IEEE TRANSACTIONS ON MAGNETICS, VOL. 50, NO. 11, NOVEMBER 2014
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Strong Perpendicular Uniaxial Magnetic Anisotropy in Tetragonal Fe0.5 Co0.5 Films of Artificially Ordered B2 State Bin Lao, Jin Won Jung, and Masashi Sahashi Department of Electronic Engineering, Tohoku University, Sendai 980-8579, Japan We studied the relationship between the magnetic anisotropy properties and crystal structure of an artificial B2 state of Fe0.5 Co0.5 epitaxial films on a Rh(001) seed layer with a tetragonal distortion structure, which exhibited large perpendicular uniaxial magnetic anisotropy over a wide range of thicknesses and c/a ratios at room temperature. We obtained the bulk anisotropy energy of the Fe0.5 Co0.5 films, which reached 2.5 × 107 erg/cm3 at c/a = 1.20. The magnetization easy axis of the Fe0.5 Co0.5 films transferred from out-of-plane to in-plane when the thickness was increased from 24 to 30 monolayers. This is due to the lattice relaxation caused by the increase of the total thickness. We observed that the cubic anisotropy component around the in-plane of the films was much smaller than the perpendicular effective magnetic anisotropy, which demonstrates that our films exhibited the uniaxial anisotropy property. Furthermore, we found the bulk anisotropy values of our artificial B2-state Fe0.5 Co0.5 to be larger than those of the A2 state under the same lattice distortion. Thus, we suggest that tuning the degree of crystal order of FeCo alloys is an effective way to experimentally enhance their perpendicular uniaxial magnetic anisotropy. Index Terms— B2 state Fe0.5 Co0.5 , critical thickness 3.4 nm, high K u without noble elements.
I. I NTRODUCTION
M
AGNETIC material films exhibiting large perpendicular uniaxial magnetic anisotropy (K u ) at room temperature (RT) are not only of fundamental interest in the field of magnetism, but also technologically important for magnetic applications, such as magnetic recording media and highfrequency magnetic devices [1], [2]. The uniaxial magnetic anisotropy property, can generally be explained in terms of spin-orbit coupling (SOC) being induced close to the Fermi level (E F ) because the energy separation of the electronic states (eg , t2g ) is reduced [3]. It is well known that a large perpendicular K u is characteristic of bi-metallic alloy systems, which consist of ferromagnetic 3d metals (Fe, Co) and 4d/5d noble elements (nonmagnetic) due to their strong SOC [4]. Much attention has recently been given to FeCo alloys with a tetragonal distortion structure; these are conventional 3d alloys without 4d/5d elements that possess a large perpendicular K u [5]–[7]. In general, body-centered cubic FeCo alloys exhibit cubic magnetic anisotropy. However, based on the first-principle theory, Burkert et al. [5] first proposed that body-centered tetragonal (bct) Fe1−x Cox alloys can possess a giant perpendicular K u because of the SOC induced strongly between the dx y and dx 2 −y 2 states of the 3d orbits at cobalt concentrations of 0.5 ≤ x ≤ 0.65 and tetragonal distortion of 1.20 ≤ c/a ≤ 1.25. Subsequently, most of the experimental realizations of bct FeCo alloys have been grown strained FeCo alloy films epitaxially on suitable nonmagnetic substrates with face-centered cubic structure or as building blocks of superlattices, such as Pd, Ir, Rh, and Pt [6], [7]. Notably, the magnetization easy axis was reported to remain out-of-plane until 17 monolayers (MLs) for a disordered Fe0.5 Co0.5 (A2 state) alloy film, which was grown on Rh(001) with a c/a Manuscript received March 7, 2014; revised April 12, 2014; accepted May 1, 2014. Date of current version November 18, 2014. Corresponding author: B. Lao (e-mail:
[email protected]). Color versions of one or more of the figures in this paper are available online at http://ieeexplore.ieee.org. Digital Object Identifier 10.1109/TMAG.2014.2322936
