Fusarium oxysporum is a soilborne fungal pathogen that causes Fusarium wilt, one of the most important diseases on tomato. During the infection of F. oxysporum, some proteins are secreted that modulate host plant immunity and promote pathogen invasion. In this study, we identify an RNase, FoRnt2, from the F. oxysporum secretome that belongs to the ribonuclease T2 family. FoRnt2 possesses an N‐terminal signal peptide and can be secreted from F. oxysporum. FoRnt2 exhibited ribonuclease activity and was able to degrade the host plant total RNA in vitro dependent on the active site residues H80 and H142. Deletion of the FoRnt2 gene reduced fungal virulence but had no obvious effect on mycelial growth and conidial production. The expression of FoRnt2 in tomato significantly enhanced plant susceptibility to pathogens. These data indicate that FoRnt2 is an important contributor to the virulence of F. oxysporum, possibly through the degradation of plant RNA.Secreted RNase proteins have been reported from only a few pathogens, and relatively little is known about their biological functions.  Fusarium oxysporum, is able to degrade host RNA and represents an important contributor to the virulence of the soilborne fungal pathogen.FoRnt2, a member of the ribonuclease T2 family secreted from  Secreted RNases have been reported in many pathogens and play crucial roles in plant–pathogen interactions. For example, Nuc1 and Nuc2, which are secreted from F. oxysporum is an important soilborne fungal pathogen that has a broad host range, causes root rot and wilting disease in the plant vascular system, and results in economic loss worldwide . This pathogen has many different formae speciales that can infect only one or a few host species. F. oxysporum can infect the vascular bundle of host plants, leading to clogged vessels, yellowing of leaves, wilting and finally death of the whole plant. The germination of dormant spores in soil results in invasion of plant roots by fungal hyphae. The movement of hyphae from root cortex to the xylem, where it produces and disseminates microconidia, is critical for disease progression. It is difficult to efficiently control this pathogen because F. oxysporum is soilborne and can produce chlamydospores in the soil. Chlamydospores are resistant to different environments, such as high temperature and drought. To control F. oxysporum, chemical fungicides are widely used. However, the resistance of this pathogen to fungicides is becoming increasingly important. To date, no effective methods to control the disease have been used in the field  proteins, which play important roles in host specificity. Six1 is required for full virulence and is responsible for avirulence in tomato plants carrying the resistance (R) gene I‐3. The effector Six6 is essential for the virulence of F. oxysporum and suppresses host cell death  in this study. In addition, the results from SignalP 5.0 predicted that FoRnt2 possesses a secretion signal peptide domain in residues 1–19 in the N‐terminal region and a cleavage site between residues 19 and 20 . The two catalytic sites of the FoRnt2 fusion protein were also mutated to other amino acids, and the FoRnt2M2 recombinant protein was then expressed under the same conditions  tag in  Figures  and S2.M2 recombinant proteins was tested by incubation with tomato total RNA. FoRnt2 could significantly degrade the tomato total RNA, while the FoRnt2M2 proteins could not degrade the RNA under the same conditions, indicating loss of RNase activity  strain  and Congo red (CR). Surprisingly, we found that compared with those of the WT and ∆FoRnt2‐C strains, the growth rate of the ∆FoRnt2 strain showed no difference under any of the stresses on PDA, implying that deletion of the FoRnt2 gene did not affect cell wall integrity or osmotic regulation  localizes in the nucleus and was used to verify the nuclear localization of N. benthamiana  alone or FoRnt2 fused with GFP at its C‐terminus  in a Figure . To furts Figure , which f2.6GFP or FoRnt2‐GFP transgenic tomato plants. Microscopic observation showed that green fluorescence signals were distributed in the tomato cells of the two transgenic plant lines  was used to analyse the differentially expressed genes (DEGs) in s Figure . The RNAs Figure . This res Figure . The tops Figure . All 12 s Figure , suggest2.7Agrobacterium‐mediated transgenic tomato seedlings were used in this assay. Immunoblot assays showed that GFP or FoRnt2 was expressed successfully in the stable Sl:GFP or Sl:FoRnt2 transformant plants, respectively  were cultured in lysogeny broth liquid or solid medium at 37°C for vector construction and protein expression. A. tumefaciens GV3101 was used to transiently express target proteins in plants. N. benthamiana and tomato (Solanum lycopersicum ‘Ailsa Craig’) were grown at 25°C under a 16 h light and 8 h dark photoperiod in an artificially controlled growth room.The WT 4.2FoRnt2 gene deletion mutant . All the transformant strains were purified by single‐spore isolation and then stored at −80°C in 30% glycerine. The primers used in the construction of different strains are listed in Table The split‐marker approach was used in this study to generate a gene replacement for the 4.3The cDNA sequences of the signal peptide regions of FoRnt2 and PsAvr1b were cloned into the pSUC2 vector, which carries a truncated invertase gene lacking both the initiation residue Met and a signal peptide  and anti‐actin antibodies .To test the secretion of the FoRnt2 protein, the FoRnt2‐FFLAG strain was cultured in potato dextrose broth to harvest the conidia. Then, the conidia were transferred into 10% YEPD liquid medium with the tomato roots and cultured at 25°C and 180 rpm for 16 h. The culture supernatants were collected after centrifugation. The total proteins in the culture supernatants were precipitated by adding 20% acetone (wt/vol) and stored at −80°C overnight. The solution was centrifuged at 13,800 × 4.5FoRnt2 gene constructs were transformed into A. tumefaciens GV3101 through heat shock treatment. The correct A. tumefaciens clone was cultured overnight in LB medium  at 28°C. The A. tumefaciens cells were collected and resuspended in infiltration buffer at an OD600 of 0.4. This experiment was carried out on 3‐week‐old N. benthamiana leaves using needleless syringes. At the same time, the same A. tumefaciens strain or cells harbouring the empty GFP vector were used as the negative control and BAX was used as a positive control. For subcellular localization, fluorescence observation was carried out after treatment for 48 h. For fluorescence observation after plasmolysis, the leaves of N. benthamiana after Agrobacterium injection were treated with 0.8 M NaCl under static conditions. The inoculation for P. capsici was performed by the same method as described previously. The total proteins in the leaves of N. benthamiana were extracted with protein extraction buffer and the proteins were detected by SDS‐PAGE or western blotting.The 4.6FoRnt2 or GFP transgenic tomato plants, the FoRnt2 gene was cloned into a GFP construct with the CaMV 35S promoter. The resulting constructs and the 35S::GFP empty vector were transformed into WT tomato plants separately using Agrobacterium‐mediated transformation as described previously  and active site mutant sequences were cloned into the expression vector pMAL‐c2x. The recombinant vectors or empty vector were transformed into E. coli BL21 (DE3) and the cells were cultured in LB liquid medium at 37°C. Then, the cells were diluted into fresh LB grown to a final OD of 0.6. IPTG (0.1 mM) was used to induce the expression of the recombinant proteins for 16 h at 16°C. The cells were harvested by centrifugation at 9,600 × g and then sonicated at 300 W for 10 min. Recombinant proteins were purified using amylose resin (BioLabs) and the protein concentration was determined using a BCA Protein Assay Kit (Solarbio). The primers used for the expression constructs are listed in Table The coding sequence of the 4.9M2 (MBP‐FoRnt2H80F/H142R) proteins at 25°C for 30 min. Five micrograms of RNase A with equal amounts of total RNA was used as a positive control in this experiment. All samples were mixed with 1× loading buffer and then run in a 0.1% agarose gel.The enzymatic activity of FoRnt2 was tested on total tomato RNA in an in vitro assay  according the manufacturer's instructions. Reverse transcription was performed using the All‐In‐One 5× RT MasterMix Kit (Abm). HiPer SYBR Premix EsTaq  was used for qPCR to analyse the expression pattern. The relative expression of each gene was calculated using the 2Ct−∆∆ method as previously described. All experiments were repeated three times. The primers used in RT‐qPCR are listed in Table Total RNA was extracted from the 4.11http://www.cbs.dtu.dk/services/SignalP/). The conserved domains of the proteins were identified using the Pfam database (http://pfam.xfam.org/).The homologous protein sequences of FoRnt2 in different pathogens were identified by querying the FoRnt2 protein sequence against the NCBI database using the BLAST tool. The protein sequence alignment of the FoRnt2 protein and other ribonuclease T2 proteins was performed using the Clustal W2 program. Phylogenetic dendrograms of FoRnt2 and its homologous proteins were generated using MEGA 5. The signal peptides of each protein sequence were predicted by the SignalP 5.0 server . The crude samples were centrifuged at 4°C for 15 min at 13,800 × Figure S1 Western blotting confirmation of FoRnt2‐overexpressing strains using anti‐FLAG antibodies. The targeted protein was detected in the #1 and #2 strainsClick here for additional data file.Figure S2 SDS–PAGE analysis of the FoRnt2 recombinant proteins used in this study. The proteins were expressed in Escherichia coli BL21 (DE3) and purified using amylose resin (BioLabs). The protein concentration was determined by a BCA kitClick here for additional data file.Figure S3 Targeted gene knockout of FoRnt2 in Fusarium oxysporum. (a) Schematic diagram of the gene deletion strategy for FoRnt2 gene in F. oxysporum using the split‐marker method. (b) PCR products of each fragment and fusion amplification of split‐marker PCR. (c) Agarose gel of PCR products from the wild type (WT) and deletion strain genomic DNA. The F. oxysporum WT strain was used as a positive control and water was used as a negative control. M represents the molecular markers of DNA fragment size; In represents the FoRnt2 gene; Out represents whether the hph gene is correctly embedded. (d) PCR confirmation of the ∆FoRnt2‐C strain. FoRnt2 genes of the same size were detected in the WT and ∆FoRnt2‐C strains but not in the ∆FoRnt2 strainClick here for additional data file.Figure S4 The numbers of differentially expressed genes (DEGs) in the FoRnt2 transgenic tomato plants. The x axis shows the difference between green fluorescent protein (GFP) and FoRnt2‐GFP expressed in tomato. The y axis represents the number of DEGs. Red and blue colours represent up‐regulated and down‐regulated DEGs, respectivelyClick here for additional data file.Figure S5 The morphology of 4‐week‐old green fluorescent protein (GFP) or FoRnt2‐GFP transgenic tomato and wild‐type tomato. All tomato seedlings were grown at 25°C under a 16 h light and 8 h dark photoperiod in an artificially controlled growth roomClick here for additional data file.Figure S6 FoRnt2 could not induce cell death in Nicotiana benthamiana. (a) The cell death‐inducing ability of FoRnt2 was determined in 3‐week‐old N. benthamiana leaves infiltrated with Agrobacterium tumefaciens carrying the target gene. BAX was used as a positive control and green fluorescent protein (GFP) was used as a negative control. The photographs were taken 4 days post‐agroinfiltration. (b) Immunoblot analysis showed the protein expression levels of GFP or GFP‐tagged protein in N. benthamiana leaves. Equal amounts of protein were confirmed by Ponceau S staining on the membraneClick here for additional data file.Table S1FoRnt2‐GFP The unique differentially expressed genes identified in FoRnt2‐GFP expression tomato plantsClick here for additional data file.Table S2 KEGG enrichment analysis of the down‐regulated genes in tomato plantsClick here for additional data file.Table S3 Primers used in this studyClick here for additional data file.


