Sublimation-based wafer-scale monolayer WS₂ formation via self-limited thinning of few-layer WS₂

Abstract

Atomically-thin monolayer WS₂ is a promising channel material for next-generation Moore's nanoelectronics owing to its high theoretical room temperature electron mobility and immunity to short channel effect. The high photoluminescence (PL) quantum yield of the monolayer WS₂ also makes it highly promising for future high-performance optoelectronics. However, the difficulty in strictly growing monolayer WS₂, due to its non-self-limiting growth mechanism, may hinder its industrial development because of the uncontrollable growth kinetics in attaining the high uniformity in thickness and property on the wafer-scale. In this study, we report a scalable process to achieve a 4 inch wafer-scale fully-covered strictly monolayer WS₂ by applying the in situ self-limited thinning of multilayer WS₂ formed by sulfurization of WO_x films. Through a pulsed supply of sulfur precursor vapor under a continuous H₂ flow, the self-limited thinning process can effectively trim down the overgrown multilayer WS₂ to the monolayer limit without damaging the remaining bottom WS₂ monolayer. Density functional theory (DFT) calculations reveal that the self-limited thinning arises from the thermodynamic instability of the WS₂ top layers as opposed to a stable bottom monolayer WS₂ on sapphire above a vacuum sublimation temperature of WS₂. The self-limited thinning approach overcomes the intrinsic limitation of conventional vapor-based growth methods in preventing the 2^nd layer WS₂ domain nucleation/growth. It also offers additional advantages, such as scalability, simplicity, and possibility for batch processing, thus opening up a new avenue to develop a manufacturing-viable growth technology for the preparation of a strictly-monolayer WS₂ on the wafer-scale.

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New concepts

In this manuscript, a sublimation-assisted and self-limited thinning approach for wafer-scale transition metal dichalcogenides (TMDCs) is demonstrated for the first time. Tungsten dichalcogenide (WS₂) – a promising TMDC-based channel material for next-generation Moore's nanoelectronics is selected as a typical example. Based on sulfurization with pulsed sulfur gas supply and constant H₂/Ar gas flow, the thinning effect coming from the volatile reduced WS_2−x phase by H₂ under a temperature substantially above the sublimation temperature of WS₂ can become self-limited. Importantly, the self-limited nature of the thinning process enables the damage-free and strictly monolayer TMDCs formation. This new thinning process simplifies the current challenges in preparing the high-quality monolayer TMDCs, where an inevitable partial oxidation and physical damages associated from conventional etching processes severely degrade the film stoichiometry and uniformity. Notably, the universality of the self-limited thinning effect is also manifested on a broader number of different TMDCs (e.g. MoS₂) and TMDCs prepared under various growth approaches (e.g. PVD sputtered WS₂ and MoS₂) and doped TMDCs system (e.g. Co-doped WS₂ prepared by sulfurization). The findings indicate that this thinning method is broadly applicable to pure and doped TMDCs systems, which will accelerate the research of TMDCs with high PL quantum yield and adjustable band gap for large-area optoelectronics and spin-electronics.

Introduction

As the feature size of silicon-based transistors is approaching their physical limit, there have been intensive efforts in searching for alternative channel materials for future ultra-scaled electronic devices.^1,2 The unique thickness-dependent properties of semiconducting 2D layered transition metal dichalcogenides (TMDCs), including WS₂, MoS₂, WSe₂, and MoSe₂, make them very attractive for applications in future electronics, optoelectronics, and smart sensors.^3–5 The lower effective mass of WS₂ among other TMDC semiconductors provides it with the highest theoretical room temperature electron mobility.^6–8 In addition, its immunity to the short channel effect, transferability for vertical stacking, and high PL quantum yield make the monolayer WS₂ as a promising channel material for sub-3 nm ultra-scaled field effect transistors (FETs) and for high-performance optoelectronics.^9,10

Over the last 5 and more years, there has been significant progress in the effort for growing large-area atomically thin 2D TMDCs, including WS₂. Among the various growth methods developed, including chemical vapor deposition (CVD), metal organic chemical vapor deposition (MOCVD), physical vapor deposition (PVD), and molecular beam epitaxy (MBE),^11–13 CVD and MOCVD are considered as the two most promising front-runners in terms of the quality of the grown TMDCs layer and the wafer-size for uniform growth. Both are able to grow monolayer TMDCs on the wafer-scale (up to 4 inch wafer for CVD and up to 8 inch wafer for MOCVD).^14,15 However, both CVD and MOCVD also have their own shortcomings for adoption in scalable mass-production. For CVD, where metal oxide (e.g., WO₃, MoO₃etc.) powders are typically used as metal precursors, it is difficult to further scale up the wafer-size due to its so-called distance-dependent growth mechanism.^16,17 In comparison, MOCVD offers a good control over the area uniformity, and therefore greatly improved scalability.¹⁸ However, an un-acceptably long growth time (more than 26 h) is required to achieve monolayer growth.¹⁹

