Herein, direct evidence was provided of the importance of Mn2+ in this type of battery by using a bespoke cell. Without pre‐addition of Mn2+, the cell exhibited an abnormal discharge–charge profile, meaning it functioned as a primary battery. By adjusting the Mn2+ content in the electrolyte, the cell recovered its charging ability through electrodeposition of MnO2. Additionally, a dynamic pH variation was observed during the discharge–charge process, with a precipitation of Zn4(OH)6(SO4)⋅5H2O buffering the pH of the electrolyte. Contrary to the conventional Zn2+ intercalation mechanism, MnO2 was first converted into MnOOH, which reverted to MnO2 through disproportionation, resulting in the dissolution of Mn2+. The charging process occurred by the electrodeposition of MnO2, thus improving the reversibility through the availability of Mn2+ ions in the solution.Poor cycling stability and mechanistic controversies have hindered the wider application of rechargeable aqueous Zn–MnO Taking charge: This work provides direct evidence of the importance of Mn2+ in rechargeable aqueous Zn–MnO2 batteries. The charging process is implemented by the electrodeposition of MnO2, thus improving the reversibility by the availability of Mn2+ in the solution. A conversion reaction between MnO2 and MnOOH occurs, rather than Zn2+ intercalation into MnO2. However, the mechanism of the Zn–MnO2 battery is far from being fully understood and remains highly controversial. Various reaction mechanisms in Zn–MnO2 batteries have been proposed. The earliest reports of rechargeable Zn–MnO2 batteries described a mechanism based on reversible Zn2+ insertion/extraction into/from the tunnels of α‐MnO2.2 and H+ without zinc ion intercalation and de‐intercalation.2 cathode underwent successive H+ and Zn2+ insertion/extraction.2 was reported without zinc intercalation into the tunnels of MnO2.3O4 has been proposed by Jiang and Ji.2 batteries is important in optimising this battery chemistry in terms of performance and lifetime and thus promoting the large‐scale application of this type of battery.Recently, rechargeable aqueous Zn–MnO2 (2×2), β‐MnO2 (1×1), δ‐MnO2 (1×∞) and todorokite‐MnO2 (3×3), which may also be responsible for the variety in reported charge‐storage mechanisms when the material is employed as an electrode.2 with 2×2 tunnels has been of particular interest for Mn‐based Zn‐ion batteries. In common with other batteries, the performance of Zn–MnO2 batteries is generally investigated in coin cells with a small amount of electrolyte (≈100 μL), which complicates in operando measurements of mechanistically relevant parameters such as local variations in electrolyte pH and composition. Consequently, in this work, a custom cell was designed and used to reveal the reaction mechanism of the rechargeable aqueous Zn–MnO2 batteries, which enabled monitoring of variations in electrode potential and pH value of the electrolyte. The effect of solution‐phase Mn2+ on cell reversibility was explored: it was found that this species significantly changes the ability to charge the cell. The potential of the electrodes and pH of the electrolyte was monitored in situ during the discharge–charge process. The morphological and structural evolution of the α‐MnO2 cathode was also investigated during discharge–charge using SEM, XRD, TEM, scanning TEM energy‐dispersive spectroscopy (STEM‐EDS) and Raman spectroscopy. This combined structural and electrochemical study sheds light on the reaction mechanism of rechargeable aqueous Zn–MnO2 batteries.As well as the differences in mechanism, various crystallographic polymorphs of manganese dioxide exist, such as α‐MnO2 batteries, which is well indexed into the crystalline phase of α‐MnO2 (JCPDS: 44‐0141). The morphology was first observed by SEM in Figure 2 nanorod with a well‐defined lattice constant of 0.310 nm for the (3 1 0) crystal plane, indicating its high degree of crystallinity. As shown in Figure 2 tunnels by the synthetic conditions owing to the KMnO4 starting material. Figure Figure 2 as the cathode, Zn foil as the anode and 2 m ZnSO4 with different concentrations of MnSO4 as the electrolyte without separator. Although the cycling stability of the MnO2 electrode has been improved by adding Mn2+ to the electrolyte, the practical function of Mn2+ is far from well understood.4, the pure 2 m ZnSO4 solution was first chosen as the electrolyte. Interestingly, we found that the cell was able to discharge with a high specific capacity of approximately 283 mAh g−1 and almost lost the ability to charge, shown in Figure −1 (discharging for 0.3 h at 30 mA g−1) without the addition of MnSO4. Also, a brown deposit was observed on the current collector after charging, indicating the formation of MnO2, shown in Figures S1–S3 in the Supporting Information. Because there is no Mn2+ in the electrolyte, it is inferred that Mn2+ dissolves into the solution during the discharge process. The dissolved Mn2+ in the first discharge process was measured by inductively coupled plasma optical emission spectroscopy (ICP‐OES). The amount of Mn2+ in the electrolyte was calculated to be 0.774 mg, which is less than half of the total mass of Mn in the MnO2 cathode (2.8 mg), shown in Figure S4 in the Supporting Information. Also, previous literature has reported that the dissolution of Mn2+ results in a rapid capacity fade.4 were added to the 2 m ZnSO4 electrolyte, as shown in Figure S5 in the Supporting Information. Significantly, there is no difference in the first discharge curve, but the charging curve begins to recover to a normal state compared with that using 2 m ZnSO4 electrolyte and the cell can then be cycled normally, as shown in Figure 2+ plays a vital role in the charging behaviour of the cell, not simply enhancing the cycling ability of the MnO2 electrode. This abnormal phenomenon requires us to reconsider the discharge and charge processes. If the discharge product is MnOOH or ZnMn2O4 as previously reported in the literature, then the cell should be charged normally to extract the inserted ions without addition of MnSO4, leading to a relatively high reversible capacity, contrary to what is observed in our study.Figure 2+ aqueous solution by the following generalized reaction  consists of stacked zinc hydroxide layers. The interlayer space is filled with zinc sulfate and different numbers of water molecules, resulting in an interlayer distance of 7–11 Å.4(OH)6(SO4)⋅5 H2O shows an upward trend, indicating its growth. During the subsequent charging process, the intensity of Zn4(OH)6(SO4)⋅5 H2O peaks gradually decreases, and finally the corresponding peaks disappear, suggesting a reversible precipitation/dissolution of Zn4(OH)6(SO4)⋅5 H2O. As shown in Figure 2 electrode exhibits a clean surface. After discharging, some large flakes emerge on the electrode surface  and start to emerge with further discharging to 1 V (vs. Zn2+/Zn). During the subsequent charging, the zinc hydroxide sulfate hydrate gradually vanishes. To confirm that the formation of zinc hydroxide sulfate hydrate is pH‐dependent, the discharged electrode was washed with acetic acid. As shown in Figure S8 in the Supporting Information, it is found that the flakes on the electrode surface can be totally removed, verifying the pH‐dependent property of zinc hydroxide sulfate hydrate. It is also observed that there are far fewer MnO2 nanorods on the surface, indicating the consumption of MnO2 during the discharge process. Also, the XRD pattern of the discharged product was compared with the pure ZnMn2O4, which does not demonstrate a good match with reported intercalated phases in the literature, shown in Figure S9 a, b in the Supporting Information. Further, the XRD pattern of the discharged product was compared with pure MnOOH, shown in Figure S9 c, d in the Supporting Information. It was previously reported that the MnOOH phase as a discharged product was submerged in the strong characteristic peaks of current collector (carbon paper) and resultant products.2O4 can be detected. Because the electrode surface is covered by the zinc hydroxide sulfate hydrate, the complex peaks complicate the determination of the structural evolution of MnO2. Therefore, the fully discharged electrode was washed with acetic acid to remove the precipitated zinc hydroxide sulfate hydrate. Surprisingly, only α‐MnO2 can be observed, and no evidence for Zn2+ or proton intercalation into α‐MnO2 can be found, as shown in Figure S10 in the Supporting Information. This is consistent with the report by Oh and coworkers.2 to form intercalated phases; this point was further investigated by TEM and Raman spectroscopy (see below). The phase and morphology evolution of α‐MnO2 with pure 2 m ZnSO4 electrolyte was also investigated, shown in Figure S11 in the Supporting Information. With regard to MnOOH, it is probably unstable owing to the Jahn–Teller effect of the Mn3+ ion, thus forming solid MnO2 and aqueous Mn2+.2+ is not pre‐added to the solution. Additionally, this is consistent with the XRD result after the zinc hydroxide sulfate hydrate flakes are removed with the weak acid. A previous report suggested that there is a phase transformation between α‐MnO2 and Zn‐buserite.2.Ex situ XRD and SEM data was acquired for the α‐MnO2 during discharge process. As shown in Figure 2 electrode is well maintained when it is first discharged to 0.8 V. The corresponding STEM‐EDS mapping reveals various elements such as Mn, O, K, Zn and S in the discharge products, shown in Figure 2 to form the spinel ZnMn2O4 or a tunnel or layered ZnxMnO2 phase.2 and only exists on the surface of the MnO2 nanorod. The Zn element shows a similar distribution to that of S, indicating the formation of zinc hydroxide