MgO Layers Grown on InAs

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Microstructural Changes of Epitaxial Fe/MgO Layers Grown on InAs(001) Substrates Kyung-Ho Kim,†,‡ Hyung-jun Kim,*,† Jae-Pyung Ahn,† Jun Woo Choi,† Jun Hyun Han,§ Rizcky Tamarany,† Seung-Cheol Lee,† Sung Ok Won,† Joonyeon Chang,† and Young Keun Kim‡ †

Korea Institute of Science and Technology, Seoul 136-791, Korea Department of Advanced Material Science and Engineering, Korea University, Seoul 136-713, Korea § Department of Nano Materials Engineering, Chugnam National University, Daejeon 305-764, Korea ‡

bS Supporting Information ABSTRACT: The microstructural evolution and the effect on in-plane magnetic properties of epitaxial Fe/MgO layers grown on InAs(001) substrates have been investigated as a function of MgO growth temperature. The Fe grows three-dimensional islands with two different in-plane textures along [010] and [110] directions on the MgO layers grown below 200
C in remarkable contrast to two-dimensional Fe layers on the MgO layers grown above 300
C. As the MgO growth temperature increases, both tensile-strained MgO and the subsequent Fe are simultaneously relaxed, and the distribution of 45
-rotated Fe lattices with [010] texture becomes dominant. The experimental results imply that the microstructural evolution of the Fe is strongly influenced by the underlying misfit strain within the MgO layers grown at different temperatures. The two different epitaxial relationships of the Fe islands lead to no magnetic anisotropy, while the Fe layer with the single epitaxial relationship of Fe[010]//MgO[110]//InAs[110] shows cubic magnetic anisotropy.

1. INTRODUCTION Efficient electrical injection of spin-polarized electrons from a ferromagnetic metal (FM) into a semiconductor (SC) has been a crucial issue for the applications of spin-based electronics.1
4 The spin injection and detection is known to be significantly influenced by the spin polarization of the FM and the interface resistance at a FM/SC junction.5 Recently, the epitaxially grown MgO tunnel barrier between the FM and SC has been widely employed because it removes a conductivity mismatch problem and serves as a diffusion barrier that facilitates crystalline FM layer growth at relatively high temperature without intermixing.2,3,6
8 Furthermore, it is predicted that highly spin-polarized wave functions in FMs match those in the MgO. As a result, there is a relatively long evanescent wave in the barrier that promotes tunnelling of the highly spin-polarized current.9,10 To date, Fe/MgO/GaAs heterostructures have been extensively attempted for spin-based electronic applications such as spin-light-emitting diodes (LEDs) because of a relatively small lattice mismatch of ∼0.7%.2,3,11,12 Recently, the use of SC substrates has been diverse with Si, Ge, and InAs for many purposes such as longer spin relaxation time.13
15 Among the alternative SC substrates, we focus on an InAs substrate because of its fascinating properties of a narrow band gap with higher electron mobility and large spin
orbit interaction (SOI), which is desirable for spintronic device applications despite the considerably larger lattice mismatch (>4.2%) to the MgO. In r 2011 American Chemical Society

particular, an InAs quantum well (QW) structure has been employed to implement a spin field effect transistor (spin-FET) due to Rashba effect.16,17 For these reasons, we believe that the epitaxial growth of two-dimensional (2D) Fe and MgO layers on InAs (001) surfaces are prerequisites for the improvement of electrical spin injection into the InAs QW. It has been reported that the MgO layers can be epitaxally grown on Ge and GaAs substrates at significantly wide growth temperature between 125 and 550
C.14,15,18
20 However, MgO growth temperatures higher than 450
C seriously degrade the InAs QW structure by selective arsenic (As) desorption, undulating interface formation of the strained InAs QW layer, or Si dopant segregation within an In0.52Al0.48As carrier supply layer.21,22 These result in significant decrease of the electron mobility as well as the loss of spin controllability in the InAs QW structure. Furthermore, it has been often claimed that Fe on MgO substrates grows in three-dimensional (3D) islands via Volmer
Weber (VW) or Stranski
Kristanov (SK) growth mode at the growth temperature between room temperature (RT) and 700
C.23,24 According to the literature, the microstructure, morphology, and magnetic properties of the 3D Fe islands are significantly affected by the experimental parameters Received: January 14, 2011 Revised: April 27, 2011 Published: May 04, 2011 2889

