Probing Evolution of Local Strain at MoS2-metal Boundaries by

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Probing Evolution of Local Strain at MoS2-metal Boundaries by Surface-enhanced Raman scattering Yan Aung Moe, Yinghui Sun, Huanyu Ye, Kai Liu, and Rongming Wang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b13241 • Publication Date (Web): 26 Oct 2018 Downloaded from http://pubs.acs.org on October 27, 2018

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ACS Applied Materials & Interfaces

Probing Evolution of Local Strain at MoS2-metal Boundaries by Surface-enhanced Raman scattering Yan Aung Moe1,#, Yinghui Sun1,#, Huanyu Ye1, Kai Liu2, Rongming Wang1,* 1

Beijing Advanced Innovation Center for Materials Genome Engineering, Beijing Key Laboratory for MagnetoPhotoelectrical Composite and Interface Science, School of Mathematics and Physics, University of Science and Technology Beijing, Beijing 100083, People’s Republic of China 2 State Key Laboratory of New Ceramics and Fine Processing, School of Materials Science and Engineering, Tsinghua University, Beijing 100084, People’s Republic of China

Abstract Strain usually exists in 2-dimensional (2D) materials and devices, and its presence drastically modulates their properties. When 2D materials interface with noble metals, local strain and surface plasmon can couple at the metal-2D materials boundaries, delivering a lot of intriguing phenomena. Current studies are mostly focused on the explanations of these strain-related phenomena based on a static point of view. Although strain can typically be relaxed in many environments, the time evolution of strain at metal-2D materials interfaces remains largely unknown. In this work, we investigate the evolution of local strain at Ag-MoS2 boundaries by surface-enhanced Raman scattering (SERS). With the split of MoS2 Raman peaks as an indicator of local strain, it is found that the originally localized strain at Ag-MoS2 boundaries evolves and relaxes with time into a delocalized strain in MoS2 plane. The time to start the strain relaxation depends on the number of layers of MoS2 flakes, suggesting that the relaxation may result from the mechanical instability of the interface between the top-most MoS2 layer and the underlying materials. The relaxation occurs in a certain period of time, i.e. ~70 days for 1L and ~30 days for 3L. Accompanying the strain relaxation, surface sulphurization of Ag also occurs, a process that reduces the strength of locally enhanced electric field. Our results not only provide deep understandings of strain evolution at metal-MoS2 interfaces, but also shed light on the optimization of MoS2-based device fabrications. 1

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Keywords: Molybdenum disulphide, Raman peaking splitting, strain relaxation, surface-enhanced Raman scattering, Ag nanoparticles

Introduction Since the discovery of graphene along with its promising properties, two-dimensional (2D) materials have been a focus of materials studies for over a decade. As a typical 2D material in the family of transition metal dichalcogenides (TMDs), MoS2 has a direct band gap of ~1.8 eV in the monolayer limit,1 which surpasses graphene that lacks a band gap. Therefore, MoS2 has attracted intensive attention and finds potential applications in nanoscale electronic and photonic devices, such as field effect transistors,2-3 sensors,4 photodetectors,5 photocatalysis,6 lithium ion batteries,7 and chemical detection.8-10 Mechanical, thermal and chemical stabilities of MoS2 have also been studied for its practical applications,11-14 as the tunability of MoS2 by those parameters has been drawing much attention.15-18 In the application of nanodevices, the electronic band structure of MoS2 can be tuned by changing layer thicknesses together with applying multiple fields. The interfaces of metal electrode/MoS2 and MoS2/substrates are of vital importance for the transport behaviors. Besides doping effect, the study of strain effects from metal electrodes and substrates is helpful for the modulation of the device interface,19-20 since the stability of the nano-devices strongly depends on the strain relaxation and structural evolution at nanoscale.21 In materials synthesis and device fabrication, strain is usually unavoidable and intriguing studies have been focused on it, in which the strain can be applied either uniaxially22-23 or biaxially.24-25 Compared to the strain effect in graphene, a much smaller amount of strain is required to vary the band gap of TMDs.26 When the strain is applied on monolayer MoS2, the direct gap turns into an indirect gap and the effective masses vary drastically.21, 27-28 A uniaxial tensile strain is known to modulate the lattice symmetry and the band gap of MoS2, thereby affecting its Raman modes and photoluminescence (PL) properties.22-23 It can be applied on MoS2 through versatile ways such as mechanically bending,23 substrate engineering,29 and metal deposition.30-31 2


