Interface Kinetics Assisted Barrier Removal in Large Area 2D-WS2 Growth to Facilitate Mass Scale Device Production

Growth of monolayer WS2 of domain size beyond few microns is a challenge even today; and it is still restricted to traditional exfoliation techniques, with no control over the dimension. Here, we present the synthesis of mono- to few layer WS2 film of centimeter2 size on graphene-oxide (GO) coated Si/SiO2 substrate using the chemical vapor deposition CVD technique. Although the individual size of WS2 crystallites is found smaller, the joining of grain boundaries due to sp2-bonded carbon nanostructures (~3–6 nm) in GO to reduced graphene-oxide (RGO) transformed film, facilitates the expansion of domain size in continuous fashion resulting in full coverage of the substrate. Another factor, equally important for expanding the domain boundary, is surface roughness of RGO film. This is confirmed by conducting WS2 growth on Si wafer marked with few scratches on polished surface. Interestingly, WS2 growth was observed in and around the rough surface irrespective of whether polished or unpolished. More the roughness is, better the yield in crystalline WS2 flakes. Raman mapping ascertains the uniform mono-to-few layer growth over the entire substrate, and it is reaffirmed by photoluminescence, AFM and HRTEM. This study may open up a new approach for growth of large area WS2 film for device application. We have also demonstrated the potential of the developed film for photodetector application, where the cycling response of the detector is highly repetitive with negligible drift.


Introduction
Lamellar two-dimensional (2D) WS 2 has immense potential for electronic applications because of its remarkable properties such as tunable direct bandgap [1][2][3], high photoluminescence intensity [4,5], high emission quantum field [6], attractive spin-orbit coupling [6], substantial exciton/trion binding energies [7][8][9], etc., which give WS 2 superior leverage over other transition-metal dichalcogenides (TMDCs). Unfortunately, the research on WS 2 is not matured yet, especially large-scale production of single-crystal monolayers [10][11][12]. Challenges lie in reliable synthesis of atomically thick 2D layers and controlled manipulation of electronic properties [13]. Though mechanical cleaving still remains a relatively facile technique to prepare high quality single to few layered WS 2 , problems such as absence of control over thickness, relatively small lateral size and low yield limit its applications and commercialization [14,15]. The most popular techniques to produce few-layered WS 2 include intercalation driven exfoliation [16], liquid phase exfoliation [17], physical vapor deposition [18], hydrothermal reaction [19], and heat treatment of W and S containing precursors [20]. However, the domain size of developed WS 2 via aforementioned methods is generally restricted to a few microns (µm), and the production of WS 2 monolayer films with desired domain size remains a distant dream. To achieve this, several strategies (e.g., atomic layer deposition [21], pulsed-laser deposition [22,23], metal/metal oxide thin films [24], suitable metal substrates [25], etc.), have been suggested Nanomaterials 2021, 11, 220 2 of 19 to achieve uniform dispersion of precursors over substrates. Presently, large scale continuous growth of TMDCs is highly complex and expensive, and results in poor quality films [9].
Chemical vapor deposition (CVD) remains a challenge for a long time for growing two-dimensional TMDC materials with high crystallinity, desired thickness, and sufficient domain size [10,12,[26][27][28]. For WS 2 growth, typical synthesis approach involves sulfidation of WO 3 powders at sufficiently high temperature. The growth conditions such as precursor type, process step, and their manipulations during growth, are found to greatly affect the product quality. Most of the reports on CVD grown WS 2 solely bank on vapor phase reaction between S and W precursors at suitably high temperatures, where a single furnace is used for both precursors. In such systems, there is no control over the temperature of individual precursor, thereby limiting the parameter space of growth reaction [28]; leading to poor uniformity and repeatability. High melting point of WO 3 powders and sulfurization rate are found to influence the growth process. Achieving large triangle shaped growth is challenging, because of lower vapor pressure of WO 3 due to its high melting temperature (1473 • C) [29]. This is a serious problem and severely lowers the partial pressure of WO 3 vapors [29]. Further, the low vapor pressure of WO 3 also hampers the availability of W atoms on substrate surface. Thus, enhancing the partial pressure of WO 3 vapors is a natural approach to obtain large sized WS 2 flakes [12]. This can be accomplished by reducing the pressure of the furnace. However, this increases the S vapor transfer speed, thereby increasing the sulfurization rate [10]. Fast sulfurization is not desirable for the transport and diffusion of precursor vapors on a substrate and may result in aggregation of precursors at certain locations on the substrate, thereby limiting the growth of large sized WS 2 crystals [10,30,31]. An efficient way involves a trade-off between the increase in partial pressure of WO 3