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of growth temperature and growth geometry. For example, Fe on MgO substrates appears in the form of either a 2D layer or 3D islands depending on the experimental parameters.25
27 Moreover, the irregular magnetization direction and diminished polarization of the 3D Fe islands result in poor electrical spin injection compared with the 2D Fe layer. Therefore, the study on the morphology and epitaxial relationship of the Fe and MgO layers is essential particularly when large lattice-mismatched InAs substrate is used. In addition, the magnetic properties of the Fe with respect to its morphology evolution are also importan from the spin injection point of view. Most of the literature has reported Fe microstructural characterization with respect to its own growth temperature governing the surface diffusion of Fe adatoms.12,26,28 In this study, we have examined the epitaxial relationships, microstructural evolutions, and corresponding magnetic properties of Fe/MgO layers grown on InAs substrates as a function of the MgO growth temperature. It should be clearly pointed out that the MgO growth temperature being a key experimental parameter in this study leads to crystalline, morphological, and in-plane magnetic anisotropy changes of the subsequent Fe layers due to the different misfit strain between the MgO and InAs substrate.

2. EXPERIMENTAL SECTION All of the Fe/MgO/InAs heterostructures used in this study were prepared using a cluster molecular beam epitaxy (MBE) system composed of a semiconductor MBE, a metal MBE, and a magnetron sputtering unit for the MgO growth. After “epi-ready” semi-insulating InAs(001) substrates were introduced into an introduction chamber without any chemical treatments, surface oxide layers were thermally desorbed by heating up to 480
C under As flux to avoid selective As desorption. The epitaxial layers of 300-nm-thick InAs buffer, 4-nm-thick MgO, and 7-nm-thick Fe were subsequently grown without vacuum break, ensuring the exclusion of native oxide formation at the interfaces. The growth of the InAs buffer layer consistently produced a streaky (2  4) As-terminated surface. The MgO layers were consecutively grown using a sintered MgO target with a rf power of 100 W without additional oxygen gas. The growth temperature of the 4-nm-thick MgO layer was changed as a key experimental parameter in the range from RT to 400
C. The growth temperature of the subsequent 7-nm-thick Fe was kept constant at 200
C. The growth rates of the MgO and Fe were 0.233 and 0.021 Å/s, respectively. All the samples were finally capped by aluminum (Al) layers at RT to prevent oxidation of Fe in the atmosphere and immediately removed from the MBE chamber without postgrowth annealing. The microstructural evolutions of the Fe and MgO were characterized by means of a high-resolution transmission electron microscope (HRTEM, Titan) operated at 300 kV. In particular, the Fe morphology was carefully observed in both plane-view and crosssection HRTEM with selected area diffraction (SAD) patterns. The inplane textures of the Fe and InAs were examined by using φ-scan mode of an X-ray diffractometer (Rigaku D/MAX 2500 V) using 8 kW with Cu KR radiation. Normal incidence X-ray diffraction (XRD) was used for qualitative analysis and for determining the interplanar spacing of the lattice planes parallel to the interface. To evaluate in-plane orientation, we use asymmetric reflection and grazing incidence diffraction (GID). The GID measurement was performed in the direction of the diffraction angle perpendicular to the direction of the incident angle (near the critical angle, 0.37
) inserting Soller slits with apertures of 0.41
and 0.45
on both the incident and receiving sides using 9 kW. The samples were tilted at 45
during the φ-scan, when the incident and diffracted X-ray beam directions were satisfied with Bragg’s diffraction condition of

Figure 1. Cross-sectional HRTEM images of the Fe/MgO/InAs structures with the MgO layers grown at (a) RT, (b) 200
C, (c) 300
C, and (d) 400
C, respectively. Note that the scales of panels a and b are different from those of panels c and d with 100 and 50 nm scale bars, respectively. Fe(110) or InAs(220). The in-plane magnetic anisotropies of Fe were ex situ measured using an alternating-gradient magnetometer (AGM) with a sensitivity up to 10
9 emu.

