Large scale 2D/3D hybrids based on gallium nitride and transition metal dichalcogenides

Abstract

Two and three-dimensional (2D/3D) hybrid materials have the potential to advance communication and sensing technologies by enabling new or improved device functionality. To date, most 2D/3D hybrid devices utilize mechanical exfoliation or post-synthesis transfer, which can be fundamentally different from directly synthesized layers that are compatible with large scale industrial needs. Therefore, understanding the process/property relationship of synthetic heterostructures is priority for industrially relevant material architectures. Here we demonstrate the scalable synthesis of molybdenum disulfide (MoS₂) and tungsten diselenide (WSe₂) via metal organic chemical vapor deposition (MOCVD) on gallium nitride (GaN), and elucidate the structure, chemistry, and vertical transport properties of the 2D/3D hybrid. We find that the 2D layer thickness and transition metal dichalcogenide (TMD) choice plays an important role in the transport properties of the hybrid structure, where monolayer TMDs exhibit direct tunneling through the layer, while transport in few layer TMDs on GaN is dominated by p–n diode behavior and varies with the 2D/3D hybrid structure. Kelvin probe force microscopy (KPFM), low energy electron microscopy (LEEM) and X-ray photoelectron spectroscopy (XPS) reveal a strong intrinsic dipole and charge transfer between n-MoS₂ and p-GaN, leading to a degraded interface and high p-type leakage current. Finally, we demonstrate integration of heterogeneous 2D layer stacks of MoS₂/WSe₂ on GaN with atomically sharp interface. Monolayer MoS₂/WSe₂/n-GaN stacks lead to near Ohmic transport due to the tunneling and non-degenerated doping, while few layer stacking is Schottky barrier dominated.

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Introduction

Beyond two dimensional (2D) heterostructures,^1–4 2D/3D hybrids enable mixed dimensional van der Waals heterostructures⁵ and provide unique opportunities towards steep slope transistors,⁶ highly efficient p–n junction solar cells,⁷ light emitters⁸ and photovoltaics.⁹ Transition metal dichalcogenide (TMD) and nitride semiconducting heterostructures are especially interesting because of the wide variety of band alignments, and the possibility to convert the nitride to a 2D form.¹⁰ Recently, p–n diodes¹¹ and room temperature Esaki diodes¹² were demonstrated with transferred 2D layers. However, transfer-induced contamination and damage often degrades the performance of the devices.¹³ Additionally, the device size is limited by the exfoliated sample size (typically < 10μm), which is not compatible to the industrial process. Our previous work using powder vaporization¹⁴ demonstrates that monolayer MoS₂ can be epitaxially grown on GaN and is electrically active in the vertical direction. However, the non-uniformity still limits its applicability to large area synthetic 2D/3D heterostructures.^15,16

In this work, we utilize metal organic chemical vapor deposition (MOCVD) to synthesize mono to few-layer MoS₂ and WSe₂ on p-type and n-type GaN, respectively, to probe the 2D/3D electrical properties. MOCVD enables large area, uniform TMDs films via layer-by-layer growth as identified by atomic force microscopy (AFM) and Raman spectroscopy, where the layer number is tuned via growth time. X-ray photoelectron spectroscopy (XPS) characterization confirms nearly stoichiometric MoS₂ (Mo : S = 1 : 1.95) and WSe₂ (W : Se = 1 : 2.03) layers are grown with negligible metal–oxide bonding compared to powder vaporized TMDs.^14,17 Furthermore, there is no structural change of the GaN substrate. Importantly, we demonstrate that the vertical transport in the 2D/3D hybrids vary with 2D thickness and the choice of the heterostructure. Monolayer TMDs on GaN exhibit clear direct tunneling, while few layer (FL) (>2 layers) TMDs lead to p–n junction behavior. Although p–n diodes based on FL MoS₂/p-GaN and WSe₂/n-GaN exhibit similar turn-on voltages under forward bias, FL MoS₂/p-GaN exhibits 100× higher current under reversed bias due to strong charge transfer and an intrinsic dipole between the MoS₂/p-GaN interface. As a result, we hypothesize that few layer WSe₂/n-GaN is the most appropriate choice for high quality synthetic 2D/3D p–n heterostructures. Finally, we demonstrate MoS₂/WSe₂/n-GaN hybrids with atomically sharp interfaces. Unlike the resonant tunneling observed on MoS₂/WSe₂/epitaxial graphene (EG) heterostructures,¹⁸ single layer MoS₂/WSe₂ on n-GaN exhibits Ohmic behavior, while FL MoS₂/WSe₂ on n-GaN is dominated by Schottky barrier transport.