Other than those limits associated with the individual growth methods, a more general issue with almost all TMDCs growth arises from the difficulty in controlling the layer number. Unlike graphene growth, which is governed by a catalytically driven self-limiting growth mechanism, the CVD/MOCVD/PVD/MBE growth of TMDCs is non-self-limiting growth.^20,21 The 2^nd layer nucleation/growth occurs after substantial growth of the 1^st layer domains, typically at the centers of domains, due to competition between lateral vs. vertical growth, or at their grain boundaries after their coalescence.^22–24 This poses a serious challenge for their consistent and robust usage in nanoelectronics since the presence of 2^nd layer domains change the electronic structure in TMDCs (since the electronic structure of the bilayer TMCDs is significantly different from monolayer TMDCs), consequently resulting in variation of the device parameters and affecting circuit yields.^25–27

Alternative approaches based on the post-growth treatments (e.g., low-power oxygen plasma,^28,29 laser thinning method,³⁰ XeF₂ gas etching,³¹etc.) were also used to achieve monolayer TMDCs. Promising thinning effects had been demonstrated by trimming down the few-layer TMDCs designated areas to atomically-thick layers. However, there are also several limitations in these approaches. First, partial oxidation happens on the post-treated TMDCs due to the thinning process being performed under ambient conditions. Such ambient air oxidation would induce uncontrollable doping in TMDCs with O₂, H₂O or N₂, and it is necessary to perform extra sulfurization on post-treated TMDCs to restore their stoichiometry. Second, the etching results are highly sensitive to the exposure time and power of the oxygen plasma, laser beam, and XeF₂ etchant gas. More importantly, these post-etching methods are not self-limiting, and it is thus very challenging to avoid the occurrence of physical damages and complete etching on the thinnest TMDCs areas.^26,31

Here, we present the use of an in situ self-limited thinning of few-layer WS₂ formed by the sulfurization of WO_x (W) for the growth of a wafer-scale fully-covered strictly-monolayer WS₂. This approach plays with the thermodynamic instability of the WS₂ top layers as opposed to the stable bottom monolayer WS₂ on suitable substrates (with strong interface coupling), especially in its reduced state, i.e., WS_2−x, above the vacuum sublimation temperature of WS₂ as revealed in Density functional theory (DFT) simulations. However, we also found that the alternating switching between thinning and non-thinning via a pulsed supply of sulfur precursor vapor in a constant flow of H₂ and Ar gas is essential for process robustness and reliability. Strictly-monolayer high quality WS₂ films have been prepared successfully on 4 inch substrates using this method with excellent uniformity. The grown monolayer WS₂ was further examined by optical microscopy, Raman microscopy, photoluminescence (PL), X-ray photoelectron spectroscopy (XPS) and transmission electron microscopy (TEM) to reveal its high quality. A number of FETs were fabricated on the synthesized monolayer WS₂ to demonstrate its potential application in electronics and optoelectronics. Other than overcoming the major intrinsic limitation of conventional vapor-based growth methods in preventing the 2^nd layer WS₂ domain nucleation and growth, additional potential advantages of the self-limited thinning approach could also include: (1) scalability (growth area is determined by the original sputter-deposited WO_x (or W) film coverage); (2) simplicity in process methodologies (the use of industry-standard sputter-deposition and conventional sulfurization); and (3) possibility for batch processing (several WO_x(W)-deposited wafers can be processed simultaneously in a large growth reactor).

Results and discussion

In our sulfurization and subsequent self-limited thinning for the formation of monolayer WS₂, a pre-deposited WO₃ (thickness of 1–3 nm) film on sapphire is placed in the downstream heating zone of a furnace with an appropriate temperature (e.g., 1000 °C) and pressure (e.g., ∼10 Torr). Diethyl sulfide (DES, (C₂H₅)₂S) (as liquid sulfur precursor) is diluted by carrier gas Ar with a precise control of the concentration through regulating the bubbler and mass flow controller, while a constant flow of Ar + H₂ mixed etching gas is also fed into the reaction chamber, as illustrated in Fig. 1a (see the Method section for the detailed growth conditions). Different from the growth of WS₂via conventional sulfurization of a pre-deposited WO₃ film with a constant sulfur source (CSS) supply, a typical self-limited thinning process for the WS₂ monolayer formation features a pulsed sulfur supply (PSS) under a constant Ar + H₂ mixed gas flow. Fig. 1b details the DES supply and process temperature temporal profiles in a typical PSS WS₂ monolayer formation process cycle, where the DES supply is switched ON and OFF 3 times under 2 stages of high temperature processing (the 1st stage involves a linear temperature ramp-up period (20 min) from 600 °C to 1000 °C and the 2^nd stage involves the annealing at 1000 °C for another 30 min). Fig. 1c shows the optical image of a fully-covered WS₂ film grown on a 4 inch (0001)-Al₂O₃ sapphire wafer by the PSS method. Micro-Raman and PL measurements were performed on different regions of the WS₂ film to assess the thickness (i.e., layer number) and PL property on the corresponding positions and analyze their area uniformity across the 4 inch wafer (Fig. 1d and e). Fig. 1d compares the Raman spectra obtained on different regions on the wafer. No noticeable variations in peak positions and intensities are observed for the two signature Raman peaks, i.e., E_2g (in-plane mode) and A_1g (out-of-plane mode). It is known that both the intensity ratio of A_1g/E_2g and the frequency difference (Δω) between the E_2g and A_1g peaks increase with the increasing number of layers in WS₂ films.¹⁹ Thus, the observation of essentially identical values of the A_1g/E_2g intensity ratio and Δω from the Raman spectra obtained on the different positions clearly indicates excellent thickness uniformity in the grown WS₂ film over the entire 4 inch wafer. The values of the A_1g/E_2g intensity ratio and Δω are ∼0.7 and ∼61.5 cm⁻¹, respectively, and are in good agreement with that reported for the monolayer WS₂,³² as well as that measured on the monolayer WS₂ single crystal (Fig. S1a and b, ESI†),¹¹ suggesting that the grown WS₂ is mostly or even completely covered by a monolayer WS₂ area. The formation of the monolayer WS₂ is further evidenced by the observation of a strong PL peak at ∼1.98 eV with the peak intensity as high as that from the single crystal monolayer WS₂ (Fig. S1c, ESI†). It should be noted that, in WS₂, only the direct bandgap monolayer WS₂ exhibits a strong PL emission (at ∼1.98 eV), as opposed to weak PL emissions from the indirect bandgap bilayer and multilayer WS₂.^33,34 Furthermore, the emission intensity of the PL peak essentially remains the same for the different positions on the wafer, again indicating excellent uniformity of the WS₂ film formed across the wafer.

