Interaction of Native Defects with Ions and Its Role in Inducing Phase

Jun 21, 2018 - In this Article, we show that defect-ion interaction is one of the key ..... in the lower frequency region of the Nyquist plot, that ac...
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Interaction of Native Defects with Ions and Its Role in Inducing Phase Transitions in p‑Type S‑Excess MoS2 Ajinkya Puntambekar,† Naveen Chandrasekeran,‡ Qi Wang,† Indroneil Roy,† Viji Premkumar,‡ and Vidhya Chakrapani*,†,$ †

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Howard P. Isermann Department of Chemical and Biological Engineering and $Department of Physics, Applied Physics & Astronomy, Rensselaer Polytechnic Institute, Troy, New York 12180, United States ‡ CSIR-Central Electrochemical Research Institute, Karaikudi, Tamil Nadu India 630006 S Supporting Information *

ABSTRACT: The effects of ion intercalation in transition metal chalcogenides like MoS2 has been well studied, although the nature of this interaction is not clearly known. In this Article, we show that defect-ion interaction is one of the key parameters that control many of the electrical, optical, structural, and electrocatalytic properties of MoS2. The results show for the first time that modulation of the concentration of intrinsic defects in MoS2 containing an excess of ‘S’ atoms in the lattice through Li+ insertion can lead to a new type of semiconductor-to-insulator-to-metal electronic phase transition with a concomitant change in the electrical conductivity from p-type to n-type, and a reversible 1T → 2H → 1T structural phase transformation. Using near-infrared photoluminescence and X-ray photoelectron spectroscopy measurements to directly monitor defect-ion interactions, it is shown that the observed changes are a direct result of changes in the electronic structure resulting from passivation of S-excess defects by Li+ and subsequently from the formation of electrochemically induced Sdeficient vacancy defects in MoS2. The effects of these broad range modulation on the catalytic rates of oxygen and hydrogen evolution reactions are shown. The structure−property−activity correlation shown here has important implications for chalcogenides-based semiconductors in general. KEYWORDS: molybdenum disulfide, ion insertion, defects, phase transitions, photoluminescence, electrocatalysis

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INTRODUCTION Electrochemical ion insertion is a facile technique for modulating many optoelectronic properties of transition metal dichalcogenides.1−3 For instance, Li+ insertion into MoS2 has been shown to cause structural phase transitions from 2H semiconducting phase to metastable 1T metallic phase. Ion insertion has also shown to increase both the visible and the infrared optical transmittance of the MoS2.4 Catalytic activity of electrochemical reactions, such as hydrogen evolution reaction (HER), have shown to markedly increase with Li+ insertion, which is believed to be due to the formation of metastable 1T phase during electrochemical charging.5−7 These results have suggested that the metallic 1T phase is catalytically the most active phase of MoS2, and attempts to stabilize this phase with chemical doping such as carbon, have been successful.8,9 While the merits of electrochemical control of electronic properties is clear, the exact nature of the ion interaction with the lattice ions and its role in affecting electrical and catalytic properties is neither clearly known nor has been investigated in detail. Intrinsic point defects, such as vacancies, edge sites, are known to markedly affect the electronic structure and catalytic properties of chalcogenides.3,10 However, the nature of interaction of intercalating © XXXX American Chemical Society

ions with the intrinsic structural defects during electrochemical tuning and in-turn its role in affecting macroscopic properties such as catalytic activity or structural 2H → 1T transition is not known. In this work, we show that a broad-range of reversible modulation of structural, electronic, electrical, and optical properties of MoS2 is possible through careful modulation of interaction of excess ‘S’ defects with Li+ in ptype MoS2. We show that progressive Li+ insertion into nonstoichiometric, S-excess MoS2 leads to a dramatic semiconductor-to-insulator-to-metal electronic phase transition with a concomitant change in the electrical conductivity from p-type to n-type, a reversible 1T → 2H → 1T structural phase transformation at room temperature, and a “volcano” type change in the optical absorbance. We show, through nearinfrared (NIR) photoluminescence (PL) and X-ray photoelectron spectroscopy (XPS) measurements, that the observed changes are a direct result of changes in the electronic structure resulting from defect-ion interactions, which consists of the passivation of ‘S’ excess defects by Li+ and subsequently from Received: February 13, 2018 Accepted: June 7, 2018

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DOI: 10.1021/acsaem.8b00215 ACS Appl. Energy Mater. XXXX, XXX, XXX−XXX

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ACS Applied Energy Materials

Figure 1. (A) Scanning electron micrograph of as-prepared MoS2 powder containing a mixture of 2H and 1T phases. (B−D) X-ray photoemission spectrum of core level of S 2p (B), Mo 3d (C), and valence band region of untreated, nonstoichiometric MoS2. The twin peaks of the S 2p envelope with BE separation of 1.16 eV is because of the S signal from S2− lattice ions of MoS2. However, the presence of broad shoulder peak at the higher BE indicates in the presence of excess S in the lower oxidation state possibly in the form of S0 or S2−, which is also evident in the Mo 3d envelope (C) that shows the presence of Mo in Mo4+ and small concentration of Mo6+ state. The presence of excess S also manifests a broad shoulder peak in the S band of the VB spectrum, as indicated in the Figure 1D. The presence of excess S renders the MoS2 p-type, which is evident from the small value of EF−EV estimated by the x-axis intercept of the rising portion of the VB spectrum. The summary of the fitting analysis of the core-level peaks and their peak energy positions, areas, and fwhm’s is given in Table S1.

