## Introduction

Surface/defect engineering of two-dimensional (2D) materials by laser exposure, ion irradiation, and plasma treatment has been studied actively because the modification of the surface can totally change the properties of 2D materials due to large surface-to-volume ratio1,2,3,4,5,6,7,8,9,10,11,12,13. Such a surface/defect-engineered 2D materials have a potential for high-performance field-effect transitor13,14,15, sensor16, and catalyst17,18,19. Among various methods for defect generation, the plasma treatment has shown great potential for transition metal dichalcogenides (TMDs), such as MoS2, MoTe2, and WSe2, due to high controllability and uniformity, leading to chemical functionalization, heteroatom doping, defect generation, and etching12,20,21,22,23. Recently, it was reported that Janus-type membrane24,25 of MoSSe and WSSe can be fabricated by selenizing the hydrogenated top surface26,27. Moreover, several studies have shown that defects in MoS2 can enhance catalytic performance for water splitting and hydrogen evolution reaction17,22,23,28,29,30,31. Nevertheless, the structural evolution of MoS2 during hydrogen plasma treatment was not studied due to the high reactivity of hydrogen plasma with MoS2 and the instability of plasma-treated MoS2. Thus, systematic study on the structure and properties of plasma-treated MoS2 is required to develop a method to slightly modify the surface or controllably generate a specific density of defects in three-atom-thick MoS232.

Here, we investigated the evolution of defect formation in monolayer MoS2 during exposure to low-energy indirect hydrogen plasma, which enables atomic precision etching. The sulfur atoms of top layer in three-atom-thick MoS2 were able to be removed with preserved molybdenum atomic layer and no damage to sulfur atoms of bottom layer. The sulfur vacancies started to form in the beginning of plasma treatment and became dominant up to 1.58 × 1014 cm−2 (for 40 min-treated sample) as the plasma treatment proceeded, resulting in plasma-treated MoS2 with preserved 2H-structure. Finally, nanocracks were generated along the zigzag orientation of MoS2 to relax the defect-induced biaxial tensile strain. These nanocracks appeared when more than 50% of sulfur atoms were removed from the top layer. The optical, electrical, and mechanical properties of the plasma-treated MoS2 were also explored to correlate them to the structural changes of MoS2. The photoluminescence (PL) intensity gradually decreased during plasma treatment and it was finally quenched. The conductivity of MoS2 also decreased by two orders of magnitude due to high density of defects, in contrast to theoretically expected metallic property of MoSH. Furthermore, the lateral tip-surface contact stiffness measured by lateral force microscopy increased by ~50% after plasma treatment. Our work shows that atomic precision etching is possible and the properties of MoS2 can be controlled by surface/defect engineering.

## Results and discussion

### Evolution of defects in plasma-treated MoS2

We used a home-built remote hydrogen plasma system, consisting of three ring-shaped electrodes and a low-frequency plasma generator, to desulfurize only the topmost atomic layer of monolayer MoS2 in an atomically precise and physically mild manner (Fig. 1a, b)32,33,34. The plasma was generated at the center of three electrodes and monolayer MoS2 crystals grown by chemical vapor deposition (CVD)35 were placed away from the plasma ignition area to avoid direct ion bombardment (see the “Methods” Section for details of the MoS2 growth and remote hydrogen plasma treatment). To investigate the structural change in the desulfurized MoS2 with hydrogen plasma treatment, we measured Raman spectra of MoS2, which is sensitive to perturbation of crystal structure induced by deformation and doping (Fig. 1c). The Raman spectra of hydrogen plasma-treated MoS2 (HP-MoS2) showed significant shifts and broadening of E′ and A1′ vibration modes, which are attributed to strain and doping10. In addition, disorder-related Raman peaks emerged as indicated by transversal optical, longitudinal optical, and out-of-plane optical modes at 342, 366, and 413 cm−1, respectively, in Fig. 1c36. As summarized in Fig. 1d, the E′ peak, in-plane vibration mode of Mo-S bond, red-shifted from 382.2 to 371.6 cm−1, meanwhile the A1′ peak, out-of-plane vibration mode, blue-shifted from 402.6 to 409.5 cm−1. As a result, the frequency difference between two peaks increased from 20.4 to 37.9 cm−1. Such notable peak shifts correspond to the formation of sulfur vacancies37. High density of sulfur vacancies at the top layer of MoS2 weakens the in-plane vibration in the Mo-S bonds, leading to a weaker restoring force constant and lower phonon frequency of E′. For frequency of A1′, restoring force constant also decreased with sulfur vacancies. However, the removal of top sulfur atoms allows out-of-plane vibration of static molybdenum atoms, then the resonance frequency of A1′ raises as we observed. Also, the intensities of two peaks gradually decreased after plasma treatment, which attribute to the degradation of MoS222,38,39,40,41.

