Microstructural and Texture Evolution of Strip Cast NdâFeâB Flake

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Microstructural and Texture Evolution of Strip Cast Nd-Fe-B Flake Hansheng Chen, Wanqiang Xu, Zhixiao Ye, Yin Yao, Jiangtao Qu, Fan Yun, Jacob Warner, Zhenxiang Cheng, Jianqiang Liu, Simon P. Ringer, Michael Ferry, and Rongkun Zheng Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.7b01213 • Publication Date (Web): 02 Nov 2017 Downloaded from http://pubs.acs.org on November 7, 2017

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Crystal Growth & Design

Microstructural and Texture Evolution of Strip Cast Nd-Fe-B Flake Hansheng Chen§,#,†, Wanqiang Xu‡,*, Zhixiao Yeǂ, Yin Yao||, Jiangtao Qu§,#,†, Fan Yun§,#,†, Jacob A Warner||, Zhenxiang Cheng⊥, Jianqiang Liu§, Simon P. Ringer#,\, Michael Ferry‡, and Rongkun Zheng§,#,†,*

§ School of Physics, The University of Sydney, NSW, 2006, Australia # Australian Institute for Nanoscale Science and Technology, The University of Sydney, Sydney, NSW, 2006, Australia † Australian Centre for Microscopy and Microanalysis, The University of Sydney, Sydney, NSW, 2006, Australia ‡ School of Materials Science and Engineering, The University of New South Wales, NSW, 2052, Australia ǂ Hengdian Group DMEGC Magnetics Co. Ltd, Zhejiang, 322118, China || Electron Microscope Unit, Mark Wainwright Analytical Centre, The University of New South Wales, New South Wales, 2052, Australia ⊥

Institute for Superconducting and Electronic Materials, University of Wollongong, New South

Wales 2522, Australia \ School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney, Sydney, NSW 2006, Australia

ACS Paragon Plus Environment

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KEYWORDS: Strip cast Nd-Fe-B flake; Microstructural and Texture evolution; X-ray diffraction; Electron backscatter diffraction; Magnetic force microscopy

ABSTRACT: In this work, the microstructure and texture evolution of the strip cast Nd-Fe-B flake has been systematically investigated by correlating multiple state-of-the-art characterization techniques. We found that: (i) besides the existence of random ultrafine equiaxed grains at the wheel side of the flake, elongated (001) textured grains were formed into V-shape zone between neighboring nucleation sites, which possibly resulted from the in-plane growth of low energy preferred growth direction of grains (a axis). Both ultrafine random equiaxed grains and elongated (001) textured grains are harmful to achieving high-performance Nd-Fe-B magnet, due to the inhomogeneous grain shape and non-uniform distribution of rare earth-rich phase within these grains or along the grain boundaries, which deteriorate the alignment of the Nd-Fe-B powders in the subsequent hydrogen decrepitation process and jet milling procedure. To overcome the issues mentioned above, two potential approaches are proposed, which are increasing the nucleation rate on the wheel side and homogenization of rare earth-rich phase within grains or along grain boundaries; (ii) columnar grains containing (Nd,Pr)2Fe14B lamellae with an average spacing of ~5 µm and discontinuous rare-earth rich phase were formed in the remaining part of the flake. Accordingly, a model in terms of the microstructure and texture evolution was proposed.


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1. INTRODUCTION The rare-earth (RE) elements have a critical role in the functional materials, among which NdFe-B permanent magnet is the most appealing, and has been broadly integrated into various applications such as the latest direct-drive generators for wind turbines by taking advantages of its extraordinary magnetic properties, such as large magnetic product energy.1-3 However, low intrinsic coercivity value (Hci), which is only ~20-30% of the theoretical value, severely restricts the applications of Nd-Fe-B magnets in high operating temperatures due to the massive thermal demagnetization.4,5 Although the incorporation of heavy rare-earth (HRE) elements (for example, dysprosium (Dy) and holmium (Ho)) enhances the magnetic properties of Nd-Fe-B magnets at higher temperatures, the scarcity and high price of HRE elements make these products non-sustainable and uncompetitive on the global market.6-8 Moreover, the magnetization behavior of these materials at high temperatures does not meet specific applications due to the antiparallel coupling between HRE elements and Fe in matrix phases.4,5

