This kind of NiCo2O4 aerogel structure shows large specific surface area (167.8 m2 g−1), high specific capacitance (903.2 F g−1 at a current density of 1 A g−1), outstanding rate performance (96.2% capacity retention from 1 to 10 A g−1), and excellent cycling stability (nearly without capacitance loss after 3000 cycles at 10 A g−1). The unique structure of the 3D hollow aerogel synergistically contributes to the high performance. For instance, the 3D interconnected porous structure of the aerogel is beneficial for electrolyte ion diffusion and for shortening the electron transport pathways, and thus can improve the rate performance. The conductive carbon joint greatly enhances the specific capacity, and the hollow structure prohibits the volume changes during the charge–discharge process to significantly improve the cycling stability. This work represents a giant step toward the preparation of high‐performance commercial supercapacitors.A biomass‐templated pathway is developed for scalable synthesis of NiCo However, for the nature of a conductor, the conductivity of NiCo2O4 still needs to be further improved to achieve higher performance. Therefore, most of the reported NiCo2O4 electrodes for high‐performance SCs so far are based on various NiCo2O4 nanostructures grown on conductive substrates by an in situ growth method.As new and promising energy‐storage devices, supercapacitors (SCs) have attracted tremendous research attention in recent years because of their high power densities, short charging/discharging time, excellent cycling stability, and electrochemical reversibility.2O4 electrode also suffers from poor ion transport kinetics and therefore, would lead to bad rate performance.2O4 hollow nanoparticles (HNPs) have been developed as electrodes for SCs to enhance the performance since they can provide high specific surface area and short ion diffusion path. However, these nanostructures always result in increased number of grain boundaries, which could increase electrical resistance in a solid electrode. Also, the synthesis of NiCo2O4 HNPs was generally dependent on hard‐template method which is limited for practical applications due to its high cost, inconvenient, and multistep synthesis.2O4 HNPs and assemble them into a matrix with interconnected porosity and excellent electrical conductivity. Aerogels, a special class of porous solid materials with hierarchical 3D morphology and adjustable porosity, precisely meet the above requirements. It could be expected that a 3D NiCo2O4 aerogel, which is composed of HNPs of NiCo2O4 connected with conductive joint, can combine the merits of each assembling unit to synergistically enhance the intrinsic properties of each component such as electrical conductivity and ion transport kinetics.The NiCo2O4@carbon aerogels assembled by HNPs of NiCo2O4 and nanocarbon joint were prepared by using sustainable sodium alginate (SA) as template and carbon source. The method reported here involves the synthesis of NiCo‐alginate hydrogel and aerogel, and subsequent pyrolysis in N2 and oxidation in air. The special “egg‐box” structure formed in SA after ion‐exchange with Ni2+ and Co2+ cations is the key to guide the synthesis.2O4 HNPs within the nanoscale. This facile and cost‐effective method is more suitable for large scale production in comparison with the complicated hard‐template synthesis. The composite aerogels were designed and synthesized by integrating the 3D porous structure and functional components. When evaluated as electrode materials for SCs, these NiCo2O4@carbon aerogels displayed high capacitance (903.2 F g−1 at a current density of 1 A g−1), outstanding rate performance (96.2% capacity retention from 1 to 10 A g−1), and superior cycling stability (nearly without capacitance loss after 3000 cycles at 10 A g−1).In this work, NiCo2O4@carbon aerogels is illustrated in Figured‐mannuronate (M) and α‐l‐guluronate (G) monomers. The G blocks of SA are active to coordinate with divalent or trivalent cations to form “egg‐box” structure.2+ and Co2+ cations with Ni/Co molar ratio of 1:2 was dropped into an aqueous solution of SA to prepare NiCo‐alginate hydrogel  were observed from the patterns , i.e., the NPs of NiCo alloy were oxidized to NiCo2O4 NPs. According to the Raman spectra shown in Figure S3 in the Supporting Information, besides the vibration peaks assigned to F2g, Eg, F2g, and A1g, modes of NiCo2O4,−1 and G band at 1592.5 cm−1 were observed.−1 are assigned to the characteristic D (defects and disorder) and G (graphitic) bands of carbon materials, respectively.ID/IG) value demonstrates the extent of structural disorder of carbon materials and the ID/IG values of the NiCo2O4 NPs here are in the range of 0.85–0.92, indicating that the formed carbon material possessed high graphitic degree and excellent electroconductivity. Field emission scanning electron microscopy (FESEM) was conducted to determine the morphology and microstructure of the obtained NiCo2O4 aerogel. As shown in Figure−1), was composed of uniform NiCo2O4 NPs with a diameter of 30–40 nm. More importantly, the sample exhibited highly interconnected 3D porous structure. A cross‐section image of the Ni‐Co‐12 aerogel revealed that the 3D porous network was throughout the whole sample  which is due to the treatment of higher concentration of Ni2+ and Co2+ ions. With the increased concentration of Ni2+ and Co2+ used during ion‐exchange, the larger NiCo2O4 particles were generated during the air oxidation process owing to the Ni2+ and Co2+ not fully utilized by the SA chains. The above formation of larger NiCo2O4 particles led to lower Brunauer‐Emmett‐Teller (BET) specific surface area as determined by the adsorption–desorption of N2.Powder X‐ray diffraction (PXRD) was carried out to monitor the evolution process of the NiCo‐alginate aerogels after pyrolysis and oxidation treatment. Figure S1 (Supporting Information) shows the XRD pattern of NiCo‐alginate aerogels after the pyrolysis process. Three diffraction peaks at 44.35°, 51.69°, and 76.12° are observed from the XRD patterns, which are ascribed to Ni–Co alloy (JCPDS No. 97‐010‐8308). The intensity of the peaks increased with increasing Ni, Co content. After oxidation in air at 300 °C, the diffraction peaks of the cubic spinel NiCo2O4 NPs suggested their typical hollow nature  and (02‐2) planes of face centered cubic NiCo2O4, respectively. In addition, we can clearly see that the NiCo2O4 HNPs were connected with graphitic carbon joint with a lattice fringe space of 0.35 nm  spectrum of Ni‐Co‐12 aerogel  shows the signals of Ni, Co, and O elements and the molar ratio between Co and Ni is nearly equal to 2. The high‐angle annular dark‐field scanning transmission electron microscopy (HADDF‐STEM) image  and high resolution TEM (HRTEM). The contrast (dark/bright) between the boundary and center of the NiCo2 adsorption/desorption isotherms and the pore size distribution of the NiCo2O4 aerogels are shown in Figure S6 in the Supporting Information. The major N2 adsorption of all the samples occurs at a low relative pressure (P/P0 < 0.02) but rises slowly thereafter to an almost plateau and at higher relative pressures shows an obvious hysteresis loop, which is typical of type IV adsorption/desorption isotherms indicating the existence of mesoporosity. The specific surface area and pore volume of all the samples are presented in Table S1 in the Supporting Information. The specific surface areas of the Ni‐Co‐3, Ni‐Co‐6, Ni‐Co‐12, and Ni‐Co‐18 aerogels are determined to be 77.5, 121.4, 167.8, and 137.6 m2 g−1, respectively. The insets of Figure S6 (Supporting Information) show the pore size distribution curves of all the samples, which indicate mesopores in the 3–4 nm and 15–40 nm ranges. The pore size distribution of Ni‐Co‐12 , in particular, shows the narrowest pore size distribution of 3–20 nm and consequently the highest specific surface area. Therefore, the Ni‐Co‐12 aerogel is selected for further investigation. The porosity derived from the aerogel structure apparently can be used as good current collector in energy storage applications and hence they are expected to be useful for electrochemical applications.The NFigure2p emission spectrum  analysis Figure3. The 2O4 aerogels, the cyclic voltammetry (CV) and galvanostatic charge–discharge (GCD) measurements were performed as shown in Figure−1 in a potential range of 0.0–0.6 V (vs Hg/HgO) in 6.0 m KOH solution. A pair of well‐defined redox peaks are clearly observed from all the CV curves, which mainly originated from the reversible faradic redox reaction related to M—O/M—O—OH (M = Ni or Co).