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I. INTRODUCTION&lt;br/&gt;A phase with extraordinarily high room-temperature coercivity in the Sm-Fe-Ti system was first observed in thin films crystallized from an amorphous precursor where μ&lt;sub&gt;0&lt;/sub&gt;&lt;i&gt;H&lt;sub&gt;c&lt;/sub&gt;&lt;/i&gt; reached 3.85 T for a Sm&lt;sub&gt;2&lt;/sub&gt;Fe&lt;sub&gt;7&lt;/sub&gt;Ti&lt;sub&gt;1&lt;/sub&gt; composition [1]. Giant room-temperature coercivities up to 5.03 T were reported later [2] in a ‘new and possibly metastable phase’ with Sm&lt;sub&gt;2&lt;/sub&gt;Fe&lt;sub&gt;7&lt;/sub&gt;Ti&lt;sub&gt;1&lt;/sub&gt; composition prepared by mechanical alloying. The reason for the maximum in coercivity at the Sm&lt;sub&gt;2&lt;/sub&gt;Fe&lt;sub&gt;7&lt;/sub&gt;Ti&lt;sub&gt;1&lt;/sub&gt; composition has been a subject of debate. The observed unit cell expansion is smaller than expected when replacing the respective Fe atoms with Ti, which led to the conclusion that the amount of Ti in the Sm&lt;sub&gt;5&lt;/sub&gt;Fe&lt;sub&gt;17&lt;/sub&gt; phase has to be smaller than weighted and thus the remaining Ti must form a secondary phase [3]. The interesting question arises then, whether the high coercivity is caused solely by the intrinsic magnetic properties of the Sm&lt;sub&gt;5&lt;/sub&gt;Fe&lt;sub&gt;17&lt;/sub&gt; phase or is a result of a multi-phase microstructure. To answer this, in this work we have performed a detailed microstructural investigation of samples with varying Ti contents and correlated our findings with the observed magnetic properties.&lt;br/&gt; &lt;br/&gt;II. EXPERIMENTAL&lt;br/&gt;Mechanical alloying of Sm-Fe-Ti samples was carried out by high energy ball milling (SPEX 8000D) under argon atmosphere from elemental Fe (74 -149 μm, 99.9%, Alfa Aesar), Ti (-74 μm, 99%, Alfa Aesar) powders and Sm pieces (3-4 mm, 99.9%, ChemPUR). The resultant powders were subsequently compacted, and heat treated in a horizontal tube furnace at different temperatures. The samples were handled in an Ar filled glovebox (p(O&lt;sub&gt;2&lt;/sub&gt;) &amp;lt; 0.1 ppm, MBraun) to avoid oxidation.&lt;br/&gt;Magnetic measurements were performed using a Quantum Design Physical Property Measurement System (PPMS) vibrating sample magnetometer (VSM). The magnetization measurements were extended up to 43 T using a non-destructive pulsed-field coil at the Dresden High Magnetic Field Laboratory. A single 1.44 MJ capacitor bank was used. When fully charged, it could produce a maximum magnetic field of 60 T with a rise time of about 7 ms and a total pulse duration of 25 ms. In our experiments, the capacitor module was charged to about two thirds. The magnetization was detected by the induction method using a coaxial pick-up coil system surrounding the sample. All pulsed-field data were calibrated against the magnetization recorded in steady fields.&lt;br/&gt; &lt;br/&gt;III. RESULTS AND DISCUSSION&lt;br/&gt;As shown in Fig. 1a, ultrahigh coercivities, 7.18 T at room temperature (8.86 T at 10 K), were achieved for the Sm&lt;sub&gt;20&lt;/sub&gt;Fe&lt;sub&gt;70&lt;/sub&gt;Ti&lt;sub&gt;10&lt;/sub&gt; alloy. The coercivity as high as 2.18 T is maintained even at 500 K. Anisotropy field μ&lt;sub&gt;0&lt;/sub&gt;&lt;i&gt;H&lt;sub&gt;a&lt;/sub&gt;&lt;/i&gt; of 20.7±0.8 T determined from high-field pulse measurements demonstrates that &lt;i&gt;H&lt;sub&gt;c&lt;/sub&gt;&lt;/i&gt; reaches 35% of the &lt;i&gt;H&lt;sub&gt;a&lt;/sub&gt;&lt;/i&gt; at room temperature. Large switching field of around 7 T is visible in the initial magnetization curve. This behavior is similar to fine-grained (below the critical single-domain grain size) Nd-Fe-B alloys.&lt;br/&gt;The Curie temperature &lt;i&gt;T&lt;sub&gt;C&lt;/sub&gt;&lt;/i&gt; of Sm&lt;sub&gt;20&lt;/sub&gt;Fe&lt;sub&gt;70&lt;/sub&gt;Ti&lt;sub&gt;10&lt;/sub&gt; is 577 K and calculated exchange stiffness parameter &lt;i&gt;A&lt;/i&gt; is 7.72 pJ/m. Because of high rare earth content as well as presence of a nonmagnetic secondary phase, the magnetization was only 60 A m&lt;sup&gt;2&lt;/sup&gt; kg&lt;sup&gt;-1&lt;/sup&gt;, which is roughly 0.58 T.&lt;br/&gt;Detailed transmission electron microscopy investigations presented in Fig. 1b show two-phase microstructure consisting of the Sm&lt;sub&gt;5&lt;/sub&gt;Fe&lt;sub&gt;17&lt;/sub&gt;-based hard magnetic matrix phase with size below 200 nm and fine, &amp;lt;100 nm, Fe&lt;sub&gt;2&lt;/sub&gt;Ti grains. Majority of the Fe&lt;sub&gt;2&lt;/sub&gt;Ti secondary phase grains are located at the grain boundaries as well as inside of the 5:17 grains. No distinct segregation of any element at the grain boundaries can be seen. Despite high areal fraction of the Fe&lt;sub&gt;2&lt;/sub&gt;Ti grains, near single-phase demagnetization loops are observed.&lt;br/&gt;Aiming at enhancing &lt;i&gt;M&lt;sub&gt;s&lt;/sub&gt;&lt;/i&gt; by realizing Sm&lt;sub&gt;5&lt;/sub&gt;Fe&lt;sub&gt;17&lt;/sub&gt;/α-Fe composite magnets, the effect of Ti content on the phase constitution, magnetic properties and microstructure was investigated in detail for Sm&lt;sub&gt;20&lt;/sub&gt;Fe&lt;sub&gt;70&lt;/sub&gt;Ti&lt;i&gt;&lt;sub&gt;x&lt;/sub&gt;&lt;/i&gt; with 4≤&lt;i&gt;x&lt;/i&gt;≤10. The fraction of Fe&lt;sub&gt;2&lt;/sub&gt;Ti phase decreases with &lt;i&gt;x&lt;/i&gt; and for x=4 Sm&lt;sub&gt;2&lt;/sub&gt;Fe&lt;sub&gt;17&lt;/sub&gt; phase appears. &lt;i&gt;M&lt;sub&gt;s&lt;/sub&gt;&lt;/i&gt; increases and &lt;i&gt;H&lt;sub&gt;c&lt;/sub&gt;&lt;/i&gt; decreases for the Ti-lean compositions. Based on the microstructural investigations, it is concluded that the reduced coercivities are the consequence of both intrinsic and extrinsic factors, such as the lower anisotropy field, larger grain size and smaller volume fraction of the Fe&lt;sub&gt;2&lt;/sub&gt;Ti precipitates.&lt;br/&gt;&lt;br/&gt;ACKNOWLEDGEMENTS&lt;br/&gt;We acknowledge funding by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation), Project ID No. 405553726-TRR 270, and Elements Strategy Initiative Center for Magnetic Materials (ESICMM), Grant Number JPMXP0112101004, through the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan. We also acknowledge the support of the HLD at HZDR, member of the European Magnetic Field Laboratory (EMFL).

