Control of Microstructure and Magnetic Properties of ... - IEEE Xplore

668

IEEE TRANSACTIONS ON MAGNETICS, VOL. 49, NO. 2, FEBRUARY 2013

Control of Microstructure and Magnetic Properties of FePt Films With TiN Intermediate Layer K. F. Dong , H. H. Li , J. F. Hu , Y. G. Peng , G. Ju , G. M. Chow , and J. S. Chen Department of Materials Science and Engineering, National University of Singapore, Singapore 117576, Singapore Data Storage Institute, A*STAR, Singapore 117608, Singapore Seagate Technology, Fremont, CA 94538 USA The effects of a TiN intermediate layer on the microstructure and magnetic properties of the FePt films were investigated. It was found that the TiN layer could effectively block the diffusion of Cr into the FePt film. The good epitaxial relationships among these layers were revealed from the transmission electron microscopy (TEM) results. With introducing TiN intermediate layer the chemical ordering and magnetic properties of FePt films significantly improved. The FePt film with 5 nm TiN exhibited a high perpendicular coercivity of 13.7 kOe and a low in-plane coercivity of 0.24 kOe, resulting from the combined contribution of TiN (200) orientation, TiN layer roughness and the effective block of Cr diffusion. Moreover, with doping C into the FePt-SiN films, the out-of-plane coercivity increased due to the decrease of the exchange coupling, the grain size of FePt films decreased, and well-separated FePt grains and uniform size were formed. By optimizing the sputtering process, the [FePt (4 nm)-SiN 40 vol %]- 20 vol % C (001) film with coercivity higher than 21.5 kOe, a single layer structure, and small FePt grain size of 5.6 nm in average diameter was obtained, which are suitable for ultrahigh density perpendicular recording. Index Terms—FePt films, microstructure, perpendicular media, TiN intermediate layer.

I. INTRODUCTION

T

HE ordered FePt alloy with high magnetocrystalline anisotropy ( erg/cc) has been intensively investigated due to its potential application for ultra-high-density perpendicular magnetic recording [1], [2]. For practical applications, it is imperative to fabricate the FePt films with good (001) texture, a large magnetocrystalline anisotropy and small -FePt films with grain size with a narrow size distribution. (001) texture and high magnetocrystalline anisotropy have been fabricated by epitaxial growth on single-crystalline MgO (001) substrates [3], [4] or on polycrystalline underlayers such as Ag (100), MgO (200), and CrRu (200) with silicon wafers or glass -FePt grown on an Ag underlayer, substrates [5]–[7]. For the rough surface caused by the Ag underlayer with low melting point was not favorable for the low flying height of recording -FePt (001) head and thus high recording density [8], [9]. films with high magnetic anisotropy and small grain size have further been fabricated on MgO underlayer/intermediate layer by C doping, which was attributed to the island growth on MgO with small surface energy [10]–[12]. However, the RF sputtering of the insulating MgO underlayer/intermediate layer is not preferred for industrial applications due to its low deposition rate and particle contamination on the media surface. Moreover, small surface energy of MgO (1.1 J/m ) comparing with FePt (2.9 J/m ) prefers to resulting in a large contacting angle between FePt grain and MgO which is not favorable for epitaxial growth and will cause large -axis distribution and/or promote the formation of in-pane variants. Relatively narrow opening-up of the in-plane hysteresis loop were reported on MgO single Manuscript received July 11, 2012; revised September 11, 2012; accepted September 27, 2012. Date of current version January 22, 2013. Corresponding author: J. S. Chen (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.2012.2223204

