Structural Instability Stimulated Heteroatoms Co‐Doping of 2D Quaternary Semiconductor for Optoelectronic Applications

Although the structural and electrical engineering of transition metal dichalcogenides using atomic doping or doping‐induced phase modulation can be used to attain high‐performance and wavelength‐tunable optoelectronic devices, accessible substitutional doping to overcome the large lattice mismatch between the host and guest atom‐related bonding states remains elusive. This study corroborates an innovative synthetic route for molybdenum disulfide (MoS2)‐derived two‐dimensional (2D) quaternary semiconductors substitutionally doped with Re and O using a solution‐based large‐area compatible approach combined with the thermal evaporation of dopants. The substitutional doping of Re into MoS2 crystals with a large lattice mismatch is effectively accomplished by adopting structurally unstable host films, resulting in the large‐scale synthesis of 2D quaternary multi‐layers with a Re doping concentration >10%. Comprehensive spectroscopic and microscopic evaluations are performed to determine the efficacy of the host films with structural instability for the synthesis of 2D RexMo(1‐x)O2yS2(1‐y) quaternary multi‐layers. The capability of the quaternary semiconductor for versatile nanophotonic devices is validated by ascertaining the simultaneous enhancement of the photoelectrical properties with wide‐range optical absorption and photoelectrochemical properties, as compared with those of their binary counterparts.


Introduction
Molybdenum disulfide (MoS 2 ), a state-of-the-art 2D transitionmetal dichalcogenide (TMD), has attracted extensive attention in DOI: 10.1002/adfm.202310178[3][4][5] From an optoelectronic application perspective, MoS 2 possesses good stability, a suitable and tunable bandgap, high electrical conductivity, and is an abundant natural mineral.8][9][10] In particular, MoS 2 can primarily be applied as an active material for photodetectors operating in the visible wavelength range; however, enhancement of the photoelectrical performance in the near-infrared (NIR) range using band engineering is required because of the limitation of the intrinsic band structure. [11]ubstitutional doping, which is an effective strategy for modulating the bandgap of MoS 2 and passivating atomic vacancies to improve the device performance has been attempted to address this issue. [11,12]ubstitutional doping leads to covalent bonding between the hosts and dopants by replacing the host transition metal or chalcogen atoms with target guest atoms, enabling more stable doping, as compared to noncovalent approaches for attaining electrically tailored MoS 2 . [13]Moreover, atomic substitution can cause phase transitions via atomic re-arrangement originating from charge transfer. [14,15]henium (Re) atoms for the atomic doping of MoS 2 host crystals were rationally considered in this study.Because of the discrepancy between the structural features of MoS 2 (trigonal prismatic lattice) and ReS 2 (rhenium disulfide, distorted tetragonal lattice), ReS 2 multi-layers behave similarly to decoupled monolayers, arising from the lack of interlayer registry and weak interlayer coupling, leading to the confinement of a Re-based valence electron in individual layers, even in bulk form.Consequently, ReS 2 has a direct bandgap of ≈1.5 eV, irrespective of the number of layers. [16]The efficacy of Re substitution correlated with the bandgap narrowing of the MoS 2 host lattices or the formation of a metastable structure dictated by the localized octahedral 1T-phase structure, can be achieved, due to these unique characteristics of Re.Thus, Re substitution facilitates increased broadband photodetection with wide-range optical adsorption can be anticipated. [17,18]Additionally, Re doping can offer additional lattice strain-based active sites, which can be accomplished by increasing the intrinsic activity of MoS 2 catalysts. [14]However, the large lattice mismatch (≈37%) between MoS 2 and ReS 2 (lattice parameters: 0.32 and 0.64 nm) is connected to that the direct atomic substitution of Re atoms into MoS 2 is considerably challenging.
This study reports a rational approach for the synthesis of MoS 2 -derived 2D quaternary semiconductors with high Re and O doping concentrations via a large-area-compatible solutionbased synthetic route using structurally unstable ammonium tetrathiomolybdate ((NH 4 ) 2 MoS 4 ) host films.The introduction of (NH 4 ) 2 MoS 4 host films with structural instability can result in a reduction of the activation barrier related to the co-doping of Re and O atoms compared with that of the utilization of structurally robust MoS 2 host films.Notably, the co-doping of O with Re accelerates structural instability to compensate for the inconsistency of the MoS 2 lattices for the viable substitution of Re atoms accompanied by a large lattice mismatch.Moreover, this approach enables large-scale thin film synthesis that surpasses micro-scale synthesis by the conventional chemical vapor deposition (CVD) method.The authors previously developed a wafer-scale synthetic platform for 2D multi-layers using a simple coating and subsequent thermal decomposition of singlesource precursors. [11,19,20]Based on this innovative strategy combined with the synthetic platform, large-scale synthesis of a 2D Re x Mo (1-x) O 2y S 2(1-y) quaternary semiconductor for wide-range optical adsorption was achieved and validated the capability of the quaternary semiconductor for industrial application in nanophotonic devices for broadband photodetector and catalysts for PEC cell.The proposed approach can significantly improve the concentration of dopants, thereby effectively tailoring the bandgap and facilitating enhanced broadband photodetection and PEC performance.