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Metal deposition is a typical way to fabricate electrodes in nanoscale electronic devices. The metalMoS2 interface is found to be crucial for the performance of electronic devices because it drastically modulates the Schottky contacts32-34 and effective carrier mobility.35-36 Lowering the contact resistance at metal-MoS2 interface has been approached by several methods including work function engineering,37 molecular doping,38-39 phase engineered contacts (i.e., semiconducting 2H phase to metallic 1T phase transition of MoS2)40 and localized doping by annealed Ag contact.41 When MoS2 interfaces with noble metals, the induced surface plasmon resonance (SPR)42 and exciton-plasmon coupling interaction43-46 in the hybrid nanostructure have enabled new functions in optoelectronic devices. It is also reported that the local strain and surface plasmon can be coupled at the metal-MoS2 boundaries, delivering some interesting optical phenomena.31, 47 However, current studies are mostly focused on the explanations of these strainrelated phenomena based on a static point of view. Although strain can typically be relaxed in many environments, the time evolution of strain at metal-2D material interfaces remains largely unknown. In this work, we systematically studied the process of strain evolution at Ag nanostructures-MoS2 interfaces by surface plasmon-enhanced Raman scattering. With the split of MoS2 Raman peaks as an indicator of local strain, it is found that the originally localized strain at Ag-MoS2 boundaries evolves and relaxes with time into a delocalized strain in the MoS2 plane. The time to start the strain relaxation depends on the number of layers of MoS2 flakes, suggesting that the relaxation may result from the mechanical instability of the interface between the top-most MoS2 layer and the underlying material. Scanning electron microscopy and X-ray photon spectroscopy further reveal the surface sulfurization of Ag nanostructures with time elapse, which reduces the strength of locally enhanced electric field as well as the Raman signals originating from the locally strained MoS2. Our results not only provide deep understandings of strain evolution at metal-MoS2 interfaces, but also shed light on the optimization of MoS2-based device fabrications.

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Results and discussion We used mechanical exfoliation to prepare ultrathin MoS2 flakes on Si substrates covered with ~300 nm thick SiO2 layer from bulk MoS2 crystal. The inset in Fig. 1a shows the optical images of monolayer (1L), bilayer (2L), and trilayer (3L) MoS2. The inset in Fig. 1b shows the Raman spectra for ultrathin MoS2 sheets with thinknesses up to hexa-layer. The Raman peak frequency can be extracted by Lorentzian fitting. The interval between Raman in-plane mode 𝐸 and out-of-plane mode 𝐴 can be used as an indicator of the layer number of ultrathin MoS2 flakes.48-49 Monolayer MoS2 is known to have the point group symmetry of D3h, and the few layer MoS2 is known to have D6h symmetry.50-51 The interval between 𝐸 and 𝐴 modes is designated as 𝛥𝜔 (where, the subscript N represents MoS2 layer number), which increases as the increase of the layer number (Fig. 1a). The values of 𝛥𝜔 are 18.6, 21.1, and 23.1 cm-1 for monolayer, bilayer, and trilayer, respectively, which are consistent with reports in literatures.31, 52-53 The full width at half maximum (FWHM) values for 𝐸

modes are ~1.88 cm-1 for monolayer, ~2.59 cm-1 for bilayer and

~2.41 cm-1 for trilayer. As the layer number increases, the long range Coulombic interaction among Mo atoms is weakened due to the increase of dielectric tensor, which therefore softens the in-plane 𝐸 mode and hardens the A

mode by the enhanced restoring force on S atoms.50, 53 The 𝐸

sensitive to strain, and the A

mode is found to be

mode is known to be sensitive to the doping level.54,55 Besides strain induced

softening of vibration modes, Ag can be considered as electron donor when contact with MoS2, and can also cause the softening of both 𝐸

and 𝐴

modes.9 The broad peak around 450 cm-1 is closely

overlapping Raman peaks assigned to second order zone-edge phonon peak (2LA(M)) and a first-order optical phonon peak 𝐴 .56 The PL spectra of monolayer, bilayer, and penta-layer (5L) MoS2 are shown in Fig. 1b, where A and B peaks result from the spin-orbit splitting at the top of the valence band, and I peak (~1.4 eV), which results from the indirect band transition, is absent for monolayer MoS2 as a consequence of inversion symmmetry breaking.57-58 The PL intensities in Fig. 1b are all normalized to the corresponding Raman peak intensity of each layer type. Accompanying to the interval between 𝐸

and 𝐴

modes (𝛥𝜔 ),

4


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the position of I peak can also be used to assign the layer number for its red-shift with the increase of layer number.

