Modification of Vapor Phase Concentrations in MoS2 Growth Using a

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Modification of Vapor Phase Concentrations in MoS Growth Using a NiO Foam Barrier 2

Yee-Fun Lim, Kumar Priyadarshi, Fabio Bussolotti, Pranjal Kumar Gogoi, Xiaoyang Cui, Ming Yang, Jisheng Pan, Shi Wun Tong, Shijie Wang, Stephen J. Pennycook, Kuan Eng Johnson Goh, Andrew T. S. Wee, Swee Liang Wong, and Dongzhi Chi ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.7b07682 • Publication Date (Web): 16 Jan 2018 Downloaded from http://pubs.acs.org on January 17, 2018

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Modification of Vapor Phase Concentrations in MoS2 Growth Using a NiO Foam Barrier Yee-Fun Lim1†, Kumar Priyadarshi1, 3†, Fabio Bussolotti1, Pranjal Kumar Gogoi2, Xiaoyang Cui1, Ming Yang1, Jisheng Pan1, Shi Wun Tong1, Shijie Wang1, Stephen J. Pennycook4, Kuan Eng Johnson Goh1, 2, Andrew T. S. Wee2, Swee Liang Wong1, 2* and Dongzhi Chi1* 1

Institute of Materials Research and Engineering, Agency for Science Technology and Research, 2 Fusionopolis Way, #08-03 Innovis, Singapore 138634

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Department of Physics, National University of Singapore, 2 Science Drive 3, Singapore 117542 3

Indian Institute of Science Education and Research, Dr.Homi Bhabha Road, Pashan Pune 411 008

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Department of Materials Science & Engineering, National University of Singapore, 9 Engineering Drive 1, Singapore 117575

*S. L. W. Email: [email protected]; *D. C. Email: [email protected] KEYWORDS: chemical vapor deposition, molybdenum disulfide, angle-resolved photoemission spectroscopy, Raman spectroscopy, electronic transport measurement

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ABSTRACT Single-layer molybdenum disulfide (MoS2) has attracted significant attention due to its electronic and physical properties with much effort invested towards obtaining large area high quality monolayer MoS2 films. In this work, we demonstrate a reactive barrier-based approach to achieve growth of highly homogeneous single layer MoS2 on sapphire through the use of a nickel oxide foam barrier during chemical vapor deposition. Due to the reactivity of the NiO barrier with MoO3, the concentration of precursors reaching the substrate and thus nucleation density is effectively reduced, allowing grain sizes of up to 170 µm and continuous monolayers on the centimeter length scale being obtained. The quality of the monolayer is further revealed by angle-resolved photoemission spectroscopy measurement through observation of a very well-resolved electronic band structure and spin orbit splitting of the bands at room temperature with only two major domain orientations, indicating the successful growth of a highly crystalline and well-oriented MoS2 monolayer.

Two-dimensional transition metal dichalcogenides (TMDCs) have been the subject of much research interest of late due to their physical and electronic properties, enabling a wide range of applications such as photodevices,1-3 field-effect transistors,4-6 biomedical applications7,8 and also as a potential material for valleytronics.9-11 MoS2, due to its indirect to direct electronic band gap transition when thinned to a monolayer,12,13 provides an avenue for thickness-based variation in its optoelectronic properties and hence control over the number of layers during synthesis is important. In addition, for applications in the aforementioned areas, both film size as well as quality is essential to the final device performance and ease of integration. Depending on the coverage and quality required, MoS2 can be synthesized using different methods, such as topdown mechanical/chemical exfoliation techniques4,14 or through bottom-up growth by chemical