3. RESULTS AND DISCUSSION Figure 1 shows the cross-sectional HRTEM images of Fe/ MgO/InAs heterostructures with the MgO layers grown at RT, 200
C, 300
C, and 400
C, respectively. It is clearly seen that the morphology of MgO appears as a 2D layer at all of the growth temperatures. In contrast, the subsequent Fe forms 3D islands below 200
C and a continuous 2D layer above 300
C. This morphological distinction of the Fe is supported by the observation that the interspace between Fe islands is filled with 2890

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Figure 2. Plane-view HRTEM images and the corresponding SAD patterns of the Fe/MgO/InAs structures with the MgO layers grown at (a) RT, (b) 200
C, and (c) 300
C.

amorphous Al capping on the MgO layers grown at RT and 200
C. Although the Fe layers on the MgO layers grown at 300 and 400
C show periodic contrast discontinuities with some notches on the top surface, they are apparently continuous 2D Fe layers in the absence of the interspaces filled with the Al capping. In addition, comparing the size and height-to-base diameter aspect ratio of the 3D Fe islands with respect to the MgO growth temperature, we found that the base diameter increases from 33.4 ( 9.4 to 51.0 ( 18.5 nm and the corresponding aspect ratio decreases from 1:(3.37 ( 0.76) to 1:(6.13 ( 2.54) as the MgO growth temperature increases from RT to 200
C. We postulate that the aspect ratio change results directly from the underlying misfit strain within the MgO layer grown at different temperature and will discuss in detail later. Figure 2 displays the plane-view HRTEM images and the corresponding SAD patterns of the Fe/MgO/InAs structures with the MgO growth temperatures of RT, 200
C, and 300
C, respectively, in which the Fe island shape can be clearly seen. The elemental mapping measured by electron energy loss spectroscopy (EELS) reveals that the Fe islands are discontinuous on the underlying MgO layers grown at RT and 200
C. In particular, the Fe islands on the MgO layer grown at RT show wider interisland spacing compared with the serpentine-shaped Fe islands on the MgO layer grown at 200
C. The corresponding SAD patterns in the insets represent the bright spots diffracted from both Fe and InAs substrates. It is difficult to find distinguishable

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spots from the MgO layers owing to very small thickness of 4 nm. As opposed to the clear diffraction spots obtained from the Fe layer on the MgO layer grown at 300
C shown in the inset of Figure 2c, the Fe islands on the MgO layers grown at RT and 200
C illustrate ring patterns and diffuse spots, respectively. Interestingly, the individual ring pattern shown in the inset of Figure 2a represents four bright regions along two perpendicular directions of [110] and [110]. This result indicates that the Fe islands are in-plane textured along the two directions despite different in-plane textures. Moreover, it is observed that the lengths of the four individual bright regions become much shorter in the inset of Figure 2b. The corresponding diffuse spots indicate that the Fe islands are more highly in-plane textured along the two directions compared with the ring pattern in the inset of Figure 2a. Now, a question is whether each Fe islands is polycrystalline or single crystalline Fe islands possessing different in-plane textures. This is a very important point to support our claim in this study, and the following zone axis analysis and XRD results clarify the crystallinity and in-plane texture of the Fe. Figure 3 shows the magnified cross-sectional HRTEM images and the corresponding fast Fourier transform (FFT) patterns of each layer in the Fe/MgO/InAs structure with the MgO layers grown at 200 and 300
C. From the zone axis analysis, we understand the in-plane epitaxial relationships of both 3D Fe islands and 2D Fe layer of which the formation strongly depends on the MgO growth temperature. First of all, examining the structure with the Fe layer on MgO layer grown at 300
C in Figure 3c, it is found that the zone axis of the Fe layer is found to be [010] while the zone axes of MgO and InAs are [110], respectively. This result elucidates the epitaxial relationship of Fe[010]//MgO[110]//InAs[110] in which the Fe lattice is inplane 45
-rotated with respect to the underlying MgO and InAs. This epitaxial relationship has been commonly reported for the structures of Fe/MgO/GaAs and Fe/MgO/Si in the literature.12,13 However, it is found that this epitaxial relationship coexists with an unusual epitaxial relationship in the structures of the 3D Fe islands on the MgO layers grown below 200
C. The FFT patterns in Figure 3a,b obtained from the spatially different Fe islands clearly represent the coexistence of the 3D Fe islands with two different epitaxial relationships on the MgO layer grown at 200
C. Figure 3a,b represents the Fe islands with [010] and [110] zone axes, respectively, while both MgO and InAs keep the identical zone axis of [110]. This experimental observation implies that the Fe islands simultaneously have two different in-plane orientations with 45
-rotation and 0
-rotation, which result in the epitaxial relationships of Fe[010]//MgO[110]// InAs[110] and Fe[110]//MgO[110]//InAs[110], respectively. Careful FFT analyses on a number of the Fe islands ensure that the zone axis of the individual Fe islands is exclusively [010] or [110], but not the mixture of both. We have claimed that the Fe morphology evolution observed in this study directly results from the underlying misfit strain in the MgO layers. In order to verify the effect, the lattice spacing of Fe and MgO layers are systematically attained by means of XRD analysis. In addition, the crystallinity of the Fe islands should be confirmed in the larger area of the samples to further clarify the effects of in-plane textures of the Fe islands and layers. Since the layer thicknesses of the Fe and MgO are very thin in the structures, both normal and grazing incidence diffractions are attempted to obtain reliable lattice spacing values. Furthermore, the diffraction patterns obtained from the normal incidence are primarily used for the lattice spacing of the MgO layers because 2891