Fig. 1 Sublimation-based wafer-scale monolayer WS₂ formation via self-limited thinning of few-layer WS₂. (a) A schematic configuration showing the position of the WO₃-coated substrate, gas flow of DES, Ar and H₂ carrier gas inside the vertical furnace tube. (b) Illustrations of the temporal profiles of the process temperature and DES supply for a typical PSS process growth cycle consisting of three DES turn-on pulse cycles. (c) Photograph of the as-grown PSS-WS₂ film on the 4-inch sapphire wafer. (d) Raman and (e) PL spectra collected on 9 different positions of the as-grown PSS-WS₂ film on the 4-inch sapphire wafer. (f) Photographs of the WS₂ sample grown in the midway (top) and completion (bottom) of the PSS process growth cycle. (g) Raman and (h) PL spectra collected on different positions of the WS₂ films, as shown in (f). Schematic drawings showing the growth mechanism of the PSS-WS₂: (i) multi-layer WS₂ growth in the early low-temperature stage, (j) reduction of the top few-layer WS₂, (k) sublimation into the volatile WS_2−x species (thinning effect), (l) sulfur diffusion to restore WS_2−x back to stable WS₂ on sapphire and (m) strictly monolayer WS₂ formation.

As mentioned earlier, the demonstration of the wafer-scale monolayer WS₂ formation shown in Fig. 1c is based on the sublimation-assisted and self-limited thinning process of few-layer WS₂. Indeed, an apparent color difference can be seen from the photographs of WS₂ films grown in the midway (top) and completion (bottom) of the PSS process growth cycle (Fig. 1f) due to the variation of the bandgap in WS₂ of different thicknesses. The large value of Δω (∼64 cm⁻¹) and small PL emission suggested that the WS₂ grown in the midway of the PSS growth cycle (the state after ramping up from 600 °C to 1000 °C in PSS process) was few-layer thick (measured on spot 5 in Fig. 1g and h). Comparatively, a monolayer WS₂ would be formed if its growth cycle went through the entire PSS process, as evidenced by the value of Δω ∼ 61.5 cm⁻¹ (spots 1–4 in Fig. 1g) and strong PL emission at 1.98 eV (spots 1–4 in Fig. 1h). In short, the self-limited thinning PSS process consisted of 5 steps: (1) multi-layer WS₂ growth in the early low-temperature stage, (2) reduction of the top few-layer WS₂, (3) sublimation into the volatile WS_2−x species (thinning process), (4) sulfur diffusion to restore WS_2−x back to stable WS₂ on sapphire, and (5) a formation of stable and strictly monolayer WS₂ (Fig. 1i–m). The detailed growth mechanism will be discussed with XPS, TEM and DFT results in the later sessions. Notably, the self-limited thinning PSS process can also be manifested on PVD-sputtered WS₂ (Fig. S2, ESI†) and MoS₂ (Fig. S3, ESI†), in which the few-layer regions are trimmed down into strictly-monolayer with full surface coverage. In addition, this PSS monolayer WS₂ method can be extended to the large area uniform formation of doped W_1−x(M_x)S₂ monolayer, as revealed by our preliminary work with Co doping (Fig. S4, ESI†). These findings are an embodiment of the universality of the self-limited thinning effect in the PSS process.

