Effect of MgO underlayer misorientation on the texture and magnetic property of FePt–C granular film

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ScienceDirect Acta Materialia 91 (2015) 41–49 www.elsevier.com/locate/actamat

Effect of MgO underlayer misorientation on the texture and magnetic property of FePt–C granular film J. Wang,a S. Hata,b Y.K. Takahashi,a H. Sepehri-Amin,a B.S.D.Ch.S. Varaprasad,a T. Shiroyama,a ⇑ T. Schreflc and K. Honoa, a

National Institute for Materials Science, 1-2-1 Sengen, Tsukuba 305-0047, Japan Department of Electrical and Material Science, Kyushu University, Kasuga 816-8580, Japan c Center for Integrated Sensor Systems, Danube University Krems, Austria

b

Received 14 January 2015; revised 22 February 2015; accepted 3 March 2015 Available online 28 March 2015

Abstract—A transmission electron microscope (TEM) based orientation mapping technique and micromagnetic simulations were applied to study the influence of easy axis distribution (EAD) on magnetic properties of FePt–C granular films which were deposited on a single crystalline MgO (0 0 1) substrate and a (0 0 1)-textured poly-crystalline MgO underlayer. The FePt–C film on the polycrystalline MgO underlayer shows smaller perpendicular coercivity, broader switching field distribution and visible in-plane minor loop compared with that deposited on the single crystalline MgO (0 0 1) substrate. Although the grain sizes and their distributions in both films look similar in TEM, orientation mapping and texture analysis revealed that the polycrystalline MgO underlayer introduces significant misorientation in the (0 0 1)-textured FePt grains. Micromagnetic simulations successfully reproduced the large hysteresis in the in-plane magnetization by introducing the specific misorientation distribution of the FePt grains obtained from the texture analysis. The misoriented FePt grains were found to be grown from misoriented MgO grains, indicating that the improvement of the (0 0 1) texture of the MgO underlayer is essential to reduce the in-plane component of FePt based recording media. Ó 2015 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Magnetic recording media; Heat assisted magnetic recording; Misorientation; FePt–C

1. Introduction FePt–C based granular films with L10-ordered FePt nano-grains have been considered as the most promising candidate for heat-assisted magnetic recording (HAMR) media for the recording density exceeding 1 Tbit/in2. For optimal recording performance, the c-axis of L10-FePt grains must be aligned normal to the film plane with a small grain size of less than 6 nm and a size distribution of less than 10% [1]. Previously, extensive studies were carried out to establish an effective way to obtain (0 0 1)-textured FePt films on single-crystalline MgO (0 0 1) substrates [2–5]. With elevated substrate temperature, L10-FePt tends to epitaxially grow on MgO (0 0 1) with easy axis for magnetization, [0 0 1]FePt, normal to the film plane. However, for industrial viability, the strongly (0 0 1)-textured FePt films must be grown on low-cost amorphous substrates, i.e., glass or thermally oxidized Si. We reported that the FePt(Ag)–C films deposited on (0 0 1)-textured MgO underlayers show the microstructure suitable for perpendicular recording with its excellent grain isolation and small size distribution [6–11]. Such films have

⇑ Corresponding author; e-mail: [email protected]

a uniform nanogranular structure consisting of well-isolated FePt particles of 5.5 nm in diameter with a size distribution of ±2.3 nm and pitch distance of 6.7 nm [6]. Due to the assembly of well-separated nano-sized L10-ordered FePt grains with high magnetocrystalline anisotropy energy (Ku  6.6 MJ/m3), l0Hc of 3.7 T was reported in FePtAg–C media [11]. A HAMR static test using the film demonstrated the areal density of 550 Gbit/in2, which was the highest in HAMR recording at that time [10]. Although the FePt–C based granular medium grown on the (0 0 1)-textured MgO underlayer is thought to be suitable for a mass production, easy axis distribution of the FePt grains remains a problem to achieve a sufficient signal-to-noise ratio (SNR) for a higher recording density. So how to estimate and control the magnetic easy axes distribution (EAD) in the HAMR recording media has become a critical issue in the magnetic recording community. For well-isolated FePt granular thin films, the individual grains are single domain, exchange-decoupled ferromagnets and non-uniform switching during magnetization reversal. In this case its magnetization behavior can be described by the Stoner–Wohlfarth (SW) model [12]. In this model, the magnetization reversal behavior and the shape of the hysteresis loop are mainly determined by the angle between

http://dx.doi.org/10.1016/j.actamat.2015.03.007 1359-6462/Ó 2015 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

