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Profound effect of substrate hydroxylation and hydration on electronic and optical properties of monolayer MoS 2
Changxi Zheng, Zai-Quan Xu, Qianhui Zhang, Mark T Edmonds, Kenji Watanabe, Takashi Taniguchi, Qiaoliang Bao, and Michael S Fuhrer Nano Lett., Just Accepted Manuscript • Publication Date (Web): 21 Apr 2015 Downloaded from http://pubs.acs.org on April 22, 2015
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Profound effect of substrate hydroxylation and hydration on electronic and optical properties of monolayer MoS2 Changxi Zheng, †,‡ Zai-Quan Xu,§ Qianhui Zhang,‡ Mark T. Edmonds,† Kenji Watanabe,∥ Takashi Taniguchi,∥ Qiaoliang Bao,§,⊥ and Michael S. Fuhrer*,† †
School of Physics, ‡Department of Civil Engineering, and §Department of Materials Engineering, Monash University, Clayton, Victoria 3800 Australia ∥
National Institute for Materials Science, 1-1 Namiki, Tsukuba 305-0044, Japan
⊥Institute
of Functional Nano and Soft Materials (FUNSOM), Jiangsu Key Laboratory for
Carbon-Based Functional Materials and Devices, and Collaborative Innovation Center of Suzhou Nano Science and Technology, Soochow University, Suzhou 215123, P. R. China
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ABSTRACT
Atomic force microscopy, Kelvin probe force microscopy, and scanning photoluminescence spectroscopy image the progressive post-growth hydroxylation and hydration of atomically flat Al2O3(0001) under monolayer MoS2, manifested in large work function shifts (100 mV) due to charge transfer (>1013 cm-2) from the substrate, and changes in PL intensity, energy, and peak width. In contrast, trapped water between exfoliated graphene and Al2O3(0001) causes surface potential and doping changes one and two orders of magnitude smaller, respectively, and MoS2 grown on hydrophobic hexagonal boron nitride is unaffected by water exposure.
KEYWORDS Molybdenum disulfide; transition metal dichalcogenides; substrate; doping; photoluminescence; boron nitride
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Hydrophilic oxides such as SiO2 and Al2O3 as substrates have played an essential role in the isolation and study of two-dimensional (2D) crystals with the implicit assumption that the substrate has negligible influence on the electronic properties of the 2D material.1 The recent interest in atomically thin materials beyond graphene necessitates revisiting this assumption and understanding the interaction between other 2D materials and hydrophilic oxides. Monolayer transition-metal dichalcogenides (TMDs) of family (Mo,W)(S,Se)2 are direct band gap semiconductors2 with an unusual two-fold valley degeneracy and large spin-orbit coupling, and are attractive for novel device applications in optoelectronics, valleytronics, and spintronics. Chemical vapour deposition (CVD) is the only technique to date to synthesize large (Mo,W)(S,Se)2 monolayers, with hydrophilic Al2O3 and SiO2, two widely used substrates.3 In addition, oxides are the prevailing dielectric materials for assembling (Mo,W)(S,Se)2 devices.4 However substrate choice and ambient environment have been found to affect the properties of MoS2 and consequently device performance.5-9 MoS2 on polymer substrates8 and h-BN9,
10
shows higher mobility and greatly enhanced p-channel
conduction compared to MoS2 on SiO2, though the underlying mechanisms are still elusive. Here we demonstrate that post-growth hydroxylation and hydration of the oxide substrate, Al2O3(0001), profoundly affect the electronic and optical properties of TMDs. Atomic force microscopy
(AFM),
Kelvin
probe
force
microscopy
(KPFM),
and
scanning
photoluminescence (PL) spectroscopy image the progressive post-growth hydroxylation and hydration of atomically flat Al2O3(0001) under MoS2, manifested in large work function shifts (100 mV), and changes in PL intensity, energy, and peak width. We indicate a large electron doping by the bare Al2O3(0001) substrate, and a reduction upon hydroxylation and a partial recovery of doping upon hydration. In contrast, exfoliated graphene on dry Al2O3(0001) prevents hydration of the substrate in ambient, and trapped water between graphene and Al2O3(0001) causes surface potential and doping changes one and two orders of
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magnitude smaller, respectively than observed for MoS2/Al2O3(0001). MoS2 grown on hydrophobic hexagonal boron nitride (h-BN) is unaffected by water exposure. Figure 1a shows an optical image of our monolayer MoS2 prepared by CVD on Al2O3(0001) with atomically flat terraces exceeding 1 micron using a novel technique with powder MoS2 and sulfur precursors (see Methods). The MoS2 crystals have lateral extent >200 microns, and the triangular shapes with straight edges indicate single crystals with preferential
molybdenum
edge
termination.11
Figures
1b-c
show
Raman
and
photoluminescence (PL) spectra taken under 532 nm laser excitation immediately (~5 min) ଵ after removal from the furnace. The in-plane (ܧଶ ) and out-of-plane (ܣଵ ) Raman peaks at
382 cm-1 and 402 cm-1, respectively, with 20 cm-1 difference, indicate that the sample is monolayer.12 The PL spectrum (Figure 1c) shows the single sharp excitonic A peak at 668 nm (1.86 eV) of monolayer MoS2. The peak width is 21 nm (61 meV), comparable to freely suspended exfoliated MoS22 and high quality CVD MoS2.11 Figures 1d-f show maps of the integrated PL intensity, peak position, and peak width, respectively. The small linewidth and uniformity of the PL indicates our MoS2 single crystals are of high quality. We now discuss the changes in the topography (measured by AFM), surface potential (measured by KPFM), and PL of our samples upon extended ambient exposure. Figures 2a-c show the topography, phase (mechanical damping of the AFM during topographic scan), and surface potential of a portion of a MoS2 crystal (shown in optical micrograph in the inset of Figure 2a) after exposure to air (~60% relative humidity, RH) for 5 hr. Figure 2a shows the atomically flat terraces of the bare Al2O3(0001) with width > 1 µm and step height 0.2 nm. The monolayer MoS2 largely follows the substrate topography, with large atomically flat regions, however raised regions are observed at the crystal boundaries and along the uphill side of the substrate steps. The height of the raised regions is 0.37 nm, strongly indicative of an ordered monolayer of water molecules intercalated between MoS2 and Al2O3, as observed