The as-grown PSS-WS₂ films were transferred onto a SiO₂/Si substrate for further examination through a conventional van der Waals 2D layer transfer method.³⁵Fig. 2a indicates the optical microscope image of a transferred monolayer WS₂ film on the substrate of SiO₂/Si, while Fig. 2b shows the mapping image of the PL peak intensity over the red dashed box region in Fig. 2a. The homogeneous color contrast shown from the mapping image suggests that the transferred WS₂ film maintains a high degree of integrity with minimum defects. This is further confirmed by AFM measurement (Fig. 2c), where a surface roughness (R_q) as low as 0.212 nm is obtained without revealing any evidence of the presence of bilayer domains. The height profile along the blue-dotted line indicates that the thickness of WS₂ is about 0.65 nm, agreeing well with the reported value (∼0.67 nm) for monolayer WS₂¹³ and thus unambiguously confirming that a strictly monolayer WS₂ is indeed prepared by the PSS sulfurization and subsequent self-limited thinning process. TEM characterization was carried out to reveal the crystalline structure of the transferred monolayer WS₂ (Fig. 2d and e). Excellent uniformity of the monolayer WS₂ is evident from the TEM image (over 60 nm × 60 nm), where neither bilayer regions nor grain boundaries are present. A well-defined atomic arrangement is observable in the High-Resolution TEM image, indicating a high crystallinity of WS₂. The measured d-spacing of the (1010) planes are about 0.27 nm, which coincides well with the reported value for 2H-WS₂.³⁶ To assess its single crystallinity, selected-area electron diffraction (SAED) was performed on four different locations (over μm-scale area) of the monolayer WS₂ (Fig. 2f). As shown in Fig. 2f, each pattern shows one set of six hexagonal dots, indicating high crystal crystallinity, while all of the patterns present an almost identical lattice orientation, suggesting the presence of large grains in μm scale.

Fig. 2 Characterizations of the transferred PSS-WS₂ on the substrates of SiO₂/Si and TEM grid. (a) Optical microscope image, (b) mapping image based on the PL intensity of PSS-WS₂ (enclosed by the red dotted-line in (a)) and (c) plane-view AFM image. The inset in (c) shows the thickness of PSS-WS₂ along the dashed line in the AFM image. TEM images acquired under (d) low-resolution and (e) high-resolution of the monolayer PSS-WS₂. (f) Representative SAED patterns collected from four typical locations of the continuous monolayer WS₂ revealing nearly the same lattice orientation; the yellow dashed line guides the orientation between the six hexagonal dots and horizontal direction.

With the formation of monolayer WS₂ by the self-limited PSS process being confirmed by the experimental results presented above, X-ray photoelectron spectroscopy (XPS) was carried out to analyze the phase evolution in the original WO_x film during/after the CSS and PSS process to gain insight of the monolayer formation mechanism (Fig. 3a). It was found that the sulfurization of the WO_x film at 600 °C during the initial stage of the CSS and PSS process (for 10 min, see Fig. 1b) resulted in a substantial, but incomplete conversion of WO_x to WS₂. This is evident from the observation of three W 4f core-level peaks (i.e., two dominant W 4f_7/2 (∼32.8 eV) and W 4f_5/2 (∼35 eV)) and a faint W 5p_3/2 (∼37.9 eV) (peaks) corresponding to WS₂, together with two weak W 4f peaks (W 4f_7/2 (∼36.1 eV) and W 4f_5/2 (∼38.2 eV)) from WO₃. Cross-section TEM of the 600 °C-sulfurized film indeed confirms the partial formation of few-layer WS₂ (2–4 layers) on top, while a substantial amount of residual WO_x remains beneath the WS₂ formed (Fig. 3c–e). It is seen that the WO_x phase remains even after further sulfurization at a much higher temperature of 1000 °C under a constant DES supply, i.e., CSS process (Fig. 3a), presumably due to the effective inhibition of S penetration by the top WS₂ multilayer.

Fig. 3 Self-limiting thinning mechanism of monolayer WS₂ on the sapphire. (a) W 4f peaks of XPS spectra of WS₂ growth under different conditions: (I) sulfur supply at 600 °C, (II) constant sulfur supply at 1000 °C, (III) pulsed sulfur supply at 1000 °C, (b) Slab structures of monolayer (right) and bilayer (left) WS₂ on Al₂O₃. (c) Cross-sectional STEM image acquired at low magnification at the interface of WS₂/sapphire prepared under sulfur supply at 600 °C. (d)–(f) The corresponding zoomed-in images were captured over three different areas enclosed by the dotted rectangles.

The multilayer nature of CSS-WS₂ is revealed by its much larger A_1g/E_2g ratio of ∼0.6, which agrees with reported results,³⁷ and significantly weaker PL intensity as compared to PSS-WS₂ (Fig. 4a and b), which may also lead to a higher roughness of the CSS-WS₂ surface (Fig. S5, ESI†). An apparent discrepancy in film color can also be observed from the photographs of the CSS-WS₂ and PSS-WS₂ films (Fig. 4c and d) due to the variation of the bandgap in WS₂ of different thicknesses. In contrast, the WO₃ phase is completely absent in the PSS processed sample. Only the three W 4f_7/2 (∼32.8 eV), (W 4f_5/2 (∼35 eV)), and W 5p_3/2 (∼37.9 eV) peaks corresponding to WS₂ phase are observed, in agreement with the confirmation of the formation of the monolayer WS₂ in the PSS sample by the experimental data presented in previous sections.