−1 scan rate  are also consistent with the specific surface areas of different aerogels . As a comparison, the SEM image and electrochemical performance of nonporous NiCo2O4 is shown in Figure S8 in the Supporting Information. It can be seen that nonporous NiCo2O4 has larger size than that of the NiCo2O4 aerogel. Its specific capacitance is ≈580.0 F g−1 at the current density of 1 A g−1, much lower than that of Ni‐Co‐12. Figure S9 (Supporting Information) shows a comparison of the specific capacitance of Ni‐Co‐12 aerogel with those of other materials previously reported.To study the electrochemical performance of all NiCo2O4 aerogel is shown in Figure 2O4 aerogels have excellent rate performance although the corresponding specific capacitance tends to decrease slightly with the increase of current density. Interestingly, the specific capacitances of the Ni‐Co‐12 were about 903.2, 898.7, 883.0, 874.5, 870.0, 869.0 F g−1 at the current densities of 1, 2, 4, 6, 8 and 10 A g−1, respectively and these results suggest that this NiCo2O4 aerogel led to an outstanding rate performance with only an ≈6.8% SC loss (from 903 to 869 F g−1) by increasing the current density from 1 to 10 A g−1. This outstanding rate performance can be attributed to the unique structure of HNP in the Ni‐Co‐12 aerogel, which provided higher specific surface area and thus, more contact interface between the electrolyte ions and the electrode materials. Such aerogel structure can also facilitate the electrolyte ion diffusion and shorten the electron transport pathways, which are beneficial to improve the rate performance of SCs. Figure 2O4 aerogels tested at the current density of 10 A g−1. Only a little capacity loss occurred with all the aerogels after 3000 cycles. When the cycle times were extended to 10 000 laps, the capacitance retention of Ni‐Co‐12 is also as high as 92.1% . Such high capacity retention can be attributed to the unique hollow nanostructures of the aerogels, which provide much space against the volume changes during the charge–discharge process.2O4 HNPs can retain their hollow structure, indicating super stability of the HNP electrode. Figure S12 (Supporting Information) gives the Nyquist plots of all the samples. Apparently, Ni‐Co‐12 shows the smaller equivalent series resistance (Rs) and charge‐transfer resistance (Rct) than other three samples, which also leads to better electrochemical performance.The rate performance of all NiCo2O4 aerogel, hybrid supercapacitor (HSC) was assembled. Figure2O4 HNPs and activated carbon were used as the positive and negative electrode, respectively. Figure 2+/Ni3+ and Co2+/Co3+. Figure −1 at the current density of 1, 2, 4, 6, 8 and 10 A g−1, respectively. The result is much better than the previous reports (≈74–95 F g−1).−1 can be achieved at a power density of 1852 W Kg−1. Furthermore, the soft package electrochemical device was prepared with NiCo2O4 aerogel as electrodes using 6.0 m KOH as electrolyte. As shown in Figure S14 in the Supporting Information, the electrochemical devices used Ni‐Co‐12 as electrode were connected in series which powered 19 LED lights (≈1.2 V) and could keep these LEDs bright enough continuously (see Movie S1 in the Supporting Information). Figure S15 (Supporting Information) shows the cycle performance of the above device. It can be seen that the capacitance retention is ≈85.1% after 4000 cycles, which is relative better than the previous report.2O4 aerogels showed superior supercapacitor behavior.To further investigate electrochemical performance of the NiCo2O4 aerogels by a simple metal ion‐exchange method from the seaweed biomass for use in high‐performance supercapacitors. Electrochemical investigations clearly demonstrated that the carbon conjoined hollow NiCo2O4 aerogels served as excellent electrode materials for electrochemical capacitors with high specific capacitance (903.2 F g−1 at 1 A g−1), improved rate capability, and excellent cyclic stability even at a high current density of 10 A g−1. This biomass‐template pathway is not only facile and cost‐effective for the controlled synthesis of the porous hollow and aerogel structure of NiCo2O4 but also more suitable for large scale production in comparison with the complicated hard‐template synthesis. So, this synthesis strategy suggests a new pathway to design electrode materials of different chemical compositions for high‐performance supercapacitor.In conclusion, we have successfully synthesized carbon supported hollow NiCoThe authors declare no conflict of interest.SupplementaryClick here for additional data file.SupplementaryClick here for additional data file.


Co3O4 activated the highly efficient oxidation of methane under mild electric energy with the help of carbonate as an oxidant, which is delivered by ZrO2. Based on the experimental results, acetaldehyde is the key intermediate product. Subsequent nucleophilic addition and free radical addition reactions accounted for the generation of 2‐propanol and 1‐propanol, respectively. Surprisingly, this work achieves a production efficiency of over 60% in the conversion of methane to produce these long‐term stable products. The as‐proposed regional electrochemical methane oxidation provides a new pathway for the synthesis of higher alcohols with high production efficiencies under ambient conditions.Due to the greenhouse effect, enormous efforts are done for carbon dioxide reduction. By contrast, more attention should be paid for the methane oxidation and conversion, which can help the effective utilization of methane without emission. However, methane conversion and utilization under ambient conditions remains a challenge. Here, this study designs a Co However, conventional CH4 conversion processes still suffer from various drawbacks, including the use of precious metal catalysts, high reaction temperatures that are far from ambient conditions and extremely low conversion efficiencies.4 to liquid alcohols, such as methanol, ethanol, and propanol, is much more economical and energy‐efficient. Among liquid alcohols, higher alcohols with high energy densities have wide applications in fabricating commodity chemicals.Methane CH is the p4 conversion, has been widely investigated. Currently, the syngas reaction is the main route for the industrial production of methanol, whereby syngas is produced through CH4 steam reforming. The two reactions for CH4 conversion to methanol are given in Equations Methanol, as a product of CH4 to methanol. Some scientists have employed a class of bacteria known as methanotrophs with methane mono‐oxygenases to convert CH4 into methanol under ambient conditions using oxygen.4 by oxygen gas (O2) is always accompanied by substantial overoxidation, which is both kinetically and thermodynamically favorable.4 to methanol with O2 is triggered by energy input and proceeds spontaneously with a substantial release of energy, which makes it exergonic for the further oxidation of methanol to formaldehyde, formic acid, carbon monoxide (CO), and CO2. Controlling CH4 oxidation to obtain methanol is extremely difficult. Thus, compared with the route of inhibiting CH4 overoxidation to obtain methanol, the control of further oxidation processes under mild conditions may enable the production of more useful higher alcohols, such as propanol. Electrochemical oxidation in aqueous electrolyte has been demonstrated to be a suitable method for the conversion of CH4 or other volatile organic compounds (VOCs) at low temperature using simple reaction instruments.From the equations, the initial reaction in syngas reforming requires more energy than that released from the second reaction, illustrating that additional energy input is necessary for the conversion of CH4 have a high dissociation energy of 104 kcal mol−1, which makes CH4 extremely inert. When O2 is used as an oxidizing agent for CH4 oxidation, the reaction of triplet O2 with singlet CH4 to form singlet methanol is a spin‐forbidden process.4 are not expected to be readily abstracted by O2 under mild conditions, such as low temperature. Therefore, other oxidizing agents are necessary to replace O2. In conventional alkaline electrochemical systems, which are always performed at room temperature, hydroxide (HO−) generally functions as the oxidizing agent. However, hydroxide has previously been demonstrated to have negligible activity for abstracting protons from CH4 under mild conditions.32−) oxidizes species by donating a charged oxygen atom accompanied by CO2 release, resulting in a large enthalpy of reaction and favorable oxidation kinetics.32− may be an attractive alternative to OH− for alkaline electrochemical reactions. In addition, to realize the donation of an oxidizing agent from CO32−, zirconia (ZrO2) should be employed to facilitate CO32− adsorption because of its surface Lewis acid sites and electron‐accepting capabilities.2 for the electrochemical oxidation of CH4 with CO32− as the oxidizing agent.4 oxidation, and the reaction mechanism was also unclear. Thus, more efficient catalysts with high activity and selectivity should be employed.The C—H bonds in CH3O4) has been demonstrated to be one of the most efficient catalysts for the oxidation of VOCs.3O4 has a strong surface adsorption capacity for formaldehyde.2 on transition metal surfaces,3O4, formaldehyde, converted from methanol, may be more active for further additional reactions to higher alcohols, illustrating the regional selectivity of Co3O4 for CH4 oxidation. Thus, we designed a Co3O4/ZrO2 composite for the electrochemical oxidation of CH4, using CO32− as the oxidizing agent source, that resulted in high selectivity for the production of 1‐propoanl and 2‐propanol with over 60% production efficiency. The reaction process is shown in SchemeAmong metal oxide catalysts, cobalt oxide . Scanning electron microscopy (SEM) was used to observe the morphology of different samples fabricated using different ZrO2/Co3O4 ratios, denoted 1–2 ZrO2/Co3O4, 1–4 ZrO2/Co3O4, and 1–6 ZrO2/Co3O4, as shown in Figure2 nanoparticles with uniform sizes were formed and were adsorbed on the surface of the Co3O4 plates. Upon increasing the amount of Co3O4, the particle size of Co3O4 gradually increased until it reached bulk  and showed large particles that were more than 10 µm in size. The size of the Co3O4 plates can be controlled through coprecipitation with ZrO2. The elemental ratios of Zr/Co for all samples were obtained through EDS (energy‐dispersive X‐ray spectroscopy) measurement . The Zr/Co ratio decreased upon increasing the amount of Co precursor. In the 1–6 ZrO2/Co3O4 sample, big Co3O4 plates were surrounded by small ZrO2 particles, which may have caused the sudden decrease in the Zr/Co ratio. The EDS measurement was focused on the surface of the samples, which would be affected much by the surface state of the materials. Thus, the ratios showed different value when compared to the stoichiometric ratio of elements calculated from the reactants described in the Experimental Section. In order to obtain the stoichiometric ratio of Zr/Co, ICP‐OES  measurement was conducted, shown in Table S1 (Supporting Information). The 1–4 ZrO2/Co3O4 sample exhibited relatively small particle sizes and a suitable amount of Co3O4 and showed