Epitaxial van der Waals Contacts between Transition-Metal

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Epitaxial van der Waals Contacts between Transitionmetal Dichalcogenide Monolayer Polymorphs Chang-Soo Lee, Seung Jae Oh, Hoseok Heo, Seung-Young Seo, Juho Kim, Yong Hyeon Kim, Donghwi Kim, Odongo Francis Ngome Okello, Hocheol Shin, Ji Ho Sung, Si-Young Choi, Jun Sung Kim, Jong Kyu Kim, and Moon-Ho Jo Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.8b04869 • Publication Date (Web): 19 Feb 2019 Downloaded from http://pubs.acs.org on February 20, 2019

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Epitaxial van der Waals Contacts between Transition-metal Dichalcogenide Monolayer Polymorphs Chang-Soo Lee,1,2, † Seung Jae Oh,2,† Hoseok Heo,1 Seung-Young Seo, 1,2 Juho Kim,1,2 Yong Hyeon Kim,1,3 Donghwi Kim, 1,2 Odongo Francis Ngome Okello, 2 Hocheol Shin, 1,2 Ji Ho Sung,1 Si-Young Choi,2 Jun Sung Kim,1,3 Jong Kyu Kim,2 and Moon-Ho Jo1,2,3,* 1Center

for Artificial Low-Dimensional Electronic Systems, Institute for Basic Science (IBS),

2Department

of Materials Science and Engineering, 3Department of Physics, Pohang University of

Science and Technology (POSTECH), 77 Cheongam-Ro, Pohang 790-784, Korea †Equal

contributors, *Corresponding author: E-mail: [email protected] (Moon-Ho Jo)

ABSTRACT We have achieved heteroepitaxial stacking of a van der Waals (vdW) monolayer metal, 1T’-WTe2 and semiconductor, 2H-WSe2, in which a distinctively low contact barrier was established across a clean epitaxial vdW gap. Our epitaxial 1T’-WTe2 films were identified as a semimetal by low temperature transport, and showed the robust breakdown current density of 5.0 x 107 A/cm2. In comparison with a series of planar metal contacts, our epitaxial vdW contact was identified to possess intrinsic Schottky barrier heights below 100 meV for both electron and hole injections, contributing to superior ambipolar field-effect transistor (FET) characteristics, i.e., higher FET mobilities and higher on-off current ratios at smaller threshold gate voltages. We discuss our observations around the critical roles of the epitaxial vdW heterointerfaces, such as incommensurate stacking sequences and absence of extrinsic interfacial defects that are inaccessible by other contact methods.


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KEYWORDS Van der Waals layered materials, Two-dimensional materials, Tungsten diselenides, Tungsten ditellurides, Vapor transport synthesis, Polymorphism

TEXT Planar metal contacts to van der Waals (vdW) semiconductors often suffer from unpredictably large contact-resistance, due to inherent vdW gaps at the contacts and metal-induced defects that act as additional series contact barriers.1,2 This unfavorable trait is critically impeding construction of atomically thin two-dimensional (2D) circuitry in higher performances based on various vdW semiconductors, i.e., monolayer (ML) transition-metal dichalcogenide (TMDC) semiconductors.3,4,5 To circumvent this limit, alternative contact methods have been proposed, including 1D edge contacts with Cr or Nb and coplanar contacts with metal-polymorphs, to achieve ultralow contact resistance, 6,7,8,9,10,11 and non-perturbing 2D metal contacts by vdW lamination to achieve predictable contact potentials.12 In this work, we report another contact architecture, namely, epitaxial vdW contacts, by vertical heteroepitaxy of metallic (1T’WTe2) and semiconducting (2H-WSe2) TMDC MLs. In other words, we established a new type of the contacts by a sequential synthesis of semiconductor channels and contact metals, utilizing the concept of polymorphism, thus providing an advantage to achieve large-area devices, rather than by conventional lithographic metals or dry transfer methods. Heteroepitaxial integration of such W-dichalcogenide MLs established distinctive metal-semiconductor interfaces, where the crystallographic stacking sequences are deterministically defined across the intrinsic vdW gap at the atomic scales. By incorporating these epitaxial vdW contacts to field-effect transistors (FETs), we have identified lower Schottky barrier heights below 100 meV for both electron and hole injections, with superior ambipolar FET characteristics, in direct comparison to those in planar metal contacts. Heteroepitaxial stacking of 2H-WSe2 and 1T’-WTe2 MLs was achieved by sequential chemical vapor deposition (CVD) at different growth temperatures (TG) for each ML, as illustrated in Fig. 1a. First, ACS Paragon Plus Environment

