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Softening and magnetic properties of ultrahigh Fe content FeSiBCuPC nanocrystalline alloy induced by low-pressure stress annealing
Scripta Materialia 179 (2020) 6–11
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Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Softening and magnetic properties of ultrahigh Fe content FeSiBCuPC nanocrystalline alloy induced by low-pressure stress annealing Jia Xu a,b, Yuanzheng Yang a,∗, Qiusheng Yan b,∗, Guihua Xiao a, Ting Luo a, Chenfeng Fan a a b
Faculty of Materials and Energy, Guangdong University of technology, Guangzhou 510006, China Faculty of Electromechanical Engineering, Guangdong University of technology, Guangzhou 510006, China
a r t i c l e
i n f o
Article history: Received 18 October 2019 Revised 23 December 2019 Accepted 29 December 2019
Keywords: Fe-based nanocrystalline alloys Low-pressure stress annealing Soft magnetic properties Hardness Young’s modulus
a b s t r a c t Due to the rapid growth of precipitated nanoparticles in ultrahigh Fe content nanocrystalline alloys, we propose a new method of low-pressure stress annealing (LPSa) to control crystallization behavior to achieve excellent magnetic properties and mechanical properties of Fe85.7 Si0.5 B9.3 Cu0.7 P3.5 C0.3 nanocrystalline alloy. It demonstrates that LPSa promotes the precipitation of α -Fe phase and suppresses the growth of nanoparticles, which obtains a higher saturation magnetization of 202.1–206.6 emu/g and a lower coercivity of 6.7–8.2 A/m. Besides, the formation of fine and uniform nanocrystalline structure induced by LPSa achieves the lower value and uniform distribution of hardness and Young’s modulus, resulting in softening.
Although nanocrystalline FeSiBPCu alloys with ultrahigh Fe content (≥ 85 at%) exhibit remarkable soft magnetic properties, brittleness severely restricts their broad application as functional materials [1–3]. Due to the absence of high melting point alloying elements and only a trace amount of Cu content, the α -Fe grains are prone to grow rapidly and non-uniformly during crystallization [4], resulting in high coercivity Hc and low plasticity. Previously, Makino et al. [5] demonstrated that the crystallization process of Fe85-86 SiBPCu alloys is very strict, requiring higher heating rate (≥ 200 °C/min) and shorter holding time (≤10 min) to achieve excellent soft magnetic properties, which seriously hinders the industrial production of high Fe content FeSiBPCu alloys. The plastic deformation ability can be tailored by changing the surface stress state of alloys [6]. Pressure as an important thermodynamic and kinetic parameter promotes the crystallization of amorphous alloys [7,8]. Unlike previous tensile stresses annealing [9,10], the crystallization of amorphous alloys by applying a compressive stress during the heat treatment promotes the precipitation of α -Fe phase and reduce the growth rate of nanoparticles [11,12]. Besides, LPSa continuously and uniformly regulates the energy state of nanocrystalline alloys, which has the advantages of fine-tuning of tissue structure, non-destructive and uniformity. In this paper, intending to suppress the rapid growth of nanocrystals and improve the ability of plastic deformation, we investigated the crystallization behavior, magnetic properties and mechanical prop∗
Corresponding authors. E-mail addresses: [email protected] (Y. Yang), [email protected] (Q. Yan).
https://doi.org/10.1016/j.scriptamat.2019.12.042 1359-6462/© 2020 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
© 2020 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
erties of ultrahigh Fe content Fe85.7 Si0.5 B9.3 Cu0.7 P3.5 C0.3 nanocrystalline alloy induced by low-pressure stress annealing (LPSa). Multi-component Fe85.7 Si0.5 B9.3 Cu0.7 P3.5 C0.3 alloy ingot was prepared by arc melting with the mixtures of raw materials in a high purity argon atmosphere. Ribbon with a width of 1 mm and a thickness of 20 μm was prepared by a single-roller melt spinning method. The diameter and width of the copper wheel are 320 mm and 50 mm, respectively. The structure and thermal stability of the ribbon was examined by X-ray diffraction (XRD) with Cu Kα radiation and differential scanning calorimetry (DSC) at a heating rate of 20 °C/min under nitrogen atmosphere, respectively. Heat treatment of amorphous ribbon was carried out in a tubular furnace with a constant heat rating of 20 °C/min. The saturation magnetization (Ms) and coercivity (Hc) were measured with a vibrating sample magnetometer (VSM) under an applied field of 12,500 Oe and a DC B-H loop tracer under a maximum applied field of 800 A/m, respectively. The hardness and Young’s modulus were measured with a Vickers hardness tester under a load of 0.49 N for 10 s and an atomic force microscope (AFM), respectively. To ensure the uniformity of the surface state of the ribbon, the free surface of the melt-spun and annealed ribbon with lower roughness was chosen for XRD, hardness and Young’s modulus measurements. The corresponding hardness distribution map is drawn by multi-point isometric hardness measurement. The sample is placed on a smooth and flat quartz glass slide as a substrate, and another glass slide is pressed in parallel on the sample to maintain a uniform force on the upper and lower sides
J. Xu, Y. Yang and Q. Yan et al. / Scripta Materialia 179 (2020) 6–11
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Fig. 1. Appearance of the Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 nanocrystalline ribbon under different applied compressive stresses (a) annealed at 360 °C and (b) annealed at 420 °C. (c) and (d) the XRD patterns of stress-annealed ribbon. (e) the DSC curves of the ribbon after annealed at 420 °C under different applied compressive stresses. (f) Compressive stress annealing dependence of the α -Fe grain size for the Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 ribbon.