ratio of 1.24 ± 0.04 [8]. However, no systematic quantitative analysis of K u for FeCo/Rh has been performed to date. The experimentally obtained K u values of FeCo alloys for similar systems were appreciably smaller than the calculated values [7], [9], [10]. Currently, a modified calculation method was employed to evaluate the anisotropy properties of FeCo alloys, which obtained optimal conditions of cobalt concentrations and tetragonal distortion for K u similar to those of [5] and [9]; however, the calculated value of K u was in agreement with the previous experimental studies [11], [12]. In addition, with this method, the crystal order degree dependence of K u in the FeCo alloys under the optimal c/a ratio was predicted to give the maximum K u values of 450–600 µeV/atom in the perfectly ordered Fe0.5 Co0.5 (B2 state) case, which is sufficiently comparable with the K u values of bi-metallic alloy systems (FePd, FePt, and CoPt), as mentioned above [4]. This result suggests that the order-induced increase of the K u value, owing to the energy separation between the related electronic states of the B2 state, is smaller than that of the A2 state of FeCo alloys. In this paper, we fabricated B2-state Fe0.5 Co0.5 films epitaxially on Rh(001) at RT. The Rh has been suggested as a promising candidate to strain the FeCo alloys in the optimal lattice distortion (c/a = 1.20 − 1.25) and give the largest K u because of the appropriate lattice mismatch between them. The Fe0.5 Co0.5 films were deposited using the alternate monatomic layered (AML) method, in which N(6–30) MLs of Fe and Co were deposited alternately to form a B2 state, which we refer to as AML[Fe/Co] N [13]. After sample fabrication, we studied the bulk anisotropy, the c/a ratio, and the critical thickness of the films and compared our results with the previous studies. II. E XPERIMENTAL P ROCEDURE The film fabrication was performed in an e-beam evaporator chamber with a base pressure better than 1 × 10−6 Pa. Our films were grown epitaxially with a stacking structure of MgO(001) single crystalline substrate/Rh seed layer of
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Fig. 2. Magnetic hysteresis loops measured at RT for AML[Fe/Co] N films of varying thickness grown on a Rh(001) seed layer. The loops display a perpendicular magnetization easy axis up to a film thickness of 24 MLs.
Fig. 1. c/a ratio as a function of AML[Fe/Co] N . The dashed line is a guide. For a thickness smaller than 12 MLs (1.7 nm), the c/a ratio remains 1.20, as this is the optimal condition for the strongest SOC. Inset: RHEED patterns of the 50-nm Rh and the AML[Fe/Co]24 MLs.
50 nm/AML[Fe/Co]6−30 MLs (in which 1 ML is 0.14 nm)/ Rh capping layer of 10 nm. The MgO(001) substrates were prepared before deposition by in situ annealing treatment at 500 °C for 1 h to obtain a clean and flat surface. For the AML[Fe/Co] N layers, low-deposition rates of 1 Å/min for Fe (99.9% purity) and Co (99.9% purity) were chosen to avoid crystal quality degradation at RT and to prevent a mixing layer being induced at the Rh/Fe interface. The thickness of the AML[Fe/Co] N varied from 0.84 to 4.2 nm, which corresponds to the numbers of MLs: N = 6, 10, 12, 18, 24 and 30 MLs. After AML[Fe/Co] N layer depositions, a 10 nm-thick Rh capping layer was deposited as a protective layer. In situ reflection high energy electron diffraction (RHEED) was performed to confirm the epitaxial growth of the films and measure the in-plane lattice spacing of AML[Fe/Co] N using the diffraction patterns. The spacing of the neighboring streaks in the diffraction patterns is in accordance with the Bragg’s diffraction condition and is proportional to the in-plane lattice spacing of the surface layer; thus, we could calculate the lattice distortion for each sample. An X-ray diffraction (XRD) was utilized to verify the results of RHEED. The in-plane and out-of-plane hysteresis loops were measured by a vibrating samples magnetometer. The in-plane cubic anisotropy component was observed by a torque magnetometer. III. R ESULTS A ND D ISCUSSION The lattice distortion (c/a ratio) as a function of AML[Fe/Co] N thickness is shown in Fig. 1. We found that the c/a ratio reached 1.20 and this value was retained up to 12 MLs. This means that an almost perfect pseudomorphic epitaxial growth was maintained up to 12 MLs, owing to the strain effect of the Rh seed layer contributing completely to the AML[Fe/Co] N films. However, for thickness thicker than 12 MLs, the in-plane lattice spacing of AML[Fe/Co] N began relaxing toward the bulk value, which resulted in a decrease of the c/a ratio. We obtained similar results for the c/a ratio from the XRD measurement. Unfortunately, we only observed the characteristic peak of the AML[Fe/Co] N films for N = 24 and 30 MLs because the measurement sensitivity of our XRD equipment was limited to ∼3 nm. The relationship between
Fig. 3. Effective magnetic anisotropy as a function of AML[Fe/Co] N . The gray dashed line is a guide. The red open squares represent the K eff for a magnetization easy axis perpendicular to the film plane; the blue open square denotes an in-plane magnetization easy axis.