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Figure 3. Fast Fourier transform (FFT) patterns and the corresponding magnified cross-sectional HRTEM images of the Fe/MgO/InAs structure showing the epitaxial relationships of (a) Fe[010]//MgO[110]//InAs[110] and (b) Fe[110]//MgO[110]//InAs[110] obtained from the spatially separated Fe islands on the MgO layers grown at 200
C and (c) Fe[010]//MgO[110]//InAs[110] obtained from an Fe layer on the MgO layers grown at 300
C.

MgO(200) and InAs(220) peaks are unfortunately overlapped at around 42.5
in grazing incidence diffraction results. The lattice spacing of the Fe with respect to the MgO growth temperature is calculated by the consideration of both diffraction pattern results. Figure 4 illustrates the representative XRD patterns and the resulting lattice parameter of the Fe and MgO in the Fe/MgO/ InAs structures at the different MgO growth temperatures from RT to 400
C. In the normal incidence XRD, Fe(002) and MgO(002) peaks are not shown in the structure with the MgO layer grown at RT, while both peaks clearly appear at the higher MgO growth temperatures. On the other hand, Fe(200) peaks are shown for all the MgO growth temperatures in the grazing incidence XRD results as shown in the inset of Figure 4a. It is understood that the appearance of the Fe peak in the structure with the MgO layer grown at RT in opposition to the normal incidence XRD results from the longer penetration depth of the grazing incidence X-ray beam. Unfortunately, the MgO peak cannot be seen in the grazing incidence XRD result due to the overlap with InAs(220) peak. Figure 4b displays the variations of the Fe and MgO in-plane lattice parameter to systematically investigate the residual misfit strain within Fe and MgO layers, which is represented by the peak shift in the XRD results. Evidently, the MgO layer grown at low temperatures such as 200
C is under tensile strain with the in-plane lattice constant much larger than the bulk value. The tensile-strained MgO becomes relaxed as the growth temperature increases. Moreover, the lattice constant of the MgO layers decreases and approaches its bulk value at MgO growth temperatures higher than 300
C. On the other hand, the highly tensile-strained Fe on the MgO layers becomes relaxed as the MgO growth temperature increases regardless of the Fe morphology. Furthermore, the lattice constant of the Fe layer comes close to its bulk value on the MgO layer grown at 400
C. The lattice spacing change of the Fe obviously supports the Fe morphology evolution and the aspect ratio change of Fe islands by the underlying misfit strain. Therefore, the MgO growth temperature used in this study is a key experimental parameter determining the underlying misfit