Results and Discussion
Figure 1a shows the proposed synthetic strategy for hybrid materials (hereinafter referred to as ReO 3 -MoS 2 ) and quaternary semiconductors (hereinafter referred to as ReMoOS) by selecting the host materials (MoS 2 and (NH 4 ) 2 MoS 4 ).A (NH 4 ) 2 MoS 4 film was used as the host material to endow structural instability for effi-cient substitutional diffusion and the subsequent reaction of Re and O atoms in the 2D quaternary system.Rhenium oxide (ReO 3 ) powder was used to simultaneously supply Re and O dopants.The authors anticipate that the use of (NH 4 ) 2 MoS 4 host films with structural instability will lead to a decrease in the activation barrier associated with the co-doping of Re and O atoms, as compared to the use of highly crystalline and structurally stable MoS 2 host films.The proof-of-concept of the presented synthetic strategy for 2D quaternary semiconductors governed by the structural instability of the host material was implemented through comprehensive microscopic evaluation.First, the surface morphologies of pristine MoS 2 , ReO 3 -MoS 2 governed by a highly crystalline MoS 2 host, and ReMoOS governed by an amorphous (NH 4 ) 2 MoS 4 host were observed using field-emission scanning electron microscopy (FE-SEM).Pristine MoS 2 had an atomically flat surface, as shown in Figure S1a (Supporting Information).Figure S1b (Supporting Information) shows that the ReO 3 -MoS 2 sample synthesized using a highly crystalline MoS 2 host was incorporated with high-density nanoparticles (NPs).Conversely, no NPs were observed in the ReMoOS sample synthesized using the amorphous host, as shown in Figure S1c (Supporting Information).These results demonstrate that there was a large discrepancy between the surface morphologies of the samples synthesized using different host materials, which suggested that the atomically formed quaternary semiconductors were caused by the structural instability of the amorphous (NH 4 ) 2 MoS 4 host.To gain further structural and chemical insights into the synthesized samples, cross-sectional transmission electron microscopy (TEM) observations combined with energy-dispersive X-ray spectroscopy (EDS) elemental mapping analyses of ReO 3 -MoS 2 induced by the highly crystalline host and ReMoOS induced by the amorphous (NH 4 ) 2 MoS 4 host were performed.Figure 1b shows a representative TEM image of ReO 3 -MoS 2 induced by a highly crystalline host (pristine MoS 2 ), revealing that NPs were distinctly formed on the surface of a 2D MoS 2 layered structure, which has an interlayer spacing of ≈0.62 nm. [21]In addition, the high-resolution TEM image of the NPs highlighted in the white box in Figure 1b and the fast-Fourier-transform (FFT) image in Figure S2 (Supporting Information) reveal an interlayer spacing of 0.37 nm.According to previous studies, this corresponds to the (100) plane of ReO 3 , supporting the identification of the surface NPs as ReO 3 . [22]The formation of crystalline ReO 3 NPs reflects that the heat-driven substitutional diffusion of Re and O into MoS 2 was restricted.This phenomenon is further supported by the EDS elemental mapping analysis of Re and Mo in Figure 1b, which indicates that Re was observed on the entire surface of MoS 2 .However, despite using a structurally stable MoS 2 host, a slight distribution of Re into the MoS 2 host was clearly observed.This can be explained by the inevitable occurrence of intrinsic vacancies in the pristine MoS 2 lattice during synthesis.The supplied Re and O atoms saturated the vacancy sites of the MoS 2 lattice and agglomerated as NPs on the MoS 2 surface.Figure 1c shows a cross-sectional TEM image and EDS maps obtained from the ReMoOS quaternary semiconductor induced by an amorphous host, indicating that a 2D layered structure with an interlayer spacing of ≈0.64 nm was clearly visible.These results can be interpreted as the lattice expansion of the MoS 2 crystal via the substitution of Re and O atoms in a localized area.Therefore, it was ascertained that Re and Mo were superimposed by the EDS mapping of the ReMoOS quaternary semiconductor induced by the amorphous host.Meanwhile, the ReMoOS quaternary semiconductor was also characterized using top-view TEM (Figure S3, Supporting Information).A previous study corroborated that the (NH 4 ) 2 MoS 4 precursor was initially converted to amorphous MoS 3 and subsequently finalized into crystalline MoS 2 during thermal decomposition. [19,23]The intermediate structure of amorphous MoS 3 contains highly reactive sites that overcome the activation energy for structural re-arrangement via substitutional doping of heteroatoms. [24,25]Therefore, in the synthesis process involving the transition from MoS 3 to MoS 2 , thermally decomposed Re and O atoms efficiently diffuse into the host material, leading to the formation of a 2D quaternary semiconductor.This process signifies the advantage of the relatively unstable (NH 4 ) 2 MoS 4 host compared to the crystalline MoS 2 host in terms of its ability to facilitate the infiltration of heteroatoms.Additionally, O atoms accelerated the structural instability to compensate for the inconsistency of the MoS 2 lattices for the viable substitution of Re atoms accompanied by a large lattice mismatch. [26,27]omprehensive spectroscopic and microscopic evaluations corroborate clear evidence that the selection of host materials determines the dimensionality of the heteromaterials.