of the ribbon. Then, it is placed in parallel in a tubular furnace for crystallization treatment. The pressure is calculated by combining the weight of glass slides and the surface area of the ribbon. The stress is adjusted by controlling the loading weight of the slides loaded on the ribbon. Based on the XRD pattern (presented in Fig. S1a), only a broad peak without any sharp diffraction peaks can be detected, indicating the typical amorphous structure. According to the DSC curve
(shown in Fig. S1b), the onset temperature of the first crystallization peak is at Tx1 ≈ 388 °C, while the secondary crystallization peak is at Tx2 ≈ 521 °C. The XRD patterns of the annealed ribbon is shown in Fig. S2. Although the ribbon has a wide ࢞T (࢞T=Tx2 Tx1 ~133 °C), the instability of nanocrystalline structure leads to premature precipitation of the secondary phase. Fig. 1(a) and (b) shows the photo images of the annealed sample with different compressive stresses applied along the ribbon
8
J. Xu, Y. Yang and Q. Yan et al. / Scripta Materialia 179 (2020) 6–11
Fig. 2. Hysteresis loops of Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 ribbon annealed at (a) 360 °C and (b) 420 °C for 15 min with different applied compressive stresses. (c) Variation in coercivity with annealing stress for the Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 ribbon.
plane during isothermal annealing. When annealed at 360 °C with stress-free, the ribbon exhibits a curled shape, which is similar to a ring. While the annealed ribbon becomes noticeably smooth with different applied low-pressure stresses, and the curvature decreases with increasing annealing stress. When annealed at 420 °C, the shape of the stress-annealed ribbon presents a straight line, indicating that the curl deformation behavior of ribbon is mainly dependent on the applied-pressure. Fig. 1(c) and (d) shows the XRD patterns of the ribbon annealed at 360 °C and 420 °C, respectively, under different applied pressure stresses. The intensity of the crystallization peak increases slightly with increasing annealing stress, which correlates with a higher volume fraction of the α -Fe phase. Fig. 1(e) shows the DSC curves of the ribbon after annealed at 420 °C with various applied-pressures. The area of exothermic peak correlates with the volume fraction of crystalline phase. Interestingly, the stress-annealed ribbon presents a straight line corresponding to the first exothermic peak of the initial crystalline phase, while the stress-free annealed ribbon still exhibits a small hill (࢞H~4.2 J/g) associated with the incomplete precipitation of α -Fe phase. Fig. 1(f) shows the applied pressure stress dependence of an average grain size (D) of the α -Fe phase estimated by Scherrer‘s equation. The D of stress-annealed ribbon is smaller than that of stress-free annealed ribbon. Besides, the D shows a significant decreasing trend with increasing annealing stress and is effectively controlled within 30 nm, indicating that PLSa contributes to the reduction of D. Consequently, LPSa effectively
promotes the precipitation of α -Fe grains and suppresses the rapid growth of nanoparticles in ultrahigh Fe content nanocrystalline alloy, which achieves the high volume fraction of α -Fe grains and maintains a finer nanocrystalline structure. Fig. 2(a) and (b) shows the typical hysteresis loops of the ribbon annealed at 360 °C and 420 °C, respectively, under different applied compressive stresses. The saturation magnetization (Ms) of the ribbon increases slightly with increasing annealing stress. We know that the Ms depends on the exchange coupling between atoms, which is directly related to the volume fraction of α -Fe grains. A large amount of precipitation of the crystalline phase induced by LPSa, resulting in a significant increase in Ms. Thus, it suggests that LPSa promotes the precipitation of the α -Fe phase causing a high volume fraction of the α -Fe grains, which is consistent with previous DSC curves. Besides, the magnetic anisotropy field Hk of the stress-annealed ribbon decreases with increasing annealing stress, which transforms into the easy axis magnetization direction. The α -Fe grains subjected to low-pressure stress induce the magnetization vector of the grains to be perpendicular to the stress direction, which generates magnetic induction anisotropy and forms an easy magnetization axis. Due to the decrease of Hk and the increase of Ms, the slope of M–H curves (magnetic susceptibility χ , χ =μ0 Ms/H) of the ribbon slightly increases with increasing annealing stress, indicating that compressive stress has a positive influence on the magnetization. Based on the principle of ferromagnetism, the anisotropy K in-
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Fig. 3. Surface state and typical AFM micrographs of stress-annealed ribbon. (a) Young’s modulus and Vickers hardness as a function of annealing stress. (b) and (c) Hardness distribution of the ribbon annealed at 360 °C and 420 °C, respectively, with different applied compressive stresses. (d)–(f) Young’s modulus and the corresponding AFM morphologies of the stress-annealed ribbon.
duced by nanocrystallization of the amorphous alloy under LPSa can be expressed as [13]:
3 K ≈ − λsσ υcr 2
magnetic anisotropy field is related to the anisotropy [14]:
Hk = (1)
where λs is the magnetostriction coefficient, σ is the applied low-pressure stress, and ν cr is the volume fraction of crystalline phase. Considering the easy axis magnetization direction, the
2Ku
μ0 Ms
(2)
where Ku is the magnetic anisotropy energy, and μ0 is the vacuum permeability. Besides, the relationship between the magnetostriction coefficient and the annealing stress can be described as [15]:
λs,σ = λs,0 − Bσ
(3)
10
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Fig. 4. Schematic of the crystallization behavior of Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 amorphous ribbon induced by (a) stress-free isothermal heat treatment and (b) low-pressure stress annealing (LPSa).
