Influence of strain and metal thickness on metal-MoS2 contacts Wissam A. Saidi Citation: The Journal of Chemical Physics 141, 094707 (2014); doi: 10.1063/1.4893875 View online: http://dx.doi.org/10.1063/1.4893875 View Table of Contents: http://scitation.aip.org/content/aip/journal/jcp/141/9?ver=pdfcov Published by the AIP Publishing Articles you may be interested in Substrate induced modulation of electronic, magnetic and chemical properties of MoSe2 monolayer AIP Advances 4, 047107 (2014); 10.1063/1.4871080 High-performance MoS2 transistors with low-resistance molybdenum contacts Appl. Phys. Lett. 104, 093106 (2014); 10.1063/1.4866340 Ballistic performance comparison of monolayer transition metal dichalcogenide MX2 (M = Mo, W; X = S, Se, Te) metal-oxide-semiconductor field effect transistors J. Appl. Phys. 115, 084506 (2014); 10.1063/1.4866872 Strain-induced magnetism in MoS2 monolayer with defects J. Appl. Phys. 115, 054305 (2014); 10.1063/1.4864015 Characterization of metal contacts for two-dimensional MoS2 nanoflakes Appl. Phys. Lett. 103, 232105 (2013); 10.1063/1.4840317

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THE JOURNAL OF CHEMICAL PHYSICS 141, 094707 (2014)

Influence of strain and metal thickness on metal-MoS2 contacts Wissam A. Saidia) Department of Mechanical Engineering and Materials Science, University of Pittsburgh, Pittsburgh, Pennsylvania 15261, USA

(Received 11 June 2014; accepted 12 August 2014; published online 5 September 2014) MoS2 and other transition metal dichalcogenides are considered as potential materials in many applications including future electronics. A prerequisite for these applications is to understand the nature of the MoS2 contact with different metals. We use semi-local density functional theory in conjunction with dispersion corrections to study the heterostructures composed of Pd and Pt monolayers with (111) orientation grown pseudomorphically on MoS2 (001). The interface properties are mapped as a function of the number of deposited overlayers, as well as a function of tensile and compressive strains. Although we show that the dependence of the contacts on strain can be fully explained using the d-band model, we find that their evolution with the number of deposited metal layers is markedly different between Pd and Pt, and at variance with the d-band model. Specifically, the Pt/MoS2 heterostructures show an anomalous large stability with the deposition of two metal monolayers for all investigated strains, while Pd/MoS2 exhibits a similar behavior only for compressive strains. It is shown that the results can be rationalized by accounting for second-nearest-neighbor effect that couples MoS2 with the subsurface metal layers. The underpinnings of this behavior are attributed to the larger polarizability and cohesive energy of Pt compared to Pd, that leads to a larger charge-response in the subsurface layers. © 2014 AIP Publishing LLC. [http://dx.doi.org/10.1063/1.4893875] INTRODUCTION

There is a growing interest in two-dimensional (2D) materials that are characterized by a thickness of a single atom such as graphene and boron-nitride, or single polyhedra such as the transition metal dichalcogenides (TMDs). These materials offer unique opportunities for applications by virtue of their exotic electrical, optical, and mechanical properties.1–4 Graphene is the most investigated material in the 2D family eventhough its realm of applications is somewhat limited owing to its zero bandgap. Recently, MoS2 from the TMD family started attracting a lot of attention considering that unlike graphene, MoS2 has a relatively large bandgap ≈1.8 eV5 allowing for field-effect devices with a low off-current.2, 6–10 A prerequisites for future electronic applications based on 2D materials is to understand and control their interfaces that are formed with different metals.11, 12 Considering that graphene is the prime example of 2D materials, it is not surprising that metal/graphene contacts have been extensively investigated. Based on previous studies, these contacts can be classified as having a weak interface energy (physiadsorbed case) such as those formed between graphene and Ir,13 Pt,14, 15 Au, or Cu,16 and heterostructures where graphene has a more significant adhesion (chemisorbed case) to the metal such as Ni,13, 17, 18 Ru,19 Co,20 and Pd.15, 21 Generally, only in the chemisorbed case, the electronic interactions with the graphene substrate significantly perturb its electronic structure such as the Dirac cones. Compared to graphene/metal systems, less is known about metal/MoS2 contacts. Recently, it was demonstrated usa) Author to whom correspondence should be addressed. Electronic mail:

[email protected].

