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Thin Solid Films 519 (2011) 8312–8316

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Thin Solid Films j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / t s f

Effect of mechanical milling and heat treatment on the structure and magnetic properties of gas atomized Mn–Al alloy powders Jung-Goo Lee a,⁎, Xiao-Lei Wang a,b, Zhi-Dong Zhang b, ChuI-Jin Choi a a

Functional Materials Division, Korea Institute of Materials Science, 531 Changwondaero, Changwon 631-831, Republic of Korea Shenyang National Laboratory for Materials Science, Institute of Metal Research and International Centre for Materials Physics, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, People's Republic of China b

a r t i c l e

i n f o

Available online 1 April 2011 Keywords: Mn–Al powders Mechanical milling Annealing Magnetic properties

a b s t r a c t Ferromagnetic Mn–Al alloy powders were fabricated by mechanical milling and heat treatment with gasatomized powders. Different processes, i.e., heat treatment before ball milling and ball milling before heat treatment, result in different microstructures and magnetic properties of the powders. It was found that Hc increased and Mr decreased with the size reduction regardless of the sequence of heat treatment and ball milling. However, tendency of the change in Hc and Mr depended on the sequence. Further annealing of the powders ball-milled after heat treatment resulted in slight decrease of Hc and large increase of Mr. The magnetic properties, Mr = 41.2 emu/g, Hc = 3.1 kOe, were obtained from the powders ball-milled for 5 h after heat treatment at 650 °C for 20 min, and subsequent annealing at 280 °C for 20 min. © 2011 Elsevier B.V. All rights reserved.

1. Introduction Recently, high-performance permanent magnets such as Nd–Fe–B, Sm–Co have attracted great attention due to the need for these magnets in new markets such as the hybrid car and the electric vehicle. However, most of them are made from rare earth elements such as Nd, Sm, and Dy which are getting more expensive. On the other hand, Mn–Al alloy is made from cheap elements but shows good magnetic properties superior to the well-known Alnicos and hard ferrites [1]. It is reported that the theoretical maximum energy product (BH)max and magnetocrystalline anisotropy field of Mn–Al alloy is 12.64 MGOe and 38 kOe, respectively [2]. In addition, Mn–Al alloy shows good machinability and corrosion resistance [1]. Accordingly, Mn–Al alloy is recently reconsidered as an important material for permanent magnets. The excellent magnetic properties of the Mn–Al system are attributed to the metastable τ-phase in which the Mn atoms exist at eight apexes in L10 superstructure [3]. In general, the τ-phase Mn–Al alloy has been produced by quenching the high temperature ε-phase followed by a proper heat treatment, or by cooling the ε-phase at a controllable rate [4,5]. It is believed that the magnetic properties of Mn–Al alloy are mostly dependent on the weight ratio of raw material, preparation approaches and sequent annealing process. Up to now, various fabrication methods have been

⁎ Corresponding author. Tel./fax: + 82 55 280 3606/3392. E-mail address: [email protected] (J.-G. Lee). 0040-6090/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2011.03.094

employed to obtain τ-phase Mn–Al alloy, such as mechanical alloying [6], magnetron sputtering [4], melt spinning [7], plasma arc discharge [8] and atomization [9]. Among these methods, the atomization method is the most economical one for mass production of alloy powders [9,10]. The atomized powders usually have a broad size distribution, which can affect not only the condition for phase transition from ε to τ-phase but also the magnetic properties [11]. And the size of the atomized powders is bigger than the single domain size, which means that the magnetic properties of the atomized powders could be improved by the size reduction. However, there have been few reports on the relationship between the particle size and the magnetic properties of the atomized Mn–Al powders. In the present study, we have investigated the change in size and magnetic properties of the gas atomized Mn–Al powders by mechanical milling and heat treatment.

