Chalcogen-vacancy group VI transition metal dichalcogenide nanosheets for electrochemical and photoelectrochemical hydrogen evolution

Abstract

Solar photoelectrochemical water splitting using the catalytic hydrogen evolution reaction (HER) is one of the most promising methods for producing hydrogen energy. Herein, we report the facile synthesis of chalcogen vacancy-rich (5–20%) transition metal dichalcogenide (TMD, i.e., MoS₂, MoSe₂, WS₂, and WSe₂) nanostructures for the (photo)electrocatalytic HER. Atomically resolved scanning transmission electron microscopy revealed the concentration of chalcogen (S and Se) vacancies, as well as the formation of double chalcogen vacancies at higher concentrations. Photocathodes consisting of a Si nanowire array sheathed in TMD layers exhibited excellent photoelectrocatalytic performance under AM 1.5 irradiation conditions. The photocurrent reached 30 mA cm⁻² (at 0 V vs. RHE) with a high faradaic efficiency (90%) and negligible degradation of the HER for 3 h. The free-standing TMD nanosheets exhibit excellent electrocatalytic HER performance (an overpotential of 162–202 mV vs. RHE for 10 mA cm⁻² and a Tafel slope of 70–84 mV dec⁻¹ in 0.5 M H₂SO₄), which can be correlated with the vacancy concentration. A detailed structural analysis suggested that the chalcogen vacancies play a major role in determining the catalytic activity toward the HER, which in turn is responsible for the enhanced PEC performance.

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1. Introduction

The drastically increased consumption of fossil fuels has caused serious environmental crises such as global warming and air pollution, prompting scientists to explore renewable and clean energy sources. Hydrogen energy is presently considered the most promising candidate to replace the use of fossil fuels by hydrogen-based energy economy. Solar photoelectrochemical (PEC) water splitting is a carbon neutral technology because it produces hydrogen using practically unlimited resources (i.e., sunlight and water). The first PEC cell was fabricated in 1972 using a TiO₂ photocatalyst as the photoanode under UV irradiation.¹ Since then, there have been tremendous efforts to develop cost-effective and stable semiconductor photoelectrodes based on earth abundant elements for efficient solar energy harvesting and catalytic water oxidation or reduction reactions.^2–10

Si has a band gap (E_g) of 1.18 eV, which is narrow enough to efficiently absorb solar energy. Its conduction band position is suitable for the hydrogen evolution reaction (HER). Furthermore, the Si technology is well established, and the material price has decreased significantly in the past few years. Therefore, there is a keen interest in using Si in both the photoanodes and photocathodes of solar PEC cells.^2,11–21 A powerful strategy is to apply low-dimensional nanostructures that can increase the surface area for light absorption.^{11,17,19–21} In particular, nanowires (NWs) can produce a desirable one-dimensional pathway for charge collection along the electrode/electrolyte interface.

Despite the significant progress in employing Si photoelectrodes for solar water splitting PEC cells, their durability and energy conversion efficiency are still far below the goal (>2000 h and 10%, respectively). The Si photoelectrode is intrinsically unstable due to its poor catalytic activity toward the water splitting reaction, which causes surface photo-corrosion. These problems could be avoided by employing efficient catalysts to increase the reaction rate and retard the surface oxidation. Recently, two-dimensional (2D) earth abundant metal dichalcogenide (TMD) nanomaterials with the chemical formula MX₂ have been extensively investigated as HER electrocatalysts. Excellent electrocatalytic activities have been reported for Group VI TMDs (i.e., MoS₂, MoSe₂, WS₂, and WSe₂).^22–24 At the same time, TMDs are often used as catalysts for Si-based photocathodes.^25–39 The HER electrocatalytic activity of TMDs can be further improved by introducing S or Se vacancies, because these chalcogen vacancies facilitate the HER by adsorbing proton or water at the metal sites.^40–53 However, the enhancement effect of chalcogen vacancies has not been verified for the Si-TMD photocathodes yet.

In this study, we synthesized four TMD nanomaterials (MoS₂, MoSe₂, WS₂, and WSe₂) with two distinctive forms: (1) shell layers on the pre-synthesized Si NW array photocathode for water splitting PEC cells (in 0.5 M H₂SO₄) and (2) freestanding nanosheets for the electrocatalytic HER. Both forms of TMDs are enriched with chalcogen vacancies and synthesized using a common sulfurization/selenization step. Atomically resolved scanning transmission electron microscopy provides evidence for the vacancies at the chalcogen sites. All the Si–MoS₂, Si–MoSe₂, Si–WS₂, and Si–WSe₂ photocathodes exhibited excellent PEC performance, with a photocurrent of 20–30 mA cm⁻² (at 0 V vs. RHE) in 0.5 M H₂SO₄ under AM 1.5 irradiation conditions. The electrocatalytic HER performance of freestanding TMD nanosheets was well correlated with the concentration of chalcogen vacancies. A detailed structural analysis shows that S/Se vacancies play a major role in enhancing the HER performance. These results demonstrate that the electrochemical and PEC catalytic efficiencies of the four Group VI TMDs could all be improved by introducing chalcogen vacancies.