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magnetic easy axis and external magnetic field. The easy axis orientation varies from grain to grain and the performance of FePt-based HAMR media such as SNR is quite sensitive to its EAD [13]. Hence, the determination of the EAD is essential for understanding the switching processes within the recording media. With such kinds of practical need and physical interests, several methods have been developed to deduce EAD. It can be roughly determined by X-ray diffraction where the dispersion is characterized by the full-width half-maximum (FWHM) of the FePt rocking curve [14]. The most widely used method was proposed by Shtrikman and Treves [15], in which EAD is deduced from the angular dependence of the remanence parallel and perpendicular to the applied field [16–19]. The EAD determined by these methods is effective EAD and is only strictly applicable to non-interacting ferromagnetic systems. Recently, high frequency electron paramagnetic resonance (HFEPR) [20] and magnetic force microscopy (MFM) [21] were also applied to investigate EAD. But these methods are limited by their localized data acquisition and limited lateral resolution, which cannot guarantee the statistic accuracy of the data analysis. Also, all of these methods are indirect way to study EAD. In this work, we employed the transmission–electron-microscope (TEM)-based orientation mapping technique, ASTARe system, and directly mapped the EAD within a micrometer scale. Using this method, we attempted to understand the origin of EAD in the media grown on glass substrates through (0 0 1)-textured polycrystalline MgO underlayer compared to the one grown on single-crystalline MgO substrates.

2. Experimental procedures

(SQUID-VSM) with an applied magnetic field of up to ±7 T. 2.2. Texture analysis The crystallographic orientation and phase maps were acquired on a FEI Tecnai G2 F20 Super Twin TEM (FEI Corporation, Hillsboro, OR) equipped with the NanoMEGAS ASTARe (Automatic crySTAl oRientation and phase mapping hardware/software package for TEM) system [23–26]. For data acquisition, the selected area was scanned by a nano-sized beam and the nano-beam diffraction pattern of each scanned point was simultaneously recorded by a charge-coupled device (CCD) camera mounted on the observation window. The system visualizes spatial distribution of crystallographic orientation and phase by comparing the experimental and calculated diffraction-spot patterns. The acquisition and analysis parameters were: step size 2.5 nm, spot size 1 nm, camera length 9.0 cm, horizontal distortion 0.89 and vertical distortion 1.15. The orientation maps were further imported within TSL OIMe software for quantification analysis. The orientation raw data were subjected to a cleanup procedure to eliminate unindexed and incorrectly indexed points. First, the “Grain Dilation” filter was applied with a tolerance angle of 3° and a minimum grain size of 7 pixels. Then, based on the first step cleaned raw date, the “Single (Average) Orientation per Grain” treatment was assigned to all of the pixels within a grain, assuming all adjacent pixels with misorientation less than 3° belonged to the same grain. Finally, the “Pseudosymmetry” cleanup was used to remove false boundaries that are created within single grains when patterns can be indexed in multiple orientations related by simple symmetry operations [27,28].