2. Experimental procedure High-purity (N99.9%) Mn–30 wt.% Al was induction melted into a master alloy, which was then atomized in nitrogen gas at a pressure of 60 kg f/mm2 to obtain ε-phase Mn–Al powders. The obtained powders were sieved to different sized powders and the powders with a diameter of 25–38 μm were used. In the present study, two different processes were employed. In the first process, the atomized powders were first annealed in vacuum at 650 °C for 20 min to obtain τ-phase Mn–Al alloy powders. The powders were then poured into a cylindrical container together with balls of 5 mm

J.-G. Lee et al. / Thin Solid Films 519 (2011) 8312–8316

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in diameter in a powder-to-ball weight ratio of 1:20. The container was filled with 40 mL hexane and sealed under an argon atmosphere to limit oxidation. The sample was milled for 5, 10, 20, and 26 h (hereafter referred to as P1). In the second process, the atomized powders were first ball milled at the same condition as mentioned above and then annealed in vacuum at 650 °C for 20 min (hereafter referred to as P2). Finally, the powders treated by P1 process were further annealed in vacuum at 280 °C for 20 min. The structure and particle size of the obtained powders were examined by X-ray diffraction (XRD) with Cu Ka radiation and scanning electron microscopy (SEM), respectively. The magnetic properties of the powders were measured with a vibrating sample magnetometer (VSM) with an external magnetic induction field up to 10 kOe. No demagnetization correction was made for the magnetic measurements. 3. Results and discussion Fig. 1 shows typical SEM image and XRD pattern of the gasatomized Mn–Al powders. The gas-atomized Mn–Al powders exhibit spherical shape and board size distribution as shown in Fig. 1(a). And it was confirmed that the gas-atomized Mn–Al powders are ε-phase which is a high temperature phase in the Mn– Al binary system as shown in Fig. 1(b). Fig. 2 shows XRD patterns of the powders obtained by the two different processes as mentioned above. From Fig. 2, it is obvious that the two different processes result in the different phase evolution in the powders. As shown in Fig. 2(a), the powders treated by P1 process mostly consist of the metastable τ-phase and there is no phase change during ball milling for 26 h. However, it is noteworthy that the diffraction peaks were slightly broadened by ball milling, indicating that ball milling process induced the strain inside the powders. On the other hand, the powders treated by P2 process were gradually decomposed into the equilibrium γ2 + β phases with an increase in ball milling time as shown in Fig. 2(b). It is because that the defects or strain induced by balling milling could facilitate the phase decomposition from the metastable τ-phase into the equilibrium γ2 + β phases during subsequent heat treatment. The typical Mn–Al powders treated by P1 process were shown in Fig. 3. It can be observed that the particle size gradually decreased with increasing the ball milling time up to 26 h. The average particle size can be estimated to be ca. 8.8, 5.8, 4.3, and 4.0 μm, corresponding to the ball milling time 5, 10, 20, and 26 h, respectively. Fig. 4 shows the dependence of coercivity (Hc) and remanence (Mr) on the ball milling time for the Mn–Al powders prepared by the

Fig. 2. XRD patterns of the gas-atomized Mn–Al alloy powders treated by different two processes: (a) annealed at 650 °C for 20 min before ball milling for different times; (b) ball milled for different times before annealing at 650 °C for 20 min.

two different processes as mentioned above. Regardless of the different experimental process, Hc increased and Mr decreased with an increase in ball milling time. It is now well known that the

Fig. 1. Typical SEM image and XRD pattern of gas-atomized Mn–Al alloy powders.

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Fig. 3. Typical SEM images of Mn–Al alloy powders annealed at 650 °C for 20 min before ball milling for different times: (a) 5 h, (b) 10 h, (c) 20 h, and (d) 26 h.