2. Experimental section

Synthesis of Si-TMD NW arrays

The Si NW array was synthesized using Ag-assisted chemical etching of a Si(100) substrate, as has been described elsewhere.⁵⁴ The deposition of TMD nanosheets on the Si NW array was carried out as follows. A substrate of Si NW array was immersed in 0.4 M MoCl₅ (or WCl₆) ethanol solution for 5 min and then rinsed with ethanol. The metal ions deposited on the Si NW substrate were then sulfurized or selenized as follows. The NW substrate was placed in a quartz tube inside a chemical vapor deposition (CVD) reactor (electrical furnace), and S or Se powder (200 mg) was located in a quartz boat 3 cm away from the Si substrate. Mixed hydrogen (H₂)/argon (Ar) gas was continuously flowed at a rate of 20/100 sccm (mL min⁻¹) during the entire process. The temperatures of the S (or Se) powder and the Si NW substrate were, respectively, 200 and 450–550 °C during the sulfurization/selenization, and the reaction time was 30 min.

Synthesis of freestanding TMD nanosheets

First, molybdenum oxide (MoO₃) and tungsten oxide (WO₃) nanocrystals were synthesized using a hydrothermal method. Sodium molybdate dihydrate (Na₂MoO₄·2H₂O, molecular weight (MW) = 241.95 g mol⁻¹, 1 mmol) or sodium tungstate dihydrate (Na₂WO₄·2H₂O, MW = 329.85 g mol⁻¹, 1 mmol) and 30 μL HCl (conc. 37%) were dissolved in deionized (DI) water (20 mL) using a bath sonicator. The pH of the mixture was 1–2. The solution was transferred to a Teflon-lined stainless-steel autoclave reactor. The hydrothermal reaction was performed at 180 °C for 6 h in an electric oven. The product (MoO₃ or WO₃) was collected by centrifugation, washed thoroughly with DI water and acetone several times, and then vacuum-dried at room temperature. For sulfurization or selenization, 20 mg of MoO₃ (or WO₃) was placed inside the CVD reactor. The reaction time was 60–120 min. The other conditions are the same as those of the Si NW samples. The characterization method and (photo)electrochemical measurements are described in the ESI.†

3. Results and discussion

Fig. 1 shows a schematic of the synthesis procedure. (1) TMD nanosheets were synthesized directly on a pre-grown Si NW array. The NW surface was coated with metal ions by soaking the substrate in a MoCl₅ or WCl₆ ethanol solution. (2) Subsequent sulfurization/selenization was performed at 450–550 °C via the chemical vapor transport of the S or Se powder under H₂/Ar gas flow. (3) Separately, the freestanding TMD nanosheets were synthesized via the sulfurization/selenization of wire-like MoO₃ or WO₃ nanocrystals under the same conditions as those for the NW sample. The MoO₃ and WO₃ were synthesized using the hydrothermal reaction of NaMoO₄ and Na₂WO₄ at 180 °C, respectively. Fig. S1 (ESI†) shows the scanning electron microscopy (SEM) and high-resolution transmission electron microscopy (HRTEM) images, X-ray energy-dispersive fluorescence spectroscopy (EDX) data, and X-ray diffraction (XRD) patterns for the wire-like MoO₃ and WO₃ nanocrystals. The XRD patterns of the two series of TMD materials are shown in Fig. S2 (ESI†). The hexagonal (2H, space group = P6₃/mmc) phase was identified for MoS₂ (JCPDS No. 87-2416, a = 3.160 Å, c = 12.290 Å), MoSe₂ (JCPDS No. 29-0914, a = 3.287 Å, c = 12.925 Å), WS₂ (JCPDS No. 08-0237, a = 3.154 Å, c = 12.362 Å), and WSe₂ (JCPDS No. 38-1388, a = 3.285 Å, c = 12.982 Å).

Fig. 1 Schematic of the synthesis of two forms of TMD (MoS₂, MoSe₂, WS₂, and WSe₂) nanosheets. (1) For the photoelectrochemical HER, the shell layers of TMD on the Si NW array (Si-TMD) were synthesized in two steps: metal ion coating and sulfurization/selenization. (2) Sulfurization/selenization step using the thermal chemical transport of S and Se powders under H₂/Ar gas flow. (3) For the electrochemical HER, freestanding TMD nanosheets were synthesized in two steps: hydrothermal synthesis of MoO₃/WO₃ nanocrystals and sulfurization/selenization.

The TMD nanosheets were synthesized as shell layers on Si NWs, and these samples were referred to as Si–MoS₂, Si–MoSe₂, Si–WS₂, and Si–WSe₂. The tilted SEM view in Fig. 2a shows that the Si-TMD NWs were grown on a large area of Si substrates via Ag-assisted etching of Si(100) wafers. The side view shows that the Si NWs had a length of approximately 4 μm (Fig. 2b). The top views confirm that TMDs were grown on the Si NWs (Fig. 2c). Moreover, HRTEM images indicate that the surface of Si NWs was coated homogeneously with MoS₂ layers with a thickness of approximately 20 nm (Fig. 2d). Lattice-resolved TEM images revealed that the MoS₂ nanoflakes assembled into a polycrystalline shell sheathing the Si NW core. The d-spacing of the (002) planes was identified as 6.1 Å, which is consistent with the reference value of 6.145 Å. Fig. 2e shows the HRTEM images of Si–MoSe₂, in which the polycrystalline MoSe₂ layers have d₀₀₂ = 6.5 Å, consistent with the reference value of 6.46 Å. Fig. 2f and g show the HRTEM images of Si–WS₂ and Si–WSe₂, respectively. These shell layers also consisted of the nanoflake assembly. The WS₂ and WSe₂ nanoflakes have uniform d₀₀₂ = 6.2 and 6.5 Å, which are also consistent with the respective reference values of 6.18 and 6.49 Å.