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for graphene on mica13 and Al2O3.14 The morphology suggests that intercalation proceeds from the sample boundaries, fastest along the uphill edges of the substrate steps. This behaviour is in stark contrast to graphene/substrate interfaces where (1) water may be trapped by exfoliated graphene but does not intercalate the dry graphene/oxide interface14 (see below) and (2) the enhancement of diffusion at substrate steps has not been seen.15 The phase image in Figure 2b shows contrast between the raised areas of MoS2 and the rest of the crystal, indicating a change in mechanical properties. Figure 2c shows the corresponding surface potential which surprisingly shows not two but three distinct regions. Figures 2d,e show higher resolution topography and surface potential images of the region indicated by the boxes in Figures 2a,c. The surface potential is highest in the center of the unintercalated terraces, surrounded by regions ~80 mV lower. The waterintercalated regions show an intermediate potential, ~20 mV lower than the terrace center. Figure 2f shows line profiles of the topography and surface potential indicating the three regions: pristine MoS2 with high surface potential (unshaded), a region of low surface potential but unmodified height in advance of the water-intercalated region (shaded yellow), and intermediate surface potential in the water-intercalated region (shaded blue). We interpret the observations as follows. MoS2 grows epitaxially with an atomically sharp boundary at the MoS2/α-Al2O3 (0001) interface.16 At the growth temperature of 940 ℃, we expect that the surface is the (1×1) reconstruction, terminated by a single Al plane with reduced coordination.17-20 This surface has a Lewis acid character, acting as an electron acceptor,20, 21 reducing the native electron doping of MoS2 resulting in the highest surface potential. Upon exposure to ambient, water first adsorbs dissociatively (hydroxylation) with hydroxyl groups binding to the Al-rich surface, and then additional water layers are physisorbed (hydration) on the hydroxyl-terminated surface.20 We expect the same sequence to occur at the Al2O3 surface beneath MoS2; thus we observe three regions: the Al-terminated
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Al2O3, and the hydroxylated and hydrated surfaces. The hydroxylation front proceeds in advance of the hydration front, due to the negligibly small barrier to surface diffusion for hydroxyl groups on α-Al2O3 (0001).22 The hydroxylated region shows no detectable height increase but a large electron doping (shift to lower workfunction) of MoS2, consistent with the Lewis base (electron donating) nature of the –OH groups on the hydroxylated surface. Hydroxylation is followed by hydration, with a height increase of about 3.7 Å, consistent with an ice-like water layer. Here, we expect the formation of an electric double layer at the hydroxylated Al2O3 interface, with H+ ions close to the interface and OH- ions close to the MoS2 with reduced electron doping and higher workfunction compared to the hydroxylated but not hydrated Al2O3 substrate.18 On further exposure to ambient (3 days; Figures 2g-i), the monolayer water region grows, occupying the first two terraces on the left, which show high surface potential. The hydroxylated area has expanded to occupy the entire centre of the third terrace, showing low surface potential, while the remaining terraces show some small regions of high surface potential indicating still unhydroxylated regions. After 30 days (RH: ~60%, Figures 2j-l) several terraces are completely hydrated, though the topography and phase images (Figures 2j,k) show that the water-intercalated regions are not uniform, and surface potential variations are still seen across the sample which correlate strongly with the differences in topography and phase, presumably reflecting differences in the structure of the water layer. We now correlate the changes in topography and surface potential with changes in the PL of MoS2. Figures 3a-c show AFM topography and phase, and surface potential of another MoS2 sample on the same substrate, similar to Figures 2j-l taken after similar ambient exposure time of 30 days. Here the surface is fully hydroxylated, and the surface potential (Figure 3c) shows only the hydroxylated and hydrated regions separated by about 40 mV. Figures 3d-f show maps of the PL intensity, peak position, and peak width. The PL is brighter
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for the hydrated regions along the edges of the sample, and in the interior the PL intensity oscillates with the same period as the terrace steps. Figure 3g shows line traces of the data in Figures 3a-f. Oscillations between two distinct regions are clear in all the traces. The hydrated regions (shaded blue) show higher topography and phase; higher surface potential; and brighter, narrower, and blue-shifted PL. The hydroxylated but unhydrated regions show lower topography and phase; lower surface potential; and less intense (by an order of magnitude), broader, and red-shifted PL. We can estimate the degree of charge transfer caused by hydroxylation, based on the change of surface potential. Assuming m* ≈ 0.4 m0 for MoS223 the density of states D(E) = 3.4 × 1014 states/eVcm2. The observed surface potential shift of ~40 mV between hydration and hydroxylation corresponds to a surface charge transfer of 1.3 × 1013 cm-2. This value is similar to the range of surface charge explored in Ref 24, where a PL intensity change of ~4×, and a shift in PL position of ~12 nm were observed, the same order of magnitude, but somewhat smaller (intensity change) and larger (PL position) than the changes reported here. We note however that the changes in PL in our experiment are probably under-resolved because the laser spot size (~500 nm) is comparable to the Al2O3 terrace size. Furthermore, we address whether similar surface potential fluctuations appear in graphene on hydroxylated vs. hydrated Al2O3. We exfoliated monolayer graphene on atomically-flat Al2O3 both in ambient (hydrated) and in a glove box (H2O < 0.1ppm, hydroxylated). Figure 4a shows the topography of graphene exfoliated in ambient. Most of the graphene/Al2O3 region is occupied by water, with dendritic shapes characteristic of diffusion limited dynamics.25 In contrast, Figure 4b shows that no water intercalates the glove box exfoliated graphene/Al2O3, even after 30 days in ambient (RH: ~60%). We next examine the doping effect of trapped water on graphene using KPFM. Figures 4c-d show the topography and surface potential of a small region of water-occupied graphene/Al2O3, and Figure 4e shows