**References:**

1. N. Kamprath, N.C. Liu, H. Hegde, F.J. Cadieu, J. Appl. Phys. 64 (1988) 5720–5722.&lt;br/&gt;2. L. Schultz, K. Schnitzke, J. Wecker, J. Magn. Magn. Mater. 83 (1990) 254–256.&lt;br/&gt;3. H. Stadelmaier, G. Schneider, E.T. Henig, M. Ellner, Mater. Lett. 10 (1991) 303–309.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/8e90a5cf8fd14c1e53d388b8def552ce.jpg)
Fig. 1 (a) Room temperature hysteresis loop and (b) the corresponding microstructure of the Sm&lt;sub&gt;20&lt;/sub&gt;Fe&lt;sub&gt;70&lt;/sub&gt;Ti&lt;sub&gt;10&lt;/sub&gt; sample. Two-phase microstructure consisting of the Sm&lt;sub&gt;5&lt;/sub&gt;Fe&lt;sub&gt;17&lt;/sub&gt;-based hard magnetic matrix phase with fine Fe&lt;sub&gt;2&lt;/sub&gt;Ti grains is visible.



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 Magnetic properties and microstructure of Sm5Fe17-based two-phase magnets

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I. INTRODUCTION A phase with extraordinarily high room-temperature coercivity in the Sm-Fe-Ti system was first observed in thin films crystallized from an amorphous precursor where μ0Hc reached 3.85 T for a Sm2Fe7Ti1 composition [1]. Giant room-temperature coercivities up to 5.03 T were reported later [2] in a ‘new and possibly metastable phase’ with Sm2Fe7Ti1 composition prepared by mechanical alloying. The reason for the maximum in coercivity at the Sm2Fe7Ti1 composition has been a subject of debate. The observed unit cell expansion is smaller than expected when replacing the respective Fe atoms with Ti, which led to the conclusion that the amount of Ti in the Sm5Fe17 phase has to be smaller than weighted and thus the remaining Ti must form a secondary phase [3]. The interesting question arises then, whether the high coercivity is caused solely by the intrinsic magnetic properties of the Sm5Fe17 phase or is a result of a multi-phase microstructure. To answer this, in this work we have performed a detailed microstructural investigation of samples with varying Ti contents and correlated our findings with the observed magnetic properties. II. EXPERIMENTAL Mechanical alloying of Sm-Fe-Ti samples was carried out by high energy ball milling (SPEX 8000D) under argon atmosphere from elemental Fe (74 -149 μm, 99.9%, Alfa Aesar), Ti (-74 μm, 99%, Alfa Aesar) powders and Sm pieces (3-4 mm, 99.9%, ChemPUR). The resultant powders were subsequently compacted, and heat treated in a horizontal tube furnace at different temperatures. The samples were handled in an Ar filled glovebox (p(O2) &lt; 0.1 ppm, MBraun) to avoid oxidation. Magnetic measurements were performed using a Quantum Design Physical Property Measurement System (PPMS) vibrating sample magnetometer (VSM). The magnetization measurements were extended up to 43 T using a non-destructive pulsed-field coil at the Dresden High Magnetic Field Laboratory. A single 1.44 MJ capacitor bank was used. When fully charged, it could produce a maximum magnetic field of 60 T with a rise time of about 7 ms and a total pulse duration of 25 ms. In our experiments, the capacitor module was charged to about two thirds. The magnetization was detected by the induction method using a coaxial pick-up coil system surrounding the sample. All pulsed-field data were calibrated against the magnetization recorded in steady fields. III. RESULTS AND DISCUSSION As shown in Fig. 1a, ultrahigh coercivities, 7.18 T at room temperature (8.86 T at 10 K), were achieved for the Sm20Fe70Ti10 alloy. The coercivity as high as 2.18 T is maintained even at 500 K. Anisotropy field μ0Ha of 20.7±0.8 T determined from high-field pulse measurements demonstrates that Hc reaches 35% of the Ha at room temperature. Large switching field of around 7 T is visible in the initial magnetization curve. This behavior is similar to fine-grained (below the critical single-domain grain size) Nd-Fe-B alloys. The Curie temperature TC of Sm20Fe70Ti10 is 577 K and calculated exchange stiffness parameter A is 7.72 pJ/m. Because of high rare earth content as well as presence of a nonmagnetic secondary phase, the magnetization was only 60 A m2 kg-1, which is roughly 0.58 T. Detailed transmission electron microscopy investigations presented in Fig. 1b show two-phase microstructure consisting of the Sm5Fe17-based hard magnetic matrix phase with size below 200 nm and fine, &lt;100 nm, Fe2Ti grains. Majority of the Fe2Ti secondary phase grains are located at the grain boundaries as well as inside of the 5:17 grains. No distinct segregation of any element at the grain boundaries can be seen. Despite high areal fraction of the Fe2Ti grains, near single-phase demagnetization loops are observed. Aiming at enhancing Ms by realizing Sm5Fe17/α-Fe composite magnets, the effect of Ti content on the phase constitution, magnetic properties and microstructure was investigated in detail for Sm20Fe70Tix with 4≤x≤10. The fraction of Fe2Ti phase decreases with x and for x=4 Sm2Fe17 phase appears. Ms increases and Hc decreases for the Ti-lean compositions. Based on the microstructural investigations, it is concluded that the reduced coercivities are the consequence of both intrinsic and extrinsic factors, such as the lower anisotropy field, larger grain size and smaller volume fraction of the Fe2Ti precipitates. ACKNOWLEDGEMENTS We acknowledge funding by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation), Project ID No. 405553726-TRR 270, and Elements Strategy Initiative Center for Magnetic Materials (ESICMM), Grant Number JPMXP0112101004, through the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan. We also acknowledge the support of the HLD at HZDR, member of the European Magnetic Field Laboratory (EMFL).

**References:**

1. N. Kamprath, N.C. Liu, H. Hegde, F.J. Cadieu, J. Appl. Phys. 64 (1988) 5720–5722. 2. L. Schultz, K. Schnitzke, J. Wecker, J. Magn. Magn. Mater. 83 (1990) 254–256. 3. H. Stadelmaier, G. Schneider, E.T. Henig, M. Ellner, Mater. Lett. 10 (1991) 303–309.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/8e90a5cf8fd14c1e53d388b8def552ce.jpg)
Fig. 1 (a) Room temperature hysteresis loop and (b) the corresponding microstructure of the Sm20Fe70Ti10 sample. Two-phase microstructure consisting of the Sm5Fe17-based hard magnetic matrix phase with fine Fe2Ti grains is visible.

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1Introduction Manufacturing of electrical machines, in general, involves cutting of electrical steel sheets to appropriate sizes. Different cutting techniques such as punching, electrical discharge machining (EDM) and laser cutting are popular. These methods cause damage inside the material close to the cutting edge, which results in deterioration of magnetic properties [1]. Existing studies are based on measurement result by Epstein frame, Ring sample and single-sheet tester. However, these methods do have problems, causing error when analyzing influence of manufacture on material properties[2]. For example, the such non-ideal factors as overlapping joint causes significant eddy current loss especially at high frequency, which may make effect of manufacture effect negligible. In order to obtain more accurate measurement results, this paper aims to propose a new method to clarify the cutting effect on magnetic properties of silicon steel, eliminating those non-ideal factors, and to find out mechanism behind the deterioration of magnetic properties based on macroscopic magnetic measurement and microstructure observation. 2Effect of Cutting Methods on Magnetic Properties of Electrical Steel 2.1Macroscopic Magnetic Measurement The effect of cutting on the electrical steel sheets is analyzed by magnetic measurement on steel strips of different widths. The magnetic test samples were 35WW270 electrical steel lamination (300 mm×30 mm×0.35 mm). Half of the samples were cut along rolling direction while the other half were in the transverse direction. Then the samples have been cut along their length into different width by shearing, EDM, laser cutting, and glued up with polyimide tapes. Magnetic measurements are usually performed on electrical steel samples with Epstein frame. Nevertheless, the Epstein frame has some intrinsic weaknesses. One of them is that the flux density and specific total loss is uneven near the corners. The Epstein frame can be divided as uniform area, transitional area and overlapping area. In order to eliminate the impact of overlapping area and transitional area, two Epstein frames were designed as Fig 1, E(25), 25cm (standard) and E(12.5), 12.5cm (non-standard) Epstein frames, which have identical corner areas and transitional areas. The only difference in size is that the uniform area of E(12.5) is shorter than E(25). The same steel sheet (300mm length *30mm width) were used in both Epstein frames and their total absolute loss was measured. Magnetic measurement are performed following standard IEC: 604004-2. It can be assumed the absolute loss of corner and affected area of E(25) and E(12.5) is equal because of the identical corners and affected area. Then the only difference in absolute power loss between the two Epstein frames is due to the difference in absolute loss of the different uniform areas. The magnetic properties of the samples were measured at 50-800 Hz, and the results at 50 Hz were shown in Fig 2. When B = 1 T and frequency is 50 Hz, compared with EDM, the iron loss of samples with shearing increased by 5.6% and that with laser cutting increased by 32.6%. It can be found that the difference decreased at a higher frequency.The iron loss can be divided into static loss and dynamic loss. With the frequency increasing, dynamic loss takes more share of total iron loss[3]. Therefore, it can be assumed that the effect of cutting stresses on static loss is greater than that on dynamic loss. 2.1.1Microstructure Observation The steel samples were prepared by grinding on 400, 800, and 1200 grit SiC papers and polishing using 3μm diamond paste to obtain a mirror-like surface. Specimen microstructure was examined by scanning electron microscopy (SEM). Fig 2 shows the microstructure of the samples with different cutting methods observed by SEM. There are different contrast effects due to the plastic deformation near the cutting edge. It can be seen that the width of damaged area of EDM is about 20 μm, smaller than that of shearing (83 μm). And it is obvious that there is melting or heat affect zone (108 μm) in the immediate proximity of the cut edge due to laser cutting procedure, which leads to the worst magnetic properties of samples with laser cutting. 3Conclusion This paper has compared and analyzed the effect of different cutting methods on the magnetic properties of electrical steel by double Epstein frames. It is found that EDM has the least influence on the sample, followed by shearing, and laser cutting has the greatest influence. The laser cutting effect is the worst, that is becacuse that there is no cooling in the cutting process, while EDM has coolant flow. The microstructure of the samples was observed by SEM. The results show that the macro measurement results are consistent with the micro observation results.