Fig. 2 SEM images of Si-TMD NWs in (a) tilted and (b) side views. (c) Magnified images of the top part, showing the homogeneous deposition of TMD nanosheets on the Si NWs. HRTEM images of (d) Si–MoS₂, (e) Si–MoSe₂, (f) Si–WS₂, and (g) Si–WSe₂. The distance between adjacent (002) planes (d₀₀₂) in the TMD is 6.1, 6.5, 6.2, and 6.5 Å, respectively.

The high-angle annular dark-field scanning TEM (HAADF-STEM) image and EDX elemental mapping spectra (using the Mo L-shell, W M-shell, S K-shell, Se L-shell, and Si K-shell peaks) revealed that only the Mo and X elements were present in the shell part, while Si was localized in the Si NWs (Fig. S3, ESI†). The EDX and X-ray photoelectron spectroscopy (XPS) analyses provided [S]/[Mo] = 1.85, [Se]/[Mo] = 1.9, [S]/[W] = 1.9, and [Se]/[W] = 1.8 (see Fig. S4, ESI†). The corresponding concentrations were 7%, 5%, 5%, and 10%, which are half of those for the freestanding TMD nanosheets (i.e., 15%, 10%, 10%, and 20%).

The electronic structures of the samples were analyzed using XPS, showing that the binding energies were lower than those in the bulk, also owing to the S/Se vacancies (Fig. S5, ESI†). The magnitude of the redshift is nearly the same for all samples, probably because the vacancy concentration varied in the narrower range compared to that of freestanding nanosheets as described later. The lower concentration of chalcogen vacancies in Si-TMD is ascribed to the shorter time of the sulfurization/selenization step. The optimized reaction time was 30 min for the TMD shell layers. In contrast, the freestanding TMD nanosheets need at least 60 min to obtain a complete conversion of the oxide.

The SEM images of freestanding TMD nanosheets show that the thin nanosheets assembled into flower-like microspheres (Fig. S6, ESI†). Fig. 3a shows the HRTEM images of WS₂ nanosheets. The average thickness of the nanosheets is 2 nm based on the statistical distribution. The lattice-resolved image (Fig. 3b) of the basal plane (magnified in the marked area) shows that the distance between adjacent (100) planes (d₁₀₀) is 2.7 Å, which is close to the value of 2H phase WS₂ (2.73 Å). The HAADF-STEM image and EDX elemental mapping of MoS₂, MoSe₂, WS₂, and WSe₂ nanosheets show uniform distributions of metal (Mo or W) and chalcogen (S or Se) atoms over the entire samples (Fig. 3c). All four types of nanosheets have significant chalcogen vacancies since [S]/[Mo] = 1.7, [Se]/[Mo] = 1.8, [S]/[W] = 1.8, and [Se]/[W] = 1.6. These ratios were supported by the XPS analysis (Fig. S7, ESI†). Therefore, the chalcogen vacancies were estimated to be 15%, 10%, 10%, and 20%, respectively, for MoS₂, MoSe₂, WS₂, and WSe₂.

Fig. 3 (a) High-resolution transmission electron microscopy (HRTEM) image of WS₂ nanosheets and (b) lattice-resolved TEM image for their basal plane. The distance between adjacent (100) planes (d₁₀₀) is 2.7 Å. (c) HAADF-STEM image; EDX elemental mapping of Mo (L shell), W (M shell), S (K shell), and Se (L shell); and EDX spectra for MoS₂, MoSe₂, WS₂, and WSe₂. Atomically resolved HAADF-STEM images of the basal plane for (d) MoS₂, (e) MoSe₂, and (f) WSe₂, and line profiles for the selected areas (i), (ii), and (iii). The single (V_1S and V_1Se) and double (V_2S and V_2Se) chalcogen vacancies are marked by circles.

Atomic-resolution HAADF STEM images for the basal planes of MoS₂, MoSe₂, and WSe₂ were obtained, and the corresponding contour maps are shown in Fig. 3d–f, respectively. They all consisted of a well-defined hexagonal pattern of 3(M–X) corresponding to the 2H phase of the TMD. Darker intensities at the atomic sites suggest that these atoms are missing. The contour map clearly distinguishes the single/double vacancies from the bright yellow S or Se atoms.⁵⁵ In MoS₂, S sites missing one and two S atoms are marked by V_1S and V_2S, respectively, and they are approximately equal in number. In contrast, MoSe₂ has only sites missing one Se atom (marked by V_1Se). In the case of WSe₂, vacancies with one and two missing Se atoms (marked by V_1Se and V_2Se, respectively) are formed in equal numbers. The production of double vacancies (V_2S and V_2Se) in MoS₂ and WSe₂ is ascribed to the higher concentration of vacancies (20%) compared to that in MoSe₂ (10%). The line profiles for regions (i)–(iii) clearly show the two types (i.e. single and double) of chalcogen vacancies. The presence of S/Se vacancies was also supported by electron paramagnetic resonance (EPR), as shown in Fig. S8 (ESI†). A stronger signal was found in MoS₂ and WSe₂ than that in MoS₂ and WS₂, which is correlated with the higher concentration of vacancies.^43,51

The electronic structures of the samples were analyzed using XPS. The binding energy was lower compared to the corresponding bulk crystals, indicating increased metallicity that is probably due to the presence of S or Se vacancies (Fig. S9, ESI†). The higher concentration of vacancies in MoS₂ and WSe₂ induced a more significant redshift compared to MoS₂ and WS₂.