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vapor deposition (CVD).15-19 Each method has its inherent advantages but CVD has shown great potential in growing MoS2 monolayers having large grain sizes with controllable layer thickness and film coverage. Currently, many efforts have concentrated on growth parameters such as operating pressure, temperature and concentrations of the chemical precursors to achieve the desired film quality and control over the deposited number of layers. Most deposition geometries involve the direct placement of the growth substrate either above the MoO3 precursor or adjacent to it.17,19 Other reaction geometries use a second quartz tube to prevent the chemical precursors from directly facing each other and thus altering the concentration gradient at the growth substrate.20 Additional methods of growth include seed promoters,21,22 pre-deposition plasma treatment of substrates23,24 and also introduction of gases such as oxygen25 during the growth process. Employment of solution based precursors26,27 to achieve a uniform growth has been reported as well. Use of precursor sources such as Mo-based hexacarbonyls15,28,29 located external to the main heating chamber has also been used to deposit monolayers on the wafer scale but are toxic and additional precursors are required to increase the grain sizes beyond a few µm. Reactor design in improving the chemical precursor flow and therefore nucleation density has also been intensively looked into. Physical parameters involved during CVD such as chemical precursor-substrate distances,30-32 multiple furnace temperature zones33 as well as segregation of precursors34,35 prior to reaction at the substrate has been reported to affect the homogeneity and coverage of the MoS2 film deposited. Orientation and placement of the substrate35-37 with respect to the precursor powder sources have also been shown to influence the grain sizes and homogeneity of the final deposited monolayer. Though much progress has been made in achievable growth coverage15 as well as quality of CVD grown MoS2 monolayer on sapphire38-40


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through tuning of process parameters, reactor geometry and choice of chemical precursors, not much work has been done in regulating the flux and concentration gradient of the chemical precursors using reactive barriers. In this article, we report the development of a reactor geometry to enable centimeter scale CVD growth of well-oriented monolayer MoS2 on sapphire with large grain sizes through addition of a nickel oxide foam as a barrier for limiting the deposition rate of chemical precursors, enabling an even distribution of reactants on the substrate. As NiO is known to react with MoO3, not only does it function as a physical barrier but also as a chemically reactive trap of MoO3, allowing greater regulation of the precursor flux during the growth process compared to an inert physical barrier. Raman and photoluminescence (PL) spectroscopy was used to identify the number of layers and its homogeneity while scanning transmission electronic microscopy (STEM) confirmed its crystallinity and atomic structure. The intrinsic electronic quality of the MoS2 monolayer and its relative orientations in the continuous film is then characterized through measuring its electronic band structure by angle-resolved photoemission spectroscopy (ARPES). RESULTS AND DISCUSSION CVD deposition was carried out at 775°C over a duration of 10 minutes in a single zone 3 inch quartz tube furnace at atmospheric pressure (additional details can be found in the Supplementary Material). A top view of the arrangement of substrate, crucible and NiO foam is shown in Figure 1a. The schematic describing the relative arrangement of the different precursors and substrate is described in Figure 1b. 10 mg of MoO3 (99.8%, Sigma Aldrich) is placed in a single open-end crucible with a piece of commercially purchased post-annealed nickel oxide foam (1 mm thickness with 400 µm average pore size) placed directly above the MoO3 powder. The substrate is then placed on top of the foam, supported by pieces of ceramic.


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The substrate used is a commercially available c-plane (0001) sapphire (Al2O3) substrate (Namiki) and N2 was employed as the carrier gas. 300 mg of Sulfur powder (99.8%, Sigma Aldrich) is placed upstream, at the periphery of the furnace heating coils, and is heated up to approximately 200°C during MoS2 growth. Prior to deposition, the sapphire is annealed ex-situ at 1050°C in air for an hour before being placed in the furnace to generate flat and well-defined surface terraces41 for MoS2 deposition. The NiO foam is formed by also annealing a Ni foam in air, at the same temperature of 1050°C to minimize any possible outgassing during the growth process. The foam serves as a physical barrier to prevent direct deposition of MoO3 onto the surface of the substrate, thereby reducing the nucleation rate of chemical precursors, leading to a slower growth rate and hence monolayer MoS2 deposition with greater homogeneity. Furthermore, NiO has been known to react with Mo oxides to form a stable NiMoO4 compound42-44 and this could greater enhance its efficacy as a barrier as the chemical reaction would serve to limit even more Mo chemical precursors from reaching the substrate. In addition, NiO was chosen as it has a very high melting point (>1900°C) and therefore would not melt or vaporize under our growth conditions (775°C). Thus, it would minimize its potential as a contamination source while retaining the reactivity of the barrier throughout the entire growth process. Shown in Figure 1c is the as-deposited MoS2 film while the inset describes the photographic image of the as-deposited MoS2 where the monolayer area (yellow region) is seen to clearly extend over several centimeters. From the micrographs, we note that the continuous layer show minimal nucleation of thicker MoS2 islands or the presence of unreacted Mo oxide particles. Under larger magnifications (Figure 1d) close to the edge of the monolayer MoS2 film, isolated monolayer MoS2 triangles of up to 170 µm in length can be observed, indicating that relatively


