Bridging the Gap Between Bulk and Nanostructured Photoelectrodes

May 2, 2013 - ... surface with Scotch tape to expose a flat and reflective (0001) face. ..... a constant number of corner sites (at the ends of the (1...
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Bridging the Gap Between Bulk and Nanostructured Photoelectrodes: The Impact of Surface States on the Electrocatalytic and Photoelectrochemical Properties of MoS2 Zhebo Chen, Arnold J. Forman, and Thomas F. Jaramillo* Department of Chemical Engineering, Stanford University, 381 North South Axis, Stanford, California 94305, United States S Supporting Information *

ABSTRACT: Molybdenite (MoS2) is a material that has been previously studied for its photoelectrochemical properties in bulk form and, more recently, for its electrocatalytic hydrogen evolution properties in various nanostructured forms. Herein, we aim to address the properties of bulk and nanostructured MoS2 in order to assess the future potential of harnessing nanostructured MoS2 as a more efficient photocathode for photoelectrochemical hydrogen production. Using an in-situ method to progressively create defects at the surface of bulk MoS2, we show that losses due to recombination of photogenerated charge carriers at surface states occur concurrently with an enhancement in catalytic activity for hydrogen evolution. We further explore how these surface defects affect the measured flat-band potential, and we discuss the consequences of the distinct electronic properties of bulk versus nanostructured MoS2. Our results for MoS2 show that differences in morphological length scales (bulk vs nanoscale) give rise to unique surface properties that can greatly impact material performance and that represent key differences that researchers can leverage in order to develop more efficient nanoscaled photoelectrodes.

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organic compounds.8−11 This work alluded to the possibility of achieving higher photovoltages with nanostructured MoS2, but earlier work by Kline et al. showed that samples of n-MoS2 with a large population of surface defects, as may be found on nanostructures, typically produced lower anodic photocurrents12 implying the presence of recombination centers which can also serve to decrease photovoltage. On the basis of earlier studies of n-WSe2, this reduced performance was attributed to high recombination at these surface defects which decreased quantum yield.13−15 The role of electronic surface states has since been evaluated extensively in semiconductor photoelectrochemistry,16,17 including studies of layered transition metal chalcogenides such as WSe2.18 The presence of electronic surface states can lead to effects such as Fermi level pinning and charge-carrier shunting, which we further discuss in this study of MoS2. Given that nanostructured MoS2, because of its geometry, often possesses a larger population of surface step and edge sites per geometric device area compared to its bulk counterpart (with the exception of spherical closed-shell fullerene structures19,20 and extended monolayers21,22 of MoS2), the question remains as to whether these sites, which are known to be catalytically active, can enhance the overall activity for MoS2 as a photocathode for hydrogen evolution by reducing kinetic overpotentials or whether they act as electronic

INTRODUCTION Molybdenite (MoS2) is a layered transition metal chalcogenide that been studied extensively over recent decades for its photoelectrochemical properties. Seminal studies originated from Tributsch and Bennett in 1977, who analyzed the dark and illuminated activity of both n-type and p-type MoS2 in various electrolytes and redox couples.1 More recently, Hinnemann et al. investigated nanostructured MoS2 for electrocatalytic hydrogen evolution by means of theory and experiment proposing that the edge sites of MoS2 could serve as catalytically active sites;2 this hypothesis was later confirmed experimentally in a combined electrochemistry and surface science study of MoS2 nanoclusters synthesized in ultrahigh vacuum by Jaramillo et al.3 The origin of this high catalytic activity was found to be due to the presence of undercoordinated sulfur atoms at the 101̅0 edge that possess a metallic character allowing for more effective adsorption of H+ for conversion into H2. Since this discovery, many of the recent efforts surrounding MoS2 have focused on its further development as an efficient catalyst for electrochemical hydrogen evolution.4−7 In the context of these developments regarding the high reactivity of metallic edge sites, we present a study of MoS2 which serves to elucidate the impact of edge sites on the photoelectrochemical properties of MoS2 with specific attention to hydrogen evolution and the challenges in developing nanostructured MoS2 as a photocathode. Previously, Wilcoxon and co-workers proposed the synthesis of nanoparticulate MoS2 which exhibited quantum confinement of its band gap leading to enhanced photooxidation of various © XXXX American Chemical Society

Received: November 18, 2012 Revised: March 30, 2013

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conductivities and photoactivities for the purpose of this study. The exact nature of the dopant atoms found intrinsically in MoS2 is not well understood, although a wide variety of elements including Bi, Si, Al, Cu, Ag, Fe, Mg, Mn, Na, and Pb have been detected as possible dopants.26,27,32 Physical characterization of our crystals was performed and indeed confirmed the MoS2 composition and crystal structure. Surface and bulk chemical composition was analyzed using X-ray photoelectron spectroscopy and Raman spectroscopy (Figures S2 and S3 of the Supporting Information), while X-ray diffraction confirmed the layered crystal structure and the high degree of orientation of well-cleaved (0001) surfaces (Figure S4 of the Supporting Information). We first present the photoactivity obtained for two samples of MoS2 with distinctly p-type or n-type conductivities. Figure 1

recombination centers that can quench photovoltage and that can degrade overall photoactivity. This is the question that we aim to address in this study.