vapors and lower the transfer speed of S vapors.
Recently, many researchers have focused on developing monolayer WS 2 via CVD. Zhang et al. employed low-pressure chemical vapor deposition (LPCVD) technique to synthesize atomically thin triangle shaped WS 2 crystals on sapphire substrate having single domain size~50 µm [32]. An improvement was introduced in the CVD process by Cong et al. to effectively increase the concentration of precursors and furnace pressure [33]. They achieved single-domain growths of~178 µm [33]. Li et al. introduced further changes in the CVD reaction by employing alkali metal halides as growth promoters [34]. The groups of Zhang [32] and Li [34] reported that due to strong reducing nature, the mixing of suitable amount of H 2 enhances the sublimation and sulfurization of WO 3 precursor. Fu et al. [35] investigated the influence of CVD reaction temperature and gas flow rate (comprising a mixture of 97% Ar and 3% H 2 ) on the growth size and morphology of WS 2 films. The optimized growth conditions yielded a domain size of~52 µm. Rong's group [28] achieved precise control over S introduction time by utilizing a two-zone furnace and obtained very large area WS 2 films of 370 µm domain size. Recently, Liu et al. [36] reported sequential synthesis of several one-, and two-dimensional nanomaterials by intentionally creating initial low sulfur conditions, to divide the growth process in two stages: the first stage is a partial reduction and the second one is sulfurization. By this, they could exploit a WO 3-x intermediate route, and were able to grow large size WS 2 (150 µm).
TMDCs growth is very susceptible to substrate treatment prior to formation [37]. To assist the nucleation by seeding the substrate with various aromatic molecules, such as -3,4,9,10-tetracarboxylic acid tetrapotassium salt (PTAS), perylene-3,4,9,10-tetracarboxylic dianhydride (PTCDA), and reduced graphene oxide (RGO), perylene encourages the lateral growth of TMDCs. However, this seeding technique helps to grow the large area mono-to few layer TMDCs but continuous film from the abovementioned technique is still limited. Although the CVD method has many advantages, the coordination among various growth parameters is highly complex and requires further elucidation.
In this work, we present a novel approach for large area synthesis of WS 2 on graphene oxide (GO) surface via CVD under high temperature conditions of 1070 • C. Upon transformation to RGO during growth process, GO becomes a mixed sp 3 -sp 2 bonded network, where nano-sized sp 2 -patches of size~3-6 nm, are randomly distributed with more than 60-70% coverage [38][39][40]. The sp 2 -patches are six-fold aromatic carbon atoms with C=C bonds, whereas major percentage of sp 3 bonds constitutes epoxides (C-O) and hydroxyl (−OH) groups in the basal plane and somewhat lesser amount of carboxyl (−COOH) and carbonyl (−C=O) functionalities in the edges [38][39][40]. These sp 2 -chemical bonds act as a seed to promote the growth of large area WS 2 film. The as-developed WS 2 layers are polycrystalline comprising mainly mono-to few layers. The single-crystal domains are triangular and hexagonal in shape, with sizes up to 60 µm on Si/SiO 2 and~15 µm on the GO coated Si/SiO 2 , respectively. The domain size of WS 2 film on GO coated Si/SiO 2 is much smaller than that on non-GO substrate, but the development of large continuous surface makes a great footprint in WS 2 growth, attributed to C=C assisted coalescence of high-density polycrystalline grain boundaries into large domain area. Suitable explanation is given how sp 2 -patches comprised of C=C bonds are pivotal in the expansion of domain size of WS 2 film. Surface roughness of the top layer of substrate is a well-recognized technique for nucleation of nanostructure growth and is cross-checked with intentionally created rough surface on silicon wafer. The proposed device fabrication technique is well suited for mass production of identical devices, and therefore bids for commercial scale development. Moreover, this technique can be extended to the fabrication of other TMDCs. We have further demonstrated the use of the developed film for photodetector applications.
The morphology, size and structure of as-synthesized WS 2 were examined via a transmission electron microscope (TEM) (JEOL JEM F-200, Tokyo, Japan) working at 200 kV. The samples for TEM investigations were kept on the conventional holey carbon Cu grid. Highresolution transmission electron microscopy (HRTEM, JEOL, Tokyo, Japan) studies reveal the uniformity and high quality of WS 2 samples grown on both Si/SiO 2 and GO coated Si/SiO 2 substrates. Basic studies of WS 2 growth and the domain size were performed on an optical microscope (Leica DM4 P, Wetzlar, Germany). Raman spectrophotometers (WITec Alpha 300RA, Ulm, Germany and LabRAM HR800 HORIBA JY, Kyoto, Japan) fitted with a Peltier cooled CCD detector and confocal microscope were operated for verifying the crystallinity, defects and number of layers. A diode laser of excitation wavelength 532 nm was used to excite the samples. A 100X objective lens with spot size of 1 µm was used to focus the excitation onto the samples. The thickness of grown sample was analyzed by atomic force microscopy (AFM) fitted with the Raman system. Surface morphology of the CVD grown WS 2 on Si/SiO 2 and GO coated Si/SiO 2 was studied on a scanning electron microscope (SEM) (Nova Nano SEM 450, FEI, Oregon, USA). The photoluminescence (PL) measurements were performed by fluorescence spectrometer (Agilent Technologies, Carry Eclipse fluorescence spectrometer, California, USA). For PL investigations, the prepared samples were dispersed in ethanol by simply ultrasonicating the growth substrate for few minutes and isolating the upper one-third solution into a quartz cuvette.

Results and Discussion
Two approaches have been adopted to synthesize monolayer WS 2 of large domain size, where CVD is employed to grow WS 2 on two substrates, i.e., Si/SiO 2 and GO-coated Si/SiO 2 (Si/SiO 2 /GO), with the aim to extend the domain size of WS 2 and uniform coverage of the entire substrate area (see Sections 3.1 and 3.2).
3.1. First Approach: Mono-to Few-Layered WS 2 Growth on Si/SiO 2 WS 2 growth was carried out inside a two-zone tube furnace having separate highand low-temperature zones, denoted as HT and LT, respectively, as shown in Figure 1a. Ultrahigh purity argon at flow rate of 100 sccm under atmospheric pressure was used as carrier gas to transport sulfur vapors to the HT zone. In a typical experiment, WO 3 powders (200 mg of pure grade 99.9%) was placed in an alumina crucible at the center of HT-zone of the furnace. Si/SiO 2 substrate (10 mm × 10 mm) was placed on the alumina crucible in inverted position so as to allow WO 3 vapor to hits the substrate and deposit on it. The sulfur powder (400 mg of pure grade 99.5%) was loaded upstream, and heated in LT-zone of the furnace. The HT-and LT-zone temperatures were controlled in such a way that the temperature in both the zones simultaneously reach at 1070 • C and 220 • C, respectively [28]. Monolayer WS 2 grains were grown on Si/SiO 2 substrate having small domain size and the growth remained scattered. The reaction occurring at the surface of the substrate can be summarized as: (1) Ultrahigh purity argon at flow rate of 100 sccm under atmospheric pressure was used as carrier gas to transport sulfur vapors to the HT zone. In a typical experiment, WO3 powders (200 mg of pure grade 99.9%) was placed in an alumina crucible at the center of HTzone of the furnace. Si/SiO2 substrate (10 mm × 10 mm) was placed on the alumina crucible in inverted position so as to allow WO3 vapor to hits the substrate and deposit on it. The sulfur powder (400 mg of pure grade 99.5%) was loaded upstream, and heated in LT-zone of the furnace. The HT-and LT-zone temperatures were controlled in such a way that the temperature in both the zones simultaneously reach at 1070 °C and 220 °C, respectively [28]. Monolayer WS2 grains were grown on Si/SiO2 substrate having small domain size and the growth remained scattered. The reaction occurring at the surface of the substrate can be summarized as: Thus, the vapor phase reaction between tungsten oxide and sulfur, produces tungsten disulfide and sulfur dioxide as end products. The optimized time duration for WS2 growth is found to be 45 min, once the temperature of HT zone reaches 1070 °C. After completing the growth process, the CVD system was allowed to cool naturally to the room temperature. Figure 1b-e shows the optical images of the WS2 sample grown on Si/SiO2 substrate. Two types of crystallites are visible: (a) equilateral triangles, and (b) hexagons; both are characteristic of WS2 single crystal formation. Most of the crystals include monolayers. However, bi-and tri-layered regions are also visible at the center of large monolayer crystals, which is supposed to be the nucleation site, even for further growth of multi-layers [41]. Formation of hexagons in case of multilayer grains originates from the random orientation of individual WS2 layers and differences in crystal orientation, where each orientation acts as nucleation site [41][42][43]. The side length of the crystals in case of equilateral triangles is found to be ~60 μm, which is reasonably large. The CVD growth of TMDCs is generally believed to be a self-limiting process [44]. The lateral size of WS2 increases with Thus, the vapor phase reaction between tungsten oxide and sulfur, produces tungsten disulfide and sulfur dioxide as end products. The optimized time duration for WS 2 growth is found to be 45 min, once the temperature of HT zone reaches 1070 • C. After completing the growth process, the CVD system was allowed to cool naturally to the room temperature. Figure 1b-e shows the optical images of the WS 2 sample grown on Si/SiO 2 substrate. Two types of crystallites are visible: (a) equilateral triangles, and (b) hexagons; both are characteristic of WS 2 single crystal formation. Most of the crystals include monolayers. However, bi-and tri-layered regions are also visible at the center of large monolayer crystals, which is supposed to be the nucleation site, even for further growth of multi-layers [41].
Formation of hexagons in case of multilayer grains originates from the random orientation of individual WS 2 layers and differences in crystal orientation, where each orientation acts as nucleation site [41][42][43]. The side length of the crystals in case of equilateral triangles is found to be~60 µm, which is reasonably large. The CVD growth of TMDCs is generally believed to be a self-limiting process [44]. The lateral size of WS 2 increases with the partial pressure of chalcogen precursors [45], while its thickness is determined by the partial pressure of transition metal precursors [46]. We have kept the weight of sulfur powders to be the twice of the weight WO 3 (chalcogen precursor). Thus, the lateral growth dominates in our case and could explain the growth of large size single crystals [46]. Furthermore, we could find WS 2 single crystals in variable sizes, ranging from 100 nm to 60 µm on the same Si/SiO 2 substrate. The growth duration is found to directly influence the area of grown crystals, before the crystal growth reaches its limit. Therefore, large sized WS 2 crystals can be formed if the growth time is kept sufficiently long.

