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Soft magnetic properties in Fe84−xB10C6Cux nanocrystalline alloys
Journal of Magnetism and Magnetic Materials 326 (2013) 22–27
Contents lists available at SciVerse ScienceDirect
Journal of Magnetism and Magnetic Materials journal homepage: www.elsevier.com/locate/jmmm
Soft magnetic properties in Fe84 xB10C6Cux nanocrystalline alloys X.D. Fan a, H. Men b, A.B. Ma a, B.L. Shen b,n a
College of Mechanics and Materials, Hohai University, Nanjing 210098, PR China Zhejiang Province Key Laboratory of Magnetic Materials and Application Technology, Key Laboratory of Magnetic Materials and Devices, Ningbo Institute of Materials Technology and Engineering, Chinese Academy of Sciences, Ningbo 315201, PR China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 14 December 2011 Received in revised form 7 June 2012 Available online 2 September 2012
The dependence of Cu addition on magnetic properties of Fe84 xB10C6Cux nanocrystalline alloys prepared by annealing the melt-spun amorphous ribbons was investigated. It was found that the coercivity (Hc) of the nanocrystalline Fe84 xB10C6Cux alloys decreased with increasing Cu addition and exhibited a minimum value with composition of x ¼ 1 after appropriate annealing, while the saturation magnetic flux density (Bs) shows an increasing trend owing to the increasing volume fraction of nanocrystalline phase. And the alloy with composition of x ¼1 exhibits excellent magnetic properties, i.e., high Bs of 1.78 T, low Hc of 5 A/m and low core loss (P) of 0.34 W/kg at 1 T and 50 Hz. & 2012 Elsevier B.V. All rights reserved.
Keywords: Soft magnetic alloy Fe-based nanocrystalline alloy High saturation magnetic flux density Low core loss
1. Introduction Recently, energy saving has been considered as one of the most important issues since energy problem has become crucial in the world. For years, Si-steels have been taken as the most widely used soft magnetic materials due to their high saturation magnetic flux density (Bs) [1,2]. However, Si-steels have a rather high core loss at high frequency compared with those of commonly used Fe-based amorphous alloys [3] and Fe-based nanocrystalline alloys [4–6]. Nanocrystalline soft magnetic alloys produced from amorphous alloys have attracted great attention due to their excellent soft magnetic properties over the past two decades. Among them, FeSiBNbCu nanocrystalline alloys (FINEMET) have been used for magnetic devices due to the high permeability and low core loss [4,7], but the Bs of most widely used FINEMET is only 1.24 T and the alloy contains the expensive metal element Nb. Recently, nanocrystalline alloy FeCuB has become a new research topic of general interest because of its excellent soft magnetic properties [8,9]. However, either the alloy FeCuB or its derivative FeCuSiB [10] contains more than 14 at% of B element, which causes the increase of the production cost. More recently, nanocrystalline alloy FeSiBPCu has been reported to exhibit a high Bs as high as 1.9 T [11,12], but it contains volatile element P which limits its industrial applications. Therefore, it is important to search and explore an Fe-based nanocrystalline alloy with high Bs, low core loss and low cost, which is also easy to be produced in industrial
n
Corresponding author. Tel.: þ86 574 87913392; fax: þ 86 574 87911392. E-mail address: [email protected] (B.L. Shen).
0304-8853/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jmmm.2012.08.045
applications. It has been reported that FeBC amorphous alloys exhibit high Bs over 1.75 T [13], but a rather high coercivity (Hc) of 10 A/m and a rather high core loss of 0.5 W/kg at 1.7 T and 60 Hz limited their applications. Therefore, with the aim of synthesizing an Fe-based nanocrystalline alloy system with good soft magnetic properties and low cost, as well as easy production, we added Cu element to this FeBC amorphous alloys. As a result, Fe-based nanocrystalline Fe84 xB10C6Cux alloys have been successfully synthesized, which exhibit high Bs of 1.74–1.78 T and low Hc of 5–8 A/m [14]. This paper intends to report in detail the microstructure, thermal properties and magnetic properties of the Fe84 xB10C6Cux alloy system. 2. Experimental procedure FeBCCu alloy ingots were prepared by arc melting the mixtures of Fe (99.99 mass%), B (99.7 mass%), Cu (99.99 mass%) and pre-alloyed Fe–C alloy in a highly purified argon atmosphere. Each ingot was melted five times to ensure compositional homogeneity. Amorphous Fe84 xB10C6Cux (x¼0, 0.5, 1, 1.3) alloy ribbons were prepared by the single roller melt-spinning method [15]. The molten alloy is pushed through the orifice of a pressurized crucible onto a rotating copper wheel with linear velocity of 40 m/s. The width and thickness of the ribbons were 1 mm and 20 mm, respectively. Melt-spun ribbons were annealed to develop nano-scale grains. The annealing was carried out by keeping the ribbons in the tubular furnace preheated to annealing temperature for the preset annealing time in argon atmosphere followed by water quenching cooling. Microstructures were identified by X-ray diffraction (XRD) with Cu Ka1 radiation and transmission
X.D. Fan et al. / Journal of Magnetism and Magnetic Materials 326 (2013) 22–27
electron microscopy (TEM). Grain size (D) was estimated by TEM observation and Scherrer’s equation from XRD profiles. The thermal stability was evaluated by a differential scanning calorimeter (DSC) at a heating rate of 0.67 1C/s. Bs was measured with a vibrating sample magnetometer (VSM) under a maximum applied field of 800 kA/m and Hc was measured with a DC B–H curve tracer under a maximum applied field of 1 kA/m. The melt-spun ribbons were wound into laminated toroids 10 mm in outer diameter, and 6 mm in inner diameter, and subsequently isothermal annealed in the same condition with the ribbon specimens to determine the AC properties of the nanocrystalline alloy. Core losses at 1 T and 50 Hz, 400 Hz and 1 kHz were measured with an AC B–H analyzer.
