APPLIED SCIENCES AND ENGINEERING Bulk layered ...Therefore, the Chevrel phase formation was understood to be the local alloying effect at the interfaces even upon an-nealing at 500°C

SC I ENCE ADVANCES | R E S EARCH ART I C L E

APPL I ED SC I ENCES AND ENG INEER ING

1Department of Energy Science, Sungkyunkwan University, Suwon 440-746, Republic ofKorea. 2Korean National Research Foundation Shinjin Scientist Program, SungkyunkwanUniversity, Suwon 440-746, Republic of Korea. 3Institute of Applied Physics, University ofHamburg, Jungiusstrasse 11, 20355 Hamburg, Germany. 4Department of Chemical andBiomolecular Engineering, Yonsei University, 50 Yonsei-ro, Seodaemun-gu, Seoul 120-749, Republic of Korea.*These authors contributed equally to this work.†Corresponding author. Email: [email protected] (C.B.); [email protected] (H.S.)

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Bulk layered heterojunction as an efficientelectrocatalyst for hydrogen evolutionChangdeuck Bae,1,2*† Thi Anh Ho,1* Hyunchul Kim,1 Seonhee Lee,1 Seulky Lim,1 Myungjun Kim,1

Hyunjun Yoo,1 Josep M. Montero-Moreno,3 Jong Hyeok Park,4 Hyunjung Shin1†

We describe the spontaneous formation of composite chalcogenide materials that consist of two-dimensional (2D)materials dispersed in bulk and their unusual charge transport properties for application in hydrogen evolution reac-tions (HERs). When MoS2 as a representative 2Dmaterial is deposited on transition metals (such as Cu) in a controlledmanner, the sulfidation reactions also occur with the metal. This process results in remarkably unique structures, thatis, bulk layered heterojunctions (BLHJs) of Cu–Mo–S that contain MoS2 flakes inside, which are uniformly dispersed inthe Cu2Smatrix. The resulting structures were expected to induce asymmetric charge transfer via layered frameworksand tested as electrocatalysts forHERs. Upon suitable thermal treatments, theBLHJ surfaces exhibited theefficientHERperformance of approximately 10 mA/cm2 at a potential of −0.1 V versus a reversible hydrogen electrode. The Tafelslope was approximately 30 to 40mV per decade. The present strategy was further generalized by demonstrating theformation of BLHJs on other transitionmetals, such as Ni. The resulting BLHJs of Ni–Mo–S also showed the remarkableHER performance and the stable operation over 10 days without using Pt counter electrodes by eliminating any pos-sible issues on the Pt contamination.

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INTRODUCTIONIn principle, electrolysis of water into oxygen and hydrogen can offer aclean, renewable energy resource. The cost-effective and efficientsplitting of water is a critical issue in technologies and economies withenergy delivery systems that use hydrogen, enabling zero emission ofgreenhouse gases (1). A key challenge in the electrocatalyst design isthe performance-cost trade-off of using the platinum-group metals asa cathode for the hydrogen evolution reaction (HER) (2–5). Two differ-ent approaches have constituted the major branches of science in thisdomain of research. One approach is to use less or optimized amountsof precious elements, such as nanoparticulate Pt on carbon supports (6).The other approach is to find new alloy compositions based on Pt. Inexploring new alloy composition, computational simulations are bene-ficial in the development of previously unknown single-phase systems(7, 8). Although more complex surface structures involving platinumhave exhibited experimental improvements (9–11), efficient electroca-talysts consisting of nonprecious elements have recently been more ac-tively studied (12–15). Transition metal chalcogenide–based systems,such as MoS2, provide many possibilities for HER because they haveunique anisotropic surface/transport properties and provide surfaceswith the desired binding energy/site for H+ (16–21). However, amaterial equivalent to platinum in terms of the onset potential, the Tafelslope, and the long-term stability has not yet been developed.

The activation processes on the catalyst surfaces were believed to bethe essential mechanisms for HER, and thus, surface energies matter inthe development ofHER catalysts with high efficacy. This would be trueas for the case of monometallic Pt (22), and seeking platinum-likesurfaces seems rational among cost-effective elements. Recent researchactivities have been headed in this direction accordingly. However,

when nonmetallic catalysts are involved, the charge transfer resistancerather than surface reactions themselves would play a significant roleand should be considered. A recent study by Voiry et al. (23) experi-mentally exhibited that the catalytically inert basal plane of 2H MoS2can be active by controlling the charge transfer resistance of the system.Beyond identifying platinum-like surfaces, we propose here a novelelectrocatalytic concept consisting of two-dimensional (2D) materialsin bulk; we refer to this concept as the inorganic bulk layered hetero-junction (BLHJ). Our approach is based on a simple fabrication techni-que for the direct growth of MoS2 on self-supported metallic substratesvia sequential gas-phase reactions, which result in the spontaneous for-mation of dense BLHJ structures via spontaneous sulfidation reactions.These structures contain MoS2 flakes, which are dispersed in the Cu2Smatrix (resembling “straw inmud plaster”). The present system not on-ly features distinctive inorganic, dense BLHJ structures, which are dif-ficult to prepare experimentally using any other methods, but also haslayered 2Dmaterials as the key anisotropic components that trigger un-usual charge transfer processes.

