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nanocrystalline hybrid materials
Powder Technology 339 (2018) 440–445
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Soft magnetic properties of Fe-based amorphous/nanocrystalline hybrid materials Yeonjoo Lee a, Jonggyu Jeon a, Seungjin Nam a, Teasuk Jang b, Hwijun Kim c, Minwoo Lee c, Yongjin Kim d, Dongyeol Yang d, Kyeongsik Min e, Hyunjoo Choi a,⁎ a
School of Advanced Materials Engineering, Kookmin University, Seoul, Republic of Korea Dept. of Advanced Materials Engineering, Sunmoon University, Asan, Republic of Korea Incheon Regional Division, Korea Institute of Industrial Technology, Incheon, Republic of Korea d Powder & Ceramics Division, Korea Institute of Materials Science, Changwon, Republic of Korea e School of Electrical Engineering, Kookmin University, Seoul, Republic of Korea b c
a r t i c l e
i n f o
Article history: Received 13 February 2018 Received in revised form 29 June 2018 Accepted 9 August 2018 Available online 11 August 2018 Keywords: Fe-based amorphous powder Ball milling Crystallization Coercivity Magnetization
a b s t r a c t In this study, we investigated the effect of crystalline structures on the magnetic properties of the Fe-based amorphous powder. We intentionally produced uniformly dispersed nanocrystallites inside the Fe-based amorphous powder using a combined process of mechanical milling and thermal annealing. In the case of powder produced after 30 min of ball milling followed by annealing at 400 °C, coercivity was decreased by 96.88% from 3698.73 to 115.49 A/m, and magnetization was increased 2.60% from 1.92 × 10−4 to 1.97 × 10−4 Wbm/kg. The simultaneous improvement of coercivity and magnetization can be attributed to the homogeneous generation of a small number of nanocrystallites as well as the release of residual stress. © 2018 Elsevier B.V. All rights reserved.
1. Introduction Soft magnetic materials have become of paramount importance in a variety of industrial fields such as electricity-generating facilities [1], power-converting parts [2], electromagnetic wave shielding materials [3], and magnetic sensors [4, 5]. As electronic devices become miniaturized and automated, the development of high-performance soft magnetic materials with good reliability is becoming urgent. In general, if the crystal size of crystalline metals is less than the exchange correlation length, a decrease in crystal size results in a decrease in coercivity and an increase in permeability, meaning an improvement in soft magnetic properties [6, 7]. Therefore, minimum coercivity can be obtained in an amorphous state. Furthermore, the disordered atomic arrangement of amorphous materials can provide excellent mechanical properties and maintain good magnetic properties even at high frequencies [6, 8]. In this regard, Fe-based amorphous materials have attracted increasing attention because of their good soft magnetic properties such as high saturation magnetization, high permeability, low coercivity, and low core loss [6]. Recently, efforts have been made to produce these amorphous materials in powder form because, differently from bulk amorphous ⁎ Corresponding author. E-mail address: [email protected] (H. Choi).
https://doi.org/10.1016/j.powtec.2018.08.037 0032-5910/© 2018 Elsevier B.V. All rights reserved.
materials, amorphous powder enables the production of parts with complex morphologies and three-dimensionally uniform magnetic properties [9, 10]. In previous work, the amorphous powder was generally produced by crushing a melt-spun ribbon [11–15], but in this case, low filling rate, high manufacturing cost, and characteristic deterioration have restricted the commercialization of the powder. Mechanical alloying has also been suggested to produce amorphous powders, whereby monolithic elemental powders are ball milled in a chamber to eventually generate an atomically mixed amorphous powder [16]. However, this process requires a long milling duration (generally longer than 48 h), which is not practical for industrial manufacturing processes. Recently, we have reported a way to produce Fe-Hf-B-Nb-P-C amorphous powder directly using gas atomization [17]. The powder revealed good soft magnetic properties such as good saturation magnetization higher than 149.7 emu/g and low coercivity below 12.8 Oe. On the other hand, the introduction of nanocrystallites into amorphous materials has been reported to reduce iron loss by preventing the generation of eddy currents at high frequencies [10]. Recently, F. Hosseini-Nasb et al. has generated nanocrystallites homogeneously in a melt-spun amorphous ribbon by controlling the wheel speed and this amorphous/nanocrystalline hybrid ribbon revealed a simultaneous improvement of saturation magnetization and coercivity [18]. An attempt has also been made to produce nanocrystallites in amorphous powder by inducing mechanical [6] or thermal [19, 20] energy into
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the powder [8, 21, 22]. However, the size and amount of the nanocrystallites generated during the process have yet to be clearly investigated. Furthermore, the simultaneous improvement of saturation magnetization and coercivity described in melt-spun ribbons has not yet been reported in amorphous powders, which may be because the crystallites formed mostly near the powder surface rather than being uniformly distributed inside the powder interior. In this study, we employed a combined treatment to introduce thermal energy after inducing mechanical energy into a gas-atomized Fe-based amorphous powder in order to generate nanocrystallites uniformly inside the powder interior. Mechanical milling may generate a number of defects inside the powder, which can act as homogeneous nucleation sites for crystallization during thermal annealing. We investigated the effect of the mechanical milling and thermal annealing processes on the microstructure and soft magnetic properties of the Fe-based amorphous powder. In general, amorphous powder is generally mixed with binders to obtain a bulk-type product because of difficulties in plastic deformation of the amorphous powder. This process does not require high temperatures or high pressures, and hence, it does not involve any significant microstructural changes (grain growth or oxidation) [23]. Therefore, the soft-magnetic properties of the products are determined in terms of the packing fraction and the inherent soft-magnetic properties of amorphous powder. Hence, in this study, the soft magnetic properties were measured in the powder state. 2. Experimental procedures 2.1. Sample preparation The Fe-B-P-C-Nb-Hf amorphous powder was produced by gas-atomization. A mixture of pure Fe, B, P, C, Nb, and Hf (≥ 99.5% purity) ingots was melted in a magnesia crucible set in a high-frequency induction furnace at 1400–1600 °C. The cast master alloy was extruded into a bar shape to easily insert and re-melt it in the gas-atomization equipment. After that, the master alloy was dissolved at 1300–1330 °C in a graphite crucible, and the melt was atomized to produce an amorphous powder by passing through a 2 mm nozzle in jets of He and N2 gases under 5–7.5 MPa pressure. The Fe-based amorphous powder was partially crystallized by mechanical milling and annealing processes. The powder was mechanically milled in a planetary mill (Pulverisette 5 classic line, Fritsch ™, Germany) at room temperature using a 500 mL stainless
