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Attractive-domain-wall-pinning controlled Sm-Co magnets overcome the coercivity-remanence trade-off
Accepted Manuscript Attractive-domain-wall-pinning controlled Sm-Co magnets overcome the coercivityremanence trade-off Hansheng Chen, Yunqiao Wang, Yin Yao, Jiangtao Qu, Fan Yun, Yuqing Li, Simon P. Ringer, Ming Yue, Rongkun Zheng PII:
S1359-6454(18)30846-2
DOI:
https://doi.org/10.1016/j.actamat.2018.10.046
Reference:
AM 14924
To appear in:
Acta Materialia
Received Date: 12 July 2018 Revised Date:
23 October 2018
Accepted Date: 24 October 2018
Please cite this article as: H. Chen, Y. Wang, Y. Yao, J. Qu, F. Yun, Y. Li, S.P. Ringer, M. Yue, R. Zheng, Attractive-domain-wall-pinning controlled Sm-Co magnets overcome the coercivity-remanence trade-off, Acta Materialia (2018), doi: https://doi.org/10.1016/j.actamat.2018.10.046. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Attractive-domain-wall-pinning controlled Sm-Co magnets overcome the coercivity-remanence trade-off
a
b
, Simon P. Ringer c, f, Ming Yue d, *, and Rongkun Zheng a, b, c, *
School of Physics, The University of Sydney, NSW, 2006, Australia
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d
The University of Sydney Nano Institute, The University of Sydney, Sydney, NSW
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2006, Australia c
Australian Centre for Microscopy and Microanalysis, The University of Sydney, NSW, 2006,
Australia d
College of Materials Science and Engineering, Beijing University of Technology, Beijing
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100124, China e
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Hansheng Chen a, b, c, #, Yunqiao Wang d, #, Yin Yao e, Jiangtao Qu a, b, c, Fan Yun a, b, c, Yuqing Li
Electron Microscope Unit, Mark Wainwright Analytical Centre, The University of New South
Wales, New South Wales, 2052, Australia
School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney,
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f
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Sydney, NSW 2006, Australia
#
These authors contributed equally to this work
*
Corresponding author. Email: [email protected] (R.Z.); [email protected]
(M.Y.)
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Abstract Traditional approaches for increasing the intrinsic coercivity of magnets typically come at the expense of remanence, a dilemma known as intrinsic coercivity-remanence trade-off, leading to a
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substantial reduction of maximum energy product. New metallurgical processing might offer the possibility of overcoming this trade-off. Here, we achieve a combination of an intrinsic coercivity of 26.9 kOe, a remanence of 11.2 kG, and a maximum energy product up to 26.6
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MGOe, which surpasses most of conventional Sm-Co based permanent magnets, by manipulating the gradient of domain wall energy landscape of constituent phases to realize the
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attractive domain wall pinning in Sm(Co,Fe,Cu,Zr)z permanent magnets. Using powerful atomicscale analysis technique known as atom probe tomography and micromagnetic simulations, we reveal that an enlarged attractive domain wall pinning strength results in the substantial coercivity enhancement with little sacrifice of remanence and maximum energy product in the
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Cu-particle-alloyed magnet. These results provide atomic-level insights into the coercivity mechanism of rare earth permanent magnets, with the methodology offering exciting possibilities for quantitative analyses and prediction between compositions and magnetic properties of other
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magnetic materials.
Keywords Sm-Co
magnets,
attractive-domain-wall-pinning,
micromagnetic simulations
2
coercivity,
atom
probe
tomography,
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1. Introduction Sm(Co,M)z (M=Cu, Fe and/or Zr) alloys are widely used as permanent magnets in satellite communications, aerospace, and defence technologies because of their capacity to retaining high
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intrinsic coercivity (Hci) and high maximum energy product [(BH)max] at elevated temperatures [1-9]. Simultaneously increasing the Hci and (BH)max in permanent magnets remains a significant scientific and technological materials design challenge, owing to the coercivity-remanence (Br)
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trade-off, which is notionally similar to the well-known strength-ductility trade-off in structural materials [10, 11]. Therefore, novel materials design and processing strategies are needed to
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circumvent this trade-off in permanent magnets, such as the Sm(Co,M)z magnets.
Fabrication of Sm(Co,M)z magnets involves a complex metallurgical process, including sintering, solutionizing for the formation of single disordered hexagonal Sm(Co,M)7 phases (1:7H),
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isothermal aging, and slow cooling down for forming rhombohedral Sm2(Co,M)17 matrix phases (2:17R) surrounded by Sm(Co,M)5 phases (1:5), and Zr-rich hexagonal platelet phases perpendicular to the rhombohedral Bravais lattice c axis [12]. Now, a simple expression for the
H ci ∝
d γ ( x) dx
(1)
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Hci of Sm(Co,M)z magnets is:
where γ(x) represents the average domain wall energy per unit area as a function of the wall position, x [13-15]. One well-accepted pinning mechanism in Sm(Co,M)z magnets is the repulsive domain wall pinning [3, 12, 14]. Magnetic domain walls unfavorably move into the 1:5 phases in the magnetic reversal process, since the magnetic domain wall energy of 1:5 phases is higher than that of the 2:17R matrix phases (i.e: γ2:17R < γ1:5) and magnetic domain walls prefer to stay at the low energy state. This is a classic example of. Now, much of the effort has been
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devoted to improving the performance of Sm(Co,M)z magnets via exploiting this repulsive domain wall pinning mechanism [10, 12, 16-20]. It was shown by Ojima et al. that a Br of up to ~11 kG and a (BH)max of ~30 MGOe can be achieved in Sm(Co,M)z magnets, while the Hci
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reached only ~6 kOe [18]. Liu and Ray developed a Sm(Co0.612Fe0.316Cu0.052Zr0.020)7.73 magnet alloy with an ultrahigh Br of up to 12 kG and a (BH)max of 31-33 MGOe, but the Hci was still below 20 kOe [17]. Given that increasing the Cu concentration of the 1:5 phases lowers the γ1:5
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[3, 13, 17, 21], the predominant factor limiting Hci in the repulsive-domain-wall-pinningcontrolled magnets is that Cu concentration cannot be reduced below a certain level in the 1:5
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phases to stabilize 2:17R matrix phases and 1:5 phases, This, in turn, sets an upper limit for the
γ1:5 and so the repulsive pinning strength that can be achieved. Therefore, the Hci-Br trade-off has still prevailed in Sm(Co,M)z magnets controlled by the repulsive domain wall pinning.
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Recently, an alternative approach has been explored, which uses an attractive domain wall pinning mechanism, and works by reducing the domain wall energy of the 1:5 phases to values whose below that of the 2:17R phases (γ2:17R > γ1:5) [22, 23]. For example, it has been shown via
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Lorentz microscopy that it was possible to have the magnetic domain walls strongly attracted by the 1:5 phases, indicating that γ2:17R > γ1:5 (implying attractive domain wall pinning), rather than
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γ2:17R < γ1:5 (implying repulsive domain wall pinning) [22, 23]. In this scenario, there is no constraint on increasing the Cu concentration in the 1:5 phases since this further lowers γ1:5, and the Hci can be enhanced substantially. However, it remains challenging to rigorously and precisely tune the microstructure and composition of constituent phases to fabricate attractivedomain-wall-pinning-controlled Sm(Co,M)z magnets.
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In this work, we manipulated the gradient of domain wall energy landscape of constituent phases (the 2:17R matrix phases and 1:5 phases) to realize the attractive domain wall pinning instead of repulsive domain wall pinning in the Sm-Co magnets (γ2:17R > γ1:5), by judiciously alloying
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copper particles in the pre-heat treatment process, to drastically enhance the Hci of ~12 kOe with ~1.8% reduction of remanence and maximum magnetic energy product in the rare earth permanent magnets. To valid this design and approach, we have used advanced microscopy and
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microanalysis techniques, such as a powerful atomic-scale analysis technique known as atom probe tomography, and micromagnetic simulations to understand the origin of exceptional
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improvement of magnetic properties caused by the Cu-particle-alloying technique.
2. Experimental 2.1. The material
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The alloy with a nominal composition of Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 was fabricated by induction melting. The crashed ingots were ball-milled into powders with an average size of ~5.5
µm [24]. Then, 1 wt.% Cu fine particles with an average size of ~2 µm prepared by inert gas
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condensation technique were alloyed into the Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 ball milled powders. The mixed powders were pressed under a pressure of 147 MPa in a magnetic field of 20 kOe,
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and then isostatically compacted under a press of 220 MPa for 30 s. The green bodies were sintered at ~1500 K for ~1h under an argon atmosphere, followed by a homogenization treatment at ~1400 K for ~3.5 h. The magnets were isothermally aged at ~1100 K for ~40 h, then slowly cooled down to ~700 K, and was aged again at ~700 K for about ~10 h [25]. Hereafter, 1 wt.% Cu fine particle alloyed Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 magnet is abbreviated to the alloyed
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magnet. To investigate the effects of the Cu-particle-alloying technique, pure magnets were
2.2. Magnetic hysteresis loops at different temperatures
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fabricated following the same procedure described above without Cu particles for comparison.
The magnetic properties of the pure and alloyed magnets were measured with a Quantum Design
2.3. Microstructural characterization
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magnetic field of up to 50 kOe at room temperature.
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physical property measurement system-vibrating sample magnetometer (PPMS-VSM) under a
The pure and alloyed magnets for scanning electron microscopy-electron backscatter diffraction (SEM-EBSD) analysis were mounted in epoxy and then mechanically polished using 220-grit, 500-grit, 1200-grit SiC papers, and standard colloidal silica suspension successively. SEM-
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EBSD mapping was carried out in a Zeiss Ultra Plus field-emission SEM equipped with an Oxford Instruments Aztec and Nordlys-nano EBSD detector. The Kikuchi patterns were recorded with an acceleration voltage of 25 kV, an aperture size of 120 µm, and a ‘high current’
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mode. The results were processed by the standard EBSD processing software. In addition, energy-dispersive X-ray spectroscopy (EDS) mapping was also obtained for revealing element
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distribution simultaneously.
The pure and alloyed magnets were cut into a small slice with the thickness of ~1 mm, polished and thinned down by 1500-grit SiC paper to ~50 µm. The samples were then attached to the copper grid and further thinned down to tens of nanometers by precision ion polishing system
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(PIPS). Transmission electron microscopy (TEM) experiments were conducted by a JEOL JEM2100F.
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The pure and alloyed magnets were milled down to a diameter of tens of nm using the tripod polishing method followed by focused ion beam (FIB) milling for atom probe tomography (APT) experiments [26]. APT measurements were carried out in a picosecond-pulse UV laser-assisted
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CAMECA local electrode atom probe (LEAP) 4000X Si. APT experiments were performed in a high vacuum, a cryogenic temperature of 50 K, under an intense electric field generated by a DC
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voltage up to 10000 V, with a 355 nm wavelength laser with pulse energy of ~50-100 pJ and a pulse frequency of ~100-200 kHz. The tomographic reconstructions were performed using Cameca’s Integrated Visualization & Analysis Software (IVAS) version 3.6.6 [27, 28]. Note that the morphology of reconstructed tips is influenced by user-defined reconstruction parameters,
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including but not limited to the detector efficiency, image compression factor (ICF) and field factor (kf), so absolute size measurement must be treated carefully [28]. If sufficient crystallographic information is present within the reconstruction process, calibration procedures
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may be used, but this was not possible in the collected APT datasets. Therefore, the default values of detector efficiency (0.50), ICF (1.650) and kf (3.30) were applied. The shank angle was
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selected as 7.5o considering the similarity between lengths of reconstructed APT datasets and actual lengths of evaporated sections of tips. The iso-concentration surfaces were defined with a voxel size of 1 nm and a delocalization value of x=3 nm, y=3 nm, z=1.5 nm. The error bars shown in 1D concentration profiles were calculated as E=(Ci(1-Ci)/N)1/2, where Ci=Ni/N, Ni represents the number of i solute ions/atoms and N represents the total number of counts with the given bin [28]. The bin width was selected as 0.2 nm in the 1D concentration profiles.
