Contacting MoS2 to MXene: Vanishing p-Type Schottky Barrier and

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C: Surfaces, Interfaces, Porous Materials, and Catalysis 2

Contacting MoS to MXene: Vanishing p-Type Schottky Barrier and Enhanced Hydrogen Evolution Catalysis Jinxuan You, Chen Si, Jian Zhou, and Zhimei Sun J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b12469 • Publication Date (Web): 21 Jan 2019 Downloaded from http://pubs.acs.org on January 22, 2019

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The Journal of Physical Chemistry

Contacting MoS2 to MXene: Vanishing p-type Schottky barrier and enhanced hydrogen evolution catalysis Jinxuan Youa , Chen Sia ,∗ Jian Zhoua , and Zhimei Suna,b† a

School of Materials Science and Engineering, Beihang University, Beijing 100191, China

b

International Research Institute for Multidisciplinary Science, Beihang University, Beijing 100191, China

Abstract It is a big challenge to make a Schottky-barrier (SB)-free hole contact to MoS2 with a high ionization of ∼6.0 eV. Here, using first-principles calculations, in a recently discovered large family of two-dimensional transition metal carbides or nitrides (MXenes), we have found six materials (V2 CO2 , Cr2 CO2 , Mo2 CO2 , V4 C3 O2 , Cr2 NO2 and V2 NO2 ) which can be used as metal contacts to monolayer MoS2 with vanishing p-type Schottky barriers at contacting interfaces, resulting in highly efficient hole injection into MoS2 . We reveal that the successful achievements of the SB-free hole contacts at these MoS2 /MXene interfaces depend on not only the high work functions of the MXenes but also the absence of the formation of interfacial gap states that usually result in strong Fermi level pinning in the midgap of semiconductor. We further propose that efficient charge injection into MoS2 facilitated by the SB-free contact could also increase the hydrogen evolution reaction (HER) activity of the 2H-MoS2 basal plane by improving its conductivity as well as its ability to adsorb hydrogen. Not only are these findings invaluable for designing high-performance MoS2 -based electronic devices, but they provide an effective route to optimize the MoS2 nanosheet catalysts for HER.

∗

Electronic address: [email protected]

†

Electronic address: [email protected]

1

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I.

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INTRODUCTION

Two dimensional (2D) transition metal dichalcogenide MoS2 has been attracting widespread attention in recent years, owing to its great application potential in future electronics[1]. A striking example is the successful fabrication of MoS2 transistors with 1-nm physical gate length, overcoming the scaling limit of 5-nm gate length for conventional Si transistors[2]. The semiconducting electronic devices based on MoS2 need to make use of fine control over the charge-carrier flow injected into MoS2 through metal contacts; hence the quantity of electric contacts directly dictates the device performances[3, 4]. However, contacting MoS2 with metals is still a problem, because it usually inclines to generate accidentally large contact barriers and resistances[5]. In this process, the Fermi level pinning arising from the strong MoS2 -metal interaction[6, 7] is uncovered to be an important factor triggering the formation of the large Schottky barrier. On the other hand, it is found experimentally that monolayer and multilayer MoS2 usually show n-type behavior when in contact with metals[3]. Common metal electrodes such as Sc, Ti, Au and Ni have been verified to form n-type contacts with MoS2 [8–10], while whether the n-type contact is formed between Pd (Pt) and MoS2 is currently under debate. First-principles calculations show the Pd/MoS2 contact is n-type[5, 6], which is supported by Kaushik et al.’s[10] and Neal et al.’s[11] experiments but opposite to Fontana et al.’s experiment which shows a p-type contact between Pd and MoS2 [9]. For Pt/MoS2 contact, it is predicted to be p-type theoretically[6] but found to be n-type experimentally by Saptarshi et al.[8]. In this context, it is desirable to develop low-resistance Schottky-barrier (SB)-free contacts to MoS2 , particularly barrierless p-type contacts. To create a p-type SB-free contact, the electrode work function (WF) should be no smaller than the ionization energy (IE) of the semiconductor. However, single-layer MoS2 possesses a quite high IE of 5.95 eV, which is larger than the work functions of all pure metals. Recently, a new family of 2D transition metal carbides or nitrides, so-called “MXenes” with a general formula of Mn+1 Xn Tx , where M is an early transition metal, X is C or N, Tx represents a surface functional group, and n=1, 2 or 3, are synthesized by selective etching of A layers from their parent layered compounds “Mn+1 AXn phases” (A=Al, Si, etc.) using a strongly acidic solution such as HF or a mixture of LiF and HCl[12, 13]. This process results in T = OH, O or/and F. Interestingly, MXenes show excellent electric conductivities, holding great