On the other hand, there is ample potential to improve the performance of traditional Nd-Fe-B magnets towards the theoretical value of the intrinsic coercivity.8-10 The fabrication process of sintered Nd-Fe-B magnets involves strip casting process, hydrogen decrepitation (HD), jet milling (JM), aligning pressing, sintering, and annealing. The desired microstructure is that matrix grains with a single magnetic domain size of ~200 nm are well magnetically separated by uniform RE-rich phases.2 The intrinsic coercivity is essentially determined by the average grain size as Hci=a-blnD,11 or Hci=a-blnD2,12 where D is the size of the matrix grains. At the same time, grain size distribution is also central to fabricate high-performance permanent magnet.13,14 Therefore, tuning the average grain size, the grain size distribution, and RE-rich phase


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distribution is the key to improve the properties of sintered Nd-Fe-B magnets. It is well known that the size of the matrix grains is mainly determined by the size of the Nd-Fe-B powders, even though limited grain growth occurs during the sintering and annealing.15,16 Therefore, controlling average powder size and the powder size distribution are central to achieving desired microstructure and corresponding magnetic properties.

One of the most paramount factors determining the average Nd-Fe-B powder size and size distribution is the microstructure and texture of the strip cast (SC) Nd-Fe-B flake. The strip casting process, as the first and most principle step in manufacturing high-performance rare earth Nd-Fe-B permanent magnets, plays a vital role in tuning the nucleation and formation of matrix grains (Nd2Fe14B) with a uniform distribution of RE-rich phases.17-21 It has been reported that undesired α-Fe precipitates in the free side and a layer of the fine equiaxed grains form near the wheel side when the thickness of SC Nd-Fe-B flake is more than 0.5 mm.21 In addition, fine equiaxed zone disappears and the formation of α-Fe is suppressed with the decreased thickness of the strip cast flake to ~0.2 mm.21 However, at the same time, lower saturated magnetization and coercivity of the final product would be achieved, due to the increased amounts of small particles after HD and JM addition and the amorphization of the Nd-Fe-B magnet.22,23 Despite considerable investigations into the SC Nd-Fe-B flakes, the underlying mechanism for the microstructural and texture evolution is not well understood yet, due to the incapability of traditional techniques on characterizing the microstructure and crystalline orientation of the matrix grains near the nucleation sites on the wheel side. It is commonly considered that the thickness of strip cast Nd-Fe-B flake in the range of 0.2 to 0.4 mm is appropriate.21 Therefore, we chose the strip cast flake with the thickness of ~0.3 mm as a representative.


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Here, we systematically investigate the SC Nd-Fe-B flake with multiple cutting-edge microanalytical techniques, including X-ray diffraction (XRD), high-resolution focused ion beam (HRFIB) assisted electron backscatter diffraction (EBSD) and energy dispersive spectroscopy (EDS), and high-resolution magnetic force microscopy (HRMFM). The experimental findings unveil that both random ultrafine equiaxed grains and elongated (001) textured grains are harmful to achieving high-performance Nd-Fe-B magnet due to the nonuniform grain shape and inhomogeneous distribution of rare earth-rich phase, which deteriorate the alignment of the Nd-Fe-B powders in the subsequent hydrogen decrepitation process and jet milling procedure. In addition, columnar grains containing (Nd,Pr)2Fe14B lamellae with an average spacing of ~5 µm and uniform rare-earth rich phase were formed in the rest of the flake. This work provides fresh insights into the mechanism of microstructure and texture evolution and optimization of the fabrication and processing of SC Nd-Fe-B flakes in a controllable manner.

2. EXPERIMENTAL SECTION 2.1. Materials Pure Fe, B-Fe alloy, Nb-Fe alloy, Dy, Nd, Al, Cu and other additives were used as raw material. The strip casting process for (Nd,Pr)-(Fe,Co)-B flakes and other similar processes are summarized in Figure S1 for comparison in the Supporting Information. The alloy was melted by induction furnace with high purity argon (Ar) atmosphere protection. The molten alloy was then poured onto a rotating copper roller. The casting rate was ~1.1 m/s. The SC Nd-Fe-B flake studied in this paper was ~300 µm thick. The chemical composition of the SC Nd-Fe-B flake measured by inductively coupled plasma (ICP) is presented in Table 1.