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EXPERIMENTAL SECTION Natural molybdenite crystals were procured from multiple vendors (Ward’s Natural Science, SPI) and were prepared by cleaving the surface with Scotch tape to expose a flat and reflective (0001) face. The crystals were then assembled into a custom compression electrochemical cell with an O-ring seal. Studies were conducted in a quiescent cell (no rotation or stirring) with no sparging of the electrolyte of residual dissolved O2. Electrical contact to the sample was made using copper foil (3M) since it provides a suitable ohmic contact for MoS2 (Figure S1 of the Supporting Information).23 Photoelectrochemical measurements were conducted using a Bio-Logic VMP3 potentiostat with 0.5 M H2SO4 as the electrolyte, a Hg/Hg2SO4 reference electrode (Hach), and a graphite foil counter electrode (Alfa). Calibration of the reference electrode (−0.7 V vs the reversible hydrogen electrode (RHE)) was performed using platinum mesh working and counter electrodes in the same electrolyte sparged with 99.999% H2 (Praxair). Adequate sparging was confirmed when the difference in open-circuit potential between the working and the counter electrodes was less than a few millivolts. The point at which the j−V curve crossed zero current was taken as 0.0 V versus RHE. A xenon lamp (Newport-Oriel) provided concentrated broad band illumination with a power density of 610 mW/cm2 between 250 and 950 nm as measured by a spectroradiometer calibrated against the NIST F-420 standard. Integration of the solar AM 1.5 G spectrum (ASTM G-173-03) over the same wavelength range yields ∼71 mW/cm2, which reveals that the illumination intensity of our lamp is approximately 8−9 suns. This concentrated light intensity allows for larger photocurrents and photocorrosion currents for simulating accelerated durability testing and is also instructive for practical performance in a real commercial-scale system that employs planar photoelectrochemical (PEC) water splitting electrodes since the use of concentrated illumination improves overall efficiency and hydrogen production and minimizes levelized balance-ofsystem costs.24,25 The synthesis of core−shell MoO3−MoS2 nanowires has been described previously.4 Briefly, MoO3−x nanowires were grown on fluorine-doped tin oxide substrates (FTO, TEC 15, Hartford Glass) using hot-wire chemical vapor deposition followed by sulfidization in a 10%/90% H2S/H2 environment at 200 °C. The synthesis of mesoporous double-gyroid MoS2 has also been described previously7 and employed an evaporationinduced self-assembled silica template on FTO followed by electrodeposition of a molybdenum bronze, sulfidization in 10%/90% H2S/H2 at 200 °C, etching of the template in 2% HF, and resulfidization in 10%/90% H2S/H2 at 200 °C.

Figure 1. Periodically illuminated j−V curves of (a) p-type and (b) ntype MoS2 crystals in 0.5 M H2SO4 under 610 mW/cm2 of illumination from a Xe lamp. Scan rate was 20 mV/s with an illumination shutter frequency of 1 Hz, 50% duty cycle.

shows the j−V curves for the p-MoS2 and n-MoS2 samples under concentrated periodic illumination. In the case of pMoS2, significant photocathodic current (∼60 mA/cm2) was generated that coincided with the visible formation of bubbles at the surface of the semiconductor. These bubbles scattered the light and also decreased the interfacial area between the pMoS2 surface and the electrolyte resulting in significant noise at current densities >40 mA/cm2. The activity for this sample remained relatively stable over the course of several hours of testing. The current generated in the dark was small across this wide potential window even at extremely large overpotentials, which is indicative of a good Schottky electrode with few bulk defects that would otherwise serve as shunts for transport of

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RESULTS AND DISCUSSION Several samples of natural molybdenite crystals were procured and were examined for their photoelectrochemical properties. Previous studies of MoS2 employed both natural crystals1,26,27 and synthetic crystals formed primarily via chemical vapor transport.28−31 We utilized natural crystals because they were readily available and because they provided a sufficient range of B