0021-9606/2014/141(9)/094707/8/$30.00

ing a wet-chemical synthesis approach that MoS2 nanosheets can be used to direct the epitaxial growth for Pd, Pt, and Ag nanostructures at ambient conditions. A wet-chemical synthesis techniques approach is very advantageous for the epitaxial growth of metallic structures on bulk substrates by virtue of its relatively low cost and high throughput compared to other conventional deposition techniques.22 On the computational side, Popov et al.23 showed using density functional theory (DFT) in conjunction with the local density approximation (LDA) that the bonding in Au and MoS2 interface is not strong enough to enable a large electron carrier mobility in Au/MoS2 metal contacts, while as the bonding in Ti/MoS2 is strong enough suggesting that Ti is a better material as an electrode in these contacts. To provide an in-depth understanding of metal-MoS2 interfaces, here we investigate the contacts formed between MoS2 and the two metals Pd and Pt that are expected to show drastic differences based on the graphene analogs.14, 15, 21 For example, Wang and Che15 employing DFT/LDA calculations showed that Pd/graphene interface is stabilized by electron transfer from electrons in graphene π orbitals to dxz + dyz orbitals of Pd that is compensated by a charge depletion in Pd dz2 orbital, while as the electron repulsion between the graphene π electrons and Pt occupied valence-band makes the Pt/graphene interface less stable. Furthermore, considering the importance of miniaturization in applications, we study the interface as a function of the number of metal layers deposited on the MoS2 substrate. Additionally, we investigated the dependence of the heterostructures on tensile and compressive strains. This is important considering that geometric strain effects due to lattice mismatch, can have a profound effect on the contact as demonstrated before in terms

141, 094707-1

© 2014 AIP Publishing LLC

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094707-2

Wissam A. Saidi

of catalytic properties of thin films. Recently, the effects of strain on reaction rates were even experimentally isolated and quantified.24 Experimentally, two types of epitaxial orientations (111) and (101) of the metallic nanoparticles Pd, Pt, and Ag are observed on MoS2 (001) substrate.22 In this study, we investigate the (111) orientation of the metals, which have the same crystalline symmetry as MoS2 . We show that the dependence of the metal/MoS2 heterostructures on strain follows the d-band model of Hammer and Nörksov,45 however, the dependence of the interface on the number of layers showed significant differences between the two metals. The results can be understood due to a charge-transfer effect, which induces charge oscillations in the sublayers that influence the interfacial properties. The underpinnings for the differences between Pd and Pt can be traced back to the larger cohesive energy and polarizability of Pt compared to Pd. COMPUTATIONAL DETAILS AND METHODOLOGY

The ab initio calculations are performed using DFT as implemented in PWSCF.25 Core electrons are replaced by norm conserving pseudopotentials. We used the generalizedgradient approximation by Perdew, Burke, and Ernzerhof26 (PBE) in all of our investigations in conjunction with dispersion DFT-TS27 corrections using the implementation of Ref. 28 and reference atomic densities obtained using selfconsistent Hartree-Fock wavefunctions.29 We employed the parameterization of Ref. 30 to account for non-local Coulomb screening within the metal bulk, which has been applied recently in several studies of metal surfaces.31–35 During geometry optimization, all atoms are relaxed using a convergence threshold of 10−4 Ry/bohr on the atomic forces, and 10−7 Ry on the energies in the self-consistent step. We included a smearing of 0.01 Ry to aid in the convergence of the electronic self-consistent step. The Kohn-Sham orbitals of the valence electrons and the corresponding electron densities are expanded in plane waves with a cutoff of 50 and 200 Ry, respectively. As a validation of the computational approach, we find that the optimum lattice constants for Pd and Pt are 3.95 and 3.94 Å that are in good agreement with experimental values 3.89 and 3.92.36 Also, our computed values are in good agreement with previous PBE calculations 3.98 and 3.99 Å for Pd and Pt, respectively.37 The integration over the Brillouin zone (BZ) for the heterostructures is performed using a 2 × 2 × 1 -centered Monkhorst-Pack (MK) grid resulting in 5 k-points in the irreducible BZ. The analysis of the electronic properties including density-of-states, electron density plots, etc., are based on Kohn-Sham wavefunctions obtained from a self-consistent solution with a 6 × 6 × 1 -centered MK grid, amounting to 20 k-points in the irreducible BZ. The DOS, and atomprojection and angular-momentum projected DOS are generated using a Lorentzian with a half-width at half-maximum value of 0.025 eV. Motivated by the recent experiment of Huang et al.,22 the metal/MoS2 overlayer structures are modeled by a metallic slab of (111) orientation grown pseudomorphically on MoS2 (001) surface. The metal slab has a thickness that ranges