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triangular-faceted 2H-WSe2 single crystals were synthesized from vapor-phase reactions of high-purity WO3 and Se precursors on SiO2/p+-Si substrates at TG = 850 ℃ – see also Fig. S1. Subsequently, at lower TG = 700 ℃, stacking growth of 1T’-WTe2 ML crystals was accomplished with WO3 and Te precursors with KI as a growth agent (Methods).13,14,15,16 We observed that the vertical stacking growth of 1T’-WTe2 MLs nucleated at the edge of the preexisting 2H-WSe2 MLs, and proceeded inwardly. A cross-sectional high-angle annular dark field scanning transmission electron microscope (HAADF-STEM) image (Fig. 1b) shows that the upper 1T’-WTe2 ML (a distorted octahedral coordinated polytype) is coherently stacked on a 2H-WSe2 ML (a trigonal prismatic one) across the vdW gap of 0.29 nm (defined as the distance between the lower Te atom of the top WTe2 ML and the upper Se atom of the lower WSe2 ML) in Fig. 1b without any atomic intermixing or damage – see also Fig. S2. Typically, the triangular-faceted 2H-WSe2 ML (0.8 nm thick and 20 μm wide) is covered on the edge areas by the 5 μm wide 1T’-WTe2 ML (1 nm in thick), as evident with images (Fig. 1c) and Raman scattering spectra (Fig. S3), where the characteristic A1g (251 cm-1) and A15 (164 cm-1) vibration modes in 2H-WSe2 and 1T’-WTe2 are spatially resolved.17,18 At the stacked region, the photoluminescence (PL) at 1.58 eV of the imbedded 2H-WSe2 ML, measured at 4K, was suppressed (Fig. 1d), while maintaining the similar full width at half maximum of ~ 77 meV with that of the bare 2H-WSe2 ML (~ 74 meV). This, in return, guarantees epitaxial ML stacking without interlayer mixing – see Fig. S4-5 and table S1 for degradation properties of epitaxial WTe2. In-plane rotations and stacking sequences of each ML were characterized by 5th order aberrationcorrected STEM under 60 kV (Fig. 1e-i) – see also Fig. S6 for energy dispersive X-ray spectroscopy mapping. The Z-contrast STEM images and the selected area electron diffraction (SAED) patterns (Fig. 1f-g), collected from the bare 2H-WSe2 region clearly verify the hexagonal ML lattice. Meanwhile the SAED patterns collected from the stacked regions (Fig. 1i) shows that the orthorhombic unit cell of the 1T’-WTe2 ML is epitaxially stacked on the hexagonal unit cell of the 2H-WSe2 ML, maintaining the (1120)2H//(200)1T’ and (-1100)2H//(020)1T’ in-plane relationships, due to the (11-20)2H edge nucleation. By nature, vdW epitaxy is characterized by absence of chemical polarity and dangling bonds at interfaces, thus the coherent epitaxial growth of the top layer is not strictly affected by lattice mismatch with the ACS Paragon Plus Environment