2.1. Film processing 2.3. Micromagnetic simulations An alternating film stack of [FePt-48 vol.% C (0.25 nm)/ FePt (0.15 nm)]25 (hereafter FePt–C granular film) was deposited by co-sputtering Fe, Pt and C at 600 °C under 0.48 Pa Ar with a deposition rate of 0.02 nm/s on two types of substrates: MgO (0 0 1) single crystalline substrate (sample A) and glass substrates through buffer layers of amorphous-NiTa (100 nm)/poly-MgO (10 nm) (sample B). The 10 nm MgO underlayer was RF sputter-deposited on the amorphous NiTa layer under an Ar pressure of 5.2 Pa with a deposition rate of 0.006 nm/s at room temperature (RT) using a MgO target [11]. The thickness of the film and the atomic fraction of Fe, Pt, and C were estimated using the pre-calibrated sputtering rates. The composition of FePt for both films is nearly 1–1. The alternating layer deposition technique applied here for FePt–C film was to suppress the growth of the randomly oriented spherical grains on the (0 0 1) textured FePt granular layer [22], and the FePt–C film thickness was optimized to ensure a single layered structure. The crystallographic structure and the degree of L10 order were examined using X-ray diffraction (XRD) with ˚ ). The microstructures were Cu Ka radiation (k = 1.542 A characterized by a spherical aberration corrected scanning transmission electron microscope (Titan G2 80–200 ChemiSTEM, FEI). The room temperature magnetic property of the films was measured by superconducting quantum interference device vibrating sample magnetometer

The effect of the easy axis distribution on the magnetization reversal was studied using micromagnetic simulations on the modeled samples with average grain size of 10 ± 1.5 nm based on the experimentally measured average grain size. The models were 400  400  10 nm3. The grains were isolated by a non-magnetic grain boundary phase with a thickness of 1.5 nm. The saturation magnetization (l0Ms), magnetocrystalline anisotropy (K1) and exchange stiffness (A) of the ordered FePt nanoparticles were chosen to be 1.43 T, 5.3 MJ/m3 and 10 pJ/m, respectively [29]. The easy axis of L10-FePt grains was assumed to be parallel to Z direction with a random misorientation which was considered based on orientation mapping results measured in transmission electron microscope for the L10FePt samples deposited on single crystalline and polycrystalline MgO substrate. Tetrahedron meshes with total element number of 100 K were generated and the Landau– Lifshitz–Gilbert (LLG) equation was solved at each node by the FEMME software [30].

3. Results Fig. 1 shows the XRD profiles of the FePt–C granular film on: (a) MgO (0 0 1) single crystalline substrate and (b) glass/NiTa (100 nm)/MgO (10 nm) polycrystalline

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Fig. 1. XRD patterns of FePt–C granular film on (a) MgO single crystal and (b) polycrystalline MgO underlayer.

underlayer. In the XRD patterns, the diffraction peaks from (0 0 1)FePt and (0 0 2)FePt were clearly observed with a missing (1 1 1)FePt peak indicating FePt grains were strongly textured to the (0 0 1) plane. This texture was induced by the epitaxial growth of FePt on MgO (0 0 1) single crystal substrate or the (0 0 2)MgO-textured 10 nm-MgO underlayer grains [9]. The superlattice (0 0 1) diffraction peak indicates that the FePt orders to the L10 structure from the as-deposited state. The degree of long-range order parameter S was evaluated from the integrated intensities of the (0 0 2) fundamental peak and the (0 0 1) superlattice peak following formula: "    #1  12 2 ðL  A  DÞf 2 I 001 Ff I 001  S¼  ¼ 0:85  ; I 002 FS ðL  A  DÞS I 002 where is Ihkl the integrated intensity, F is the structure factor, L is the Lorentz polarization factor, A is the absorption factor, D is the temperature factor, and the subscripts f and s refer to the fundamental peak and super-lattice peak, respectively [31]. The values of the degree of order were 0.74 for the samples grown on the MgO (0 0 1) substrate and 1.0 for the sample grown on the poly-MgO underlayer. The apparent high degree of order in the sample grown on the poly-MgO underlayer may be an artifact due to the strong broad amorphous peak around 23° that comes from the glass substrate below the thin MgO underlayer. In order to estimate the quality of the (0 0 1) texture of the L10-FePt grains, the rocking curves of (0 0 1)FePt of the two films are measured and shown in Fig. 2. The full-width at half-maximum (FWHM) of the FePt (0 0 1) rocking curve of the sample grown on the single MgO (0 0 1) substrate is about 4.15°, while it is 8.78° for the sample grown on the poly-MgO underlayer. This indicates that the epitaxial growth of the FePt film on the MgO (0 0 1) single crystal substrate provides much improved (0 0 1) texture than the FePt film on the polycrystalline MgO underlayer, suggesting the presence of angular distributions in the polycrystalline MgO underlayer. Fig. 3 presents in-plane and cross-sectional TEM brightfield images and the corresponding size distribution (insets) of 10-nm-thick FePt–C granular film on (a) MgO single crystal and (b) polycrystalline MgO underlayer. Both of the films show well-isolated uniform granular structure with the average grain sizes of 10.8 ± 1.6 nm and 10.1 ± 1.3 nm, respectively. From the cross-sectional TEM images, one can confirm that the FePt grains have a columnar shape with a single layered structure without small particles covering on the top [32]. This single layered