To elucidate the host-material-dependent doping of Re and O atoms in the 2D MoS 2g mode, originating from the in-plane displacement of Mo and S atoms, and the A 1g phonon mode at 406.3 cm −1 originating from the vibration perpendicular to the basal plane (Figure 2f). [28,29]The relative frequency difference extracted from the wavenumber difference between the E 1 2g and the A 1g modes was estimated to be 23.9 cm −1 for pristine MoS 2 , confirming the spectroscopic evidence of the presence of a multi-layer structure, which replicates the TEM results. [21]he Raman spectra of ReO 3 -MoS 2 and ReMoOS were compared with that of pristine MoS 2 , as shown in Figure 2g,h.These results provided a clear indication of the spectral alterations after the formation of ReO 3 -MoS 2 and ReMoOS, that is i) a decrease in the full-width half maximum (FWHM) of the A 1g phonon mode, as compared with that of pristine MoS 2 , and ii) a blueshift in the A 1g phonon modes of MoS 2 via co-doping with Re and O atoms (Figure 2i).The FWHM of the A 1g phonon mode of ReO 3 -MoS 2 corresponded to 8.4, as compared with that of 10.8 in pristine MoS 2 , demonstrating a significant reduction.This reduction can be attributed to the saturation of the intrinsic vacancies in Mo and S with the supplied Re and O atoms, that is, the effective heating of the defective sites.The FWHM of the A 1g phonon mode of Re-MoOS was 9.6, which is greater than that of ReO 3 -MoS 2 , presumably because of the local distortion caused by the increase in the Re and O contents within the multi-layer structure.Meanwhile, no distinct peaks related to Re-S, Re-O, and Mo-O were observed in the total Raman spectrum (Figure S5, Supporting Information).This suggests that the MoS 2 phase still dominates in the sample, making it challenging to observe the presence of a single phase other than MoS 2 .For instance, in the case of Re-MoS 2 , it has been reported that a substantial Re percentage in the range of 40-50% is required to exhibit new peaks. [30]Nevertheless, it was noted that the substitution of Re and O led to the structural alteration of MoS 2 , as ascertained by the Raman features.Both MoS 2 and ReS 2 exhibit a similar arrangement, where the Re and Mo atoms are surrounded by six S atoms.When Re atoms with a relatively smaller atomic radius substitute Mo atoms, they cause internal strain in the quaternary semiconductor and the crystal structure of MoS 2 undergoes local distortion in the quaternary semiconductor.In particular, the substitution of Re atoms at the Mo sites led to changes in the length of the initial Mo─S bonds.This alteration served as the primary driving force behind the tensile strain observed in the 2D alloy, ultimately leading to a blue shift in the phonon modes associated with MoS 2 , as shown in Figure 2i,j. [30,31]In addition, the introduction of oxygen atoms can also lead to changes in MoS 2 , causing a blue shift in the A 1g mode. [27]o investigate the effect of the formation of quaternary semiconductors by the atomic substitution of Re and O on the optical properties, UV-vis absorption spectroscopy was performed.Figure 2k shows the UV-vis absorption spectra of the various samples (MoS 2 , ReO 3 -MoS 2 , and ReMoOS), showing that all samples exhibited the two characteristic absorption peaks of MoS 2 , known as A and B excitons, respectively. [32,33]The absorption properties of ReMoOS in the visible light region were clearly enhanced, as compared to those of MoS 2 and ReO 3 -MoS 2 .The absorption edges of ReO 3 -MoS 2 and ReMoOS shifted to longer wavelength regions, which indicated a discernible decrease in the optical bandgap (E g ) energy due to atomic substitution-induced band engineering. [17]Figure 2l and Figure S6 (Supporting Information) show the indirect and direct bandgap energy levels of MoS 2 , ReO 3 -MoS 2 , and ReMoOS.MoS 2 and ReO 3 -MoS 2 exhibited in-direct bandgap energies of 1.5 and 1.47 eV, respectively, as estimated from the intercept of the linear portion of the plot of (h) 1/2 as a function of energy (h).Notably, ReMoOS exhibited a narrower indirect bandgap energy (1.36 eV) than those of the other samples, and the direct bandgap also decreased after the formation of the quaternary semiconductor (Figure 2l,m).Considering these comprehensive spectroscopic evaluation results, the proposed approach is indicated to effectively modulate the bandgap and enhance optical absorption properties, which can potentially lead to improvements in the performance of photode-tectors and PEC.Meanwhile, ReO 3 exhibits a metallic nature, unlike other stoichiometric metal oxides that offer wide-bandgap semiconducting characteristics.These unique properties of ReO 3 enable the possibility of plasmon resonance at the interface with MoS 2 , and ReO 3 is known to exhibit a plasmon absorption band at ≈530-590 nm in its absorption spectrum. [34,35]Although the properties of ReO A two-terminal device for visible or NIR photodetection was fabricated to investigate the photoelectrical properties of the various samples (pristine MoS 2 , ReO 3 -MoS 2 , ReMoOS), as depicted in Figure 4a.Electrical contact was formed via thermal evaporation using a shadow mask, defining a channel length and width of 50 and 500 μm, respectively (Figure 4b). Figure 4c-e  A linear correlation between the photocurrent and the applied bias voltage was observed with both lasers, which was attributed to the increase in the carrier drift velocity. [11,36]The ReMoOSbased photodetector exhibited a significantly enhanced photocurrent and rapid response/recovery times, as compared with those of the MoS 2 -and ReO 3 -MoS 2 -based photodetectors, as shown in Figure 4c-e.The photocurrent of ReMoOS at 532 nm was The enhanced photoelectrical performance and reduced response/recovery times in the quaternary semiconductor systems can be attributed to the following effects: i) Improved crystalline and enhanced suppression of recombination achieved by anchoring heteroatoms in the vacancy sites.ii) Bandgap narrowing induced by substitutional doping.The theoretically and experimentally confirmed bandgap narrowing through substitutional doping leads to a decrease in the metastable chargetrapping centers, [11] consequently enhancing the photodetector performance.Figure 4f,g shows the extracted photocurrents as a function of the voltage and illumination