electrochemically induced formation of ‘S’ deficient vacancy defects in MoS2. The effects of these broad range modulations on the catalytic rate of HER and oxygen evolution reaction (OER) are shown. A structure−property−activity correlation is proposed that has important implications for chalcogenidebased heterogeneous catalysts and field-effect devices. Lastly, the results presented here demonstrate that type of intrinsic defects and nature of intercalating ions are important parameters for enabling a broad range of structural and electronic phase transitions. MoS2 powder was purchased from Alfa-Aesar. Raman spectrum (shown later in the text) showed that the powder consisted of a mixture of 2H semiconducting and 1T metallic phase of MoS2. The powder consisted of mixture of large hexagonal crystallites of ∼20 μm of 2H phase and smaller nanoparticles in the size range of 40−100 nm of predominantly 1T phase, as seen from the scanning electron microscopy (SEM) image shown in Figure 1A. X-ray diffraction pattern, shown in Figure S1, conforms to the reference spectrum of 2H phase of MoS2. Results of X-ray photoelectron spectroscopy (XPS) suggest that as-purchased MoS2 contains excess sulfur in the lattice. Figure 1B and 1C shows the core level XPS spectra of S 2p and Mo 3d envelopes. The S 2p envelope shows twin peaks at the binding energies (BE) of 162.2 and 163.4 eV that are characteristic of the spin−orbit splitting of peaks into S 2p3/2 and S 2p1/2, with an energy difference of 1.16 eV, and an area intensity ratio of 2:1, respectively. These peaks arise from the sulfur present as S2− atoms in MoS2. In addition to the two main peaks, each of the S 2p3/2 and S 2p1/2

envelope also contain a broad peak at a higher BE of 162.7 and 163.9 eV, respectively, from the main peaks, which is characteristic of ‘S’ species present at the higher oxidation state such as S0 or S2−. By fitting each of the S 2p peaks with two Gaussian functions and estimating the integrated peak area, the composition of as-purchased sample was found to be MoS2.25. A small peak was observed near 169 eV corresponds to the BE of sulfur in a sulfate, SO42− state (shown in Figure S2). The exact nature of excess S in the bulk MoS2 used in the present study is not known, though both the core-level and valence band (VB) spectrum show some resemblance to the spectrum of amorphous MoS3, which contain sulfur in a combination of S22− and S2− groups, as reported in the literature.11−13 The presence of some excess S in MoS2 leads to the increase in the oxidation state of surrounding Mo atoms from +4 to +6, which gives rise to changes in the Mo 3d emission line. Figure 1C shows the Mo 3d envelope, which consists of Mo 3d5/2 and 3d3/2 at the BEs of 229.4 and 232.6 eV, respectively, from Mo4+ atoms, along with two additional peaks at BEs of 226.6 and 235.7 eV, respectively, from Mo6+ atoms surrounding excess S sites. It is also noted here that Mo 3d and S 2p core-level peaks have previously been used to identify 1T and 2H phases of MoS2. However, changes in the stoichiometry of the sample also affects the core-level spectrum that generally leads to broadening of the Mo 3d and S 2p peaks. The relationship between stoichiometry and type of structural phases (1T or 2H) is not clearly known. The presence of excess S in the sample may be the reason for stabilization of the metastable 1T phase. Therefore, the coreB

DOI: 10.1021/acsaem.8b00215 ACS Appl. Energy Mater. XXXX, XXX, XXX−XXX

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Figure 2. (A) Nyquist plot of real and imaginary components of electrochemical impedance of MoS2 at various charging potentials in 0.1 M LiClO4 in propylene carbonate at an AC amplitude of 5 mV. The inset shows the equivalent circuit used for modeling the measured electrochemical response. (B) Changes in the electrode resistance, which is plotted as the sum of charge transfer resistance (RCT) and grain boundary resistance (RGB), during various stages of electrochemical charging. The data shows a semiconductor-to-insulator-to-metal transition during electrochemical reduction, (C and D) Mott−Schottky plots of MoS2 electrode at OCP (C) and 1.0 V (D) in an electrolyte solution consisting of 0.1 M TBAP in propylene carbonate showing p-type and n-type behavior in the untreated and heavily Li+ intercalated state, respectively.

level spectra was fitted to account for the different compositional states rather than different structural phases. The valence band (VB) electronic structure of MoS2 has been extensively studied both experimentally and theoretically.14−16 The VB structure derived from photoemission studies is generally in good agreement with the theoretical calculations. The dominant features of VB spectrum are the observation of three regions that correspond to the different bonding nature: (i) dz2 band at the BE of ∼2 eV which originates mainly from the nonbonding 4d electrons of Mo; (ii) the d-p band in the BE range of 3−9 eV arising from the bonding states of hybridized Mo 4d and S 3p electrons; (iii) the S band at ∼13.5 eV from the nonbonding S 3s electrons.11,12 The VB spectrum of as-synthesized nonstoichiometric MoS2 in the energy range of 0−20 eV is shown in Figure 1C. Comparison of this spectrum with that reported for stoichiometric MoS2 shows similar bonding characteristics in the energy range between 0 and 10 eV. However, the major difference is the presence of an additional shoulder peak in the higher BE of S band, indicating the presence of two different types of S in the nonstoichiometric sample. Similar VB structure has been observed for amorphous MoS3.12 Interestingly, the VB spectrum of both MoS3 reported in the literature, as well as the MoS2.25, used in the present study indicates p-type electronic conduction. Density functional theory (DFT) calculations of Zhou et al.17 have predicted that nonstoichiometric, monolayer MoS2 with excess S related defects lead to the generation of midgap electronic states whose energy lies close to the VB maximum, while nonstoichiometric, monolayer MoS2 with S-deficient vacancy defects have electronic states closer to the conduction band (CB) minimum, which leads to n-type conductivity. From the results of scanning tunneling microscopy (STM) on natural bulk MoS2 crystals, Addou et al.18 concluded that

