Optimization of synthesis parameters to obtain single phase NbC NPs has been successfully done. The phase identification, morphology and nature of carbon were determined with the help of X-ray diffraction, transmission electron microscopy (TEM) and Raman spectroscopy. X-ray photoelectron spectroscopy (XPS) suggested the presence of multiple oxidation states of Nb associated to NbC and NbCxOy centers on the surface of NPs. Due to the presence of NbCxOy on the surface of NPs, absorption under visible region of EM spectrum has been observed by UV-visible spectroscopy. Different organic dyes  were used to study the effect of holding time on the photocatalytic performance of as-synthesized samples. RhB dye was found to be the most sensitive organic molecule among all the considered dyes and degraded 78% in 120 min.A facile carbothermal route was adopted to obtain niobium carbide nanoparticles (NPs) embedded in carbon network from Nb Moreover, dyes are colored and cause hindrance to the sunlight penetration and reduction in dissolved oxygen, which affects aquatic living organisms3. About 15% of industrial dyes are being disposed without any treatment4. Among different types of treatment methods , photocatalysis has emerged as a green technology due to the complete mineralization of pollutants to water, CO2, and mineral acids even at room temperature6.The extensive spillage of organic pollutants in water from industries has led to severe environmental problems. These organic pollutants are not only creating aquatic pollution, but also disturbing the ecological balance. The required biochemical oxygen demand (BOD) in river streams of most Asian rivers has decreased due to organic pollutants released by industries2, ZnO and carbon allotropes  are the common photocatalysts used for degradation of different dyes under UV-visible irradiation13. Dissimilar to these photocatalysts, metallic photoconductors contain either small band gaps or absence of energy difference between highest occupied and lowest unoccupied bands14. Metallic photocatalysts exhibit the generation of electron-hole pair due to inter-band transitions  and higher carrier density than semiconducting photocatalysts make them suitable as photocatalysts15. Transition metal carbides (TMCs), excellent high-temperature stable compounds exhibiting high corrosion and tribological properties, are now being exploited for other applications like catalyst supports in electrochemical cells19. Among all the TMCs, niobium carbide (NbC) is a highly stable compound and is mostly used in nuclear and thermal engines22. Different morphology like thin films, micro-cage shaped, nano wires and fibers of NbC have been synthesized to extract the required properties suitable for different applications such as corrosion resistance, catalyst supports, rotors of micro ceramics engines, conduction and electrodes for electrochemical cells, respectively27. As a function of microstructure, Coy et al.28 studied the electrocatalytic activity of NbC thin films for hydrogen evolution reaction (HER). Further, NbC thin film has also been found as a electrocatalyst support for both alkaline and acidic mediums for HERs29. Carbon coated TMCs have drawn the attention of researchers for their unique physicochemical properties for electrocatalytic H2 production, targeted drug delivery, flexible electronics, magnetic properties and electro-oxidation34. Encapsulation of NPs help to overcome the agglomeration and oxidation limitations particularly for high temperature applications of nano TMCs. Among all the elements, carbon is preferred for encapsulation due to good electrical conductivity and stability in different chemical environments and has a high specific surface area in different forms36. Carbon as coating hinders the oxidation of NPs as reported by Wang et al.37 and Singla et al.32 for tungsten carbide NPs which enables these nanomaterials to be used in various electro oxidation reactions38.Metal oxides viz. TiO47. To overcome the high temperature and time consuming procedures, carbothermal synthesis in an autoclave has been practiced at low temperature34. With the help of an autoclave, various reactants were subjected to carburization of the transition metal  source in the presence of a reducing agent to obtain nano TMCs at relatively low temperature48.Various researchers have practiced encapsulation of carbon on NPs through different synthesis techniques and chemical reagents50, the developing process is not so economical. Zhang et al.42 dsynthesized core-shell NbC/C nano spheres by laser ablation methods in which they used a Nb target in a carbon rich liquid environment (ethanol) for ablation. As far as photocatalytic applications are concerned, except for WC and Mo2C, other carbides are less studied52. Chen et al.52 have reported the photocatalytic degradation of RhB with the help of NbC in almost 12 h under a 300 W mercury lamp. A detailed analysis of photocatalytic characteristics of NbC and coated NbC is not available. In the present approach, NbC/C nano-composite has been synthesized via a carbothermal route. A mechanism responsible for the formation of NbC NPs has been proposed from the data obtained by various crystallographic, thermal, microstructural, spectroscopic, and thermodynamic calculations. The photo-enhanced catalytic efficiency of NbC/C nano-composite has been studied by using RhB, MB, and MO dye organic model compounds, frequently used as a tracer, bright colorant of wool, Nylon fabric, metal detection, and as a food colorant54.Although a significant amount of work has been reported on the synthesis of carbon encapsulated TMCs2 where a = lattice parameter in Å and x = atomic ratio C/Nb and listed in Table 155. For the optimization of synthesis parameters (temperature and holding time) single phase NbC NPs were synthesized using experiments discussed below.For the synthesis of single phase carbon coated NbC NPs, synthesis temperatures and holding times were optimized on the basis of XRD results which are shown in Fig. 2O5 to NbO (sample 10C600). While, at 700 and 800 °C, mixtures of (NbC + NbO2) and NbC were obtained respectively. The absence of the NbC phase at 600 °C  resulted in the complete transformation of Nb2O5 to NbC without showing any signature of lower oxides of Nb.XRD patterns of samples synthesized at low temperatures  with 10 h holding is shown in Fig.  °C Fig.  may be d2O5 with optimized temperature which is shown in Fig. 2O5 after carburization and the formation of a mixture of NbO2 and NbC. Sample 7C800 showed contrary results in which carburization exceeded the reduction process as shown by NbO2 as a minor phase and NbC as a major phase. In the case of 10 h, complete reduction followed by carburization of Nb2O5 was observed in sample 10C800  beyond 10 h at 800 °C57. While the formation of NbO2 as a result of decarburization at longer holding (11 h) instead of NbO (as observed in the case of lower temperature) may be attributed to the higher stability of NbO2 than NbO58.After obtaining single phase NbC at 800 °C, the holding time has been varied to study the behavior of reduction and carburization of Nb800 Fig.  resultinOptimization of synthesis parameters to obtain single phase NbC also affects the constituted crystallographic distortion. In order to understand all the distortional transitions that occurred in NbC due to different synthesis temperatures and holding times, stress-strain analysis  of the obtained samples has been done. Detailed theory and assumptions of Williamson-Hall analysis has been provided in Appendix-II.2O5 to NbC, while presence of NbO2 at 700 °C (sample 10C700) distorts the crystallites which shows higher strain in crystallites. As a function of holding time at 800 °C, the induced strain tends to decrease as carburization occurs. Figure 2) to 7 h (NbC + NbO2) then 10 h (pure NbC)) leading to a large reduction in strain. The diffusion of carbon in the Nb lattice is increased by increasing the holding time which supports the enhancement of lattice parameters by reducing the strain and supporting the attainment of equilibrium positions of Nb and C atoms in the NbC unit cell. Beyond 10 h (11C800), decarburization of NbC led to increase the distortion which might be associated to incorporation of oxygen resulting NbO2.Table 5/2 and 3d3/2 as shown in Figs XPS analysis of all the samples synthesized at 800 °C was carried out to understand the surface chemical variations  that reactants undergo during reduction and carburization processes. The survey spectrum confirms the presence of Nb, C and O on the surface of as-synthesized samples which is shown in Figs xOy, NbO2 and Nb2O5) even after obtaining single phase XRD pattern of 10C800 as shown in Fig. 4+ (NbO2) and Nb5+ (Nb2O5) has increased  as the holding time was shifted from 10 h (81.72%). Moreover, binding energy associated to NbCxOy (~204.8 eV) was not observed in 5C800. HR-XPS data for Nb3d suggest that direct conversion of Nb2O5 to NbC involved reduction and carburization simultaneously which was responsible for the multiple transitions of oxidation states of Nb in the lattice as earlier discussed for XRD analyses. The HR-XPS spectrum of C 1 s as shown in Figs 2 and C-C sp3 bonds in all the samples at 282.8, 284.4 and 285.1 eV, respectively63. Further, peaks observed around 286.4, 288.1 and 290.1 eV might be associated to hydroxyl, carbonyl groups and π → π* transitions, respectively63. HR-XPS spectra of O 1 s  and adsorbed water/carbonyl (C=O) group, respectively65.Figure 1 s Figs  and S1 sTEM micrographs of 10C800 are shown in Figs x or NbCxOy. Further, elemental line profile (STEM) was also taken to observe the linear distribution of each elements from periphery to core which is also shown in Fig. xOy (NbOx centers) as observed from XPS analysis.Further, to observe the distribution of Nb, C and O in the synthesized NbC NPs (10C800), elemental profile was also carried out as shown in Fig. 2 gas for the samples synthesized at 800 °C as shown in Fig. 2/g with 0.3697 cm3/g pore volume for 10C800 while, the presence of oxide and Nb metal with NbC in other samples induced the reduction in SSA and pore volume as shown in the inset of Fig. et al.49. Such variation of SSA and pore volume of synthesized samples might have occurred due to evolution of CO or CO2 as a result of in situ reduction and carburization of niobium oxide by encapsulated carbon. Further, pore size distribution suggested the contribution of high surface area and mesopores of as-synthesized samples as shown in Fig. BET specific surface areas (SSA) were calculated from adsorption-desorption curves using N66 which is higher as compared to the present work. The transition of Nb2O5 to NbC at relatively low temperature (800 °C) in a high pressure closed chamber is a temperature driven transition. The mixture of reactants contains Mg and charcoal, which act as reducing agents providing MgO, CO and CO2 as byproducts. The gaseous byproducts (CO and CO2) of the reactions lead to increase the pressure inside the autoclave. Various chemical reactions occurring inside the autoclave have been listed in the supplementary information.Normally NbC is synthesized at >1000 °C2O5 forming NbO2, NbO and (NbO2 + NbO) mixture in the presence of reducing agents (Mg and C) and their mixture. All the reduction reactions suggested that Mg  reduced Nb2O5 more efficiently to obtain NbO2, NbO and (NbO2 + NbO) mixture as compared to C and (C + Mg) mixture as represented by reaction  in Fig. 2O5 equivalently to obtain similar products as reaction  and , respectively, which are shown in Fig. 2O5.All the possible chemical reactions are categorized into two categories; (i) reduction and (ii) carburization. Reactions (S1–S24) represents the reduction of Nb2  which is also supported by the XRD pattern of sample 10C600. Further, the formation