sulfate, which is consistent with results of SEM and XRD. To further confirm the localized elemental distribution, the corresponding linear scanning position is shown in the highlighted rectangle. As shown in Figure 2. HRTEM was further used to reveal the lattice distance of the discharged products, shown in Figure 2 and a smaller lattice distance of 0.311 nm is consistent with the (3 1 0) plane of α‐MnO2. The TEM results showed that the lattice fringes from the discharged products do not match the reported intercalated phases in the literature. Also, the synthesized pure ZnMn2O4  was investigated by TEM. The corresponding STEM‐EDS mapping shows a homogeneous distribution of Zn, Mn and O, shown in Figures S14 and S15. The lattice distance was revealed by HRTEM, showing a 0.485 nm corresponding to the (1 0 1) plane of ZnMn2O4, shown in Figure S16 in the Supporting Information. Furthermore, the electrochemical performance of the ZnMn2O4 was investigated by galvanostatic charge–discharge method. It is found that the ZnMn2O4 does not exhibit electrochemical activity with an extremely low specific capacity of 15 mAh g−1, as shown in Figure S17 in the Supporting Information. The charge–discharge profile, showing a capacitor‐like behaviour, suggests that the zinc ions in the spinel ZnMn2O4 cannot be extracted from the tunnel, which provides an indirect proof to support our conclusion that the mechanism does not involve the insertion of zinc ions. Steingart and co‐workers reported that ZnMn2O4 is electrochemically inert, which is consistent with our results.2+ has a relatively small ionic radius of 0.75 Å, it is still difficult to find ideal cathode materials to accommodate Zn2+ owing to the strong electrostatic interaction between divalent Zn2+ and solid framework of host materials as well as the difficulty in redistributing the multiple charges of the inserted cations in the host.2+ (Table S1 in the Supporting Information) will add an additional barrier for the intercalation process despite the partial charge shielding of the solvation shell.2.TEM was used to gain further insight into the structural evolution of α‐MnO2, the composition of the electrode and measurement conditions such as discharge current and electrolyte, the capacity obtained at a low current density is more likely to approach the theoretical capacity of the electrode materials. The practical specific capacity of MnO2 reported by most researchers is less than the theoretical capacity of 308 mAh g−1,2/a‐CNT (a‐CNT: acid‐treated carbon nanotubes) nanocomposites displayed an ultrahigh capacity of 665 mAh g−1 at 0.1 Ag−1, the improved capacity was attributed to the numerous oxidative functional groups in the a‐CNT, which could act as additional storage sites.4++e−→Mn3+) in this battery system. It means that the MnO2 will not directly transform into Mn2+ although we have directly observed the dissolution of Mn2+ into the solution. The intermediate is likely to be MnOOH. Previous observations showed that MnOOH dissolved significantly below pH 6 through an acidic disproportionation owing to the high‐spin electronic configuration.2+ and solid MnO2 [2 Mn3+→Mn2+(aq)+Mn4+(s)] owing to the Jahn–Teller effect.2+ and MnO2. This is probably the reason for the difficulty in detecting the intermediate MnOOH. It should be noted that this disproportionation involves a localized electron transfer without contributing any capacity.Although the test capacity may be affected by various factors such as the properties  of MnO3+ can described as follows:Based on the above analysis, the reaction mechanism of this battery during the discharge process including the intermediate MnReductive half‐reaction [Eqs. (2)–Oxidative half‐reaction Eq. :(5)Zn(s)Overall Eq. :(6)4MnO2During the charge process, the reaction mechanism can be formulated as follows:Oxidative half‐reaction [Eqs. (7) and Reductive half‐reaction Eq. :(9)Zn2+4Mn22+ in the solution is crucial to improving the reversibility of this battery, and the pH of the electrolyte needs to be controlled. The precipitation reaction plays a vital role in dynamically tuning the pH of the solution. If the pH is higher than approximately 7 , this type of battery functions as an alkaline primary battery. If the pH is lower than approximately 3 (acid solution), the zinc metal is unstable with respect to the generation of hydrogen gas. This fundamental research is expected to provide useful guidance in promoting the practical application of rechargeable aqueous Zn–MnO2 batteries.Although this aqueous battery is safe and environmentally friendly compared with organic‐based batteries, the key conversion reaction is more complicated than that in intercalation chemistry because it is accompanied by MnOOH disproportionation and a precipitation reaction. The accessibility of Mn2+ in rechargeable aqueous Zn–MnO2 batteries by studying their discharge–charge behaviour. The reaction mechanism was studied with a home‐made cell containing a large amount of electrolyte (≈12 mL) without separator. In the absence of pre‐added Mn2+ in the electrolyte, the dissolved Mn2+ ions are not easily returned to the electrode, leading to the instantaneous failure of the cell. This provides a direct way to identify the degradation mechanism of the rechargeable aqueous Zn–MnO2 batteries, thus helping to reveal the underlying reaction mechanism. By increasing the content of Mn2+ in the electrolyte, the loss of MnO2 from the electrode can be replenished by electrodeposition of MnO2 from the pre‐addition Mn2+ in the electrolyte, leading to a normal charging behaviour. This degradation and charging mechanism also applies to other manganese oxides such as β‐MnO2. More interestingly, the in‐depth investigation of morphology and structure of the discharged cathode reveals that there is a conversion reaction between MnO2 and MnOOH without zinc ion intercalation into the tunnel of α‐MnO2 to form the ZnMn2O4 or ZnxMnO2 phase. The dissolution of Mn2+ is caused by the acidic disproportionation of MnOOH. The precipitation of zinc hydroxide sulfate hydrate is essential to buffer the pH value of the solution. These findings shed light on the degradation and reaction mechanism of rechargeable aqueous Zn–MnO2 batteries and provide useful guidance in designing high‐performance rechargeable aqueous Zn–MnO2 batteries.We have provided direct evidence of the importance of MnThe authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


Though many studies have been conducted, it is necessary to systematically summarize and introduce the research on photocatalytic O2 evolution, and thus a systematic review is needed. First, the corresponding principles about O2 evolution and some urgently encountered issues based on the fundamentals of photocatalytic water splitting are introduced. Then, several types of classical water oxidation photocatalysts, including TiO2, BiVO4, WO3, α‐Fe2O3, and some newly developed ones, such as Sillén–Aurivillius perovskites, porphyrins, metal–organic frameworks, etc., are highlighted in detail, in terms of their crystal structures, synthetic approaches, and morphologies. Third, diverse strategies for O2 evolution activity improvement via enhancing photoabsorption and charge separation are presented, including the cocatalysts loading, heterojunction construction, doping and vacancy formation, and other strategies. Finally, the key challenges and future prospects with regard to photocatalytic O2 evolution are proposed. The purpose of this review is to provide a timely summary and guideline for the future research works for O2 evolution.Photocatalytic water splitting has attracted a lot of attention in recent years, and O 2 evolution reaction is the rate‐determining step in the water splitting process. This review provides a comprehensive summary on the progress of photocatalytic O2 evolution from water splitting, involving the fundamentals and mechanism of O2 evolution, several typical photocatalysts in terms of their crystal structures, synthetic approaches and morphologies, and diverse strategies for O2 evolution activity improvement.Photocatalytic O A large variety of semiconductor photocatalysts have been developed, such as TiO2, ZnO, g‐C3N4, sulfides, bismuth‐based compounds, black phosphorus, etc., which show various applications in the areas of degradation of domestic sewage, industrial dyes and pesticides, water splitting for H2, O2, and H2O2, carbon dioxide conversion for generating CO, CH4, HCHO, CH3OH, etc., nitrogen fixation, removal of harmful indoor gases, solar cell, etc. The photocatalytic process generally includes three stages: i) the generation of the electron–hole pairs in the bulk phase of photocatalysts under light irradiation (hv ≥ Eg); ii) the migration of electrons and holes from bulk to the surface of photocatalysts, and some of them will be recombined during this process; and iii) the residual electrons and holes on the surface of photocatalysts will react with the adsorbed reactants for reduction and oxidation reactions.Over the last few decades, the problems of energy shortage and environmental pollution have become the major issues that need to be solved for the sustainable development of modern society and economics.1 SincHϕ = 285.5 KJ mol−1), which cannot occur spontaneously. It is composed of two half reactions, namely, the hydrogen evolution reaction (HER) and the oxygen evolution reaction (OER), both of which are the processes of Gibbs free with. It can be realized through the solar light irradiation over photocatalysts, which has become one of the most promising ways to produce clean energy due to the advantages of cleanness, low cost and high physicochemical stability, the equations are as followsPhotocatalytic