2


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potential as electrode materials[14, 15]. More importantly, various elemental compositional and surface functional possibilities and fine thickness controllability of MXenes make their work functions be able to vary in a wide range[16, 17], which yields a high possibility of finding electrode materials with higher WFs than the IE of MoS2 in the large MXene family. In addition, it should be stressed that a high work function for the electrode alone is not sufficient for attaining a vanishing p-type Schottky barrier height (SBH), and a moderate interfacial coupling to avoid strong Fermi level pinning is another indispensable factor. We notice that a recent theoretical work proposed that some MXenes such as Nb2 CO2 can realize barrierless hole injection into WSe2 [18]. However, the IE of MoS2 exceeds that of WSe2 by as large as 0.84 eV, and meanwhile, the interfacial electronic coupling between MoS2 and MXene would differ from that between WSe2 and MXene, leaving whether SB-free hole contacts to MoS2 can be formed by using MXenes as electrodes remaining an open question. In this paper, based on first-principles calculations, we sift out seven O-terminated MXenes (V2 CO2 , Cr2 CO2 , Mo2 CO2 , V4 C3 O2 , Cr2 NO2 , V2 NO2 and Ti3 C2 O2 ) possessing WFs larger than the IE of single-layer MoS2 from the large MXene family. A further study of the electronic and interfacial properties of the heterostructures constructed by monolayer MoS2 adsorbed on top of MXene shows that the former six MXenes mentioned above can form SB-free contacts to MoS2 without creating interfacial states in the midgap of MoS2 . By contrast, when Ti3 C2 O2 is in contact with MoS2 , their interfacial electric coupling produces gap states which pin the Fermi level of the heterostructure, giving rise to a sizable SBH. We further propose that the formation of SB-free contact could also enhance the catalytic activity of the 2H basal plane of MoS2 for the hydrogen evolution reaction (HER). This is because efficient charge injection into MoS2 facilitated by the SB-free contact will significantly increase the electric conductivity of the 2H basal plane and its ability to adsorb hydrogen. Our work demonstrates the great potential of MXenes as favorable p-type SB-free contacts to MoS2 which can facilitate not only the realizations of MoS2 field-effect transistors with both polarities but also the optimizations of the HER activities of MoS2 nanosheet catalysts.

II.

COMPUTATIONAL METHODS

All the calculations are performed in the framework of density functional theory (DFT)[19] with the projected augmented wave (PAW)[20] formalism as implemented in the 3



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FIG. 1: (Color online) (a) Work functions (denoted by dots) of MXenes with O terminations compared with the ionization energies (denoted by dashed lines) of monolayer MoS2 , MoSe2 , MoTe2 , WS2 and WSe2 . (b) Representative structure of MXene shown by using Cr2 NO2 as an example.