3. Results and discussion The XRD patterns of Fe84 xB10C6Cux melt-spun ribbons are shown in Fig. 1. All samples were examined on free surface. It can be seen that a sharp crystallization peak corresponding to the Fe3C phase exists for Fe84B10C6 alloy. However, with increase of Cu addition, the crystallization peak disappeared and all samples exhibit halo patterns, indicating the amorphous structure. It can also be concluded from the XRD patterns of Fe84 xB10C6Cux alloys that the glass-forming ability of the Fe84B10C6 alloy is poor but can be improved by Cu addition, that is consisted with the former results reported by Liu et al. [16], who pointed out that Cu has a flux effect to the base alloy; thus the impurities such as oxides which are usually harmful to the glass formation were distinctly decreased by Cu addition. The DSC curves for Fe84 xB10C6Cux melt-spun ribbons are shown in Fig. 2. It can be seen that the crystallization process of Fe84 xB10C6Cux alloy includes two stages. The first crystallization
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onset temperature (Tx1) decreases slightly with increasing Cu addition from x¼0 to 1. When x¼ 1.3, Tx1 decreases dramatically, whereas the second crystallization onset temperature (Tx2) shows little variation. Hence, the difference between Tx1 and Tx2 increases with increasing Cu addition. In order to investigate the crystallization behavior, XRD measurement was carried out for the Fe83B10C6Cu1 alloy annealed at different temperatures. Fig. 3 shows the XRD patterns of the Fe83B10C6Cu1 alloy subjected to annealing for 10 min at 340 1C, which is far lower than the temperature Tx1, and 10 min at 430 1C and 520 1C, respectively, corresponding to a temperature between Tx1 and Tx2, and a temperature just higher than that of the second exothermic peak. The XRD pattern of the as quenched alloy is also shown for comparison. As a result, a crystalline phase which is identified as a-Fe superimposes on that of the amorphous phase for the sample annealed at 340 1C, which means the precipitation amount of the a-Fe phase is small. As the annealing temperature (Ta) increases to 430 1C, the precipitation of a-Fe single phase is obvious. When annealed at 520 1C, the precipitation phases are identified as a-Fe, Fe3B and Fe2B phases. Therefore, it is confirmed that Tx1 and Tx2 represent the crystallization temperature of a-Fe phase and the precipitation temperature of Fe–B compounds, respectively. It has been reported that the achievement of good soft magnetic properties is closely related to the nanoscale grain size of a-Fe phase, while the formation of Fe2B and Fe3B phases with large particle size, which have a large magnetocrystalline anisotropy, causes deterioration in the softness of the materials [17]. Therefore, in order to obtain good soft magnetic properties, the proper annealing temperature should be selected between Tx1 and Tx2. Meanwhile, according to DSC results, the crystallization peak is sharp in the alloy with composition x¼0–1, indicating that the crystallization process of a-Fe phase is quick. The temperature difference between Tx1 and Tx2 enlarges with composition x¼ 1, which is advantageous for synthesizing nanocrystallites of this alloy. But in the alloy with the composition x¼1.3, the crystallization peak becomes broad and the crystallization behavior of a-Fe phase occurs within a wide temperature range of about 100 1C. Therefore, the crystallization process is difficult to control, so it will not be advantageous for precipitating the nanocrystalline phase. According to former research results [14], Hc for Fe83B10C6Cu1 nanocrystalline alloy decreases with increasing annealing temperature up to 430 1C, then increases rapidly as Ta is increased, so that 430 1C was chosen as the annealing temperature. Fig. 4 shows the dependence of Hc on annealing time (t) for Fe83B10C6Cu1 alloy annealed at 430 1C. Here the values of Hc are calculated by averaging the three specimens annealed in the same condition. It is seen that Hc decreases gradually, reaches its
Fig. 1. XRD patterns of Fe84 xB10C6Cux melt-spun ribbons.
Fig. 2. DSC curves for Fe84 xB10C6Cux melt-spun ribbons.
Fig. 3. XRD patterns of Fe83B10C6Cu1 alloy ribbons annealed for 10 min at 340, 430, and 520 1C. The pattern of the as quenched alloy is also shown for comparison.
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X.D. Fan et al. / Journal of Magnetism and Magnetic Materials 326 (2013) 22–27
minimum value at t ¼3 min, and then increases rapidly as t is increased. Therefore, in this study, we determined the annealing time as 3 min. Fig. 5 shows XRD patterns of Fe84 xB10C6Cux alloys annealed at 430 1C for 3 min. The alloy with composition x ¼0.5 is almost an amorphous structure after annealing. For the alloys with composition of x¼1 and 1.3, the precipitation of a-Fe phase can be observed obviously. The average grain size (D) estimated by Scherrer’s equation according to the (110) reflection peak is 15 nm for x ¼1 and 25 nm for x¼1.3, indicating that the nanocrystalline a-Fe phase was obtained. Hence, Cu addition has a vital impact on the Fe84 xB10C6Cux alloy system and favors the precipitation of a-Fe phase. According to the previous results
[4,18], Cu has a tendency to segregate from Fe during the annealing process and form Cu clusters, which act as the nucleation sites of a-Fe primary crystals that formed during melt-spun progress. Fig. 6 shows the bright-field TEM images and selected area electron diffraction (SAED) patterns of (a) Fe83B10C6Cu1 and (b) Fe82.7B10C6Cu1.3 nanocrystalline alloys. The TEM image shows that nanocrystalline grains of a-Fe phase precipitated from Fe83B10C6Cu1 amorphous alloy after annealing, but the volume fraction of nanocrystalline phase is low and the grains seem to distribute heterogeneously. However, Fe82.7B10C6Cu1.3 alloy shows a uniform nanocrystallized structure with grain size of 20–30 nm and obvious higher volume fraction of nanocrystalline phase. The a-Fe grain size distributions for Fe83B10C6Cu1 and Fe82.7B10C6Cu1.3 nanocrystalline alloys determined by TEM are shown in Fig. 7. The distribution of the alloy with more Cu addition is much narrower, and the standard deviation decreases from 4.88 to 3.48 with increasing x from 1 to 1.3. The average D for Fe83B10C6Cu1 and Fe82.7B10C6Cu1.3 nanocrystalline alloys is about 18 nm and 26 nm, respectively. According to the distribution statistical results, the values of D by means of TEM are quite uniform compared to that estimated by XRD. Fig. 8 shows the HRTEM images of (a) Fe83B10C6Cu1 and (b) Fe82.7B10C6Cu1.3 nanocrystalline alloys. The HRTEM observation reveals that the microstructure consists of residual amorphous phase and fine a-Fe grains embedded in the amorphous matrix phase. The grain size is about 15 nm and 25 nm,
Fig. 4. Dependence of Hc on annealing time for Fe83B10C6Cu1 alloy annealed at 430 1C.