Fig. 4 Structural characterization of PSS-WS₂. (a) Raman feature, (b) PL spectrum and (c) photograph of PSS-WS₂, compared with those measured on CSS-WS₂. (e) PL spectra of PSS-WS₂ prepared under various processing temperatures during the PSS process. (f) The intensities of the PL peaks as plotted in (e) with the error bars showing the standard mean. Insets are the AFM images of PSS-WS₂ prepared under 900 °C and 1000 °C. The temperature-dependent optical/structural transition is highlighted into two colored zones.

As opposed to the multilayer WS₂ in the CSS-WS₂ sample, the formation of monolayer PSS-WS₂ certainly indicates that effective thinning of the multilayer WS₂ occurs only during the DES turn-off periods. This is presumably due to the etching and/or enhanced sublimation of the top WS₂ layers by H₂. In understanding this thinning process, we note that there are several previous reports relevant to the thermal and chemical stability of WS₂, which, together with our own experimental observations, may help to reveal the underlying mechanism. First, it was previously reported that a post-growth thermal treatment under H₂ + Ar mixed gas is not effective in etching CVD-grown WS₂ at temperatures up to 900 °C.²⁴ This is in agreement with our experimental finding that monolayer WS₂ formation by PSS is possible only for temperatures above 950 °C. As shown in Fig. 4e and f, there is a significant increase in the PL intensity of PSS-WS₂ without the presence of bilayer domains upon preparing the film above 950 °C. This temperature-dependent optical/structural transition is highlighted into two colored zones showing that the multilayer WS₂ and strictly-monolayer WS₂ were prepared at low temperature range (850–925 °C) and high temperature range (950–1025 °C), respectively. This finding validates that a high enough temperature (larger than the vacuum sublimation temperature ∼870 °C of WS₂)³⁸ is crucial for providing the thermodynamic driving force to trim down the overgrown WS₂ top layers from the bottom monolayer WS₂/SAP. Second, a study of the thermal stability of WS₂ and MoS₂ in vacuum set the sublimation and decomposition temperatures for bulk WS₂ around 870 °C and 1040 °C, respectively, as observed from the compact WS₂ beginning to have a detectable weight loss rate at 870 °C and sulfur peaks appearing at 1040 °C.³⁸ Third, Wu et al. demonstrated the vapor-solid growth of high optical quality MoS₂ monolayers by the sublimation of the MoS₂ powder source at 900 °C.³⁹ Meanwhile, a vacuum sublimation temperature of ∼930 °C was determined for bulk MoS₂, a temperature higher than the vacuum sublimation temperature ∼870 °C of bulk WS₂.³⁸

Based on the experimental observations summarized above and knowing the critical role of H₂ in the PSS process (note: the formation of WS₂ monolayer by PSS failed in the absence of H₂ gas; see Fig. S6, ESI†), we explain the self-limited thinning of multilayer WS₂ by PSS for the formation of a stable monolayer WS₂ as follows. (1) At a temperature substantially higher than the vacuum sublimation temperature of ∼870 °C of WS₂, there exists a certain thermodynamic driving force for multilayer WS₂ to sublimate, with a drastically accelerated sublimation for highly volatile reduced WS_2−x (Fig. 1j). (2) At a temperature (e.g., 1000 °C) much higher than ∼870 °C and in the presence of H₂, the top few-layer WS₂ (formed after the initial 600 °C sulfurization) would undergo a reduction process via H₂(g) + WS₂(s) → H₂S(g) + WS_2−x(s), leading to a drastically enhanced sublimation in the volatile WS_2−x and thus layer-by-layer thinning, during DES turn-off periods (Fig. 1k). (3) This thinning process would be inhibited or drastically minimized in the presence of the DES supply (i.e., during DES turn-on period), due to the massive annihilation of S vacancies (created either during previous DES turn-off period or current DES turn-on period) by the readily available S atoms and, consequently, the restoration of the volatile WS_2−x back to more stable WS₂. During this DES turn-on period, further conversion of underlying remaining WO_x, if there is any, into WS₂ also takes place due to enhanced S diffusion through the top WS₂ multilayer, the thickness of which was trimmed down during the previous turn-off period. (4) Once the multilayer WS₂ finally reduces to a monolayer limit after several DES turn-on/turn-off cycles, the sublimation would stop effectively, even for the volatile WS_2−x phase, most likely due to the stronger interface coupling of the bottom monolayer WS₂ with sapphire substrate, as compared to that for WS₂–WS₂ interface. It should be noted that a more negative binding energy (E_b) of −2.23 eV is obtained for the monolayer WS₂–sapphire interface as compared to E_b = −2.01 eV for the WS₂–WS₂ interface by DFT calculations (Fig. 3b), which means a significantly increased sublimation temperature for the monolayer WS₂ on the sapphire surface. The effective suppression of sublimation in monolayer WS₂ is also evident from the observation of the persistent existence of the PSS-WS₂ monolayer after high temperature annealing under the steady flow of the H₂ + Ar mixed gases without a DES supply (Fig. S7, ESI†). (5) The final pulse DES supply in a complete PSS process cycle would restore the WS_2−x phase to a stable WS₂ monolayer via a reaction of WS_2−x(s) + xS(g) → WS₂(s), thus realizing the self-limited thinning of multilayer WS₂ for monolayer WS₂ formation (Fig. 1l and m). (6) On the other hand, the failure in the formation of monolayer WS₂ by the CSS process is due to the overwhelming restoration process (WS_2−x(g) + xS(g) → WS₂(s)) against the thinning process involving (1) reduction (H₂(g) + WS₂(s) → H₂S(g) + WS_2−x(s)) and (2) subsequent sublimation (WS_2−x(s) → WS_2−y(s) + volatile WS_2−x(g)) under an environment with a constant sulfur supply. Here, we would like to highlight that the self-limited thinning process can also be employed on other suitable substrates, as demonstrated by the formation of the stable monolayer WS₂ on SiO₂/Si substrates (Fig. S8, ESI†). However, it is found that the formation of the monolayer MoS₂ fails on SiO₂/Si (i.e., complete stripping-off of MoS₂), as shown in Fig. S9 (ESI†). In contrast, a continuous monolayer MoS₂ is successfully prepared on sapphire substrates under the same PSS processing condition (Fig. S3, ESI†), clearly indicating the crucial effect of substrate on the PSS process and its extendibility. The instability of monolayer MoS₂ on SiO₂/Si is most likely due to the weaker binding energy of the MoS₂–SiO₂ interface than that for the MoS₂–sapphire interface.