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large grain sizes are also achievable using this method. We note that this value is an estimate of the upper limit of the domain sizes deposited. Atomic force microscopy (AFM) measurements (Figure S3, Supplementary Information) of grain boundaries within the continuous monolayer film which was aged in ambient humidity (>50%) yields an estimated range of grain sizes from 6 µm to 55 µm. We observe that the majority of MoS2 growth takes place on the top surface of the sapphire substrate even though it is placed directly above the MoO3 precursor, albeit with a NiO foam between the two. A much smaller coverage of MoS2 ( 95% of the counts) showing characteristic Raman peak separation


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for monolayer MoS2, while noting the formation of a small MoS2 bilayer triangle that shows a larger separation of 22 cm-1. This indicates formation of a fairly homogenous monolayer MoS2. The crystallinity and atomic structure of the grown MoS2 sample is also confirmed by Zcontrast high angle annular dark field (HAADF) STEM images acquired at 60 kV after transfer to a TEM grid. The representative raw and deconvolved45 STEM images (Fig. 2e and f) show good quality of the grown MoS2 crystal at atomic levels with minimum defects. Particularly, the intensity profiles of the Mo and S sites (Fig. 2h) matches well when compared with a simulated STEM image (Fig. 1g) obtained using the same parameters as the experimental imaging conditions. As can be seen in Fig. 2e and f, there is no evidence of molybdenum vacancies in the image. The comparison of the simulated and experimental intensity profiles of the STEM images allows the identification of monosulphur vacancy sites (See Supplementary Fig. S5). We further quantify the sulphur deficiency to be ~4.45% from a representative larger area STEM image (See Supplementary Fig. S6) based on the concentration of these monosulphur vacancies. To further confirm the intrinsic electronic quality of the CVD grown MoS2, ARPES has also been performed at room temperature with a 21.2 eV He Iα source of 800 µm spot size and is shown in Figure 3. Note that the measured sample was first transferred from the insulating growth substrate, sapphire, onto a conducting HOPG substrate before ARPES measurements are conducted. Unlike previous work which were primarily done on exfoliated flakes or MoS2 grown on graphitic surfaces, the resolution of the bands on our deposited are much higher, with a clear spin-orbit splitting of the valence bands which was not observed in previous work. The band structure of the monolayer MoS2 is shown in Figure 3c, with the spin split bands having an energy separation of ≈ 0.18 ± 0.02 eV at valley point K1 due to spin orbit coupling effect,


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agreeing with theoretical and experimental reports. This ascertains the crystallinity and quality of the monolayer that is deposited using our process. We also note the absence of any additional bands in the measured energy range at Γ (at least in an 800 µm probe area), indicating that the number of areas with > 1 MoS2 layer is negligible and that a homogenous growth is achieved using our improved method. Though previous reports have shown ARPES results on monolayer MoS2 using micrometer spot sizes,46-48 our samples yield the best resolved spin-split bands and band dispersion at room temperature to date. Considering that the beam size involved is approximately 800 micrometers in diameter, our CVD grown samples represent a high degree of crystallinity over a substantially large area. Note also that the band dispersions along Γ – M1 and Γ – K1 can be observed simultaneously due to preferred orientations of the MoS2 monolayer grown on sapphire which can be further probed through a constant energy plot. As the electronic band dispersion is dependent on the individual crystal grain orientation present in the monolayer, a constant energy plot is used to reveal the number of orientations present in the continuous MoS2 layer and is described in Figure 3a. The plot yields the relative orientation of the different domains present in the MoS2 layer as it is taken close to the valence band maximum of MoS2, which is located at the K point in the respective Brillouin zones of each domain that differ according to their individual orientations. Our measurements reveal a preferential orientation adopted by these domains, as evident from the high photoemission intensities being centered about specific points, labeled K1 and K2 in Figure 3a. K1 and K2 are equidistant (1.33 ± 0.01 Å-1) from the Γ point while the ΓK1 and ΓK2 vectors are separated by an angle of ± 30°. Note that there are still trace intensities distributed along arbitrary directions but the majority of the measured dispersion is centered about these points. An intensity profile


