Highly Oriented Atomically Thin Ambipolar MoSe2 Grown by Molecular Beam Epitaxy.

Transition metal dichalcogenides (TMDCs), together with other two-dimensional (2D) materials, have attracted great interest due to the unique optical and electrical properties of atomically thin layers. In order to fulfill their potential, developing large-area growth and understanding the properties of TMDCs have become crucial. Here, we have used molecular beam epitaxy (MBE) to grow atomically thin MoSe2 on GaAs(111)B. No intermediate compounds were detected at the interface of as-grown films. Careful optimization of the growth temperature can result in the growth of highly aligned films with only two possible crystalline orientations due to broken inversion symmetry. As-grown films can be transferred onto insulating substrates, allowing their optical and electrical properties to be probed. By using polymer electrolyte gating, we have achieved ambipolar transport in MBE-grown MoSe2. The temperature-dependent transport characteristics can be explained by the 2D variable-range hopping (2D-VRH) model, indicating that the transport is strongly limited by the disorder in the film.

Atomically thin two-dimensional (2D) materials have shown great potential because of their interesting electrical and optical properties.[1−4] Potential applications in flexible electronics and the possibility to further extend their range of applications by integrating them into heterostructures[5] motivate scientists to develop a reliable way to grow large-area 2D materials.[6−9] Whereas chemical vapor deposition (CVD) can on one hand result in high-quality 2D materials and 2D heterostructures with sharp and clean interfaces,[10] the toxicity of some of the precursors introduces challenges, and the lack of in situ characterization and monitoring during growth can lead to nonreproducible results. On the other hand, using molecular beam epitaxy (MBE) to grow atomically thin transition metal dichalcogenides (TMDCs) and other 2D materials via van der Waals epitaxy[11−16] has several potential advantages. The deposition can be well controlled using high-purity elemental sources; direct heterostructure growth can be achieved, and the growth in an ultrahigh vacuum environment limits the amount of impurity atoms. In addition, the quality of the MBE-grown film can be monitored during growth in situ using reflection high-energy electron diffraction (RHEED). The growth of different TMDCs on various substrates has been recently demonstrated,[17−19] and the band structures and thin film morphology have been intensively investigated using angle-resolved photoemission spectroscopy and scanning tunneling microscopy.[20−22] However, reports on the electrical properties of atomically thin layers grown by MBE[23] are rare and indicate that the material quality needs to be improved further. Starting from this point, it is crucial to optimize epitaxial growth and to investigate the optical and electrical properties of MBE-grown atomically thin TMDCs.

Results

Here, we report on the use of MBE to grow atomically thin MoSe2 on GaAs(111)B down to nominal monolayer (ML) thickness, exhibiting high optical quality confirmed by photoluminescence (PL) and Raman spectroscopy. We find that using GaAs(111)B as the growth substrate results in a high degree of control over lattice orientation of the domains forming the polycrystalline film. The features of van der Waals epitaxy were thoroughly investigated by X-ray photoelectron spectroscopy (XPS). Using scanning transmission electron microscopy (STEM), we have observed that depending on the growth temperature, polycrystalline MoSe2 films can be composed of grains with either a 15° or a 60° misorientation angle. Both configurations are predicted to be energetically stable by density functional theory (DFT) calculations. We further address the electrical properties of highly oriented nominal ML MoSe2 and demonstrate ambipolar transport behavior using polymer electrolyte gating. Temperature-dependent measurements indicate that the transport is limited by the disorder and can be explained using the 2D variable-range hopping (2D-VRH) model.

