Growth and Characterization of Epitaxial Fe Nanoislands on Si(001): Size Effects on Ferromagnetic Anisotropy Shu-Fang Chen, Hung-Chin Chung, and Chuan-Pu Liu* Department of Materials Science and Engineering and Center for Micro/Nano Science and Technology, National Cheng Kung UniVersity, Taiwan
CRYSTAL GROWTH & DESIGN 2008 VOL. 8, NO. 11 3885–3888
ReceiVed NoVember 2, 2007; ReVised Manuscript ReceiVed July 29, 2008
ABSTRACT: Self-assembled Fe islands were successfully grown on Si(001) by ion-beam sputtering at room temperature. We present a combined study on the growth, structure, and magnetic properties of Fe islands on Si(001). Nanometer-scale iron islands, ranging from 25 to 100 nm with narrow size distributions, can be achieved with a silicide interfacial layer analogous to the S-K growth mode. Highresolution transmission electron microscopy results confirm the epitaxial nature of the Fe islands with Si. Furthermore, magnetic anisotropy of different-sized islands has been studied by superconducting quantum interference device measurements. The perpendicular magnetic anisotropy falls away with increasing island size. We found that surface effects from the morphologies of three-dimensional islands are mainly responsible for the spin reorientation. Heteroepitaxial growth of self-assembled magnetic nanoparticles has been a subject of considerable studies for decades, since they are expected to be applied in magnetic nanodevices and high-density magnetic recording media.1,2 A key advantage in using selfassembled nanoparticle arrays in data storage media is the potential for a high degree of magnetic uniformity, which would improve the signal-to-noise ratio, and could lead to a new recording paradigm with a single particle per bit. The perpendicular magnetization properties appear to be of great importance for such applications,3 and are determined by the competition between different anisotropy energies. Therefore, investigation of material structure and surface morphology at the nanometer scale is necessary in order to understand the magnetic properties of ferromagnetic nanostructures. From previous investigations, it is clear that the composition, shape, and size of particles all have strong influences on reversal behavior.3,4 Although numerous efforts have been made to study the magnetic properties of nanostructures, the correlation between microstructure and magnetism in an epitaxial nanoisland system has still not been well studied. Besides, the anisotropic magnetic behaviors of magnetic nanostructures have not been particularly emphasized, especially for Fe nanoislands. Previous reports focused mainly on granular films5 and two-dimensional islands with coverage lower than 1 ML.6 In this paper, we show that selfassembled Fe islands on Si(001) with a narrow size distribution could be successfully grown by ultrahigh vacuum ion-beam sputtering (UHV-IBS) at room temperature. The relationship between microstructure and magnetic properties of epitaxial Fe islands was then studied. We found that even at high coverage, the magnetization direction of a Fe nanoisland does not completely rotate toward the film plane. The experiments were performed in an UHV-IBS system with a base pressure of 1 × 10-10 torr. Si(001) substrates were cleaned with a standard RCA procedure, followed by heating at 900 °C for 30 min to remove surface oxide in the growth chamber. After surface cleaning, Fe (99.99% purity) was deposited on Si(001) at room temperature with a beam voltage of 2.25 keV and argon gas flux of 4 sccm. Fe coverage was calibrated in pseudomorphic monolayers (ML) using atomic force microscopy (AFM). X-ray photoelectron spectroscopy (XPS) with monochromatic Al KR radiation was conducted to identify the chemical bonding states of Fe islands. The samples were cleaned by Ar+ ion sputtering at 1.5 keV energy for 5 min prior to XPS measurement to remove the surface contamination. Structure and morphologies of the islands were investigated by a JEOL 2100F transmission electron micro* Corresponding author. E-mail:
[email protected].
Figure 1. Typical Fe 2p3/2 core level spectrum of 13 ML Fe nanoislands on Si(001). The sample is after ordinary Ar ion sputtering at 1.5 keV energy.
