Comparison of the non‐resonant and resonant XES for a series of high‐spin ferrous and ferric complexes shows that oxidation state assignments that were previously unclear are now easily made. The present study spans iron tetrachlorides, iron sulfur clusters, and the MoFe protein of nitrogenase. While 1s3p RXES studies have previously been reported, to our knowledge, 1s3p RXES has not been previously utilized to resolve questions of metal valency in highly covalent systems. As such, the approach presented herein provides chemists with means to more rigorously and quantitatively address challenging electronic‐structure questions.The ability of resonant X‐ray emission spectroscopy (XES) to recover physical oxidation state information, which may often be ambiguous in conventional X‐ray spectroscopy, is demonstrated. By combining Kβ XES with resonant excitation in the XAS pre‐edge region, resonant Kβ XES (or 1s3p RXES) data are obtained, which probe the 3d Assigning transition‐metal physical oxidation states is a major goal of X‐ray spectroscopy. However, competing influences of covalency and d‐count on Kβ XES spectra often make assignments ambiguous. It is now shown that resonant Kβ X‐ray emission spectroscopy (RXES) yields unambiguous oxidation‐state determinations for iron monomers, dimers, cubanes, and metalloenzymes. The concept of oxidation state is fundamental to how chemists communicate about electronic structure and chemical reactivity. We distinguish formal oxidation states, which are effectively used for bookkeeping of electrons, from physical oxidation states, which are assigned based on experimental observables. Spectroscopic methods play a key role in the identification of physical oxidation states. X‐ray spectroscopy is frequently utilized to assign oxidation states in transition metal (TM) complexes, largely due to its element selectivity and involvement of core electron excitations which provide an experimental means to probe the changes in effective nuclear charge at the selected photoabsorbing atom. Metal K‐edge X‐ray absorption spectroscopy (XAS) serves as a reasonable fingerprint for changes in metal oxidation state, with the 1s to 4p edge features of first‐row transition metals (TMs) shifting up by ∼1 eV per unit change in oxidation state.10, 11III complexes have been used to both supportIII physical oxidation state assignment.Metal L‐edge XAS (2p→3d)25, 26n3p5 final state (FS). Hence, Kβ mainlines are dominated by 3p–3d exchange coupling that splits the mainline into Kβ1,3 and Kβ′ features and provides information about the number of unpaired d‐electrons (where a larger number of unpaired d‐electrons results in larger 3p–3d exchange coupling).2+/1+/0 complexes spanning three oxidation states, we showed that the rising edges of the diferrous and mixed valent iron dimers are nearly superimposable and that the Kβ XES mainlines of the full [Fe2S2]2+/1+/0 series are essentially identical. Despite this, crystallographic data, Mössbauer, and calculations of spectroscopic data supported diferrous, mixed‐valent, and diferric oxidation states assignments for these complexes, respectively. By combining spectroscopic results with computation, we were able to highlight the relative strengths and weaknesses of each spectroscopic approach in assessing the oxidation state.Consequently, it appears that there are many ways for X‐ray spectroscopic measurements to lead to ambiguous conclusions. In our view, this issue is best addressed by taking a holistic view of all available spectroscopic data (both X‐ray based and other methods) to arrive at a consistent picture of the electronic structure. In a recent studyThis however, raises an important question: is there a spectroscopic approach that allows our concepts of metal valency, covalency, and oxidation state to be more rigorously tested and experimentally assessed? It is here that we draw inspiration from the words of Linus Pauling in his 1948 Liversidge lecture where he stated with regard to our understanding of valency, “…and we may hope that powerful methods of investigation that are not yet known will be discovered”.Below, we present 1s3p RXES (Kβ RXES) as a means to recover physical oxidation state information, which may be lost in classical non‐resonant emission (Kβ XES) measurements due to the countering effects of covalency and metal d‐count Figure , was stu53dn+1 FS. By performing both resonant and non‐resonant