Evolution of Surface and Interface Structures in Molecular-Beam

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Evolution of Surface and Interface Structures in MolecularBeam Epiaxy of MoSe on GaAs(111)A and (111)B 2

Akihiro Ohtake, and Yoshiki Sakuma Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.6b01605 • Publication Date (Web): 13 Dec 2016 Downloaded from http://pubs.acs.org on December 19, 2016

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Crystal Growth & Design

Evolution of Surface and Interface Structures in Molecular-Beam Epiaxy of MoSe2 on GaAs(111)A and (111)B Akihiro Ohtake∗ and Yoshiki Sakuma National Institute for Materials Science (NIMS), Tsukuba 305-0044, Japan E-mail: [email protected] Phone: +81 (0)29 860 4198. Fax: +81 (0)29 860 4753

∗

To whom correspondence should be addressed

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Abstract We have systematically studied the atomistic growth processes of monolayer MoSe2 on GaAs(111)A and (111)B substrates. A combination of complementary techniques of reflection high-energy electron diffraction, scanning tunneling microscopy, and xray photoelectron spectroscopy allows us to study the evolution processes of surface and interface structures during the MoSe2 growth. Highly-oriented MoSe2 films are epitaxially grown in two steps: the amorphous Mo islands are initially formed on Setreated GaAs surfaces, which is followed by the crystallization into MoSe2 under the supply of a Se molecular beam. While the initial Mo deposition leads to the segregation of Se atoms from the Se-treated GaAs surface, the subsequent supply of the Se beam induces the reconstruction of the Se-terminated GaAs structure beneath the MoSe2 film.

Keywords Transition metal dichalcogenides, Molecular-beam epitaxy, Scanning tunneling microscopy, Surface Reconstructions

1.Introduction Transition metal dichalcogenides (TMDs) have attracted considerable interest, because of their functional properties complementary to graphene. 1–3 While bulk TMDs with the general formula MX2 (M=Mo, W; X=S, Se, Te) are indirect band-gap semiconductors, monolayer TMDs have direct band gaps, making these materials suitable for optoelectronic applications. 4–9 The bulk TMDs have a layered structure: covalently bonded MX2 layers, consisting of three two-dimensional (2D) atomic layers in a sequence of X-M-X, are held together by weak van der Waals (vdW) forces. Thus, as with grapheme, flakes of thin TMDs can be easily exfoliated from bulk, and, so far, numerous experimental studies have been performed 2


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on such exfoliated flakes. 1,2 However, this method has the difficulty in controlling the thickness, size, and uniformity in the TMDs flakes, that prevents large-scale device fabrication. Thus, alternative approach is required for large-scale synthesis of TMDs films with monolayer precision. Molecular-beam epitaxy (MBE) 6,10–15 and chemical vapor deposition 16–18 have recently received increased attention for the growth of TMDs. In particular, MBE has been expected to produce highly-oriented TMDs monolayer over the entire wafer, because of its superiority in controlling film thickness and the availability of in-situ characterization during the growth. This paper reports on a systematic study of the MBE growth of monolayer MoSe2 on GaAs(111)A and (111)B substrates. Early studies 19,20 have shown that MoSe2 films are epitaxially grown on GaAs(111)A and (111)B substrates despite a large lattice mismatch of -17.4%. It has been proposed that the termination of the GaAs surface by the S 19 and Se 20 atoms weaken the interaction between the film and the substrate, allowing for the vdW epitaxy of MoSe2 . However, atomistic processes at surfaces and interfaces in MoSe2 heteroepitaxy are far from being completely understood. Here, we present a detailed description of the growth process of MoSe2 , especially focusing on the evolution of the atomic structures at the surface and the interface. For this purpose, MoSe2 films were grown in two basic steps: first, amorphous Mo islands were formed on the Se-treated GaAs surface, which were, then, crystallized into epitaxial MoSe2 films by supplying a Se molecular beam. Complementary experimental techniques of reflection high-energy electron (RHEED) diffraction, scanning tunneling microscopy (STM), and x-ray photoelectron spectroscopy (XPS) have been used to study the detailed growth processes at each step. We show that the initial Mo deposition is accompanied by the segregation of Se atoms from the substrates, and that the subsequent supply of the Se beam allows the Se atoms to intrude into the MoSe2 /GaAs interface, resulting in the reconstruction of the Se-terminated GaAs structure beneath the MoSe2 film.

