The in‐depth mechanism behind this significant improvement is revealed through a combined experimental and theoretical analysis for QDs/TiO2–Au:CNTs hybrid network and demonstrates the multifaceted impact of plasmonic Au NPs and CNTs: i) hot‐electron injection from Au NPs into CNTs and TiO2; ii) near‐field enhancement of the QDs absorption and carrier generation/separation processes by the plasmonic Au NPs; iii) enhanced photoinjected electron transport due to the highly directional pathways offered by CNTs. These results provide fundamental insights on the properties of QDs/TiO2–Au:CNTs hybrid network, and highlights the possibility to improve the performance of other solar technologies.Here, a facile approach to enhance the performance of solar‐driven photoelectrochemical (PEC) water splitting is described by means of the synergistic effects of a hybrid network of plasmonic Au nanoparticles (NPs) decorated on multiwalled carbon nanotubes (CNTs). The device based on TiO 2–Au:CNTs hybrid photoanode yields a saturated photocurrent density of 16.10 ± 0.10 mA cm−2 [1.0 V vs reversible hydrogen electrode (RHE)]. This approach offers a promising solution to improve the performance of optoelectronic devices.The performance of solar‐driven photoelectrochemical (PEC) devices is enhanced through the synergistic effects of plasmonic Au nanoparticles and multiwalled carbon nanotubes (CNTs). A PEC device based on alloyed core/shell quantum dot sensitized TiO A typical solar‐driven water splitting system exploits two‐half chemical reactions such as the reduction of water to H2 at the cathode and oxidation of water to oxygen (O2) at the anode. For this purpose, the starting point is to optimize the band alignment of a semiconductor photoanode and photocathode with suitable reduction and oxidation potential of the electrolyte. A photon absorbed by the semiconductor photoanode with energy higher than its bandgap generates an electron–hole pair (exciton), which is dissociated at the semiconductor–electrolyte interface and leads to hole oxidation and electron transfer through an external circuit to the cathode for H2 generation.Solar‐driven hydrogen , the absorption of TiO2 is limited to the ultraviolet (UV) spectral region (which represents only a 3–5% of the full solar spectrum). In addition, the low electron mobility (1 cm2 V−1 s−1) and short carrier diffusion length (10–100 nm) of TiO2 result in an overall low efficiency of H2 generation.Following the pioneering work by Fujishima and Honda,2 seve2 anode with colloidal quantum dots (QDs) is a promising approach to extend the absorption spectrum toward the visible and near‐infrared (NIR) region. The properties of QDs such as size/shape/composition‐tunable absorption edge, their high absorption coefficient and the possibility of multiple exciton generation (MEG) by single photon absorption allow to significantly enhance the performance of photoelectrochemical (PEC) cells for H2 generation. Similarly, the efficient transport of photogenerated electrons from the QDs to the collecting electrode surface through the network of wide bandgap semiconductor mesoporous film without quenching at grain boundaries is another key aspect. In this context, the design and synthesis of one‐dimensional (1D) , branched nanorods and hierarchical nanotubes of TiO2 have been extensively studied to further improve performance.Multidisciplinary approaches have been explored to address these challenges. For example, sensitizing the TiO) region.12 The2) is a promising approach to enhance the light harvesting efficiency of solar technologies. These plasmonic NPs work as nanoantennas that localize the impinging light's energy, to then transfer a fraction of it to its environment through different optical and charge‐transfer mechanisms. However, plasmonic effects only originate from nanoscale noble metal crystals, and their proximity to other systems also causes charge trapping, thereby hampers device performance.Recently, it has been shown that combining localized surface plasmon resonance (LSPR) of noble metal  nanoparticles (NPs) with wide‐bandgap semiconductors , graphene, graphene nanoribbons, and graphene oxide may be useful for enhancing charge carrier transport and collection. In particular, CNTs could be the most suitable candidates due to their unique transport and structural properties. The addition of a precise amount of CNTs in TiO2 mesoporous paste boosts the performance of solar technologies thanks to improved electron transport and reduced carrier recombination within the photoanode. Recently, Lee et al. demonstrated that plasmonic metal NPs–CNTs hybrid networks lead to significant suppression of charge trapping and related nonradiative recombination that are typical of plasmonic NPs. To the best of our knowledge, there is no report exploring the use of plasmonic metal NPs–CNTs hybrid networks in colloidal QD based PEC devices for H2 production. In particular, the synergistic effect of plasmonic metal NPs–CNTs in mesoporous wide bandgap semiconductors sensitized with QDs on the light harvesting efficiency and the carrier transport/collection efficiency is as yet not understood. In addition, the optimization of plasmonic metal NPs–CNTs hybrid network is critical to optimize performance in solar energy conversion devices.To address this issue, we propose to use 1D materials capable of promoting the efficient separation and transport of photogenerated charge carriers.23 In 2–Au:CNTs hybrid network based PEC devices for H2 generation. We decorated Au NPs on the surface of CNTs by using a simple hydrothermal approach with different relative ratio of Au and CNTs. Subsequently we prepared TiO2–Au:CNTs hybrid anodes with different amount (wt%) of Au:CNTs hybrid network using the simple and cost‐effective doctor blade approach. A specially designed colloidal alloyed CdSe/(CdSexS1−x)5/(CdS)1 QDs “giant” core/shell QDs (denoted as “g‐QD”) with broad absorption as well as efficient exciton separation were applied as light harvester. The PEC device based on “g‐QD” sensitized TiO2–Au:CNTs hybrid photoanode with an optimized amount of Au:CNTs (0.10:0.014 wt%) yielded a ≈26% higher saturated photocurrent density than that of the control PEC device. This improvement is attributed to enhanced light harvesting efficiency, hot‐electron injection and near‐field enhancement induced by the Au NPs and the improved electron transport provided by CNTs directional path through the mesoporous film. In addition, the long‐term stability of the PEC device based on a QDs/TiO2–Au:CNTs photoanode is better than the PEC devices based on QDs/TiO2 photoanode.Here we demonstrate for the first time a facile synthetic approach for the fabrication of QDs/TiO22.1Figure  Selected area electron diffraction (SAED) pattern of Au NPs is displayed in Figure Furthermore, energy‐dispersive X‐ray spectroscopy (EDS) measurements of the Au:CNTs hybrid network confirms the presence of Au and C (from CNTs) Figure .2 paste and deposited on the fluorine doped tin oxide (FTO) by using doctor blade approach followed by annealing at 500 °C for 30 min. Figure 2–Au:CNTs hybrid anode with Au NPs highlighted by white dotted circles and CNTs by white dotted lines. The HR‐TEM of the TiO2–Au:CNTs hybrid anode shown in Figure 2 confirms the anatase phase of TiO2 and 3.40 Å for CNTs and 2.35 Å for Au. There is no change in the size, morphology, and crystal structure of both Au NPs and CNTs during the preparation of TiO2–Au:CNTs hybrid anode, confirming the thermal and structural stability of CNTs and Au NPs. Figure 2–Au:CNTs hybrid anode, which confirms the presence of Au, C (from Au:CNTs) and Ti, O (from TiO2). Raman spectroscopy was used to further investigate the nature of interaction of cationic PEI coated Au NPs with CNTs. Figure −1, respectively. The later bands are most commonly observed peaks for the carbonaceous materials. The Raman spectrum of the Au:CNTs hybrid network shows a slightly shift in G bands from 1578 to 1583 cm−1, and D band from 1345 to 1347 cm−1, which confirms binding of PEI coated Au NPs to CNTs  is accepted as an indication of binding at the sidewalls or changes in the defects of CNTs. The ID/IG ratio was reduced from 1.23 to 0.98 for pristine CNTs to Au:CNTs, respectively (see Table S1 in the Supporting Information). This suggests that PEI coated Au NPs are attached to the sidewall of the CNTs via physisorption and electrostatic interaction as expected.The Au:CNTs hybrid network solution was then mixed with an appropriate amount of TiOaterials.28 Thes Figure . The rel2, TiO2–CNTs, TiO2–Au, and TiO2–Au:CNTs hybrid mesoporous film deposited on the FTO glass substrate are shown in Figure S3 in the Supporting Information. The inset of each figure displays a high‐resolution SEM image of the corresponding mesoporous films. The CNTs and Au NPs are hard to distinguish within the hybrid mesoporous films due to the relatively low concentration of CNTs (0.014 wt%) and Au (0.010 wt%) and their conformal coverage by the TiO2 NPs. However the TiO2–Au:CNTs hybrid mesoporous film exhibits cracks at higher concentration of the Au:CNTs (0.015:0.014) wt%  thereby hindering electron transport within the mesoporous film and acting as recombination sites for the photogenerated carrier at the electrolyte/FTO interface. The absorption spectra of bare TiO2 and TiO2–Au, and TiO2–Au:CNTs hybrid mesoporous films with different content of Au:CNTs (wt%) are shown in Figure S4 in the Supporting Information. Although the content of Au NPs (0.010 wt%) is quite low and conformal coverage of Au NPs by the TiO2 NPs, there is still a small photoabsorption in 500–600 nm range in TiO2–Au, and TiO2–Au:CNTs hybrid mesoporous films due to the surface plasmonic effect of Au NPs, which is absent in bare TiO2.The morphology of the mesoporous anode is a crucial component for PEC device performance. Scanning electron microscopy (SEM) images of TiO2 and TiO2–Au, TiO2–CNTs and TiO2–Au:CNTs hybrid mesoporous films are shown in Figure θ = 25.3°, 37.9°, and 48.1° of bare TiO2 and TiO2–Au, TiO2–CNTs, and TiO2–Au:CNTs hybrid mesoporous films, which correspond to the (101), (004) and (200) crystal planes of TiO2 (JCPDS No. 21‐1272). This confirms the anatase phase of TiO2 and TiO2–Au:CNTs hybrid mesoporous anodes and consistent with previous work. However, it is difficult to distinguish the characteristic peak of CNTs (2θ = 26.0°) from the TiO2–CNTs and TiO2–Au:CNTs hybrid mesoporous anodes due to the overlap with the characteristic peak of anatase TiO2 (2θ = 25.3°). There is a characteristic peak of Au NPs observed at 2θ = 