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2.4. Magnetic domain observation Magnetic force microscopy (MFM) measurements were conducted with a Bruker Dimension
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Icon (Santa Barbara, CA, USA). MFM is a secondary imaging mode of atomic force microscopy (AFM) using a two-pass technique. The initial pass measures the surface topography using tapping mode, then cantilever was lifted at a certain distance (~40 nm) above the sample surface
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in the second pass, thereby measuring the long-range magnetic forces between the probe and sample. Magnetic probes MESP-V2 (from Bruker AFM probes) were used to perform the
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magnetic measurements. 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 was recorded via changes in the phase and amplitude of the probe oscillation. They were further processed in the
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Gwyddion software to acquire two/three-dimensional (2/3D) MFM images.
2.5. Micromagnetic simulations
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Sandwich models (2:17R matrix phase/1:5 phase/2:17R matrix phase) with 80 nm × 30 nm × 30 nm in size were used to simulate the demagnetization curves of the pure and alloyed magnets.
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The thickness of the interfaces between the 1:5 phases and 2:17R matrix phases was set to 1 nm. Since simulated Hci values are almost identical no matter whether the magnetic domain wall is introduced in 2:17R matrix phases or 1:5 phases, the magnetic domain wall was introduced in the right 2:17R matrix phase in the model. The initial magnetization directions of the left 2:17R matrix phase, the 1:5 phase and the region left to the domain wall of the right 2:17R matrix phase were set upward, while the initial magnetization direction of the region right to the domain wall
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of the right 2:17R matrix phase was set downward. The external field was applied downwards from 0 to 40 kOe. The models were discretized by cubic nodes with a size of 1 nm, which is smaller than the domain wall width (δW=π(A/K1)1/2=~3.84-7.35 nm) and exchange length
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(δW=(A/µ 0MS2)1/2=~3.37-10.90 nm). The Landau-Lifshitz-Gilbert equation at each node was calculated by the 3D NIST Object Oriented MicroMagnetic Framework (OOMMF) software [29]. The exchange stiffness constant (A), magnetocrystalline anisotropy (K1) and saturation
3. Results and Discussion 3.1. Macromagnetic properties
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and Fe were calculated and summarized in Table 3.
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magnetization (MS) of 2:17R matrix phases and the 1:5 phases with various compositions of Cu
We measured the initial magnetization curves and hysteresis loops of the pure and Cu-particle-
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alloyed magnets. The initial magnetization curves manifest that the domain wall motion and domain wall pinning exist in both samples, but the alloyed magnet exhibits a larger pinning field (Fig. 1). Table 1 summarizes the corresponding magnetic properties of the pure and alloyed
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magnets. The Hci has been dramatically enhanced from 14.9 kOe in the pure magnet to 26.9 kOe in the alloyed magnet with ~1.8% reduction of the Br (from 11.4 kG to 11.2 kG) and (BH)max
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(from 26.4 MGOe to 26.6 MGOe). The Hci of the alloyed magnet is ~35% higher than that of the Sm(CobalFevCu0.056Zr0.020)7.546 sintered magnets while maintaining a comparably high (BH)max in Sm-Co based permanent magnets [17]. In addition, the squareness (Sr) of hysteresis curves has been improved from 81.2% to 85.8% by the Cu-particle-alloying technique. To summarize, we have achieved Sm(Co,M)z permanent magnets with exceptionally high overall magnetic performance (OMP) of Hci (kOe)+(BH)max (MGOe)=53.5, which has surpassed most of
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conventional Sm-Co based permanent magnets, as reported in the literatures [2, 3, 13, 16, 17, 30,
3.2. Microstructural observation by SEM and TEM
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31].
To fully unveil the origin of such substantial improvement of magnetic properties, global microstructure
and
microchemistry
of
pure
and
Cu-particle-alloyed
magnets
were
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comprehensively investigated by SEM-EBSD/EDS technique. Fig. 2 shows the global microstructure and elemental distributions of Sm, Co, Fe, O, Zr, and Cu of the pure and alloyed
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magnets. The bright and dark regions correspond to the Sm oxides and matrix grains (a combination of the 2:17R phases, 1:5 phases, and Zr-rich phases), respectively (Fig. 2a and b). Fig. 2c and d show the matrix grains with various grain shapes, crystallographic orientations, and grain sizes ranging from ~10 µm to ~30 µm of the pure and alloyed magnets. In addition, some
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large Zr-rich pockets were observed in both samples as well, in agreement with previous work done by Yosuke Horiuchi et al.[32]. Table 2 exhibits that the Cu concentration of matrix grains increases from ~5.20% in the pure magnet to ~5.82 at.% in the alloyed magnet, indicating that
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the additive Cu atoms were diffused into the 1:5 cell boundary phases through grain boundaries
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(GBs) in the alloyed magnet since Cu is not soluble in the 2:17R matrix phases.
We further investigated the constituent phases of pure and Cu-particle-alloyed magnets by using TEM. The typical TEM images show the cellular 2:17R matrix phases and 1:5 cell boundary phases of the pure and alloyed magnets at the nanoscale (Fig. 3a and b). The cellular sizes of 2:17R matrix phases are comparably similar (~100-150 nm). However, some discontinuous 1:5 cell boundary phases were observed in the pure magnet marked by white arrows (Fig. 3a). In the
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alloyed magnet, we found the improved continuity of 1:5 cell boundary phases (Fig. 3b), which is believed to contribute to the Hci enhancement [10]. Furthermore, the steady formation of the 1:5 cell boundary phases and cellular 2:17R matrix phases in the alloyed magnet leads to a high
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Br and a high (BH)max as well [10].
3.3. Distribution of individual atoms and compositions of the constituent phases
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We utilized APT to investigate the 3D distributions of various atom species of the cellular 2:17R matrix phases and 1:5 cell boundary phases of the pure and alloyed magnets at the atomic scale.
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To assure the reliability, we performed two separate APT experiments for each sample (One of them is shown in the main context and the other one is exhibited in the Supplemental Materials). Fig. 4 shows the 3D atom distribution maps of Fe, Sm, Zr, and Cu in the pure and alloyed magnets at the atomic scale, respectively. Iso-concentration surfaces of the Cu atoms (12.0 at.%
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for both pure and alloyed magnets) reveal the interfaces between the 1:5 cell boundary phases and cellular 2:17R matrix phases and iso-concentration surfaces of Zr atoms (1.0 at.% for the pure magnet/3.7 at.% for the alloyed magnet) illustrate the interfaces between the cellular 2:17R
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matrix phases and Zr-rich phases.
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As a representative, 1D concentration profiles from the black box (Fig. 4a to c) reveals that Cu and Sm are enriched in the ~6-nm-wide 1:5 cell boundary phase in the pure magnet, while Fe and Co are depleted. The Cu (Fe) concentrations are as high (low) as ~30 at.% (~12 at.%) at the maximum (minimum) value in the 1:5 cell boundary phase in the pure magnet. 1D concentration profiles from the black box (Fig. 4d to f) reveal that Cu and Sm are also enriched in the ~6-nmwide 1:5 cell boundary phase, while Fe and Co are depleted as well. However, the Cu (Fe)
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concentrations are as high (low) as ~43 at.% (~9 at.%) at the maximum (minimum) value in the 1:5 cell boundary phase in the alloyed magnet. Owing to the inhomogeneous distribution of these species in/across the 1:5 cell boundary phases, around twenty different regions were selected for
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identifying the distribution of maximum Cu concentration of the 1:5 cell boundary phases on the basis of two APT datasets of the pure and alloyed magnets, respectively (Supplemental Text, Fig. S1-S4, and Table S1-S2). Fig. 4g exhibits the corresponding probability distribution of the
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maximum Cu concentration of the 1:5 cell boundary phases in the pure and alloyed magnets. It should be noted that (1) the maximum Cu concentration of the 1:5 cell boundary phases ranges
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from ~18-38 at.% in the pure magnet, and ~26-50 at.% in the alloyed magnet, respectively, which lessens the Sr of the pure and alloyed magnets; (2) given that the Hci is a manifestation of the pinning strength of individual 1:5 cell boundary phases and represents the state in which the positive and negative magnetization of the Sm-Co magnets are equivalent to each other, the Hci is
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approximately equal to the pinning strength of the 1:5 cell boundary phases whose maximum Cu concentration is the value (~30 at.% for pure magnet and ~43 at.% for the alloyed magnet) in
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which the cumulative probability is 0.5.
3.4. Parameter estimation for micromagnetic simulations
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We bridged the aforementioned experimental observations and theoretical simulations by estimating the saturation magnetization (MS), magnetocrystalline anisotropy constant (K1), and exchange stiffness (A) of the cellular 2:17R matrix phases and 1:5 cell boundary phases in the pure and alloyed magnets on the basis of the APT results via polynomial fitting methods.
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Here,
we
selected
Sm10.55Co58.50Fe26.71Cu1.29Zr2.88Other0.07
and
Sm10.51Co57.82Fe27.52Cu1.52Zr1.61Other1.02 as the compositions of cellular 2:17R matrix phases in pure and alloyed magnets, respectively. To simplify compositions of cellular 2:17R matrix
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phases in pure and alloyed magnets, Cu and Zr atoms were neglected due to their little contribution to the magnetic properties of cellular 2:17R matrix phases. Under the circumstance, the compositions of cellular 2:17R matrix phases were determined to be Sm2(Co0.69Fe0.31)17 for
pure
and
alloyed
magnets,
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the pure magnet and Sm2(Co0.68Fe0.32)17 for the alloyed magnet, respectively. To simulate Hci of Sm16.89Co42.92Fe8.51Cu30.61Zr0.50Other0.57
and
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Sm16.41Co31.93Fe5.91Cu43.45Zr0.41Other1.89 were selected as compositions of 1:5 cell boundary phases in pure and alloyed magnets in which the cumulative probability is 0.5. To simplify the compositions of 1:5 cell boundary phases in the pure and alloyed magnets, Zr atoms were neglected due to its little contribution to magnetic properties of 1:5 cell boundary phases. Under
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the circumstance, the compositions of 1:5 cell boundary phases were determined to be Sm(Co0.52Fe0.11Cu0.37)5 and Sm(Co0.39Fe0.07Cu0.54)5, respectively.