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2.Experiments The experiments were performed using a multi-chamber ultra-high vacuum (UHV) system consisting of MBE chambers for the III-V growth and the Se treatments and an electron-beam (EB) evaporation chamber for the Mo deposition. The system is equipped with the STM and XPS apparatuses for on-line characterization . 21,22 The clean surfaces of GaAs(111)A and (111)B were prepared by MBE. 23 While only a (2×2) reconstruction is observed on √ √ the (111)A surface under the conventional MBE condition, Ga-rich ( 19 × 19)R23.4◦ and As-rich (2×2) reconstructions are observed on the (111)B surface depending on the preparation conditions. 24 We found that high density of one bilayer-height steps was formed √ √ on the ( 19 × 19)R23.4◦ surface after the Se treatment (Supporting Information Figure S1). Thus, in the following the results on the As-rich (2×2) surfaces will be shown for the (111)B orientation. Clean GaAs substrates were transferred to another UHV chamber for the Se treatments via UHV transfer modules (< 2×10−10 Torr). The GaAs surfaces are exposed to the Se beam at 300◦ C, and then annealed at 530◦ C under the Se flux. The samples were then transferred to the EB evaporation chamber for the deposition of Mo atoms. Mo atoms were deposited at room temperature (RT) at a rate of 0.3∼0.4 monolayer (ML)/min. Here, 1 ML is defined as 1.06×1015 atoms/cm2 , which corresponds to the number density of the MoSe2 (0001) unit cells. The growth rate was measured by quartz-crystal microbalance, and was calibrated by cross-sectional transmission electron microscopy observations for amorphous Mo films. The pressure during the Mo deposition was in the range of 8×10−10 Torr∼2×10−9 Torr. After the Mo deposition, the samples were exposed to the Se beams at RT and were annealed at 600-630◦ C to form MoSe2 . The beam-equivalent pressure of Se was 7∼9×10−8 Torr for both the Se treatments of the initial surfaces and the selenidation of Mo films. All the STM images were collected at RT in the constant current mode with a tunneling current of 0.1 nA and a sample voltage of -3 V. XPS measurements were performed using monochromatic Al Kα radiation (1486.6 eV). Photoelectrons were detected at an angle of 4


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35◦ from the surface. The Se 3d and Mo 3d spectra were fitted using a Voigt function. A peak separation of 0.85 eV (3.2 eV) is assumed for the 5/2 and 3/2 spin-orbit components of Se 3d (Mo 3d) and the branching ratio is kept at ∼1.5.

3.Results and discussion Figure 1a shows a series of RHEED patterns during the formation of MoSe2 films on GaAs(111)A-(2×2) substrates. Well-ordered (2×2) reconstruction of the (111)A surface (I) became disordered after the Se treatment (II). Figure 2a shows the corresponding STM image of the Se-treated GaAs(111)A surface: high densities (8×1012 ∼1×1013 cm−2 ) of small islands are observed. In the magnified image (Figure 2b), bright features are arranged with a (2×2) periodicity. This is in marked contrast with the observation of (1×1) RHEED patterns on S-terminated GaAs(111)A. 25,26 On the other hand, a (2×2) reconstruction has been observed for S on InAs(111)A. 27,28 While a possible structure model has been proposed for S-terminated InAs(111)A, 28 disordered surface features and the existence of small islands make it difficult to determine the detailed atomic structure of the Se-treated GaAs(111)A surface. As shown in Figure 1b-II, the Se-treated (111)B surface shows a (1×1) RHEED pattern, similarly to the case for the S-terminated (111)B surfaces of GaAs 26 and InAs. 27 Previous studies have shown that the S-terminated GaAs(111)B surface has a simple structure, in which the S atom replaces the first-layer As atom. 26,29 However, high-resolution STM image (Figure 2d) shows highly disordered array of bright features with a density of 0.75 per (2×2) unit. This is incompatible with the simple structure model proposed for the S-terminated GaAs(111)B surface. We note that the (1×1) structures having S-Ga and Se-Ga surface bilayers 26,29 do not satisfy the electron-counting requirement. 30 The deposition of Mo (nominal coverage of 1 ML) at RT increased the background intensity in RHEED patterns and weakened the reflections from GaAs substrates (Figure 1a-III