Fig. 2. Rocking curve of FePt (0 0 1) peak from FePt–C film on MgO single crystal and polycrystalline MgO underlayer.

structure is necessary for the texture orientation mapping to extract the orientation mapping from the FePt grains on the MgO (0 0 1) substrate and poly-MgO underlayer. The room temperature magnetic hysteresis loops of the films grown on MgO (0 0 1) and MgO underlayer were measured both in perpendicular and in-plane directions as shown in Fig. 4. l0Hc were 4.3 and 3.7 T for the films deposited on MgO (0 0 1) and MgO underlayer, respectively. Obviously, both of the films show strong perpendicular magnetic anisotropy. However, compared with the one grown on MgO (0 0 1), the film grown on the MgO underlayer presents a slanting loop or poor squareness with smaller perpendicular coercivity, broader switching field distribution (SFD), DM/DH, and visible in-plane minor loop. Since both films hold similar morphology and grain size, the degradation of perpendicular coercivity, broader SFD and large in-plane component in the film grown on the MgO underlayer can be attributed to the [0 0 1] easy axis misorientation. In order to clarify the effect of the EAD on the magnetic properties, we carried out orientation mapping with ASTARe (NanoMEGAS, Brussels, Belgium) system. Fig. 5 shows the orientation maps of the FePt–C granular films grown on (a and b) MgO (0 0 1) single-crystal and (c and d) poly-MgO underlayer obtained from the plane-view specimens. The out-of-plane orientation map for the film grown on the single crystalline MgO (0 0 1)

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Fig. 3. In-plane & cross-sectional bright-field TEM images of FePt–C granular film on (a) MgO single crystal and (b) polycrystalline MgO underlayer. (Inset: FePt grain size distribution, unit in nm.)

Fig. 4. In-plane and out-of-plane magnetization curves of FePt–C granular film on (a) MgO single crystal and (b) polycrystalline MgO underlayer.

substrate (Fig. 5a) demonstrates almost homogeneous red index color, indicating a prefect [0 0 1] texture of FePt grains along the film normal direction. Moreover, the inplane orientation map (Fig. 5b) shows homogeneous [1 1 0] blue index color, indicating all the FePt grains hold the same direction along both film normal and in-plane directions. It confirms the good epitaxial growth of FePt nanograins on the MgO (0 0 1) single crystal substrate. On the other hand, the FePt–C film grown on the poly-MgO underlayer (Fig. 5c) shows a distinctly different orientation map. While the majority of the grains are oriented to the [0 0 1] direction along the normal direction to the film plane, variants such as [1 1 0] and even [1 0 0] (90° misorientation) poles were also detected. From the gradient and contrast of the red index color, one can detect a distribution of [0 0 1] orientation along the film normal direction. In addition, the in-plane orientation distribution is random as shown in Fig. 5d, indicating the [0 0 1] fibrous texture on the poly-MgO underlayer. The misorientation of the (0 0 1) texture of the FePt grains mapped by the ASTARe system was further quantified with TexSEM Laboratories Orientation Imaging Microscopy (TSL OIMe, EDAX, Mahwah, NJ, USA)