### Atomic structure and defects in MoS2 after treatment

To further examine the atomic structure of HP-MoS2, we utilized transmission electron microscopy (TEM) and high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM). The CVD-grown MoS2 samples were transferred onto the graphene-coated holey carbon TEM grid, followed by the hydrogen plasma treatment. The graphene supporting film enables high-resolution imaging of the transferred MoS2 by TEM due to the dissipation of electrostatic charging and heating under e-beam (see details in the “Methods” Section)42. In the HAADF-STEM, the image contrast is determined by the Z-contrast mechanism, IZn, where I and Z are image contrast and atomic number. The index (n) is the number between 1 and 2, which depends on the scattering angles43. In the MoS2 supported on graphene-coated TEM grid, the molybdenum atoms with Z = 42 look brighter than the sulfur atoms with Z = 16 and carbon atoms with Z = 12, so that the graphene membrane can be ignored in our images. The HAADF-STEM images of the Fig. 2a, b show that the as-grown MoS2 has the hexagonal structure of 2H-phase with small number of sulfur vacancies as indicated by yellow circles. The estimated density of sulfur vacancies in the as-grown MoS2 was 1.60 × 1013 cm−2, which shows the high quality of the sample, compared to MoS2 samples used in other works44,45,46. The density of vacancies was estimated by the direct counting in obtained images to determine a quantitative value. Meanwhile, the HP-MoS2 treated for 40 min showed significantly increased density of sulfur vacancies (1.58 × 1014 cm−2) as shown in Fig. 2c, d. The HP-MoS2 showed mainly two types of vacancies at sulfur positions, monosulfur (VS) and disulfur vacancies (VS2, indicated by green circles). Note that the calculated formation energy of VS2 is relatively high as 4.14 eV45. Therefore, VS2 is formed by the generation of VS in topmost layer while VS already resided in the bottom layer, not directly generated by hydrogen plasma treatment (See Supplementary Fig. 1 for schematic diagram of VS2 formation). In addition, the sulfur vacancies are aligned preferentially along the zigzag orientation (indicated by red arrows in Fig. 2c) as expected by theoretical caculation47,48,49. The higher magnification color-contrasted HAADF-STEM images of the as-grown MoS2 (Fig. 2b) and HP-MoS2 (Fig. 2d) clearly show that the number of sulfur vacancies increased after hydrogen plasma treatment. From the line intensity profiles of the HAADF-STEM images (Fig. 2e), the local sites and types of sulfur vacancies were identified by using the different local intensities of Vs and VS2. For statistical analysis, we extract the local intensities at the atomic sites of Mo and S as shown in the histograms of Fig. 2f. The Mo and S atoms showed two distributions around 1.5 and 1.0, respectively. The local intensity distributions of the Mo atoms in the as-grown MoS2 and HP-MoS2 showed negligibly small difference, meanwhile the local intensity of two S atoms (S2) showed a significant decrease due to formation of the sulfur vacancies. The increased distribution peak of VS is located at 0.6 in the histograms. Even though a few VS2 were observed (very small intensity at 0.2), it is not clear that those are generated during hydrogen plasma treatment since the as-grown MoS2 probably has some missing sulfur atoms at bottom layer before the hydrogen plasma treatment. Nevertheless, it should be noted that VS significantly increased after hydrogen plasma treatment. However, it is still difficult to verify the presence of hydrogen atoms bonded to Mo atoms (Mo-H) or sulfur atoms (S-H), which were reported as a result of hydrogenation26,50.