•
Electron Microscope Analysis Figure 2a-c shows the FESEM images of WS 2 samples deposited on Si/SiO 2 substrate, corroborating the findings of optical microscopy, indicating both equilateral triangles and hexagons. In case of WS 2 , the growth of equilateral triangles suggests single crystallinity; whereas the hexagonal or star-like structures result from several rotationally symmetric grains [4,9,26,46,47]. The size distribution histogram of hexagonal and triangle flakes of WS 2 is shown in Figure 2d-e. Where the average size of the hexagonal and triangular flake is found to be 71.5 µm and 44.5 µm, respectively. The ratio of the hexagon to the triangular flakes over the substrate is 2:1. Notably, the growth products, though consisting of monolayers, are scattered over the substrate and do not cover the entire substrate surface.
the partial pressure of chalcogen precursors [45], while its thickness is determined by the partial pressure of transition metal precursors [46]. We have kept the weight of sulfur powders to be the twice of the weight WO3 (chalcogen precursor). Thus, the lateral growth dominates in our case and could explain the growth of large size single crystals [46]. Furthermore, we could find WS2 single crystals in variable sizes, ranging from 100 nm to 60 μm on the same Si/SiO2 substrate. The growth duration is found to directly influence the area of grown crystals, before the crystal growth reaches its limit. Therefore, large sized WS2 crystals can be formed if the growth time is kept sufficiently long.

•
Electron Microscope Analysis Figure 2a-c shows the FESEM images of WS2 samples deposited on Si/SiO2 substrate, corroborating the findings of optical microscopy, indicating both equilateral triangles and hexagons. In case of WS2, the growth of equilateral triangles suggests single crystallinity; whereas the hexagonal or star-like structures result from several rotationally symmetric grains [4,9,26,46,47]. The size distribution histogram of hexagonal and triangle flakes of WS2 is shown in Figure 2d-e. Where the average size of the hexagonal and triangular flake is found to be 71.5 μm and 44.5 μm, respectively. The ratio of the hexagon to the triangular flakes over the substrate is 2:1. Notably, the growth products, though consisting of monolayers, are scattered over the substrate and do not cover the entire substrate surface.  3a-c shows a low magnification bright field TEM imaging of WS2 crystals suspended on a holey carbon Cu grid. For TEM imaging, first the WS2 grown onto the Si/SiO2 substrate was ultrasonicated in ethanol for 20 min to get the WS2 flakes. After this, the suspension was centrifuged at a speed of 3000 rpm for 15 min to remove unwanted contamination or to allow settling down of heavy particles. Finally, top one-third supernatant was used for TEM analysis. We used HRTEM and selected area electron diffraction (SAED) to characterize the WS2 domains. Figure 3a shows the hexagonal flakes with lateral dimension of about ~2.5 μm. In Figure 3b, the distorted cum truncated WS2 triangle is clearly observable. At higher resolutions (Figure 3c-d), the crystalline nature of the  After this, the suspension was centrifuged at a speed of 3000 rpm for 15 min to remove unwanted contamination or to allow settling down of heavy particles. Finally, top onethird supernatant was used for TEM analysis. We used HRTEM and selected area electron diffraction (SAED) to characterize the WS 2 domains. Figure 3a shows the hexagonal flakes with lateral dimension of about~2.5 µm. In Figure 3b, the distorted cum truncated WS 2 triangle is clearly observable. At higher resolutions (Figure 3c,d), the crystalline nature of the flakes is clearly displayed. In Figure 3e, the selected area electron diffraction (SAED) pattern is indicated, which further confirms the highly crystalline morphology of as-grown WS 2 . flakes is clearly displayed. In Figure 3e, the selected area electron diffraction (SAED) pattern is indicated, which further confirms the highly crystalline morphology of as-grown WS2. •

Raman Analysis
The E-k diagram and the phonon dispersion in Brillouin zone for WS2 are shown in Figure 4a-c. Raman spectroscopy was employed to detect both the in-plane (E2g 1 ) and outof-plane (A1g) phonon modes. Figure 4d shows Raman spectrum of the WS2 flakes deposited over the Si/SiO2 substrate, indicating a strong peak at 354 cm −1 (2LA(M)), signifying a longitudinal acoustic mode at M point. This signals the monolayer growth of WS2 [9,26]. The peak at 421 cm −1 corresponding to the A1g mode, although weak, is also observed. As is widely reported, the separation between these two peaks increases from 67 to 72 cm −1 with increasing the number of layers from single-to tri-layer due to a red-shifting of 2LA(M) mode [26,[48][49][50][51]. However, the observed peak separation of ~67 cm −1 is the fingerprint of WS2 monolayer growth. Its topography, revealing a triangular shape probed via Atomic Force Microscopy (AFM), is shown in Figure 4e,f. The whole crystal is almost atomically flat and clean. The height profile in Figure 4f indicates the as-grown WS2 on Si/SiO2 having a step size of 0.89 nm; reaffirming the growth of a WS2 monolayer. To verify the uniformity in the flakes, Raman mapping studies were performed from one edge to the other (Figure 5a,b). The synthesized WS2 flakes were further analyzed by recording multiple Raman spectra across the area of post WS2 grown substrate. Typically, seven measurement sites were selected on the top surface, and it can be concluded that we have achieved the monolayered WS2 flakes. As seen in Figure 5a,b, the spots 1 and 7 only show •