Atomically thin MoSe2 films were grown on GaAs(111)B substrates by MBE, and the growth was monitored in situ using a RHEED camera. Atomic hydrogen was used to effectively remove native oxide at low temperature without As desorption,[24,25] and a clean GaAs(111)B surface with sharp RHEED streaks along the GaAs[11-2] azimuth was observed (see Figure a). As the growth progresses, GaAs-related streaks gradually fade out and are replaced by a new set of faint streaks with a smaller spacing, indicating the formation of MoSe2. A RHEED pattern from the MoSe2 film with a half-layer coverage is shown on Figure b, with GaAs and MoSe2 streaks having similar intensities. GaAs-related streaks completely vanish as a full film is grown, while MoSe2 streaks become more intense Figure c. The sharp streaks in Figure c,d indicate the epitaxy with MoSe2[10-10]//GaAs[11-2] and MoSe2[11-20]//GaAs[10-1], respectively. The van der Waals epitaxy thus provides a reliable way for developing large-scale 2D materials on different substrates.[19,26]

Figure 1

Growth of atomically thin MoSe2 by MBE. (a) RHEED patterns of GaAs(111) along the GaAs[11–2] azimuth at growth start. (b) Half–half transition along the GaAs[11–2] azimuth. (c,d) Nominal monolayer (ML) MoSe2 observed along MoSe2[10–10] and MoSe2[11–20] azimuths at growth end after 22 min. (e) Mo 3d and (f) Se 3d core-level spectra in XPS. (g) Comparison of Raman spectra from MBE-grown and exfoliated ML MoSe2. (h) Photoluminescence of transferred ML MoSe2 and exfoliated ML MoSe2. Inset shows optical image of the transferred film. The sample was grown at 470 °C.

Films grown on GaAs(111)B were examined ex situ using X-ray photoelectron spectroscopy (XPS). Figure e shows the core-level spectrum of the Mo 3d range. The binding energies of Mo 3d5/2 and Mo 3d3/2 peaks are 229.2 and 232.4 eV, respectively. An additional AsLMM peak with a lower binding energy of 227.2 eV is also observed. The Se 3d core-level spectrum in Figure f shows an Se 3d5/2 peak at 54.9 eV and an Se 3d3/2 peak at 55.8 eV, demonstrating the existence of Mo–Se bonds. A consistent binding energy shift is observed in the core-level spectra of Ga 3d and As 3d between pristine GaAs substrate and as-grown MoSe2, implying that charge transfer takes place at the interface (see Supplementary Section 1 for more details). On the other hand, the peak positions and the high quality of fitting show that no intermediate compounds exist at the interface, as expected from van der Waals epitaxy.

The Raman spectrum of nominal ML MoSe2/GaAs(111) shown on Figure g clearly shows the MoSe2 A1g mode at 238.7 cm–1, which is comparable with that of exfoliated ML MoSe2 on 270 nm SiO2. Photoluminescence spectroscopy was also used to investigate the optical properties of ML MoSe2. In order to avoid the optical quenching effect from the substrate,[18,27] the as-grown film was transferred onto a 270 nm SiO2/Si chip, with the optical image shown in the inset of Figure h and Figure S2. The 1.58 eV peak recorded at room temperature is clearly shown in Figure h and is comparable with the exfoliated ML flake in terms of the peak positon. Despite inevitable existence of wrinkles and folded regions due to the imperfect transfer process which might introduce defects and peak broadening, the full width at half-maximum (fwhm) of 37 meV is comparable to that of exfoliated ML flakes (fwhm = 27 meV).

The films were transferred onto the TEM grids and investigated using Cs-corrected STEM in order to examine their morphology. Figure a shows a low-magnification high-angle annular dark-field (HAADF)-STEM image of MoSe2 grown at the temperature of 470 °C. The patch-like islands with a typical size of a few nm with different brightness represent regions with varying thicknesses. Surprisingly, the corresponding fast Fourier transform (FFT) image calculated from this image and shown in the inset of Figure a shows only two sets of spots with a six-fold symmetry, rotated by ∼15° with respect to each other. The high-magnification HAADF-STEM image in Figure b shows corresponding domains with a relative rotation of 15°, labeled ML(A) and ML(B). These are sometimes also stacked on top of each other, forming bilayers (BLs). A more detailed analysis of the grain orientation is presented on Figure S3 and shows that both orientations appear with roughly the same frequency in the ML film. Most of the BL area (∼70%) is, however, of the ML(A) + ML(B) type, composed of two MoSe2 layers with an interlayer twist of 15° and could present an interesting material system for studying the effect of interlayer twist on electrical and optical properties of 2D semiconductors.[28]