power of the 532 and 1064 nm lasers, respectively; the laser power used in the experiment was set at 5, 10, 15, and 20 mW at a constant voltage of 20 V.As mentioned previously, the linear correlation between the photocurrent and voltage can be attributed to a progressive increase in the drift velocity of the photoexcited carriers.The ReMoOS-based photodetector conformed to a linear relationship between the photocurrent and laser power, whereas the MoS 2based photodetector exhibited slight sub-linearity under 532 and 1064 nm light illumination.This tendency can be attributed to the formation of localized deep-level defect states (DLDSs) within the bandgap caused by the presence of intrinsic vacancies. [37,38]LDSs can act as recombination centers for photoexcited carriers, leading to a reduction in the carrier lifetime.This effect is dominated by the higher power density of the photons. [39]he doping of Re and O at high concentrations facilitated the effective elimination of these unintentional vacancies associated with the DLDSs, resulting in a positive impact on photodetection properties through enhanced recombination suppression and improved lifetime of photoexcited carriers.Figure 4h shows the photocurrent enhancement factor (I doped-MoS2 /I MoS2 ) under 532 and 1064 nm laser illumination, indicating that the enhancement factor of the ReMoOS-based photodetector was greater than that of the ReO 3 -MoS 2 -based photodetector.Consequently, the ReMoOS-based photodetector exhibited a significantly improved photocurrent ratio exceeding 125 and 400 times that of MoS 2 at 532 and 1064 nm, respectively.These photodetector performance results emphasize the efficacy of the proposed strategy for synthesizing a quaternary-semiconductor-based photodetector with exceptional broadband sensitivity spanning from the visible to NIR region.
To further demonstrate the potential of ReMoOS as a multifunctional material for various applications, its capabilities for PEC hydrogen production were extensively investigated using a Si photocathode as a reference.Despite the potential of the MoS 2 /Si photocathode, pristine MoS 2 with a 2H stable phase has not been fully exploited because the catalytically active sites are primarily confined to the edge rather than the basal plane. [40]herefore, activation of the basal plane and the implementation of an innovative approach to enhance the PEC performance can unlock the true potential of these materials, making them even more compelling as photocathodes.Figure 5a shows the linear sweep voltammetry of various photocathodes obtained from the simulated AM 1.5 G illumination, clearly indicating a pronounced photoresponse attributed to the PEC reaction, in contrast to the dark current density-potential curves. [29]The Si photocathode exhibited a slow negative onset potential of ≈−0.2 V versus reversible hydrogen electrode (RHE), resulting in an inferior PEC performance due to poor catalytic activity and sluggish hydrogen evolution reaction (HER) kinetics.The incorporation of a MoS 2 thin film on the Si photocathode induced a positive shift in the onset potential (≈−0.05V vs RHE), accompanied by an improved PEC performance due to the enhanced catalytic activity of the photocathode.Notably, the ReMoOS corresponding to quaternary semiconductors, including Re and O, demonstrated a significantly enhanced PEC photocurrent and an accelerated onset potential of ≈0.24 V versus RHE, surpassing the PEC performance of both the Si and MoS 2 photocathodes.The improved PEC properties of ReMoOS can be attributed to the following effects: i) Re doping serves to not only enhance catalytic activity by activating the basal plane but also enables efficient light harvesting. [14,41]i) The O doping within the MoS 2 basal plane accelerates the HER kinetics by increasing the intrinsic conductivity of MoS 2 and activating inert sites on the basal planes. [9]iii) Co-doping of Re and O in ReMoOS synergistically enhanced light absorption, suppressed recombination, and allowed for bandgap engineering, all of which collectively contributed to the improved PEC performance.Therefore, the Re and O co-doped ReMoOS photocathode exhibited a significantly enhanced PEC reactivity.Meanwhile, the PEC performance improvement observed in the ReO 3 -MoS 2 photocathode was relatively unsatisfactory, underscoring the utility of the proposed approach for fabricating a quaternary semiconductor (ReMoOS) for PEC applications.The Tafel plot in Figure 5b shows that ReMoOS exhibited faster catalytic reaction kinetics (Tafel slope of 157 mV dec −1 ) than that of ReO 3 -MoS 2 (184 mV dec −1 ) or MoS 2 (245 mV dec −1 ).The smallest Tafel slope of ReMoOS indicated an enhancement in the HER kinetics, which is consistent with the observed improvements in PEC performance. [42]Figure 5c shows the long-term current density versus time curve of the ReMoOS photocathode to demonstrate the PEC stability, and the inset of Figure 5c presents a schematic representation of the solar hydrogen generation with the Re-MoOS photocathode.The PEC photocurrent density remained stable throughout 18 h of continuous illumination.An electrochemical impedance spectroscopy (EIS) analysis was conducted to gain further insight into the photogenerated charge transport properties and to specifically examine the charge transfer resistance.Figure 5d and the inset of Figure 5d illustrate the Nyquist plots obtained from EIS measurements conducted under illumination.These Nyquist plots were then fitted using a simplified Randles circuit (inset of Figure 5e), which consists of elements such as the solution resistance (R s ), charge transfer resistance (R ct ), and constant phase elements (C).Previous studies consistently reported that a smaller arc diameter in EIS measurements corresponds to a lower R ct . [28,43]As shown in Figure 5e, MoS 2 exhibited a smaller arc radius, indicating a lower R ct,2 value than that of Si.This observation indicates the improved catalytic activity and electron transport properties of MoS 2 .In particular, the ReMoOS photocathode exhibited a significantly decreased R ct,2 value (≈25.59Ω), as compared with the ReO 3 -MoS 2 (≈39.31Ω), MoS 2 (≈120.1 Ω), and Si (≈1068 Ω) photocathodes.These results suggest that rapid carrier transport and activation of the basal plane can be achieved by substituting Re and O.