38.6° both for TiO2–Au and TiO2–Au:CNTs hybrid anodes corresponding to (111) plane, which is consistent with SAED measurements of Au NPs and confirms the presence of Au NPs.X‐ray diffraction (XRD) patterns of bare TiO2, TiO2–Au, TiO2–CNTs, and TiO2–Au:CNTs hybrid photoanodes. There is a peak at 1120 cm−1, which corresponds to the presence of Au NPs in TiO2–Au hybrid photoanode (lower image). Similarly, the presence of CNTs in TiO2–CNTs hybrid photoanode is confirmed by the D and G bands at 1350 and 1580 cm−1, respectively, which are the specific bands position for the carbonaceous materials. The Raman spectra of TiO2–Au:CNTs hybrid anodes show the Au peak at 1120 cm−1 and the D and G bands at 1350 and 1580 cm−1, respectively, which indicate the presence of Au and CNTs. These results demonstrate that the Au NPs and CNTs did not undergo any structural changes while annealing the TiO2–Au, TiO2–CNTs, and TiO2–Au:CNTs hybrid photoanodes at 500 °C for 30 min under ambient conditions. The relative peak intensities of the D and G bands are lower compared to the same peaks before annealing. This lower intensity of D and G bands is due to a small amount of Au:CNTs hybrid network in TiO2 NPs mesoporous film and conformal coverage of CNTs by TiO2 NPs.Figure aterials.28b Th7/2 and f5/2 at 84.6 and 88.2 eV, respectively. A slight shift was found in the Au 4f characteristic peaks toward lower binding energy (BE) in the Au:CNTs hybrid network such as Au 4f7/2 and f5/2 at 83.8 and 87.4 eV, respectively. This shift confirms the interaction between the Au NPs and CNTs and the metallic nature of the Au NPs. At the same time, N 1s BE peaks of Au–PEI, which located at 400.06 and 401.74 eV with a ratio of 13.28 are slightly shifted to 400.43 and 401.97 eV with a ratio of 4.93 in the case of the Au:CNTs hybrid network . A higher BE peak is usually associated with protonation and/or binding, so decreasing ratio suggests a significant electron donation from N to CNTs . This observation is consistent with the Raman measurement of Au:CNTs hybrid network  spectra of the Au NPs and Au:CNTs hybrid network are shown in Figure e Au NPs.33 At k Figure . The XPS2.2 In brief, CdSe core QDs with average diameter of 3.33 ± 0.02 nm were synthesized via a hot injection approach. A typical TEM image and size distribution of as‐synthesized CdSe core QDs are shown in Figure S5a,b in the Supporting Information. CdSexS1−x graded alloyed interfacial layers were further grown over the CdSe core QDs by tailoring the Se:S ratio during in situ growth of each CdSexS1−x (x = 0.9–0.1) via a successive ionic layer adsorption and reaction (SILAR) at 240 °C under N2 flow. The final diameter of the “g‐QDs” after the growth of five graded alloyed interfacial layers of CdSexS1−x (x = 0.9–0.1) with one monolayer of CdS reaches 7.75 ± 0.10 nm (estimated from the TEM image) . In the bilinker approach, a bifunctional molecule is linked between the QDs to surface TiO2 NPs mesoporous film and the electron transport rate at the QD/TiO2 interface is directly related to the length of bifunctional molecule. In the EPD approach, colloidal QDs are directly anchored to surface of the TiO2 NPs mesoporous film, yielding enhanced electron transfer rate compared to bilinker approach as confirmed in our previous reports. Therefore, we used the EPD approach to deposit the as‐synthesized graded alloyed core/shell QDs into the TiO2 NPs mesoporous film. To verify the homogenous distribution of colloidal QDs into the TiO2 mesoporous film, we carried out EDS measurements. Figure 2 NPs mesoporous film and EDS mapping of the selected green rectangular area are reported in Figure 2 NPs mesoporous film coated with ZnS and SiO2 confirm the homogeneous distribution of “g‐QDs” within the TiO2 NPs mesoporous film and confirm the presence of Ti, O, Cd, Se, S, Zn, and Si. In addition, homogenous dispersion of the colloidal “g‐QDs” on the surface of the TiO2 NPs is also confirmed from the HR‐TEM images of QD sensitized TiO2 NPs mesoporous film. Figure 2 NPs mesoporous film, which confirms that there is no aggregation of QDs during their deposition via EPD. The EDS spectra of “g‐QD” sensitized TiO2 NPs mesoporous film coated with ZnS and SiO2 also confirm the presence of Ti, O, Cd, Se, S, Zn and Si .37 In i Figure .2.32 and TiO2–Au:CNTs hybrid mesoporous film with different amounts (wt%) of Au:CNTs hybrid network as working electrode to highlight the effect of Au:CNTs. The current density measurements were carried out under dark, chopped and continuous 1 sun light illumination  by using a typical three electrodes configuration system with Platinum as the counter electrode and Ag/AgCl as the reference electrode. Figure 2–Au:CNTs hybrid photoanodes with different relative concentration of Au and CNTs. The calculated saturated photocurrent densities of all PEC devices at 1.0 V versus RHE are reported in Table 2 mesoporous film with varying weight ratios of Au:CNTs (0.0:0.0 to 0.15:0.014 wt%). In brief, the highest saturated photocurrent density of the PEC device based on QDs/TiO2 photoanode under 1 sun illumination (100 mW cm−2) is 8.20 ± 0.11 mA cm−2  wt% due to enhanced electron transport and collection efficiency of the photoinjected electrons by the directional path of CNTs as confirmed in our previous report , while maintaining the same content (0.014 wt%) of CNTs, which is ≈40% higher than that of the PEC device based on the bare QDs/TiO2 photoanode . In addition, the presence of this small amount of Au NPs (0.10 wt%) decorated CNTs (0.014 wt%) in TiO2 NPs mesoporous film does not alter the structural morphology and optical transparency, which is confirmed by SEM images  and UV–visible absorption .Solar‐driven PEC devices for hydrogen generation were fabricated by using colloidal “g‐QD” sensitized TiO2 Figure , which it Figure .26]262 an2 Figure , then ree Figure . This si−2 , which act as recombination center during device operation and reduce overall electron transport. In addition, it starts partially absorbing the light that is available for exciton generation, which reduces the overall performance of PEC devices. To highlight the effect of Au NPs, the PEC device was also fabricated using QDs/TiO2–Au with 0.10 wt% of Au NPs in TiO2, yielding a saturated photocurrent density of 10.60 ± 0.15 mA cm−2 under 1 sun illumination, which is ≈29% and ≈10% higher than that of the PEC devices based on QDs/TiO2 and QDs/TiO2–CNTs (0.014 wt%), respectively. The variation of the saturated photocurrent density of PEC devices based on QDs/TiO2–Au:CNTs hybrid photoanodes with varying content (wt%) of Au NPs and CNTs at 1.0 V versus RHE under 1 sun illumination is shown in Figure However, if we further increase the concentration of Au:CNTs (0.15:0.014 wt%), the saturated photocurrent density decreases from 11.50 ± 0.24 to 10.15 ± 0.11 mA cm2 Figure . This is2–Au:CNTs photoanodes with Au:CNTs of 0.10:0.014 wt% yields the highest saturated photocurrent density compared to all other types of QDs/TiO2–Au:CNTs photoanodes with different relative content (wt%) of Au and CNTs.Thus, the PEC device based on the QDs/TiO2.42–Au:CNTs hybrid and QDs/TiO2 photoanodes, we used transient photovoltage decay spectroscopy. To measure the transient photovoltage decay, both PEC devices based on QDs/TiO2–Au:CNTs hybrid and QDs/TiO2 photoanodes were illuminated under 1 sun simulated sunlight  until reaching a steady voltage and then shut down and recoded the voltage decay with time under dark conditions. Figure Voc) response of PEC devices based on QDs/TiO2 and QDs/TiO2–Au:CNTs hybrid photoanodes under illumination followed by measurements without illumination. Under illumination, the photogenerated electron injected to the conduction band (CB) of TiO2 from the photoexcited QDs leads to the shift of the Fermi level to more negative potential and enhances the Voc. When the electron accumulation equal to charge recombination and then the Voc value reaches to the maximum value and attains a steady state. Under dark conditions, the Voc decay rate corresponds to the recombination of the accumulated electrons with electron acceptor in the electrolyte and within the TiO2 network.To gain further insights into the carrier transport properties of QDs/TiOVoc rate is faster for the PEC device based on QDs/TiO2 photoanodes than the PEC device based on QDs/TiO2–Au:CNTs hybrid photoanodes  of the corresponding PEC devices was calculated from Voc decay rate by using the following equationkB is Boltzmann's constant, T is the absolute temperature, and e is the electron's charge.The s Figure . The eleτe depends on the nonradiative carrier recombination at the photoanode/electrolyte interface. The calculated τe value is higher for PEC devices based on QDs/TiO2–Au:CNTs hybrid photoanode than that of PEC devices based on the QDs/TiO2 photoanode . An equivalent circuit composed of the series resistance (Rs) followed by the series of the charge transfer resistance (Rct) and Warburg element and then parallel by a constant phase element (CPE) was used to fit the mid‐frequency range data . Figure 2–Au:CNTs and TiO2 photoanodes. The semicircle of the Nyquist plots defines the carrier dynamics at the photoanode/electrolyte interface. The diameter of the semicircle for a PEC device with TiO2–Au:CNTs is smaller than that of PEC device with TiO2, which confirms reduced Rct for the photogenerated carrier in TiO2–Au:CNTs. Thus the incorporation of a precise amount of Au:CNTs hybrid network (0.10:0.014 wt%) in TiO2 mesoporous film reduces the Rct compared to TiO2 as well. The Rct values of PEC devices based on TiO2–Au:CNTs and TiO2 photoanodes were calculated by fitting the Nyquist plots at different bias potential versus RHE. At 0.6 V versus RHE, Rct of the PEC device based on the TiO2–Au:CNTs (Rct = 6.3 kΩ) is lower than that of PEC device based on TiO2 (Rct = 9.9 kΩ), which confirm the better electron transport and injection rate in TiO2–Au:CNTs compared to the TiO2 photoanode.In addition, EIS measurements were carried out to understand the effect of Au:CNTs hybrid network on the carrier dynamics within TiO2–Au:CNTs hybrid and TiO2 anodes as an electron and electrolyte containing 0.25 m sodium sulfide (Na2S) and 0.35 m sodium sulfite (Na2SO3) (pH = 13) as hole scavengers. QD/ZrO2 photoanodes were used as a reference, in which there is no electron transfer from QDs to ZrO2 due to the unfavorable electronic band alignment of QDs and ZrO2. For efficient carrier transfer of the photogenerated carriers in the QDs to their scavengers depends on the optimization of the band alignment of the conduction band edge of the QDs with