Vienna ab initio simulation package (VASP)[21]. The exchange and correlation effects are described with the generalized gradient approximation (GGA) of Perdew-Burke-Emzerhof (PBE)[22]. A kinetic energy cutoff of 500 eV is set for the plane wave expansion. A vacuum region thicker than 15 ˚ A is employed to avoid the interaction between periodic replicas along the z direction. All the structures are fully relaxed till the residual force on each atom is less than 0.01 eV/˚ A and an energy convergence threshold of 10−5 eV is adopted for electronic optimization. Γ-centered k-point meshes with sufficient k-point densities are used for the Brillouin-zone integrations. Van der Waals interaction is taken into account by the DFTD2 approach[23]. Considering the asymmetric geometries of the hetero-bilayers, the dipole correction is added to cancel the errors of electrostatic potential, total energies and forces introduced by the periodic boundary conditions[24].

III.

RESULTS AND DISCUSSION

MoS2 , WS2 , MoSe2 , WSe2 and MoTe2 are the most-common five semiconducting 2D transition metal dichalcogenides (TMDCs). All of them are thermodynamically stabilized in the 2H structure which is a sandwich of triple layers of 2D hexagonally packed atoms, X (chalcogen)-M (transition metal)-X, forming Bernal (ABA) stacking[1]. Their calculated ionization energies are displayed in Figure 1a. It is seen that as the atomic number of M or 4


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X decreases (i.e., M changes from W to Mo, or X changes from Te to Se to S), the IE of MX2 increases. As a consequence, among the five TMDCs, MoS2 possesses the largest IE of 5.95 eV, even larger than the WF of Pt[18] (Pt has the highest WF in all elemental metals). On the other hand, MoS2 also has the highest electron affinity in the above five TMDCs (data not shown)[25]. These results explain why in experiments MoS2 is usually n-type[3] while other TMDCs such as WSe2 can show n-type[26, 27] or p-type[28, 29] transport depending on the choice of the contacting metal. Next, we are going to perform high-throughput calculations of work functions for all MXenes to find out candidate electrode materials whose WFs are larger than the IE of MoS2 . Given that MXenes are usually fabricated by extracting “A” layers from their parent MAX phases, here we focus on the MXenes that have existing MAX precursors[30], and in fact most of them have been synthesized recently[15]. As the MXenes with fully saturated surfaces are thermodynamically more favorable than those with partially saturated surfaces[31], we consider the MXenes with surfaces fully terminated by the possible functional groups which have the formula Mn+1 Xn T2 (T = OH, O or F). Figure 1b displays a representative structure of Mn+1 Xn T2 , where, M and X layers alternate with each other in a ABCABC· · · stacking order, and the most stable adsorption sites of T are determined by comparing the total energies of all possible adsorption configurations. After a global screening of WFs for all MXenes, we discover that with the change of the surface functional group, WFs show the following trend: for a given Mn+1 Xn , WF(Mn+1 Xn (OH)2 ) < WF(Mn+1 Xn F2 ) < WF(Mn+1 Xn O2 ). We find seven MXenes with WFs larger than IE of MoS2 , which are V2 CO2 , Cr2 CO2 , Mo2 CO2 , Ti3 C2 O2 , V4 C3 O2 , Cr2 NO2 and V2 NO2 (see Figure 1a, where, for clarity only the WFs of Mn+1 Xn O2 are given, but those of Mn+1 Xn F2 and Mn+1 Xn (OH)2 are not given because they are all smaller than IE of MoS2 ). Noted in Figure 1a, Ti2 CO2 , Zr2 CO2 and Hf2 CO2 are semiconductors[16], and the work function values shown for them are defined as the energy differences between the vacuum levels and the valence band maximums, which reflect the work functions in the cases of small p-type doping of these MXenes. With the seven candidate MXene materials Mn+1 Xn O2 at hand, we can construct the corresponding heterostructures which consist of a monolayer MoS2 adsorbed on a monolayer Mn+1 Xn O2 so as to investigate the band alignments and Schottky barriers at interfaces. Considering that MoS2 and Mn+1 Xn O2 possess different lattice constants, particular care is required in the building of the structural model for the heterostructure to make the lattice 5