### Chemical characterization of the plasma-treated MoS2

We performed chemical composition analysis using Auger electron spectroscopy (AES) and time-of-flight secondary ion mass spectroscopy (ToF-SIMS) to confirm the relationship between crack formation and the chemical composition. As shown in Fig. 3a, the AES features around 156 (S LVV) and 190 eV (Mo MVV) for the MoS2 samples. The intensity of the peak at S decreases while that of Mo peak is constant, implying that sulfur vacancies were produced without molybdenum vacancies, as observed in TEM and STEM. In the AES result, chemical compositions were extracted based on peak-to-valley height proportional to the atomic concentration and sensitivity factor51. Figure 3b shows that the calculated atomic ratio (S/Mo) of the as-grown MoS2 is ~2, close to the ideal stoichiometry of MoS2. However, after treatment, the S/Mo ratio gradually decreased to 1.60 and 1.15 for HP-MoS2 (10 min) and HP-MoS2 (30 min), respectively. The ToF-SIMS depth profiling of the as-grown MoS2 and HP-MoS2 (30 min) was performed with Cesium (Cs) sputtering (0.5 keV) to reveal the vertical position of sulfur vacancies. Figure 3c shows the ToF-SIMS negative spectra of sulfur (S) as a function of the elapsed sputtering time. The intensity of sulfur in the as-grown MoS2 is higher than HP-MoS2 (30 min) at early sputtering time and difference between them becomes negligible, indicating that the indirect hydrogen plasma removes only top sulfur layer of the three-atom-thick MoS2 as depicted in the inset of Fig. 3c.

### Formation of nanocracks in the plasma-treated MoS2

In the HP-MoS2 (30 min), we observed the formation of nanocracks as shown in Fig. 4. Although the surface of HP-MoS2 (30 min) looked smooth and clean in the optical microscopic image of Fig. 4a, many line patterns with a hexagonal shape were observed in the dark-field optical microscopic image as shown in Fig. 4b. These line patterns were investigated by using atomic force microscopy (AFM), scanning electron microscope (SEM), and TEM as shown in Fig. 4c–h. The line patterns were not detected in the AFM topography image of Fig. 4c, probably due to the smaller width of the line rather than AFM tip radius and roughness of the SiO2 substrate. However, the AFM friction image of Fig. 4d showed clear line patterns due to the lower friction of line regions. The SEM image of Fig. 4e clearly shows that the hexagonal patterns are 1 μm in size and angle distribution of the lines relative to the edge of the MoS2 crystal has preferential orientations of 0°, 60°, and 120° as summarized in Fig. 4f, which indicates that the line patterns are associated with the crystallographic orientation of the MoS2. For TEM imaging of HP-MoS2, we treated hydrogen plasma followed by transferring MoS2 sample onto a silicon nitride TEM grid. The TEM image of Fig. 4g shows that a narrow lines (indicated by yellow arrows) are between two regions (termed ‘A’ and ‘B’). The selected area electron diffraction pattern in the inset of Fig. 4g shows that two regions of A and B have the same hexagonal symmetry, indicating that two regions are separated from the MoS2 grains, not misaligned grains. On the other hand, two-line patterns are aligned along the zigzag direction of the MoS2 as indicated by red arrows. The HAADF-STEM image of Fig. 4h clearly shows that the line between A and B regions is a nanocrack with a width of ~10 nm. However, the edges of the crack are rough, compared to highly sharp edges of chemically etched MoS222.