RESULTSWe selected MoS2 as the representative 2D component with the feasi-bility of anisotropic transport properties (17). Incorporating the targetlayered materials into the bulk chalcogenide host with secure contactinterfaces between suitable nanoscale junctions in a controlled manneris difficult (Fig. 1B). We used the sequential gas-phase surface reactiontechnique for which the reactants, such as Mo or S, are independentlydelivered into the substrates so that only the surface-limited reactionsoccur. Because the substratemetal used for the self-supporting electrodeis simultaneously sulfidizable, this growthmode is expected to incorpo-rate a layered system into the bulk chalcogenide host. In the design ofthe BLHJ, another important consideration is to select two chalcogenidesystems that are thermodynamically immiscible at given temperatures.For example, copper as the substrate material and Mo for sulfidationwill lead to the desired immiscible phase separation to form an in-organic BLHJ with MoS2 at relatively lower temperatures (


systems, it is Earth-abundant and cheap. Moreover, the quality of Cu asthe electrochemical electrode has been verified in the field of battery,and thus, the possible side effects could be ruled out.

First, we established the sequential gas-phase surface reaction tech-nique for the precursors on SiO2/Si substrates in a flow-type reactor(about 150 mm in diameter), analogous to the atomic layer deposition(ALD) procedure. To obtain a preliminary result, for example, we usedMoCl5 andH2S as the reactants and separated their alternating exposuresbyN2 purging (fig. S1). The nucleation and growth along the basal planes


were strongly affectedby the gas flow fromthe inlet of the chamber (fig. S2).Remarkably, the uniformity of thickness was not maintained across thesubstrates under the given process conditions (fig. S2). This result hasnot been reported previously in the context of ALD studies (25, 26), butit is understandable on the basis of the strong structural anisotropy ofthe resulting materials. The resulting MoS2 layers have distinctivefeatures that are dependent on the distance from the chamber inletnot only in thickness but also in terms of the morphologies (fig. S2D).Larger uniformgrowth zoneswere optimizedby increasing thepulse time

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Fig. 1. The anisotropic charge transport in the bulk and the spontaneous formation of Cu–Mo–S–baseddense nanocomposites during deposition ofMoS2 directly onCu foils. (A) Illustration of a typical bulk heterojunction (BHJ). A bulk heterojunction consists of particulate/granular structures connected to the electrodes with different chemicalpotentials (green and red plates). When either electrons or holes are generated, they are then separated isotropically (black arrows). (B) Illustration of the anisotropic chargetransport in bulk. The BLHJ concept uses layered chalcogenides inside the bulk chalcogenides. When charges are injected into the layered chalcogenides from one electrode (forexample, red plate), they would undergo fast transfer because of the strong anisotropy in the transport properties. (C) Image of the as-grown CMS layers on Cu foil. (D) Plan-viewSEM of the as-deposited CMS. (E) Cross-sectional scanning transmission electron microscopy (TEM) image of the resulting CMS layer. (F) Bright-field TEM image of the cross-sectioned TiO2/CMS/Cu, where the topmost layers consist of Pt particles and carbon deposited as protection for focused ion beam sectioning. (G) High-resolution TEM (HR-TEM)image of the Chevrel phase of Cu2.76Mo6S8 formed at the interface of MoS2 and Cu2S. The inset shows that the electron diffraction patterns collected in the dashed orange circleconfirm formation of the Chevrel phase. ZA, zone axis.

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(tp) for MoCl5 (fig. S3), and we used this phenomenon in the presentstrategy for preparing BLHJ structures.

On the basis of the chemical reaction of MoCl5 and H2S in the gasphase, we suspect that chloride ions are inevitably left in the films duringthe reaction; thus, the resulting MoS2 nanoflakes incorporate Cl as animpurity. Evidence for the presence of Cl was confirmed by the chem-ical analysis techniques of energy-dispersive x-ray (EDX) spectroscopyand x-ray photoelectron spectroscopy (XPS) (fig. S4, A to E), althoughthe layered nature was preserved (figs. S4F and S7). Approximately 5 to10 atomic % of chloride ions were detected. Mott-Schottky measure-ments revealed the significant amount of negatively charged species,such as Cl ions, that had a density of ~1023/cm3 (fig. S8). This orderof magnitude is consistent with the chemical analysis results. The dop-ing density was measured using the Hall measurement technique withvan der Pauw geometry and indicates that the Cl dopants are not fullythermally activated (with an electron density of ~1019/cm3; fig. S9). Thisresult is simply a consequence of the low growth temperature (that is,250°C). In a recent study byYang et al. (27), this Cl dopingwas reportedto decrease the contact resistance through a significant reduction of theSchottky barrier width. Therefore, the fact that our methodology is in-trinsically capable of Cl doping should be beneficial in the design of thepresent BLHJ structures with efficient charge transport characteristics,as will be discussed later.