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FIG. 2: The structure diagram of supercell heterojunction MoS2 /Cr2 NO2 . (a) Top view of 3 × 3 √ √ MoS2 on the top of 13× 13 R13.6◦ Cr2 NO2 as modeled in the supercell approach. The construction of the supercell basis vectors ma1 + na2 and pb1 + qb2 are illustrated and the parallelogram indicates the surface supercell. (b) Side view of the MoS2 /Cr2 NO2 heterostructure.

strain minimized. Let us label the base vectors of the MoS2 and Mn+1 Xn O2 unitcells as {a1 , a2 } and {b1 , b2 }, respectively. Their supercells basis vectors can be built severally as {ma1 + na2 , -na1 + (m + n)a2 } and {pb1 + q b2 , -qb1 + (p + q)b2 }, where, m, n, p, q are integers. Then we search for a series of integers so that the lattice constant of the MoS2 supercell can well match that of the Mn+1 Xn O2 : |ma1 + na2 | ≈ |pb1 + qb2 |. In practice, the commensurable smallest supercell satisfying that the lattice mismatch is smaller than 1% is chosen. As the electronic properties of the semiconducting MoS2 are sensitive to the external strain, we fix the lattice constant of MoS2 and slightly adjust the lattice of metallic Mn+1 Xn O2 to compensate the lattice mismatch. The selected lattice models for the seven MoS2 /Mn+1 Xn O2 heterostructures are given in Table 1 together with the rotation angles (θ) between the two constituent monolayers in the heterostructures as determined by θ = cos−1 ( √

√

mp+nq+(mq+np)/2

m2 +n2 +mn×

p2 +q 2 +pq

). Taking the MoS2 /Cr2 NO2 heterostructure as an example,

{m = 3, n = 0} and {p = 3, q = 1} are sought out, i.e., a 3 × 3 MoS2 lattice is put on top √ √ of a 13 × 13 Cr2 NO2 lattice with a rotation angle of 13.6◦ (see Figure 2). The binding energy Eb between MoS2 and Mn+1 Xn O2 in the heterostructure is defined as Eb =

E(MoS2 ) + E(Mn+1 Xn O2 ) − E(MoS2 /Mn+1 Xn O2 ) . A

Here, E(MoS2 ) and E(Mn+1 Xn O2 ) are the energies of the isolated MoS2 and Mn+1 Xn O2 monolayers, respectively, and E(MoS2 /Mn+1 Xn O2 ) is the total energy of the bilayer 6


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FIG. 3: (a) (Color online) Band structure of the MoS2 /Cr2 NO2 heterostructure with the Fermi level set to zero. The red dots represent the projection of the wave function on the MoS2 orbitals. (b) The plane averaged electron density difference ∆ρ(z) = ρMoS2 /Cr2 NO2 (z) − ρMoS2 (z) − ρCr2 NO2 (z). (c) The plane-averaged electrostatic potential of MoS2 /Cr2 NO2 along the z direction. The potential step ∆V across the vacuum is shown. The interface is denoted by the vertical dashed line. (d) The schematic illustration of the band alignment of the MoS2 /Cr2 NO2 heterostructure.

MoS2 /Mn+1 Xn O2 heterostructure. A positive value for Eb indicates energy gain upon the formation of the heterostructure. The calculated binding energies for the seven heterostructures we study here are listed in Table 1. It is seen that the binding energy between MoS2 and Mn+1 Xn O2 ranges from 15 to 24 meV/˚ A2 , lying in the range of van der Waals (vdW) interaction. In a metal-semiconductor (MS) heterostructure, the SBH is the most important physical quantity controlling the transport of charge carriers across the MS interface[32]. It is defined as the energy difference between the Fermi level and the band edge of the semiconductor in the junction: Φh = EF - EVBM , Φe = ECBM - EF , where Φh and Φe are SBHs for holes and electrons injected from the metal to the semiconductor, respectively, EF is the Fermi-level