To gain insight into the origin of the photocurrent and establish a direct correlation between the PEC photocurrent and H 2 production, further investigations on the evolution of hydrogen were conducted.As shown in Figure 5f, the remarkable enhancement in the hydrogen evolution rate for the ReMoOS photocathode, reaching 206.7 μmol h −1 , was more that 2.4-fold higher than that of the MoS 2 photocathode (83.4 μmol h −1 ).Moreover, Re-MoOS exhibited the highest hydrogen generation rate among the various samples tested, surpassing ReO 3 -MoS 2 (147.4 μmol h −1 ), MoS 2 (83.4 μmol h −1 ), and pristine Si (16.9 μmol h −1 ).These findings correspond with the observed PEC performance, further confirming the efficacy of ReMoOS as a superior photocathode material.Consequently, the hydrogen evolution performance not only implies that the photocurrent is primarily attributed to PEC hydrogen production, but also demonstrates that the produced amount of hydrogen was most significantly enhanced in the Re-MoOS sample owing to the synergetic effect of activating the basal plane and accelerating the photocatalytic reactivity.The average estimated Faradaic efficiency of the ReMoOS photocathode during the hydrogen evolution period was 96%, suggesting that the majority of the photogenerated carriers actively contributed to hydrogen generation without engaging in side reactions.
Density functional theory (DFT) calculations were performed to evaluate the effects of co-doping Re and O on the catalytic function of ReMoOS, by comparing Re-O-doped MoS 2 corresponding to ReMoOS with pristine MoS 2 .The DFT-estimated diagram of the proton adsorption free energy and the morphologies of the proton binding of the calculated models are shown in  ative but small ΔG H * of doped MoS 2 indicates that the state of the adsorbed proton (H * ) is energetically equivalent to that of the proton in the (H + + e − ; 0 eV) and the H in the product (H 2 ; 0 eV); therefore, the HER on doped MoS 2 was kinetically feasible (Figure 5g).Interestingly, the most thermodynamically stable proton binding sites for the three types of catalysts used for the DFT calculations (pristine, non-pair 4ReO doped, and pair 4ReO doped MoS 2 ; Figure 5i) were different.The proton binds to S in pristine MoS 2 and in the vicinity of the anion in doped MoS 2. A charge analysis was conducted to determine the electronic and structural features of each proton-binding site.As shown in Table S1 (Supporting Information), the Bader charge analysis results indicated that the anion in the vicinity of Re was more negatively charged than that in pristine MoS 2.Moreover, the 2D contour diagram shows that the charge density of Re was lower than that of Mo, indicating that Re donates more electrons to the surroundings (Figure S12, Supporting Information).The theoretical results of this study showed that protons prefer to bind to electron-rich sites, such as the anions near Re and O.

Conclusion
This study reports a successful solution-based synthesis strategy for MoS 2 -based 2D multi-layer quaternary semiconductor films co-doped with Re and O.The structural instability of the host film and the co-doping of O atoms enabled the efficient doping of Re with a large lattice mismatch, and the formation of a MoS 2 -derived quaternary semiconductor was demonstrated.Moreover, the quaternary semiconductor exhibited a significant performance improvement over MoS 2 in photodetector and PEC applications.These results not only provide excellent potential for broadband photodetection and PEC applications but also suggest exciting opportunities for expanding the scope of 2D TMD semiconductor synthesis.

Experimental Section
Synthesis of MoS 2 , ReO 3 -MoS 2 , and ReMoOS Layers: The large-scale synthesis of MoS 2 multilayers was performed using a well-designed two-step thermal decomposition of (NH 4 ) 2 MoS 4 as a single-source precursor. [11](NH 4 ) 2 MoS 4 (5 wt.%, ACROS, 99.95%), as a single source precursor, was stirred in ethylene glycol at room temperature for 60 min.The resulting solution was then spin-coated onto hydrophilic-treated substrates at 2000 rpm for 20 s and 3000 rpm for 40 s.After coating, the samples were immediately annealed at 130 °C for 1 min to remove the solvent.The resultant (NH 4 ) 2 MoS 4 films were annealed at 280 °C (first step) by introducing Ar (1000 sccm) at a pressure of 1.5 Torr for 30 min and subsequently annealed at 600 °C (second step) for 30 min to complete the synthesis of pristine MoS 2 .A synthetic procedure identical to that described above was used for the synthesis of the ReMoOS layers; however, ReO 3 powder (0.01 g) was added together with the (NH 4 ) 2 MoS 4 host film.The (NH 4 ) 2 MoS 4 films and 0.01 g of ReO 3 powder were placed in a furnace and annealed at 600 °C by introducing Ar (1000 sccm) at a pressure of 1.5 Torr for 60 min.The ReO 3 -MoS 2 synthesis was implemented using a MoS 2 host crystal under the same procedure.