the conduction band edge of the metal oxide (TiO2–Au:CNTs hybrid and TiO2 anodes) and the redox potential of the sulfur‐based electrolyte with the valence band of the QDs. In our previous reports, we confirmed that there is no possibility of the energy transfer from the QDs to the carrier scavengers as there is no spectral overlap between the PL of the QDs and absorption of the carrier scavengers. The PL decay of the QDs was measured under pulse laser of 444 nm and all curves were fitted using an exponential decay. The average lifetime (〈τ 〉) was extracted by using the following equationia  are the fitting coefficients of the PL decay and τi  are the characteristic lifetimes, respectively.In addition, we used transient PL spectroscopy to investigate the carrier dynamics of QDs coupled with TiOavengers.42 The2, electrolyte/QDs/ZrO2, QDs/TiO2–Au:CNTs and QDs/TiO2. These results demonstrate that the PL decay is faster for QDs/TiO2–Au:CNTs than the QDs/ZrO2 due to different band offset potential between the QDs and ZrO2 or TiO2–Au:CNTs. This confirms the efficient transfer of the electrons from the QDs to the TiO2–Au:CNTs with suitable band alignment. Similarly, the PL decay of QDs with ZrO2/electrolyte is faster than only ZrO2 due to efficient transfer of the holes from QDs to the electrolyte in suitable band alignment of the QDs with the redox potential of the electrolyte  is mainly due to enhanced carrier extraction and injection kinetics by plasmonic‐electrical effect of Au NPs. In QDs/TiO2–Au:CNTs hybrid system, Au NPs induce the local field enhancement that improves not only light absorption, but also leads to better carrier separation and transport. In addition, the improved injection kinetics is also attributed to slight negative shift in the CB minimum (CBM) with the incorporation of Au:CNTs hybrid network in TiO2 as confirmed by the ultraviolet photoelectron spectroscopy (UPS) measurements.Figure e Figure . To high2 Figure . The PL Ket or Kht) of QDs coupled with the different carrier scavengers were calculated from the difference in the lifetime values of the corresponding carriers, by using the following equationτ〉QDs/e(TiO2 − Au: CNTs) or h(electrolyte) and 〈τ〉QDs/ZrO2 are the average lifetimes of the QDs with carrier scavengers and QDs with ZrO2, respectively. The calculated electron transfer rate for QDs/TiO2–Au:CNTs (Ket = 4.4 ± 0.8 × 107 s−1) is faster than that of QDs/TiO2 (1.4 ± 0.08 × 107 s−1) and hole transfer rate (Kht = 2.1 ± 0.05 × 107 s−1) values, confirming efficient transfer from the QDs to the carrier scavengers (TiO2–Au:CNTs and electrolyte).The carrier transfer rates . UPS with He I radiation (21.22 eV) was used to estimate the Fermi level and valance band maximum (VBM) level energy of TiO2 and TiO2–Au:CNTs. Figure S7b in the Supporting Information displays full UPS spectrum of both TiO2 and TiO2–Au:CNTs with high and low binding energy cut‐off, which determines the Fermi level and VBM level energy, respectively. The CBM level value for TiO2 and TiO2–Au:CNTs can be calculated from the VBM and bandgap values . These results show that the bandgap of TiO2 is slightly reduced with the incorporation of the Au–CNTs hybrid network consistent with the combined effect of Au and CNTs, and having negative shift (eV) in the CBM level. This shift has positive effect on electron injection rate from the CB of QDs to TiO2–Au:CNTs and suppresses the charge recombination during the device performance, which is consistent with transient PL spectroscopy and photovoltage decay measurements  spectra were acquired to further investigate the effect of the incorporation of Au:CNTs hybrid network on the electronic bandgap alignment of TiOwn value,43 wheIn summary, the results of the carrier dynamics measurements are consistent with the obtained photocurrent density of PEC devices based on the respective photoanodes.2.5Figure 2 occurs at longer wavelengths than those covered by the absorption band of TiO2, their main contribution to the photocurrent will be the injection of plasmonic hot electrons . Briefly, the mechanism responsible for introducing these additional carriers into the generated photocurrent proceeds in two steps, carrier excitation in the metal and injection through the metal–semiconductor interface . Given that the plasmon resonance of the small Au NPs surrounded by TiOs Figure 6a.[46e Figure . HoweverEb. This magnitude, with units of 1/s, also depends on the Fermi energy of the metal, EF, the photon's energy, ℏω, and the integral over the NPs surface, SNC, of the squared absolute value of the complex amplitude of the field normal to the inner surface of the metal particle.We can calculate the number of high‐energy hot electrons generated in a given plasmonic nanostructure by using the formalism described in refs. , 49, whi2 and CNTs. In the context of this model, these two semiconductors will induce different injection rates by virtue of their different Schottky barrier heights when interfacing with the Au NPs. Thus, and as schematically depicted in Figure 2 interface is modelled with ΔEb = 1 eV, while the Au/CNT is estimated as ΔEb = 0.5 eV. A lower energy barrier at the Au/CNT interface would thus lead us to expect an increase in the plasmon‐induced photocurrent if we compare a system including CNTs with one in which the Au NPs only interacts with TiO2. The quantitative expression of this concept is displayed in Figure In our case, the injection of plasmonic hot electrons occurs at the interfaces with two different semiconducting systems, TiO2 and CNTs is greatly surpassed in magnitude by the sensitization of the g‐QDs. Nonetheless, the data in Figure  This contribution to the QD photoabsorption can be interpreted as the Au NPs increasing the photonic density of states at the neighboring QDs at the Au NPs’ resonant frequency, thus increasing the probability of light absorption by the QDs. Supporting the interpretation that this is what occurs in our QD‐sensitized photoanode, we show that the plasmon band of 10 nm Au NPs immersed in mesoporous TiO2 overlaps with the absorption spectrum of the QDs. This is depicted in Figure V is the volume of the space surrounding the Au NP , E0 is the electric field amplitude of the incoming radiation, and E(r) is the electric field at a given point in space r. Finally, the Au NPs in this system do not contribute to increasing the photon travel path in the photoanode via scattering events, as it is the case for other designs aiming to optimize solar capture. This is so because the plasmon excited in small NPs decay almost exclusively nonradiatively, as shown by results on the interaction cross sections of 10 nm Au NPs in TiO2 .In the case of the system sensitized with QDs, the previous perspective still holds, but the injection of plasmonic hot electrons into TiOn Figure 52 Thi2.6): i) 300 nm thin layer of 20 nm sized TiO2 NPs spin‐coated between the blocking and active layer; ii) 4–5  µm scattering layer of 150–200 nm sized TiO2 NPs deposited over the active layer. The PEC device based on QDs/TiO2–Au:CNTs (0.10:0.014 wt%) and QDs/TiO2 photoanodes with spin‐coated and scattering TiO2 layer exhibits a higher saturated photocurrent density than that of the respective configuration of photoanodes without spin‐coated and scattering TiO2 layers. Figure 2 and QDs/TiO2–Au:CNTs (0.10:0.014 wt%) hybrid photoanodes with spin‐coated and scattering layers under dark, chopped and continuous 1 sun light illumination . The calculated saturated photocurrent density of the respective PEC devices at 1.0 V versus RHE under 1 sun illumination is reported in Table −2 for QDs/TiO2–Au:CNTs (0.10:0.014 wt%), which is ≈26% higher than the corresponding values obtained from the control device (QDs/TiO2). This significant improvement in the saturated photocurrent densities is mainly attributed to: i) improved adhesion between the compact blocking layer and active layer of TiO2 NPs with the addition of spin coated TiO2 NPs layer; ii) enhanced light scattering with the addition of TiO2 scattering layer, which increases the time span by the light within the active layer and the possibility of better light absorption by the QDs. To the best of our knowledge, the highest saturated photocurrent density of 16.10 ± 0.10 mA cm−2 (at 1.0 V vs RHE) under 1 sun illumination  obtained in this work is higher than the photocurrent density value reported for PEC devices based on hybrid network of plasmonic NPs and wide‐bandgap semiconductor nanostructures sensitized with QDs. For instance, Gong and co‐workers fabricated a PEC device based on TiO2 branched nanorod arrays decorated with Au NPs and reported a photocurrent density of 2.32 ± 0.1 mA cm−2. Wu and co‐workers designed a sandwich‐structured CdS QD sensitized Au–TiO2 hybrid nanorod array photoanode and reported the highest photocurrent density of 4.07 mA cm−2. Later, Wang and co‐workers fabricated a PEC device by exploring novel Au@TiO2/Al2O3/Cu2O (p‐n junction) configuration and reported a photocurrent density of −4.34 mA cm−2, which is ≈20 times higher compared to that obtained from a TiO2‐P25/Cu2O photoelectrode. A significant improvement in photocurrent density is obtained by replacing the TiO2 mesoporous film with a novel matchlike ZnO nanostructure decorated with plasmonic Au NPs. The resulting PEC device based on ZnO/Au heterostructures yields a photocurrent density of 9.11 mA cm−2, which is much higher than that of the pristine ZnO nanorods array (0.33 mA cm−2). In addition, our methodology adopted in this work to improve the performance of the PEC device is simple, low‐cost, and large‐area scalable compared to the above‐mentioned literature.To further improve the performance of PEC devices, the photoanode architecture was optimized by introducing two additional layers (details reported in our previous work26): i2–Au:CNTs and QDs/TiO2 measured at 0.6 V versus RHE under continuous 1 sun illumination  are displayed in Figure 2–Au:CNTs maintaining 64% of its initial value of photocurrent density after 2 h of continuous illumination, whereas the PEC device based on QDs/TiO2 maintained only 50%. This enhancement in long‐term stability of the PEC device with the incorporation of Au:CNTs hybrid network in TiO2 mesoporous film is ascribed to several physical and mechanical phenomena occurring in the TiO2 mesoporous film. As discussed in previous reports, under 1 sun illumination, mesoporous TiO2 films undergo structural defects, which affects the carrier‐transport properties and hence reduces the overall long‐term stability of PEC devices. In this TiO2–Au:CNTs hybrid mesoporous film, CNTs