of NbC can be obtained via in-situ reduction carburization of oxide products (NbO2 and NbO) as Nb metal was not obtained in any of the synthesized sample  mixture, all are feasible with negative ΔG values. Unlike reduction reactions, reduction-carburization of NbO2 and NbO follow different paths as shown in Fig. 2 + NbO) mixture is represented by reaction (S37–S42) which are more spontaneous than previous chemical reactions (S25–S36) with high difference of ΔG values. As single phase NbC was obtained at 800 °C, reduction-carburization reactions follow the following order according to ΔG at 800 °C; S37 > S41 > S38 > S42 > S25 > S27 > S29 > S33 > S36 > S32 > S40 > S35 > S31 > S39 > S34 > S26 > S28 > S30.As a result of reduction reaction, feasibility of the formation of NbO is higher than that of NbO2 + NbO mixture is also evidenced in the XRD pattern  were obtained at 600 °C. This is further converted to NbC (major) + NbO2 (minor) at 700 °C and NbC at 800 °C. In this sequence, charcoal encapsulated NbO particles which is followed by reduction-carburization reactions. As a result, CO and CO2 gases evolved through coated carbon resulting to porous particles. Moreover, the transition Thermodynamically higher feasibility of the formation of NbOern Fig.  of 10C60−1 and 1589.8 nm−1 respectively. Ferrari and Robertson67 suggested that the disordered scattering and bond stretching of sp2 carbon atoms contribute to the formation of D- and G- band respectively. The position of the G- band near to 1600 nm−1 suggests nano-crystalline graphite. The observed higher intensity of the D- band than the G- band in these samples depicts the amorphous nature of carbon as has also been observed from TEM analysis.Raman spectroscopy was used to analyze the nature of carbon in samples 5C800 and 10C800. Figure 2 and Nb2O5) or continuous carbon network (highly disordered graphic carbon) in the nanocomposite powder samples70. The band gap for all the samples was calculated with the help of Tauc’s relationship between photon energy (hν) and absorption coefficient (α) as71.Figure 2 vs. hν curve as shown in Fig. 70 and Nb-O centers69, respectively. Table The values of band gap were calculated by extrapolating the linear portion of the (αhν)Moreover, to observe the photoemission characteristics of the synthesized samples photoluminescence (PL) spectroscopy was conducted, which is shown in Fig. 72. Moreover, the presence of NbCxOy on the surface also promotes the photoluminescence quenching.For 11 h of holding time (11C800) more intense photoemission is observed as compared to 10C800. This is because of the presence of some functional groups, which reduce the efficiency of non-radiative decay. The XPS spectra showed the energy corresponding to π → π* transition which also promotes the photoluminescence quenching as observed in Fig. Photocatalytic activity of prepared nanocomposites was monitored by observing the degradation of RhB, MB and MO dyes. All of them are cationic dyes, which acquire electrons from donor excited sites of a photocatalyst. Figure Moreover, Fig. 0, C, K1 and K2 are the solution concentrations at irradiation time t = 0, at different times and pseudo first and second order rate constant (per minute), respectively71. A plot of −ln(C/C0) vs. t (min) and t/C vs. t (min) as shown in Fig. 1 and K2; min−1) and regression coefficients (R2P) have been listed in Table 1) of 0.0143, 0.0102, and 0.0075 min−1 for RhB, MB, and MO dyes, respectively, following pseudo first order kinetics. Moreover, based on quality of fitting (R2), it can be suggested that 10C800 and 11C800 followed pseudo first order kinetics while, 5C800 and 7C800 followed second order kinetic  and superoxide anion radicals  such as hydroxyl radicals (OH2/g (5C800) to 506 m2/g (10C800) and then decreased to 475 m2/g (11C800). Further, PL spectroscopy data show that recombination decreases with the decreased impurities in the synthesized samples resulting in better mineralization of dye with 10C800 than other samples. Further, the reusability of the nanocomposite photocatalyst is a significant parameter which maneuvers the practical usefulness of photocatalyst. To observe the reusability and stability of as-prepared C coated NbC NPs, recycle reactions were carried out for the degradation of all the dyes over 10C800 as shown in Fig. rd cycle of photodegradation reaction as shown in Fig. The photocatalytic activity of the synthesized samples was increased with increasing holding time at 800 °C up to 10 h, which again was decreased at 11 h. As observed from XRD analysis, carbon content in the lattice of NbC has increased from 0.79 (5C800) to 0.959 (10C800) resulting in less distorted crystallites Table  which is−, h+, OH· and To establish the mechanism responsible for such a good photocatalytic behavior of synthesized photocatalyst, the generation of reactive oxygen species (ROS) has to be confirmed by using eet al.75 and Ishihara et al.76 also suggested the enhanced generation and transportation of charge carriers by inducing oxygen centers  in the carbide nanoparticles. Figure 77. Further, VB potential edge of 77. Being metallic in nature, NbC supported the charge transfer towards the surface of NPs where the reaction with the hydroxyl anion and dissolved oxygen in water produces hydroxyl free radicals , π→π* transitions and oxygen vacancies present on the surface. Among all of these parameters, the presence of NbCxOy and π→π* governs whole catalytic reaction . Moreover, oxygen vacancies promotes the delayed recombination (PL emission peak at 468 nm) which in turn provide more excitons78. But, excess vacancies can also quench the photocatalytic reaction which might be responsible for the observed reduced photodegradation in 7C800 as compared to 10C800. While, 11C800 possess insufficient vacancies providing lesser degradation efficiency than 10C800.For all the samples, degradation efficiency is a function of phase content  and 5+ (Nb2O5) oxidation states in which Nb5+ tries to attain a more stable state, i.e. Nb4+ by scavenging an electron to retard the photodegradation. Sample 10C800 consists of Nb in the 3+  and 4+  states in which Nb3+ will try to achieve Nb4+ by ejecting an electron to enhance the degradation efficiency which can be expressed by the following equations:Except for 10C800, both 7C800 and 11C800 contain higher carbonyl (C=O) group, considered as hole scavengers81. With the help of results of detection of ROS species, the generation of ROS species is a function of both photocatalyst and dye solution. In the case of MO dye, all the species are being generated among which hydroxyl radical is the most active oxidizing agent. But, degradation of MB dye occurred by the hole, electron, and hydroxyl radical equivalently where, holes dominated the degradation of RhB dye as compared to other ROS species. Thus, smaller crystallite size, high pore volume and specific surface area, smaller recombination, and fine tuning of bandgaps are considered to be prominent factors for the enhanced photocatalytic properties of single phase carbon coated NbC NPs. Further, to check the toxicity of the best degraded dye solutions where 10C800 photocatalyst was used, a toxicity test using IS: 6582-2001 method was done. In the adopted method, 100% survival of fish species was observed after 96 h indicating the non-toxicity of the degraded solution as shown in Fig. As the degradation can be achieved by the transformation of chromophore groups and the generation of energetic electrons in visible light, photocatalysis is thermodynamically and kinetically limited which might cause the partial degradation of RhB, MB and MO dyes even with the visible range band gap of photo catalysts based on MS and TOC resultsx NPs has been successfully synthesized by using Nb2O5, Mg, and charcoal at relatively low temperature (800 °C). However, the holding time at 800 °C for the synthesis has been optimized as 10 h to obtain carbon coated NbCx NPs in which x = 0.95 is obtained. The distorted lattice due to simultaneous transportation of carbon and oxygen has been supported by W-H analyses having the least distortion in 10C800. The as synthesized NPs are thermally stable up to 400 °C in air atmosphere  along with the both sp2 and sp3 hybridized carbon, oxygen vacancies confirming the removal of oxygen from Nb2O5 (reactant). Raman spectroscopy also confirmed the presence of disordered graphitic carbon as coatings of NPs. UV-visible spectroscopy affirmed the absorption of visible radiation corresponding to small bandgaps. The higher O-vacancies, less π → π* transition and lower NbC/NbCxOy content are responsible for the increased PL emission intensity in 5C800, 7C800, and 11C800 as compared to 10C800. Small crystallites, high SSA, and less PL emission intensity gave enhanced photocatalytic properties of optimized sample (10C800) that follow pseudo first order kinetics. Further, visible light photocatalysis is very complex phenomenon comprising of decolorization and degradation of dye solution with the help of photocatalyst. For visible light photocatalysis with synthesized single phase NbC (10C800), rate of decolorization is faster than that of degradation of dye solution in 120 min under visible irradiation, which is much better than that reported by Chen et al.52.From a series of experiments, carbon coated NbCere Fig. . The exi2O5 (1.329 g), metallic Mg powder (2.0 g) and activated charcoal (1.5 g) were used as the niobium source, reducing agent, and carbon source respectively. The mixture of reactants was heated to different temperatures  at a heating rate of 5 °C/min for different holding times  in a specially designed stainless steel autoclave50. The autoclave was allowed to cool within the furnace to room temperature. Thereafter, the obtained black powder was leached in 50% diluted HCl (v/v) and then washed with distilled water. The washed powder sample was further dried at 120 °C. The obtained powder samples were assigned as 10C600, 10C700, 5C800, 7C800, 10C800, and 11C800 corresponding to their synthesis temperature and holding time as listed in Table To synthesize carbon coated NbC nano-composite, Nb2O5 .329 g, m5/2:d3/2) with doublet separation of 2.74 eV. Morphologies of synthesized powder samples were observed with transmission electron microscopy (TEM) . Line profile of nanoparticle was also taken using Cu grid in STEM mode to study the linear distribution of elemental concentration. Thermal analysis (TG) of the samples were carried out with a NETZSCH STA 449F3 at a heating rate of 5 °C/min up to 850 °C in air atmosphere to predict the thermal stability as a function of temperature. The surface area and pore size distribution calculations were done with the help of N2 adsorption-desorption analyses (Tristar 3000). Optical absorption spectra of the samples were recorded on a double beam UV-Visible spectrophotometer (Hitachi U-3900H) in the range of 350–700 nm. Raman spectroscopic signals were recorded with a Renishaw Raman Microscope at 50X magnification with 785 nm (wavelength) laser source over a range of 100–1800 cm−1.For the identification of phases, X ray diffraction (XRD) analysis of synthesized samples was performed on a PANalytical Xpert-Pro diffractometer with Cu-Kα radiation  over 20 ≤ θ ≤ 80° and X’pert High Score was used to match the diffraction peaks with standard ICDD cards. Surface chemical compositions of synthesized samples (synthesized at 800 °C) were analyzed by X-ray photoelectron spectroscopy (XPS)  using Al-Kα radiation (1486.7 eV) and C1s (284.5 eV) for the calibration of binding energies of all the elements. All the high resolution XPS spectra were analyzed with the help of XPSPEAK41 software by considering the equal FWHM of doublet peaks of Nb3d keeping area ratio of 3:2 , sodium sulphate , ammonium oxalate  and ascorbic acid  were added to dye solution in a 1:10 volume ratio prior to irradiation in separate systems.For photocatalytic studies, a concentration of 1 mg/L of dye  was used. The photocatalyst (0.002 g) was suspended in the 100 mL aliquot after homogenizing the dye solution. For the establishment of adsorption-desorption equilibria, the suspended solution was stirred for 60 min in a dark chamber. Thereafter, the solution was exposed to visible radiation of around 8900 lux intensity for 120 min in which the sample aliquots were extracted at an interval of 30 min. The photo reactor was open to air atmosphere to avoid the depletion of dissolved oxygen during the irradiation. Each sample was centrifuged at 4500 rpm to extract photocatalyst from the suspension and supernatant was characterized with a double beam UV-Visible spectrophotometer (Hitachi U-3900H) to observe the absorption spectra of dye solution at regular time interval. To confirm the variation of absorbance of dye solution due to decolorization or degradation, the filtered aliquots  after 120 min irradiation were characterized by mass spectroscopy (MS) and total organic carbon (TOC). MS was carried out on UPLC-XEVO-G2-XS/QTOF mass spectrometer equipped with electron ionization source operated in positive ion mode. The mass spectrometer ion source and desolvation temperatures were set to 120 °C and 350 °C, respectively. Further, TOC was measured by MULTI N/C 3100 (N3-800/O) with the gas flow rate 160 mL/min and furnace temperature 800 °C. Further, the generation of reactive oxidative species (ROS) in the photodegradation of dyes were observed with Supplementary Information