Fig. 5. XRD patterns of Fe84 xB10C6Cux alloys annealed at 430 1C for 3 min.
Fig. 7. a-Fe grain size distributions for Fe83B10C6Cu1 and Fe82.7B10C6Cu1.3 nanocrystalline alloys.
Fig. 6. Bright-field TEM images and SAED patterns of (a) Fe83B10C6Cu1 and (b) Fe82.7B10C6Cu1.3 nanocrystalline alloys.
X.D. Fan et al. / Journal of Magnetism and Magnetic Materials 326 (2013) 22–27
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Fig. 8. HRTEM images of (a) Fe83B10C6Cu1 and (b) Fe82.7B10C6Cu1.3 nanocrystalline alloys.
Fig. 9. EDS analyses on a-Fe grains and amorphous phase for Fe82.7B10C6Cu1.3 nanocrystalline alloy.
respectively, for Fe83B10C6Cu1 and Fe82.7B10C6Cu1.3 nanocrystalline alloys, which is completely consistent with the XRD results. Fig. 9 shows the EDS analyses on a-Fe grain and amorphous phase for Fe82.7B10C6Cu1.3 nanocrystalline alloy. It clearly shows that the spectrum of C detected from the residual amorphous phase is slightly stronger than that detected from the a-Fe grain, which means the residual amorphous phase is enriched with C element. The C-rich amorphous phase may inhibit grain growth, similarly to the Nb-, Zr- or B-rich amorphous phase in traditional nanocrystalline alloy systems [4,19–21]. In Fig. 10, the dependences of (a) magnetic flux density at 800 kA/m (B800) and (b) coercivity (Hc) on Cu content for Fe84 xB10C6Cux alloys annealed at 430 1C for 3 min are shown. Note that in this alloy system, B800 EBs. According to the results, Bs shows an increasing trend with increasing Cu addition. It is reported that Bs can be calculated by the equation Bs ¼BscVc/ VþBsaVa/V, where Vc/V and Va/V stand for the volume fraction of the crystalline phase and the amorphous phase with V¼Vc þVa; and Bsc and Bsa are the saturation magnetic flux densities in the crystalline and amorphous phases, respectively [11]. The value of Bsc of a-Fe phase is more than 2 T, whereas Bsa is lower than 1.7 T for this alloy system. Therefore, the higher the Vc/V, the higher the Bs. Vc/V is 15% larger for x ¼1.3 than x¼1. That is why the alloy with composition of x ¼1.3 exhibits a higher Bs. Meanwhile, Hc shows a decreasing trend with increasing Cu content x. The value of Hc for the alloy with x¼0 is as large as 13 A/m. With increasing Cu addition x from 0.5 to 1, Hc decreases from 8 to 5 A/m. Then Hc increases to 15 A/m for x ¼1.3. The increasing Hc is due to the increasing average grain size of a-Fe
Fig. 10. Dependences of (a) magnetic flux density at 800 kA/m (B800) and (b) coercivity (Hc) on Cu content for Fe84 xB10C6Cux alloys annealed at 430 1C for 3 min.
Fig. 11. Dependence of core loss (P) on maximum magnetic flux density (Bm) for Fe84 xB10C6Cux nanocrystalline alloys at 50 Hz, together with those of commercial oriented 3 mass% Si-steel.
phase. According to Herzer’s reports, when the grain size D is smaller than the exchange length Lex, anisotropy density /KS goes close to zero; thus excellent soft magnetic properties are obtained with the reduction in D. Coercivity is a function of grain size; the rate of increase in Hc is almost proportional to D6 up to D E50 nm [22,23]. The average grain size D estimated by XRD is 15 nm for x ¼1 and 25 nm for x ¼1.3. Therefore, the alloy with composition of x¼ 1 exhibits a lower Hc, which means excellent soft magnetic properties are obtained for Fe83B10C6Cu1 nanocrystalline alloy. Fig. 11 shows the dependence of core loss (P) on maximum magnetic flux density (Bm) for Fe84 xB10C6Cux nanocrystalline
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X.D. Fan et al. / Journal of Magnetism and Magnetic Materials 326 (2013) 22–27
that of Fe73.5Cu1Nb3Si13.5B9 nanocrystalline alloy. Core loss of P10/50 for this alloy is 0.34 W/kg, which is about 80% that of oriented 3 mass% Si-steel. The core losses at high frequencies, such as P10/400 and P10/1k is 4.3 W/kg and 12.5 W/kg, respectively, which is much lower than those of oriented 3 mass% Si-steel. Though the optimum annealing temperature range for this alloy is relatively narrow compared to that of traditional nanocrystalline alloy [4], the excellent soft magnetic properties should promise their application in soft magnetic industry. Moreover, this nanocrystalline alloy system has a large economical application value of low material cost and convenient productivity because of the absence of easily oxidizable metal elements and volatile elements.
Fig. 12. Dependence of core loss (P) on maximum magnetic flux density (Bm) for Fe84 xB10C6Cux nanocrystalline alloys at 400 Hz, together with those of commercial oriented 3 mass% Si-steel.
Table 1 Bs, Hc and P of Fe84 xB10C6Cux nanocrystalline alloys under several conditions are compared with those of Fe73.5Cu1Nb3Si13.5B9, Fe83.7Cu1.5B14.8 and Fe83.3Si4B8P4Cu0.7 nanocrystalline alloys and commercial oriented 3 mass% Si-steel. Alloy
Bs (T)
Hc (A/m)
P10/50a (W/kg)
P10/400a (W/kg)
P10/1ka (W/kg)
Fe83.5B10C6Cu0.5 Fe83B10C6Cu1 Fe82.7B10C6Cu1.3 Fe73.5Cu1Nb3Si13.5B9 Fe83.7Cu1.5B14.8 Fe83.3Si4B8P4Cu0.7 Oriented 3 mass% Si-steel (CGO, 0.23 mm)
1.74 1.78 1.83 1.24 1.82 1.88 1.92
8 5 15 0.53 7 7 9.8
0.37 0.34 0.51 P2/20k ¼2.1 P15/50 ¼0.38 0.12 0.41
4.7 4.3 6.9
14.2 12.5 20.3
7.9
27.1
a Here, the symbols P10/50, P10/400 and P10/1k stand for core losses at 1 T and 50 Hz, 400 Hz and 1 kHz, respectively.