excess-S leads to generation of structural defects and strain that gives rise to p-type conductivity, while S-deficiency leads to ntype conductivity. The presence of S-excess related defect states within the band gap was confirmed by NIR PL measurements, wherein a defect related PL peak at the energy of ∼0.78 eV can be seen that is not present in stoichiometric MoS2, which is discussed in detail later in the text. We note here that the p-type conductivity could be caused by the presence of metallic impurities, such as Nb, Zr, Y, that are present below the limit of XPS detection, and therefore not detected. However, the observation of dominant S peak of higher oxidation state, suggests that excess ‘S’ may be the dominant p-type dopant in this sample. Reversible electrochemical insertion of Li+ into MoS2 was performed in an Ar-filled glovebox using a three-electrode setup with the application of various constant potentials and is described in detail in the Methods section. In a typical experiment, a step voltage was applied for a duration of time needed for the equilibrium to be established, which was indicated when the charging current dropped to a low value in the range of 10−7 Amperes. Changes in the electrical resistivity of MoS2 was measured in situ, while changes in the optical and structural characteristics were recorded ex situ using Raman, absorption and photoluminescence spectroscopy. The results are summarized in Figures 2−6. Changes in the electrical resistivity, as well as the sign of the majority charge carriers, at various stages of the intercalation were determined by electrochemical impedance spectroscopy. Experiments were performed by charging the electrode at a given DC potential until equilibrium was established. The total impedance of the cell was then measured in the frequency range of 10 Hz to 50 kHz at an AC amplitude of 5 mV superimposed on the DC charging potential. To determine the C

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observed before. The resistance of the untreated film increases with increase in Li+ concentration and reaches a maximum at the charging voltage of 1.7 V versus Li/Li+. This is followed by a decrease in the resistance to lower values at a higher Li+ concentration in the potential range of 1.7 to 1 V versus Li/Li+. This change in the electrical resistance is accompanied by the switch in the polarity of majority charge carriers from p-type to n-type with Li+ insertion and occurs at a charging voltage of 1.7 V versus Li/Li+. The n-type conductivity of the film is evident in the M-S plot shown in Figure 2D that was obtained at charging voltage of 1.0 V. The measured shift in the flat band potential of ∼1.4 V from 1.33 V at open circuit potential (OCP) to −0.19 at 1.0 V is consistent with the movement of the Fermi level, EF, near VB in the p-type MoS2 to near CB in the n-type MoS2, which roughly corresponds to the bulk band gap (EG = 1.3 eV) of MoS2. Ex situ measurements of Hall voltage of the heavily lithiated sample emersed from the electrolyte confirmed the n-type nature of the film. Note that the resistance of MoS2 in the insulating phase was ∼500 ohms, which is lower than what would be expected for an insulating material. Several factors may be responsible for the decreased resistance. Because of the high surface-area and nanoporous nature of the electrode, not all regions of the electrode is uniformly intercalated with Li+. As a result, despite the formation of an insulating phase, channels with high conductivity persists in undoped regions (regions not lithiated) that cause the overall resistance to be lower than the film that is uniformly doped. It was also observed that sample in electrolyte had an order of magnitude lower resistance than in air. The observed parabolic trend in the resistance change with Li+ insertion is in contrast to the previous reports that showed only a decrease in the electrical resistance of MoS2 with insertion.4 But the stoichiometry of the starting material was not reported. Our results suggest that if the untreated MoS2 is S-excess to begin with, then electrochemical insertion of Li+ and e− (from the back contact) leads to passivation of acceptor-type states of excess S, which leads to a metal-toinsulator transition in the voltage regime between OCP and 1.7 V. However, insertion of Li+ at a concentration greater than the passivation limit results in n-type conductivity and an insulator-to-metal transition. Thus, high Li+ insertion must give rise to a new type of donor-type defects closer to the CB of MoS2. This suggests that at smaller reduction potentials, Li+ initially binds to excess ‘S’ species such as S0 or S2− to form Li−S bonds; a process that is similar to the formation of Li−O bonds during Li+ insertion in transition metal oxide.21−23 This would lead to passivation of acceptor-type S excess defects states and increase in electrical resistivity. It is known that S vacancy in MoS2 leads to n-type electrical conductivity.17 This is because removal of sulfur as neutral atom will result in the electron transfer to neighboring Mo atoms, which will result in the change in the oxidation state from +4 to +3 (polaron formation). In the electrochemical reduction process, insertion of e− and Li+ at a concentration above the passivation limit results in binding of Li+ to the lattice S2− sites. This would result in the formation of localized Li−S bonds (as LiS or Li2S) and a change in the oxidation state of surrounding Mo atoms from +4 to +3; a case similar to the formation of S vacancy. Thus, electrochemically induced S-vacancy formation also leads to the increase in the n-type conductivity and formation of donor-type (such as Mo3+) defect states. To confirm the above hypothesis, the interaction of vacancy defects with Li+ was directly monitored using NIR PL