crystal substrate which may be due to the good crystallinity of MgO single crystal and the large FePt grain size where the non-wetting effect can be neglected [3], [4]. With FePt grain size decreasing to below 5 nm required for high density magnetic recording, the ratio of the contacting region between one grain and substrate (underlayer) to the lateral grain size due to the wetting effect become smaller. Therefore, any small deviation of texture of underlayer and roughness changes will cause deviation of the crystal orientation of FePt overlayer from film normal (001) orientation, which became predominant for polycrystalline MgO underlayer/intermediate layers. The large opening-up of in-plane hysteresis loop—large in-plane coercivity were indeed observed on MgO intermediate layer/underlayer. The opening-up of in-plane hysteresis loop may broaden the switching field distribution and thus reduce signal-to-noise J/m ) inratio. Although FePt films on Pt (surface energy termediate layer showed the narrow opening-up of in-plane hysteresis, the inter diffusion between Pt and FePt caused a thick dead magnetic layer [13]–[15]. Based on previous investigations, one can find that an intermediate layer or underlayer with smaller surface energy is required for the granular FePt films with small grain size and an intermediate layer or underlayer with larger surface energy is required for good (001) texture and thus less opening-up of in-plane hysteresis loop. Therefore it is desirable to search for conductive underlayers/intermediate layers with appropriate surface energy for the growth -FePt (001) films with good (001) texture. TiN is conof ductive and a good diffusion barrier used in semiconductor. It also has larger surface energy (1.28 J/m than MgO. The lat//TiN (001) is about tice mismatch of FePt (001) 9.5%. Either Ti or N deficiency will reduce the lattice constant of TiN and thus lattice mismatch. It was recently reported that -FePt (001) thin films were grown on a TiN underlayer, which was epitaxially grown on a Si (100) wafer or on a RuAl underlayer. However, the FePt films exhibited poor perpendicular anisotropy and large in-plane coercivity [16]–[19].

0018-9464/$31.00 © 2013 IEEE

DONG et al.: CONTROL OF MICROSTRUCTURE AND MAGNETIC PROPERTIES OF FEPT FILMS WITH TIN INTERMEDIATE LAYER

669

In this paper, the investigation on the microstructure and magnetic properties of FePt (001) films grown on CrRu underlayer with TiN intermediate layer were carried out. Moreover, FePt-SiN -C films were fabricated and the microstructure and the magnetic properties of those films were systematically investigated. II. EXPERIMENTAL DETAIL FePt (10 nm/ TiN/ CrRu (30 nm)/ glass films with various TiN thicknesses (0, 2, 3, 4, 5 and 7 nm), and [FePt (4, 6 and 10 nm)-SiN 40 vol.%]-C/ TiN 5 nm/ CrRu (30 nm)/ glass films with different C volume concentration were deposited using an AJA sputtering system with a base pressure better than 2 10 Torr. The CrRu, TiN, FePt and C layers were all fabricated by dc-sputtering of CrRu, TiN, FePt and C targets, while the Si N layer was fabricated by RF-sputtering of the Si N target. The Ar working pressure was 2 mTorr for CrRu and 10 mTorr for TiN and FePt-SiN -C. The crystallographic texture was examined with X-ray diffraction (XRD) using Cu K radiation. The microstructure of the films was characterized by transmission electron microscopy (TEM). The morphologies of the samples were examined by atomic force microscopy (AFM) and scanning electron microscopy (SEM). The magnetic properties were measured by the vibrating sample magnetometry (VSM) and the superconduction quantum inference device (SQUID) at a maximum applied field of 60 kOe at room temperature. The depth profiling of the films was carried out by x-ray photoelectron spectroscopy (XPS) with 1 kV etching. III. RESULTS AND DISCUSSIONS A. Microstructure and Magnetic Properties of FePt Films With TiN Intermediate Layer Fig. 1(a) shows the XRD 2 -scans of the samples with various TiN thicknesses deposited at 500 C. All the FePt films exhibited (001) preferred orientation. Without the TiN intermediate layer, the FePt (002) peak shifted to lower angles, implying that the fcc phase in the film increased. It has been reported that the CrRu underlayer caused a very thick initial growth layer in FePt thin films due to the diffusion of Cr into the magnetic layer that resulted in the degradation of the (001) texture and the magnetic properties of -FePt [10]–[12]. With increasing TiN thickness from 2 to 5 nm, the -FePt (001) peak intensity increased and the FePt (002) peak shifted to higher angle, indicating that the fcc phase in the film decreased and fct phase increased, which should be attributed to the suppression of Cr diffusion. This was further proved by the results of XPS in-depth profiling (Fig. 2). The XPS in-depth profiling showed that significant Cr diffusion occurred for the sample without TiN layer and was effectively suppressed for the sample with a 5 nm TiN layer. With further increasing TiN thickness to 7 nm, the -FePt (001) peak intensity decreased and the FePt (002) peak shifted slightly to lower angles. The chemical ordering of -FePt can be qualitatively represented by the square root of the ratio of the integrated peak intensities of the (001) superlattice peak to the (002) fundamental peak, [20], and of the FePt films is summarized in Fig. 1(b). increased with increasing TiN layer thickness, reached its maximum at 5 nm