The aforementioned processes were used to directly deposit MoS2onto copper foil (~20 mm in thickness). The corresponding surface re-actions were visible across a large area (approximately 10 cm × 10 cm;Fig. 1C). As shown in Fig. 1D, scanning electron microscopy (SEM)image is the representative morphology of the resulting materials onCu (see fig. S11 for more details). Well-dispersed nanoparticles (brightcontrast) embedded in the matrix were observed. A representativescanning transmission electron micrograph shows that the layeredstructures were densely incorporated into thematrix (that is, CuSx) dur-ing the reactions (Fig. 1E). The embodiment ofMoS2 and the sulfidationof Cu simultaneously occurred during the fabrication ofMoS2, resultingin the formation of CuxS and MoS2 (hereafter termed as CMS), whichare immiscible according to their phase diagram at the given tempera-ture (24). The sulfidation of Cu can be understood in the context of theadsorbate-induced surface band bending (28), charge transfer by theH2S molecules, and the subsequent diffusion processes of sulfur inside.We observed a much higher growth rate of CuSx/CMS than that ofMoS2 by comparing the resulting thicknesses of MoS2 and CMS layers(fig. S12). The degree of sulfidation of Cu was also affected by the flowduring the reaction, as in the case of MoS2, whereas the amounts ofMoS2 (determined by detecting Mo) inside the CMS layers remainednearly constant (fig. S12C). This independence offers valuable benefitsto the rapid optimization of the MoS2 concentration within the bulklayers when the present system is applied to the HER.Within the thick-ness gradient investigated, the HER activities showed few differences(fig. S13). A comparison of the resulting thicknesses of the depositedhybrid composites confirmed that we could generate BLHJs of Cu–Mo–S with controlled amounts of MoS2.

The subsequent thermal annealing at 500°C (for ~1 hour under N2flow; see Materials and Methods for more details) allowed further con-trol over the heterojunction structures of CMS by local alloying at theinterfaces between CuxS and MoS2 (fig. S14). Investigation by TEM re-vealed the formation of Chevrel CuxMo6S8 clusters at the Cu2S/MoS2interfaces (see the superlattices marked as blue arrows in fig. S15). No-tably, according to the bulk phase diagram of the Cu–Mo–S system, al-loyingCuxMo6S8 is expected at higher temperatures, that is, temperatures


greater than ~700°C (24). Therefore, the Chevrel phase formation wasunderstood to be the local alloying effect at the interfaces even upon an-nealing at 500°C (Fig. 1G). Although their presence was evident in theTEM investigation, the clear detection of the smaller portions of bothMoS2 and CuxMo6S8 was limited to the x-ray diffraction (XRD) patterns(fig. S16) (29). Upon being subjected to further heat treatment processes,the Chevrel phase was unambiguously observed to be Cu2.76Mo6S8 in theXRD patterns (fig. S16). Note that the overall distribution of Mo wasuniform across the CMS layers (figs. S17 to S19). These observations fur-ther imply that the CMS surfaces were also partly terminated by both theMoS2 nanoflakes and the Chevrel clusters, elucidating the CMS/TiO2interface in detail (fig. S20).

We constructed a conventional three-electrode measurement setupin homemade Teflon cells (fig. S21) using a CMS/Cu cathode, a Pt an-ode, an Ag/AgCl reference electrode, and a sulfuric acid electrolyte. Thecathode was first protected with amorphous TiO2, which was grown byALD to achieve initial protection (fig. S22, A and B) (30). Notably, thisTiO2 coating can initially function as an efficient electronic transportlayer for HER (31) and serve the initial activation processes (for exam-ple, electrical annealing) of the CMS layers, as will be discussed later.HER activity was investigated at different processing steps and com-pared with control samples (Fig. 2, A to C). Although Cu and Cu2Sthemselves exhibit little HER activity under the given measurementconditions, the as-grown CMS that was exhibited substantiallyenhanced the HER activity. The annealed CMS shows further improve-ment in the HER. Notably, although the surfaces of Cu2S showed neg-ligible HER activities, the presence of the MoS2 inside certainlycontributed to the enhancement of the charge transfer characteristics.After one more junction of TiO2 was added not only as a protectivelayer but also as an efficient charge transport layer, remarkably, ourCMS material exhibited a nearly zero onset potential and a Tafel slopeof ~39 mV/decade (dec) (Fig. 2C). These values are comparable withthose of our sputtered Pt films and the literature values of Pt (22, 32).To the best of our knowledge, the present study shows the bestperformance record for HER activities among any other reportedmaterials’ systems using nonprecious elements in terms of both theonset potential and the Tafel slope. Note that the resulting high currentdensity did not stem from the large specific surface area and representsthe material’s property itself. Normalizing the porosity in the HERperformance simplifies comparisons among intrinsic properties (32).The roughness factor of our system was estimated to be ~1.3 byelectrochemically active surface area with double-layer capacitancemeasurements (fig. S23). Surface roughness of our planar structure is5 to 8 nm in root mean square roughness (RRMS) measured by atomicforce microscopy (AFM) (fig. S22C). Moreover, the TiO2/CMS/Cusamples were operated stably as HER cathodes for well over a 10-dayrun at −0.1 V (Fig. 2E). Possible issues on the contamination of Pt wereruled out by carefully analyzing XPS results before and after stabilitytests (figs. S24 to S26). These superior performances that are compara-ble with that of Pt were reproducibly observed in samples from fourdifferent batches (fig. S27). Nonetheless, a lingering suspicion on thesubtle amounts of Pt contamination would be the electrochemical re-deposition of Pt during HER under a harsh condition (33) (that is,cycling at larger voltages than −0.1 V versus reversible hydrogenelectrode (RHE) used for stability tests, for example, similar to fig.S25). Therefore, we present the HER results by both the conventionalthree-electrode measurements using Pt counter electrodes as the stan-dard methodology (that is, BLHJ-Ag/AgCl-Pt) and the three-electrodemeasurements with BLHJ as monolithic electrodes [that is, BLHJ–SCE