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(Figure 3b) taken along the directions (indicated in Figure 3a), ΓA, ΓK1 and ΓK2 further corroborates our conclusion as the photoemission intensities at other points at angles that are not at ± 30° from the K point are much lower. Constant energy plots taken from other randomly selected spots about 1mm apart of the transferred film on HOPG also describes a similar situation (Fig. S7). This indirectly alludes to the homogeneity of the film. Our observation matches that of Ji et al.,40 in which two distinct orientations of deposited MoS2 on sapphire at 0° and 30° are preferred due to their relatively lower binding energies with sapphire while a minority of the monolayers are distributed at arbitrary angles. Measuring the relative spread of crystal orientations in Figure 1d yields a similar observation (Figure S8, Supplementary Information), with majority of the crystals oriented similarly and some deviating at multiples of 30 degrees away. Although previous studies39,40,49,50 showing similar preferential crystal orientations have been reported on different substrates, epitaxy is observed either through statistical counting of isolated MoS2 crystals prior to their merging and formation of a continuous layer or selected area electron diffraction TEM measurements of smaller areas (couple of µm). Ours present a definitive confirmation of such an orientation preference over a large area of a continuous MoS2 monolayer (≈ 1 mm2) through verification of a well-resolved electronic band dispersion. Our results are different to that reported by Dumcenco et al.,39 in which almost only a single grain orientation is observed over a centimeter square of MoS2. This is likely due to the higher growth temperatures employed (775°C compared to 700°C), resulting in provision of the necessary thermal energy for the MoS2 domains to adopt other orientations. We attributed the observed dominant crystal orientations to the annealed sapphire surface which provides a template for preferred alignment during growth. Nevertheless, having dominant preferences


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instead of the commonly observed random orientations in other reports would enable higher quality electronic transport as there will be less angular mismatch between domain boundaries. Furthermore, our observed grain sizes are larger than that described in previous reports.39,49 Hence, not only does our method provides large area growth (up to a few cm2), it also allows semi-epitaxial growth of MoS2 on a hexagonally symmetric sapphire (0001) surface. To validate the effectiveness of the annealed NiO foam in improving the quality of the deposited film, a solid SiO2/Si wafer barrier of the same size is used in place of the NiO foam and growth is carried out under the same temperature profile, gas flow and chemical quantities (Figure S4b). When a solid barrier was used instead of a porous foam, even though a continuous monolayer can be achieved, there is a much higher nucleation density with smaller grain sizes (inset of Figure S4b) prior to merging of a continuous monolayer. Such an observation is attributed to the limited effectiveness of the SiO2/Si barrier to reduce the deposition rate of Mo chemical precursors onto the sapphire surface. Even though the SiO2/Si barrier can serve as a physical obstruction, the reduced MoO3-x precursors are able to desorb from its surface and reach the sapphire at these growth temperatures. Thus, this results in a higher flux density of reduced MoO3-x precursors that reach the sapphire surface, resulting in much higher nucleation densities and small grain sizes. To confirm that the presence of NiO does not result in any elemental contamination of our MoS2 monolayers, XPS is performed ex-situ without any prior annealing of the sample to check for the presence of Ni and we see no presence of Ni related core level peaks (not shown), confirming that our method is suitable for CVD synthesis of MoS2 and does not introduce additional contaminants to the product. Hence, our hypothesis is as such: Due to the presence of a NiO foam, the Mo-based precursors are well-adsorbed on the NiO barrier and do not desorb readily at high temperatures. As a result,