**References:**

&lt;!--[if supportFields]&gt;&lt;spanlang=EN-US&gt;&lt;span style='mso-element:field-begin'&gt;&lt;/span&gt;&lt;spanstyle='mso-spacerun:yes'&gt; &lt;/span&gt;ADDIN EN.REFLIST &lt;span style='mso-element:field-separator'&gt;&lt;/span&gt;&lt;/span&gt;&lt;![endif]--&gt;[1] H. Naumoski, "Investigation of the influence of different cutting procedures on the global and local magnetic properties of non-oriented electrical steel," Journal of Magnetism and Magnetic Materials, vol. 392, pp. 126-133, 2015. [2] Z. Xiaoyang, "Study on Measurement and Simulation Method of Magnetic Properties of Electrical Steel Sheets Considering Influence of Cutting Method," Shenyang University of Technology, 2016. [3] W. D. Chen J, Cheng S, "Modeling of Temperature Effects on Magnetic Property of Nonoriented Silicon Steel Lamination," IEEE Transactions on Magnetics, vol. 51, no. 11, pp. 1-4, 2015. &lt;!--[if supportFields]&gt;&lt;span lang=EN-US style='font-size:12.0pt;font-family:"Times New Roman",serif;mso-fareast-font-family:宋体;mso-bidi-theme-font:minor-bidi;mso-ansi-language:EN-US;mso-fareast-language:ZH-CN;mso-bidi-language:AR-SA'&gt;&lt;spanstyle='mso-element:field-end'&gt;&lt;/span&gt;&lt;/span&gt;&lt;![endif]--&gt;

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![](https://assets.underline.io/uploads/markdown_image/1/image/4befe7ab34ce100a7257c4a4004b0ad8.jpg)

![](https://assets.underline.io/uploads/markdown_image/1/image/a40e5a0db3f0e78ca813dc19a706c8a2.jpg)

 Effect of Cutting Methods on Magnetic Properties of Electrical Steel

Fe-Ga alloy is an advanced magnetostrictive material for actuators and sensors in terms of combining excellent mechanical and magnetostrictive properties. It is desirable to produce η(&lt;100&gt;//RD) textured Fe-Ga sheets due to the magnetostriction anisotropy with the largest coefficient along &lt;100&gt; direction and the serious eddy current loss in high-frequency use 1,2. The secondary recrystallization Goss ({110}&lt;001&gt;) texture and large magnetostriction coefficient were mainly achieved in Fe-Ga aaloy thin sheet with the thickness of 0.3~0.5mm, by the combination of micron-sized NbC particles and surface energy effect in high-temperature annealing 3,4, as well as the initial sharp &lt;100&gt;-oriented columnar grains by directional solidification 5. Since the eddy current loss is proportional to the square of the sheet thickness, the thinning of the sheet thickness can improve the energy conversion efficiency significantly. However, the secondary recrystallization of Goss texture significantly deteriorated in the Fe-Ga alloy sheet with a thickness of less than 0.3 mm due to the insufficient pinning force from the micron-sized NbC particles 6. The introduction of surface energy effect from H2S or S element, as well as the initial columnar crystals by directional solidification 7, cannot guarantee the complete secondary recrystallization in Fe-Ga alloy sheet with the thickness less than 0.3 mm. Qi et al 8 supposed the 0.10-mm Fe-Ga alloy ultra-thin sheets were prepared by rolling the 0.3 mm secondary recrystallized sheets, the strong cube texture and magnetostriction of 140 ppm were obtained after annealed for 2 h at 1200 °C under a flowing argon–hydrogen mixture. However, the preparation process is complicated and the incomplete secondary recrystallization restricts the magnetostriction coefficient. Therefore, it is necessary for the realization of complete secondary recrystallization in the Fe-Ga alloy sheet with a thickness of less than 0.3 mm. In the present study, the Fe-Ga ultra-thin sheets with the thickness of 0.10–0.25 mm were prepared via two-stage cold rolling method. The composite inhibitors composed of nanometer-sized sulfide and carbonitride are dispersedly precipitated during hot rolling and subsequent annealing, which provides effective pinning force for the normal growth of matrix grains. The gradually dissipated inhibition force during high-temperature annealing promotes a few exact Goss grains to grow abnormally by consuming primary recrystallized matrix grains textured by γ-fiber. Centimeter-sized secondary recrystallization Goss grains and large magnetostriction coefficient as high as 240 ppm is obtained in Fe-Ga alloy ultra-thin sheet without surface energy effect. The result indicates that complete secondary recrystallization can be realized based on the precise design and control of nano-sized composite precipitates and primary recrystallization microstructure/texture in Fe-Ga alloy ultra-thin sheet by conventional rolling and annealing methods.

**References:**

1. S. M. Na and A. B. Flatau, Scripta Mater. Vol., 66, p.307-310 (2012). 2. J. H. Li, X. X. Gao, J. Zhu, T. Xia and M. C. Zhang, Scripta Mater. Vol., 63, p.246-249 (2010). 3. C. Yuan, J. H. Li, X. Q. Bao and X. X. Gao, J Magn Magn Mater. Vol., 391, p.145-150 (2015). 4. S. M. Na, K. M. Atwater and A. B. Flatau, Scripta Mater. Vol., 100, p.1-4 (2015). 5. C. Yuan, J. H. Li, X. G. Bao and X. X. Gao, J Magn Magn Mater. Vol., 374, p.459-462 (2015). 6. S. M. Na and A. B. Flatau, AIP Adv. Vol., 7, p.056406 (2017). 7. J. H. Li, C. Yuan, Q. L. Qi, X. Q. Bao and X. X. Gao, Metals. Vol., 7, p.36 (2017). 8. Q. Qi, J. Li, X. Mu, Z. Ding, X. Bao and X. Gao, J Mater Sci. Vol., 55, p.2226–2238 (2019).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/326557a3327ec6e86a0e54601fff5ed3.jpg)
Orientation image maps of Fe-Ga ultra-thin sheets annealed at 950°C (a), 1000°C (b), 1100°C (c), and macrostructure after final annealing (d), as well as the variation of magnetostriction coefficients (λ//, λ⊥) with magnetic field annealed at 1000°C and 1100°C (e)

 The secondary recrystallization texture and magnetostriction in Fe-Ga alloy ultra-thin sheet