Fig. 1. (a) XRD 2 spectra, (b) the full width at half maximum (FWHM) of the rocking curves of the CrRu (200), TiN (200) and FePt (001) peaks, and the ratio of the integrated peak intensities of FePt (001) to (002) of the FePt films deposited on different TiN thicknesse and (c) the rocking curve of CrRu (200), TiN (200) and FePt (001) peaks at 5 nm TiN layer.

Fig. 2. XPS depth profiles of the FePt films with (a) 2 nm and (b) 5 nm TiN.

and then slightly decreased with further increasing TiN thickness to 7 nm. The full width at half maximum (FWHM) of the rocking curves of CrRu (200), TiN (200) and FePt (001) peaks at different TiN thicknesses, as shown in Fig. 1(b), were measured. In order to avoid the misleading by the inaccuracy caused by the small thickness, only of TiN (200) with 4, 5 and 7 nm thickness were measured. The typical rocking curves of CrRu (200), TiN (200) and FePt (001) peaks with 5 nm TiN layer are shown in Fig. 1(c). The of CrRu (200) remained around 4.5 for all samples. The addition of the 2 nm

670


Fig. 3. Average roughness of TiN/CrRu films with different TiN thicknesses.

TiN intermediate layer led to a dramatic decrease of the of the FePt (001) peak. Further increase of the TiN thickness up to 5 nm, the of FePt (001) showed slight decrease to 6.5 . With comparison of the of CrRu (200), TiN (200) and FePt (001) peaks, it was found that the of TiN (200) was larger than that of CrRu (200) and comparable with that of FePt (001) peak. This suggested that the diffusion of CrRu into FePt layer would deteriorate the FePt (001) texture and the FePt (001) texture was governed by TiN (200) texture when the diffusion effect was small. However, a small discrepancy between the TiN (200) and FePt (001) texture was observed when the TiN layer was 7 nm, where although of TiN (200) decreased, of FePt (001) increased. This suggested that the FePt (001) texture might be affected by other factors, e.g., the surface morphology of TiN layer. In order to evaluate the effect of TiN surface morphology, TiN films with different thicknesses were fabricated on CrRu underlayer with same fabrication conditions as the samples FePt/TiN/CrRu. The average roughness as the function of the TiN thickness is shown in Fig. 3. It was clearly found that the average roughness ( linearly increased with the thickness of TiN layer. Therefore, the change in the FePt-fct (001) preferred orientation and the degree of the ordering can be explained as the combined effects of the degree of TiN (200) orientation, the roughness of TiN intermediate layer and Cr diffusion, i.e., the rougher surface and Cr diffusion would deteriorate the FePt (001) orientation. For the TiN layer with 7 nm thickness, the slight deterioration of FePt (001) texture was caused by the increase of the TiN surface roughness although the Cr diffusion effect can be ignored. The out-of-plane and in-plane hysteresis loops of FePt films with different TiN layer thicknesses are shown in Fig. 4(a)–(c). It can be seen that the sample without TiN intermediate layer was magnetically soft with in-plane magnetic anisotropy which was due to the Cr diffusion [Fig. 4(a)]. With the addition of a 2 nm TiN intermediate layer [Fig. 4(b)], the out-of-plane and in-plane coercivity were 2.1 and 0.4 kOe, respectively, indicating that the easy axis of magnetization changed to the perpendicular direction. It was also observed that the FePt film with 5 nm TiN exhibited high perpendicular anisotropy [Fig. 4(c)], with a high out-of-plane coercivity of 13.7 kOe and a low in-plane coercivity of 0.24 kOe. These VSM results were consistent with the XRD results. The perpendicular anisotropy can be calculated from ,