J. Chem. Phys. 141, 094707 (2014)

between 1 and 5 monolayers (ML), and with a hexagonal ABCABC stacking sequence with an inter-layer spacing set √ initially to a/ 3, where a is the lattice constant of the hexagonal cell. In our approach that employs 3D periodic boundary conditions, the 2D slab is realized using a supercell, where the fictitious interactions between images along the non-periodic directions are mitigated by using more than 15 Å of vacuum in the direction perpendicular to the slab and in conjunction with the dipole-correction.38 The MoS2 layer is represented using a 3 × 3 surface supercell (27 atoms),√where √ as the (111) metal surface is represented using a (2 3 × 2 3) R30◦ supercell (12 atoms per layer). The supercell used to describe the metal/MoS2 interface is similar to the one used to describe heterostructures formed between Pd and Pt with graphene.15 With this supercell, and from the optimum lattice constants of 3.18, 3.95, and 3.94 Å for MoS2 , Pd, and Pt, this would result in a lattice mismatch of 2.7% and 3.0% for Pd and Pt, respectively. Additionally, to study the effects of the lattice mismatch, the surface slab is subjected to a biaxial in-plane strain. The lattice constant for the strained supercells are obtained by multiplying the MoS √ 2 equilibrium lattice constant a0 = 3.18 Å by a factor of 1 + e where e = −0.03, −0.015, 0.015, and 0.02. Because the number of metal layers is small, it is expected that the overlayers will adapt the lattice structure of the substrate inducing either tensile or compressive strain in the surface metal.39 The stability of the heterostructure is gauged using the interface adhesion energy defined as E = EMoS + Emetal − Emetal/MoS . 2

2

(1)

Here, EMoS and Emetal are the energies of the free-standing 2 optimized MoS2 and metal slabs, and Emetal/MoS is the en2 ergy of the heterostructure. The contributions from van der Waal (vdW) interactions are assessed by partitioning E into a short-range ESR as described by the PBE exchangecorrelation functional, and long-range vdW attraction EvdW , i.e., E = EvdW + ESR .

(2)

The SR contribution is computed by using the optimum DFTTS geometries of the interface, MoS2 , and metal systems, and calculating E via Eq. (1) using self-consistent energies from the PBE only. The long-range vdW contribution is treated semi-empirically using a damped sum of pair-wise dispersion terms between atom pairs. Nevertheless, the bonding hybridization state of the atoms are accounted for by using a C6 coefficient that depends on their chemical environment as measured with a Hirshfeld partitioning40 scheme of the electronic density.27 This approach has been applied to several systems yielding results that are in good agreement with experiment for semiconducting41–43 and metallic systems.31–35 RESULTS AND DISCUSSION

Figures 1(a) and 1(b) show top and side views of (Pd)1 /MoS2 (the heterostructures will be refereed to as (M)n /MoS2 where n is the number of layers of the metal M). Other interfaces are very similar except that the rumpling of

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J. Chem. Phys. 141, 094707 (2014) TABLE I. dMS is the distance between the top metal surface and the closest sulfur layer. δ is the rumpling of the top metal layer. dMM is the average distance between metal layers (shown for 3 ML case only). Results for 5 ML are not shown as these are very close to the 3 ML values. All distances are in Å. 1 ML

FIG. 1. Top (a) and side (b) views for the (Pd)1 /MoS2 , and (c) side view for (Pt)1 /MoS2 where all the heterostructures are under 0% strain. The position of the Mo atoms in registry with the metal highly symmetric sites top (T), fcc (F), and hcp (H) are indicated on the Mo atoms. Note the displacements δ of the metal atoms that are atop of the Mo atoms which are towards Mo layer as in (b) and away from the Mo layer as in (c). The solid lines show the hexagonal supercell. The S, Mo, Pd, and Pt atoms are shown as gold, red, green, and blue spheres, respectively.