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and 1b-III). After being completely covered by amorphous Se at RT, the sample was heated under the Se molecular-beam. The broad and weak streaks appeared at 500◦ C, and became sharp and intense as the substrate temperature is increased as shown in Figure 1a-IV and 1b-IV. The spacings of streaks in Figure 1a-IV and 1b-IV are 21% larger than those of GaAs in both [101] and [121] directions, which correspond to the in-plane lattice constant of 0.329 nm. Thus, it turns out that the crystalline MoSe2 films were epitaxially grown with the epitaxial relationship of (0001)MoSe2 //(111)GaAs and [1120]MoSe2 //[110]GaAs, similar to the case for the simultaneous deposition of Mo and Se. 19 MoSe2 films on both substrates show symmetric RHEED patterns along the [101] direction, in contrast to those of GaAs substrates (Figure 1a-I and 1b-I), indicating the coexistence of domains rotated by 180◦ with respect to each other. It is also suggested that MoSe2 films consist of rather small domains, because elongated streaks are observed in Figure 1a-IV and 1b-IV instead of spots on the zeroth-order Laue zone. Figure 3a and 3b show filled-state STM images of MoSe2 films (nominal coverage of 1 ML) on GaAs(111)A and (111)B substrates, respectively. The MoSe2 growth on both substrates proceeds with the nucleation of 2D islands of irregular shape: the edges of islands are not aligned in a specific crystallographic direction. In the magnified image (inset of Figure 3a), bright features show hexagonal close packed structure, which corresponds to the Moiré pattern. The spacing of bright features is 0.18∼0.19 nm, in good agreement with the value calculated from the in-plane lattice constants of GaAs (0.3997 nm) and MoSe2 (0.3299 nm). As shown in Figure 3c and 3d, the height of islands is 0.7∼0.8 nm, which is slightly larger than the interlayer spacing of MoSe2 (0.647 nm). The size of MoSe2 islands increases with selenidation temperature from 600 to 630◦ C, but beyond which the surface morphology significantly degrades, because of the roughening of the GaAs substrates, as we will show below. Small regions located at a lower level than the surrounding area exist in both images.

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In particular, for (111)B, the islands having different terminating heights are connected, as indicated by arrow heads in Figure 3b. As shown in Figure 3c and 3d, the height difference is close to the step height on GaAs surface (0.326 nm), and much smaller than the interlayer spacing of MoSe2 . Thus, it is likely that such a height difference is not originated from the coexistence of 1 ML- and 2 ML-thick MoSe2 islands, and that lower islands are located at the 2D holes of one BL depth on the GaAs substrate. Since such small holes are hardly observed on the initial and Se-treated GaAs surfaces, it is suggested that the crystallization processes of MoSe2 at high temperatures causes the height undulation. We confirmed that the selenidation at temperatures higher than 630◦ C usually results in a highly multilayered morphology of the GaAs substrate. There also exist second-layer MoSe2 islands (arrows in Figure 3b) and gaps between the first-layer 2D islands, which might be ascribed to the insufficient surface diffusion due to the kinetics of growth processes. It is also possible that the actual MoSe2 coverage is less than 1 ML. Although the gap between the 2D islands becomes narrower as the film thickness is increased, small holes are still observed in the first layer even at 1.5 ML (Supporting Information Figure S2). Carefully comparing surface morphologies on (111)A (Figure 3a) and (111)B (Figure 3b), we found that MoSe2 islands on (111)B form more continuous films. To obtain details on such a polar surface dependence, now we focus on the changes in surface morphologies of Mo films induced by the selenidation. Figure 4 compares the STM images before and after the selenidation of Mo films. Mo atoms form small islands on the Se treated surfaces. The height and diameter of the islands at 0.2 ML are typically 0.3∼0.8 nm and 1∼3 nm, respectively, for both surface orientations. The size and density of the islands are increased with Mo coverage, which is naturally accompanied by the decrease in the RHEED intensity from the GaAs substrates (inset of Figure 4). While the Mo islands are crystallized into MoSe2 islands by the irradiation of the Se molecular beam at a high temperature of 620◦ C on both substrates, the surface morphologies