Figure 1. (a) The Raman peak position of 𝐸 and 𝐴 modes of pure MoS2 with varying thicknesses from monolayer to hexa-layer. The red axis at the right depicts the intervals between 𝐸 and 𝐴 modes. Inset shows the optical images of 1L, 2L, and 3L MoS2 on SiO2/Si substrate with the white scale bar of 5 µm. (b) Photoluminescence (PL) of pristine 1L, 2L, and 5L MoS2 on SiO2/Si substrate where the magnified PL peak (×4) with dotted line originates from the indirect band transition of 5L MoS2. (Note that the PL intensities are all normalized to the corresponding Raman peak intensity of each layer type). Inset shows the Raman spectra of assorted layers. (c) The comparison of Raman spectra of pristine (black dashed lines) and 3-nm-Ag deposited (red lines) 1L, 2L, and 3L MoS2 on SiO2/Si substrate.

Ag was deposited by e-beam evaporation at a rate of 0.2 Å/sec for 150 sec. The nominal thickness of the deposited Ag was about 3 nm. Following the Volmer-Weber island growth mode, Ag forms nanoparticles on MoS2 instead of a continuous film.31 In our experiment, the deposited Ag forms dendritelike nanostructures, as illustrated in the SEM images in Fig. 3c. The morphology of Ag on MoS2 is identical 5



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and independent on the layer number. Figure 1c shows the comparison of Raman spectra with arbitrary intensity profiles for 1L, 2L, and 3L MoS2 to clearly depict the difference in spectra profiles before and after Ag deposition. The SERS enhancement after Ag deposition compared to pristine sample can be observed in Fig. S1. Compared with those of pristine MoS2 (before Ag deposition), the Raman spectra of Ag deposited MoS2 reveal distinct changes. While the peak positions of in-plane 𝐸 modes remain the same after deposition, another redshifted 𝐸 𝐸

and 𝐴

and out-of-plane 𝐴

and 𝐴 ′ peaks emerge near the pristine

peaks. We have reported that the Raman signals for 1L MoS2 after Ag deposition (upper red

spectral line in the top panel of Fig. 1c) originate from two planar regions: one region of pure unstrained MoS2 which is not covered by the Ag dendrites, and the other region of strained MoS2 around the Ag-MoS2 boundary. Illustration of Ag deposited MoS2 surface illuminated by the 532 nm laser line is provided in Fig. S2. The un-shifted 𝐸 shifted 𝐸

and 𝐴

and 𝐴 ′ peaks result from the metal-MoS2 boundary where local strain is mainly induced.30, 31 and 𝐴 ′ peaks for few-layer MoS2 are weaker compared to the respective 𝐸

The relative intensity of 𝐸 and 𝐴

peaks can be ascribed to the region of pure unstrained MoS2, and the red-

peaks. The intensity ratios [I(𝐸 )/I(𝐸 )] of the 𝐸

peak to the pristine 𝐸

peak just after Ag

deposition are 1.08 for monolayer, 0.55 for bilayer, and 0.23 for trilayer MoS2.The strongest relative intensity for 1L MoS2 is due to the strain mainly localized in the top-most layer of MoS2. For bilayer to bulk MoS2, the splitting of 𝐴

mode is almost unconceivable. The splitting of MoS2 Raman peaks indicates

the presence of local strain induced by Ag dendrites, which is consistent with our and other group’s previous work.30-31 Then the splitting of the strain-sensitive 𝐸