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electrons through the p-type semiconductor into the electrolyte. A small gradual rise in the dark current occurs at reverse biases either because of charge-carrier shunting or because of Schottky barrier height lowering.33 The n-MoS2 sample (Figure 1b) generated anodic photocurrents of a similarly large magnitude (∼60 mA/cm2) as the ptype sample, but no bubbles rose from the n-MoS2 surface and the data remained noise-free. This indicates that the evolution of O2, the desired anodic reaction for photoelectrochemical water-splitting, did not proceed because current densities much greater than ∼1 mA/cm2 corresponding to the evolution of gaseous products typically result in visible bubble formation. Instead, the n-MoS2 sample suffered severe photocorrosion and rapidly lost activity as the surface became visibly pitted by eye. The observations for p-MoS2 and n-MoS2 are consistent with earlier studies by Fujishima et al., who used rotating ring disk electrode experiments to determine that the primary reaction products in aqueous electrolyte with no sacrificial reactants from n-MoS2 include SO42− and Mo6+ (the latter in the form of HMoO4− or MoO42−), while the primary reaction product from the cathodic activity of p-MoS2 consists of H2.27 This illustrates why MoS2 is more naturally suited for use as a photocathode rather than as a photoanode in aqueous environments unless protective coatings can be developed as has been shown for nSi.34 These results show that the photocurrent generated by MoS2 under illumination can be significant and that photogenerated charge carriers can be effectively extracted albeit requiring large reverse biases. At 8−9 suns of illumination intensity, the 60 mA/cm2 of photocurrent corresponds to an incident photonto-current conversion efficiency (IPCE) or external quantum efficiency (EQE) of 9−11%, which includes significant losses because of reflection from the shiny MoS2 surface. For MoS2 to become a viable material for photoelectrochemical energy storage, it must generate its photocurrent at a significant underpotential (forward bias) relative to the redox reaction of interest. In particular, p-MoS2 must exhibit a photocurrent onset potential (Eonset) positive of 0.0 V versus the reversible hydrogen electrode (RHE). One possible means to shift the pMoS2 onset to more anodic potentials is by improving surface catalysis such that photogenerated electrons are more efficiently transferred to protons in solution to form H2. We thus turned our attention toward investigating the photoactivity of two nanostructured forms of MoS2 both of which have previously demonstrated excellent catalytic activity for hydrogen evolution in the dark. The nanostructured samples consisted of core−shell nanowires of MoO3−MoS2 and a mesoporous double-gyroid MoS2.4,7 Neither form of nanostructured MoS2, however, was able to generate a significant amount of photocurrent as shown in Figure 2. Applying a large external bias in an effort to obtain greater photocurrents only served to increase electrolytic dark current, which will be discussed further below. Even under concentrated illumination, current densities reached only a few μA/cm2, 2−3 orders of magnitude lower than what we observed on bulk MoS2 samples. Earlier work by Kiesewetter et al. investigated the use of small MoS2 microcrystals as the absorber in a dye-sensitized solar cell.35 Unfortunately, few crystallites were found to be active, and overall photoactivity remained low. Both our results and those of Kiesewetter et al. are consistent with previous studies by Kline et al. which found lower photoactivity for highly structured MoS2.12 This naturally led to the question as to whether nanostructured MoS2 is

Figure 2. Periodically illuminated j−V curves of double-gyroid MoS2 and core−shell MoO3−MoS2 nanowires taken in the cathodic direction. The dark current has been subtracted to more clearly distinguish the photocurrent response.

intrinsically prohibited from generating significant photoactivity because of electronic surface states that arise from physical surface defects (e.g., edges, corners, steps). Given that bulk single crystals possess a surface consisting of large, atomically flat domains of the (0001) plane while highly nanostructured morphologies exhibit a surface with a plethora of undercoordinated edge sites as illustrated in Figure 4, we developed a strategy to address this question by bridging the gap between these two extremes. To do so, we started with a freshly cleaved, bulk single-crystal MoS2 and progressively etched the surface anisotropically so as to deliberately populate the surface with the kinds of sites encountered on the nanostructured surfaces. A multitude of methods can be employed to partially etch the MoS2 surface to generate surface defects. For example, McIntyre et al. showed that argon ion bombardment of MoS2 produced significant surface defects that resulted in Fermi level pinning.36 However, such ex-situ methods may lead to a convolution of effects that may arise from handling procedures such as damage during sample mounting or surface oxidation upon exposure to air. An in-situ method of progressively generating surface defects while simultaneously measuring current−voltage characteristics electrochemically enables more consistent and extensive analysis of the effects of edge-site formation on the ability of MoS2 to generate electrolytic hydrogen in the dark and photocatalytic hydrogen under illumination. By simultaneously studying the effects of photogenerated carrier recombination and catalytic activity, we aim to provide greater insight into the fundamental physical chemistry of the processes at play. For this particular study, we investigated a third bulk MoS2 sample which was determined to be predominantly n-type with donor concentrations of 1018−1019 cm−3 by Mott−Schottky analysis (see the Supporting Information). This sample was capable of generating significant dark current under reverse bias (>1.1 V vs RHE) likely because of low shunt resistance which produced an ohmic rather than a rectifying response. This effect likely arose from the high concentration of donors or from the presence of bulk defects which additionally served as bulk recombination centers resulting in decreased photoactivity compared to the samples in Figure 1. Additionally, these states likely served to increase hole conductivity enabling holes to be C