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bowl. Powder was poured into the bowl with 200 mL ethanol and 5 mm-diameter stainless steel balls at a powder-to-ball weight ratio of 1:20. The milling was carried out at a constant rotating speed of 100 rpm for 30 min. Afterwards, as-atomized or ball-milled powders were annealed at 400 °C using a tube furnace (PTF-15/100/610, Lenton, Inc.) in argon; samples were heated to the target temperature at a heating rate of 10 °C/min, maintained for 1 h, and then furnace-cooled to room temperature. 2.2. Characterization The lattice structure of the Fe-based amorphous/nanocrystalline hybrid powders was characterized by X-ray diffraction (XRD) with a CN2301 (Rigaku, Japan) using CuKα (λ = 1.5405 Å). Samples were scanned over an angular range of 30–60° with a step size of 0.05° and a scan speed of 0.02°/min. During the XRD analysis, the sample was rotated to observe the general lattice. At that time, the diffractometer was set at 40 kV working voltage and 200 mA working current with a 2 mm slit size. The degree of crystallinity was examined by differential scanning calorimetry (DSC) with a TGA/DSC 1 (Mettler Toledo, USA) in an argon atmosphere. The sample was heated to 600 °C at a heating rate of 30 °C/min, and was then cooled to room temperature using circulated water. The same amount of powder used in this analysis was (300 ± 10 mg) for all samples to obtain comparable results. The morphology of the powders was observed using scanning electron microscopy (SEM) with a JEM-7610F (JEOL Ltd., Japan). Prior to the SEM analysis, the powders were coated with platinum (Pt) and attached using carbon tape to improve imaging. The microstructure of the Fe-based amorphous powder and amorphous/nanocrystalline hybrid powders were observed using transmission electron microscopy (TEM) with a Tecnai F20 G2 (FEI, USA) operating at 200 kV; TEM specimens were prepared using a focused ion beam technique. Fast Fourier transform (FFT) patterns and an inverse FFT (IFFT) images were obtained using Gatan digital microscope software for detailed structure analysis. Soft magnetic properties (i.e. saturation magnetization and coercivity) of the prepared powders were measured using a vibration sample magnetometer (VSM; Lake Shore 7400) with a maximum applied field of 10 kG and a B\\H loop tracer at room temperature. For all samples, 0.07 ± 0.01 g powder was loaded in a quartz crucible for the test, and prior to the test, the place of sample holder was adjusted using a standard Ni specimen.
Fig. 1. SEM images of powders after ball milling for (a) 0 h (as-atomized), (b) 15 min, (c) 30 min, (d) 1 h, (e) 15 h, and (f) 72 h.
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3. Results and discussion Fig. 1(a-f) shows the morphology of the powders after ball milling for a variety of time periods. As-atomized powder exhibited a spherical shape with a mean diameter of 10.88 μm, which became flattened as ball milling proceeded. In general, the initial impact of the grinding ball induced the soft crystalline metal powder to flatten and become work-hardened. With continued milling, the work-hardened powder tended to fracture and become refined. Although the starting material was in an amorphous state in this study, the morphological evolution of the powder during milling was similar to that of crystalline metal powders. Although it proceeded much more slowly, the amorphous powder started to become flattened after 15 h of milling whereas crystalline metal powders generally became flattened within 3 h under similar milling conditions [24]. This demonstrate that, similar to crystalline metal powders, the amorphous powder also experienced significant plastic deformation during milling [25]. Unlike bulk amorphous materials, a number of defects were generated from the large specific surface in the powder state, and these defects facilitated the plastic deformation [26]. Temperature rise during milling might also have helped plastic deformation of the amorphous powder as N90% of the mechanical energy imparted to the powders during milling was transformed into heat [27]. In addition, a certain proportion of the amorphous phase was crystallized during the process, which also contributed to the plastic deformation [6]. Therefore, in this study, the milling duration was fixed at 30 min to avoid significant crystallization and morphological changes. Fig. 2 shows XRD patterns of as-atomized powder and powders annealed and/or ball milled under different condition. Regardless of ball milling time and annealing temperature, all samples exhibited a typical amorphous pattern at 2θ ranging from 40° to 50° [6, 28]. Crystallization of the amorphous powder was not thought to have notably occurred during a short time period with low-energy milling combined with annealing at 400 °C (81.87% of crystallization temperature of the amorphous powder (i.e. 488.59 °C)). A peak indicating crystalline α-Fe phases was observed at 44.5° in 2θ when the powder was annealed at temperatures higher than 400 °C. This result is comparable to previous reports in that crystallization of the Fe-based amorphous alloy resulted in the formation of the α-Fe phase followed by the Fe2B phase [29]. On the other hand, ball milling was found to stimulate the formation of the Fe2B phase (the peak detected at 42.5° in 2θ). Compared with general heat treatment, it has been found that ball milling of Fe-based powders can accelerate crystallization of this phase due to excessive generation of vacancy-type defects [6]. According to previous research, when Fe81B13.5Si13.5C2 was heat-treated at a heating rate of 5 °C/min, the
Fig. 2. XRD patterns of the amorphous powder: as-atomized, after annealing at 400 °C, after 30 min ball milling, and after 30 min ball milling and annealing at 400 °C, in ascending order.
Fig. 3. (a) DSC traces and (b) degree of crystallinity for the amorphous powder: asatomized, after annealing at 400 °C, after 30 min ball milling, and after 30 min ball milling and annealing at 400 °C.
Fig. 4. TEM images of the amorphous powder: (a) as-atomized, (b) after annealing at 425 °C, (c) after 1-h ball milling and (d) after 1-h ball milling and annealing at 425 °C The insets in the top right-hand corners of (a)-(d) are selected area diffraction (SAD) patterns.
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Fig. 5. Schematic of the crystallization of the amorphous powder via (a) thermal processing and (b) thermal processing after mechanical processing.