All samples crystallize from the melt in mixtures of α and γ forms and the fraction of γ form increases with decreasing cooling rate. At high cooling rates, the crystallization of the α form is always favored, even for samples that contain high total concentration of defects that should crystallize in the γ form. The results demonstrate that in iPPs containing significant concentrations of defects, such as stereodefects and comonomeric units, the γ form is the thermodynamically stable form of iPP and crystallizes in selective conditions of very slow crystallization, whereas the α form is the kinetically favored form and crystallizes in conditions of fast crystallization.Random isotactic propene-butene copolymers (iPPC4) of different stereoregularity have been synthesized with three different homogeneous single center metallocene catalysts having different stereoselectivity. All samples crystallize from the polymerization solution in mixtures of α and γ forms, and the relative amount of γ form increases with increasing concentrations of butene and of  The crystallization of α and γ forms of isotactic polypropylene (iPP) depends on the conditions of crystallization and on the molecular structure of polypropylene; the latter depends on the polymerization conditions and catalysis . DiffereThe α form is considered the most stable form of iPP and crystallizes usually in the iPP homopoymer prepared with heterogeneous Ziegler-Natta catalysts in common crystallization conditions from the melt or from solution, and in stretched fibers ,7,8. In The crystallization of the γ form in chains containing defects is due to the fact that the γ form crystallizes when iPP chains are characterized by short regular propene sequences, therefore, it occurs when iPP chains contain any type of defect that interrupts and shortens the regular propene sequences ,54,62,79The crystallization of α form is, instead, favored when the regular propene sequences are very long, which are generated when the concentration of defects is low or when defects are segregated in blocks of the macromolecules ,54,62,79Understanding the conditions of crystallization of α and γ forms of iPP is of particular relevance since the two different polymorphs exhibit significant differences in mechanical behavior ,100,101.In this paper we report a study of the crystallization of α and γ forms of iPP in propene-butene copolymers in different crystallization conditions to analyze the thermodynamic and kinetics effects on the crystallization of the two polymorphic forms. Propene-butene copolymers of different stereoregularity synthesized with different metallocene catalysts have been chosen for this study because they crystallize easily in the γ form, thanks to the incorporation in the iPP chains of butene comonomeric units that shorten the length of the regular propene sequences . Hence, 2 and passing over activated Al2O3 , and stored under nitrogen. The MAO cocatalyst was used as received . The catalyst mixture was prepared by dissolving the desired amount of the metallocene with the proper amount of the MAO solution, obtaining a solution which was stirred for 10 min at 25 °C before being injected into the reactor. All copolymerizations were run at 25 °C in a 250 mL Pyrex reactor, agitated with magnetic stirrer, containing toluene (100 mL) and MAO (2.0 mL). Gas mixtures of propene and 1-butene at the appropriate composition, prepared with vacuum line techniques in a gas cylinder at pressure of 4–5 bar and standardized by gas chromatography, were bubbled through the liquid phase at atmospheric pressure and a flow rate of 0.3 L/min. The polymerizations started by syringing in the toluene solution of the catalyst (2–3 mg) and proceeded under a constant flow of the gas mixture. Under such conditions, total monomer conversions were lower than 15%, this ensuring a nearly constant feeding ratio. The copolymers were coagulated with excess methanol acidified with enough HCl  to prevent the precipitation of alumina from MAO hydrolysis, filtered, washed with further methanol, and vacuum-dried. Typical yields were 2–5 g with a 120 min reaction time. Propylene-butene isotactic copolymers (iPPC4) were synthesized with the metallocene catalysts of 2-symmetric metallocene Rf is highly isospecific [rr triads) and small amount of regiodefects around 0.1–0.5 mol%, represented by secondary 2,1-erythro units . The two C1-symmetric metallocenes Z4 and Z9 are instead fully regioselective but introduce significant amounts of rr stereodefects [rr stereodefects, whereas samples synthesized with the catalyst Z9 (samples iPPC4Z9-x) contain 2.5–3.4 mol% of rr defects [The Cspecific ,105 and odefects ,108,109. defects ,108. Con defects . The com13C NMR analysis (2 at 120 °C. The carbon spectra were acquired with a 90° pulse and 15 s of delay between pulses and CPD (WALTZ 16) to remove 1H-13C coupling. About 1500–3000 transients were stored in 32K data points using a spectral window of 6000 Hz. For all copolymer samples, the peak of the propylene methine carbon atoms was used as internal reference at 28.83 ppm. The 13C NMR spectra of two samples of iPPC4 copolymers are reported in rP × rB ≈ 1, calculated using the The composition and comonomer distribution were determined by analysis . All spe6. The molecular masses and the dispersity were determined by gel permeation chromatography (GPC), using a Polymer Laboratories GPC220 apparatus equipped with a differential refractive index (RI) detector and a Viscotek 220R viscometer , on polymer solutions in 1,2,4-trichlorobenzene at 135 °C of 2 mg/mL concentration. The injection volume was 300 μL with a flow rate of 1.0 mL/min. The GPC apparatus was calibrated with 12 standard samples of polystyrene having narrow dispersity and molecular masses in the range 580 and 13.2 × 102 atmosphere and scanning rate of 10 °C/min.The calorimetry measurements were carried out with differential scanning calorimeter (DSC) Mettler Toledo DSC-822  performing scans in a flowing NX-ray powder diffraction profiles were recorded with Ni filtered Cu Kα radiation by using an Empyrean diffractometer .All samples of iPPC4 copolymers were crystallized in DSC by cooling the melt at 180 °C down to 25 °C at different cooling rates from 1 to 40 °C/min. After the crystallization in DSC, the samples were analyzed by X-ray diffraction.fγ, with respect to the α form, was calculated from the intensities of the (117)γ and (130)α reflections at 2θ = 20.1° of the γ form and at 2θ = 18.6° of the α form, respectively, as the ratio: In samples that crystallized in mixtures of α and γ forms, the relative fraction of the γ form α and (117)γ reflections at 2θ = 18.6 and 20.1° of the α and γ forms, respectively, indicating that all as-prepared samples of iPPC4 copolymers are crystallized in mixtures of α and γ forms. The intensity of the (117)γ reflection increases with increasing butene concentration and, at the same butene content, is higher in the low stereoregular samples synthesized with the catalysts Z4 and Z9  calculated from the intensities of the (117)γ and (130)α reflections in the diffraction profiles of fγ = 0 for butene concentration equal to zero. It is apparent that the amount of γ form increases with increasing concentration of butene and of rr stereodefects [fγ ≈ 80%) for butene concentration of nearly 9 mol%  crystallizes at about 6 mol% of butene  samples of iPPC4 copolymers of odefects  γ reflection at 2θ = 20.1° of the γ form increases with decreasing cooling rate. The amount of γ form is reported in rr defects concentrations) that tend to crystallize normally in the γ form. In fact, even the less isotactic samples iPPC4Z4-3 and iPPC4Z9-3 with the highest butene concentrations of 8.2 and 6.4 mol%, respectively, that crystallize in the as-prepared samples in the almost pure γ form  These results demonstrate that in iPP samples containing significant concentrations of defects, such as stereodefects and butene comonomeric units, the γ form is the thermodynamically stable form of iPP and crystallizes normally from polymer solution or in selective conditions of very slow crystallization, as isothermal crystallizations from the melt at high crystallization temperatures  or by co2-symmetric metallocene produces highly isotactic iPPC4 copolymers containing negligible amounts of stereodefects, whereas the two C1-symmetric metallocenes produce less isotactic copolymers containing on average 2.0–2.4 mol% and 2.5–3.4 mol% of rr stereodefects, respectively.Random isotactic propene-butene copolymers of different stereoregularity have been prepared with three different metallocene catalysts having different stereoselectivity. The Crr stereodefects.All as-prepared samples crystallize in mixtures of α and γ forms from the polymerization solution and the relative amount of γ form increases with increasing concentrations of butene and of The samples have been crystallized from the melt by cooling the melt from 170 °C down to 25 °C at different cooling rates. All samples crystallize from the melt in mixtures of α and γ forms and the amount of the two forms strongly depends on the cooling rate. The amount of γ form increases with decreasing cooling rate. Correspondingly, the γ form almost disappears and the almost pure α form crystallizes in all samples at the highest cooling rate of 40 °C/min. The highest concentration of γ form is always obtained at low cooling rates, whereas at high cooling rates the crystallization of the α form is always favored, even for samples that contain high total concentration of defects that should crystallize in the γ form.These results demonstrate that in iPP samples containing significant concentrations of defects, such as stereodefects and butene comonomeric units, the γ form is the thermodynamically stable form of iPP and crystallizes normally from polymer solution or in selective conditions of very slow crystallization, such as by cooling from the melt at very low cooling rates, whereas the α form is the kinetically favored form and crystallizes in conditions of fast crystallization, such as fast cooling from the melt.