nature of majority charge carriers, Mott−Schottky (M-S) measurement was performed on the electrode after charging in the Li+ electrolyte by transferring the electrode to a different cell with an electrolyte consisting of 0.1 M tertiary butyl ammonium perchlorate (TBAP) in propylene carbonate (PC) and silver metal foil serving as the reference and platinum as the counter electrode. The capacitance of the electrode was determined at a fixed AC frequency of 500 Hz as a function of the scan voltage and the flat band potential was estimated through the M-S analysis. Figure 2C shows the M-S plot of the measured capacitance of as-synthesized sample, which indicates that nonstoichiometric MoS2 with excess S is a ptype semiconductor, as indicated by the negative slope of the C2− versus V curve. This is consistent with the p-type nature of the sample deduced from the VB XPS curve shown in Figure 1D. The flat band potential was estimated to be 1.13 V versus Ag/AgCl or 1.33 V versus the standard hydrogen electrode (SHE). Thus, the value of Fermi level of ‘S’-excess MoS2 was calculated to be −5.77 eV with respect to the vacuum energy scale, which is close to the value of −5.3 eV reported previously for p-type natural crystal of MoS2 through kelvin probe measurements in air.14 The p-type nature of the material was also independently confirmed by Hall effect measurements of the MoS2 film in air prior to electrolyte exposure in a Van der Pauw geometry. Measurements were done by pressing four Ag clips to the four corners of the square MoS2 film deposited on a glass substrate. Untreated samples showed a small positive Hall voltage of 4.4 mV, which indicates that the samples were p-type. The nanostructured nature of the film prevented the estimation of the charge carrier concentration. The electron mobility of the film was calculated to be 5.7 cm2 V−1 s−1. Electrical resistance of the MoS2 film at various charging potentials of Li+ insertion was extracted from the measured total impedance by fitting it to the equivalent circuit proposed by Passerini et al.19 and Ho et al.20 with slight modification to account for the nanostructure nature of the electrode. As a typical example, Figure 2A illustrates the Nyquist plots of the measured real and imaginary components of the total impedance of MoS2 in the untreated and Li+ intercalated states at various electrochemical potentials. The data were fitted with an equivalent circuit, which is shown in the inset of Figure 2A, which consisted of two combinations of resistive and capacitive elements along with a third resistor in series with these elements. The various electrode parameters corresponding to these elements are (i) MoS2/electrolyte interface with charge transfer resistance and bulk electronic resistance, RCT, and Warburg impedance, W, in parallel with the double layer capacitance, Cdl; (ii) grain boundary resistance, RGB, in parallel with the grain boundary capacitance, CGB, to account for the nanostructure nature of the electrode; (iii) solution resistance, RS; (iv) a finite diffusion element, FD, which manifests as a straight line in the lower frequency region of the Nyquist plot, that accounts for the finite thickness effect. The reversible evolution of various resistances, RCT, RGB, and RS, with progressive Li+ insertion, is shown in Figure S3. Changes in the electrical resistance of the MoS2 film, as represented by the sum of RCT and RGB, as a function of charging potential is shown in Figure 2B. All changes in the resistance are reversible within the potential window reported here. As can be seen, the resistance curve shows a “volcano”type behavior that indicates the occurrence of a semiconductor-to-insulator-to-metal electronic phase transition in S-excess MoS2 during Li+ insertion, which has not been D

DOI: 10.1021/acsaem.8b00215 ACS Appl. Energy Mater. XXXX, XXX, XXX−XXX

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Figure 3. (A) Ex-situ photoluminescence spectra of nonstoichiometric MoS2 in the energy range of 0.7−1.6 eV measured after lithiation at various electrochemical potential. The spectra shows two emission peaks related to the band edge (Eg) emission at 1.3 eV and a defect related peak emission at ∼0.75 eV. (B) Changes in the intensity of Eg emission as a function of charging voltage. The emission intensity shows a maximum at potential of 1.7 V when the material behaves as an insulator. (C) Changes in the intensity of NIR emission measured at 0.8 eV occurring as a result of Li+ insertion into the MoS2 lattice. The NIR emission seen in the untreated MoS2 quenches upon Li+ insertion at 1.7 V, whereas subsequent Li+ insertion at potential less than 1.7 V (versus Li/Li+) leads to an increase in the NIR emission, which is due to the formation of new defect-related emission peak that occurs at nearly similar energy range as the peak seen in untreated sample. (D) VB spectrum of MoS2 before (black) and after (red) Li+ insertion at a charging voltage of 2.0 V. The change in the conductivity from p-type to n-type is evident from the shift in the x-axis intercept, which gives the value of EF − EV.

measurements. Our prior work on transition metal oxides22−24 and CdSe25 have shown that most point defects, which give rise to electronic midgap states closer to the band edges can be probed by NIR PL spectroscopy through their characteristic defect-related luminescence. For instance, Figure 3A shows the PL emission from untreated S-excess MoS2 in the spectral range of 0.7 to 1.6 eV. MoS2 is an indirect band gap semiconductor with a band gap of 1.3 eV. However, at room temperature, the band gap emission is weak, as seen by the broad weak peak at ∼1.3 eV in the PL spectrum. The emission at room temperature is dominated by a strong sub-band gap emission that has a peak intensity at ∼0.75 eV. This emission arises most probably from the electronic transition between excess-S related acceptor states and the VB of MoS2. Changes in the PL emission with Li+ insertion was measured ex situ at various charging potentials on emersed electrodes. After it was charged at each potential, the sample was washed with propylene carbonate (PC) and dried under vacuum. The PL spectrum was recorded immediately to avoid bleaching in air. The changes in the spectral intensity measured at different voltages is shown in Figure 3A. The complex changes in the trend are however best observed from Figure 3B and 3C, which shows the changes in the intensity of band gap, EG, emission (measured at 1.3 eV) and the defect-related emission (measured at 0.8 eV) as a function of charging potential, respectively. As can be seen in Figure 3B, the intensity of EG emission increase with an increase in Li+ concentration until the charging voltage reaches 1.7 V, which is also the potential at which MoS2 shows the maximum value of resistance (Figure