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the line traces of height and surface potential corresponding to the white lines in Figures 4c-d. As indicated, the thickness of trapped water is 0.37 nm, equal to the value of monolayer ice on mica.13 Similarly, this layer is stable and only slightly reduced in area after air exposure over 30 days (RH: ~60%). The surface potential of graphene on Al2O3 increases only ~5 mV with trapped water, an order of magnitude smaller change than observed for MoS2. Raman spectroscopy (Figure 4f) of graphene on Al2O3 with partial water and without water show a broad G peak (17.1 and 15.5 cm-1 respectively) consistent with low doping (Fermi level 1013 cm-2). The details are listed in Table 1. While the imaging of the successive hydroxylation/hydration of the Al2O3 substrate under MoS2 was made possible by the use of atomically flat Al2O3(0001), the hydroxylation and hydration of high-temperature treated oxide substrates on ambient exposure is a generic phenomenon and occurs for e.g. amorphous SiO2. Our results can explain the observation by other groups of differences in PL between MoS2 sample edges and center as due to water intercalation, and not a different edge structure or reactivity. Notably, differences in PL between sample edge and center have also been observed for (Mo,W)(S,Se)2 on SiO228, 29 & Al2O3.30 The observed doping shifts of order 1013 cm-2 are large, exceeding the maximum achievable gate-induced doping in many device geometries. The surface potential inhomogeneity observed here for even the highest quality atomically flat substrates is likely significantly worse for less ordered or amorphous substrates, and is likely a major source of disorder in MoS2 on oxides. This work thus has strong implications for fabrication of MoS2 devices including field effect transistors, and optoelectronic or valleytronic devices, where uniformity of doping and electronic properties is essential. It points out the special advantage of hydrophobic dielectrics such as PMMA8 and BN9,
10
in avoiding disorder caused by inhomogeneous hydroxylation and water
intercalation, and helps to understand the higher mobility,8, 9 greatly reduced hysteresis10 and improved environmental stability for such substrates.
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METHODS Sample preparation. Monolayer MoS2 single crystals were grown by a novel method, sulfur-assisted CVD. The difference between this and other CVD techniques is that the precursors are sulfur powder (Sigma-Aldrich, 84683) and MoS2 powder (Sigma-Aldrich, 69860), instead of the most commonly used MoO3 powder.3 Growth substrates of Al2O3(0001) were cleaned by 10 min acetone and isopropanol sonication.
Substrates
(sapphire and exfoliated h-BN), MoS2 (0.35 g) and sulfur (0.8 g) were loaded into three separated quartz crucibles. The growth was performed in a 1 inch diameter quartz tube flowing ultra-high purity argon at atmospheric pressure. During growth, MoS2 and substrates were located at the centre and downstream of the furnace, respectively. The sulfur powder was loaded upstream of the MoS2 and substrates and independently heated. The furnace and the sulfur were heated to 940 ℃ and 225 ℃ respectively in 30 min. The furnace was cooled to 500 ℃ over 10 min, followed by rapid cooldown to room temperature (furnace opened for maximum cooling rate achievable). The flow rate of argon was kept at 500 sccm until the furnace temperature initially reached 500 ℃. The flow rate was then decreased to 60 sccm during growth, until the furnace temperature returned to 500 ℃, when the flow rate was increased again to 500 sccm. Hexagonal-BN crystals were obtained in solvent growth process under high pressure by using belt type high pressure apparatus.31 In order to reduce impurities of oxygen and carbon, Ba-B-N solvent were employed. The graphene flakes were exfoliated on atomically flat Al2O3(0001) surface in ambient environment (RH: 60%) and glove box (H2O < 0.1 ppm). Before exfoliation, the sapphire pieces were cleaned by acetone and isopropanol sonication for 10 min. The substrates next were annealed at 1300 ℃ for 2h under ambient environment and then allowed to cool down to room temperature naturally. The substrates were further cleaned by oxygen annealing at
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1050 ℃ for 1 hr. Before loading to the main chamber of glove box (H2O < 0.1ppm), sapphire substrates and scotch tape sticking with graphene flakes were cleaned by repeating the cycle of pumping-flushing nitrogen (purity: 99.999%) three times in the load lock. For glove box exfoliation, the substrates were additionally heated on a hot plate at 200 ℃ for 1 hr soon after being loaded into the glove box filled with nitrogen. Optical measurements. Raman/PL measurements were performed using a confocal microscope system (WITec alpha 300R) with a 100× objective with a numerical aperture (NA) of 0.90. A 532 nm laser with spot diameter of ~500 nm was used to excite samples placed on a piezocrystal-controlled scanning stage. The best mapping spatial resolution was ~250 nm. The spectra were collected using a 600 line/mm grating. To avoid sample damage, low laser power (50 µW) and short integration time (0.1 s) were applied during PL mapping. The Raman spectra of both MoS2 and graphene were measured using much longer integration time (2s). The PL intensity images were obtained by summing the PL intensity from 630 nm to 690 nm, and the PL peak position and full width of half maximum (FWHM) were estimated from fits to a single Lorentzian lineshape. The Raman spectral parameters of graphene were obtained by Lorentzian fitting as well. AFM & KPFM measurements. The AFM and KPFM images were obtained using Bruker Dimension Icon SPM through LiftMode, a dual-pass technique. The topography line profile of the sample was achieved in amplitude-modulated tapping mode during the first pass. The second pass traced the acquired topography line at a set lift height to allow surface potential measurement. A chrome/platinum coated silicon tip was used as the (conductive) tip. For MoS2 measurements, the lift height and the AC bias voltage were 20 nm and 2 V, respectively. While for graphene measurements, the parameters were 10 nm and 0.5 V.