Grain-oriented (GO) steel sheets increasingly face novel applications, like the ones enforced by the emerging smart grid technologies and the related medium-to-high frequencies conversion devices and those connected with the development of highly efficient low-noise rotating machines [1-3]. Comprehensive knowledge of the material behavior, going beyond the standard assessment of its properties at power frequencies along the rolling direction (RD), is therefore appreciated in the practice of machine design. This requires both experiments made upon broader than usual range of frequencies [4] and the investigation of the complex anisotropic response of the GO sheets and the related 2D magnetization process [5]. Such a process is shown to consists in the balanced contributions by moving [001] directed 180° domain walls (dws) and 90° dw transitions towards the [100] and [010] phases. It is then observed that, with the applied field leaving the RD, not only the magnetic losses suffer a sharp increase, but also a strong geometry dependent effect, under both static and dynamic regimes. This descends from the interference of the demagnetizing field, invariably arising in Fe-Si single crystals, but for closed samples cut along the &lt;100&gt;, &lt;110&gt;, and &lt;111&gt; axes. Consequently, the popular approach to GO characterization using Epstein strips cut at different angles θ to RD can provide the intrinsic properties of the material only for θ = 0° (RD) and θ = 90° (TD). This is ill-suited to straightforward use in those cases, like the rotating machine cores, where directions other than RD and TD are involved. However, the experimental knowledge of the geometry-independent RD and TD magnetization curves can be usefully exploited, following the lines of the Néel’s phase theory, to reconstruct the same curves measured on Epstein strips cut along a generic angle θ [6]. The same applies to the DC hysteresis loops and their area, once the fundamental role of the demagnetizing field is recognized. One can also emulate an infinitely extended sheet, the ideal configuration leading to the intrinsic material response, by adopting an X-stacked sheet configuration, where flux continuity across the sheet edges can be achieved, resulting in a net magnetization variation both along the applied field (i.e., the strip edges) and orthogonal to it. Also in this case the DC properties become predictable by starting from the RD and TD magnetic properties [6]. Any other sample geometry will then be fully represented by or will be intermediate between the Epstein and the X-stack configurations. In this work, we provide a general approach to DC and AC (1 Hz - 200 Hz) magnetic hysteresis loop and losses in Epstein and X-stacked HGO sheets. The experiments are made using either the standard Epstein frame or a single sheet tester, where 120 mm wide 300 mm long X-stacked sheet samples are inserted in a laminated flux-closing yoke. A phenomenological predicting method is devised, which relies on the pre-emptive knowledge of the behavior of hysteresis loop and energy loss independently measured under sinusoidal induction in the RD- and TD-cut Epstein strips. We treat, in accordance with magneto-optical observations, the magnetization process unfolding for generic angle θ as the result, dictated by applied, demagnetizing, and eddy current fields, of two mechanisms: 1) Displacement of the 180° dws, the process pertaining to the RD cut sheets and strips; 2) Motion of 90° dws, converting the antiparallel slab-like [001] domains into [010] and [100] domains, and vice versa. The frames in Fig. 1 qualitatively show an example of domain structure in an X-stacked sheet at two instants of time along a major hysteresis loop (peak polarization Jp = 1.45 T). The antiparallel [001] domain structure observed at J = 0, largely disappears upon attaining J = 0.87 T. With this configuration, only a faint demagnetizing field is expected to arise at the sample edges and a net 2D magnetization variation is observed. By taking then, for a given Jp, the experimental energy loss and the hysteresis loops pertaining to the RD and TD samples, obtained at JpRD = Jp cos(θ) and JpTD = Jp sin(θ), and realizing that the fractional sample volume v180(t) occupied by the [001] phase is related to the same quantity v90(t) = √2 JTD /Js associated with the [010] and [100] domains as v180(t ) = 1 – v90(t), we calculate loss and loop for any given θ value. An example of so-obtained energy loss W(f) at 50 Hz across a range of Jp values on 0.288 mm thick high-permeability GO Epstein strips cut at different angles θ to RD is provided in Fig. 2. The experiments show that W(θ) peaks around θ = 75°. The predicting approach to W(θ) is based on the experimental knowledge of the energy loss measured at the same frequency in the RD and TD strips for JpRD = Jp cos(θ) and JpTD = Jp sin(θ). This associates with the calculation of the hysteresis loop. A more involved approach is required by the W(θ) prediction in the X-stacked samples, where the simplifying assumption of zero magnetization across the sheet width does not hold.

**References:**

[1] T. Belgrand, R. Lemaître, A. Benabou, J. Blaszkowski et C. Wang, AIP Advances 8, 047611 (2018). [2] G. Parent, R. Penin, J. P. Lecointe, J. F. Brudny, and T. Belgrand, IEEE Trans. Magn. 49, 1977 (2013). [3] S. Lopez, B. Cassoret, J. Brudny, L. Lefebvre, and J. Vincent, IEEE Trans. Magn. 45, 4161(2009). [4] M. A. Shamsuddin, F. Rojas, R. Cardenas, J. Pereda, M. Diaz, and R. Kennel, Energies 13, 2319 (2020). [5] W. Pluta, IEEE Trans. Magn. 52, 6300912 (2016). [6] F. Fiorillo, L. R. Dupré, C. Appino, and A. M. Rietto, IEEE Trans. Magn. 38, 1467 (2002).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/6b5cb05f0e84b50643df6d3ceb985c44.jpg)
Fig. 1. X-stacked GO sheet, cut at 75° with respect to the RD. The [001] antiparallel slab-like domain structure found in the demagnetized state (J = 0) is converted by 90° dw motion into symmetric [100] and [010] structures under increasing applied field (Jp = 1.45 T).

![](https://assets.underline.io/uploads/markdown_image/1/image/853d1d39b26e5ad4a2c3c620a36c162b.jpg)
Fig. 2. Energy loss at 50 Hz measured versus Jp and θ in Epstein strips. Solid lines and symbols: experiments. Dashed lines: prediction by a weighted combination of the 50 Hz RD and TD loss figures.

 The 2D Magnetization Process in HGO Fe-Si Sheets

Increasing the efficiency of electrical machines is a global objective. One way to contribute to its fulfilment is to reduce iron losses of soft magnetic materials within their magnetic circuits (stator, rotor, core, …). Improvement of soft magnetic material properties, especially Grain-Oriented Electrical Steels (GOES) based on iron and 3% silicon, is possible using surface laser treatments through the magnetic domain refinement technic and introduction of favorable stress [1-2]. These improvements are classically obtained for unidirectional excitation fields parallel to the rolling direction, the pronounced easy magnetic axis of GOES. However, in appliances such as rotating electrical machines, the flux lines are not always parallel to this preferred configuration (ex: in the teeth) and can also rotate (ex: in the yoke). In the literature, papers address the development of the expression and measurements of 2-D-magnetic properties [3], but the behavior of laser treated GOES under rotational and parallel to transverse direction fields remains slightly explored [4] and can be investigated further. Indeed, advanced pulsed laser technologies have been proven to improve magnetic [5-6] and magneto-mechanical [7] properties significantly under unidirectional field along the rolling direction. They include irradiation and scribing processes but also a new laser process called ablation with ultra-short laser pulses [5]. Therefore, in this research work, the effects of pulsed laser irradiation, scribing and ablation on GOES under transverse and rotational fields are presented. The objective is to optimize the vector behavior of GOES. The study focuses on power losses and B-H curves. The considered material is a laboratory grade of 0,28mm thickness with a CARLITE coating of 3-4µm thickness. The experimental plan includes 6 laser sets of parameters. The selection was made to enable comparisons between the types of laser processes (irradiation, scribing or ablation) and regarding the power loss reductions in the rolling direction. The main configurations were based on reference [5] in which the study was conducted with unidirectional excitation field parallel to the rolling direction on 150x150mm2 samples. In the present work, power losses have been measured in both rolling and transverse directions with a Single Sheet Tester (SST) and in rotational field with a Rotational Power loss Tester (RPT). An accuracy study will be detailed in the extended paper. So as to avoid any sample dispersion, measurements were carried out systematically on the same samples before and after laser treatment with the same setups (SST for 150x150mm2 samples and RPT for 60x60mm2 samples). The rotational power losses were obtained by considering the average of Clock-Wise (CW) and Counter-Clock-Wise (CCW) measurements and compared to the sum of unidirectional power losses. Thanks to loss model and B-H curve representation, while separating each direction but also the static and dynamic contributions, the impact of various laser parameters will be analyzed separately on scalar and vector magnetic properties as a function of working conditions. Additionally, all samples have been observed at each step (before and after laser treatment) with a Magneto-Optic Indicator Film technic (MOIF), which is very adapted for GOES’s magnetic domains observations. This will be illustrated in the final paper. Both macroscopic and microscopic analysis will be helpful to understand the reasons for opposite or simultaneous improvements in case of unidirectional or rotational fields. Concerning RPT measurements, rotational power losses reductions are reported for several laser sets of parameters (to be compared) and for several Polarization/Frequency couples as shown in figures 1 and 2. By using the pre-selected laser treatments, magnetic properties have been improved with iron losses reduction rate up to -12% at 50Hz&1T and up to -7% at 500Hz&0.5T in comparison with non-laser treated material for one of the selected sets of parameters. Further measurements will be carried out by the authors by improving the laser parameters (laser power and patterns), which might enhance the actual results and reduce even more significantly power losses. In the extended paper, measurement results (SST, RPT) will be presented with more details. Comparisons and discussions about the impact of considered laser treatments on GOES measured in both unidirectional and rotational fields on magnetic properties will be proposed as well. Acknowledgments: This work was supported by JEUMONT Electric and was carried out with the help of financing from the Association Nationale Recherche Technologie (ANRT). It was carried out in parallel with the ESSIAL project, which the authors would like to thank. The ESSIAL project is funding from the European Union’s Horizon 2020 research and innovation program under grant agreement No 766437. The authors would also like to thank the Aperam-Brazil company for having supplied the samples.