Material Characterization: The surface morphologies of the samples were characterized using SEM (Hitachi S-4800) and cross-sectional TEM (JEM-ARM200F, JEOL).The microstructural and structural properties of the samples were investigated using micro-Raman spectroscopy at an excitation wavelength of 532 nm with a charge-coupled device detector (inVia Raman microscope, Renishaw).The chemical states and sample compositions were determined using XPS (K-Alpha, Thermo Scientific) equipped with a monochromatic Al K X-ray source (h = 1486.6eV) operating in the constant analyzer energy mode.Each spectrum was deconvoluted using a Gaussian-Lorentzian function and Shirley background subtraction.The optical properties were characterized using UV-vis spectroscopy (Shimadzu UV-2600).
Fabrication of Photodetector and Photoelectrical Properties Measurement: Photodetectors with a Si/SiO 2 /active layer/Cr/Au architecture were fabricated to examine the photoelectrical properties of various active layers (MoS 2 , ReO 3 -MoS 2 , and ReMoOS).Briefly, a 70 nm-thick Au layer was deposited to form the electrical contacts of the devices, and the inserted 5 nm-thick Cr layer served as a functional adhesion and ohmic contact layer.Electrical contacts were deposited by thermal evaporation using a shadow mask, defining channel lengths and widths of 50 and 500 μm, respectively.Two laser sources (MGL-S-532 and MIL-S-1064, CNI Lasers) and a source meter (Keithley 2612 B) were used to evaluate the photoelectric properties of the photodetectors.The incident photon wavelengths were adjusted to 532 nm (5-20 mW cm −2 ) and 1064 nm (230-280 mW cm −2 ).
Photoelectrode Preparation and PEC Measurements: PEC photoelectrodes were fabricated on 1 × 2 cm 2 Si substrates.The p-type (100) Si wafer was washed with 10 vol.% hydrofluoric acid to remove impurities and native oxides, and Ti was deposited on the Si substrate using thermal evaporation at ≈0.8 Å s −1 .The Ti layer (≈10 nm) served as a functional adhesion and protection layer and was attached to the Si substrate through thermal evaporation.Various active layers (MoS 2 , ReO 3 -MoS 2 , and ReMoOS) were formed on the Si/Ti photoelectrodes in the same manner as previously mentioned in the thin-film synthesis section.To achieve ohmic contact, the back sides of the prepared photoelectrodes were scratched using a blade and coated with a Ga-In eutectic alloy.Subsequently, the photoelectrodes were passivated using nonconductive epoxy to fix the working area for subsequent PEC measurements.PEC characterization was performed using a three-electrode system and an electrochemical analyzer (potentiostat/galvanostat 263A).A Pt plate and Ag/AgCl electrode were used as the counter and reference electrodes, respectively.The electrolyte solution comprised 0.5 m H 2 SO 4 .A 150 W Xe arc lamp that delivered an intensity of 100 mW cm −2 was used as the light source, simulating AM 1.5 G irradiation.The current-voltage characteristics were recorded using a source meter (Keithley 2400).The potentials versus the Ag/AgCl reference electrode were converted to potentials versus RHE using the Nernst equation (E RHE = E Ag/AgCl + 0.197 + 0.059 × pH).EIS analyses were performed under constant light illumination (100 mW cm −2 ) at a bias of −0.35 V while varying the AC frequency from 100 kHz to 1 Hz.The H gas products were analyzed using a YL 6500 gas chromatograph (Young In Chromass Co., Ltd.) equipped with flame ionization and thermal conductivity detectors.
Density Functional Theory Calculations: All spin-polarized DFT calculations were performed using the vienna ab initio simulation package (VASP) code and the Perdew-Burke-Ernzerhof functional. [44,45]The projector-augmented wave method describes the interaction between an ionic core and valence electrons. [46]The DFT+U scheme, [47] with U eff = 4 eV, [48] was applied to the Mo ions to appropriately consider the localized Mo-d orbitals.Valence electron wave functions were expanded on a plane-wave basis up to an energy cutoff of 400 eV.The first Brillouin zone was sampled at the Γ point with dimensions of 12 × 12 × 1 for geometry optimization and spin-polarized band structure.The convergence criteria for the electronic structure and atomic geometry were 10 −4 eV and 0.05 eV Å −1 , respectively.To improve the convergence of states near the Fermi level, a Gaussian smearing function with a finite temperature width of 0.05 eV was used to improve the convergence of states near the Fermi level.The DFT-optimized lattice constants of the 2H-phase MoS 2 were a = b = 3.20 Å and c = 13.65 Å (Figure S13, Supporting Information).A (4 × 4) supercell was constructed with a few layers of free-standing MoS 2 to describe the monolayer and bilayer MoS 2 .The vacuum layer was applied at >10 Å to avoid electronic interference between periodic supercells.The electrical band structures of pristine MoS 2 based on the number of layers are shown in Figure S9 (Supporting Information).As layers of MoS 2 were exfoliated down to bilayers or monolayers, the bandgap increased to 1.24, 1.36, and 1.55 eV in the bulk, bilayer, and monolayer MoS 2 , respectively.The bandgaps of the three types of pristine MoS 2 were characterized by a gradual shift from indirect to direct.[51] According to the crystal field theory, the trigonal coordination of Mo-S leads to the splitting of the 4dorbital in 2H-MoS 2 (Figure S14, Supporting Information).Among these separated orbitals, the energy states of d x 2 −y 2 and d xy were ascribed to the semiconducting properties of 2H-MoS 2 . [52,53]o evaluate the HER activity of the pristine and doped MoS 2 , the DFTcalculated H adsorption energy to the Gibbs free energy of H binding

Figure 1 .