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effectively extracted through the rear contact and allowing the flow of cathodic photocurrent toward the front semiconductor−electrolyte contact under large forward bias (negative potentials, see Figure 3). This provided a unique

Figure 3. Periodically illuminated j−V curve of bulk MoS2 sample with both anodic and cathodic photocurrent at 20 mV/s with an illumination shutter frequency of 1 Hz, 50% duty cycle.

opportunity since employing our in-situ method to generate edge sites on this sample (described below) allowed us to track anodic and cathodic electrolytic currents in the dark as well as photoanodic and photocathodic photocurrents under illumination all as a function of an increase in the density of surface defects. While the sample studied here is n-type, the ability to observe the effect of surface states on the charge transfer of photogenerated electrons into solution at the semiconductor− electrolyte interface may provide some insight into the behavior of p-type samples as well which are more relevant for development of a H2 evolving photocathode. Additional studies that evaluate p-MoS2 directly would provide great value in this respect. As mentioned above, anodic photocurrents in MoS2 proceed toward oxidative corrosion evidenced by the formation of pits and edge sites in previous work. Both Kautek and Gerischer and Sakamaki et al. showed that the photocorrosion process of n-MoS2 is anisotropic and occurs preferentially along the direction,29,37,38 which is the same behavior observed earlier by Bahl et al. for the case of thermal oxidation of MoS2.39 Corrosion along the direction subsequently leads to the formation of new catalytically active (101̅0) edge sites as visualized in Figure 4. Our in-situ method for generating edge sites and for continuously evaluating the light and dark activities for cathodic hydrogen evolution and anodic corrosion consisted of continuous cyclic voltammograms between −0.8 and +1.7 V versus RHE under chopped, periodic illumination (Figure 3) over the course of over 11 h. The dark and light activities for this MoS2 sample are presented in Figure 5a and b, respectively, normalized to initial activity (plots of non-normalized photocurrent are shown in Figure S5 of the Supporting Information). We focus first on a description of the sample’s activity in the dark. As the sample is subjected to progressive anodic photocorrosion over the course of the experiment, Figure 5a shows that the cathodic dark current rises steadily over time. This is consistent with recent observations that edge sites provide catalytic centers for hydrogen evolution confirming the ability of this approach to generate surface defects that represent the highly catalytically

Figure 4. Top-down schematic of the (0001) basal plane of MoS2. Labeled are the catalytically active (101̅0) Mo edges, minimally active (1̅010) S edges, and preferential direction of anodic photocorrosion. The undercoordinated S atoms along the (101̅0) edge are responsible for efficiently binding H to drive the hydrogen evolution reaction. Schematic of anodic corrosion process for a single MoS2 (0001) plane (a) before and (b) after corrosion along the direction revealing additional catalytically active sites along the (101̅0) Mo edge.

active sites observed in nanostructured MoS2. In contrast, the anodic dark current remains steady throughout the course of the 11 h experiment. This occurs because the mechanism of the corrosion process proceeds in the direction parallel to the (1010̅ ) edge resulting in a constant number of corner sites (at the ends of the (101̅0) edges) that are oxidized as visualized in Figure 4. Regarding activity under illumination, the anodic photocurrent decreases markedly over time consistent with previous observations in which the anodic photocurrent of MoS2 decreased in the presence of edge sites which served as electronic recombination centers. More interestingly, however, the cathodic photocurrent did not experience the same level of degradation and retained nearly all of its initial activity throughout the experiment. Consideration must also be given to the photoconductive behavior of MoS2,26,40,41 which can contribute to the absolute majority carrier (electron) photocurrent. Although photoconductivity is a bulk property and should not significantly contribute to the observed change in photocurrent over time as surface states are introduced, it may mask the surface-state effects if it is a primary contributor to the apparent photocurrent. In such a scenario, one should expect that the apparent cathodic photocurrent should continuously rise with increasing bias, whereas the observed saturation in photocurrent implies a photon-to-electron limited process. D