2B). Below 1.7 V, the intensity of EG emission quenches completely. In contrast, the intensity of defect-related emission seen in untreated samples decreases with increasing Li+ concentration in the sample and quenches completely at 1.7 V, which is consistent with our hypothesis that Li+ initially binds specifically to the excess S defect sites, which could be edge sites. However, with the passivation of all excess S sites (S0 or S2− sites) by Li+ at 1.7 V, the material behaves as a stoichiometric insulator with relatively high resistance. This also leads to an intense EG emission at room temperature, which is due to the suppression of nonradiative recombination occurring at defect sites, a process that otherwise dominates during PL measurements at room temperature. These results show that optical emission from indirect band gap semiconductor is strongly affected by the density of defects states, which in turn is controlled by the stoichiometry of the material. Interestingly, a new defect-related emission that has a peak at nearly the same energy as the emission peak seen in the untreated sample is also seen in MoS2 at higher concentration of Li+. The intensity of this peak progressively increases with increasing concentration of Li+ in the sample even down to the potential of 0 V, where irreversible conversion to Li2S is known to occur.26 Given that the conductivity of the sample switches to n-type at potential less than 1.7 V, the defect-related emission must be a result of electronic transitions between newly formed donor-type defects that have a density of states closer to the CB edge. Comparison of VB XPS spectrum of MoS2 before and after Li+ insertion at a charging voltage of 2.0 V, as shown in Figure 3D, also confirms the change in the E

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ACS Applied Energy Materials conductivity from p-type to n-type. This is evident from the shift in the x-axis intercept, which is obtained from the extrapolation of the straight line fitted to the rising portion of VB curve, that gives the value of EF − EV from ∼0.1 to ∼1.3 eV after Li+ insertion. As suggested above, these states probably arise due to the formation of S vacancy defects as a result of binding of Li+ at the lattice S2− sites. Figure 4 shows the band

Figure 4. Schematic of energy band diagrams of (A) nonstoichiometric MoS2 containing excess S-related defects that gives rise to midgap electronic states closer to the valence band (VB) and consequently a p-type conductivity. (B) Li+-passivated insulating MoS2 (LixMoS2) at a potential of 1.7 V. (C) Heavily Li+ intercalated n-type MoS2 (LiMoS2−x) at a potential below 1.7 V with the midgap energy states related to S-vacancy defects that gives rise to n-type conductivity. Also shown are the various radiative transitions resulting from the formation and passivation of defects.

Figure 5. (A) Optical absorbance of MoS2 powder measured ex situ in the reflectance mode after Li+ insertion at various charging potential with respect to Li/Li+ reference electrode. (B) Changes in the value of absorbance measured at the wavelengths of 650 and 1500 nm as a function of charging potential.

diagram of MoS2 and LixMoS2 corresponding to the untreated p-type semiconducting, insulating and n-type semiconducting phases formed during various stages of intercalation along with the major electronic transitions during optical excitations. Also shown are the positions of EF and band edges determined from the M-S measurements. Similar to the changes in the PL, the passivation and formation of vacancy defect states as a result of ion intercalation also strongly influences adsorption/transmission properties of MoS2. The results are shown in Figure 5. The effect of Li+ insertion on the optical transmittance of MoS2 sheets have been studied before. For instance, Xiong et al.4 showed that Li+ insertion leads to an increase in the optical transmittance in both visible and NIR spectral regions from OCP down to the charging potential of 1 V. In this work, we observe two regions of opposite trend. Ex situ UV−vis−NIR absorption measurements were done in the diffuse reflectance mode by dispersing a known amount of MoS2 powder scrapped from the electrode post electrochemical charging onto a BaSO4 pellet. Measurements were done immediately after charging to minimize bleaching in air. Figure 5 shows that the optical absorbance of MoS2 powder after various stages of electrochemical charging in the ultraviolet to NIR spectral range. There are two major features of each of the spectra: (i) three excitonic peaks at 340, 600, and 650 nm arising as a result of band edge transition in the K point and (ii) a defectrelated absorption tail that can be seen at the longer wavelength between 700 and 1500 nm. Figure 5B shows the

changes in the measured absorbance at spectral wavelength of 650 and 1500 nm as a function of the charging potential. Between OCP and 2 V, the optical absorbance corresponding to the excitonic peak as well as the defect-related absorbance decreases with increasing Li+ concentration. This trend is similar to the observation of Xiong et al.4 mentioned earlier. However, further insertion of Li+ at higher reducing voltages between 2 and 1 V leads to a strong increase in both the visible and NIR absorbance. The observation that the absorbance changes parabolically with charging potential is similar to the trend seen with defect-related emission intensity. This suggests that the same defects are responsible for both electronic transitions. Thus, we conclude that the decrease in the absorbance with Li+ insertion is due to the passivation of Sexcess defects, which enhances optical transmittance throughout the UV-to-NIR spectral range. Moreover, the absorbance increase seen at higher Li+ insertion is probably is a due to the optical excitations from midgap electronic states formed as a result of electrochemical S-vacancy defects. Importantly, these results show the importance of the correlation between stoichiometry and optical and electronic properties of chalcogenide semiconductors. While the trend of defect-related absorbance changes with Li+ is similar to that seen with defectrelated emission, the changes in excitonic absorbance measured at 650 nm show a different trend compared to the band edge emission. The indirect band gap of MoS2 is 1.3 eV. The three peaks at 340, 600, and 650 nm arise as a result of direct excitonic transition in the K point. The underlying F