analysis software. Fig. 6 shows the [0 0 1] texture (easy axis) distribution according to the film normal direction of the FePt–C granular films grown on the MgO (0 0 1) single crystal and the poly-MgO underlayer. (0 0 1)-textured FePt grains in both of the samples show mainly misorientation angle of around 5°. However, unlike the sample grown on the MgO (0 0 1) single crystal, we found total 23% of FePt grains in the film grown on the MgO underlayer have 45° or even 90° misoriented easy axis. In the above, we have shown that the FePt–C granular film deposited on poly-MgO underlayer has a larger EAD compared to the one grown on the MgO (0 0 1) single crystal. Then, what is the origin of the misorientation? To clarify this question, detailed microstructure studies were implemented on the poly-MgO underlayer. Fig. 7 displays in-plane and cross-sectional TEM bright-field images and the corresponding size distribution (inset) of a 10-nm-thick polycrystalline MgO film. The diffraction contrast causes black and gray contrast in the in-plane TEM image which helps to distinguish individual MgO grains, and the average grain size was determined to be about 16 nm. The cross-sectional image also reveals a polycrystalline feature with uniform film thickness and flat surface. To clarify the effect of

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Fig. 5. Orientation maps of FePt–C granular films deposited on (a and b) MgO single crystal and (c and d) polycrystalline MgO underlayer. In these maps, (a and c) are the inverse pole figures along film normal direction (point out of the screen), while (b and d) are the inverse pole figures along inplane direction (within the screen). The orientation of each grain is given by the reference triangle in which (0 0 1, 1 1 0, and 1 0 0) are directions in the FePt crystal reference frame.

Fig. 6. [0 0 1] texture (easy axis) distribution related to film normal direction of FePt–C granular film on (a) MgO single crystal and (b) polycrystalline MgO underlayer.

the poly-MgO underlayer on the misorientation of the subsequently deposited FePt grains, orientation mapping was also obtained from the plan-view specimen of the polyMgO underlayer film. Fig. 8 presents orientation maps of the polycrystalline MgO underlayer film. Fig. 8a is the inverse pole figure along the film normal direction, while Fig. 8b is in-plane orientation map. As can be seen from the gradient and contrast of the index color in Fig. 8a, one can clearly detect a wide distribution of MgO [0 0 1] zone axis along the film normal direction. Furthermore, one can see a random crystal orientation along the in-plane direction in Fig. 8b. The distinct contrast of the index color

and many variants indicate a weak out-of-plane [0 0 1] textured ploy-MgO film. The MgO [0 0 1] distribution (inset in Fig. 8a) along the film normal direction was plotted within 0–45° considering its symmetry fcc structure. Compared with the single crystalline MgO substrate, the poly-MgO underlayer has misorientations around 7.5° (95%) and 40° (5%). This confirms that the misorientation in the FePt–C granular film is mainly originated from the misorientations in the polycrystalline MgO underlayer. Fig. 9 shows orientation maps of cross-sectional view of FePt–C granular films grown on the MgO (0 0 1) single crystal and the polycrytalline MgO underlayer along film

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boundaries that was reported by Wicht et al. [29]. These misorientated FePt grains were further confirmed with cross-sectional HRTEM observation (Fig. 10). From the TEM images, we found that the FePt nanoparticles with misorientation angle smaller than 5° mainly come from the surface roughness of single crystalline MgO substrate or poly-MgO underlayer. However, the FePt nanoparticles with misorientation angle larger than 35° can be attributed to their epitaxial growth on misorientated MgO grains. It is also in good agreement with the cross-sectional plane orientation mapping results.

4. Discussion

Fig. 7. In-plane & cross-sectional bright-field TEM images and size distribution of polycrystalline MgO film. (Inset: MgO grain size distribution, unit in nm).

normal direction (Y direction in the TSL OIMe reference frame). It is clear that the FePt grains grows epitaxially on the single crystal MgO substrate with a perfect [0 0 1] texture along the film normal direction, However, the FePt film grown on the poly-MgO underlayer displays visible [0 0 1] misorientations along the film normal direction. With clear color contrast, one can easily identify the MgO grain boundaries, which are quite difficult to distinguish using bright field TEM images. It seems that the misorientated FePt nanoparticles mainly originate from misorientated MgO grains rather than on the MgO grain