where B is the positive coefficient of order 10−10 MPa, λs ,0 is the magnetostriction constant under stress-free annealing, and λs,σ is the magnetostriction constant under stress annealing. It is noted that the λs of the sample decreases with increasing annealing stress. Consequently, LPSa is beneficial to reduce the K and low the λs , which contributes to the enhancement of μ and the decrease of Hc. Fig. 2(c) shows the Hc of the Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 ribbon as a function of annealing stress. At lower Ta ~340 °C, the Hc decreases slightly with increasing annealing stress, which is ascribed to the relaxation of internal stress. Under the stress-free annealing treatment, the Hc shows an increasing trend with increasing Ta from 360 °C to 500 °C. Interestingly, the Hc of the stress-annealed ribbon significantly decreases with increasing annealing stress and reaches the minimum value at 420 °C with an applied stress of 5 kPa. The magnetization process of materials is mainly achieved by domain wall displacement and magnetic moment rotation magnetization. This suggests that the Hc depends on the magnetocrystalline anisotropy and the internal stress fluctuation induced by the content and distribution of impurities. Due to the low density of nucleation sites and ultrahigh Fe content, the structural instability of the annealed ribbon causes the precipitation of multisized α -Fe nanoparticles during stress-free heat treatment, which greatly increases the magneto-crystalline anisotropy. In combination with Fig. 2(a) and (b), LPSa not only promotes the reduction of K and achieves a low λs, but also suppresses the rapid growth of nanocrystals, thus reducing Hc. This phenomenon is consistent with the stress-induced magnetic anisotropy reported by Lachowicz et al. [16], who pointed out that the residual amorphous phase is subjected to elastic strain inducing the sliding of the α -Fe grains, which results in a certain degree of deformation. This deformation changes the original orderly arrangement of atoms and causes the orientation of the similar atoms along the direction perpendicular to the stress. Based on the former result reported by Herzer [17], due to the magnetoelastic coupling effect, the volume shrinkage generated by the stress-induced crystallization leads to internal compressive stress in the α -Fe grains, which causes the magnetization vector to rotate in a direction perpendicular to the applied pressure stress, thus contributing to the improvement of soft magnetic properties.
Fig. 3(a) shows the Vickers hardness (Hv) and Young’s modulus (E) of the samples as a function of annealing stress. Obviously, the Hv and E of stress-annealed ribbon are significantly lower than that of stress-free annealed ribbon. Also, the Hv and E show a sharp downward trend with an increase in annealing stress. Compared with stress-free annealed ribbon, the Hv and E of stressannealed ribbon (~at 420 °C with applied 5 kPa) reach the minimum value of about ~821 Hv and 33.9 GPa, respectively, which are drastically reduced by about 33% and 51%, respectively. Fig. 3(b) and (c) shows the Hv distribution with different applied compressive stresses annealed at 360 °C and 420 °C, respectively. The Hv in the middle of the stress-free annealed ribbon is significantly greater than that on both sides, indicating the non-uniformity of the surface hardness. Besides, the value of Hv gradually decreases from the middle to the both sides, showing the phenomenon in which the middle is harder and the sides are softer. Interestingly, the higher Hv of the middle portion decreases sharply with an increase in compressive stresses during isothermal annealing, which promotes the uniform distribution of Hv. Fig. 3(d)–(f) shows the Young’s modulus distribution of the ribbon and the corresponding AFM morphology annealed at 420 °C with different applied compressive stresses. The E of the middle portion is slightly higher than that of the edge portion, which reflects the non-uniformity of the surface state in the stress-free annealed ribbon. The E significantly decreases with increasing annealing stress and the surface state of the stress-annealed ribbon transforms from inhomogeneity to homogenization. Consequently, LPSa is considered to be an effective way to adjust the inhomogeneity of nanocrystalline alloys during crystallization, which effectively reduces the Hv and E and promote homogenization distribution, thereby achieving a low value and uniform distribution of surface state. Based on the Hall–Petch relationship [18], LPSa promotes the formation of smaller α -Fe grains with a size of around 30 nm, which contributes to the achievement of a lower Hv and E [19]. Each part of the ribbon participates in the deformation induced by LPSa and exhibits a uniform viscous flow at the macroscopic level, which reduces the localization degree of the non-uniform deformation, thus promoting the uniform distribution of the surface state [20]. In the crystallization process, LPSa reduces the barrier of nucleation and promotes the precipitation of the α -Fe phase, which
J. Xu, Y. Yang and Q. Yan et al. / Scripta Materialia 179 (2020) 6–11