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Firstly, the MS of cellular 2:17R matrix phases in the pure and alloyed magnets were polynomial fitting as follows [33, 34]:
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y = 12.60 + 13.99x − 62.71x2 + 164.41x3 − 183.17 x4 + 65.70x5
(2)
where x and y represent the concentration of Fe in Sm2Co17 phases and their corresponding MS, respectively. Therefore, the MS of Sm2(Co0.69Fe0.31)17 was determined to be 14.30 kG in the pure magnet. The MS of Sm2(Co0.68Fe0.32)17 was obtained as 14.34 kG in the alloyed magnet as well. If only Sm, Co, and Fe exist in 1:5 cell boundary phases, the MS of 1:5 cell boundary phases in the pure and alloyed magnets were polynomial fitting as follows [35]:
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y = 9.45 + 16.46 x − 41.15x2 + 163.41x3 − 229.49 x4 + 96.95x5
(3)
where x and y represent the concentration of Fe in SmCo5 phases and their corresponding MS,
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respectively. Therefore, the MS of 1:5 cell boundary phases (only considering Sm, Co, and Fe) in pure and alloyed magnets was determined to be 11.68 kG and 11.44 kG, respectively. However, the incorporation of Cu severely reduces the MS of Sm(Fe,Co)5 phases. The relation between the Cu content in the Sm(Fe,Co,Cu)5 phase (x) and the corresponding MS (y) was fitted with
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experimental datasets in [36]:
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y = 11.47 − 2.26x − 111.07 x2 + 304.14x3 − 255.94x4 − 23.67 x5
y = 11.25 − 2.92x − 117.07 x2 + 380.07 x3 − 464.70x4 + 152.05x5
(4) (5)
Therefore, the MS of Sm(Co0.52Fe0.11Cu0.37)5 was determined to be 5.87 kG in the pure magnet.
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The MS of Sm(Co0.39Fe0.07Cu0.54)5 was obtained as 2.85 kG in the alloyed magnet as well.
Secondly, the exchange stiffness (A) of cellular 2:17R matrix phases in the pure and alloyed magnets were calculated. Curie Temperature (TC) of cellular 2:17R matrix phases in the pure and
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alloyed magnets were polynomial fitting as follows [34]: y = 1189.00 + 246.59x − 2584.71x2 + 5108.61x3 − 5592.88x4 + 2023.60x5
(6)
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where x and y represent the concentration of Fe in Sm2Co17 phases and their corresponding TC, respectively. Therefore, the TC of Sm2(Co0.69Fe0.31)17 phase in the pure magnet was calculated as ~1123 K. The TC of Sm2(Co0.68Fe0.32)17 phase in the alloyed magnet was calculated as ~1119 K. The relationship between the A and TC can be shown as [37]: A=
3k BTC 2 ZS ( S + 1)
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(7)
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where kB, Z, and S correspond to Boltzmann constant, the coordination number, spin quantum number of the material. Here, kB, Z, and S were assumed nearly constant in the same system even with little chemical variation. Therefore,
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A ∝ TC
(8)
The A of Sm2Co17 was reported as 19.60 pJ/m [2]. According to the equation (8), the A of Sm2(Co0.69Fe0.31)17 phase in the pure magnet was determined to be 18.50 pJ/m. The A of
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Sm2(Co0.68Fe0.32)17 phase in the alloyed magnet was calculated as 18.45 pJ/m.
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If only Sm, Co, and Fe exist in 1:5 cell boundary phases, TC remains nearly constant with Fe concentration lower than 40 at.% [38]. Therefore the A of 1:5 cell boundary phases (only considering Sm, Co, and Fe) are the same as that of pure SmCo5 phases, which was reported as 15.10 pJ/m [2]. However, the incorporation of Cu severely reduces TC of Sm(Fe,Co)5 phases.
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The relationship between the Cu content in the Sm(Fe,Co,Cu)5 phase (x) and the corresponding TC (y) was fitted with experimental datasets in [36]: y = 977.00 − 536.97 x − 2643.65x2 + 10410.68x3 − 17622.54x4 + 9584.99x5
(9)
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Therefore, the TC of Sm(Co0.52Fe0.11Cu0.37)5 phase was determined to be 679.93 K in the pure
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magnet. The TC of Sm(Co0.39Fe0.07Cu0.54)5 phase was obtained as 497.11 K in the alloyed magnet. Thus, the A of Sm(Co0.52Fe0.11Cu0.37)5 phase was determined to be 10.51 pJ/m in the pure magnet. The A of Sm(Co0.39Fe0.07Cu0.54)5 phase was obtained as 7.68 pJ/m in the alloyed magnet.
Finally, the magnetocrystalline anisotropy constant (K1) of cellular 2:17R matrix phases in pure and alloyed magnets were polynomial fitting as follows [39]: y = 4.99 + 9.88x − 46.28x2 − 239.80 x3 + 1173.02 x4 − 1336.83x5
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(10)
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where x and y represent the concentration of Fe in Sm2Co17 phases and their corresponding K1, respectively. Therefore, the K1 of Sm2(Co0.69Fe0.31)17 phase in the pure magnet and Sm2(Co0.68Fe0.32)17 phase in the alloyed magnet was calculated as 3.47 MJ/m3 and 3.37 MJ/m3,
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respectively.
The K1 of 1:5 cell boundary phases remains challenging to determine due to various values of
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anisotropy field HA (317 kOe and 415 kOe) of SmCo5 phases reported by researchers [40, 41]. Therefore, the average value of HA was selected, which is 366 kOe. If only Sm, Co, and Fe exist
were polynomial fitting as follows [40]:
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in 1:5 cell boundary phases, the HA of 1:5 cell boundary phases in the pure and alloyed magnets
y = 373.92 − 210.30x + 3674.83x2 − 12547.92x3 + 13932.20x4 − 5082.05x5
(11)
where x and y represent the concentration of Fe in Sm2Co17 phases and their corresponding HA,
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respectively. Therefore, the HA of Sm(Co0.52Fe0.11Cu0.37)5 phase was determined to be 393.62 kOe in the pure magnet. The HA of Sm(Co0.39Fe0.07Cu0.54)5 phase was obtained as 389.38 kOe in
EP
the alloyed magnet as well.
If only Sm, Co, and Fe exist in 1:5 cell boundary phases, the MS of 1:5 cell boundary phases in
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pure and alloyed magnets are 11.68 kG and 11.44 kG, respectively. The relation between the HA and K1 can be shown as [37]: K1 =
H AM S 2
(12)
Therefore, the K1 of 1:5 cell boundary phases in pure and alloyed magnets were calculated as 18.30 and 17.72 MJ/m3, respectively.
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However, the incorporation of Cu severely reduces K1 of 1:5 cell boundary phases in pure and alloyed magnets. The relation between the Cu content in the Sm(Fe,Co,Cu)5 phase (x) and the
y=−
a 0.6
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corresponding K1 (y) was fitted based on experimental datasets in [36]: x+a
(13)
where a corresponds to the K1 without any Cu incorporation. In the pure magnet, a=18.30
SC
(MJ/m3), and x=0.37. Therefore, the K1 of Sm(Co0.52Fe0.11Cu0.37)5 in the pure magnet was determined to be 7.02 MJ/m3. In the alloyed magnet, c=17.72 (MJ/m3), and x=0.54. Therefore,
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the K1 of 1:5 Sm(Co0.39Fe0.07Cu0.54)5 in the alloyed magnet was determined to be 1.77 MJ/m3. The calculated values are in good agreement with the results published by the Toshikazu et al.[40].
Here,
we
selected
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3.5. Micromagnetic simulations
Sm10.55Co58.50Fe26.71Cu1.29Zr2.88Other0.07
[Sm2(Co0.69Fe0.31)17]
and
Sm10.51Co57.82Fe27.52Cu1.52Zr1.61Other1.02 [Sm2(Co0.68Fe0.32)17] as the compositions of the cellular
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2:17R matrix phases in the pure and alloyed magnets, respectively. To simulate the Hci values of the pure and alloyed magnets, Sm16.89Co42.92Fe8.51Cu30.61Zr0.50Other0.57 [Sm(Co0.52Fe0.11Cu0.37)5]
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and Sm16.41Co31.93Fe5.91Cu43.45Zr0.41Other1.89 [Sm(Co0.39Fe0.07Cu0.54)5] were selected as the compositions of the 1:5 cell boundary phases in the pure and alloyed magnets owing to the previously mentioned correlation between the Hci and the individual pinning strength of the 1:5 cell boundary phases. Strikingly, we found that the repulsive domain wall pinning (γ2:17R < γ1:5) predominates in the pure magnet, whereas the attractive domain wall pinning (γ2:17R > γ1:5) in the alloyed magnet (Table 3), in consistency with previous observations by Lorentz TEM microscopy [22, 23]. In this scenario, there is no constraint on increasing the Cu concentration in 17
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the 1:5 phases, which further lowers the γ1:5, and the Hci can be enhanced substantially. More precisely, ∆γ between the cellular 2:17R matrix phases and 1:5 cell boundary phases in the alloyed magnet is 16.79 mJ/m2, which is around sevenfold larger than that between these phases
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in the pure magnet (~|-2.31| mJ/m2), which leads to the substantial Hci enhancement.
We performed micromagnetic simulations on the pure and alloyed magnets based on the
SC
parameters estimated from our experimental observations (Fig. 5a to d). Under the increasing external field, the magnetic domain wall moves leftwards and is pinned at the 1:5 cell boundary
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phase. Fig. 5b shows the simulated demagnetization curves of the pure and alloyed magnets, in which the plateaus in the curves results from the domain wall pinning. Our micromagnetic simulation results show that the Hci is enhanced by ~18.0 kOe from the pure magnet (~9.0 kOe) to the alloyed magnet (~27.0 kOe), in excellent agreement with our macromagnetic properties
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(Fig. 1). In the pure magnet, the magnetic domain wall moves in the right 2:17R matrix phase towards the 1:5 cell boundary phase under a small external field of 1 kOe. With the increasing external field, the center of magnetic domain wall is repulsively pinned from the largest domain
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wall energy gradient at the interface between the right 2:17 phase and the 1:5 cell boundary phase. When the external field is up to 9.0 kOe, the domain wall is repulsively depinned from the
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1:5 cell boundary phase and moved leftwards continuously into the 1:5 cell boundary phase and then the left 2:17R matrix phase (Fig. 5d). Similarly, the domain wall moves in the right 2:17R matrix phase towards the 1:5 cell boundary phase under a small external field of 1.0 kOe in the alloyed magnet. However, the domain wall moves across the 1:5 cell boundary phase and is attractively pinned at the largest domain wall energy gradient at the interface between the left 2:17R matrix phase and the 1:5 cell boundary phase with increasing external field, since the
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magnetic domain wall is easily trapped in the 1:5 cell boundary phase. When the external field is up to 27.0 kOe, the domain wall is attractively depinned from the 1:5 cell boundary phase and moved leftwards continuously into the left 2:17R matrix phase (Fig. 5d). The actual position of
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the domain wall is determined by the interplay between the domain wall energy and magnetostatic energy [37].
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We also quantified the effects of the domain wall energy gradient and the thickness of the 1:5 cell boundary phases on the Hci of the alloyed magnet. The smallest domain wall energy gradient
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only gives rise to a smaller Hci up to 20 kOe in the alloyed magnet (Fig. 5e and f). In contrast, the largest domain wall energy gradient results in a larger Hci of 27 kOe in the alloyed magnet, even with the same domain wall energy difference between the 2:17R matrix phase and 1:5 cell boundary phase (Fig. 5e and f). Moreover, the Hci of the alloyed magnet is reduced from 27 kOe
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to 20 kOe with the thickness of the 1:5 cell boundary phases reducing from 7 nm to 3 nm (Fig. 5g and h). because the domain wall may not be fully trapped into these 1:5 cell boundary phases if the domain wall width in the 1:5 cell boundary phases (~6.54 nm) is larger than the thickness
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of 1:5 cell boundary phases. Furthermore, the Hci of the alloyed magnet remains constant with a larger thickness of the 1:5 cell boundary phases (>7 nm) since the domain wall has already been
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fully trapped in 1:5 cell boundary phases (Fig. 5g and h). To summarize, the pinning strength of the 1:5 cell boundary phases in Sm-Co based permanent magnets is not only correlated to the domain wall energy difference of the cellular 2:17R matrix phases and 1:5 cell boundary phases, but also associated with the domain wall energy gradient and the thickness of the 1:5 cell boundary phases.