the best electrochemical catalytic performance for CH4 oxidation, which will be discussed next.Coprecipitation and hydrothermal methods were employed to fabricate the Co2 and Co3O4, 1–4 ZrO2/Co3O4 was examined by transmission electron microscopy (TEM)  image  facet, and that of Co3O4 was 0.238 nm, corresponding to the (022) facet. In addition, elemental mapping was performed for Co, Zr, and O atoms  and X‐ray photoelectron spectroscopy (XPS) spectra , and the peaks of Co3O4 were indexed to the cubic structure (JCPDS No. 42‐1467). In the XRD patterns, the typical (001), (100), (011), (−111), and (022) planes of ZrO2 were observed at ≈17.5°, 24.2°, 24.6°, 28.3°, and 50.3°. All related peaks gradually decreased in intensity as the amount of Co3O4 increased. The typical (111), (311), and (440) planes of Co3O4 were observed at ≈19.0°, 36.9°, and 65.2°, which also showed the same change in peak intensity with an increase in ZrO2.The crystalline structures of the ZrO2/Co3O4 nanocomposite. Therefore, we next set out to determine the surface state, which is responsible for the unique regional selectivity of CH4 oxidation, by XPS, as shown in Figure 2/Co3O4 ratios. The binding energies of the Co and Zr signals did not greatly shift upon changing the component ratio. However, the intensity of the Co signals clearly changed as the amount of Co decreased. The Zr signals also changed, but not as much as the Co signals, which may be due to ZrO2 being on top of the Co3O4 surface. Additionally, the O 1s spectroscopic signals were affected. The O 1s peaks at ≈530 and 532 eV are ascribed to lattice oxygen and nonuniform surface sites, respectively.3O4 is known to be readily covered with a monolayer of negatively charged chemisorbed oxygen.3O4, shown in Figure 3O4. In addition, after coprecipitation with ZrO2, the O 1s signals of all ZrO2/Co3O4 composites showed enhanced peaks at ≈532 eV, which can be observed more clearly in the overlapping plots in Figure S4b (Supporting Information). This result shows the surface electron accepting capabilities of ZrO2, as reported previously,32− electrolyte during the electrochemical oxidation of methane.The abovementioned data explained the microscopic and crystalline structures of the ZrO2.24 oxidation, a glassy carbon disc was employed to load the catalyst, forming the working electrode. Details of the preparation and measurement process are provided in the Experimental Section. As mentioned above, NiO is not an efficient catalyst for CH4 oxidation. Thus, we prepared a ZrO2/NiO composite and compared its catalytic ability with that of ZrO2/Co3O4, as shown in Figure S5 (Supporting Information). In the measurement of the current–voltage (J–V) curves and Nyquist plots , argon (Ar) saturation was used to analyze the water oxidation and CH4 saturation was used to analyze the CH4 oxidation competing with water oxidation. Additional anodic activity under CH4 saturation indicated that both NiO and Co3O4 have CH4 oxidizing activity. However, CH4 oxidation performed better on the surface of Co3O4 than NiO. The same results were observed via electrochemical impedance spectroscopy (EIS) , which illustrated the superior activity of Co3O4 for CH4 electrochemical oxidation.To measure the electrochemical performance of CHFigure2/Co3O4 samples with different ratios in CH4‐saturated carbonate electrolyte. The 1–4 ZrO2/Co3O4 sample showed the highest electrochemical current density for CH4 oxidation, which agrees with the sample morphology results after adjusting the component ratio, as 1–4 ZrO2/Co3O4 had a relatively suitable size and amount of Co3O4. The current density of the 1–2 ZrO2/Co3O4 sample was relatively low even after increasing the potential to a high value, which illustrated that a low amount of Co3O4 provided poor catalytic activity in CH4 oxidation. By contrast, ZrO2 showed extremely weak oxidizing activity under both H2O and CH4 saturation due to its high bandgap of more than 5 eV.2/Co3O4 sample showed the same J–V curve trend as the 1–4 ZrO2/Co3O4 sample but a lower current density value, illustrating the weaker surface adsorption ability toward the carbonate electrolyte due to the lower amount of ZrO2. In addition, J–V curves of the optimized sample in Ar‐ and CH4‐saturated electrolyte were recorded to demonstrate its ability for CH4 oxidation, as shown in Figure S6a (Supporting Information). In addition, pure Co3O4 powder was also prepared to examine the electrochemical oxidation of CH4 . Unfortunately, the Co3O4 sample without ZrO2, which aids the adsorption of oxygen donors, showed no additional anodic activity in CH4‐saturated electrolyte and even worse performance, which confirmed that the outstanding electrochemical oxidation of CH4 resulted from the synergistic effects of ZrO2 and Co3O4, i.e., carbonate adsorption and CH4 oxidation, respectively.2.34 conversion, the catalyst was loaded on carbon paper to form the anode, carbonate was the electrolyte and platinum foil was the counter electrode. The preparation and measurement processes are shown in the schematic diagram in Figure S7 (Supporting Information), and further details are provided in the Experimental Section. CH4 conversion using the Co3O4 catalyst was performed to obtain more stable and useful higher alcohols. Therefore, a sealed reactor instrument is more suitable for the conversion and for tracking the CH4 oxidation progress in real time. To determine the optimal potential for long reaction times, the difference in the current density of the optimized sample in Ar‐ and CH4‐saturated electrolyte  was calculated, and the curves are provided in Figure S8 (Supporting Information). Based on the curves, 2.0 V was selected as the suitable potential for CH4 electrochemical oxidation to obtain a relatively high current density and low competition with water oxidation. The products were collected after reaction for 3, 6, and 12 h with vigorous stirring. I–t curves for 12 h measurement was shown in Figure S9 (Supporting Information). To identify the products, proton nuclear magnetic resonance (1H‐NMR) spectroscopy was performed. Figure1H‐NMR spectrum of the products obtained after 12 h of reaction. The main products were 1‐propanol and 2‐propanol.1H‐NMR peak of methanol is located at ≈3.3–3.5 ppm,1H‐NMR peaks of ethanol appear at the same positions as the 2‐propanol peaks.4. The 1H‐NMR result of the products from control experiment was shown in Figure S10a (Supporting Information). In addition, the pure carbonate electrolyte before reaction was also detected with 1H‐NMR spectrum for comparison . To confirm the amount of CH4 consumed and products generated, gas chromatography (GC) and GC‐mass spectrometry (MS) were performed. The amount of CH4 remaining in the reactor after 3, 6, and 12 h of reaction is shown in Figure 4, measured by GC, is shown in Figure S11 (Supporting Information). CH4 gas was mostly consumed, and the amount decreased gradually with the reaction time. After 12 h of reaction, almost 40% of the CH4 gas was converted. Meanwhile, the amount of various products measured by the GC‐MS system is shown in TableIn the measurement of the CH4 oxidization. As detailed in a previous study, formaldehyde should be the methanol oxidation product.4 and formaldehyde. Moreover, the amount of acetaldehyde decreased with the reaction time, which illustrates that acetaldehyde plays a pivotal role in the production of 1‐propanol and 2‐propanol, indicating that 1‐propanol and 2‐propanol were converted from acetaldehyde. Then, after 12 h of reaction, 1‐propanol, and 2‐propanol were the main stable products of CH4 oxidation, which agrees with the 1H‐NMR results. The conversion efficiencies for acetaldehyde, 1‐propanol and 2‐propanol were calculated and are shown in Figure As described in Table 2.4As observed in the results in Table 4 oxidation, CH4 was oxidized by the oxidant (carbonate in this work) to form CH3OH, which was subsequently oxidized to HCHO. The production mechanism has been investigated in detail, and there are several ways to achieve the reaction.4, methanol, and formaldehyde. Thus, the generation and accumulation of acetaldehyde is rapid and large, which is in accordance with the results in Table In the primary reaction involved in CH4 conversion. The mechanism involves a type of addition reaction, as shown in Figure4 acts as a nucleophilic reagent and attacks the carbonyl carbon in acetaldehyde. Then, a nucleophilic addition reaction occurs to form 2‐propanol as one of the main products in this CH4 conversion reaction. However, considering the addition reaction mechanism, the formation of 1‐propanol from acetaldehyde and CH4 is impossible. Therefore, we considered all the reaction conditions to determine a possible route for 1‐propanol production and found that the free radical addition reaction is suitable  is not the most stable state. However, when the methyl radical attacks carbon 2, the resulting 1‐propanol radical (free electron on carbon 1) is more stable than the 2‐propabol radical, illustrating 1‐propanol as the main product. 