partially absorb the UV light and act as blocking agent as well as provide better connectivity between the TiO2 NPs, whereas the presence of Au NPs also reduces carrier recombination. These promising synergetic effects of CNTs and Au NPs lead to an improved long‐term stability of QDs/TiO2–Au:CNT based PEC devices compared to QDs/TiO2.The long‐term stability of the QD‐based PEC devices is an important factor to be considered for commercialization. The photocurrent density versus time curves of PEC device based QDs/TiO reports,56 und2–Au:CNTs hybrid photoanodes were used as photoanodes to fabricate QD sensitized solar cells (QDSCs). Figure S10a in the Supporting Information displays the J–V curves of QDSCs based on QDs/TiO2 and QDs/TiO2–Au:CNTs hybrid photoanodes under 1 sun illumination ; the calculated photovoltaic parameters of QDSCs are reported in Table S3 in the Supporting Information. QDSCs based on QDs/TiO2–Au:CNTs hybrid photoanodes yield the PCE of 3.56%, which is 40% higher than the PEC of QDSCs based on the QDs/TiO2 fabricated under identical experimental condition and components (Pt counter electrode and iodide redox couple electrolyte).In addition, QDs/TiO2–Au:CNTs/electrolyte interface by transient photovoltage decay. As shown in Figure S10b in the Supporting Information, Voc decay is faster in the QDSC based on QDs/TiO2 than that of QDSC based on QDs/TiO2–Au:CNTs. The calculated τe by using Equation , the decay rate of the calculated τe for QDSC based on QDs/TiO2 is higher than that of τe for QDSC based on QDs/TiO2–Au:CNTs, which is consistent with the obtained photovoltaic performances of the respective QDSCs . This suggests the versatility of TiO2–Au:CNTs hybrid anode for emerging solar technologies.We also studied the carrier dynamics at the QDs/TiO32 generation. PEC devices based on QDs/TiO2–Au:CNTs hybrid photoanode with optimized amount of Au:CNTs (wt%) yields a saturated photocurrent density of 16.10 ± 0.10 mA cm−2 (at 1.0 V vs RHE) under standard 1 sun illumination , which is ≈26% higher than the control PEC device. Carrier dynamics studies highlight that the addition of a precise amount of Au:CNTs hybrid network in the TiO2 mesoporous films enhances the electron lifetime as well electron injection rate and reduces the Rct compared to bare TiO2 mesoporous films. This is consistent with the obtained photocurrent density of PECs devices based on the respective photoanodes. In addition, we presented theoretical calculations supporting the role of the Au NPs in increasing the device photocurrents by injecting hot electrons into the TiO2–CNT hybrid matrix and, primarily, enhancing the “g‐QD” absorption rates, while showing that their size is too small to contribute as scattering centers in the photoanode.In conclusion, we demonstrated that the synergetic effects of plasmonic Au NPs and CNTs to boost the performance of PEC device for HThese results provide fundamental insights about the hybrid network of QDs/Au:CNTs and offer a promising solution to improve the performance of PEC and other optoelectronic devices. Future directions will focus on the optimization of Au:CNTs hybrid networks with different size and shape of Au NPs  to understand the dynamics of hot‐electrons and to further improve the performance of PEC water splitting and other optoelectronic devices. The hybrid anode could be used for other optoelectronic devices, such as dye‐sensitized solar cells (DSSCs), QDSCs, photocatalytic water splitting, etc.42S, sodium hydroxide, Na2SO3, toluene, methanol, acetone, ethanol, and isopropanol (IPA) were obtained from Sigma‐Aldrich Inc. Au (III) chloride solution 30 wt% in dilute HCl (HAuCl4) was purchased from Aldrich (Germany). Titania paste (code18NR‐AO) consisting of a blend of active anatase particles (20 nm in diameter) and larger anatase scattering particles (up to 450 nm in diameter) was supplied by Dyesol . FTO coated glass substrates with sheet resistance of 10 Ω square−1 were purchased from Pilkington glasses. All chemicals were used as received.Sulfur (100%), oleylamine (OLA) , cadmium oxide (99%), cadmium nitrate tetra hydrate (≥99%), oleic acid (OA), Rhodamine 6G and octadecene (ODE), selenium pellet (≥99.999%), trioctyl phosphine oxide (TOPO), trioctyl phosphine (TOP) (97%), hexane, zinc acetate dihydrate (98%), PEI, , Na Briefly, 6.8 × 10−2 × 10−3m HAuCl4 was mixed with 10 mL 25 kDa bPEI solution at a concentration of 10 mg mL−1. The reaction pH was adjusted to 9.5 with 0.5 m sodium hydroxide and was mixed for 2 min at room temperature. Then, the reaction mixture was placed in an oil bath at 95 °C and stirred with a mechanical stirrer at a constant speed of 350 rmp for 35 min. After cooling down the solution to room temperature, Au NPs were washed with deionized water with 30 kDa molecular weight cut‐off ultracentrifuge filter at 4500 rpm. The resulting GNP dispersions were stored in a vial covered with aluminum foil and kept at 4 °C.Au NPs coated with bPEI were synthesized as reported previously.59 Bri−1) and CNTs dispersion . After 1 h of ultrasonication, mixture of Au NPs and CNTs dispersion was heated to 90 °C for overnight, until the whole solvent was evaporated. Then, the precise amount of ethanol was added and again ultrasonicated for 1 h.Au:CNTs hybrid network solutions with different relative ratios of Au NPs and CNTs were prepared by hydrothermal synthesis. Briefly, a homogeneous dispersion of CNTs (an average length of 10 µm) in ethanol was prepared by mixing 6 mg of CNTs in 15 mL of ethanol and sonicated for 3 h. Then, a precise amount of aqueous solution of Au NPs was mixed with ethanolic dispersion of CNTs with varying relative volume contents of Au  and large (up to 450 nm in diameter) size anatase particles (18NR‐AO) to prepare TiO2–Au:CNTs hybrid pastes. Subsequently, the above prepared TiO2–Au:CNTs hybrid pastes with different wt% of Au:CNTs, were deposited on top of the compact TiO2 layer by tape casting. A drying process was followed for 15 min at ambient conditions and then placed on a hot plate for 6 min at 120 °C. A second layer was then deposited on the top, following the same procedure. All the photoanodes were then annealed at 500 °C for 30 min under ambient conditions. For a systematic comparison, bare TiO2 photoanodes were also prepared under the same conditions. The thickness of all photoanodes was measured using a profilometer.A thin and compact TiO2–Au:CNTs hybrid and TiO2 mesoporous films grown on FTO substrate were vertically immersed in the QDs dispersion in such a way that the deposited films were facing each other. The distance between them was adjusted at around 1 cm and a direct current (DC) bias of 200 V was applied for 2 h. To wash off unabsorbed QDs after the EPD process, the samples were rinsed several times with toluene and dried with N2. Prior to ZnS capping, the photoanodes underwent three SILAR cycles of ligand exchange of methanolic solution cetyl‐trimethyl ammonium bromide (CTAB) and toluene for 1 min dipping. After CTAB treatment, 1 min dipping in methanol was applied to wash and remove the chemical residuals from the surface and dried with N2. Then, 1 min dipping in toluene was applied and dried with N2 to complete one SILAR cycle. Then, ZnS capping layer was applied by using the SILAR. In a typical SILAR deposition cycle, first Zn2+ ions were deposited from a methanolic 0.1 m solution of Zn(OAc)2. Then, sulfide precursor was 0.1 m solution of Na2S in the mixture of methanol and water. One complete SILAR cycle consisted of 1 min of dip‐coating the TiO2 photoanode into the Zn2+ precursors, and subsequently into the S2− solutions. After each bath, to remove the unabsorbed chemical residuals, the photoanode was thoroughly rinsed with corresponding solvent (methanol or mixed solution), respectively, and dried with N2. There were two SILAR cycles applied to form the ZnS capping layer.TiOα radiation source (λ = 0.15 418 nm). TEM and HR‐TEM images of Au, Au:CNTs, and TiO2–Au:CNTs film were collected by using a JEOL 2100F TEM system. Field emission scanning electron microscopy  was applied to study the morphology of the TiO2 and TiO2–Au:CNTs hybrid photoanodes. The chemical composition mapping of QD sensitized TiO2 photoanodes was carried out by EDS. The Raman spectra of the photoanode were measured with a 514 nm excitation source. The UV–vis absorption spectra were recorded with a Cary 5000 UV‐Vis‐NIR spectrophotometer (Varian) with a scan speed of 600 nm min−1. Fluorescence spectra were taken with a Fluorolog‐3 system (Horiba Jobin Yvon). The PL lifetime of the QDs was measured in the time‐correlated single‐photon counting (TCSPC) mode with a 444 nm laser. XPS was performed in a VG Escalab 220i‐XL equipped with hemispherical analyzer, applying a Twin Anode X‐Ray Source. UPS measurement was performed on a VG ESCALAB 3 Mark II high vacuum system.Small angle XRD of extensively purified QDs was carried out with a Philips X′pert diffractometer using a Cu K2–Au:CNTs photoanode as a working electrode, a Pt counter electrode, and a saturated Ag/AgCl reference electrode. An insulating epoxy resin was used to cover the sample's surface except for the active area, to avoid any direct contact between the electrolyte and the conducting back‐contact and/or the connecting wire. Subsequently, the sample was fully immersed in the electrolyte containing 0.25 m Na2S and 0.35 m Na2SO3 (pH ≈13) as the sacrificial hole scavenger. All potentials, measured with respect to the reference electrode of Ag/AgCl during the PEC measurements, were converted to scale according to the following equation VRHE = VAg/AgCl + 0.197+ pH × (0.059). The photoresponse was measured by using a 150 W Xenon lamp as a light source with an AM 1.5G filter. The sample was placed at a distance of 2 cm from the lamp case . Prior to each measurement, light intensity was monitored by a thermopile and adjusted to 100 mW cm−2. All the current versus potential measurements were carried out at a 20 mV s−1 sweep rate.The PEC performance of the photoanodes was evaluated in a typical three‐electrode configuration, consisting of a QDs/TiOElectrochemical impedance spectroscopy (EIS) was carried out under dark conditions by using a SOLARTRON 1260 A Impedance/Gain‐Phase Analyzer with a bias voltage from 550 to 750 mV. The tests were recorded over a frequency range between 10 mHz and 300 kHz, and the AC signal was 10 mV in amplitude. All impedance measurements were analyzed using an appropriate equivalent circuit model with Z‐View software .The authors declare no conflict of interest.Supporting InformationClick here for additional data file.