Raman Analysis
The E-k diagram and the phonon dispersion in Brillouin zone for WS 2 are shown in Figure 4a-c. Raman spectroscopy was employed to detect both the in-plane (E 2g 1 ) and out-of-plane (A 1g ) phonon modes. Figure 4d shows Raman spectrum of the WS 2 flakes deposited over the Si/SiO 2 substrate, indicating a strong peak at 354 cm −1 (2LA(M)), signifying a longitudinal acoustic mode at M point. This signals the monolayer growth of WS 2 [9,26]. The peak at 421 cm −1 corresponding to the A 1g mode, although weak, is also observed. As is widely reported, the separation between these two peaks increases from 67 to 72 cm −1 with increasing the number of layers from single-to tri-layer due to a redshifting of 2LA(M) mode [26,[48][49][50][51]. However, the observed peak separation of~67 cm −1 is the fingerprint of WS 2 monolayer growth. Its topography, revealing a triangular shape probed via Atomic Force Microscopy (AFM), is shown in Figure 4e,f. The whole crystal is almost atomically flat and clean. The height profile in Figure 4f indicates the as-grown WS 2 on Si/SiO 2 having a step size of 0.89 nm; reaffirming the growth of a WS 2 monolayer. To verify the uniformity in the flakes, Raman mapping studies were performed from one edge to the other (Figure 5a,b). The synthesized WS 2 flakes were further analyzed by recording multiple Raman spectra across the area of post WS 2 grown substrate. Typically, seven measurement sites were selected on the top surface, and it can be concluded that we have achieved the monolayered WS 2 flakes. As seen in Figure 5a,b, the spots 1 and 7 only show the Raman-active LO mode of Si, while spots 2, 3, 4, 5, and 6 express the Raman features of WS 2 monolayers, and interestingly, the Raman result is also found supplemented by AFM studies. the Raman-active LO mode of Si, while spots 2, 3, 4, 5, and 6 express the Raman features of WS2 monolayers, and interestingly, the Raman result is also found supplemented by AFM studies.  •

Photoluminescence (PL) Analysis
The PL properties of 2D-TMDCs are mainly governed by their excitonic features [1,[7][8][9]52]. In the limiting case of 2D materials, as shown in Figure 6a, the lines of electric field between the hole and electron start extending beyond the material, thereby reducing the dielectric screening and enhancing coulomb interactions [36,38,39]. A strong coulombic force is the main reason behind strong photoemission efficiency. This enhanced coulomb interaction suggests formation of bound e-h pairs or "excitons", having greater the Raman-active LO mode of Si, while spots 2, 3, 4, 5, and 6 express the Raman features of WS2 monolayers, and interestingly, the Raman result is also found supplemented by AFM studies.

Photoluminescence (PL) Analysis
The PL properties of 2D-TMDCs are mainly governed by their excitonic features [1,[7][8][9]52]. In the limiting case of 2D materials, as shown in Figure 6a, the lines of electric field between the hole and electron start extending beyond the material, thereby reducing the dielectric screening and enhancing coulomb interactions [36,38,39]. A strong coulombic force is the main reason behind strong photoemission efficiency. This enhanced coulomb interaction suggests formation of bound e-h pairs or "excitons", having greater •

Photoluminescence (PL) Analysis
The PL properties of 2D-TMDCs are mainly governed by their excitonic features [1,[7][8][9]52]. In the limiting case of 2D materials, as shown in Figure 6a, the lines of electric field between the hole and electron start extending beyond the material, thereby reducing the dielectric screening and enhancing coulomb interactions [36,38,39]. A strong coulombic force is the main reason behind strong photoemission efficiency. This enhanced coulomb interaction suggests formation of bound e-h pairs or "excitons", having greater binding energies making them stable even at room temperature [7,[52][53][54]. In case of monolayer 2D material, significant deviation is expected from the hydrogenic model, prevalent to describe Wannier-excitons in inorganic semiconducting materials [7,51]. Therefore, the calculation of binding energy of excitons in case of single-layered WS 2 needs to consider modified potential [54]. High intense peak is observed at 630 nm (1.968 eV) in the as-synthesized monolayer WS 2 , which mainly originates from neutral-exciton (X A ) emission due to direct transition between the lowest conduction band (CB) and the highest valence band (VB) at the same K-point in the Brillouin zone [7,49,[55][56][57][58], and schematically represented in Figure 4c. The fundamental bandgap of a single layer WS 2 is~2.1 eV, but the large excitonic binding energy of 0.032 eV renders the optical bandgap measured at 1.968 eV [8,[50][51][52] as shown schematically in Figure 6a,c shows the PL spectra of WS 2 flakes on the Si/SiO 2 substrate. Figure 6b represents the Neutral A-excitonic transition (XA) in the energy level diagram. The FWHM values of the PL spectra lie between 25 to 35 meV for both the approaches which is in accordance with the literature and further establish the fact of monolayer formation [59,60]. The presence of enhanced coulombic interaction at higher distant points suggests more likelihood of formation of higher-order excitons, i.e., trions and bi-excitons [51].
binding energies making them stable even at room temperature [7,[52][53][54]. In case of monolayer 2D material, significant deviation is expected from the hydrogenic model, prevalent to describe Wannier-excitons in inorganic semiconducting materials [7,51]. Therefore, the calculation of binding energy of excitons in case of single-layered WS2 needs to consider modified potential [54]. High intense peak is observed at 630 nm (1.968 eV) in the as-synthesized monolayer WS2, which mainly originates from neutral-exciton (XA) emission due to direct transition between the lowest conduction band (CB) and the highest valence band (VB) at the same K-point in the Brillouin zone [7,49,[55][56][57][58], and schematically represented in Figure 4c. The fundamental bandgap of a single layer WS2 is ~2.1 eV, but the large excitonic binding energy of 0.032 eV renders the optical bandgap measured at 1.968 eV [8,[50][51][52] as shown schematically in Figure 6a,c shows the PL spectra of WS2 flakes on the Si/SiO2 substrate. Figure 6b represents the Neutral A-excitonic transition (XA) in the energy level diagram. The FWHM values of the PL spectra lie between 25 to 35 meV for both the approaches which is in accordance with the literature and further establish the fact of monolayer formation [59,60]. The presence of enhanced coulombic interaction at higher distant points suggests more likelihood of formation of higher-order excitons, i.e., trions and bi-excitons [51].