**References:**

[1] Y. Huang, W. Ming, M. Li et al., The International Journal of Advanced Manufacturing Technology, vol. 70, p. 1–9 (2014) [2] I. Petryshynets, F. Kováč, V. Puchý et al., AIP Advances, vol. 8, p. 047604 (2018) [3] Y. Maeda, H. Shimoji, T. Todaka and M. Enokizono, Journal of Magnetism and Magnetic Materials, vol. 320, p. e567–e570 (2008) [4] T. Kajiwara and M. Enokizono, IEEE Transactions on Magnetics, vol. 50, p. 1-4 (2014) [5] M. Nesser, O. Maloberti, J. Dupuy et al., Journal of Magnetism and Magnetic Materials, vol. 504, p. 166696 (2020) [6] E. Salloum, O. Maloberti, M. Nesser et al., Journal of Magnetism and Magnetic Materials, vol. 503, p. 166613 (2020) [7] E. Salloum, M. Nesser, O. Maloberti et al., conference MMM2019, Las Vegas, USA (2019)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/a2ce0b6b2f6b9c37b7f5c444b51e925b.jpg)
Fig.1 - Rotational power loss variation [%] after laser treatment at 50Hz for various polarizations (0.1T, 0.5T, 1T and 1.4T)

![](https://assets.underline.io/uploads/markdown_image/1/image/58aea06788288312a5bfb3fb25ed90e4.jpg)
Fig.2 - Rotational power loss variation [%] after laser treatment at 1T for various frequencies (50Hz, 150Hz, 500Hz and 1500Hz)

 Impact of Pulsed Laser Irradiation, Scribing and Ablation on 2-D Scalar and Vector Magnetic Losses and General Properties of Grain-Oriented Electrical Steels

Due to their texture, the magnetic characteristics of grain oriented (GO) steels depend on the angle between the applied magnetic field and the rolling direction RD. These GO steels exhibit advantageous magnetic characteristics, such as high permeability in the RD and relatively poor properties in the transverse direction TD. In order to obtain a better electromagnetic modeling of the electrical energy conversion systems using GO steels, the dissipative phenomena and the anisotropy must be reliably represented. One example is the case of the GO electrical steels used in stators of turbo-generators. In order to create an appropriate model, the associated magnetic losses have to be taken into account for any angle between the applied field and the RD. However, the problem of computing iron losses whatever the angle between the applied field and the RD has remained unsolved especially within the finite element method (FEM) framework. In this paper, a phenomenological iron loss model, recently proposed in [1], taking into account the anisotropy and nonlinearity of GO sheets has been successfully implemented in the L2EP FEM software baptized code_Carmel [2]. The model relies on the ‘loss separation’ approach. In this approach, which is based on the traditional decomposition proposed by Bertotti [3], the total losses are decomposed into three terms (hysteresis, classical and excess losses). The originality of the model concerns the first loss term i.e., the quasi-static losses dissipated per magnetization cycle, depending on the polarization level Jp and on the angle θ between the applied field and RD. The proposed quasi-static loss model considers the magnetization process and the evolution of the domain wall structure (dws) with the angle θ [4]. It is built from the two sets of W-Jp quasi-static experimental data obtained in the principal directions RD and TD and is based on the interactions between the main classes of magnetic domains, 180° and 90° dws, known as “Néel phases” [5] through a single analytical relationship involving the volumic fraction occupied, for any direction of the magnetic field, by these classes. This same principle is applied for the determination of the excess losses. At this stage, the corresponding input data are the excess losses measured in both RD and TD. Concerning the classical loss term, it simply takes the expression resulting from the classical eddy current theory which is independent of the angle θ. Indeed, the case of an infinite sheet of thickness e, with a linear constitutive law (µ independent of H) and conductivity σ, under a sinusoidal field of frequency f, allows an entirely analytical resolution of the magnetic diffusion equation. Finally, we get a description of the total losses W(f,Jp) as a function of the angle θ, without any adjustable parameter. The model has been validated on industrial sheets of EDF turbo-generators by means of simple study cases under periodic (cyclic) waveforms. Thus, we dispose of an efficient and reliable tool for modeling the magnetic losses of GO sheets within the FEM simulation environment. In order to validate the implementation of the developed model, a finite element analysis (FEA) has been performed using code_Carmel. A 3D linear magnetostatic simulation has been done using the magnetic vector potential formulation. A diagonal permeability tensor has been used to represent the anisotropic properties of the GO material. This tensor is formed assuming non-diagonal terms. The values of the tensor for the axes x and y were obtained from the characterization of the industrial GO steel C150-35S in the RD and TD, while for the axis z a very low relative permeability of ≈ 30 is assumed, as measured in [6]. Furthermore, the analyzed geometry corresponds to a GO sheet with dimensions of 10 cm x 10 cm and a thickness of 0.35 mm, surrounded by a coil along the x-axis. The iron loss calculation is performed in a section at the center of the sheet with an area of 1 cm x 1 cm. The elements at the center are subjected to a sinusoidal waveform of Jp of 0.5 T. Thirteen values of θ have been evaluated in this study (0°, 5°, 10°, 15°, 20°, 25°, 30°, 42°, 54.7°, 65°, 75°, 82.5° and 90°). They correspond to the angles characterized experimentally using the Epstein frame method with Epstein strips cut at different angles from the industrial GO steel sheet. Finally, from the distribution of H and J obtained in the FEA, the iron losses have been calculated in post-treatment. The main entries for the iron loss model were the hysteresis and excess losses in the RD and TD. In Table 1, an example comparing the static and dynamic losses obtained from the FEA against the experimental losses is given for some angles at 0.5 T. Further validations, analysis and discussions are envisaged in the extended version of the paper.

**References:**

[1] Carlo Appino, Enzo Ferrara, Fausto Fiorillo, et al., Journal of Magnetism and Magnetic Materials, vol. 500, (2019). [2] code_Carmel. [Online]. Available: https://code-carmel.univ-lille.fr/?q=fr/ node/15, accessed 2021 [3] G. Bertotti, IEEE Transactions on Magnetics, vol. 24, no. 1, pp. 621–630, (1988) [4] F. Fiorillo, L. R. Dupré, IEEE Trans. on Mag., vol. 38, no. 3, (2002) [5] L. Néel, J. Phys. Rad., vol. 5, pp. 241, (1944) [6] N. Hihat, J. P. Lecointe, O. Ninet, et al., International Journal of Applied Electromagnetics and Mechanics, vol. 46. pp. 349 - 354, (2014)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/0e5efc22699b2e5d1252a3fef02af167.png)
Table 1. Comparison of computed and experimental losses at 0.5 T

 Iron Loss Modeling of Anisotropic Soft Magnetic Steels in FEM Simulation Environment