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3.6. Summary We illustrated the microstructural changes occurring in the pure and Cu-particle-alloyed magnets (Fig. 6). In the pure magnet, normal 1:7H matrix grains are formed after the sintering process
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and homogenization (Fig. 6a). Moreover, discontinuous 1:5 cell boundary phases and cellular 2:17R matrix phases are formed after the aging process and furnace cooling. Using the Cuparticle-alloying-technique (Fig. 6b), Cu-rich 1:7H matrix grains and GBs are formed after the
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sintering process and homogenization. In addition, continuous 1:5 cell boundary phases and cellular 2:17R matrix phases are formed after the aging process and furnace cooling. Most
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importantly, the Cu concentration in 1:5 cell boundary phases is increased significantly, which is not only driven by the thermal diffusion during the slow cooling process, but also the concentration gradient of residual Cu located near the GBs. This results in the attractive domain pinning instead of repulsive domain wall pinning in Cu-particle-alloyed Sm(Co,M)z permanent
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magnets.
In Cu-particle-alloyed Sm-Co magnets, not only the intrinsic coercivity but also the squareness is
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improved significantly. For the coercivity enhancement, the addition of Cu lowers the domain wall energy of 1:5 cell boundary phases, which enlarges the difference of domain wall difference
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between 1:5 cell boundary phases and 2:17 matrix phases and therefore leads to the enhancement of the domain wall pinning strength at 1:5 cell boundary phases.
For the improved squareness, the addition of Cu is likely to increase the fraction of 1:5 cell boundary phases near grain boundaries. In the pure Sm-Co magnets, the grain boundaries are Culean [32]. Most importantly, few 1:5 cell boundary phases are formed in the regions near these
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Cu-lean grain boundaries [32]. Furthermore, the concentration of Cu in these inchoate 1:5 cell boundary phases varies, which leads to the low squareness of pure Sm-Co magnets. In the Cuparticle-alloyed Sm-Co magnets, the concentration of Cu at grain boundaries is increased, which
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also facilitates the formation of 1:5 cell boundary phases near grain boundaries. Since the number of Cu-lean or inchoate 1:5 cell boundary phases near grain boundaries is reduced in the Cu-particle-alloyed Sm-Co magnet, the number of 1:5 cell boundary phases with low domain
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wall pining strength is reduced (In other words, the distribution of maximum Cu concentration in
particle-alloyed Sm-Co magnet.
4. Conclusions
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1:5 cell boundary phases becomes narrow), which leads to the improved squareness of Cu-
In conclusion, our work here demonstrates high iron content Sm(Co,Fe,Cu,Zr)z permanent
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magnets with a combination of a intrinsic coercivity of 26.9 kOe, a remanence of 11.2 kG, and a maximum energy product up to 26.6 MGOe simultaneously via the Cu-particle-alloying technique, which outperforms most of conventional Sm(Co,Fe,Cu,Zr)z permanent magnets. By
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integrating multiple cutting-edge microscopy techniques and micromagnetic simulations, we find that (1) attractive domain wall pinning, instead of widely accepted repulsive domain wall pinning,
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predominates in the Cu-particle-alloyed magnet, (2) the exceptionally high Hci is attributed to the continuity and enlarged pinning strength of the 1:5 cell boundary phases in the Cu-particlealloyed magnet. (3) High (BH)max results from the stable formation of the 1:5 cell boundary phases and cellular 2:17R matrix phases in high content Sm(Co,Fe,Cu,Zr)z permanent magnets. Our work not only demonstrates the transformation from repulsive pinning to attractive pinning by alloying Cu particles in Sm-Co based permanent magnets but also provides insights into the
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grain boundary engineering for enhancing the Hci of rare earth permanent magnets. Due to the versatility of the Cu-particle-alloying technique, further performance optimization is possible, for example, by tuning the Cu particle size to achieve uniform Cu diffusion from the grain
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boundaries to the 1:5 cell boundary phases, which will further improve the magnetic properties of Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 permanent magnets. Moreover, the methodology applied in this work may also offer exciting possibilities for quantitative analysis and prediction between
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compositions and magnetic properties of other magnetic materials.
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Acknowledgments
The authors would like to thank Ms. Katie Levick, Dr. Patrick Trimby, Dr. Anna Ceguerra, Dr. Takanori Sato, Mr. Adam Sikorski, and Dr. Jacob Warner for their technical support. The authors also would like to Dr. Yisheng Chen and Dr. Yahui Cheng for fruitful discussion. The
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authors would like to thank the Australian Microscopy and Microanalysis Research Facility (AMMRF) and Australian Centre for Microscopy and Microanalysis (ACMM). This work was
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partly supported by the Australian Research Council (DP150100018).
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[36] E. Lectard, C. Allibert, R. Ballou, Saturation magnetization and anisotropy fields in the Sm(Co1-xCux)5 phases, J. Appl. Phys. 75 (1994) 6277-6279. [37] J. M. Coey, Magnetism and magnetic materials, Cambridge University Press, 2010. [38] T. Saito, D. Nishio-Hamane, Magnetic properties of SmCo5-xFex (x=0-4) melt-spun ribbon, J. Alloys Compd. 585 (2014) 423-427. [39] H. Mildrum, M. S. Hartings, K. J. Strnat, J. Tront, Magnetic properties of the intermetallic phases Sm2(Co,Fe)17, AIP Conf. Proc. 10 (1973) 618-622. [40] T. Katayama, T. Shibata, Magneto-crystalline anisotropy constant in Sm(Co,Cu)5 base alloy, Jpn. J. Appl. Phys. 12 (1973) 762. [41] T.-S. Chin, C.-H. Lin, H. Bai, Y. Huang, K. Cheng, S. Huang, Effect of nitrogen content on magnetic properties of Sm2Fe17Nx and SmFe5Ny alloys, IEEE Trans. Magn. 28 (1992) 2587-2589.
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Table
1.
Magnetic
properties
of
the
pure
and
Cu-particle-alloyed
Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 permanent magnets. Hci (kOe)
11.4
14.9
1.0 (Cu particles)
8.4
11.2
26.9
(BH)max (MGOe) 26.4
Sr (%)
26.6
85.8
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81.2
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Br (kG)
0
Density (g/cm3) 8.4
t (wt. %)
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Table 2. Experimentally determined compositions of the matrix grains (a combination of
particle-alloyed magnets by SEM-EDS. Sample
Composition (at.%) Sm 11.32 11.74
Cu 5.20 5.82
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Fe 23.04 23.46
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Pure Cu particlealloyed
Co 58.62 57.10
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2:17R matrix phases, 1:5 cell boundary phases, and Zr-rich phases) in pure and Cu-
Zr 1.82 1.88
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Table 3. The parameter estimation of the cellular 2:17R matrix phases and 1:5 cell boundary phases in the pure and alloyed magnets. Phas e
Composition
K1 (MJ /m3)
MS (kG)
A (pJ/ m)
γw (mJ/m 2 )
Pure
2:17 R 1:5
Sm2(Co0.69Fe0.31)17
3.47
14.30 18.50
32.05
Sm(Co0.52Fe0.11Cu 0.37)5 Sm2(Co0.68Fe0.32)17
7.02
5.87
10.51
34.36
3.37
14.34 18.45
31.54
Sm(Co0.39Fe0.07Cu 0.54)5
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2:17 R 1:5
1.77
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Cuparti clealloy ed
2.85
Dom Pinnin ain g Type wall width (nm) 7.25 Repuls ive 3.84 Pinnin g 7.35 Attrac tive 6.54 Pinnin g
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Sam ple
7.68
14.75
Fig.
1.
The
magnetization
curves
of
pure
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Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 permanent magnets.
the
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and
Cu-particle-alloyed
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Fig. 2. The microstructure and microchemistry of the pure and Cu-particle-alloyed magnets. (a) Secondary electron (SE) image, (c) Backscattered electron (BSE) image, (e) The corresponding all Euler map, EDS maps of (g) Sm, (i) Co, (k) Fe, (m) Zr, (o) O and (q) Cu of the pure magnet.
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(b) SE image, (d) BSE image, (f) The corresponding all Euler map, EDS maps of (h) Sm, (j) Co,
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(l) Fe, (n) Zr, (p) O and (r) Cu of the doped magnet. The scale bars are 50 µm.
Fig.
3.
Bright
field
TEM
images
of
(a)
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pure
and
(b)
Cu-particle-alloyed
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Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 permanent magnets. c axis is out of the plane. The scale bars are
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100 nm.
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FIG. 4. APT data of the pure and Cu-particle-alloyed magnets at the atomic scale. The pure magnet: (a) 3D atom distribution maps of Fe, Sm, Cu, Zr, iso-concentration surfaces of Cu (12.0 at.%) and Zr (1.0 at.%). (319 nm × 71 nm × 71 nm). Movie S1 shows the 3D rotation of Fe, Sm,
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Cu and Zr atoms. (b) A rotational view of (a) (without Zr iso-concentration surfaces). (c) 1D concentration profiles of Co, Fe, Sm, Cu, and Zr of the black box in( b). The alloyed magnet: (d), 3D atom distribution maps of Fe, Sm, Cu, Zr, iso-concentration surfaces of Cu (12.0 at.%) and
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Zr (3.7 at.%). (301 nm × 68 nm × 68 nm). Movie S2 shows the 3D rotation of Fe, Sm, Cu and Zr atoms. (e) A rotational view of d) (without Zr iso-concentration surfaces). (f) 1D concentration
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profiles of Co, Fe, Sm, Cu, and Zr of the black box in (e). (g) The probability distribution of the maximum Cu concentration of 1:5 cell boundary phases in the pure and Cu-particle-alloyed
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magnets based on around twenty atom probe tomography datasets for each sample.
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Fig. 5. Micromagnetic simulation results of the pure and Cu-particle-alloyed magnets. (a) The parameters used for simulating the demagnetization curves of the pure and alloyed magnets. (b) The simulated demagnetization curves of the pure and alloyed magnets. (c) Schematic diagram
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of the initial state of the micromagnetic model containing two 2:17R matrix phases and one 1:5 cell boundary phase. (d) The simulated repulsively and attractively depinned situations of the pure and alloyed magnets. (e) The parameters used for simulating the effects of the domain wall
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energy gradient on the Hci of the alloyed magnet. (f) The simulated demagnetization curves of the alloyed magnet with various domain wall energy gradients. (g) The parameters used for simulating the effects of the thickness of the 1:5 cell boundary phases on the Hci of the alloyed
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thicknesses of the 1:5 cell boundary phases.