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FIGURES Figure 1. Large-area monolayer MoS2 crystals. (a) Optical image of MoS2 monolayer crystals on Al2O3(0001) grown by sulfur-assisted CVD. (b,c) Typical Raman and photoluminescence (PL) spectra taken from the same position on a MoS2 crystal. (d-f) PL mapping image of a portion of a large crystal showing the integrated PL intensity (d), peak position (e), and peak width (f). Scale bar in (d) applies to (e, f). The PL and Raman measurements were done right after removing the samples from the furnace. Figure 2. Atomic force microscopy (AFM) and Kelvin probe force microscopy (KPFM) of water intercalation. (a-c) AFM height and phase, and KPFM surface potential images of atomically flat MoS2 after 5hr air exposure. Inset in (a) shows the optical image of the MoS2 crystal. (d,e) Height image (d) and KPFM image (e) of the square region shown in (a,c). (f) Correlated height and surface potential profiles corresponding to the white lines shown in (d,e). Shading indicates three regions of surface potential corresponding to bare Alterminated Al2O3 (unshaded), hydroxylated Al2O3 (yellow), and hydrated Al2O3 (blue). (g-i) AFM height, phase and surface potential images of the MoS2 crystal after air exposure for 3 days. (j-l) AFM height, phase and surface potential images of the MoS2 crystal after air exposure for 30 days. Scale bar in (a) applies to (b, c, g-l). Figure 3. Correlated atomic force microscopy (AFM) and photoluminescence (PL) imaging. (a-c) AFM height, phase and surface potential images of a MoS2 crystal after air exposure for 30 days. (d-f) PL mapping images of the flake represented in the form of sum intensity (d), peak position (e), and peak width (f). (g) Line traces of the data corresponding to the lines shown in (a-f). Shading indicates two regions of surface potential corresponding to hydroxylated Al2O3 (yellow), and hydrated Al2O3 (blue).
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Figure 4. Graphene on atomically flat Al2O3(0001). (a,b) atomic force microscopy topographs of graphene on Al2O3(0001) exfoliated in ambient (RH: ~60%) and glove box (H2O < 0.1 ppm). (c,d) Detail images of height (c) and surface potential (d) of graphene exfoliated in ambient on an atomically flat sapphire surface. (e) Line traces of height (black dash dot) and surface potential (blue solid dot) corresponding to the white lines shown in panels (d) and (e). (f) Raman spectra of graphene with partial intercalation of water (blue circles) and without water (red circles). Lines are Lorentzian fits to the data. Figure 5. (a,b) Photoluminescence (PL) images showing the integrated PL intensity of a MoS2 crystal grown on exfoliated h-BN immediately post-growth (a) and after ambient exposure for 3 weeks (b). (c) PL spectra corresponding to (a) and (b). (d) Atomic force microscopy (AFM) topograph of the MoS2 crystal on h-BN. (e) AFM height profile following the black line shown in (d). Table 1. Summary of interactions between ambient water and MoS2/Al2O3, graphene/Al2O3 and MoS2/h-BN interfaces.
AUTHOR INFORMATION Corresponding Author * [email protected]
ACKNOWLEDGMENT M.S.F, M.T.E., and C.Z. acknowledge support from ARC (DP150103837 and FL120100038). C.Z. acknowledges support from ARC DECRA (DE140101555). Q.B. acknowledges support from the 863 Program (Grant No. 2013AA031903), the youth 973
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program (2015CB932700), the NSFC grants (Grant No. 51222208, 51290273), ARC DECRA (DE120101569), DP (DP140101501) and Engineering Seed Funding Scheme (2014) at Monash University. This work was performed in part at the Melbourne Centre for Nanofabrication (MCN) in the Victorian Node of the Australian National Fabrication Facility (ANFF). K.W. and T.T. acknowledge support from the Elemental Strategy Initiative conducted by the MEXT, Japan. T.T. acknowledges support from a Grant-in-Aid for Scientific Research on Grant 262480621 and on Innovative Areas “Nano Informatics” (Grant 25106006) from JSPS.
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15. Sutter, P.; Sadowski, J. T.; Sutter, E. A. Journal of the American Chemical Society 2010, 132, (23), 8175-8179. 16. Ma, L.; Nath, D. N.; Lee, E. W.; Lee, C. H.; Yu, M.; Arehart, A.; Rajan, S.; Wu, Y. Applied Physics Letters 2014, 105, (7), 072105. 17. Elam, J. W.; Nelson, C. E.; Cameron, M. A.; Tolbert, M. A.; George, S. M. The Journal of Physical Chemistry B 1998, 102, (36), 7008-7015. 18. Franks, G. V.; Gan, Y. Journal of the American Ceramic Society 2007, 90, (11), 3373-3388. 19. Felice, R. D.; Northrup, J. E. Physical Review B 1999, 60, (24), R16287-R16290. 20. Hass, K. C.; Schneider, W. F.; Curioni, A.; Andreoni, W. Science 1998, 282, (5387), 265-268. 21. Eng, P. J.; Trainor, T. P.; Brown Jr., G. E.; Waychunas, G. A.; Newville, M.; Sutton, S. R.; Rivers, M. L. Science 2000, 288, (5468), 1029-1033. 22. Nelson, C. E.; Elam, J. W.; Cameron, M. A.; Tolbert, M. A.; George, S. M. Surface Science 1998, 416, (3), 341-353. 23. Peelaers, H.; Van de Walle, C. G. Physical Review B 2012, 86, (24), 241401. 24. Mak, K. F.; He, K.; Lee, C.; Lee, G. H.; Hone, J.; Heinz, T. F.; Shan, J. Nat Mater 2013, 12, (3), 207-211. 25. Severin, N.; Lange, P.; Sokolov, I. M.; Rabe, J. P. Nano Letters 2012, 12, (2), 774779.