Figure 1.a) Schematic depiction illustrating a procedure for the synthetic strategy by the selection of host materials (i) pristine MoS 2 and ii) (NH 4 ) 2 MoS 4 ).Cross-sectional TEM and EDS mapping images of b) ReO 3 -MoS 2 and c) ReMoOS.
2 crystal, X-ray photoelectron spectroscopy (XPS) with conventional monochromatic Al K radiation (hv = 1486.6eV) and a pass energy of 50 eV was used to chemically identify MoS 2 , ReO 3 NPs on MoS 2 , and ReMoOS.The survey spectra acquired from pristine MoS 2 , ReO 3 -MoS 2 , and ReMoOS are presented in Figure S4 (Supporting Information).The survey spectrum of pristine MoS 2 shows the presence of Mo, S, O, and C. Five chemical elements, namely Mo, S, O, C, and Re, were distinctly observed in the survey spectra of ReO 3 -MoS 2 and ReMoOS.Figure 2a-d shows the evolution of the deconvoluted core-level spectra for Mo 3d, Re 3f, S 2p, and O 1s obtained from the various samples (pristine MoS 2 , ReO 3 -MoS 2 , and Re-MoOS).Each spectrum was fitted using a Gaussian-Lorentzian function and Shirley background, enabling a detailed analysis of the individual chemical components and their characteristic peaks.Figure 2a shows the Mo 3d core-level spectra of pristine MoS 2 , ReO 3 -MoS 2, and ReMoOS.The Mo 3d core-level spectrum of pristine MoS 2 exhibited three pronounced peaks attributed to the doublet peaks originating from the Mo 4+ bonding states of Mo 3d 5/2 and Mo3d 3/2 centered at 232.4 and 229.2 eV, respectively, and the S 2s state at 226.4 eV was associated with the trigonal prismatic lattice of MoS 2 .After the formation of ReO 3 NPs on MoS 2 synthesized using the MoS 2 host, doublet peaks correlated with the Mo 6+ bonding states in MoO 3 (Binding energy (E B ) = 232.4and 235.5 eV) were observed.This phenomenon indicated the substitutional doping of O atoms in the trigonal prismatic MoS 2 lattice.Intriguingly, the Mo 4+ bonding states in Mo-S-Re, induced by the substitutional doping of Re atoms, were distinctly observed after the formation of the ReMoOS quaternary semiconductors.The Re 4f and Mo 4p core-level spectra were simultaneously observed in a similar binding energy region, enabling an efficient comparison of the relative atomic concentrations of Re and Mo atoms in the quaternary system.Figure 2b shows the Re 4f core-level spectra obtained from pristine MoS 2 , ReO 3 -MoS 2, and ReMoOS, which could be deconvoluted into distinct peaks attributed to the Re-and Mo-related bonding states.The doublet peaks correlating to the Mo 4p bonding states in pristine MoS 2 were located at E B = 36.4and 38.1 eV.After the formation of ReO 3 -MoS 2 , the Re 4f spectrum exhibited Mo 4p, S-Re-S (E B = 41.4 and 43.8 eV), and O-Re-S (E B = 42.1 and 44.4 eV) originating from the presence of ReMoOS, and ReO 2 (E B = 43.2 and 45.6 eV) and ReO 3 (E B = 45.8 and 48.2 eV) originating from the presence of Re-O NPs.This result indicates that the supplied Re and O atoms saturated the vacancy sites of the MoS 2 lattice and then agglomerated as ReO x (x = 2 and 3) NPs on the MoS 2 surface, which is consistent with the comprehensive microscopic evaluation results.After the formation of ReMoOS, the S-Re-S and O-Re-S states were predominant, and the ReO x (x = 2 and 3) NPinduced bonding states were marginal.Figure 2c shows the S 2p core-level spectra of pristine MoS 2 , ReO 3 -MoS 2, and ReMoOS.The doublet peaks of pristine MoS 2 were located at 162.0 and 163.2 eV, corresponding to the S 2p 3/2 and S 2p 1/2 states of 2H-MoS 2 , respectively.The formation of ReO 3 -MoS 2 causes the S-Re-S bonding states (E B = 162.5 and 163.7 eV) induced by the formation of ReMoOS at a localized area determined by the intrinsic vacancy sites of the MoS 2 crystal.The intensity of the S-Re-S states distinctly increased with the formation of the ReMoOS quaternary semiconductor using (NH 4 ) 2 MoS 4 host films with structural instability.As shown in Figure 2d, the O 1s core-level spectrum of pristine MoS 2 exhibited weak-intensity peaks correlated with the co-existence of an adsorbed hydroxyl group and C─O bonding states from surface residues.In contrast, a distinctive O 1s peak with a higher intensity was discernible in the ReO 3 -MoS 2 and ReMoOS O 1s spectra, which could be deconvoluted into five states associated with MoO 3 (E B = 530.4eV), ReO (E B = 531.0eV), ReO 2 (E B = 531.8eV), and ReO 3 (E B = 532.7 eV) chemical bonding, indicating the formation of ReO x NPs and is direct evidence of atomic substitution of O atoms into MoS 2 crystals.A quantitative analysis of the Re and O in the MoS 2 crystal extracted from the XPS core-level spectra was performed, as shown in Figure 2e.The Re, Mo, O, and S atomic percentages in ReO 3 -MoS 2 and ReMoOS were 5.5, 26.5, 29.6, and 38.4%, and 13.6, 21.8, 15.5, and 49.1%, respectively.The chemical composition of the quaternary semiconductor was estimated from these results to be Re 0.38 Mo 0.61 O 0.48 S 1.51 .It should be noted that the proposed synthesis route enables the facile synthesis of quaternary semiconductors with high concentrations of guest atoms using host films with structural instability.Raman analysis was performed to explore the structural evolution of MoS 2 after the formation of ReO 3 -MoS 2 and ReMoOS, as shown in Figure 2f-h.Two distinct peaks were observed at 382.4 and 406.3 cm −1 in the Raman spectrum of pristine MoS 2 , which was attributed to the representative phonon modes of 2H-MoS 2 .Specifically, the vibration mode at ≈382.4 cm −1 corresponded to the E 1