overall water splitting (OWS) is a thermodynamically uphill process 2 producing photocatalyst, the most common method is to test the turnover frequency (TOF), turnover number (TON), and apparent quantum efficiency (AQE) of the reaction system, and the corresponding formulas are in the followingIn order to evaluate the photocatalytic activity and physicochemical stability of Ofollowing28(6)2 per unit reaction time to the incident photons absorption number at a certain monochromatic wavelength, whereas the conversion efficiencies for O2 evolution photocatalysts are distinct based on their different kinds of bandgaps (direct and indirect). Compared with indirect photocatalysts, the direct semiconductor materials have little change in the momentums during the migration process of photogenerated charge carriers after the absorption of incident photons, leading to the faster water oxidation rates than those of indirect ones.In particular, the AQE is the ratio of the moles of generated Oect ones.282 evolution from water splitting mainly focused on some specific class of materials, such as BiVO4, metal–organic frameworks (MOFs), cobalt complexes, etc., or limited to the theoretical calculations, defects design from crystal structures, etc. So far, there is a lack of a systematic overview of the literatures on the photocatalytic O2 evolution. Herein, we provide a review for the advances on the researches of photocatalytic O2 evolution by starting with the fundamentals of photocatalysis and photocatalytic O2 evolution, and then introduce various types of water oxidation photocatalysts based on different crystal structures, compositions, and morphologies. Then, we focus on the diverse strategies that improve the performance of photocatalytic O2 evolution. Finally, the challenges and future prospects in this field are proposed. The summary contents are shown in Figure Previously, the published reviews that are related to the photocatalytic Oes, etc.,31 or 22 evolution reaction is the rate‐determining step during the process of photocatalytic OWS ; ii) to make more use of solar energy, bandgaps of the photocatalysts should be in the range of 1.23–3 eV; iii) the separation of the photogenerated electron–hole pairs within the photocatalysts should be fast under illumination, and then the corresponding redox reactions can occur in time on the surface of photocatalysts. In addition, the water oxidation active sites should be sufficient on the surface of the photocatalysts in order to restrain the rapid recombination of the photogenerated electrons and holes on the surface. Nowadays, the heterogeneous photocatalysts are still the most commonly researched photocatalysts.As Oatalysts.35 In 2.12) can be mainly divided into four types: TiO2(B), brookite, anatase, and rutile, all of whose crystal structures are composed of TiO6 octahedra  is the layered titanate, and the brookite TiO2 belonging to rhombic system has a low physicochemical stability. As anatase TiO2 has more oxygen defects inside the crystal and greater band bending to capture the electrons and promote charge separation than the rutile, the photocatalytic activity of the anatase TiO2 is usually better than that of the rutile phase. Commercial TiO2 (P25) is composed of both anatase and rutile, which has a higher photocatalytic activity than the monocomponent one, due to the formation of energy barrier at the interface that improves the separation efficiency of the photogenerated charge carriers.Titanium dioxide   to anatase, and finally some of them were transformed into rutile, showing the 1D nanowires morphology with photocatalytic activity higher than that of the P25 nanoparticles due to the large specific surface area  on the surface , 2D nanosheets (by adding hydrofluoric acid)  Figure ,f. For ia Figure .[44]e Figure –j. This e Figure ,l.4 Figure ,h. direction, making the crystal of α‐Fe2O3 multifaceted. As seen from the band structure and DOS of α‐Fe2O3  in the lattice of α‐Fe2O3 can promote the photocatalytic activity of α‐Fe2O3 and have little effect on its bandgap and band edge energy at the same time, which can be regarded as a good strategy.From the perspective of crystal structure Figure 9a,b, theifaceted.78 As 3 Figure ,d, the C2 evolution performance of α‐Fe2O3 can be greatly influenced by the morphology and size, a large number of investigations on the microstructure regulation of α‐Fe2O3 were conducted by using different synthetic techniques under various conditions. The 0D α‐Fe2O3 quantum dots were obtained by the facile microwave‐assisted reverse micelle route  plane  and surfactant (PVP) as the precursors. The α‐Fe2O3 hollow nanocages were composed of nanorods with rough surfaces. The hematite phase was confirmed by the exposed crystal facets  and commercial Fe2O3 nanoparticles (14 µmol h−1 g−1)  Figure ,i, emphas Figure .[93]2.52O2]2+ metal oxide layers  separated by the halide layers and proposed the formula of Xn  to represent the number of halide layers that separated the metal oxide layers, such as BiPbO2X  and BiOX  are the X1 and X2 type of structure, respectively. Later, Aurivillius discovered in 1949 that some layered bismuth oxide compounds consist of the alternately stacked [Bi2O2]2+ units and perovskite layers, like Bi2MoO6, Bi4Ti3O12, Bi3TiNbO9, etc., the general formula can be described as (Bi2O2)2+(An−1BnO3n+1)2−, where A and B represent the cations. In recent years, it has been discovered that some compounds possess both of the above two structural characteristics, and they are called Sillén–Aurivillius perovskites; the crystal structures can be expressed as [(Bi2O2)2X]2+(An−1BnO3n+1)2−, where A and B represent the cations and X represents the halide anions. This type of material has recently been found to demonstrate excellent performance of photocatalytic O2 evolution from water splitting.Sillén et al. found a series of compounds that were made up of the 6− perovskite units and [(Bi2O2)2Cl]6− blocks. Different from some simple oxyhalides, such as BiOCl and BiOBr, the VB maximum of Bi4NbO8Cl is mainly occupied by the highly dispersive O 2p orbitals and is not prone to self‐oxidation due to the stable oxygen anions, so it has great potential for the catalytic water oxidation , which led to the morphology of irregular chunks of the as‐prepared samples  planes were obtained when calcination temperature reached above the melting points of CsCl and NaCl like 650 °C   Figure .2+    with the quantum efficiency of 17.1% at 473 nm  Figure , so the ) Figure , the No.m Figure ,(NO3)2 and Na2S2O8 served as the photosensitizer and sacrificial agent, respectively 2OH has a monoclinic crystal structure, in which the copper‐polyhedral layers and copper‐triazole layers were alternately stacked along [001] direction to form a 3D structure 2OH with the potential of superior water oxidation 2OH reached ≈4.2 mL after 1 h photoirradiation with the addition of sulfite as sacrificial agent.As a kind of organic–inorganic hybrid material, MOFs are generally constructed by the self‐assembly of organic ligands and metal ions/clusters through coordination bonds, featuring the advantages of low density, high porosity, and large specific surface area.112 Inn Figure . As seen3+ ions easily coordinate with the organic ligands, so the Bi‐based MOFs are ideal candidates worth to be explored. For instance, a orthorhombic bismuth‐based metal–organic framework (denoted as Bi‐mna) was obtained by solvothermal method with Bi(NO3)3·5H2O and 2‐mercaptonicotinic acid as the precursors , indicating that Bi‐mna was able to more efficiently separate photogenerated charge carriers , noble metal oxides , and transition metal oxides (TMOs) .In the photocatalytic water splitting process, oxidation reaction requires more rigorous thermodynamic and kinetic conditions. As a small amount of cocatalysts can provide the active sites and trap the photogenerated charge carriers, they can reduce the activation energy and enhance the photocatalytic performance.128 Th3  further boosted the O2 evolution performance of PtOx/WO3, and the most effective one was RuO2 with a O2 evolution rate of 41 µmol h−1  effect  by the impregnation route was also explored. As shown in Figure 2 evolution of TaON was Rh species, and the rate reached ≈43 µmol h−1 when the calcination temperature was 500 °C. In addition, the O2 evolution rate of TaON loaded with two cocatalysts was higher than those of solely loaded samples, demonstrating the superiority of dual‐loading of cocatalysts  on the surface of the TiO2 nanosheets compared to pristine TiO2 nanosheets  was higher than that of Ru/Ir‐loaded counterparts, reflecting the superiority of TMO cocatalysts , it is significant to explore the distributions of different types of cocatalysts on the crystal facets. Under photoirradiation, the Pt and CoOx were respectively accumulated on {010} and {110} facets of BiVO4   ,141 it4 Figure –o, which) Figure ,q and du) Figure . Notablyg Figure , attribuFigure 4 by electrostatic interaction, leading to the O2 evolution rate of 67.2 µmol h−1 with 0.5% loading amount of FeOOH nanosheets hydroxides, 2D ultrathin FeOOH nanosheets with a thickness of ≈1.8 nm were prepared by a bottom‐up strategy, which had abundant OVs compared to the bulk FeOOH and FeOOH nanoparticles Figure 20a–c. Ths Figure .10− (denoted as CoPOM), the MOF MIL‐101 was applied as the support to encapsulate CoPOM inside the cavity. TEM and energy dispersive X‐ray (EDX) mapping images  as the electron transfer platform  to construct TiO2@CoAl‐LDH hollow nanospheres with average size of ≈250 nm, which achieved the utilization of the full sunlight spectrum  was much higher than that of the pristine TiO2 hollow nanospheres (0.27 mmol h−1 g−1)  was grown on the surface of TiOm Figure 22a–c. Un) Figure . The DOSs Figure .