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magnet. (h) The simulated demagnetization curves of the alloyed magnet with various
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Fig. 6. Schematic illustration of the microstructural evolution of the pure and Cu-particle-alloyed
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magnets. (a) The microstructure characteristics of the pure magnet after mixing the matrix powers, after the sintering process and homogenization, and after the aging process and furnace cooling. (b) The microstructure characteristics of the alloyed magnet after mixing the matrix
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powers and Cu particles, after the sintering process and homogenization, and after the aging process and furnace cooling. The scale bars in the left and middle columns are 5 µm. The scale
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S1359-6454(18)30846-2
DOI:
https://doi.org/10.1016/j.actamat.2018.10.046
Reference:
AM 14924
To appear in:
Acta Materialia
Received Date: 12 July 2018 Revised Date:
23 October 2018
Accepted Date: 24 October 2018
Please cite this article as: H. Chen, Y. Wang, Y. Yao, J. Qu, F. Yun, Y. Li, S.P. Ringer, M. Yue, R. Zheng, Attractive-domain-wall-pinning controlled Sm-Co magnets overcome the coercivity-remanence trade-off, Acta Materialia (2018), doi: https://doi.org/10.1016/j.actamat.2018.10.046. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Attractive-domain-wall-pinning controlled Sm-Co magnets overcome the coercivity-remanence trade-off
a
b
, Simon P. Ringer c, f, Ming Yue d, *, and Rongkun Zheng a, b, c, *
School of Physics, The University of Sydney, NSW, 2006, Australia
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d
The University of Sydney Nano Institute, The University of Sydney, Sydney, NSW
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2006, Australia c
Australian Centre for Microscopy and Microanalysis, The University of Sydney, NSW, 2006,
Australia d
College of Materials Science and Engineering, Beijing University of Technology, Beijing
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100124, China e
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Hansheng Chen a, b, c, #, Yunqiao Wang d, #, Yin Yao e, Jiangtao Qu a, b, c, Fan Yun a, b, c, Yuqing Li
Electron Microscope Unit, Mark Wainwright Analytical Centre, The University of New South
Wales, New South Wales, 2052, Australia
School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney,
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f
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Sydney, NSW 2006, Australia
#
These authors contributed equally to this work
*
Corresponding author. Email: [email protected] (R.Z.); [email protected]
(M.Y.)
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Abstract Traditional approaches for increasing the intrinsic coercivity of magnets typically come at the expense of remanence, a dilemma known as intrinsic coercivity-remanence trade-off, leading to a
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substantial reduction of maximum energy product. New metallurgical processing might offer the possibility of overcoming this trade-off. Here, we achieve a combination of an intrinsic coercivity of 26.9 kOe, a remanence of 11.2 kG, and a maximum energy product up to 26.6
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MGOe, which surpasses most of conventional Sm-Co based permanent magnets, by manipulating the gradient of domain wall energy landscape of constituent phases to realize the
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attractive domain wall pinning in Sm(Co,Fe,Cu,Zr)z permanent magnets. Using powerful atomicscale analysis technique known as atom probe tomography and micromagnetic simulations, we reveal that an enlarged attractive domain wall pinning strength results in the substantial coercivity enhancement with little sacrifice of remanence and maximum energy product in the
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Cu-particle-alloyed magnet. These results provide atomic-level insights into the coercivity mechanism of rare earth permanent magnets, with the methodology offering exciting possibilities for quantitative analyses and prediction between compositions and magnetic properties of other
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magnetic materials.
Keywords Sm-Co
magnets,
attractive-domain-wall-pinning,
micromagnetic simulations
2
coercivity,
atom
probe
tomography,
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1. Introduction Sm(Co,M)z (M=Cu, Fe and/or Zr) alloys are widely used as permanent magnets in satellite communications, aerospace, and defence technologies because of their capacity to retaining high
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intrinsic coercivity (Hci) and high maximum energy product [(BH)max] at elevated temperatures [1-9]. Simultaneously increasing the Hci and (BH)max in permanent magnets remains a significant scientific and technological materials design challenge, owing to the coercivity-remanence (Br)
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trade-off, which is notionally similar to the well-known strength-ductility trade-off in structural materials [10, 11]. Therefore, novel materials design and processing strategies are needed to
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circumvent this trade-off in permanent magnets, such as the Sm(Co,M)z magnets.
Fabrication of Sm(Co,M)z magnets involves a complex metallurgical process, including sintering, solutionizing for the formation of single disordered hexagonal Sm(Co,M)7 phases (1:7H),
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isothermal aging, and slow cooling down for forming rhombohedral Sm2(Co,M)17 matrix phases (2:17R) surrounded by Sm(Co,M)5 phases (1:5), and Zr-rich hexagonal platelet phases perpendicular to the rhombohedral Bravais lattice c axis [12]. Now, a simple expression for the
H ci ∝
d γ ( x) dx
(1)
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Hci of Sm(Co,M)z magnets is:
where γ(x) represents the average domain wall energy per unit area as a function of the wall position, x [13-15]. One well-accepted pinning mechanism in Sm(Co,M)z magnets is the repulsive domain wall pinning [3, 12, 14]. Magnetic domain walls unfavorably move into the 1:5 phases in the magnetic reversal process, since the magnetic domain wall energy of 1:5 phases is higher than that of the 2:17R matrix phases (i.e: γ2:17R < γ1:5) and magnetic domain walls prefer to stay at the low energy state. This is a classic example of. Now, much of the effort has been
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devoted to improving the performance of Sm(Co,M)z magnets via exploiting this repulsive domain wall pinning mechanism [10, 12, 16-20]. It was shown by Ojima et al. that a Br of up to ~11 kG and a (BH)max of ~30 MGOe can be achieved in Sm(Co,M)z magnets, while the Hci
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reached only ~6 kOe [18]. Liu and Ray developed a Sm(Co0.612Fe0.316Cu0.052Zr0.020)7.73 magnet alloy with an ultrahigh Br of up to 12 kG and a (BH)max of 31-33 MGOe, but the Hci was still below 20 kOe [17]. Given that increasing the Cu concentration of the 1:5 phases lowers the γ1:5
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[3, 13, 17, 21], the predominant factor limiting Hci in the repulsive-domain-wall-pinningcontrolled magnets is that Cu concentration cannot be reduced below a certain level in the 1:5
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phases to stabilize 2:17R matrix phases and 1:5 phases, This, in turn, sets an upper limit for the
γ1:5 and so the repulsive pinning strength that can be achieved. Therefore, the Hci-Br trade-off has still prevailed in Sm(Co,M)z magnets controlled by the repulsive domain wall pinning.
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Recently, an alternative approach has been explored, which uses an attractive domain wall pinning mechanism, and works by reducing the domain wall energy of the 1:5 phases to values whose below that of the 2:17R phases (γ2:17R > γ1:5) [22, 23]. For example, it has been shown via
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Lorentz microscopy that it was possible to have the magnetic domain walls strongly attracted by the 1:5 phases, indicating that γ2:17R > γ1:5 (implying attractive domain wall pinning), rather than
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γ2:17R < γ1:5 (implying repulsive domain wall pinning) [22, 23]. In this scenario, there is no constraint on increasing the Cu concentration in the 1:5 phases since this further lowers γ1:5, and the Hci can be enhanced substantially. However, it remains challenging to rigorously and precisely tune the microstructure and composition of constituent phases to fabricate attractivedomain-wall-pinning-controlled Sm(Co,M)z magnets.
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In this work, we manipulated the gradient of domain wall energy landscape of constituent phases (the 2:17R matrix phases and 1:5 phases) to realize the attractive domain wall pinning instead of repulsive domain wall pinning in the Sm-Co magnets (γ2:17R > γ1:5), by judiciously alloying
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copper particles in the pre-heat treatment process, to drastically enhance the Hci of ~12 kOe with ~1.8% reduction of remanence and maximum magnetic energy product in the rare earth permanent magnets. To valid this design and approach, we have used advanced microscopy and
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microanalysis techniques, such as a powerful atomic-scale analysis technique known as atom probe tomography, and micromagnetic simulations to understand the origin of exceptional
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improvement of magnetic properties caused by the Cu-particle-alloying technique.
2. Experimental 2.1. The material
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The alloy with a nominal composition of Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 was fabricated by induction melting. The crashed ingots were ball-milled into powders with an average size of ~5.5
µm [24]. Then, 1 wt.% Cu fine particles with an average size of ~2 µm prepared by inert gas
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condensation technique were alloyed into the Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 ball milled powders. The mixed powders were pressed under a pressure of 147 MPa in a magnetic field of 20 kOe,
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and then isostatically compacted under a press of 220 MPa for 30 s. The green bodies were sintered at ~1500 K for ~1h under an argon atmosphere, followed by a homogenization treatment at ~1400 K for ~3.5 h. The magnets were isothermally aged at ~1100 K for ~40 h, then slowly cooled down to ~700 K, and was aged again at ~700 K for about ~10 h [25]. Hereafter, 1 wt.% Cu fine particle alloyed Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 magnet is abbreviated to the alloyed
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magnet. To investigate the effects of the Cu-particle-alloying technique, pure magnets were
2.2. Magnetic hysteresis loops at different temperatures
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fabricated following the same procedure described above without Cu particles for comparison.
The magnetic properties of the pure and alloyed magnets were measured with a Quantum Design
2.3. Microstructural characterization
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magnetic field of up to 50 kOe at room temperature.
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physical property measurement system-vibrating sample magnetometer (PPMS-VSM) under a
The pure and alloyed magnets for scanning electron microscopy-electron backscatter diffraction (SEM-EBSD) analysis were mounted in epoxy and then mechanically polished using 220-grit, 500-grit, 1200-grit SiC papers, and standard colloidal silica suspension successively. SEM-
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EBSD mapping was carried out in a Zeiss Ultra Plus field-emission SEM equipped with an Oxford Instruments Aztec and Nordlys-nano EBSD detector. The Kikuchi patterns were recorded with an acceleration voltage of 25 kV, an aperture size of 120 µm, and a ‘high current’
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mode. The results were processed by the standard EBSD processing software. In addition, energy-dispersive X-ray spectroscopy (EDS) mapping was also obtained for revealing element
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distribution simultaneously.
The pure and alloyed magnets were cut into a small slice with the thickness of ~1 mm, polished and thinned down by 1500-grit SiC paper to ~50 µm. The samples were then attached to the copper grid and further thinned down to tens of nanometers by precision ion polishing system
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(PIPS). Transmission electron microscopy (TEM) experiments were conducted by a JEOL JEM2100F.
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The pure and alloyed magnets were milled down to a diameter of tens of nm using the tripod polishing method followed by focused ion beam (FIB) milling for atom probe tomography (APT) experiments [26]. APT measurements were carried out in a picosecond-pulse UV laser-assisted
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CAMECA local electrode atom probe (LEAP) 4000X Si. APT experiments were performed in a high vacuum, a cryogenic temperature of 50 K, under an intense electric field generated by a DC
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voltage up to 10000 V, with a 355 nm wavelength laser with pulse energy of ~50-100 pJ and a pulse frequency of ~100-200 kHz. The tomographic reconstructions were performed using Cameca’s Integrated Visualization & Analysis Software (IVAS) version 3.6.6 [27, 28]. Note that the morphology of reconstructed tips is influenced by user-defined reconstruction parameters,
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including but not limited to the detector efficiency, image compression factor (ICF) and field factor (kf), so absolute size measurement must be treated carefully [28]. If sufficient crystallographic information is present within the reconstruction process, calibration procedures
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may be used, but this was not possible in the collected APT datasets. Therefore, the default values of detector efficiency (0.50), ICF (1.650) and kf (3.30) were applied. The shank angle was
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selected as 7.5o considering the similarity between lengths of reconstructed APT datasets and actual lengths of evaporated sections of tips. The iso-concentration surfaces were defined with a voxel size of 1 nm and a delocalization value of x=3 nm, y=3 nm, z=1.5 nm. The error bars shown in 1D concentration profiles were calculated as E=(Ci(1-Ci)/N)1/2, where Ci=Ni/N, Ni represents the number of i solute ions/atoms and N represents the total number of counts with the given bin [28]. The bin width was selected as 0.2 nm in the 1D concentration profiles.