DOI: 10.1021/acsaem.8b00215 ACS Appl. Energy Mater. XXXX, XXX, XXX−XXX

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at room temperature through chemical doping with ‘C’,9 substitutional electron doping of Re, Tc, and Mn atoms,8 and through electron beam exposure,32 the stabilization of 1T phase seen in undoped, nonstoichiometric MoS2 with excess sulfur seen in the present study has not been reported before. It may be that the presence of excess S as S2− or S0 helps in stabilizing 1T phase. This assertion is born from the observation that passivation of this excess S through Li+ insertion induces a conversion of 1T to 2H phase. This is seen from the complete disappearance of the Raman peaks at 146.8, 196.4, 235.5, and 335.9 cm−1 corresponding to the vibrational stretches of 1T phase (indicated by *) when untreated MoS2 is charged between OCP and 1.7 V. This is the first observation of reversible 1T to 2H phase transition at room temperature achieved via ion insertion. However, charging at higher reduction potentials that results in the formation of electrochemical S-vacancy defects leads to a reversible conversion of 2H to 1T phase transformation. This result is similar to prior reports on the transformation to 1T MoS2 at potential less than 1.0 V versus Li/Li+ by various researchers. Thus, the results of Raman measurements together with the NIR PL results shown in Figure 3C show that the reversible 1T → 2H → 1T phase transformation that occurs as a result of ion insertion crucially depends on the nature of defect-ion interactions. While the 1T → 2H transformation is driven by the passivation of S-excess sites by Li+, the 2H → 1T transformation is induced by the formation of electrochemical S-vacancy defects. Changes in the structural phase transition have been studied using XPS. Our attempts to obtain a reliable ex-situ XPS spectrum from lithiated MoS2 was hampered by the formation of Li(OH)3 on the surface as a result of reaction of Li+ with ambient air during brief exposure doing sample transfer. Prior results by Lin et al.32 showed that exposure of MoS2 to an electron beam and/or high temperature annealing can induce the nucleation of metastable 1T or 1T′ phases within the 2H semiconducting phase. Our results here show that electrochemically generated vacancy-defects created as a result of ion insertion can induce similar phase transitions at room temperature. Thus, electrochemical modulation of defect equilibria can enable precise control of electronic and structural properties of transition metal chalcogenides. Finally, we studied the correlation between structural and electronic properties, and electrocatalytic properties. MoS2 is an important electrocatalyst for a number of electrochemical reactions, such as HER, CO2 reduction.13,34 It is well-known that the edge sites of MoS2 are the active sites for HER reaction, while the basal sites are electrochemically inactive.35,36 However, recent studies have shown that S vacancy37 and/or strain can also lead to an improvement in the HER kinetics.10,38 Vacancy defects in both transition metal oxides and sulfides have long been known to play an important role in catalyzing heterogeneous reactions by acting as “hot spots” for adsorptive binding. Both theoretical and experimental results of Li et al.10 showed that the electrocatalytic activity of HER shows a volcano-type dependence on the concentration of Svacancy defects created in the basal plane of MoS2. They reported that the highest activity of HER occurs at the Svacancy concentration of 12.5%. Their results suggest that the presence of vacancy defects enhances activity by affecting the binding energy of S−H bond. Thus, concentration of Svacancy defects greater that 12.5% leads to much stronger binding of H+ such that it lowers the rate of desorption of H2 from active sites. Here, we studied the catalytic effects of the

reason for the increase in the absorbance of excitonic transitions is not immediately clear. Detailed band structure calculations have to be performed to understand how various transitions are affected by Li+ insertion. Insertion of alkali metal ions, such as Li+, Na+, K+, into 2H MoS2 at room temperature is well-known to induce reversible phase transformation from 2H phase to metastable 1T phase.1,2,6,7,27 From the results of electrocatalytic HER studies on various dichalcogenides surfaces of MoS2, MoSe2, WS2, and WSe2, Ambrosi et al.7 concluded that the efficiency of the 2H →1T transition that occurs upon Li+ insertion determines the efficiency of the HER. Their results showed that WS2 and MoSe2 are the best material for HER. We studied the effect of modulation of stoichiometry through lithiation and defect formation on the structural phase transformation between the 2H and 1T phases of MoS2. Ex-situ Raman spectroscopic measurements were performed on MoS2 electrode at various electrochemical potentials and the results are shown in Figure 6. Raman spectrum of untreated, S-excess MoS2 showed the

Figure 6. Changes in the ex-situ Raman spectrum of MoS2 electrode during various potentials of electrochemical charging. The characteristic in-plane (E1g, E12g) and the out-of-plane (A1g) vibrational modes of the 2H phase are indicated by (+), which occurs at 280.9, 375.6, and 402.8 cm−1, respectively. The characteristic vibrations of 1T phase are indicated by (*). Data suggest that the nonstoichiometric Sexcess MoS2 contains mixture of 1T and 2H phases, which upon low Li+ intercalation (between OCP to 2 V) undergoes a 1T → 2H phase transformation. Upon Li+ insertion at higher reducing potential (below 2 V), the material undergoes a 2H → 1T phase transformation.