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Great research efforts have been devoted to the recycling of these disilanes into monosilanes to allow reintroduction into the siloxane production chain. In this work, disilane cleavage by using alkali and alkaline earth metal salts is reported. The reaction with metal hydrides, in particular lithium hydride (LiH), leads to efficient reduction of chlorine containing disilanes but also induces disproportionation into mono‐ and oligosilanes. Alkali and alkaline earth chlorides, formed in the course of the reduction, specifically induce disproportionation of highly chlorinated disilanes, whereas highly methylated disilanes (n>3) remain unreacted. Nearly quantitative DPR conversion into monosilanes was achieved by using concentrated HCl/ether solutions in the presence of lithium chloride.The industry‐scale production of methylchloromonosilanes in the Müller–Rochow Direct Process is accompanied by the formation of a residue, the direct process residue (DPR), comprised of disilanes Me Too valuable for disposal or incineration: Simple recycling of the Müller–Rochow Direct Process residue with LiH yields monosilanes suitable for the reintroduction into the silicone production chain. In search of alternative hydride sources, we reacted methylchlorodisilane mixtures mimicking the DPR with lithium hydride. To our surprise, we found that most disilanes are efficiently cleaved, resulting in formation of mostly bifunctional monosilanes in high yields (see below). These findings prompted us to conduct further investigations on the disilane hydrogenation and cleavage reactions with alkali‐ and alkaline earth hydrides. The results of these studies are reported in the following.Recently, we have reported on the competitive chlorination and cleavage of methylhydridodisilanes with ether/HCl solutions to yield bifunctional monosilanes in excellent yields.4,44.24 and NaBH4 for reduction of, for example, SiCl4, Me2SiCl2 or GeCl4.4A number of earlier studies by others have shown that the hydrogenation of chlorosilanes can efficiently be achieved by using complex reducing agents, such as LiAlH2SiCl2 to Me2SiHCl with organotin hydrides in the presence of phosphonium chlorides or amine bases.2 as hydrogen source have been reported.2SiHCl and RSiHCl2 has been established by selective chlorination of hydridomonosilanes with HCl in the presence of catalytic amounts of Lewis acidsBifunctional monosilanes represent fundamentally important building blocks in silicone technology. Utilizing 1) the Si−H functionality for hydrosilylation reactions to create silicon–carbon bonds2H6 with alkali metal chlorides and hydrides to yield SiH4, ‐(SiH2)n‐ polymers and metal silanides,The efficient cleavage of silicon‐silicon bonds with alkali metal salts has been first reported by Ring and co‐workers.2MeSi−SiMeCl2  and ClMe2Si−SiMeCl2  are referred to in the literature as “cleavable fraction” of the DPR, because they can be thermally cleaved by amine catalyzed reaction with HCl.2Si−SiMe2Cl , Me3Si−SiMeCl2 , Me3Si−SiMe2Cl  and Me3Si−SiMe3  as their cleavage requires harsh conditions.3,Highly chlorinated disilanes such as Cl4 maintaining the Si−Si and the Si−C−Si backbone.Table 3) as a representative model compound for the “uncleavable” fraction of the DPR. 3 was quantitatively reduced at room temperature (RT) to tetramethyldisilane 18 with two equivalents of lithium hydride [Eq. To study the reduction and cleavage of disilanes with lithium hydride, we chose tetramethyldichlorodisilane (ride Eq. .(1)ClMe218 to form dimethylsilane (94 %) together with the NMR‐detectable oligosilanes HMe2Si−(SiMe2)n‐SiMe2H .1618 even at elevated temperatures. Also in contrast to earlier reports,16 afforded MeSiH3 (44 %) along with oligosilanes (4 %), 52 % of disilane 16 remained uncleaved. Upon increasing the reaction temperature to 60 °C, only MeSiH3 was identified by 29Si NMR spectroscopy. The reaction of 18 with excess LiH was studied further in variable‐temperature NMR experiments: Me2SiH2 was, apart from traces of oligosilanes, the only product detectable up to 140 °C.Increase of the reaction temperature to 140 °C with excess of LiH led to cleavage of 2Cl6,18 , which subsequently undergoes Si−Si bond cleavage to give Me2SiH2 (M) and the silanide anion HMe2Si− (A−).A− can then abstract a protonD to yield monosilane Me2SiH2 (M) along with the higher silanide anion HMe2Si−SiMe2− (−B). A quantum chemical assessment of this step at the SMD(THF)‐M062X/6‐31+G level reveals a moderate exoergicity  and an activation barrier of Δ≠G=28 kcal mol−1 n−SiMe2H results, which eventually become insoluble and escape NMR spectroscopic identification (for characterized species with n=1–4 see Supporting Information). Overall, this scenario is in line with the previous work of Ring and co‐workers3 nor 18 react with LiCl, even at temperatures as high as 220 °C.Based on earlier detailed studies on the chloride‐induced aufbau of higher perchlorinated oligosilanes from Si1), chosen as a representative model for the “cleavable” fraction of the DPR, does not form dimethyltetrahydridodisilane 16 upon reaction with 4 equiv LiH but undergoes quantitative cleavage into monosilanes. Most notably, the industrially important bifunctional silanes MeSiHCl2 and MeSiH2Cl comprise almost 80 % of the product mixture obtained at RT. We optimized their yield by using substoichiometric amounts of LiH: The reaction with 1.3 equiv LiH yields 8 and 9 in almost 90 %, with 8 in significant excess .In contrast to our observations for 1 with LiH initially leads to partial reduction and kinetically favored Si−Si bond cleavage sets in, once sufficient amounts of LiCl have formed. The resulting monosilanes, in turn, are then partially reduced by LiH to yield the bifunctional monosilanes observed in the experiments. These findings complement our related study on the disilane cleavage with phosphonium chloride saltsWe thus conclude that reaction of disilane 1 (69 mol %), 2 (26 %), 3 (4 %), and 4 (1 %) dissolved in diglyme, with LiH . After about 30 min at 60 °C disilanes were almost quantitatively consumed and the bifunctional monosilanes MeSiH2Cl (33 %), MeSiHCl2 (21 %), and Me2SiHCl (9 %) were formed along with MeSiCl3, Me2SiCl2, and Me3SiCl (together 23 %) and MeSiH3 (13 %). As detailed in the Supporting Information, use of substoichiometric amounts of LiH leads to predominant formation of bifunctional monosilanes in up to 70 %, whereas use of 400 mol % LiH results in complete reduction to Me2SiH2 (6 %) and MeSiH3 (78 %) together with formation of hydridodisilanes 16–19 (16 %) that remained stable under the reaction conditions. Treatment of a mixture of the highly methylated chlorodisilanes 3–4 with different molar amounts of LiH, in turn, led to partial disilane reduction while cleavage reactions were not detected in significant amounts in most cases. Only the reaction with a high excess of LiH (350 mol %) led to 50 % disilane cleavage at 140 °C (cf. Table S10).For further scrutiny, we investigated the reaction of an authentic sample of a highly chlorinated DPR fraction, comprised of 30 (45 %), 31 (31 %), 32 (14 %), 34 (10 %) and 35 (1 %) was reacted with excess LiH . Heating the sample to 180 °C led to carbodisilane reduction and Si‐C cleavage to give MeSiH3 (37 %) and Me2SiH2 (31 %) as main products, along with the hydridocarbodisilanes 36–39 (32 %).Generally, the disilane fraction of the DPR is contaminated with carbodisilanes. A representative mixture of 2, and MgCl2 is impeded by their lower solubility. Acceptable reaction rates, however, were found in tetraglyme at 140 °C and above. In exemplary reactions performed at a preparative scale with both, LiCl and KCl, the DPR mixture was efficiently converted: with a maximum theoretical yield of 50 % the disproportionation led to 42 % monosilanes . The residue remaining after distillation of the monosilanes consists of highly methylated disilanes, carbodisilanes and oligosilanes. A broad signal at +35 ppm in the 29Si NMR spectrum of the sample was assigned to branched oligosilanes with terminal Cl2MeSi groups. Slightly higher conversion ratios were obtained with LiCl at 220 °C.The suitability of other alkali and alkaline earth metal chlorides to induce disilane disproportionation was investigated in reactions with an industrial DPR mixture . Although disilane conversion was found most effective with LiCl in diglyme, the use of NaCl, KCl, CaCl4. We have further shown that LiCl, formed in the course of the reduction of chlorinated disilanes with LiH, acts as an efficient catalyst to trigger disproportionation of disilanes bearing SiMeX2 groups  into the corresponding mono‐ and higher oligosilanes.3, Me2SiH2, and Me3SiH and oligosilanes.In summary, we have shown that chlorosilane reduction is possible with lithium hydride, which thereby is established as economically favorable alternative to LiAlHnSi2Cln6‐ (n=2–6) were isolated from the DPR by fractional distillation and investigated as pure model compounds or in complex mixtures. The reactants for example, HCl/ether solutions, catalysts and solvents were placed in an NMR tube under nitrogen atmosphere and cooled to −196 °C, subsequently the disilanes were added and frozen. Then the NMR tube was evacuated (at −196 °C) and sealed in vacuo to avoid losses of low boiling monosilanes, such as MeSiH3 (b.p. −58 °C), Me2SiH2 (b.p. −20 °C), MeSiH2Cl (b.p. −46 °C), MeSiHCl2 (b.p. 41 °C) and Me2SiHCl (b.p. 35 °C). After warming the mixture to RT the reaction temperatures were increased, and the course of reaction was followed by NMR spectroscopy, especially by 29Si NMR. The molar ratios of products formed were determined by integration of product specific NMR signals o f the resulting mixtures. According to the optimum reaction conditions evaluated from the NMR investigations, upscaling was performed with larger amounts of starting materials in closed reaction ampules. Filling of reactants was similar as described for the experiments in sealed NMR tubes. Alternatively, upscaling was performed in open systems. This procedure is described in the Supporting Information.For the elucidation of the reaction conditions, disilanes MeThe 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.