### Relaxation of defect-induced strain and formation of nanocracks

To investigate the strain effect on the HP-MoS2 (30 min), we analyzed the Raman spectra of cracked and uncracked regions as shown in Fig. 5a. The peak positions of E′ and A1′ modes are influenced by doping and strain52,53,54. Both regions showed the increased separation between the E′ and A1′ (Δ ~ 33 cm−1) after plasma treatment due to the formation of vacancies as we discussed. The cracked region showed blue shifts of two Raman peaks compared to those of the uncracked region as shown in Fig. 5b. This peak shift in defective MoS2 has been explained by strain effect53,54. The blue shift of the Raman peaks can be attributed to a decrease of bond length under compressive strain and increase of the force constant. On the other hand, red shift can be result from increasing tensile strain, attributing to the reduction of the force constant in MoS2. Thus, cracked region undergoes less tensile strain than uncracked region, meaning that nanocraks generates during the relaxation of tensile strain. The alignment of vacancies (Fig. 2c) is an intermediate step between vacancies and cracks, and transition occurs when the strain or defect size exceeds the certain value. We assumed that the crack generates after more than 50% of sulfur atoms in the topmost layer is etched (S/Mo value between 1.5 and 1.0, which is calculated from AES result), as a result of relaxation of strong defect-induced strain. According to our observations, the etching process for the top-sulfur atoms and the following crack formation are speculated as depicted in Fig. 5c. First, atomic defects are generated randomly in topmost surface of monolayer MoS2 after hydrogen plasma treatment. As we discussed above, the defect is thought to dominantly consist of sulfur vacancies (VS, VS2) After longer treatment, sulfur vacancies preferentially generate along zigzag orientation as shown in the third image. Sequentially, when the half of top sulfur atoms is etched, cracks generate along zigzag orientation during relaxation of defect-induced strain. Furthermore, the HP-MoS2 (180 min) was separated into small islands with scrolled edges (for optical and AFM image see Supplementary Fig. 2). The low S/Mo ratio of 0.64 indicates that even bottom sulfur atoms of MoS2 are etched (Supplementary Fig. 3 contains AES spectra).

### Evolution of optical and electrical properties

To investigate the effect of defects on the optical properties, the PL spectra of the MoS2 were measured after plasma treatment (Fig. 6a). Two emission peaks at ~1.84 and ~2.0 eV, corresponding to A and B excitonic transitions, were rapidly decreased and finally queched as shown in the inset of Fig. 6a as the plasma treatment proceeded. To understand the origin of PL spectral change, we used Gaussian fits to extract the contribution of excitons as shown in Fig. 6b. The PL spectra were deconvoluted into four peaks of A-exciton (A), negatively charged trion (A), B-exciton (B), and defect-induced exciton (XD)55. The A-peak and A--peak were dominant in the PL spectra, and B-peak was also evident in emission, while XD was absent for as-grown MoS2. By contrast, B-peak became dominant and XD appeared after treatment. This behavior results from the increase of non-radiative recombination due to a high density of defects56,57. In as-grown MoS2, radiative recombination dominantly occurs at ground state, a strong A-peak and weak B-peak are expected. In HP-MoS2, non-radiative recombination becomes common due to defects and reduces effective A-exciton lifetime, while B-exciton lifetime is moderately impacted57. Also, A-peak decreaes together as A-peak decreases. The energy of excitons shifts to lower energy level with increasing treatment time. The PL shift during treatment can be explained based on strain effect as we observed in Raman spectroscopy58. The energy states of the valence band and conduction band originate from the 3p orbital of S atoms and 4d orbital of Mo atoms. The strain effect leads interatomic distance change between the Mo and S atoms, and further modifies the overlap between those orbitals, which results in energy shift in band gap. Tensile strain causes the reduction of the distance between the S atoms and the increase of the distance between Mo and S atoms, which induces a rise of the energy at the top of the valence band near the Г point and at the bottom of the conduction band between the K and the Г points and results reducing band gap59. The evolution of spectral weight and photon energy of each emission following treatment time is summarized in Fig. 6c. The change of electrical properties was investigated by plasma-treating the MoS2 FET as shown in the inset of Fig. 6d. We fabricated PMMA window only over the channel area to prevent damage to the metal contact area. The output curves (Ids–Vds) show that the treated MoS2 becomes more resistive with increasing treatment time (Fig. 6d). Note that the hydrogen-based bandgap modulation of MoS2, such as hydrogenation and vacancy generation, has been studied by experimental and computational method50,60. The hydrogenation of MoS2 includes the formation of S-H bond by surface adsorption and atomic substitution of sulfur to hydrogen, showing bandgap reduction or semiconductor to metal transition. Similarly, it has been reported that sulfur vacancies act as an electron donors and induce n-doping of MoS261,62. However, unlike the previous results, we observed the decreasing conductivity of MoS2-FET, which is presumed to be due to the convolution of defect and strain. The sulfur vacancies and strain-induced cracks in HP-MoS2 act as electron-hole scattering sites, which decrease the conductivity of carriers, and finally show insulating behavior after degradation of MoS2.