2‐Propanol can be directly converted from acetaldehyde and CH4 through a nucleophilic addition reaction, leading to more 2‐propanol being produced than 1‐propanol at short oxidation times. However, after long reaction times, the amount of 1‐propanol exceeds that of 2‐propanol, even though 2‐propanol is more thermodynamically stable, illustrating the unique regional selectivity of 1‐propanol production through radical addition with the participation of a Co3O4 catalyst and carbonate electrolyte. In summary, after comprehensive analysis with the above content, the complete reaction pathways for electrochemical oxidation of methane with ZrO2/Co3O4 anode and carbonate electrolyte were proposed, shown in Figure S12 (Supporting Information). The reaction process may help other researchers to have an overall understanding of the one carbon related reactions.The production of 1‐propanol and 2‐propanol from acetaldehyde is the most important reaction step in CHe Figure b1–b3. Fi32/Co3O4 nanocomposite that aids in the regional selective oxidation of CH4 to 1‐propanol and 2‐propanol via an electrochemical method. We expected stable products, such as higher alcohols, to be formed due to the strong surface adsorption ability of Co3O4 and the participation of carbonate, which has a soft oxidizing ability, delivered by ZrO2. After tracking the products for different reaction times, acetaldehyde was found to be the key intermediate. To convert acetaldehyde to 1‐propanol with the participation of CH4, a free radical addition reaction was conducted by an electrochemical reaction, with Co3O4 as the catalyst and carbonate as the electrolyte. Finally, as a result of competition between reactions, both 1‐propanol and 2‐propanol were the main products. This electrochemical partial oxidation of CH4 may aid in the synthesis of other oxygenates and long‐chain hydrocarbons.In summary, we designed a ZrO4Fabrication of Electrocatalyst Materials: All reagents were used as received without further treatment. The ZrO2/Co3O4 nanocomposite was synthesized using precipitation and a hydrothermal method. For the 1–2 ZrO2/Co3O4 sample, 0.1611 g ZrOCl2·8H2O , 0.291 g Co(NO3)2·6H2O , and 9.6 g NaOH  were dissolved in 40 mL deionized (DI) water with vigorous stirring for 30 min. Then, the solution was transferred to a 60 mL autoclave container and heated at 180 °C for 24 h. After that, the powder was collected by centrifuge and washed with DI water 3 times. Finally, the 1–2 ZrO2/Co3O4 sample was obtained after thermal annealing at 500 °C for 3 h. For the ZrO2/Co3O4 nanocomposites with different ratios, the amount of Co(NO3)2·6H2O was adjusted to 0.582 g for 1–4 ZrO2/Co3O4 and 0.873 g for 1–6 ZrO2/Co3O4, with no changes in the other conditions. For comparison, pure Co3O4 and ZrO2/NiO samples were prepared by the same method as 1–4 ZrO2/Co3O4, without the addition of ZrOCl2·8H2O and with the addition of 0.582 g Ni(NO3)2·6H2O  instead of Co(NO3)2·6H2O, respectively.Electrochemical Test: LSV and EIS tests for the comparison of ZrO2/Co3O4 and ZrO2/NiO samples were conducted in a three‐electrode system using a potentiostat  with a glassy carbon electrode as the working electrode, Ag/AgCl as the reference electrode, a Pt foil as the counter electrode and 0.5 m Na2CO3 solution as the electrolyte. Meanwhile, LSV tests for ZrO2/Co3O4 samples with different component ratio were conducted in a two‐electrode system without the utilization of Ag/AgCl reference electrode. All working electrodes were prepared by dispersing the powder samples in DI water in a concentration of 3 mg mL−1 with vigorous stirring for 30 min and then dropping 20 µL of the dispersed solution on the surface of the glassy carbon electrode (area = 0.07 cm2), followed by drying at room temperature. Next, 10 µL of a 5% Nafion 117 solution (Aldrich) was deposited on the surface of the glassy carbon electrode to cover the sample films, followed by drying at room temperature. Ultrahigh‐purity argon gas  and methane gas  were used. Before each electrochemical test, the electrolyte was bubbled with Ar or CH4 for 1 h to prepare the Ar‐ or CH4‐saturated electrolyte.4 Conversion TestCH: The long‐term electrochemical oxidation of CH4 was conducted in a two‐electrode system with a closed reaction instrument, employing carbon paper (Alfa) as the working electrode, Pt foil as the counter electrode and 30 mL 0.5 m Na2CO3 solution (pH around 12.0 before reaction and about 11.9 after 12 h reaction) as the electrolyte. The working electrode was prepared by dispersing the powder sample in DI water in a concentration of 3 mg mL−1 with vigorous stirring for 30 min and then dropping 5.7 mL of the dispersed solution on the surface of the carbon paper (area = 20 cm2), followed by drying at room temperature. Next, 3 mL of a 5% Nafion 117 solution was deposited to cover the sample film on the carbon paper, followed by drying at room temperature. Before the electrochemical reaction, the electrolyte was bubbled with CH4 for 1.5 h to remove the oxygen and fill the space in the reaction instrument. In this case, after the consumption of saturated CH4 in aqueous solution, the gas phase CH4 could dissolve in the electrolyte continually, guaranteeing the adequate reactant. Electrochemical oxidation was conducted at 2.0 V versus Pt for 3, 6, or 12 h.Characterization and Products Analysis: Morphology analyses of the samples were carried out using field‐emission scanning electron microscopy  and a JEOL JEM‐2100F (Japan) electron microscope. EDS spectra and ICP‐OES measurements were employed for the elements ratio detection. XRD measurements were conducted using a Siemens diffractometer D500/5000 in a Bragg–Brentano geometry. XPS data were obtained from a K‐alpha instrument . 1H‐NMR was conducted using an Avance III HD 400 FT‐NMR instrument (Bruker Biospin), where the sample was prepared by mixing 0.4 mL of the product solution with 0.2 mL D2O. The amount of methane was determined using a 7890B GC instrument (Agilent Technologies). The product was examined using a 7890B‐5977A GC‐MS instrument (Agilent Technologies).The authors declare no conflict of interest.SupplementaryClick here for additional data file.


Mycobacterium tuberculosis (Mtb) DprE1, an essential isomerase for the biosynthesis of the mycobacterial cell wall, is a validated target for tuberculosis (TB) drug development. Here we report the X‐ray crystal structures of DprE1 and the DprE1 resistant mutant (Y314C) in complexes with TCA1 derivatives to elucidate the molecular basis of their inhibitory activities and an unconventional resistance mechanism, which enabled us to optimize the potency of the analogs. The selected lead compound showed excellent in vitro and in vivo activities, and low risk of toxicity profile except for the inhibition of CYP2C9. A crystal structure of CYP2C9 in complex with a TCA1 analog revealed the similar interaction patterns to the DprE1–TCA1 complex. Guided by the structures, an optimized molecule was generated with differential inhibitory activities against DprE1 and CYP2C9, which provides insights for development of a clinical candidate to treat TB. Mtb Decaprenylphosphoryl‐β‐d‐ribofuranose 2‐oxidase (DprE1) is a flavin adenine dinucleotide (FAD)‐dependent enzyme, which together with decaprenylphosphoryl‐d‐2‐ketoerythropentose reductase (DprE2), converts decaprenylphosphoryl‐beta‐d‐ribose (DPR) to decaprenylphosphoryl‐beta‐d‐arabinofuranose (DPA), an essential cell wall component.dprE1 results in a strong bactericidal effect in vitro.Mtb, further validating DprE1 as an attractive anti‐tuberculosis (TB) drug target.dase DprE is a flaWe previously reported the identification of TCA1 from a cell‐based phenotypic screen, and demonstrated it is a DprE1 inhibitor.Mtb, and selectively targeting DprE1 over CYP2C9.Here we report the X‐ray crystal structures of WT DprE1, a TCA1‐resistant DprE1 mutant, and CYP2C9 in complex with TCA1 analogs. Along with structure activity relationship studies, the mechanism of resistance by the mutant DprE1 enzyme and the tight correlation between DprE1 and CYP2C9 inhibition was elucidated. We further generated analogs with excellent in vitro and in vivo activities against In the DprE1–TCA1 structure, TCA1 showed a planar conformation and tight fit in the DprE1 active site. The benzothiazole core is oriented parallel to the FAD isoalloxazine ring and the thiophene carboxamide resides in small hydrophobic pocket. Besides substantial hydrophobic interactions, multiple hydrogen bonding interactions are formed between TCA1 and residues Lys418, His132, and Ser228 . Guided by the structure, our medicinal chemistry effort focused on the benzothiazole core (red), thiophene carboxamide (green) and the acylcarbamate (blue) of TCA1  compared to TCA1 (48±14 nm) . Only a fluoro substituent on the thiophene afforded an analog (TCA481) that is 5‐fold more potent than TCA1 based on MIC . Replacement of thiophene with other heterocycles or adding bulky substitutions resulted in reduced activity against m, IC50=370±133 nm)   and the mutant strain . To elucidate the molecular basis of the resistance mechanism, the structure of DprE1 Y314C bound with TCA1 was determined to 2.2 Å , suggesting a low risk of cardiotoxicity. In addition, we observed no genetic toxicity in a mini‐Ames mutagenicity test. TCA007 showed no inhibition of four major CYPs, 1A2, 3A4, 2D6, 2C19, but a strong inhibitory effect for CYP2C9 (SI methods).A typical treatment of TB lasts for more than six months, and therefore, the safety profile of anti‐TB drugs is of great importance. TCA007 is not cytotoxic against Vero and HepG2 cell lines at up to a concentration of 100 μThe potent inhibition of P450 2C9 could reflect oxygenation of thiophene moiety of TCA analogs to an electrophilic S‐oxide or epoxidePhe100, Leu102, and Phe114 form a hydrophobic pocket together with Phe476 where the carbamate group of TCA007 is bound  for TCA582 and TCA711 were 33 and 1515, respectively, which, compared with TCA007 (ratio=1.8), represents a significant improvement in selectivity   Figure  D. The iIn summary, we demonstrate the mode of inhibition of DprE1 by the thiophene scaffold via enzymatic and structural studies. The SAR studies not only helped us to optimize the potency of the inhibitors but revealed an unconventional resistant mechanism of DprE1. The optimized lead compound TCA007 showed excellent in vitro and in vivo activities, PK properties, and low risk of toxicity profile except for the strong inhibition of CYP2C9. Guided by the structures of CYP2C9 in complex with TCA1 analogs, an optimized molecule, TCA711, was generated with differential inhibitory activities against DprE1 and CYP2C9, which paved the road for further development of a clinical candidate to treat TB.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.