Second Approach: Uniform Growth of WS2 Flakes on Spin-Coated GO
In this approach, we have spin coated pre-synthesized graphene oxide (GO) solution over the Si/SiO2 substrate prior to CVD growth. Hummers' method [38][39][40] has emerged as a well-established technique to prepare graphite oxide which can then be ultrasonicated to yield graphene oxide (GO). Briefly, natural graphite powders (5.0 g), H2SO4 (120 mL) and NaNO3 (2.5 g) were mixed via continuous stirring. During stirring, the temperature was kept below 10 °C in order to check overheating of the mixture. After obtaining a uniform mixing, KMnO4 (15.0 g) was added to this mixture by stirring for 24 h at room temperature. A condensed light brown colored slurry was obtained, which was diluted with deionized water (150 mL) and stirring for 2 h at 98 °C. The resulting mixture was vigorously washed with HCl and deionized water, to yield graphite oxide. Subsequently, graphite oxide was ultrasonicated (37 kHz, 500 W) for 3 h. The resulting suspension was centrifuged at 5000 rpm and GO was obtained by collecting the top one-third supernatant from the centrifugation product. A suspension of GO was prepared by mixing 5 mg GO in deionized water via ultrasonication. The substrate for WS2 growth was prepared by spin coating (2500 rpm) 100 μL GO suspension on a Si/SiO2 wafer. Figure 7a includes the SEM micrograph of the GO film coated over Si/SiO2 substrate. The GO flakes appear uniformly dispersed over the surface of the substrate, and are expected to act as a seed layer for the growth of WS2 film. In Figure 7b, the Raman spectrum of the GO film is displayed in 500-3500 cm −1 range, in which the D-, G-and 2D-band are clearly defined at 1354, 1592,

