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Direct Growth of MoS2/h-BN Heterostructures via a Sulfide-Resistant Alloy lei fu, Yangyong Sun, Nian Wu, Rafael G. Mendes, Linfeng Chen, Zhen Xu, Tao Zhang, Mark H. Rümmeli, Bernd Rellinghaus, Darius Pohl, Lin Zhuang, and Lei Fu ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.5b06254 • Publication Date (Web): 12 Jan 2016 Downloaded from http://pubs.acs.org on January 13, 2016

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Direct Growth of MoS2/h-BN Heterostructures via a Sulfide-Resistant Alloy Lei Fu†, Yangyong Sun†, Nian Wu†, Rafael G. Mendes‡, Linfeng Chen†, Zhen Xu†, Tao Zhang†, Mark H. Rümmeli‡, Bernd Rellinghaus‡, Darius Pohl‡, Lin Zhuang† and Lei Fu†,* †

College of Chemistry and Molecular Science, Wuhan University, Wuhan 430072, China

‡

IFW Dresden, P.O. Box 270116, 01069 Dresden, Germany

*Address correspondence to [email protected] RECEIVED DATE Corresponding Author Footnote: Prof. Lei Fu College of Chemistry and Molecular Science, Wuhan University No.16 Luojiashan Road, Wuchang District, Wuhan, China 430072 Tel: (+)86-27-6875-5867 Fax: (+)86-27-6875-5867 E-mail: [email protected]

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ABSTRACT: Improved properties arise in transition metal dichalcogenides (TMDCs) materials when they are stacked onto insulating hexagonal boron nitride (h-BN). Therefore, the scalable fabrication of TMDCs/h-BN heterostructures by direct chemical vapor deposition (CVD) growth is highly desirable. Unfortunately, to achieve this experimentally is challenging. Ideal substrates for h-BN growth, such as Ni, become sulfides during the synthesis process. This leads to the decomposition of the pre-grown h-BN film and the thus no TMDCs/h-BN heterostructure forms. Here, we report a thoroughly direct CVD approach to obtain TMDCs/h-BN vertical heterostructures without any intermediate transfer steps. This is attributed to the use of a nickel-based alloy with excellent sulfide-resistant properties and a high catalytic activity for h-BN growth. The strategy enables the direct growth of single-crystal MoS2 grains of up to 200 µm2 on h-BN, which is approximately one order of magnitude larger than that in previous reports. The direct band gap of our grown single-layer MoS2 on h-BN is 1.85 eV, which is quite close to that for free-standing exfoliated equivalents. This strategy is not limited to MoS2-based heterostructures, and so allows the fabrication of a variety of TMDCs/h-BN heterostructures, suggesting the technique has promise for nanoelectronics and optoelectronic applications.

KEYWORDS: MoS2/h-BN heterostructures, direct CVD growth, sulfide-resistant alloy, optical properties

Two-dimensional (2D) atomic crystals have emerged as versatile materials with the potential to be applied in future optoelectronic devices owing to their unique planar structure and distinct electrical and optical properties.1,2 The controlled stacking of different 2D materials will greatly expand their family and broaden their applications.3,4 The vertical stacking of 2D materials and hexagonal boron nitride (h-BN) provides a unique platform for exploring unique phenomena in condensed matter physics and electrical

properties.5–9

The

stacking

of

transition

metal

dichalcogenides

(TMDCs)/h-BN

heterostructures were first created through sequential mechanical transfer procedures, which are severely limited by interface contamination, the poor interlayer contact and a lack of production 2 Environment ACS Paragon Plus