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bottom layer, i.e., the unit cell symmetries and lattice constants.19,20 It is true of our case of epitaxial stacking of the orthorhombic 1T’-WTe2 on the hexagonal 2H-WSe2, where there is large difference in lattice constants. Indeed, we have measured the interplanar spacings of the (200) of 1T’-WTe2 and the (11-20) of 2H-WSe2 to be 1.730 Å and 1.670 Å in the (11-20)2H//(200)1T’ relationships, and those of the (020) of 1T’-WTe2 and the (-1100) of 2H-WSe2 to be 3.140 Å and 2.893 Å in the (-1100)2H//(020)1T direction, all of which are consistent with the bare 1T’-WTe2 ML and 2H-WSe2 ML.21,22 Interlayer atomic conformations in the epitaxial 2H-WSe2/1T’-WTe2 stacks were further investigated with a HAADF image simulation (Dr. Probe, open simulator, Ernst Ruska Center Juliech), where the three types of the representative “domains” as large as ~ 6.3 nm × ~ 6.6 nm, identified in the experimental HAADF-STEM image (Figure 2a-c) – see also Figure S7 for such moiré fringes in details.23 They can be described as regions, where (d) the W atoms in the top layer (1T’-WTe2) sitting on the W atoms in the bottom layer (2H-WSe2), i.e., Wtop-Wbottom, (e) the Te atoms in the top layer on the W in the bottom layer, Tetop-Wbottom, and (f) the W atoms in the top layer on the Se atoms in the bottom layer, WTop-Sebottom. Each stacking pair is characterized by stronger atomic contrasts forming the unique patterns: the zig-zag pattern (blue circles), the short zig-zag pattern (purple circles), and the hexagonal pattern (the purple and green circles) in Fig. 2d, where there is a close agreement between the simulations and the experiments. Indeed, such simulated patterns were periodically observed in the HAADF-STEM image of Fig. 2c, as the blue, green, and purple boxes. These observations are direct evidence of the heteroepitaxial stacking of 1T’WTe2/2H-WSe2 MLs as incommensurate, where the crystallographic and electronic structures of each ML as a metal and a semiconductor are distinctly intact without strong interlayer coupling, dictated by any specific atomic conformation. Before we investigated heteroepitaxial vdW contact properties between 1T’-WTe2 and 2H-WSe2 ML polymorphs, we first examined the metallicity of epitaxial 1T’-WTe2 layers as contact electrodes. For that, we have grown 1T’-WTe2 films in large-areas at TG = 680 oC with a prolonged growth time, where individual facet crystals started to consequently coalesce into continuous polycrystalline films of mono to ACS Paragon Plus Environment

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few layers, spanning over several centimeters in the lateral size. Such centimeter-scale 1T’-WTe2 films were patterned into arbitrary shapes with conventional lithography in a scalable manner. Figure 3a is a patterned 1T’-WTe2 film with multiple Ti/Au electrodes, defined by lithography, with which we measured sheet resistance (Rs) using transfer length methods. We indeed confirmed that the continuous and uniform channels were formed over the large areas, where the channel resistance was found to be linearly proportional to the channel length, to produce the Rs of 12.1 k and the finite contact resistance (Rc) of 4.2 k·μm (Fig. 3b) – see also Fig. 3c for the film thickness dependent metallicity. As we discuss below, this finite Rc between between Ti/Au and 1T’-WTe2 can be largely negligible, compared to the Rc between 1T’-WTe2 and 2H-WSe2 (~ M·μm), when extracting the Schottky barrier heights. The distinct metallic properties of 1T’-WTe2 films were further characterized with low temperature transport measurements. Figure 3d shows the temperature dependent resistivity for thin (2.1 nm) polycrystalline films and exfoliated crystals (4.8 and 40 nm). With lowering temperatures, the resistivity decreases (40 nm) or weakly grows (4.8 and 2.1 nm), but it is never diverging in any case. This is a signature of a semimetal, in which there is a finite density of state at the Fermi level. We found that the crossover from the metallic to the weakly-insulating behaviors occurs, when the sheet resistivity is approximately h/e2. Thus, the weakly-insulating behavior of the thin films can be attributed to disorder induced Anderson localization, which introduces qualitatively different magnetoresistance (MR).24,

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For the exfoliated crystals, we

observed quadratic field dependence, as expected in the classical transport regime. The MR, (H)/(0), however, reaches only ~ 5 % at 5 T, far smaller than ones found in the bulk 1T’-WTe2.26 In bulk semimetal 1T’-WTe2, the magnitude of the positive MR becomes extremely large when electron- and hole-carriers are almost perfectly-compensated and have large mobilities. For our 4.8 nm thick (exfoliated) crystal, the positive MR is observed up to 40 % with a sharp cusp near zero magnetic field, which is a signature of weak antilocalization. In this case, the reduced thickness, significant amounts of disorders, and strong spin-orbit coupling introduce quantum correction to carrier conduction, dominating over the classical transport. Similar behavior is also found in our epitaxial 1T’-WTe2 films (2.1 nm), where the even larger