### Mechanical properties of the plasma-treated MoS2

To verify the tribological change in the HP-MoS2, the local frictional properties of the HP-MoS2 were characterized using AFM. Fig. 7a–c presents the high-resolution topography and lateral force images of the as-grown MoS2 and HP-MoS2 samples treated for 10 and 30 min, respectively. The atomic-scale sawtooth pattern of the as-grown MoS2 was clearly observed in both topography and lateral force images, indicative of high crystallinity and uniformity of the as-grown MoS2. As shown in the fast Fourier transform (FFT) of the lateral force image of Fig. 7d, the as-grown MoS2 exhibits a set of hexagonal spots with a distance of ~0.3 nm, close to the lattice parameter of MoS263. This indicates that the single stick-slip occurs by periodic potential wells consisting of adjacent top-sulfur atoms64,65. The onset of stick-slip occurs when the value, $${{{\rm{\eta }}}}=\frac{2\pi {E}_{0}}{k{a}^{2}}$$ becomes larger than 166. Here, E0, k, and a denote the barrier of potential well, effective contact stiffness, and lattice space, respectively. The E0 reflects interaction between tip and sample made from every lattice point. Therefore, the emergence of regular sawtooth signal in Fig. 7a implies that the as-grown MoS2 contains constant energy barrier and lattice spacing together with low contact stiffness, meeting the condition of η > 1. On the other hand, HP-MoS2 contains irregular energy barrier and large lattice spacing in defective sites, meeting the condition of η < 1 and losing periodic sawtooth signal. Meanwhile, the FFT images of the HP-MoS2 show just two broad strips with the reduced intervals as shown in Fig. 7e, f. The line profile shows losing periodicity between lattice point in HP-MoS2 (Fig. 7g). Sharp peaks were observed in as-grown MoS2, while periodic distance became broader in HP-MoS2 (10 min) and nearly disappeared in HP-MoS2 (30 min). Nevertheless, the surface roughness of the HP-MoS2 is similar to that of the as-grown MoS2 irrespective of the plasma treatment time (Fig. 7h). Even though the surface roughness increased somewhat from ~0.003 nm (for as-grown) to ~0.03 nm (for 10- and 30 min-treated ones), the topographical change is still negligible compared to the nominal step height of MoS2 (~0.65 nm).

In addition, the lateral stiffness of the as-grown MoS2 measured from the regular sawtooth patterns of the friction loop was 15 ± 2.3 mV/nm (Fig. 7i). However, the prolonged plasma treatment increased the lateral stiffness to 22.9 ± 16.7 mV/nm for the HP-MoS2 treated for 10 min, finally leading to no detectable sawtooth pattern for the HP-MoS2 treated for 30 min.