The development of selective electrocatalysts for CO2 reduction in water offers a sustainable route to carbon based fuels and feedstocks. 2 reduction in water offers a sustainable route to carbon based fuels and feedstocks. However, molecular catalysts are typically studied in non-aqueous solvents, in part to avoid competitive H2 evolution. [Ni(cyclam)]2+ (1) is one of the few known electrocatalysts that operate in water and 30 years after its report its activity remains a rarely surpassed benchmark. Here we report that [Ni(cyclam-CO2H)]2+ ) shows greatly enhanced activity versus1 for CO production. At pHs < pKa of the pendant carboxylic acid a large increase in catalytic activity occurs. Remarkably, despite the high proton concentration (pH 2), 2 maintains selectivity for CO2 reduction and is believed to be unique in operating selectively in such acidic aqueous solutions.The development of selective electrocatalysts for CO To enable practical utilisation, CO2 reduction electrocatalysts will need to be used in tandem with a sustainable oxidation reaction, such as water splitting  = 1.23–0.059 pH) making the development of low cost, selective CO2 reduction catalysts that operate in water at a range of pHs an imperative goal. However the majority of studies to date using molecular catalysts have been carried out in aprotic solvents such as dimethylformamide (DMF) and acetonitrile (CH3CN) with Brønsted acids added. Careful control of the acid concentration, coupled to the relatively high solubility of CO2 in these solvents minimises competitive H2 production  = 0–0.059 pH). A further complication is that any CO2 electrolyser will require the cathode and anode to be separated by a membrane. To date the most effective membranes are proton exchange materials3 making the study of CO2 reduction in acidic conditions of particular interest. [Ni(cyclam)]2+ (1) is a low cost, highly selective CO2 reduction catalyst producing solely CO in water at pHs 7–4. Since the initial reports over 30 years ago,6 numerous attempts have been made to develop nickel cyclam catalysts with improved rate constants and onset potentials.7 However to the best of our knowledge only two reports observed an increase in the catalyst performance,9 with functionalisation of both the amines and carbon backbone typically causing losses in selectivity and excessive hydrogen production.The discovery of catalysts for the conversion of carbon dioxide ]+ adsorbs on to some metal electrodes including Sn, Pb,14 and Hg,5 and adsorption onto the electrode is key for efficient CO2 reduction.111 has also been shown to act as a homogeneous CO2 reduction catalyst when used with a glassy carbon electrode (GCE)15 however the level of activity was significantly lower than can be achieved on Hg, which in part may be due to suppression of catalyst degradation pathways on Hg.16 Indeed Hg remains a common electrode for fundamental studies such as that presented here. At pH 5 adsorption initiates at potentials positive of the formal NiII/I couple in solution (–1.3 VNHE) and a monolayer is formed at ca. –1 VNHE. The adsorbed NiI complex is predicted to bind in an η1-C mode to CO2, prior to the transfer of a second electron to the catalyst centre.13 Computational studies18 indicate that the structure of the adsorbed complex is a trans-I conformation,19 with the axial amine hydrogens aiding CO2 binding. In solution 1 can adopt five different isomeric forms11 and for clarity structures of 1 and 2 are drawn in a planar geometry , the turnover frequency per adsorbed catalyst, for 2 compared to 1, at –0.99 VNHE, near the foot of the catalytic wave. We also note an extremely high kobs = 3.4 (±1.0) × 103 s–1 at –1.25 VNHE. Perhaps most remarkable is that catalyst 2 operates in acidic conditions whilst maintaining selectivity towards CO2.Only a limited number of other classes of molecular CO2, a derivative of 1 with a carboxylate group on the carbon backbone  of 1 and 2 on a Hg/Au electrode at pH 5, higher than the pKa of the carboxylic acid group of 2 are in line with past reports, II/I couple is present under argon at –1.30 VNHE (1) and –1.33 VNHE (2),  of 2 at –0.99 VNHE with a very good selectivity towards CO production >4 : 1 CO : H2 (faradaic efficiency (FE), total = 81%, H2 = 15 ± 5%, CO = 66 ± 9%, with errors being the result of 3 experiments). This corresponds to an average bulk turnover number of 591 for CO in 1 hour. In contrast 1 passes only 2.0 ± 0.2C in 1 hour with a lower selectivity 0.2 : 1 CO : H2 , and a bulk turnover number of ca. 45 for CO production. No liquid phase products were detected by NMR. [Ni(cyclam)]2+ and its derivatives are known to form inactive species in the presence of CO,34 however activity can be maintained through constant CO2 purging and experiments with 2 over a 7.5 hour period show activity being maintained, Fig. S8.Bulk electrolysis experiments Fig. S7 confirm 2 at low pH we have examined the electrochemical response of 1 and 2 over a wide pH (6–2) range using rotating disc electrode (RDE) voltammetry,  growing in under CO2 as the pH decreases, which is shown above to be due to catalytic CO2 reduction. This leads to a large decrease in the potential necessary for catalysis between pH 5 and 2 of ca. 240 mV versus the normal hydrogen electrode. In the RDE measurements we define the potential necessary for catalysis as being when the current density exceeds 2 mA cm–2.35 In contrast with 1 we only measure a very small shift (ca. 50 mV) in the potential necessary for catalysis between pH 6 and 2, which will be at least in part due to the increased level of H2 production at low pHs. By pH 2 there is minimal separation of the RDE curves of 1 in the presence and absence of CO2, 2 but not 1 is indicative of a change in catalytic mechanism for 2 between pH 3 and 2. Furthermore whilst the variation in overpotential for CO2 reduction brought about by a change in pH (0.18 V) is equivalent for both catalysts only 2 shows a significant change in potential necessary for catalysis. The lack of a pH dependence for 1 is further explored in the ESI  spectroscopy in solution  at 1575 cm–1 and νs(CO2–) modes at 1375 cm–1 in line with literature reports for similar complexes.36 Titration of a 0.1 M solution of 2 in D2O  with DCl showed the clear emergence of the carboxylic acid form of 2 with νas(CO) at 1706 cm–1 in D2O, with pKa ∼ 2.6. Deuterated solvents are required to avoid the δ(HOH) mode of H2O masking the spectral window of interest. There is an excellent correlation between the relative concentration of the protonated carboxylic acid in solution and the current density for 2 measured under CO2 using RDE at –0.99 VNHE  × 10–10 mol cm–2.34 At pH 2 the surface concentration of both 1 and 2 are changed by a minimal amount (2.2 (±0.2) × 10–10 mol cm–2 and 1.6 (±0.2) × 10–10 mol cm–2 respectively at –0.99 VNHE), indicating that the large increase in activity of 2 cannot be attributed to a change in the surface coverage of the catalyst with pH.It has been shown that for 1 and 2 at pH 2 can be obtained from the RDE measurements carried out at different rotation rates , i.e. the turnover frequency (TOF) per adsorbed catalyst using eqn (1).38n is the number of electrons transferred (2), Γ the surface coverage (mol cm–2) and A the electrode area (cm2). We obtain kobs values of 3.5 (±1.0) × 101 s–1 and 1.9 (±0.2) × 102 s–1 for 1 and 2 respectively at pH 2, –0.99 VNHE. It is apparent that at pH 2 complex 2 turns over approximately five times faster than 1 at –0.99 VNHE and the activity of 2 exceeds 1 at all potentials examined, Table S4.1 primarily produces H2 in bulk electrolysis experiments therefore the measured kobs for 1 at pH 2 under CO2 is expected to have a significant contribution from proton reduction. In contrast 2 is shown to be selective towards CO2 and at potentials positive of –1.3 VNHE there is a large difference in kobs obtained in the presence and absence of the substrate (CO2), NHE, we calculate a very large rate constant under CO2, kobs = 3.4 (±1.0) × 103 s–1 compared to only kobs = 3.3 (±0.4) × 102 s–1 under N2. This kinetic control between proton and CO2 reduction offers a rationalisation of the very high selectivity of 2 even in the presence of a high proton concentration.Using the limiting current obtained from the intercepts of Koutecký–Levich plots  × 103 s–1 for 2 at pH 2 under CO2, measured at a single potential (–1.25 VNHE) exceeds the reported TOF of the majority of known water soluble CO2 reduction catalysts,34 including 1 (6.3 × 101 s–1).34 To the best of our knowledge there has only been one reported water soluble catalyst that operates at a greater rate, the recently reported iron porphyrin catalyst WSCAT.27 We also note that the measured rate constant for 2 under CO2 also exceeds that of many of the most commonly studied CO2 reduction catalysts operating in aprotic solvents,1 which is perhaps surprising given the significantly lower dissolved CO2 concentration in water (0.28 M (CH3CN), 34 mM (H2O)).Comparison of the catalyst performance with existing benchmarks is ideally carried out by comparison of the overpotential dependence of the catalytic rate constant.2 is its selectivity towards CO2 even under acidic conditions. All previously reported derivatives of 1 have shown predominantly hydrogen production outside of a small pH window11 and we note that the majority of CO2 reduction catalysts are reported at pHs close to neutral (5–7),27 making the ability of 2 to operate at pHs as low as 2 unusual. The correlation between the current density under CO2 and the protonation state of the carboxylic acid group of 2 suggests that the protonation state of the catalyst is an important factor in the enhanced TOF, and hence selectivity towards CO2 of 2 in acidic solutions  under identical conditions, something that has been rarely achieved in over 30 years of research. 2 is also found to have be amongst the most active aqueous CO2 reduction catalysts and we believe that these characteristics make it of great significance to the field of electrocatalytic CO2 reduction.The development of selective and efficient molecular catalysts for electrocatalytic CO