[153t Figure ,h. X‐rayy Figure . When th3 Figure . The imps Figure .[154e Figure . As show4 Figure . The synn Figure .[1553N4 was deposited on the dual‐phases BiVO4  to form the p–n heterojunction  than the monoclinic‐phase BiVO4 (321 µmol h−1 g−1) and tetragonal‐phase BiVO4 (256 µmol h−1 g−1) , which was attributed to the formation of p–n heterojunction  Figure ,e. Furthn Figure .[1563.2.2 As a result, the concept of Z‐scheme has been proposed by recombination of the photogenerated holes of PS I and the electrons of PS II  was much higher than that of the pristine BiVO4  were applied as the bridge to construct BiVO4/g‐C3N4 Z‐scheme junction for accelerating the charge separation and water oxidation . The weaker peak in XRD patterns of α‐Fe2O3/C‐C3N4 attributed to the exfoliation and partial carbonization of g‐C3N4 after calcination  was 30‐fold higher than that of the pristine g‐C3N4 (0.7 µmol h−1)  Figure ,k.[1622MoO6 has a great potential in the field of photocatalysis due to the layered crystal structure and suitable bandgap. The heterojunction of Bi2MoO6 hybridized with the good electron transporters, such as g‐C3N4, has shown excellent photocatalytic performances in terms of the degradation of organic pollutants, water splitting for H2 evolution, etc. To further improve the activity or stability, it is significant to introduce appropriate electron transfer media, like Au, Rh, and Ru. Recently, the Bi2MoO6/Ru/g‐C3N4 catalyst was fabricated by the solvothermal method combined with the reduction of precursor  has gradually become the mainstream, as they bring more flexible tunability in the improvement of photocatalytic activity. Nowadays, a generally accepted viewpoint is that the effect of metallic/nonmetallic doping for photocatalysts can be affected by the types of element, doping methods, concentration of dopant, and doping position.Since 1982, it has been found that the incorporation of a certain amount of transition metals, such as Cr and Ru, into TiOnductors.169 At2, the concept of gradient doping with nonmetallic heteroatoms in TiO2 to improve the electronic structure has been proposed. As shown in Figure 2 microspheres doped with boron were synthesized by hydrothermal method and heat treatment with TiB2 as the precursor of titanium and boron. The X‐ray photoelectron spectroscopy (XPS) spectra with argon ion sputtering revealed that the binding energy of B 1s in the TiO2 microspheres changed from 187.9 to 192.2 eV, indicating the transition of substitutional boron (Bδ−) to interstitial boron (Bσ+) during the heat treatment process , the electron will increase in concentration and be captured into the Ti4+ lattice, tuning the latter into Ti3+ species. Then, the Fermi energy level was shifted upward and a built‐in electric field was formed, promoting the separation of photogenerated charge carriers  was synthesized by the impregnation approach. After the introduction of cesium, an amorphous layer appeared on the surface of Cs–WO3 with slightly changed size   Figure . Contrars Figure ,r. Baseds Figure ,[1752 can be adjusted through the nitridation, and its water oxidation performance can be improved by the donor doping, thus the synergistic effect of metal and nonmetal codoping on the water oxidation activity of rutile TiO2 was investigated. Kazuhiko's group reported that the rutile TiO2 nanorods codoped with Ta/N species (denoted as TiO2:Ta/N) were synthesized by the microwave‐assisted hydrothermal route followed by the subsequent calcination treatment in the ammonia atmosphere  as the O2 evolution component in the Rh‐doped SrTiO3‐based Z‐scheme system.The band position of TiOr doping,178 ths Figure ,e. DRS aN Figure . Based os Figure ,j. Impor3 Figure ,l. Furthe Figure , and a s3.3.22 evolution performance based on defect modulation is the vacancy creation. Vacancy is one of the intrinsic defects, which can introduce the defect energy level into the forbidden band of semiconductors, thereby enhancing the light absorbance ability and charge separation.In addition to doping, another strategy for improving the Oparation.185 In recent years, 2D monocrystalline nanosheets are the ideal candidates to create OVs. For instance, the OVs were in situ generated on the surface of BiOCl single‐crystalline nanosheets with the exposed {010}/{001} facets by the hydrothermal method followed by UV‐light irradiation. The water molecules adsorbed over OVs on the {010} facets in the dissociated manner were more easily to be oxidized than that on {001} facets in the molecular manner, thus leading to an enhanced O2 production activity. Yan et al. fabricated WO3 single‐crystalline nanosheets with OVs introduced by the exfoliation and the following calcination treatment  peaks in the infrared region of the DRS spectra also proved the successful creation of OVs on the outside surface of WO3 nanosheets , which was attributed to the LSPR effect induced by the OVs that promoted the utilization of the solar energy , ≈2.3‐fold higher than the pristine WO3 (32.6 µmol h−1), benefiting from the formation of the HxWO3 layer around the edge of WO3  that were closely loaded on the surface of Ta3N5  when 3% Ba was doped, ≈4.79 and 1.97‐fold higher than that of pure SrTiO3 and BaTiO3, respectively. The enhanced O2 production activity was mainly attributed to the optimized energy level of SrTiO3 by the Ba doping , featuring the crystalline structures, synthesis approaches, and morphologies. On this basis of the critical issues that hinder the photocatalytic activity of photocatalysts, such as the low utilization of visible light and fast recombination of photogenerated charge carriers, the corresponding effective solutions, including the cocatalyst loading, heterojunction construction, doping and vacancy formation, and other strategies, are summarized.Under the social background of advancing the sustainable development of energy resources, the photocatalytic water splitting has been gradually become the focus due to the features of abundant resources, environmental‐friendly, etc. As the rate‐determining step of the water splitting reaction, water oxidation is the key bottleneck that restricts the efficiency during this process. This review summarizes the latest research progresses of photocatalytic water oxidation. The content includes the introduction of several classical water oxidation photocatalysts 2 evolution. First, the VB of these perovskites such as Bi4NbO8Cl is mainly occupied by the O 2p orbitals. It makes them not be easily corroded by the photogenerated holes, thus demonstrating high photochemical stability. Second, the synthesis methods of this kind of materials are diverse, including 1SSR, 2SSR, 2PC, and flux, which allows the preparation process to be flexible and easy to optimize the photocatalytic performance. Additionally, the Sillén–Aurivillius perovskites have diverse compositions, which enables the adjustable light absorption or bandgap and photocatalytic O2 evolution activity. Therefore, the rational design strategies based on crystal structure and band structure are expected to yield high‐performance perovskites for O2 evolution in the future.The development of efficient water oxidation photocatalysts is still the present research focus, and the Sillén–Aurivillius perovskites will show significant potential for photocatalytic Oii)2 evolution performance enhancement are mainly achieved by improving the light absorption ability and charge separation efficiency, whereas the researches on the surface catalytic reaction are rarely involved. Actually, the reactive sites of photocatalysts are closely related to the adsorption of reactants and the reaction activation energy. For example, the construction of the surface defects can obviously enrich the active sites of photocatalysts. Therefore, the exploration of the catalytic active sites for O2 evolution is expected to be one of the research priorities in the future.At present, the strategies for Oiii)2 evolution performance. However, less attention has been paid to the physicochemical stability of photocatalysts during the photocatalytic reaction process, whereas the actual situation is that the physicochemical stability of most of as‐synthesized materials is often affected by many factors, such as the synthetic routes and conditions. In addition, some O2 evolution photocatalysts are prone to be self‐poisoned by the photogenerated holes, which results in the inactivation after photoreactions. Thus, developing effective tactics for improving the chemical and physicochemical stability of O2 evolution photocatalysts are necessary.Most of the reported works mainly focused on promoting the separation efficiency of photogenerated charge carriers to achieve the purpose of improving the OIn the last ten years, although a series of oxygen evolution photocatalysts have been developed, there is still a long way to go before the practical industrial applications. The photocatalysts for oxygen evolution still suffer from low efficiency or poor physicochemical stability, and especially most of them require the presence of sacrificial agents and cocatalysts, which also undoubtedly increase the economic costs to the industrial applications. In the future, more efforts are in need: and in which we believe more breakthroughs will be achieved.Besides, the spatial separation for the occurrence of reduction and oxidation reactions should be considered, which can effectively inhibit the inverse reaction that usually occurs on the surface of photocatalyst. For instance, the hydrogen farm strategy proposed by Li's group is a promising direction,222 anThe authors declare no conflict of interest.