This article reports a compatible strategy to enhance the efficiency of planar n–i–p Sb2Se3 solar cells through Sb2Se3 surface modification and an architecture with oriented 1D van der Waals material, trigonal selenium (t‐Se). A seed layer assisted successive close spaced sublimation (CSS) is developed to fabricate highly crystalline Sb2Se3 absorbers. It is found that the Sb2Se3 absorber exhibits a Se‐deficient surface and negative surface band bending. Reactive Se is innovatively introduced to compensate the surface Se deficiency and form an (101) oriented 1D t‐Se interlayer. The p‐type t‐Se layer promotes a favored band alignment and band bending at the Sb2Se3/t‐Se interface, and functionally works as a surface passivation and hole transport material, which significantly suppresses interface recombination and enhances carrier extraction efficiency. An efficiency of 7.45% is obtained in a planar Sb2Se3 solar cell in superstrate n–i–p configuration, which is the highest efficiency for planar Sb2Se3 solar cells prepared by CSS. The all‐inorganic Sb2Se3 solar cell with t‐Se shows superb stability, retaining ≈98% of the initial efficiency after 40 days storage in open air without encapsulation.Environmentally benign and potentially cost‐effective Sb 2Se3 solar cells through Sb2Se3 surface modification and an architecture with oriented 1D van der Waals material, trigonal selenium (t‐Se). The p‐type t‐Se layer functionally works as a surface passivation and hole transport material. The all‐inorganic device structure enables high efficiency and superb stability.This work reports a compatible strategy to enhance the efficiency of planar n–i–p Sb The earth‐abundant low‐toxic materials, together with scalable and cost‐effective fabrication processes promise Sb2Se3 solar cells a competitive technology for high‐efficiency, low‐cost photovoltaics. The unique Q1D structure of Sb2Se3 enables intrinsically benign grain boundaries if the crystals are properly aligned. To date, the highest reported power conversion efficiencies of both superstrate‐planar and substrate‐nanorod type Sb2Se3 solar cells are 7.6% and 9.2%, respectively.Quasi‐1D(Q1D) V–VI antimony chalcogenide‐based solar cells have drawn intensive attention as a result of the meteoric rise in power conversion efficiencies over the past few years.14 In 2Se3 and high interface recombination loss represent the two most urgent obstacles to further improvement of the device efficiency. The intrinsic defects in the low‐symmetry Sb2Se3 semiconductors are complicated and limit their p‐type doping level. To obtain high‐quality Sb2Se3 absorbers, several approaches have been developed, such as rapid thermal evaporation, close spaced sublimation (CSS), vapor transport deposition. Early work mainly focused on improving the quality of the absorber  and on engineering junction interfaces. To further improve device performance, increasing attention has been paid to the trapping and carrier transport at the Sb2Se3 rear surface and related interface. Tang and co‐workers reported that, due to the electrical loss induced by rear surface/interface defects and the presence of a contact barrier, the additional deficit of open‐circuit voltage of Sb2Se3 solar cells can exceed 0.2 V. Nevertheless, little attention has been paid to the properties of the Sb2Se3 rear surface, despite their role in limiting device performance. In this regard, effective surface treatment or passivation methods are yet to be developed to further advance the device efficiency.Nevertheless, the low carrier density of Sbficiency.17 The2Se3 cell structures do not have a wide‐bandgap barrier/passivating layer at the absorber interface to provide carrier confinement and reduce interface recombination velocity. Furthermore, considering the quasi‐intrinsic nature of the Sb2Se3 absorber layer, incorporation of a p‐type interfacial layer to construct n–i–p configuration is a feasible way to overcome the bottleneck of the low doping concentration  of the absorber layer and improve the carrier extraction efficiency. Exploration of efficient charge‐extraction material or hole‐transporting material (HTM) is the key to construct n–i–p Sb2Se3 solar cells. Recently, organic materials  and inorganic materials  have been investigated as HTMs to improve device performance . Especially, PbS quantum dots have been successfully used to prepare record‐efficiency (7.62%) Sb2Se3 solar cells. However, the toxicity of Pb, the high‐mobility Cu impurity and the instability of organic materials limit their practical applications. An ideal interfacial structure should feature compatible material composition, matched energy band level, low interface defect density, good conformal coverage, and low‐contact barrier. Meeting all these requirements is not easy but essential to realize efficient Sb2Se3 devices.In addition to the lack of well‐established surface modification methods, most of the existing superstrate Sbficiency.24 Exp2Se3 surface and construct n–i–p Sb2Se3 solar cells. Different from the ring structure of selenium clusters, trigonal selenium is a 1D van der Waals solid with helical atomic chains. The unique properties of wide bandgap (1.8–2.0 eV), high work function (>5.0 eV), p‐type conductivity (≈1014 cm−3), and high mobility (≈60 cm2 V−1 s−1) make t‐Se a promising p‐type interfacial layer for constructing efficient Sb2Se3 solar cells. We developed a seed layer assisted successive CSS process to fabricate high‐crystalline Sb2Se3 films. The chemical and electrical properties of the Sb2Se3 surface were carefully analyzed. To reduce surface Se deficiency and the related downward band bending, reactive Se was introduced to compensate the surface Se deficiency and form an oriented 1D trigonal Se interlayer on the Sb2Se3 surface. It is found that the prepared t‐Se, which has (101) preferred orientation matches well with the crystal orientation of the Sb2Se3 absorber. Further, the p‐type t‐Se functionally worked as surface passivation and p‐type hole transport layer, facilitating the formation of n–i–p structured Sb2Se3 solar cells. As a consequence, an efficiency of 7.45% was achieved, which is the highest efficiency ever reported for superstrate Sb2Se3 solar cells prepared by CSS.In this work, we present a strategy that uses oriented trigonal selenium (t‐Se) to modify the Sbc chains.26 The22.12Se3 films were prepared by the CSS method according to our previous work with some modifications. A ceramic light visor was inserted between source and substrate to prevent overheating the graphite substrate holder by interactive light radiation  was first deposited onto the CdS surface, then the main Sb2Se3 absorber layer (≈500 nm) was deposited on the seed layer. The two deposition processes were conducted in a successive manner without breaking the vacuum while keeping the substrate temperature higher than 270 °C. This successive high‐temperature process without low‐temperature stage can eliminate repeated nucleation on the grain surface and favor continuous grain growth. Figure 2Se3 film. It can be seen that the Sb2Se3 film has an excellent crystallinity with Sb2Se3 grains as tall as the thickness of the film (through‐thickness grains). The grains are densely packed and have a lateral size of 0.5–1 µm, which is about two times larger than that of the grains in films made by the conventional single‐step CSS process. The XRD patterns of the Sb2Se3 films deposited by seed layer assisted successive CSS process and single‐step CSS process are shown in Figure 2Se3 film prepared by seed layer assisted successive CSS is much stronger than that of the film by single‐step CSS, indicating that seed layer assisted successive CSS process could strongly enhances the crystallinity. Furthermore, the film by seed layer assisted successive CSS shows a more (221) preferred orientation. As Sb2Se3 is composed of 1D (Sb4Se6)n ribbons stacked together via weak van der Waals interaction, strong (221) preferred orientation growth with through‐thickness grains is highly favored for efficient carrier transport and benign grain boundaries.We developed a seed layer assisted successive CSS process to enhance preferred orientation and grain growth. As shown in Figure n growth.29 Fig2.2Figure 2Se3 film. It was observed that the shoulder peak with lower binding energy at 528.7 eV was identified, which is representative of the reduced Sb elemental state or Se‐deficient component at the Sb2Se3 surface. It has been reported that Sb2Se3 is slightly decomposed during high‐temperature process, leading to the formation of Se‐deficient Sb2Se3‐x because of the high vapor pressure of Se. We further conducted XPS depth‐profile measurements to investigate the chemical distribution of Sb2Se3 absorber. Figure S2 of the Supporting Information shows XPS scans of the Sb 3d at various depths. The intensity of the shoulder peak decreased and disappeared with increasing depth, indicating that the detectable Se deficiency only exists within several nanometers from the surface. Because of the high vapor pressure of Se, the high‐temperature deposited Sb2Se3 film is prone to deviate from the stoichiometric ratio. Theoretically, Se‐related defects produce a series of intrinsic doping levels in the bandgap and make Sb2Se3 p‐type under the Se‐rich condition and n‐type under the Se‐poor condition. We assume that the surface Se deficiency is related to the change of chemical environment during the whole CSS process. When the source temperature is high (510 °C), due to the higher vapor pressure of Se, source vapor provides a Se‐rich environment around the substrate, leading to a less Se‐deficient condition. A sufficiently high Se content is critical to guarantee intrinsic p‐type conductivity in Sb2Se3 films. During the cooling process, however, the Se vapor around the hot substrate tends to diminish with decreasing source temperature. Under such circumstances, the deposited Sb2Se3 film may decompose or re‐evaporate, which lead to surface Se deficiency because of the high vapor pressure of Se.In addition to structural properties, the chemical composition was characterized by high‐resolution X‐ray photoelectron spectroscopy (XPS).  surface.31 It 2Se3, we deposited an ultrathin Se layer  on the Sb2Se3 surface by high‐vacuum thermal evaporation. The reactivity of Se on the Sb2Se3 surface was regulated by altering the substrate temperature. It is found that, when the substrate temperature was lower than 80 °C, the device performance changes were negligible. When the substrate temperature was elevated to 100 °C, enhanced spectral response at long wavelengths and enhanced device performance could be observed . However, when the substrate temperature was higher than 110 °C, it was difficult to deposit Se on the Sb2Se3 surface. To understand the mechanism of the performance enhancement, the Se‐modified Sb2Se3 surface was characterized. In order to characterize the change in chemical composition with respect to the pure Sb2Se3 surface, the samples with Se coating were cleaned to make sure that no superfluous Se residue was present on the surface . As shown in Figure 3+ in Sb2Se3. These observations provide evidence that the Se modification process is fairly effective in compensating surface Se deficiency of the Sb2Se3 absorber.In order to tackle the Se deficiency at the rear surface of Sbn Sb2Se3.35 The2Se3 films, Kelvin probe force microscopy (KPFM) was employed to map the surface morphology and potential of the as‐deposited Sb2Se3 surface and the Se‐modified Sb2Se3 surface. To calibrate the potential value of the probe, an Au sample was measured together with the Sb2Se3 samples . The atomic force microscopy (AFM) images . The surface potential measurements indicate a work function increase of 125 mV for the Se‐modified Sb2Se3 surface as compared to the as‐deposited Sb2Se3 surface. This is a considerable change for a 1.2 eV bandgap semiconductor. The reason for the work function variation of the Sb2Se3 surface with different chemical composition can be explained by means of the schematic band diagram in Figure Φ = Ea + ΔEfn, where Ea is the electron affinity and ΔEfn is the Fermi level position referenced to the bottom of the conduction band. In case of the as‐deposited Sb2Se3 surface, the Se‐deficiency related defects, such as Se vacancies, usually produce deep‐donor levels in the bandgap which may act as carrier recombination centers. The deep‐donor states near the