Photoluminescence spectra of DY1/MoS2 show quenching of the DY1 emission and enhancement of the MoS2 emission, indicating a strong electronic interaction between these two materials. Nanoscopic mappings of the light‐induced contact potential difference changes rule out the CT process at the interface. Using femtosecond transient absorption spectroscopy, the rapid interfacial ET process from DY1 aggregates to MoS2 and a fourfold extension of the exciton lifetime in MoS2 are elucidated. These results suggest that the integration of 2D inorganic semiconductors with fluorescent molecules can provide versatile approaches to engineer the physical characteristics of materials for both fundamental studies and novel optoelectronic device applications.Integration of distinct materials to form heterostructures enables the proposal of new functional devices based on emergent physical phenomena beyond the properties of the constituent materials. The optical responses and electrical transport characteristics of heterostructures depend on the charge and exciton transfer (CT and ET) at the interfaces, determined by the interfacial energy level alignment. In this work, heterostructures consisting of aggregates of fluorescent molecules (DY1) and 2D semiconductor MoS 2 monolayers provides energy level alignment appropriate for exciton transfer. Optical characterizations and probe‐based surface charge distribution measurements reveal rapid transfer of photogenerated excitons at the interface. This work highlights a strategy to enable exciton funneling and boosts visible light emission in 2D inorganic semiconductors.Integration of blue‐fluorescent organic molecules and atomically thin MoS After dissociation, photogenerated electrons and holes in type‐I (II) heterojunctions move in identical (opposite) directions. Thus, spontaneous transfer of photogenerated excitons under light illumination can occur from a high bandgap layer to a low bandgap layer in type‐I heterojunctions. Such exciton funneling leads to enhanced light emission. Type‐II heterojunctions, in which the transfer of net charge from one layer to the other occurs, are beneficial for the collection of photogenerated carriers in solar cells.Proper combination of constituent materials is key to tuning the photophysical properties of semiconductor heterostructures, which depend on exciton formation and harvesting at the interfaces.6, 72 compounds, in which the transition metal atom, M, is sandwiched between two chalcogen atoms, X. TMDs have attracted growing research attention for optoelectronic device applications since they have sizable bandgap energies (≈1.2–1.8 eV) and optical characteristics tunable by external stimuli, such as strain and electric/magnetic fields. TMD‐based heterostructures have been fabricated similar to conventional semiconductors. Energy band structure engineering at heterointerfaces enables control of the electrical transport and recombination of charge carriers in heterostructures. Among numerous TMDs, MoS2 has been regarded as one of the strongest candidates for novel optoelectronic device applications.Transition metal dichalcogenides (TMDs) are 2D layered inorganic semiconductors that have the form of MXc fields.20 TMD Their utility benefits from the high versatility in the design and synthesis of organic materials. For instance, recent advances have been made in OLEDs by taking advantage of the ability of molecules to accommodate transition metals or to spatially separate frontier orbitals. These molecules are routinely deposited onto semiconductors and metals to form solid‐state devices. Accumulated knowledge indicates that the intrinsic electronic properties of organic compounds are perturbed at interfaces with semiconductors and metals. Exploitation of the full potential of molecules thus requires an understanding of the electronic processes that occur at the heterojunction between organic aggregates and other materials.Conjugated organic compounds have emerged as valuable active materials in a variety of devices, such as organic thin film transistors, solar cells, and organic light‐emitting devices (OLEDs).31 The For example, fabrication of hybrid structures can improve the electrical transport and reliability of OLEDs due to the high carrier mobility and chemical stability of the TMD layers. Additionally, stacking of organic layers on TMDs is beneficial for improving the light‐matter interaction in TMDs. TMDs usually form type‐II band alignment, and charge transfer (CT) rather than exciton transfer (ET) is dominant at the interfaces. Insertion of insulating layers between two TMD layers can suppress the interlayer CT process, and hence, the ET process enables enhanced light emission in TMD heterostructures. Recent works have shown that combinations of suitable materials allow fabrication of both type‐I and type‐II organic/TMD heterojunctions. Qiao et al. fabricated heterostructures consisting of MoSe2 monolayers  and pentacene (Eg: 1.8 eV). They reported 3.7‐times‐enhanced photoluminescence (PL) intensity from MoSe2 monolayers at the emission wavelength of ≈800 nm. Park et al. fabricated a hybrid structure composed of MoS2 monolayers (Eg: 2.11 eV) and perylenetetracarboxylic dianhydride (Eg: 2.55 eV), which exhibited a twofold‐enhanced visible‐light PL yield of MoS2 monolayer. Since many visible‐light‐fluorescent organic materials are available, efforts to find materials for type‐I organic/TMD heterojunctions and their careful characterizations should deserve considerable attention for both fundamental science and device applications. Also, systematic investigations to study interfacial excitonic behaviors in such hybrid systems should be established.In organic/TMD hybrid systems, the device performance can be improved with the help of synergistic effects of the two materials.6, 7 DY1 is based on a planar pyrazinoquinoxaline core; the planar structure enables the formation of self‐assemblies of DY1. The highly symmetrical molecular structure is anticipated to facilitate the formation of molecular films of DY1 on flat substrates. The propensity for intermolecular packing and lateral layering would provide a valuable opportunity to investigate the behavior of molecular excitons at heterojunctions. In this work, we prepared fluorescent molecular aggregates of DY1 on MoS2 monolayer flakes to form hybrid heterostructures. Steady‐state photophysical measurements were carried out to investigate the interfacial excitonic processes at the DY1/MoS2 interface. Light‐induced contact potential difference (CPD) changes were studied to examine the photoinduced modifications of the surface charge distribution at the heterointerface. Furthermore, femtosecond transient absorption techniques were employed to study the dynamics of the photogenerated excitons at the heterojunctionWe envisioned that understanding the ET and CT processes in heterostructures consisting of fluorescent organic materials and TMDs would provide novel insights into the construction of future devices involving organic/TMD heterostructures. DY1, a blue‐fluorescent molecule recently developed in our group, exhibits excellent photochemical and electrochemical stabilities.29 DY12Figure and a schematic diagram of a DY1/MoS2 heterostructure on a quartz substrate. MoS2 monolayer flakes were grown on quartz substrates by chemical vapor deposition (CVD)  image , is as high as several tens of nm ” and DY1(F)’. The DY1‐coated MoS2 flake is called “DY1/MoS2,” which contains high‐density DY1 aggregates.e Figure  shows th2/quartz obtained using a bright‐field optical microscope . A white LED source  with a 550‐nm short‐pass filter was used as an excitation source, and the emission image was taken by a color camera  with a 600‐nm long‐pass filter. The PL intensity of DY1(F) is weaker than that of DY1(A) due to the aforementioned difference in thickness. In the case of DY1/MoS2, the PL intensity near the MoS2 flake edge is much stronger than that in the other regions. As shown in the SEM and AFM images .33 A w2 regions in DY1/MoS2/quartz and MoS2 flakes on quartz substrates obtained with an excitation wavelength (λex) of 532 nm (2.33 eV). The beam diameter of the excitation laser light was ≈1 µm, which was small enough to collect the spectra from selected regions: DY1(A), DY1(F), and DY1/MoS2. The PL spectrum of DY1(A) exhibits significantly distinct features compared with that of DY1/MoS2. DY1(A) and DY1(F) exhibit very broad emission. The PL intensity of the former is higher than that of the latter due to the aforementioned difference in their thickness. DY1/MoS2 and the pristine MoS2 flakes show strong emission peaks at a wavelength (λ) of 670 nm, which are attributed to the A exciton resonance of the MoS2 monolayers (bandgap energy: 1.8 eV). The PL intensity of DY1/MoS2 at an emission wavelength (λem) of 570 nm, where a broad peak appears for DY1(A), is only 20% of that of DY1(A), which reveals that the atomically thin MoS2 monolayers can drastically suppress the light emission from the DY1 aggregates. In sharp contrast, the MoS2 excitonic resonance peak at λem = 670 nm of DY1/MoS2 is 1.7 times larger than that of the pristine MoS2 flake, even though the DY1 layer on the MoS2 flake reflects and absorbs part of the incident excitation light. The differential reflectance measurements show that the emission spectrum of DY1 overlaps with the absorption spectrum of MoS2  using a confocal laser scanning microscope (CLSM). The laser beam was focused on the sample surface and the emission spectra at specific pixels in the CLSM image were reconstructed from the PL intensity recorded at each channel of a 32‐channel detector. Typical PL spectra of the distinct sample regions are shown in Figure 2 monolayers is too weak to be observed at a λex of 405 nm, and hence, the exciton emission from the MoS2 monolayer is hardly seen in Figure 2 relative to DY1(A). This result is qualitatively consistent with the PL spectra obtained under a λex of 532 nm  and minor (λem = 570 nm) emissions of DY1/MoS2 remain at only 70% and 40% of those of DY1(A), respectively. The greater decrease in the major peak can be attributed to the conformeric heterogeneity or the incomplete aggregate formation of DY1 molecules. The CSLM image in Figure λem of 494 nm is very close to the major emission peak of DY1 . All these results indicate a notable interaction at the MoS2/DY1 interface.m Figure . The PL 1 Figure . The emi2/quartz were performed using scanning confocal microscopy and a picosecond 405‐nm‐wavelength laser, as shown in Figure 2/quartz  is negligibly small, which enables estimation of the intensity‐weighted average lifetime, 〈τ〉, ofthe DY1 emission. As shown in Figure τ〉 of DY1/MoS2 (4.2 ns) is longer than those of DY1(A) (2.5 ns) and DY1(F) (1.7 ns)  measurements of DY1/MoSz Figure . The emi2 monolayers were transferred onto a DY1/quartz sample, as illustrated in the schematic diagram in Figure2 flake can be as large as 100 nm in DY1/MoS2. Even though light‐induced CT occurs at the interface, the resulting CPD change at the DY1/MoS2 surface could be too small to be measured due to the electric field screening. An AFM image near the edge of a MoS2 flake on a DY1 surface is shown in Figure 2 surface without serious concerns about the topographic artifacts. The CPD maps and histograms obtained from the region in Figure  Thus, the CPD measurements enable nanoscale mapping of light‐induced changes in the surface potential and charge distributions. The polarity and magnitude of the light‐induced CPD changes can reveal the types and amount of transferred charges. The results in Figure 2 in the dark. This relative CPD difference between MoS2 and DY1 (50 mV) is almost invariant under light illumination. This suggests that charging by exciton scission does not occur at the DY1/MoS2 heterointerface. Consequently, there is no signature of light‐induced CT at the DY1/MoS2 interface.For CPD measurements, MoS2 heterojunction, femtosecond transient absorption (TA) spectroscopy was performed at a pump wavelength of 450 nm (2.76 eV) and an excitation fluence of ≈30 µJ cm−2. As written above, DY1/MoS2 indicates the DY1‐coated MoS2 flake, containing high‐density DY1 aggregates  were obtained from pristine MoS2 and DY1/MoS2  peaks of the B exciton at 610 nm and the A exciton at 660 nm. In pristine MoS2, the PB peak of the B exciton is observed at 610 nm, but the PB peak of the A exciton is mixed with the broadband of positive absorption (photoinduced absorption). This can be attributed to the excitonic transitions via band filling. The position of each peak in the TA map is similar to that in the PL spectra  of the excitation photon energy, free carriers can be generated in MoS2. The free carriers in MoS2 form neutral excitons (A0) and trions (A–), of which TA signals are located near the energy around 1.9 eV (650 nm) and below, respectively. The TA spectrum of pristine MoS2  molecular orbital level of the evaporated DY1 molecules can be higher (lower) than the conduction  band edge of the MoS2 monolayers. Thus, a type‐I heterojunction is supposed to be formed at the DY1/MoS2 interface, as illustrated in Figure 2 under light illumination, which agrees with the CPD results . Immediately after light illumination, the density of A excitons in MoS2 exhibits a more rapid increase for DY1/MoS2 than for pristine MoS2. A excitons are generally formed after the generation of B excitons via hole cooling in monolayer MoS2. In DY1/MoS2, as described above, the energy is directly transferred from DY1 to MoS2 in the form of excitons. The transferred excitons instantaneously occupy excitonic states in MoS2. Consequently, the rise time of the A exciton peak for DY1/MoS2 (≈0.4 ps) is smaller than that for pristine MoS2, confirming the excitonic ET process across the interface , which originates from charge trapping mostly by hole trap states. The slow component (τslow), contributing 27% of the decay, corresponds to the radiative recombination time of excitons because the position of the steady‐state PL emission peak well matches the position of the A exciton peak in the TA spectrum , which contributes 27% of the decay, can be explained by radiative recombination. At elevated densities of the A exciton via ET from DY1 to MoS2, a strong interaction between carriers leads to rapid depopulation of the A exciton (τ1 ≈ 0.5 ps). This faster component, contributing 48% of the decay, can be attributed to the fast Auger scattering. Nonradiative relaxation via electron‐phonon and slow Auger processes are consistent with the component of τ2 ≈ 18 ps (25%).The decay kinetics for pristine MoSm Figure .