Fig. 4. Out-of-plane and in-plane M-H loops of FePt films with different TiN thickness (a) 0 nm, (b) 2 nm and (c) 5 nm; (d) the out-of-plane and in-plane coercivities and the perpendicular anisotropies of FePt films with different TiN thicknesses.

where is the magnetic anisotropy field (estimated by extrapolating the hard axis loop) [21], [22]. The out-of-plane and in-plane coercivities and the with different TiN thicknesses are summarized in Fig. 4(d). The in-plane coercivity remained low values (less than 0.5 kOe) for all the samples. first increased with increasing TiN thickness and reached the maximum of erg/cc at 5 nm TiN. Further increase of TiN thickness caused a slight decrease of the . By comparison of chemical ordering with different TiN thickness, the change of of the FePt films with TiN layer thickness was consistent with the chemical ordering of the FePt layer. The change of out-of-plane coercivity as a function of the TiN thickness had the same trend with the . It is known that the coercivity was not only related to magnetic anisotropy but also the magnetization reversal mechanism [23]. The initial magnetization curves were measured to elucidate the magnetization reversal mechanism, as shown in Fig. 5. It is known that when the magnetization reversal was dominated by pinning, the domain wall would not move until the external field was larger than the pinning field. The initial magnetization curve was convex downward. The same tendency would appear if there was an S–W model rotation, in which single domain grains only reversed its magnetization when applied field exceeded the anisotropy energy. When the magnetization reversal was dominated by nucleation, the initial magnetization curve was convex upward [24], [25]. Fig. 5 clearly showed that the initial magnetization curves of FePt films were dependent on the TiN thickness. The magnetization reversals of FePt films with TiN thickness of 0 and 2 nm were dominated by the nucleation mechanism. For FePt films with TiN thickness above 3 nm, the magnetization reversal was dominated by the pinning or S–W rotational mechanism. Therefore, the increase of the magnetic anisotropy and the change of magnetization reversal mechanism from nucleation mode to pinning mode/S–W rotation led to the increase in coercivity when TiN thickness increased from 0 to


671

Fig. 5. Initial magnetization curves with different TiN thicknesses.

Fig. 7. (a) Planar view, (b) low magnification cross-sectional and (c) the high resolution TEM images of the FePt film with 5 nm TiN layer. Inset to (c) shows the corresponding selected area electron diffraction (SAED) patterns. Fig. 6. SEM images of the FePt films with different TiN thicknesses (a) 2 nm and (b) 5 nm. The inset is the corresponding grain size distribution of FePt film with 5 nm TiN.

5 nm. For the sample with 7 nm TiN layer the decrease in coercivity could be attributed to the decrease in magnetic anisotropy caused by the decrease in the chemical ordering. The SEM images of the FePt films with different TiN thicknesses are shown in Fig. 6. The evolution of film growth mechanism and magnetization reversal mechanism with different TiN thickness could be understood also by the microstructure of the films. When TiN layer was 2 nm [Fig. 6(a)], the grain boundary was not distinct, and the particles that consisted of several grains were predominant. This might be due to the deterioration of the island growth of FePt grains caused by the Cr diffusion into the FePt layer. On the other hand, almost continuous FePt film on 2 nm TiN was easy for domain wall motion. Therefore, the magnetization reversal of FePt films with 2 nm TiN was dominated by the nucleation mechanism. With increasing TiN thickness to 5 nm [Fig. 6(b)], the grain boundary became more distinct, and island growth of FePt grains was formed due to the effective blocking of Cr diffusion (confirmed by TEM images in Fig. 7 later). The inset of Fig. 6(b) is the corresponding grain size distribution of FePt film with 5 nm TiN. The average grain size was nm, but many grains larger than 100 nm were observed. Considering the critical radius for a plate-like single domain FePt grain with thickness , the energy to create a domain wall to bisect the grain, namely, , should equal to the magnetostatic energy reduction by changing from single domain state to two equal domain state, . Therefore, critical single domain size was obtained as (1)