Results and discussion

Transition metal dichalcogenide films are directly deposited on (p- and n-)GaN/c-sapphire substrates via MOCVD. The hole concentration for p-type GaN is 10¹⁷ cm⁻³ and the electron concentration for n-type GaN is and 10¹⁸ cm⁻³. The MoS₂ films are grown in a hot wall reactor at 650 °C using molybdenum hexacarbonyl (Mo(CO)₆) and diethyl sulfide (DES) precursors with NaCl powder¹⁶ placed upstream from the reactor hot zone for nucleation control and carbon sequestration (Fig. S1†). The WSe₂ films are prepared using a cold wall reactor at 650 °C with tungsten hexacarbonyl (W(CO)₆) and hydrogen selenide (H₂Se) (Fig. S2†). Layer-by-layer growth has previously been demonstrated via MOCVD,^15,16 and here we tune the layer thickness via growth time. Atomic force microscopy (Fig. 1a–c) demonstrates controlled formation of 1L, 2L, and FL (3.5 nm) MoS₂ with increasing growth time from 30 min to 60 min to 120 min, and the topography of monolayer MoS₂ (Fig. 1a) closely matches the roughness of the GaN substrate.¹⁴ An additional adhesion map (inset of Fig. 1a) is provided to visualize the distinction between 1L MoS₂ and p-GaN. Monolayer, 2L and FL (4.7 nm) WSe₂ (Fig. 1d–f, respectively) is realized using growth times of 5 min, 10 min and 30 min. Raman spectra (see ESI† for measurement details) of MoS₂ (Fig. 1g) exhibits the typical 384 cm⁻¹ (E_2g) and 403 cm⁻¹ (A_1g) vibration modes for 1L MoS₂, 384 cm⁻¹ (E_2g) and 405 cm⁻¹ (A_1g) for 2L MoS₂, and 383 cm⁻¹ (E_2g) and 407 cm⁻¹ (A_1g) for FL MoS₂, confirming the thickness difference.¹⁹ A similar Raman shift trend is also observed on WSe₂ (Fig. 1h) for E_2g peak (250 cm⁻¹ for 1L and 249 cm⁻¹ for 6L), agreeing with previous reports.²⁰

Fig. 1 Thickness controlled growth of TMDs: a–c. AFM image of 1L–FL MoS₂ grown on p-GaN. To identify the thickness, the MoS₂ film is intentionally scratched. The height profile in b and c clearly show 2L (1.3 nm) and ∼5L (3.5 nm) of MoS₂ on p-GaN. It is difficult to visualize the height monolayer in a. due to the roughness of p-GaN, but the adhesion map (inset of a) demonstrates the sub-monolayer flakes on p-GaN; d–f. AFM image of 1L–FL WSe₂ grown on n-GaN. The height profile indicates the different thicknesses of WSe₂; g. Raman spectra for 1L–FL MoS₂. A_1g peak exhibits a significant red shift as thickness decreases, while E_2g peak remains nearly the same, confirming the thickness reduction. h. Raman spectra of 1L–FL WSe₂. The observed E_2g peak shows a slight red shift as thickness increases. i. XPS spectrum of Mo 3d core level, clear Mo–S bonding is detected and the Mo–O bonding is below the detection limit. j. XPS spectrum of W 4f core level, W–Se bonding is identified and the W 5p is also observed. k. Valence band edge of the FL MoS₂/p-GaN (blue) and FL WSe₂/n-GaN (red). The Fermi level is 1.6 eV and 0.4 eV above the valence band maxima of MoS₂ and WSe₂, respectively, indicating that the MoS₂ is n-type and WSe₂ is p-type.