Second Approach: Uniform Growth of WS 2 Flakes on Spin-Coated GO
In this approach, we have spin coated pre-synthesized graphene oxide (GO) solution over the Si/SiO 2 substrate prior to CVD growth. Hummers' method [38][39][40] has emerged as a well-established technique to prepare graphite oxide which can then be ultrasonicated to yield graphene oxide (GO). Briefly, natural graphite powders (5.0 g), H 2 SO 4 (120 mL) and NaNO 3 (2.5 g) were mixed via continuous stirring. During stirring, the temperature was kept below 10 • C in order to check overheating of the mixture. After obtaining a uniform mixing, KMnO 4 (15.0 g) was added to this mixture by stirring for 24 h at room temperature. A condensed light brown colored slurry was obtained, which was diluted with deionized water (150 mL) and stirring for 2 h at 98 • C. The resulting mixture was vigorously washed with HCl and deionized water, to yield graphite oxide. Subsequently, graphite oxide was ultrasonicated (37 kHz, 500 W) for 3 h. The resulting suspension was centrifuged at 5000 rpm and GO was obtained by collecting the top one-third supernatant from the centrifugation product. A suspension of GO was prepared by mixing 5 mg GO in deionized water via ultrasonication. The substrate for WS 2 growth was prepared by spin coating (2500 rpm) 100 µL GO suspension on a Si/SiO 2 wafer. Figure 7a includes the SEM micrograph of the GO film coated over Si/SiO 2 substrate. The GO flakes appear uniformly dispersed over the surface of the substrate, and are expected to act as a seed layer for the growth of WS 2 film. In Figure 7b, the Raman spectrum of the GO film is displayed in 500-3500 cm −1 range, in which the D-, G-and 2D-band are clearly defined at 1354, 1592, and 2712 cm −1 , respectively. Further, XRD analysis was performed to confirm the quality of prepared GO. The XRD pattern in Figure 7c shows the peak at 2θ ≈ 11 • , which validates the high quality of the GO flakes. and 2712 cm −1 , respectively. Further, XRD analysis was performed to confirm the quality of prepared GO. The XRD pattern in Figure 7c shows the peak at 2θ ≈ 11°, which validates the high quality of the GO flakes. The WS2 growth was performed similar to that described in Section 3.1 (i.e., WS2 grown at Si/SiO2 substrate). The only difference is the coating of Si/SiO2 substrate with GO, which is expected to get converted into RGO due to in-situ heating encountered in the HT-zone during the growth process. The synthesis process is schematically illustrated in Figure 8a-c. Figure 8d-g shows the Optical images of the WS2 grown on the GO coated Si/SiO2, where in Figure 8d clear interface is shown between WS2 film and the Si substrate. Figure 8h-k shows the FESEM micrographs at different magnification. In this approach, sufficient amount of sulfur vapor is taken into consideration for the reaction, similar to the first approach. Here, it is noteworthy to say that the sp 2 -carbon nanostructures in the RGO matrix fulfill three-pronged objectives: (i) they improve the surface adhesion, thereby enhancing local adsorption of WO3 molecules; (ii) they reduce WO3 to yield WO3x; and finally, (iii) they also act as active sites for heterogeneous nucleation, as demonstrated in Figure 9.  The WS 2 growth was performed similar to that described in Section 3.1 (i.e., WS 2 grown at Si/SiO 2 substrate). The only difference is the coating of Si/SiO 2 substrate with GO, which is expected to get converted into RGO due to in-situ heating encountered in the HT-zone during the growth process. The synthesis process is schematically illustrated in Figure 8a-c. Figure 8d-g shows the Optical images of the WS 2 grown on the GO coated Si/SiO 2 , where in Figure 8d clear interface is shown between WS 2 film and the Si substrate. Figure 8h-k shows the FESEM micrographs at different magnification. In this approach, sufficient amount of sulfur vapor is taken into consideration for the reaction, similar to the first approach. Here, it is noteworthy to say that the sp 2 -carbon nanostructures in the RGO matrix fulfill three-pronged objectives: (i) they improve the surface adhesion, thereby enhancing local adsorption of WO 3 molecules; (ii) they reduce WO 3 to yield WO 3-x ; and finally, (iii) they also act as active sites for heterogeneous nucleation, as demonstrated in Figure 9. and 2712 cm −1 , respectively. Further, XRD analysis was performed to confirm the quality of prepared GO. The XRD pattern in Figure 7c shows the peak at 2θ ≈ 11°, which validates the high quality of the GO flakes. The WS2 growth was performed similar to that described in Section 3.1 (i.e., WS2 grown at Si/SiO2 substrate). The only difference is the coating of Si/SiO2 substrate with GO, which is expected to get converted into RGO due to in-situ heating encountered in the HT-zone during the growth process. The synthesis process is schematically illustrated in Figure 8a-c. Figure 8d-g shows the Optical images of the WS2 grown on the GO coated Si/SiO2, where in Figure 8d clear interface is shown between WS2 film and the Si substrate. Figure 8h-k shows the FESEM micrographs at different magnification. In this approach, sufficient amount of sulfur vapor is taken into consideration for the reaction, similar to the first approach. Here, it is noteworthy to say that the sp 2 -carbon nanostructures in the RGO matrix fulfill three-pronged objectives: (i) they improve the surface adhesion, thereby enhancing local adsorption of WO3 molecules; (ii) they reduce WO3 to yield WO3x; and finally, (iii) they also act as active sites for heterogeneous nucleation, as demonstrated in Figure 9.   Upon heating, WO3 vapors reach the surface of the substrate. A portion of these vapors diffuses and reacts with sp 2 -bonded carbon patches of RGO, resulting in the formation of an intermediate phase, WO3-x, along with the release of CO2. Thus, an increase in the concentration of WO3 vapors, results in simultaneous increase in the CO2 and a corresponding decrease in the size of sp 2 -carbon patches. Further, at elevated temperatures, WO3-x molecules gain the tendency to migrate and aggregate around the sp 2 -bonded carbon patches. These carbon patches or 'carbon nanostructures' are expected to act as heterogeneous nucleation sites; where nucleation takes place via sulfurization of adjoining WO3-x molecules. The synthesis kinetics involve heterogeneous and homogeneous nucleation reactions, occurring simultaneously and competing with each other, as the main growth processes [61,62]. Initially, due to a lower surface free-energy barrier, heterogeneous nucleation outpaces the homogeneous nucleation [59]. In addition, the local high concentration of WO3-x facilitates the formation of bilayer nuclei. In case of perfectly overlapping bilayer formation, the nucleation shifts from the bottom layer to the second layer almost instantly. If growth does not stop at this point, few-or multi-layer WS2 flakes are formed. Nonetheless, a complete understanding of growth mechanics warrants further investigations. The growth process can be summarized in the following chemical reactions [61,62]: The use of GO produces substantially different growth characteristics than the isolated islands observed during the first approach. The WS2 islands now coalesce to form nearly continuous films on centimeter scale as is evident in the optical images and SEM micrograph (Figure 8). Apart from the monolayers, multilayered growth is also visible. The multilayered growth is more favored at either the nucleation sites or grain boundaries (the locations shown by pointing arrows). Continuous monoto-few layer growths on centimeter scale are representative of the samples grown via the second approach, where the Si/SiO2 substrate is coated with GO. We have found the size of the individual WS2 crystals less than 15 μm, which is much larger than those reported for h-BN substrates, where a domain size of less than 1 μm is observed [63][64][65]. Although it is smaller than the WS2 grown on pristine Si/SiO2 substrate, its continuous expansion on the entire substrate has strong merits and fulfils the objectives of this work, as well as the requisite need for device applications. Figure 10a shows the Raman spectrum of a WS2 film grown on GO coated substrate. The peaks present at 351.2 and 421.4 cm −1 is relatable with the and A1g Raman modes. Besides, the and A1g modes are found to be slightly red-and blue-shifted, respectively. Our concern is to study the shift in and A1g peaks due to thickness variation in a mixed multi-layered WS2 flake. The increase in number of layers strongly enhances the out-of-plane vibrations, while coulomb interactions tend to decrease the frequency of in- Thus, an increase in the concentration of WO 3 vapors, results in simultaneous increase in the CO 2 and a corresponding decrease in the size of sp 2 -carbon patches. Further, at elevated temperatures, WO 3-x molecules gain the tendency to migrate and aggregate around the sp 2 -bonded carbon patches. These carbon patches or 'carbon nanostructures' are expected to act as heterogeneous nucleation sites; where nucleation takes place via sulfurization of adjoining WO 3-x molecules. The synthesis kinetics involve heterogeneous and homogeneous nucleation reactions, occurring simultaneously and competing with each other, as the main growth processes [61,62]. Initially, due to a lower surface free-energy barrier, heterogeneous nucleation outpaces the homogeneous nucleation [59]. In addition, the local high concentration of WO 3-x facilitates the formation of bilayer nuclei. In case of perfectly overlapping bilayer formation, the nucleation shifts from the bottom layer to the second layer almost instantly. If growth does not stop at this point, few-or multi-layer WS 2 flakes are formed. Nonetheless, a complete understanding of growth mechanics warrants further investigations. The growth process can be summarized in the following chemical reactions [61,62]: The use of GO produces substantially different growth characteristics than the isolated islands observed during the first approach. The WS 2 islands now coalesce to form nearly continuous films on centimeter scale as is evident in the optical images and SEM micrograph (Figure 8). Apart from the monolayers, multilayered growth is also visible. The multilayered growth is more favored at either the nucleation sites or grain boundaries (the locations shown by pointing arrows). Continuous monoto-few layer growths on centimeter scale are representative of the samples grown via the second approach, where the Si/SiO 2 substrate is coated with GO. We have found the size of the individual WS 2 crystals less than 15 µm, which is much larger than those reported for h-BN substrates, where a domain size of less than 1 µm is observed [63][64][65]. Although it is smaller than the WS 2 grown on pristine Si/SiO 2 substrate, its continuous expansion on the entire substrate has strong merits and fulfils the objectives of this work, as well as the requisite need for device applications. Figure 10a shows the Raman spectrum of a WS 2 film grown on GO coated substrate. The peaks present at 351.2 and 421.4 cm −1 is relatable with the E 1 2g and A 1g Raman modes. Besides, the E 1 2g and A 1g modes are found to be slightly red-and blue-shifted, respectively. Our concern is to study the shift in E 1 2g and A 1g peaks due to thickness variation in a mixed multi-layered WS 2 flake. The increase in number of layers strongly enhances the out-ofplane vibrations, while coulomb interactions tend to decrease the frequency of in-plane vibrations, thereby causing monotonous increase in frequency separation between E 1 2g and A 1g peaks [58]. While there is little variation in peak positions, small modifications in the relative intensities between E 1 2g and A 1g modes are clearly observed. Samples synthesized via first approach have a slightly higher E 2g 1 intensity, with the ratio E 1 2g /A 1g of~3.2. Conversely, the A 1g intensity is higher for the second approach, where E 1 2g /A 1g ratio is of 1.25. Moderate changes in the Raman intensity ratio are believed to indicate changes in the number of layers for exfoliated WS 2 [46,66]. We have performed Raman mapping analysis at multiple spots of the same sample to confirm the uniformity of WS 2 across the substrate area as shown in Figure 10b,c. It is evident that there is slight variation in the separation as well as intensity ratio