the c/a ratio and the thickness is qualitatively consistent with the results of disordered FeCo alloys grown on Rh [6]. The inset of Fig. 1 shows the RHEED diffraction patterns for the surface of the 50 nm Rh and the AML[Fe/Co]24 MLs. Streaky patterns are clearly present for the (001) surface of the Rh seed layer in [100] azimuth and of AML[Fe/Co]24MLs in [110] azimuth, caused by a 45° rotation of the FeCo bulk pseudomorphically grown on the Rh for atomic matching. This indicates that good epitaxial films with Frank-van der Merwe growth and flat surfaces were obtained for our samples [14]. Furthermore, the in-plane lattice spacing of the Rh seed layer was confirmed as the intrinsic lattice constant length, which can strain AML[Fe/Co] N to the optimal lattice distortion. The hysteresis loops of the samples show a perpendicular magnetization easy axis, which was retained up to 24 MLs. A spin reorientation transition from perpendicular to in-plane magnetization occurred between 24 and 30 MLs. The characteristic loops are shown in Fig. 2. The effective magnetic anisotropy (the energy required to change the orientation of the magnetization from easy to hard axis, K eff = K u − K d , where K d is the shape anisotropy of the film) of each film were determined by measuring the enclosed area of the hysteresis loop by applying in-plane and out-of-plane field orientations. The K eff data plotted as a function of AML[Fe/Co] N thickness are shown in Fig. 3. The K eff of AML[Fe/Co] N remained around 3 × 106 erg/cm3 for N ≤ 12 MLs and decreased between N = 18 and 30 MLs. Therefore, the reduction in the c/a ratio for N ≥ 18 MLs leads to the decrease of K eff values
LAO et al.: STRONG PERPENDICULAR UNIAXIAL MAGNETIC ANISOTROPY IN TETRAGONAL Fe0.5 Co0.5 FILMS
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Fig. 4. In-plane cubic anisotropy increase corresponding to the decrease of uniaxial magnetic anisotropy of AML[Fe/Co] N , displaying the weakening of the SOC for the c/a ratio’s departure from the optimal region when increasing the total thickness of AML[Fe/Co] N . The dashed lines are eye guides.
Fig. 5. Bulk anisotropy of our artificial B2-state Fe0.5 Co0.5 -AML[Fe/Co] N and A2-state Fe0.5 Co0.5 as a function of c/a ratio [12]. The value of K V for the B2-state Fe0.5 Co0.5 is larger than that for the A2 state at the same c/a ratio.
owing to the complete strained film portion (below 12 MLs) contributing less to the total K u . The thickness boundary of the positive and negative K eff is defined as the critical thickness. In this case, the thickness reached 24 MLs (3.4 nm), which is larger than that of the disordered Fe50 Co50 alloy [8]. We determined the in-plane cubic anisotropy (K cubic ) component using a torque magnetometer with an applied field rotated around the in-plane of the samples to confirm the isotropic property of the plane. As shown in Fig. 4, for a AML[Fe/Co] N thickness smaller than 12 MLs, K cubic is estimated to be around 103 erg/cm3, which means that almost no energy difference exists between any in-plane directions. In contrast, a higher K cubic was found for the AML[Fe/Co]18−30 MLs, of the order of 104 erg/cm3. These results demonstrate that the c/a ratio is related to the magnitude of K cubic . For c/a < 1.20, K cubic increased owing to the energy separation between in-plane 3d orbits becoming relatively large. Similar results for c/a > 1.25 have been obtained previously [5]. However, compared with the magnitude of K u , the K cubic values are extremely small, which indicates that the bct-AML[Fe/Co] N exhibits the uniaxial magnetic anisotropy property. Since the K u of the films has both a contribution of bulk anisotropy from the volume and a contribution of the interface anisotropy from the two interfaces enclosing the bulk of the magnetic layer, it can be expressed by the following relation:
results from similar analyses, which found K V = 1.87 × 107 erg/cm3 for a [FeCo]3 /Pt7 multilayer with c/a = 1.18, and K V = 1.26 × 107 erg/cm3 for a Fe0.5 Co0.5 alloy with c/a = 1.13 [9], [10]. The K i , which consists of the interface anisotropy from the bottom Rh(001)/Fe(001) and the top Co(001)/Rh(001) interfaces, was obtained from the vertical axis intercept at tFM = 0 and was found to be 0.68 ± 0.34 erg/cm2. We assumed that the interface anisotropy is insensitive to the different in-plane lattice spacing and can be defined as a constant. The K V values for the different c/a ratios of the N = 18, 24, and 30 MLs were estimated by (1). Fig. 5 shows the relationship between the estimated K V for each c/a ratio of AML[Fe/Co] N and the bulk magnetic anisotropy of the A2-state Fe0.5 Co0.5 , as determined by the first-principle calculation [12]. The behavior of K V versus the c/a ratio for AML[Fe/Co]N is in qualitative agreement with the theoretical calculation of the A2-state Fe0.5 Co0.5 , indicating that the c/a ratio is an essential parameter to achieve large uniaxial magnetic anisotropy. In addition, for the same c/a ratio, the K V of bulk AML[Fe/Co]N is larger than that of the A2 state, which demonstrates that improving the crystal order of Fe0.5 Co0.5 is an effective way to further enhance the uniaxial magnetic anisotropy under a certain lattice distortion.