the metal layer δ is reduced significantly (see Table I). δ is defined as the difference between the smallest and largest displacements from the two-dimensional nuclear framework of the metal. The optimum structure of the interface can be conveniently described using the high symmetry sites top (T), hcp (H), fcc (F), and bridge (B) of the fcc (111) lattice. In the equilibrium configuration, the nine Mo atoms in the supercell are ¯ direction that are in registry with arranged in rows along [110] H-T-F, F-H-T, T-F-H sites of the metal lattice. The S atoms are located nearly between a T and a B site with respect to the fcc (111) lattice. MoS2 in the heterostructure retains, to a very large degree, its structure in the free-standing form, where each Mo atom has a trigonal prismatic coordination with the nearby S atoms. This is not surprising considering the strong covalent forces that hold the S-Mo-S trilayer system. The distance dMS between the metal surface and the closest sulfur layer (denoted (S)2 while the other layer is denoted (S)1 ) decreases as the strain develops from compressive to tensile (see Table I). The values of dMS for the 0% strain cases in Pd and Pt contacts

2 ML

3 ML

Strain

dMS

δ

dMS

δ

dMS

δ

dMM

Pd −3% −1.5% 0% 1.5% 2%

2.38 2.33 2.01 1.98 1.97

0.67 0.57 0.83 0.80 0.79

2.35 2.31 2.28 2.25 2.23

0.17 0.13 0.11 0.10 0.09

2.32 2.27 2.23 2.17 2.15

0.12 0.06 0.02 0.05 0.07

2.39 2.36 2.33 2.30 2.29

Pt −3% −1.5% 0% 1.5% 2%

2.45 2.40 2.35 2.30 2.28

0.52 0.48 0.44 0.39 0.37

2.40 2.36 2.32 2.28 2.27

0.24 0.19 0.17 0.17 0.17

2.48 2.43 2.37 2.32 2.31

0.20 0.17 0.16 0.15 0.15

2.39 2.36 2.33 2.29 2.28

are similar to the LDA values 2.0 and 2.6 Å reported before for Ti and Au contacts with MoS2 .23 With increasing strain in (Pd)1 /MoS2 , dMS shows a large decrease by more than 0.3 Å as strain changes from −3% to 0%, while the decrease for other heterostructures is smaller. On the other hand, the distance between the two metal layers dMM decreases as the strain increases from tensile to compressive as shown in the table. This out-of-plane relaxation along the [001] direction is expected as a response for the inplane strain. The rumplings δ in Table I show that the metal layer bonded to MoS2 is relatively flat, where the rumpling magnitude generally decreases as more layers are deposited or as the strain increases from compressive to tensile. The exception to these trends is the 1 ML case, which has a significantly larger rumpling because the three metal atoms in registry with Mo show a larger out-of-plane displacements from their nominal positions. By relaxing the heterostructures keeping the metal monolayer fixed, the binding energy decreases by ≈0.3 eV for the unstrained case. This decrease is a result of two competing factors: an energy gain from optimizing the interactions of the metal layer with the substrate that leads to its rumpling, and an energy penalty due to the deformation of the metal monolayer from its optimum unrumpled structure. Interestingly, for the Pt contacts, the displacement are away from the Mo layer, while for the Pd contacts, the displacements are towards Mo for zero and tensile strains, and away from Mo for compressive strains, compare Figures 1(b) and 1(c). This behavior has an important impact on the properties of the contacts (vide infra). As seen in the table, Pd has a larger rumpling compared to Pt, which can be understood because the cohesive energy of Pd is smaller than that of Pt (experimental values36 are 3.89 eV for Pd and 5.84 eV for Pt). Figure 2 shows the total interfacial energies E, as defined in Eq. (1), plotted as a function of the number of metal layers. As reflected in the figure, the adhesion energy of Pd/MoS2 heterostructures is ≈60% larger than that of the corresponding Pt/MoS2 interface. This correlates with the

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094707-4

Wissam A. Saidi (a) Pd/MoS2

Short-Range (eV)

Total (eV)

8.0

6.0

7.0

5.0

6.0

4.0

5.0

3.0

4.5 4.0 3.5 3.0 2.5 2.0

3.0 2.5 2.0 1.5 1.0 0.5

3.2 vdW (eV)