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are strongly dependent on the surface polarity of the substrates. For the (111)B surface, small Mo islands are evolved into larger 2D MoSe2 islands after the selenidation, irrespective of the Mo coverage. This is in accordance with the appearance of rather sharp reflections from MoSe2 islands in the RHEED patterns. On the other hand, much smaller MoSe2 islands are formed on the (111)A surface at Mo coverages of 0.2 ML and 0.5 ML, in good agreement with the broad MoSe2 reflections in the RHEED patterns, and are coalesced into larger islands at 1 ML coverage. Thus, it is plausible to consider that the surface diffusion of MoSe2 is less enhanced on the (111)A surface, which could be ascribed to the difference in the surface structures of Se-treated (111)A and (111)B surfaces. XPS measurements were carried out to study the electronic states before and after the selenidation of Mo films. Figure 5a and 5b show the high-resolution Se 3d and Mo 3d spectra for (111)A and (111)B, respectively. The Se 3d spectrum for the Se-treated (111)A surface (Figure 5a-I) is composed of two components denoted Se1 (54.7.eV) and Se2 (53.9 eV), while that for the (111)B consists of a single component Se1 (55.2 eV). The Se1 and Se2 peaks have been assigned to Se atoms at the surface layer and those at the subsurface layers. 31 The intensity ratio of Se1 components of (111)A to (111)B is 1:0.7, roughly proportional to the density of the bright spots observed in the STM images of the Se-treated GaAs surfaces (Figure 2b and 2d). After the Mo deposition, Se1 intensities for both surfaces are significantly decreased (Figure 5a-II and 5b-II). Simultaneously, a new component Se3 appeared at a lower binding energy of 54.5 eV for (111)A and 54.4 eV for (111)B. Their peak positions are close to the values of bulk 2H-MoSe2 32 and monolayer-MoSe2 . 12 From these results, it is suggested that the majority of Se atoms were taken away from the GaAs substrate to form covalent bonds with Mo atoms after the Mo deposition. On the other hand, as shown in Figure 5a-IV and 5b-IV, the Mo 3d spectra for the as-deposited Mo films are located at 228.2 eV and 228.0 eV for (111)A and (111)B, respectively. These values are in good agreement with that reported for metallic Mo (227.8 eV). 33 Thus, we conclude that Mo islands have metallic properties, and

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Mo atoms at the bottom of the islands are bonded to Se atoms. An important implication of the present results is that the Se-treated surface is not effectively passivated, and is unstable against the formation of Mo-Se bonds. This motivated us to check whether the initial Setreatments of GaAs surfaces are necessary for the formation of MoSe2 films. As shown in Figure 3e and 3f, we confirmed that MoSe2 films having morphologies comparable to those in Figure 3a and 3b are formed on the GaAs(111)A and (111)B surface without the initial Se treatments. Here we note that the Se2 component for (111)A completely disappeared after the Mo deposition (Figure 5a-II and 5b-II). This means that the Mo islands attract Se atoms from deeper layers of the GaAs(111)A substrate, but such an interpretation is somewhat awkward. On the other hand, as shown in Figure 2a, our STM observations have revealed the formation of high densities of small islands (typical diameter and height are 2 nm and 0.4 nm, respectively) on the Se-treated (111)A surface, and such islands are hardly observed for the (111)B surface (Figure 2c). If we assume that small islands consist of Se atoms, the total amount of Se in the islands could be estimated to be 0.3±0.1 ML in coverage, which is enough to be detected by XPS. Since the Se2 component were not detected from the Setreated (111)B surface (Figure 5b-I), it appears likely that the Se2 components comes from Se islands. After the Mo films were exposed to the Se beam, Mo 3d peaks are shifted to higher binding energies, so that their peak positions coincidence with those of bulk 2H-MoSe2 32 and monolayer- MoSe2 12 (Figure 5a-V and 5b-V), while leaving the Se 3d peak positions almost unchanged (Figure 5a-III and 5b-III). Another interesting finding is that the Se1 components for the Se-treated surface reappeared after the selenidation. Thus, it turns out that the supplied Se molecules are consumed to the formation of MoSe2 and to the rearrangement of the Se-terminated GaAs surface beneath the growing MoSe2 film, giving rise to the weakening of the interaction at the MoSe2 /GaAs interface (pseudo vdW interaction). It is also likely that the interfaces are similarly formed for the MoSe2 growth without the

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initial Se treatments.

4.Conclusions We have studied the growth processes of MoSe2 on the GaAs(111)A and (111)B-(2×2) surfaces. Metallic Mo islands are first formed on the Se-treated surfaces, which are crystallized into highly-oriented 2D MoSe2 islands by supplying the Se molecular beam. The MoSe2 islands on the (111)B surface are larger and more continuous than those on the (111)A surface. The initial surfaces are not effectively passivated by Se atoms: the deposition of Mo atoms leads to the segregation of Se atoms from Se-treated GaAs substrates, resulting in the formation of Mo-Se bonds. However, when the Mo films are exposed to the Se beam, the atomic rearrangement occurs to form Se-terminated GaAs structure beneath the MoSe2 film. The present findings lead us to a more complete understanding of the growth processes of TMDs, and are expected to stimulate further studies to realize large-scale synthesis of monolayer-TMDs for future applications.

Acknowledgement Helpful discussion with Dr. T. Mano is gratefully acknowledged.

Supporting Information Available √ √ STM images of Se-treated GaAs(111)A-(2×2), (111)B-( 19× 19)R23.4◦ , and (111)B-(2×2) surfaces; STM images of 1.5 ML- MoSe2 on GaAs(111)A and (111)B substrates. This material is available free of charge via the Internet at http://pubs.acs.org/.

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