where is exchange stiffness constant, is magnetic anisotropy constant and is saturation magnetization. For FePt films on 5 nm TiN, was estimated by this method to be nm, which means a 80 nm in diameter, larger than average grain size but smaller than many big grains appeared in the FePt film with 5 nm TiN. Note that the demagnetizing factor for the grains in this experiment should be smaller than 1 in ideal infinite plate, the magnetostatic energy in this experiment is smaller than calculation and hence should be a little larger than 80 nm. Because the grain size in the FePt film with 5 nm TiN distributed in a wide range around calculated , single domain grains and multi-domain grains coexisted with each other. Moreover, distinct grain boundaries in FePt film on 5 nm TiN played a role as obstacle to domain wall motion, thus the magnetization reversal mechanism will include S–W rotation and pining model domain wall motion. A planar view TEM image and a low magnification cross-sectional TEM image of the FePt film with 5 nm TiN intermediate layer are shown in Fig. 7(a) and (b). Well-segregated grains were observed in the planar view TEM image [Fig. 7(a)]. The low magnification cross-sectional TEM image showed that FePt grains were grown on the top of TiN polycrystalline intermediate layer with island growth mode [Fig. 7(b)], which might be due to the difference in surface energy between FePt (2.9 J m ) layer and TiN (1.28 J m ) layer. It is known that island growth (Vomer-Weber growth) was preferred when the sum of the surface energy of film materials and the interface energy was larger than the surface energy of the substrate [26]. Therefore, the TiN intermediate layer promoted the three-dimensional island mode epitaxial growth of FePt film. The high resolution TEM (HR-TEM) images as well as the selected area electron diffraction (SAED) patterns of the FePt film with 5 nm TiN intermediate layer are illustrated in Fig. 7(c). It showed the matching of

672


Fig. 9. M-H loops of [FePt-SiN 40 vol.%]-20 vol.%C films with different FePt thickness (a) 4 nm, (b) 6 nm and (c) 10 nm, the inset is the corresponding rocking curves of FePt (001) peak. Fig. 8. (a) XRD 2 spectra and (b) the ratio of the integrated peak intensities of FePt (001) to (002) peaks of [FePt (4, 6 and 10 nm)-SiN 40 vol.%]-C films with different C volume concentration (the lines are drawn for visual guide only).

the atomic planes across the TiN and FePt interface. The interface was sharp and clear. As can be seen from the SAED patterns (d) and (e), (001) and (111) axis of FePt aligned very well along TiN (002) and (111) axis, respectively, confirming the epitaxial relationship of FePt (001) //TiN (001) . Therefore, TiN can be used to replace MgO to induce ordering and (001) texture of FePt films. The lattice constant of FePt was measured to be 3.67 Å, slightly smaller than the reported value. This might be due to the tensile stress resulted from the large lattice mismatch between the FePt and TiN layers. B. Microstructure and Magnetic Properties of FePt-SiN -C films With TiN Intermediate Layer From the results above, FePt films with enhanced orientation and narrower open-up of in-plane hysteresis loop by using the TiN intermediate layer were formed. It is expected that well isolated FePt films with small grain size, narrow grain size distribution and open-up of in-plane hysteresis loop can be obtained by using this new CrRu/TiN underlayer. In this section, the systematic investigation on the microstructure and magnetic properties of FePt-SiN -C (001) textured composite films grown on the CrRu underlayer with TiN intermediate layers was carried out. Fig. 8(a) shows the dependence of crystallographic texture of FePt-SiN -C films on doping concentration of C and FePt thickness. Only FePt (001) and (002) peaks were observed, indicating that all the FePt-SiN -C films exhibited good (001) texture. The ratios of the integrated intensity of FePt (001) to FePt (002) peaks were measured to evaluate the C volume concentration and FePt thickness on the chemical ordering of FePt films, as shown in Fig. 8(b). It was found that when FePt thickness was fixed at 4 nm, the ratio of increased with increasing the C concentration. However, when FePt thickness were fixed at 6 and 10 nm, the ratio of increased with 10 vol.% C doping. With further increase of C concentration to 20 vol. % caused the decrease of . The mechanism of the improvement of chemical ordering by 10 vol.% C doping is still under investigation. Moreover, the decrease of the chemical ordering by 20 vol.% C doping might be due to the change of the microstructure of FePt layer similar to the results of FePt-C films grown on MgO interlayer [15]. The out-of-plane and in-plane M-H loops of [FePt-SiN 40 vol.%]- 20 vol.%C films with different FePt thickness (a) 4 nm, (b) 6 nm and (c) 10 nm are shown in Fig. 9. It is observed that