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2.2. Microstructure & texture evolution XRD was performed on the wheel side and the free surface of the SC Nd-Fe-B flake using Cu Kα radiation (λ = 1.5418 Å) for the sake of phase identification. The peak shape analysis was performed using the Topas 3.0 software, fitted using the simple Rietveld refinement.

The EBSD/EDS experiments were performed on the wheel side and cross-section of the SC NdFe-B flake. Samples were mechanically polished using 220-grit, 500-grit, and 1200-grit SiC papers successively and finally mechanically polished using standard colloidal silica suspension. Firstly, the microstructure and elemental distribution of the flake were investigated using a combination of secondary electron imaging and EDS in a Zeiss Ultra Plus field-emission scanning electron microscope (SEM) operating at 15 kV equipped with the Oxford Instruments Aztec integrated EBSD/EDS system. Secondly, the EBSD mapping was performed on the cross section and plain section using 20 kV with the map analyzed using the CHANNEL5 software.

For the FIB assisted EBSD/EDS, the wheel side was coarsely milled by ~2 µm under 30 kV, 2.5 nA and then was polished using 10 kV, 200 pA (in order to minimize the Ga damage) in a Zeiss Auriga FIB/SEM for obtaining accurate EBSD mapping.

2.3. Magnetic domain structure The HRMFM experiments were carried out using a Bruker Dimension Icon scanning probe microscopy (SPM). HRMFM is a secondary imaging mode, which uses a two-pass technique. The initial pass measures the height data, while the second pass at a prescribed ‘lift height’ (~40


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nm) above the sample surface thereby measures the longer-range magnetic forces between the probe and sample. The probes employed for imaging were MESP probes (Bruker) with spring constants ~2.8 N/m. The probes were tuned to their resonant frequencies and then driven slightly below resonances. The free-air oscillation amplitude was set between 30-50 nm and the scan rate was around 0.4 to 0.65 Hz depending on the requirement. The data were processed in the Gwyddion software to acquire two-dimensional (2D) atomic force microscopy (AFM) and MFM images.

3. RESULTS AND DISCUSSION 3.1. Phase identification During rapid solidification, the side of flakes directly contacted with the copper wheel is referred as the wheel side. The other side is referred as the free surface. Figure 1 shows various phases, including (Nd,Pr)2Fe14B, Nd-Pr, and Nd-Pr-O, were formed at the wheel side and free surface of untreated samples with NdPrFe14B phase having a tetragonal structure (P42/mnm space group) and lattice constants a=8.816 Å and c=12.230 Å, Nd0.5Pr0.5 phase having a hexagonal structure (P63/mmc space group) and lattice constants a=3.665 Å and c=11.818 Å, and NdPrO3 phase having a cubic structure and a lattice constant of a=10.986 Å, respectively.

3.2. Microstructural and texture evolution Besides the phase identification by XRD techniques, strong (004), (006), and (008) peaks shown in Figure 1(a) indicate that a pronounced (001) texture of the matrix grains was formed at the wheel side during the rapid solidification,17 in agreement with the work done by H. Q. Liu and etc.19 In addition, several less strong (220), (330), and (440) peaks suggest the existence of a


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relatively weak (110) texture. However, the (001) (major) or (110) (minor) peaks gradually disappeared and the (410) peak was gradually enhanced from the wheel side towards the free surface (Figures 1(a) and (d)).

Figure 2 shows the microstructural and texture evolution of the SC Nd-Fe-B flake from the wheel side (upper) to the free surface (bottom), parallel to the direction of rapid solidification. In Figure 2(a), dark and bright regions correspond to the (Nd,Pr)2Fe14B matrix phase and RE-rich phase, respectively (Figure S2 shows the elemental distributions of Nd, Pr, and Fe of the same area). Figure 2(b) shows the cleaned inverse pole figure (IPF)-Y map of the (a) with black lines highlighting the high angle (>10°) grain boundaries. It should be noted that RE-rich phases exist as not only intergranular phases but also intragranular platelet. In addition, the crystalline orientations of matrix grains forming on the wheel side were not random, but the c axis of matrix grains was preferentially parallel along the normal direction of the surface, in agreement with the observation in the melt spun Nd-Fe-B magnets.24 However, this preferred c axis texture disappeared towards the free surface gradually with the occurrence of large columnar dendritic grains, in agreement with the observation by Takashi Hattori and etc.25 Furthermore, the black arrows mark out the growth direction of the columnar dendritic grains which is not parallel but deviated ~14-25° to the normal direction of the wheel side or free surface, suggesting that not only the intrinsic easy growth direction of the Nd2Fe14B but also other factors, such as the presence of the fluid flow, play an important role in determining the growth direction of the columnar grains, which has been observed in the steel fabricated by the twin-roll casting technique.26 To sum up, the SC Nd-Fe-B flake can be classified into two zones according to ascast microstructural and texture evolution: (i) the (001) (major) and (110) (minor) textured fine