Neutral and negatively charged substitutional Cr impurities CrTi0 and CrTi1− as well as Cr-oxygen vacancy complex 2CrTi + VO are studied by the density functional theory (DFT) within the generalized gradient approximation (GGA) of Perdew-Burke-Ernzerhof (PBE) functional. Experimental results based on X-Ray absorption spectroscopy (XAS) and X-Ray photoelectron spectroscopy (XPS) performed on Cr doped TiO2 at the Synchrotron facility were compared to the theoretical results. It is shown that the electrons of the oxygen vacancy tend to be localized at the tg2 states of the Cr ions in order to reach the stable oxidation state of Cr3+. Effects of Cr on crystal field (CF) and structural distortions in the rutile TiO2 cell were analyzed by the DFT calculations and XAS spectra revealing that the CF and tetragonal distortions in TiO2 are very sensitive to the concentration of Cr.The effects of Cr on local environment and electronic structure of rutile TiO Study of cationic dopants in TiO2 is of great interest because it allows to modify and/or enhance physical and chemical properties of TiO2, thus advance the respective device performance and extend its application range10. Transition metal impurities such as Cr, V, Mn or Fe, can be easily incorporated into TiO2 lattice, which aroused interest since these dopants commonly improve physicochemical reactions at the surface that is important for degradation of pollutants in waste water treatment or in energy related applications where the oxidation state of the dopant plays a key role11. Nevertheless, achieving a controlled oxidation state of the transition metals in TiO2 is not a simple task, due to their characteristic multivalence. Many factors should be accounted for in the analysis of doping process, such as the temperature, atmosphere and precursors used during the synthesis, the dimensions and morphology of the material, and the final concentration of dopants, among others, which usually affect the position of the dopant in the lattice and its ground state nearby other intrinsic defects13. Among the other transition metals, Cr has frequently been the matter of study because of the controversy related to its influence on the physical and chemical properties of TiO2. As an example, Fan et al.3 observed that mesoporous Cr-doped TiO2 presents lower photocatalytic activity for acetaldehyde photodecomposition under the UV light radiation and its performance varies as a function of the Cr concentration and the irradiation wavelength, generally increasing below the critical Cr doping level. Similar observations were reported by Wei et al.14, Tian et al.4 On the contrary, Wilke et al.6 and recently Mittal et al.15 did not observe any photocatalytic improvement by Cr doping despite higher visible light absorption was reported. The magnetic properties of Cr doped TiO2 has been also investigated due to potential applications for spintronics and novel information processing16. In contrast to non-magnetic TiO2 observed by Matsumoto et al.13, ferromagnetic behavior at room temperature has been observed by other authors, however the variables involved in the magnetic properties are still unclear and remains under debate17. Therefore, a deeply understanding of electronic structure of doped systems like TiO2 is still needed, and the parameters such as the dopant oxidation state and the site where the dopant is located in the host lattice, as well as the influence of the dopant on the oxygen related defects should be further investigated in order to understand the dopant-induced physical and chemical effects18.Due to many technologically important optical, electronic and mechanical properties transition metal oxides have attracted increased research attention during the last years. Their optical, electronic and mechanical properties are appropriate for a wide range of applications. Titania  based calculations were carried out to investigate the effect of Cr on electronic density of states (DOS) of ideal and oxygen-deficient crystalline TiO2. The electronic structure has been experimentally studied by resonant and non-resonant X-ray Photoelectron Spectroscopy (XPS) and X-ray Absorption Spectroscopy (XAS) performed at a Synchrotron beamline that have been compared to the above theoretical results. The morphological and crystalline characteristics of the present microstructures, as well as the doping homogeneity achieved, make them an useful reference material for theoretical studies of Cr-doped TiO2. Moreover, this study could serve as a reference model for other metallic oxides that also crystallize in the rutile-like structure such as VO2, CrO2, MnO2, SnO2 or RuO2.The present study focuses on modifications of electronic structure of rutile TiO2 microtubes analyzed in this work have been synthesized by a vapor-solid method at temperatures of 1300 °C during 15 h, using 5 cat.% and 10 cat.% Cr-doped nanoparticles as precursor, as reported in a previous work19. These microtubes exhibit high crystallinity and chromium homogeneity, up to 3–4 cat.% Cr. A combined Raman spectroscopy in a confocal microscope and electron back scattered diffraction (EBSD) in a SEM19 confirm the growth direction and the lateral planes forming the tubes, which correspond to the [001] direction and the {110} family planes, respectively. These techniques also revealed the high crystalline quality of the microtubes, which lateral faces can be considered as single-crystals. In addition, undoped rutile TiO2 and Cr doped samples up to 5 cat.% Cr have been analyzed for comparison and commercial α-Cr2O3 polycrystalline powder , were employed as reference samples. Table The Cr-doped rutile TiO20 has been employed for calibration of the XPS spectra. Chromium quantification was performed by energy dispersive spectroscopy (EDS) with a Bruker AXS 4010 detector mounted in a Leica 440 SEM.XAS experiments in total electron yield (TEY) mode, as well as resonant and non-resonant XPS measurements were carried out at the BACH beamline at the Elettra Synchrotron light source facility  using a monochromatic photon energy in the range from 400 to 600 eV, with energy resolution of 40–200 meV. The spectra were acquired in normal incidence geometry. C(1 s) peak at 284.6 eV from adventitious carbonab initio simulation package (VASP)23 together with the potential projector augmented-wave (PAW) method26. The core and valence electronic states are expanded with a plane-wave basis set, with an optimal energy cutoff of 500 eV. Standard PAW-PBE27 pseudopotentials were employed to describe the Ti(3s3p3d4s), Cr(4s3p3d) and O(2s2p) valence states, allowing spin polarization during all calculations.All the calculations were performed using the density functional theory (DFT) implemented in the Vienna Ti0, CrTi1−) and also for the combined neutral defect 2CrTi0 + VO0. A 4 × 4 × 6 mesh for the k-point sampling centered at the Gamma point was found optimal for the calculations. Additionally, calculations using hybrid functional with a 25% portion of Hartree-Fock exchange using the standard HSE0628 screening parameter of 0.2 Å−1 were performed to analyze the 2CrTi0 + VO0 case using a 2 × 2 × 2 supercell and a 2 × 2 × 2 mesh for the k-point sampling. The structural optimization for point defects in the supercell were performed starting from a pre-converged unit cell and allowing ionic relaxation with energy convergence of 10−7 eV per atom and minimizing the forces on all atoms less than 10−2 eV·Å−1. The calculation of the effective charges were performed using the Henkelman’s grid-based algorithm for the Bader electron decomposition method29. The effective charge is defined as QX = ZX-qBader,X  similar to that calculated in Vasquez et al.30, where ZX is the number of valence electrons and qBader,X is the calculated Bader charge for the corresponding X atom. Graphical illustrations were drawn using the software VESTA31.The defect calculations were performed using a 2 × 2 × 4 (96 atoms) supercell for the neutral and negative charged Cr defects  and two apical (D0) Ti-O bond lengths. The basal bonds form an angle of 98.8° whereas the apical bonds form a right angle with respect to the basal plane32. Therefore, each octahedron is slightly tetragonal and trigonal distorted. In a previous work30 we have compared electronic properties of bulk rutile as calculated by PBE and hybrid functional (HSE06), as well as the effect of VO2+ on the TiO2 lattice and conduction band, so the complete list of rutile TiO2 parameters, as well as our experimental values and data from other authors, are summarized in the Table The rutile structure is formed by an infinite chain of TiOg), which is underestimated by the PBE method. Janotti et al.33 reported that this effect in rutile TiO2 can be attributed to a reduced self-interaction for the O(2p)-derived states at the upper part of the valence band. Bader charge analysis calculated by HSE06 indicates a slightly higher ionic character in the Ti-O bond as compared with the PBE calculations30. Taking into account that, apart from Eg underestimation by the PAW-PBE funtional, comparable values are obtained by both computational methods, in this work the PBE method has been employed in most of the analyzed simulations, as it allows to spread the type of defects to be simulated while keeping calculation time lower than HSE06.Our results are in good agreement with those commonly reported by other authors. Small differences, below 1%, are found between the results obtained from the calculations obtained within PBE and HSE. Large differences are found in the calculated band gap  and basal (d) Cr–O and Ti–O bond lengths in CrO6 and TiO6 octahedra are indicated for both CrTi1− and CrTi0 defects in rutile TiO2.After the ionic relaxation we have observed that the symmetry of the CrOD-d)/D percentages and the calculated Bader effective charge values are also included in Table 6 octahedron, the tetragonal distortion is significantly reduced, where the relative difference between calculated