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the flux density of reactants is reduced and in-turn, the nucleation density is lowered, allowing development of larger grain sizes together with formation of a well-oriented continuous monolayer. In order to elucidate the mechanism behind the effectiveness of the NiO foam, XPS and X-ray diffraction measurements were carried out to identify the compound formed on the foam after the growth process. Figure 4a -c describes the evolution in Mo 3d and Ni 2p XPS peaks before and after being used as a barrier while Figure 4d shows the XRD spectrum of the NiO foam post growth. Prior to the CVD process, characteristic Ni 2p3/2 and 2p1/2 peaks belonging to Ni2+ present in NiO can be seen in Figure 4a at binding energies of 853.5 and 872.3 ± 0.7 eV, with a representative spin-orbit splitting of 17.8 eV, as reported by others.51,52 The associated multiplet splitting (855.6 eV, 873.1 eV) due to single unpaired electron interactions between orbitals and also the satellite peaks can be clearly observed. After the NiO foam has been used as a barrier in the CVD process, the Mo 3d related peaks can now be observed in Figure 4c and there is a distinct shift in the binding energies of the Ni 2p peaks (Figure 4b). As seen in Figure 4b, the Ni 2p3/2 and 2p1/2 peaks now reside at binding energies of 856.5 and 874.3 ± 0.5 eV respectively while possessing the same extent of spin-orbit splitting as before together with the satellite peaks. In addition, the multiplet split peaks are now absent. Such a shift in binding energy and absence of multiplet split peaks are well-documented for Ni 2p peaks related to NiMoO442,53,54 and occurs due to the change in chemical environment of the Ni2+ species. Furthermore, the Mo 3d5/2 and Mo 3d3/2 core level peaks located at binding energies of 232.8 and 235.9 ± 0.5 eV with a spin orbit split of 3.1 eV is indicative of the presence of Mo in the +6 oxidation state, which was similarly reported for NiMoO4.55 The other doublets fitted within the Mo 3d envelope with the 3d5/2 peak located at binding energies of 229.4 and 230.9 ± 0.5 eV correspond to respective Mo oxidations states of +4 and +5 which have lower binding


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energies. Such oxidation states could be present due to reduced MoO2 and MoO3-x species remaining on the barrier after the growth process. X-ray diffraction (XRD) data was used to further identify the oxide phases produced in the NiO foam during the CVD growth and are shown in Figure 4d. Unsurprising, the XRD spectrum of the oxidized Ni foam was dominated by peaks corresponding to Ni and NiO (see Figure S9a in Supporting Information). There were also a number of smaller peaks, as plotted in Figure 4d. These peaks were attributed to NiMoO4, MoO3 and MoO2 by comparison with standard database values. The presence of these oxide phases are consistent with the XPS data, and further supports the proposed mechanism. Raman spectroscopy of the same NiO foam after use as a barrier bears signature vibrational modes of the aforementioned NiMoO4 chemical species, confirming our observation55 (Figure S9b). Thus, this validates our initial hypothesis: Not only does the NiO foam act as a physical barrier, it further lowers the Mo related chemical precursors reaching the substrate surface by reacting with them, forming NiMoO4 in the process. In addition, as NiMoO4 has been demonstrated as a catalyst in other works,56-58 its presence could potentially serve to facilitate the ease in formation of MoS2. These combination of factors allows NiO foam to function as a highly effective barrier to facilitate large area CVD growth of well-aligned MoS2 monolayers through control of the chemical precursor diffusion as well as reducing the flux density reaching the substrate surface. Top-gate field-effect transistors (FET) were fabricated to investigate the electronic properties of as-grown MoS2 film. Ti/Au (10 nm/ 20 nm) source and drain electrodes were patterned onto as deposited MoS2 on sapphire by standard photolithographic techniques. After photoresist removal, the sample was annealed at 200 °C in H2/Ar atmosphere to remove any remaining resist residue,4 following which 35 nm of Al2O3 high-κ dielectric was deposited by Atomic Layer


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Deposition (ALD) in the presence of a pre-deposited Al seed layer.59 Finally, patterned top gate electrodes were deposited onto the dielectric. A schematic of the FET device structure is depicted in Figure 5a, while Figure 5b shows an optical microscope image of the top view of the completed device. Our device channel lengths range from 3 - 15 µm, with a common channel width of 40 µm. MoS2 FET mobility was calculated using the expression µ = [dID/dVG] × [L/(WCOXVD)], where ID, VD and VG are the drain current, drain voltage and gate voltage, L is the channel length, W = 40 µm is the channel width, and COX is the oxide dielectric capacitance per unit area. We obtained a value of 1.9 × 10-3 F/m2 for the capacitance using the formula COX = ε0κ/d, where d = 35 nm is the thickness of the dielectric, ε0 = 8.854 × 10-12 F/m is the permittivity of free space and κ is the oxide dielectric constant, for which we use a value of κ~ 7.5 that is typical for ALD grown Al2O3.60-62 Our FET devices show typical mobilities in the range of µ = 0.5 – 2.0 cm2/Vs, with our best device giving µ = 2.5 cm2/Vs. ID-VG and ID-VD transfer curves of this device are plotted in Figures 5c and 5d respectively. Our values compare quite favorably to those reported in the literature for monolayer MoS2 with typical mobilities of 0.01 – 4.0 cm2/Vs.6,63-67 This is indicative of the good quality of our MoS2 film. The mobility of exfoliated MoS2 flakes has been reported to be as high as 200 cm2/Vs,2 which may be attributed to their higher crystallinity as the conducting channel is through a single flake as well as the presence of a high k HfO2 dielectric. Devices on a continous CVD grown MoS2 layer exhibit lower mobilities due to possible grain boundaries in the channel unlike films with single orientation.39 Our reported carrier mobilities here should be taken as a lower limit of its intrinsic carrier mobility as there is still potential for improvement, perhaps through optimization of ALD oxide growth,65 using bottom gated configuration5 or choice of dielectric material4 (HfO2 instead of Al2O3) while the contact resistance can also be reduced