α-Fe phase was formed at 485 °C, the Fe3B phase was formed at 515 °C, and the Fe3C phase was formed at 630 °C [29, 30]. In this study, a small amount of α-Fe and Fe2B phases were formed at a relatively low temperature of 425 °C, much lower than that in previous reports [31]. In addition, the crystal size was found to be very small based on the broad XRD spectrum [6], the significance of which is discussed along with the TEM observations later. All of the samples also revealed two peaks at 33.8 and 34.3°, which indicate the presence of an Fe3O4 layer possibly on the powder surface. The intensity of the peaks did not significantly vary as milling and heating progressed. When as-atomized powder was
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annealed at 425 °C, the thickness of the oxide layer was found to increase from 70 to 90 nm according to the SEM analysis in a previous study [32]. The growth rate of the Fe3O4 layer of Fe-based amorphous powder is much slower than that of pure iron [32, 33] because diffusion of iron atoms in the amorphous state of a powder is much suppressed compared to in a crystalline state. Furthermore, the presence of Nb and P in this amorphous alloy might have interrupted oxidation of the material [21, 34]. Therefore, in this study, we think that the growth of the oxide layer did not occur at a serious level during the ball milling and annealing processes and the influence of the oxide layer on the soft magnetic properties might have been weak. Fig. 3(a) shows DSC traces of the Fe-based amorphous powder either as-atomized or annealed and/or ball milled under different conditions, which revealed the glass transition (Tg) and crystallization (Tx) temperature of each sample. Tx did not vary significantly among the samples whereas Tg tended to decrease after both the ball milling and annealing processes. This is because Tx is mainly determined by the chemical composition of materials, and so the negligible variation of Tx among the samples indicates that impurities were not significantly introduced through the milling and annealing processes [35]. On the other hand, Tg decreased as crystallization progressed [36], thus we considered that the powder was notably crystallized after the combined processes of ball milling and annealing. Fig. 3(b) exhibits the calculated degree of crystallinity of each sample based on the DSC results according to a method reported elsewhere [32], for which the crystallinity of the as-atomized Fe-based amorphous powder was assumed to be 0%. The crystallinity of each powder after annealing, ball milling, and combining the two increased to 1.16%, 2.81%, and 4.62%, respectively. These results are consistent with those predicted from the change in Tg in Fig. 3(a). Both annealing and ball milling induced crystallization of the amorphous powder but the effect of each process on crystallization behavior
Fig. 6. Hysteresis loops of the amorphous powder: (a) as-atomized, (b) after annealing at 400 °C, (c) after 30 min ball milling, and (d) after 30 min ball milling and annealing at 400 °C.
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Fig. 7. (a) Coercivity and (b) magnetization of the amorphous powder: as-atomized, after annealing at 400 °C, after 30 min ball milling, and after 30 min ball milling and annealing at 400 °C.
of the amorphous powder was somewhat different. This is discussed with the TEM analysis later. Fig. 4 reveals TEM images of the various samples. As shown in the SAD pattern in Fig. 4(a) and Fig. 4(c), which are the Fe-based atomized powder and the powder after ball milling respectively, manufactured by gas atomization, has a complete amorphous phase. A crystalline layer was formed underneath the oxide layer after annealing at 425 °C, as shown in Fig. 4(b), while small crystallites were uniformly distributed in the interior of the powder when 1 h ball milling was followed by annealing at 425 °C, as shown in Fig. 4(d) (the SAD pattern in Fig. 4(d) also confirms the presence of crystallites). A schematic of the crystallization behavior of the amorphous powder after thermal annealing and the hybrid process (thermal annealing after ball milling) are shown in Fig. 5(a) and (b), respectively. As can be seen in the TEM images, the interface between the oxide layer and Table 1 The average and standard deviation values of coercivity and saturation magnetization of the amorphous powder: as-atomized, after annealing at 400 °C, after 30 min ball milling, and after 30 min ball milling and annealing at 400 °C. Sample
Coercivity (A/m)
Magnetization (Wbm/kg)
Average Std. Average Std. Deviation Deviation As atomized powder powder after annealing at 400 °C powder after 30 min ball milling powder after 30 min ball milling and annealing at 400 °C
3698.73 162.61 766.8 115.49
9.14 3.9 1.43 2.2
0.192 0.19 0.196 0.197
0.001 0.004 0.004 0.003
amorphous powder can be act as a heterogeneous nucleation sites for crystallization during thermal annealing [37]. However, the introduction of mechanical energy prior to thermal annealing might have generated a number of crystal defects that can act as homogeneous nucleation sites at interior of the amorphous powder [38]. Further into the ballmilling process, the amorphous powder was crystallized to some extent without thermal annealing, as revealed by the DSC analysis in Fig. 3. When thermal annealing followed the ball-milling process, the crystallization was much accelerated; the degree of crystallization of the amorphous powder after the hybrid process was higher than solely ball milling or thermal annealing. However, the size of the crystallites in the amorphous powder was smaller after the hybrid process than after thermal annealing only. Since the number of nucleation sites was much larger than with thermal annealing only, crystalline growth was much suppressed in the case of the hybrid process. This observation is contrary to previous work that reported heterogeneous nucleation of crystallites at the grain boundaries after ball milling for much longer durations (10 to 60 h) [39]. Fig. 6 reveals hysteresis loops of each sample under various conditions measured with a VSM, while the measured values of coercivity and saturation magnetization of each sample are presented in Fig. 7(a) and (b), respectively. The corresponding average and standard deviation of coercivity and magnetization value of four samples are summarized in Table 1. Coercivity decreased by 95.60%, 79.27%, and 96.88% (from 3698.73 to 162.61, 766.80, and 115.49 A/m) after annealing, ball milling, and the combined process, respectively. On the other hand, saturation