effectively controls the growth of crystal nuclei and reduces the growth rate of grains, thus contributing to the formation of finer and uniformly distributed nanoparticles [21]. Besides, LPSa induces the directional movement of the α -Fe phase and decreases the volume of nanovoids at the interface, which reduces the resistance of dislocation motion, thereby decreasing Hv and E [22]. Hence, LPSa plays a significant role in refining the grains and adjusting the nonuniform distribution of the surface state, which achieves the lower value and uniform distribution of Hv and E. Fig. 4(a) and (b) shows the schematic diagram of crystallization behavior in the ultrahigh Fe content Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 alloy induced by stress-free isothermal heat treatment and LPSa, respectively. For stress-free isothermal heat treatment, the crystallization behavior is mainly dependent on the interaction between atoms. Due to the presence of low-density nucleation sites (caused by low Cu content) [4] along with the absence of the diffusion barrier (influenced by large-sized Nb atoms), ultrahigh Fe content causes the rapid growth of nanoparticles and accelerates the crystallization process, which results in the formation of the largesized and non-uniform nanocrystalline structure. Considering that the high and inhomogeneity distribution of Hv and E originates from multi-sized mixed structure induced by the stress-free annealing, which also results in a higher Hc [23]. Under low-pressure stress, PLSa induces a large amount of the α -Fe grains to precipitate uniformly from the amorphous matrix and suppresses the rapid growth of nanoparticles, which reduces the number of holes and dislocations, thus increasing the total area of grain boundaries and resulting in softening [24]. Consequently, LPSa effectively improves soft magnetic properties to achieve high Ms and low Hc, and promotes the transformation of the surface state of the ribbon from higher value and inhomogeneity distribution of Hv and E to lower value and homogenization distribution of Hv and E. In summary, by applying a compressive stress during heat treatment, we demonstrate that PLSa promotes the precipitation of the α -Fe phase and refines the grain size in ultrahigh Fe content Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 nanocrystalline alloy, which contributes to the achievement of a high volume fraction, fine and uniform nanocrystalline structure. Besides, LPSa induces the transformation of hysteresis loop, which results in the reduction of magnetic anisotropy and the lower magnetostriction coefficient, thus decreasing Hc. The formation of finer and uniformly distributed α Fe nanoparticles without internal defects caused by LPSa promotes an increase in the total area of grain boundaries, which causes a lower value and uniform distribution of Hv and E, thereby resulting in softening of the ribbon. Through optimal LPSa, the nanocrystalline ribbon exhibits excellent soft magnetic properties (low Hc of ~6.7–8.2 A/m and high Ms of ~202.1–206.6 emu/g) and good manufacturability (uniform distribution and low Hv ~80 0–90 0 Hv and low E ~ 34–43 GPa). Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Acknowledgment This work is supported by National Natural Science Foundation of China (Nos: 50971046, 51575112), Natural Science Foundation of Guangdong (No: 2019A1515010886), Major Projects of Guangdong Basic and Applied Basic Research, NSFC-Guangdong Joint Fund Project (No: U1801259), and the Foshan Science and Technology Innovation Project (No: 2018IT100242). Supplementary material Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2019.12. 042. References [1] Y. Yoshizawa, S. Oguma, K. Yamauchi, J. Appl. Phys. 64 (1988) 6044–6046. [2] A. Makino, IEEE Trans. Magn. 48 (2012) 1331–1335. [3] A. Makino, T. Kubota, K. Yubuta, A. Inoue, A. Urata, H. Matsumoto, S. Yoshida, J. Appl. Phys. 109 (2011) 07A302. [4] L.X. Jiang, Y. Zhang, X. Tong, T. Suzuki, A. Makino, J. Magn. Magn. Mater. 471 (2019) 148–152. [5] A. Makino, H. Men, K. Yubuta, T. Kubota, J. Appl. Phys. 105 (2009) 07A308. [6] C.A. Schuh, T.C. Hufnagel, U. Ramamurty, Acta Mater. 55 (2007) 4067–4109. [7] M. Kuhnt, M. Marsilius, T. Strache, C. Polak, G. Herzer, Scr. Mater. 130 (2017) 46–48. [8] H.R. Lashgari, J.M. Cadogan, C. Kong, C. Tang, C. Doherty, D. Chu, S. Li, J. Magn. Magn. Mater. 456 (2018) 62–70. [9] P. Corte-Leon, V. Zhukova, M. Ipatov, J.M. Blanco, J. Gonzalez, M. Churyukanova, J.M. Bararibar, S. Taskaev, A. Zhukov, J. Alloy Compd. 789 (2019) 201–208. [10] X.Z. Fan, X.W. He, R.K. Nutor, R.M. Pan, J.J. Zheng, H.Q. Ye, F.M. Wu, J.Z. Jiang, Y.Z. Fang, J. Magn. Magn. Mater. 469 (2019) 349–353. [11] J.C. Baret, D. Vandembroucq, S. Roux, Phys. Rev. Lett. 89 (2002) 195506. [12] X.K. Xi, D.Q. Zhao, M.X. Pan, W.H. Wang, Y. Wu, J.J. Lewandowski, Phys. Rev. Lett. 94 (2005) 125510. [13] G. Herzer, IEEE Trans. Magn. 30 (1994) 4800–4802. [14] A. Siemko, H.K. Lachowicz, IEEE Trans. Magn. Mag. 23 (1987) 2563–2565. [15] J.M. Barandiaran, A. Hernando, V. Madurga, O.V. Nielsen, M. Vazquez, M. Vazquez-Lopez, Phys. Rev. B 35 (1987) 5066–5071. [16] H.K. Lachowicz, A. Neuweiler, F. Poplawski, E. Dynowska, J. Magn. Magn. Mater. 173 (1997) 287–294. [17] G. Herzer, Acta Mater. 61 (2013) 718–734. [18] S. Takaki, Mater. Sci. Forum 654 (2010) 11–16. [19] K. Lu, W.D. Wei, I.T. Wang, Scr. Metall. 24 (1991) 2319–2323. [20] Z. Wang, W.H. Wang, Natl. Sci. Rev. 6 (2018) 304–323. [21] L.Z. Zhao, R.J. Xue, Z.G. Zhu, Z. Lu, E. Axinte, W.H. Wang, H.Y. Bai, J. Appl. Phys. 116 (2014) 103516. [22] J.C. Qiao, Q. Wang, J.M. Pelletier, H. Kato, R. Casalini, D. Crespo, E. Pineda, Y. Yao, Y. Yang, Prog. Mater. Sci. 104 (2019) 250–329. [23] E. Lopatina, I. Soldatov, V. Budinsky, M. Marsilius, L. Schultz, G. Herzer, R. Schäfer, Acta Mater. 96 (2015) 10–17. [24] K. Kosiba, D. Sopu, S. Scudino, L. Zhang, J. Bednarcik, S. Pauly, Int. J. Plast. 119 (2019) 156–170.