Raman peak can be regarded as the dominant

indicator of local strain at the metal-MoS2 boundary. Immediately after the Ag deposition, the interval between 𝐸 ′ peak and 𝐸

peak are observed to be ~8.0 cm-1 for monolayer, ~6.5 cm-1 for bilayer, and 5.9

cm-1 for trilayer, which is similar to the splitting resulted from the uniaxial tensile strain achieved by mechanically bending the underlying substrate.22-23 According to the relation between Raman shift and tensile strain, the observed Raman splitting of 𝐸

and 𝐸

peaks at ~8.0 cm-1 in monolayer MoS2 with Ag

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nanostructures suggests an effective strain of ~2.0%.22-23 The peak splitting in 1L MoS2 is nearly the same as our previous work where the strain is induced by the deposition of Ag nanoparticles (NPs) with different sizes, which indicates that the induced strain by Ag is not dependent on the specific morphology of Ag nanostructures. Although the Ag induced strain in monolayer and few layer MoS2 is distinct enough to be detected, the Raman spectra of bulk MoS2 before and after Ag deposition are almost identical, as shown in Fig. S3. The phenomenon indicates that the weak interlayer interaction leading to inefficient load transfer between MoS2 layers. Ag dendrites can induce the local strain which distorts the lattice structure of MoS2 along the metalMoS2 boundary. Due to the presence of metal NPs, the resultant strain changes the symmetry of monolayer (bilayer) to Cs (C2h) from D3h (D6h) along the Ag-MoS2 boundary.22, 50 The change in the lattice symmetry introduces new Raman modes with the peak position in accordance with the strain being applied. This strain resulting from the Ag-MoS2 interactions is highly inhomogeneous and distributes locally along the AgMoS2 boundaries, and it is found to be tensile, therefore leading to the softening of in-plane phonon modes at the boundaries.30-31 Away from the Ag-MoS2 boundaries, the strain will gradually disappear due to the load transfer between the top-most MoS2 layer and the substrate (for monolayer) or the underlying MoS2 layers (for bilayer or more layers).31 Another important effect induced by Ag nanostructures along with the strain is localized surface plasmon resonance (LSPR) under the incidence of visible light. The enhanced electric field due to LSPR of Ag NPs is localized within ~1 nm around the circular edge of the Ag NPs, spatially overlapping with the strain centralized zone at the Ag-MoS2 boundaries and enhancing the Raman signals from there. Therefore, Raman scattering from the local area right at the Ag NP-MoS2 boundaries can be detected distinctly, although the strained MoS2 along the Ag-MoS2 boundary occupies only a small proportion compared with the unstained MoS2.31 The coupling of LSPR and local strain delivers a split of Raman modes, as shown in Fig. 1c as well as in our previous study,31 which are contributed by the strained MoS2 region along the Ag-MoS2 boundaries (i.e., 𝐸 ) and the unstrained MoS2 region away from the boundaries (e.g., 𝐸 ), respectively (Fig. S2). The profile of 𝐸

Raman peak is observed to be sharp despite 7



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the highly inhomogeneous strain induced by Ag, which may be attributed to the lightning rod effect59 and the hot spots of local electric field around sharp edges.31 Our previous work indicates that the simulated electric field is highly localized and enhanced within ~1 nm around the circular edge of the semisphere contacting with the substrate.31 Raman scattering from the local area right at the Ag-MoS2 boundary dominates the overall Raman signal. It thereby makes us observe only one value of splitting in the Raman spectra. In this scenario, the relative intensity of the shifted Raman peak [I(𝐸 )/I(𝐸 )] suggests the strength of LSPR, and the splitting interval between 𝐸

and 𝐸

modes (Δω ) is an indicator of the strain difference

between the strained and the unstrained MoS2 regions. In the following, we will analyze the strain evolution at Ag-MoS2 interfaces through the changes in Raman spectra based on this scenario. Figure 2 shows the Raman spectra for the Ag/MoS2 with 1L, 2L, and 3L MoS2 over 72 days after Ag deposition. For all layer numbers, the relative intensity of either redshifted 𝐸

or 𝐴 ′ peak decreases with time elapse. The speed

at which the intensity decreases is found to be dependent on the layer number, i.e., the 𝐸 MoS2 decays fastest, the 𝐸

peak for 3L

peak for 1L MoS2 decays in the slowest speed. After 48 days, the 𝐸

peak in

3L MoS2 totally disappeared, while that in 2L MoS2 become a shoulder peak after 72 days. However, the 𝐸

peak is still evident in 1L MoS2 after 72 days. The intensity enhancement in 𝐴 +2LA(M) peak around

450 cm-1 can simply be addressed to the SERS effect of freshly deposited Ag dendrites, and the enhancement also fade away following along the Ag dendrites aging.60

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Figure 2. The Raman spectra of pristine MoS2 (before Ag deposition) and their evolution with time after Ag deposition for the 1L (a), 2L (b), and 3L (c) MoS2 respectively. The 𝐸 Raman peak for 3L-MoS2 totally disappeared after 48 days, while 2L-MoS2 and 1L-MoS2 preserve the 𝐸 Raman peak for longer time.