[44, 32 monolayers. The hybrid heterostructures were fabricated by evaporating DY1 aggregates on CVD‐grown MoS2 flakes on quartz substrates. The PL intensity from the DY1 aggregates on the MoS2 monolayers was only 20% of that from the DY1 aggregates on the quartz substrate under 532‐nm light illumination. Sub‐nm‐thick MoS2 flakes drastically suppressed the light emission from the DY1 aggregates with thicknesses of several tens of nm. In contrast, the MoS2 flakes with evaporated DY1 exhibited 1.7‐fold enhanced PL emission compared to pristine MoS2 flakes. The negligible light‐induced CPD changes at the DY1/MoS2 interface ruled out the interlayer CT process and indicated the formation of a type‐I heterojunction. Femtosecond TA investigations enabled a detailed understanding of the kinetics of the photogenerated excitons in the DY1/MoS2 heterostructures. The TA spectra showed a rapid increase in the exciton density and a fourfold extension of the exciton lifetime in the MoS2 monolayers. All these results suggested transfer of charge‐neutral excitons from the higher bandgap DY1 aggregates to the lower bandgap MoS2 flakes. This knowledge will be instrumental in the design of high‐performance organic/inorganic hybrid devices. Specifically, our results suggest that integration of fluorescent organic materials and TMDs will provide a new strategy to explore emergent physical phenomena and propose novel optoelectronic devices.This work successfully demonstrates the transfer of photogenerated excitons from fluorescent DY1 molecules to 2D semiconducting MoS42/quartz samples were fabricated and used for most of the optical characterizations in this work. The synthesis of DY1 was previously reported. MoS2 flakes were grown on quartz substrates by the CVD technique. A bubbler containing ammonium sulfide solution was used as the sulfur precursor. DY1 fluorescent molecules were deposited on MoS2/quartz using a thermal evaporator (base pressure: 10–8 mbar) following a prior report. For CPD measurements, MoS2/DY1/quartz samples were prepared. The CVD‐grown MoS2 monolayers on SiO2/Si were transferred onto quartz substrates with evaporated DY1 by polymer‐based capillary‐force‐assisted stamp transfer. Only water was used to separate the MoS2 flakes from the quartz substrates since the organic molecules of DY1 could be dissolved in organic solvents.DY1/MoSreported.32 MoS2 atmosphere to avoid artifacts caused by ambient gas adsorption. Pt/Ir‐coated Si cantilevers were used as the tip, and the work function was calibrated against a highly ordered pyrolytic graphite reference sample. The dark states were characterized after storing the sample for more than 3 h in a light‐blocked glove box. A laser diode  was used as a light source to measure the illumination‐induced CPD changes.The surface topography and CPD maps were obtained by an AFM system  with a glove box. All the measurements were carried out in a NMicro‐PL spectra under 532‐nm‐excitation were measured by a multifunctional optical microscopy system using a 532‐nm laser with a power of 5 µW and an acquisition time of 1 s . Under 405‐nm‐excitation, confocal PL scanning images were obtained using a multiphoton CLSM  and PL spectra at specific pixels in the image were acquired with a 32‐channel spectral detector. The time‐resolved PL decay time was measured using an inverted‐type scanning confocal microscope  with a 20× objective lens (beam diameter: ≈660 nm), a picosecond 405‐nm laser source, and a hybrid photon detector. FLIM images, consisting of 512 × 512 pixels, were recorded using a galvo‐stage and the time‐correlated single‐photon counting technique. All data manipulations for the obtained fluorescence decays were performed using Leica suite software (LAS X Ver. 3.5.2).−1 and a pulse duration of ≈30 fs. The other beam was focused onto a nonlinear crystal to generate a broadband white light continuum in the visible range as a probe beam, which was used to detect changes in absorption as a function of the time delay between the pump and probe pulses. TA data were acquired using a confocal TA microscope system with a spatial resolution of ≈2 µm .For TA spectroscopy, the fundamental light from a 1‐kHz Ti:sapphire amplifier with a 790 nm center wavelength  was divided into two beams. One was sent to an optical parametric amplifier  to generate a pump beam with a certain wavelength. The pump wavelength for photocarrier excitation was tuned to 450 nm, with an energy per pulse of 0.1 nJ pulseThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


In this mechanistic study, we systematically address the redox processes related to the reverse water‐gas shift reaction (rWGSR) over In2O3 nanoparticles, both at the surface and in the bulk. Based on temperature‐dependent operando UV/Vis spectra and a novel operando impedance approach for thermal powder catalysts, we propose oxidation by CO2 as the rate‐determining step for the rWGSR. The results are consistent with redox processes, whereby hydrogen‐containing surface species are shown to exhibit a promoting effect. Our findings demonstrate that oxygen/hydrogen dynamics, in addition to surface processes, are important for the activity, which is expected to be of relevance not only for In2O3 but also for other reducible oxide catalysts.In 2 activation over In2O3 nanoparticle catalysts is consistent with redox processes, including oxidation by CO2 as the rate‐determining step, whereby hydrogen‐containing surface species are shown to exhibit a promoting effect. Besides surface processes, subsurface/bulk dynamics are shown to be essential for the activity in reverse water gas‐shift reaction, as revealed by operando UV/Vis and impedance spectroscopies.CO Such an approach would allow the increasing energy demand to be met while reducing the large amounts of CO2 emitted by burning fossil fuels.Catalysts based on cubic In2 hydrogenation over In2O3, highlighting the importance of oxygen vacancies,2 dissociation is thermodynamically and kinetically favored.2Ox3−(OH)y(111) surface has been proposed to be responsible for H2 dissociation as well as for CO2 reduction, with the latter being the rate‐limiting step.[2 adsorbate (bt‐CO2*) to a carboxylate intermediate (COOH*) is the rate‐limiting step for CO formation, which is followed by COOH* decomposition into CO and hydroxide on the surface.2O3 was investigated.Theoretical studies have addressed CO2O3, which is associated with the formation of surface indium hydrides and hydroxyl groups.2O3‐CeO2 catalysts highlight the importance of oxygen vacancies in In2O3 for CO2 conversion.2O3 has been shown to be more suitable for rWGSR than hexagonal In2O3, by enhancing the dissociative adsorption of H2, facilitating the formation of oxygen vacancies and increasing the ability to adsorb and activate CO2.2O3 has been employed as a photocatalyst for rWGSR,[On the experimental side, surface indium oxo species have been proposed to be responsible for the heterolytic dissociation of hydrogen on alumina‐supported c‐Inor rWGSR,19, 202O3 during the thermal rWGSR (≥200 °C) is still very limited, in particular on the surface chemistry and subsurface/bulk dynamics, which may be related to experimental challenges for IR and Raman spectroscopy resulting from the In2O3 absorption properties.[2O3 catalysts have been reported. Nevertheless, (sub)surface/bulk processes seem to play an important role in c‐In2O3 catalysts, as our previous studyMechanistic data on Inoperties.21, 222O3 during rWGSR. Bulk properties are directly probed by operando impedance spectroscopy supported by ex situ X‐ray diffraction (XRD) analysis. We demonstrate, on one hand, the readiness of the In2O3 surface for reduction by H2 and oxidation by CO2, and, on the other, the participation of the subsurface/bulk in the reaction, thus enhancing the understanding of the mechanism, in particular, the CO2 activation process, the role of oxygen vacancies, and the participation of hydrogen‐related adsorbates.In this mechanistic study, we employ (quasi) in situ and operando spectroscopies ) to explore the surface and subsurface dynamics of In2O3 sample used in this study have already been published.2 g−1 and our ex situ XRD results show that only cubic (Ia3‾) In2O3 is present. In addition, transmission electron microscopy (TEM) measurements show that the particles are present as sheets terminating with an In2O3(222) surface. Contaminations caused by the synthesis, e.g. nitrogen, can be excluded within the sensitivity of our XPS measurements.Details of the characterization of the In2O3 was analyzed for its activity as a function of temperature, while we simultaneously monitored the electronic structure using operando UV/Vis spectroscopy. Our previous UV/Vis studies on In2O3 during 10 vol % H2 exposure have revealed a strong dependence of the absorption in the visible range on the oxygen defect density.2 to H2 leads to an increase in absorption in the visible range and H2O formation in the gas phase, which illustrates that In2O3 undergoes reduction even at just 130 °C. Exposure to reaction conditions shows a further increase in absorbance, resulting from the longer residence time under H2 flow, since CO2 is not yet activated at this temperature, as evidenced by the absence of CO in the gas phase (not shown).First, the In2O in the gas phase. In this context, H2‐temperature programmed desorption (TPD) experiments from previous studiesIncreasing the temperature to 160 °C leads to a strong increase in absorption, which continues even after 160 °C is reached. This behavior can be explained by a strong reduction of the catalyst, as supported by the absence of CO but increasing concentration of H2 conversion and shows only a weak increase or even a decrease in visible absorption. The absorption reaches its maximum during the first measurement at 220 °C and then decreases again over the next 30 min at 220 °C. The changes in absorption behavior can be explained in terms of counteracting contributions to the indium reduction state from temperature‐dependent reduction and CO2 conversion leading to oxidation. At 220 °C the maximum degree of reduction is reached under reaction conditions but at this temperature CO2 is significantly converted, which oxidizes the surface and thus results in an overall decrease in visible absorption. With time, a stationary state is established by balancing the contributions from reduction by H2 and oxidation by CO2. At temperatures >190 °C, a similar absorption behavior is observed, consisting of an initial increase and a subsequent decrease in absorption, which clearly shows that reduction by H2 is faster than oxidation by CO2. Comparison of the different equilibrated states reveals that the absorption is at its highest at 190 °C and decreases with temperature, strongly suggesting that the oxidation of the catalyst by CO2 increases with temperature while the reduction remains at about the same level, because the maximum amount of oxygen vacancies has been reached. This could be a first indication that CO2 activation with subsequent formation of CO is the rate‐determining step. The same experiment was performed with a CO2 to H2 ratio of 4 : 2 . The results are similar, but the reduction is much slower due to the lower H2 content, which is why the further reduction of the material ends only after 250 °C. These findings underline our previous results.Starting at a temperature of 190 °C, the catalyst becomes active towards CO2O3,[3)3) and the fact that it is the most thermally stable. However, this should be viewed with caution, as previously lower flow rates were used, resulting in higher residence times, and higher reactant concentrations in the gas feed. For this reason, we will refrain from a more detailed comparison of the activities. Other factors leading to higher conversions may be the smaller crystallite size and thus the larger specific surface area of the In2O3 particles used here. Besides, the particle shape or the surface termination may have an influence on the activity, as has been shown in the context of CO oxidation,2 or CO2 activation over In2O3(111) or In2O3(110) have been demonstrated by theoretical studies (DFT).