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Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Softening and magnetic properties of ultrahigh Fe content FeSiBCuPC nanocrystalline alloy induced by low-pressure stress annealing Jia Xu a,b, Yuanzheng Yang a,∗, Qiusheng Yan b,∗, Guihua Xiao a, Ting Luo a, Chenfeng Fan a a b
Faculty of Materials and Energy, Guangdong University of technology, Guangzhou 510006, China Faculty of Electromechanical Engineering, Guangdong University of technology, Guangzhou 510006, China
a r t i c l e
i n f o
Article history: Received 18 October 2019 Revised 23 December 2019 Accepted 29 December 2019
Keywords: Fe-based nanocrystalline alloys Low-pressure stress annealing Soft magnetic properties Hardness Young’s modulus
a b s t r a c t Due to the rapid growth of precipitated nanoparticles in ultrahigh Fe content nanocrystalline alloys, we propose a new method of low-pressure stress annealing (LPSa) to control crystallization behavior to achieve excellent magnetic properties and mechanical properties of Fe85.7 Si0.5 B9.3 Cu0.7 P3.5 C0.3 nanocrystalline alloy. It demonstrates that LPSa promotes the precipitation of α -Fe phase and suppresses the growth of nanoparticles, which obtains a higher saturation magnetization of 202.1–206.6 emu/g and a lower coercivity of 6.7–8.2 A/m. Besides, the formation of fine and uniform nanocrystalline structure induced by LPSa achieves the lower value and uniform distribution of hardness and Young’s modulus, resulting in softening.
Although nanocrystalline FeSiBPCu alloys with ultrahigh Fe content (≥ 85 at%) exhibit remarkable soft magnetic properties, brittleness severely restricts their broad application as functional materials [1–3]. Due to the absence of high melting point alloying elements and only a trace amount of Cu content, the α -Fe grains are prone to grow rapidly and non-uniformly during crystallization [4], resulting in high coercivity Hc and low plasticity. Previously, Makino et al. [5] demonstrated that the crystallization process of Fe85-86 SiBPCu alloys is very strict, requiring higher heating rate (≥ 200 °C/min) and shorter holding time (≤10 min) to achieve excellent soft magnetic properties, which seriously hinders the industrial production of high Fe content FeSiBPCu alloys. The plastic deformation ability can be tailored by changing the surface stress state of alloys [6]. Pressure as an important thermodynamic and kinetic parameter promotes the crystallization of amorphous alloys [7,8]. Unlike previous tensile stresses annealing [9,10], the crystallization of amorphous alloys by applying a compressive stress during the heat treatment promotes the precipitation of α -Fe phase and reduce the growth rate of nanoparticles [11,12]. Besides, LPSa continuously and uniformly regulates the energy state of nanocrystalline alloys, which has the advantages of fine-tuning of tissue structure, non-destructive and uniformity. In this paper, intending to suppress the rapid growth of nanocrystals and improve the ability of plastic deformation, we investigated the crystallization behavior, magnetic properties and mechanical prop∗
Corresponding authors. E-mail addresses: [email protected] (Y. Yang), [email protected] (Q. Yan).
https://doi.org/10.1016/j.scriptamat.2019.12.042 1359-6462/© 2020 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
© 2020 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
erties of ultrahigh Fe content Fe85.7 Si0.5 B9.3 Cu0.7 P3.5 C0.3 nanocrystalline alloy induced by low-pressure stress annealing (LPSa). Multi-component Fe85.7 Si0.5 B9.3 Cu0.7 P3.5 C0.3 alloy ingot was prepared by arc melting with the mixtures of raw materials in a high purity argon atmosphere. Ribbon with a width of 1 mm and a thickness of 20 μm was prepared by a single-roller melt spinning method. The diameter and width of the copper wheel are 320 mm and 50 mm, respectively. The structure and thermal stability of the ribbon was examined by X-ray diffraction (XRD) with Cu Kα radiation and differential scanning calorimetry (DSC) at a heating rate of 20 °C/min under nitrogen atmosphere, respectively. Heat treatment of amorphous ribbon was carried out in a tubular furnace with a constant heat rating of 20 °C/min. The saturation magnetization (Ms) and coercivity (Hc) were measured with a vibrating sample magnetometer (VSM) under an applied field of 12,500 Oe and a DC B-H loop tracer under a maximum applied field of 800 A/m, respectively. The hardness and Young’s modulus were measured with a Vickers hardness tester under a load of 0.49 N for 10 s and an atomic force microscope (AFM), respectively. To ensure the uniformity of the surface state of the ribbon, the free surface of the melt-spun and annealed ribbon with lower roughness was chosen for XRD, hardness and Young’s modulus measurements. The corresponding hardness distribution map is drawn by multi-point isometric hardness measurement. The sample is placed on a smooth and flat quartz glass slide as a substrate, and another glass slide is pressed in parallel on the sample to maintain a uniform force on the upper and lower sides
J. Xu, Y. Yang and Q. Yan et al. / Scripta Materialia 179 (2020) 6–11
7
Fig. 1. Appearance of the Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 nanocrystalline ribbon under different applied compressive stresses (a) annealed at 360 °C and (b) annealed at 420 °C. (c) and (d) the XRD patterns of stress-annealed ribbon. (e) the DSC curves of the ribbon after annealed at 420 °C under different applied compressive stresses. (f) Compressive stress annealing dependence of the α -Fe grain size for the Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 ribbon.