magnetization was increased by 2.08% and 2.60% (from 0.192 to 0.196 and 0.197 Wbm/g) after ball milling and the combined process, respectively, while it was decreased by 1.04% (from 0.192 to 0.19 Wbm/g) after annealing at 400 °C. Although the formation of non-magnetic layers (i.e. oxide layers) on the powder surface might have decreased magnetization [40], their influence on the powder's magnetic properties were considered to be minimal in this study because of their negligible amount compared to the volume of the powder interior. Therefore, the improvement of soft magnetic properties (a decrease in coercivity and an increase in saturation magnetization) mainly originated from the reduced residual stress and formation of nanocrystallites [41, 42] (amorphous powder is essentially subjected to residual stress generated during a rapid cooling process [43]). Although magnetic anisotropy is essentially absent in amorphous materials because of their non-crystalline structure, it can be formed when the residual stress of the amorphous powder interacts with magnetostriction. As a result, coercivity is increased and permeability is decreased because magnetic anisotropy interrupts magnetization [8]. Residual stress can be effectively removed through annealing, so coercivity can be decreased. However, according to a previous study, saturation magnetization is also decreased due to a decrease in the lattice constant and the unit cell volume when the residual stress in the amorphous powder is removed by the annealing process [22]. In this study, the saturation magnetization was rather increased, possibly because of the reduction in free volume; amorphous materials possess a more disordered atomic arrangement and larger free volume compared to crystalline materials, and partial crystallization may reduce the free volume so as to enhance the saturation magnetization [44]. 4. Conclusions In this study, we produced a Fe-based amorphous powder using gas atomization and investigated the effect of mechanical milling and annealing processes on its microstructure and soft magnetic properties. During thermal annealing, crystallites with a mean size larger than 10 nm were formed at the interface between the oxide layer and amorphous powder because the interface acts as heterogeneous nucleation sites for crystallization. However, the introduction of mechanical energy prior to thermal annealing might have generated a number of crystal
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defects that were able to act as homogeneous nucleation sites in the interior of the amorphous powder, thereby generating small crystallites (b5 nm in mean size) homogeneously inside the powder. As a result, the introduction of nanocrystals during the combined process (30 min ball milling followed by annealing at 400 °C) simultaneously improved the powder's soft magnetic properties; it decreased coercivity by 96.88% and also increased saturation magnetization by 2.60% compared to the as-atomized amorphous powder. This may open up potential applications of these amorphous/nanocrystalline hybrid powders in the electric/electronic, IT, and electric vehicle industries that need power-converting materials with high-energy efficiency. Acknowledgments This research was supported by Civil-Military Technology Cooperation Program (18-CM-MA-15). This work was also supported by the National Research Foundation of Korea(NRF) Grant funded by the Korean Government(MSIP) (No. Grant Number - 2015R1A5A7037615). References [1] H. Shokrollahi, K. Janghorban, Soft magnetic composite materials, J. Mater. Process Tech. 189 (2007) 1–12. [2] A.M. Leary, P.R. Ohodnicki, M.E. Mchenry, Soft magnetic materials in high-frequency, high-power conversion applications, J. Miner. Met. Mater. Soc. 64 (2012) 772–781. [3] S.W. Kim, Y.W. Yoon, S.J. Lee, G.Y. Kim, Y.B. Kim, Y.Y. Chun, K.S. Lee, Electromagnetic shielding properties of soft magnetic powder-polymer composite films for the application to suppress noise in the radio frequency range, J. Magn. Magn. Mater. 316 (2007) 472–474. [4] L.V. Panina, K. Mohri, High-Frequency Giant Magneto-Impedance in Co-Rich Amorphous Wires and Films, J. Mag. Soc. Japan. 19 (1995) 265–268. [5] H. Lee, K.J. Lee, Y.K. Kim, T.K. Kim, C.O. Kim, S.C. Yu, Magnetoimpedance effect in the nanocrystalline Fe–Zr–Cu–B–Al alloy system, J. Appl. Phys. 87 (2000) 5269–5271. [6] M.A. Nowroozi, H. Shokrollahi, The effects of milling time and heat treatment on the micro-structural and magnetic behavior of Fe42Ni28Zr8Ta2B10C10 synthesized by mechanical alloying, J. Magn. Magn. Mater. 335 (2013) 53–58. [7] M. Pekala, D. Oleszak, E. Jartych, J.K. Zurawicz, Structural and magnetic study of crystalline Fe80Ni20 alloys with nanometer-sized grains, J. Non-Cryst. Solids 250-252 (1999) 757–761. [8] G. Herzer, Grain size dependence of coercivity and permeability in nanocrystalline ferromagnets, IEEE. T. Magn. 26 (1990) 1397–1402. [9] Y. Liu, Y. Yi, W. Shao, Y. Shao, Microstructure and magnetic properties of soft magnetic powder cores of amorphous and nanocrystalline alloys, J. Magn. Magn. Mater. 330 (2013) 119–133. [10] G. Herzer, Modern soft magnets: Amorphous and nanocrystalline materials, Acta Mater. 61 (2013) 718–734. [11] Y.Y. Zheng, Y.G. Wang, G.T. Xia, Amorphous soft magnetic composite-cores with various orientations of the powder-flakes, J. Magn. Magn. Mater. 396 (2015) 97–101. [12] T.H. Kim, K.K. Jee, Y.B. Kim, D.J. Byun, J.H. Han, High-frequency magnetic properties of soft magnetic cores based on nanocrystalline alloy powder prepared by thermal oxidation, J. Magn. Magn. Mater. 322 (2010) 2423–2427. [13] J. Zhou, Y.F. Cui, H.S. Liu, W. Wang, K. Peng, Y.D. Xiao, Magnetic properties of Fe78.4Si9.5B9Cu0.6Nb2.5 nanocrystalline alloy powder cores, J. Mater. Sci. 46 (2011) 7567–7572. [14] X. Wang, C. Lu, F. Guo, Z. Lu, D. Li, S. Zhou, New Fe-based amorphous compound powder cores with superior DC-bias properties and low loss characteristics, J. Magn. Magn. Mater. 324 (2012) 2727–2730. [15] T. Ma, M. Yan, W. Wang, The evolution of microstructure and magnetic properties of Fe-Si-Al powders prepared through melt-spinning, Scripta. Mater. 58 (2008) 243–246. [16] P. Schaaf, G. Rixecker, E. Yang, C.N.J. Wagner, U. Gonser, Study of nanocrystalline and amorphous powders prepared by mechanical alloying, Hyperfine Interac. 94 (1994) 2239–2244. [17] J.W. Jeong, D.Y. Yang, K.B. Kim, J. Lee, Y.J. Kim, T.S. Lim, S. Yang, M.H. Lee, H.J. Kim, Y.J. Kim, Investigation on Fe-Hf-B-Nb-P-C Soft Magnetic Powders Prepared by HighPressure Gas Atomization, J. Korean Powder Metall. Inst. 23 (2016) 391–396.