When synthesizing the freestanding TMD nanosheets or the TMD shell layers of Si NWs, a H₂ gas flow was essential because the sulfurization/selenization was not efficient. However, the H₂ gas flow also inevitably depleted the S/Se, owing to the reaction H₂ + S (or Se) → H₂S (or H₂Se) during/after the growth of nanosheets. In fact, H₂ annealing is a common method for producing vacancies in TMD nanomaterials.^45,50 Guo et al. reported the vacancy formation energy for monolayer MoS₂, MoSe₂, WS₂, and WSe₂ to be 2.35, 2.82, 3.38, and 2.81 eV, respectively.⁵⁶ Lee et al. calculated these values as 2.21, 2.64, 2.37, and 2.70 eV.⁵⁷ A lower formation energy implies more favored vacancy formation. Thus, these formation energies can explain the higher vacancy concentration of MoS₂ nanosheets than that of MoSe₂ nanosheets. However, there is a discrepancy in the values of WS₂ and WSe₂, which needs further studies to explain our results.

Now, we discuss the I–V curves measured for water splitting PEC cells using photocathodes of Si, Si–MoS₂, Si–MoSe₂, Si–WS₂, and Si–WSe₂ in 0.5 M H₂SO₄ (pH 0) under solar irradiation (AM 1.5G, 100 mW cm⁻²). For the first and second scanned linear sweep voltammetry (LSV) curves, the current density (mA cm⁻²) vs. applied potential (V vs. RHE) is plotted in Fig. 4a. A quick photocurrent response was observed during light irradiation in 2 s on/off cycles (first scan). The second scan was performed under the light-on condition. The onset potential (=open circuit voltage, V_OC) was 0.25, 0.22, 0.24, and 0.15 V, respectively, for Si–MoS₂, Si–MoSe₂, Si–WS₂, and Si–WSe₂, which was approximately 0.2 V for all Si-TMD samples. The photocurrent at 0 V vs. RHE (J_SC, short circuit current) was 30, 26, 23, and 22 mA cm⁻², respectively, for Si–MoS₂, Si–MoSe₂, Si–WS₂, and Si–WSe₂, and 0 for pristine (bare) Si NWs. The maximum current reaches 32 mA cm⁻². The PEC performance of the four TMDs is much enhanced compared to the bare Si NWs and is summarized in Table 1.

Fig. 4 (a) Current density (mA cm⁻²) vs. potential (V vs. RHE) for Si, Si–MoS₂, Si–MoSe₂, Si–WS₂, and Si–WSe₂ photocathodes, measured in 0.5 M H₂SO₄ (pH 0) under AM 1.5G irradiation (100 mW cm⁻²), with the light-on/off cycles (first scan) and light-on (second) scan. The scan rate is 20 mV s⁻¹. E⁰ represents the 0 V vs. RHE. (b) CA response of photocurrent and (c) H₂/O₂ evolution vs. time (min). Applied potential of 0 V vs. RHE. The line represents the calculated H₂/O₂ evolution from current generation with FE = 100%, and the circle symbol for the data.

Table 1 Summary of the PEC performance of Si–MoS₂, Si–MoSe₂, Si–WS₂, and Si–WSe₂

Samples	S or Se vacancies	PEC				R ct	E fb
		J SC	V OC	Degradation for 3 h^c	FEH2^d
a Short circuit current: current density (mA cm^−2) at 0 V vs. RHE. b Open circuit voltage (V vs. RHE). c Degradation of current density at 0 V vs. RHE for the 3 h CA test. d Average faradaic efficiency for 3 h H2 gas generation. e Charge transfer resistance (Ω) determined using Nyquist plots at 0 V (vs. RHE) under light irradiation. f Flat band potential (V vs. RHE) determined using Mott–Schottky plots.
Si–MoS2	7%	30	0.25	0%	90%	172	0.17
Si–MoSe2	5%	26	0.22	3%	88%	112	0.23
Si–WS2	5%	23	0.24	0%	86%	109	0.25
Si–WSe2	10%	22	0.18	8%	85%	203	0.20

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Fig. 4b shows the chronoamperometric (CA) response of the photocurrent (at 0 V vs. RHE) with a degradation of 3% over 3 h. In order to confirm water splitting in the PEC cell at pH 0, we monitored the evolution of hydrogen (H₂) and oxygen (O₂) gases at 0 V (vs. RHE). Fig. 4c displays the H₂ and O₂ evolution data (in mmol). The corresponding faradaic efficiencies (FEs) for 2H⁺ + 2e⁻ → H₂(g) and 2H₂O → O₂(g) + 4H⁺ + 4e⁻ were calculated using the equations and where N_H2 and N_O2 are the amounts (mol) of H₂ and O₂, respectively, and Q is the total amount of generated charge in Coulombs (photocurrent × time). The average FE was 90% for H₂ and 85% for O₂ (the line in the figure represents the calculated H₂/O₂ evolution from the photocurrents assuming FE = 100%). Table S1 (ESI†) compares the PEC performance with previously reported data for Si-TMD PEC cells, showing that the J_SC value of Si–MoS₂ is close to that obtained by the Shen group using the MoS₂ layer on the pyramid shaped Al₂O₃/n⁺p-Si photocathode in 1 M HClO₄ (34 mA cm⁻²).³²

Fig. 5 HER performance of freestanding MoS₂, MoSe₂, WS₂, and WSe₂ nanosheets. (a) LSV curves (scan rate: 2 mV s⁻¹) of the HER in H₂-saturated 0.5 M H₂SO₄ (pH 0) using TMD nanosheets (with bulk and oxide samples) and commercial Pt/C as catalysts. (b) Tafel plots derived from the LSV curves at low potentials that correspond to the activation-controlled region, using the Tafel equation η = b log(J/J₀), where η is the overpotential, b is the Tafel slope (mV dec⁻¹), J is the measured current density (mA cm⁻²), and J₀ is the exchange current density. The b values are given in parentheses. (c) TOF (s⁻¹) curves and (d) comparison of η_J=10, b, and TOF values at 0.2 V. (e) CA responses at η_J=10 for 12 h.