alloys at 50 Hz, together with that of a toroidal core of oriented 3 mass% Si-steel. Here the Si-steel that is used for comparison is the conventional grain-oriented (CGO) steel 0.23 mm in thickness. Compared to the oriented 3 mass% Si-steel, this Fe-based nanocrystalline alloy system shows a lower P for Bm o0.8 T. Although the alloy with x ¼1.3 shows a higher value and quick increasing trend with increasing Bm, core loss for Fe83B10C6Cu1 nanocrystalline alloy exhibits a much lower value than that of oriented 3 mass% Si-steel. Fig. 12 shows the dependence of core loss (P) on maximum magnetic flux density (Bm) for Fe84 xB10C6Cux nanocrystalline alloys at 400 Hz, together with those of a toroidal core of commercial oriented 3 mass% Si-steel. The increasing rate of P against Bm for this nanocrystalline alloy system is small when Bm o0.8 T, and the values of P show much lower than those of oriented 3 mass% Si-steel over the whole magnetic flux density range. P for the alloy with composition x¼1 at 1 T is 4.3 W/kg, which is about 55% that of commercial oriented 3 mass% Si-steel. For this alloy, the value of P keeps a low value with increasing Bm. Therefore, these nanocrystalline alloys are suitable for commercial use at a higher frequency. In Table 1, the Bs, Hc and P of Fe84 xB10C6Cux nanocrystalline alloys under several conditions are compared with those of Fe73.5Cu1Nb3Si13.5B9, Fe83.7Cu1.5B14.8 and Fe83.3Si4B8P4Cu0.7 nanocrystalline alloys [4,8,12] and commercial oriented 3 mass% Si-steel. Here, the symbols P10/50, P10/400 and P10/1k stand for core losses at 1 T and 50 Hz, 400 Hz and 1 kHz, respectively. These Fe84 xB10C6Cux nanocrystalline alloys exhibit low Hc of less than 15 A/m. Coercivity for Fe83B10C6Cu1 nanocrystalline alloy is 5 A/m, which is lower than those of Fe83.7Cu1.5B14.8 and Fe83.3Si4B8P4Cu0.7 nanocrystalline alloys and commercial oriented 3 mass% Si-steel. Although the Bs shows a slight decrease for these materials, it is more than 40% higher than
4. Conclusions The improvement of soft magnetic properties in Fe84 xB10C6Cux nanocrystalline alloy system was investigated in this study. The main results are summarized as follows. 1. Cu addition suppressed the crystallization in Fe84 xB10C6Cux alloy system and improved the glass-forming ability of this alloy system. 2. Cu addition favors the precipitation of nanocrystalline a-Fe phase. The average grain size is 15 and 25 nm for the alloy with composition x¼1 and 1.3, respectively. 3. The C element is enriched in the residual amorphous phase for this nanocrystalline alloy system, which may inhibit grain growth. 4. Hc decreases with increasing Cu addition and exhibits a minimum value for composition x¼1, while Bs shows an increasing trend owing to the increasing volume fraction of nanocrystalline a-Fe phase. 5. The Fe83B10C6Cu1 nanocrystalline alloy exhibits excellent soft magnetic properties with a high Bs of 1.78 T, low Hc of 5 A/m and low core loss of 0.34 W/kg at 1 T and 50 Hz.
Acknowledgments This work was supported by the National 863 Project (Grant no. 2009AA03Z214), the National Science Fund of China for Distinguished Young Scholars (Grant no. 50825103), the National Natural Science Foundation of China (Grant no. 51001112), the Hundred Talents Program (Grant no. KGCX-2-YW-803) by Chinese Academy of Sciences and the Fundamental Research Funds for the Central Universities (2010B15414). References [1] D.F. Binns, A.B. Crompton, A. Jaberansari, IEE Proceedings Part C: Generation, Transmission and Distribution 133 (1986) 451. [2] M. Abe, Y. Takada, T. Murakami, Y. Tanaka, Y. Mihara, Journal of Materials Engineering 11 (1989) 109. [3] Y. Ogawa, M. Naoe, Y. Yoshizawa, R. Hasegawa, Journal of Magnetism and Magnetic Materials 304 (2006) e675. [4] Y. Yoshizawa, S. Oguma, K. Yamauchi, Journal of Applied Physics 64 (1988) 6044. [5] K. Suzuki, A. Makino, N. Kataoka, A. Inoue, T. Masumoto, Materials Transactions, The Japan Institute of Metals 32 (1991) 93. [6] M.A. Willard, M.Q. Huang, D.E. Laughlin, M.E. McHenry, J.O. Cross, V.G. Harris, C. Franchetti, Journal of Applied Physics 85 (1999) 4421. [7] J. Petzold, Scripta Materialia 48 (2003) 895. [8] M. Ohta, Y. Yoshizawa, Japanese Journal of Applied Physics 46 (2007) 477. [9] M. Ohta, Y. Yoshizawa, Materials Transactions 48 (2007) 2378. [10] M. Ohta, Y. Yoshizawa, Applied Physics Letters 91 (2007) 062517. [11] A. Makino, H. Men, T. Kubota, K. Yubuta, A. Inoue, Materials Transactions 50 (2009) 204. [12] A. Makino, H. Men, T. Kubota, K. Yubuta, A. Inoue, IEEE Transactions on Magnetics 45 (2009) 4302. [13] S. Hatta, T. Egami, C.D. Graham Jr., Applied Physics Letters 34 (1979) 113.
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[14] X.D. Fan, A.B. Ma, H. Men, G.Q. Xie, B.L. Shen, A. Makino, A. Inoue, Journal of Applied Physics 109 (2011) 07A314. [15] M.E. McHenry, M.A. Willard, D.E. Laughlin, Progress in Materials Science 44 (1999) 291–433. [16] Y.Z. Jia, S.Y. Zeng, S.F. Shan, L.Y. Zhang, C.Z. Fan, B.Q. Zhang, Z.J. Zhan, R.P. Liu, W.K. Wang, Journal of Alloys and Compounds 440 (2007) 113. [17] A. Makino, A. Inoue, T. Masumoto, Materials Transactions, The Japan Institute of Metals 36 (1995) 924. [18] M. Ohta, Y. Yoshizawa, Applied Physics Express 2 (2009) 023005.
