig. 1. A schematic illustration of the cryo-milling process used in this study.

![](https://assets.underline.io/uploads/markdown_image/1/image/ad03f42558598a49f705daf6a85b3f5b.jpg)
Fig. 2. Hysteresis loops for the cryo-milled Mn54Bi46 powder measured parallel to the alignment direction.

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