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futher. Nonetheless, our transport measurements indicate that a semiconducting MoS2 monolayer film has been successfully deposited. CONCLUSION In summary, we have demonstrated a simple method of achieving large area uniform monolayer MoS2 through the addition of a NiO foam barrier as well as design of reaction geometry. The achieved monolayer displays signs of preferred orientation on sapphire as measured from constant energy plots by ARPES. Furthermore, the as-grown monolayer MoS2 is also of high crystallinity as observed by TEM and evident by the clear resolution of the ARPES measurement of well-defined electronic band dispersions at room temperature as well as confirmation of a continuous semiconducting monolayer from electronic transport measurements on the as-deposited MoS2. As such, we believe our method presents an avenue for future scaling up of MoS2 production for large-area devices and electronics integration on well-oriented monolayer crystal domains. Furthermore, demonstrating that a chemically reactive barrier is effective in lowering the concentration of precursors provides an additional parameter to consider in future growth setups to improve growth coverage and quality. METHODS Chemical vapor deposition of MoS2. The detailed process profile is as such: 300°C for 10 mins with 200 sccm N2 flow followed by a direct ramp (approx. 70° C min-1) to 775°C. It is then allowed to dwell at 775°C with 10 sccm N2 flow for 10 mins, after which the tube is allowed to naturally cool down (10 sccm N2 flow) to 600°C before the furnace heaters are open to allow rapid cooling to room temperature through ambient exposure while the N2 flow is increased to 200 sccm. All steps described are performed at atmospheric pressure. 10 mg of MoO3 (99.98%, Sigma Aldrich) is placed in a single open-end crucible with a piece of nickel foam (size 3.5 cm x


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3 cm, 1 mm thickness with 400 µm average pore size) placed directly above the MoO3 powder. 300 mg of S (99.998%, Sigma Aldrich) was placed at the periphery of the furnace. The substrate is then placed on top of the foam, supported by pieces of ceramic. The substrate used is a commercially bought c-plane (0001) sapphire (Al2O3) substrate (Namiki Inc). Relative positions of the chemical precursors are shown in Fig. S1. Raman and photoluminescence Spectroscopy. Raman and PL measurements were carried out at room temperature with a 532 nm laser using a WiTEC© system in a backscattering geometry. Step size of mapping achieved in Fig. 2d is 0.2 µm with the plot taken over 100 x 100 points of 20 x 20 µm2 area with an integration time of 0.23s per step. Beam spot size is approximately 500 nm. The peak positions are corrected according to a reference Si Raman peak (520 cm-1). Transfer of MoS2 onto HOPG. First, a hydrophobic polymer support, Polymethyl Methacrylate (PMMA, A4 950K in anisole, MicroChem©) is spin coated (4000 rpm, 60s) onto the MoS2/sapphire sample and cured at 175°C for 3 mins. Following which, the edges of the sample is scratched off to expose the MoS2 – sapphire interface to facilitate infiltration of water. DI water is then introduced to these exposed areas using a dropper. Subsequently, a polydimethylsiloxane stamp (GelPak©) is then used to stamp transfer the PMMA supported MoS2 onto a freshly cleaved HOPG. The PMMA on MoS2is then heated at 150°C for 30s to allow greater adhesion of the layer as well as evaporation of water before being washed off using acetone and isopropanol. For transfer of MoS2 onto TEM grid, PMMA is first coated using the same procedure. It is then floated on a NaOH (3M) etching agent before being fetched by the TEM grid and then dried in atmosphere. The polymer is then washed off using acetone and IPA.