The measured total lateral stiffness (ktot) is described as

$$\frac{1}{{k}_{{tot}}}=\frac{1}{{k}_{{lever}}}+\frac{1}{{k}_{{cont}}}$$
(1)

where klever, and kcont are the lateral stiffness of cantilever and tip-sample contact, respectively67. When the tip approaches, the surface of MoS2, kcont is approximated as the out-of-bending stiffness of 2D layer (K2D, bend), so that ktot is governed by the weakest spring68. As shown in Fig. 7i, we can clearly observe the stick-slip from as-grown MoS2 at atomic scale, meaning that the effect of tip-sample contact is dominant factor as we discussed. However, in the HP-MoS2, the stick-slip was weaknened as η decreased below 1 due to reduced contribution of kcont to ktot. This indicates that HP-MoS2 has the increased out-of-bending stiffness. Instead, the effect of klever becomes dominant in Eq. (1), leading to dramatic increase in the average friction of MoS2. The friction ratio between MoS2 and SiO2 (see Supplementary Fig. 4a for friction ratio of MoS2 to the surrounding SiO2), shows that as-grown MoS2 is less sticky than SiO2 substrate, while the drastic reversion was observed in HP-MoS2 (30 min) due to enhanced cantilever bending effect at defective surface. In addition, the energy dissipation calculated from the area of the friction loop shows increasing fluctuation after treatment as depicted in Supplementary Fig. 4b, which is in agreement with that such high fluctuation occurs in the vicinity of sulfur vacancy of monolayer MoS269.

According to structural characterization, we confirmed that the indirect hydrogen plasma primarily generates sulfur vacancies in the topmost surface of MoS2 without harming its basal plane. Furthermore, defects and defect-induced strain in MoS2 affect optical, electrical, and mechanical properties.

## Conclusion

We observed the evolution of defect formation during the desulfurization process, especially, one-atom-thick removal of sulfur atoms in three-atom-thick monolayer MoS2. We exposed low-energy remote hydrogen plasma to demonstrate atomic structure and chemical composition of monolayer MoS2 with a mild manner, which allow us to observe changes systematically with atomic precision. The detailed analysis revealed that low-energy plasma leads to generate sulfur vacancies at the topmost surface of MoS2, and both molybdenum and sulfur atoms in lower layers remained intact. In addition, cracks along the zigzag edge of MoS2 appeared during the relaxation of defect-induced biaxial tensile strain when sulfur vacancies possessed half of the topmost sulfur layer. We believe that our observations provide information on the evolution of atomic structure during surface engineering process and potential of plasma-based precise surface modification techniques of TMDs. Such modulation of defect could be further extended potential of MoS2 and other TMDs for future electronics and catalytic reactions.

## Methods

### Sample synthesis

Monolayer MoS2 on SiO2 (285 nm)/Si substrate was synthesized using an APCVD (atmospheric pressure CVD) system. Before the growth process, substrates (2.3 × 2.3 cm2) were cleaned via sonication in acetone and isopropyl alcohol for 10 min each. Molybdenum trioxide powder (MoO3, 3 mg, Sigma-Aldrich) loaded quartz boat was heated to 750 °C at center of the furnace. Another quartz boat filled with sulfur powder (S, 180 mg, Sigma-Aldrich) was placed at 23 cm away from the center toward the upstream (temperature of sulfur is measured as 220 °C). The substrate was positioned upside-down on MoO3 powder filled quartz boat. The temperature of furnace was ramping up to the reaction temperature of 750 °C with a ramping rate of 50 °C /min. Reaction maintained for 15 min with supplying Ar gas (300 sccm) for carrying sulfur and MoO3-x vapors to the substrate. After growth, the furnace was cooled naturally to 600 °C, followed by rapid cooling.

### Plasma treatment

The hydrogen plasma treatment was performed using home-built indirect plasma system (L-Gen; Femto Science), which is embedded in conventional CVD system. The plasma generator is composed of three electrodes and power generator, whose adjustable parameters are generation power and time. In three electrodes configuration, electrodes at both ends are connected to ground, and center electrode is worked as low-frequency electrode for plasma generation, respectively. This plasma generator is movable along quartz tube to control plasma-to-sample distance. Owing to the plasma-to-sample distance, plasma generation area and target sample are separated along the reaction chamber. Thus, MoS2 is exposed to the reactive plasma species with a minimized bombardment effect, which realizes ‘indirect’ hydrogen plasma treatment. The target sample was located 3 cm away from plasma ignition for indirect plasma treatment. The plasma generation power was 20 W with 50 kHz of fixed frequency, and hydrogen gas flow was 20 sccm (1.2 × 10−1 Torr).