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(saturated calomel electrode)–BLHJ]. The HER results of the CMSlayers with a non-Pt counter electrode showed similar performance(fig. S28). How to intrinsically avoid the usage of Pt electrodes willalso be demonstrated below with the development of Ni–Mo–S (NMS)systems.

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DISCUSSIONAquestion naturally arises as to what the presentmechanisms are. Notethat this experimental value of the Tafel slope (around 30 mV/dec) hasbeen observed on Pt(111), assuming extreme coverage of hydrogen andtheTafel reaction (34–36). Our observationmight indicate that the rate-determining steps are not simply based on the surface species adsorbed(that is, binding energy matters), but the electron transfer kineticsthrough the bulk should be seriously considered. One clue is that theimpedance analysis exhibited a systematic reduction in the resistanceof the whole system as a result of thermal treatment and the additionof distinctive junction interfaces (Fig. 2D).We ascribe the observed highperformance forHER to not only themacroscopicmechanismbased onband diagrams but also the microscopic structures of our unique sys-tem. To measure the work functions of our CMS layers and other con-trol samples, we first complementarily used ultraviolet photoelectronspectroscopy (UPS) (fig. S29). The order of the work function values(Fs) was MoS2 > Cu2S > CMS. This implies the reduction of energybarriers when electrons transfer from the Cu electrode and is consistentwith the impedance analysis. The presence of Cl in the MoS2 and itsdiffusion in the resulting BLHJ structures should be addressed. If the


MoS2 flakes are to be formed inside the BLHJs in a similar way, thentheCl-dopedMoS2 nanoflakes should contribute to improved transportproperties (figs. S5 and S6). The estimated small energy gap from anArrhenius plot also supports easy processes for the thermal activationof electrons (fig. S10). Therefore, the presence of Cl and the resultingelectrical doping effect will play a key role in the charge transfer viathe inner frameworks consisting of the MoS2 flakes. Moreover, theXPS results on the annealed CMS would testify a possibility of Cl dif-fusion intoCu2S if larger amounts ofClwere detected out of the surfacesof CMS than those of MoS2. We observed the higher portions of Clwhen compared with those expected in theMoS2 nanoflakes. From thisobservation, we could therefore conclude the presence of Cl even in theCu2S matrix, and similar electrical effects, such as a reduction of theSchottky barrierwidth by doping, are expected. At the interface betweenTiO2 and electrolytes, both scenarios are possible. One is the transfer viathe surface Cu2S because the conduction level of TiO2 is well below thatof Cu2S. The other is the transport through the conductive TiO2. Theoptical absorptionmeasurementswere carried out to estimate the band-gap of our TiO2 layers (fig. S30). Although whether the ALD-grown,amorphousTiO2 has direct or indirect transitions is not clear, we believethat the charge transfer mechanism occurs below the conduction bandofTiO2 through the defect sites (fig. S31). The detailed transportmecha-nisms of the ALD-grown, amorphous TiO2 are under debate (37–39).Nonetheless, the coating of TiO2 overlayers resulted in the formation ofideal interfaces with our CMS layers in terms of electron transfer andinitial activation. The former is further understood as a framework of theformation of Cu2S/TiO2 ohmic junctions. Notably, all the other control

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Fig. 2. Electrochemical analysis and HER of the CMS layers. (A) Linear sweep voltammetry. (B) Onset HER potentials. (C) Tafel curves of the CMS layers (as-grown, cyan;annealed, orange; annealed with TiO2, red) together with those of the control materials (Cu, black; MoS2, green; Pt, blue). (D) Nyquist plots. (E) Chronogalvanometry of MoS2 andCMS/TiO2. The legend colors in (E) are common to all.