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26. DasA; PisanaS; ChakrabortyB; PiscanecS; Saha, S. K.; Waghmare, U. V.; Novoselov, K. S.; Krishnamurthy, H. R.; Geim, A. K.; Ferrari, A. C.; Sood, A. K. Nat Nano 2008, 3, (4), 210-215. 27. Yan, J.; Zhang, Y.; Kim, P.; Pinczuk, A. Physical Review Letters 2007, 98, (16), 166802. 28. Ling, X.; Lee, Y.-H.; Lin, Y.; Fang, W.; Yu, L.; Dresselhaus, M. S.; Kong, J. Nano Letters 2014, 14, (2), 464-472. 29. Wang, X.; Gong, Y.; Shi, G.; Chow, W. L.; Keyshar, K.; Ye, G.; Vajtai, R.; Lou, J.; Liu, Z.; Ringe, E.; Tay, B. K.; Ajayan, P. M. ACS Nano 2014, 8, (5), 5125-5131. 30. Zhang, Y.; Zhang, Y.; Ji, Q.; Ju, J.; Yuan, H.; Shi, J.; Gao, T.; Ma, D.; Liu, M.; Chen, Y.; Song, X.; Hwang, H. Y.; Cui, Y.; Liu, Z. ACS Nano 2013, 7, (10), 8963-8971. 31. Taniguchi, T.; Watanabe, K. Journal of Crystal Growth 2007, 303, (2), 525-529.
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Figure 1.
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Figure 2.
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Figure 3.
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Figure 4.
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Figure 5.
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Table 1. Water intercalation Water-induced work from ambient function shifts MoS2/Al2O3 Yes ~80 mV graphene/Al2O3 No* ≤5mV MoS2/h-BN No n/a *water trapped upon exfoliation in ambient Interface
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Charge transfer >1.3 × 1013 cm-2 ≤1.5 × 1011 cm-2 n/a
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Table of Contents Graphic
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Communication
Profound effect of substrate hydroxylation and hydration on electronic and optical properties of monolayer MoS 2
Changxi Zheng, Zai-Quan Xu, Qianhui Zhang, Mark T Edmonds, Kenji Watanabe, Takashi Taniguchi, Qiaoliang Bao, and Michael S Fuhrer Nano Lett., Just Accepted Manuscript • Publication Date (Web): 21 Apr 2015 Downloaded from http://pubs.acs.org on April 22, 2015
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Profound effect of substrate hydroxylation and hydration on electronic and optical properties of monolayer MoS2 Changxi Zheng, †,‡ Zai-Quan Xu,§ Qianhui Zhang,‡ Mark T. Edmonds,† Kenji Watanabe,∥ Takashi Taniguchi,∥ Qiaoliang Bao,§,⊥ and Michael S. Fuhrer*,† †
School of Physics, ‡Department of Civil Engineering, and §Department of Materials Engineering, Monash University, Clayton, Victoria 3800 Australia ∥
National Institute for Materials Science, 1-1 Namiki, Tsukuba 305-0044, Japan
⊥Institute
of Functional Nano and Soft Materials (FUNSOM), Jiangsu Key Laboratory for
Carbon-Based Functional Materials and Devices, and Collaborative Innovation Center of Suzhou Nano Science and Technology, Soochow University, Suzhou 215123, P. R. China
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ABSTRACT
Atomic force microscopy, Kelvin probe force microscopy, and scanning photoluminescence spectroscopy image the progressive post-growth hydroxylation and hydration of atomically flat Al2O3(0001) under monolayer MoS2, manifested in large work function shifts (100 mV) due to charge transfer (>1013 cm-2) from the substrate, and changes in PL intensity, energy, and peak width. In contrast, trapped water between exfoliated graphene and Al2O3(0001) causes surface potential and doping changes one and two orders of magnitude smaller, respectively, and MoS2 grown on hydrophobic hexagonal boron nitride is unaffected by water exposure.
KEYWORDS Molybdenum disulfide; transition metal dichalcogenides; substrate; doping; photoluminescence; boron nitride
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Hydrophilic oxides such as SiO2 and Al2O3 as substrates have played an essential role in the isolation and study of two-dimensional (2D) crystals with the implicit assumption that the substrate has negligible influence on the electronic properties of the 2D material.1 The recent interest in atomically thin materials beyond graphene necessitates revisiting this assumption and understanding the interaction between other 2D materials and hydrophilic oxides. Monolayer transition-metal dichalcogenides (TMDs) of family (Mo,W)(S,Se)2 are direct band gap semiconductors2 with an unusual two-fold valley degeneracy and large spin-orbit coupling, and are attractive for novel device applications in optoelectronics, valleytronics, and spintronics. Chemical vapour deposition (CVD) is the only technique to date to synthesize large (Mo,W)(S,Se)2 monolayers, with hydrophilic Al2O3 and SiO2, two widely used substrates.3 In addition, oxides are the prevailing dielectric materials for assembling (Mo,W)(S,Se)2 devices.4 However substrate choice and ambient environment have been found to affect the properties of MoS2 and consequently device performance.5-9 MoS2 on polymer substrates8 and h-BN9,
10
shows higher mobility and greatly enhanced p-channel
conduction compared to MoS2 on SiO2, though the underlying mechanisms are still elusive. Here we demonstrate that post-growth hydroxylation and hydration of the oxide substrate, Al2O3(0001), profoundly affect the electronic and optical properties of TMDs. Atomic force microscopy
(AFM),
Kelvin
probe
force
microscopy
(KPFM),
and
scanning
photoluminescence (PL) spectroscopy image the progressive post-growth hydroxylation and hydration of atomically flat Al2O3(0001) under MoS2, manifested in large work function shifts (100 mV), and changes in PL intensity, energy, and peak width. We indicate a large electron doping by the bare Al2O3(0001) substrate, and a reduction upon hydroxylation and a partial recovery of doping upon hydration. In contrast, exfoliated graphene on dry Al2O3(0001) prevents hydration of the substrate in ambient, and trapped water between graphene and Al2O3(0001) causes surface potential and doping changes one and two orders of