### Structural characterizations

The optical properties (Raman and PL) of the MoS2 were measured using Raman spectroscope (inVia; Renishaw) with a wavelength of 532 nm and a spot size of 1 μm. To prevent the degradation of samples during measurement, the laser power (2.2 mW) was carefully chosen. Both Raman and PL spectra at the same position on MoS2 before and after plasma treatment were obtained. Observation of the surface morphology of MoS2 and HP-MoS2, SEM (VEGA 3; Tescan) was used under 30 kV of acceleration voltage, and AFM (NX-10; Park Systems) was used under contact mode. To investigate tribological variation of MoS2 with different treatment time, high-resolution AFM (Cypher ES; Oxford Instruments) was utilized under contact mode. In order to make more accurate comparison, the loading force of tip was maintained at 2 nN for every measurement, and the fast scan direction was oriented parallel to one edge of each triangular MoS2 single crystal and the scan velocity was fixed at 20 nm/sec. The root-mean square roughness values (Rrms) were determined by scanning more than a few different regions for each sample. Considering the atomically smooth surface of samples, ~100 nm × ~100 nm areas were scanned.

### Transmission electron microscopy measurements

To prepare graphene support film for TEM and STEM measurement, CVD-grown graphene was transferred on holey carbon film deposited TEM grid (Quantifoil) by a direct transfer method70. TEM grid was placed on graphene grown Cu foil, and a droplet of isopropyl alcohol was dropped on the surface of TEM grid. After evaporation of isopropyl alcohol naturally, the sample was floated on copper etchant (0.1 M ammonium persulfate solution) for 2–3 h. The copper etched graphene/TEM grid was sequentially moved to deionized water to remove residual chemical several times and dried naturally. The CVD-grown MoS2 on SiO2 substrate was located on as-prepared graphene-coated holey carbon TEM grid for Fig. 2a–d or silicon nitride membrane (PELCO) for Fig. 3g, h using PDMS stamp method71. TEM and STEM imaging were performed by an atomic resolution TEM (JEM-ARM200F; JEOL) equipped with image and probe aberration correctors operated at 80 kV. HAADF-STEM images were obtained with a convergence angle of 23 mrad and the collection semi-angles from 40 to 160 mrad. ImageJ and Gatan Digital Micrograph were used to process the HAADF images.

### Electrical characterizations

To fabricate MoS2-based FETs using CVD-grown MoS2 on SiO2 substrate, source and drain electrodes were defined by an e-beam lithography using a nanolithography system (PIONEER 2; Raith). We deposited metal electrodes (Al 20 nm/Cr 5 nm/Au 30 nm) with an e-beam evaporator (Korea Vacuum Tech.) in a high vacuum of 10−8 Torr. The lift-off process was performed in acetone. Electrical measurements were carried out by semiconductor parameter analyzer (Keithley 4200; Keithley) at room temperature under ambient condition.

### Chemical characterizations

The chemical composition of MoS2 and HP-MoS2 was performed by AES and ToF-SIMS. The AES data were obtained by Auger nanoprobe (PHI 700; ULVAC-PHI) equipped with cylindrical mirror analyzer (energy resolution < 0.5%). The base pressure is under 5 × 10−10 Torr. To avoid the charging effect of SiO2 substrate, the samples were tilted during measurement process. The electron beam current of 10 nA with an acceleration voltage of 3 kV was used. The quantitative analysis of sulfur and molybdenum atoms was performed using sensitive factors derived from as-synthesized MoS2 to directly compare the effect of sulfur vacancy generation. ToF-SIMS (TOF.SIMS-5; ION-TOF) was performed in negative mode using a Bismuth primary ion beam at 30 keV of beam energy with 1 pA of current. The sputtering was processed using a Cs at 0.5 keV of beam energy with 40 nA of current to achieve depth profile. The measurements were performed using facilities of the Korea Institute of Science and Technology (Seoul, South Korea).