of the ribbon. Then, it is placed in parallel in a tubular furnace for crystallization treatment. The pressure is calculated by combining the weight of glass slides and the surface area of the ribbon. The stress is adjusted by controlling the loading weight of the slides loaded on the ribbon. Based on the XRD pattern (presented in Fig. S1a), only a broad peak without any sharp diffraction peaks can be detected, indicating the typical amorphous structure. According to the DSC curve
(shown in Fig. S1b), the onset temperature of the first crystallization peak is at Tx1 ≈ 388 °C, while the secondary crystallization peak is at Tx2 ≈ 521 °C. The XRD patterns of the annealed ribbon is shown in Fig. S2. Although the ribbon has a wide ࢞T (࢞T=Tx2 Tx1 ~133 °C), the instability of nanocrystalline structure leads to premature precipitation of the secondary phase. Fig. 1(a) and (b) shows the photo images of the annealed sample with different compressive stresses applied along the ribbon
8
J. Xu, Y. Yang and Q. Yan et al. / Scripta Materialia 179 (2020) 6–11
Fig. 2. Hysteresis loops of Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 ribbon annealed at (a) 360 °C and (b) 420 °C for 15 min with different applied compressive stresses. (c) Variation in coercivity with annealing stress for the Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 ribbon.
plane during isothermal annealing. When annealed at 360 °C with stress-free, the ribbon exhibits a curled shape, which is similar to a ring. While the annealed ribbon becomes noticeably smooth with different applied low-pressure stresses, and the curvature decreases with increasing annealing stress. When annealed at 420 °C, the shape of the stress-annealed ribbon presents a straight line, indicating that the curl deformation behavior of ribbon is mainly dependent on the applied-pressure. Fig. 1(c) and (d) shows the XRD patterns of the ribbon annealed at 360 °C and 420 °C, respectively, under different applied pressure stresses. The intensity of the crystallization peak increases slightly with increasing annealing stress, which correlates with a higher volume fraction of the α -Fe phase. Fig. 1(e) shows the DSC curves of the ribbon after annealed at 420 °C with various applied-pressures. The area of exothermic peak correlates with the volume fraction of crystalline phase. Interestingly, the stress-annealed ribbon presents a straight line corresponding to the first exothermic peak of the initial crystalline phase, while the stress-free annealed ribbon still exhibits a small hill (࢞H~4.2 J/g) associated with the incomplete precipitation of α -Fe phase. Fig. 1(f) shows the applied pressure stress dependence of an average grain size (D) of the α -Fe phase estimated by Scherrer‘s equation. The D of stress-annealed ribbon is smaller than that of stress-free annealed ribbon. Besides, the D shows a significant decreasing trend with increasing annealing stress and is effectively controlled within 30 nm, indicating that PLSa contributes to the reduction of D. Consequently, LPSa effectively
promotes the precipitation of α -Fe grains and suppresses the rapid growth of nanoparticles in ultrahigh Fe content nanocrystalline alloy, which achieves the high volume fraction of α -Fe grains and maintains a finer nanocrystalline structure. Fig. 2(a) and (b) shows the typical hysteresis loops of the ribbon annealed at 360 °C and 420 °C, respectively, under different applied compressive stresses. The saturation magnetization (Ms) of the ribbon increases slightly with increasing annealing stress. We know that the Ms depends on the exchange coupling between atoms, which is directly related to the volume fraction of α -Fe grains. A large amount of precipitation of the crystalline phase induced by LPSa, resulting in a significant increase in Ms. Thus, it suggests that LPSa promotes the precipitation of the α -Fe phase causing a high volume fraction of the α -Fe grains, which is consistent with previous DSC curves. Besides, the magnetic anisotropy field Hk of the stress-annealed ribbon decreases with increasing annealing stress, which transforms into the easy axis magnetization direction. The α -Fe grains subjected to low-pressure stress induce the magnetization vector of the grains to be perpendicular to the stress direction, which generates magnetic induction anisotropy and forms an easy magnetization axis. Due to the decrease of Hk and the increase of Ms, the slope of M–H curves (magnetic susceptibility χ , χ =μ0 Ms/H) of the ribbon slightly increases with increasing annealing stress, indicating that compressive stress has a positive influence on the magnetization. Based on the principle of ferromagnetism, the anisotropy K in-
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Fig. 3. Surface state and typical AFM micrographs of stress-annealed ribbon. (a) Young’s modulus and Vickers hardness as a function of annealing stress. (b) and (c) Hardness distribution of the ribbon annealed at 360 °C and 420 °C, respectively, with different applied compressive stresses. (d)–(f) Young’s modulus and the corresponding AFM morphologies of the stress-annealed ribbon.
duced by nanocrystallization of the amorphous alloy under LPSa can be expressed as [13]:
3 K ≈ − λsσ υcr 2
magnetic anisotropy field is related to the anisotropy [14]:
Hk = (1)
where λs is the magnetostriction coefficient, σ is the applied low-pressure stress, and ν cr is the volume fraction of crystalline phase. Considering the easy axis magnetization direction, the
2Ku
μ0 Ms
(2)
where Ku is the magnetic anisotropy energy, and μ0 is the vacuum permeability. Besides, the relationship between the magnetostriction coefficient and the annealing stress can be described as [15]:
λs,σ = λs,0 − Bσ
(3)
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Fig. 4. Schematic of the crystallization behavior of Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 amorphous ribbon induced by (a) stress-free isothermal heat treatment and (b) low-pressure stress annealing (LPSa).