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junctions, such as MoS2/TiO2, exhibited no improvement in chargetransfer during HER (fig. S32). A very recent report demonstrated thatTiO2 layers grown via the same ALD chemistry exhibited significantimprovement in electron collection efficacy, with a shift (greater than0.2 V) of the open-circuit voltage in InP heterojunction solar cells(40). The Cu2S/TiO2 control junctions solely exhibited clear improve-ment for HER (fig. S32). Note that the amorphous titania layers werenot kept on the CMS layers upon long-term operation by detecting noXPS signal of titania (fig. S26). Nevertheless, this layer functioned forinitial activation of the underlying CMS layers for long-term stability.We believe that the presence of TiO2 surface layer contributes to theinitial charge extraction processes from the CMS BLHJs as follows.We suspect that the long-term operation at the high current density(about tens of microampere per square centimeter) resulted in the elec-trical annealing of the CMS layers (41), and the performance would be-come stable possibly by improving the electrical connectivity of theframeworks of MoS2. This argument could be supported by comparingthe XPS results after ~100-hour testing (fig. S26) and the chronogalva-nometry results for ~10 days (Fig. 2). Moreover, these currents could befurther localized for their pass only through theMoS2 frameworks [seethe results of conductive AFM (C-AFM) below].

Second, the unique microstructures inside our system are synergisti-cally responsible for the observedHER activity. The local probemeasure-ments on the resulting CMS layer confirmed that the MoS2 nanoflakeswere likely manifested in the transport mechanisms. We carried outC-AFM and Kelvin probe forcemicroscopy (KPFM)measurements onthe surfaces of the CMS. The potential maps that were thus monitoredindicated the successful formation of nanoscale spatial junctions on thesurfaces, indicative of the direct modification of the surface energy byembedding 2D materials inside (fig. S33). The local conduction resultsalso imply the presence of selective pathways in the conductionthroughout the CMS layers, probably via the MoS2 nanoflakes (Fig. 3,A to C). Notably, these measurements are sensitive to the ideality factorof the junctions between the tip and the sample (42). Our systematic


measurements also demonstrate the formation of conductive junctioninterfaces even at very small loading forces (fig. S34). In addition to thepresence of an anisotropic trigger for charge transport, the inhomo-geneity in the energy landscape might help evolve H2 because of theasymmetric charge injection for HER. The ingeniously arranged localnucleation in a nanometer/micrometer spatial regions would be bene-ficial for HERs (43). All these findings result in synergetic contributionsin determining the charge transfer kinetics for HER.

Therefore, we propose operation principles that might be account-able for the efficient HER performance observed in this study. The pre-sent mechanisms have been proposed first by recognizing the fact thatthe surfaces were mostly terminated by electrocatalytically nonactivematerials. Therefore, the classical Volmer-Heyrovsky-Tafel mechanismfor HER activity at metal interfaces cannot explain the high HER activ-ity observed here. Although the Tafel reaction having the ideal surfacecoverage resulted in the theoretical kinetic description to be an idealTafel slope of ~30 mV/dec, the estimated Tafel slope of ~39 mV/decfrom our BLHJs could not directly assign the rate-determining step(36). As a result, we suggested that the present BLHJ systems have lo-calized transport paths with reduced charge transfer resistance, whichserve as the true active site for HER on the basis of the abovementionedobservations. By spontaneous sulfidation reactions, followed by heattreatment, the dense BLHJ structures were prepared, and the resultingmicrostructure exhibited the successful formation of internal networksconsisting of the MoS2 and the Chevrel phases. The presence of MoS2nanoflakes with strong anisotropy inside the bulk can boost the internaltransport of charge carriers across the CMS layers inside during the sur-face reactions. That is, the metal substrate serves as a self-supportedcathode, injecting electrons into the CMS layer. In general, the driftand diffusion processes are responsible for the major transport mecha-nisms duringHER. Because theMoS2 is rather confined in the BLHJ, thestrong anisotropy should lead to an enhancement in the spatial chargetransport via different mechanisms, thereby enabling the design ofunique electrocatalysts. Moreover, the mobility of MoS2 is strongly

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Fig. 3. Selectivity of transport properties of the resulting CMS BLHJs. (A) AFM height image of CMS. (B) Corresponding current map by C-AFM at a sample bias of −0.5 V.(C) Simplifiedmodel of BLHJ structures, highlighting localized transport paths of 2Dmaterials (blue). (D) Proposedmechanisms on energy (e) barrier lowering by the localizedelectric fields.

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dependent on the dielectric screening of its surroundings (44). The re-sulting MoS2 nanoflakes were embedded in the Cu2S matrix; thus, thissituation should be advantageous to the anisotropic charge transport.As previously discussed, the Cl doping of MoS2 would modify thestructures, lower the energy barriers, and shorten the barrier widthswith its contact materials, if any (45). Moreover, the Chevrel CuxMo6S8clusters formed between the MoS2 and Cu2S should facilitate the trans-fer of charge carriers, operating as gradient junctions. The nanoscaledimensions of the resulting MoS2 flakes in the bulk CMS should leadto the environmental pinning at the Fermi levels inside the bulk and thefurther lowering of the barrier heights (45). Despite the complex geo-metries in defining the contact interfaces, the energy band diagrams


suggest ideal junction interfaces with negligible barrier heights (fig.S31) (46). As shown in Fig. 3D, the embeddedMoS2 flakes can functionas sharp tips at the near surfaces and are additionally responsible forenergy barrier lowering by the localized fields as proposedmechanisms.The present strategy reported here differs from conventional electroca-talytic concepts that seek the optimum surface binding energy to H+,not only experimentally but also theoretically (table S1). To date, nosuch electrocatalysts exist, satisfying both the ideal catalyst surface (thatis, the onset potential and the Tafel slope) and the long-term stability(fig. S35).