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disorder strength at grain-boundaries may spoil the balance of electron- and hole-carriers and also reduce their mobilities. Nevertheless the qualitatively similar transport behaviors unambiguously suggest that our epitaxial 1T’-WTe2 films is comparable with exfoliated crystals in crystal qualities as a contact metal. We have also measured the breakdown strength by ramping up the high dc bias voltage across the epitaxial 1T’-WTe2 films, and found the maximum current density of 5.0 x 107 A/cm2 at 12 V (Fig. S8). This value is comparable to CVD-graphene metal films.27 Finally, we have investigated our epitaxial vdW contact properties in a back-gated field-effect transistor (FET), in direct comparison to ones of planar electrical contacts with various metals, fabricated on the identical channels (Fig. S9). Figure 4a illustrates such contact schemes, fabricated on a heteroepitaxial crystal, as we labeled as FET-1 (epitaxial vdW 1T’-WTe2 contact) and FET-2 (lithographic planar Ti contact). In both cases, the work functions of Ti (Ti ~ 4.3 eV) and 1T’-WTe2 (WTe2 ~ 4.4 eV) are expected to be similarly aligned in the midgap of 2H-WSe228 (Fig. 4b, the inset for a conceptual band alignment) - see Fig. S10 for the photocurrent spectrum to confirm the energy band gap of 1.61 eV. Current-bias voltage (I-Vb) characteristics of the two FETs in Fig. 4b (at Vg = -20 V) immediately show that the FET-1 carries much higher current than FET-2; we found the channel conductance of the FET-1 is 905 nS at 300 K, which is much higher than 6.9 nS of the FET-2 by more than two orders of magnitude. This finding is also evident in the I-Vg characteristics at Vb = 50 mV (Fig. 4c), where both FETs exhibit ambipolar characteristics. We found more efficient gate tunability in FET1, where the maximum on/off current ratio is ~107 (~105 for FET-2) at 300 K. We tested more than 20 devices to quantify the statistical variation in the hole mobility of our epitaxial vdW contacts, and it is ranged to 5 - 20 cm2V-1s-1 at 300 K, which is systematically higher than that of FET-2 by more than a factor of ten (Fig. S11). Both I-Vb and I-Vg characteristics were weakly temperature-dependent in FET-1, compared to those of FET-2, both of which can be fitted by thermally activated transport. Therein, the FET-1 requires much smaller activation energy by a factor of 4.5 (-0.013 S/eV vs. -0.059 S/eV) for hole transport from the fit at 150 - 300 K (Fig. 4d) – see Fig. S12 for the temperature dependent FET mobility. ACS Paragon Plus Environment

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The observed higher sheet conductivity at 300 K and lower thermal activation energy barrier of FET-1 in Fig. 4 can be altogether attributed to the lower contact barrier, which is in turn responsible for the enhanced FET characteristics. Schottky barrier height (SBH) in our 2H-WSe2 ML FETs can be quantitatively determined within the thermionic emission model. Figure 5a is variation in the qΦbi, where q is the elementary electric charge and Φbi is the built-in potential at the contacts, extracted from the temperature dependent I measurements of FET-1 and FET-2. At the flat band condition, where variation of qΦbi starts to deviate from the linear fits in the qΦbi - Vg relations, we found the SBH (qΦSB) for hole injection in FET-1 to be ~ 71 meV at VFB of -10 V, where VFB is the flat band gate voltage, which is much smaller than ~ 126 meV at -21 V of FET-2 – see also the SBH of ~ 99 meV for the electron injection, measured on another device, in Figure S13. For another set of comparison, we also measured the similar FET-2, separately fabricated with other metals of different work functions, to be ~ 98 meV, and ~ 187 meV for the Pd- and Sc-contacted FET-2 (Fig. 5b), which are consistent with values in literature,29,30 ranged 100 ≤ qΦSB ≤ 200 meV. Notably, we have achieved the lowest qΦSB of ∼70 meV in our epitaxial vdW contact of FET-1. Systematic variation in the qΦSB of the FET-2 type with different metals allows us to estimate the effects of Fermi level pinning (FLP) within the framework of Schottky-Mott model on interfacial states in our 2H-WSe2 channels. Figure 5c is variation in the qΦSB as a function of the  of our employed metals, from which a pinning factor (S) and charge neutral level (CNL) can be extracted in the linear fit, i.e., S = ΦSB/ and CNL = ( + b)/(1- S), where  is the electron affinity, and b is the y-intersect: the S = 1 is the ideal Schottky-Mott limit.12 We found that S and CNL to be - 0.065 and 4.84 eV in our work, implying the interfacial states are strongly pinned closer to the valence band edges, thus to promote hole transport, consistent with our FET characteristics in Fig. 4. Strong FLP in vdW TMDCs (i.e., a small pinning factor of - 0.065 in our work) can be ascribed to various factors, such as interfacial d-orbital mixing between metals and TMDCs across the vdW gaps, atomic point vacancies in TMDCs, and extrinsic defects formed during device fabrications (liftoff and evaporation).28,31,32,33 Considering that our heteroepitaxial stacking of 1T’ACS Paragon Plus Environment