conduction band uplift the Fermi energy, resulting in a downward band bending, smaller ΔEfn, and decreased Ф at the surface .37 Thee Figure .2.42Se3/t‐Se interface. The quantitative valence band offset (VBO) can be obtained by VBO=EVBb−EVBa+Vbb, where EVBb and EVBa are the energy positions of the valence band edges of bulk Sb2Se3 and bulk t‐Se, respectively, and Vbb is the band bending at the interface. The Vbb can be estimated from the core level energy shift between the bulk area and the interface, and can be calculated by Vbb=(Ecla−Ecla(i))−(Eclb(i)−Eclb), where Eclb and Ecla are the core level energies of Sb 3d5/2 and Se 3d5/2 in the bulk Sb2Se3 and bulk t‐Se, respectively, and Eclb(i) and Ecla(i) are the core level energies of Sb 3d5/2 and Se 3d5/2 at the Sb2Se3/t‐Se interface, respectively. Note that the values in the equation are calculated under the assumption that binding energy below Fermi edge is negative. The XPS curves and the binding energies are shown in Figure 2Se3 and t‐Se, we chose higher peak position (≈55.4 eV) as the core level energy of Se 3d5/2, which corresponds to the binding energy of Se 3d5/2 in t‐Se . The difference of valence band maxima (VBM) of Sb2Se3 layer and t‐Se layer is 0.10 eV. The total Vbb and VBO at the Sb2Se3/t‐Se interface are 0.11 and 0.21 eV, respectively. To confirm the valence band alignment and band bending at the interface, ultraviolet photoelectric spectroscopy (UPS) was used to examine the energy levels . The calculated VBM and work functions are 0.56 and 4.55 eV for the Sb2Se3 film, and 0.70 and 5.18 eV for the t‐Se film, respectively. The VBM values are in good agreement with those shown in Figure 2Se3 to t‐Se, which corresponds to an upward band bending at the interface of Sb2Se3/t‐Se. Given that the bandgaps of bulk Sb2Se3 and t‐Se remain unchanged, the conduction band offset (CBO) can be calculated by CBO=(Egapb−Egapa)+VBO. The bandgaps of the Sb2Se3 and t‐Se were measured to be 1.20 and 1.98 eV, respectively , and the CBO was calculated to be −0.57 eV.To provide insight into the impact of t‐Se layer on the interfacial band structure, XPS measurements including core‐level spectroscopy and valence band (VB) spectroscopy were conducted at the Sb2Se3/t‐Se/Au. The Sb2Se3/t‐Se forms a favorable band alignment at the contact interface. The upward band bending at the Sb2Se3 surface along with the close VBM levels between the Sb2Se3 and t‐Se will facilitate the efficient extraction of holes from Sb2Se3 to electrode. At the same time, due to the higher conduction band maximum, the t‐Se layer acts as a back field reflector, which blocks photogenerated electrons from injecting into the back electrode and consequently reduces carrier recombination at the back contact. Furthermore, the p‐type conductivity  and the preferred orientation of 1D t‐Se favors hole injection and transport in the t‐Se layer. Based on the discussions above, it can be concluded that the t‐Se layer functionally works as a p‐type hole transport material in the n–i–p device, by passivating the rear surface, facilitating both hole extraction and electron reflection. The matching interfacial properties enable to effectively suppress the interface recombination and reduce the series resistance of the devices, which is expected to deliver a significantly enhanced photovoltaic performance.Figure 2.5Figure 2Se3 solar cells with and without 15 nm thick t‐Se layer. It is obvious that, compared to the control device, solar cells modified with t‐Se layers exhibit noticeably higher overall performance in terms of all photovoltaic parameters. The average power conversion efficiency (PCE) dramatically increases from ≈5.2% to over 7%. Figure J–V) and external quantum efficiency (EQE) curves of the best devices with and without t‐Se measured under standard test conditions. The detailed device parameters are summarized in Table RS and RSH were extracted from the J–V curves using Site's method. Compared with the control devices, the open‐circuit voltage (VOC), short‐circuit current density (JSC), and fill factor (FF) of the device with t‐Se layer are enhanced by 8.7%, 13.8%, and 13.1%, respectively, resulting in a 40% improvement in PCE, from 5.32% to 7.45%. To the best of our knowledge, the obtained efficiency of 7.45% in this study is the highest efficiency value among all reported superstrate Sb2Se3 solar cells fabricated by CSS, and very close to the record efficiency (7.62%) of the superstrate Sb2Se3 solar cells, which is employed with toxic PbS hole transport layer. In addition, no hysteresis between forward and reverse scans was observed in the Sb2Se3 solar cell with t‐Se layer . The EQE characterizations were carried out to verify the reliability of the J–V measurement. The integrated JSC values from EQE curves are 25.2 and 28.6 mA cm−2 for Sb2Se3 solar cells without and with t‐Se layer, respectively, which are in good agreement with the values derived from the J–V measurements . The EQE at the wavelength of 1000 nm is still about 50%, which is higher than most previously reported values (≈30%). The EQE improvement in the long wavelength region (600–1100 nm) contributes to an integrated photocurrent increase of ≈3 mA cm−2. Due to the photons at the short‐wavelength region (<625 nm) have smaller absorption depth, the t‐Se layer does not absorb the photons passing through the Sb2Se3 absorber and contribute to the EQE enhancement . It is known that the photons with long wavelength are generally absorbed at a deep position in the absorber layer, the increased EQE response at long wavelengths is mainly due to the sufficient generation and collection of carriers close to the deep or back side. Considering that the limited thickness of the Sb2Se3 layer (≈500 nm), the transport and recombination of carriers generated at deeply in the absorber layer can be easily affected by the rear surface or interface properties. With the introduction of t‐Se layer, the upward band bending in the surface region promotes holes drifting as well as injection via the back field. The passivated deep‐donor surface defects and the blocking effect of the t‐Se layer for electrons reduce the carrier recombination at the back interface. In addition, the depletion width (xd) of the main junction is found to be increased from 223 to 238 nm for the cell with the t‐Se layer, which strengthens the charge carrier separation and transport deeply in the Sb2Se3 absorber layer  of the main recombination mechanism. The EA linearly extrapolated from the curve of VOC versus temperature at 0 K shows a high value of 1.08 eV for the solar cells with a t‐Se layer, which is close to the bandgap of the Sb2Se3 absorber (1.20 eV), suggesting a dominant Schottky–Read–Hall recombination process in the space charge region. On the other hand, for the device without t‐Se, the EA is calculated to be 0.99 eV, which is about 0.21 eV lower than the bandgap, indicating additional interface recombination and thus the VOC is also limited as well. Figure JSC on the light intensity for both types of solar cells, in which both slopes of the lines are close to 1, indicating that both devices suffer negligible from space‐charge limited photocurrent. The relatively higher value of the slope of the device with t‐Se layer indicates that a better back contact leads to a much easier charge transport.The carrier transport behavior was analyzed by temperature‐dependent VOC decay measurements. In a typical solar cell under open‐circuit conditions, the decay time of VOC upon removal of light is representative of the recombination time constant of the solar cell. Figure VOC decay curves. The calculated decay times are determined to be 2.67 and 0.96 µs for Sb2Se3 solar cells with and without t‐Se layer, respectively. The longer VOC decay time and the better retardation of carrier recombination confirm the reduction of interface defects and the related electronic recombination loss for the device with t‐Se layer.The carrier recombination dynamics was quantitatively analyzed by transient C−2–V) measurements were conducted to understand the origin of the VOC improvement from the built‐in potential (Vbi) variation. As shown in Figure Vbi of the device with t‐Se layer increased from 0.391 to 0.452 V. As Vbi is positively correlated to VOC, the increased Vbi theoretically contributes to the improvement of VOC. Considering the identical CdS/Sb2Se3 junction structure and the same slopes of the linear regions, the increased Vbi can be attributed mainly to the band bending at the Sb2Se3/t‐Se interface. Further, the diode properties are determined from the corresponding dark J–V curves of the two cells, as shown in Figure RSH) and the decreased series resistance (RS) . The diode ideality factors (A) are calculated to be 1.58 and 1.91 for the cells with and without t‐Se layer. The variation in A between 1 and 2 depends on the energies of the deep defects that act as dominant trap states or the tunneling contribution. The lower value of A affirms the reduced trap‐assisted recombination in the device with t‐Se layer.Mott–Schottky C−2–V meas2Se3 solar cell, and compared this to the cell with the commonly used organic Spiro‐OMeTAD HTL. The two devices have the same structure except the HTL materials. The details of the Sb2Se3 solar cell with Spiro‐OMeTAD HTL used in the present stability studies can be found in our previous work. Both devices were stored and measured in ambient environment without encapsulation. As shown in Figure 2Se3 solar cell with t‐Se retains ≈98% of its initial value after storage for 40 days. By contrast, the efficiency shows about 40% decline for the device with the Spiro‐OMeTAD HTL after the same storage duration. These results suggest that the chemically compatible t‐Se based all‐inorganic structure is an effective approach to fabricate highly efficient and stable Sb2Se3 solar cells.In addition, we studied the long‐term stability of the Sbous work.48 Bot32Se3 surface and construct efficient n–i–p Sb2Se3 solar cells. A seed layer assisted successive CSS process was developed to fabricate high‐crystalline Sb2Se3 film. The surface Se deficiency and negative surface band bending were discovered. Reactive Se was innovatively introduced to compensate surface Se deficiency and form an oriented 1D t‐Se interlayer on the Sb2Se3 surface. It is found that the t‐Se layer has a matched (101) preferred orientation, and functionally works as surface passivation and p‐type hole transport material. Benefiting from the suppressed recombination and enhanced band structure, the device performance was significantly improved. An efficiency as high as 7.45% was achieved, which is the highest efficiency reported for superstrate Sb2Se3 solar cells prepared by CSS. Besides, the all‐inorganic Sb2Se3 solar cell shows excellent device stability. This work provides a compatible component and device architecture design for highly efficient and stable Sb2Se3 solar cells.In summary, we have demonstrated a compatible approach to modify the Sb4−1 was cleaned by sonication with by acetone, ethanol, and deionized water for 10 min, respectively. A resistive 30 nm SnO2 buffer layer was deposited on the FTO surface by the RF magnetron sputtering technique at the substrate temperature of 200 °C. CdS layers with a thickness of ≈60 nm were prepared on glass/FTO/SnO2 substrates by chemical bath deposition at 85 °C from a solution composed of deionized water, cadmium acetate, ammonium acetate, and thiourea. The substrate with deposited CdS layer was then immersed into saturated CdCl2 methanol solution to dip‐coat a thin CdCl2 layer, and heat treated in a furnace in air atmosphere at a temperature of 380 °C for 10 min. The Sb2Se3 absorber layers were deposited by a seed layer assisted successive CSS process in a modified CSS facility, as sketched in Figure 2Se3 powder with 99.999% purity. The deposition of Sb2Se3 layer includes two processes: seed layer deposition and main absorber layer deposition. The CSS process was conducted at a vacuum pressure of ≈0.1 Pa. There is no inert gas or reactive gas in the chamber. First, seed layer deposition: set source and substrate temperatures at 310 °C for 120 s, raise source temperature to 510 °C in 60 s, keep the substrate temperature at 310 °C and actually the temperature increases to ≈330 °C after 60 s, keep the source temperature at 510 °C for 10 s; then the source and substrate temperatures naturally cooling to 270 °C and begin the main