The addition of Ti3C2 makes more balanced charge carrier transport in LEDs by changing the energy level structure and electron mobility of ETL. Moreover, lower surface roughness is obtained for the ETL, thus guaranteeing uniform distribution of the perovskite NCs layer and further reducing leakage current. As a result, a 17.4% external quantum efficiency (EQE) with low efficiency roll‐off is achieved with 10% Ti3C2, which is a 22.5% improvement compared to LEDs without Ti3C2.2D transition metal carbides, nitrides, and carbonitrides called MXenes show outstanding performance in many applications due to their superior physical and chemical properties. Herein, a ZnO–MXene mixture with different contents of Ti 3C2 into the ZnO electron transport layer (ETL) to change the energy level structure and electron mobility of the ETL and further make the carrier transport balanced is explored. The ZnO/TiC ETL‐based red emitting perovskite nanocrystals light‐emitting diodes exhibit an external quantum efficiency of 17.4%.Adding an appropriate amount of Ti In addition, they have wider color gamut and higher color purity than II–VI NCs and organic fluorophores. Based on the above advantages, their application in light‐emitting diodes (LEDs) is also progressing rapidly. So far, the external quantum efficiency (EQE) of CsPbX3 NC LEDs has exceeded 20%, which makes them very competitive in the application of next‐generation display technology.All‐inorganic metal halide perovskite nanocrystals (NCs) have very recently emerged as interesting materials for optoelectronic applications because of their superior properties , tunable bandgaps over the entire visible spectral range, and facile synthesis).1 In a EQE = ηoutχγηPL, where ηout is the optical outcoupling coefficient, γ is the fraction of excitons that can radiatively decay, ηPL is the PL QY of emitting layer (EML), χ is the balance of charge injection. Based on the above formula, when only the internal factors of the device are considered, the factors affecting the EQE of LEDs can be mainly divided into two aspects: i) the influence of the light‐emitting layer materials; ii) the design of the device structure. For the former, the strategy of ligand engineering or element doping is basically used to increase the radiative recombination probability of electrons and holes, thereby improving the PL QY. For the latter, a lot of research works have been focused on device structure optimization to guarantee balanced charge injection, which needs appropriate energy band alignment and equivalent carrier mobility of charge transport layer (CTL) materials. For instance, Kim et al. utilized a perfluorinated polymeric acid ) to modify PEDOT:PSS, which forms a self‐assembled gradient work function between PEDOT:PSS and perovskite, thereby reducing the hole injection barrier between them. Similarly to PFI, ethanolamine (EA)‐modified TiO2, polyethyleneimine (PEI)‐modified ZnO, polyethylenimine ethoxylated (PEIE)‐modified ZnO, and PFI modified poly‐TPD have been applied in perovskite LEDs (PeLEDs) to achieve a balanced charge injection and further improving the device performance. In addition, Wu et al. and Zhang et al. used Mg‐doped ZnO and Li‐doped TiO2 as the electron transport layers (ETLs) to reduce the electron injection barrier and the carrier mobility, making the electron and hole injection more balanced. Therefore, the device performances were significantly improved. Sim et al. prepared amorphous Zn–Si–O ETL by magnetron sputtering at room temperature, which had sufficiently shallow electron affinity (≈3.2 eV)  to confine excitons and sufficiently high electron mobility (≈0.8 cm2 V−1 s−1) to transport electrons. Consequently, the very low operating voltage and high power efficiency were achieved for CsPbX (X=Br3 or BrI2) perovskite LEDs.LED's EQE can be expressed as4 EQE n+1XnTx  show outstanding performance in many applications such as supercapacitors, catalysts, conducting thin films, sensors, and antennas due to their superior physical and chemical properties. In addition, MXenes offer relatively high conductivity, high light transmittance in the visible range, and the possibility to tailor their electronic structure such as work function (WF) or bandgap. All of these make MXenes possible to be the transport material in optoelectronic applications. Ti3C2Tx MXenes have been incorporated into the perovskite absorber or into the ETL, modifying the chemical and/or physical properties of the interface and providing superior charge transfer paths, therefore, resulting in an enhanced power conversion efficiency in perovskite solar cells. Considering the interoperability of them, it is also possible to apply the MXenes to LEDs.Recently, 2D transition metal carbides, nitrides, and carbonitrides called MXenes with a formula of M3C2 was incorporated into the ZnO ETL for CsPb0.64Zn0.36I3 NC LEDs. ETL films of ZnO with different Ti3C2 contents were prepared to explore its impact on device performance. It can be seen that the addition of Ti3C2 significantly changes the energy level structure of the ETL and affects the carrier transport balance. The highest EQE up to 17.4% was achieved for 10% Ti3C2 based LEDs, which is 22.5% improvement compared with the ZnO‐based device (14.2% EQE). In addition, the efficiency roll‐off in high current density was also effectively improved due to the more balanced carrier recombination. To our best knowledge, this is the first work to study the influence of Ti3C2 MXene on device performances of perovskite LEDs.Herein, sheet‐like Ti3C2 MXene is illustrated in Figure 3C2 MXene was obtained through etching bulk Ti3AlC2 ceramics by a mixed solution of LiF salt and HCl acid. The accordion‐like Ti3C2 was further exfoliated into Ti3C2 sheets via ultrasound (details in the Supporting Information). Figure 3C2 MXene, respectively. The sheet‐like Ti3C2 MXene has a rather flat and smooth surface with an average edge length of several hundred nanometers  and Ti3C2 MXene are shown in Figure 3AlC2, disappeared after etching out the Al layer. Meanwhile, the (002) plane at 9.54° and (004) plane at 19.1° of Ti3AlC2 were broadened and shifted to smaller angles, indicating the formation of Ti3C2 sheets.The fabrication route of Tis Figure . The edg2 sheets.11b0.64Zn0.36I3 perovskite NCs was based on our previous work using a 0.5 feed ratio of ZnI2 to PbI2 (details in the Supporting Information). Figure S2a (Supporting Information) shows the UV−vis absorption and photoluminescence (PL) spectra of the as‐synthesized CsPb0.64Zn0.36I3 perovskite NCs whose first excitonic absorption peak and PL peak are 675 nm  and 688 nm,  respectively. The corresponding XRD patterns and TEM images are shown in Figure S2b,c (Supporting Information), proving the perovskite NCs with a cubic phase.For perovskite EML, the synthesis of CsPb3C2 (designated as ZTC) mixture are made up of the different amounts of Ti3C2 MXene dispersed into the ZnO NC solution. As shown in Figure S2 (Supporting Information), they formed a uniform dispersion and the color of the solution gradually darkened as the amount of Ti3C2 increased. The XRD patterns of different samples are shown in Figure S3 (Supporting Information). Figure S4a (Supporting Information) shows the TEM image of the ZTC. It can be seen that ZnO is attached to or scattered around the surface of Ti3C2. Figure S4b (Supporting Information) is the diffraction pattern without the Ti3C2 area and the diffraction rings of ZnO can be seen. Figure S4c (Supporting Information) is the diffraction pattern of the mixed region. Compared with Figure S4b (Supporting Information), there are two more obvious diffraction rings, corresponding to the (101) and (110) planes of Ti3C2, indicating that there is indeed Ti3C2 in the mixed region. SEM images of unheated and heated ZTC are shown in Figure S4d,e (Supporting Information), respectively. Figure S4d (Supporting Information) shows that a small amount of sheet‐like Ti3C2 is mixed on the ZnO film, and Figure S4e (Supporting Information) shows that the annealed film has uniform grains. In the process of device preparation, ZTC mixture were spin‐coated onto an ITO substrate as ETL and a light output medium. In order to guarantee a high optical coupling output of the device, the light transmittance of the spin‐coated ZTC mixture film cannot be greatly reduced. Subsequently, the transmittances of pristine ZnO film and ZTC film with different amounts of Ti3C2 were characterized and presented in Figure S5 (Supporting Information). There is no significant change in transmittance between ZTC and pristine ZnO film over the entire visible range, demonstrating that the introduction of Ti3C2 has a negligible effect on the optical coupling output of the device.The ZnO–Tiof Ti3C2,12 ind3C2 on the electronic properties of ETL, the