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WTe2/2H-WSe2 ML crystals is intrinsically incommensurate at the contacts in the micrometer scales, as discussed above, the interfacial orbital mixing is not dictated by specific interlayer atomic conformations in the nanometer scales. Rather, we expect minimal defect formations at heteroepitaxial 1T’-WTe2/2HWSe2 MLs, as it is identified by the full width at half maximum of the 1.58 eV PL peak in the FET-1 structure, which is almost identical in the bare 2H-WSe2 MLs. Crystal polymorphism in TMDC MLs is often accompanied by the distinct “electronic” polymorphism as either ultra-thin metals or semiconductors, and we demonstrated vdW epitaxial integration of metallic (1T’-WTe2) and semiconducting (2H-WSe2) TMDC MLs, which then serves as a basic circuit unit of an electrical contact, namely “epitaxial vdW contacts”. We report that these epitaxial vdW contacts establish lower Schottky barrier heights to achieve efficient carrier injections for both electrons and holes, thereby showing superior ambipolar field-effect transistor characteristics, in direct comparison to those in conventional metal contacts by vacuum deposition. Our results may suggest an alternative design rule of the 2D semiconductor circuitry by polymorphism engineering.


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FIGURES

Figure 1. (a) Sequential growth scheme for heteroepitaxial stacking of 2H-WSe2 and 1T’-WTe2 MLs. Inset: optical microscope (OM) images of monolayer WSe2 (left) and bilayer WTe2/WSe2 (right) (b) Atomic-resolution cross-sectional HAADF-STEM image with corresponding atomic model for 1T’-WTe2 (upper layer) and 2H-WSe2 (bottom layer) (c) OM image of heteroepitaxial stacked WTe2/WSe2 crystal. Inset: atomic force microscopy (AFM) height profile obtained from the crystal in image (d) PL spectra measured from the ML WSe2 region (navy) and stacked region (orange) at 4K. Inset: PL intensity mapping images from bare WSe2 (left upper) and WTe2/WSe2 heterostructures (right lower). (e) Highmagnification HAADF-STEM image at the step edge. (f) Atomic-resolution HAADF-STEM image and (g) SAED pattern from WSe2 ML region. (h) Atomic-resolution HAADF-STEM image and (i) SAED pattern from WTe2/WSe2 stacked region


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Figure 2. (a) Atomic configuration of the heteroepitaxial 2H-WSe2 and 1T’-WTe2 ML stacks (blue, red, and yellow spheres represent W, Te, and Se atoms, respectively) and (b) the simulated Z-contrast image based on the built atomic configuration of (a). (c) The experimental HAADF-STEM image from the stacked layers. The simulation and experimental HAADF-STEM images show the several interference patterns arising from two different lattices of 2H-WSe2 and 1T’-WTe2, which can be identified as representative “domains” in the nanometer scales. (d-f) Atomic models and simulated and experimental HAADF-STEM images with Z-contrast intensity profiles corresponding to each domains described in (d), (e), and (f), where the coherency between the top and bottom layers is induced by Wtop-Wbottom, TetopACS Paragon Plus Environment

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Wbottom, and Wtop-Sebottom as well as Tetop-Wbottom, respectively. Note that the stronger contrast results from the atomic channeling effect and thus the well-aligned array between the top and bottom layers gives rise to the higher yield of electron scattering in a STEM.


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Figure 3. (a) False-color SEM image of patterned WTe2 film with conventional lithography, into FET device configuration. (b) Channel length dependent resistance of 1T’-WTe2 film. The contact resistance and sheet resistance are 4.2 kΩ·μm and 12.1 kΩ respectively. (c) Thickness dependent channel sheet resistances. Inset: thickness dependent contact resistances. (d) (left) Temperature dependence of the resistivity of 1T’-WTe2 film of thickness (t) of 2.1 nm (red) and crystals of different t = 40 nm, (blue) and 4.8 nm (navy), taken at 2 K. The classical quadratic field dependence is observed in the thick sample, whereas the weak antilocalization effect, i.e. quantum correction to the 2D conductivity, is dominant for the thin sample.