coercivity increases, goes through a maximum with decreasing the particle size and then tends toward zero [12]. According to Kittel's approach, the critical single domain size, Dc is 1.4γ/M2s , where γ is the domain wall energy per unit area and Ms is the saturation magnetization [13]. With regard to the spherical τ-phase particle, the critical size for single-domain is about 710 nm [14]. In the present study, the particle size gradually decreased down to 4.0 μm after ball milling, which is bigger than the single-domain size. So it is easily expected that Hc would increase by ball milling. It should be noticed that Hc of the powders treated by P2 process is higher than that of the powders treated by P1 process as shown in Fig. 4(a). It is because that Hc is dependent on not only the particle size but also its microstructure. The powders treated by P2 process have nonmagnetic γ2 and β phases which could raise coercivity by domain walls pinning [1,15–17]. On the other hand, the powders treated by P1 process show higher Mr than that of the powders treated by P2 process as shown in Fig. 4 (b). It is because the magnetization is greatly dependent on the amount of τ-phase and the powders treated by P2 process were gradually decomposed into the equilibrium γ2 + β phases with ball milling time as shown in Fig. 2(b). It is expected that Mr would slightly increase with decreasing the particle size because the reduction of the particle size improves magnetic anisotropy [18]. However, the Mr of the powders treated by P1 process still decreased although no phase decomposition took place. One possible reason is that a lot of strain

and/or defect could be accumulated inside the powders after ball milling which deteriorate squareness of hysteresis loop and decrease Mr [19]. To make clear this point, the powders treated by P1 process were subjected to further annealing at a relatively low temperature. Fig. 5 shows the dependence of Hc and Mr on the ball milling time for the Mn–Al powders treated by P1 process with or without further annealing at 280 °C for 20 min. As shown in Fig. 5, it is obvious that further annealing at 280 °C for 20 min has a significant influence on the magnetic properties of the powders. Hc slightly decreased but Mr greatly increased by further annealing. The increase in Mr is ca. 7, 10.2, 10, and 10.1 emu/g corresponding to ball milling time 5, 10, 20, and 26 h, respectively. These results indicate that further annealing may release the strain or defect from the powder, which is beneficial to the enhancement of Mr. The optimal magnetic properties, M r = 41.2 emu/g, H c = 3.1 kOe, (BH) m a x = 1.59 MGOe, were obtained from the powders ball-milled for 5 h and further annealing at 280 °C for 20 min. The corresponding M–H loop is shown in Fig. 6. Magnetic properties of the Mn–Al alloy prepared by different methods are shown in Table 1. Although the Mr of the powders increased by the further annealing process, it is still lower than that without ball milling except for one point. It can be expected that the degree of strain and/or defect accumulation is dependent on the milling time, which means that different annealing conditions should be applied to the powder depending on the milling time. The


Fig. 4. Dependence of Hc and Mr on the ball milling time for the Mn–Al alloy powders treated by P1 and P2 processes.

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Fig. 5. Dependence of Hc and Mr on the ball milling time for the Mn–Al alloy powder annealed at 650 °C for 20 min before ball milling with or without further annealing at 280 °C for 20 min.

more detailed experiments on the further annealing with the ballmilled Mn–Al powders are in progress. 4. Conclusions The gas-atomized Mn–Al powders were subjected to ball milling and heat treatment. Depending on the sequence of ball milling and heat treatment, different microstructures and magnetic properties were confirmed. Regardless of the sequence, Hc increased and Mr decreased with an increase in ball milling time. However, Hc of the powders heat-treated after ball milling was higher than that of the powders ball-milled after heat treatment. With regard to Mr the tendency was reversed. Further annealing of the powders ball-milled after heat treatment resulted in slight decrease of Hc and large increase of M r . The magnetic properties, M r = 41.2 emu/g, Hc = 3.1 kOe, were obtained from the powders ball-milled for 5 h after heat treatment at 650 °C for 20 min, and subsequent annealing at 280 °C for 20 min. Acknowledgment This research was supported by a grant from the Fundamental R&D Program for Core Technology of Materials funded by the Ministry of Knowledge Economy, Republic of Korea.

Fig. 6. Typical M–H loop of the samples ball-milled for 5 h after heat treatment at 650 °C for 20 min and subsequently annealed at 280 °C for 20 min.

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Table 1 Magnetic properties of the Mn–Al alloy prepared by different methods. Mr (emu/g)

Hc (kOe)

Preparation method

Source

45 49 8.8 41.2

4.8 1.78 5.6 3.1

Mechanical milling Melt spinning Plasma arc discharge Gas-atomization and ball milling

[1] [7] [8] In this work

[5] [6] [7] [8] [9] [10] [11] [12] [13] [14]

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