K u tFM = K V tFM + K i .
(1)
Here, K V , K i , and tFM are the bulk anisotropy, interface anisotropy, and total thickness of the AML[Fe/Co] N films, respectively. We estimated the K V of our samples to distinguish the bulk contribution from the total K u of the film. Since K V is expected to depend on the degree of tetragonal distortion, we plotted K u tFM as a function of tFM by a linear fit for the N = 6, 10, and 12 MLs to obtain the K V at c/a ratio of 1.20. The bulk anisotropy is given by the slope of the fitting line and we determined that K V = (2.50 ± 0.34) × 107 erg/cm3. For comparison, we quote some previous
IV. C ONCLUSION Strong perpendicular uniaxial magnetic anisotropy was achieved in tetragonal Fe0.5 Co0.5 films of an artificially ordered B2 state. The isotropy plane property was confirmed in these films. The magnetization easy axis remained out-ofplane up to the thickness of 24 MLs. The anisotropy energy of bulk AML[Fe/Co]N reached (2.50 ± 0.34) × 107 erg/cm3 and remained constant up to 12 MLs at a c/a ratio of 1.20, which is larger than that of disordered Fe0.5 Co0.5 under the same lattice distortion. This indicates that the degree of crystal order of FeCo alloys is a decisive factor in the perpendicular uniaxial anisotropy under the optimal conditions of FeCo composition and lattice distortion.
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ACKNOWLEDGMENT This work was supported by the Storage Research Consortium (SRC). R EFERENCES [1] H. N. Bertram, H. Zhou, and R. Gustafson, “Signal to noise ratio scaling and density limit estimates in longitudinal magnetic recording,” IEEE Trans. Magn., vol. 34, no. 4, pp. 1845–1847, Jul. 1998. [2] J.-G. Zhu and Y. Wang, “Microwave assisted magnetic recording utilizing perpendicular spin torque oscillator with switchable perpendicular electrodes,” IEEE Trans. Magn., vol. 46, no. 3, pp. 751–757, Mar. 2010. [3] P. Bruno, “Tight-binding approach to the orbital magnetic moment and magnetocrystalline anisotropy of transition-metal monolayers,” Phys. Rev. B, Condens. Matter, vol. 39, no. 1, pp. 865–868, 1989. [4] D. Weller et al., “High Ku materials approach to 100 Gbits/in2 ,” IEEE Trans. Magn., vol. 36, no. 1, pp. 10–15, Jan. 2000. [5] T. Burkert, L. Nordström, O. Eriksson, and O. Heinonen, “Giant magnetic anisotropy in tetragonal FeCo alloys,” Phys. Rev. Lett., vol. 93, no. 2, pp. 027203-1–027203-4, 2003. [6] F. Yildiz, M. Przybylshi, X.-D. Ma, and J. Kirschner, “Strong perpendicular anisotropy in Fe1−x Cox alloy films epitaxially grown on mismatching Pd(001), Ir(001), and Rh(001) substrates,” Phys. Rev. B, Condens. Matter, vol. 80, no. 6, pp. 064415-1–064415-6, 2009. [7] P. Warnicke, G. Andersson, M. Björck, J. Ferré, and P. Nordblad, “Magnetic anisotropy of tetragonal FeCo/Pt(001) superlattices,” J. Phys., Condensed Matter, vol. 19, no. 22, p. 226218, 2007.
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