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Figure 4. (a) Illustration of our device scheme for comparison vdW epitaxial metal (1T’-WTe2) contact (FET-1) with planar metal (Ti) contact (FET-2) and optical microscope image of vdW epitaxial contact devices, where upper 4 metals for FET-1 and lower 3 metals for FET-2. (b) Current-voltage (I-Vb) curves at RT for epitaxial 1T’-WTe2 contacted (navy) and Ti contacted (orange) 2H-WSe2 channel, and Ti contacted 1T’-WTe2 channel (black dotted line). Inset: Schematic energy band diagram for 2H-WSe2 with diverse metal including 1T’-WTe2. (c) Temperature-dependent Vg modulation of sheet conductivity (σ) at Vb = 50 mV for FET-1 (color scale) and FET-2 (grey scale) (d) Arrhenius plot of σ for FET-1 (navy) and FET-2 (orange) and Ti-contacted 1T’-WTe2 (grey). Solid dots : Vg = -30 V, hollow dots : Vg = 30 V.


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Figure 5. (a-b) Built-in potential energy (qΦbi) as a function of Vg for (a) FET-1 (1T’-WTe2 contact, navy), FET-2 (Ti contact, orange) and (b) FET-2 type on bare WSe2 with other metals, Sc (red) and Pd (blue). (c) Hole SBH of 2H-WSe2 channel for various metal work function. Epitaxial vdW metal contact corresponds to solid square (navy), planar metal contact corresponds to hollow red dots and line, and Schottky-Mott rule fitting for black line.


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Acknowledgments This work was supported by Institute for Basic Science (IBS), Korea under the Project Code (IBS-R014G1).

Supporting Information Available This material is available free of charge via the Internet at http://pubs.acs.org.

Methods Gas-phase growth of ML stack of 2H-WSe2 and 1T’-WTe2. ML stack of 2H-WSe2 and 1T’-WTe2 crystals were grown by sequential vapor transport synthesis from solid powder precursors. 2H-WSe2 was firstly synthesized from WO3 (99.9%, sigma aldrich) and Se (99.5%, sigma aldrich), and 1T’-WTe2 was secondly grown at other furnace. WO3 was uniformly ground on half opened quartz tube with faced down SiO2 (100 nm)/Si wafers as the growth substrates and placed at the center of a hot-walled quartztube furnace (diameter 3.81 cm, length 30.48 cm). Selenium (99.5%, 5 g) powder in an alumina boat was 24 cm placed away from the center of furnace, the 240 ℃ region during growth. Prior to heating the furnace, the furnace was evacuated to 5 × 10-3 Torr and purged by 100 s.c.c.m Ar (99.9999 %) (standard cubic centimeters per minute, in the standard temperature and pressure condition of 0 ℃ and 1 atm) for 10 min. And then, the furnace was heated to 850 ℃ for 38 min, and the temperature was kept for 5 min, with a flow of 24 s.c.c.m Ar and 2 s.c.c.m H2 (99.9999 %). At the last of growth, the final chamber pressure was about 800 torr. After the growth process, the furnace was rapidly cooled to RT. After taking out WSe2 crystals from the vacuum furnace, the as-grown WSe2 substrate was immediately transferred to another furnace for growing 1T-WTe2 to maintain minimum exposure to ambient conditions. For the stacking of ML 1T’-WTe2, the growth scheme is almost same with WSe2 growth except additional KI (99.9 %, 3 mg, sigma aldrich) with WO3. The furnace was heated to 700 ℃ for 30 min and maintained for 3 min, with a flow of 20 s.c.c.m N2. After that, the furnace was cooled to 680 ℃ ACS Paragon Plus Environment

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in 1 min and gas flow was changed to 8 s.c.c.m N2 and 2 s.c.c.m H2 simultaneously. Final conditions for growth are maintained for 10 min, then the furnace was rapidly cooled to RT.