[The activities we measured compare favorably to those of previous studies on polycrystalline In2O3,9, 11 2 to start from an oxidized state. During heating to 250 °C under O2 exposure, desorption of H2O and CO2 was observed (not shown), indicating that the surface is being cleaned from adsorbates. Upon exposure to O2 flow, the absorption at 532 nm increased significantly surface and possibly also to an increase in reflectivity through small metallic indium domains. This is supported by the strong presence of H2O in the gas phase. Since the same trends in absorption can be observed at both wavelengths (and the entire remaining visible range), only the absorption at 700 nm will be discussed in the following.On switching to H2, there is again a strong decrease in absorption, which is in the range of the absorption in O2, demonstrating that the sample can be oxidized again by CO2 and that the process of reduction seems to be completely reversible on the basis of UV/Vis spectra and their penetration depth. At the same time, the gas phase initially shows a small increase in CO  induces a decrease in absorption despite the same H2 concentration, which, however, does not reach the level detected in CO2. This behavior confirms the findings of Figure 2 to CO2 or vice versa. In this context, previous photocatalytic studies have shown that the surface hydroxide concentration is also higher under reaction conditions than under pure H2.2 after reaction conditions leads to a renewed increase in absorption and thus a reduction of the surface.Exposure to reaction conditions H/CO2 indu2O3 is exposed to O2 flow to check the reversibility of the system. The gas‐phase data reveals a strong H2O signal decaying with time, which implies the presence/formation of hydroxides/adsorbed hydrogen during H2 exposure. During this O2 exposure, the absorbance is lower than during the first O2 phase, and major differences were detected for 532 nm absorption, which is located in the absorption edge. Since in both cases the experimental conditions were the same, we attribute the observed drop in absorption to a decrease in the band gap resulting from a larger crystallite size. To test this hypothesis, we performed ex situ XRD measurements after different gas exposures until the first reaction phase, which clearly show that the crystallite size increases from 12 nm after the first O2 phase to 27 nm after the subsequent H2 . Note that all following gas atmospheres have no significant influence on the crystallite size. Notably, after the second H2 exposure, additional reflections are detected, which originate from metallic indium (JCPDS 85‐1409), showing that metallic indium is formed during reduction with 4 vol % H2 at 250 °C. To this end, it should be noted that small domains of metallic indium, formed during the first H2 phase and under reaction conditions, may have oxidized back under air exposure and thus cannot be excluded by the present ex situ XRD analysis.Finally, In2O3, we recorded the XRD pattern after CO exposure (4 vol %) after a prior oxygen treatment, analogous to the other measurements. For CO exposure, significantly more metallic indium was formed and the crystallite size grew only to 21 nm, in contrast to the observed 27 nm for H2. Thus, hydrogen is a weaker reducing agent towards In2O3 than CO, which is consistent with the literature.2 treatment . Interestingly, the sample absorbance almost returns to its original value (measured in O2) on subsequent exposure to CO2. Thus, In2O3 can be re‐oxidized by CO2, releasing CO regardless of the reducing agent (H2 or CO). Moreover, no surface hydrogen species are required for the oxidation by CO2, indicating that no hydrogen‐containing intermediates are involved in the re‐oxidation process.To further probe the reducibility behavior of In2O3 nanoparticles, operando potentiostatic electrochemical impedance spectroscopy (p‐EIS) spectra were acquired. It should be mentioned that due to the sample shape (no pellet) a Mott–Schottky analysis is disturbed by gas inclusion and an increased number of grain boundaries and is not completely reproducible. For this reason, such an analysis was not performed in this work. The non‐uniform sample distribution also has an impact on the p‐EIS spectra, which deviate from ideal behavior by ca. 10 %, as validated by Kramers‐Kronig‐relations (see Supporting Information). All spectra .To elucidate the charge and mass transport in In2O3 is a semiconductor with a band gap of approximately 3 eV,2O3 during the oxygen treatment. The spectrum in O2 atmosphere may therefore be treated as a response of bare In2O3 nanoparticles and will serve as a reference for the discussion of changes observed during the following gas treatments.As In2 at room temperature.2O3 surface. The voltage applied to the sample causes external polarization due to the non‐conductive character of In2O3 at room temperature, and the activation energy from the temperature dependent conductivity behavior of 137 kJ mol−1 will not be reached by simple polarization but requires elevated temperatures.Regarding surface charging, a similar behavior was previously observed for CeO2O3 was observed during the hydrogen treatment . The pronounced contribution of the inductive part (L=2.03×10−5 H during the first H2 treatment) observed in the operando p‐EIS spectrum is attributed to the percolation of In2O3 nanoparticles with hydrogen species,2O3, which have demonstrated the diffusion of interstitial hydrogen to increase conductivity.[2O3 nanoparticles is proposed to proceed according to the reaction In2O3+3 H2→2 In0+3 H2O. Thus, the first two H2 phases show a purely resistive behavior at the beginning of the Nyquist plot, indicating the formation of metallic indium. In the third phase, this behavior is less pronounced, but this is also consistent with our UV/Vis results (lowest absorption at 532 nm).The major reduction of Inuctivity.24, 252 into protons within In2O3 was speculated to occur by a homolytic or a heterolytic splitting process, as discussed by García‐Melchor and López.The decomposition of H2O3 was exposed to CO2 and an operando p‐EIS spectrum was recorded , which was found to be in the ohm range (see Table S2), speaking for improved conductivity compared to the pristine material (with R1 in the megohm range); nevertheless, remaining vacancies induced by the H2 phase remain active, judging by our Raman spectra below. In the second part, the reorganization of vacancies occurs.2O3, while at low frequencies (L1R5 region) gas diffusion takes place, as, for example, observed for polycrystalline tin oxide.[Following the hydrogen treatment, In2O3, carries the CO2 gas with it, as judged by the inductance contribution (L=1.03×10−5 H), which differs from the value obtained for the first H2 treatment (L=2.03×10−5 H). The reaction occurs in the bulk due to gas percolation, rather than on the surface as in the case of pure CO2 exposure. In this context, we propose that the higher inductance is caused by the reaction CO2+H2→CO+H2O and thus by the presence of CO, which is also detected in the gas phase. Therefore, the respective parts of the triple equivalent circuits of Figure 2O3 with H2 (right part of the circuit) and CO2 (left part of the circuit). Please note that the EIS spectrum is reversed compared to what it was in the case of H2 treatment, implying that the high frequency points are in the region of high imaginary and real impedances, originating from gas percolation and the induction of the gas impedance.2 atmosphere.During reaction conditions , ν(InO6), In−O−In, and ν(InO6) vibrations of cubic In2O3, respectively, which is in accordance with the literature.[2 and H2/CO2 treatment, Figure −1 also originates from cubic In2O3 and is attributed to a vibration with Eg symmetry.2, as a result of defect formation, consistent with our previous study.−1) are related to oxygen defects, as confirmed by DFT.−1 shows a shoulder towards larger wavenumbers, which is caused by reduction and was also observed in previous studies.[Figure terature.16, 412 and reaction conditions, which is caused by fluorescence, self‐absorption of the laser, and an increased reflectivity by metallic indium . Nevertheless, closer inspection of the spectra recorded after O2 and CO2 exposure shows that the 415 cm−1 band is less pronounced for O2, implying that re‐oxidation by CO2 is not as efficient as by oxygen. This behavior is in agreement with previous impedance studies,2O3 is best re‐oxidized with O2, followed by H2O and CO2. In the UV/Vis measurements such a difference was not observed due to the lower penetration depth, implying that the surface is completely oxidized, but the bulk is not. Thus, a certain number of defects remain present even under O2.Comparison of the different gas phases reveals a significantly higher background for H−1 , which originate from bridging hydroxyl groups[−1, which is assigned to the C−H stretching vibration of a formate‐like species.The high‐wavenumber region of the Raman spectra shows bands at 3647 and 3675 cmyl groups44, 452O3 using He I radiation after exposure to the different gas atmospheres, analogous to the XP spectra. All UP spectra show the typical valence band features at 4.4, 6.1 and 8.7 eV.[2 exposure, but most prominently for exposure to CO2 and reaction conditions. Previous studies have proposed the contributions at 5.7 and 10.5 eV to be determined by modified valence band (VB) states subjected to adsorption of molecules.2 spectrum, but the presence of carbonaceous adsorbates appears to be more likely, considering the increased intensity after exposure to CO2 and the reaction phase. This is further supported by the IR gas‐phase measurements, which show the presence of hydrogen‐containing adsorbates on the surface after H2 treatment  was applied (for XPS data and discussion see Supporting Information). Figure d 8.7 eV.46, 472) and reductive (H2) atmospheres, while the overall changes in the adsorbate species  are rather distinct, somewhat resembling the observed subsurface/bulk behavior  newly developed operando impedance spectroscopy, as well as quasi in situ spectroscopies . The new mechanistic insight into the In2O3 mode of operation was supported by a systematic investigation of the related reduction and oxidation processes and their comparison to reaction conditions. For example, our study shows that despite surface reduction, diffusion of oxygen vacancies leads to bulk reduction. While reduced c‐In2O3 nanoparticles can be re‐oxidized by CO2 on the surface, complete oxidation of the bulk is only possible with O2. This behavior is independent of the reducing agent (H2 or CO), although CO leads to stronger In2O3 reduction.In this study, we have generated a deeper understanding of the mechanism of rWGSR over c‐In2 is identified as the rate‐determining step for the rWGSR, which is supported by operando impedance spectra, showing a predominance of the hydrogen contribution. The significant evolution of CO and H2O, also after separate treatment with CO2 or H2, is fully consistent with redox processes but because conversions are somewhat higher under reaction conditions, hydrogen‐containing surface species  are proposed to have a promoting effect on CO2 conversion and/or hydrogen containing intermediates (e.g. COOH*) may play an role. Such a behavior would be in accordance with results from theoretical studies, which associate CO2 conversion with frustrated Lewis pairs.[Based on temperature‐dependent UV/Vis analysis under reaction conditions, oxidation by COis pairs.14, 152O3 nanoparticle is involved in the reaction, which is only accessible by using complementary in situ/operando techniques. Combining our new experimental findings with results from theory allows us to develop a consistent mechanistic picture of rWGSR over c‐In2O3. Our findings are of immediate relevance for a detailed understanding of CO2 hydrogenation and other related catalytic processes over In2O3, but they are also expected to be of more general impact, considering the growing importance of catalysts based on reducible oxide materials.The above analysis of the (sub)surface and bulk properties during the rWGSR demonstrates that the entire InThe authors declare no conflict of interest.1As 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.Supporting InformationClick here for additional data file.