We examined the structure and morphology of the samples after PEC H₂ evolution for 3 h. The XRD patterns showed that all the crystals retained their crystal phases (Fig. S10, ESI†). The XPS and EDX data indicate that the electronic structures also remained the same afterwards. Electrochemical impedance spectroscopy (EIS) measurements provided the charge transfer resistance (R_ct) under light irradiation, as shown in Fig. S11 (ESI†). The R_ct values of Si–MoSe₂ (112 Ω) and Si–WS₂ (109 Ω) are smaller than those of Si–MoS₂ (172 Ω) and Si–WSe₂ (203 Ω), which is consistent with the freestanding nanosheets. Using Mott–Schottky (MS) plots, the flat-band potential (E_fb) was estimated to be approximately 0.2 V (Fig. S12, ESI†). Overall, the PEC performance of Si-TMD is nearly the same because they all have 5–10% chalcogen vacancies, which are optimal according to the HER catalytic activity for freestanding nanosheets with 10–20% chalcogen vacancies.

The extent of band bending (E_b) that occurred in the surface-electrolyte interface can be determined by E − E_fb, where E is the applied potential. If E_fb is more positive, E_b becomes more negative at the cathodic potential E, which provides a driving force for electron transfer to the surface and benefits the catalytic HER at the active sites of TMDs. Because E_fb = 0.2 V for all considered Si-TMDs, a similar downward band bending would be expected. The metallicity of the TMDs due to the chalcogen vacancies increases the carrier concentration to induce band bending. The good catalytic activity of TMD nanosheets with chalcogen vacancies resulted in the efficient separation of the photoexcited electron–hole pairs. Therefore, their excellent HER catalytic activity gives rise to excellent PEC performance.

As a next step, the electrocatalytic HER performances of freestanding MoS₂, MoSe₂, WS₂, and WSe₂ nanosheets were investigated via LSV, using these samples as working electrodes in a typical three-electrode setup. Fig. 5a shows the LSV curves measured at pH 0, where the potentials are referenced to the reversible hydrogen electrode (RHE). The overpotential required to obtain a current density of 10 mA cm⁻² (η_J=10) is 185, 169, 162, and 202 mV for MoS₂, MoSe₂, WS₂, and WSe₂ nanosheets, respectively. For comparison, the LSV curves measured for commercially available bulk powders (purchased from Alfa Aesar) provided the following values: η_J=10 = 274, 261, 272, and 269 mV, respectively, for MoS₂ (purity 99%), MoSe₂ (99%), WS₂ (99.8%), and WSe₂ (99.8%). As shown in the earlier XPS and EPR data (see Fig. S7 and S8, ESI†), these bulk powders have no S/Se vacancies. MoO₃ and WO₃ also showed much lower catalytic activity (η_J=10 = 326 and 336 mV, respectively). A commercial 20 wt% Pt/C catalyst exhibited η_J=10 = 27 mV.

Fig. 5b displays the Tafel plots, η(V) vs. log[J (mA cm⁻²)], in the low potential region. The slope of the linear fit provides the Tafel slope (b). The b values are 81, 80, 70, and 84 mV dec⁻¹, respectively, for MoS₂, MoSe₂, WS₂, and WSe₂ nanosheets. Respective bulk powders show b = 91, 106, 125, and 109 mV dec⁻¹. In comparison, Pt/C, MoO₃, and WO₃ showed b = 30, 64, and 69 mV dec⁻¹, respectively. The catalytic activity can be compared using the turnover frequency (TOF) per active site. We quantified the active sites using cyclic voltammetry (CV) curves at pH 7 (phosphate buffer solution) in the region of −0.2 to 0.6 V vs. RHE (see “Experimental details” in the ESI†). Fig. 5c shows the TOF curves derived from LSV data by normalizing the active sites in the samples. The TOF values at η = 0.2 V are 2.3, 4.4, 5.2, and 1.8 s⁻¹ for MoS₂, MoSe₂, WS₂, and WSe₂, respectively. Fig. 5d compares the values of η_J=10, b, and TOF for the four samples, showing that MoSe₂ and WS₂ display better HER performance than MoS₂ and WSe₂.

Electrochemical impedance spectroscopy (EIS) measurements provided the charge transfer resistance (R_ct), as shown in Fig. S13 (ESI†). Its values at −0.2 V (vs. RHE) are 41.0, 20.0, 9.1, and 66.4 Ω for MoS₂, MoSe₂, WS₂, and WSe₂, respectively, in the reverse order of their HER performance. Since a smaller R_ct value implies more facile electron transfer kinetics that enhances the catalytic activity, the R_ct values are well correlated with the relative HER performance. As shown in Fig. S14 (ESI†), the double-layer capacitance (C_dl) was measured using CV scans in the non-faradaic region (0.1–0.2 V vs. RHE) to be 10.7, 11.6, 14.7, and 8.2 mF cm⁻² for MoS₂, MoSe₂, WS₂, and WSe₂, respectively. Thus, the sample with higher catalytic activity consistently exhibits a larger charge capacitance value. Fig. 5e shows the CA responses of the samples at η_J=10 for 12 h. There was no current attenuation for MoS₂ after 12 h, indicating that it had the best stability among the samples. In contrast, WSe₂ showed a 12% current decrease, but it was only 3% until 3 h. In the cases of MoSe₂ and WS₂, no degradation occurred until 6 h and the decrease was 3% after 12 h.