To investigate the electrical properties of the monolayer WS₂ film grown via the self-limited thinning PSS process, bottom-gated FET devices were fabricated, with a channel length L = 25 μm and a channel width W = 200 μm, respectively. Fig. S10 (ESI†) shows the schematic of the device structure. Details of the device fabrication steps are described in the Methods. The fabricated WS₂ FET devices exhibited typical n-channel FET characteristics, as shown in Fig. 5, similar to previous reports for WS₂ FETs either with exfoliated or CVD-grown monolayer WS₂ as a channel.^13,40,41 The WS₂ FET mobility was calculated using the expression,

where I_DS and V_G, are the drain current in the saturation region (under a constant drain voltage V_DS = 2 V), gate voltage and threshold voltage, respectively; C_OX is the gate oxide dielectric capacitance per unit area. We obtained a value of ∼229 nF cm⁻² from the capacitance measurement from the metal–insulator–metal (MIM) capacitor device. Our FET devices show typical mobilities in the range of ∼0.35–0.65 cm² V⁻¹ s⁻¹ and ON/OFF ratios of I_DS ∼ 6 × 10⁴–6 × 10⁵ (Fig. 5a and Fig. S11, ESI†).

Fig. 5 Electrical characteristics of PSS-WS₂ FETs. (a) Transfer characteristic curves at a fixed bias voltage of 2 V. (b) I_DSversus V_DS output curves at a step V_GS from 0.4 to 2 V with 0.4 V intervals.

Our values compare rather favorably to those reported in the literature (μ_eff = 0.16–1.97 cm² V⁻¹ s⁻¹) for the exfoliated single crystalline monolayer WS₂ FETs fabricated with similar source and drain metal contacts,^9,42,43 implying the good quality of our monolayer WS₂ film. It must be pointed out that, while the intrinsic mobility is a material-related parameter that is determined by intrinsic material properties such as crystallinity, crystal-defects/impurities, and grain sizes etc., μ_eff is also profoundly influenced by extrinsic effects, such as the contact resistance, gate-dielectric layer used, and integrity of the interface between the channel and gate-dielectric layer. For atomically-thin WS₂ FETs, a μ_eff as high as 214 cm² V⁻¹ s⁻¹ at 300 K was reported on single layer WS₂ sandwiched between CVD-grown h-BN films.⁴⁴ A μ_eff up to ∼115 cm² V⁻¹ s⁻¹ at 300 K was also achieved for the exfoliated few-layer WS₂ through monolayer-h-BN/Cr/Au tunneling barrier contact. In contrast, a low mobility of ∼1.44 cm² V⁻¹ s⁻¹ at 300 K was extracted on the same few-layer WS₂ film with direct Indium contact.⁴² Therefore, our reported carrier mobilities here should be taken as a lower limit of its intrinsic carrier mobility as there is still potential for improvement through various approaches, e.g., optimization of ALD oxide growth,⁴⁵ choices of dielectric material (HfO₂ instead of Al₂O₃),⁴⁶ improvements in interface quality,^44,47 while the contact resistance can also be further reduced.^42,48 Nonetheless, our transport measurements indicate that a semiconducting WS₂ monolayer film has been successfully grown uniformly by self-limited process on a 4 inch sapphire wafer.