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scalability compared to chemical vapor deposition (CVD).1,10,11 Thus, the scalable fabricating TMDCs/h-BN heterostructures by direct CVD growth is highly attractive. Unfortunately, it is very challenging to directly synthesize the TMDCs on h-BN by CVD without any intermediate transfer steps owing to the tendency for sulfide formation of the metal catalysts and substrates (e.g., Cu or Ni) used for h-BN growth.12,13 The strong sulfur-metal bond causes the rearrangement of the metal surface structure which leads to the decomposition of CVD grown h-BN and thus fabrication of TMDCs/h-BN fails.14 Recently, Warner et al. presented the fabrication of TMDCs/h-BN by all CVD process.15 They transferred the CVD grown h-BN films from Cu substrates to SiO2/Si and then conducted the CVD growth of MoS2. The PMMA-assisted transfer played a key role in avoiding sulfidation of the Cu substrate during subsequent MoS2 growth. However, the transfer step will obviously introduce impurities and strain at the heterostructure interface and thereby degrade layer-layer interactions. Thus far, the stacking of TMDCs/h-BN layers in a controllable manner by direct CVD growth without any intermediate transfer steps has not been achieved and is a highly desirable procedure. To address this challenge, we designed a strategy for the fabrication of TMDCs/h-BN heterostructures via sulfide-resistant metal, such as Ni–Ga alloys, which possess high catalytic activity for h-BN growth and excellent sulfide-resistant properties as confirmed by density functional theory (DFT) calculations. Interestingly, the Mo foil acts as the support substrate beneath the Ni–Ga alloy and also serves as the Mo source for MoS2 growth. This avoids the introduction of molybdenum oxide or other contaminations. The strategy enables the direct growth of single-crystal MoS2 grains up to 200 µm2 in area over h-BN, which is approximately one order of magnitude larger than quoted in previous reports. The direct band gap of our grown single-layer MoS2 on h-BN is quite close to that found for free-standing exfoliated systems. This is because our strategy benefits from cleaner interfaces, smaller lattice strain, and lower doping levels than similar heterostructures formed using transfer procedures. RESULTS AND DISCUSSION Synthesis and characterizations of direct grown MoS2/h-BN heterostructures. The fabrication 3 Environment ACS Paragon Plus

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procedure for the MoS2/h-BN heterostructures is illustrated in Figure 1. The Ni–Ga alloy was utilized as a model system to demonstrate our strategy. The Ni–Ga alloy was formed by thermal annealing using Mo foil as a support base (see Figure S1, the Supporting Information). After that, the few-layer h-BN film was initially grown on the Ni–Ga surface by CVD using an ammonia borane precursor under atmospheric pressure. Then H2S was introduced into the CVD system to grow MoS2 on the top of h-BN. The MoS2/h-BN heterostructures were fabricated by direct all CVD growth without any intermediate transfer steps. Figure 2a schematically illustrates the h-BN growth over a Ni–Ga substrate, where the monomeric aminoborane (BH2NH2) decomposed from the ammonia borane precursor, reacts with Ni atoms to form a Ni–B phase which then further reacts with ammonia.16 The Ni–Ga alloy reduces the reaction barrier for the borazane oligomer molecules to form a honeycomb lattice along the h-BN domain edge during growth. Macroscopic uniformity of the h-BN films, catalyzed by the Ni–Ga alloy at 1000 °C, is achieved (Figure S2, Supporting Information). Figure 2b shows typical XPS data of the h-BN film with the B 1s and N 1s peaks are located at 190.46 and 397.96 eV, respectively, and the B:N atomic ratio calculated from the XPS data is 1.07:1, which is close to the 1:1 stoichiometry found in h-BN.17 The resistance of the h-BN film was measured to determine insulating property. The optical image in the inset of Figure 2c shows the device configuration of h-BN that has been transferred onto the SiO2/Si substrate. As shown in Figure 2c, the current–voltage (I–V) curve shows a measured current of h-BN is below 1×10–12 A, indicating the excellent insulating nature of our h-BN.16 By controlling the growth temperature, growth time and precursor heating temperature, single layer to several layer h-BN films could be synthesized. The selected area electron diffraction pattern (SAED) presented in the inset of Figure 2d reveals a distinct hexagonal structure (inset of Figure 4c). The atomic positions of the B and N atoms in the monolayer h-BN can be directly identified via HRTEM images with N atoms being slightly brighter than B,18,19 further confirming the high crystallinity of the obtained h-BN (Figure 2d). More characterizations of the h-BN films, including optical image, atomic force microscope (AFM),