Figure 2

Morphology of MBE-grown MoSe2 films. (a) Low-magnification HAADF-STEM image of MoSe2 grown at 470 °C. Inset is the corresponding FFT image showing two sets of spots. (b) High-magnification image of an incomplete bilayer (BL) composed of ML domains with two orientations denoted as ML(A) and ML(B). (c) Low-magnification HAADF-STEM image of MoSe2 grown at 530 °C. Inset shows the corresponding FFT image with a single set of diffraction patterns. (d) High-magnification image of the ML region with a schematic of atom positions. Bright spots correspond to double Se atoms. (e) Intensity histogram of the image shown in (c). (f) Intensity profile along the slice shown in (c). (g) Intensity profile of the slice from (d).

We have gained further insight into the stability of both types of MoSe2/GaAs(111)B superlattices by considering several models and determining their formation energies through DFT calculations (see Supplementary Section 3). The picture that emerges suggests that a twisting angle not only dramatically reduces the strain in the MoSe2 lattice but also leads to a slightly more favorable interaction with the GaAs(111)B substrate, irrespectively of the relative MoSe2–GaAs orientation. Overall, formation energies of oriented and misoriented MoSe2/GaAs(111)B superlattices are very similar, indicating that both cases are likely to form. Nevertheless, it is also worth mentioning that the growth dynamics should also play a role in the migration of adatoms prior to in-plane bond formation, which could influence the growth, resulting in large regions with single orientation. Possible changes in the surface reconstruction at these temperatures close to GaAs decomposition are not taken into account either, which could explain the difference in the morphology of films grown at these two temperatures.

A slight increase of growth temperature, into the 500–530 °C range, results in increased order in the film, with the grains no longer showing the 15° misorientation. The HAADF-STEM image of the film is shown on Figure c. The corresponding FFT image now shows only one set of peaks. Since ML MoSe2 does not possess inversion symmetry, we cannot exclude the presence of grains with a 60° relative orientation at this point as these would result in a set of diffraction peaks at the same positions in Fourier space. The presence of voids in the film indicates that BLs start to form before the first MLs complete. Similar morphologies of MBE-grown MoSe2 on highly oriented pyrolytic graphite, graphene, and SiC(0001) were also observed by STM by other groups.[18,21,27,29,30]

Because the intensities recorded in HAADF-STEM images are related to Rutherford scattering, which increases with the atomic number (Z),[31,32] different intensities in the image can be attributed to different layer thickness. Figure b shows the intensity histogram with dashed lines that correspond to intensities in ML and BL regions. The intensity profile along the slice of interest is plotted in Figure c, showing the thickness varying from ML to BL. Figure d shows a high-magnification image of a ML region with the intensity profile along a slice of interest shown in Figure f, where the positions of Mo and Se atoms can be assigned to periodically varying intensities. The ML has a 2H structure with a lattice constant estimated to be 3.29 ± 0.03 nm, which is in line with the bulk lattice constants reported in literature.[33,34]

To further confirm the long-range uniformity of the MoSe2 film, we have measured second harmonic generation (SHG) from as-grown MoSe2/GaAs(111)B and suspended MoSe2 on TEM grids (Figure S7). Polar plots of SHG intensity show six-fold symmetry, whereas the PL intensity and Raman A1g peak position and intensity maps indicate a high degree of uniformity (see Supplementary Section 4).