The THz scanning near‐field optical microscopy experimentally reveals the strong light–matter interaction in the THz band of EI state in the Ta2NiSe5. Benefiting from the strong light–matter interaction, the Ta2NiSe5‐based photodetectors exhibit superior THz detection performances with a detection sensitivity of ≈42 pW Hz−1/2 and a response time of ≈1.1 µs at 0.1 THz at room temperature. This study provides a new avenue for realizing novel high‐performance THz photodetectors by exploiting the emerging EI materials.Despite the interest toward the terahertz (THz) rapidly increasing, the high‐efficient detection of THz photon is not widely available due to the low photoelectric conversion efficiency at this low‐energy photon regime. Excitonic insulator (EI) states in emerging materials with anomalous optical transitions and renormalized valence band dispersions render their nontrivial photoresponse, which offers the prospect of harnessing the novel EI properties for the THz detection. Here, an EI‐based photodetector is developed for efficient photoelectric conversion in the THz band. High‐quality EI material Ta 2NiSe5 are reported. The THz scanning near‐field optical microscopy experimentally reveals the strong light‐matter interaction in the THz band of EI state in Ta2NiSe5. Benefitting from the strong interaction, the photodetectors exhibit superior THz detection performances with an ultra‐high sensitivity and fast response speed at room temperature.High‐performance terahertz (THz) photodetectors based on excitonic insulator (EI) material Ta The main challenge and technical bottleneck of THz photodetection is the large mismatch between the nanoscale active region and the extremely long wavelength of the incident THz wave, resulting in low photoelectric conversion efficiency. In order to break through this technical bottleneck in THz detection, it is crucial to develop novel materials with strong light–matter interactions for efficient photoelectric conversion of low‐energy photons.Nowadays, high‐efficient photodetection in the long‐wave infrared especially in far infrared up to terahertz (THz) range (30 µm to 3 mm) at room temperature, is particularly rare but becomes more and more important. THz wave is located in the cross transition region of microwave electronics and infrared photonics, which promises it a wide range of applications, including security check, biomedical, wireless communication, and remote sensing.6, 7 While, excitonic insulator (EI) states as the novel quantum states with anomalous optical transitions and renormalized valence band dispersion has gained rediscovery due to its unique band structure and accompanied photoelectric characteristics, which can be realized in narrow band gap semiconductors or small band overlapped semimetals at finite temperatures. Specifically, when the exciton binding energy (Eb) exceeds the band gap (Eg), the excitons spontaneously generate and condense into a completely new ground state due to the weakly screened Coulomb interaction between electrons and holes. However, due to the difficulty in finding suitable compound materials with an EI state, the EI state has been considered as an unconfirmed quantum state. In recent years, the EI state has been experimentally confirmed in a handful of strongly correlated systems , which provide a unique platform for the study of quantum many‐body physics and optoelectronic devices due to their exotic properties.Quantum states in emerging materials, such as topological states, superconductivity, and unconventional excitonic states, have inspired innovations in the fields of electronic and optoelectronic devices.22 Whi Ta2NiSe5, a ternary layered TMC compound with a narrow direct energy gap structure has been identified as a stable EI material. Over the past few years, the EI state of Ta2NiSe5 at room temperature has been widely studied and characterized. The electrical transport tests on bulk and few‐layer samples have confirmed that there is a transition from semiconductor to EI state in Ta2NiSe5 at phase transition temperature Tc ≈ 326 K. At the same time, the crystal structure of Ta2NiSe5 undergoes an orthorhombic to monoclinic distortion across in the second‐order phase transition. Angle‐resolved photoelectron spectroscopy (ARPES) experiments showed that the valence band edges are unusually flat while the energy band gap enlarges as the temperature decreases below the Tc. In addition, giant Fano resonance and pump‐probe time‐resolved measurements further revealed strong band hybridization and interband interactions. The special EI state of Ta2NiSe5 provides an unprecedented opportunity to study the role of electronic states in the exciton formation on electronic and optoelectronic devices. However, the current research on Ta2NiSe5 mostly focuses on its physical properties, such as how to demonstrate and modulate the EI state, its applications including photodetections are rarely researched. The expected strong light–matter interaction in the EI material Ta2NiSe5 is still elusive and the harnessing of the EI properties in the THz photodetection is far unexplored.Recently, 2D transition metal chalcogenide compounds (TMC) with narrow bandgap or semi‐metallicity have also shown promise as hosts of the novel EI ground state.36 Ta22NiSe5 and utilize the novel EI state to achieve sensitive THz detection with highly efficient low‐energy photon harvesting. First, the high‐quality Ta2NiSe5 crystals are synthesized by the chemical vapor transport (CVT) method, and the existence of the EI state at room temperature and the second‐order phase transition at Tc ≈ 328 K are confirmed by the combination of various characterization methods. Furthermore, we reveal that the EI state of Ta2NiSe5 has strong light–matter interaction in the THz band through near‐field THz microscopy. Benefiting from the strong light–matter interaction at the THz band, the EI‐based photodetectors achieve a high sensitivity of ≈42 pW Hz−1/2 and a fast response speed of ≈1.1 µs at room temperature. This work opens up the possibility to explore new fundamental quantum states for the future applications in THz optoelectronic devices.In this work, we demonstrate a photodetector based on the EI material Ta22NiSe5 is a layered structure stacked by a weak van der Waals force along the b‐axis, which possesses a monoclinic atomic structure at room temperature. As shown in Figurec axis. The top view of the crystallographic structure of Ta2NiSe5 is depicted in Figure 6]2 dimer chains and NiSe4 single chains are arranged periodically along the c‐axis and run along the a‐axis to form the quasi‐one‐dimensional structure. To investigate the crystallographic structure of our synthesized Ta2NiSe5, X‐ray diffraction (XRD) pattern was performed, as displayed in Figure c (PDF#78‐0279), which is consistent with the EI phase structure. Meanwhile, these strong peaks illustrate the high‐purity phase and highly preferred orientation along the (0l0) direction of Ta2NiSe5 single crystal. Furthermore, the scanning electron microscopy (SEM) with energy dispersive X‐ray spectroscopy (EDX) was adopted to characterize the surface morphology and elemental composition of Ta2NiSe5, as shown in Figure 2NiSe5 crystal. As depicted in Figure 7/2 and Ta 4f5/2 of the Ta–Se bonding. Figure 3/2 and Ni 2p1/2 of the Ni–Se bonding. In addition, two peaks at 54.1 eV and 54.9 eV of Se 3d spectrum displayed in Figure 2− of Ta2NiSe5. Furthermore, the atomic ratio Ta, Ni, and Se was calculated to be around 1.95:1:5.05 from the XPS spectrum, which is also close to the ideal composition of Ta2NiSe5.The ternary compound Ta2NiSe5. Figure 2NiSe5 nanosheets on the gold substrate. The Ta2NiSe5 shows an absorption edge of ≈1337 cm−1, corresponding to a narrow band gap of ≈0.165 eV, which is consistent with the previous reports. The infrared absorption spectrum is a reproducible result for the other sample  images of the Ta2NiSe5 nanoflake are depicted in Figure g vibration mode, which matches well with the previous reports. Among them, two strong Raman peaks at 97 cm−1 and 122 cm−1 are the characteristic peaks of monoclinic structure (EI state) without obvious shift in all samples  spectrometer was deployed to investigate the optical band gap of Ta reports.28 The2NiSe5 were examined by the spherical‐aberration corrected scanning transmission electron microscopy (STEM). Figure2NiSe5 crystal. In addition, the cross‐section STEM image , (001), and (010) planes of Ta2NiSe5 crystal, respectively. All of the interplanar spacings can be well matched to the monoclinic crystal structure of Ta2NiSe5, which further illustrates the existence of EI state at room temperature. The periodic arrangement of Ta, Ni, and Se atoms can be seen from the atomic intensity line profile shown in Figure 2NiSe5 nanoflake, which shows that the atomic ratio of as‐prepared Ta2NiSe5 is close to 2:1:5. The low‐magnification high‐angle annular dark‐field (HAADF) STEM image and element mapping images are depicted in Figure 2NiSe5. In the meanwhile, the uniform element distribution is also shown in the cross‐section element mapping images of Ta2NiSe5, as depicted in Figure 2NiSe5 crystals at room temperature, which plays an important role in the subsequent THz photodetections.To further evaluate the crystalline quality, atomic structure and chemical composition of our synthesized Tae Figure  shows a 2NiSe5‐based field‐effect transistors (FET) and Hall bar devices were fabricated on the degenerately p‐type doped silicon (Si) substrate  to study the electrical properties of Ta2NiSe5. Figure 2NiSe5 FET, which displays a weak gate‐control ability and a good ohmic contact between