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Raman spectroscopy and optical absorption spectrum presented in Supplementary Figure S2. To have a better understanding the formation of h-BN over Ni–Ga alloys, we carried out DFT calculations for the decomposition energy of (BH2NH2)n=1 on substrates with different compositions (Figure 2e–f, and Figure S3, Supporting Information).20 The dehydrogenation energy of BH2NH2 on substrates and the decomposition energy of BN on substrates to form separate B, N on substrates was taken into account. The former calculation showed that dehydrogenation energy of BH2NH2 on Ni–Ga alloy is 1.00 eV, which is 0.57 eV lower than that for pure Ni (1.57 eV), which confirms that the dehydrogenation of BH2NH2 is preferable over Ni–Ga alloy. Furthermore, through the geometry of BH2NH2 adsorbed on Ni and Ni–Ga alloys, B–N was tilting on substrates with B closing to substrates (Figure S3, Supporting Information), N atoms do not obviously interact with substrates in terms of chemical electronic effects, because of the large distance. This is consistent with the XPS experiments which indicated only Ni–B present viz. no Ni–N was observed,16,20 which indicates that the growth of h-BN requires BN in the form of a molecule rather than individual B and N atoms. Therefore, we developed further decomposition calculations for BN; the decomposition energy of BN on Ni–Ga alloys was found to be 0.39 eV higher than that on Ni. This implies Ni–Ga alloys inactively catalyzed BN to decompose in contrast to Ni, that is to say, BN adsorbed on Ni–Ga alloy surfaces, as compared to pure Ni surfaces, is more stable. All aspects evidently demonstrate that Ni–Ga alloys are a superior catalyst to promote the growth of two dimensional h-BN (see DFT details in Methods). Figure 3a shows the scheme for the follow up growth of MoS2/h-BN heterostructures on Ni–Ga alloys. Mo foil beneath Ni–Ga acts as the Mo source and the highly reactive sulfur precursor (H2S) facilitates sufficient sulfurization (see experimental details in Methods). The growth mode of MoS2 on h-BN would be a Frank van der Merwe mechanism, where small 2D nuclei are formed and grow to be large 2D crystals. In view of the difference in nucleation and growth rates on h-BN compared with SiO2, the precise reaction temperature for the growth of crystalline of MoS2 with large size is preferred at ∼680 °C. In addition, the clean and smooth surface of h-BN is beneficial to the diffusion of additional


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MoS2 molecules obtained by sulfidation. Through that, the direct growth of single-crystal MoS2 on h-BN substrates form with grain sizes up to 200 µm2 (Figure 3b), which is approximately one order of magnitude larger than the previous reports.7,14 To further characterize MoS2 grown on h-BN, TEM and SAED measurements were conducted (Figure 3c). In the SAED patterns, the six spots of the inner hexagon (denoted by green lines) correspond to the MoS2 (110) plane with the spacings of 1.6 Å and the _

six spots of the outer hexagon (denoted by blue lines) correspond to the BN (1010) plane with the spacings of 2.5 Å.10,11 This data demonstrates two sets of hexagonal diffraction patterns in our MoS2/h-BN sample, providing a strong evidence of the coexistence of single crystal MoS2 and h-BN. HRTEM analysis also agreed with the spacings obtained from SAED. Raman spectra can investigate lattice strain effect, doping levels, and stacking interaction of heterostructures.21 The in-plane E2g1 mode is associated with the built-in strain of MoS2 with a blue shift for a smaller biaxial lattice tensile strain.22–24 An E2g1 peak position variation of ∼2.8 cm–1 (blue shift) was observed as shown in Figure 3d. This indicates that using a CVD-grown h-BN film as a substrate can make lattice strain easier to release. An A1g peak position variation of ∼1.2 cm‒1 (blue shift) for our sample compared with that on SiO2, can indicate the reduced doping levels because of less electron density between the interface of heterostructures.25,26 For the transferred ones on h-BN, the A1g peak position remained almost unchanged compared to the ones on SiO2, which means the use of a transfer process can increase doping effects, spoil the original interface and reduce the interlayer interaction between the stacking layers. The inset of Figure 3d provides the measured characteristic Raman peak of CVD-grown h-BN under MoS2 domains, which is located at ∼1372 cm‒1, confirming the existence of the h-BN film after the growth of TMDCs. XPS was used to measure the binding energy of the Mo 3d and S 2p Mo and S, with 229.5 and 232.6 eV attributed to the Mo 3d5/2 and Mo 3d3/2 levels and 162.5 and 163.7 eV assigned to the S 2p1/2 and S 2p3/2 levels, respectively. The positions of these XPS peaks suggested that the valence of Mo is +4, which is evidence for the formation of a pure MoS2 phase (Figure 3e–f).27 We also fabricated a field effect transistor (FET) device of MoS2/h-BN using 10 nm Ti/50 nm Au 6 Environment ACS Paragon Plus