Furthermore, it could offer a general tool to explore other efficientcomposite electrocatalysts with 2D chalcogenide materials inside the


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Fig. 4. Electrochemical analysis and HER of the NMS layers. (A) Polarization curves. (B) Onset HER potentials. (C) Tafel analysis (NMS, red) together with those of the controlmaterials (bare Ni form, gray; sulfurizedNi form, green; Pt, blue). (D) Nyquist plots. The legend colors in (A) are common to all, except for (E). (E) Stability tests of NMSmaterials bothat 10 mA/cm2 and at −0.1 V versus RHE when using Pt as counter electrode for about 100 hours. (F) Long-term stability of our NMS samples for about 1 month where Pt wasreplaced by the other sheet of NMS (see the inset image).

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bulk. We applied the present strategy to the porous Ni form as anothertransitionmetal where the similar incorporation is expected aswell as tothe BLHJs of NMS. The difference between Cu and Ni is the cohesiveenergy (47), and the degree of sulfidation should be different whenMoS2 was grown under identical conditions. Because Ni has greaterbonding strength than Cu, the formation of NiSx was suppressed andthe relative amounts ofMoS2were found to be larger by its spontaneousincorporation as in the CMS case (fig. S36). The BLHJs of NMS alsoshowed the remarkable HER performance without postthermal treat-ments (Fig. 4). The resulting NMS plates were tested without using Ptcounter electrodes to eliminate any possible issues on the trace amountof Pt contamination. Three-electrode measurements [that is, NMS as aworking electrode, NMS as a counter electrode, and SCE as a reference]were carried out and exhibited similar performance and superior stabil-ity (Fig. 4F). These results explicitly testify that the outstanding HERperformance undoubtedly comes from our BLHJ structures. Further-more, any issues on the contamination of Pt can be completely ruledout by substituting the counter Pt electrode as our NMS, indicative ofthe technical feasibility of direct utility of the present systems for HERs.The simultaneous incorporation of 2D chalcogenides during sulfidationof desiredmetals should also open a new path for research in nanocom-posites that are currently difficult or not possible to synthesize.


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MATERIALS AND METHODSALD of MoS2 and TiO2MoS2 was grown on SiO2/Si substrates at 250°C using a commerciallyavailable ALD system (Lucida D100, NCD). MoCl5 (99.6%; StremChemicals) andH2S (3.99%, balanceN2; JCGas) were used as reactants.The MoCl5 was kept in a stainless canister at 140°C and delivered intothe chamber. Ultrahigh-purity N2 (5 N; JC Gas) was used as both thecarrier and the purging gas. The total flow rate was 200 standard cubiccentimeters per minute (sccm). A full cycle consisted of tp (MoCl5) of0.5 to 5 s and tp (H2S) of 1 s, followed by N2 purging for 30 s. The CMSwas also prepared using the same configuration with MoS2 on Cu foil(I2B ~20 mm, battery grade, ILJIN Materials). After ALD, the chamberwas cooled to room temperature, and the samples were removed fromthe chamber. Using MoCl5 allowed for sufficient vapor pressure at thelow temperature used (that is, 140°C); however, it also resulted in dop-ing byCl via the following chemical reactions:MoCl5 +H2S⇄MoSCl3 +2HCl and MoSCl3 + H2S⇄MoS2Cl + 2HCl.

Once the conditions for growing MoS2 have been established, wefurther optimized the number of ALD cycles. The progressive improve-ment was shown in linear sweep voltammetry (fig. S37), and we took3000-cycle CMS layers in the present study. The annealing conditionshave been studied in the temperature range of 300° to 700°C and chosento give an optimumBLHJ structure (figs. S38 and S39). As-grown CMSsamples were annealed in a tube furnace in a flow of inert gas (HTF-Q50, Hantech Co.). The furnace was first evacuated down to less than3 × 10−2 torr for more than 30min before introducing 250 sccm of N2(5 N; JC Gas), and the N2 flow was continuously maintained duringannealing. The ramping rate was 5°C/min until the desired tempera-ture increased, and the temperature was kept for 1 hour. The furnacewas then cooled down to room temperature naturally.