J. Chem. Phys. 141, 094707 (2014)

(b) Pt/MoS2

3.2

3.0

-3% -1.5% 0% 1.5% 2%

2.8 2.6 2.4 1

2

3

4

3.0 2.8 2.6

2.4 5 1 2 Number of Metal Layers

3

4

5

FIG. 2. Interfacial energies for Pd/MoS2 (a) and Pt/MoS2 (b) as a function for the number of metal layers for different strains. Top, middle, and bottom rows show, respectively, total, short-range, and vdW energies, as defined in Eqs. (1) and (2).

adsorption distance dMS shown in Table I where Pd contacts are closer to MoS2 than Pt contacts. This is also consistent with the results obtained with a graphene substrate where the Pd/graphene interface showed a larger stability compared to Pt/graphene.14, 15, 21 For the 0% strain case, the dependence of E on the number of layers shows a drastic difference between Pd and Pt. For Pd, E is the largest 6.25 eV for (Pd)1 /MoS2 and decreases as more layers are deposited up to 6.0 eV with 5 ML. On the other hand, (Pt)1 /MoS2 has the smallest interfacial energy E = 4.0 eV, which then increases as more layers are deposited with a value of 4.5 eV for 5 ML. However, as shown in the figure, (Pt)2 /MoS2 is the most stable showing an anomalous adhesion energy, E = 5 eV. For both Pd and Pt, E for 3 ML and 5 ML are not significantly different showing that the effects of the MoS2 substrate on the metallic slab is only limited up to 3–4 ML, and that the interfacial energies with 5 ML are nearly converged with slab thickness. Strain is in general an important factor in heterostructures due to the lattice mismatch between the parent structures. To show whether strain can affect the properties of metal/MoS2 contacts, we have additionally investigated the interface under compressive and tensile strains. Figure 2 summarizes the adhesion energies, and as shown, regardless of the imposed strain, the interfacial binding energies of Pt/MoS2 exhibit generally similar behavior as the 0% strain case except that the 3 ML and 5 ML energies are similar in energy for compressive strains, while for tensile strains E for 3 ML is larger than that of 5 ML. In contrast to the Pt/MoS2 contacts, strain has a huge effect on Pd/MoS2 heterostructures where the interfaces with compressive strain show similar trends as those with Pt, while the interfaces with tensile strain show a behavior similar to the Pd/MoS2 0% strain case. To shed light into the energetics and the differences between the Pd and Pt contacts with MoS2 , we examine the short-range and vdW contributions to E, as defined in Eq. (2). As reflected in Figure 2 (middle panel), the trends of ESR are similar to those of the total energy suggesting that

the observed trends in the interfacial energies are more likely to exist with local and semi-local functionals without dispersion corrections. Also, comparing the total and short-range contributions, we can infer that the extra stability of the Pd heterostructures compared to the Pt ones is mostly because of ESR . vdW interactions, as expected, enhances the stabilization of the interfaces and as seen EvdW is of appreciable magnitude that is comparable to ESR . The vdW contribution EvdW increases with the number of layers, as shown in Figure 2 (bottom panel), which is because of the additivity of the vdW interactions. The slow variations of vdW energy with distance as 1/r6 explain why EvdW saturates rather quickly with the number of layers. Also, it is noteworthy that EvdW is similar for 1 ML overlayer of Pd and Pt, and are larger in the Pd heterostructures with thicker metal slabs than the Pt ones. At first, this is surprising because Pt (van der Waal coefficient in Hartree×Bohr6 C6 = 60) is more polarizable than Pd (C6 = 51), and thus it would have been expected that EvdW for Pt is larger than that for Pd, and not the opposite as reflected in Figure 2 (bottom panel). However, these results can be rationalized because the adsorption distance dMS is smaller for Pd than Pt (vide supra), leading to stronger EvdW for Pd heterostructures. The larger dMS value for Pd/MoS2 under compressive strains also explain why EvdW is larger for −3% and −1.5% strain cases compared to the un-strained and tensile-strained slabs. The adhesion energy of the metal/MoS2 interface correlates with the imposed strain. This result can be understood using the d-band model of Hammer and Nörskov,45 which stipulates that as the strain proceeds from compression to tension, the d-band becomes narrower due to the decrease in the overlap between the d metal orbitals on neighboring atoms. This accordingly causes an up-shift in d-orbitals energies to maintain a constant d-band occupancy in Pd and Pt with partially filled d bands.44 The center of the d band,  d , reflects these changes and has been demonstrated to be a key parameter influencing the interaction strength between adsorbates and metal surfaces.44, 45 To validate that this is indeed the case for the metal/MoS2 interface, we show in Figure 3  d of the top layer in the free standing metal slab. The d-band center as the averaged energy of the d-band:  d  is calculated  = dEEρ(E)/ dEρ(E). As expected, with lattice expansion, the center of the d-band moves closer to the Fermi energy, making the metallic slab more reactive and explaining the trends seen in the adhesion energies. However, based on the d-band and for the same inplane strain value,  d downshifts as the number of layers of the metal slab increases. This indicates that the reactivity of the metallic slab, and accordingly E for the metal/MoS2 interface, would decrease with the number of deposited layers. We find that this is indeed the case for Pd in the case of 0% and tensile strains. However, for the compressive strains in Pd/MoS2 interface, as well as for Pt/MoS2 for all the investigated strain values, the adhesion energies are inconsistent with the trends anticipated from the d-band model. These results are similar to what was observed before for the adsorption of CO and H on Pd/Au layered systems with (100) and (111) orientations.46, 47