strain, leading to the different Fe morphology and epitaxial relationship. The in-plane epitaxial relationships and quantitative ratio of the Fe islands between the two different zone axes with respect to the MgO growth temperature are examined by means of XRD φscan analysis. Figure 5 depicts the corresponding XRD φ-scan results obtained from the Fe/MgO/InAs structures with the MgO layers grown at RT, 200
C, and 300
C. For the Fe islands on the MgO layers grown at RT and 200
C as shown in Figure 5a,b, there are eight characteristic peaks of Fe(110) reflections being 45
apart. Among them, four peaks match the peaks of InAs(220) reflections being 90
apart whereas the other four peaks are located in the middle of the peaks of InAs(220) reflections. The relative intensity of the latter four peaks of Fe(110) reflections becomes higher as the MgO growth temperature increases from RT to 200
C. The appearance of Fe(110) reflections in the structure with the MgO layer grown at RT presumably results from the smaller incident angle of X-ray for φ-scan as for grazing incidence results in Figure 4. Moreover, the crystallinity of the Fe islands on the MgO grown at RT is not inferior to those on the MgO layers grown at higher temperature. On the other hand, in the Fe/MgO/InAs structure with the MgO layer grown at 300
C as shown in Figure 5c, there are only four characteristic peaks of Fe(110) reflections being 90
apart, which are located in the middle of the peaks of InAs(220) because the whole 2D Fe layer exclusively keeps the single eiptiaxial relationship of Fe[010]//MgO[110]//InAs[110]. This experimental result elucidates that the Fe islands of two different in-plane textures along Æ100æ and Æ110æ directions coexist in the structure with the MgO layer grown below 200
C. The quantitative analysis for the ratio of Fe islands with the in-plane textures along between Æ100æ and Æ110æ are simply investigated in the structures with the MgO layers grown at RT and 200
C shown in Figure 5a, b. The area of each Æ100æ and Æ110æ texture peak of the Fe islands from the XRD φ-scan results is calculated when the background elimination of the peaks is included. From the total areas of the four Æ100æ and four Æ110æ peaks, the Fe island ratio of the Æ100æ 2892

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Figure 5. X-ray φ-scan diffraction patterns of the Fe/MgO/InAs structures with the MgO layers grown at (a) RT, (b) 200
C, and (c) 300
C.

Figure 4. (a) Normal incidence XRD patterns of the Fe/MgO/InAs structures with respect to the MgO growth temperature. The inset is grazing incidence diffraction patterns showing exclusively Fe(200) peaks. (b) Calculated in-plane lattice spacing of Fe and MgO with respect to the MgO growth temperature deduced from both normal and grazing incidence XRD patterns.

and Æ110æ textures is estimated. The analysis result reveals that Æ100æ-to-Æ110æ textured Fe island ratio on the MgO layer grown at RT is approximately 0.33:1. The ratio drastically changes to 2.17:1 in the structure with the MgO layer grown at 200
C. In other words, the number of the Fe islands along Æ100æ texture becomes dominant as the MgO growth temperature increases. Our XRD and φ-scan results evidently prove that the Fe islands

on the MgO layer grown at 200
C are perfect single-crystalline and are the assembly of different in-plane textured islands. It is expected that the Æ110æ-textured Fe gradually disappears and Æ100æ-textured Fe exclusively forms the 2D layer at the MgO growth temperature higher than 300
C as shown in Figure 5c. According to the experimental results, Fe islands of 0
rotation and 45
-rotation coexist on the MgO layers grown below 200
C. However, the formation of the Fe islands of 0
-rotation is surprising because it has been well-known that the preferential site of Fe adatoms are exclusively on oxygen atoms in the MgO layer due to the strong affinity between Fe and O. In order to compare the relative stability between 0
- and 45
rotation, model calculations based on density functional theory are performed. We assume that two Fe atoms serve as a nucleus for thin film growth, which is reasonable when the growth temperature is sufficiently low. One Fe atom is located on the oxygen on-top site since this site is the most energetically favorable in the case of a single Fe atom on MgO surface. The second Fe determines the distinction of 0
- and 45
-rotation. If a second Fe prefers the on-top site of neighboring oxygen atom, the formed Fe dimer serve as a nucleus of 45
-rotation. If a second Fe atom favors the Mg on-top site, a nucleus of 0
-rotation is formed. Based on the calculation of bulk terminated MgO(001) surface, the calculation reveals that the energy of the 45
configuration is 0.3056 eV lower than the 0
configuration. This calculation result evidently supports the 45
-rotated Fe lattice with respect to the MgO layer grown above 300
C. 2893

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Figure 6. AGM hysteresis loops obtained from the Fe/MgO/InAs structures with the MgO layers grown at (a) 200 and (b) 300
C. Red and blue lines correspond to the hysteresis loops when the magnetic field is applied along Fe[100] and Fe[110] directions, respectively. Note that the scale of the magnetic field in panel a is 10 times larger than that in that in panel b. The insets are the angle dependence of Mr/Ms ratio.