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matrix grains near the wheel side; (ii) the coarse columnar dendritic matrix grains in the remaining part of the SC flake.

For a detailed understanding of the local microstructural and texture formation of the SC Nd-FeB flake at the wheel side, high-resolution EBSD experiments were conducted. Figure 3(a) shows the cleaned IPF-Y map of the SC Nd-Fe-B flake near the wheel side (cross-section). High angle (>10°) grain boundaries are marked by black lines. Figure 3(b) shows the corresponding (100), (110) and (001) pole figures, indicating a pronounced (001) texture of the matrix grains was formed at the wheel side, in agreement with our XRD results. It should be mentioned that most of the matrix grains in the V-shape zone were elongated and not normal to the wheel side, indicating the existence of the temperature gradient between the nucleation sites. In addition, besides strong (001) (major) and (110) (minor) textured grains marked by yellow dotted lines, ultrafine random equiaxed grains with size of ~1-5 µm marked by violet dotted lines were formed at the wheel side as well, which was observed in the SC Nd-Fe-B flake with the thickness over ~0.5 mm.21

Figure 3(c) shows magnetic domain structure of the SC Nd-Fe-B flake (cross-section). The phase shift ( δφ ) shown in Figure 3(d) is proportional to the force gradient ( δ F ′t ), which can be formulated as: δφ = Q / k × δ F ′ t , where Q and k are a quality factor and cantilever spring constant.27 In Figure 3(c), small random equiaxed grains marked by violet dotted lines were found with smaller domain size with different phase contrast. In addition, V-shape zone marked by yellow dotted lines was also observed at the wheel side at the demagnetized state. The black line marks through four stripe domains and five magnetic domain walls normal to the wheel side,


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indicating the c axis of those matrix grains was in plane and perpendicular to the wheel side in the V-shaped (001) zone because c axis is the easy magnetization axis for uniaxial Nd2Fe14B. According to the average domain size calculated from D1, D2, D3, and D4, the arrow penetrates through the magnetic domains with ~1.5 µm wide and throughout the area, indicating that magnetic domain walls penetrate across the RE-rich grain boundary phases.

Figure 4 shows secondary electron image, backscatter electron image, cleaned IPF-Y map, and corresponding EDS maps of another region near the wheel side (cross-section). First of all, strong (001) and (110) textured elongated matrix grains with inhomogeneous grain shape and grain size ranging from ~1 µm to ~10 µm were formed near the wheel side, which are harmful to achieving the high performance of the sintered Nd-Fe-B magnet. Secondly, the etched regions in the Figure 4(a), the brighter regions in Figure 4(b), and the EDS maps of the Fe, Nd, and Pr indicate some large RE-rich pockets were formed inside the grains and along the GBs, suggesting the non-uniform distribution the intragranular and intergranular RE-rich phases. This also plays a detrimental role in the efficiency of the following fabrication procedures, such as HD and JM. Besides, Cu was found to segregate in the RE-rich phase, as shown in the white arrows in Figure 4(g).

Figure 5 shows the plain-view EBSD results of the SC Nd-Fe-B flake in the columnar zone. Figure 5(a) shows the BSE image of the sample with brighter regions corresponding to the RErich phases. Figure 5(b) shows an enlarged BSE image of the white rectangles in Figure 5(a), indicating (Nd,Pr)2Fe14B dendritic lamellae with an average spacing of ~5 µm were formed in the columnar zone. Figure 5(c) shows the IPF-Z map of the same area in Figure 5(a). High angle


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(>10°) boundaries are marked by black lines. The number of the detected grains and the grain size are 589, and ~16.0 µm, respectively, which confirms the existence of the intragranular RErich phase. Figure 5(d) shows that the strongest intensity point in the (100) pole figure deviates to the center at ~14°, indicating that 〈100〉 direction was not parallel to the normal direction of the free surface, in agreement with our previous EBSD results in Figure 2(b).