Paradoxically, it evolved independently in numerous plant lineages, despite requiring specialised leaf anatomy. The anatomical modifications underlying C4 evolution have previously been evaluated through interspecific comparisons, which capture numerous changes besides those needed for C4 functionality. Here, we quantify the anatomical changes accompanying the transition between non‐C4 and C4 phenotypes by sampling widely across the continuum of leaf anatomical traits in the grass Alloteropsis semialata. Within this species, the only trait that is shared among and specific to C4 individuals is an increase in vein density, driven specifically by minor vein development that yields multiple secondary effects facilitating C4 function. For species with the necessary anatomical preconditions, developmental proliferation of veins can therefore be sufficient to produce a functional C4 leaf anatomy, creating an evolutionary entry point to complex C4 syndromes that can become more specialised.C Leaf samples from 44 of the 50 accessions were collected from their original field site and preserved until embedding was possible. For the remaining six accessions, leaf samples were taken from plants grown under controlled environment conditions as in Lundgren et al. , outer sheath and vein areas were measured separately for secondary, tertiary and any minor veins. The area of M tissue was calculated as the total area remaining after accounting for all other tissue types. In addition, the cross‐sectional area of individual M and BS cells (hereafter ‘size’) was measured . Although the depth of individual cells can vary, it is their cross‐sectional areas, and not their three‐dimensional volumes, that primarily influence the proportion of each tissue in the leaf.The cross‐sectional area of the whole segment, combining M, BS, epidermis and bulliform cells, extraxylary fibres and BS extensions as well as any transverse veins or tear spaces was measured. For all accessions, the total BS . We chose predictor variables that were likely to influence the M : BS ratio, including the number of M cells between major veins, average size of individual M cells, number of minor veins per segment, leaf thickness and the average size of BS cells on tertiary veins.We used a linear discriminant analysis (LDA) to explain the variation between photosynthetic types (i.e. the test maximises between‐group variance while minimising within‐group variance). We performed the LDA on the 50 accessions with leave‐one‐out cross‐validation and then bootstrapping over 100 runs, using the MASS package in R  and 69 non‐C4  accessions distributed across the species’ geographic range.To determine whether the pattern of vein density observed in the main data set was maintained across a larger sample of 3, C3‐C4 and C4 types. Recently expanded mature leaf tissue was prepared for transmission electron microscopy and processed for immunodetection of the large subunit of Rubisco (RBCL) and glycine decarboxylase H subunit (GLDH) as previously described  after growing for a minimum of 3 months in a common growth chamber, with conditions as described in Lundgren et al. To determine the degree to which the various leaf anatomical phenotypes arose from plastic development responses to their differing native growth environments, we compared field phenotypes to those obtained from live tillers of 17 4‐compatible phenotypes could emerge from plastic responses to the environment, as previously suggested  and one C3  plant were initially grown from seed in a controlled environment chamber set as described in Lundgren et al. −1 CO2 concentration. Both plants were split into five replicate cuttings and kept in the same growth chamber conditions to re‐establish for four months. From each replicated clone, one fully expanded, mature leaf was sampled and fixed in 4 : 1 ethanol : acetic acid solution. The growth chamber was then set to 180 μmol mol−1 CO2 concentration for the next four months to promote photorespiration, while maintaining the other environmental conditions, and one new fully expanded leaf was again sampled and fixed. All leaf samples from the 400 (i.e. ambient) and 180 (i.e. low) CO2 treatments were embedded, sectioned and imaged as described above. To determine whether the plants used different photosynthetic pathways in the two CO2 growth environments, we determined CCP and carboxylation efficiency, as described in Lundgren et al. , the M area is over 10 times larger than the BS area in C3 accessions, but less than five times larger in C4 accessions (Data set S1). As expected, C3‐C4 accessions are intermediate in their overall leaf anatomy, with 5–10 times more M than BS. These M : BS ranges are consistent with those measured in other C3 and C4 grasses , C3  and C3‐C4  accessions is also statistically indistinguishable from 1.0. The analysis, therefore, confirms that leaf anatomy varies among photosynthetic types in a statistically predictable manner.Variation in M : BS ratios can arise via changes to several underlying traits  and the number of minor veins , compared to C3‐C4 (3–6) and C3 (5–11) plants . However, the number of M cells between major veins overlaps between the C4 and non‐C4 groups  than is found in non‐C4 accessions . As the area of extraxylary fibres per vein does not differ between the photosynthetic types , the increased fibre area in C4 plants derives entirely from their greater vein density.The proliferation of minor veins explains a number of patterns associated with C4 and non‐C4 accessions and is therefore linked to the emergence of a fully functioning C4 physiology from a C3‐C4 intermediate state. Evolutionary changes that happened once this C4 physiology was in place would be restricted to some, but not all, C4 individuals. In our data set, such changes include further reductions to the M : BS ratio, potentially achieved via contractions to M airspace . The BS cell enlargement was therefore involved in the adaptation of C4 physiology after it had emerged, possibly to accommodate more or larger organelles for a more efficient C4 cycle, rather than being involved in its origin. Occasional hybridisation between C4 and non‐C4 individuals could affect the distribution of trait values; however, non‐C4A. semialata individuals are restricted to Africa, so that hybridisation outside of Africa is unlikely. Yet, Asia and Australian accessions exhibit some of the smallest BS cells among C4 accessions .The development of minor veins explains the overall anatomical difference between Cce Byott  or incre4 accessions are also present in C3‐C4 individuals, but not C3 ones, indicating that they are not associated with the transition to fully functional C4 physiology, but might have facilitated it. These include a small increase in BS cell sizes in C3‐C4 compared with C3 plants and a decrease in outer sheath cell size, with C3‐C4 accessions bridging the anatomical gap between C3 and C4 outer sheath cell sizes  in the common environment. Moreover, C3‐C4 plants produced thicker leaves (P = 0.040), such that leaf thickness of the three photosynthetic types converged in the common environment, which is likely a result of the non‐limiting light, nutrients and water available in these conditions. However, the other traits were not influenced by growth conditions and leaf anatomy of the three photosynthetic types remained distinct when grown in the common environment.3 and C4A. semialata under current ambient (400 ppm) and the Pleistocene minimum (180 ppm) CO2 concentrations. Plants grown under the low CO2 concentration experience elevated photorespiration rates, which might have induced a more C4‐like anatomy. However, we found that plants did not shift photosynthetic state under the differing CO2 conditions . Both C3 and C4 plants produced thinner leaves in the low CO2 environment (P = 0.0049 C3/0.0065 C4), and C4 plants developed smaller BS cells , probably because the lower carbon supply restricted development , larger BS cells or fewer M cells between veins when grown under this high photorespiration condition. These results show that, even when photorespiration is high, a C4‐like phenotype is not plastically induced in C3A. semialata.We further verified that historical changes in atmospheric composition did not influence the leaf phenotype by comparing C2 fixed via the C4 cycle. The emerging model of C4 evolution involves gradual and overlapping phenotypic changes , although the small sample of populations prevents us from differentiating ultrastructural changes linked to the transition to C4 from those that happened later.Photosynthetic types form a continuum, along which multiple biochemical, anatomical and ultrastructural alterations increase the proportion of CO4 photosynthesis in A. semialata may have a number of physiological and ecological consequences. First, the increase in vein frequency is accompanied by an enhancement of unpigmented extraxylary fibres, which improves light transmission to the BS, and thus ATP production in these cells, facilitating photosynthetic carbon reduction (Bellasio & Lundgren et al. et al. et al. et al. et al. The change in venation inferred during the evolution of C4 to C4 forms of Flaveria (McKown & Dengler 3 and C4Flaveria species (Kümpers et al. 4 grasses that primarily localise Rubisco within the mestome sheath. The insertion of minor veins could occur via relatively few developmental changes, likely underpinned by changes to auxin, brassinosteroids, SHORTROOT/SCARECROW and/or INDETERMINATE DOMAIN transcription factors (Kumar & Kellogg et al. et al. 4 plants (Nelson et al. et al. Our results complement those from previous comparisons among species, which show that an additional order of minor veins develops during the evolutionary transition from non‐C4 to C4 states can be caused by a single developmental alteration is a plausible explanation for the recurrent origins of C4 leaf anatomy and helps to resolve the paradox of how this complex trait emerged so many times. We also show that organelle number and size differ among photosynthetic types of A. semialata, but, here too, recent work indicates that one gene can control multiple ultrastructural modifications (Wang et al. 4 to C4 in A. semialata (Dunning et al. 4 state to the form of C4 photosynthesis observed in A. semialata involved relatively few genetic mutations.The possibility that a transition from non‐C4 anatomy in A. semialata is partially explained by the presence of relatively enlarged BS in the C3‐C4 accessions, while C3A. semialata BS size is similar to C3 species from other grass lineages (Lundgren et al. et al. 3‐C4A. semialata are also characterised by fewer M cells compared to C3 accessions and higher BS organelle abundance (Figs 3‐C4 physiology eased the subsequent transition to a full C4 state, but it is important to note that the physiology and anatomy of C3‐C4A. semialata are typical for C3‐C4 plants in general (Lundgren et al. 3 grasses (Hattersley et al. et al. 4 anatomy evolved in A. semialata is therefore not exceptional.The limited number of changes involved in the emergence of C4 leaf anatomy can arise from one key developmental modification is apparently incompatible with the great anatomical specialisation of other C4 lineages as well as the large phenotypic gaps separating them from their closest C3 relatives (Dengler et al. et al. 3 and C4 sister lineages are separated by long periods of evolution, and comparing these groups therefore captures all of the changes that happened after the origin of C4 photosynthesis to improve the efficiency of C4 physiology and adapt it to various organismal and ecological contexts (Christin & Osborne 4A. semialata than in species from some older C4 lineages (Lundgren et al. et al. 4 photosynthesis in A. semialata may represent a rudimentary version of the physiological trait (Ueno & Sentoku 4 cycle in A. semialata may be one reason for this (Bräutigam et al. A. semialata populations (Lundgren et al. A. semialata, enlargements of the BS cells beyond those seen in non‐C4 individuals are restricted to a subset of C4 populations (Fig. 4 physiology. Over time, accumulated modifications will move C4 leaf anatomy far beyond that realised via a single developmental change. However, the fact that an initial C4 phenotype and the associated physiology can be accessed via a single modification likely placed multiple groups on a selective highway to highly specialised and successful variants of the C4 syndrome.Our conclusion that Cons Fig.  and thusMRL, PAC and CPO designed the study. MRL produced and analysed the data, with the help of LTD, JJMV and JWB. TS and RK performed the immunolocalisations and TEM imaging. SS assisted with immunolocalisation sample preparation. MS assisted with tissue fixation. MRL, LTD, JKO, BR, MSV, GB, CA, NC, AM, MB, CML, PAC and CPO contributed plant material. MRL, PAC and CPO interpreted the results and wrote the paper, with the help of all the authors. Click here for additional data file. Click here for additional data file. Click here for additional data file.