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biases, ηEQE should be voltage independent and should become limited entirely by ηe−/h+ and ηrecombination. Consequently, one should expect that the saturated photocurrent decreases over time since the generation of surface states lowers ηrecombination. However, this was not observed for the cathodic photocurrent suggesting that there was no change in ηrecombination for electrons despite an increase in surface states. To further explore the complexity in the observed behavior, we present here a framework for analyzing the role of surface states as possible recombination centers and as catalytically active sites in photoelectrochemical systems. The rate of surface recombination is governed by the following equations for electrons (n) under forward bias and for holes (p) under reverse bias for a n-type semiconductor: dn dt

surface

dp dt

surface

= −cnNssΔn

= −cpNssΔp

where cn and cp are the electron and hole capture coefficients, Nss is the number of recombination centers at the surface, and Δn and Δp are the number of photogenerated electrons and holes. Because Nss at any given time is a constant and because Δn = Δp for a given illumination intensity at large biases (both the cathodic and the anodic photocurrents achieve 1.8 mA/cm2 initially as shown in Figure S5 of the Supporting Information), differences in the cn and cp govern the differences in recombination rate. The different values for the two capture coefficients could arise from a couple of different reasons. One hypothesis has to do with the catalytic nature of these surface states since they promote cathodic H2 evolution. As we observe an enhancement in electrocatalytic hydrogen evolution (in the dark) along the Mo (1010̅ ) edge, this means that electron transfer from the semiconductor into the electrolyte could be facilitated with more surface states present thereby decreasing the likelihood of recombination; this would result in a smaller cn than it would otherwise have been without a catalytic effect (i.e., ηcatalysis and ηrecombination are coupled) thereby offsetting the increase in Nss. Another possibility is that the energy levels of the surface states (Ess) might lie closer to the conduction band (EC) than to the valence band (EV). In this hypothesis, such surface states would serve as shallow traps for electrons (resulting in a low cn) and deep traps for holes (resulting in a high cp) as illustrated in Figure 6. Either explanation could be responsible for the observed stability in cathodic photocurrent throughout the experiment as more and more surface states are formed in situ. To determine whether the stability in cathodic photocurrent was due to a lack of recombination at potentially shallow traps for electrons or was due to an increase in electron-transfer rate to the electrolyte because of increased catalysis, we analyzed the cathodic dark current and photocurrent in a regime that was kinetically limited (−0.2 V vs RHE, where ηcatalysis < 1) for H2 evolution rather than absorption limited (−0.8 V vs RHE, where ηcatalysis ∼ 1) as was presented earlier. It is not straightforward to disentangle these two effects in the absorption-limited regime, as shown in Figure 5, because the cathodic photocurrent at −0.8 V versus RHE is not limited by the number of catalytically active sites but rather by the photon flux. This is not the case at −0.2 V versus RHE, however, near the onset of photocurrent. If the surface states were shallow

Figure 5. (a) Plot of cathodic dark current at −0.8 V vs RHE and anodic dark current at +1.7 V vs RHE. (b) Plot of cathodic dark current and photocurrent at −0.8 V vs RHE and anodic photocurrent at +1.7 V vs RHE.

Upon initial inspection, the stability of the cathodic photocurrent seems surprising since it does not mirror the decreases observed in the anodic photocurrent. Only the anodic photocurrent decreases with time, and if degradation of the electronic quality of the semiconductor is responsible for the decline in photoactivity, one may expect similar effects to be measured for both the anodic and the cathodic photocurrent. The cathodic dark current shows a different trend and exhibits a significant increase in performance over time, rising more than 30% above its initial value, suggesting that the loss in electronic quality in the semiconductor because of recombination of photogenerated carriers at newly created surface edge sites may be offset by enhanced catalysis for hydrogen evolution. However, this is not necessarily the case at large biases. The external quantum efficiency can be described as a combination of several contributing factors: ηEQE = ηe−/ h+ × ηtransport × ηinterface where ηe−/h+ is the efficiency of absorbing photons and generating electron−hole pairs, ηtransport is the efficiency of separating and transferring charges within the semiconductor bulk, and ηinterface is the ability to transfer charges from the semiconductor into solution. One can further describe ηinterface = ηcatalysis × ηrecombination, where ηcatalysis is the catalytic ability to drive a particular electrochemical reaction, and where ηrecombination accounts for losses due to surface-state recombination. The addition of greater bias enhances both ηtransport by providing a larger electric field and ηcatalysis by providing a larger overpotential. Once ηtransport and ηcatalysis approach ∼1 at large E