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FIGURE CAPTIONS Figure 1. (a) Sequential growth scheme for heteroepitaxial stacking of 2H-WSe2 and 1T’-WTe2 MLs. Inset: optical images of monolayer WSe2 (left) and bilayer WTe2/WSe2 (right) (b) Atomic-resolution cross-sectional HAADF-STEM image with corresponding atomic model for 1T’-WTe2 (upper layer) and 2H-WSe2 (bottom layer) (c) Optical images of heteroepitaxial stacked WTe2/WSe2 crystal. Inset: Atomic force microscopy (AFM) height profile obtained from the crystal in image (d) PL spectra measured from the ML WSe2 region (navy) and stacked region (orange) at 4K. Inset: PL intensity mapping images from bare WSe2 (left upper) and WTe2/WSe2 heterostructures (right lower). (e) High-magnification HAADFSTEM image at the step edge. (f) Atomic-resolution HAADF-STEM image and (g) SAED pattern from WSe2 ML region. (h) Atomic-resolution HAADF-STEM image and (i) SAED pattern from WTe2/WSe2 stacked region Figure 2. (a) Atomic configuration of the heteroepitaxial 2H-WSe2 and 1T’-WTe2 ML stacks (blue, red, and yellow spheres represent W, Te, and Se atoms, respectively) and (b) the simulated Z-contrast image based on the built atomic configuration of (a). (c) The experimental HAADF-STEM image from the stacked layers. The simulation image and experimental HAADF-STEM image show the several interference patterns arising from two different lattices of 2H-WSe2 and 1T’-WTe2, which can be identified as representative “domains” in the nanometer scales. For clarity, the domains are described in (d), (e), and (f), where the coherency between the top and bottom layers is induced by Wtop-Wbottom, TetopWbottom, and Wtop-Sebottom as well as Tetop-Wbottom, respectively. Note that the stronger contrast results from the atomic channeling effect and thus the well-aligned array between the top and bottom layers gives rise to the higher yield of electron scattering in scanning transmission electron microscopy. Figure 3. (a) False-color SEM image of patterned WTe2 film with conventional lithography, into FET device configuration. (b) Channel length dependent resistance of 1T’-WTe2 film. The contact resistance and sheet resistance are 4.2 kΩ·μm and 12.1 kΩ respectively. (c) Thickness dependent channel sheet resistances. Inset: thickness dependent contact resistances. (d) (left) Temperature dependence of the resistivity of 1T’-WTe2 film of thickness (t) of 2.1 nm (red) and crystals of different t = 40 nm, (blue) and ACS Paragon Plus Environment

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4.8 nm (navy), taken at 2 K. The classical quadratic field dependence is observed in the thick sample, whereas the weak antilocalization effect, i.e. quantum correction to the 2D conductivity, is dominant for the thin sample. Figure 4. (a) Illustration of our device scheme for comparison vdW epitaxial metal (1T’-WTe2) contact (FET-1) with planar metal (Ti) contact (FET-2) and optical microscope image of vdW epitaxial contact devices, where upper 4 metals for FET-1 and lower 3 metals for FET-2. (b) Current-voltage (I-Vb) curves at RT for epitaxial 1T’-WTe2 contacted (navy) and Ti contacted (orange) 2H-WSe2 channel, and Ti contacted 1T’-WTe2 channel (black dotted line). Inset: Schematic energy band diagram for 2H-WSe2 with diverse metal including 1T’-WTe2. (c) Temperature-dependent Vg modulation of sheet conductivity (σ) at Vb = 50 mV for FET-1 (color scale) and FET-2 (grey scale) (d) Arrhenius plot of σ for FET-1 (navy) and FET-2 (orange) and Ti-contacted 1T’-WTe2 (grey). Solid dots : Vg = -30 V, hollow dots : Vg = 30 V. Figure 5. (a-b) Built-in potential energy (qΦbi) as a function of Vg for (a) FET-1 (1T’-WTe2 contact, navy), FET-2 (Ti contact, orange) and (b) FET-2 type on bare WSe2 with other metals, Sc (red) and Pd (blue). (c) Hole SBH of 2H-WSe2 channel for various metal work function. Epitaxial vdW metal contact corresponds to solid square (navy), planar metal contact corresponds to hollow red dots and line, and Schottky-Mott rule fitting for black line.


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Heo, H.; Sung, J. H.; Cha, S.; Jang, B.-G.; Kim, J.-Y.; Jin, G.; Lee, D.; Ahn, J.-H.; Lee, M.-J.; Shim, J. H.; Choi, H.; Jo, M.-H. Nat. Commun. 2015, 6, 7372. 15

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Tonndorf, P.; Schmidt, R.; Böttger, P.; Zhang, X.; Börner, J.; Liebig, A.; Albrecht, M.; Kloc, C.; Gordan, O.; Zahn, D. R. T.; Vasconcellos, S. M.; Bratschitsch, R. Opt. Express 2013, 21, 4908− 4916. ACS Paragon Plus Environment