The band structure and photoelectrochemical stability of photoelectrodes are found to be key factors determining the charging performance of sunlight-promoted zinc–air batteries. The introduction of photoelectrode as an air electrode opens a facile way for developing integrated single-unit zinc–air batteries that can efficiently use solar energy to overcome the high charging overpotential of conventional zinc–air batteries.Directly harvesting solar energy for battery charging represents an ultimate solution toward low-cost, green, efficient and sustainable electrochemical energy storage. Here, we design a sunlight promotion strategy into rechargeable zinc–air battery with significantly reduced charging potential below the theoretical cell voltage of zinc–air batteries. The sunlight-promoted zinc–air battery using BiVO The authors here report a sunlight-promoted rechargeable zinc–air battery in which photoelectrode is used as the air electrode to substantially lower the charge potential under illumination. Notably, the battery can be initially charged with an extremely low voltage of ~1.20 V. Rechargeable metal–air battery, on the other hand, can store (and release) energy repeatedly as needed via oxygen evolution reaction (OER) but this process is essentially plagued by the large overpotential caused by sluggish OER kinetics2. Therefore, a combination of these two battery systems would theoretically guarantee a clean and energy-efficient functional integrity, which is however, poorly explored so far.Solar cell and rechargeable metal–air battery are two totally different systems converting energy of different forms into electricity. Intrinsically, solar cell can only convert light into instant electricity for immediate usage but cannot store energy7, lithium–oxygen battery has been the mostly focused system to be successfully loaded with a photo-responsive electrode where the in situ generated holes possessing positive charges directly5 or indirectly4 catalyze the OER. Therefore, the notorious large overpotential of the OER has been efficiently suppressed via this synergistic process where solar energy is nominally stored into the rechargeable battery system and further consumed to compensate the energy needed for battery charge. Yet, the high cost and poor safety issues of lithium–oxygen battery make it unsuitable to be coupled with the widely commercialized solar cells aiming at low cost and reliable safety, which thus drives the search for appropriate battery systems.Among the very few reported studies−1 in theory), environmental friendliness, earth-abundance of Zn, low cost, and excellent safety enabled by the use of aqueous electrolytes9. Although primary zinc–air batteries have been commercialized11, the development of rechargeable zinc–air batteries is still in the early stages with several major challenges that should be addressed. One of the most significant challenges is the high charge overpotential of the battery, which is caused by the practical charge voltage typically above 2 V despite the calculated theoretical voltage of 1.65 V8. Therefore, the energy efficiency is generally lowered below 60%12. Besides, the high charge voltage results in the possible degradation of air electrode due to the side reactions during the charging process, greatly reducing the cycle life of rechargeable zinc–air batteries14. The large charge overpotential is essentially caused by the sluggish kinetics of oxygen evolution reaction (OER) process that proceeds through four proton-coupled electron transfer steps at the air electrode during the charging process15. Therefore, the improvement of the OER kinetics is the central scheme for the further development and commercialization of rechargeable zinc–air battery. In this sense, it should be beneficial to effectively harvest the solar energy using the zinc–air battery system and meanwhile to reduce the charge potential with input electric energy.Among the candidacy of various rechargeable battery technologies, rechargeable zinc–air batteries have attracted great interest thanks to their high energy density (1086 Wh kg4 and α-Fe2O3) with different band structures, the effect of the band structure of the semiconductor photoelectrode on the charging performance of the zinc–air battery has been clearly demonstrated and the corresponding sunlight-promoted charging mechanism has been proposed. BiVO4- and α-Fe2O3-based sunlight-promoted zinc–air batteries exhibit an ultra-low initial charge potential of 1.20 and 1.43 V, respectively, even lower than the theoretical cell voltage of zinc–air batteries. Sunlight-promoted zinc–air batteries with α-Fe2O3 photoelectrode show robust cycling stability due to the high photoelectrochemical stability of α-Fe2O3 electrode. The concept of sunlight-promoted OER in rechargeable zinc–air battery offers an effective strategy to store the in situ converted energy from the solar cell system and more importantly, to enhance the OER kinetics of zinc–air battery toward its wide application in future.In present work, we report a sunlight-promoted rechargeable zinc–air battery by integrating semiconductor photoelectrode as air electrode to significantly lower the charge potential by ~0.5–0.8 V under sunlight illumination. By utilizing two typical semiconductor photoelectrodes  as an air electrode assembled in alkaline electrolyte . The following step is the reaction of OH– with the adsorbed OH species to produce H2O and adsorbed atomic O and the release of an electron (reaction 2). After that, the step in the sequence involves the reaction of an OH– with an adsorbed O atom to form adsorbed OOH species (reaction 3), which then undergo reaction with additional OH– (reaction 4), leading to the formation of adsorbed O2 and H2O and the release of an electron. Adsorbed O2 then desorbs in the last step of the sequence (reaction 5).Figure yte Fig. . The dis−. The photoexcited holes simultaneously migrate to the photoelectrode surface to oxidize water to oxygen. The reactions for a sunlight-promoted zinc–air battery during discharging and charging are illustrated in Eqs.  of semiconducting photoelectrode and further transferred to the Zn electrode through the external circuit, resulting in the reduction of mentclasspt{minima19Discharge process:Charge process:− to O2 should fulfill the essential condition that the valence band (VB) potential lies higher than the O2/OH− couple potential (0.401 V vs. standard hydrogen electrode (SHE)). As illustrated in the proposed sunlight-promoted charging mechanism  of electrons, which, at its most negative situation, approaches to the conduction band minimum (CBM), in the semiconductor photoelectrode. Based on the above discussion, the band structure  of the photoelectrode significantly affects the charging performance of the sunlight-promoted zinc–air battery. To better understand the mechanism of sunlight-promoted charging process, for the first time, two typical semiconductor photoelectrodes with different band structures  were investigated simultaneously as air electrodes for zinc–air batteries.This mechanism is expected to reduce the charge potential of a zinc–air battery through the photooxidation reaction by photogenerated holes to facilitate the OER reaction in the sunlight-promoted charging process. The requirement of photoelectrode in the sunlight-promoted rechargeable zinc–air battery for conversion of OHism Fig. , the pho4 and Fe2O3 photoelectrodes are analyzed in Fig. 4 photoelectrode exhibits a structure identical to monoclinic scheelite BiVO4 (JCPDS Card No. 14–0688)20. The Fe2O3 photoelectrode is indexed to α-phase Fe2O3 (JCPDS Card No. 33–0664)21. Strong (110) diffraction peak implies preferential growth of hematite nanorods in the [110] direction22. All the other peaks are indexed to be SnO2 (JCPDS Card No. 46–1088) from the fluorine-doped tin oxide (FTO) substrate. The surface chemical compositions of both BiVO4 and α-Fe2O3 are confirmed by X-ray photoelectron spectroscopy (XPS). Apparently, the characteristic spin–orbit split of the Bi 4f5/2 and Bi 4f7/2 signals  reveals that the obtained BiVO4 product consists of interconnected columnar-like particles with an average particle size of ~110 nm, forming a three-dimensional (3D) porous network  above ~1.4 V for water oxidation in the dark condition. In sharp contrast, the BiVO4 and α-Fe2O3 electrodes show very low onset potential below ~0.9 V under simulated sunlight illumination. Furthermore, the photocurrent increases linearly with the increasing voltage and achieves high values of 0.79 and 0.60 mA cm−2 at 1.23 V vs. RHE, respectively, which are more than one order of magnitude larger than those under the dark condition. This indicates that the application of photoelectrodes dramatically accelerate water oxidation kinetics under illumination. The charge-transfer properties of the photoanodes were further investigated by electrochemical impedance spectroscopy (EIS)  analysis was performed at 1 kHz  and donor densities (ND) of two electrodes are calculated based on the slopes of M–S plots28, which are 0.17 V (vs. RHE) and 2.18 × 1020 cm−3 for BiVO4 and 0.47 and 1.25 × 1019 cm−3 for α-Fe2O3. It should be noted that the M–S analysis is derived from a planar electrode model29 and therefore these values are only for the purpose of relative comparison. Compared to α-Fe2O3 the negative shift of EFB and larger ND of BiVO4 electrode indicate that the Fermi level is closer to the CB31, leading to more significant band bending in the space charge region than that of α-Fe2O3 due to a larger difference between the Fermi level of BiVO4 photoelectrode and the redox potential of the electrolyte. This facilitates the charge separation efficiency of the photogenerated charge carriers33, which is consistent with the higher photocurrent density of BiVO4 electrode compared to α-Fe2O3 electrode  absorption spectra as shown in Fig. 4 and α-Fe2O3 extends to 504 nm and 590 nm, corresponding to the band gap energy (Eg) of 2.46 and 2.10 eV, respectively  can be estimated by linear extrapolation of the leading edges in the spectra34, which are deduced to be 2.11 and 1.70 eV for BiVO4 and α-Fe2O3, respectively.The optical absorption properties of BiVO2O3 and BiVO4 was measured using