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Angle-Resolved Photoemission Spectroscopy (ARPES). The ARPES measurements were conducted in a custom-designed PREVAC© ARPES system, with a hemispherical electron analyser (SCIENTA DA30L) and monochromatized He Iα(hν=21.218 eV) radiation source (SCIENTA VUV5k). A schematic description of the experimental geometry is shown in Fig. S2. Thanks to the special analyser lens system design, data acquisition is possible in (i) “normal” ARPES mode, where the emission angle θx is defined in the photoemission incidence plane and (ii) “deflection” ARPES mode, where full photoemission cone is accessible [i.e. both θx and θy are simultaneously measured (see Fig. S2)] within a range of ±15o with respect to the surface normal direction. Higher angular limits, to reach the boundaries of the SBZ, were obtained by proper adjustment of sample surface orientation, defined by the Θx and Θy angles of the surface normal with respect with respect to the analyser lens entrance axis z (Fig. S2). In both the ARPES acquisition modes the total energy resolution was set to 20 meV, the angular resolution being better than 0.2o. The binding energy scale was referred to the Fermi level (EF) as measured for a clean gold substrate. Scanning Transmission Electron Microscopy (STEM). The transfer of the MoS2 layer to quantifoil grid was performed in the clean room (Class 1000) in Singapore Synchrotron Light Source (SSLS), NUS. The scanning transmission electron microscopy (STEM) images were aquired using abberation corrected JEOL-ARM200 F microscope also located in SSLS, NUS. High angle annular dark field (HAADF) imaging provides Z-contrast images where the image intenisites depend on the atomic number. The lower accelerating voltage of 60 kV ensures negligible beam damage and the abberation correction is crucial for high quality and high resolution images to reveal the atomic strcuture. The microscope has a resolution of 80 picometers at 200 kV with a demonstrated information transfer of 95 pm at 40 kV. The


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simulations were performed using QSTEM software.68 For simulations, the convergence angle was set at 30 mrad and the HAADF detector angles were from 68 to 280 mrad. These are same as experimental imaging conditions and similarly all other parameters (aberrations) in the simulations are as in experiment. X-ray Photoelectron Spectroscopy (XPS) of NiO barrier after deposition. For Figure 4, XPS measurement was performed ex-situ on NiO foam barriers using Al-Kα source (photon energy hυ=1486.7 eV) and a beam spot size of about 200 µm with an energy resolution of 0.7 eV at 100 eV pass energy with the photoelectrons collected at normal emission angle and the light incident at 60 degrees to the surface normal. Binding energy were calibrated according to the C1s peak position (288.5 eV) as a reference against the standard (284.5 eV). X-Ray diffraction (XRD) of NiO foam after deposition. X-ray diffraction was performed on the NiO foam after CVD growth using a Bruker D8 Advance X-ray diffractometer. The scan was done using Cu K-α radiation (1.54 Å wavelength), at a 2θ range of 10° to 70°, step size of 0.02° and scan time of 0.5 seconds per step. FET device fabrication and measurement. Ti/Au source and drain electrodes were fabricated on as-grown MoS2 film on sapphire using standard photolithographic techniques. AZ 7127 photoresist was spin-coated onto MoS2 at 3000 rpm for 45 seconds, and then baked at 95 °C for 1 minute. 10 nm of Ti and 20 nm of Au were then deposited by E-beam evaporation using a Denton Explorer system. Different device channel lengths of 3 µm, 6 µm, 10 µm, and 15 µm were obtained by varying the gap between the source and drain electrodes. Prior to dielectric deposition, the sample was annealed in a 1:9 H2/Ar gas mixture for 2 hours at 200 °C followed by E-beam evaporation of 1 nm seed layer of Al. Al2O3 high-κ oxide dielectric was then deposited using a Beneq TFS 200 Atomic Layer Deposition (ALD) system. A


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35 nm layer of Al2O3 was obtained from 300 ALD cycles of alternate introduction of trimethylaluminum (TMA) and water as ALD precursors; film thickness was measured using a Woollam VB-400 VASE Ellipsometer. Finally, top gate Ti/Au electrodes were patterned on the Al2O3 dielectric using a similar photolithographic lift-off process as described above. The top gate electrodes defined a channel width of 40 µm. FET device characterization was performed at room temperature under vacuum (~10-4 Torr) inside a probe station. An Agilent HP 4156B Semiconductor Parameter Analyzer was used to source voltage and measure current. A computer with LabVIEW software installed was used for instrument control and data acquisition.