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as source and drain electrodes, p++ Si as the back gate by electron beam lithography and electron beam deposition (Figure 3g, see in Methods). The insets display the corresponding schematic diagram (top) and the optical image of the real FET device (bottom). The Ids−Vgs curve was measured at Vds = 0.5 V, with the back-gate voltage sweeping from −100 to 100 V. The transfer characteristics of the MoS2 on h-BN measured at room temperature in atmosphere exhibited n-type conduction with an on/off current ratio of ∼106, comparable with those of backgated FETs made with mechanically exfoliated MoS2 flakes.28,29 The current minimum of MoS2 on h-BN substrate presents an increase in current level compared with the ones on SiO2 28,30 and an n-type shift of the minimum current level towards negative Vgs (red curve) down to ‒78 V. As a consequence of better contact by reducing the doping level, the mobility of the MoS2 FET on the h-BN substrate was up to 22 cm2 V‒1 s‒1 on h-BN substrate (blue curve) calculated using the equation µ = [dId/dVg] [L/WCiVd], where the channel length L = 2.6 µm, channel width W = 5.5 µm, back gate capacitance perunit area Ci = εrε0/d (εr = 3.9, d = 300 nm) and Vds = 0.5 V. The mobility of our devices are about 2 times higher than those reported for monolayer CVD grown MoS2 on other substrates.30,31 These results prove that our method can be used for enhancing the device performance of TMDCs materials. Importantly, the CVD growth method developed is versatile as a general strategy for fabricating other TMDCs/h-BN heterostructures. For example, high-quality WS2/h-BN heterostructures were fabricated using W foils in place of the Mo foils during the growth process. Related characterizations for the WS2/h-BN heterostructures are presented in Supplementary Figure S4. The photoluminescence performance of the direct grown TMDCs/h-BN heterostructures. To investigate

the

optical

performance

towards

optoelectronic

applications,

we

explore

the

photoluminescence (PL) properties of the directly grown MoS2/h-BN heterostructures. As shown in Figure 4a, the PL spectra of direct-grown MoS2 on h-BN/Ni–Ga, Ni–Ga and SiO2, as well as transferred MoS2 on h-BN were performed. The strong PL spectra of MoS2 direct grown on h-BN/Ni–Ga is centered at 1.85 eV (Figure 4a). This measured direct band gap is quite close to the exfoliated 7 Environment ACS Paragon Plus

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free-standing monolayer MoS2 (1.86 eV),32 and is superior to the transferred monolayer MoS2 on h-BN (1.83 eV) as well as the MoS2 grown on SiO2 (1.80 eV). The PL spectra of the MoS2 grown on Ni–Ga without h-BN film shows almost no features in the wavelength range from 550 to 750 nm, which is probably due to a fluorescence quenching effect from the Ni–Ga substrate. As we know, the doping effects can affect the population of neutral excitonic and charged species in 2D electron systems through a chemical equilibrium:32 A + e– ←→ A− Charged excitons (A–) are produced by the binding of a free electron (e–) to a neutral exciton (A). We extract three different types of samples’ PL intensity ratio between A and A− (A/A−) to indicate the relevant doping level. As depicted in Figure 4b–d, the MoS2 grown on h-BN shows the highest A/A− ratio, followed by the MoS2 transferred on h-BN, and the MoS2 grown on SiO2 showed the lowest ratio. The reduction of A− peak in PL could be explained by the low level of charged impurities at the interface and the minute quantity of electron transfer from MoS2 to the underlying metallic substrate. As seen in Figure 4d, The PL exciton emission of MoS2 grown on h-BN is 2 times greater than that of MoS2 on SiO2. The MoS2 grown on h-BN has a narrower full width at half-maximum (FWHM), with the value of 26.39 nm, than both the transferred sample (36.61 nm) and the sample on SiO2 (44.96 nm). These results indicate the direct grown MoS2/h-BN heterostructures with a tight interlayer contact and low level of charged impurities. The transfer process might spoil the original interface and increase the doping level, leading to a lower A/A− and wider FWHM.33 The charged impurities of SiO2 results in the highest doping level, consistent with our study of the Raman A1g mode.27,34 This observation further highlights the superiority of direct grown TMDCs/h-BN heterostructures. High sulfide-resistant performance of Ni–Ga to fabricate TMDCs/h-BN heterostructures. The success in fabricating TMDCs/h-BN heterostructures with adjacent pristine direct band gaps is attributed to the high sulfide-resistant performance of the Ni–Ga alloy substrate. Characterization of the annealed and cooled Ni–Ga substrate was performed. The cooling during the study was rapid (quenched)