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energy of the MS heterostructure, EVBM and ECBM denote the energies of the valence band maximum (VBM) and conduction band minimum (CBM) of the semiconductor[32, 33]. Thus, the value of SBH can be predicted through the band structure calculation for the MS heterostructure, where the band edges of the semiconductor can be identified by the projections of total electronic states onto given atomic orbitals. When SBH approaches zero or is even negative, the MS contact could be regarded as Ohmic. Using the MoS2 /Cr2 NO2 heterostructure as an example, its calculated band structure is shown in Figure 3a, where the sizes of red circles represent the weights of MoS2 orbitals. Remarkably, the Fermi level of the heterostructure is located below the VBM of MoS2 , resulting in a negative Φh of -0.102 eV. This suggests holes are spontaneously injected from Cr2 NO2 to MoS2 upon their contact. Additionally, in the heterostructure the direct band gap of MoS2 is preserved, yet the band edges (CBM and VBM) are folded onto the Γ point in the reciprocal space of the 3 × 3 superlattice of MoS2 from the original K point of the 1 × 1 unitcell of isolated MoS2 . At the same time, the size of band gap is decreased from 1.67 eV in isolated MoS2 to 1.47 eV in the heterostructure. This is because the hybridization between the electronic states of Cr2 NO2 with those of MoS2 induces the shifts of the band edges of MoS2 . It is noted that though the interfacial vdW interaction is weak, it can rearrange the charge densities at interface. Figure 3b shows the differential charge density ∆ρ(z) along the z direction perpendicular to the MoS2 /Cr2 NO2 interface, which is defined as ∆ρ(z) = ρMoS2 /Cr2 NO2 (z) − ρMoS2 (z) − ρCr2 NO2 (z). Here, ρMoS2 /Cr2 NO2 (z), ρMoS2 (z), and ρCr2 NO2 (z) are the plane-averaged charge densities of the MoS2 /Cr2 NO2 heterostructure, the isolated MoS2 and Cr2 NO2 , respectively. ∆ρ(z) unambiguously reveals an accumulation of electrons at the interface region close to MoS2 and an electron depletion at the interface region near Cr2 NO2 . Such charge redistribution results in the electron wave polarization and thus creates an interface dipole pointing from MoS2 to Cr2 NO2 . This interface dipole is accompanied by a potential step (∆V ) formed across the interface, which will induce the shift of the vacuum level. From the plane-averaged electrostatic potential plotted in Figure 3c, one can indeed observe a discontinuity of 1.05 eV between the vacuum levels on the MoS2 and Cr2 NO2 sides. The presence of ∆V will significantly modify the band alignment at interface (see Figure 4d), resulting in the p-type SBH approximately following the relation Φh = IE − W F + ∆V [32, 34], where IE (W F ) is the ionization energy (work function) of the semiconductor 8


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TABLE I: The supercell models (the supercells for MoS2 and MXene layers and the rotation angles (θ) between the supercells) and the calculated potential steps ∆V (eV), binding energies (eV·˚ A−2 ) and p-type SBHs Φh (eV) for all the considered MoS2 /MXene heterostructures. heterostructure supercell 1 supercell 2