apical (D) and basal (d) Cr–O bond length is less than 0.55% in comparison to the value obtained for Ti–O bonds of 2.4%. This observation indicates that the Cr ion symmetry in rutile TiO2 is nearly octahedral (Oh) rather than tetragonal distorted octahedral symmetry as the TiO6 octahedron in perfect rutile.In addition, the corresponding  in Fig. Ti1−, CrTi0) influences the Cr-O bonding, as expected. The Cr d-orbital is occupied by two electrons in the case of a Cr4+ ion (CrTi0 defect), and three in the case of a Cr3+ ion (CrTi1−). This implies that the additional electron in the CrTi1− defect belongs to the Cr valence electrons, expanding slightly the octahedron volume by increasing the Cr-O bond length , corresponding to both types of Cr defects. In Fig. Ti1- defect. However, the apical bonds are different depending on the Cr oxidation state. For the neutral CrTi0 defect (Cr4+) the ELF value is higher in the apical Cr–O bonds, contrary to CrTi1− (Cr3+) where the electron density is still minimum as can be observed in Figs Ti1− and CrTi0 defects are shown in the Figs Ti1− defect is distributed forming a cube-like volume, whereas the electron localization around the CrTi0 defect shows a more complex distribution with localized electrons in the regions around the apical bonds. The relatively small difference in the effective charges between the CrTi1− and CrTi0 defects  crystal field. In the case of the CrTi0, the higher electron localization along the apical bonds indicates that the electrons can partially occupy the In both cases the electron density around the basal Cr-O bonds Figs  is practCrTi1− Cr+ where tp), Ti(3d) and Cr(3d), for the CrTi1− and CrTi0 defects in rutile TiO2 are shown in Fig. s+ and s- indicate the spin-up and spin-down components of the DOS, respectively.In order to complete the DFT study, the density of electronic states (DOS) has also been analyzed. The total DOS and partial DOS (pDOS), including O and Cr(3d) orbitals for CrO6 octahedron compared with a single TiO6 octahedron (dashed line) for both CrTi1− and CrTi0. In the case of CrTi1− [top of the Fig. 6 octahedron are similar to the calculated contributions for the TiO6 octahedron. Within the band gap region, the O(2p) orbitals are slightly hybridized with Cr(3d) orbitals (tg2 related states) in the CrTi1− defect. On the contrary, for the CrTi0 defect [down in Fig. p) and Cr(3d) orbitals show stronger hybridization states at about 5.2 eV below the EF partially overlapped with the O(2p)-σ bonds of TiO234. In the conduction band (CB) the Cr contributions are partially mixed with the empty tg2 states of the TiO2 CB33, showing localized states at about 3 eV above the EF. Hereinafter our analysis will be focused on the VB region, thus results comparing the PAW-PBE and HSE06 functional are shown in Supplementary Fig. The pDOS reveals that the Cr defect states within the TiO3+ defects in a neutral rutile lattice could be achieved by the electron transfer from neutral oxygen vacancies, VO0, or ionized VOn+ defects , to nearby Cr4+ defects. To perform and simulate this possibility, a single VO0 is created in the proximity of two CrTi0 defects resulting in the combined defect 2CrTi + VO studied in this work. Figure Ti + VO. The dotted lines indicate broken Cr–O and Ti–O bonds due to the removed oxygen atom (VO).Other way to obtain Crd orbitals. The presence of the VO defect generally results in displacements of the Ti and O atoms outward and inward, respectively, from the VO site35 In the combined defect, the Cr atoms [Cr(1) and Cr(2) in Fig. O0, but only a small amount (1.04d0) as compared to the nearest Ti atoms from the VO2+ defect . However, the next nearest Ti atom [Ti(1) in Fig. D0, equivalent to that calculated for VO2+ . Despite the fact that charge density in the bonded O atom [between Cr(1) and Cr(2) in the Fig. O, only small distortions can be observed around Cr atoms because of the ionic relaxation as compared to single VO2+. The isosurfaces represented in the Fig. Ti1− defect [Fig. 3+ ions. However, in this case the electrons from the VO state are bounded with Cr atoms indicating the presence of a mixed Cr(3d)–VO electronic state. The Total DOS and pDOS corresponding to O(2p), Ti(3d) and Cr(3d) states for the combined defect 2CrTi + VO are represented in Fig. Ti1− defect [top in Fig. d) nature at about 0.75 eV above the VBM. For the combined defect there are also partially occupied states close to the EF which increase the energy splitting between the VBM and the EF from 0.95 eV, for CrTi1−, to 1.2 eV, for the combined 2CrTi + VO. In this case the electronic states of the VO0 are mixed with the nearest Cr tg2 orbitals resulting in a high spin polarized state, similar to that observed for the CrTi1−, which means that electrons from VO contribute to the formation of a pair of Cr3+ defects.After ionic relaxation, the trapped electrons from the removed O atom are localized in the region between the Cr and Ti dangling ect Fig. . Therefo2. The combination of these experimental results with the theoretical study will extend the understanding of the Cr doping process in rutile TiO2. In this way, XAS spectra have been acquired on the Cr doped microtubes, as well as on undoped TiO2 and Cr2O3 reference samples. The Fig. 2,3, O-K and Cr-L2,3 absorption edges, respectively, from the doped Cr03 microtubes. Spectra from undoped TiO2 and Cr2O3 are also included for comparison. The Ti-L2,3 edge is splitted in two regions due to the spin-orbit coupling forming the L3 and L2 edges, as marked in Fig. XAS and XPS measurements have been performed at BACH beamline at the Elettra synchrotron facility in order to investigate the electronic structure and the effect of Cr in rutile TiO3 and L2 bands are subdivided in two contributions as a result of the crystal field (CF), that originates the fivefold d-orbital splitting into two degenerated levels of tg2 (three-fold degenerate states) and eg (two-fold degenerate states) symmetry by an energy amount of 10 Dq eV, also known as CF splitting36. According to the atomic coordinates calculated by DFT, TiO6 octahedron in the rutile phase presents both trigonal and tetragonal distortions. TiO2 tetragonal distortions, affecting directly to apical Ti–O bonds, are more sensitive in the L2,3-edge. Under tetragonal distortion the eg states are splitted in two states of bg1 and ag1 symmetry, whereas tg2 states are splitted in three states of bg2 and eg (two-fold degenerated) symmetry37. The high energy resolution achieved in the acquisition of the XAS spectra (<0.1 eV) enables the observation of the eg splitting at the L3 edge, resulting in the characteristic rutile TiO2 fingerprint with a maximum at 460 eV and a shoulder around 459 eV37, as observed in Fig. 2 edge, although the absorption bands in the latter region are broader because of Auger decay effects and vibrational dispersions38. It can be noticed that in Fig. eg band presents slight, but not negligible, modifications in Cr doped samples such as a less defined eg splitting at the L3 edge or variations in the eg relative intensity at the L2 edge.Both Ls → 2p, also shows features that are sensitive to the crystal environment. The hybridization of the Ti(3d)–O(2p) orbitals in TiO2 makes able the observation of the tg2 and eg bands in the O-K edge32, observed at 530 and 533 eV, respectively. Complex contributions at higher energies with two maxima at 539 and 542 eV and a shoulder at around 545 eV, are also characteristic of the rutile phase of TiO239. In this case variations related to the eg relative intensity can be observed in Fig. 2 edge in Fig. On the other hand, the O-K edge Fig. , corresp2,3 edge [Fig. 3+ or Cr4+ oxidation states, making difficult the analysis of the Cr-L edge. Thus, the Cr-L2,3 edge from an α–Cr2O3 sample has been used as a reference and compared to the Cr-L2,3 from Cr03 microtubes. The main features in the corresponding XAS spectra in Fig. 3–L2 separation), indicate that the most probable oxidation state in the Cr doped TiO2 microtubes is Cr3+ under octahedral coordination, in agreement with previous CL results19. However, the characteristic spectral features could be associated with the different octahedral distortions or local environment, as can be observed from the ball and stick diagrams for a CrO6 unit from α–Cr2O3 [Fig. 6 octahedron for a Cr3+ defect in rutile TiO2 [Fig. The Cr Ldge Fig.  shows a 2O3 Fig.  and the iO2 Fig. .2 sample used as a reference. VB-XPS spectra acquired with energy of 450 eV are shown in Fig. XPS spectra at the VB region (VB-XPS) from microtubes with different amount of Cr –σ and O(2p)–π states that are partially hybridized with Ti(3d)40, can be also observed, although less defined, in the Cr doped samples Cr02 and Cr03 [marked with arrows in Fig. 2, the VB is composed by a broad band, wider than that from undoped TiO2, centered at about 6 eV that extends approximately from 9 to 1 eV. Moreover, a shoulder close to the VBM, around 2.5–2.6 eV, is observed differently from the undoped TiO2. This band can be attributed to Cr(3d) states10, as it is higher as the amount of Cr increases. At the same time, the VBM shifts towards the EF as the amount of Cr increases in the microtubes, which should involve a less n-type behavior because of Cr doping at the surface of the probed microtubes.These contributions, attributed to O using a photon energy off-resonance (Eν = 450 eV) and on-resonance (Eν = 457.2 eV). The XPS signal from Ti states in the VB increases under on-resonance conditions, showing peaks at 7.3 eV and 0.7 eV below EF, marked with arrows in Fig. p)–Ti(3d) σ-bonding states40, whereas the band at 0.7 eV is associated to Ti(3d) states as the result of reduced Ti3+ ions due to presence of oxygen vacancies and structural defects35. The same experiment was carried out for Cr doped microtubes (Cr03), as shown in Fig. ν = 457.2 eV) which can be also related to O(2p)–Ti(3d) hybridized states. However, selecting a photon energy on-resonance for Cr (Eν = 577.5 eV) a peak at 2.7 eV clearly dominates the VB spectrum. This band at 2.7 eV was previously