ultraviolet photoelectron spectroscopy (UPS). The work function of BiVO4 and α-Fe2O3 is 5.05 and 5.30 eV, respectively . Based on the charging mechanism mentioned above, the CB edge combined with To further analyze the band structure, the surface work function of α-Feely Fig. . Further4 and α-Fe2O3 photoelectrodes are shown in Fig. 4-based zinc–air battery achieves a very low initial charging voltage of ~1.20 V corresponding to the voltage reduction of ~0.76 V . Based on the estimated CBM value  and the discussion in Fig. 4 and α-Fe2O3 is expected to be 0.71 and 0.91 V, respectively. The initial charge potential of BiVO4-based zinc–air battery is ~0.23 V lower than the α-Fe2O3-based one, which matches well with the theoretical value of 0.20 V. Considering the factors including energy gap between the photoelectrode CBM and the EF of electrons under working condition, the charge recombination occurring in the photoelectrode–electrolyte interface, and the internal series resistance within the device, it is reasonable to obtain a higher actual charging voltage than the theoretical value.To demonstrate the potential application of photoelectrode as an air electrode, we assembled a single-body-structured zinc–air battery using the semiconductor photoelectrode. The charging curves of zinc–air battery in the dark and under illumination based on BiVO6 V Fig. , which tode Fig. . For α-F4 V Fig. , corresp2O3 Fig. . Both th2O3 Fig.  and the −2, the initial charging voltage (~1.35 V) of BiVO4-based battery is still much lower than the theoretical working voltage (1.65 V) of the zinc–air system; α-Fe2O3-based battery gives a larger charging potential of ~1.64 V, which is very close to the theoretical working voltage. The discharge–charge polarization curves for zinc–air battery using BiVO4 and α-Fe2O3 as the air electrode with and without light illumination are shown in Supplementary Fig. 4 delivers a lower charging voltage compared to α-Fe2O3 due to the higher OER photoelectrocatalytic activity of BiVO4. BiVO4- and α-Fe2O3-based batteries exhibit similar discharge behaviors in dark and under illumination, which is consistent with the oxygen reduction reaction (ORR) performance of the two photoelectrodes –(7)), which will increase the resistance of the zinc–air battery and make it failure to discharge. Therefore, the limiting factor for the long-term cycling of zinc–air battery is the passivation of zinc electrode in alkaline electrolyte. The structure and morphology of two samples after the cycling test are further explored. The transmission electron microscopy (TEM) images for α-Fe2O3 and BiVO4 materials before and after cycling tests are analyzed. It can be seen that the surface morphology of BiVO4 electrode exhibits significant dissolution after cycling test compared to the initial electrode, but the high-resolution TEM (HRTEM) images of the BiVO4 electrode remains a lattice spacing of 0.46 nm, which is well consistent with the (011) plane of crystalline BiVO4  plane of α-Fe2O3. The TEM results for α-Fe2O3 and BiVO4 materials after long-term cycling tests further indicate the high stability of α-Fe2O3. In addition, XPS tests for photoelectrodes before and after cycling test are conducted to understand the possible chemical changes of the materials. The XPS for BiVO4 before and after the cycling performance is provided in Supplementary Fig. 4 electrode during the cycling test. In contrast, there is no noticeable change in XPS spectra of α-Fe2O3 after stability test  as an example. The “light-response” charging feature of the Fe–air battery is shown in Supplementary Fig. 2O3. Once the sunlight is cut off, it can be seen that the applied voltage increases immediately. The example of the sunlight-promoted Fe–air batteries further demonstrates that the concept of utilizing solar energy to improve the oxygen evolution reaction kinetics can be extended to other metal–air batteries for addressing the charging overpotential issue.To further analyze the limiting factor for the long-term cycling of α-Fe4 and α-Fe2O3) with different band structures. BiVO4-based zinc–air battery can be initially charged with an extremely low voltage of ~1.20 V corresponding to ~0.76 V reduction compared to that in dark. However, it increases quickly due to severe photocorrosion of BiVO4. In contrast, α-Fe2O3 exhibits promising potential as stable and efficient air photoelectrode with a low charge potential of ~1.43 V and high cycling stability during 50 h discharge–charge test. It is also interesting to note that the difference in the actual charge potential (~0.23 V) of sunlight-promoted zinc–air batteries based on two different photoelectrodes matches well with the predicted theoretical value difference (0.20 V). This work demonstrates that the introduction of photoelectrode with more negative CB and appropriate VB positions as well as high photoelectrochemical stability is promising for the development of zinc–air batteries with high energy efficiency.In conclusion, we have successfully incorporated solar energy capturing into the charging process of a rechargeable zinc–air battery which integrates the semiconducting photoelectrode as air electrode. Under sunlight illumination, the photoelectrode undergoes photoexcitation and generates holes and electrons. The holes transfer to the photoelectrode surface to facilitate OER process, thus reducing the charging voltage. The proposed sunlight-promoted charging mechanism has been demonstrated using two typical photoelectrodes  substrate followed by thermal conversion to the porous BiVO4 film20. Briefly, 0.4 M KI solution was prepared by dissolving 20 mmol potassium iodide  into 50 ml deionized water and adjust its pH to 1.7 by nitric acid . Then, 2 mmol bismuth(Ш) nitrate pentahydrate (Bi(NO3)3·5H2O, ≥99.99%, Sigma-Aldrich) was added into the solution and stirred by 15 min. Afterwards, this solution was mixed with 20 ml of absolute ethanol containing 0.4 M p-benzoquinone , and was vigorously stirred for 30 min resulting in a dark transparent black solution with a pH of 2.6. A typical three-electrode system was carried out for the electrodeposition, with the prepared solution as electrolyte, a FTO substrate as the working electrode (WE), a platinum foil as the counter electrode (CE) and the saturated Ag/AgCl as the reference electrode. After potentiostatically cathodic deposition at an optimized condition of –0.1 V vs. Ag/AgCl for 5 min, a uniform BiOI film was deposited on the FTO substrate followed by rinsing with deionized water and drying in ambient air. BiVO4 electrodes were prepared by dropping 0.2 ml of a dimethyl sulfoxide  solution containing 0.2 M vanadyl acetylacetonate (VO(acac)2, 99.98%, Sigma-Aldrich) onto the formed BiOI layer, followed by annealing in a muffle furnace at 450 °C (ramping rate 2 °C min−1) for 2 h. To remove excess V2O5, the obtained BiVO4 electrodes were soaked in 1 M potassium hydroxide  after cooling to room temperature. Finally, the electrodes were rinsed with deionized water  and dried in ambient air to obtain pure BiVO4.The porous BiVO2O3 photoelectrode with one dimensional nanorod arrays on the FTO substrate was prepared by a hydrothermal method and two-step annealing. Typically, a cleaned FTO glass was immersed in a Teflon-lined stainless steel autoclave containing aqueous solution of 0.15 M ferric chloride  and 1 M sodium nitrate  at pH 1.25 adjusted by hydrochloric acid . The hydrothermal reaction was conducted at 95 °C for 6 h and naturally cooled down to room temperature. The obtained products were cleaned with deionized water and subsequently sintered in air at 550 °C for 2 h to form Fe2O3 nanorod arrays. Afterwards, the formed Fe2O3 product was further annealed at 800 °C for 10 min for facilitating the doping of Sn from the FTO into Fe2O3 to increase the donor density in the electrode.α-Feα radiation and operating in a 2θ range of 20–80° at a scan rate of 5° min−1. The surface morphology of the samples was characterized by field-emission SEM  and TEM . The Optical absorption measurements of as-prepared samples were performed using a spectrophotometer . VB analysis was carried out by XPS . UPS was measured with a monochromatic He I light source (21.2 eV) and a VG Scienta R4000 analyzer . A sample bias of −5 V was applied to observe the secondary electron cutoff. The work function can be determined by the secondary electron cutoff at the high kinetic energy region.The crystal structures of the obtained samples were determined by XRD  with Cu K-electrode configuration using BiVO4 or α-Fe2O3 as the working electrode, saturated calomel electrode (SCE) as reference electrode, graphite rod as counter electrode and 1 M KOH with pH of 13.6 as electrolyte. In order to ensure the dark environment, the quartz electrolytic cell is wrapped in tin foil. Simulated solar illumination was provided by a 500 W Xe lamp equipped with an AM 1.5 G filter  and the incident photo intensity was calibrated to 100 mW cm−2 by a solar power meter. Liner sweep voltammetry (LSV) was carried out in O2-saturated electrolyte with scanning rate of 20 mV s−1 in dark and under simulated solar illumination, respectively. Potentials were calibrated to a RHE using the Nernst equation:The photoelectrochemical and electrochemical performances of the photoelectrode were measured with potentiostat  in a three-electrode configuration using the electrochemical station with bias of 1.23 V vs. RHE in the dark and under illumination , respectively. Mott–Schottky measurements were carried out in 1 M KOH aqueous solution at a frequency of 1 kHz with a scan rate of 10 mV s−1 in the dark. The donor densities were calculated with the following equation:Nd is the donor concentration, e the electron charge (1.60 × 10−19), C the electron charge, ε the dielectric constant , C the capacitance of the space charge region, and V is the electrode applied potential.The photocurrent stability tests were conducted under continuous illumination at 1.23 V vs. RHE. Electrochemical impedance spectroscopy (EIS) measurements were performed in a three4 or α-Fe2O3) as the air electrode. The electrolyte used was 1 M KOH with 0.03 M zinc acetate (Zn(CH3COO)2, 99.99%, Sigma-Aldrich)  and the energy density (mWh g−1) based on the weight of consumed Zn are calculated according to Eqs. The zinc–air battery tests were performed by the home-made liquid zinc–air battery which was assembled using a polished Zn foil (0.25 mm thickness) as the metal electrode and the semiconductor photoelectrode (i.e., BiVOSupplementary Information