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094707-5

Wissam A. Saidi

J. Chem. Phys. 141, 094707 (2014)

-1.0

1.2 -3% -1.5% 0% 1.5% 2%

-1.4 -1.6

(b)

0.4 0.0 -0.4 -0.8

-1.8 -2.0 (b) Pt

-1.4 -1.6

Q (-e)

d-band center (eV)

(a)

0.8 Δ ρ (-e/Å)

d-band center (eV)

(a) Pd -1.2

-1.8 -2 -2.2 -2.4 -2.6 1

2

3 Number of Metal Layers

4

5

FIG. 3. Evolution of the d-band center  d of the top metal layer for (a) Pd and (b) Pt slabs as a function of the number of layers, and for different imposed strains. The Fermi energy is set at 0.

The d-band model is based on a second-order perturbation theory where the whole d-band is assumed to act as a single electronic level (center of d-band) that is interacting with the adsorbate. It is known that the d-band is not applicable in systems where there is a strong interaction between the adsorbate and substrate as observed before for hydrogen interacting with Pd overlayers on Au single crystals.47 Thus, it is conceivable that the d-band is also not applicable for Pt/MoS2 systems because of the strong coupling between Pt and MoS2 . However, this argument is not valid because the d-band explains successfully the Pd/MoS2 results eventhough the interaction energies are larger than those of the corresponding Pt systems, and eventhough the occupation of the d-band for Pd is more than that for Pt. We will show below that our results can nevertheless be rationalized using the d-band including charge transfer and effects of the metal subsurface layers as argued before for H and CO adsorption on Pd layers that are deposited on Au substrate.46, 47 First, to gain further insight into the charge rearrangements, we inspect the plane-averaged charge densities that can provide a quantitative picture of the charge distribution due to bonding and especially the extent of the charge rearrangements in the metallic slab. This can be conveniently done by examining ρ(z) = ρmetal/MoS − ρMoS − ρmetal where z 2 2 is perpendicular to the interface. Here, ρmetal/MoS is the 2 plane-averaged density of the metal/MoS2 heterostructure, and ρMoS and ρ metal are, respectively, the plane-averaged 2 densities of the isolated MoS2 and metal subsystems adapting their structures as in the optimized   metal/MoS2 system. The plane-averaging is defined as x y dxdy ρ(x, y, z) where the integration of the electronic density ρ(x, y, z) is over the surface unit cell. Positive (negative) values in ρ(z) indicate accumulation (reduction) of the electron charge density. Furthermore, to quantify the amount of charge transfer, we also  z examined the cumulative charge-rearrangement Q(z) = 0 dz ρ(z ).48 A negative value of Q(z) specifies the number of electrons transferred from left to right of a plane perpendicular to the surface plane at z, a value of 0 indicates

-1.2 0.4 0.3 0.2 0.1 0.0 -0.1 -0.2 -0.3 -0.4

(c)

(d)

1ML 2ML 3ML 5ML S1MoS2

Pd

Pd

Pd

Pd

Pd

S1Mo S2

Pt

Pt

Pt

Pt

Pt

FIG. 4. Plane-integrated charge rearrangements for (a) Pd/MoS2 and (b) Pt/MoS2 with 0% strain for different number of metallic layers. For comparison purposes, we used the same y-axis scale for both metals. The approximate locations of the different layers are indicated by the dotted lines. Negative values in ρ correspond to a reduction of the electron density, while positive values indicate an increase in the electron density. (c) and (d) show the accumulative charge transfer Q(z) for Pd and Pt, as defined in the text.