Synthetic MoS₂ and WSe₂ is observed to be nearly stoichiometric (Mo : S = 1 : 1.95; W : Se = 1 : 2.03) with negligible metal–oxygen bonding and no structural degradation. In the case of MoS₂, the molybdenum 3d doublet peaks are observed at 230.0 eV and 233.1 eV together with a sulfur 2s peak at 227.2 eV (Fig. 1i). Following peak fitting, the only detected bonding is between Mo and S. Importantly, Mo–O bonding is significantly reduced for MoS₂ grown by MOCVD compared to PV.¹⁴ This is likely the result of the elimination of the oxygen precursor, MoO₃, often used as the Mo source in the PV method. X-ray photoelectron spectroscopy of WSe₂ (Fig. 1j) exhibits three W peaks at 32.3 eV (W 4f_7/2), 34.4 eV (W 4f_5/2) and 37.7 eV (W 5p_3/2), indicating no W–O bonding or other impurities within the detection limits of XPS.²¹ Finally, valence band edge measurements (Fig. 1k) of FL MoS₂ and FL WSe₂ indicate the valence band maximum (VBM) of MoS₂ is 1.6 eV below the Fermi level, suggesting that MoS₂ is electron doped (n-type). In contrast, the VBM of WSe₂ is 0.4 eV below the Fermi level, indicating hole doping (p-type) behavior of WSe₂.²² This is typical of MoS₂ and WSe₂, and is attributed to native defects in the TMD layer.^16,23,24 Further XPS characterization of the GaN before and after growth also indicates no detectable structural degradation of the GaN after the TMD growth (Fig. S3†).

Vertical charge transport through the 2D/3D hybrid is highly sensitive to the 2D layer properties. Here, we use conductive AFM (CAFM) with a Pt/Ir tip to probe the nanoscale vertical transport of p–n junction TMD/GaN heterostructures. Based on fitting the local current–voltage (I–V) characteristics with the Fowler–Nordheim tunneling equations,^25,26 the carrier transport can be described to be dominated by direct tunneling through the TMDs for 1L MoS₂/p-GaN and 1L WSe₂/n-GaN (Fig. S4†). Thus, monolayer TMDs are considered nearly electrically transparent due to the full depletion region overlap²⁷ (see ESI† for depletion width calculation). Additionally, the TMD layer between the Pt/Ir metal tip and GaN surface leads to a reduction in the Schottky barrier due to Fermi level pinning at the Tip/TMD interface, thereby reducing the contact resistance.^14,27,28 As the 2D layer thickness increases, the carrier transport becomes dominated by the barrier at the 2D/GaN interface (Fig. S4†). This is evident when analyzing the I–V characteristics of FL n-MoS₂/p-GaN and p-WSe₂/n-GaN (Fig. 2a). For simplicity, positive bias is applied on the p-type semiconductor (p-GaN or WSe₂). The two heterostructures show slightly different turn-on voltage at approximately 1 V for WSe₂/n-GaN and 1.4 V for MoS₂/p-GaN under forward bias, with an ideality factor (n) of 30 for MoS₂/p-GaN and 15 for WSe₂/n-GaN_. The high ideality factor is likely due to the Schottky barrier between the tip and TMDs,^14,29 leading to a combination of p–n and Schottky barrier transport. Improved Ohmic contacts to the TMDs will enable lower ideality factors. We note that under reverse bias, the MoS₂/p-GaN structure exhibits high current, even at low voltages. To understand this behavior, Kelvin Probe Force Microscopy (KPFM) and Low Energy Electron Microscopy/Reflectivity (LEEM/LEER) measurements are employed (see Fig. S5† for KPFM calibration). KPFM measurements of FL MoS₂ (Fig. 2b) and FL WSe₂ (Fig. 2c) reveal distinct work function differences between the TMDs and (p-, n-)GaN. The work function of WSe₂ is measured to be ∼300 meV higher than n-GaN, and the work function of MoS₂ is ∼150 meV higher than p-GaN. This work function difference is confirmed by LEEM/LEER measurements, where multiple spots are evaluated on FL MoS₂/p-GaN (LEEM, Fig. 2d) and FL WSe₂/n-GaN (LEEM, Fig. 2e). The corresponding LEER spectra (Fig. 2f) suggest that the work function difference between FL MoS₂ and p-GaN is ∼220 meV and ∼340 meV between FL WSe₂ and n-GaN. Based on the measured work function (Φ_WSe2 = 4.8 eV; Φ_GaN = 4.5 eV), XPS measurements³⁰ (see ESI† for details) and reported electron affinity (χ_WSe2 = 3.9 eV (ref. 31); χ_GaN = 4.1 eV (ref. 14)), the band diagram of WSe₂/n-GaN can be established (Fig. 2g), which agrees with literature reported staggered band alignment for WSe₂/GaN.³⁰ However, counter to the reported values,^13,32 KPFM and LEER in this work both demonstrate that the work function of MoS₂ is higher than that of p-GaN (Fig. 2b and d). Note that the GaN surface is exposed by scratching the continuous TMD, and it results in some roughness on the GaN. We hypothesize that this could be due to modification of the GaN Mg dopant properties. First, hydrogen present during the growth process may passivate the Mg dopant, leading to a reduced hole concentration of p-GaN.^33–36 Second, the strong intrinsic dipole and charge transfer at the semiconductor/semiconductor interface may be responsible for the observation.^37,38 In either case, this would shift the Fermi level (and hence work function) toward higher values in the GaN bandgap, leading to a modification of the MoS2/p-GaN behavior. Additional XPS measurements (Fig. S6†) reveal the Type I band alignment between MoS₂ and p-GaN, agreeing with the recent study (Fig. 2h).³⁷ Combining this information with the aforementioned KPFM and LEEM/LEER measurement, the band bending under different biases is shown in Fig. 2h. Conventional electron conduction is responsible for the forward biased current, but unlike WSe₂, the hole conduction due to the tunneling can introduce reversed bias current.³⁹ Furthermore, material modifications (below the detection limits of XPS) in the GaN due to high temperature synthesis of the MoS₂ could also be the source of high leakage currents in reverse bias observed in Fig. 2a,⁴⁰ however, additional theoretical work is needed to understand the aforementioned scenario in depth. This suggests that the utilization of n-MoS₂/p-GaN could suffer from performance and reliability challenges due to the required elevated temperature of the MoS₂ growth, thereby pointing to WSe₂/n-GaN the preferred choice for 2D/3D hybrids for p–n junction applications.