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Soft magnetic properties of Fe-based amorphous/nanocrystalline hybrid materials Yeonjoo Lee a, Jonggyu Jeon a, Seungjin Nam a, Teasuk Jang b, Hwijun Kim c, Minwoo Lee c, Yongjin Kim d, Dongyeol Yang d, Kyeongsik Min e, Hyunjoo Choi a,⁎ a
School of Advanced Materials Engineering, Kookmin University, Seoul, Republic of Korea Dept. of Advanced Materials Engineering, Sunmoon University, Asan, Republic of Korea Incheon Regional Division, Korea Institute of Industrial Technology, Incheon, Republic of Korea d Powder & Ceramics Division, Korea Institute of Materials Science, Changwon, Republic of Korea e School of Electrical Engineering, Kookmin University, Seoul, Republic of Korea b c
a r t i c l e
i n f o
Article history: Received 13 February 2018 Received in revised form 29 June 2018 Accepted 9 August 2018 Available online 11 August 2018 Keywords: Fe-based amorphous powder Ball milling Crystallization Coercivity Magnetization
a b s t r a c t In this study, we investigated the effect of crystalline structures on the magnetic properties of the Fe-based amorphous powder. We intentionally produced uniformly dispersed nanocrystallites inside the Fe-based amorphous powder using a combined process of mechanical milling and thermal annealing. In the case of powder produced after 30 min of ball milling followed by annealing at 400 °C, coercivity was decreased by 96.88% from 3698.73 to 115.49 A/m, and magnetization was increased 2.60% from 1.92 × 10−4 to 1.97 × 10−4 Wbm/kg. The simultaneous improvement of coercivity and magnetization can be attributed to the homogeneous generation of a small number of nanocrystallites as well as the release of residual stress. © 2018 Elsevier B.V. All rights reserved.
1. Introduction Soft magnetic materials have become of paramount importance in a variety of industrial fields such as electricity-generating facilities [1], power-converting parts [2], electromagnetic wave shielding materials [3], and magnetic sensors [4, 5]. As electronic devices become miniaturized and automated, the development of high-performance soft magnetic materials with good reliability is becoming urgent. In general, if the crystal size of crystalline metals is less than the exchange correlation length, a decrease in crystal size results in a decrease in coercivity and an increase in permeability, meaning an improvement in soft magnetic properties [6, 7]. Therefore, minimum coercivity can be obtained in an amorphous state. Furthermore, the disordered atomic arrangement of amorphous materials can provide excellent mechanical properties and maintain good magnetic properties even at high frequencies [6, 8]. In this regard, Fe-based amorphous materials have attracted increasing attention because of their good soft magnetic properties such as high saturation magnetization, high permeability, low coercivity, and low core loss [6]. Recently, efforts have been made to produce these amorphous materials in powder form because, differently from bulk amorphous ⁎ Corresponding author. E-mail address: [email protected] (H. Choi).
https://doi.org/10.1016/j.powtec.2018.08.037 0032-5910/© 2018 Elsevier B.V. All rights reserved.
materials, amorphous powder enables the production of parts with complex morphologies and three-dimensionally uniform magnetic properties [9, 10]. In previous work, the amorphous powder was generally produced by crushing a melt-spun ribbon [11–15], but in this case, low filling rate, high manufacturing cost, and characteristic deterioration have restricted the commercialization of the powder. Mechanical alloying has also been suggested to produce amorphous powders, whereby monolithic elemental powders are ball milled in a chamber to eventually generate an atomically mixed amorphous powder [16]. However, this process requires a long milling duration (generally longer than 48 h), which is not practical for industrial manufacturing processes. Recently, we have reported a way to produce Fe-Hf-B-Nb-P-C amorphous powder directly using gas atomization [17]. The powder revealed good soft magnetic properties such as good saturation magnetization higher than 149.7 emu/g and low coercivity below 12.8 Oe. On the other hand, the introduction of nanocrystallites into amorphous materials has been reported to reduce iron loss by preventing the generation of eddy currents at high frequencies [10]. Recently, F. Hosseini-Nasb et al. has generated nanocrystallites homogeneously in a melt-spun amorphous ribbon by controlling the wheel speed and this amorphous/nanocrystalline hybrid ribbon revealed a simultaneous improvement of saturation magnetization and coercivity [18]. An attempt has also been made to produce nanocrystallites in amorphous powder by inducing mechanical [6] or thermal [19, 20] energy into
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the powder [8, 21, 22]. However, the size and amount of the nanocrystallites generated during the process have yet to be clearly investigated. Furthermore, the simultaneous improvement of saturation magnetization and coercivity described in melt-spun ribbons has not yet been reported in amorphous powders, which may be because the crystallites formed mostly near the powder surface rather than being uniformly distributed inside the powder interior. In this study, we employed a combined treatment to introduce thermal energy after inducing mechanical energy into a gas-atomized Fe-based amorphous powder in order to generate nanocrystallites uniformly inside the powder interior. Mechanical milling may generate a number of defects inside the powder, which can act as homogeneous nucleation sites for crystallization during thermal annealing. We investigated the effect of the mechanical milling and thermal annealing processes on the microstructure and soft magnetic properties of the Fe-based amorphous powder. In general, amorphous powder is generally mixed with binders to obtain a bulk-type product because of difficulties in plastic deformation of the amorphous powder. This process does not require high temperatures or high pressures, and hence, it does not involve any significant microstructural changes (grain growth or oxidation) [23]. Therefore, the soft-magnetic properties of the products are determined in terms of the packing fraction and the inherent soft-magnetic properties of amorphous powder. Hence, in this study, the soft magnetic properties were measured in the powder state. 2. Experimental procedures 2.1. Sample preparation The Fe-B-P-C-Nb-Hf amorphous powder was produced by gas-atomization. A mixture of pure Fe, B, P, C, Nb, and Hf (≥ 99.5% purity) ingots was melted in a magnesia crucible set in a high-frequency induction furnace at 1400–1600 °C. The cast master alloy was extruded into a bar shape to easily insert and re-melt it in the gas-atomization equipment. After that, the master alloy was dissolved at 1300–1330 °C in a graphite crucible, and the melt was atomized to produce an amorphous powder by passing through a 2 mm nozzle in jets of He and N2 gases under 5–7.5 MPa pressure. The Fe-based amorphous powder was partially crystallized by mechanical milling and annealing processes. The powder was mechanically milled in a planetary mill (Pulverisette 5 classic line, Fritsch ™, Germany) at room temperature using a 500 mL stainless