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specifically, the coupling of ηcatalysis and ηrecombination may explain the behavior observed in this study, although additional studies are required to further explore this possibility. This highlights the interplay among multiple effects arising from surface sites (effects which may be advantageous in one respect and detrimental in another) and complicates the ability to make simple and broad conclusions regarding the factors that contribute to material performance in photoelectrochemistry. The presence of surface states at multiple energy levels may further complicate the behavior of a nanostructured PEC device. In nanostructured MoS2, a high concentration of both Mo (1010̅ ) and S (10̅ 10) edges exists in tandem. Previous density functional theory calculations have shown that the hydrogen-binding energy at the Mo edge is ΔGH = 0.08 eV and at the S edge is ΔGH = 0.18 eV.6 The disparate binding energies may be indicative of a multitude of surface-state energetics leading to multiple charge carrier lifetimes that may further complicate recombination. Time-resolved optical spectroscopy of nanostructured MoS2 performed by Parsapour et al. revealed that, indeed, several electronic midgap states are likely involved.42,43 However, it may be possible to favorably tune the energy levels of these surface states and carrier lifetimes by the addition of promoter atoms like Ni or Co, which have been shown to change the electronic properties of the S edge.6,44 At this time, further investigation is necessary to evaluate the exact energetic nature of states responsible for catalytic hydrogen evolution versus oxidative corrosion. Typical high quality Schottky photodiode behavior under small forward bias (Figure 8a) and large reverse bias (Figure

Figure 6. Band diagrams illustrating the impact of surface-state energetics (Ess) on the recombination of charge carriers. A given surface state can represent (a) a shallow trap for electrons and (b) a deep trap for holes giving rise to a disparity in cn vs cp.

electron traps that did not play a significant role in recombination of photogenerated electrons, one would expect the cathodic photocurrent in the kinetically limited regime to increase with the creation of more surface states since they are catalytically active (Figure 7a). As explained above, no such

Figure 8. Band diagram of an n-type semiconductor Schottky junction (a) under small forward bias and (b and c) under reverse bias . In a and b, the barrier height (ΦB) and the lack of surface states prevent the flow of dark current producing rectifying behavior. In c, the presence of surface states at ESS serves as recombination centers (pathways depicted by arrows) as well as electronic shunts that allow majority carrier electrons to flow under reverse bias.

Figure 7. (a) j−V curves in the dark (dotted) and under illumination (solid) for hypothetical photocathodes with good surface catalysis as a result of more edge sites (blue) and with poor surface catalysis as a result of fewer edge sites (green). (b) Cathodic dark current and photocurrent at −0.2 V vs RHE plotted as % of initial current vs time.

8b) should exhibit rectifying behavior characterized by absorption limited photocurrent and low dark current because of the presence of the Schottky barrier at the semiconductor− electrolyte interface. However, semiconducting materials of low electronic quality can often produce significant dark current and decreased photocurrent at these potentials (Figure S7 of the Supporting Information).45 These effects often arise because of the presence of defect states such as grain boundaries or atomic vacancies which serve as electronic

enhancement would be expected in the photon absorption limited regime as presented in Figure 5. In Figure 7b, we show distinctly that the cathodic photocurrent at −0.2 V versus RHE remained steady and did not increase within the kinetically limited regime. This strongly suggests that, in fact, the recombination of photogenerated electrons must play a significant role in limiting overall photoactivity. More F

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Figure 9. (a) Periodically illuminated j−V curve of a bulk MoS2 sample which displays both photoanodic and photocathodic current zoomed-in to the region of the flat-band potential marked by the gray region. (b) Flat-band potential measured from photocurrent onset for the duration of the study. The top and bottom lines represent the onset measured on the cathodic and anodic sweeps, respectively, while the line in the middle represents the average value. (c) Band diagram of n-MoS2 electrode at negative potentials relative to EFB (accumulation), at 0.0 V vs EFB (flat-band), and at positive potentials relative to EFB (depletion).