Conclusions

In summary, a sublimation-based self-limited thinning process is demonstrated for the preparation of strictly monolayer WS₂ on a 4-inch sapphire wafer. At a temperature substantially above the vacuum sublimation temperature of WS₂, the multilayer WS₂ formed by the sulfurization of the WO_x precursor film is effectively trimmed down to the monolayer limit via the pulsed supply of sulfur precursor vapor under a continuous flow of H₂. This self-limited thinning process is attributed to the following combined effect: (1) an effective reduction of stable WS₂ to volatile reduced WS_x phase by H₂, and (2) significantly higher sublimation temperature of monolayer WS₂ on sapphire as compared to multilayer or bulk WS₂, as suggested by DFT calculations. The monolayer WS₂ film showed high crystallinity and excellent property uniformity across the entire 4-inch wafer, as revealed by transmission electron microscopy (TEM), Raman microscopy, photoluminescence (PL), and X-ray photoelectron spectroscopy (XPS). The self-limited thinning process overcomes the intrinsic limitation of conventional vapor-based growth methods in preventing the 2^nd layer WS₂ domain nucleation/growth, and thus provides us with the strictly-monolayer WS₂ on the wafer-scale. This approach opens a new avenue for a feasible manufacturing technology of 2D channel material for high efficiency optoelectronics.

Methods

Sputtering growth of the WO₃ precursor films

The WO₃ were magnetron-sputtered on sapphire at a substrate temperature of 450 °C in pulsed DC mode using a W metal target. The base pressure of the system was 1.0 × 10⁻⁷ Torr. During the sputtering process, the working pressure was 2.6 × 10⁻³ Torr with Ar and O₂ gas flow rates of 6 and 5 sccm, respectively. The pulsed DC power generator was operated at 5 kHz with reverse pulse widths of 0.4 μs. The thickness of WO₃ is increased by increasing the sputtering time. For preparing WO₃ with thicknesses of ∼1–3 nm, the sputtering duration was 12–30 s under the same power of 60 W.

Formation and self-limited thinning process of multilayer WS₂ for monolayer formation

For the growth of the WS₂ monolayer, the wafer-sized sapphire substrates (c-plane, Namiki Precision Singapore Pte., Ltd) deposited with 1–3 nm WO₃ precursor films were placed into a vertical 2-zone furnace. The chamber pressure was maintained at 10 Torr under the steady flow of mixed gases (78% volume Ar/22% volume H₂). During the growth, 4 sccm diethyl sulfide (DES, (C₂H₅)₂S, 98%, Sigma-Aldrich) was supplied from the bubbler at 30 °C and regulated precisely with a mass flow controller. Details of the pulsed supply of the DES and the process temperature temporal profile (600–1025 °C) are shown in Fig. 1b. Similarly, for the growth of the Co-doped WS₂ monolayer, the substrate was firstly pre-deposited with a bilayer of Co (0.2 nm)/W (0.6 nm), and followed by the pulsed DES supply at 1000 °C. After the growth process, the quartz tube was naturally cooled down to room temperature under 200 sccm Ar gas.

Materials characterization

The morphologies of WS₂ and Co-doped WS₂ were analyzed via Bruker's dimension icon AFM system. The WITec alpha300S Raman instrument was used to measure the micro-Raman and photoluminescence (PL) spectra in order to confirm the quality and crystallinity of the grown WS₂ and Co-doped WS₂ film under a laser source with 532 nm wavelength. The XPS measurements for the pre-deposited tungsten oxides, WS₂ and Co-doped WS₂ were conducted by a VG ESCALAB 220i-XL system with a monochromatic Al Kα source and a pass energy of 10 eV. The optical absorption spectra were measured with the Shimadzu UV-VIS-NIR spectrophotometer UV-3150. To determine the atomic structure and orientations of the monolayer WS₂, the FEI Titan 80/300 Scanning/TEM (200 kV) and FEI Tecnai F20 with fast Fourier transform (FFT) measurements were used for imaging the high-resolution TEM images after general transferring of the WS₂ by poly(methyl methacrylate) (PMMA 950 in anisole, A5, KAYAKU Advanced Materials, Inc.) onto the 200-mesh TEM Cu grid. The focused ion beam (FIB) technique was employed to prepare a lamella, which was used for obtaining a cross-sectional image of the WS₂ film.

Simulation study

To execute the DFT simulations, the Vienna Ab initio Simulation Package (version 5.4) using the Perdew–Burke–Ernzerhof (PBE)⁴⁹ was applied. The geometry was optimized under DFT simulation within the generalized gradient approximation.^50,51 The expansion of the plane-wave had a kinetic energy cutoff at 450 eV. The slab structures consisted of Gamma-centered grids with 5 × 5 × 1 k-point grids were employed. Under the correction of the dipole and the DFT-D3 method,⁵² the interactions between the adjacent images and the dispersions were reduced respectively under the vacuum setting of no less than 14 Å.