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bowl. Powder was poured into the bowl with 200 mL ethanol and 5 mm-diameter stainless steel balls at a powder-to-ball weight ratio of 1:20. The milling was carried out at a constant rotating speed of 100 rpm for 30 min. Afterwards, as-atomized or ball-milled powders were annealed at 400 °C using a tube furnace (PTF-15/100/610, Lenton, Inc.) in argon; samples were heated to the target temperature at a heating rate of 10 °C/min, maintained for 1 h, and then furnace-cooled to room temperature. 2.2. Characterization The lattice structure of the Fe-based amorphous/nanocrystalline hybrid powders was characterized by X-ray diffraction (XRD) with a CN2301 (Rigaku, Japan) using CuKα (λ = 1.5405 Å). Samples were scanned over an angular range of 30–60° with a step size of 0.05° and a scan speed of 0.02°/min. During the XRD analysis, the sample was rotated to observe the general lattice. At that time, the diffractometer was set at 40 kV working voltage and 200 mA working current with a 2 mm slit size. The degree of crystallinity was examined by differential scanning calorimetry (DSC) with a TGA/DSC 1 (Mettler Toledo, USA) in an argon atmosphere. The sample was heated to 600 °C at a heating rate of 30 °C/min, and was then cooled to room temperature using circulated water. The same amount of powder used in this analysis was (300 ± 10 mg) for all samples to obtain comparable results. The morphology of the powders was observed using scanning electron microscopy (SEM) with a JEM-7610F (JEOL Ltd., Japan). Prior to the SEM analysis, the powders were coated with platinum (Pt) and attached using carbon tape to improve imaging. The microstructure of the Fe-based amorphous powder and amorphous/nanocrystalline hybrid powders were observed using transmission electron microscopy (TEM) with a Tecnai F20 G2 (FEI, USA) operating at 200 kV; TEM specimens were prepared using a focused ion beam technique. Fast Fourier transform (FFT) patterns and an inverse FFT (IFFT) images were obtained using Gatan digital microscope software for detailed structure analysis. Soft magnetic properties (i.e. saturation magnetization and coercivity) of the prepared powders were measured using a vibration sample magnetometer (VSM; Lake Shore 7400) with a maximum applied field of 10 kG and a B\\H loop tracer at room temperature. For all samples, 0.07 ± 0.01 g powder was loaded in a quartz crucible for the test, and prior to the test, the place of sample holder was adjusted using a standard Ni specimen.
Fig. 1. SEM images of powders after ball milling for (a) 0 h (as-atomized), (b) 15 min, (c) 30 min, (d) 1 h, (e) 15 h, and (f) 72 h.
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3. Results and discussion Fig. 1(a-f) shows the morphology of the powders after ball milling for a variety of time periods. As-atomized powder exhibited a spherical shape with a mean diameter of 10.88 μm, which became flattened as ball milling proceeded. In general, the initial impact of the grinding ball induced the soft crystalline metal powder to flatten and become work-hardened. With continued milling, the work-hardened powder tended to fracture and become refined. Although the starting material was in an amorphous state in this study, the morphological evolution of the powder during milling was similar to that of crystalline metal powders. Although it proceeded much more slowly, the amorphous powder started to become flattened after 15 h of milling whereas crystalline metal powders generally became flattened within 3 h under similar milling conditions [24]. This demonstrate that, similar to crystalline metal powders, the amorphous powder also experienced significant plastic deformation during milling [25]. Unlike bulk amorphous materials, a number of defects were generated from the large specific surface in the powder state, and these defects facilitated the plastic deformation [26]. Temperature rise during milling might also have helped plastic deformation of the amorphous powder as N90% of the mechanical energy imparted to the powders during milling was transformed into heat [27]. In addition, a certain proportion of the amorphous phase was crystallized during the process, which also contributed to the plastic deformation [6]. Therefore, in this study, the milling duration was fixed at 30 min to avoid significant crystallization and morphological changes. Fig. 2 shows XRD patterns of as-atomized powder and powders annealed and/or ball milled under different condition. Regardless of ball milling time and annealing temperature, all samples exhibited a typical amorphous pattern at 2θ ranging from 40° to 50° [6, 28]. Crystallization of the amorphous powder was not thought to have notably occurred during a short time period with low-energy milling combined with annealing at 400 °C (81.87% of crystallization temperature of the amorphous powder (i.e. 488.59 °C)). A peak indicating crystalline α-Fe phases was observed at 44.5° in 2θ when the powder was annealed at temperatures higher than 400 °C. This result is comparable to previous reports in that crystallization of the Fe-based amorphous alloy resulted in the formation of the α-Fe phase followed by the Fe2B phase [29]. On the other hand, ball milling was found to stimulate the formation of the Fe2B phase (the peak detected at 42.5° in 2θ). Compared with general heat treatment, it has been found that ball milling of Fe-based powders can accelerate crystallization of this phase due to excessive generation of vacancy-type defects [6]. According to previous research, when Fe81B13.5Si13.5C2 was heat-treated at a heating rate of 5 °C/min, the
Fig. 2. XRD patterns of the amorphous powder: as-atomized, after annealing at 400 °C, after 30 min ball milling, and after 30 min ball milling and annealing at 400 °C, in ascending order.
Fig. 3. (a) DSC traces and (b) degree of crystallinity for the amorphous powder: asatomized, after annealing at 400 °C, after 30 min ball milling, and after 30 min ball milling and annealing at 400 °C.
Fig. 4. TEM images of the amorphous powder: (a) as-atomized, (b) after annealing at 425 °C, (c) after 1-h ball milling and (d) after 1-h ball milling and annealing at 425 °C The insets in the top right-hand corners of (a)-(d) are selected area diffraction (SAD) patterns.
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Fig. 5. Schematic of the crystallization of the amorphous powder via (a) thermal processing and (b) thermal processing after mechanical processing.