In this work, we have investigated the microstructure and magnetic properties of FePt–C granular films grown on a MgO (0 0 1) single crystalline substrate and on a (0 0 1)-textured MgO underlayer that was predeposited on a glass substrate through an amorphous NiTa buffer layer. The results have shown that the presence of misoriented grains in the FePt–C granular films grown on the polyMgO underlayer is one of the major reasons for the inplane coercivity and degradation of SFD. Wicht et al. discussed the reason of such texture dispersion in FePt–C granular media by observing the in-plane and cross-sectional TEM images. They concluded the misorientation in FePt nanograins as a consequence of the nucleation of the FePt grain growth at the step edges of the MgO seed layer [29]. Dong et al. pointed out that due to the differences in the surface energies between FePt (2.9 J/m2) and MgO (1.1 J/m2), the contact angle between FePt and MgO becomes too large to allow the epitaxial growth on MgO underlayer, resulting in broad EAD [33]. However, our experimental results on the FePt–C film grown on a MgO (0 0 1) single crystalline substrate have convincingly shown that [0 0 1]-oriented L10-FePt grains grow with excellent epitaxy with MgO (0 0 1). Instead, the large contact angle between FePt and MgO is rather favorable for the formation of well-separated FePt grains. This merit can be clearly understood by comparing the FePt–C granular films grown on the MgO underlayer and the (Mg0.2Ti0.8)O underlayer [34]. In the latter, well-isolated

Fig. 8. Orientation maps of polycrystalline MgO underlayer film. (a) Is the inverse pole figure along film normal direction (point out of the screen), while (b) is the inverse pole figure along in-plane direction (within the screen). The orientation of each grain is given by the reference triangle in which (0 0 1, 1 1 1, and 1 0 1) are directions in the MgO crystal reference frame. (Inset: MgO grain [0 0 1] texture distribution related to film normal direction.)

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Fig. 9. Orientation maps of FePt–C granular film on (a) MgO single crystal and (b) polycrystalline MgO underlayer from cross-sectional specimens. The orientation of each grain is given by the reference triangle in which (0 0 1, 1 1 0 and 1 0 0) and (0 0 1, 1 0 1 and 1 1 1) are directions in the FePt and MgO crystal reference frame, respectively. X, Y and Z are directions in the TSL OIMe reference frame. (The black line in figures indicates the FePt grain boundary.)

Fig. 10. Cross-sectional bright field TEM image of FePt–C granular film on polycrystalline MgO underlayer with different misorientated FePt grains: (a) 45° and (b) 90°.

granular structure did not develop because of the small contact angle between FePt and (Mg0.2Ti0.8)O. To sum up, MgO underlayer is excellent choice for the growth of [0 0 1]-oriented L10-FePt granular film with its moderate mismatch to allow single variant epitaxial growth of the L10-FePt (0 0 1). A drawback of the poly-MgO underlayer is its poor texture when it was grown on amorphous substrate. The texture analysis has clearly shown that the high quality misorientation-free underlayer is necessary for achieving the small in-plane component as seen in the film grown on the single crystalline MgO (0 0 1) substrate. Although the mechanism of the large misorientation in the polycrystalline MgO underlayer is not clear, the growth of strongly (0 0 1)-textured MgO underlayer is the key to reduce the inplane component. Kodzuka et al. reported the orientation of polycrystalline MgO varies depending on the chemical composition of the amorphous CoFeB underlayer [35]. In this sense, intensive research on the way to improve the (0 0 1) texture of MgO underlayer appears to be essential for improving the (0 0 1) texture of the L10-FePt grains and SNR of HAMR media. More recently, Ho et al. reported that more than 15% of FePt grains grown on MgO underlayer are multi-variant, which induces the inplane component in the FePt–C based HAMR media [36]. In our study, we did not confirm the presence of multi-variant particles, but if that is the case, the grain size of the MgO underlayer must also be increased to reduce the chance of the FePt-particles crossing MgO grain boundaries.