scope (TEM) operated at 200 keV. TEM samples were prepared by mechanical thinning, followed by ion milling. In addition, the magnetic behaviors of Fe dots were analyzed by superconducting quantum interference device (SQUID) magnetometry (maximum magnetic field of (5000 Oe) at room temperature. The chemistry of the resulting surface islands was studied by means of the Fe 2p XPS signals, as shown in Figure 1 for the sample of 13 ML Fe on Si (001) (1 ML is equivalent to the nominal surface atomic density of bcc Fe(110), 1.7 × 1015 atoms/cm2). The spectrum shows metallic Fe features, where the asymmetrical shape of the Fe 2p3/2 peak and the binding energy at 706.8 eV are in agreement with those from bulk Fe.7 The Fe 2p3/2 and Fe 2p1/2 signals in Figure 1 are pronounced and only trace amounts of iron oxide can be detected from the Fe 2p3/2 peak at the binding energy of 711.1 eV,7 caused by natural oxidation when exposed to air. As a result, the resulting surface islands synthesized by this method are confirmed to be made of pure metallic Fe. Figure 2a-c shows plan-view bright-field TEM images of the Fe islands with Fe coverage of 5, 8, 13 ML, respectively, where the sample is tilted to the two-beam condition, g ) 220, so that only (220) planes satisfy Bragg diffraction. These islands are randomly distributed on the surface of the Si(100) substrate. The morphology of the Fe islands with increasing Fe coverage gradually changed from circular (e8 ML) to irregular (>8 ML), resulting from the coalescence of small particles. Figure 2 reveals coherently strained islands without defects as indicated by the characteristic dark/bright contrast related to an inhomogeneous lattice deformation
10.1021/cg701083e CCC: $40.75 2008 American Chemical Society Published on Web 09/18/2008
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Figure 2. TEM plan-view images of Fe islands grown on Si(001) with different Fe coverage of (a) 5 ML (b) 8 ML, where the inset is the corresponding diffraction pattern, and (c) 13 ML; (d) TEM cross-sectional image of the 8 ML Fe islands; (e) STEM-HAADF image of the 13 ML sample with a series of EELS spectra for the Fe M2.3 edge and Si L2,3 edge taken from a line scan as indicated in the image; (f) enlarged image of (d); and (g) HRTEM image after inverse Fast Fourier transform processing of (f). The dotted circle and the white arrows show the interface and dislocations, respectively.
in a three-dimensional structure. The diffraction pattern, in the inset of Figure 2b, indicates that iron grows in the body-centered cubic (bcc) structure with the epitaxial relations of Fe(001)|Si(001) and Fe100|Si100. The cross-sectional TEM image of the 8 ML sample, as shown in Figure 2d, shows that the island lies on a thin “wetting layer”. The thickness of the wetting layer is approximately 0.75 nm (equivalent to 2 ML), and the height (width) of the island is 4.2 nm (47.2 nm). The presence of the silicide layer has also been observed in previous studies,8,9 indicating that a strong intermixing occurs at the Fe/Si interface, even at room temperature. The Fe-Si bonding (binding energy ) 707.3 eV7) could not be detected in the XPS measurement, mainly due to the relatively lower amount of silicide. Therefore, we used electron energy-loss spectroscopy (EELS) to investigate the composition of the thin silicide layer. Figure 2e is the scanning TEM (STEM)-high-angle annular darkfield (HAADF) image of the 13 ML sample with a series of EELS spectra for Fe M2,3 edges and Si L2,3 edges, as the STEM probe scans from the top of the island to the Fe/Si interfacial layer as indicated. The Si L2.3 edge is only detected within the interfacial layer. By EELS quantification, the composition of the silicide layer is close to Fe3Si. So we demonstrated that Fe islands were formed on Si(001) via a pathway analogous to the Stranski-Krastanow (SK) growth mode. When the deposition amount exceeded 2 ML, the following deposited materials (Fe) tended to nucleate in 3D islands on the silicide layer. Figure 2f is the high-resolution TEM image of a Fe island, showing that the Fe island exhibits single crystallinity. The Fe islands are coherently matched to the Si(001) surface, with the lattice parameter larger (aFe ) 2.86 Å) than half that of Si (aSi ) 5.43 Å), causing in-plane compression of the Fe lattice at the Fe/Si(001) interface, as well as expansion of the Fe lattice normal to the interface. Thus, the Fe lattice undergoes a tetragonal
distortion in the thin layer, but relaxes toward its natural bcc form in the thick layers. Because of the large misfit (5.6%) between Fe and Si, lattice distortion and structure defects spontaneously appeared to reduce the strain energy. In addition, the oxidation layer is marked with the white dotted lines, shown in Figure 2f, where the thickness is about 0.8 nm, thinner than the previous study.10 The behavior of less oxidation may be due to the large strain induced from the epitaxial growth and different growth methods. Figure 2g is the inverse fast Fourier transform (IFFT) image taken with the 220g vector, exhibiting the presence of dislocations at the interface, marked by the white arrows, and with no defects inside the Fe islands. However, dislocations were formed in larger islands. Figure 3a is the dot size distribution with coverage, along with the corresponding logarithmic normal fits, showing that the peak is located at 27.3, 45.3, and 97.2 nm for the coverage of 5, 8 and 13 ML, respectively. The width of size distribution broadened as coverage increased, which is the result of island coalescence, indicating that dots with uniform size could be obtained at lower coverage. Figure 3b is the plot of density and average size of the Fe dots as a function of Fe coverage. The island density for 5, 8, 13 ML is 1.5 × 1010, 1.2 × 1010, and 3 × 109/cm2, respectively. Apparently, the island size increases while density decreases, with increasing Fe coverage. To investigate the relationship between microstructure and magnetism in Fe nanostructures, each sample was analyzed by SQUID magnetometry, measuring hysteresis loops at room temperature, with the field applied both perpendicular and parallel to the film plane. Figure 4a-c shows the evolution of the M-H loops with increasing deposition coverage, where the insets display both the enlarged loop of the central portion and magnetization axis of each sample. Clearly, all three samples are ferromagnetic with a canted orientation (neither in-plane, nor out-of-plane). The angle
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Figure 3. (a) Diameter distributions of the islands shown in Figure 1 with the corresponding logarithmic normal fits. (b) Average size and density of Fe islands as a function of Fe coverage.