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[19] Y. Yoshizawa, K. Yamauchi, Materials Science and Engineering A 133 (1991) 176. [20] K. Suzuki, M. Kikuchi, A. Makino, A. Inoue, T. Masumoto, Materials Transactions, The Japan Institute of Metals 32 (1991) 961. [21] K. Suzuki, A. Makino, A.P. Tsai, A. Inoue, T. Masumoto, Materials Science and Engineering A 179 (1994) 501. [22] G. Herzer, IEEE Transactions on Magnetics 26 (1990) 1397. [23] G. Herzer, Handbook of Magnetic Materials, Elsevier Science, Amsterdam, 1997.
Contents lists available at SciVerse ScienceDirect
Journal of Magnetism and Magnetic Materials journal homepage: www.elsevier.com/locate/jmmm
Soft magnetic properties in Fe84 xB10C6Cux nanocrystalline alloys X.D. Fan a, H. Men b, A.B. Ma a, B.L. Shen b,n a
College of Mechanics and Materials, Hohai University, Nanjing 210098, PR China Zhejiang Province Key Laboratory of Magnetic Materials and Application Technology, Key Laboratory of Magnetic Materials and Devices, Ningbo Institute of Materials Technology and Engineering, Chinese Academy of Sciences, Ningbo 315201, PR China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 14 December 2011 Received in revised form 7 June 2012 Available online 2 September 2012
The dependence of Cu addition on magnetic properties of Fe84 xB10C6Cux nanocrystalline alloys prepared by annealing the melt-spun amorphous ribbons was investigated. It was found that the coercivity (Hc) of the nanocrystalline Fe84 xB10C6Cux alloys decreased with increasing Cu addition and exhibited a minimum value with composition of x ¼ 1 after appropriate annealing, while the saturation magnetic flux density (Bs) shows an increasing trend owing to the increasing volume fraction of nanocrystalline phase. And the alloy with composition of x ¼1 exhibits excellent magnetic properties, i.e., high Bs of 1.78 T, low Hc of 5 A/m and low core loss (P) of 0.34 W/kg at 1 T and 50 Hz. & 2012 Elsevier B.V. All rights reserved.
Keywords: Soft magnetic alloy Fe-based nanocrystalline alloy High saturation magnetic flux density Low core loss
1. Introduction Recently, energy saving has been considered as one of the most important issues since energy problem has become crucial in the world. For years, Si-steels have been taken as the most widely used soft magnetic materials due to their high saturation magnetic flux density (Bs) [1,2]. However, Si-steels have a rather high core loss at high frequency compared with those of commonly used Fe-based amorphous alloys [3] and Fe-based nanocrystalline alloys [4–6]. Nanocrystalline soft magnetic alloys produced from amorphous alloys have attracted great attention due to their excellent soft magnetic properties over the past two decades. Among them, FeSiBNbCu nanocrystalline alloys (FINEMET) have been used for magnetic devices due to the high permeability and low core loss [4,7], but the Bs of most widely used FINEMET is only 1.24 T and the alloy contains the expensive metal element Nb. Recently, nanocrystalline alloy FeCuB has become a new research topic of general interest because of its excellent soft magnetic properties [8,9]. However, either the alloy FeCuB or its derivative FeCuSiB [10] contains more than 14 at% of B element, which causes the increase of the production cost. More recently, nanocrystalline alloy FeSiBPCu has been reported to exhibit a high Bs as high as 1.9 T [11,12], but it contains volatile element P which limits its industrial applications. Therefore, it is important to search and explore an Fe-based nanocrystalline alloy with high Bs, low core loss and low cost, which is also easy to be produced in industrial
n
Corresponding author. Tel.: þ86 574 87913392; fax: þ 86 574 87911392. E-mail address: [email protected] (B.L. Shen).
0304-8853/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jmmm.2012.08.045
applications. It has been reported that FeBC amorphous alloys exhibit high Bs over 1.75 T [13], but a rather high coercivity (Hc) of 10 A/m and a rather high core loss of 0.5 W/kg at 1.7 T and 60 Hz limited their applications. Therefore, with the aim of synthesizing an Fe-based nanocrystalline alloy system with good soft magnetic properties and low cost, as well as easy production, we added Cu element to this FeBC amorphous alloys. As a result, Fe-based nanocrystalline Fe84 xB10C6Cux alloys have been successfully synthesized, which exhibit high Bs of 1.74–1.78 T and low Hc of 5–8 A/m [14]. This paper intends to report in detail the microstructure, thermal properties and magnetic properties of the Fe84 xB10C6Cux alloy system. 2. Experimental procedure FeBCCu alloy ingots were prepared by arc melting the mixtures of Fe (99.99 mass%), B (99.7 mass%), Cu (99.99 mass%) and pre-alloyed Fe–C alloy in a highly purified argon atmosphere. Each ingot was melted five times to ensure compositional homogeneity. Amorphous Fe84 xB10C6Cux (x¼0, 0.5, 1, 1.3) alloy ribbons were prepared by the single roller melt-spinning method [15]. The molten alloy is pushed through the orifice of a pressurized crucible onto a rotating copper wheel with linear velocity of 40 m/s. The width and thickness of the ribbons were 1 mm and 20 mm, respectively. Melt-spun ribbons were annealed to develop nano-scale grains. The annealing was carried out by keeping the ribbons in the tubular furnace preheated to annealing temperature for the preset annealing time in argon atmosphere followed by water quenching cooling. Microstructures were identified by X-ray diffraction (XRD) with Cu Ka1 radiation and transmission
X.D. Fan et al. / Journal of Magnetism and Magnetic Materials 326 (2013) 22–27
electron microscopy (TEM). Grain size (D) was estimated by TEM observation and Scherrer’s equation from XRD profiles. The thermal stability was evaluated by a differential scanning calorimeter (DSC) at a heating rate of 0.67 1C/s. Bs was measured with a vibrating sample magnetometer (VSM) under a maximum applied field of 800 kA/m and Hc was measured with a DC B–H curve tracer under a maximum applied field of 1 kA/m. The melt-spun ribbons were wound into laminated toroids 10 mm in outer diameter, and 6 mm in inner diameter, and subsequently isothermal annealed in the same condition with the ribbon specimens to determine the AC properties of the nanocrystalline alloy. Core losses at 1 T and 50 Hz, 400 Hz and 1 kHz were measured with an AC B–H analyzer.