of in-plane and out-plane vibrational peaks as a function of sample spot location; and a summary of the peak position, difference in Raman peak frequency and intensity ratio of the peaks is tabulated in Table 1. The shift in the frequency of 2LA (M) and A 1g modes has been employed as a measure of the number of layers. However, this approach to determine the number of layers is erroneous due to the close proximity of 2LA (M) and E 1 2g phonon modes. It has been found that with increasing the number of layers, the intensity ratio (I 2LA(M) /I A1g ) decreases from 2.2 (for single layer) to 0.47 (for bulk) [64]. At the first instance, it looks the possible growth of mono-to few-layered WS 2 . However, in case of monolayer MoS 2 , such small intensity variations in E 2g 1 and A 1g peaks were correlated to the differences in electronic doping levels or strain as WS 2 and RGO effectively combine to form heterostructures [67]. Generally, TMDCs are highly prone to significant number of chalcogenide vacancies [59,64,65]. Sulphur vacancies exist in WS 2 , thus introducing localized donor states deep inside the bandgap, which can be subsequently filled by environmental impurities to make it doped. As of now, various competitive approaches exist to find out the origin of anomaly in the interpretation of peak intensity ratio and shift in peak position of Raman modes. Nonetheless, the contribution to the intensity changes due to sulfur content in WS 2 films as a dopant, and local strain in between WS 2 and RGO, cannot be ignored [66]. plane vibrations, thereby causing monotonous increase in frequency separation between and A1g peaks [58]. While there is little variation in peak positions, small modifications in the relative intensities between and A1g modes are clearly observed. Samples synthesized via first approach have a slightly higher E2g 1 intensity, with the ratio /A1g of ~3.2. Conversely, the A1g intensity is higher for the second approach, where /A1g ratio is of ~1.25. Moderate changes in the Raman intensity ratio are believed to indicate changes in the number of layers for exfoliated WS2 [46,66]. We have performed Raman mapping analysis at multiple spots of the same sample to confirm the uniformity of WS2 across the substrate area as shown in Figure 10 b,c. It is evident that there is slight variation in the separation as well as intensity ratio of in-plane and out-plane vibrational peaks as a function of sample spot location; and a summary of the peak position, difference in Raman peak frequency and intensity ratio of the peaks is tabulated in Table 1. The shift in the frequency of 2LA (M) and A1g modes has been employed as a measure of the number of layers. However, this approach to determine the number of layers is erroneous due to the close proximity of 2LA (M) and phonon modes. It has been found that with increasing the number of layers, the intensity ratio (I2LA(M)/IA1g) decreases from 2.2 (for single layer) to 0.47 (for bulk) [64]. At the first instance, it looks the possible growth of mono-to fewlayered WS2. However, in case of monolayer MoS2, such small intensity variations in E2g 1 and A1g peaks were correlated to the differences in electronic doping levels or strain as WS2 and RGO effectively combine to form heterostructures [67]. Generally, TMDCs are highly prone to significant number of chalcogenide vacancies [59,64,65]. Sulphur vacancies exist in WS2, thus introducing localized donor states deep inside the bandgap, which can be subsequently filled by environmental impurities to make it doped. As of now, various competitive approaches exist to find out the origin of anomaly in the interpretation of peak intensity ratio and shift in peak position of Raman modes. Nonetheless, the contribution to the intensity changes due to sulfur content in WS2 films as a dopant, and local strain in between WS2 and RGO, cannot be ignored [66].   The PL spectra, shown in Figure 10d, comprise with neutral A-exciton (X A ) at 1.968 eV as also observed in case of WS 2 growth on Si/SiO 2 substrate. A striking difference observed in WS 2 growth on GO-coated Si/SiO 2 substrate, is the appearance of a low intensity peak at 2.03 eV (611 nm), attributable to the weak transition of carriers from disorder induced states within π-π * band structure of RGO [67] (schematically shown in Figure 10e). Our laser excitation visible wavelength does not cover the intense emission peak around 440 nm. A comparative study in terms of PL spectral positions, FWHM and shift in peak positions are put in Table 2. Surface roughness, in particular, the sharp edges, is reported to help in catalyst-free CNT growth [68]. To ascertain the role of surface roughness on the continuous growth of WS 2 , we conducted CVD growth in identical conditions: first, on polished surface, marked with scratches on polished side of silicon wafer; and second, on backside, the unpolished surface of the same wafer. In the polished wafer side, WS 2 debris having no shape and sizes of flakes, are observed on and around the scratch marks. SEM images (Figure 11a-d) show few vertically aligned and flower petal shaped flakes with some horizontal flakes on the scratch marked region having negligible presence of triangular as well as hexagonal WS 2 crystallites. In contrast, we have achieved horizontally grown WS 2 crystallites, majority in hexagonal shape on the unpolished surface (Figure 11e-g). In both cases, thickness of WS 2 flake(s) is multilayered. It may be inferred that surface roughness does also play an important role in the WS 2 growth. The optical profilometer image (surface microscopy) of the GO coated sample surface before WS 2 deposition is shown in Figure 11h. It is to be noted that surface of RGO film is highly rough in nature; and this may be one of the sources of nucleation sites too, in addition to the major contribution from high density sp 2 -bonded patches, to enhance the joining of grain boundaries, resulting in continuous WS 2 film growth over GO coated Si/SiO 2 substrate. Therefore, it may be understood that continuous growth is possible in GO coated substrate, in which sp 2 -patches of carbon purely act as a seed or catalyst that enhance growth origins. sources of nucleation sites too, in addition to the major contribution from high density sp 2bonded patches, to enhance the joining of grain boundaries, resulting in continuous WS2 film growth over GO coated Si/SiO2 substrate. Therefore, it may be understood that continuous growth is possible in GO coated substrate, in which sp 2 -patches of carbon purely act as a seed or catalyst that enhance growth origins. The present work has shown the successful growth of mono to few layer WS2 film of square-centimeter size and triangular shape on graphene-oxide-coated Si/SiO2 substrate using CVD technique. The real issue addressed in this work is the large area coverage of the WS2 continuous film; however, many of the researchers have also tried to resolve this issue by using various approaches such as atomic layer deposition, pulse laser deposition, and metal/metal oxide thin film or noble metal substrates, etc. Although these strategies yield large area continuous TMDCs, these techniques also suffer from high costs, increased complexity, and sometimes poor quality. Further, to be useful, WS2 grown over inert metal needs etching via acidic or basic (e.g., concentrated HF, KOH and NaOH) solutions, before transferring to Si or flexible substrates. Usually, the transfer processes not only end up in damaging the grown material (WS2), but produce substrate residues and negative environment impact, as well. We successfully grew the continuous WS2 film at very low cost by simply coating the Si/SiO2 substrate with graphene oxide (GO) and this technique is highly reproducible. Table 3 enlists the optical properties of WS2 nanoflakes deposited on different substrates. However, HRTEM and Raman spectroscopy confirm the high crystallinity and quality of the grown WS2 flakes. Note that the Raman peak frequencies difference (66-72 cm −1 ) and the PL peak position (2.00-1.92 eV) as well as intensity fluctuates over the WS2 triangles as enlisted in Table 2. The change of frequency and/or width of Raman modes of 2H-type transition-metal dichalcogenides can be used to evaluate the sample quality because they are strongly sensitive to external electrostatic doping (A1g mode) and lattice strain (E2g 1 mode) [69,70]. Wang et al. found that E2g 1 mode exhibits obvious red-shift when increasing strain [70]. In contrast, Chakraborty demonstrated that a softening and broadening of the A1g mode occur with electron doping, whereas the E2g 1 mode remains essentially inert [69]. The present work has shown the successful growth of mono to few layer WS 2 film of square-centimeter size and triangular shape on graphene-oxide-coated Si/SiO 2 substrate using CVD technique. The real issue addressed in this work is the large area coverage of the WS 2 continuous film; however, many of the researchers have also tried to resolve this issue by using various approaches such as atomic layer deposition, pulse laser deposition, and metal/metal oxide thin film or noble metal substrates, etc. Although these strategies yield large area continuous TMDCs, these techniques also suffer from high costs, increased complexity, and sometimes poor quality. Further, to be useful, WS 2 grown over inert metal needs etching via acidic or basic (e.g., concentrated HF, KOH and NaOH) solutions, before transferring to Si or flexible substrates. Usually, the transfer processes not only end up in damaging the grown material (WS 2 ), but produce substrate residues and negative environment impact, as well. We successfully grew the continuous WS 2 film at very low cost by simply coating the Si/SiO 2 substrate with graphene oxide (GO) and this technique is highly reproducible. Table 3 enlists the optical properties of WS 2 nanoflakes deposited on different substrates. However, HRTEM and Raman spectroscopy confirm the high crystallinity and quality of the grown WS 2 flakes. Note that the Raman peak frequencies difference (66-72 cm −1 ) and the PL peak position (2.00-1.92 eV) as well as intensity fluctuates over the WS 2 triangles as enlisted in Table 2. The change of frequency and/or width of Raman modes of 2H-type transition-metal dichalcogenides can be used to evaluate the sample quality because they are strongly sensitive to external electrostatic doping (A 1g mode) and lattice strain (E 2g 1 mode) [69,70]. Wang et al. found that E 2g 1 mode exhibits obvious red-shift when increasing strain [70]. In contrast, Chakraborty demonstrated that a softening and broadening of the A 1g mode occur with electron doping, whereas the E 2g 1 mode remains essentially inert [69]. To check the continuity of the grown film, I-V characteristics were monitored in the voltage range of −12 V to +12 V using Keithley SCS 4200 for variable channel length (from 0.5 to 5.0 mm). The inset of Figure 12a shows the digital photographs of the samples, where the electrical probes are set at different channel lengths. The observed I-V curves indicate that flakes grown over the GO modified substrate are interconnecting and form a continuous network. Figure 12b demonstrates the possible use of developed film as photodetector where the change in photocurrent under illumination of laser wavelength of 635 nm, keeping the laser power constant at 40 mW, is shown. The cyclic response is highly reproducible with negligible drift. The photoresponsivity and specific detectivity are found to be 0.9 µA W −1 and 0.15 × 10 5 jones.  To check the continuity of the grown film, I-V characteristics were monitored in the voltage range of −12 V to +12 V using Keithley SCS 4200 for variable channel length (from 0.5 to 5.0 mm). The inset of Figure 12a shows the digital photographs of the samples, where the electrical probes are set at different channel lengths. The observed I-V curves indicate that flakes grown over the GO modified substrate are interconnecting and form a continuous network. Figure 12b demonstrates the possible use of developed film as photodetector where the change in photocurrent under illumination of laser wavelength of 635 nm, keeping the laser power constant at 40 mW, is shown. The cyclic response is highly reproducible with negligible drift. The photoresponsivity and specific detectivity are found to be 0.9 μA W −1 and 0.15 × 10 5 jones.