α-Fe phase was formed at 485 °C, the Fe3B phase was formed at 515 °C, and the Fe3C phase was formed at 630 °C [29, 30]. In this study, a small amount of α-Fe and Fe2B phases were formed at a relatively low temperature of 425 °C, much lower than that in previous reports [31]. In addition, the crystal size was found to be very small based on the broad XRD spectrum [6], the significance of which is discussed along with the TEM observations later. All of the samples also revealed two peaks at 33.8 and 34.3°, which indicate the presence of an Fe3O4 layer possibly on the powder surface. The intensity of the peaks did not significantly vary as milling and heating progressed. When as-atomized powder was
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annealed at 425 °C, the thickness of the oxide layer was found to increase from 70 to 90 nm according to the SEM analysis in a previous study [32]. The growth rate of the Fe3O4 layer of Fe-based amorphous powder is much slower than that of pure iron [32, 33] because diffusion of iron atoms in the amorphous state of a powder is much suppressed compared to in a crystalline state. Furthermore, the presence of Nb and P in this amorphous alloy might have interrupted oxidation of the material [21, 34]. Therefore, in this study, we think that the growth of the oxide layer did not occur at a serious level during the ball milling and annealing processes and the influence of the oxide layer on the soft magnetic properties might have been weak. Fig. 3(a) shows DSC traces of the Fe-based amorphous powder either as-atomized or annealed and/or ball milled under different conditions, which revealed the glass transition (Tg) and crystallization (Tx) temperature of each sample. Tx did not vary significantly among the samples whereas Tg tended to decrease after both the ball milling and annealing processes. This is because Tx is mainly determined by the chemical composition of materials, and so the negligible variation of Tx among the samples indicates that impurities were not significantly introduced through the milling and annealing processes [35]. On the other hand, Tg decreased as crystallization progressed [36], thus we considered that the powder was notably crystallized after the combined processes of ball milling and annealing. Fig. 3(b) exhibits the calculated degree of crystallinity of each sample based on the DSC results according to a method reported elsewhere [32], for which the crystallinity of the as-atomized Fe-based amorphous powder was assumed to be 0%. The crystallinity of each powder after annealing, ball milling, and combining the two increased to 1.16%, 2.81%, and 4.62%, respectively. These results are consistent with those predicted from the change in Tg in Fig. 3(a). Both annealing and ball milling induced crystallization of the amorphous powder but the effect of each process on crystallization behavior
Fig. 6. Hysteresis loops of the amorphous powder: (a) as-atomized, (b) after annealing at 400 °C, (c) after 30 min ball milling, and (d) after 30 min ball milling and annealing at 400 °C.
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Fig. 7. (a) Coercivity and (b) magnetization of the amorphous powder: as-atomized, after annealing at 400 °C, after 30 min ball milling, and after 30 min ball milling and annealing at 400 °C.
of the amorphous powder was somewhat different. This is discussed with the TEM analysis later. Fig. 4 reveals TEM images of the various samples. As shown in the SAD pattern in Fig. 4(a) and Fig. 4(c), which are the Fe-based atomized powder and the powder after ball milling respectively, manufactured by gas atomization, has a complete amorphous phase. A crystalline layer was formed underneath the oxide layer after annealing at 425 °C, as shown in Fig. 4(b), while small crystallites were uniformly distributed in the interior of the powder when 1 h ball milling was followed by annealing at 425 °C, as shown in Fig. 4(d) (the SAD pattern in Fig. 4(d) also confirms the presence of crystallites). A schematic of the crystallization behavior of the amorphous powder after thermal annealing and the hybrid process (thermal annealing after ball milling) are shown in Fig. 5(a) and (b), respectively. As can be seen in the TEM images, the interface between the oxide layer and Table 1 The average and standard deviation values of coercivity and saturation magnetization of the amorphous powder: as-atomized, after annealing at 400 °C, after 30 min ball milling, and after 30 min ball milling and annealing at 400 °C. Sample
Coercivity (A/m)
Magnetization (Wbm/kg)
Average Std. Average Std. Deviation Deviation As atomized powder powder after annealing at 400 °C powder after 30 min ball milling powder after 30 min ball milling and annealing at 400 °C
3698.73 162.61 766.8 115.49
9.14 3.9 1.43 2.2
0.192 0.19 0.196 0.197
0.001 0.004 0.004 0.003
amorphous powder can be act as a heterogeneous nucleation sites for crystallization during thermal annealing [37]. However, the introduction of mechanical energy prior to thermal annealing might have generated a number of crystal defects that can act as homogeneous nucleation sites at interior of the amorphous powder [38]. Further into the ballmilling process, the amorphous powder was crystallized to some extent without thermal annealing, as revealed by the DSC analysis in Fig. 3. When thermal annealing followed the ball-milling process, the crystallization was much accelerated; the degree of crystallization of the amorphous powder after the hybrid process was higher than solely ball milling or thermal annealing. However, the size of the crystallites in the amorphous powder was smaller after the hybrid process than after thermal annealing only. Since the number of nucleation sites was much larger than with thermal annealing only, crystalline growth was much suppressed in the case of the hybrid process. This observation is contrary to previous work that reported heterogeneous nucleation of crystallites at the grain boundaries after ball milling for much longer durations (10 to 60 h) [39]. Fig. 6 reveals hysteresis loops of each sample under various conditions measured with a VSM, while the measured values of coercivity and saturation magnetization of each sample are presented in Fig. 7(a) and (b), respectively. The corresponding average and standard deviation of coercivity and magnetization value of four samples are summarized in Table 1. Coercivity decreased by 95.60%, 79.27%, and 96.88% (from 3698.73 to 162.61, 766.80, and 115.49 A/m) after annealing, ball milling, and the combined process, respectively. On the other hand, saturation magnetization was