The conformations of AIE ligands in 1 are switchable by the absorption/desorption of guest molecules. As a result, 1 exhibited not only intense but also guest molecule switched luminescent properties. More importantly, the switching rate is tunable by using different guest molecules. 1 provides a unique visualized prototype to understand the mechanism of guest‐triggered aggregation‐induced emission in MOFs.Utilizing aggregation‐induced emission luminogens (AIEgens) as ligands has proven to be an effective strategy for constructing metal–organic frameworks (MOFs) with intense luminescent properties. However, highly luminescent AIEgen‐based MOFs with adjustable emission properties are rarely achieved because of the rigid conformation of AIEgens in the crystalline state. Here, a dual‐node 3D silver chalcogenolate cluster MOF ( Metal–organic frameworks (MOFs) with good luminescent properties have great potential in application areas such as chemical sensors and light‐emitting diodes.1⊃DMAC, DMAC = dimethylacetamide) was successfully synthesized 6(CF3CO2)6]0.5[Ag8(StBu)4(CF3CO2)4](tppe)2(DMAC)10}n . This formula was further confirmed by thermogravimetric analysis (TGA), elemental analysis, and 1H‐NMR spectroscopy . As shown in Figure 1⊃DMAC is built of Ag8 cluster  and Ag12 cluster  nodes. The Ag8 cluster can be considered as a deformed trigonal prism with two silver atoms hanging outside a prismatic edge, which are united by four StBu and four CF3CO2− auxiliary ligands . Each Ag8 cluster acts as a 6‐connected node, in which each of six silver atoms on the triangular planes is extended by one tppe via Ag—N bonding  bound by six StBu ligands and six CF3CO2− ligands ; four of the six silver atoms on the waist are connected by four different tppe ligands, making Ag12 a 4‐connected node , upon removal of the guest molecules.Single‐crystal X‐ray analysis reveals that Figure1⊃DMAC exhibited strong blue emission at 470 nm at room temperature (quantum yield 51.1%). When 1⊃DMAC were exposed in the atmosphere, its fluorescence gradually changed from blue to green and the maximum emission wavelength shifted to 532 nm (quantum yield 59.7%). Interestingly, this fluorescence change was reversible, as it turned back to blue immediately when the atmosphere exposed 1⊃DMAC was treated with DMAC again . To investigate its fatigue resistance, 1⊃DMAC was alternatively exposed to the atmosphere and treated with DMAC for 1 min. As shown in Figure 1⊃DMAC was investigated. As shown in Figure S5 in the Supporting Information, the emission intensity at 470 nm of 1⊃DMAC barely changed after being irradiated with 365 nm light for 3 h, indicating its good photostability.As shown in 1⊃DMAC was investigated first. As shown in Table S1 in the Supporting Information, the nitrogen content of the atmosphere exposed 1⊃DMAC was significantly lower than that of the as‐synthesized 1⊃DMAC. The composition of the atmosphere exposed 1⊃DMAC agreed well with the general formula ({[Ag12(StBu)6(CF3CO2)6]0.5[Ag8(StBu)4(CF3CO2)4](tppe)2}n, i.e., the composition of empty framework of 1. This result suggested that DMAC molecules have escaped from 1⊃DMAC by exposing in the atmosphere, producing 1. This result was supported by infrared spectroscopic measurements. As shown in Figure S6 in the Supporting Information compared with 1⊃DMAC, the peaks belonging to DMAC disappeared in the infrared spectra of 1. TGA results also confirmed that there was no DMAC in 1 . More direct evidence for the loss of DMAC from 1⊃DMAC in the atmosphere was obtained from 1H‐NMR spectroscopy. The as‐synthesized 1⊃DMAC and 1 were dissolved in deuterated dimethylsulfoxide (DMSO‐d6) and their 1H‐NMR spectra were recorded. As shown in Figure S8 in the Supporting Information, for the as‐synthesized 1⊃DMAC, the proton resonance peaks of tppe and DMAC could be well assigned. Integration of the proton resonance peaks demonstrated that the ratio of tppe and DMAC in 1⊃DMAC was 2:10, which fitted well with its composition. On the contrary, there was no signal of DMAC in the 1H‐NMR spectrum of 1, while the proton resonance peaks of tppe could still be observed. These results demonstrated that the atmosphere exposed 1⊃DMAC was actually 1. Meanwhile, from 1⊃DMAC to 1, the morphology of crystal remained unchanged even though its fluorescence emission changed significantly . As shown in Figure 1⊃DMAC transformed to 1 and the diffraction peaks decreased and even disappeared. This result suggested that 1 has a relatively disorder structure, which might originate from the rotations of phenyl and pyridyl groups since the loss of guest molecules released space around the tppe ligands. When 1 was treated with DMAC for 1 min, the diffraction peaks appeared again, which suggested the recovery of ordered structure of 1⊃DMAC . The limited adsorption capacity of 1 for N2, O2, and CO2 also hints that the framework structure of 1 might lose high order and become partially disorder. The loss of guest molecules released space around the tppe ligands in 1, which endowed it with more flexible twisting angles than that in 1⊃DMAC. As a result, the tppe ligands in 1 exhibit multiple absorption/emission characteristics which thus gives rise to broader absorption/emission spectra. As shown in Figure 1 in the emission spectrum was 105 nm, which was much broader than that of 1⊃DMAC (69 nm). Interestingly, the absorption spectrum of 1 was not only broader than that of 1⊃DMAC, but also red‐shifted, which indicated its better conjugated structure than 1⊃DMAC  was larger than that in 1 (path b). Thus, the emission wavelength of 1⊃DMAC was shorter than that of 1. The proposed decay paths are also supported by fluorescence lifetime measurements. As shown in Figure 1⊃DMAC and 1 are 0.59 and 3.28 ns, respectively. The steric hindrance between tppe and DMAC prevents the excited state distortion of tppe in 1⊃DMAC, resulting in a faster exciton energy release via fluorescence with a shorter lifetime. These results also suggested that the luminescence of the MOFs are not originated from the silver chalcogenolate clusters, whose luminescence lifetimes are normally tens to hundreds of nanoseconds.]Possible fluorescence decay paths for tppe in the MOFs are proposed in 1⊃DMAC were restricted by the incorporated guest molecules, resulting in rigid high‐energy excited state conformations for blue emission. As shown in Figure S10 in the Supporting Information, C‐H···O and C‐H···π weak interactions are found between DMAC and tppe ligands, indicating that there was interaction between solvent molecules and tppe backbone, which may hinder the rotation of the phenyl rings in tppe ligand to trigger the AIE property in the MOFs.As shown in Figure S3 in the Supporting Information, the determined DMAC molecules were mainly distributed nearby tppe ligands, which indicated that tppe in 1⊃DMAC suggested that tppe exhibits two conformations . It was noticed that the angles between the two adjacent C—C single bonds of tetraphenylethene moiety in Type III are 114.5° and 114.4°, respectively, which are highly similar to that in Type I (115.1° and 114.3°) and Type II (115.9° and 114.6°). This result suggested that the rotation of phenyl and pyridyl groups in tppe barely changed their extending directions. That is why the morphologies of 1⊃DMAC and 1 crystal are analogous  with different polarity were used for exchange with DMAC from 1⊃DMAC. First, 1⊃DMAC was soaked in the selected solvents separately for 1 h. Then the crystals were filtered out to obtain 1⊃THF, 1⊃Toluene, and 1⊃DMF. The 1H‐NMR spectra of the exchanged crystals were shown in Figure S12 in the Supporting Information. It was noted that the signals of DMAC disappeared after exchanging with the solvents, while the peaks of exchange solvents could be clearly observed. These results suggested that the DMAC in 1⊃DMAC had been completely replaced by THF, toluene, and DMF. Infrared spectra also confirmed the formation of 1⊃THF, 1⊃Toluene, and 1⊃DMF. As shown in Figure S13–S15 in the Supporting Information, the peaks of guest molecules could be observed in the corresponding infrared spectra of solvated MOFs. Interestingly, there was no distinct difference observed in the PXRD patterns for 1⊃DMAC, 1⊃THF, 1⊃Toluene, and 1⊃DMF , which demonstrated that their frameworks are uniform. Meanwhile, highly analogous emission spectra and fluorescence lifetimes were found in all these MOFs  based on AIE ligand was designed and synthesized. 1 exhibited intense and unique guest molecule‐switched fluorescence emission. The easily reversible transformation between the empty framework of 1 and the solvated 1⊃DMAC led to a ratiometric fluorescence change with good fatigue resistance. In addition, different guest molecules did not vary the fluorescence, but significantly changed the switching rate between 1 and guest‐absorbed 1. Mechanism study showed that the rotations of phenyl and pyridyl groups in tppe were restricted by the incorporated guest molecules in the pore of 1, resulting in rigid high‐energy excited state conformations for blue emission; otherwise, the free space for the rotors of tppe in the empty framework of 1 contributed to the other end of the switch– a yellow emission. In all, the reported AIEgen‐based MOFs provide a visualized prototype to understand the mechanism of guest‐triggered aggregation‐induced emission in crystalline materials. Currently, efforts toward developing more luminescent MOFs for sensing and light‐emitting materials are under investigation in our laboratories.In summary, a dual‐node 3D silver chalcogenolate cluster MOF (www.ccdc.cam.ac.uk/data_request/cif.][CCDC 1821839 contains the supplementary crystallographic data for this paper. These data can be obtained free of charge from The Cambridge Crystallographic Data Centre via The authors declare no conflict of interest.SupplementaryClick here for additional data file.SupplementaryClick here for additional data file.SupplementaryClick here for additional data file.