TiO2 was deposited using a separate, commercial ALD chamber(Ozone, ForALL) at 120°C. Titanium(IV) isopropoxide (TTIP; 99.99%;UP Chemical) and deionized H2O were used as the metal reactantand oxygen source, respectively. TheTTIP andwater were kept in stain-less bubblers at 70°C and room temperature, respectively. A cycle con-


sisted of tp of 2 s for both reactants andAr (5N) purging for 8 s at a totalflow rate of 200 sccm. The growth rate of TiO2was determined to be 0.4Å per cycle.

NMS was synthesized in the same way with CMS except on porousNi foam (0.8 mm thick; Wellcos Corporation). First, Ni foams were so-nicated in ethanol for 30min and then thermally treated at 800°C for2 hours under the reducing atmosphere with a mixture gas of 5%H2 inAr. Three thousand ALD cycles of MoS2 were applied on cleaned Nifoam at 250°C under the identical conditions with the case of CMS.

Structural characterization and surface analysisThe physical dimensions and morphologies were observed by field-emission SEM (JSM7500F, JEOL). The structures were investigatedby HR-TEM (JEM 2100F, JEOL) equipped with an EDX spectrometer(AZtec, Oxford Instruments). Thin sections of the samples for TEMwere prepared by focused ion beam etching (SMI3050TB, SII). Surfacechemical compositions were analyzed using XPS (ESCA Sigma Probe,Thermo VG Scientific), and surface work functions were estimated byUPS (AXIS Ultra DLD, Kratos Inc.) using a He I photon source (hv =21.2 eV).

Electrochemical measurementsElectrochemical measurements were carried out with a three-electrodesystem. Pt wires were used as the counter electrode, Ag/AgCl (MF-2052for H2SO4; Bioanalytical Systems Inc.) as the reference electrode, andthe CMS (as controls, Pt, Cu, and MoS2) as working electrodes. Theelectrochemical properties were recorded using a commercially availa-ble potentiostat (VMP-300, Bio-Logic). Cyclic voltammogramswere re-corded at a scan rate of 5 mV/s either in 0.5 or 1 M H2SO4 electrolyte.The virgin curves were discarded, and the third results were routinelydisplayed, unless otherwise specified. The polarization curves were re-plotted as the overpotential (h) versus logarithmic current density, log|j|,to obtain the Tafel slopes. The reference electrode was calibrated to RHEpotential in the electrolyte of 1MH2SO4, that is,E(RHE)=E(Ag/AgCl) +0.21V. Electrochemical impedance analysis was carried out at a bias volt-age of −0.2 versus Ag/AgCl, in the frequency range from 100 mHz to100 kHz, at a voltage amplitude of 6 mV, and at room temperature.

NMS was used as both working and counter electrode, and SCE foralkaline electrolytes was used as a reference. The three-electrode con-figuration with Pt counter electrode was also used for comparison.The reference electrode was calibrated to RHE in the electrolyte of1 M KOH (pH 14), according to E(RHE) = E(SCE, alkaline) + 0.971 V.Electrochemical impedance analysis was carried out at a bias voltageof−0.2 versus RHEwith a voltage amplitude of 10mV, in the frequencyrange from 100 mHz to 100 kHz.

Chronoamperometry and chronopotentiometry were carried out ata constant potential of −0.1 V versus RHE and a constant current den-sity of−10mA/cm2, respectively.Mott-Schottky plots were attained at afrequency of 7.8Hz, from−0.6 to 0V versus RHE. The impedance spec-troscopy results were analyzed and fitted using EC-Lab software. All themeasurements were carried out under dark conditions and mildmagnetic stirring (200 rpm, unless otherwise specified).

Local probe measurementsA commercial atomic force microscope (SPA-400, SII) was used forC-AFM and KPFM measurements using Au-coated tips [SI-DF3-A,with a spring constant (C) of ~0.2 N/m; SII NanoTechnology Inc.]and Pt/Ir-coated cantilevers [CONTPt-W, with a resonance frequency( f ) of 25 to 27 kHz andC = 1.9 N/m; NanoWorld), respectively. KPFM

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images were acquired at scan rates of 0.1 to 0.2 Hz with an applied volt-age of 1 V (peak-to-peak) and alternating current frequencies near thef value of the cantilevers. We simultaneously measured the topographyand surface potential/conduction map of the samples under ambientconditions and in the dark.