Fig. 10. (a) Uut-of-plane coercivity and (b) the loop slope parameter of [FePt (4, 6 and 10 nm)-SiN 40 vol.%]-C films with different C volume concentration (the lines are drawn for visual guide only).

all samples exhibited obvious perpendicular anisotropy, with a high out-of-plane coercivity and low in-plane coercivity. These VSM results are consistent with the XRD results. The squareness ratio of all loops was close to 1. The inset shows the corresponding rocking curves of the FePt (001) peak. For the three samples, the FWHM of the FePt (001) peak was 5.1 , 7.0 and 6.9 , respectively. As compared to the FWHM of the FePt-SiN films without C doping (not shown here), it is observed that the doping of C did not deteriorate the FePt (001) texture. The out-ofplane coercivity and the slope at coercivity of [FePt (4, 6 and 10 nm)-SiN 40 vol.%]-C films as a function of C concentrations are illustrated in Fig. 10(a) and (b). As seen in Fig. 10(a), the out-of-plane coercivity increased with increasing C concentration in the FePt-SiN films. In addition, the slope of the hysteresis loop at coercivity decreased [Fig. 10(b)] with increasing C doping. The decrease in the slope, together with the well-isolated grains shown in the TEM results, indicated that the exchange coupling decreased with C doping. Therefore, the increase in the coercivity with C doping can be largely attributed to the decrease in the lateral exchange coupling of the FePt grains. The planar view of [FePt (4, 6 and 10 nm)-SiN 40 vol.%]-C films with different C volume concentration, are shown in Fig. 11. The inset is the corresponding histogram of grain size distribution. As seen from planar view TEM images, by introducing C into the FePt-SiN films, the grain boundaries became more distinct, and the grain shape changed from maze-like to circular. Moreover, grain size decreased and grain size distribution became more uniform. The grain size distribution matched very well to the simulated lognormal distribution. Compared with undoped samples, grain aggregation decreased in the samples with C doping. Therefore, upon C doping, the grain isolation was improved, which was consistent with the results of the slope of hysteresis. The high resolution cross-sectional TEM images of [FePt (4, 6, 10 nm)-SiN 40 vol.%]-C films with 20 vol.% C doping was illustrated in Fig. 12. It revealed


673

Fig. 11. Planar view TEM images of [FePt (4, 6 and 10 nm)-SiN 40 vol.%]-C films with different C volume concentration; the inset is the corresponding statistical grain size distribution.

Fig. 13. Grain sizes of [FePt-SiNx 40 vol.%]- 20 vol.% C films with different FePt thickness (the lines are drawn for visual guide only).

Fig. 12. High resolution cross-sectional TEM images of [FePt-SiN vol.%]-20 vol.% C films with different FePt thickness.

40

that for all the three samples, the FePt grains in the layer 1 with (001) orientation were epitaxially grown on the (200) textured

TiN intermediate layer and the atomic planes across the TiN and FePt interface matched well with each other. Moreover, with increasing FePt thickness to 6 nm and 10 nm, despite the FePt grains with small grain size and clear grain boundary were obtained, the FePt grains began to form double or multiple layers due to excess C diffusing to the surface and thus FePt renucleation occurred. The FePt particle “1” nucleated on FePt particles without C coverage and the fcc FePt particle “2” nucleated on C coverage were also observed in layer 2 for FePt thickness of 6 and 10 nm, which would cause