θ

∆V

Eb

Φh

MoS2 /V2 CO2

√ √ 13 × 13

4×4

13.9◦

0.80

0.018

0.052

MoS2 /Cr2 CO2

√ √ 3× 3

2×2

29.9◦

1.75

0.025

0.005

MoS2 /Mo2 CO2

√ √ 13 × 13

4×4

13.9◦

1.69

0.024

-0.185

√ √ MoS2 /Ti3 C2 O2 2 3 × 2 3

3×3

16.1◦

0.30

0.015

0.245

MoS2 /V4 C3 O2

4×4

√ √ 19 × 19

23.4◦

0.70

0.019

0.027

MoS2 /Cr2 NO2

3×3

√ √ 13 × 13

22.4◦

1.05

0.019

-0.102

MoS2 /V2 NO2

4×4

√ √ 19 × 19

23.4◦

0.18

0.016

0.022

(metal) constituting the heterostructure. The sign and size of the ∆V are consistent with the orientation and magnitude of the interface dipole. Here, an interface dipole with the orientation pointing from MoS2 to MXene corresponds to a positive ∆V . Similar charge redistribution and ∆V formation have also been observed in some vdW heterostructures such as the 2H-MoS2 /1T-MoS2 heterostructure where the 1T-MoS2 forms a p-type contact with 2H-MoS2 , with a small Φh [35], well explaining the low contact resistance in the MoS2 field effect transistors with 1T-MoS2 electrodes. We further calculate the band structures of the other six MoS2 /Mn+1 Cn O2 heterostructures and determine the corresponding Φh which are listed in Table 1. Interestingly, it is found that p-type SB-free contacts are also achieved at MoS2 /V2 CO2 , MoS2 /Cr2 CO2 , MoS2 /Mo2 CO2 , MoS2 /V4 C3 O2 and MoS2 /V2 NO2 interfaces with Φh quite close to zero or even negative. In these vdW heterostructures, the interface charge redistributions which are constituted by the charge depletion in the interface region near MoS2 and the charge accumulation in the interface region near MXene are universally observed, causing the formation of the positive potential step ∆V whose value ranges from 0.18 to 1.75 eV depending on the type of MXene. It should be mentioned that the p-type SBH values listed in Table 1 are calculated in the framework of PBE. The PBE functional is well-known to be suffered from band-gap underestimation problem, which gives a band gap of 1.67 eV for monolayer MoS2 , 9



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smaller than the 1.8 eV obtained experimentally[36]. Although generally PBE cannot give reliable electronic properties and energies of the excited states, it is believed to be able to well describe the valence electronic states. Therefore, the p-type SBH, the energy difference between the VBM of the semiconductor and the Fermi level of the metal, calculated by PBE, is considered to be relatively reliable[37, 38]. We also considered the effects of spinorbital coupling (SOC) on the p-type SBH. We find that the SOC-triggered energy splitting of valence band maximums of MoS2 in the MoS2 /MXene heterostructures is minor, and accordingly, the inclusion of SOC only introduces slight modifications to the SBH. Taking MoS2 /Cr2 NO2 and MoS2 /Cr2 CO2 as examples, their p-type SBHs without (with) SOC are calculated to be -0.102 (-0.091) and 0.005 (0.007) eV, respectively. The conclusion about the the formation of p-type SB-free contacts are not changed after taking SOC into account. Additionally, in Table 1 it is noticed that MoS2 /Ti3 C2 O2 heterostructure peculiarly shows a sizeable Φh of 0.25 eV at interface. To explore why Ti3 C2 O2 fails to form a SB-free contact with MoS2 , we carefully analyze the electronic properties of the MoS2 /Ti3 C2 O2 heterostructure, as shown in Figure 4a. The CBM and VBM of MoS2 in this heterostructure are both located at the Γ point, with a gap size of 1.69 eV, quite close to the gap size of 1.67 eV in the isolated MoS2 . Meanwhile, it is interesting to observe the spilling of the valence states of MoS2 into the gap region, forming gap states. The presence of gap states results in the pinning of the Fermi level in the band gap, producing a sizable Schottky barrier. A further real-space charge density analysis shows that the gap states around the Γ point near the Fermi level are composed of both the electronic states of MoS2 and those of Ti3 C2 O2 (see Figure 4b), in marked contrast to the VBM of MoS2 that is totally contributed by the MoS2 electronic states (see Figure 4c). This indicates that the gap states arise from the interfacial electronic states coupling between MoS2 and Ti3 C2 O2 . To clearly illustrate the role of gap states, we have drawn a simplified sketch for the band alignment of the MoS2 /Ti3 C2 O2 heterostructure in Figure 4d. The holes are transferred from the interface region near Ti3 C2 O2 to the interface region near MoS2 , and the gap states serve as a reservoir for holes, thus finally pinning the Fermi level. As is widely known, the gap states are commonly formed in the traditional heterojunctions with interfacial covalent bonding[32]. Here, the case of MoS2 /Ti3 C2 O2 suggests that the possibility of forming gap states in the vdW heterostructures can not be completely excluded. Since the gap states are originated from the interfacial electronic coupling, they may be 10