where B is the positive coefficient of order 10−10 MPa, λs ,0 is the magnetostriction constant under stress-free annealing, and λs,σ is the magnetostriction constant under stress annealing. It is noted that the λs of the sample decreases with increasing annealing stress. Consequently, LPSa is beneficial to reduce the K and low the λs , which contributes to the enhancement of μ and the decrease of Hc. Fig. 2(c) shows the Hc of the Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 ribbon as a function of annealing stress. At lower Ta ~340 °C, the Hc decreases slightly with increasing annealing stress, which is ascribed to the relaxation of internal stress. Under the stress-free annealing treatment, the Hc shows an increasing trend with increasing Ta from 360 °C to 500 °C. Interestingly, the Hc of the stress-annealed ribbon significantly decreases with increasing annealing stress and reaches the minimum value at 420 °C with an applied stress of 5 kPa. The magnetization process of materials is mainly achieved by domain wall displacement and magnetic moment rotation magnetization. This suggests that the Hc depends on the magnetocrystalline anisotropy and the internal stress fluctuation induced by the content and distribution of impurities. Due to the low density of nucleation sites and ultrahigh Fe content, the structural instability of the annealed ribbon causes the precipitation of multisized α -Fe nanoparticles during stress-free heat treatment, which greatly increases the magneto-crystalline anisotropy. In combination with Fig. 2(a) and (b), LPSa not only promotes the reduction of K and achieves a low λs, but also suppresses the rapid growth of nanocrystals, thus reducing Hc. This phenomenon is consistent with the stress-induced magnetic anisotropy reported by Lachowicz et al. [16], who pointed out that the residual amorphous phase is subjected to elastic strain inducing the sliding of the α -Fe grains, which results in a certain degree of deformation. This deformation changes the original orderly arrangement of atoms and causes the orientation of the similar atoms along the direction perpendicular to the stress. Based on the former result reported by Herzer [17], due to the magnetoelastic coupling effect, the volume shrinkage generated by the stress-induced crystallization leads to internal compressive stress in the α -Fe grains, which causes the magnetization vector to rotate in a direction perpendicular to the applied pressure stress, thus contributing to the improvement of soft magnetic properties.
Fig. 3(a) shows the Vickers hardness (Hv) and Young’s modulus (E) of the samples as a function of annealing stress. Obviously, the Hv and E of stress-annealed ribbon are significantly lower than that of stress-free annealed ribbon. Also, the Hv and E show a sharp downward trend with an increase in annealing stress. Compared with stress-free annealed ribbon, the Hv and E of stressannealed ribbon (~at 420 °C with applied 5 kPa) reach the minimum value of about ~821 Hv and 33.9 GPa, respectively, which are drastically reduced by about 33% and 51%, respectively. Fig. 3(b) and (c) shows the Hv distribution with different applied compressive stresses annealed at 360 °C and 420 °C, respectively. The Hv in the middle of the stress-free annealed ribbon is significantly greater than that on both sides, indicating the non-uniformity of the surface hardness. Besides, the value of Hv gradually decreases from the middle to the both sides, showing the phenomenon in which the middle is harder and the sides are softer. Interestingly, the higher Hv of the middle portion decreases sharply with an increase in compressive stresses during isothermal annealing, which promotes the uniform distribution of Hv. Fig. 3(d)–(f) shows the Young’s modulus distribution of the ribbon and the corresponding AFM morphology annealed at 420 °C with different applied compressive stresses. The E of the middle portion is slightly higher than that of the edge portion, which reflects the non-uniformity of the surface state in the stress-free annealed ribbon. The E significantly decreases with increasing annealing stress and the surface state of the stress-annealed ribbon transforms from inhomogeneity to homogenization. Consequently, LPSa is considered to be an effective way to adjust the inhomogeneity of nanocrystalline alloys during crystallization, which effectively reduces the Hv and E and promote homogenization distribution, thereby achieving a low value and uniform distribution of surface state. Based on the Hall–Petch relationship [18], LPSa promotes the formation of smaller α -Fe grains with a size of around 30 nm, which contributes to the achievement of a lower Hv and E [19]. Each part of the ribbon participates in the deformation induced by LPSa and exhibits a uniform viscous flow at the macroscopic level, which reduces the localization degree of the non-uniform deformation, thus promoting the uniform distribution of the surface state [20]. In the crystallization process, LPSa reduces the barrier of nucleation and promotes the precipitation of the α -Fe phase, which