Additive manufacturing (AM) via 3-D printing technologies have become a frontier in materials research, including its application in the development and recycling of permanent magnets1,2. The increasing criticality in availability of the rare earth minerals in magnets present a demanding challenge to the losses from conventional or subtractive manufacturing. AM allows for a minimal waste production and re-utilization of recovered rare earths in the magnet processing. In-situ alignment during 3D printing of magnetic materials has opened new horizons for manufacturing of complex permanent magnets3,4. The introduction of functionalized magnetic 3D printing (Fig. 1a) as an AM process yields several unique advantages including rapid prototyping of multi-directionally aligned magnetic systems, and reduced energy requirement due to the elimination of post-production alignment3. In this work, we have developed a mechanism to align magnetic materials using fused filament 3D printer with a novel magnetic field source architecture (Fig. 1b). The work addresses opportunities to integrate magnetic field into 3D printing process to enable printing, and alignment of anisotropic permanent magnets, without requiring further processing. The principle behind the analysis is rotation of magnetic particles in a molten matrix under low magnetic field. The alignment architecture was designed and optimized for weight and applied field. An in-depth finite element modeling and CAD was performed followed by prototyping of the magnetic field source (Fig. 1c-d). Several electromagnetic core models were developed using finite element modeling to obtain a field strength of &gt; 1 T with a magnetic material in the nozzle. The electromagnetic simulations were coupled with thermal analysis to obtain realistic transient thermal profiles. Design modifications were incorporated to improve the thermal performance of the models without compromising the electromagnetic output. The results predict models that could be operated over prolonged durations (&gt; 10 hours) at less than the rated temperature (80 oC) and at the rated current (5 A), while producing field strength of &gt; 1 T. As a future prospective, the model would be implemented as a prototype to be incorporated with the 3D printer. Magnetization vs. field measurements of extruded and printed Sm-Co (15 vol.%) in PLA and Nd-Fe-B/Sm-Fe-N composite (65 vol.%) in Nylon-12, prepared with and without magnetic field, confirmed alignment of magnetic particles (Fig. 2a-b). Field vs. alignment and printer temperature vs. alignment analyses were performed to understand the effects of process variables on the degree of alignment of the samples. Finally, a multiphysics model was developed, that couples fluid dynamics and electromagnetic interactions, to predict the degree of alignment (DoA) of a polymer bonded magnet printed under applied field. The model predicts the flow of magnetic particles in a viscous fluid through a nozzle under applied magnetic field. A particle-fluid interactive flow simulation was performed to model the flow regime of molten bonded magnet. The interactions between the drag, inertial, and magnetophoretic forces were analyzed to predict the particle trajectory. A torque balance was performed between the drag torque, magnetic torque on particles from the applied field, and torque from particle-particle interactions. The analysis was used to predict the degree of rotation of magnetic particles during printing. The simulated DoA of 0.68 compares well with the experimentally obtained for 65 vol.% NdFeB+SmFeN in Nylon-12. For 15 vol.% Sm-Co in PLA, the simulated DoA of 0.86 also compared well with 0.83 obtained experimentally (Fig.2c-d). This process can be used for any permanent magnet composition or hybrid magnets and presents a benchmark for application of this technology in industrial and commercial applications. This work is supported by the Critical Materials Institute (CMI), an Energy Innovation Hub funded by the U.S. Department of Energy (DOE), Office of Energy Efficiency and Renewable Energy, Advanced Manufacturing Office. Ames Laboratory is operated for the U.S. Department of Energy by Iowa State University of Science and Technology under Contract No. DE-AC02-07CH11358.

**References:**

1. C. Huber et al., Appl. Phys. Lett., 109, 162401 (2016) http://arxiv.org/abs/1605.07309. 2. L. Li, B. Post, V. Kunc, A. M. Elliott, and M. P. Paranthaman, Scr. Mater. (2017). 3. A. Sarkar et al., Addit. Manuf., 34, 101289 (2020) https://linkinghub.elsevier.com/retrieve/pii/S2214860420306618. 4. Y. Kim, H. Yuk, R. Zhao, S. A. Chester, and X. Zhao, Nature, 558, 274–279 (2018) http://www.nature.com/articles/s41586-018-0185-0.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/0f1fa7ee3edeb633b194223095ca2154.jpg)
Figure 1. a) Schematic of in-situ alignment mechanism for magnetic 3D printing, b) Representation of the prototyped 3D printer, c) Novel EM-core design with advanced cooling technology, and d) FEM simulation of c-core @ 5A current.

![](https://assets.underline.io/uploads/markdown_image/1/image/a0cec09893008e380b67cbc258cb5916.jpg)
Figure 2. a-b) MH loop for aligned magnetic samples, and c-d) Simulation validation.

 Advancing Additive Manufacturing of Bonded Permanent Magnets via in-situ Magnetic Field Alignment During 3D Printing

Today, the 3D-printing technology faces a new challenge - the Additive Manufacturing (AM) of functional magnetic materials, especially permanent magnets. AM should provide very good functional properties of the printed magnets, namely locally tailored remanent magnetization Mr and coercivity Hc. The highest magnetic performance is reached for Nd-Fe-B magnets based on the hard magnetic Nd2Fe14B phase, and on extrinsic properties (micro and nanostructure) mastered during the production. Ideal microstructure consists of magnetically decoupled grains of the hard magnetic RE2Fe14B (RE = Rare earth element) phase. The main challenge for Nd-Fe-B magnets, which arises during direct AM techniques, consists of an unsuitable microstructure appearing after re-melting of Nd-Fe-B precursor by the laser, followed by a fast solidification of the alloy. This usually leads to a very low coercivity in magnets made from commonly used commercial Nd-Fe-B alloy powders. Thus, one of the main obstacles in successful 3D printing of full-dense hard magnetic materials, is the engineering of a desired phase composition and the mastering of an appropriate microstructure. In the contrast to conventional Nd-Fe-B-based bulk alloys, a large HC has been found in annealed Pr-Fe-Cu-B compounds. The Pr2Fe14B phase is isostructural to Nd2Fe14B and the Pr-Fe-B and Nd-Fe-B ternary phase diagrams exhibit the same phases. By doping the system with Cu, additional phases can be formed in the sample and some of them affecting drastically the hard-magnetic properties of the magnet. Kajitani et al. [4] found a new grain-boundary phase with tetragonal structure (space group I4/mcm) and Pr6Fe13Cu stoichiometry in hot-pressed Pr17Fe76B5.5Cu1.5 alloys. Annealing at 500 °C leads to an increase in HC of 50 % compared to the hot-pressed state, which was associated to the formation of the antiferromagnetic grain-boundary phase. Further investigations by Marcondes et al. [5] proved the occurrence of the antiferromagnetic Pr3Fe13Cu phase and it was found that annealing at 500 °C improves the HC drastically due to the formation of the Pr3Fe13Cu phase at the expense of the Pr,Cu-rich grain-boundary phase. Recently, the occurrence of the RE6(Fe,Ga)14 within the grain boundary [6]–[8] of Nd-Fe-B permanent magnets was reported. The magnets show good hard- magnetic properties, but the role of the RE6Fe13(Cu,Ga) phase is not clear yet. Our studies explore the applicability of Pr-Fe-Cu-B based alloys for the use in Additive Manufacturing (AM) techniques to produce fully dense permanent magnets. Induction molten Pr17Fe76B5.5Cu1.5 alloys were annealed at high temperature (1000 °C for 5 hours), followed by a low temperature annealing (500 °C for 3 hours). The samples show the formation of soft magnetic Pr2Fe17 phase at high temperature accompanied by low coercivity. After a low temperature annealing, the mentioned Pr6Fe13Cu is forming (see Fig. 1) and an increase of coercivity can be observed. Based on these experiments, the composition has been optimized regarding the coercivity. For this, three pre-alloys (Pr6Fe13Cu, Pr2Fe14B, FeB) were prepared by induction melting and two compositional series by mixing the Pr6Fe13Cu with FeB or Pr2Fe14B were prepared by arc melting. The heat treatments were the same for all samples and the development of coercivity was characterized. The shape of the hysteresis is drastically changing within the sample series. High amount of the Pr6Fe13Cu phase leads to a shoulder and decreasing saturation magnetization, which indicates the non-ferromagnetic properties of the phase. The combination of Pr6Fe13Cu + FeB shows the largest coercvity of ~1250 kA/m which is comparable to a sintered magnet (without heavy Rare Earth elements). In both sample series, the formation of the reported Pr6Fe13Cu phase could be observed by SEM-EDX. The large coercivity in the bulk state makes this material interesting to use as stable precursors for AM of fully dense magnets. The material needs to be in powder form for techniques like SLM. Therefore, the hydrogen decrepitation (HD) process was successfully applied on high coercive samples. Afterwards, the material was suction casted which simulates the remelting and rapid solidification during the SLM process. After the aforementioned annealing, the coercivity is equal to the as-cast condition which demonstrates that the magnetic hardening mechanism is stable throughout the different preparation steps and the material will exhibit large coercivity after SLM. We thank the German Research Foundation (DFG) and the Collaborative Research Center/Transregio 270 HoMMage, Project ID No. 405553726, TRR 270, for making this work possible.