Although considering the effect of the tensile strain in the MgO layers, which corresponds to the growth below 200
C, the reversal of the stability between two configurations is not believed to have happened in order to explain the formation of Fe islands of 0
configuration. Another possible reason can be found by the existence of Mg or O vacancies on the top surface of the MgO layer, which will be investigated in the future. Next, the in-plane magnetic anisotropies of the two different Fe morphologies of 3D islands and 2D layers are measured by AGM. Figure 6 represents the magnetic hysteresis (M-H) loops of the Fe/MgO/InAs structures with the MgO layers grown at 200 and 300
C where 3D Fe islands and 2D Fe layers exist, respectively. There is a significant contrast in the M-H loops between the 3D Fe islands and 2D Fe layer. The hysteresis loops are measured with the applied magnetic field in four major InAs in-plane directions of [100], [110], [110], and [110], respectively. For the Fe islands on the MgO layers grown at 200
C, the hysteresis loops are similar regardless of the applied field directions with nonsquare hysteresis loops. On the other hand, the Fe layers on the MgO layers grown at 300
C unambiguously illustrate the strong cubic magnetic anisotropy with the easy and hard axes along InAsÆ110æ and Æ100æ directions, respectively. These directions correspond to the Fe having magnetic easy and

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hard axes along FeÆ100æ and FeÆ110æ directions because the epitaxial relationship of the structure is Fe[010]//MgO[110]//InAs[110]. In order to verify the Fe morphology-dependent magnetic anisotropy, the in-plane angular-dependent Mr/Ms ratio is measured as illustrated in the insets of Figure 6 in which all data points are taken at 9
intervals; 0
and 45
correspond to the applied magnetic field being swept along the InAs[110] and InAs[100] directions, respectively. It should be pointed out that the Mr/Ms ratio of the Fe islands is constant in the structure of the Fe islands, which is clear evidence of no magnetic anisotropy. In contrast, the 2D Fe layer shows noticeably large variation in the Mr/Ms ratio with the in-plane angle measured. The large fluctuation of the Mr/Ms ratio along the Fe[100]//InAs[110] (0
) and Fe[110]//InAs[100] (90
) indicates the cubic anisotropy of the Fe layer. One possible reason for no anisotropy in the Fe islands is that each Fe island may simultaneously have the epitaxial relationship of either Fe[010]//MgO[110]//InAs[110] or Fe[110]//MgO[110]//InAs[110] as mentioned earlier. When the magnetic field is applied along the InAs[110] direction, some Fe islands are magnetized along the easy axis (Fe[010]), while others are magnetized along the hard axis (Fe[110]), therefore resulting in randomization effect of overall magnetocrystalline anisotropy. Also, since the irregularly shaped islands as seen in the plane-view TEM image of Figure 2b are oriented in random directions, the shape magnetic anisotropy of each island would be in different directions, so there would be no net shape anisotropy. Furthermore, since small magnetic particles usually tend to have single domain magnetic structure, each Fe islands of ∼50 nm in diameter shown in Figure 2b could contain magnetic single domains.29 Then the magnetization process of the Fe islands would proceed via magnetization rotation of each Fe island rather than other mechanism such as domain wall motion.30 This is evidenced by the large coercivity of the Fe islands compared with that of the Fe layers as shown in Figure 6. Magnetic domain imaging techniques with high spatial resolution would be needed to confirm this. The results presented in this study lead to two important and interrelated observations when the growth temperature of the MgO layer is lowered in the Fe/MgO/InAs structures. First, the tensile strain in the MgO layer increases due to pseudomorphic growth. Second, the enhanced underlying misfit strain gives rise to the formation of the 3D Fe islands with two different in-plane textures, resulting in no in-plane magnetic anisotropy. The formation of the 3D Fe islands is evidently contrasted with the 2D Fe layer in the structures of Fe/MgO/GaAs of similar Fe and MgO growth temperatures.12 It is well-known that the reason is the considerably smaller lattice mismatch in the layer stack due to the use of GaAs substrates. Since the growth temperature of the Fe remains constant at 200
C for all of the samples in this study, it is certain that the thermally assisted surface diffusion of Fe adatoms does not much work on the morphology change of the Fe. Instead, our previous work manifested that the formation of the 3D Fe islands is closely associated with the underlying misfit strain.31 As the lattice mismatch between MgO and SC substrates increases by use of the substrates of the larger lattice constants, the underlying misfit strain becomes larger so that the formation of 3D Fe islands is favorable for strain energy reduction. The MgO growth temperature employed in this study is a consequential experimental parameter controlling the underlying misfit strain by the partial strain relaxation of the MgO layers. The misfit strain between the MgO and InAs substrates is noticeably large at 4.2% or
7.7% assuming when the MgO lattice is fully strain-relaxed with the lattice registry of 2InAs/3MgO 2894