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effectively controls the growth of crystal nuclei and reduces the growth rate of grains, thus contributing to the formation of finer and uniformly distributed nanoparticles [21]. Besides, LPSa induces the directional movement of the α -Fe phase and decreases the volume of nanovoids at the interface, which reduces the resistance of dislocation motion, thereby decreasing Hv and E [22]. Hence, LPSa plays a significant role in refining the grains and adjusting the nonuniform distribution of the surface state, which achieves the lower value and uniform distribution of Hv and E. Fig. 4(a) and (b) shows the schematic diagram of crystallization behavior in the ultrahigh Fe content Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 alloy induced by stress-free isothermal heat treatment and LPSa, respectively. For stress-free isothermal heat treatment, the crystallization behavior is mainly dependent on the interaction between atoms. Due to the presence of low-density nucleation sites (caused by low Cu content) [4] along with the absence of the diffusion barrier (influenced by large-sized Nb atoms), ultrahigh Fe content causes the rapid growth of nanoparticles and accelerates the crystallization process, which results in the formation of the largesized and non-uniform nanocrystalline structure. Considering that the high and inhomogeneity distribution of Hv and E originates from multi-sized mixed structure induced by the stress-free annealing, which also results in a higher Hc [23]. Under low-pressure stress, PLSa induces a large amount of the α -Fe grains to precipitate uniformly from the amorphous matrix and suppresses the rapid growth of nanoparticles, which reduces the number of holes and dislocations, thus increasing the total area of grain boundaries and resulting in softening [24]. Consequently, LPSa effectively improves soft magnetic properties to achieve high Ms and low Hc, and promotes the transformation of the surface state of the ribbon from higher value and inhomogeneity distribution of Hv and E to lower value and homogenization distribution of Hv and E. In summary, by applying a compressive stress during heat treatment, we demonstrate that PLSa promotes the precipitation of the α -Fe phase and refines the grain size in ultrahigh Fe content Fe85.7 Si0.5 B9.3 P3.5 Cu0.7 C0.3 nanocrystalline alloy, which contributes to the achievement of a high volume fraction, fine and uniform nanocrystalline structure. Besides, LPSa induces the transformation of hysteresis loop, which results in the reduction of magnetic anisotropy and the lower magnetostriction coefficient, thus decreasing Hc. The formation of finer and uniformly distributed α Fe nanoparticles without internal defects caused by LPSa promotes an increase in the total area of grain boundaries, which causes a lower value and uniform distribution of Hv and E, thereby resulting in softening of the ribbon. Through optimal LPSa, the nanocrystalline ribbon exhibits excellent soft magnetic properties (low Hc of ~6.7–8.2 A/m and high Ms of ~202.1–206.6 emu/g) and good manufacturability (uniform distribution and low Hv ~80 0–90 0 Hv and low E ~ 34–43 GPa). Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Acknowledgment This work is supported by National Natural Science Foundation of China (Nos: 50971046, 51575112), Natural Science Foundation of Guangdong (No: 2019A1515010886), Major Projects of Guangdong Basic and Applied Basic Research, NSFC-Guangdong Joint Fund Project (No: U1801259), and the Foshan Science and Technology Innovation Project (No: 2018IT100242). Supplementary material Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2019.12. 042. References [1] Y. Yoshizawa, S. Oguma, K. Yamauchi, J. Appl. Phys. 64 (1988) 6044–6046. [2] A. Makino, IEEE Trans. Magn. 48 (2012) 1331–1335. [3] A. Makino, T. Kubota, K. Yubuta, A. Inoue, A. Urata, H. Matsumoto, S. Yoshida, J. Appl. Phys. 109 (2011) 07A302. [4] L.X. Jiang, Y. Zhang, X. Tong, T. Suzuki, A. Makino, J. Magn. Magn. Mater. 471 (2019) 148–152. [5] A. Makino, H. Men, K. Yubuta, T. Kubota, J. Appl. Phys. 105 (2009) 07A308. [6] C.A. Schuh, T.C. Hufnagel, U. Ramamurty, Acta Mater. 55 (2007) 4067–4109. [7] M. Kuhnt, M. Marsilius, T. Strache, C. Polak, G. Herzer, Scr. Mater. 130 (2017) 46–48. [8] H.R. Lashgari, J.M. Cadogan, C. Kong, C. Tang, C. Doherty, D. Chu, S. Li, J. Magn. Magn. Mater. 456 (2018) 62–70. [9] P. Corte-Leon, V. Zhukova, M. Ipatov, J.M. Blanco, J. Gonzalez, M. Churyukanova, J.M. Bararibar, S. Taskaev, A. Zhukov, J. Alloy Compd. 789 (2019) 201–208. [10] X.Z. Fan, X.W. He, R.K. Nutor, R.M. Pan, J.J. Zheng, H.Q. Ye, F.M. Wu, J.Z. Jiang, Y.Z. Fang, J. Magn. Magn. Mater. 469 (2019) 349–353. [11] J.C. Baret, D. Vandembroucq, S. Roux, Phys. Rev. Lett. 89 (2002) 195506. [12] X.K. Xi, D.Q. Zhao, M.X. Pan, W.H. Wang, Y. Wu, J.J. Lewandowski, Phys. Rev. Lett. 94 (2005) 125510. [13] G. Herzer, IEEE Trans. Magn. 30 (1994) 4800–4802. [14] A. Siemko, H.K. Lachowicz, IEEE Trans. Magn. Mag. 23 (1987) 2563–2565. [15] J.M. Barandiaran, A. Hernando, V. Madurga, O.V. Nielsen, M. Vazquez, M. Vazquez-Lopez, Phys. Rev. B 35 (1987) 5066–5071. [16] H.K. Lachowicz, A. Neuweiler, F. Poplawski, E. Dynowska, J. Magn. Magn. Mater. 173 (1997) 287–294. [17] G. Herzer, Acta Mater. 61 (2013) 718–734. [18] S. Takaki, Mater. Sci. Forum 654 (2010) 11–16. [19] K. Lu, W.D. Wei, I.T. Wang, Scr. Metall. 24 (1991) 2319–2323. [20] Z. Wang, W.H. Wang, Natl. Sci. Rev. 6 (2018) 304–323. [21] L.Z. Zhao, R.J. Xue, Z.G. Zhu, Z. Lu, E. Axinte, W.H. Wang, H.Y. Bai, J. Appl. Phys. 116 (2014) 103516. [22] J.C. Qiao, Q. Wang, J.M. Pelletier, H. Kato, R. Casalini, D. Crespo, E. Pineda, Y. Yao, Y. Yang, Prog. Mater. Sci. 104 (2019) 250–329. [23] E. Lopatina, I. Soldatov, V. Budinsky, M. Marsilius, L. Schultz, G. Herzer, R. Schäfer, Acta Mater. 96 (2015) 10–17. [24] K. Kosiba, D. Sopu, S. Scudino, L. Zhang, J. Bednarcik, S. Pauly, Int. J. Plast. 119 (2019) 156–170.