**References:**

[1] O. Gutfleisch, “Controlling the properties of high energy density permanent magnetic materials by different processing routes,” J. Phys. D. Appl. Phys., vol. 33, no. 17, pp. R157–R172, Sep. 2000. [2] T. G. Woodcock et al., “Understanding the microstructure and coercivity of high performance NdFeB-based magnets,” Scripta Materialia, vol. 67, no. 6. pp. 536–541, 2012. [3] G. Hrkac et al., “Impact of different Nd-rich crystal-phases on the coercivity of Nd-Fe-B grain ensembles,” Scr. Mater., vol. 70, no. 1, pp. 35–38, 2014. [4] T. Kajitani, K. Nagayama, and T. Umeda, “Microstructure of Cu-added Pr-Fe-B magnets: Crystallization of antiferromagnetic Pr6Fe13Cu in the boundary region,” J. Magn. Magn. Mater., vol. 117, no. 3, pp. 379–386, Dec. 1992. [5] P. V. P. Marcondes and R. N. Faria, “Microstructural studies on Pr-Fe-B-Cu magnets produced by upset forging of cast ingot,” Mater. Sci. Eng. A, vol. 272, no. 2, pp. 245–249, 1999. [6] K. Niitsu et al., “Magnetization measurements for grain boundary phases in Ga-doped Nd-Fe-B sintered magnet,” J. Alloys Compd., vol. 752, pp. 220–230, 2018. [7] G. Ding et al., “Coercivity enhancement in Dy-free sintered Nd-Fe-B magnets by effective structure optimization of grain boundaries,” J. Alloys Compd., vol. 735, pp. 795–801, 2018. [8] T. T. Sasaki et al., “Formation of non-ferromagnetic grain boundary phase in a Ga-doped Nd-rich Nd-Fe-B sintered magnet,” Scr. Mater., vol. 113, pp. 218–221, 2016.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/254a76c7d50f328f1a476ac33e1bda97.jpg)
Fig. 1: Microstructural evolution during the annealing treatment of Pr-Fe-Cu-B alloys. The sample is shown in as-cast condition (top) and annealing (bottom).

 High coercivity in bulk Pr-Fe-Cu-B alloys as a stable precursor for permanent magnets by Additive Manufacturing

Current production of dense Nd-Fe-B permanent magnets is restricted to the powder metallurgical way consisting of powder production, pressing and magnet field alignment, sintering and post processing. This method allows the production of simple shaped magnets with superior magnetic performance. Bonded magnets show a higher geometrical flexibility on the expense of magnetic performance due the mixing with non-magnetic polymer with a volume content of up to 50%. Additive Manufacturing by means of laser powder bed fusion (LPBF) offers the opportunity to overcome design restrictions of conventional production techniques and allows new strategies for the application-oriented product development. LPBF is characterized by a layer-by-layer local melting of powder on a built platform by a focused laser beam, which is scanned across the powder. New powder is applied on top of melted material afterwards and the melting process is continued. Local melting, rapid solidification, directional heat transfer and several re-melting and re-heating cycles control the microstructural development of metals during LPBF. For this reason, achievable microstructures and properties differ in general significantly from such ones, which can be obtained from conventional fabrication processes. We applied laser powder bed fusion to commercial Nd-lean MQPTM-S powder from Magnequench for the additive manufacturing of Nd-Fe-B bulk permanent magnets. Samples were manufactured on a “M2 cusing” machine from Concept Laser GmbH under protective Argon atmosphere and an Oxygen content below 0.2%. Nd-Fe-B behaves very different compared to established materials during LPBF and the resulting magnetic performance is mainly controlled by the energy input from the laser beam and depends on the processing parameter laser power PL, laser scan velocity vL and hatch distance hy. The latter one represents the offset of neighbouring laser scan lines [1]. The impact of processing parameter laser power and scan velocity on coercivity is shown in Figure 1 (a). It is obvious, that Hc is enhanced, when laser power increases or scan velocity decreases – in other words, if the energy input raises. A similar behaviour was observed for remanence Br and maximum energy product (BH)max. However, the enhancement of magnetic performance is limited by a material specific maximum allowed limit of processibility, at which the samples will be increasingly destroyed. For optimized processing parameter a coercivity of 920 kA/m (1.16 T), a remanence of 0.63 T and a maximum energy product of 63 kJ/m3 is obtained. Thereby, Br and (BH)max represent the highest reported values for additively manufactured permanent magnets so far. The demagnetization branch of the hysteresis curve for a magnet with optimized magnetic properties is shown in Figure 1 (b) and compared to the initial MQPTM-S powder and a polymer bonded magnet from the powder, which is produced by injection molding with a loading factor of 60 vol.%. The coercivity of the powder is 700 kA/m and is slightly reduced by injection molding. However, LPBF leads to an unexpected enhancement of coercivity and the value of the powder is exceeded by 30% without the addition of any rare earth containing eutectics or other post processes. The composition of the used material exhibits 20 wt.% of rare earth, which is equivalent to 8 at.% and is in the range of α-Fe/Nd2Fe14B nanocomposites [2]. These are typically prepared by melt spinning and show coercivities of 400 – 600 kA/m for comparable compositions [2,3], e.g rare earth content between 8 and 9 at.%. These values are clearly exceeded by our LPBF-processed Nd-Fe-B magnets. The link between LPBF-processing and enhanced magnetic performance is the microstructure of the magnets and the microstructural development is controlled – as described above – by several factors. In the case of LPBF-processed Nd-Fe-B magnets we obtained a unique fine-grained microstructure with a grain size between 60 and 600 nm. Our results demonstrate, that improved magnetic performance can be achieved by LPBF of Nd-Fe-B bulk permanent magnets with an optimized coercivity, which is unexpected high for the given Nd-lean composition. This offers new perspectives for the development of new manufacturing processes for the production of improved Nd-Fe-B permanent magnets.

**References:**

[1] F. Bittner, J. Thielsch and W.-G. Drossel: Progress in Additive Manufacturing, Vol. 5, p. 3-9 (2020) [2] G.C. Hadjipanayis: Journal of Magnetism and Magnetic Materials, Vol. 200, p. 379-391 (1999) [3] C. Wang, M. Yan and Q. Li: Materials Science and Engineering B, Vol. 150, p. 77-82 (2008)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/0747ee065f8f32f43ee73c42c10486b5.jpg)
Figure 1: Magnetic performance of additively manufactured Nd-Fe-B permanent magnets produced by laser powder bed fusion: (a) dependence of Hc on laser power and scan velocity and (b) demagnetization curve of additively manufactured magnet (“AM Magnet”, solid line), initial MQPTM-S powder (“powder”, dashed line) and injection molded magnet from MQPTM-S with loading of 60 vol.% (“injection molding”, dotted line).