length scale of nanostructured systems is often on the same order (or less) than the depletion layer width. Additionally, one must also consider the crystal phase of MoS 2 ; only the 2H and 3R polytypes of MoS2 are semiconducting, whereas the 1T phase is suggested to be metallic.46,47 Even though the 1T phase is a metastable phase, it has been found in nanostructured forms of MoS248−50 and may also serve to decrease shunt resistance. Increasing the conductivity of MoS2 (e.g., by doping)51 could serve as a route toward further enhancing its properties as an electrocatalyst, but the impact to overall photoactivity could be detrimental. To address this challenge of low shunt resistance in nanostructured MoS2, it will be important to develop selective passivation strategies that can inhibit charge-transport processes at the relevant surface states. Fujishima et al. showed that I− could suppress the anodic dissolution of n-MoS2 by serving as a hole acceptor.27 Kubiak et al. found that other hole acceptors, such as Cl− or Br−, could also serve to limit corrosion in aqueous solutions, but very high concentrations of these acceptors were needed to outcompete surface oxidation.52 These studies demonstrated that redox couples can be used to kinetically outcompete corrosion at anodic potentials although, ideally, the charge-transfer process for majority carrier holes in p-MoS2 would be halted entirely to prevent photocurrent degradation at all anodic potentials. Further development of surface passivation chemistries or charge-selective contacts will be necessary to reach this goal. If applied optimally, performance can be greatly enhanced. In one particular example, Bard et al. demonstrated that passivation of n-WSe2

shunts (increasing dark current) and recombination centers (degrading photocurrent). In contrast to bulk crystalline MoS2, the highly structured nature of nanoscaled MoS2 gives rise to a higher density of surface states which can form an electronic percolation network across the entire surface of the structure that extends to the rear electrical contact. While the enhanced electronic conductivity due to the surface states lowers resistive losses and allows for more facile catalytic turnover (a feature desirable for an electrocatalyst), these same surface states will limit performance for a photoabsorber by increasing recombination as well as dark current (Figure 8c). In the dark, nanostructured MoS2 produces high-current densities for hydrogen evolution at low cathodic overpotentials3 (−0.15 V vs RHE) and rapid corrosion current6 at anodic potentials (>0.6 V vs RHE). Because the observed potential for the onset of anodic corrosion in nanostructured MoS2 lies well within the band gap, an effective shunt pathway must be present to allow flow of charge carriers. This low shunt resistance in nanostructured MoS2 can negatively impact the voltage window of operation by allowing majority carriers to flow readily under forward bias thereby limiting the maximum achievable underpotential for a MoS2 photoelectrode. More specifically in the context of a p-MoS2 photocathode, a device with low shunt resistance (i.e., redox active surface states) will allow majority hole carriers to access the corrosion reaction preventing Eonset from being achieved at more positive potentials than the corrosion redox potential will allow. Note that the utility of band diagrams, such as those presented in Figure 6 and Figure 8, is limited to bulk systems because the G

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by Cl− ions could enhance photooxidation of thianthrene because of surface-state passivation.18 Last, we examined the effects of the in-situ surface-state formation on the flat-band potential (EFB) of bulk MoS2. Bard et al. discussed the role of surface states in Fermi level pinning of p-WSe2 showing that these states could cause a shift in the EFB that corresponded to a shift in the band structure enabling the photoreduction of Ph-NO2 which was otherwise an inaccessible reaction because of its redox potential existing outside of the band gap of WSe2.18 Gobrecht et al., Tributsch et al., and Kautek and Gerischer showed how the EFB of various layered chalcogenides including MoS2 could be shifted in the presence of an I−/I2 redox couple,30,53−55 which Fujishima et al. later confirmed.27 Because the band structure is a fundamentally determining factor in the ability of the photoelectrode to drive a particular reaction of interest, measuring the EFB as surface states are generated is paramount to understanding their effects. Traditionally, the Mott−Schottky technique has been applied to assess the EFB of MoS2.27,30,53 However, accurate application of this technique is dependent on the ability to fit the data to a representative equivalent circuit model in order to determine the space-charge capacitance. Typically, the simple model of a resistor and capacitor in series is utilized to determine EFB along with selection of an appropriate frequency to determine the space-charge capacitance, which is typically materials-dependent. Kautek and Gerischer noted that this model could be applied to MoS2 with a frequency of approximately 4 kHz.30 We show in Figure S8 of the Supporting Information the application of this model to the bulk MoS2 sample which was analyzed for its dual photoactivity. Depending on the exact frequency and value of relative permittivity, we measured an EFB of approximately 0.2−0.3 V versus RHE and decreasing capacitance toward more positive potentials characteristic of an n-type semiconductor as shown in Table S1 of the Supporting Information. The EFB measured for this sample using the Mott−Schottky approach is consistent with previous results for n-MoS2 by Kautek and Gerischer at approximately 0.2 V versus normal hydrogen electrode (NHE),30 by Fujishima et al. at approximately 0.4 V versus NHE,27 and by Gobrecht et al. at approximately 0.3−0.4 V versus NHE.53 In our study, the NHE and RHE scales are approximately equivalent considering the electrolyte used (RHE = −17 mV vs NHE for pH 0.3). To complement the conventional Mott−Schottky method of determining EFB, we utilized the technique of photocurrent onset56 in which the EFB is assigned to the potential of the photocurrent switch from cathodic to anodic in direction as highlighted by Figure 9a and as explained with a band diagram in Figure 9c. This method allowed us to dynamically measure EFB over the course of the study. Because the exact transition point depends on the direction of the potential sweep (anodic or cathodic) as a result of the capacitive arrangement of ions at the semiconductor−electrolyte interface, we present the average between anodic and cathodic sweeps for a given cycle in Figure 9b. The EFB measured using this technique yielded values of 0.4−0.6 V versus RHE with an average of 0.5 V versus RHE that remained constant over the 11 h of the study. This value is consistent with the value obtained by Schneemeyer and Wrighton for n-MoS2 using the same technique of photocurrent onset at approximately 0.5 V versus NHE.57 The difference between these EFB measurements by photocurrent onset versus those from Mott−Schottky (0.2−0.3 V vs RHE) may be partly attributed to deviations from the equivalent