absorber layer deposition: keep the source and substrate temperatures at 270 °C for 120 s, raise source temperature to 510 °C in 60 s, and then keep at 510 °C for 150 s, the substrate temperature is set at 270 °C the whole process and actually increases to ≈350 °C when the process is over, then turn off the power to stop the deposition and finally take the film out when it is cooled to 150 °C. After that, the Sb2Se3 film was rinsed with carbon disulfide  for 1 min and then rinsed by deionized water. The Se films were deposited by thermal evaporation with 99.999% pure Se particles. The chamber was pumped to a pressure of less than 5 × 10−4 Pa before evaporation. The substrates were rear heated at an average rate of 10 °C min−1 from room temperature to the designed temperature. The deposition rate was monitored by crystal oscillator. The Deposition rate was kept at ≈2 Å s−1. Subsequently, gold back contacts (≈100 nm) were deposited by thermal evaporation in a vacuum chamber and the active area of individual device was defined as 0.126 cm2.FTO‐coated glass with sheet resistance of 8 Ω sq2Se3 surface measurements: the as‐deposited Sb2Se3 after rinsed by CS2 solution for 1 min was defined as as‐deposited Sb2Se3; the Sb2Se3/t‐Se (2 nm) was first re‐evaporated in an RTP furnace for 60 s at 150 °C under a vacuum pressure below 0.1 Pa, then immersed in CS2 solution for 5 min and rinsed by deionized water, the obtained sample with pure Sb2Se3 surface without Se residue was defined as Se‐modified Sb2Se3. Samples for interface band alignment measurements: a 50 nm thick t‐Se film deposited on top of the Sb2Se3 substrate was defined as bulk t‐Se; a 3 nm thick t‐Se film on the Sb2Se3 substrate was defined as Sb2Se3/t‐Se interface; the Sb2Se3/t‐Se(3 nm) was re‐evaporated at 150 °C and immersed in CS2 solution, the sample with clean Sb2Se3 surface without Se residual was defined as bulk Sb2Se3. Sample t‐Se for GIXRD measurement: a 50 nm thick t‐Se film on the Sb2Se3 substrate.Samples for Sbα as the radiation source. GIXRD pattern was performed with incidence angle of 0.5° and five consecutive measurements were collected and averaged into single spectra. TEM and HRTEM: Samples for TEM were prepared by ablating the device using a focused ion beam microscope. Pt layers were deposited on top of the device for protection. TEM and HRTEM images were taken on a JEOL JEM 3200FS TEM at 300 kV. AFM and KPFM: The AFM and KPFM measurements were performed on a Bruker Dimension Icon with Nanoscope V controller. All the KPFM images were collected in lift mode by using SCM–PIT probes with a spring constant of 2.8 N m−1 and a resonant frequency of 78.3 kHz. During all KPFM measurements, the optimal lift height is 30 nm. UPS and XPS: UPS and XPS characterizations were obtained with a Thermo Fisher Scientific K‐Alpha+, using the HeI (21.22 eV) emission line and Al Kα radiation (1486.6 eV). X‐ray photoelectron depth profiling spectra were recorded by using Ar+ sputtering gun operated at 3 keV, rastering a 2.0 × 2.0 mm2 area with 5 s postetch delays.The film morphologies were characterized by SEM (FEI Apreo LoVac). The optical properties were measured by UV–vis spectrophotometer . XRD and GIXRD: XRD measurements were performed by a Bruker D8 Advance X‐ray diffractometer with Cu KJ–V curves were measured using a Keithley 2400 source meter. The illumination was provided by a Newport Oriel 92192 solar simulator with an AM1.5G spectrum, operating at 100 mW cm−2, which was calibrated by a standard silicon solar cell (91150 V) from Newport. The EQE spectra were measured in AC mode on a spectrum corresponding system (Enlitech QE‐R), calibrated by Si reference solar cell. Capacitance–voltage (C–V) was measured with the bias voltage varying from −0.5 to 0.5 V at room temperature under near‐dark condition. Besides, the measurement frequency was set to 100 kHz and the AC modulation voltage was set to 20 mV. The voltage dependent depletion width was calculated according the formulation xd = εA/C(V), where xd is the depletion width, ɛ is the permittivity of Sb2Se3, A is the device area, and C(V) is voltage dependent capacitance, respectively. For low temperature J–V measurements, the devices were placed in a closed‐cycle nitrogen cryostat, and the temperature was varied from 100 to 300 K with a step size of 20 K. A LabVIEW control program was used to operate the temperature controller and the Keithley 2400 source measurement unit for current and voltage data acquisition. Light intensity J–V measurements and transient photovoltage VOC decay curves were measured by an electrochemical workstation . The light source is the standard wideband white LED light source for ZAHNER CIMPS. The light intensity changed from 1000 to 100 W m−2 in light intensity J–V measurements. The light intensity used in transient photovoltage VOC decay measurements was 1000 W m−2.The The authors declare no conflict of interest.Supporting InformationClick here for additional data file.


Here, we report that PCBP2, which strongly resembles PCBP1, can stabilize sortilin transcripts in macrophages using the same mechanism employed by PCBP1. PCBP2 recognized the C‐rich element in the 3′ UTR of sortilin mRNA, and PCBP2 knockdown decreased sortilin transcripts, indicating that PCBP2 stabilizes sortilin mRNA by binding to its 3′ UTR. Zn2+ reversibly inhibited the nucleotide binding ability of PCBP2 in vitro. These findings suggest that both PCBP2 and PCBP1 may control the stability of sortilin transcripts by sensing intracellular Zn2+ levels in immune cells.Post‐transcriptional regulation of cytokine production is crucial to ensure appropriate immune responses. We previously demonstrated that poly‐rC‐binding protein‐1 (PCBP1) can act as a  2+ interferes with this interaction, suggesting that both PCBP1 and PCBP2 may regulate the stability of sortilin mRNA by sensing the intracellular Zn2+ levels.We analyzed the function of poly‐rC‐binding protein‐2 (PCBP2), a paralog of PCBP1, in stabilizing the mRNA of sortilin, a cytokine trafficking mediator. PCBP2 binds to C‐rich elements in the 3′ UTR of sortilin mRNA; Zn They recognize C‐rich elements (CREs) of single‐stranded RNA and have diverse functions affecting RNA processing, translation, and stability through their binding to CREs 2+ to PCBP1 reversibly interfered with the nucleotide binding activity of PCBP1 in vitro2+ affects the nucleotide binding activity of PCBP1 paralogs. In this study, we measured the protein expression levels of PCBP1 and PCBP2 in multiple types of immune cells and examined the involvement of PCBP2 in the stabilization of sortilin mRNA. We describe the interaction of PCBP2 with the CRE of the 3′ UTR of sortilin mRNA, which leads to the stabilization of the sortilin mRNA. Zn2+ reversibly interferes with PCBP2 nucleotide binding to the CRE. We previously observed that loading primary macrophages with Zn2+ leads to a reduction of sortilin mRNA 2+ levels in immune cells. Our observations provide insights into the roles of PCBP1 and PCBP2 in mRNA regulation, metal homeostasis, and innate immunity via the post‐transcriptional control of sortilin expression.We previously demonstrated that PCBP1 is involved in the regulation of sortilin, which mediates cytokine trafficking via stabilization of its mRNA + and CD8+ splenocytes , plasmacytoid DCs (pDCs), and eosinophils  on yeast cell lysates expressing FLAG‐PCBP2 in the presence of Zn2+. As expected, the shifted band derived from the complex of FLAG‐PCBP2 with radiolabeled probe was decreased when the lysate was incubated with zinc, whereas the signals were unchanged when the lysate was incubated with iron, indicating the Zn2+ specificity of PCBP2 binding  binding.We previously reported reversible inhibition of nucleotide binding ability of PCBP1 by Zning Fig. . Incubating Fig. , indicatPathogen infection and cellular stimuli induce PRR signaling, resulting in cytokine and chemokine production by immune cells. These inflammatory responses protect the body against harmful stimuli; however, chronic inflammation and/or autoimmune diseases are primarily attributed to excessive immune responses caused by cytokine overproduction. Therefore, the appropriate regulation of cytokine production is crucial for accurate immune responses. Here, we demonstrate that PCBP2, a paralog of PCBP1, can control the stability of sortilin mRNA through the binding to the CRE in 3′ UTR, along with PCBP1. So far, it has been reported that sortilin plays a key role in exocytic trafficking of cytokines in multiple types of immune cells Toll‐like receptor (TLR) stimulation leads to a decrease in the amount of sortilin mRNA that is stabilized by PCBP1 2+ affected the poly(C)‐binding ability of PCBP1 2+. We have also demonstrated that cells, which were loaded with Zn2+, have decreased sortilin levels 2+ content 2+ to PCBP1/2, destabilizing sortilin mRNA to attenuate the immune response by preventing cytokine overproduction. We infer that both PCBP1 and PCBP2 may function as sensors of intracellular Zn2+ levels to control multiple genes, in addition to Sort1, by the binding to their CREs; however, further studies are required to better understand the underlying mechanisms. In summary, we show that the post‐transcriptional regulation of sortilin mRNA by PCBP2 and PCBP1 is important for the post‐translational control of cytokine production.Both PCBPs are iron chaperones that can bind and deliver iron, in addition to their functions as RBPs C57BL/six mice  were purchased from Sankyo Labo Service Corporation . All animal experimental protocols were approved by the Animal Research Committee of the Medical Research Institute, Kanazawa Medical University. All animal experiments were carried out in accordance with institutional guidelines.2 atmosphere. Media for primary culture were described previously −1 penicillin, and 100 U·mL−1 streptomycin (Wako Chem.).All cells were cultured at 37 °C in a humidified, 5% COPCBP1 depletion and PCBP2 depletion in RAW264.7 cells were performed by transfection of siRNA  using ScreenFect siRNA (Wako Chem.) according to the manufacturer's instructions. A nontargeting sequence siRNA pool  was used as a control. The preparation of whole‐cell lysate and immunoblotting was carried out as previously described 5 cells was isolated with ReliaPrep RNA Cell Miniprep System , and cDNA was synthesized using 500 ng of total RNA with ReverTra Ace qPCR RT Master Mix , according to the manufacturers' instructions. qRT‐PCR was performed as previously described Gapdh. The primer set for the detection of Pcbp1, Sort1, and Gapdh was described previously https://primer3plus.com/cgi-bin/dev/primer3plus.cgi): forward, 5′‐TATGCCATTCCACAGCCAGA‐3′; reverse, 5′‐CTGCCCAATAGCCTTTCACC‐3′.Total RNA from 5 × 10m) was incubated with cell lysate prior to incubation with probes. Samples were separated on 6% polyacrylamide DNA gels  and analyzed by the Odyssey system or phosphorimaging.RNA EMSA with macrophage cell lysate and EMSA with a yeast pep4Δ strain transformed with pYES2 FLAG‐PCBP2 were performed as described previously RNA immunoprecipitation was carried out as described previously 7 macrophages using a RiboCluster Profiler RiboTrap kit (MBL), according to the manufacturer's instructions. For in vitro transcription with T7 RNA polymerase, the 3′ UTR of sortilin mRNA was amplified from a plasmid containing full‐length sortilin cDNA , and the resulting amplicon was cloned into a pGEM‐4Z vector (Promega) using the In‐Fusion HD Cloning Kit . The BrUTP‐labeled 3′ UTR of sortilin mRNA was prepared using an RNA Riboprobe System‐T7 kit (Promega). Magnetic Beads Protein A/G (Merck Millipore) bound to anti‐BrdU (MBL) or rabbit IgG (Merck Millipore) were incubated with the mixture of cytosol and BrUTP‐labeled 3′ UTR at 4 °C for 3 h. The BrUTP‐labeled RNA–protein complex bound to magnetic beads was eluted with PBS containing BrdU, followed by immunoblotting.The cytosolic fraction of macrophages for RiboTrap analysis was prepared from 5 × 10graphpad prism software . One‐way and two‐way ANOVA were used for experiments with one and two variables, respectively. A P value of < 0.05 was considered statistically significant.Statistical analysis was performed by ANOVA using The authors declare no conflict of interest.TY‐W conceived the project. TY‐W, CCP, and NO supervised the project. TY‐W performed the experiments and analyzed the results. TY‐W, CCP, and NO wrote and edited the manuscript.