Table 2 summarizes the measured HER performance for MoS₂, MoSe₂, WS₂, and WSe₂ nanosheets and bulk powders. We highlight the following results: (i) The HER catalytic activity of the bulk powders is much lower than that of the nanosheets. The difference between them is negligible compared to the nanosheets. (ii) MoSe₂ and WS₂ nanosheets exhibit higher HER performance than MoS₂ and WSe₂. (iii) MoS₂ has the best stability. (iv) WSe₂ possesses the worst catalytic capability, especially in terms of stability. These trends in HER performance can be simply explained using the following model. The data of bulk powders suggest that the intrinsic catalytic activity of four TMDs may not be much different. Therefore, the enhanced HER performance of nanosheets can be correlated with their S/Se vacancies. Based on previous studies, the S/Se vacancies enhance the HER performance by producing active sites that favorably form the Mo–H or W–H bond.^40–53 When the vacancy concentration exceeds the optimum value, the activity of active sites usually decreases.^40,45,50 Therefore, the higher HER performance of MoSe₂ and WS₂ compared to MoS₂ and WSe₂ is due to the optimum S/Se vacancies (10% vs. 15% and 20%). The worst performance of WSe₂ can be similarly ascribed to the rich Se vacancies. The double chalcogen vacancies at a higher concentration can act as a defective trapping site for the adsorbed protons. Nevertheless, the dependence of the HER performance on the concentration of chalcogen vacancies over the range of 10–20% is not significant. It is also possible that the optimal HER performance occurs at below 10%, which was not feasible to examine in this study. On the other hand, the highest stability of MoS₂ may originate from its inherent nature, which needs further investigation.

Table 2 Electrocatalytic HER performance of freestanding MoS₂, MoSe₂, WS₂, and WSe₂ nanosheets and bulk powders

Nanosheets									Bulk powders
TMD	[X]/[M]^a	S or Se vacancies^b	η J=10	b	TOF^e	R ct	C dl	CA^h	η J=10	b
a M = Mo or W, X = S, Se. b Determined using EDX and XPS data. c Overpotential (mV vs. RHE) at J = 10 mA cm^−2. d Tafel slope (mV dec^−1). e TOF (s^−1) at η = 0.2 V. f Charge transfer resistance (Ω) obtained using the Nyquist plot of EIS at η = 0.2 V. g Double layer capacitance (mF cm^−2). h Degradation of current during the CA test for 12 h.
MoS2	1.7	15%	185	81	2.3	41.0	10.7	0%	274	91
MoSe2	1.8	10%	169	80	4.4	20.0	11.6	3%	261	106
WS2	1.8	10%	162	76	5.2	9.1	14.7	3%	272	125
WSe2	1.6	20%	202	84	1.8	66.4	8.2	12%	269	109

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4. Conclusions

We synthesized MoS₂, MoSe₂, WS₂, and WSe₂ nanosheets with chalcogen vacancies as excellent catalysts for solar-driven water splitting PEC cells and the electrocatalytic HER. Using the shared step of gas-phase sulfurization/selenization under H₂ flow, these TMDs were prepared in two forms: freestanding nanosheets and shell layers on a pre-grown Si NW array. This process produced 5–20% chalcogen vacancies. High-resolution STEM images revealed that double chalcogen vacancies were formed at higher concentrations while single vacancies formed at lower concentrations. The XPS analysis showed that a higher chalcogen vacancy concentration induces a larger redshift of the binding energy. The Si-TMD NW array with optimal chalcogen vacancies (5–10%) exhibited excellent PEC performance in 0.5 M H₂SO₄; the photocurrent is 22–30 mA cm⁻² (at 0 V vs. RHE) and the onset potential is 0.18–0.25 V. In terms of stability, nearly 100% of the HER performance is retained after 3 h, with a high FE of H₂ generation (90%). The enhanced catalytic activity due to the chalcogen vacancies strongly boosts the performance of PEC cells. Meanwhile, we evaluated the electrocatalytic activity of the freestanding nanosheets toward the HER, showing that the catalytic performance in 0.5 M H₂SO₄ was optimized when there were 10% chalcogen vacancies. For the four TMDs having 10–20% chalcogen vacancies, η_J=10 = 162–202 mV (vs. RHE) and b = 76–84 mV dec⁻¹. These results demonstrated that the introduction of S/Se vacancies in TMDs can improve the (photo)electrocatalytic activity toward the HER.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This study was supported by Korea University. The HVEM measurements were supported by the KBSI R&D program (Project No. C030440).