 Unexpected coercivity enhancement > 1 T for Nd-Fe-B permanent magnets with 20 wt.% Nd produced by laser powder bed fusion

Excellent hard magnetic properties have been achieved in rare earth transition metal films (Nd-Fe-B, Sm-Co) produced by sputtering, pulsed laser deposition and molecular beam epitaxy. Such films can serve as model systems to probe intrinsic magnetic properties or to study the link between microstructure and extrinsic magnetic properties. They can also be used to produce miniaturized magnetic field sources for applications. In the latter case, the films must be patterned at a length scale comparable to the film thickness, because of the demagnetising field effect. The choice of film thickness will depend on the targeted application, nm-sized field sources require nm-thick films while µm-sized sources require µm-thick films. The film deposition rate becomes very relevant as the film thickness is increased, especially if the ultimate aim is for commercial production. What is more, deposition on large surface areas is needed to allow batch fabrication and integration of micro-magnets. We have previously reported on triode sputtering of 5 µm thick Nd-Fe-B films at deposition rates of up to 18 µm/h [1]. We then demonstrated how such films could be physically patterned at the micro-scale using a topographic approach based on pre-patterning of the substrate or post-deposition etching of the magnetic layer [2]. We also showed that localised heating by laser irradiation through a mask can be exploited to produce magnetisation reversal at the µm-scale in the so-called thermo-magnetic patterning approach [3]. Small arrays of micro-magnets produced using these approaches have already been used in a range of bio-medical experiments [4]. In this work we will demonstrate the potential that triode sputtering holds for batch fabrication of even thicker micro-magnets, for future integration into micro-scaled devices. Deep Reactive Ion Etching (DRIE) was used to anisotropically etch 55 µm deep trenches into Si substrates of diameter 100 mm and thickness 525 µm. Patterns consisting of squares with lateral size varying from 20 to 400 μm, and 4 mm long stripes with width varying from 20 to 400 μm, were used as test structures (Fig 1). Ta (100 nm) / Nd-Fe-B (50 µm) / Ta (100 nm) trilayers were deposited onto these topographically patterned substrates heated to a temperature of 500°C or 700°C. The NdFeB layer deposited at 500°C was amorphous, while that deposited at 700°C was crystallised. Post-deposition crystallisation of the former was achieved by annealing the full wafer in a Rapid Thermal Annealing (RTA) furnace, under secondary vacuum (10-6 mbar) at 600 °C, with a heating rate of 20°C/s and a dwell time of 10 minutes. Hereafter the samples will be referred to as 1-step processed (Tdep = 700°C) or 2-step processed (Tdep = 500°C + Tanneal = 600°C). Despite the relatively large thickness of the NdFeB layers, film peel-off does not occur. It should be noted that NdFeB films deposited onto non-patterned Si substrates start to peel off the substrate at a film thickness greater than 10 µm. The stress relief afforded by topographic patterning of the substrate depends on the dimensions of the pattern features and the film thickness, and some peel off was observed in 50 µm thick NdFeB layers deposited on square structures of width 500 µm prepared in another study. For the samples detailed here, some cracking of the deposit was observed in expansive regions far from etched features. SEM images of fractured cross-sections of stripes of two different width (200 µm and 30 µm) of each sample type are compared in Fig 2. Deposition on the sidewalls of the etched trenches leads to the deposition on the “hills” being wider than the underlying Si pillars. A shadowing effect results in the profiles of the deposit in the “valleys” being different to that on the neighbouring “hills”, and the effect becomes stronger as the feature size is reduced. The shadowing effect can be enhanced by increasing the depth of the trenches, as will be shown elsewhere. SEM imaging at higher resolution reveals that the NdFeB grains of the 1-step processed sample are columnar and have widths as large as 1 µm while those of the 2-step processed film are almost equiaxed with diameters in the range of 200 – 500 nm (not shown). Hysteresis loops measured on pieces taken from each sample are shown as insets to Fig 2. The 1-step sample has a coercivity of 1.2 T and an M0T/M7.5T of 0.79 while the corresponding values for the 2-step sample are 2.2 T and 0.85. The higher coercivity of the latter is due to its much finer microstructure and the 2-stage initial magnetization curve may be explained by easy domain wall motion in grains coarser than 300 nm, the single domain particle size of Nd2Fe14B, and domain wall pinning in finer grains. An in-house developed Scanning Hall Probe Microscope (SHPM) equipped with a Hall cross of active area size 4 x 4 µm2 [5], was used to measure the stray magnetic field produced at different heights above out-of-plane magnetized pieces of each sample. In this presentation we will provide full details of structural and magnetic characterization and will compare measured field profiles with simulated profiles. We will finish by discussing future prospects for the fabrication, integration and use of high-performance NdFeB micro-magnets.

**References:**

[1] N.M. Dempsey et al., “High performance hard magnetic NdFeB thick films for integration into Micro-Electro-Mechanical-Systems”, Appl. Phys. Lett. 90 092509 (2007) [2] A. Walther et al., “Micro-patterning of NdFeB and SmCo magnet films for integration into Micro-Electro-Mechanical-Systems”, J. Magn. Mag. Mat. 321 (2009) 590–594 [3] F. Dumas-Bouchiat et al., “Thermo-magnetically patterned micro-magnets”, Appl. Phys. Lett. 96, 102511 (2010). [4] N.M. Dempsey, “Magnetism and Biology” in " Handbook of Magnetic Materials", J.M.D. Coey et al. (eds), Springer, in-press [5] Gorky Shaw et al., “A Scanning Hall Probe Microscope for high resolution, large area, variable height Magnetic Field Imaging”, Rev Sci Inst 87 (2016) 113702

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/71d2432dd00b4772234c43bc8e6f423d.jpg)
Fig 1: Schematic of the pre-patterned substrates.

![](https://assets.underline.io/uploads/markdown_image/1/image/d17c28d143c41c74d8f313da1a848cdf.jpg)
Fig 2: SEM images of fractured cross-sections of stripes of width 200 µm (top) and 30 µm (bottom) prepared using 1-step (left) and 2-step processing. Insets: M(H) loops of pieces taken from 1-step (left) and 2-step processed samples.

 Batch fabrication of 50 µm thick anisotropic NdFeB micro-magnets

Revived interest on SmFe12-based compounds of the ThMn12 type as the permanent magnet materials has brought into focus the Sm1-xZrx(Fe,Co)11+δTi1-δ alloys remarkable by their high Curie temperature, large magnetization and strong magnetic anisotropy. However, attempts to develop practically useful coercivity in these alloys has encountered considerable difficulties. As most of the ongoing efforts rely on pulverization or on creation of a nanostructure in powders or melt-spun ribbons, relatively unexplored remain the massive cast alloys and "thick" ribbons melt-spun at a low solidification rate, even though they can be utilized as magnets without the need for costly compaction. In this study, alloys of Sm1.1-xZrx(Fe0.8Co0.2)11.3-yTi0.7By with 0 ≤ x ≤ 0.6 and 0 ≤ y ≤ 0.9 were subjected to structural, microscopic, thermomagnetic and isothermal magnetic characterization in samples immediately after arc-melting and after melt-spinning at a low wheel speed of 0.75–6.0 m/s. Slower solidification rates and Zr were found to generally promote the ThMn12-type crystal structure, whereas higher solidification rates and B tend to suppress the 1:12 structure in favor of the TbCu7-type one. When added simultaneously, B and Zr can alter the morphology of the 1:12 phase thus noticeably affecting the coercivity (Fig. 1). In particular, this alloying was found to generate in the arc-melted ingots a microstructure of isolated 1:12 crystallites 1–3 μm in size (Fig. 2); a better control of the minority phases, at least some of which are ferromagnetic, would be needed, however, to obtain in the alloy with such microstructure a Hc greater than 0.5–0.7 kOe. A moderately accelerated solidification can refine the 1:12 crystallites and increase the Hc; a Sm0.7Zr0.4(Fe,Co)10.8Ti0.7B0.5 alloy exhibited a Hc of 1.5 kOe and a maximum energy product (BH)max of 2.5 MGOe when it was melt-spun at 1.5 m/s into a 0.26-mm-thick ribbon. However, a more rapid solidification was found to suppress the 1:12 phase in favor of the magnetically softer 1:7 phase. After annealing at 800–900 °C, the alloys with the large B and Zr substitutions develop reasonably high Hc and (BH)max values even if melt-spun at the moderate wheel speed of 6 m/s. The respective values for the Sm0.7Zr0.4(Fe,Co)10.8Ti0.7B0.5 alloy were 4.1 kOe and 7.8 MGOe, whereas for a very Sm-lean Sm0.5Zr0.6(Fe,Co)10.6Ti0.7B0.7 alloy they were 3.6 kOe and 7.4 MGOe. In comparison, the parent Sm1.1(Fe,Co)11.3Ti0.7 alloy subjected to the same processing developed a Hc of only 0.4 kOe. The work was supported by the U.S. Department of Energy under Grant DE-FG02-90ER45413.

**References:**



**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/c2bc2978a39790ea8a94a07f85548392.jpg)
Fig. 1. Coercivity of bulk arc-melted alloys controlled through B and Zr substitutions.

![](https://assets.underline.io/uploads/markdown_image/1/image/3204a80ef4adcec70392f0d44b5abc93.jpg)
Fig. 2. Microstructure of arc-melted Sm0.5Zr0.6(Fe0.8Co0.2)10.8Ti0.7B0.5 alloy features fine well-separated crystallites of the 1:12 phase.

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