increased by 2.08% and 2.60% (from 0.192 to 0.196 and 0.197 Wbm/g) after ball milling and the combined process, respectively, while it was decreased by 1.04% (from 0.192 to 0.19 Wbm/g) after annealing at 400 °C. Although the formation of non-magnetic layers (i.e. oxide layers) on the powder surface might have decreased magnetization [40], their influence on the powder's magnetic properties were considered to be minimal in this study because of their negligible amount compared to the volume of the powder interior. Therefore, the improvement of soft magnetic properties (a decrease in coercivity and an increase in saturation magnetization) mainly originated from the reduced residual stress and formation of nanocrystallites [41, 42] (amorphous powder is essentially subjected to residual stress generated during a rapid cooling process [43]). Although magnetic anisotropy is essentially absent in amorphous materials because of their non-crystalline structure, it can be formed when the residual stress of the amorphous powder interacts with magnetostriction. As a result, coercivity is increased and permeability is decreased because magnetic anisotropy interrupts magnetization [8]. Residual stress can be effectively removed through annealing, so coercivity can be decreased. However, according to a previous study, saturation magnetization is also decreased due to a decrease in the lattice constant and the unit cell volume when the residual stress in the amorphous powder is removed by the annealing process [22]. In this study, the saturation magnetization was rather increased, possibly because of the reduction in free volume; amorphous materials possess a more disordered atomic arrangement and larger free volume compared to crystalline materials, and partial crystallization may reduce the free volume so as to enhance the saturation magnetization [44]. 4. Conclusions In this study, we produced a Fe-based amorphous powder using gas atomization and investigated the effect of mechanical milling and annealing processes on its microstructure and soft magnetic properties. During thermal annealing, crystallites with a mean size larger than 10 nm were formed at the interface between the oxide layer and amorphous powder because the interface acts as heterogeneous nucleation sites for crystallization. However, the introduction of mechanical energy prior to thermal annealing might have generated a number of crystal
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defects that were able to act as homogeneous nucleation sites in the interior of the amorphous powder, thereby generating small crystallites (b5 nm in mean size) homogeneously inside the powder. As a result, the introduction of nanocrystals during the combined process (30 min ball milling followed by annealing at 400 °C) simultaneously improved the powder's soft magnetic properties; it decreased coercivity by 96.88% and also increased saturation magnetization by 2.60% compared to the as-atomized amorphous powder. This may open up potential applications of these amorphous/nanocrystalline hybrid powders in the electric/electronic, IT, and electric vehicle industries that need power-converting materials with high-energy efficiency. Acknowledgments This research was supported by Civil-Military Technology Cooperation Program (18-CM-MA-15). This work was also supported by the National Research Foundation of Korea(NRF) Grant funded by the Korean Government(MSIP) (No. Grant Number - 2015R1A5A7037615). References [1] H. Shokrollahi, K. Janghorban, Soft magnetic composite materials, J. Mater. Process Tech. 189 (2007) 1–12. [2] A.M. Leary, P.R. Ohodnicki, M.E. Mchenry, Soft magnetic materials in high-frequency, high-power conversion applications, J. Miner. Met. Mater. Soc. 64 (2012) 772–781. [3] S.W. Kim, Y.W. Yoon, S.J. Lee, G.Y. Kim, Y.B. Kim, Y.Y. Chun, K.S. Lee, Electromagnetic shielding properties of soft magnetic powder-polymer composite films for the application to suppress noise in the radio frequency range, J. Magn. Magn. Mater. 316 (2007) 472–474. [4] L.V. Panina, K. Mohri, High-Frequency Giant Magneto-Impedance in Co-Rich Amorphous Wires and Films, J. Mag. Soc. Japan. 19 (1995) 265–268. [5] H. Lee, K.J. Lee, Y.K. Kim, T.K. Kim, C.O. Kim, S.C. Yu, Magnetoimpedance effect in the nanocrystalline Fe–Zr–Cu–B–Al alloy system, J. Appl. Phys. 87 (2000) 5269–5271. [6] M.A. Nowroozi, H. Shokrollahi, The effects of milling time and heat treatment on the micro-structural and magnetic behavior of Fe42Ni28Zr8Ta2B10C10 synthesized by mechanical alloying, J. Magn. Magn. Mater. 335 (2013) 53–58. [7] M. Pekala, D. Oleszak, E. Jartych, J.K. Zurawicz, Structural and magnetic study of crystalline Fe80Ni20 alloys with nanometer-sized grains, J. Non-Cryst. Solids 250-252 (1999) 757–761. [8] G. Herzer, Grain size dependence of coercivity and permeability in nanocrystalline ferromagnets, IEEE. T. Magn. 26 (1990) 1397–1402. [9] Y. Liu, Y. Yi, W. Shao, Y. Shao, Microstructure and magnetic properties of soft magnetic powder cores of amorphous and nanocrystalline alloys, J. Magn. Magn. Mater. 330 (2013) 119–133. [10] G. Herzer, Modern soft magnets: Amorphous and nanocrystalline materials, Acta Mater. 61 (2013) 718–734. [11] Y.Y. Zheng, Y.G. Wang, G.T. Xia, Amorphous soft magnetic composite-cores with various orientations of the powder-flakes, J. Magn. Magn. Mater. 396 (2015) 97–101. [12] T.H. Kim, K.K. Jee, Y.B. Kim, D.J. Byun, J.H. Han, High-frequency magnetic properties of soft magnetic cores based on nanocrystalline alloy powder prepared by thermal oxidation, J. Magn. Magn. Mater. 322 (2010) 2423–2427. [13] J. Zhou, Y.F. Cui, H.S. Liu, W. Wang, K. Peng, Y.D. Xiao, Magnetic properties of Fe78.4Si9.5B9Cu0.6Nb2.5 nanocrystalline alloy powder cores, J. Mater. Sci. 46 (2011) 7567–7572. [14] X. Wang, C. Lu, F. Guo, Z. Lu, D. Li, S. Zhou, New Fe-based amorphous compound powder cores with superior DC-bias properties and low loss characteristics, J. Magn. Magn. Mater. 324 (2012) 2727–2730. [15] T. Ma, M. Yan, W. Wang, The evolution of microstructure and magnetic properties of Fe-Si-Al powders prepared through melt-spinning, Scripta. Mater. 58 (2008) 243–246. [16] P. Schaaf, G. Rixecker, E. Yang, C.N.J. Wagner, U. Gonser, Study of nanocrystalline and amorphous powders prepared by mechanical alloying, Hyperfine Interac. 94 (1994) 2239–2244. [17] J.W. Jeong, D.Y. Yang, K.B. Kim, J. Lee, Y.J. Kim, T.S. Lim, S. Yang, M.H. Lee, H.J. Kim, Y.J. Kim, Investigation on Fe-Hf-B-Nb-P-C Soft Magnetic Powders Prepared by HighPressure Gas Atomization, J. Korean Powder Metall. Inst. 23 (2016) 391–396.
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