3.3. Mechanism of nucleation and growth of the flake Here, we proposed the underlying mechanism of the microstructural and texture evolution of SC Nd-Fe-B flake described above.

3.3.1. Nucleation at the wheel side Figure 6 shows the morphology, microstructure and the composition distribution of element Fe, Nd, Pr, O, Dy, Ho, and Al near the nucleation sites at the wheel side. The nucleation sites are indicated by arrows and the distance between the nucleation sites is tens of µm (Figure 6(a)), in agreement with the results shown in Figure 2(b). The microstructure and EDS results show the petal-like (Nd,Pr)2Fe14B phase and Nd-Pr-O phase. The Nd-Pr-O phases could be the precipitates mainly forming in the boundaries of the (Nd,Pr)2Fe14B phase in this dual phase eutectic solidification structure during rapid solidification, similar to the Nd-Fe-B alloys.17 The coexistence of oxygen with Nd-Pr rich phase indicates that the RE-rich precipitate was oxidized when the melt was exposed to air during strip casting process, in agreement with our XRD results. The other elements such as Dy, Ho, and Al were relatively uniform solid solute into both matrix grains and the RE-rich phase.


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3.3.2. In-plane grain growth at the wheel side Figure 7 shows the local microstructure and texture of the slightly polished wheel side by FIB. In Figure 7(a), red and yellow regions are matrix phases and Nd-Pr-O phases, respectively. It should be noted that white arrows highlight the nucleation sites (Figure 7(b)). Smaller (100) matrix grains and relatively larger (001) matrix grains were both formed at the nucleation sites. At the nucleation process, the size of (001) and (100) matrix grains should be nearly identical. However, the favorite growth direction of Nd2Fe14B is a axis (〈100〉 direction).28 Therefore, the reason for why (001) matrix grain is larger than (100) grain is that (001) matrix grain grows fast along 〈100〉 direction at the wheel side, confirmed by the 〈100〉 pole figure of the (001) grains at the nucleation sites which exhibit that the 〈100〉 direction of each grain was parallel to the elongated direction of the (001) grains.

Figure 8 shows another local microstructure and texture analysis of the slightly polished wheel side by FIB, similar to Figure 7. Figure 8(a) shows the microstructure of the SC Nd-Fe-B flake. The gray and bright regions correspond to matrix grains and RE-rich phases, respectively. The majority of the Nd-Pr rich phase(s) at the wheel side may be intergranular lamella or equiaxed precipitate in the boundaries of (Nd,Pr)2Fe14B grains in this dual phase eutectic solidification structure while the minority of the Nd and Pr precipitates were formed within grains. The local EBSD analysis was performed on the dashed rectangles in Figure 8(a). Matrix grains and Nd-PrO phases are marked by red and yellow color, as shown in Figure 8(b). Figure 8(c) shows the crystalline orientation of each matrix grain near the near the nucleation sites. Similar to pole figures in Figure 7, the 〈100〉 pole figure of the G1 and G2 are shown in Figure 8(d), exhibiting that the 〈100〉 direction of each grain was parallel to the elongated directions (marked by white


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arrows in Figure 8(a)) of the grains, confirming the in-plane grain growth on the wheel side, which is possibly due to existence of the temperature gradient between the nucleation sites.