Fig. 2 The electronic properties of TMD/GaN heterostructures: a. I–V measurement of FL MoS₂/p-GaN and WSe₂/n-GaN. The turn on voltage for FL MoS₂/p-GaN and FL WSe₂ is 1.4 V and 1.0 V under forward bias, respectively. However, FL MoS₂/p-GaN exhibits a higher leakage current; b–c. KPFM characterization of FL MoS₂/p-GaN (b) and FL WSe₂/n-GaN (c), the work function of WSe₂ and MoS₂ is 300 mV and 150 mV higher than n-GaN and p-GaN respectively; d–e. LEEM images of FL MoS₂/p-GaN acquired at 1.9 eV (d) and FL WSe₂/n-GaN acquired at 2.3 eV (e). The scratch of GaN is clearly visible by LEE. The LEER analysis points are labeled with different colors in Fig. 3d and e; f. the LEER spectra of FL MoS₂/p-GaN (left) and FL WSe₂/n-GaN (right) corresponding to points labeled in (d) and (e), which quantifies the electrostatic surface variation and hence the variation of vacuum level; g. the band alignment of WSe₂/n-GaN; h. the band alignment of MoS₂/p-GaN. It is clear that hole conduction can be responsible for the reversed bias current.

Utilizing the same synthesis techniques, it is possible to realize more complex 2D/3D hybrids such as MoS₂/WSe₂/n-GaN (MWnG). Here, WSe₂ films are deposited on n-GaN, and subsequently MoS₂ films are deposited on WSe₂/n-GaN to form the MWnG hybrid. The MoS₂ is grown at 650 °C, resulting in nanocrystalline sub-monolayers on 1L WSe₂ (Fig. 3a) and uniform FL MoS₂ on FL WSe₂ (Fig. 3b). Raman spectroscopy (Fig. 3c) confirms the presence of MoS₂ and WSe₂, with no evidence of intermediated phases (WS₂, MoSe₂, etc.). Cross-sectional high resolution transmission electron microscopy (HRTEM) and energy dispersive spectroscopy (EDS) (Fig. 3d and e) reveal that the FL hybrid structure is 6–7 layers, evenly split between MoS₂ and WSe₂. The EDS mapping (Fig. 3e) indicates that the layers of MoS₂ and WSe₂ are well-defined with a sharp interface, and no evidence of alloying. The electrical properties (characterized by CAFM, Fig. 3f) are shown to vary greatly between monolayer heterostructures and FL heterostructures. When the film contains FL MoS₂ and FL WSe₂, a rectifying conductivity is observed. This rectification is attributed to the forward biased Schottky barrier between the Pt/Ir tip and MoS₂.^14,41 Interestingly, unlike the reported resonant tunneling of 1L MoS₂/1L WSe₂ heterostructures on graphene, we find nearly Ohmic behavior of 1L MoS₂/1L WSe₂ on n-GaN, with no negative differential resistance (NDR) at room temperature. Such behavior is likely due to the thermionic emission at room temperature and non-degenerate doping concentration of our MoS₂ and WSe₂ films.^18,42