3. Results and discussion The XRD patterns of Fe84 xB10C6Cux melt-spun ribbons are shown in Fig. 1. All samples were examined on free surface. It can be seen that a sharp crystallization peak corresponding to the Fe3C phase exists for Fe84B10C6 alloy. However, with increase of Cu addition, the crystallization peak disappeared and all samples exhibit halo patterns, indicating the amorphous structure. It can also be concluded from the XRD patterns of Fe84 xB10C6Cux alloys that the glass-forming ability of the Fe84B10C6 alloy is poor but can be improved by Cu addition, that is consisted with the former results reported by Liu et al. [16], who pointed out that Cu has a flux effect to the base alloy; thus the impurities such as oxides which are usually harmful to the glass formation were distinctly decreased by Cu addition. The DSC curves for Fe84 xB10C6Cux melt-spun ribbons are shown in Fig. 2. It can be seen that the crystallization process of Fe84 xB10C6Cux alloy includes two stages. The first crystallization
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onset temperature (Tx1) decreases slightly with increasing Cu addition from x¼0 to 1. When x¼ 1.3, Tx1 decreases dramatically, whereas the second crystallization onset temperature (Tx2) shows little variation. Hence, the difference between Tx1 and Tx2 increases with increasing Cu addition. In order to investigate the crystallization behavior, XRD measurement was carried out for the Fe83B10C6Cu1 alloy annealed at different temperatures. Fig. 3 shows the XRD patterns of the Fe83B10C6Cu1 alloy subjected to annealing for 10 min at 340 1C, which is far lower than the temperature Tx1, and 10 min at 430 1C and 520 1C, respectively, corresponding to a temperature between Tx1 and Tx2, and a temperature just higher than that of the second exothermic peak. The XRD pattern of the as quenched alloy is also shown for comparison. As a result, a crystalline phase which is identified as a-Fe superimposes on that of the amorphous phase for the sample annealed at 340 1C, which means the precipitation amount of the a-Fe phase is small. As the annealing temperature (Ta) increases to 430 1C, the precipitation of a-Fe single phase is obvious. When annealed at 520 1C, the precipitation phases are identified as a-Fe, Fe3B and Fe2B phases. Therefore, it is confirmed that Tx1 and Tx2 represent the crystallization temperature of a-Fe phase and the precipitation temperature of Fe–B compounds, respectively. It has been reported that the achievement of good soft magnetic properties is closely related to the nanoscale grain size of a-Fe phase, while the formation of Fe2B and Fe3B phases with large particle size, which have a large magnetocrystalline anisotropy, causes deterioration in the softness of the materials [17]. Therefore, in order to obtain good soft magnetic properties, the proper annealing temperature should be selected between Tx1 and Tx2. Meanwhile, according to DSC results, the crystallization peak is sharp in the alloy with composition x¼0–1, indicating that the crystallization process of a-Fe phase is quick. The temperature difference between Tx1 and Tx2 enlarges with composition x¼ 1, which is advantageous for synthesizing nanocrystallites of this alloy. But in the alloy with the composition x¼1.3, the crystallization peak becomes broad and the crystallization behavior of a-Fe phase occurs within a wide temperature range of about 100 1C. Therefore, the crystallization process is difficult to control, so it will not be advantageous for precipitating the nanocrystalline phase. According to former research results [14], Hc for Fe83B10C6Cu1 nanocrystalline alloy decreases with increasing annealing temperature up to 430 1C, then increases rapidly as Ta is increased, so that 430 1C was chosen as the annealing temperature. Fig. 4 shows the dependence of Hc on annealing time (t) for Fe83B10C6Cu1 alloy annealed at 430 1C. Here the values of Hc are calculated by averaging the three specimens annealed in the same condition. It is seen that Hc decreases gradually, reaches its
Fig. 1. XRD patterns of Fe84 xB10C6Cux melt-spun ribbons.
Fig. 2. DSC curves for Fe84 xB10C6Cux melt-spun ribbons.
Fig. 3. XRD patterns of Fe83B10C6Cu1 alloy ribbons annealed for 10 min at 340, 430, and 520 1C. The pattern of the as quenched alloy is also shown for comparison.
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minimum value at t ¼3 min, and then increases rapidly as t is increased. Therefore, in this study, we determined the annealing time as 3 min. Fig. 5 shows XRD patterns of Fe84 xB10C6Cux alloys annealed at 430 1C for 3 min. The alloy with composition x ¼0.5 is almost an amorphous structure after annealing. For the alloys with composition of x¼1 and 1.3, the precipitation of a-Fe phase can be observed obviously. The average grain size (D) estimated by Scherrer’s equation according to the (110) reflection peak is 15 nm for x ¼1 and 25 nm for x¼1.3, indicating that the nanocrystalline a-Fe phase was obtained. Hence, Cu addition has a vital impact on the Fe84 xB10C6Cux alloy system and favors the precipitation of a-Fe phase. According to the previous results
[4,18], Cu has a tendency to segregate from Fe during the annealing process and form Cu clusters, which act as the nucleation sites of a-Fe primary crystals that formed during melt-spun progress. Fig. 6 shows the bright-field TEM images and selected area electron diffraction (SAED) patterns of (a) Fe83B10C6Cu1 and (b) Fe82.7B10C6Cu1.3 nanocrystalline alloys. The TEM image shows that nanocrystalline grains of a-Fe phase precipitated from Fe83B10C6Cu1 amorphous alloy after annealing, but the volume fraction of nanocrystalline phase is low and the grains seem to distribute heterogeneously. However, Fe82.7B10C6Cu1.3 alloy shows a uniform nanocrystallized structure with grain size of 20–30 nm and obvious higher volume fraction of nanocrystalline phase. The a-Fe grain size distributions for Fe83B10C6Cu1 and Fe82.7B10C6Cu1.3 nanocrystalline alloys determined by TEM are shown in Fig. 7. The distribution of the alloy with more Cu addition is much narrower, and the standard deviation decreases from 4.88 to 3.48 with increasing x from 1 to 1.3. The average D for Fe83B10C6Cu1 and Fe82.7B10C6Cu1.3 nanocrystalline alloys is about 18 nm and 26 nm, respectively. According to the distribution statistical results, the values of D by means of TEM are quite uniform compared to that estimated by XRD. Fig. 8 shows the HRTEM images of (a) Fe83B10C6Cu1 and (b) Fe82.7B10C6Cu1.3 nanocrystalline alloys. The HRTEM observation reveals that the microstructure consists of residual amorphous phase and fine a-Fe grains embedded in the amorphous matrix phase. The grain size is about 15 nm and 25 nm,
Fig. 4. Dependence of Hc on annealing time for Fe83B10C6Cu1 alloy annealed at 430 1C.