Detailed scanning transmission electron microscopy images clearly demonstrate that MnO6 octahedral rotations in La0.9Ba0.1MnO3 close to the interface are strongly suppressed. As determined from in situ X‐ray photoemission spectroscopy, O K‐edge X‐ray absorption spectroscopy, and density functional theory, the realization of the FMI state arises from a reduction of Mn eg bandwidth caused by the quenched MnO6 octahedral rotations. The emerging FMI state in La0.9Ba0.1MnO3 together with necessary coherent interface achieved with the perovskite substrate gives very high potential for future high performance electronic devices.Ultrathin epitaxial films of ferromagnetic insulators (FMIs) with Curie temperatures near room temperature are critically needed for use in dissipationless quantum computation and spintronic devices. However, such materials are extremely rare. Here, a room‐temperature FMI is achieved in ultrathin La 0.9Ba0.1MnO3 films, which could serve as promising candidates for future oxide‐based electronic devices.Ultrathin ferromagnetic insulators with Curie temperatures above room temperature are critically needed for developing dissipationless quantum electronic and spintronic devices. Unfortunately, such materials are extremely rare in nature. Room‐temperature ferromagnetic insulating states are successfully synthesized by an interfacial octahedral proximity effect in ultrathin La Some spinel ferrimagnetic insulators such as NiFe2O4 (TC ≈ 850 K),2O4 (TC ≈ 796 K),2NiMnO6 (TC ≈ 280 K)TC above or very close to room temperature. However, their complex structure, tendency to form antisite defects, and poor chemical compatibility are obstacles to achieving high‐quality interfaces with electrodes, which significantly reduce the performance of devices. Consequently, ferromagnetic insulators with a higher TC need to be further explored for realizing the spin‐polarization at room temperature.Ultrathin ferromagnetic insulators  with Curie temperatures above room temperature and forming coherent interface with electrodes are strongly needed for generating pure spin‐polarized currents in next‐generation dissipationless quantum electronic and spintronic devices.Room‐temperature ferromagnetic metallic manganite perovskites are excellent electrode materials and are widely used for the spin detection layer.0.9Ba0.1MnO3 (LBMO) films grown on SrTiO3 (STO) (001) substrates. Bulk LBMO (pseudocubic lattice parameter ≈3.89 Å) is a ferromagnetic insulator (TC ≈ 185 K) whereas thin films of LBMO grown on STO substrates typically show a ferromagnetic metallic state with TC of 250–295 K.K‐edge X‐ray absorption spectroscopy (XAS)) and density functional theory (DFT) calculations, we identified that, when MnO6 octahedral rotations in LBMO close to the interface are strongly suppressed, the Mn eg bandwidth is reduced and magnetic interactions are enhanced simultaneously. This leads to the emergence of ferromagnetic insulating state in an ultrathin (5 unit cell) LBMO film with a remarkable TC above room temperature.In this study, we explore the utilization of the interfacial octahedral proximity effect to achieve a room‐temperature ferromagnetic insulating state in ultrathin La2‐terminated STO (001) substrates. The film thickness was controlled by in situ monitoring of the reflection high‐energy electron diffraction (RHEED) oscillations.Figurea−a−c+ rotation pattern (Pbnm) with Mn–O–Mn angles of 159.20°(in‐plane) and 162.18°(out‐of‐plane),a0a0a0 pattern (Pm‐3m).We used pulsed laser deposition to fabricate high‐quality LBMO epitaxial thin films on TiO3 perovskites, the relatively heavier cation sublattices can be directly visualized in the high angle annular dark field (HAADF) mode, while oxygen sublattices can be observed using annular bright‐field (ABF) mode. Therefore, BO6 octahedral tilt angle and full 3D rotation pattern can be determined using the combination of STEM images taken in HAADF and ABF modes.6 tilting was observed. The plane averaged and projected tilt angle determined from the ABF STEM image is close to 0°  have been quenched. Combined with octahedral rotations observed in the 40 uc film, it can be envisaged that, with decreasing the film thickness, the MnO6 octahedral rotations are significantly modified by the interfacial octahedral proximity effect.For STEM images of ABO° Figure e. In ordFigureTC of around 291 K  detection on the 5 uc and 40 uc films, as shown in Figure µB/atom) and element‐specific magnetic moment, thus excluding the contribution from magnetic impurities.L2,3 edges measured at 2 K under 4 T are shown in Figure 3 and La0.7Sr0.3MnO3.3 with a thickness of 9 uc and hole‐doped La0.7Sr0.3MnO3 with a thickness of 40 nm.TC1 ≈ 100 K; Figure TC2 > 400 K; Figure L2,3 edges under 0 T were carried out and XMCD spectra are displayed in Figure TC , suggesting the La0.9Ba0.1O layer donates electrons to the interfacial Ti. Furthermore, by applying sum rules for Ti and Mn L2,3 edges,TC1 ≈ 100 K; Figure S6, Supporting Information).We found that the Ti close to the LBMO/STO interface in the 5 uc film exhibits a mixed 3+/4+ oxidation state, indicating the charge transfer also occur at the interface similar to the LaAlOL2,3 edges were performed at 300 K without a magnetic field in grazing incidence geometry. The XLD spectra provide information on the empty Mn 3d orbital states. Specifically, the smaller (larger) absorption for in‐plane polarization suggests more out‐of‐plane (in‐plane) empty states in the eg band and thus a higher occupancy in the in‐plane (out‐of‐plane) orbitals.L2 absorption peak display a positive sign, implying a preferential occupation of the out‐of‐plane d3z2−r2 orbitals.3z2−r2 occupancy in eg orbital shown in the 40 uc film under tensile strain can be explained by a vacuum‐interface induced extra contribution of the orbital configuration, favoring occupancy of the out‐of‐plane d3z2−r2 states. This has been demonstrated in La0.67Sr0.33MnO3 films under tensile strain.L2 absorption peak (648–660 eV) in the XLD spectra was calculated by integration, as presented in Figure 3z2−r2 orbital occupancy increases with a decrease of the film thickness, in agreement with an enhanced occupation of d3z2−r2 orbital in the La0.67Sr0.33MnO3 films caused by surface effect.TC2) observed in the 5 uc film.To determine the orbital occupancy, XLD measurements at the Mn Figure0.7Sr0.3MnO3 films,2g (≈2.4 eV) and eg (≈1.05 eV) states at the top of VB. The VB maximum (VBM) for the 5 uc film is located at 0.42 eV higher binding energy than that for the 40 uc film. Also, no visible intensity at the Fermi level (Ef) is observed for the 5 uc film . These results further confirm the films change from a metallic state to an insulating state with decreasing film thickness, consistent with the transport measurement.To further elucidate the origin of room‐temperature ferromagnetic insulating state in the 5 uc films, we performed detailed investigations on the electronic structures. K‐edge XAS spectra are shown in Figure K‐edge XAS probes the transition from O 1s to the unoccupied states derived mainly from the hybridization of O 2p states with Mn 3d states, La 5d states, and Mn 4sp states. The spectrum can be qualitatively related to the unoccupied density of states above Ef.1−xSrxMnO3,g band with the occupancy of the out‐of‐plane d3z2−r2 states, which can be assigned to be the low‐energy eg↑orbital. Consequently, in the O K‐edge XAS spectra, the first peak feature just below 530 eV is attributed to an accidental superposition of majority eg↑ with dx2−y2 symmetry and minority t2g↓bands. The first peak feature shares the same peak position in the 5 uc and 40 uc films, suggesting no change of the oxidation state of the Mn atoms. However, comparing with the 40 uc film, the pre‐edge position of the first peak in the 5 uc film moves toward higher photon energies . Combined with the VBM shift toward higher binding energies in the 5 uc film, it can be concluded that the eg bandwidth is reduced and a bandgap emerges with decreasing film thickness.g bandwidth increases as the Mn—O—Mn bond angles increase and the Mn—O bond lengths decrease. The reduction of the eg bandwidth in the 5 uc film can be explained by the increase of both of the Mn—O—Mn bond angles and Mn—O bond lengths caused by the tensile strain and the suppressed MnO6 octahedral rotations.O 6 octahedral rotations are suppressed at the interface , consistent with suppression of interfacial MnO6 octahedral rotations observed in the 5 uc and 40 uc films  emerges in the 5 uc film  were aligned via autocorrelation and displayed as sum image, which were wiener filtered to reduce noise. Image analysis was done via Digital Micrograph and ImageJ scripts.Magnetic and Transport Characterization: The magnetization measurements for LBMO films were performed with a Quantum Design MPMS3 SQUID‐VSM magnetometer. Resistance was measured by a Quantum Design Physical Property Measurement System (PPMS).X‐Ray Photoemission Spectroscopy: In situ X‐ray photoemission spectroscopy was undertaken by a monochromatic Al Kα X‐ray source (hν = 1486.6 eV) using a SPECS PHOIBOS 150 electron energy analyzer with a total energy resolution of 0.5 eV. The measurements were performed at 300 K. The Fermi level of the films was calibrated using a polycrystalline Au foil.X‐Ray Absorption Spectroscopy: X‐ray absorption spectroscopy measurements with polarization dependence were performed at the I06 beamline of the Diamond Light Source. The XMCD measurements were carried out using a TEY detection in grazing incidence geometry. XMCD measurements were performed between 2 and 300 K in a 4 T magnetic field applied in the a–b plane of the films, parallel to the beam propagation direction. To ensure that the XMCD signal is of magnetic origin, the magnetic field was applied in the opposite direction to verify the sign of the XMCD reversed. XLD measurements were carried out at 300 K without a magnetic field and a TEY detection was used in grazing incidence geometry. The XLD spectra were obtained by the intensity difference (Iv − Ih) between the spectra measured with horizontal (Eh) and vertical (Ev) linear polarizations. Due to the second transition temperature observed in the 5 uc films above 300 K, the XLD spectra might be slightly affected by the magnetic contribution.Density Functional Theory Calculations: DFT calculations were carried out using the plane wave pseudopotential code CASTEP.U correction was used with U = 3 eV for the Mn 3d electrons.The authors declare no conflict of interest.Supporting InformationClick here for additional data file.