energy levels of ZnO/PEI and ZTC/PEI were explored by UPS measurement and the Tauc plots, the corresponding data are given in Figure 3C2 loading, illustrating that Ti3C2 played a role of energy level regulation. Here, we designed ITO/ZTC/PEI/CsPb0.64Zn0.36I3/TCTA/MoO3/Au device structure . Typically, electrons are injected into the emitting layer through the CBM of the ETL. The energy level difference between the ETL and EML is an important factor affecting the electron injection efficiency. Previous reports have used PEI as an interface modification layer to improve electron injection efficiency while blocking electron transport. The thickness of PEI as an insulating material must be moderate, otherwise, an excessive thickness will increase device resistance, causing unnecessary thermal effects, thereby reducing the device efficiency. Therefore, balanced charge transfer in the device may not be completely carried out only through adding a PEI layer.In order to explore the effect of Tie Figure , and theannealing13  /ETL (70 nm)/CsPb0.64Zn0.36I3 (50 nm)/TCTA (70 nm) /MoO3/Au (60 nm)  were fabricated . The normalized electroluminescence (EL) and PL spectra for all sample devices are shown in Figure S8 (Supporting Information). It is noted that the EL spectra are almost consistent with the PL spectra, illustrating that the EL is indeed from the perovskite NCs without any noticeable contribution from other charge transport materials.Figure 0.64Zn0.36I3 NC LEDs with different ETLs. It can be seen that the turn‐on voltages and maximum luminance of these devices show a slight increase as the additional content of Ti3C2 from 0% to 10%. When the content of Ti3C2 continuously increased to 20%, the turn‐on voltage became larger and the maximum luminance significantly decreased. The increase in turn‐on voltage is attributed to the fact that electron injection gradually changes from a potential well to a barrier, which means that the electron injection capability is lowered. The increase in luminance is attributed to a more balanced charge carrier recombination at low Ti3C2 concentrations. When the content of Ti3C2 increased, the barrier height between ETL and EML further increased, the recombination of charge carriers was again unbalanced, resulting in a sharp drop in luminance. Similar changes were observed in EQE versus current density characters of these four devices .In order to determine the influence of Tiaterials.16Figs Figure , and the−3 cm2 V−1 s−1 which is much higher than that of TCTA (2 × 10−5 cm2 V−1 s−1). The large electron mobility can cause excess electron accumulation at the interface between the EML and the ETL, thereby forming a space charge and causing luminescence quenching, which ultimately leads to a reduction of device efficiency. To further confirm the effect of the addition of Ti3C2 (electron mobility of 2.23 × 10−5 cm2 V−1 s−1) on the carrier transport balance, the electron‐only devices with device structure of ITO/ZTC/PEI (0% and 10%) /perovskite NCs /LiF/Al and the hole‐only devices with device structure of ITO/PEDOT:PSS/perovskite NCs/TCTA/MoO3/Au were fabricated. As shown in Figure S9a (Supporting Information), the current density of the electron‐only device with 10% ZTC became lower than that of the electron‐only device without Ti3C2 and was much closer to that of the hole‐only device, proving that the carrier transport between electron and hole are more balanced. In addition, the electron mobility of different ETL films are estimated by fitting the space‐charge‐limited‐current region (SCLC). As shown in Figure S9b (Supporting Information), the mobility of the added ZnO film with 0% to 20% Ti3C2 content are 1.98 × 10−3 cm2 V−1 s−1, 1.61 × 10−3 cm2 V−1 s−1, 1.16 × 10−3 cm2 V−1 s−1, 9.03 × 10−4 cm2 V−1 s−1, illustrating the addition of Ti3C2 reduces the electron mobility of ZnO and the current density decreases with the increase of Ti3C2 addition. As discussed above, we conclude that the effect of Ti3C2 on carrier transport balance can be divided into two aspects. On the one hand, the addition of Ti3C2 increases the work function of ETL. On the other hand, the electron mobility of ETL decreases due to the lower electron mobility of Ti3C2, thus leading to more matched energy levels and carrier mobility between ETL and HTL. Furthermore, the ZTC layer also offers lower surface roughness than that of the ZnO layer. The AFM image of the ZnO film is shown in Figure S10a (Supporting Information), exhibiting root‐mean‐square (RMS) roughness of 3.42 nm.  The RMS roughness is further reduced to 1.49 nm  when 10% of Ti3C2 was added in the ZnO . The uniform ZTC films ensure low leakage current in LEDs and the semilogarithmic curve illustrated in Figure  The main works on the red emitting CsPbI3 perovskite NC LEDs are summarized in Table S1 (Supporting Information). For now, our devices have achieved the highest EQE. We also tested the repeatability and stability of these devices, as shown in Figure S11 (Supporting Information). The average EQEmax of the four types of devices added with 0, 5%, 10%, and 20% Ti3C2 are 14.2%, 15.3%, 17.3%, and 11.5%, respectively. It can be seen that the LED devices with different Ti3C2 contents have good repeatability. In addition, the times for LED luminances without Ti3C2 and with 10% Ti3C2 drop to half of their initial values are 9 and 23 min,  respectively, illustrating a better charge transport balance can reduce the damage to the device.The electron mobility of ZnO films was 2.21 × 10V−1 s−1).17 The3C2 Mxene was prepared and mixture of Ti3C2 and ZnO NCs were first utilized as the ETL for perovskite NCs LEDs. On the one hand, the addition of Ti3C2 achieves charge carrier balance between ETL and HTL by changing the work function and electron mobility of ZnO. On the other hand, the ZTC film also offers lower surface roughness, thus guaranteeing uniform distribution of the perovskite NCs layer for leakage current reduction. As a result, a relatively high EQE of 17.4% and lower efficiency roll‐off were achieved by using ZTC (10%) as the ETL, which is 22.5% improvement compared with the pristine ZnO‐based device (14.2% EQE), showing a great promise for the applications in future display and lighting.In conclusion, conductive sheet‐like TiThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


Theoretical calculation shows that, with ∼20 meV resolved Co 2p3d resonant inelastic X-ray scattering (RIXS), the excited states of the isotropic 1Ag1(Oh) ground state are split by 3d spin–orbit coupling, which can be distinguished via their angular dependence. However, strong self-absorption and saturation effects distort the spectra of the LaCoO3 single crystal and limit the observation of small angular dependence. In contrast, the RIXS on 55 nm LaCoO3 shows less self-absorption effects and preserves the angular dependence of the excited states.Angle-dependent 2 The angular dependence of RIXS has been examined by Michel van Veenendaal using crystal field calculations for L- and M-edges in the case of a single valence hole   or external magnetic field, the system symmetry is further reduced  is a developing technique that is used to investigate the ground and excited states of transition metals, in particular when combined with the angular degrees of freedom (photon-in) process and an emission (photon-out) process. The absorption process excites the electrons to higher energy levels and the emission process emits photons when the electrons decay to the lower energy levels. Two geometrical effects have to be taken into account here: (i) the probing depth is dependent on the X-ray absorption spectroscopy (XAS) cross-section (saturation) and (ii) the emitted photons can be re-absorbed (self-absorption). Consequently, the emission intensity might be distorted according to the photon energy and the experimental geometry, which is well known in fluorescence yield XAS (FY-XAS)  were compared to illustrate these effects. Furthermore, we show that the fine structure of the spin–orbit coupling and exchange interaction within the 3d6 manifold of a Co3+ ion can potentially be revealed using RIXS angle-dependent measurements.In this article, we raise awareness that the 22.2.1.Af(ωin) is the scattering amplitude which is defined asin and ℏωout are the energies of the incident and the emitted photons; Eg, Ef and Em are the eigenvalues of the ground state |g〉, the final state |f〉 and the intermediate state |m〉, respectively; Γ gives the lifetime broadening of the intermediate state. VI(VE) is the transition operator induced by the electro-magnetic field of the incident (emitted) photon and can be expanded as ∊I (∊E) and kI (kE) and rI (rE) give the wavevector and position of the electro-magnetic wave. For 2p3d RIXS, the 2p3d (3d2p) transition operator is approximated as an electric