Fig. 5. XRD patterns of Fe84 xB10C6Cux alloys annealed at 430 1C for 3 min.
Fig. 7. a-Fe grain size distributions for Fe83B10C6Cu1 and Fe82.7B10C6Cu1.3 nanocrystalline alloys.
Fig. 6. Bright-field TEM images and SAED patterns of (a) Fe83B10C6Cu1 and (b) Fe82.7B10C6Cu1.3 nanocrystalline alloys.
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Fig. 8. HRTEM images of (a) Fe83B10C6Cu1 and (b) Fe82.7B10C6Cu1.3 nanocrystalline alloys.
Fig. 9. EDS analyses on a-Fe grains and amorphous phase for Fe82.7B10C6Cu1.3 nanocrystalline alloy.
respectively, for Fe83B10C6Cu1 and Fe82.7B10C6Cu1.3 nanocrystalline alloys, which is completely consistent with the XRD results. Fig. 9 shows the EDS analyses on a-Fe grain and amorphous phase for Fe82.7B10C6Cu1.3 nanocrystalline alloy. It clearly shows that the spectrum of C detected from the residual amorphous phase is slightly stronger than that detected from the a-Fe grain, which means the residual amorphous phase is enriched with C element. The C-rich amorphous phase may inhibit grain growth, similarly to the Nb-, Zr- or B-rich amorphous phase in traditional nanocrystalline alloy systems [4,19–21]. In Fig. 10, the dependences of (a) magnetic flux density at 800 kA/m (B800) and (b) coercivity (Hc) on Cu content for Fe84 xB10C6Cux alloys annealed at 430 1C for 3 min are shown. Note that in this alloy system, B800 EBs. According to the results, Bs shows an increasing trend with increasing Cu addition. It is reported that Bs can be calculated by the equation Bs ¼BscVc/ VþBsaVa/V, where Vc/V and Va/V stand for the volume fraction of the crystalline phase and the amorphous phase with V¼Vc þVa; and Bsc and Bsa are the saturation magnetic flux densities in the crystalline and amorphous phases, respectively [11]. The value of Bsc of a-Fe phase is more than 2 T, whereas Bsa is lower than 1.7 T for this alloy system. Therefore, the higher the Vc/V, the higher the Bs. Vc/V is 15% larger for x ¼1.3 than x¼1. That is why the alloy with composition of x ¼1.3 exhibits a higher Bs. Meanwhile, Hc shows a decreasing trend with increasing Cu content x. The value of Hc for the alloy with x¼0 is as large as 13 A/m. With increasing Cu addition x from 0.5 to 1, Hc decreases from 8 to 5 A/m. Then Hc increases to 15 A/m for x ¼1.3. The increasing Hc is due to the increasing average grain size of a-Fe
Fig. 10. Dependences of (a) magnetic flux density at 800 kA/m (B800) and (b) coercivity (Hc) on Cu content for Fe84 xB10C6Cux alloys annealed at 430 1C for 3 min.
Fig. 11. Dependence of core loss (P) on maximum magnetic flux density (Bm) for Fe84 xB10C6Cux nanocrystalline alloys at 50 Hz, together with those of commercial oriented 3 mass% Si-steel.
phase. According to Herzer’s reports, when the grain size D is smaller than the exchange length Lex, anisotropy density /KS goes close to zero; thus excellent soft magnetic properties are obtained with the reduction in D. Coercivity is a function of grain size; the rate of increase in Hc is almost proportional to D6 up to D E50 nm [22,23]. The average grain size D estimated by XRD is 15 nm for x ¼1 and 25 nm for x ¼1.3. Therefore, the alloy with composition of x¼ 1 exhibits a lower Hc, which means excellent soft magnetic properties are obtained for Fe83B10C6Cu1 nanocrystalline alloy. Fig. 11 shows the dependence of core loss (P) on maximum magnetic flux density (Bm) for Fe84 xB10C6Cux nanocrystalline
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that of Fe73.5Cu1Nb3Si13.5B9 nanocrystalline alloy. Core loss of P10/50 for this alloy is 0.34 W/kg, which is about 80% that of oriented 3 mass% Si-steel. The core losses at high frequencies, such as P10/400 and P10/1k is 4.3 W/kg and 12.5 W/kg, respectively, which is much lower than those of oriented 3 mass% Si-steel. Though the optimum annealing temperature range for this alloy is relatively narrow compared to that of traditional nanocrystalline alloy [4], the excellent soft magnetic properties should promise their application in soft magnetic industry. Moreover, this nanocrystalline alloy system has a large economical application value of low material cost and convenient productivity because of the absence of easily oxidizable metal elements and volatile elements.
Fig. 12. Dependence of core loss (P) on maximum magnetic flux density (Bm) for Fe84 xB10C6Cux nanocrystalline alloys at 400 Hz, together with those of commercial oriented 3 mass% Si-steel.