The changes of 𝐸

peak intensity and peak position, which were extracted by Lorentzian peak fitting

after the subtraction of the baseline in each spectrum, were used to analyze the evolution of the local strain. In Figs. 2a-2c, the change in the peak intensity for the splitting 𝐸

mode is more noticeable than the change

in the peak shift. Figure 3a shows the intervals or splitting between 𝐸

and 𝐸

modes (Δω ) shortening

over time for the Ag deposited 1L, 2L, and 3L MoS2. Immediately after Ag deposition, the interval Δω is observed to be ~8.0 cm-1 for monolayer and decreases to be ~7.0 cm-1 after 72 days, which is still larger than Δω for bilayer just after metal deposition (~6.5 cm-1). The intensity ratio [I(𝐸 )/I( 𝐸 )] also decreases with time elapse. For monolayer, the relative intensity ratio decreases from ~1.0 to ~0.6 (Fig. 3b). 9



The data for 3L MoS2 were cut off at 48 days because the 𝐸 the appearance of 𝐸 evolution of 𝐸

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peak totally disappeared after that. Because

mode can be ascribed to the induced strain by Ag nanostructures on MoS2, the

peak in intensity and position may reflect the change of interfacial strain. Monotonic

decrease in intensity ratio and separation is designated as the consequence of the Ag nanostructures aging with time. We characterized the morphology change of Ag/MoS2 heterostructures by SEM imaging. The overall morphology of the Ag dendrite-like nanostructure is found to be consistent through time, as shown in Figs. 3c-3d.

Figure 3. (a) Raman peak splitting of 𝐸

mode decreasing with time for 1L, 2L, and 3L MoS2 (the dotted lines

are added as guide-line for eyes). (b) Intensity ratio of 𝐸

to 𝐸

Raman mode decreasing with time for 1L, 2L,

and 3L MoS2. (c-e) SEM images of Ag nanostructures on MoS2 and the SiO2/Si substrates for the different time periods of 0 days, ~24 days, and over 70 days after Ag deposition. (Note that image (c) and (d) are taken from the same sample and image (e) is taken from another sample.)

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As discussed above, the change of intensity ratio of [I(𝐸 )/I(𝐸 )] in Fig. 3b reveals that the strength of LSPR decreases gradually with time elapse. The splitting interval between 𝐸

and 𝐸

modes (Δω ) is

an indicator of the strain difference between the strained and the unstrained MoS2 regions. The local strain mainly induces a redshifted Raman peak of 𝐸 increase of the Raman shift of 𝐸 The Raman intensity of the 𝐸

mode. The decrease of the local strain can cause the

peak, and consequently results in the decrease of Raman splitting (Δω ). peak can be ascribed to two factors: (1) The relative area of strained MoS2;

(2) the strength of the local electric field. Because the morphology change of Ag is negligible according to SEM images (Figs. 3c-3d), the area proportion of strained MoS2 to un-strained bare MoS2 remains almost identical. In the case where the strain is applied by mechanically bending a polymer substrate, the decrease in strain from 1.6% to 0.8% changes the intensity ratio of the 𝐸 Therefore, the change of Raman intensity of 𝐸