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FIG. 4: (Color on line) The electronic properties of MoS2 /Ti3 C2 O2 heterostructure. (a) Band structure with the Fermi level set to zero. (b)-(c) The real-space charge-density distributions of the gap states (b) and the VBM (c) at the Γ point. The isosurface value is set to 0.005 e·A−3 . (d) The schematic illustration of the band alignment.

suppressed by increasing the separation between metal and semiconductor. If the interlayer distance is increased to an artificially large distance d = 6 ˚ A, the interfacial hybridization is negligible, and the gap states disappear. However, the interfacial electron transfer from the interface region near MoS2 to the interface region near Ti3 C2 O2 is not forbidden, as the work function of Ti3 C2 O2 of 6.1 eV is larger than the MoS2 ionization energy of 5.95 eV. This charge transfer equilibrates the Fermi level, generating a nearly zero Φh . So far, we have sifted out six potential electrode materials, V2 CO2 , Cr2 CO2 , Mo2 CO2 , V4 C3 O2 , Cr2 NO2 and V2 NO2 from the MXene family, which could form p-type SB-free contacts to MoS2 . It is worth mentioning that such contacts are not only beneficial to the optimization of the electronic devices based on MoS2 but also able to help to increase the electrocatalytic activity of the 2H basal plane of MoS2 for the hydrogen evolution reaction. MoS2 which thermodynamically favors the 2H phase has been identified as a promising catalyst for HER in acidic media[39]. Nevertheless, only the exposed metallic edge sites of 2H-MoS2 are catalytically active, while its basal plane is rather inert[40]. A main factor limiting the activity of the 2H basal plane is its semiconducting nature, which hinders the charge-transfer kinetics required in HER. Recent works have shown that the HER activity of the basal plane can be significantly improved by converting the semiconducting 2H phase 11



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FIG. 5: (a) The schematic of the MoS2 hydrogen evolution reaction (left panel) and the metallic energy-band feature of MoS2 induced by the p-type SB-free contact (right panel). The shaded areas in the right panel represent the occupied electronic states of MoS2 . (b) The hydrogen adsorption energies ∆EH for 2H-MoS2 , 1T-MoS2 and 2H-MoS2 /MXene heterostructures. The numbers in the brackets above the x axis are corresponding H coverage rates.

into the metallic 1T phase to enhance the conductivity[41]. However, the 2H-to-1T phase transition is usually local. Here, using a suitable substrate to achieve a low-resistance contact to 2H-MoS2 with a vanishing Schottky barrier offers another feasible way to increase the electric conductivity of the basal plane. When MXene forms a p-type contact to 2H-MoS2 with Φh ≤ 0, as shown by the band structure schematics of 2H-MoS2 in Figure 5a, the Fermi level will be just at or below the VBM of 2H-MoS2 , resulting in a metallic nature for 2H-MoS2 which allows rapid charge transport to reaction sites and thus ensures efficient HER catalysis. Besides facilitating charge-transfer kinetics, the metallic nature of MoS2 induced by the SB-free contact at MoS2 /MXene interface enables hydrogen to adsorb easily on the basal plane of 2H-MoS2 . In the HER process, the binding energy of the H atom on the catalyst 12