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Crystal Growth & Design or 3InAs/4MgO. Therefore, the morphology of Fe is subject to change from the 3D islands on the MgO layers grown below 200
C to the 2D layer above 300
C. Additional evidence supporting the fact that the underlying misfit strain is the direct cause of the Fe morphology change is the size and aspect ratio change of the 3D Fe islands with respect to the MgO growth temperature between RT and 200
C. The size and aspect ratio of the 3D Fe islands become larger as the MgO growth temperature increases from RT to 200
C as described before. Comparable morphological evolution has been extensively examined in the growth of semiconductor self-assembled quantum dots (SAQDs). For example, InAs SAQDs of ∼7.1% misfit strain grown on GaAs substrates display much smaller base diameter and larger aspect ratio than Ge SAQDs of ∼4.2% misfit strain grown on Si substrates.32,33 Additionally, the GexSi1
x alloy SAQDs become 2D layers when grown on Si substrates at lower Ge concentration (x).34 Because the misfit strain can be relieved via the 3D island formation to reduce the total strain energy, it is a well-known that the 3D islands with smaller size and higher aspect ratio relieve strain more effectively. Therefore, our experimental observation evidently reveals that the Fe morphological evolution directly results from the underlying misfit strain in the MgO layers grown on InAs substrates at different temperatures. All the explanations above are understandable based on the fact that the MgO layers are epitaxially grown at all the growth temperatures used in this study. It has been reported that Mg and O molecules have enough surface diffusion mobility to form long-range order at extremely low growth temperatures.20,35 Although some polycrystalline components in the MgO are sparsely observed in zoom-in HRTEM images of the Fe/ MgO/InAs structure with the MgO layer grown at RT, the effect on the microstructural evolution of the Fe morphology might be negligibly small. From the results observed in this study, the MgO growth temperature is optimized between 300 and 450
C for successful spin-FET applications using InAs QW structures. In the future, studies with respect to the Fe growth temperature and the thicknesses of Fe and MgO layers will reveal more useful information about the effect of Fe microstrucutral evolution on underlying misfit strain dependent surface diffusion.

4. CONCLUSIONS We have investigated the microstructural evolution and the corresponding magnetic properties of the Fe/MgO/InAs structures with respect to the MgO growth temperature. As the MgO growth temperature increases, tensile-strained MgO layer and the subsequent Fe become gradually relaxed. The 3D Fe islands form on the MgO layers grown below 200
C because of the large underlying misfit strain originating from the pseudomorphic MgO growth. On the other hand, the 2D Fe layers grow on the strain-relaxed MgO layers grown above 300
C with the epitaxial relationship of Fe[010]//MgO[110]//InAs[110]. It is found that there is a clear contrast in in-plane magnetic anisotropy between 3D Fe islands and 2D Fe layers. The Fe islands present no magnetic anisotropy due to the two perpendicular inplane textures along Æ100æ and Æ110æ directions, while the Fe layer shows cubic magnetic anisotropy. ’ ASSOCIATED CONTENT

bS

Supporting Information. Model calculation details based on density functional theory. This material is available free of charge via the Internet at http://pubs.acs.org.

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’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected]. Telephone: þ82-2-958-5736. Fax: þ82-2-958-6851.