In order to understand the influence of the grain misorientation on the magnetic properties (Fig. 4), the magnetization curves of L10-FePt granular films deposited on a single crystalline MgO (0 0 1) substrate and a polycrystalline MgO underlayer were simulated using finite element micromagnetic simulations. Fig. 11(a) shows the geometry of the modeled L10-FePt film with average grain size of 10 ± 1.5 nm. The L10-FePt grains were assumed to be exchange decoupled and only magnetostatic coupling was considered in these simulations. Easy axis distribution was taken into account in the model based on the experimental result in Fig. 6. Fig. 11(b and c) shows simulated out-of-plane and in-plane magnetization curves of the samples deposited on a MgO (0 0 1) single crystal and a polycrystalline MgO underlayer. Simulated in-plane coercivity of the modeled sample with the misorientations smaller than ±5° is 0.45 T, which is only 6% of the simulated out-of-plane coercivity (7.0 T). A large in-plane coercivity of 1.35 T appeared in the modeled sample on the polycrystalline MgO underlayer, which is 19.5% of the out-of-plane coercivity (6.92 T). Note that a small kink is observed at around 4 T in the simulated out-of-plane magnetization curve for the polycrystalline MgO underlayer. The origin of the kink is 45° misaligned grains, which cause a lower switching field following the Stoner–Wohlfarth behavior. However, the experimentally observed demagnetization curve of the FePt–C film grown on the poly-MgO underlayer does not show such a kink (Fig. 4b). One possible reason for this discrepancy is the switching field distribution of the L10-FePt grains that have different values of anisotropy

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Fig. 11. (a) Modeled geometry of L10-FePt media with average grain size of 10 ± 1.5 nm and thickness of 10 nm. Simulated magnetization curves of FePt–C granular films on (b) MgO single crystal and (c) polycrystalline MgO underlayer. Distribution of easy axis direction was introduced based on the experimental data (Fig. 6).

Fig. 12. Simulated magnetization curves of FePt–C granular films deposited on (a) MgO single crystal and (b) polycrystalline MgO underlayer with varied effective anisotropy field of the FePt grains between 10% and 40%.

field. Hence, we varied the distribution of the effective anisotropy field (l0Ha) of the L10-FePt grains from 10% to 40% by changing the effective magnetocrystalline anisotropy energy of each grain. Note that if the anisotropy field of the L10-FePt grains is 7.4 T, 10% distribution of anisotropy field causes random distribution of l0Ha among L10-FePt grains between 7.4 and 6.7 T. Fig. 12(a and b) shows the effect of the distribution of anisotropy field to the out-of-plane magnetization curves of the samples deposited on MgO (0 0 1) single crystal and MgO polycrystalline underlayer, respectively. The kink in the second quadrant of demagnetization curve disappears in the model with 30% distribution of anisotropy field in good agreement with the experimentally observed magnetization curves, Fig. 4b.

5. Conclusions The influence of MgO underlayers on the misorientation of (0 0 1) texture in FePt–C granular films and corresponding magnetic properties was studied. Compared with the FePt–C film grown on a MgO (0 0 1) single crystal, the FePt grains in the FePt–C film deposited on polycrystalline

MgO underlayer showed a smaller perpendicular coercivity, a large switching field distribution and broader in-plane magnetization loop. Texture analysis using the TEM-based orientation mapping technique indicated that the polycrystalline MgO underlayer introduces significant (0 0 1) texture misorientations in FePt grains (77% ±5°, 23% ±45° and ±90o). Micromagnetic simulations reproduced the large hysteresis in the in-plane magnetization by introducing the misorientation of the FePt grains observed from the texture analysis. In addition, the existence of at least 30% distribution of anisotropy field in L10-FePt grains reproduced the experimentally observed magnetization curve. Cross-sectional HRTEM observation of the misorientated FePt grains confirmed that they were originated from their epitaxial growth on misorientated MgO grains. These indicate that the improvement of the (0 0 1) texture of the MgO underlayer is essential to improve the in-plane component of FePt–C based HAMR media.

Acknowledgments This work was supported in part by the Advanced Storage Technology Consortium (ASTC) of the International Disk Drive

J. Wang et al. / Acta Materialia 91 (2015) 41–49

Equipment and Materials Association (IDEMA), USA, and the Grant-in-Aid (B) (Grant No. 26289232). The MEXT/JSPS Kakenhi (22102002, 25286027). ASTARTM analysis was performed at the Ultramicroscopy Research Center, Kyushu University.

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