3.3.3. Out-of-plane grain growth from the wheel side to the free surface Although the strong (001) texture was formed on the wheel side, the crystals preferred to grow with a favorite crystal orientation anti-parallel to the heat flow direction as further solidification goes on.29 The selection of the favorite orientation mainly depends on the crystal structure of solidified phase and can be affected by the temperature gradient distribution and the flow movement of melted alloy during solidification. The dendritic grains that have aligned well with the favorite orientation, gain growth advantages over other oriented ones which will be eliminated eventually, to form a columnar zone. Figure 9 shows the cross-section of the SC NdFe-B flake. Even though a relatively uniform distribution of RE-rich in the columnar zone compared with the (001) textured zone, there are still some discontinuous RE-rich phases marked out by black arrows. Figure 10 shows the dendritic structure and elemental distribution of Nd, Pr, Ho, Dy, Fe, O on the free surface. Columnar grains containing (Nd,Pr)2Fe14B lamellae with an average spacing of ~5 µm and continuous rare-earth rich phase were formed at the free surface. It should be noted that Dy and Ho prefer to enter the 2:14:1 phases rather than the RE-rich phases in the free surface. X. B. Liu30 found that the substitution energy of Dy in 2:14:1 phases is negative, while the substitution energy of Dy in NdO has a large positive value, indicating that Dy prefers to enter 2:14:1 phases rather than NdO phases. However, Dy and Ho are still found in some of the RE-rich phases, as shown in Error! Reference source not found.(e-f).

3.3.4. Summary


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Figure 11 summarizes the mechanism for the microstructural and texture evolution as discussed before. First of all, nucleation occurs at the agents, such as impurities, marked by gray circles, as shown in Figure 11(a). Secondly, matrix grains grow with their 〈100〉 direction in-plane due to the heat gradient between the nucleation sites, resulting in the formation of elongated (001) textured grains between the nucleation sites, as shown in Figures 11(b-c). In addition, ultrafine random equiaxed matrix grains are formed with high nucleation density simultaneously (not shown in Figure 11). Meantime, some matrix grains grow with their 〈100〉 direction out-of-plane from the wheel side to the free surface, therefore the columnar grains grow with RE-rich phases existing as not only intergranular phases, but also intragranular platelets.

3.4. The effects of the microstructural and texture evolution on the magnetic properties of the final Nd-Fe-B magnet Firstly, solidification starts with heterogeneous nucleation at the surface of the mold wall and liquid metal. Subsequently, randomly oriented equiaxed grains form from the nucleation sites and the dendritic grains with preferred crystallographic directions dominate at the rest of the flake to form the columnar zone.31 However, strong (001) textured grains were observed on the wheel side of the flake and gradually disappeared from the wheel side to the free surface previously by X-ray diffraction in the Nd-Fe-B SC alloys.25 These (001) textured grains were not consistent with the direction of the heat flow direction, which reduced the possibilities of the nucleation of (100) grains in the equiaxed zone and degraded the alignment of Nd-Fe-B powders in the JM and HD and therefore the coercivity and remanence of final magnets.21


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Secondly, both ultrafine random equiaxed grains and elongated (001) textured grains are formed near the wheel side with inhomogeneous grain shape and nonuniform distribution of rare earthrich phase, as shown in Figure 4. Ultrafine grains are easily incorporated into the neighboring larger grains in the sintering and annealing process, resulting in the larger average grain size and non-uniform grain size distribution.14 In addition, Nd-rich phases are not well distributed inside the grains and along the grain boundaries and segregated into large pockets, as shown in Figure 4, which is harmful to the HD process. If large dispersive RE-rich pockets can be converted into the fine continuous RE-rich phases with the smaller columnar grain size of the SC Nd-Fe-B flake, smaller average powder size and uniform powder size distribution can be achieved by HD and JM.

3.5. Potential optimization approaches The core ideas are to increase the nucleation rate on the wheel side, which can reduce the possibilities of the nucleation of (100) grains in the equiaxed zone, and homogenize the distribution of RE-rich intragranular platelets and intergranular phases. For increasing the nucleation rate on the whee side, there are two approaches reported by Michael Miller, which are (i) increasing the cooling rate and (ii) inoculation, respectively.29

(i) For the homogeneous nucleation, the relationship between the nucleation rate and undercooling can be expressed as follows:29 •

N hom = k1

3  16π  γ SL Tm2 DL  exp  −  3 L2 ( ∆T )2 kT  DLM f  

(1)