Conclusions
Efforts have been made to understand the growth mechanics of large domain size of mono-to-few layer growth of WS2. Nature of substrate is utilized to understand the intricate mechanism behind continuous film growth. Large area (domain size of around 60 μm), mono-to-few layer WS2 crystals of triangular shape, were grown on pristine Si/SiO2 and GO-coated Si/SiO2 substrates, using CVD technique with a two-zone furnace. For GO coated Si/SiO2 substrate, uniform mono-to-few layer growth of WS2 crystals is achieved, although the size of WS2 crystals is smaller than those grown on pristine Si/SiO2 substrate. In case of growth on Si/SiO2 substrate, monolayer WS2 grain boundaries are in isolation, and highly scattered on the substrate surface. Further tests were conducted to examine the role of surface roughness, as reported in the past for catalyst-free CNT growth on pristine silicon surface. We conducted CVD growth in identical conditions, one on polished surface, marked with scratches; and the other one, on the unpolished surface of the same wafer. Strikingly, WS2 growth exclusively on rough surface in both the cases, confirmed

Conclusions
Efforts have been made to understand the growth mechanics of large domain size of mono-to-few layer growth of WS 2 . Nature of substrate is utilized to understand the intricate mechanism behind continuous film growth. Large area (domain size of around 60 µm), mono-to-few layer WS 2 crystals of triangular shape, were grown on pristine Si/SiO 2 and GO-coated Si/SiO 2 substrates, using CVD technique with a two-zone furnace. For GO coated Si/SiO 2 substrate, uniform mono-to-few layer growth of WS 2 crystals is achieved, although the size of WS 2 crystals is smaller than those grown on pristine Si/SiO 2 substrate. In case of growth on Si/SiO 2 substrate, monolayer WS 2 grain boundaries are in isolation, and highly scattered on the substrate surface. Further tests were conducted to examine the role of surface roughness, as reported in the past for catalyst-free CNT growth on pristine silicon surface. We conducted CVD growth in identical conditions, one on polished surface, marked with scratches; and the other one, on the unpolished surface of the same wafer. Strikingly, WS 2 growth exclusively on rough surface in both the cases, confirmed that surface roughness does play an important role. In addition, high density sp 2 -bonded patches enhance the joining of grain boundaries, to enable continuous WS 2 film growth over GO coated Si/SiO 2 substrate. It is concluded that the combination of high-temperature CVD with reduced graphene oxide surface is supposed to be a favorable condition for TMDCs growth with full coverage of the substrate. The idea of GO coated Si/SiO 2 substrate may be employed for mass production of identical devices, and therefore bids for commercial scale development and can be extended to other TMDCs.

Data Availability Statement:
The data presented in this study are available on request from the corresponding author.