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SUPPLEMENTARY MATERIALSSupplementary material for this article is available at http://advances.sciencemag.org/cgi/content/full/3/3/e1602215/DC1fig. S1. Schematic illustration of the deposition system.fig. S2. Sequential gas-phase reaction of MoS2 on Si wafers.fig. S3. Influence of tp of MoCl5 on the uniform growth zone.fig. S4. Elemental and structural analyses of thin MoS2 films.fig. S5. Elemental analysis of annealed CMS layers.fig. S6. High-resolution XPS spectra of the annealed CMS layers in Mo (3d), S (2p), Cu (2p), andCl (2p) regions.fig. S7. Raman spectra of MoS2 (300 cycles) grown on Au.fig. S8. The Mott-Schottky measurement of MoS2 on an Au/Si substrate.fig. S9. The Hall effect measurements of MoS2 on 500-nm-thick SiO2/Si.fig. S10. Arrhenius plot of the resistivity from Hall effect measurements on ALD-grown MoS2/SiO2 (500 nm)/Si.fig. S11. SEM images of the as-grown CMS layers.fig. S12. Thickness dependence of ALD films and Mo contents as a function of position.fig. S13. HER activities from the CMS samples with thickness gradient.fig. S14. Schematic illustration of the structural evolution of BLHJs upon annealing.fig. S15. Low-magnification TEM micrograph of our CMS layers upon annealing (500°C for ~1hour under N2 flow) to give an overview of the structures that consist of layered MoS2 and thesuperstructures of Chevrel clusters (marked by yellow and blue arrows, respectively).fig. S16. XRD patterns of our CMS on Cu subjected to different thermal treatments.fig. S17. EDX elemental analysis of our TiO2/CMS/Cu structures.fig. S18. EDX line scan results for the TiO2/CMS/Cu structures.fig. S19. Detailed elemental maps of our CMS layers shown in Fig. 1E (main text), indicative ofthe origin of local variations in the detected elements.fig. S20. HR-TEM image of annealed CMS to give an overview of the local surface termination.fig. S21. Schematic of our three-electrode cell used for HER experiments.fig. S22. Surface morphology of TiO2-coated annealed CMS.fig. S23. Estimation of electrochemically active surface area of our CMS material by double-layer capacitance measurements.fig. S24. XPS analyses of CMS materials to check possible contamination of noble metals.fig. S25. Stability against scanning of the present CMS system (10,000 times).fig. S26. XPS analyses of TiO2/CMS after stability tests.fig. S27. Reproducibility tests in the electrocatalytic performance of CMS and NMS materials.fig. S28. Electrochemical analysis and HER of the CMS layers with a non-Pt counter electrode(that is, graphite).fig. S29. UPS spectra.fig. S30. Optical absorption of amorphous TiO2 grown on quartz glass.fig. S31. Energy band diagrams and the corresponding circuit models of various structures.fig. S32. HER measurements of our 40-nm-thick TiO2 with different control samples.fig. S33. KPFM study of our CMS on Cu.fig. S34. Local transport study of our annealed CMS/Cu samples.fig. S35. Comparison between our CMS/TiO2 and various HER materials (32) in theelectrocatalytic performance.fig. S36. TEM image of the NMS layer as prepared to give an overview of the structures thatdensely consist of layered MoS2.fig. S37. ALD cycle–dependent HER performance of annealed CMS.fig. S38. HR-TEM image of a CMS layer annealed at 700°C.fig. S39. Linear sweep voltammetry of the CMS layers annealed at different temperatures.table S1. Summary of the catalytic performance of various materials for the hydrogenevolution reaction, reported in the literature.Reference (48)

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Acknowledgments: We thank J. Hwang and S. W. Kim for the XRD measurements. Funding:We acknowledge the grant by the Samsung Science and Technology Foundation (SRFC-MA1502-09). Author contributions: C.B. and H.S. conceived the project. C.B. developed thepresent strategy. C.B., H.K., and T.A.H. prepared the samples. C.B., T.A.H., S. Lee, M.K., J.M.M.-M.,J.H.P., and H.S. analyzed the structures. C.B., J.M.M.-M., and T.A.H. carried out theelectrochemical experiments. C.B., S. Lim, and H.Y. performed the local probe experiments andanalysis. C.B. and H.S. coadvised the research. C.B. wrote the manuscript. All authorsreviewed the paper. Competing interests: The authors declare that they have no competinginterests. Data and materials availability: All data needed to evaluate the conclusionsin the paper are present in the paper and/or the Supplementary Materials. Additional datarelated to this paper may be requested from the authors.

Submitted 12 September 2016Accepted 10 February 2017Published 31 March 201710.1126/sciadv.1602215

Citation: C. Bae, T. A. Ho, H. Kim, S. Lee, S. Lim, M. Kim, H. Yoo, J. M. Montero-Moreno, J. H. Park,H. Shin, Bulk layered heterojunction as an efficient electrocatalyst for hydrogen evolution. Sci.Adv. 3, e1602215 (2017).

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Bulk layered heterojunction as an efficient electrocatalyst for hydrogen evolution

Montero-Moreno, Jong Hyeok Park and Hyunjung ShinChangdeuck Bae, Thi Anh Ho, Hyunchul Kim, Seonhee Lee, Seulky Lim, Myungjun Kim, Hyunjun Yoo, Josep M.

DOI: 10.1126/sciadv.1602215 (3), e1602215.3Sci Adv

ARTICLE TOOLS http://advances.sciencemag.org/content/3/3/e1602215

MATERIALSSUPPLEMENTARY http://advances.sciencemag.org/content/suppl/2017/03/27/3.3.e1602215.DC1

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APPLIED SCIENCES AND ENGINEERING Bulk layered ...Therefore, the Chevrel phase formation was understood to be the local alloying effect at the interfaces even upon an-nealing at 500°C

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