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surface is an important energy scale determining the overall HER efficiency[42]. Using metallic 1T-MoS2 possessing good catalytic activity as a reference system, we compared the H adsorption on free-standing 2H-MoS2 and MoS2 /MXene heterostreuctures with that on 1T-MoS2 . Experimentally it has been found that the active sites of 1T-MoS2 nanosheets are mainly located on the basal plane rather than at the edges[43], and subsequent theoretical studies further indicate that catalytic activity of the 1T basal plane is mainly originated from its affinity for binding H atoms on the surface sulfur sites[44]. So for comparison and simplicity, here we considered a single H atom adsorption on the surface S site of 1T-MoS2 , 2H-MoS2 and 2H-MoS2 /MXene systems with same supercell sizes. The H adsorption energy is defined as ∆EH = E(catalyst + H) + E(catalyst) − E(H), where E(catalyst + H) and E(catalyst) are the total energies of the catalyst (e.g. the MoS2 /MXene heterostructure) with and without a H adatom on the surface, and E(H) represents the energy of an isolated H atom. The obtained ∆EH for the heterostructure, free-standing 2H-MoS2 and 1T-MoS2 are shown in Figure 5b. For the six SB-free heterostructures that we considered here, as they have different supercell sizes, a single H adsorption corresponds to different coverage ranging from 1/16 to 1/9. However, a common feature for them is that the values of ∆EH for the heterostructures are substantially reduced from the values for the free-standing MoS2 , becoming comparable to those for 1T-MoS2 (see Figure 5b). This indeed indicates that the ability of the 2H basal plane to adsorb hydrogen is dramatically improved in the MoS2 /MXene heterostructures. In fact, a recent experimental work has demonstrated that the 2H basal plane of monolayer MoS2 can be as catalytically active as the 1T phase by reducing the contact resistance to facilitate charge injection from the electrode to MoS2 [45]. Additionally, it was found that the S vacancies could be formed in the basal plane of 2HMoS2 under highly reducing potentials[46] and the HER activity of the basal plane can also be enhanced by the introduction of S vacancies[46, 47]. So we further consider the influences of MXene contacts on the H adsorption on the S vacancy site of the basal plane. Taking the MoS2 /Cr2 CO2 contact as an example, the EH for the H adsorption on a single S vacancy of MoS2 basal plane is reduced from -2.06 eV before contact to -3.08 eV after contact, which indicates that the H adsorption ability of the S-vacancy site are also strengthened in the heterostructure. Similar decreasing trends of EH for the H adsorption on the S vacancy site are also found for other five MoS2 /MXene contacts considered in Figure 5b.

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IV.

Page 14 of 19

CONCLUSIONS

In summary, on the basis of first-principles calculations, we have found six O-terminated MXenes (V2 CO2 , Cr2 CO2 , Mo2 CO2 , V4 C3 O2 , Cr2 NO2 and V2 NO2 ) from the large MXene family, which could be used as metal contacts to monolayer MoS2 with vanishing p-type Schottky barriers at the contacting interfaces, leading to highly efficient hole injection into MoS2 . The high work functions of these MXenes which are larger than the ionization energy of MoS2 , together with the absence of the formation of interfacial gap states which usually strongly pin the Fermi level in the midgap of semiconductor, ensure the realizations of the SB-free hole contacts. We further point out that the SB-free contact is also conductive to enhance the catalytic activity of the 2H basal plane of MoS2 for HER, because it can significantly increase the conductivity of the basal plane as well as its ability to adsorb hydrogen by facilitating charge injection into MoS2 . Our work suggests that the 2D metallic MXenes with widely tunable work functions are good supplements to the conventional 3D metals on which MoS2 usually exhibits n-type behaviors with large SBHs.

V.

ACKNOWLEDGMENTS

This work is financially supported by the National Research and Development Program of China (Grant No. 2017YFB0701700) and the National Natural Science Foundation of China (11504015, 11874079 and 51871009).

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TABLE OF CONTENTS (TOC)

19


Contacting MoS2 to MXene: Vanishing p-Type Schottky Barrier and

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