no charge is shifted across that plane, and a positive value of Q denotes a transfer from right to left. Figure 4 shows ρ(z) and Q(z) for metal/MoS2 systems for 0% strain case. The systems with other strain values show similar behavior but with few notable differences, as will be discussed below in the context of 1 ML contacts. First, the charge rearrangements for both metals are fairly similar, where there is charge depletion in the interface region above (S)2 , and concomitantly a pronounced charge accumulation above the top metal layer. The direction of charge transfer is expected because the work-functions of the metals are larger than that of MoS2 . Moreover, the charge depletion is asymmetric with respect to (S)2 , where it is more pronounced on the MoS2 side. The asymmetry in the profile is in accordance with a Pauli push-back mechanism because the electron density depletion is more enhanced on the side where the charge density of the tailing metal electron cloud is significantly larger. This is verified by examining the charge densities of the separate systems in their free standing form. A similar behavior was also observed for graphene/MoS2 interface.49 The Q(z) profiles in Figures 4(c) and 4(d) show that the charge rearrangement in the interface region is oscillatory resulting in a complex Q(z) profile with more than one local minima. Pt contacts show a larger charge response than Pd contacts, where the minima of Q(z) are −0.3 electron for Pd/MoS2 and −0.4 electron for Pt/MoS2 . The amount of charge transfer is similar to what was observed before for the adsorption of the CuPc and ZnPc metal phthalocyanines on noble metal surfaces.35 Additionally, the profile for the Pd contacts is more complex than the profile for the Pt contacts. Namely, in the interfacial region of the Pd contacts, there are locations in the interface where Q(z) = 0, which indicate that no charge-transfer is taking place across these planes. This succession of regions with depletion and accumulation of electron density can be viewed as a series of dipole layers. On the other hand for Pt, the charge transfer profile show that the heterostructures have a larger interfacial dipole than Pd as

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Wissam A. Saidi

 can be verified by computing dzzρ(z). The electrostatic interactions between these dipoles are repulsive, which at least partially explain why the Pt/MoS2 interfaces are less stable than the Pd/MoS2 interfaces (see Figure 3). It is interesting to note that eventhough the Mo layer is ≈4 Å far from the metal layer, there are substantial charge arrangements near the Mo layer that are of a similar or even larger magnitude compared to those near the (S)2 layer. There is no direct interaction between Mo and the top metal layer as shown from the negligible overlap between their electron densities. Nevertheless, there is still an indirect interaction, as a response to the charge depletion of (S)2 , mediated by the strong Mo–S bond. The orbital-decomposition of the PDOS (not shown) confirms what is inferred from charge analysis, and show additionally that the orbitals involved in the hybridization are the s and pz orbitals of (S)2 and the dz2 and dx 2 − dy 2 orbitals of Mo, with the dz2 metal orbitals. It is expected that the strain dependence of the metal/MoS2 contacts is because of the Mo dx 2 − dy 2 orbitals that are involved in the bonding. The PDOS on (S)1 confirms that this layer is not involved in the bonding, and thus validating the negligible charge rearrangement seen in Figure 4. On the metal side, the charge rearrangements are also non-negligible below the top metal surface that is closest to MoS2 , and with an oscillatory decaying profile, as Figure 4 shows. Examining (Pt)5 /MoS2 , we see that the charge rearrangement is of appreciable magnitude in the second and third layers below the top metal layer. This charge response in turn rationalizes our findings shown in Figure 2 for the dependence of the adhesion energies on the number of layers, and specially the anomalous behavior with 2 ML. In these cases, the subsurface layers, and chiefly the second nearest-neighbor layer, that provide additional binding to the MoS2 substrate are absent in (M)1 /MoS2 , which makes (M)1 adhere weakly to MoS2 in comparison to (M)n > 1 . In fact, the d-band centers of the subsurface layers are of a similar proximity to the Fermi energy as the top surface layer, and thus are expected to exhibit high reactivity. The evolution of the Pd/MoS2 interface with the strain nicely corroborates the trends that are characteristics of the Pt interfaces. For both Pd and Pt, Figure 4 shows that the charge rearrangement for the 1 ML contacts are larger than those of the thicker layers. The evolution of the charge response for the 1 ML case, measured using Q(z), are depicted in Figure 5. Contrasting the charge response for Pd with compressive and tensile stresses, we see that the charge over-spill into the vacuum is more substantial for compressive than for tensile strains. This is because in these cases, the metal atoms that are atop of the Mo atoms are displaced away from the Mo layer (vide supra). This is energetically not favorable. However, this charge spill out will cause a coupling between MoS2 and subsurface metal layers for thicker metal slabs. The coupling will in turn leads to the anomalous behavior that is seen in the energetics of Figure 2. The fact that this happens for Pt and not for Pd is because Pt is more polarizable than Pd as evidenced by examining their static polarizabilities, α Pd = 13.9 and α Pt = 14.5 bohrs6 . The same argument also explains why the stabilization energy of the 5 ML case is larger