(2 × 2) Al₂O₃ and (3 × 3) WS₂ were employed in building the model of WS₂ on Al₂O₃. The lattice mismatch was smaller than 0.7% with around 9.3 Å thick Al₂O₃ slab. In the optimization stage, we have frozen the bottom four layers (total nine layers). The optimized lattices of the sapphire substrate (bulk Al₂O₃) were a = 4.785 Å and c = 13.053 Å, while the bulk WS₂ (2H) were a = 3.169 Å and c = 12.423 Å. The oxygen-terminated slab models were constructed based on the bulk structures.

Considering the interaction along the interface will affect the thermal stability significantly, the theoretical binding energy (E_b) at the WS₂–WS₂ interface and WS₂–sapphire interface (Fig. 3b) was calculated by eqn (2):

E_b = E_{substrate+xWS2} − (E_{substrate+(x−1)WS2} + E_WS2) (2)

where E_{substrate+xWS2}, E_{substrate+(x−1)WS2} and E_WS2 are the sum of energies in the system of sapphire (bulk Al₂O₃) and x-layer of WS₂, the sum of energies in the system of sapphire and (x − 1) WS₂ layer (x is 1 or 2), and a WS₂ monolayer, respectively.

Device fabrication and electrical measurements

SiO₂/Si substrates and AlO_x/p⁺-Si substrates were prepared by atomic layer deposition (ALD) method. A 30 nm of AlO_x was grown by using 300 ALD cycles (Picosun) at 150 °C with the precursor trimethylaluminum (TMA) and oxidant H₂O. Under a N₂ carrier gas flow rate of 150 sccm, the pulse and purging time for TMA and H₂O were set to 100 ms and 20 s, respectively, in every ALD cycle. The capacitance of AlO_x was characterized by Keithley 4200. The 1 × 1 mm² Au electrodes were evaporated onto AlO_x to fabricate the metal–insulator–metal (MIM) capacitor device. Across the sweeping frequency range from 1000 to 10⁶ Hz, C_i remains almost constant at ∼229 nF cm⁻². Both SiO₂/Si and AlO_x/p⁺-Si substrates were cleaned by acetone and isopropanol (IPA) sequentially, and were subjected to 20 min O₂ plasma (Harrick Plasma Cleaner) treatment to hydroxylate the oxide surface.

To characterize the electrical performance, the monolayer WS₂ was transferred onto the fresh target substrate by a conventional PMMA method.³⁵ In short, PMMA was spun onto the as-grown WS₂/sapphire with 3000 rpm for 60 s, and then heated at 150 °C for 2 min. Next, the PMMA-coated WS₂ was separated from sapphire by dipping into 1 M KOH solution slowly. It was then picked up with a PDMS sheet, followed by rinsing in DI water several times to remove the residuals. Subsequently, the new PDMS stamp (DGL-45 × 45-0065-X4, GelPak) was then used to transfer the resultant film onto the target substrate and dried inside a fume hood overnight. Next, the PDMS was carefully detached from PMMA/WS₂ at 150 °C, followed by a rinsing process to remove PMMA in the acetone and IPA baths. Finally, the sample was annealed at 300 °C for 30 min under Ar atmosphere.

The WS₂ devices were fabricated based on the classical bottom-grid-top-electrode structure. The Ti/Au (10 nm/40 nm) electrodes were patterned by shadow mask under electron beam evaporation. The channel width and length were determined by Field Emission Scanning Electron Microscopy (FESEM, JEOL JSM7600F). Electrical measurements were carried out in air condition by a semiconductor device analyzer (B1500A, Agilent) equipped with probe stations (Probing Solutions, Inc.).

Author contributions

M. C.: writing – original draft. H. M and S. W. T: conceptualization. M. C., S. W. T., H. Z. and H. M. conducted the experiments including the growth and the transfer of WS₂ for device fabrication and TEM study. J. C. conducted XPS measurements. M. C. and J. W. fabricated and characterized the two-terminal FETs. W. W. and M. L. performed the TEM measurements. J. L. and H. M. proposed and prepared the Co-doped WS₂ samples. M.-F. N. performed and analyzed the DFT simulation results. S. W. T. and M. C. conducted the Raman, PL, AFM and UV-VIS measurements. J.-Y. K. and J. K. prepared the PVD sputtered MoS₂. D. C., S. W. T., and H. M. supervised the work. All authors reviewed, edited, and approved the manuscript.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

The authors gratefully acknowledge financial supports from National Research Foundation Competitive Research Program (NRF-CRP24-2020-0002), the Institute of Materials Research and Engineering (IMRE) under the Agency for Science, Technology, and Research (A*STAR) via A*STAR Science and Engineering Research Council (SERC Grant No 152-70-00015, SERC Grant No A20G9b0135) and Career Development (Fund/C210112032). J. W. acknowledges the Advanced Manufacturing and Engineering Young Individual Research Grant (AME YIRG Grant No.: A2084c170) and the SERC Central Research Fund (CRF). We also acknowledge the National Supercomputing Centre (NSCC) Singapore and A*STAR Computational Resource Centre (A*CRC) of Singapore through the use of its high-performance computing facilities.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d3nh00358b

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