dipole operator i.e.Eg) is taken as zero. Note that we can ignore the scattering coefficient in ≃ ωout) and the classical electron radius re is constant.The RIXS cross-section is determined by the Kramers–Heisenberg equation  p3/2 and 2p1/2 edges  TQq. Pdipole is the reduced matrix element between the 2p (core) and 3d  levels. With this simplification, the beam polarization and the electronic configuration can be decoupled from each other and they therefore can be linked directly to the experimental geometry.For a better insight into the angular dependence of RIXS, many authors have already reformulated the scattering amplitude by using the spherical tensor expansion et al., 2014f〉 were obtained from the cluster calculation using the program Quanty  which is neglected in the current simulation and discussion since it shows only tiny effects on the spectra. The electronic configurations of Co3+ ions in the LaCoO3 crystal and the LaCoO3/SrTiO3 film were suggested to be the pure 1Ag1(Oh) configuration and a mixture of the 1Ag1(Oh) and the 5Bg2(Dh4) configurations, respectively  ground state configuration are: d = 0.055 eV, 10Dq = 0.595 eV, U = 6.5 eV, Q = 7.5 eV, Δ = 1.5 eV, 5Bg2(Dh4) ground state configuration, Ds = −0.12 eV and Dt = −0.018 eV are applied in addition to the 1Ag1(Oh) case. The calculated excited states are indicated by the term symbol notations in Fig. 1(b)5Bg2(Dh4) state is the subgroup of the 5Tg2(Oh) state in the tetragonal distortion . The emission spectrum becomes the summation of the emitted photon energy (ωout) at the oxygen K edge, which leads to μO(ωin)SO that is approximated as constant with respect to the incident energy of the cobalt L2,3 edge. It also implies that the absorption factors of the emitted photons [μCo(ωout) and μO(ωout)] are approximately constant. Then the emission intensity is simplified asC1, C2 and μO(ωin) are constant with respect to the incident energy of the cobalt L2,3 edge. Thus the emission intensity is mainly dominated by the absorption factor μCo(ωin) which provides a bulk-sensitive absorption spectrum free from self-absorption and saturation effects.If the sample is thick in) is fixed at one and μ(ωout) is scaled from zero to a hundred (infinity) to simulate the trends. The scattering angle is 140°. Depending on the competition between μ(ωin) and μ(ωout), three cases can occur:For RIXS, equation (13)out) ≫ μ(ωin) (Iems ≃ 0). In this condition, the absorption factor of the emitted photons is much bigger than the incident beam absorption. The scattered photons can more easily be reabsorbed in the emission process, which implies that the emission intensity tends to vanish [the line μ(ωout) ≃ 100 in Fig. 2(b)(i) μ(ωout) value. In other words, the spectrum distortion is angle-independent [the line μ(ωout) ≃ 0 in Fig. 2(b)(ii) out) ≃ μ(ωin). In this case, none of the factors can be omitted, as shown by the lines μ(ωout) ≃ 1. The self-absorption effect is dependent on the geometry and the photon energy of both the incident and emitted beams. From the equation, we specify the two extreme conditions. The first case is that the incident beam is propagating along the sample surface plane (grazing incidence), where α ≃ 0. Here, the emission intensity X(ωin) and the total absorption factor μ(ωin). The second case describes the propagation of the emitted beam being along the sample surface plane (grazing exit). It implies that the scattered photons are totally reabsorbed and no scattering intensity should be expected (Iems ≃ 0).(iii) μ(ω3+1Ag1(Oh) ground state as an example for bulk LaCoO3. The theoretical μCo(ωout) and μCo(ωin) can be obtained by calculating the XAS spectra. μCo(ωout) is approximately equal to the total μCo(ω). For a fixed incident photon energy, the incident absorption factor μCo(ωin) is always the same. Note that the total μ(ω) includes not only the contribution of the target element but also the contributions from other elements. We assume that the other absorption channels result in a constant background μB of about 5% of the total μ(ω) . It was grown under a 0.2 mbar O2 background pressure and at a deposition temperature of 750°C. The laser flux was adjusted to 1.9 J cm−2. The sample thickness was 55 nm which is determined by X-ray reflectivity (XRR) measurement. More sample details are given by Tomiyasu et al.  at 20 K. The experimental resolution of RIXS was calibrated as ∼90 meV and the recorded incident energy broadening was ∼1000 meV (FWHM). The calculations of 2p3d RIXS spectra were performed using the program Quanty with intrinsic lifetime broadenings of ∼300 meV and 20 meV (FWHM) for the intermediate and final states, respectively 3 single crystal at the maximum of the Co L3 edge. The elastic peak shows a maximum intensity at the specular angle (β = α) due to strong reflection. In contrast, the intensity of the dd excitations decreases when α increases. In comparison, we show in Fig. 4(b)1Ag1(Oh) ground state, where the angular dependence can be observed. The 3Tg1(Oh) state slightly changes its intensity and shifts its position due to the competition of spin–orbit fine structures, which will be discussed later. However, the trends are not in good agreement with the experimental results because the saturation and the self-absorption effects should be considered in a concentrated sample. By including the saturation and the self-absorption effects, the multiplied results show good agreement for the trends of the angular dependence except that the elastic peak is overestimated [Fig. 4(c)3Tg1 excited states are still preserved theoretically. Unfortunately, better statistics and energy resolution are required to observe this tiny shift in the experiment.Fig. 4 is observed  because of the combination of two dipole transitions (rank 1 transition). On the one hand, the incident absorption channel of RIXS excites the ground state to intermediate states which preserve a mixture of different symmetries. On the other hand, due to the spin–orbit coupling, the final excited state splits into different branches with a small energy difference. In combination, the effective transition preserves spin and orbit characteristics and the polarization weighting factor TQq provides the intensity of states as a function of geometry  and 5(b)p3d RIXS spectra of three geometries with the scattering angle (α + β) set to 90°. With a finer broadening, the energy splittings caused by spin–orbit coupling are better resolved [Figs. 5(a) and 5(b)3Tg1 states split into the Eg, Tg2, Tg1 and Ag1 states, where we have indicated the double-group labels in italic. In the case of vertically-polarized incident photons (V-polarization), the elements in the polarization tensor equally contribute in all geometries. Thus, no angular dependence is expected [cf. Fig. 5(b)c) and 5(d)L3 edge as a function of rotation angle. Here the Eg and Tg2 groups of the 3Tg1 state are selected for comparison. We note that the broadening of the spectra in Fig. 4For an isotropic cf. equation (10)dd excitations can be represented properly. We examine the angle-dependent 2p3d RIXS of the 55 nm LaCoO3 film grown on the SrTiO3 substrate. This thickness is on the limit of the attenuation length of the Co L3 edge (25–125 nm) but can still be considered to be free from self-absorption effects  ground state (gray arrows). As has been discussed, the ground state of the Co3+ ions in the LaCoO3/SrTiO3 film shows a mixture of the low-spin 1Ag1(Oh) state and high-spin 5Bg2(Dh4) state  state, the angular dependence of RIXS is negligible for the 1Ag1(Oh) state  state are indicated by black arrows, but are almost invisible compared with the 5Bg2(Dh4) ground state. Some discrepancies remain, for example the 1.3 eV feature indicated with the red arrow in the 5Bg2(Dh4) calculation is not evident in the experimental data. This could be because: (i) the mixture of the spin states suppresses the intensity of the 1.3 eV feature  – experimentally, not all the CoO6 clusters should be aligned exactly along the pseudo-cubic orientation; (iii) the discussions above do not take the trigonal symmetry reduction into account. The CoO6 clusters are naturally aligned in a trigonal field for a relaxed LaCoO3 crystal  ground state shows angle-dependent 2p3d RIXS spectra, despite the absence of angular dependence in the 2p XAS spectra (Merz et al., 2010et al., 20193 single crystal distorts the RIXS spectra due to the saturation and self-absorption effects which limits the distinguishability. By applying a model to consider these effects, the trends are explained as a function of both the energy and the rotation angle. In contrast, for the LaCoO3/SrTiO3 film, the self-absorption effect is small. Thus the LaCoO3/SrTiO3 film shows better agreement with the angular trend although some features are not reproduced well. The possible reason for this mismatch could be a mixture of ground states, non-perfect empirical parameters and/or reduced symmetry.We have presented the angular dependence of 2