characteristic in-plane (E1g, E12g) and the out-of-plane (A1g) vibrational modes of the 2H phase at 280.9, 375.6, and 402.8 cm−1, respectively.28,29 In addition, the material also shows the characteristic vibrations of 1T phase at 146.8, 196.4, 235.5, and 335.9 cm−1, which indicates that the starting material contains a mixture of 1T and 2H phases of MoS2. The 1T MoS2 is metastable at room temperature. While the transition from the 2H to the stable 1T phase is known to occur at high temperature, the metastable 1T phase obtained though Li+ insertion at room temperature is known to revert back to the 2H phase upon bleaching in the air or upon mild laser annealing.30,31 Lin et al.32 showed the conversion between the 2H to 1T phase occurs via the gliding of intralayer atomic plane that involves a transversal displacement of one of the S planes, which is induced strain. Wang et al.33 suggested that strain caused by volumetric expansion during Na+ insertion is responsible for the nucleation of metastable 1T phase. While there have been recent reports on the stabilization of 1T phase G

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Figure 7. (A) Cyclic voltammogram of p-type, untreated S-excess MoS2 electrode in the potential range of −0.5 to 1.7 V obtained at a scan rate of 5 mV/s in 0.5 M H2SO4 aqueous solution. The peak potentials of various reactions are indicated. (B) Comparison of cyclic voltammograms of MoS2 before after Li+ insertion at 1.7 V. Insertion of Li+ into MoS2 leads to an enhancement in both the HER and OER currents. (C) Tafel plots for OER on MoS2 intercalated with Li+ at various charging voltages. (D) Tafel plots of S-excess MoS2 showing the effect of Li+ insertion on the HER kinetics. A strong reduction in the Tafel slope, expressed in millivolts per decade (mV/dec), is seen for both HER and OER at all Li+ intercalation voltages.

HOR currents. This result is striking because prior studies on MoS2, which are dominantly stoichiometric or S-deficient, have studied only for the HER activity. Our results show that Sexcess MoS2 can serve as excellent bifunctional catalysts and that stoichiometry of the material is another crucial knob for modulating the electrocatalytic activity of transition metal chalcogenides. Figure 7B shows the comparison of the voltammograms of MoS2 before and after Li+ insertion at 1.7 V of charging potential. A clear enhancement in both OER and HER currents can be observed with Li+ insertion. Tafel polarization curves recorded for OER and HER on untreated MoS2 and on MoS2 after electrochemical Li+ insertion at various charging potentials of OCP, 1.7 and 1.0 V, which corresponds to the p-type semiconducting, insulating, and ntype semiconducting phases of MoS2, are shown in Figure 7C and 7D. As can be seen, Li+ insertion drastically improves both the HER and OER kinetics with the marked reduction in the Tafel slopes at all charging potentials. The above results together with the results of PL, Raman, electrochemical Mott−Schottky, and electrical conductivity experiments during Li+ charging provide important insights on the reasons for electrocatalytic enhancement. First, the polarization results together with in situ electrical measurements show that electrocatalytic activity does not always directly scale with electrical conductivity. Otherwise, one would have expected to observe lowest HER and OER currents at 1.7 V of Li+ charging voltage when the electrical resistance of MoS2 is maximum (Figure 2B). Second, the enhancement in the HER current is also not directly related to the metallic 1T phase formation, as commonly believed, while it may play a secondary role in HER. According to the Raman results shown in Figure 6, there is an interconversion of 1T to 2H phase of untreated S-excess MoS2 between OCP and 1.7 V as a result of

changes in the electronic properties resulting from electrochemical modulation of defect equilibria. Prior works have shown that Li+ insertion in transition metal dichalcogenides, such as MoS2, MoSe2, WSe2, and WS2, leads to a marked increase in the current density of HER. Wang et al.5,6 observed increases in the HER current density at all concentration of Li+ i.e. at all charging voltages. The enhancement in HER current was attributed to the formation of metallic 1T phase of low electrical resistivity. However, in addition to electrical resistivity changes, alkali ion insertion leads to other marked changes in the electronic and structural properties, such as vacancy formation, passivation of midgap defects states, all of which can affect catalytic activities. Thus, the question of whether Li+ truly catalyzes HER reaction kinetics by lowering the activation barrier for charge transfer, such as that observed from S-vacancy formation, or simply leads to better electronic conductivity for charge transport is not resolved. In the present work, we evaluated the effect of modulation of stoichiometry achieved through Li+ intercalation on the catalytic activity of water dissociation reactions consisting of oxygen evolution reaction (OER), oxygen reduction reaction (ORR), hydrogen oxidation reaction (HOR), and HER. Electrodes for catalytic studies were prepared by drop casting a mixture of MoS2 powder in dimethylformamide (DMF) on a fluorine-doped tin oxide (FTO) glass slide. All electrocatalytic testing was done in air-saturated 0.5 M H2SO4 aqueous solution at room temperature. Figure 7A shows the cyclic voltammogram in the potential range of −0.5 to 1.7 V of untreated, S-excess MoS2 electrode obtained at a scan rate of 5 mV/s. The peak currents of various reactions are indicated in the figure. Also shown is the voltammogram recorded on bare FTO electrode. As can be seen, p-type S-excess MoS2 exhibits excellent OER and HER kinetics but relatively poor ORR or H