Fig. 3 2D/2D/3D heterostructure of TMDs/GaN: a–b. AFM images of 1L MoS₂/1L WSe₂/n-GaN (a) and FL MoS₂/FL WSe₂/n-GaN (b). Sub-monolayer domains can be seen on the top of WSe₂ for 1L MoS₂/1L WSe₂/n-GaN. The surface is uniform for FL heterostructures with a total thickness of ∼4.6 nm. The Raman spectra (c) confirms the co-existence of MoS₂ and WSe₂; d–e the HRTEM image of the FL MoS₂/WSe₂/n-GaN heterostructure (d), ∼7 atomic layers is clearly observed. HRSTEM image colored by EDS mapping (e) shows a sharp interface between MoS₂/WSe₂; f. the I–V curves of 1L MoS₂/WSe₂/n-GaN and FL MoS₂/WSe₂/n-GaN. Unlike the Schottky-dominated rectifying characteristic shown on FL sample, 1L sample exhibits nearly linear tunneling characteristic without negative differential resistance.

Conclusions

This work demonstrates the scalable growth of TMD/GaN hybrid structures, where film thickness is controlled by tuning the growth time. The process results in high quality MoS₂ and WSe₂, and enables a route to 2D heterostructures on GaN for next generation electronics. Transport vertically through the hybrid structures is dominated by the 2D layer thickness and materials choice, where monolayer films exhibit strong direct tunneling characteristics and few layer 2D enables p–n junction formation at the 2D/3D interface. Compared to FL WSe₂, FL MoS₂ films exhibit high leakage current in reversed bias. Our KPFM and LEEM/LEER characterization reveals a counter-intuitive work function difference between the MoS₂ and p-GaN, potentially due to strong interface charge transfer and possible non-structural degradation (Mg passivation) in p-GaN. Furthermore, FL MoS₂/FL WSe₂ heterostructures on n-GaN exhibits rectifying behavior due to the presence of the metal/MoS₂ Schottky barrier. In contrast, we find Ohmic behavior when the thickness of the film is reduced to one atomic layer. This study elucidates that the thickness and materials choice is critical towards developing high quality 2D/3D heterostructures.

Author contributions

J. A. R and K. Z conceived the idea, and J. A. R., T. G. I., R. M. F., and M. H directed the research; K. Z. and B. J synthesized the 2D/3D hybrids; K. Z. carried out AFM, Raman, and electrical measurements; N. C. B. conducted XPS characterization and K. Z. analyzed XPS data and proposed the band alignment; J. L. performed LEEM/LEER measurements; R. A. B. and D. R. prepared GaN sample and analyzed electrical data; B. W. carried out TEM sample preparation and characterization; J. O. L. performed TOF-SIMS characterization. All authors discussed the results. K. Z. and J. A. R., wrote the paper with significant input from J. L., R. M. F. and R. A. B. All authors have read and approved the manuscript.

Conflicts of interest

Authors declare no conflicts of interest.

Acknowledgements

Research was conducted as part of the Center for Atomically Thin Multifunctional Coatings (ATOMIC), sponsored by the National Science Foundation (NSF) division of Industrial, Innovation & Partnership (IIP) under award #1540018. The work at Carnegie Mellon University is supported by the center for Low Energy Systems Technology (LEAST), one of the six STARnet centers, sponsored by MARCO and DARPA.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c7nr07586c

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