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to obtain the closest possible condition to that during the synthesis. Figure 5a–b shows an energy dispersive X-ray (EDX) spectral mapping of the Ni and Ga element, filled with a Ni (green) and Ga (red) compound uniformly. The Ni and Ga peak is very prominent (Figure 5c) in EDX spectroscopy showing the stable composition of the structure. Complimentary X-ray diffraction (XRD) analysis was performed to confirm the chemical composition. For this, a gallium pellet was deposited on Ni to get the material, resulting in the formation of a Ni5Ga3 phase that matched well with the JCPDS card (Figure 5d), but possibly other phases can be present under different synthesis conditions. In our experimental condition, the Ni–Ga intermetallic compounds, Ni5Ga3, could be prepared with a pure phase, which can be attributed to the high formation energy of the different phases and the very sharp lines in the Ni–Ga phase diagram at elevated temperature 1000 °C. To examine the sulfide-resistant performance offered by Ni–Ga, some commonly used substrates, such as Cu and Ni, were chosen for comparison. We first sulfurized all four samples (pristine Ni–Ga foil, Cu foil, Ni foil and Ni–Ga foil with h-BN, respectively) under 4 sccm H2S and 100 sccm Ar, 10 sccm H2 at different temperatures ranging from 200 °C to 1,000 °C (Figure 5e). The sulfurization time was kept the same at 3 min. The net weight gains due to sulfurization were ~29.02% for pristine Cu foil and ~22.41% for pristine Ni foil at 800°C. Furthermore, the sulfurization duration versus weight gain is presented in Figure 5f, at 700 °C and under Ar+H2S gaseous atmosphere. The gained weight saturates after sulfurization for 30 min and the Ni–Ga alloy dramatically reduced this value from 35.0% (pristine Ni foil) to 0.57%. To carefully realize the sulfide-resistant mechanism of the Ni–Ga alloy, XPS depth profiles were used to map the elemental distribution across surface and interior before and after the sulfurization treatment. For Ni–Ga after treatment, the intensity evolution during etching is shown in Figure 6a, for Ni, Ga, B, N and S, respectively. The main peaks at ~852.6, 870.4 eV are attributed to Ni 2p3/2, Ni 2p1/2 and that at ~18.6 eV is assigned to the Ga 3d5/2 indicating the formation of a Ni–Ga interaction, with no Ni-S or Ga-S peaks detected,12 and is consistent with the sample without any sulfurization treatment


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(Figure S5, Supporting Information). After sulfurization treatment, the peaks for the B1s (190.9 eV) and N 1s (398.1 eV) are still prominent, indicating h-BN is well-preserved. For sulfur, the only peak located at ~162.8 eV, means that it mainly exists as an elementary deposit rather than as a sulfide compound. In sharp contrast, sulfide regions are prominent in Ni substrates and moreover they form deep into the Ni substrate (Figure 6b). The two strong peaks at around 856.4 and 874.2 eV are attributed to Ni 2p3/2 and Ni 2p1/2, respectively. The peak at 161.7 eV can be assigned to S 2p. Taking into account the atomic sensitivity factor of S and Ni, the atomic ratio of Ni:S is approximately 3:2 (Figure S6, Supporting Information), revealing the pure Ni forms a Ni+Ni3S2 eutectic melt under the same condition,35 with BN peaks not detected indicating the decomposition of BN. To better understand sulfur desorption from Ni5Ga3, we investigated the adsorption of S atoms on Ni (111) and Ni5Ga3 (111) surfaces using DFT simulations (Figure 6c). After DFT simulation optimization, we gain stable geometries for S atoms adsorbed on the substrates. The binding energy of an S atom on Ni (111) is 0.563 eV stronger than for Ni5Ga3 (111). In addition, we analyze the electronic structure according to charge density difference plots, which reflect electron transfer between the substrates and S atoms. There is an obvious difference between the two plots. For Ni5Ga3, charge accumulation (red color) appears in the middle of Ni and S atom locations and charge depletion (green color) is found around both Ni and S atoms. This means that electrons transfer from both Ni and S atoms to their middle regions. They both share electrons together and dominate a covalent interaction. As for Ni, charge accumulation apparently occurs around three Ni atoms, charge depletion around S atoms, this illustrates that electron transfer from an S atom to Ni atoms, which points to an ionic interaction, rather than a covalent interaction as found with Ni5Ga3. Moreover, we choose an isosurface value based on the criteria of located density cloud to be as large as possible. It turns out that the isosurface of charge density difference is significantly lower for Ni5Ga3. Therefore, we can draw the conclusion that Ni5Ga3 has a weaker interaction with S atoms and prefers S desorption while Ni tends to inhibit sulfurization.