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ACS Applied Energy Materials Li+ insertion. Yet, the electrocatalytic activity of both OER and HER is higher for MoS2 charged with Li+ at 1.7 V compared to the values recorded on untreated sample. Third, the catalytic activity also does not correlate with the density of vacancy defects in MoS2. Untreated MoS2 used in this study contained excess sulfur that most likely was present at the edge sites. The presence of excess S results in midgap defect states closer to the VB edge. These S-excess sites can act as binding sites for both H+ and Li+, which explains the high HER and OER activity seen in untreated samples. Li+ insertion between charging voltage of OCP and 1.7 V leads to the quenching of these catalytic defects sites, as seen from the PL results shown in Figure 3A and 3C, by forming Li−S bonds. Yet, the electrocatalytic activity as a result of H+ insertion during HER and OER is higher at 1.7 V than the value recorded at OCP. These arguments lead to the conclusion that there must be another reason for electrocatalytic enhancement as a result of Li+ insertion. Our hypothesis is that current enhancement is a result of an increase in the net electrocatalytic area as a result of the volumetric expansion caused by the insertion of Li+ into the MoS2 lattice. However, the enhancement in the HER and OER current seen with MoS2 intercalated with Li+ at 1.0 charging voltage could be due to synergetic effects of electrochemical S-vacancy formation, lowering of the electrical resistivity with concomitant formation of the 1T phase, as well as increase in the surface area. Finally, we note that though the electrochemical properties were stable for the first two cycles, long-term stability of the electrode against bleaching of Li+ from the electrode surface has to be taken into account while considering ion insertion for modulating electronic and catalytic properties. Li+ will most likely deintercalate at positive potentials or oxidizing conditions. Thus, stability of the electrode is currently being evaluated. In conclusion, the results presented in this study have shown that defect-ion interactions in MoS2 is one of the key parameters that controls many of the electrical, optical, structural and electrocatalytic properties. Electrochemical modulation of defect equilibria in untreated MoS2 with excess of ‘S’ through Li+ insertion can lead to a semiconductor-toinsulator-to-metal electronic phase transition with a concomitant change in the electric conductivity from p-type to n-type, a reversible 1T → 2H → 1T structural phase transformation, and a volcano-type behavior of optical absorbance. Through NIR PL emission and X-ray photoelectron spectroscopy measurements we have shown that the observed changes are a result of changes in the electronic structure resulting from passivation of S excess defects by Li+ and subsequently from the formation of S deficient vacancy defects in MoS2. The effects of these broad range modulation on the catalytic rate of oxygen and hydrogen evolution reactions is shown. The structure−property−activity correlation shown here has important implications for chalcogenides-based semiconductors in general.

electrodes in a nonaqueous electrolyte consisting of 0.1 M LiClO4 in propylene carbonate. Electrochemical charging was done by applying a constant potential step using a CH Instrument 660E electrochemical workstation and allowing the electrode to reach equilibrium, which was until the charging current reached a minimum value. Mott−Schottky analysis was used to determine the flat band potentials and the type of majority charge carriers under various stages of Li + intercalation using a silver metal foil as the reference and platinum metal as the reference electrodes in 0.1 M TBAP in propylene carbonate. All electrical and electrochemical measurements except the electrocatalytic studies were carried out in an argon-filled glovebox. Optical Testing. Absorbance spectra of the MoS2 powder were recorded using a UV−vis−NIR spectrophotometer (Shimadzu UV-3600) in the range of 300−2600 nm. Photoluminescence and Raman measurements were performed using HORIBA Scientific LabRAM HR Evolution spectrometer in the energy range from 1.8 to 0.7 eV using an InGaAs CCD detector at an excitation wavelength of 633 nm. The laser power used for sample excitation was 6.7 mW. To avoid spot damage with laser-induced heating, scan acquisition was limited to 45 s. X-ray Studies. The composition and structure of MoS2 film before and after electrochemical charging were characterized by XPS (Physical Electronics PHI 5000 VersaProbe). XPS measurements were carried out at room temperature in a UHV system with a base pressure of 10−9 mbar. Al Kα radiation (hν = 1486.6 eV) from a monochromatic X-ray source was used for excitation. All spectra are normalized in energy to the C 1s peak. Survey scans were collected over the range from 1100 to 0 eV with a pass energy of 117.4 eV. Typical surveys were collected for 10 min. Higher-resolution scans were collected over a range of 20 eV around the peak of interest with 23.5 eV pass energy. Typical close-up scans were detected for 10−20 min per peak of interest.

METHOD Electrochemical Charging. MoS2 electrodes were prepared by sequential layer-by-layer deposition by drop casting a slurry of MoS2 powder in dimethylformamide (DMF) on a FTO-coated glass slide that was heated at 80 °C. This process resulted in a continuous, uniform porous film that was several microns in thickness. Lithium ion intercalation of MoS2 was performed in a standard three-electrode electrochemical cell configuration with Li metal as both the counter and reference

The authors declare no competing financial interest.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsaem.8b00215.

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XRD, resistance changes, and detailed XPS analysis of the data reported in the text (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Naveen Chandrasekeran: 0000-0002-7565-1065 Vidhya Chakrapani: 0000-0003-2682-3833

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Notes

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ACKNOWLEDGMENTS The authors would like to thank Prof. Gamini Sumanasekera of University of Louisville for the helpful discussions and Rensselaer Polytechnic Institute (RPI) for the financial support. A.P, Q.W, and I.R gratefully acknowledge the partial support of Howard P. Isermann fellowship provided by the Department of Chemical and Biological Engineering at RPI. I

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