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Figure 1. (a) Sequential growth scheme for heteroepitaxial stacking of 2H-WSe2 and 1T’-WTe2 MLs. Inset: optical microscope (OM) images of monolayer WSe2 (left) and bilayer WTe2/WSe2 (right) (b) Atomicresolution cross-sectional HAADF-STEM image with corresponding atomic model for 1T’-WTe2 (upper layer) and 2H-WSe2 (bottom layer) (c) OM image of heteroepitaxial stacked WTe2/WSe2 crystal. Inset: atomic force microscopy (AFM) height profile obtained from the crystal in image (d) PL spectra measured from the ML WSe2 region (navy) and stacked region (orange) at 4K. Inset: PL intensity mapping images from bare WSe2 (left upper) and WTe2/WSe2 heterostructures (right lower). (e) High-magnification HAADF-STEM image at the step edge. (f) Atomic-resolution HAADF-STEM image and (g) SAED pattern from WSe2 ML region. (h) Atomic-resolution HAADF-STEM image and (i) SAED pattern from WTe2/WSe2 stacked region 110x72mm (300 x 300 DPI)


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Figure 2. (a) Atomic configuration of the heteroepitaxial 2H-WSe2 and 1T’-WTe2 ML stacks (blue, red, and yellow spheres represent W, Te, and Se atoms, respectively) and (b) the simulated Z-contrast image based on the built atomic configuration of (a). (c) The experimental HAADF-STEM image from the stacked layers. The simulation and experimental HAADF-STEM images show the several interference patterns arising from two different lattices of 2H-WSe2 and 1T’-WTe2, which can be identified as representative “domains” in the nanometer scales. (d-f) Atomic models and simulated and experimental HAADF-STEM images with Zcontrast intensity profiles corresponding to each domains described in (d), (e), and (f), where the coherency between the top and bottom layers is induced by Wtop-Wbottom, Tetop-Wbottom, and Wtop-Sebottom as well as Tetop-Wbottom, respectively. Note that the stronger contrast results from the atomic channeling effect and thus the well-aligned array between the top and bottom layers gives rise to the higher yield of electron scattering in a STEM. 170x171mm (300 x 300 DPI)


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Figure 3. (a) False-color SEM image of patterned WTe2 film with conventional lithography, into FET device configuration. (b) Channel length dependent resistance of 1T’-WTe2 film. The contact resistance and sheet resistance are 4.2 kΩ·μm and 12.1 kΩ respectively. (c) Thickness dependent channel sheet resistances. Inset: thickness dependent contact resistances. (d) (left) Temperature dependence of the resistivity of 1T’WTe2 film of thickness (t) of 2.1 nm (red) and crystals of different t = 40 nm, (blue) and 4.8 nm (navy), taken at 2 K. The classical quadratic field dependence is observed in the thick sample, whereas the weak antilocalization effect, i.e. quantum correction to the 2D conductivity, is dominant for the thin sample. 169x52mm (300 x 300 DPI)


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Figure 4. (a) Illustration of our device scheme for comparison vdW epitaxial metal (1T’-WTe2) contact (FET-1) with planar metal (Ti) contact (FET-2) and optical microscope image of vdW epitaxial contact devices, where upper 4 metals for FET-1 and lower 3 metals for FET-2. (b) Current-voltage (I-Vb) curves at RT for epitaxial 1T’-WTe2 contacted (navy) and Ti contacted (orange) 2H-WSe2 channel, and Ti contacted 1T’-WTe2 channel (black dotted line). Inset: Schematic energy band diagram for 2H-WSe2 with diverse metal including 1T’-WTe2. (c) Temperature-dependent Vg modulation of sheet conductivity (σ) at Vb = 50 mV for FET-1 (color scale) and FET-2 (grey scale). (d) Arrhenius plot of σ for FET-1 (navy) and FET-2 (orange) and Ti-contacted 1T’-WTe2 (grey). Solid dots : Vg = -30 V, hollow dots : Vg = 30 V. 170x55mm (300 x 300 DPI)


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Figure 5. (a-b) Built-in potential energy (qΦbi) as a function of Vg for (a) FET-1 (1T’-WTe2 contact, navy), FET-2 (Ti contact, orange) and (b) FET-2 type on bare WSe2 with other metals, Sc (red) and Pd (blue). (c) Hole SBH of 2H-WSe2 channel for various metal work function. Epitaxial vdW metal contact corresponds to solid square (navy), planar metal contact corresponds to hollow red dots and line, and Schottky-Mott rule fitting for black line. 169x64mm (300 x 300 DPI)


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TOC 89x33mm (600 x 600 DPI)


Epitaxial van der Waals Contacts between Transition-Metal

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