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where k1 depends on the critical nucleus size and surface energy, DL and DLM are the diffusivities of the liquid at T and Tm, respectively, Tm is the equilibrium melting temperature of the liquid, γSL is the surface energy of the particle/liquid interface, Lf is the latent heat of the fusion per unit volume, and ∆T represents the undercooling for the nucleation. From the equation (1), the nucleation rate is proportional to the undercooling for the nucleation.29 One practical way to increase the undercooling for the nucleation is to increase the cooling rate.32 It has been reported that there is a linear increase in ln ∆T with ln q, where q represents the fixed rate of the solution cooling.33 Therefore, increasing the cooling rate can increase the nucleation rate during the solidification process. However, two potential issues may occur when increasing the cooling rate. The first issue is that extremely high cooling rate may produce submicron matrix grains, which are difficult to mill into the single crystal in the following hydrogen decrepitation (HD) process.34 The second and most serious issue is that metastable and amorphous phases may be synthesized at high undercooling,35 which is harmful to achieving high-performance permanent magnet. Shumpei Ozawa and etc. have reported that the decrease in intrinsic coercivity in the high cooling rate is mainly due to the amorphization of the Nd-Fe-B materials.36 Therefore, the different cooling rate should be selected based on the compositions of the alloys. (ii) For the inoculation, there are many successful cases in increasing the nucleation rate and therefore refining the grain size, such as Fe in copper and TiN in ferrite iron.29 However, fewer studies have been reported in refining the grain size of Nd-Fe-B magnets by inoculation.

For the homogenization of the distribution of RE-rich intragranular platelets and intergranular phases, grain boundary diffusion process (GBDP) should be considered. Grain boundary diffusion process (GBDP) has been regarded as one of the most effective approaches to enhance


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the coercivity of final sintered Nd-Fe-B magnet. Therefore, it may be useful to utilize GBDP for homogenization of the RE-rich phase in the SC Nd-Fe-B flake.

4. CONCLUSIONS In conclusion, HRFIB-assisted EBSD in conjunction with XRD, EDS, and HRMFM offers fresh insights into the microstructural and texture evolution of the SC Nd-Fe-B flake. Besides the existence of ultrafine random equiaxed matrix grains at the wheel side of the flake, elongated (001) textured matrix grains were found forming into V-shape zones near the wheel side of the flake, mainly due to the preferential in-plane growth along 〈100〉. Both ultrafine random equiaxed and elongated (001) textured matrix grains deteriorate the alignment of the Nd-Fe-B powders in the subsequent hydrogen decrepitation process and jet milling procedure due to the inhomogeneous grain shape and non-uniform distribution of rare earth-rich phase. In the remaining part of the SC flake, columnar grains containing (Nd,Pr)2Fe14B lamellae with an average spacing of ~5 µm were formed. Our work proposed a new mechanism of microstructural and texture evolution of SC Nd-Fe-B flake and could have significant implications for manufacturing controllable high-performance SC Nd-Fe-B flakes.


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FIGURES

Figure 1. Phase identification on the wheel side and free surface. The XRD spectra of the wheel side (a)(b)(c) and the free surface (d) of the SC Nd-Fe-B flake.


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Figure 2. Microstructural and texture evolution (cross-section). The SEM-EBSD analysis on the whole cross-section of the SC Nd-Fe-B flake. (a) Back-scattered electron image. (b) Cleaned inverse pole figure (IPF)-Y map of the sample. The color of the grains represents the orientation. The color code is shown in the standard triangle (inset). (c) (100), (001) and (110) pole figures.


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Figure 3. Microstructural and texture formation at the wheel side (cross-section). The SEMEBSD and HRMFM analysis on the local regions at the wheel side (a) Cleaned IPF-Y map. The color of each grain represents the orientation. The color code is shown in the standard triangle (inset). (b) (100), (001) and (110) pole figures. (c) Magnetic domain structure of the SC Nd-Fe-B flake (cross-section) (d) Phase shift information extracted from the black line in (c).


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Figure 4. The SEM-EBSD analysis on the grains near the wheel side (cross-section). (a) Secondary electron (SE) image. (b) Backscattered electron (BSE) image. (c) Cleaned IPF-Y map for the sample. EDS maps of (d) Fe, (e) Nd, (f) Pr, and (g) Cu measured from the same region in (a).


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Figure 5: SEM-EBSD analysis of the columnar zone (plain view) (a) Backscatter electron image. (b) Enlarged backscatter electron image. (c) Cleaned IPF-Z map of the sample. High angle boundaries (>10°) are marked by black lines. The color of the grains represents the orientation. The color code is shown in the standard triangle (inset). (d) (100) pole figure.