of the magnetization direction with respect to the film plane for Fe islands with 8, 13, and 29 ML is 50.2°, 37.6°, and 29.2°, respectively (as shown in the insets of Figure 4a-c). As Fe thickness increases, the easy-magnetized axis gradually rotates toward the Si substrate. Besides, the shapes of the M-H loops of all the samples are not square shaped, revealing that the Fe islands are of multidomain structures. It is well-known that for bulk Fe films, the easy axes, due to the magneto-crystalline anisotropy, are along the (100) of the bcc structure. Nevertheless, there are studies11,12 suggesting that perpendicular magnetization occurs in ultrathin magnetic films and is only stable in an intermediate thickness range. For the bcc Fe/ Si(111) studied in ref 13, they showed three different magnetization orientations: (1) No ferromagnetic response is detected for nominal thickness below 4.5 ML; (2) for thickness between 4.5-6.3 ML, the magnetization lies perpendicular to the film surface; and (3) for film thickness above 7 ML, the magnetic easy direction rotates toward the film plane. Although the Fe thickness for the spin reorientation has been shown to occur at ∼6 ML,11-13 we show that the critical thickness is larger than 29 ML for our case. In our case here, since the morphology of the Fe nanostructures as shown in Figure 2 is not a 2D film but 3D islands, it is necessary to consider the contribution of the surface characteristics. As illustrated in Figure 5, the morphology of the Fe islands in this study is lenslike; that is, the height is smaller than the diameter, indicating that the shape anisotropy favors in-plane behavior. On the other hand, the surface anisotropy resulting from the broken symmetry at the surface, favors out-of-plane behavior. In addition, the 0.75-nmthick wetting layer (Fe3Si) as mentioned previously is far below the critical thickness of about 1.8 nm at room temperature14 for the ferromagnetic onset of Fe3Si. Therefore, the Fe3Si layer is not responsible for the perpendicular anisotropy. Qualitatively, the competition between surface and shape anisotropies leads to the canted magnetization of Fe islands. From a calculation that assumes the island shape as lens-like, the interface to volume (surface to volume) ratio decreases from 0.67 (0.71) to 0.27 (0.32), as Fe coverage increases from 8 to 29 ML. Therefore, the contributions from surface and interface are relatively lower for larger islands; thus, surface/interface anisotropy is not strong enough to support perpendicular magnetization. Besides, Hjortstam et al.15 have shown that if the cubic symmetry of the crystal is broken by a tetragonal distortion, an important contribution may come from the volume
Figure 4. SQUID M-H hysteresis loops of different Fe coverage: (a) 8 ML, (b) 13 ML, (c) 29 ML with the insets for further enlargement. The loops were measured in a magnetic field of up to ( 5 kOe at room temperature, and the field was applied perpendicular (open circles) and parallel (filled circles) to the film plane.