J. Chem. Phys. 141, 094707 (2014) 0.5

Q(-e)

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(a) Pd

0.5

0.4

0.4

0.3

0.3

0.2

0.2

0.1

0.1

0.0

0.0

-0.1

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-0.2 -3% -1.5% 0% 1.5% 2%

-0.3 -0.4 -0.5

-0.3 -0.4 -0.5

S1 Mo S2

Pd

S1 Mo S2

Pt

FIG. 5. Accumulative charge transfer Q(z) for (a) (Pd)1 /MoS2 and (b) (Pt)1 / MoS2 with different strain values. The dotted line shows the location of the different layers for the 0% strain only to avoid clutter.

than the 3 ML case in Pt/MoS2 for the compressive strains. As can be seen from Figure 5(b) the charge spill out into the vacuum for the 1 ML is the largest for compressive strain. The interaction of MoS2 with metal Pd and Pt slabs show several similarities with previous study for the interaction of CO and H with (100) and (111) substrates of Pd grown pseudomorphically on Au.46, 47 However, there are some notable differences. First, the interactions between Pd and Au are mostly due to hybridization between sp orbitals, which is because Au has a deep-lying filled d-band that makes it an inert noble metal. Here in contrast, the interaction of Pd or Pt with MoS2 is stronger where both sp bonding and d bonding is involved as discussed before. Additionally, the stronger anomalous stabilization of the (M)2 /MoS2 interface is due to the interaction between an extended system, MoS2 trilayer, with the metals, while previously H and CO are introduced as local probes to measure the reactivity of the Pd/Au system. In the latter case, the strength of the interaction between H and the substrate is relatively strong compared to (M)n /MoS2 , as it involves a coupling between s orbitals of H, and the Pd d3z2 −r 2 orbital. CONCLUSION

Using semi-local density functional theory in conjunction with semi-empirical dispersion corrections, we have characterized the Pd/MoS2 and Pt/MoS2 interfaces as a function of the number of layers deposited on the substrate and in terms of their dependence on the inplane strain. Our results show that as the lattice expands, the metal adhesion to MoS2 decreases for all metal thicknesses. This dependence of the contacts on strain can be fully understood using the d-band model. However, we find that evolution of the contacts with the number of metal layers showed markedly different behavior between Pd and Pt that was at variance with the d-band model. Specifically, the Pt/MoS2 contacts show an anomalous large stability with the deposition of two metal monolayers for all investigated strains, although the d-band model stipulates that the stability of the interface will decrease as more metal layers are deposited. On the other hand, the Pd/MoS2 exhibit

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the expected behavior based on the d-band model for unstrained and tensile strained lattices, but heterostructures that are under compressive strains showed a similar behavior as in Pt/MoS2 contacts. The results are rationalized by accounting for second-nearest-neighbor effect that couples MoS2 with the subsurface metal layers, as validated using charge decomposition. Additionally, the underpinnings of the differences between Pd and Pt are attributed to the larger polarizability and cohesive energy of Pt compared to Pd, that leads to a larger charge-response in the subsurface layers. Previously, it was shown that Pt/MoS2 hybrid nanomaterials exhibit much higher electrocatalytic activity towards hydrogen evolution reaction compared to Pt catalysts with the same Pt loading.22 Based on the present results, it is anticipated that the catalytic properties of MoS2 on supported metal layers can be tailored using either strain or metal slabs with different thicknesses.

ACKNOWLEDGMENTS

Calculations are performed in part at the University of Pittsburgh Center for Simulation and Modeling. This work used the Extreme Science and Engineering Discovery Environment (XSEDE), which is supported by National Science Foundation (NSF) Grant No. OCI-1053575. 1 A.

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