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Figure 6. Characteristics of the analyzed wheel side. (a)Low- and (b) high-magnification secondary electron SEM images of the wheel side. EDS maps of (c) Fe, (d) Nd, (e) Pr, (f) O, (g) Dy, (h) Ho and (i) Al measured from the black rectangular region in (b).


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Figure 7: FIB assisted EBSD analysis at the slightly polished wheel side (plain view). (a) phase map (b) Cleaned IPF-Z map of the sample. The color of the grains represents its orientation. The color code is shown in the standard triangle (inset). The pole figures are 〈100〉 and (001) pole figures of grains marked by black lines.


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Figure 8: Another FIB assisted EBSD analysis at the slightly polished wheel side (plain view). (a) Backscattered electron image of the wheel side after slightly polishing by FIB. (b) Cleaned phase map. Phase and grain boundaries are marked by green and black lines, respectively. (c) Cleaned IPF-Z map of the sample. The color of the grains represents its orientation. The color code is shown in the standard triangle (inset). (d) The corresponding 〈100〉 and 〈001〉 pole figures of G1 and G2 in (c).


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Figure 9. Characteristics of the cross-section of the SC Nd-Fe-B flake. (a) Backscatter electron image. EDS maps of (b) Fe, (c) Nd, and (d) Pr measured from (a).


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Figure 10. Characteristics of the analyzed free surface. (a)Low- and (b) high-magnification secondary electron SEM images of the free surface. EDS maps of (c) Nd, (d) Pr, (e) Ho, (f) Dy, (g) Fe, and (h) O measured from (b).


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Figure 11. The underlying mechanism for the abnormal texture formation. (a) Nucleation at the wheel side, (b) In-plane growth of low energy preferred growth direction (〈100〉 direction) of grains, (c) Out-of-plane growth of low energy preferred growth direction (〈100〉 direction) of grains, (d) the formation of the coarse columnar zone.


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TABLES

Table 1. The chemical composition of the SC Nd-Fe-B flake measured by inductively coupled plasma (ICP).

wt.%

Nd

Pr

Dy

Ho

Gd

Co

Nb

Cu

Al

Zr

Fe

B

21.06

6.67

2.32

0.60

0.06

1.14

0.32

0.23

0.36

0.05

66.29

0.90


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ASSOCIATED CONTENT Supporting information The Supporting Information is available free of charge on the ACS Publications website. SI#1: Cast strip fabrication techniques. SI#2: Elemental distribution of the strip cast Nd-Fe-B flake.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] *E-mail: [email protected]

ORCID Rongkun Zheng: 0000-0002-7860-2023

AUTHOR CONTRIBUTIONS The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

FUNDING SOURCES This research was supported by Australian Research Council (DP150100018).

NOTES The authors declare no competing financial interest


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ACKNOWLEDGMENT The authors would like to thank Dr. Hongwei Liu, Mr. Steve Moody, Dr. Patrick Trimby, Mr. Adam Sikorski, Dr. Matthew Foley for their technical support. The authors would like to thank Australian Microscopy and Microanalysis Research Facility (AMMRF), Australian Centre for Microscopy and Microanalysis (ACMM), the University of Sydney for their support as well.

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For Table of Contents Use Only

Microstructural and Texture Evolution of Strip Cast Nd-Fe-B Flake Hansheng Chen§,#,†, Wanqiang Xu‡,*, Zhixiao Yeǂ, Yin Yao||, Jiangtao Qu§,#,†, Fan Yun§,#,†, Jacob A Warner||, Zhenxiang Cheng⊥, Jianqiang Liu§, Simon P. Ringer#,\, Michael Ferry‡, and Rongkun Zheng§,#,†,* We report a comprehensive study of microstructural and texture evolution of strip cast Nd-Fe-B flake by correlating high-resolution focused ion beam (HRFIB) assisted electron backscatter diffraction (EBSD), X-ray diffraction (XRD) and high-resolution magnetic force microscopy (HRMFM).

SYNOPSIS


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Microstructural and Texture Evolution of Strip Cast NdâFeâB Flake

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Microstructural and Texture Evolution of Strip Cast NdâFeâB Flake