Ta2NiSe5 nanosheet and the metal electrodes. As displayed in Figure 2NiSe5 FET exhibit a typical n‐type semiconductor characteristic. The temperature dependence of the longitudinal resistance RXX was performed from 50 to 200 K at B = 0 T to understand the transport properties of the Ta2NiSe5, as shown in Figure RXX increases with the decreasing temperature, exhibiting typical semiconductor behavior. In the meanwhile, Figure R as a function of temperature T from 310 to 342 K. The resistance R displays an anomaly at T ≈ 327 K, corresponding to the second‐order phase transition, which illustrates the existence of an EI state in Ta2NiSe5 below the characteristic temperature Tc ≈ 327 K. Below Tc, the rapid increase in resistance is related to the opening of band gap under the EI state. These results are consistent with the previous reports.Subsequently, the Ta2NiSe5, it may bring about strong light–matter interaction. The lattice structure and band structure of the EI state in Ta2NiSe5 were depicted in Figure2NiSe5 crystal exhibits a standard orthogonal crystal structure and a narrow gap semiconductor band structure. During the formation of EI states below Tc, the lattice distortion of the Ta/Ni chain leads to the structure transition from orthorhombic to monoclinic, resulting in a symmetry breaking, which benefits for enhancing the light–matter interaction in THz band. In addition, a large number of excitons condensation between the valence and conduction bands results in the opening of the band gap and the flattening of the valence band. Meanwhile, there is a strong interaction between excitons and phonons in the special EI state due to the coherent phonon excitation, which has a resonant absorption in the THz band. Furthermore, the THz time‐domain spectroscopy (TDS)‐based scattering‐type scanning near‐field optical microscope (THz s‐SNOM) system was adopted to demonstrate the interaction of the Ta2NiSe5 sample with incident THz photons. Figure 2NiSe5 on Si/SiO2 substrate are displayed in Figure 2 substrate, the Ta2NiSe5 nanoflake has an obvious and stronger THz field distribution in all mapping images. As depicted in Figure 2NiSe5 nanoflake (marked as 4–6) and the Si/SiO2 substrate (marked as 1–2) over the THz frequency range. These results illustrate the strong light–matter interaction in the THz band of EI state in the Ta2NiSe5 at room temperature, which can couple more THz radiation with high efficient harvesting of the low‐energy photons, and produce an excellent THz response.Considering the spontaneous crystal structure symmetry breaking and exotic valence band flattening of the EI state in Ta2NiSe5, the THz photodetector was designed as a sub‐wavelength planer metal‐Ta2NiSe5‐metal (MTM) structure with a symmetric bow‐tie antenna, which can convert the incident THz radiation into the localized oscillating electric field  was calculated from the R–T curve . As shown in Figure 2NiSe5 photodetector at different incident radiation  are also depicted in Figure Iph of the photodetector grows approximately linearly with increasing bias voltage, which might be driven by the photoconductive effect. Figure I–V curve of the photodetector under pulse modulated radiation at 0.03 THz, which also shows the same growing trends. It should be noted that the Iph at 0.3 THz is small due to the small incident power. Figure Iph versus incident radiation power at different bias voltages, indicating the large dynamic range and second‐order nonlinear behavior of the photodetector. The response/recovery time is defined as the time for the Iph increases/decreases from 10%/90% of the maximum Iph to 90%/10% at single impulse, respectively, which is recorded by a fast oscilloscope. Figure Next, the room‐temperature THz response performances of the TaRi) and noise‐equivalent power (NEP) of the photodetector at 0.1 and 0.3 THz were calculated from the bias‐dependent Iph to further evaluate the THz photoresponse performances, as shown in Figure Ri is calculated by the relation Ri = Iph/P . It can be seen that Ri grows with increasing bias voltage for all incident THz radiation. When the bias voltage reaches 1.5 V, the values of Ri are ≈0.36 A W−1 and ≈0.036 A W−1 at 0.1 and 0.3 THz, respectively. As another important parameter, NEP, can be obtained from the formula NEP = in/Ri, where in is the noise current density for the photodetector. In general, the dominant noise in our test system contains the Johnson−Nyquist noise (vt), shot noise (vb), and flick noise (1/f noise). Among them, the 1/f noise is related to the change of electronic state, which usually occurs at low modulation frequency (below 1 kHz). During our test, the modulation frequency is over 1 kHz, so it can be ignored in our system. Therefore, the main noise sources are vt and vb in our system, which originate from the thermal motion of carriers due to the nonzero ohmic resistance and the generation of carriers under bias voltage, respectively. Inherently, the total noise current density can be obtained from the electrical characteristic of the detector by the formula in = (vt2 + vb2)1/2/r = (4kBT/r + 2qIDS), where kB is Boltzmann constant, T is the temperature, r is the resistance of the detector, q is elementary charge, and IDS is the dark current of the detector. Figure −1/2 at 0.1 THz and ≈417 pW Hz−1/2 at 0.3 THz, respectively. To better evaluate the performance of the Ta2NiSe5 photodetectors, a comparative survey of recently reported room‐temperature THz detectors based on different 2D materials was carried out, as displayed in Table2NiSe5 photodetector has obvious advantages in sensitivity and response speed.The photocurrent responsivity Ri and noi32NiSe5 single crystals and exploited their novel EI state for THz detection. The existence of the EI state at room temperature and the second‐order phase transition at Tc ≈ 328 K were confirmed by XRD, Raman spectroscopy, ARPES, STEM, and electrical transport tests. The EI state with spontaneous crystal structure symmetry breaking and exotic valence band flattening exhibits strong light–matter interaction in the THz band, which was underpinned by strong THz field distribution on the Ta2NiSe5 surface illustrated by the THz s‐SNOM. Based on the local electric field of the antenna‐integrated THz photodetector enhancement in MTM structure and the strong interaction of incident THz radiation and EI state, the Ta2NiSe5‐based photodetectors display a wide‐band (0.03–0.30 THz) response with a responsivity of ≈0.36 A W−1, an NEP of ≈42 pW Hz−1/2, and a response time of ≈1.1 µs at room temperature. Considering the high sensitivity and fast response speed of our Ta2NiSe5 photodetectors, this work lays the groundwork for future research into the EIs for THz photodetections.In summary, we synthesized high‐quality Ta42NiSe5 bulk crystals were obtained by the CVT method, the same as that of the authors’ previously reported work. The XRD (Bruker D8 Advance) was adopted to characterize the crystal structure of the synthesized Ta2NiSe5. The infrared absorption spectroscopy was taken from FTIR spectrometer (Bruker Vertex 70) equipped with a microscope system. A micro‐Raman system (LABRAM HR) with a 532 nm laser was performed to obtain the Raman spectrum. The morphology and thickness of the nanoflakes were characterized by optical microscope (Nikon Eclipse LV100ND) and AFM . The atomic structure, chemical composition, and elemental mapping of Ta2NiSe5 were characterized by the STEM (FEI Titan Themis Z), SEM (Quanta FEG 250), and XPS (PHI 5000 VersaProbeIII). The energy band structure of Ta2NiSe5 single crystal was measured on the ARPES system (DA30L). The incident photon energy, energy resolution, and angular resolution were ≈21.2 eV, ≈1.8 meV, and ≈0.1°, respectively.The high‐quality Tated work.54 The2NiSe5 nanoflakes on the Si/SiO2 substrate were obtained from as‐synthesized bulk crystals by the mechanical exfoliation method. It should be noted that the substrate used in the THz detectors was high‐resistance (ρ ≈ 20 000 Ω cm) intrinsic silicon. Subsequently, the 2‐terminal configuration, planer bow‐tie, and Hall bar electrode structures were determined by electron‐beam lithography . Then, Cr/Au (10/70 nm) films were deposited by an electron‐beam evaporator  system, followed by a lift‐off process to form the THz photodetectors and Hall devices. The structure of THz detector with a bow‐tie antenna was characterized by the SEM. The electrical transport properties of Ta2NiSe5 devices were measured in a semiconductor parameter analyzer (Keithley 4200) and a physical property measurement System .The TaIph) signal was collected at the drain electrode by means of lock‐in amplifier technique. A microwave signal source  with different VDI frequency multiplier was selected to the THz emission source, which could tune up to 0.07–0.12 THz (WR 9 Tripler) and 0.24–0.30 THz (WR 2.8 Tripler). The THz incident radiation power intensity of 0.1 and 0.3 THz measured by the TK100 power meter were 8.6 and 10 µW mm−2, respectively. The photocurrent responsivity Ri of the detector could be retrieved from the formula Ri = Iph/Pin×Sd, where Pin and Sd is the incident radiation power intensity and the effective area of detector, respectively. The whole area of device Sd (Sd = 200 µm × 140 µm = 28 000 µm2 = 2.8 × 10−2 mm2) was much smaller than the diffraction‐limited area λS = λ2/4 . So Sd was generally taken as λS in the previous reports.In this work, the photocurrent mode was adopted to measure the THz response measurements, in which the source electrode was grounded and the photocurrent (The authors declare no conflict of interest.Supporting InformationClick here for additional data file.