and 𝐸

peaks only from ~1.20 to ~1.00.23

peak in our case can be mainly attributed to the change of

the local electric field along the boundary between Ag and MoS2, where the peak intensity ratio decreases from 1.1 to 0.6. The reduction of local electric field is possibly due to the change of the Ag surface. Therefore, XPS characterizations are further used to probe the change in the composition of Ag nanostructures. The XPS spectra were taken for three different conditions, one for pristine MoS2 without any Ag and the other two for fresh and aging Ag/MoS2 after Ag deposition. The corresponding XPS spectra are shown in Fig. 4 with peak fitting, respectively. The associated binding energy values for each constituent element and the changes in binding energy (BE) values measured for each condition are listed in Table 1. The BE value for each single separate measurement is referenced by setting that of the adventitious carbon C1s to 284.8 eV. For the pristine MoS2, the Mo 3d5/2 and Mo 3d3/2 doublet peaks are found to be at 229.40 eV and 232.55 eV. The S 2p doublet peaks are found to be at 162.27 and 163.47 eV, corresponding to S 2p3/2 and S 2p1/2 peaks. For the freshly Ag deposited sample, the doublet peaks at 368.30 and 374.30 eV are assigned to Ag, corresponding to Ag 3d5/2 and Ag 3d3/2, respectively. The Mo 3d5/2 and Mo 3d3/2 doublet 11



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peaks are found at 229.53 and 232.68 eV. The S 2p doublet peaks for the freshly Ag-deposited sample are found to be at 162.39 and 163.58 eV, corresponding to S 2p3/2 and S 2p1/2 peaks. The S 2p spectra acquired for the samples has doublet structure due to the presence of S 2p3/2 and S 2p1/2 peaks and could be fit using 2:1 peak area ratio and a splitting of ~1.2 eV. Compared with pristine MoS2, we found a synchronous blue shift in BE of ~0.12 eV for both Mo 3d and S 2p species after Ag deposition. It can be attributed to the charge transfer process and the doping of Ag over MoS2.9, 39 After the Ag deposition for 20 days, the Ag nanostructures are in an aging state through XPS spectra. The Ag peaks are shifted to lower energy and can be fitted by 𝐴𝑔 species and another doublet peaks (Fig. 4a). The 𝐴𝑔 species compose of the Ag 3d5/2 and Ag 3d3/2 peaks at 368.27 and 374.27 eV. The fitted doublet peaks at 368.07 and 374.07 eV can be assigned as Ag cation (𝐴𝑔

).61 The XPS analysis indicates the gradual formation of a layer of Ag compound on the

surface of Ag nanostructures as the sample aging.62-63 For the Mo species during aging process, the BE are found at 229.51 and 232.65 eV, corresponding to Mo 3d5/2 and Mo 3d3/2, respectively. Being sandwiched between S atoms, Mo atoms in MoS2 are structurally protected from the outside environment, and the BE for Mo species are thought to be not changing, considering the energy accuracy of the spectrometer, which is 0.05 eV for non-insulating sample. For the sulphur during the aging process, another set of red-shifted peaks designated as 𝑆 𝑆

species are found to evolve at 161.36 and 162.44 eV along with the original

species at 162.36 and 163.55 eV, corresponding to S 2p3/2 and S 2p1/2 peaks, respectively. This

suggests that the Ag compound on the surface is related to the chemical bonding between the sulphur atoms and Ag at the metal-MoS2 interface61, 64-65. As mentioned before, the decrease of intensity ratio of [I(𝐸 )/I(𝐸 )] in Fig. 3b can be attributed to the gradual decease of the strength of LSPR with time elapse. The reason is that the local electric field at the metal-MoS2 boundary is reduced. It is known that Ag nanostructures have prominent surface plasmon excitations and can lead to local electric field hot spots with strongly enhanced Raman scattering upon laser illumination.66 Another important factor determining the high local electric field at the metal-MoS2 boundary is the lightning rod effect, where the field is dramatically enhanced around sharp edges.59 The 12


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electric field is highly localized and enhanced within ~1 nm around the circular edge of the metal nanostructure contacting to the substrate.31 When the surface sulphurization of Ag occurs as indicated by the XPS results, the dielectric constant and the curvature of the surface of Ag nanostructures are changed, which would reduce the local electric field.60 Furthermore, the Raman signal of MoS2 in our experiment is not only related to the strength of local enhanced electric field, but also related to the distance of MoS2 to the local electric field hot spots. The sulphurized layer of Ag dendrites gradually grows thicker with time. When the Ag have been sulphurized by 2~3 atomic layers, the pure Ag part gets away from the underlying MoS2 by about 0.5~1 nm. This means that the MoS2 become far away from the hot spots of local electric field. Therefore, the Raman signal of MoS2 is reduced significantly.