Figure 5. Schematic representation of the contribution from the shape and surface magnetic anisotropies of Fe islands.
contribution, leading to a perpendicular easy axis, which accords with the case of the epitaxial Fe islands grown on Si substrate. Interestingly, the strain would be largely relaxed by introducing defects in larger islands as having been demonstrated, but the magnetization direction did not completely rotate to the film plane, even with the high coverage of 29 ML here. Hence, we suggested that the strain-induced perpendicular anisotropy would not be the dominant factor. It is well-known that the shape anisotropy given by -RMs2, where R is the demagnetization factor, is responsible for in-plane magnetization. The R is 2π for a perfectly smooth film, but it is structure sensitive and becomes smaller than 2π for a rough surface. The Fe islands discussed here can be regarded as a rough surface due to growth in 3D, resulting in a smaller R. Thus, the morphology of the Fe islands causes a reduced tendency for inplane magnetization as compared to the reported continuous films.
3888 Crystal Growth & Design, Vol. 8, No. 11, 2008 Therefore, we proposed that the delay of the spin reorientation for film thickness correlates directly to the morphological effect. The understanding of the magnetic behavior of Fe islands may help in the development of magnetic dots in the application of high-density magnetic recording media. In conclusion, we report on the morphological evolution, microstructure and magnetic anisotropy of epitaxial Fe islands on Si(001). The relation between the magnetic properties and interfacial microscopic structure was clearly investigated. We have shown that under proper growth conditions, the self-assembly of iron islands on Si (001) leads to the formation of nanometer-scale magnetic iron islands, with narrow size distributions. Furthermore, the surface anisotropy resulting from a Fe island with high surface to volume ratio causes the easy axis of magnetization to be canted toward out of plane, with low Fe coverage.
Acknowledgment. The work has been supported by the National Science Council of Taiwan (Grant Nos. NSC95-2221-E006-079-MY3 and NSC95-2221-E-006-080-MY3).
References (1) Tian, D.; Jona, F.; Marcus, P. M. Phys. ReV. B 1992, 45, 11–216. (2) (a) Kadau, K.; Meyer, R.; Entel, P. M. Surf. ReV. Lett. 1999, 6, 35. (b) Ohno, Y.; Young, D. K.; Beschoten, B.; Matsukura, F.; Ohno, H.; Awschalom, D. D. Nature 1999, 402, 790.
Communications (3) Datta, S.; Das, B. Appl. Phys. Lett. 1990, 56, 665. (4) Cowburn, R. P.; Adeyeye, A. O.; Welland, M. E. Phys. ReV. Lett. 1998, 81, 5414. (5) Xiao, G.; Wang, J.-Q. J. Appl. Phys. 1994, 75, 6604. (6) Tsay, J. S.; Fu, T. Y.; Lin, M. H.; Yang, C. S.; Yao, Y. D. Appl. Phys. Lett. 2006, 88, 102506. (7) R¨ uhrnschopf, K.; Borgmann, D.; Wedler, G. Thin Solid Films 1996, 280, 171. (8) Gallego, J. M.; Garcia, J. M.; Alvarez, J.; Miranda, R. Phys. ReV. B 1992, 46, 13–339. (9) Alvarez, J.; Vazquez de Parga, A. L.; Hinarejos, J. J.; de la Figuera, J.; Michel, E. G.; Ocal, C.; Miranda, R. Phys. ReV. B 1993, 47, 16– 048. (10) Wang, C. M.; Baer, D. R.; Thomas, L. E.; Amonette, J. E.; Antony, J.; Qiang, Y.; Duscher, G. J. Appl. Phys. 2005, 98, 094308. (11) Stampanoni, M.; Vaterlaus, A.; Aeschlimann, M.; Meier, F. Phys. ReV. Lett. 1987, 59, 2483. (12) Qiu, Z. Q.; Pearson, J.; Bader, S. D. Phys. ReV. Lett. 1993, 70, 1006. (13) Garreau, G.; Hajjar, S.; Bubendorff, J. L.; Pirri, C.; Berling, D.; Mehdaoui, A.; Stephan, R.; Wetzel, P.; Zabrocki, S.; Gewinner, G.; Boukari, S.; Beaurepaire, E. Phys. ReV. B 2005, 71, 094430. (14) Herfort, J.; Schonherr, H.-P.; Jenichen, B. J. Appl. Phys. 2008, 103, 07B506. (15) Hjortstam, O.; Baberschke, K.; Wills, J. M.; Johansson, B.; Eriksson, O. Phys. ReV. B 1997, 55, 15026. (16) Suzuki, T.; Notarys, H.; Dobertin, D. C.; Lin, C.-J.; Weller, D.; Miller, D. C.; Gorman, G. IEEE Trans. Magn. 1992, 28, 2756.
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