Their structural and electrochemical characteristics were investigated by SEM, XRD, XPS, charge–discharge tests and EIS. XPS data indicate that the substitution of Al3+ for Zn2+ causes manganese to be in a mixed valence state by a charge compensation mechanism. Moreover, the presence of this charge compensation significantly improves the electrochemical performance of Znx1−AlxMn2O4, such as increasing the initial coulombic efficiency, stabilizing the cycleability as well as improving the rate capability. The sample with 2% Al doping shows the best performance, with a first cycle coulombic efficiency of 69.6% and a reversible capacity of 597.7 mA h g−1 after 100 cycles. Even at the high current density of 1600 mA g−1, it still retained a capacity of 558 mA h g−1.ZnMn 2O4. This is a new approach to improve the electrochemical performance of ZnMn2O4 through a charge compensation mechanism using the substitution of Al3+ for Zn2+.This work reports the nonequivalent substitution of ZnMn Thus, the demand for next-generation LIBs with higher capacity has stimulated efforts to develop new materials6,7 such as graphene, carbon nanotubes, silicon, tin oxide, etc.Recently, rechargeable lithium ion batteries (LIBs), as energy storage devices, have been wildly applied in hybrid electric vehicles (HEVs), pure electric vehicles (EVs) and plug-in hybrid vehicles (PHEVs). However, in these cases, consumers want the vehicles to have a long cruising mileage and good safety, which requires LIBs with enhanced performance in terms of their energy density and safety.8 The first is the intercalation/de-intercalation type, including many carbonaceous materials like hard carbons,9–11 carbon nanotubes12–15 and graphene,16 and some titanium oxides, such as TiO2  were prepared by a typical spray drying method using a citrate ligand to complex the metal ions. For this purpose, Zn(Ac)2·H2O , Al(Ac)3·9H2O , Mn(Ac)2·4H2O (0.1 mol) and citric acid (0.05 mol) were dissolved in 125 ml of deionized water with vigorous agitation. The mixed solution was spray-dried as the outlet temperature was set to 185 °C. Then the obtained white acetate powder was pre-sintered at 300 °C in air for 3 hours to decompose the organic constituents. After grinding into powder, the precursor was sintered at 600 °C in air for 20 hours to obtain target products.ZnMnθ range of 10–90°. The morphologies and sizes of the samples were directly examined by scan electron microscopy (SEM). The electronic states of Mn in the samples were determined X-ray photoelectron spectroscopy .The phase composition and structure of the sample were performed using an X-ray diffractometer equipped with Cu Kα radiation  over the 2Φ 15 × 1 mm) was used as a negative electrode (counter electrode and reference electrode). A Celgard 2400 porous polypropylene film was served as a separator. The electrolyte was 1 M LiPF6 dissolved in a compound of diethyl carbonate/ethylene carbonate . The galvanostatic charge–discharge cycling and cyclic voltammetry were tested at room temperature (RT ∼ 25 °C) by means of a computer-controlled battery evaluation system . And the electrochemical impedance spectroscopy (EIS) measurements were performed via PARSTAT 2273 electrochemical workstation system  over frequency range of 100 kHz to 10 mHz with amplitude of 10 mV.CR2032 coin-tape cell was used to investigate the electrochemical performances of the synthesized materials. The active material was mixed in slurry containing 10 wt% of Super P as a conductive agent and 10 wt% NaCMC as a binder. The homogeneous slurries were cast onto copper foils to obtain electrode laminate which were dried at 110 °C in the vacuum drying oven all over the night. The working cell was assembled in a glove box filled with dry argon. A lithium disk . No additional peaks are observed, indicating that the samples are pure. The lattice parameters calculated by JADE are shown in x1−AlxMn2O4a/b are gradually decreased from 5.722 Å to 5.6965 Å when x is increased from 0 to 10%. In the case of lattice parameter c, it is initially increased from 9.236 Å to 9.2793 Å when x is increased from 0 to 2%, then decreased to 5.6965 Å when x is increased 10%. This variation of lattice parameters should be ascribed to the substitution of Al3+ to Zn2+ since the radius of the Al3+ is smaller than that of the Zn2+ . These results also indicate that Al3+ has been successfully introduced into the ZnMn2O4 lattice.XRD patterns of the Zn2O4 and Znx1−AlxMn2O4 samples. All samples have an aggregated morphology of some irregular primary particles whose size is around 150 nm. There is no change in terms of the particle size and shape after the substitution of Al3+ for Zn2+. The EDX images of the ZAMO2 are shown in 2O4 lattice.x1−AlxMn2O4 electrode are provided in −1 in the voltage range of 0.01–3.0 V versus Li/Li+. In the case of the ZnMn2O4, there are three cathodic peaks and two anodic peaks in the first cycle. According to the discussion in literatures, the first peak located at 1.2 V is attributed to the reduction of Mn3+ to Mn2+.36 The second one located at 0.8 V is ascribed to the electrolyte decomposition along with the formation of solid-electrolyte interface (SEI) on the surface of the electrode. Both of those two peaks disappear in the following cycles. The sharp and intense peak at 0.13 V is attributed to reduction of Mn2+ and Zn2+ to metallic Mn and Zn nanoparticles, respectively.37 Simultaneously, the Li–Zn alloy is also formed at this low potential. There are two peaks in the first charge process, one is at 1.3 V, and the other is at 1.5 V, corresponding to the oxidation of metallic Mn and Zn nanoparticles to MnO and ZnO.38 In the second and the third cycles, there is only one peak located at 0.5 V in the discharge process. They are resulted from the reduction of MnO and ZnO to Mn and Zn. The similarity of the subsequent CV curves indicates the high electrochemical reversibility of the sample. The cyclic voltammograms of Znx1−AlxMn2O4 electrodes are shown in 2O4, implying that the Al3+ doping has no effect on the charge–discharge mechanism of ZnMn2O4.The cyclic voltammograms of Zn2O4 and Znx1−AlxMn2O4 are shown in 2O4 and Znx1−AlxMn2O4 are all at about 0.4 V. The specific capacities of the Znx1−AlxMn2O4 between 0.4 V and 1.2 V slightly decrease after the Al3+ doping. It is probably related to the decrease in Mn3+ concentration in Znx1−AlxMn2O4. Thus the total discharge capacities are decreased. The initial discharge capacities of five samples are much higher than the theoretical capacity of 1024 mA h g−1 corresponding to the reaction of ZnMn2O4 + 9Li+ + 9e− → ZnLi + 2Mn + 4Li2O. The extra capacity could result from the SEI layer on the surface of the electrode.39–42 The coulombic efficiency (CE) is an important parameter of LIBs. It reflects the degree of the Li+ going back to the lattice. The charge retention capacity increases along with CE increasing. The CEs of ZnMn2O4 and Znx1−AlxMn2O4 are about 70% . This value is higher than CEs of other ZnMn2O4-based anodes reported in literature.43–46The first discharge–charge curves of ZnMn2O4 and Znx1−AlxMn2O4 are depicted in −1 in the voltage range of 0.01–3.0 V in room temperature, using Li metal as an anode. The specific discharge capacity of ZnMn2O4 is gradually decreased till the 40th cycle and then increases till the 120th cycle. This abnormal increase in the discharge capacity between the 40th and 120th cycle have been ascribed to two factors. One is the formation of a polymer organic layer on the electrode during the cycling which can reserve Li reversibly.47,48 The other is the particles cracked during the cycling so that the contact area between the electrode and electrolyte solution increases.49 This phenomenon is not always good. When the nanostructure completely crumbled, the active materials will be separated from the conductive agent. This is bad for charge and ion transmission. As expected, following the rise, there is a quick decline in the capacity.The cyclic performances of ZnMnx1−AlxMn2O4 electrodes are shown in −1.The discharge/charge curves of Znx1−AlxMn2O4 exhibits a much higher capacity at a current density of 1600 mA g−1. The capacities of ZAMO2 and ZMO are 557.8 mA h g−1 and 379.8 mA h g−1. When the current density goes back to 100 mA g−1, the capacity of ZAMO2 is also higher than ZMO. It indicates that ZAMO2 is more appropriate for charge/discharge at high current density.The rate capabilities of five samples are shown in In order to clarify the mechanism of this improved performances after the introduction of Al, the X-ray Photoelectron Spectroscopy (XPS) and Electrochemical Impedance Spectroscopy (EIS) were measured and the results were provided in 2p, Mn3s and O1s. The binding energy of Mn3/22p, multiplet splitting of the Mn3s level, the binding energy of O1s and the atomic ratio of Mn3+/Mn2+ are list in 3/22p for the samples of Znx1−AlxMn2O4 slightly shifts from 641.1 eV to 641.2 eV, while the binding energy of O1s shifts from 529.8 eV to 529.9 eV, when x increases from 0 to 0.1. These results implied that the substitution of Al for Zn increases the interaction between Mn and O. As to the ΔE3s of Mn, it increases from 4.5 eV to 5.1 eV when x is from 0 to 2%. According to previous research,50 the higher ΔE3s value indicated a lower valence of Mn. It is well known that Zn in ZnMn2O4 exhibits +2 and Mn is +3, after the substitution of Al3+ for Zn2+, the valence of Mn should be lower so that the sample displays electric neutrality. Furthermore, the atomic ratio of Mn3+/Mn2+ calculated through XPS analysis decreases from 6.5 to 4.5 when x is from 0 to 2%. Those results prove that some Mn3+ in Znx1−AlxMn2O4 is partly transformed to Mn2+ in order to balance the Al3+ doping.x1−AlxMn2O4 before cycling. The depressed semicircles in the high frequency region are related to the ohmic resistance (Re) and the charge transfer resistance (Rct) through the electrode/electrolyte interface.51–53 Low resistance is good for the charge and ion diffusion. From the figure, we can easily find that ZAMO2 has the least diameter of the semicircles.x1−AlxMn2O4 after 10 cycles. The curves are almost lines with a slope of 45 degrees. This indicates that the batteries are controlled by the diffusion step of Li+ in solid phase.According to the result of XPS and EIS, we believed that the introduction of Al makes Mn into a mixed valence state which can shorten the band gap width of manganese oxide and result in materials with higher conductivity. SEM shows that the particle size and morphology have no change after the doping, so that this improvement is all result from the charge compensation after the nonequivalent substitution.2O4 was successfully realized by spray drying process following with annealing treatment. Results of XRD, EDX and XPS indicate that the materials are pure and the aluminum ions are successfully introduced into the ZnMn2O4 lattice. Meanwhile, the valance of Mn in Znx1−AlxMn2O4 is in a mixed state because of charge compensation. ZAMO2 exhibits excellent stability, which maintains a high reversible of 597.7 mA h g−1 at the current density of 100 mA g−1 after 100 cycles. Even at high current density of 1600 mA g−1, the reversible capacity is still kept at 557.8 mA h g−1, which is higher than the capacity of theoretical commercial graphite. The electrochemical results demonstrated that the nonequivalent-substitution is a successful try for development of advanced anode material for high-performance LIBs.In summary, the nonequivalent substitution of ZnMnThere are no conflicts to declare.