Notes and references

1. A. A. Fujishima and K. Honda, Nature, 1972, 238, 37–38.
2. R. N. Dominey, N. S. Lewis, J. M. Bruce, D. C. Bookbinder and M. S. Wrighton, J. Am. Chem. Soc., 1982, 104, 467–482.
3. A. J. Bard and M. A. Fox, Acc. Chem. Res., 1995, 28, 141–145.
4. O. Khaselev and J. A. Turner, Science, 1998, 280, 425–427.
5. R. Asahi, T. Morikawa, T. Ohwaki, K. Aoki and Y. Taga, Science, 2001, 293, 269–271.
6. M. Grätzel, Nature, 2001, 414, 338–344.
7. M. G. Walter, E. L. Warren, J. R. McKone, S. W. Boettcher, Q. Mi, E. A. Santori and N. S. Lewis, Chem. Rev., 2010, 110, 6446–6473.
8. X. Chen, S. Shen, L. Guo and S. S. Mao, Chem. Rev., 2010, 110, 6503–6570.
9. A. Paracchino, V. Laporte, K. Sivula, M. Grätzel and E. Thimsen, Nat. Mater., 2011, 10, 456–461.
10. D. Bae, B. Seger, P. C. K. Vesborg, O. Hansen and I. Chorkendorff, Chem. Soc. Rev., 2017, 46, 1933–1954.
11. Y. Hou, B. L. Abrams, P. C. K. Vesborg, M. E. Björketun, K. Herbst, L. Bech, A. M. Setti, C. D. Damsgaard, T. Pedersen, O. Hansen, J. Rossmeisl, S. Dahl, J. K. Nørskov and I. Chorkendorff, Nat. Mater., 2011, 10, 434–438.
12. Y. W. Chen, J. D. Prange, S. Dühnen, Y. Park, M. Gunji, C. E. D. Chidsey and P. C. McIntyre, Nat. Mater., 2011, 10, 539–544.
13. S. Y. Reece, J. A. Hamel, K. Sung, T. D. Jarvi, A. J. Esswein, J. J. H. Pijpers and D. G. Nocera, Science, 2011, 334, 645–648.
14. M. J. Kenney, M. Gong, Y. Li, J. Z. Wu, J. Feng, M. Lanza and H. Dai, Science, 2013, 342, 836–840.
15. K. Sun, S. Shen, Y. Liang, P. E. Burrows, S. S. Mao and D. Wang, Chem. Rev., 2014, 114, 8662–8719.
16. L. Ji, M. D. McDaniel, S. Wang, A. B. Posadas, X. Li, H. Huang, J. C. Lee, A. A. Demkov, A. J. Bard, J. G. Ekerdt and E. T. Yu, Nat. Nanotechnol., 2015, 10, 84–90.
17. Q. Ding, J. Zhai, M. Cabán-Acevedo, M. J. Shearer, L. Li, H.-C. Chang, M.-L. Tsai, D. Ma, X. Zhang, R. J. Hamers, J.-H. He and S. Jin, Adv. Mater., 2015, 27, 6511–6518.
18. J. C. Hill, A. T. Landers and J. A. Switzer, Nat. Nanotechnol., 2015, 14, 1150–1155.
19. H. Zhang, Q. Ding, D. He, H. Liu, W. Liu, Z. Li, B. Yang, X. Zhang, L. Lei and S. Jin, Energy Environ. Sci., 2016, 9, 3113–3119.
20. D. Liu, J. Ma, R. Long, C. Gao and Y. Xiong, Nano Today, 2017, 17, 96–116.
21. Z. Luo, T. Wang and J. Gong, Chem. Soc. Rev., 2019, 48, 2158–2181.
22. T. F. Jaramillo, K. P. Jørgensen, J. Bonde, J. H. Nielsen, S. Horch and I. Chorkendorff, Science, 2007, 317, 100–102.
23. J. Kibsgaard, Z. Chen, B. N. Reinecke and T. F. Jaramillo, Nat. Mater., 2012, 11, 963–969.
24. D. Voiry, H. Yamaguchi, J. Li, R. Silva, D. C. B. Alves, T. Fujita, M. Chen, T. Asefa, V. B. Shenoy, G. Eda and M. Chhowalla, Nat. Mater., 2013, 12, 850–855.
25. P. D. Tran, S. S. Pramana, V. S. Kale, M. Nguyen, S. Y. Chiam, S. K. Batabyal, L. H. Wong, J. Barber and J. Loo, Chem. – Eur. J., 2012, 18, 13994–13999.
26. Q. Ding, F. Meng, C. R. English, M. C. Acevedo, M. J. Shearer, D. Liang, A. S. Daniel, R. J. Hamers and S. Jin, J. Am. Chem. Soc., 2014, 136, 8504–8507.
27. L. Zhang, C. Liu, A. B. Wong, J. Resasco and P. Yang, Nano Res., 2015, 8, 281–287.
28. K. C. Kwon, S. Choi, K. Hong, C. W. Moon, Y.-S. Shim, D. H. Kim, T. Kim, W. Sohn, J.-M. Jeon, C.-H. Lee, K. T. Nam, S. Han, S. Y. Kim and H. W. Jang, Energy Environ. Sci., 2016, 9, 2240–2248.
29. C.-J. Chen, K.-C. Yang, C.-W. Liu, Y.-R. Lu, C.-L. Dong, D.-H. Wei, S.-F. Hu and R.-S. Liu, Nano Energy, 2017, 32, 422–432.
30. Y. Hou, Z. Zhu, Y. Xu, F. Guo, J. Zhang and X. Wang, J. Hydrogen Energy, 2017, 42, 2832–2838.
31. S. Oh, J. B. Kim, J. T. Song, J. Oh and S.-H. Kim, J. Mater. Chem. A, 2017, 5, 3304–3310.
32. R. Fan, J. Mao, Z. Yin, J. Jie, W. Dong, L. Fang, F. Zheng and M. Shen, ACS Appl. Mater. Interfaces, 2017, 9, 6123–6129.