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2.4. Magnetic domain observation Magnetic force microscopy (MFM) measurements were conducted with a Bruker Dimension
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Icon (Santa Barbara, CA, USA). MFM is a secondary imaging mode of atomic force microscopy (AFM) using a two-pass technique. The initial pass measures the surface topography using tapping mode, then cantilever was lifted at a certain distance (~40 nm) above the sample surface
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in the second pass, thereby measuring the long-range magnetic forces between the probe and sample. Magnetic probes MESP-V2 (from Bruker AFM probes) were used to perform the
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magnetic measurements. 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 was recorded via changes in the phase and amplitude of the probe oscillation. They were further processed in the
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Gwyddion software to acquire two/three-dimensional (2/3D) MFM images.
2.5. Micromagnetic simulations
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Sandwich models (2:17R matrix phase/1:5 phase/2:17R matrix phase) with 80 nm × 30 nm × 30 nm in size were used to simulate the demagnetization curves of the pure and alloyed magnets.
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The thickness of the interfaces between the 1:5 phases and 2:17R matrix phases was set to 1 nm. Since simulated Hci values are almost identical no matter whether the magnetic domain wall is introduced in 2:17R matrix phases or 1:5 phases, the magnetic domain wall was introduced in the right 2:17R matrix phase in the model. The initial magnetization directions of the left 2:17R matrix phase, the 1:5 phase and the region left to the domain wall of the right 2:17R matrix phase were set upward, while the initial magnetization direction of the region right to the domain wall
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of the right 2:17R matrix phase was set downward. The external field was applied downwards from 0 to 40 kOe. The models were discretized by cubic nodes with a size of 1 nm, which is smaller than the domain wall width (δW=π(A/K1)1/2=~3.84-7.35 nm) and exchange length
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(δW=(A/µ 0MS2)1/2=~3.37-10.90 nm). The Landau-Lifshitz-Gilbert equation at each node was calculated by the 3D NIST Object Oriented MicroMagnetic Framework (OOMMF) software [29]. The exchange stiffness constant (A), magnetocrystalline anisotropy (K1) and saturation
3. Results and Discussion 3.1. Macromagnetic properties
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and Fe were calculated and summarized in Table 3.
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magnetization (MS) of 2:17R matrix phases and the 1:5 phases with various compositions of Cu
We measured the initial magnetization curves and hysteresis loops of the pure and Cu-particle-
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alloyed magnets. The initial magnetization curves manifest that the domain wall motion and domain wall pinning exist in both samples, but the alloyed magnet exhibits a larger pinning field (Fig. 1). Table 1 summarizes the corresponding magnetic properties of the pure and alloyed
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magnets. The Hci has been dramatically enhanced from 14.9 kOe in the pure magnet to 26.9 kOe in the alloyed magnet with ~1.8% reduction of the Br (from 11.4 kG to 11.2 kG) and (BH)max
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(from 26.4 MGOe to 26.6 MGOe). The Hci of the alloyed magnet is ~35% higher than that of the Sm(CobalFevCu0.056Zr0.020)7.546 sintered magnets while maintaining a comparably high (BH)max in Sm-Co based permanent magnets [17]. In addition, the squareness (Sr) of hysteresis curves has been improved from 81.2% to 85.8% by the Cu-particle-alloying technique. To summarize, we have achieved Sm(Co,M)z permanent magnets with exceptionally high overall magnetic performance (OMP) of Hci (kOe)+(BH)max (MGOe)=53.5, which has surpassed most of
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conventional Sm-Co based permanent magnets, as reported in the literatures [2, 3, 13, 16, 17, 30,
3.2. Microstructural observation by SEM and TEM
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31].
To fully unveil the origin of such substantial improvement of magnetic properties, global microstructure
and
microchemistry
of
pure
and
Cu-particle-alloyed
magnets
were
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comprehensively investigated by SEM-EBSD/EDS technique. Fig. 2 shows the global microstructure and elemental distributions of Sm, Co, Fe, O, Zr, and Cu of the pure and alloyed
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magnets. The bright and dark regions correspond to the Sm oxides and matrix grains (a combination of the 2:17R phases, 1:5 phases, and Zr-rich phases), respectively (Fig. 2a and b). Fig. 2c and d show the matrix grains with various grain shapes, crystallographic orientations, and grain sizes ranging from ~10 µm to ~30 µm of the pure and alloyed magnets. In addition, some
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large Zr-rich pockets were observed in both samples as well, in agreement with previous work done by Yosuke Horiuchi et al.[32]. Table 2 exhibits that the Cu concentration of matrix grains increases from ~5.20% in the pure magnet to ~5.82 at.% in the alloyed magnet, indicating that
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the additive Cu atoms were diffused into the 1:5 cell boundary phases through grain boundaries
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(GBs) in the alloyed magnet since Cu is not soluble in the 2:17R matrix phases.
We further investigated the constituent phases of pure and Cu-particle-alloyed magnets by using TEM. The typical TEM images show the cellular 2:17R matrix phases and 1:5 cell boundary phases of the pure and alloyed magnets at the nanoscale (Fig. 3a and b). The cellular sizes of 2:17R matrix phases are comparably similar (~100-150 nm). However, some discontinuous 1:5 cell boundary phases were observed in the pure magnet marked by white arrows (Fig. 3a). In the
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alloyed magnet, we found the improved continuity of 1:5 cell boundary phases (Fig. 3b), which is believed to contribute to the Hci enhancement [10]. Furthermore, the steady formation of the 1:5 cell boundary phases and cellular 2:17R matrix phases in the alloyed magnet leads to a high
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Br and a high (BH)max as well [10].
3.3. Distribution of individual atoms and compositions of the constituent phases
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We utilized APT to investigate the 3D distributions of various atom species of the cellular 2:17R matrix phases and 1:5 cell boundary phases of the pure and alloyed magnets at the atomic scale.
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To assure the reliability, we performed two separate APT experiments for each sample (One of them is shown in the main context and the other one is exhibited in the Supplemental Materials). Fig. 4 shows the 3D atom distribution maps of Fe, Sm, Zr, and Cu in the pure and alloyed magnets at the atomic scale, respectively. Iso-concentration surfaces of the Cu atoms (12.0 at.%
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for both pure and alloyed magnets) reveal the interfaces between the 1:5 cell boundary phases and cellular 2:17R matrix phases and iso-concentration surfaces of Zr atoms (1.0 at.% for the pure magnet/3.7 at.% for the alloyed magnet) illustrate the interfaces between the cellular 2:17R
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matrix phases and Zr-rich phases.
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As a representative, 1D concentration profiles from the black box (Fig. 4a to c) reveals that Cu and Sm are enriched in the ~6-nm-wide 1:5 cell boundary phase in the pure magnet, while Fe and Co are depleted. The Cu (Fe) concentrations are as high (low) as ~30 at.% (~12 at.%) at the maximum (minimum) value in the 1:5 cell boundary phase in the pure magnet. 1D concentration profiles from the black box (Fig. 4d to f) reveal that Cu and Sm are also enriched in the ~6-nmwide 1:5 cell boundary phase, while Fe and Co are depleted as well. However, the Cu (Fe)
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concentrations are as high (low) as ~43 at.% (~9 at.%) at the maximum (minimum) value in the 1:5 cell boundary phase in the alloyed magnet. Owing to the inhomogeneous distribution of these species in/across the 1:5 cell boundary phases, around twenty different regions were selected for
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identifying the distribution of maximum Cu concentration of the 1:5 cell boundary phases on the basis of two APT datasets of the pure and alloyed magnets, respectively (Supplemental Text, Fig. S1-S4, and Table S1-S2). Fig. 4g exhibits the corresponding probability distribution of the
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maximum Cu concentration of the 1:5 cell boundary phases in the pure and alloyed magnets. It should be noted that (1) the maximum Cu concentration of the 1:5 cell boundary phases ranges
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from ~18-38 at.% in the pure magnet, and ~26-50 at.% in the alloyed magnet, respectively, which lessens the Sr of the pure and alloyed magnets; (2) given that the Hci is a manifestation of the pinning strength of individual 1:5 cell boundary phases and represents the state in which the positive and negative magnetization of the Sm-Co magnets are equivalent to each other, the Hci is
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approximately equal to the pinning strength of the 1:5 cell boundary phases whose maximum Cu concentration is the value (~30 at.% for pure magnet and ~43 at.% for the alloyed magnet) in
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which the cumulative probability is 0.5.
3.4. Parameter estimation for micromagnetic simulations
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We bridged the aforementioned experimental observations and theoretical simulations by estimating the saturation magnetization (MS), magnetocrystalline anisotropy constant (K1), and exchange stiffness (A) of the cellular 2:17R matrix phases and 1:5 cell boundary phases in the pure and alloyed magnets on the basis of the APT results via polynomial fitting methods.
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Here,
we
selected
Sm10.55Co58.50Fe26.71Cu1.29Zr2.88Other0.07
and
Sm10.51Co57.82Fe27.52Cu1.52Zr1.61Other1.02 as the compositions of cellular 2:17R matrix phases in pure and alloyed magnets, respectively. To simplify compositions of cellular 2:17R matrix
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phases in pure and alloyed magnets, Cu and Zr atoms were neglected due to their little contribution to the magnetic properties of cellular 2:17R matrix phases. Under the circumstance, the compositions of cellular 2:17R matrix phases were determined to be Sm2(Co0.69Fe0.31)17 for
pure
and
alloyed
magnets,
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the pure magnet and Sm2(Co0.68Fe0.32)17 for the alloyed magnet, respectively. To simulate Hci of Sm16.89Co42.92Fe8.51Cu30.61Zr0.50Other0.57
and
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Sm16.41Co31.93Fe5.91Cu43.45Zr0.41Other1.89 were selected as compositions of 1:5 cell boundary phases in pure and alloyed magnets in which the cumulative probability is 0.5. To simplify the compositions of 1:5 cell boundary phases in the pure and alloyed magnets, Zr atoms were neglected due to its little contribution to magnetic properties of 1:5 cell boundary phases. Under
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the circumstance, the compositions of 1:5 cell boundary phases were determined to be Sm(Co0.52Fe0.11Cu0.37)5 and Sm(Co0.39Fe0.07Cu0.54)5, respectively.