674


the chemical ordering decrease [15]. The average grain size of this series of films is summarized in Fig. 13 together with their standard deviation. It can be seen that the grain size showed obviously linear decrease with decreasing the FePt thickness and increasing the C concentration. Especially for the FePt 4 nm films co-doped with 40 vol.% SiN and 20 vol.% C, well isolated grains with size of nm were obtained. IV. SUMMARY The effects of TiN intermediate layer on the microstructure and magnetic properties of FePt films were investigated. The experimental results showed that the TiN thickness had significant effects on the fct-FePt (001) preferred orientation and the better (001) texture was obtained with 5 nm TiN. The FePt film with 5 nm TiN exhibited a high perpendicular coercivity of 13.7 kOe and a low in-plane coercivity of 0.24 kOe, resulting from the combined contribution of TiN (200) orientation, TiN layer roughness and the effective block of Cr diffusion. With doping C into the FePt-SiN films, the grain isolation was improved. Moreover, the out-of-plane coercivity increased and the grain size decreased with increasing C doping concentration. The microstructure of FePt-SiN -C films changed from one-layer structure to two-layer structure when FePt thickness increased from 4 to 10 nm. ACKNOWLEDGMENT This work was supported in part by the Agency of Science, Technology and Research (A*STAR), Singapore, SERC grant-092-156-0118, Ministry of Education, Singapore, Tier 1 funding-T11-1001-P04 and Seagate Technology. REFERENCES [1] V. B. M. Lairson and B. M. Clemens, “Enhanced magneto-optic Kerr rotation in epitaxial PtFe(001) and PtCo(001) thin films,” Appl. Phys. Lett., vol. 63, pp. 1438–1440, 1993. [2] M. R. Visokay and R. Sinclair, “Direct formation of ordered CoPt and FePt compound thin films by sputtering,” Appl. Phys. Lett., vol. 66, pp. 1692–1694, 1995. [3] T. Shima, K. Takanashi, Y. K. Takahashi, and K. Hono, “Preparation and magnetic properties of highly coercive FePt films,” Appl. Phys. Lett., vol. 81, pp. 1050–1052, 2002. [4] G. Li, H. Saito, S. Ishio, T. Shima, K. Takanashi, and Z. Xiong, “MorFePt elongated phology and domain pattern of epitaxially grown particles,” J. Magn. Magn. Mater., vol. 319, pp. 73–79, 2007. [5] L. Zhang, Y. K. Takahashi, K. Hono, B. C. Stipe, J. Y. Juang, and M. -ordered FePtAg-C granular thin film for thermally asGrobis, “ sisted magnetic recording media,” J. Appl. Phys., vol. 109, pp. 07B7031–07B703-4,, 2011. [6] C. C. Chiang, C. H. Lai, and Y. C. Wu, “Low-temperature ordering FePt by PtMn underlayer,” Appl. Phys. Lett., vol. 88, pp. of 152508–152510, 2006.