3 and the potential for plasmon resonance at the interface with MoS 2 are intriguing, a comparison of the absorption spectra of ReO 3 -MoS 2 and ReMoOS without ReO 3 revealed that the contribution of the plasmon resonance of ReO 3 was negligible.Density functional theory (DFT) calculations for the Re-O doping in 2H-phase MoS 2 were performed to gain further insight into the adjustment of the electronic structure induced by doping.Bilayer 2H-MoS 2 was used to investigate the Re-O-induced modulation of the band structure.Moreover, we have constructed two-type models: paired-and non-paired-Re-O doped in MoS 2 , representing the DFT-modeled morphologies of doped MoS 2 to clarify the effect of Re and O doping.The two conformations of Re-O denote whether the Re is in direct contact with oxygen (Figure S7, Supporting Information).Two configurations of the Re-O formed pairs and non-pairs are presented in Figure 3c and Figure S8 (Supporting Information).The pristine bilayer of 2H-MoS 2 exhibited an indirect bandgap of 1.36 eV with the edges of the conduction band (CB) and valence band (VB) located at the Γ-and K-points, respectively (Figure S9, Supporting Information).Although Re-O-doped MoS 2 retained a similar location of the band edges to pristine MoS 2 , the bandgap of the quaternary system (ReMoOS) gradually decreased with the dopant concentrations, regardless of the conformation of Re-O in MoS 2 (Figure 3a,b,d).The variation in the chemical bonding of Re-S, Re-O, and Mo-O emerged at an energy state near the Fermi level, entailing modulation of the band structure.The Re-S coupling became more prominent at the Fermi level with an increasing dopant concentration (Figure S10, Supporting Information).The distinguished energy states of Re-S in ReMoOS downshifted the original VB edge of pristine MoS 2 .
shows the photocurrent dynamics of the devices based on pristine MoS 2 , ReO 3 -MoS 2 , and ReMoOS with varying applied voltages (3, 5, 10, 15, and 20 V) under periodic 532 and 1064 nm laser illumination to evaluate the visible and NIR photodetection performance.The devices were irradiated at the incident laser powers of 20 and 280 mW, corresponding to 532 and 1064 nm, respectively.The insets of Figure 4c-e show the regions of the response time (90% of maximum photocurrent) and recovery time (10% minimum photocurrent) for the various photodetectors at 20 V.

Figure 3 .
Figure 3.The DFT-calculated electronic band structure of doped MoS 2. Series of a) pair Re-O and b) non-pair Re-O doped MoS 2 .The green circles denoted the band edges of valence band maximum (VBM) and conduction band minimum (CBM).c) The morphologies of pristine MoS 2 and Re-O doped MoS 2 , in which the Re-O is incorporated and separated.The blue area is the forbidden energy region to electrons.d) In the plotted reduction of the bandgap of doped MoS 2, the red and blue lines represented pair or non-pair doping of the Re-O unit.The pink arrow indicates that the bandgap is gradually reduced as increasing the concentration of dopants regardless of the Re-O configuration.

Figure 5 .
Figure 5. a) Photocurrent density-potential curves and b) Tafel slopes of various photocathodes (Si, MoS 2 , ReO 3 -MoS 2 , and ReMoOS).c) Photocurrenttime plots for ReMoOS sample.d) EIS plot and e) bar graph of R ct,2 based on: Si and Si/Ti coupled with MoS 2 , ReO 3 -MoS 2 , and ReMoOS.f) Hydrogen production amounts of PEC cells with various working electrodes.g) The hydrogen binding free energy for the models estimated: purple, red, and green lines represent the pristine 2H-phase MoS 2 , pair 4Re-O doped, and non-pair 4Re-O doped MoS 2 , respectively.i) The morphologies of proton binding on each model.The red circles indicate the binding site of protons.

Figure
Figure 5g-i.The Gibbs free energy of proton adsorption can generally act as a vital reaction descriptor representing the HER activity because optimal catalysts for the HER should adsorb adsorbates and reaction intermediates that are neither too strong nor too weak, as stated in the Sabatier principle.Because the pristine MoS 2 exhibited a positive free energy for proton binding, ΔG H * , of 0.46 eV, the reaction was prohibited on pristine MoS 2 .In contrast, the Re-O-doped MoS 2 systems exhibited a negative ΔG H * of −0.16 eV (pair 4ReO) and −0.11 eV (non-pair 4ReO).The neg-