circuit model which has been noted by Tributsch.58 In spite of the possible differences in the exact value of the EFB, we observed that the EFB of the bulk MoS2 did not change significantly as the number of surface states increased over time as shown in Figure 9b suggesting that the band structure was not affected. However, other factors could affect this result. For example, it is possible that the sample was already Fermi level pinned from the start of the study or that the number of surface states generated during the course of this study was insufficient to induce Fermi level pinning. Notwithstanding, these data suggest that the progressive generation of additional surface defects did not additionally modify the interfacial energetics and that their influence on the band structure of nanostructured MoS2 may be minimal. This supports our earlier conclusions with respect to anodic and cathodic photocurrents being influenced by recombination and catalysis as opposed to changes in barrier height or band alignment.

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CONCLUSIONS

To develop improved photoelectrodes, it is important to understand the fundamental physical and chemical phenomena occurring at the photoelectrode surface. Nanostructured forms of MoS2 have been shown to be excellent catalysts for electrochemically driven H2 evolution; however, to the best of our knowledge, these materials as photoelectrodes have yet to achieve practically relevant photocurrent densities on the order of mA/cm2. This suggests that many of the factors concerning surface-state effects may play a significant role in determining overall photoactivity and must be taken into consideration in any attempt to develop efficient MoS2 photoelectrodes. Consequently, in this study, we aim to understand the role of surface defects in MoS2; we developed a controlled, in-situ method to create surface defects (electronic surface states) during electrochemical and photoelectrochemical studies of single crystal MoS2 producing surfaces with features that bridge the gap between nanostructured MoS2 and pristine single crystal MoS2. We show that the decrease in photoactivity as a result of the in-situ surface-defect generation may be largely balanced by an enhancement in catalytic activity for the hydrogen evolution reaction from edge sites which are known to be highly catalytic for that particular reaction. The 11 h of stability testing conducted here may not be sufficient to fully account for the behavior of MoS2, and further defecting of the sample surface may reach a point in which losses in photoactivity significantly outweigh gains in catalytic activity. Regardless, it may yet be possible to achieve greater photovoltages from nanostructured MoS2 by leveraging quantum confinement effects as previously proposed, but careful attention must be paid to the chemistry and electronic properties of the edge sites, which may need to undergo chemical transformation or passivation to yield optimal properties. The tunability of the electronic structure of these edge sites demonstrated in the literature offers a path forward to potentially improve photogenerated carrier lifetimes, while the development of passivating treatments may provide a way to enable MoS2 as well as other types of semiconductors to become viable photoelectrodes for applications in energy storage on the basis of chemical transformations. H

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ASSOCIATED CONTENT

S Supporting Information *

Additional XPS, Raman, XRD, and PEC data are available. This material is available free of charge via the Internet at http:// pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was primarily supported by the U.S. Department of Energy, Office of Energy Efficiency & Renewable Energy, under subcontract NFT-9-88567-01 under Prime Contract No. DEAC36-08GO28308. Partial support for physical and electrochemical characterization was provided by the Center on Nanostructuring for Efficient Energy Conversion (CNEEC) at Stanford University, an Energy Frontier Research Center funded by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under award number DESC0001060. We thank Prof. Mahendra Sunkara, Mr. Dustin Cummins, and Mr. Ezra Clark (University of Louisville) for providing materials used in this study. We also thank Dr. Jakob Kibsgaard for helpful discussions and Mr. Jesse D. Benck for technical assistance.

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