Figure 4. High resolution XPS spectra and corresponding peak fittings for (a) Ag 3d, (b) Mo 3d & S 2s and (c) S 2p orbitals in pristine MoS2, fresh and aging Ag-MoS2 samples. Upper spectral lines of (a) represent the Ag 3d species of freshly deposited Ag dendrites just after deposition, and the lower spectral lines represent the Ag 3d species after ~20 days of aging. In (b) and (c), the top-most spectral lines correspond to the pristine MoS2 without Ag deposition, and the middle ones belong to the freshly Ag deposited MoS2 sample, while the bottom ones represent the aging sample in each panel.

Table 1. Binding energy (BE) of each constituent element for pristine MoS2, fresh Ag/MoS2 and aging Ag/MoS2 heterostructures.

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Figs. 5a-5c reveal the changes of Raman peak positions for 𝐸

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and 𝐸

modes with the increase of

time for Ag deposited on three different layer numbers of MoS2 extracted over 72 days and >120 days after Ag deposition. For all the Ag-MoS2 samples, a distinct trend is shown that 𝐸

peak red-shifts and 𝐸

peak blue-shifts with time elapse, leading to the decreasing intervals between 𝐸

and 𝐸

and 𝐸

modes. The 𝐸

peaks finally appear to merge into one after more than 150 days (Figs. 5d-5f). As we discussed

above, 𝐸

and 𝐸

modes originate from the strained MoS2 region along the Ag-MoS2 boundaries, and the

unstrained MoS2 region away from the boundaries, respectively. Therefore, the hardening of the 𝐸

mode

indicates that the tensile strain around Ag-MoS2 boundaries reduces with time, while the softening of the 𝐸

mode reveals the emergence of a tensile strain at the originally unstrained MoS2 region. The newly

merged 𝐸 𝐸

peak in the aged Ag-MoS2 sample generally has a slight redshift compared with the original

peak of pristine MoS2, which is more evident for 1L MoS2 but much less for 2L and 3L MoS2 (Fig. 5d).

These phenomena suggest that the originally localized strain along the Ag-MoS2 boundary is ultimately relaxed into a delocalized strain in the MoS2 planar directions. The process of strain relaxation is depicted in Figs. 5g-i. Fig. 5g shows the fresh Ag deposited 1L MoS2, where the strain is highly localized along the boundary of Ag/MoS2. Fig. 5h shows the slight sulphurization of Ag surface with time elapse, but the most local strain still remains. In the sample of Ag deposited 1L MoS2, the strain persists for more than 60 days (Fig. 3a), while the intensity of 𝐸

peak is reduced from the beginning (Fig. 3b), indicating the LSPR effect

is gradually decreased because of sulphurization. Finally, with continuous sulphurization of Ag surface (indicated by the thicker shell on Ag), the originally localized strain along the Ag-MoS2 boundary (shown as red color) is ultimately relaxed into a delocalized strain in the plane (Fig. 5i). The representing color for strain is changed from red to purple with an extended range in MoS2. We noted that the relaxation of the localized strain appears not to synchronize with the sulphurization of Ag, because the sulphurization starts immediately after the Ag deposition as revealed by the monotonic decrease of [I(𝐸 )/I(𝐸 )] with time (Fig. 3b). However, the strain relaxation occurs in a certain period of time after Ag deposition (see Fig. 3a, 14


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~70 days for 1L, ~50 days for 2L, and ~30 days for 3L). It suggests that the sulphurization of Ag is probably not the dominant factor to drive the strain relaxation. The relaxation of strain may result from the mechanical instability at the MoS2-substrate interface (for 1L) or the top-most MoS2-underlying MoS2 interface (for 2L or more layers). It has been reported that the MoS2-SiO2 interface has a relatively stronger mechanical interaction compared to the MoS2-MoS2 interlayer interface in our previous studies.67 For Ag/1L-MoS2, the stronger interaction at the MoS2-substrate (SiO2/Si) interface pins the 1L MoS2 more tightly, which offers a better mechanical stability of the interface and slows down the strain relaxation process. For Ag on 2L and few-layer MoS2, however, the local strain relaxes more easily because of the weaker van der Waals interaction between the top-most MoS2 layer and the underlying MoS2 layers. These results suggest that the substrate plays a great role in the redistribution of localized strain for monolayer MoS2, while the interlayer interaction does for two or more layers of MoS2.