[7] L. S. Huang, J. F. Hu, G. M. Chow, and J. S. Chen, “Deposition temperature induced magnetic anisotropy variation in FePt-C soft/hard multilayer films,” J. Appl. Phys., vol. 109, pp. 063910–063917, 2011. [8] Y. N. Hsu, S. Jeong, D. N. Lambeth, and D. E. Laughlin, “In situ ordering of FePt thin films by using Ag/Si and Ag/Mn Si/Ag/Si templates,” IEEE Trans. Magn., vol. 36, pp. 2945–2947, 2000. [9] Y. N. Hsu, S. Jeong, D. E. Laughlin, and D. N. Lambeth, “Effects of Ag underlayers on the microstructure and magnetic properties of epitaxial FePt thin films,” J. Appl. Phys., vol. 89, pp. 7068–7070, 2001. [10] J. S. Chen, B. C. Lim, Y. F. Ding, and G. M. Chow, “Low-temperFePt films for ultra-high density magnetic ature deposition of recording,” J. Magn. Magn. Mater., vol. 303, pp. 309–317, 2006. [11] J. S. Chen, Y. F. Xu, and J. P. Wang, “Effect of Pt buffer layer on structural and magnetic properties of FePt thin films,” J. Appl. Phys., vol. 93, pp. 1661–1665, 2003. [12] J. S. Chen, B. C. Lim, J. F. Hu, Y. F. Ding, G. M. Chow, and G. Ju, FePt-C (0 0 1) tex“Microstructural and magnetic properties of tured nanocomposite films grown on different intermediate layers,” J. Phys. D: Appl. Phys., vol. 41, pp. 205001-1–205001-6, 2008. [13] J. S. Chen, B. C. Lim, J. F. Hu, Y. K. Lim, B. Liu, and G. M. Chow, FePt films with perpendicular anisotropy de“High coercivity posited on glass substrate at reduced temperature,” Appl. Phys. Lett., vol. 90, pp. 042508–042510, 2007. [14] B. C. Lim, J. S. Chen, J. F. Hu, Y. K. Lim, B. Liu, G. M. Chow, and G. Ju, “Improvement of chemical ordering of FePt (001) oriented films by MgO buffer layer,” J. Appl. Phys., vol. 103, pp. 07E143-1–07E143-3, 2008. [15] J. S. Chen, B. C. Lim, J. F. Hu, B. Liu, G. M. Chow, and G. Ju, “Low FePt-C (001) films with high coercivity and temperature deposited small grain size,” Appl. Phys. Lett., vol. 91, pp. 132506-1–132506-3, 2007. ordered epitaxial FePt [16] G. R. Trichy, J. Narayan, and H. Zhou, “ (001) thin films on TiN/Si (100) by pulsed laser deposition,” Appl. Phys. Lett., vol. 89, pp. 132502–132504, 2006. [17] Y. Tsuji, S. Noda, and Y. Yamaguchi, “Structure and magnetic property -FePt nanoparticles on TiN/a-Si underlayers,” J. of c-axis oriented Vac. Sci. Technol. B., vol. 25, pp. 1892–1895, 2007. [18] Y. Tsuji, S. Noda, and S. Nakamura, “Nanostructure and magnetic -FePt nanoparticles and nanocrysproperties of c-axis oriented talline films on polycrystalline TiN underlayers,” J. Vac. Sci. Technol. B., vol. 29, pp. 031801–031810, 2011. [19] E. Yang, S. Ratanaphan, J. Zhu, and D. E. Laughlin, “Structure and -FePt thin films on TiN/RuAl underlayers,” magnetic properties of J. Appl. Phys., vol. 109, pp. 07B770-1–07B770-3, 2011. [20] K. Barmak, J. Kim, L. H. Lewis, K. R. Coffrey, M. F. Toney, A. J. Kellock, and J. U. Thiele, “On the relationship of magnetocrystalline CoPt (001) and FePt anisotropy and stoichiometry in epitaxial (001) thin films,” J. Appl. Phys., vol. 98, pp. 033904–033913, 2005. [21] T. Shima, K. Takanashi, Y. K. Takahashi, and K. Hono, “Coercivity exceeding 100 kOe in epitaxially grown FePt sputtered films,” Appl. Phys. Lett., vol. 85, pp. 2571–2573, 2004. [22] C. J. Jiang, J. S. Chen, J. F. Hu, and G. M. Chow, “FePt-TiO exchange coupled composite media with well-isolated columnar microstructure for high density magnetic recording,” J. Appl. Phys., vol. 107, pp. 123915–123921, 2010. [23] P. L. Kim and J. C. Lodder, “Thickness dependence of structural, magnetic properties and reversal mechanism of Co-Cr-Ta/Cr longitudinal recording media,” J. Magn. Magn. Mater., vol. 242, pp. 395–397, 2002. [24] G. Bertotti, Hysteresis in Magnetism: For Physicists, Materials Scientists, and Engineers. New York: Academic, 1998. [25] K. Srinivasan, S. N. Piramanayagam, R. Sbiaa, and R. W. Chantrell, “Thermal stability and the magnetization process in CoCrPt-SiO perpendicular recording media,” J. Magn. Magn. Mater., vol. 320, pp. 3041–3045, 2008. [26] Y. C. Feng, D. E. Laughlin, and D. N. Lambeth, “Formation of crystallographic texture in rf sputter-deposited Cr thin films,” J. Appl. Phys., vol. 76, pp. 7311–7316, 1994.