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magnitude smaller, respectively than observed for MoS2/Al2O3(0001). MoS2 grown on hydrophobic hexagonal boron nitride (h-BN) is unaffected by water exposure. Figure 1a shows an optical image of our monolayer MoS2 prepared by CVD on Al2O3(0001) with atomically flat terraces exceeding 1 micron using a novel technique with powder MoS2 and sulfur precursors (see Methods). The MoS2 crystals have lateral extent >200 microns, and the triangular shapes with straight edges indicate single crystals with preferential
molybdenum
edge
termination.11
Figures
1b-c
show
Raman
and
photoluminescence (PL) spectra taken under 532 nm laser excitation immediately (~5 min) ଵ after removal from the furnace. The in-plane (ܧଶ ) and out-of-plane (ܣଵ ) Raman peaks at
382 cm-1 and 402 cm-1, respectively, with 20 cm-1 difference, indicate that the sample is monolayer.12 The PL spectrum (Figure 1c) shows the single sharp excitonic A peak at 668 nm (1.86 eV) of monolayer MoS2. The peak width is 21 nm (61 meV), comparable to freely suspended exfoliated MoS22 and high quality CVD MoS2.11 Figures 1d-f show maps of the integrated PL intensity, peak position, and peak width, respectively. The small linewidth and uniformity of the PL indicates our MoS2 single crystals are of high quality. We now discuss the changes in the topography (measured by AFM), surface potential (measured by KPFM), and PL of our samples upon extended ambient exposure. Figures 2a-c show the topography, phase (mechanical damping of the AFM during topographic scan), and surface potential of a portion of a MoS2 crystal (shown in optical micrograph in the inset of Figure 2a) after exposure to air (~60% relative humidity, RH) for 5 hr. Figure 2a shows the atomically flat terraces of the bare Al2O3(0001) with width > 1 µm and step height 0.2 nm. The monolayer MoS2 largely follows the substrate topography, with large atomically flat regions, however raised regions are observed at the crystal boundaries and along the uphill side of the substrate steps. The height of the raised regions is 0.37 nm, strongly indicative of an ordered monolayer of water molecules intercalated between MoS2 and Al2O3, as observed
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for graphene on mica13 and Al2O3.14 The morphology suggests that intercalation proceeds from the sample boundaries, fastest along the uphill edges of the substrate steps. This behaviour is in stark contrast to graphene/substrate interfaces where (1) water may be trapped by exfoliated graphene but does not intercalate the dry graphene/oxide interface14 (see below) and (2) the enhancement of diffusion at substrate steps has not been seen.15 The phase image in Figure 2b shows contrast between the raised areas of MoS2 and the rest of the crystal, indicating a change in mechanical properties. Figure 2c shows the corresponding surface potential which surprisingly shows not two but three distinct regions. Figures 2d,e show higher resolution topography and surface potential images of the region indicated by the boxes in Figures 2a,c. The surface potential is highest in the center of the unintercalated terraces, surrounded by regions ~80 mV lower. The waterintercalated regions show an intermediate potential, ~20 mV lower than the terrace center. Figure 2f shows line profiles of the topography and surface potential indicating the three regions: pristine MoS2 with high surface potential (unshaded), a region of low surface potential but unmodified height in advance of the water-intercalated region (shaded yellow), and intermediate surface potential in the water-intercalated region (shaded blue). We interpret the observations as follows. MoS2 grows epitaxially with an atomically sharp boundary at the MoS2/α-Al2O3 (0001) interface.16 At the growth temperature of 940 ℃, we expect that the surface is the (1×1) reconstruction, terminated by a single Al plane with reduced coordination.17-20 This surface has a Lewis acid character, acting as an electron acceptor,20, 21 reducing the native electron doping of MoS2 resulting in the highest surface potential. Upon exposure to ambient, water first adsorbs dissociatively (hydroxylation) with hydroxyl groups binding to the Al-rich surface, and then additional water layers are physisorbed (hydration) on the hydroxyl-terminated surface.20 We expect the same sequence to occur at the Al2O3 surface beneath MoS2; thus we observe three regions: the Al-terminated
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Al2O3, and the hydroxylated and hydrated surfaces. The hydroxylation front proceeds in advance of the hydration front, due to the negligibly small barrier to surface diffusion for hydroxyl groups on α-Al2O3 (0001).22 The hydroxylated region shows no detectable height increase but a large electron doping (shift to lower workfunction) of MoS2, consistent with the Lewis base (electron donating) nature of the –OH groups on the hydroxylated surface. Hydroxylation is followed by hydration, with a height increase of about 3.7 Å, consistent with an ice-like water layer. Here, we expect the formation of an electric double layer at the hydroxylated Al2O3 interface, with H+ ions close to the interface and OH- ions close to the MoS2 with reduced electron doping and higher workfunction compared to the hydroxylated but not hydrated Al2O3 substrate.18 On further exposure to ambient (3 days; Figures 2g-i), the monolayer water region grows, occupying the first two terraces on the left, which show high surface potential. The hydroxylated area has expanded to occupy the entire centre of the third terrace, showing low surface potential, while the remaining terraces show some small regions of high surface potential indicating still unhydroxylated regions. After 30 days (RH: ~60%, Figures 2j-l) several terraces are completely hydrated, though the topography and phase images (Figures 2j,k) show that the water-intercalated regions are not uniform, and surface potential variations are still seen across the sample which correlate strongly with the differences in topography and phase, presumably reflecting differences in the structure of the water layer. We now correlate the changes in topography and surface potential with changes in the PL of MoS2. Figures 3a-c show AFM topography and phase, and surface potential of another MoS2 sample on the same substrate, similar to Figures 2j-l taken after similar ambient exposure time of 30 days. Here the surface is fully hydroxylated, and the surface potential (Figure 3c) shows only the hydroxylated and hydrated regions separated by about 40 mV. Figures 3d-f show maps of the PL intensity, peak position, and peak width. The PL is brighter
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for the hydrated regions along the edges of the sample, and in the interior the PL intensity oscillates with the same period as the terrace steps. Figure 3g shows line traces of the data in Figures 3a-f. Oscillations between two distinct regions are clear in all the traces. The hydrated regions (shaded blue) show higher topography and phase; higher surface potential; and brighter, narrower, and blue-shifted PL. The hydroxylated but unhydrated regions show lower topography and phase; lower surface potential; and less intense (by an order of magnitude), broader, and red-shifted PL. We can estimate the degree of charge transfer caused by hydroxylation, based on the change of surface potential. Assuming m* ≈ 0.4 m0 for MoS223 the density of states D(E) = 3.4 × 1014 states/eVcm2. The observed surface potential shift of ~40 mV between hydration and hydroxylation corresponds to a surface charge transfer of 1.3 × 1013 cm-2. This value is similar to the range of surface charge explored in Ref 24, where a PL intensity change of ~4×, and a shift in PL position of ~12 nm were observed, the same order of magnitude, but somewhat smaller (intensity change) and larger (PL position) than the changes reported here. We note however that the changes in PL in our experiment are probably under-resolved because the laser spot size (~500 nm) is comparable to the Al2O3 terrace size. Furthermore, we address whether similar surface potential fluctuations appear in graphene on hydroxylated vs. hydrated Al2O3. We exfoliated monolayer graphene on atomically-flat Al2O3 both in ambient (hydrated) and in a glove box (H2O < 0.1ppm, hydroxylated). Figure 4a shows the topography of graphene exfoliated in ambient. Most of the graphene/Al2O3 region is occupied by water, with dendritic shapes characteristic of diffusion limited dynamics.25 In contrast, Figure 4b shows that no water intercalates the glove box exfoliated graphene/Al2O3, even after 30 days in ambient (RH: ~60%). We next examine the doping effect of trapped water on graphene using KPFM. Figures 4c-d show the topography and surface potential of a small region of water-occupied graphene/Al2O3, and Figure 4e shows
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the line traces of height and surface potential corresponding to the white lines in Figures 4c-d. As indicated, the thickness of trapped water is 0.37 nm, equal to the value of monolayer ice on mica.13 Similarly, this layer is stable and only slightly reduced in area after air exposure over 30 days (RH: ~60%). The surface potential of graphene on Al2O3 increases only ~5 mV with trapped water, an order of magnitude smaller change than observed for MoS2. Raman spectroscopy (Figure 4f) of graphene on Al2O3 with partial water and without water show a broad G peak (17.1 and 15.5 cm-1 respectively) consistent with low doping (Fermi level 1013 cm-2). The details are listed in Table 1. While the imaging of the successive hydroxylation/hydration of the Al2O3 substrate under MoS2 was made possible by the use of atomically flat Al2O3(0001), the hydroxylation and hydration of high-temperature treated oxide substrates on ambient exposure is a generic phenomenon and occurs for e.g. amorphous SiO2. Our results can explain the observation by other groups of differences in PL between MoS2 sample edges and center as due to water intercalation, and not a different edge structure or reactivity. Notably, differences in PL between sample edge and center have also been observed for (Mo,W)(S,Se)2 on SiO228, 29 & Al2O3.30 The observed doping shifts of order 1013 cm-2 are large, exceeding the maximum achievable gate-induced doping in many device geometries. The surface potential inhomogeneity observed here for even the highest quality atomically flat substrates is likely significantly worse for less ordered or amorphous substrates, and is likely a major source of disorder in MoS2 on oxides. This work thus has strong implications for fabrication of MoS2 devices including field effect transistors, and optoelectronic or valleytronic devices, where uniformity of doping and electronic properties is essential. It points out the special advantage of hydrophobic dielectrics such as PMMA8 and BN9,
10
in avoiding disorder caused by inhomogeneous hydroxylation and water
intercalation, and helps to understand the higher mobility,8, 9 greatly reduced hysteresis10 and improved environmental stability for such substrates.
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METHODS Sample preparation. Monolayer MoS2 single crystals were grown by a novel method, sulfur-assisted CVD. The difference between this and other CVD techniques is that the precursors are sulfur powder (Sigma-Aldrich, 84683) and MoS2 powder (Sigma-Aldrich, 69860), instead of the most commonly used MoO3 powder.3 Growth substrates of Al2O3(0001) were cleaned by 10 min acetone and isopropanol sonication.