Figure 5. Raman peak position changes with time for 𝐸 and 𝐸 modes up until fully stabilized state for Ag deposited (a) 1L, (b) 2L, and (c) 3L MoS2. The data for 3L is cut off at 48 days because the 𝐸 peak totally disappeared after that. (d-f) The comparison of Raman spectra for pristine (as exfoliated) MoS2 and Ag deposited MoS2 after a long period of time (>150 days). Illustrations of strain evolution in Ag-1L MoS2 heterosystems for (g) fresh sample, (h) aging sample where Ag going through sulphurization and forming thin sulphurized layer at the outer surface, and (i) sample with thicker sulphurized Ag layer with a significant strain delocalization. The extent of the localized strain along the Ag/MoS2 boundary is depicted by the gradient color scheme (i.e., red color refers to the maximum strain and the blue color refers to the minimum strain in the MoS2 layer). 15



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Conclusion In summary, we analysed and reported the kinetics of metal induced strain relaxation process on assorted ultra thin MoS2 sheets with different thicknesses following the metal aging through time elapse. The gradual changes in Raman spectra with time evolution provide insight understanding to the behaviour of strain relaxation and interface interaction of supported Ag-MoS2 heterostructures. The speed of the strain relaxation process is found to be dependent on the number of MoS2 layer where monolayer MoS2 bearing the longest time of relaxation and thicker layer MoS2 hosting the faster relaxation process. The metal induced localized strain at the Ag-MoS2 boundary on assorted MoS2 layers gradually relaxes up to a point where the strain is distributed more or less homogenously across the MoS2 sheet as delocalized strain and theAg-MoS2 heterostructure inhereits the stabilized strain residue throught the entire surface. Our results are helpful to understand the evolution of properties at metal-MoS2 interfaces, and to optimize the MoS2based device fabrications.

Experimental Atomically thin MoS2 samples of well-defined crystallographic orientation were exfoliated from bulk MoS2 crystals onto Si substrates covered with a 300-nm-thick SiO2 layer. Single- and few-layer MoS2 films were first identified by the optical contrast, and then confirmed by the Raman spectra, where the Raman shifts of E

and A

depended on the layer thickness.48,52 Raman and PL measurements were performed in

ambient condition using Horiba-JY LabRAM 800 micro-Raman system with 532 nm laser excitation. The Raman spectral resolution was ~1.5 cm-1. The optical beams were focused on the sample with a spot diameter of ~2 μm. A low laser power of ~200 μW was used to prevent overheating of MoS2. SEM images were taken with SUPRA 55VP high resolution SEM system and XPS data were obtained by using X-ray photoelectron spectroscopy (Thermoﬁsher, ESCALAB 250 system, Al Kα). The operating X-ray spot size

16


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is ~500 µm, which can cover tens of MoS2 flakes with varying thicknesses on the SiO2/Si substrate and the detection depth is 1~3 nm. The setup enables the XPS signals collected to originate from multiple Ag dendrite-MoS2 structures regardless of the MoS2 thickness under the X-ray spot. As such, our XPS measurement could detect the extent of the surface sulphurization of Ag nanostructures, and the data collected by changing the sample or moving the X-ray spot position would give consistent results. The samples were kept in a dry nitrogen box for storage before and after performing any measurement.

Supporting Information (1) SERS enhancement for monolayer MoS2 after Ag deposition. (2) Illustration of experimental condition while in the Raman measurement for the Ag deposited MoS2 sample. (3) Illustration of Raman spectra for bulk MoS2 before and after Ag deposition.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] ORCID Rongming Wang: 0000-0003-4075-6956 Kai Liu: 0000-0002-0638-5189 Author Contributions #

Y. A. M. and Y. H. S. contributed equally. The manuscript was written through contributions of

all authors. All authors have given approval to the final version of the manuscript. Notes The authors declare no competing financial interest. 17



Page 18 of 22

ACKNOWLEDGMENT We thank Fangtao Li for his assistance in Ag deposition and Prof. W. M. Lau for valuable discussions on XPS data interpretations. This work was supported by the National Natural Science Foundation of China (Grant No. 11604010, 11674023), 111 Project (No. B170003), Open Research Fund Program of the State Key Laboratory of Low-Dimensional Quantum Physics (KF201611) and the Fundamental Research Funds for the Central Universities (FRF-BD-17004A).

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