AUTHOR INFORMATION Corresponding Author *Wong S. L. Email: [email protected]; *Chi D. Email: [email protected] Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. †Both Y. L. and K. P. have equal contributions. Y. L. and K. P. performed and designed the experiments. B. F. performed the ARPES experiments and J. P. assisted in the XPS setup used on the NiO foam barrier. P. K. G. carried out TEM measurements. M. Y. did XPS for contamination study with input from S. J. W. T. S. W. carried out AFM measurements. X. C. carried out the transfer of MoS2 monolayers from substrate to target. A. T. S. W., K. E. J. G., and S. J. P. provided technical input for the experiments. S.L.W. designed and supervised the experiments. D.C. supervised the experiments.


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ACKNOWLEDGMENT This research is supported by A*STAR Science and Engineering Research Council Pharos 2D Program (SERC Grant No 152-70-00012). The authors would like to acknowledge the Singapore Synchrotron Light Source (SSLS) for the use of their STEM measurement facility. The laboratory is a National Research Infrastructure under the National Research Foundation Singapore. The authors would also like to thank V. S. F. Lim for his assistance in the device fabrication. S. J. P. and A. T. S. W. are grateful to NUS for support. Y. L., S. L. W., K. P., B. F., M. Y., X. C., K. E. J. G. and D.C. gratefully acknowledge support from the Institute of Materials Research and Engineering (IMRE) under the Agency for Science, Technology, and Research (A*STAR). ASSOCIATED CONTENT Supporting Information. Supplementary material detailing experimental methods and additional experimental data (.doc). These include additional ARPES constant energy plots on CVD grown monolayer MoS2. HAADF STEM intensity profile and quantification of mono sulphur vacancies. Atomic force microscopy measurements as well as statistical counts of relative isolated crystal orientations. XRD and Raman spectroscopy of NiO foam barrier are also included. This material is available free of charge via the Internet at http://pubs.acs.org. The authors declare no competing financial interests. REFERENCES 1. Wang, Q.H.; Kalantar-Zadeh, K.; Kis, A.; Coleman, J.N. and Strano, M.S. Electronics and Optoelectronics of Two-Dimensional Transition Metal Dichalcogenides, Nat. Nanotechnol. 2012, 7, 699-712.


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Figure 1. a) Photograph of the actual layout of crucible, substrate and NiO foam used during CVD growth of MoS2. b) Schematic side view describing relative position in the tube furnace. Optical micrograph c) showing continuous MoS2 monolayer deposited on sapphire with small areas of exposed sapphire surface outlined in blue to highlight the optical contrast. (Inset: photograph of as-deposited MoS2 on sapphire) and d) isolated MoS2 monolayer triangles found at the edge of the continuous film. 139x92mm (300 x 300 DPI)


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Figure 2. a) Raman and b) PL spectra of as-deposited MoS2 monolayer on sapphire. c) Optical image at edge of MoS2 film where mapping is carried out. d) Raman mapping showing difference between A1g and E2g Raman peak positions. Color bar range is indicated for the panel. e) Raw with f) deconvolved and smoothed HAADF STEM images of transferred MoS2 taken at 60 kV. e) and f) share the same scale. The brighter and darker atomic site intensities are due to one molybdenum and two sulfur atoms respectively. g) Simulated HAADF STEM image. h) Line profiles of the intensity of the experimental and simulated images. e), f), and g) share the same scale. The brighter and darker atomic site intensities are due to one molybdenum and two sulfur atoms respectively. 139x218mm (300 x 300 DPI)


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ACS Nano

Figure 3. ARPES results describing the band structure of the filled electronic states of CVD grown MoS2 transferred onto HOPG. a) Constant energy plot taken around the valence band maximum at K (±10 meV of energy integration range) Lines act as a guide to the eye describing the position of the K1 and K2 points. The intensity map was obtained by symmetrizing the experimental data at ky

Modification of Vapor Phase Concentrations in MoS2 Growth Using a

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