Table 1 Bs, Hc and P of Fe84 xB10C6Cux nanocrystalline alloys under several conditions are compared with those of Fe73.5Cu1Nb3Si13.5B9, Fe83.7Cu1.5B14.8 and Fe83.3Si4B8P4Cu0.7 nanocrystalline alloys and commercial oriented 3 mass% Si-steel. Alloy
Bs (T)
Hc (A/m)
P10/50a (W/kg)
P10/400a (W/kg)
P10/1ka (W/kg)
Fe83.5B10C6Cu0.5 Fe83B10C6Cu1 Fe82.7B10C6Cu1.3 Fe73.5Cu1Nb3Si13.5B9 Fe83.7Cu1.5B14.8 Fe83.3Si4B8P4Cu0.7 Oriented 3 mass% Si-steel (CGO, 0.23 mm)
1.74 1.78 1.83 1.24 1.82 1.88 1.92
8 5 15 0.53 7 7 9.8
0.37 0.34 0.51 P2/20k ¼2.1 P15/50 ¼0.38 0.12 0.41
4.7 4.3 6.9
14.2 12.5 20.3
7.9
27.1
a Here, the symbols P10/50, P10/400 and P10/1k stand for core losses at 1 T and 50 Hz, 400 Hz and 1 kHz, respectively.
alloys at 50 Hz, together with that of a toroidal core of oriented 3 mass% Si-steel. Here the Si-steel that is used for comparison is the conventional grain-oriented (CGO) steel 0.23 mm in thickness. Compared to the oriented 3 mass% Si-steel, this Fe-based nanocrystalline alloy system shows a lower P for Bm o0.8 T. Although the alloy with x ¼1.3 shows a higher value and quick increasing trend with increasing Bm, core loss for Fe83B10C6Cu1 nanocrystalline alloy exhibits a much lower value than that of oriented 3 mass% Si-steel. Fig. 12 shows the dependence of core loss (P) on maximum magnetic flux density (Bm) for Fe84 xB10C6Cux nanocrystalline alloys at 400 Hz, together with those of a toroidal core of commercial oriented 3 mass% Si-steel. The increasing rate of P against Bm for this nanocrystalline alloy system is small when Bm o0.8 T, and the values of P show much lower than those of oriented 3 mass% Si-steel over the whole magnetic flux density range. P for the alloy with composition x¼1 at 1 T is 4.3 W/kg, which is about 55% that of commercial oriented 3 mass% Si-steel. For this alloy, the value of P keeps a low value with increasing Bm. Therefore, these nanocrystalline alloys are suitable for commercial use at a higher frequency. In Table 1, the Bs, Hc and P of Fe84 xB10C6Cux nanocrystalline alloys under several conditions are compared with those of Fe73.5Cu1Nb3Si13.5B9, Fe83.7Cu1.5B14.8 and Fe83.3Si4B8P4Cu0.7 nanocrystalline alloys [4,8,12] and commercial oriented 3 mass% Si-steel. Here, the symbols P10/50, P10/400 and P10/1k stand for core losses at 1 T and 50 Hz, 400 Hz and 1 kHz, respectively. These Fe84 xB10C6Cux nanocrystalline alloys exhibit low Hc of less than 15 A/m. Coercivity for Fe83B10C6Cu1 nanocrystalline alloy is 5 A/m, which is lower than those of Fe83.7Cu1.5B14.8 and Fe83.3Si4B8P4Cu0.7 nanocrystalline alloys and commercial oriented 3 mass% Si-steel. Although the Bs shows a slight decrease for these materials, it is more than 40% higher than
4. Conclusions The improvement of soft magnetic properties in Fe84 xB10C6Cux nanocrystalline alloy system was investigated in this study. The main results are summarized as follows. 1. Cu addition suppressed the crystallization in Fe84 xB10C6Cux alloy system and improved the glass-forming ability of this alloy system. 2. Cu addition favors the precipitation of nanocrystalline a-Fe phase. The average grain size is 15 and 25 nm for the alloy with composition x¼1 and 1.3, respectively. 3. The C element is enriched in the residual amorphous phase for this nanocrystalline alloy system, which may inhibit grain growth. 4. Hc decreases with increasing Cu addition and exhibits a minimum value for composition x¼1, while Bs shows an increasing trend owing to the increasing volume fraction of nanocrystalline a-Fe phase. 5. The Fe83B10C6Cu1 nanocrystalline alloy exhibits excellent soft magnetic properties with a high Bs of 1.78 T, low Hc of 5 A/m and low core loss of 0.34 W/kg at 1 T and 50 Hz.
Acknowledgments This work was supported by the National 863 Project (Grant no. 2009AA03Z214), the National Science Fund of China for Distinguished Young Scholars (Grant no. 50825103), the National Natural Science Foundation of China (Grant no. 51001112), the Hundred Talents Program (Grant no. KGCX-2-YW-803) by Chinese Academy of Sciences and the Fundamental Research Funds for the Central Universities (2010B15414). References [1] D.F. Binns, A.B. Crompton, A. Jaberansari, IEE Proceedings Part C: Generation, Transmission and Distribution 133 (1986) 451. [2] M. Abe, Y. Takada, T. Murakami, Y. Tanaka, Y. Mihara, Journal of Materials Engineering 11 (1989) 109. [3] Y. Ogawa, M. Naoe, Y. Yoshizawa, R. Hasegawa, Journal of Magnetism and Magnetic Materials 304 (2006) e675. [4] Y. Yoshizawa, S. Oguma, K. Yamauchi, Journal of Applied Physics 64 (1988) 6044. [5] K. Suzuki, A. Makino, N. Kataoka, A. Inoue, T. Masumoto, Materials Transactions, The Japan Institute of Metals 32 (1991) 93. [6] M.A. Willard, M.Q. Huang, D.E. Laughlin, M.E. McHenry, J.O. Cross, V.G. Harris, C. Franchetti, Journal of Applied Physics 85 (1999) 4421. [7] J. Petzold, Scripta Materialia 48 (2003) 895. [8] M. Ohta, Y. Yoshizawa, Japanese Journal of Applied Physics 46 (2007) 477. [9] M. Ohta, Y. Yoshizawa, Materials Transactions 48 (2007) 2378. [10] M. Ohta, Y. Yoshizawa, Applied Physics Letters 91 (2007) 062517. [11] A. Makino, H. Men, T. Kubota, K. Yubuta, A. Inoue, Materials Transactions 50 (2009) 204. [12] A. Makino, H. Men, T. Kubota, K. Yubuta, A. Inoue, IEEE Transactions on Magnetics 45 (2009) 4302. [13] S. Hatta, T. Egami, C.D. Graham Jr., Applied Physics Letters 34 (1979) 113.
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