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Firstly, the MS of cellular 2:17R matrix phases in the pure and alloyed magnets were polynomial fitting as follows [33, 34]:
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y = 12.60 + 13.99x − 62.71x2 + 164.41x3 − 183.17 x4 + 65.70x5
(2)
where x and y represent the concentration of Fe in Sm2Co17 phases and their corresponding MS, respectively. Therefore, the MS of Sm2(Co0.69Fe0.31)17 was determined to be 14.30 kG in the pure magnet. The MS of Sm2(Co0.68Fe0.32)17 was obtained as 14.34 kG in the alloyed magnet as well. If only Sm, Co, and Fe exist in 1:5 cell boundary phases, the MS of 1:5 cell boundary phases in the pure and alloyed magnets were polynomial fitting as follows [35]:
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y = 9.45 + 16.46 x − 41.15x2 + 163.41x3 − 229.49 x4 + 96.95x5
(3)
where x and y represent the concentration of Fe in SmCo5 phases and their corresponding MS,
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respectively. Therefore, the MS of 1:5 cell boundary phases (only considering Sm, Co, and Fe) in pure and alloyed magnets was determined to be 11.68 kG and 11.44 kG, respectively. However, the incorporation of Cu severely reduces the MS of Sm(Fe,Co)5 phases. The relation between the Cu content in the Sm(Fe,Co,Cu)5 phase (x) and the corresponding MS (y) was fitted with
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experimental datasets in [36]:
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y = 11.47 − 2.26x − 111.07 x2 + 304.14x3 − 255.94x4 − 23.67 x5
y = 11.25 − 2.92x − 117.07 x2 + 380.07 x3 − 464.70x4 + 152.05x5
(4) (5)
Therefore, the MS of Sm(Co0.52Fe0.11Cu0.37)5 was determined to be 5.87 kG in the pure magnet.
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The MS of Sm(Co0.39Fe0.07Cu0.54)5 was obtained as 2.85 kG in the alloyed magnet as well.
Secondly, the exchange stiffness (A) of cellular 2:17R matrix phases in the pure and alloyed magnets were calculated. Curie Temperature (TC) of cellular 2:17R matrix phases in the pure and
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alloyed magnets were polynomial fitting as follows [34]: y = 1189.00 + 246.59x − 2584.71x2 + 5108.61x3 − 5592.88x4 + 2023.60x5
(6)
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where x and y represent the concentration of Fe in Sm2Co17 phases and their corresponding TC, respectively. Therefore, the TC of Sm2(Co0.69Fe0.31)17 phase in the pure magnet was calculated as ~1123 K. The TC of Sm2(Co0.68Fe0.32)17 phase in the alloyed magnet was calculated as ~1119 K. The relationship between the A and TC can be shown as [37]: A=
3k BTC 2 ZS ( S + 1)
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(7)
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where kB, Z, and S correspond to Boltzmann constant, the coordination number, spin quantum number of the material. Here, kB, Z, and S were assumed nearly constant in the same system even with little chemical variation. Therefore,
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A ∝ TC
(8)
The A of Sm2Co17 was reported as 19.60 pJ/m [2]. According to the equation (8), the A of Sm2(Co0.69Fe0.31)17 phase in the pure magnet was determined to be 18.50 pJ/m. The A of
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Sm2(Co0.68Fe0.32)17 phase in the alloyed magnet was calculated as 18.45 pJ/m.
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If only Sm, Co, and Fe exist in 1:5 cell boundary phases, TC remains nearly constant with Fe concentration lower than 40 at.% [38]. Therefore the A of 1:5 cell boundary phases (only considering Sm, Co, and Fe) are the same as that of pure SmCo5 phases, which was reported as 15.10 pJ/m [2]. However, the incorporation of Cu severely reduces TC of Sm(Fe,Co)5 phases.
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The relationship between the Cu content in the Sm(Fe,Co,Cu)5 phase (x) and the corresponding TC (y) was fitted with experimental datasets in [36]: y = 977.00 − 536.97 x − 2643.65x2 + 10410.68x3 − 17622.54x4 + 9584.99x5
(9)
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Therefore, the TC of Sm(Co0.52Fe0.11Cu0.37)5 phase was determined to be 679.93 K in the pure
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magnet. The TC of Sm(Co0.39Fe0.07Cu0.54)5 phase was obtained as 497.11 K in the alloyed magnet. Thus, the A of Sm(Co0.52Fe0.11Cu0.37)5 phase was determined to be 10.51 pJ/m in the pure magnet. The A of Sm(Co0.39Fe0.07Cu0.54)5 phase was obtained as 7.68 pJ/m in the alloyed magnet.
Finally, the magnetocrystalline anisotropy constant (K1) of cellular 2:17R matrix phases in pure and alloyed magnets were polynomial fitting as follows [39]: y = 4.99 + 9.88x − 46.28x2 − 239.80 x3 + 1173.02 x4 − 1336.83x5
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(10)
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where x and y represent the concentration of Fe in Sm2Co17 phases and their corresponding K1, respectively. Therefore, the K1 of Sm2(Co0.69Fe0.31)17 phase in the pure magnet and Sm2(Co0.68Fe0.32)17 phase in the alloyed magnet was calculated as 3.47 MJ/m3 and 3.37 MJ/m3,
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respectively.
The K1 of 1:5 cell boundary phases remains challenging to determine due to various values of
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anisotropy field HA (317 kOe and 415 kOe) of SmCo5 phases reported by researchers [40, 41]. Therefore, the average value of HA was selected, which is 366 kOe. If only Sm, Co, and Fe exist
were polynomial fitting as follows [40]:
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in 1:5 cell boundary phases, the HA of 1:5 cell boundary phases in the pure and alloyed magnets
y = 373.92 − 210.30x + 3674.83x2 − 12547.92x3 + 13932.20x4 − 5082.05x5
(11)
where x and y represent the concentration of Fe in Sm2Co17 phases and their corresponding HA,
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respectively. Therefore, the HA of Sm(Co0.52Fe0.11Cu0.37)5 phase was determined to be 393.62 kOe in the pure magnet. The HA of Sm(Co0.39Fe0.07Cu0.54)5 phase was obtained as 389.38 kOe in
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the alloyed magnet as well.
If only Sm, Co, and Fe exist in 1:5 cell boundary phases, the MS of 1:5 cell boundary phases in
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pure and alloyed magnets are 11.68 kG and 11.44 kG, respectively. The relation between the HA and K1 can be shown as [37]: K1 =
H AM S 2
(12)
Therefore, the K1 of 1:5 cell boundary phases in pure and alloyed magnets were calculated as 18.30 and 17.72 MJ/m3, respectively.
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However, the incorporation of Cu severely reduces K1 of 1:5 cell boundary phases in pure and alloyed magnets. The relation between the Cu content in the Sm(Fe,Co,Cu)5 phase (x) and the
y=−
a 0.6
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corresponding K1 (y) was fitted based on experimental datasets in [36]: x+a
(13)
where a corresponds to the K1 without any Cu incorporation. In the pure magnet, a=18.30
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(MJ/m3), and x=0.37. Therefore, the K1 of Sm(Co0.52Fe0.11Cu0.37)5 in the pure magnet was determined to be 7.02 MJ/m3. In the alloyed magnet, c=17.72 (MJ/m3), and x=0.54. Therefore,
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the K1 of 1:5 Sm(Co0.39Fe0.07Cu0.54)5 in the alloyed magnet was determined to be 1.77 MJ/m3. The calculated values are in good agreement with the results published by the Toshikazu et al.[40].
Here,
we
selected
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3.5. Micromagnetic simulations
Sm10.55Co58.50Fe26.71Cu1.29Zr2.88Other0.07
[Sm2(Co0.69Fe0.31)17]
and
Sm10.51Co57.82Fe27.52Cu1.52Zr1.61Other1.02 [Sm2(Co0.68Fe0.32)17] as the compositions of the cellular
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2:17R matrix phases in the pure and alloyed magnets, respectively. To simulate the Hci values of the pure and alloyed magnets, Sm16.89Co42.92Fe8.51Cu30.61Zr0.50Other0.57 [Sm(Co0.52Fe0.11Cu0.37)5]
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and Sm16.41Co31.93Fe5.91Cu43.45Zr0.41Other1.89 [Sm(Co0.39Fe0.07Cu0.54)5] were selected as the compositions of the 1:5 cell boundary phases in the pure and alloyed magnets owing to the previously mentioned correlation between the Hci and the individual pinning strength of the 1:5 cell boundary phases. Strikingly, we found that the repulsive domain wall pinning (γ2:17R < γ1:5) predominates in the pure magnet, whereas the attractive domain wall pinning (γ2:17R > γ1:5) in the alloyed magnet (Table 3), in consistency with previous observations by Lorentz TEM microscopy [22, 23]. In this scenario, there is no constraint on increasing the Cu concentration in 17
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the 1:5 phases, which further lowers the γ1:5, and the Hci can be enhanced substantially. More precisely, ∆γ between the cellular 2:17R matrix phases and 1:5 cell boundary phases in the alloyed magnet is 16.79 mJ/m2, which is around sevenfold larger than that between these phases
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in the pure magnet (~|-2.31| mJ/m2), which leads to the substantial Hci enhancement.
We performed micromagnetic simulations on the pure and alloyed magnets based on the
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parameters estimated from our experimental observations (Fig. 5a to d). Under the increasing external field, the magnetic domain wall moves leftwards and is pinned at the 1:5 cell boundary
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phase. Fig. 5b shows the simulated demagnetization curves of the pure and alloyed magnets, in which the plateaus in the curves results from the domain wall pinning. Our micromagnetic simulation results show that the Hci is enhanced by ~18.0 kOe from the pure magnet (~9.0 kOe) to the alloyed magnet (~27.0 kOe), in excellent agreement with our macromagnetic properties
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(Fig. 1). In the pure magnet, the magnetic domain wall moves in the right 2:17R matrix phase towards the 1:5 cell boundary phase under a small external field of 1 kOe. With the increasing external field, the center of magnetic domain wall is repulsively pinned from the largest domain
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wall energy gradient at the interface between the right 2:17 phase and the 1:5 cell boundary phase. When the external field is up to 9.0 kOe, the domain wall is repulsively depinned from the
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1:5 cell boundary phase and moved leftwards continuously into the 1:5 cell boundary phase and then the left 2:17R matrix phase (Fig. 5d). Similarly, the domain wall moves in the right 2:17R matrix phase towards the 1:5 cell boundary phase under a small external field of 1.0 kOe in the alloyed magnet. However, the domain wall moves across the 1:5 cell boundary phase and is attractively pinned at the largest domain wall energy gradient at the interface between the left 2:17R matrix phase and the 1:5 cell boundary phase with increasing external field, since the
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magnetic domain wall is easily trapped in the 1:5 cell boundary phase. When the external field is up to 27.0 kOe, the domain wall is attractively depinned from the 1:5 cell boundary phase and moved leftwards continuously into the left 2:17R matrix phase (Fig. 5d). The actual position of
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the domain wall is determined by the interplay between the domain wall energy and magnetostatic energy [37].
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We also quantified the effects of the domain wall energy gradient and the thickness of the 1:5 cell boundary phases on the Hci of the alloyed magnet. The smallest domain wall energy gradient
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only gives rise to a smaller Hci up to 20 kOe in the alloyed magnet (Fig. 5e and f). In contrast, the largest domain wall energy gradient results in a larger Hci of 27 kOe in the alloyed magnet, even with the same domain wall energy difference between the 2:17R matrix phase and 1:5 cell boundary phase (Fig. 5e and f). Moreover, the Hci of the alloyed magnet is reduced from 27 kOe
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to 20 kOe with the thickness of the 1:5 cell boundary phases reducing from 7 nm to 3 nm (Fig. 5g and h). because the domain wall may not be fully trapped into these 1:5 cell boundary phases if the domain wall width in the 1:5 cell boundary phases (~6.54 nm) is larger than the thickness
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of 1:5 cell boundary phases. Furthermore, the Hci of the alloyed magnet remains constant with a larger thickness of the 1:5 cell boundary phases (>7 nm) since the domain wall has already been
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fully trapped in 1:5 cell boundary phases (Fig. 5g and h). To summarize, the pinning strength of the 1:5 cell boundary phases in Sm-Co based permanent magnets is not only correlated to the domain wall energy difference of the cellular 2:17R matrix phases and 1:5 cell boundary phases, but also associated with the domain wall energy gradient and the thickness of the 1:5 cell boundary phases.