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CONCLUSION In summary, we have developed a Ni–Ga alloy, with excellent sulfide-resistant properties, as a model system to achieve the direct CVD fabrication of TMDCs/h-BN stacking heterostructures without any intermediate transfer steps. The directly grown heterostructures, have a tight interlayer contact, cleaner and sharper interfaces, smaller lattice strain, and lower doping levels allowing the direct band gap of single-layer MoS2 on h-BN up to 1.85 eV, which is close to that for free-standing exfoliated MoS2. The strategy enables the direct growth of single-crystal MoS2 grains up to 200 µm2 on h-BN, which is approximately one order larger than that shown in previous reports. Importantly, our versatile CVD process allows the fabrication of other high-quality 2D TMDCs/h-BN heterostructures, such as WS2/h-BN, which confirms it as a universal approach for creating various novel 2D heterostructures. This synthesis strategy can provide further possibilities for the investigation of many exciting properties and intriguing applications that are complementary to those of existing 2D heterostructures between h-BN, graphene and TMDCs. METHODS AND MATERIALS Direct CVD growth of MoS2/h-BN heterostructures. For the growth of h-BN, the Ni–Ga/Mo foil was placed in a CVD chamber and was gradually heated up to 1000 °C for 20 min in a mixed gas Ar/H2 (4:1) with a flow rate of 100 sccm. The main growth was carried out in a mixed Ar/H2 gas flow of 75 sccm for 30 min at 1000 °C with NH3-BH3 as the source material sublimated at 110−130 °C. The h-BN films on Ni–Ga were subsequently used as the substrate to grow MoS2. The synthesis of MoS2 was accomplished with precursors of Mo foil (Mo>99.98%, Sigma Aldrich) and H2S in APCVD using Ar as the carrier gas. The substrate was placed in the center of furnace. After the system was flushed for 60 min with Ar, 4 sccm H2S was pre-introduced into the growth region for 15 min. Then the temperature was increased to ∼680 °C at a speed of 40 °C/min and kept for 10 min under 500 sccm Ar flow. Next, the Ar flow rate was reduced to 10 sccm and maintained for 25 min, followed by a fast cooling process.


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Device fabrication. The MoS2/h-BN heterostructures were transferred onto 300 nm SiO2/Si substrates via PMMA transfer method. This was followed by acetone and isopropyl alcohol (IPA) cleaning to remove the PMMA residues. FET devices were made via standard e-beam lithography procedure. The electrodes (10 nm Ti/50 nm Au) were deposited by e-beam evaporator with rate deposition of 1 and 3 Å/s, respectively. Lift off process was performed with acetone followed by IPA cleaning. DFT calculation. All total energy calculations are implemented using VASP code with the projector augmented wave method (PAW). The interaction between valence electrons and ion cores is described by plane waves, and the electronic exchange and correlation effects are treated by Perdew-Burke-Ernzerhof (PBE) approximation with the van der Waals correction. Spin polarization is taken into account and the corresponding kinetic energy cut-off of 450 eV is applied to converge the total energy within a 1E–4 eV precision. The smearing scheme in the electronic relaxation was from the second order Methfessel-Paxton method. The GGA-PBE method optimized equilibrium bulk lattice constants with a = 3.737 Å, b = 5.062 Å, c = 5.062 Å for Ni5Ga3, and a = 3.513Å for Ni were used. In order to simulate dissociation and adsorption on a metal surface we choose both crystals in the (111) direction. Ni5Ga3(111) surface is modeled with 2x3 repeating unit cells which corresponds to a slab thickness of approximately 7.5 Å, and separated by a vacuum region of 15Å. And Ni(111) surface is modeled with 4×8 repeating unit cells and four layers which has a similar size as Ni5Ga3(111) surface in terms of surface area, slab thickness and vacuum region. To balance the computational cost with simulation accuracy, the bottom atoms under 3.6 Å of Ni5Ga3 surface model and the bottom two layers atoms of Ni surface model are fixed, and the remaining atoms were fully relaxed for optimizing geometry with the conjugate-gradient algorithm until the forces are less than 0.02 eV/Å. The Brillouin zone are sampled on a 1×1×1 Monkhorst-Pack grid. The dehydrogenation energies, Edh, of each dehydrogenated borazane on the surfaces are calculated by Edh = E(Sub-BN) + 2×E(H2) – E(Sub) – E(BH2NH2), the decomposition energies are calculated by Edc = E(Sub-B–N) – E(Sub-BN), the adsorption energies of S atom on the surfaces are defined as Eads(S) =