The use of transient absorption spectroscopy showed that a complex ping‐pong energy‐transfer mechanism, namely from the porphyrin to MoS2 and back to the porphyrin, operated. This study reveals the potential of transition‐metal dichalcogenides in photosensitization processes.Molybdenum disulfide nanosheets covalently modified with porphyrin were prepared and fully characterized. Neither the porphyrin absorption nor its fluorescence was notably affected by covalent linkage to MoS Two‐way street: A hybrid porphyrin‐MoS2 material has been shown to undergo complex ping‐pong energy transfer from the porphyrin to the MoS2 and back to the porphyrin. This finding illustrates the potential of transition‐metal dichalcogenides in photosensitization processes. In general, the structure of such dichalcogenides is analogous to that of graphene, with an atomic layer of a transition metal sandwiched between two layers of chalcogen atoms. More specifically, each Mo atom is bound to six S atoms and forms a three‐atom‐thick monolayer.2 materials have received enormous attention in recent years because of their extraordinary optoelectronic and electrocatalytic properties, especially in the area of energy applications and catalysis.2. Most notable is the employment of a) strong intercalants such as Li+, which result in phase transfer to an octahedral metallic 1T polytype,Molybdenum disulfide MoS is a typ2. In fact, incorporation of molecular dopants results in fine‐tuning of the electronic as well as optical properties of MoS2 which, in turn, enables broadening the spectrum of its applications. In MoS2, the S atoms in the basal plane are, however, rather inert as a result of saturation. Furthermore, Mo atoms are placed between S layers and, thereby have challenging chemical reactivity. However, the covalent functionalization of MoS2 has been accomplished.2 polytype has been reacted with organoiodides,2.2 nanosheets.2 to yield interesting hybrids. Pyrene and phthalocyanines have, for example, been coupled with MoS2.2 and WS2, which facilitated the electrostatic association of carbon dots2 with photoactive porphyrins.Additional efforts have been placed on developing strategies for chemically modifying transition‐metal dichalcogenides.2 with a 1,2‐dithiolane derivative 1 featuring a porphyrin (H2P). The newly prepared H2P‐MoS22 was comprehensively characterized by spectroscopic, thermal, and microscopy means. We also gathered insight into the electronic interactions between the porphyrin and MoS2 on the femto‐ to nanosecond timescales upon photoexcitation.Herein, we report on the modification of the edges of exfoliated MoS1 was synthesized by a condensation reaction between α‐lipoic acid and 5‐(4‐aminophenyl)‐10,15,20‐(triphenyl)porphyrin. In parallel, treatment of bulk MoS2 with chlorosulfonic acid allowed exfoliation of semiconducting nanosheets.2 with 1 yielded H2P‐MoS22, as summarized in Figure 1, as revealed by UV/Vis spectroscopy . Purified 2 showed reasonable solubility in DMF, benzonitrile, and 2‐propanol, while it is completely insoluble in dichloromethane and water.1,2‐Dithiolane‐based porphyrin 2 with H2P came from vibrational spectroscopy. Attenuated total reflectance infrared (ATR‐IR) measurements on H2P‐MoS22  an amide carbonyl vibration at 1660 cm−1, and c) an aromatic C=C bending at 1590 cm−1. We further validated the origin of the IR bands related to H2P within the hybrid by performing a reference experiment in which a mixture of tetraphenylporphyrin and exfoliated MoS2 was processed under the same experimental conditions as those for obtaining 2. The presence of physisorbed tetraphenylporphyrin was not detected in the UV/Vis and emission spectra of the reference material, and no contribution in the mass loss detected by TGA was observed .Evidence for the success in the covalent modification of MoS2 came from Raman spectroscopy. Excitation at λ=514 nm showed the existence of the C=C and C=N bonds of H2P in the 1000–1600 cm−1 region . Characteristic modes corresponding to MoS2 were observed in the lower wavenumber region, especially upon excitation at λ=633 nm. These modes become visible as a result of coupling with the A1 excitonic transition, which in turn produces resonance Raman enhancement of the first and second order vibrational modes. Exfoliated semiconducting MoS2 gave rise to bands corresponding to 2LA(M) at 447 cm−1, A1g at 406 cm−1, and E2g1at 382 cm−1  was found to be considerably higher for 2, that is, 3.6 versus 1.9 for exfoliated MoS2. Complementary Raman mapping assays revealed an increased IA1g/I2LA(M) intensity ratio for 2 compared to that of exfoliated MoS2  maps for the reference material prepared by physisorption . Overall, this is indicative of a “healing” of the S‐vacancies by the 1,2‐dithiolane species in 2.−1, indicative of the metallic polytype, highlights the semiconducting nature of MoS2 in 2.Additional structural information regarding 2 functionalization in 2. Exfoliated MoS2 was found to be thermally stable under N2 from 100 to 800 °C. Notably, the mass loss observed in the modified MoS2 is directly related to the thermal decomposition of the organic addends. Hence, the mass loss of 9 % for 2  provided information related to the degree of MoS2. Figure 2 flake . Figure 2,3, Mo‐M4,5, C‐K, and Mo‐M2,3 edges are observed in both spectra. The energy‐loss near‐edge structure (ELNES) investigations of the C‐K edge are displayed in the right panel of Figure S7. Two signatures at 291.2 and 293.4 eV can be observed in the σ* region (ca. 290–325 eV) for 2. Such signals are not discerned in the spectra of other carbon‐containing materials employed as references. Some differences can also be observed in the main contribution to the π* signal at ≈284–290 eV. Here, the C‐K edge is dominated by an aromatic C contribution  contributions of H2P.2P because the Mo‐M2,3 edge is in the same energy range. For all these reasons, the features highlighted in the C‐K edge of 2 correspond to a porphyrin contribution. Furthermore, we conclude from the EELS analyses that H2P is homogeneously distributed across the surface of the MoS2 flakes in 2 .Spatially resolved electron energy loss spectroscopy (SR‐EELS) using a scanning transmission electron microscope (STEM) enables the morphology and chemical composition of complex hybrid nanostructures to be investigated at the local scale.2, the UV/Vis spectrum was recorded in DMF and compared with that of 1. In particular, the UV/Vis spectrum of 1 showed a sharp absorption at λ=418 nm that corresponds to the Soret band, accompanied by weaker Q‐band absorptions at λ=516, 550, 596, and 648 nm , 626 (B‐exciton), and 681 nm (A‐exciton) are discernable in the UV/Vis spectrum of exfoliated MoS2  was determined by time‐correlated single photon counting (TCSPC) measurements on 1 with DMF as the solvent. Photoexcitation of 2 at λ=420 nm, however, resulted in no significant quenching of the porphyrin‐centered fluorescence relative to that of 1 .Excited‐state electronic interactions could also be used to probe interactions between the Hm Figure , right, 1, exfoliated MoS2, and 2 by means of femtosecond transient absorption spectroscopy (fsTAS) using an excitation wavelength of 420 nm . Global analysis revealed the presence of three exponential decays on the fsTAS timescale. These three decays are followed by a much slower decay, whose dynamics are outside the timescale of our experiments. Therefore, a target model was applied to fit the transient absorption data ], consistent with excitation into the Soret‐band absorption, which internally converts into the first singlet excited state [S1(H2P)] within 16 ps. The latter undergoes intersystem crossing within 10 ns to the triplet state [T1(H2P)], which, in turn, decays back to the ground state on the microsecond timescale. Additionally, aggregates of 1 are present in the sample and they were found to decay to the ground state within 142 ps. These dynamics resemble very closely those already reported for a variety of porphyrins.As a complement to the aforementioned method, we probed 2, ground‐state bleaching evolved at λ=476, 623, and 683 nm together with positive absorptions at λ=591, 655, and 743 nm . Global analysis of the data revealed five exponential decays after the absorption of light. The first three decays, which take place within 2.7, 7.3, and 250 ps, involve biexciton [be(MoS2)] and trion [tr(MoS2)] formation as well as the decay of these many‐body particles into single excitons [se(MoS2)]. Single excitons then diffuse within 1.8 ns across the layers in the z direction [de(MoS2)], before they recombine to recover the ground state within 21.2 ns  from S2, which is populated upon excitation at λ=420 nm, to MoS2. However, since the energy transfer is faster than the time resolution of our instrument, we are unable to measure the population of the S2(H2P). Global analysis showed four exponential decays on the fsTAS timescale and an additional decay, which is, however, outside of our detection range. Our interpretation is based on a kinetic target model, which is depicted in the middle panel of Figure 2, with a minor contribution from a MoS2‐centered excited state. A nonquantitative ultrafast energy transfer from H2P to MoS2 then occurs, which results in comparable populations of the H2P‐centered (70 %) and MoS2‐centered (30 %) excited states. Afterwards, these three species, which govern the deactivation of exfoliated MoS2 and which result in the formation of single excitons, are discernible. The times for the interconversions are 2.5, 7.5, and 204 ps. Once populated, the single exciton of MoS2 is subject to a second energy transfer. This time, however, single excitons are transferred from the MoS2 to the S1 state of H2P. This energy transfer takes 1.75 ns. From here, the porphyrin decay proceeds through intersystem crossing to the corresponding triplet excited state within 14 ns, and recovery of the ground state on the microsecond timescale. This mechanism is also based on the energetics of 2. The S1 state of 1, with its fluorescence at 1.73 eV (λ=715 nm) and the MoS2 fluorescence at 1.82 eV,1 and 2 upon photoexcitation.For e Figure . Selecti1 onto exfoliated MoS2 to afford novel H2P‐MoS22, which was fully characterized. Thorough photophysical investigations based on steady‐state and time‐resolved measurements corroborate that the decay of the photoexcited porphyrin involves a ping‐pong energy transfer to and from MoS2. Our findings suggest that transition‐metal dichalcogenides have great potential in photosensitization. Such hybrid materials may be useful in energy‐conversion applications.In summary, we accomplished the covalent grafting of porphyrin The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.