’ ACKNOWLEDGMENT This work was supported by the Korea Institute of Science and Technology (KIST) Institutional Program (Contract Nos. 2E21672 and 2E21580), the Degree and Research Center (DRC) Program funded by the Korea Research Council of Fundamental Science and Technology (KRCF), the National Research Foundation of Korea (NRF) grant funded by the Korea government (MEST) (No. 2010-0017457), and the IT R&D program of MKE/KEIT (2009-F-004-01, Technology Development of 30 nm level High Density Perpendicular STTMRAM). The calculations were performed on the KIST grand supercomputer. ’ REFERENCES (1) Hanbicki, A. T.; Jonker, B. T.; Itskos, G.; Kioseoglou, G.; Petrou, A. Appl. Phys. Lett. 2002, 80, 1240. (2) Jiang, X.; Wang, R.; Shelby, R. M.; Macfarlane, R. M.; Bank, S. R.; Harris, J. S.; Parkin, S. S. P. Phys. Rev. Lett. 2005, 94, No. 056601. (3) Zhu, H. J.; Ramsteiner, M.; Kostial, H.; Wassermeier, M.; Sch€onherr, H.-P.; Ploog, K. H. Phys. Rev. Lett. 2001, 87, No. 016601. (4) Koo, H. C.; Yi, H. J.; Ko, J. B.; Chang, J. Y.; Han, S. H.; Jung, D. H.; Huh, S. G.; Eom, J. H. Appl. Phys. Lett. 2007, 90, No. 022101. (5) Hammar, P. R.; Bennett, B. R.; Yang, M. J.; Johnson, M. Phys. Rev. Lett. 1999, 83, 203. (6) Wang, R.; Jiang, X.; Shelby, R. M.; Macfarlane, R. M.; Parkin, S. S. P.; Bank, S. R.; Harris, J. S. Appl. Phys. Lett. 2005, 86, No. 052901. (7) Chromik, S.; Spankova, M.; Vavra, I.; Liday, J.; Vogrincic, P.; Lobotka, P. Appl. Surf. Sci. 2008, 254, 3635. (8) Rashba, E. I. Phys. Rev. B 2000, 62, R16267. (9) Butler, W. H.; Zhang, X.-G.; Schulthess, T. C. Phys. Rev. B 2001, 63, No. 054416. (10) Mathon, J.; Umerski, A. Phys. Rev. B 2006, 74, No. 140404. (11) Manago, T.; Sinsarp, A.; Akinaga, H. J. Appl. Phys. 2007, 102, No. 083914. (12) Martínez Boubeta, C.; Navarro, E.; Cebollada, A.; Briones, F.; Peiro, F.; Cornet, A. J. Cryst. Growth 2001, 226, 223. (13) Miao, G. X.; Chang, J. Y.; van Veenhuizen, M. J.; Thiel, K.; Seibt, M.; Eilers, G.; M€unzenberg, M.; Moodera, J. S. Appl. Phys. Lett. 2008, 93, No. 142511. (14) Jeon, K. R.; Park, C. Y.; Shin, S. C. Cryst. Growth Des. 2010, 10, 1346. (15) Han, W.; Zhou, Y.; Wang, Y.; Li, Y.; Wong, J. J. I.; Pi, K.; Swartz, A. G.; McCreary, K. M.; Xiu, F.; Wang, K. L.; Zou, J.; Kawakami, R. K. J. Cryst. Growth 2009, 312, 44. (16) Koo, H. C.; Kwon, J. H.; Eom, J. H.; Chang, J. Y.; Han, S. H.; Johnson, M. Science 2009, 325, 1515. (17) Kim, K. H.; Kim, H. J.; Koo, H. C.; Chang, J. Y.; Han, S. H. Appl. Phys. Lett. 2010, 97, No. 012504. (18) Tarsa, E. J.; De Graef, M.; Clarke, D. R.; Gossard, A. C.; Speck, J. S. J. Appl. Phys. 1993, 73, 3276. (19) Hung, L. S.; Zheng, L. R.; Blanton, T. N. Appl. Phys. Lett. 1992, 60, 3129. (20) Hsu, W.-Y.; Raj, R. Appl. Phys. Lett. 1992, 60, 3105. (21) Ferhat, M.; Yoh, K. Appl. Phys. Lett. 2007, 90, No. 112501. (22) Sagisaka, H.; Kitada, T.; Shimomura, S.; Hiyamizu, S.; Watanabe, I.; Matsui, T.; Minura, T. J. Vac. Sci. Technol. B 2006, 24, 2668. (23) Wedler, G.; Schneider, C. M.; Trampert, A.; Koch, R. Phys. Rev. Lett. 2004, 93, No. 236101. 2895

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