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E(Sub) + E(B) – E(Sub-S). Where E(Sub-BN) is the total energy of fully dehydrogenated borazane adsorbed on the substrate; E(Sub) is the energy of clean substrate; E(BH2NH2), E(S) and E(H2) are the energies of BH2NH2, S and H2, respectively in the vacuum; E(Sub-B–N) is the energy of separated B, N atoms adsorbed on the substrate. Characterization. Optical images were taken with an optical microscope (Olympus DX51, Olympus), and Raman spectroscopy and photoluminescence (PL) with an excitation wavelength 532 nm was carried out using Renishaw inVia, Renishaw. The AFM images were taken with NT-MDT Ntegra Spectra with h-BN transferred onto the 300 nm SiO2/Si. The TEM images were taken with an aberration-corrected, high-resolution TEM (AC-HRTEM, FEI Titan3) operating at 80 kV with h-BN samples directly transferred onto a copper TEM grid. The X-ray photoelectron spectroscopy (XPS, Thermo Scientific, ESCALAB 250Xi) depth profiling was performed by Ar ionic bombardment to gradually remove the surface layers until 9 nm downward into the bulk phase. The current (I)–voltage (V) data were collected in the cryogenic probe station (Lakeshore) under vacuum (4×10−3 torr). SEM energy-dispersive X-ray (EDX) mapping was performed on a Focused Ion Beam Scanning Electron Microscope (JIB 4500 MultiBeam). Acknowledgement. The research was supported by the Natural Science Foundation of China (Grants 51322209, 21473124), the Sino-German Center for Research Promotion (Grants GZ 871) and the Ministry of Education (Grants 20120141110030). We thank Prof. Lei Liao for the FET fabrication and Prof. Hongxing Xu for the PL characterization. Supporting Information Available: Further details and discussions on the typical microscope images and layer analysis by Raman spectroscopy of the TMDCs/h-BN and the XPS depth profiles in the Ni–Ga and Ni foil before sulfurization. This material is available free of charge via the Internet at http://pubs.acs.org.


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Figure 1. Schematic showing the preparation of TMDCs/h-BN heterostructures. (a) The Ni–Ga alloy (Mo foil serving as the support substrate) is prepared as the growth substrate. (b) Continuous h-BN films is obtained under an atmospheric CVD process. (c) H2S is introduced into the CVD system to grow MoS2 on the top of h-BN (the Mo foil serves as the Mo source for MoS2 growth).


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Figure 2. Characterizations of the h-BN films and growth mechanism on Ni–Ga. (a) Schematic illustration of various dynamic processes involved in B atoms diffusion on the Ni–Ga system. (b) XPS spectra of the as-grown h-BN film, the binding energies of B 1s and N 1s in h-BN are identified. (c) The current–voltage curve of h-BN film shows its excellent insulating nature. The inset is the optical image of the device configuration. (d) High-resolution TEM image of h-BN film. The inset shows the corresponding SAED pattern. (e–f) The molecular dynamic simulation (top view) of H2BNH2 dehydrogenation and BN decomposition on the surface of Ni–Ga alloy (e) and Ni substrate (f).