33. L. A. King, T. R. Hellstern, J. Park, R. Sinclair and T. F. Jaramillo, ACS Appl. Mater. Interfaces, 2017, 9, 36792–36798.
34. G. Huang, J. Mao, R. Fan, Z. Yin, X. Wu, J. Jie, Z. Kang and M. Shen, Appl. Phys. Lett., 2018, 112, 013902.
35. R. Fan, G. Huang, Y. Wang, Z. Mi and M. Shen, Appl. Catal., B, 2018, 237, 158–165.
36. M. Alqahtani, S. Sathasivam, F. Cui, L. Steier, X. Xia, C. Blackman, E. Kim, H. Shin, M. Benamara, Y. I. Mazur, G. J. Salamo, I. P. Parkin, H. Liua and J. Wu, J. Mater. Chem. A, 2019, 7, 8550–8558.
37. S. Seo, S. Kim, H. Choi, J. Lee, H. Yoon, G. Piao, J. C. Park, Y. Jung, J. Song, S. Y. Jeong, H. Park and S. Lee, Adv. Sci., 2019, 6, 1900301.
38. A. Hasani, Q. V. Le, M. Tekalgne, M.-J. Choi, T. H. Lee, S. H. Ahn, H. W. Jang and S. Y. Kim, ACS Appl. Mater. Interfaces, 2019, 11, 29910–29916.
39. R. Fan, J. Zhou, W. Xun, S. Cheng, S. Vanka, T. Cai, S. Ju, Z. Mi and M. Shen, Nano Energy, 2020, 71, 104631.
40. H. Li, C. Tsai, A. L. Koh, L. Cai, A. W. Contryman, A. H. Fragapane, J. Zhao, H. S. Han, H. C. Manoharan, F. Abild-Pedersen, J. K. Nørskov and X. Zheng, Nat. Mater., 2016, 15, 48–54.
41. G. Ye, Y. Gong, J. Lin, B. Li, Y. He, S. T. Pantelides, W. Zhou, R. Vajtai and P. M. Ajayan, Nano Lett., 2016, 16, 1097–1103.
42. H. Li, M. Du, M. J. Mleczko, A. L. Koh, Y. Nishi, E. Pop, A. J. Bard and X. Zheng, J. Am. Chem. Soc., 2016, 138, 5123–5129.
43. Y. Yin, J. Han, Y. Zhang, X. Zhang, P. Xu, Q. Yuan, L. Samad, X. Wang, Y. Wang, Z. Zhang, P. Zhang, X. Cao, B. Song and S. Jin, J. Am. Chem. Soc., 2016, 138, 7965–7972.
44. L. Lin, N. Miao, Y. Wen, S. Zhang, P. Ghosez, Z. Sun and D. A. Allwood, ACS Nano, 2016, 10, 8929–8937.
45. Y. Sun, X. Zhang, B. Mao and M. Cao, Chem. Commun., 2016, 52, 14266–14269.
46. G. Li, D. Zhang, Q. Qiao, Y. Yu, D. Peterson, A. Zafar, R. Kumar, S. Curtarolo, F. Hunte, S. Shannon, Y. Zhu, W. Yang and L. Cao, J. Am. Chem. Soc., 2016, 138, 16632–16638.
47. D. Gao, B. Xia, Y. Wang, W. Xiao, P. Xi, D. Xue and J. Ding, Small, 2018, 14, 1704150.
48. L.-B. Huang, L. Zhao, Y. Zhang, Y.-Y. Chen, Q.-H. Zhang, H. Luo, X. Zhang, T. Tang, L. Gu and J.-S. Hu, Adv. Energy Mater., 2018, 8, 1800734.
49. S. Park, J. Park, H. Abroshan, L. Zhang, J. K. Kim, J. Zhang, J. Guo, S. Siahrostami and X. Zheng, ACS Energy Lett., 2018, 3, 2685–2693.
50. Q. Zhu, W. Chen, H. Cheng, Z. Lu and H. Pan, ChemCatChem, 2019, 11, 2667–2675.
51. L. Li, Z. Qin, L. Ries, S. Hong, T. Michel, J. Yang, C. Salameh, M. Bechelany, P. Miele, D. Kaplan, M. Chhowalla and D. Voiry, ACS Nano, 2019, 13, 6824–6834.
52. J. Lee, C. Kim, K. S. Choi, J. Seo, Y. Choi, W. Choi, Y.-M. Kim, H. Y. Jeong, J. H. Lee, G. Kim and H. Park, Nano Energy, 2019, 63, 103846.
53. C. Hu, Z. Jiang, W. Zhou, M. Guo, T. Yu, X. Luo and C. Yuan, J. Phys. Chem. Lett., 2019, 10, 4763–4768.
54. S. Lee, S. Cha, Y. Myung, K. Park, I. H. Kwak, I. S. Kwon, J. Seo, S. A. Lim, E. H. Cha and J. Park, ACS Appl. Mater. Interfaces, 2018, 10, 33198–33204.
55. J. Hong, Z. Hu, M. Probert, K. Li, D. Lv, X. Yang, L. Gu, N. Mao, Q. Feng, L. Xie, J. Zhang, D. Wu, Z. Zhang, C. Jin, W. Ji, X. Zhang, J. Yuan and Z. Zhang, Nat. Commun., 2015, 6, 6293.
56. Y. Guo, D. Liu and J. Robertson, Appl. Phys. Lett., 2015, 106, 173106.
57. J. Lee, S. Kang, K. Yim, K. Y. Kim, H. W. Jang, Y. Kang and S. Han, J. Phys. Chem. Lett., 2018, 9, 2049–2055.

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Footnotes

† Electronic supplementary information (ESI) available: Additional data. See DOI: 10.1039/d0tc04715e

‡ I. H. K. and I. S. K equally contributed.

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