We now focus on the electrical properties of 0 and 60° MoSe2 with nominal ML thickness. We use electric double-layer transistors (EDLT) in order to access a wide range of electrostatically induced doping levels and to reduce Schottky barrier heights at contacts, allowing efficient electron and hole injection using the same contact material.[35]Figure shows the schematic of the device. The polymer electrolyte PS–PMMA-PS:[EMIM]-[TFSI] is spin-coated on top. It allows us to reach high carrier densities and increases the efficiency of carrier injection, making it ideal for achieving ambipolar regime of operation.[35,36] We also include a back gate, allowing charge carrier modulation in the semiconducting film at temperatures below the freezing point of the polymer electrolyte (T ∼ 200 K). Devices without a polymer electrolyte show a very poor transistor behavior (see Supplementary Section 5). The dependence of the EDLT on the polymer electrolyte voltage VPE exhibits a clear ambipolar behavior close to room temperature (T = 280 K) (Figure b). We find a current Ion/Ioff ratio of ∼104 and ∼102 for the n and p sides, respectively. The maximum current density on the n side, with the value of ∼1 μA/μm, is 2 orders of magnitude larger than that on the p side, possibly due to intervalley scattering of electrons, while the off current remains at pA levels. The subthreshold swing calculated from the linear region is ∼390 mV/dec for both sides. Field-effect mobilities can be extracted from four-contact devices by freezing the polymer electrolyte at 200 K and performing a back-gate sweep (see Supplementary Section 6), which allows the charge carrier concentration in the 2D semiconductor to be modulated around the value set by VPE prior to freezing the electrolyte. The extracted electron mobility μe is ∼0.05 cm2 V–1 s–1, and hole mobility μh is ∼0.28 cm2 V–1 s–1. The mobility values are significantly lower than those of CVD-grown MoSe2.[37,38] These results indicate that charge carrier transport is strongly influenced by the disorder in the film.[23,39]

Figure 3

Ambipolar transport in MoSe2 EDLT. (a) Schematic of the MoSe2 EDLT in a dual-gate geometry. The back gate is used to modulate the charge density in the 2D channel around the value set by the reference electrode with the applied voltage VPE. (b) Channel current as a function of VPE showing ambipolar behavior.

In order to elucidate the dominant transport mechanism in MBE-grown atomically thin MoSe2, we have performed electrical measurements as a function of temperature, polymer electrolyte, and back-gate voltage. Figure a shows the sheet conductivity Gsh as a function of VPE based on the four-contact device shown in the inset. The VPE applied for the p side needs to be pushed to VPE < −4 V to reach the same value as on Gsh of the n side, indicating strong electron doping or Fermi level pinning to the conduction band. A drop in Gsh takes place at VPE > 3 V and can be attributed to the electrolyte-induced disorder which is commonly observed in experiments involving EDLTs.[40,41] The sweep at 280 K can then provide a reference curve for different doping levels that can be achieved by changing VPE. Once the doping level at a given VPE is stabilized at 280 K, we freeze the electrolyte down to 200 K with a cool-down rate of 1 K/min. The VPE is then disconnected at 200 K after the electrolyte was completely frozen so that Gsh can be stabilized (see Supplementary Section 7). The Gsh during each cool down was recorded down to 12 K with a cooling rate of 0.5 °C/min. The process was reversible after a mild annealing by ramping the temperature to 333 K.

Figure 4

Two-dimensional VRH transport mechanism in MoSe2 EDLT. (a) Sheet conductivity Gsh as a function of VPE at 280 K. Inset: Optical image of the four-contact device. (b,c) Gsh as a function of T–1/3 on the hole and electron sides for different values of VPE filled lines corresponds to fits to the VRH model. (d) Evolution of the localization length ξloc with VPE, extracted from fits to the VRH model on the hole branch. (e) Dependence of ξloc on VPE for the electron branch.

The Gsh of MoSe2 at a given VPE monotonically decreases with decreasing temperature, showing semiconducting behavior for both sides. The dependence above 80 K follows the 2D-VRH model[42] with the relation Gsh ∝ exp[−(T0/T)1/3], where T0 is the characteristic temperature. A linear fit of ln Gsh to T–1/3 is plotted at each given VPE, demonstrating the validity of the 2D-VRH mechanism for both sides (Figure b,c). The VRH mechanism is usually observed in disorded systems,[23,40] and the results imply that the transport of MoSe2 EDLT is strongly influenced by the voids and the nanometer-scale grains in the film.