observed as a shoulder near the VBM in all the VB-XPS spectra measured for Cr doped samples [Fig. d) states10.Resonant XPS helps for the identification of valence states from Ti and Cr species. Figure tra Fig. . Figure les Fig. , confirm2-np and Cr03 samples, respectively. The insets of the Fig. F region. In the undoped sample the signal associated with Ti3+ defects at 0.7 eV below EF is clearly observed, indicating higher concentration of these defects in comparison to Cr doped TiO2, in which the signal is practically negligible as shown in the Fig. 3+ defects can be observed on-resonant conditions as a very weak band. This can explain the observed EF shift towards the VBM [Fig. 41.Figures VBM Fig.  as a res2 have been compared, as shown in Fig. Theoretical DFT calculations and experimental VB-XPS results acquired on Cr doped TiOO defect using the PBE functional. In the calculated combined defect the difference between the pDOS maxima related to Ti(3d) and Cr(3d) states is about 5.5 eV for HSE [Fig. d nature are located within the Eg of TiO2. The deviation between the Ti(3d)-Cr(3d) states separation calculated by DOS and experimental VB may be originated by the resultant structural relaxation of the Cr defects42, as well as the differences between the DFT model and the measurement conditions of the actual system.Among the considered defects, the best fit between experimental VB and calculated data is achieved for the 2Cr + VHSE Fig. , whereasHSE Fig. . Thus, t2 system, and more frequently at the surface43. According to our DFT results, in order to achieve charge neutrality, Cr atoms could trap electrons from VO0 defects and hence be reduced to the most stable oxidation state (Cr3+). Therefore, two Cr defects could be compensated by the creation of a single VO0. The presence of Ti3+ defects are related to the presence of VO0, in which two electrons can occupy Ti(3d) states reducing two nearby Ti4+ ions44. Kim et al.45 stated that Cr3+ defects in rutile TiO2 are promoted when the samples are grown under oxygen poor conditions and determined that the formation energy of VO2+ and Cr3+ is smaller than the formation energy of Cr4+. Taking into account that the formation energy of VO2+ is also smaller than VO1+ and VO0 35, it is expected that the preferential oxidation state of Cr in TiO2 should be Cr3+. Thus, a single VO0 could be changed into VO2+ by the reduction of Cr4+ ions into Cr3+, leading to higher concentration of coordinated Cr3+–Vo2+–Cr3+ defects and hence reducing the presence of Vo0. This is consistent with the observed quenching of the Ti3+ related emissions in the luminescence spectra of Cr doped TiO2 favoring the Cr3+ luminescent emissions, reported in previous works19. However, experimental VB spectrum shows that the EF location is about 2.6 eV above the Cr resonant peak, which means that, considering that the Eg of crystalline rutile is about 3.1 eV, the EF localization is still close to the CB of rutile TiO2.Oxygen defects and other structural defects, such as Ti interstitial, can be found naturally in a real TiOp) orbitals in TiO2) are very sensitive to local distortions in the rutile lattice32. A complementary study on the CF splitting and tetragonal distortions associated with the presence of Cr in rutile TiO2 has been also performed by a deeper analysis of the Ti-L3 absorption edge. The tetragonal distortion, which is due to the elongation of the apical Ti–O bonds, has a strong influence on the Ti-L2,3 edge features rather than trigonal distortions, which can be neglected in TiO238. This implies that a simplified Dh4 symmetry is enough to describe the Ti4+ ion in the rutile cell38. Figure 3 edge for undoped TiO2 and Cr-doped TiO2 microtubes (Cr02 and Cr03), where slight variations can be appreciated as the concentration of Cr increases. As aforementioned, Cr incorporation modifies slightly the L3 edge, and in a greater extent the eg related band. In this case, by Gauss-Lorentz (G-L) band deconvolution, the L3 edge can be decomposed into four bands labeled as A, B, C and D. It can be observed in Fig. A and B bands are nearly overlapped for tg2 states. Thereby, tg2 states are identified by the A + B maximum. C and D bands, located at higher energies and related to the eg states, show an absolute energy difference |EC − ED| of about 1 eV. Information related to the crystal environment of Ti ions can be extracted from the G-L deconvolution of experimental Ti L3 edge, in which ED − EA+B could be related to the variations of the CF splitting (10Dq) and the difference |EC – ED| could be associated with variations of the tetragonal distortion parameter defined in absolute value as |E(bg1) − E(ag1)| = |δ|38.The effect of the crystal field (CF) is directly related to the coordination symmetry of Ti atoms, which implies that the hybridized metal-ligand bonds (O(2Dq) and the |δ| parameter as a function of the amount of Cr, in cationic fraction, calculated from the Ti-L3 absorption edge corresponding to samples Cr02, Cr03 and samples with variable content of around 4 cat.% and 5 cat.% Cr, composed by sintered grains and precursor powder respectively. In both cases, the CF splitting decreases when Cr dopants are incorporated up to 3 cat.% Cr, and tends to increase when the amount of Cr is above this value. According to our previous results, the Cr solubility limit observed for rutile TiO2 microtubes grown by vapor-solid method is in the order of 3 cat.% Cr, which could indicate that the Cr defect site in the rutile lattice varies for concentrations either below or above this value. The |δ| parameter, associated with tetragonal distortion, also presents a slight reduction up to 3 cat.% Cr. Analyzing the calculated atomic coordinates of the Ti and O atoms surrounding the CrTi1− defect, a tetragonal distortion reduction has been observed in some of the TiO6 octahedra. The apical (D) and basal (d) Ti–O bonds of the in-plane TiO6 octahedra which are sharing corners with the CrO6 octahedron are D = 2.000 Å and d = 1.967 Å respectively, so the relative difference |D − d|/D is reduced from 2.4% (bulk rutile TiO2) to 1.6% using the calculated D and d, indicating a reduction of the tetragonal distortions. This is in agreement with the reduction of the |δ| parameter calculated by the G-L deconvolution of experimental Ti-L3 spectra.Figures 2, the physical behavior of the Cr defects should be also dependent on their local environment. The relatively small tetragonal distortion calculated for both CrTi0 and CrTi1− defects indicates that the CrO6 octahedra in TiO2 present Oh symmetry independently of the Cr oxidation state. XAS, XPS and DFT results indicate a predominant Cr3+ ion in our system. Previous work19 reported a characteristic emission related to Cr3+, also called R-lines, using CL spectroscopy as an evidence of the presence of octahedral coordinated Cr3+ ions in our Cr-doped TiO2 microtubes. However, that emission was observed at low temperature (T = 110 K). Taking into account the energy of the Cr3+ characteristic emission (1.79 eV)19 and the corresponding Tanabe-Sugano diagram for d3 systems46, the minimum crystal field value (Dq) necessary to observe the R-lines is about 1.8 eV. On the other hand, Urusov and Taran47 reported the evolution of 10Dq as a function of the Cr–O distance (dCr-O) for a large number of compounds with octahedral coordinated Cr3+ ions obtaining the relation 10Dq ∝ (dCr-O)−5, revealing the high sensitivity of the CF to the local environment. In accordance to our calculated Cr–O distances (1.99–2.00 Å) and applying a correction factor of 0.99 to arrange the calculations to experimental TiO2 lattice parameters, the 10Dq parameter could vary from 2.05 to 2.12 eV. Therefore, according to the Tanabe-Sugano diagrams, it should be possible to observe the R-lines in Cr-doped TiO2. However, the energy separation between Ti3+ (0.9 eV below EF) and Cr3+ (2.6 eV below EF) defect states obtained from the VB spectrum [Fig. O plays an important role on the recombination process of Cr3+ ions in rutile TiO2. As the Cr concentration increases the charge neutrality can be achieved by creating Cr4+ defects or by inducing VO that, according to our results, in the most stable structures tends to transfer electrons to the Cr atoms resulting in an increment of VO2+ defects.As the local distortions induced by Cr affects directly to the crystal environment of TiOrum Fig.  is about2 micro-tubes have been employed as model material to success-fully compare theoretical simulations and experimental results on the Cr doping of rutile TiO2. As a result, a deeper comprehension on the Cr incorporation in the rutile structure has been obtained. Aspects regarding the incorporation of chromium by forming a complex defect with two Cr atoms and one O vacancy (2CrTi + VO) have been elucidated. The electrons from the oxygen vacancy tend to be localized at the tg2 states of the Cr ions in order to reach the stable oxidation state of Cr3+. These results, showing that oxygen defects play a crucial role in the stabilization of Cr3+ in the rutile lattice, have been confirmed both theoretically and experimentally, as the separation between the theoretical states due to Ti(3d) and Cr(3d), calculated by the pDOS, is in agreement with the one obtained experimentally by XPS resonant measurements. Moreover Cr effects on the crystal field and tetragonal distortion have been studied both from DFT simulations as well as by fitting experimental XAS measurements. The results indicate a decrease in the value of the 10Dq parameter and the tetragonal distortion |δ| in samples with Cr content lower than 3 cat.%, whereas these parameters increase for samples with higher Cr concentrations. Cr doping of rutile TiO2 leads to the generation of an energy level 0.55 eV over the VBM of the TiO2 as obtained by DFT simulations, which is in agreement to the Cr-resonant XPS measurements of the VB. These Cr related shallow levels behave competitively with the Ti3+ defect related level, as measured by luminescence and XPS.In summary, single crystalline Cr doped TiOSupplementary Information