Substrates
(sapphire and exfoliated h-BN), MoS2 (0.35 g) and sulfur (0.8 g) were loaded into three separated quartz crucibles. The growth was performed in a 1 inch diameter quartz tube flowing ultra-high purity argon at atmospheric pressure. During growth, MoS2 and substrates were located at the centre and downstream of the furnace, respectively. The sulfur powder was loaded upstream of the MoS2 and substrates and independently heated. The furnace and the sulfur were heated to 940 ℃ and 225 ℃ respectively in 30 min. The furnace was cooled to 500 ℃ over 10 min, followed by rapid cooldown to room temperature (furnace opened for maximum cooling rate achievable). The flow rate of argon was kept at 500 sccm until the furnace temperature initially reached 500 ℃. The flow rate was then decreased to 60 sccm during growth, until the furnace temperature returned to 500 ℃, when the flow rate was increased again to 500 sccm. Hexagonal-BN crystals were obtained in solvent growth process under high pressure by using belt type high pressure apparatus.31 In order to reduce impurities of oxygen and carbon, Ba-B-N solvent were employed. The graphene flakes were exfoliated on atomically flat Al2O3(0001) surface in ambient environment (RH: 60%) and glove box (H2O < 0.1 ppm). Before exfoliation, the sapphire pieces were cleaned by acetone and isopropanol sonication for 10 min. The substrates next were annealed at 1300 ℃ for 2h under ambient environment and then allowed to cool down to room temperature naturally. The substrates were further cleaned by oxygen annealing at
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1050 ℃ for 1 hr. Before loading to the main chamber of glove box (H2O < 0.1ppm), sapphire substrates and scotch tape sticking with graphene flakes were cleaned by repeating the cycle of pumping-flushing nitrogen (purity: 99.999%) three times in the load lock. For glove box exfoliation, the substrates were additionally heated on a hot plate at 200 ℃ for 1 hr soon after being loaded into the glove box filled with nitrogen. Optical measurements. Raman/PL measurements were performed using a confocal microscope system (WITec alpha 300R) with a 100× objective with a numerical aperture (NA) of 0.90. A 532 nm laser with spot diameter of ~500 nm was used to excite samples placed on a piezocrystal-controlled scanning stage. The best mapping spatial resolution was ~250 nm. The spectra were collected using a 600 line/mm grating. To avoid sample damage, low laser power (50 µW) and short integration time (0.1 s) were applied during PL mapping. The Raman spectra of both MoS2 and graphene were measured using much longer integration time (2s). The PL intensity images were obtained by summing the PL intensity from 630 nm to 690 nm, and the PL peak position and full width of half maximum (FWHM) were estimated from fits to a single Lorentzian lineshape. The Raman spectral parameters of graphene were obtained by Lorentzian fitting as well. AFM & KPFM measurements. The AFM and KPFM images were obtained using Bruker Dimension Icon SPM through LiftMode, a dual-pass technique. The topography line profile of the sample was achieved in amplitude-modulated tapping mode during the first pass. The second pass traced the acquired topography line at a set lift height to allow surface potential measurement. A chrome/platinum coated silicon tip was used as the (conductive) tip. For MoS2 measurements, the lift height and the AC bias voltage were 20 nm and 2 V, respectively. While for graphene measurements, the parameters were 10 nm and 0.5 V.
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FIGURES Figure 1. Large-area monolayer MoS2 crystals. (a) Optical image of MoS2 monolayer crystals on Al2O3(0001) grown by sulfur-assisted CVD. (b,c) Typical Raman and photoluminescence (PL) spectra taken from the same position on a MoS2 crystal. (d-f) PL mapping image of a portion of a large crystal showing the integrated PL intensity (d), peak position (e), and peak width (f). Scale bar in (d) applies to (e, f). The PL and Raman measurements were done right after removing the samples from the furnace. Figure 2. Atomic force microscopy (AFM) and Kelvin probe force microscopy (KPFM) of water intercalation. (a-c) AFM height and phase, and KPFM surface potential images of atomically flat MoS2 after 5hr air exposure. Inset in (a) shows the optical image of the MoS2 crystal. (d,e) Height image (d) and KPFM image (e) of the square region shown in (a,c). (f) Correlated height and surface potential profiles corresponding to the white lines shown in (d,e). Shading indicates three regions of surface potential corresponding to bare Alterminated Al2O3 (unshaded), hydroxylated Al2O3 (yellow), and hydrated Al2O3 (blue). (g-i) AFM height, phase and surface potential images of the MoS2 crystal after air exposure for 3 days. (j-l) AFM height, phase and surface potential images of the MoS2 crystal after air exposure for 30 days. Scale bar in (a) applies to (b, c, g-l). Figure 3. Correlated atomic force microscopy (AFM) and photoluminescence (PL) imaging. (a-c) AFM height, phase and surface potential images of a MoS2 crystal after air exposure for 30 days. (d-f) PL mapping images of the flake represented in the form of sum intensity (d), peak position (e), and peak width (f). (g) Line traces of the data corresponding to the lines shown in (a-f). Shading indicates two regions of surface potential corresponding to hydroxylated Al2O3 (yellow), and hydrated Al2O3 (blue).
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Figure 4. Graphene on atomically flat Al2O3(0001). (a,b) atomic force microscopy topographs of graphene on Al2O3(0001) exfoliated in ambient (RH: ~60%) and glove box (H2O < 0.1 ppm). (c,d) Detail images of height (c) and surface potential (d) of graphene exfoliated in ambient on an atomically flat sapphire surface. (e) Line traces of height (black dash dot) and surface potential (blue solid dot) corresponding to the white lines shown in panels (d) and (e). (f) Raman spectra of graphene with partial intercalation of water (blue circles) and without water (red circles). Lines are Lorentzian fits to the data. Figure 5. (a,b) Photoluminescence (PL) images showing the integrated PL intensity of a MoS2 crystal grown on exfoliated h-BN immediately post-growth (a) and after ambient exposure for 3 weeks (b). (c) PL spectra corresponding to (a) and (b). (d) Atomic force microscopy (AFM) topograph of the MoS2 crystal on h-BN. (e) AFM height profile following the black line shown in (d). Table 1. Summary of interactions between ambient water and MoS2/Al2O3, graphene/Al2O3 and MoS2/h-BN interfaces.
AUTHOR INFORMATION Corresponding Author * [email protected]
ACKNOWLEDGMENT M.S.F, M.T.E., and C.Z. acknowledge support from ARC (DP150103837 and FL120100038). C.Z. acknowledges support from ARC DECRA (DE140101555). Q.B. acknowledges support from the 863 Program (Grant No. 2013AA031903), the youth 973
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program (2015CB932700), the NSFC grants (Grant No. 51222208, 51290273), ARC DECRA (DE120101569), DP (DP140101501) and Engineering Seed Funding Scheme (2014) at Monash University. This work was performed in part at the Melbourne Centre for Nanofabrication (MCN) in the Victorian Node of the Australian National Fabrication Facility (ANFF). K.W. and T.T. acknowledge support from the Elemental Strategy Initiative conducted by the MEXT, Japan. T.T. acknowledges support from a Grant-in-Aid for Scientific Research on Grant 262480621 and on Innovative Areas “Nano Informatics” (Grant 25106006) from JSPS.
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Figure 1.
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Table 1. Water intercalation Water-induced work from ambient function shifts MoS2/Al2O3 Yes ~80 mV graphene/Al2O3 No* ≤5mV MoS2/h-BN No n/a *water trapped upon exfoliation in ambient Interface
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