Extensive DFT calculations provided evidence for stereocontrolled proton transfer from catalyst to substrate as the rate‐determining step, and showed the importance of steric interactions from the catalyst's alkyl groups in enforcing the high enantioselectivity. Crystal structure data show the dominance of noncovalent interactions in the core structure conformation, enabling modulation of the conformational landscape. Ramachandran‐type analysis of conformer distribution and Protein Data Bank mining indicated that benzylic fluorination by this approach has the potential to improve the potency of several marketed drugs.A catalytic enantioselective synthesis of heterocyclic vicinal fluoroamines is reported. A chiral Brønsted acid promotes aza‐Michael addition to fluoroalkenyl heterocycles to give a prochiral enamine intermediate that undergoes asymmetric protonation upon rearomatization. The reaction accommodates a range of azaheterocycles and nucleophiles, generating the C−F stereocentre in high enantioselectivity, and is also amenable to stereogenic C−CF Chiral Brønsted acid catalysis enables the synthesis of heterocyclic vicinal fluoroamines by asymmetric protonation. The reaction accommodates a range of azaheterocycles and nucleophiles, generating the C−F stereocentre in high enantioselectivity, and allows the formation of chiral C−CF3 bonds. DFT calculations provided insight into the reaction mechanism, and conformational control of the products is elucidated. Several observations relating to optimization were noted . Ethereal solvents  were particularly effective (entries 1 and 2), with other solvents affording good to excellent conversion but with notably poorer enantioselectivity (entries 3 and 4). Lowering the reaction temperature from −10 to −20 °C had little effect on enantioselectivity but impacted reaction efficiency (entry 5). A similar effect was observed by lowering catalyst loading, where 10 mol % was less efficient but maintained selectivity (entry 6). Control experiments supported a catalyst‐promoted reaction that lacked background reactivity (entry 7).Based on this hypothesis, a benchmark process was established whereby ortho, meta, para) was typically accommodated while maintaining selectivity (Scheme 7 m). It should be noted that the choice of solvent was important for conversion, due to solubility: product precipitation as the reaction progressed became problematic for certain substrates; however, changing ethereal solvent based on substrate (THF or CPME) resolved this issue.The scope of the reaction was investigated Scheme . A rangey Scheme  a. Addit7 a, 7 n–q) and the reaction tolerated quinoxaline (7 r), benzothiazole (7 s), and pyridine (7 t), with the latter a significantly more challenging substrate due to its higher dearomatization barrier, hence requiring a high temperature for the reaction to proceed, which negatively affects enantioinduction. Significantly, the reaction also allows enantioselective formation of stereogenic C−CF3 bonds (11); however, catalyst loading had to be increased and reaction temperature decreased to overcome a significant nonselective background reaction observed for this substrate (see the Supporting Information).A range of vinylheterocycles was also generally well accommodated Scheme  b. Varia5 a) by the catalyst to provide LUMO‐lowered intermediate complex 12 and enabling reversible aza‐Michael using PhNH2 (8) to deliver key intermediate 13. Two mechanistic pathways are then possible from this intermediate: pathway 1 proceeds via direct stereocontrolled proton transfer from the anilinium via TS1 and delivers the product‐catalyst complex 15, which subsequently liberates the product (7 a). Alternatively, in pathway 2 proton transfer from the anilinium of 13 to the phosphate (via TS2) delivers 14, which undergoes stereocontrolled proton transfer via TS3 to deliver 15. In our previous report,TS1 leading to the observed enantiomer. These purely quantum mechanical studies did not yield transition states that would have supported pathway 2 . A series of kinetic isotope effect experiments were conducted via the use of 15N‐aniline and PhND2. However, these proved inconclusive, with independent rate experiments (see the Supporting Information) resulting in observed 14/15N KIE of approximately 0.8 and H/D KIE of approximately 1.8, which might be affected by the pre‐RDS equilibrium associated with this reaction.Two main mechanistic pathways are possible for the key asymmetric protonation event Scheme . The ini9 a, can be studied effectively and efficiently by QM/MM ONIOM calculations where the quantum mechanical aspects are described by B3LYP/6‐31G** and the molecular mechanics by UFF.Goodman and others have shown that catalysis by BINOL‐derived catalysts, such as 5 a with 8 (see the Supporting Information).Experimentally, no background reaction was observed, which was consistent with DFT calculations that indicated a prohibitively high barrier for direct reaction of 5 a with 9 a to give 12 is moderately favourable, with the preferred dipole‐induced s‐trans conformation of 5 a also retained in 12. This initial complex is held together by a H‐bond (OH⋅⋅⋅N=1.63 Å) and a weak peri CH⋅⋅⋅O=P interaction (2.40 Å). Complex 12 then undergoes dearomatizing aza‐Michael addition to deliver 13, where the loss of aromaticity is compensated for by the formation of a tightly bound ionic interaction between the anilinium NH and phosphate (P=O⋅⋅⋅HNHPH=1.38 Å and P=O⋅⋅⋅HNquin=1.71 Å). Rearrangement within this complex by proton transfer from the anilinium to the phosphate involves a low barrier and yields a complex of the enamine (14) that is higher in energy than reactants. All of these steps are therefore strongly reversible and no significant concentration of any of the intermediates subsequent to quinoline complexation would be expected—this was confirmed by parallel NMR experiments (see the Supporting Information).Complexation of 15‐(R) and 15‐(S) were optimized leading to an array of conformations for each of TS1 and TS3. Consistent with our previous report,TS1). This process has a significant barrier  but is predicted to be highly stereoselective  in favour of the experimentally observed enantiomer. Although this rationalizes the stereoselectivity of the process, it is not consistent with the experimental rate of reaction.Transition states leading to each low‐energy conformation of complex 14 via TS3 and exhibits a clear preference for the experimentally observed enantiomer , which arises from geometrical restrictions between the catalyst and enamine in the developing transition states TS3‐(R) and TS3‐(S) (vide infra). The iPr substituents of catalyst 9 a are also particularly important for imposing this geometrical restriction  and ‐phenyl (9 c) analogues , resulting in the observed diminished conversion. This is accompanied by a sharp erosion in ΔΔG≠ between TS3‐(R) and TS3‐(S)—the energy associated with TS3‐(R) remains similar for all three, and this erosion is principally driven by a change in energy of TS3‐(S). This is highlighted in the preferred conformation of each of the three structures equivalent to TS3‐(S) ; this interchanges with an interaction between the aniline NH and a phosphate oxygen . The combination of the steric clashing and this change in hydrogen bonding pattern clearly disfavours TS3‐(S) compared to TS3‐(R) for 9 a; this difference is significantly reduced for 9 b and 9 c.With regards to the optimal catalyst, a catalyst survey demonstrated the superior level of asymmetric induction using 5 t, associated with the larger dearomatization barrier (see the Supporting Information). The observed significant background reaction for 10 was also investigated computationally, confirming the accelerating role of the LUMO‐lowering CF3 unit as previously observed for other Michael acceptors.In line with experimental observations, the computational model also confirms a lower rate of catalysed reaction for 7 a shows an anti relationship between C−F and the aniline nitrogen , which is likely preferred in comparison to the gauche due to a favourable Npz⋅⋅⋅σ*C‐F interaction . The dihedral angle between the C−F bond and quinoline nitrogen is 174°, explicitly affected by the C−F/NAr dipole minimization and no longer modulated by the other effects described.To reduce the impact of hydrogen bonding and any N6 Figure  a. The p7 a, 16, and the parent 2‐pyridylethylamine (not shown) were computed and reveal that the introduction of the benzylic fluoride has a profound effect on the overall conformational landscape  should be addressed to the authors.SupplementaryClick here for additional data file.