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Figure 3. Synthesis and characterizations of MoS2/h-BN heterostructures. (a) Schematic of the synthesis process for heterostructures. (b) SEM image of the directly grown single-crystal MoS2 on h-BN, Insert: MoS2 crystal with grain up to 200 µm2 and the scale bar is 5 µm. (c) TEM characterizations of MoS2/h-BN heterostructures. SAED patterns corresponding to the heterostructures (insert) and the scale bar is 0.2 nm. The spots in the green dashed hexagons indicate (110) plane of _

MoS2, and the spots in the blue dashed hexagons indicate (1010) plane of BN. (d) Raman spectra collected from different MoS2 domains grown on three types of substrates, grown on a SiO2 surface (green), transferred (blue) and grown (red) on h-BN film, respectively. The inset shows the Raman spectra measured from the h-BN film after direct growth of MoS2. (e-f) XPS spectra of Mo 3d and S 2p for CVD MoS2. The two peaks at 229.5 and 232.6 eV are attributed to the Mo 3d5/2 and Mo 3d3/2 levels, respectively. The binding energy at 162.5 and 163.70 eV can be assigned to the S 2p1/2 and S 2p3/2 levels, respectively. (g) Electrical properties of Ids−Vgs curves for the FET devices made on MoS2/h-BN heterostructures. Inset: optical image of the device and the scale bar is 5 µm.


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Figure 4. Photoluminescence performance of the direct grown TMDCs/h-BN heterostructures. (a) PL spectra of directly grown MoS2 on h-BN/Ni–Ga, Ni–Ga or SiO2 and transferred MoS2 on h-BN, respectively, which indicates the different band gap compared to freestanding MoS2 flake (1.86 eV, pointed out by the arrow). The black annotations indicate the positions of the three main peak emissions: B excitonic transition (B), the neutral excitonic transition of A (A), and the charged excitonic transition of A (A–). (b‒d) Fits to the PL peak of strain-free MoS2 with three sub-peaks corresponding to A, A–, and B.

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Figure 5. The structural composition and sulfide-resistant performance of Ni–Ga substrates. (a-b) EDX map of the Ni and Ga peak, showing the substrate uniformly filled with a Ni (green) and Ga (red) compound. (c) An EDX spectrum of the Ni–Ga substrate showing the Ni and Ga peak are prominent. (d) XRD spectrum of the substrate, confirming the phase as Ni5Ga3. (e) Weight gain versus temperature of pristine Ni foils, Cu foils, Ni–Ga and Ni–Ga/h-BN. Sulfurization duration is 5 min in an Ar+H2S gaseous atmosphere. (f) Weight gain versus sulfurization duration at 700 °C in an Ar+H2S gaseous atmosphere.


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Figure 6. The sulfide-resistant mechanism of Ni–Ga alloys compared with pure Ni. (a) The XPS depth spectra of Ni–Ga alloy after sulfurization treatment, shown with Ni (2p3), Ga (3d), B (1s), N (1s) and S (1s), respectively, acquired after each etching. The arrows indicate the evolution of the spectra. The sulfurization treatment is accomplished at 700 °C in an Ar+H2S gaseous atmosphere for 5 min. (b) The XPS depth spectra of Ni foil after sulfurization treatment shows the sulfurization region propagates deep into the Ni substrate. (c) The three-dimensional density difference plots for S atom adsorption on Ni5Ga3 and Ni substrates. The isosurface value is +0.081 a.u. and -0.0463 a.u. for Ni5Ga3, while +0.3855 a.u. and –0.11 a.u. for Ni. Charge accumulation is in red and depletion in green. Ni atom in blue, Ga in pink and S in yellow.


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TOC Direct Growth of MoS2/h-BN Heterostructures via a Sulfide-Resistant Alloy

Improved properties arise in transition metal dichalcogenides (TMDCs) materials when they are stacked onto insulating hexagonal boron nitride (h-BN). Therefore, the scalable fabrication of TMDCs/h-BN heterostructures by direct chemical vapor deposition (CVD) growth is highly desirable. Unfortunately, this is experimentally challenging. Most substrates for h-BN growth, such as Ni, form a sulfide during the synthesis of TMDCs, leading to the decomposition of any pre-grown h-BN film and thus TMDCs/h-BN layers do not form. Here, we report a full CVD approach to obtain TMDCs/h-BN heterostructures without any intermediate transfer steps. This is attributed to the nickel-based alloy which has excellent sulfide-resistant properties and high catalytic activity for h-BN growth. The strategy enables the direct growth of single-crystal MoS2 grains up to 200 µm2 on h-BN, approximately one order larger than that in previous reports. The direct band gap of our grown single-layer MoS2 on h-BN is 1.85 eV, which is close to that for free-standing (exfoliated) MoS2. This strategy is not limited to MoS2-based heterostructures, which allows the fabrication of versatile TMDCs/h-BN heterostructures, suggesting its promising potential in nanoelectronics and optoelectronic applications.


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