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3.6. Summary We illustrated the microstructural changes occurring in the pure and Cu-particle-alloyed magnets (Fig. 6). In the pure magnet, normal 1:7H matrix grains are formed after the sintering process
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and homogenization (Fig. 6a). Moreover, discontinuous 1:5 cell boundary phases and cellular 2:17R matrix phases are formed after the aging process and furnace cooling. Using the Cuparticle-alloying-technique (Fig. 6b), Cu-rich 1:7H matrix grains and GBs are formed after the
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sintering process and homogenization. In addition, continuous 1:5 cell boundary phases and cellular 2:17R matrix phases are formed after the aging process and furnace cooling. Most
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importantly, the Cu concentration in 1:5 cell boundary phases is increased significantly, which is not only driven by the thermal diffusion during the slow cooling process, but also the concentration gradient of residual Cu located near the GBs. This results in the attractive domain pinning instead of repulsive domain wall pinning in Cu-particle-alloyed Sm(Co,M)z permanent
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magnets.
In Cu-particle-alloyed Sm-Co magnets, not only the intrinsic coercivity but also the squareness is
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improved significantly. For the coercivity enhancement, the addition of Cu lowers the domain wall energy of 1:5 cell boundary phases, which enlarges the difference of domain wall difference
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between 1:5 cell boundary phases and 2:17 matrix phases and therefore leads to the enhancement of the domain wall pinning strength at 1:5 cell boundary phases.
For the improved squareness, the addition of Cu is likely to increase the fraction of 1:5 cell boundary phases near grain boundaries. In the pure Sm-Co magnets, the grain boundaries are Culean [32]. Most importantly, few 1:5 cell boundary phases are formed in the regions near these
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Cu-lean grain boundaries [32]. Furthermore, the concentration of Cu in these inchoate 1:5 cell boundary phases varies, which leads to the low squareness of pure Sm-Co magnets. In the Cuparticle-alloyed Sm-Co magnets, the concentration of Cu at grain boundaries is increased, which
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also facilitates the formation of 1:5 cell boundary phases near grain boundaries. Since the number of Cu-lean or inchoate 1:5 cell boundary phases near grain boundaries is reduced in the Cu-particle-alloyed Sm-Co magnet, the number of 1:5 cell boundary phases with low domain
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wall pining strength is reduced (In other words, the distribution of maximum Cu concentration in
particle-alloyed Sm-Co magnet.
4. Conclusions
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1:5 cell boundary phases becomes narrow), which leads to the improved squareness of Cu-
In conclusion, our work here demonstrates high iron content Sm(Co,Fe,Cu,Zr)z permanent
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magnets with a combination of a intrinsic coercivity of 26.9 kOe, a remanence of 11.2 kG, and a maximum energy product up to 26.6 MGOe simultaneously via the Cu-particle-alloying technique, which outperforms most of conventional Sm(Co,Fe,Cu,Zr)z permanent magnets. By
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integrating multiple cutting-edge microscopy techniques and micromagnetic simulations, we find that (1) attractive domain wall pinning, instead of widely accepted repulsive domain wall pinning,
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predominates in the Cu-particle-alloyed magnet, (2) the exceptionally high Hci is attributed to the continuity and enlarged pinning strength of the 1:5 cell boundary phases in the Cu-particlealloyed magnet. (3) High (BH)max results from the stable formation of the 1:5 cell boundary phases and cellular 2:17R matrix phases in high content Sm(Co,Fe,Cu,Zr)z permanent magnets. Our work not only demonstrates the transformation from repulsive pinning to attractive pinning by alloying Cu particles in Sm-Co based permanent magnets but also provides insights into the
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grain boundary engineering for enhancing the Hci of rare earth permanent magnets. Due to the versatility of the Cu-particle-alloying technique, further performance optimization is possible, for example, by tuning the Cu particle size to achieve uniform Cu diffusion from the grain
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boundaries to the 1:5 cell boundary phases, which will further improve the magnetic properties of Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 permanent magnets. Moreover, the methodology applied in this work may also offer exciting possibilities for quantitative analysis and prediction between
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compositions and magnetic properties of other magnetic materials.
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Acknowledgments
The authors would like to thank Ms. Katie Levick, Dr. Patrick Trimby, Dr. Anna Ceguerra, Dr. Takanori Sato, Mr. Adam Sikorski, and Dr. Jacob Warner for their technical support. The authors also would like to Dr. Yisheng Chen and Dr. Yahui Cheng for fruitful discussion. The
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authors would like to thank the Australian Microscopy and Microanalysis Research Facility (AMMRF) and Australian Centre for Microscopy and Microanalysis (ACMM). This work was
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partly supported by the Australian Research Council (DP150100018).
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[36] E. Lectard, C. Allibert, R. Ballou, Saturation magnetization and anisotropy fields in the Sm(Co1-xCux)5 phases, J. Appl. Phys. 75 (1994) 6277-6279. [37] J. M. Coey, Magnetism and magnetic materials, Cambridge University Press, 2010. [38] T. Saito, D. Nishio-Hamane, Magnetic properties of SmCo5-xFex (x=0-4) melt-spun ribbon, J. Alloys Compd. 585 (2014) 423-427. [39] H. Mildrum, M. S. Hartings, K. J. Strnat, J. Tront, Magnetic properties of the intermetallic phases Sm2(Co,Fe)17, AIP Conf. Proc. 10 (1973) 618-622. [40] T. Katayama, T. Shibata, Magneto-crystalline anisotropy constant in Sm(Co,Cu)5 base alloy, Jpn. J. Appl. Phys. 12 (1973) 762. [41] T.-S. Chin, C.-H. Lin, H. Bai, Y. Huang, K. Cheng, S. Huang, Effect of nitrogen content on magnetic properties of Sm2Fe17Nx and SmFe5Ny alloys, IEEE Trans. Magn. 28 (1992) 2587-2589.
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Table
1.
Magnetic
properties
of
the
pure
and
Cu-particle-alloyed
Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 permanent magnets. Hci (kOe)
11.4
14.9
1.0 (Cu particles)
8.4
11.2
26.9
(BH)max (MGOe) 26.4
Sr (%)
26.6
85.8
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81.2
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Br (kG)
0
Density (g/cm3) 8.4
t (wt. %)
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Table 2. Experimentally determined compositions of the matrix grains (a combination of
particle-alloyed magnets by SEM-EDS. Sample
Composition (at.%) Sm 11.32 11.74
Cu 5.20 5.82
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Fe 23.04 23.46
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Pure Cu particlealloyed
Co 58.62 57.10
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2:17R matrix phases, 1:5 cell boundary phases, and Zr-rich phases) in pure and Cu-
Zr 1.82 1.88
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Table 3. The parameter estimation of the cellular 2:17R matrix phases and 1:5 cell boundary phases in the pure and alloyed magnets. Phas e
Composition
K1 (MJ /m3)
MS (kG)
A (pJ/ m)
γw (mJ/m 2 )
Pure
2:17 R 1:5
Sm2(Co0.69Fe0.31)17
3.47
14.30 18.50
32.05
Sm(Co0.52Fe0.11Cu 0.37)5 Sm2(Co0.68Fe0.32)17
7.02
5.87
10.51
34.36
3.37
14.34 18.45
31.54
Sm(Co0.39Fe0.07Cu 0.54)5
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2:17 R 1:5
1.77
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Cuparti clealloy ed
2.85
Dom Pinnin ain g Type wall width (nm) 7.25 Repuls ive 3.84 Pinnin g 7.35 Attrac tive 6.54 Pinnin g
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Sam ple
7.68
14.75
Fig.
1.
The
magnetization
curves
of
pure
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Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 permanent magnets.
the
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and
Cu-particle-alloyed
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Fig. 2. The microstructure and microchemistry of the pure and Cu-particle-alloyed magnets. (a) Secondary electron (SE) image, (c) Backscattered electron (BSE) image, (e) The corresponding all Euler map, EDS maps of (g) Sm, (i) Co, (k) Fe, (m) Zr, (o) O and (q) Cu of the pure magnet.
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(b) SE image, (d) BSE image, (f) The corresponding all Euler map, EDS maps of (h) Sm, (j) Co,
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(l) Fe, (n) Zr, (p) O and (r) Cu of the doped magnet. The scale bars are 50 µm.
Fig.
3.
Bright
field
TEM
images
of
(a)
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pure
and
(b)
Cu-particle-alloyed
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Sm(Co0.665Fe0.25Cu0.06Zr0.025)7 permanent magnets. c axis is out of the plane. The scale bars are
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100 nm.
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FIG. 4. APT data of the pure and Cu-particle-alloyed magnets at the atomic scale. The pure magnet: (a) 3D atom distribution maps of Fe, Sm, Cu, Zr, iso-concentration surfaces of Cu (12.0 at.%) and Zr (1.0 at.%). (319 nm × 71 nm × 71 nm). Movie S1 shows the 3D rotation of Fe, Sm,
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Cu and Zr atoms. (b) A rotational view of (a) (without Zr iso-concentration surfaces). (c) 1D concentration profiles of Co, Fe, Sm, Cu, and Zr of the black box in( b). The alloyed magnet: (d), 3D atom distribution maps of Fe, Sm, Cu, Zr, iso-concentration surfaces of Cu (12.0 at.%) and
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Zr (3.7 at.%). (301 nm × 68 nm × 68 nm). Movie S2 shows the 3D rotation of Fe, Sm, Cu and Zr atoms. (e) A rotational view of d) (without Zr iso-concentration surfaces). (f) 1D concentration
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profiles of Co, Fe, Sm, Cu, and Zr of the black box in (e). (g) The probability distribution of the maximum Cu concentration of 1:5 cell boundary phases in the pure and Cu-particle-alloyed
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magnets based on around twenty atom probe tomography datasets for each sample.
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Fig. 5. Micromagnetic simulation results of the pure and Cu-particle-alloyed magnets. (a) The parameters used for simulating the demagnetization curves of the pure and alloyed magnets. (b) The simulated demagnetization curves of the pure and alloyed magnets. (c) Schematic diagram
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of the initial state of the micromagnetic model containing two 2:17R matrix phases and one 1:5 cell boundary phase. (d) The simulated repulsively and attractively depinned situations of the pure and alloyed magnets. (e) The parameters used for simulating the effects of the domain wall
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energy gradient on the Hci of the alloyed magnet. (f) The simulated demagnetization curves of the alloyed magnet with various domain wall energy gradients. (g) The parameters used for simulating the effects of the thickness of the 1:5 cell boundary phases on the Hci of the alloyed
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thicknesses of the 1:5 cell boundary phases.
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magnet. (h) The simulated demagnetization curves of the alloyed magnet with various
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Fig. 6. Schematic illustration of the microstructural evolution of the pure and Cu-particle-alloyed
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magnets. (a) The microstructure characteristics of the pure magnet after mixing the matrix powers, after the sintering process and homogenization, and after the aging process and furnace cooling. (b) The microstructure characteristics of the alloyed magnet after mixing the matrix
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powers and Cu particles, after the sintering process and homogenization, and after the aging process and furnace cooling. The scale bars in the left and middle columns are 5 µm. The scale
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