Charge carriers are strongly scattered and tend to hop between different conductive paths. The localization length can be changed in a small range by changing VPE, that is, changing the doping level. The dependency is evident by extracting T0 from the fits to the 2D-VRH model with the value of slopes s extracted from Figure b,c, where T0= s3. The values of T0 decrease by more than 1 order of magnitude with the increase of |VPE| because of the increase of carrier densities in the material, thus screening the disorder along the conductive paths. The localization length ξloc can be extracted using the expression , where kB is the Boltzmann constant and D the density of states (see Supplementary Section 8). The results are plotted in Figure d,e for p and n sides, respectively. Holes have a slightly lower ξloc, with ξloc reaching a maximum value of ∼4 nm at VPE = −4.2 V. The n side, on the other hand, shows tunable ξloc up to ∼9 nm with VPE = 2.45 V. All values of ξloc have the same order of magnitude as the grain size shown in STEM images, indicating that the 2D-VRH transport is linked to the disorder in the as-grown MoSe2. Future MBE-based growth efforts will have to concentrate on increasing the grain size in order to improve the film quality.

Conclusion

In conclusion, we have grown atomically thin MoSe2 using MBE. Films show a high degree of alignment due to the van der Waals interaction with the GaAs substrate and can be transferred to insulating substrates for further optical and electrical transport studies. We realize electrolytically gated transistors based on transferred ML MoSe2 films. Electrical transport follows the 2D-VRH model due to disorder in the film, with localization length comparable to the grain size.

Methods

MBE Setup and Material Growth

The growth was carried out in an Omicron MBE (Lab 10) with a ∼10–10 mbar base pressure. Cleaved 1 × 1 cm2 GaAs(111)B substrates were outgassed at up to 500 °C for at least 30 min. The native oxide was removed from the surface of GaAs(111)B by heating it to 350 °C under a flux of atomic hydrogen. Hydrogen molecules were dissociated by a tungsten filament with Joule heating at 70 W and were introduced into the chamber via a leak valve. The procedure lasted 30 min or more at base pressure of ∼3 × 10–7 mbar, resulting in sharp streaks in RHEED. A Kundsen cell and an electron beam source (EFM-3 from Omicron) were used for Se and Mo evaporation, respectively. The flux rates were calibrated using a quartz crystal microbalance, and the flux ratio of Se/Mo was optimized to be ∼40 for growth. A RHEED camera (Staib Co.) was used to monitor the growth in situ. The growth temperature was optimized in the 470–530 °C temperature range. Post-annealing at up to 550 °C was performed in Se atmosphere. Higher temperature leads to GaAs decomposition and increased surface roughness.

XPS and Raman Spectroscopy

The XPS spectra were obtained ex situ in a commercial KRATOS AXIS ULTRA system, and a C 1s core-level peak at 284.8 eV was used for the reference. Peak identification and fitting were performed in PHI MultiPak processing software. Raman analysis was performed using a Horiba LabRAM HR800 system using a 532 nm wavelength green laser with spot size ∼4 μm. The laser power was kept below 4 mW during all measurements. We used an 1800 lines/mm grating and have calibrated the system using the polycrystalline Si peak at 520 cm–1. The PL was measured in a home-built setup using a 488 nm laser (Coherent) for excitation.

STEM Microscopy and Analysis

The STEM experiments were performed on a FEI Titan Themis 300 double Cs-corrected microscope at an acceleration voltage of 80 kV in order to minimize beam damage. The scanning probe had a 28 mrad semiconvergence angle, resulting in a resolution close to 1 Å. The data were acquired under annular dark-field conditions using an annular detector with a collection half-angle between 40 and 200 mrad.

Material Transfer and Device Fabrication

As-grown films were coated with PMMA and immersed into 30% KOH(aq) at 90 °C. The detached PMMA layer with the as-grown film was then transferred to a beaker with deionized water several times to remove excess KOH(aq) and was transferred onto a degenerately doped n++ Si chip covered by 270 nm SiO2 or TEM grids. PMMA was removed in acetone. Palladium was used for electrical contacts in a standard PMMA-based e-beam lithography process. A second e-beam lithography was performed, followed by O2/SF6 plasma etching in order to define the device geometry. To fabricate an EDLT, the PS–PMMA-PS:[EMIM]-[TFSI] electrolyte was spin-coated onto the device and soft-baked at 60 °C for 10 min. A more detailed description is available elsewhere.[43,44] Electrical measurements were carried out using an Agilent 5270B SMU and Keithley 2000 DMM. Cryogenic measurements were performed in a Janis closed-cycle cryogen-free cryostat.