XES, we are able to experimentally differentiate between different 3dn counts despite the changes in covalency. As such, the presented 1s3p RXES approach provides chemists with a tool for more rigorously assigning physical oxidation states. To our knowledge the ability to use 1s3p RXES to assess oxidation states in highly covalent systems has not previously been reported. The extension of 1s3p RXES to the MoFe protein of nitrogenase further establishes the viability of this method for studying electronic structural questions in biological systems.In a Kβ RXES experiment, the 1s electron is excited into specific unoccupied orbitals by tuning the incident excitation energy. For instance, when resonantly exciting into the pre‐edge region of the XAS spectra (dominated by 1s to 3d transitions), 1s3p RXES spectra are obtained, with 3pIIICl4]1− and [FeIICl4]2−.In order to discuss the advantages of 1s3p RXES, it is helpful first to illustrate how changes in oxidation state manifest in Fe Kβ HERFD‐XAS and Kβ XES, and how these changes are modulated by covalency. Figure 2+ to Fe3+. In contrast, the changes in the Kβ mainlines are more subtle, with a modest decrease in the intensity of the Kβ′ feature (at ca. 7045 eV) and a slight decrease in the energy of the Kβ1,3 maximum (by −0.4 eV) on going from d6S=2 Fe2+ to d5S=5/2 Fe3+. In a simple picture, one would expect the splitting between the Kβ1,3 and Kβ′ to increase with an increasing number of unpaired d electrons. The fact that the splitting of the Kβ mainline features is somewhat smaller in the S=5/2 ferric case than the S=2 ferrous case indicates that the increase in covalency (resulting from the Fe−Cl bonds contracting from (2.30±0.01) Å to (2.18±0.01) Å upon oxidation) counters the increase in spin.The XAS spectra show the classic changes that one expects to see upon oxidation of the metal center. Namely, the Fe K‐edge shifts up in energy by ∼1.2 eV, reflecting the increase in effective nuclear charge on going from Fe2FeIIFeIIS, mixed valent [L2FeIIFeIIIS2]1− and diferric L2FeIIIFeIIIS2. The assignment of oxidation states in this series is based on both structural changes and Mössbauer parameters . While the diferrous complex has the expected lowest energy rising edge, the mixed‐valent and diferric complexes are nearly superimposable. Consequently, in this case, the use of the Fe K‐edge as an isolated measure of oxidation state could lead to an incorrect assignment of the physical oxidation state. For the Kβ mainlines, the changes are even smaller. Upon going from the diferrous to the diferric, the Kβ1,3 maxima shift by <0.3 eV and the Kβ′ for all three complexes remains constant at 7059.3 eV. This is once again attributed to the canceling effects of covalency and changing d‐count, as has been previously established.2+ (S=3/2), [Fe4S4]2+ (S=0), and the MoFe protein of nitrogenase. The MoFe protein contains 8 Fe2+ ions in its P‐cluster and 4 Fe2+/3 Fe3+ in its iron–molybdenum cofactor FeMoco,3S4]2+ contains 2 Fe2+/Fe3+ with a delocalized mixed‐valence pair,4S4]2+ contains 2 Fe2+/2 Fe3+. Hence, one would expect the MoFe, containing the highest fraction of FeII should have the lowest energy edge, however as shown here , and a 3p53dn+1 final state (FS). Thus, the contrasting electronic configurations accessed via resonant (dn+1) and non‐resonant (dn) XES will result in distinct multiplet structures, analysis of which could allow for unambiguous assignment of physical oxidation states.As shown in Figure IICl4]2−, and [FeIIICl4]1−, at three different incident excitation energies: one in the pre‐edge and two in the rising edge . Two clear features appear in the Kβ1,3 region with maxima at ∼7059.3 and ∼7061.4 eV, in addition, a weak shoulder is observed at ∼7051.4 eV. The differences in the RXES spectra in the pre‐edge region relative to those in the rising edge clearly demonstrate that a different FS (3p53d7 for [FeIICl4]2−) has been reached.In order to illustrate this hypothesis and show the dependence of RXES on the excitation energy, we present 1s3p RXES of FeIICl2−, and [IIICl4]1− with resonant excitations in both the edge and pre‐edge region. Once again, the data obtained with resonant excitation into the edge region are essentially identical to the Kβ XES indicating that, at these energies, the FS is equivalent to the non‐resonant process. Interestingly, for resonant excitation into the pre‐edge region, one observes a Kβ1,3 maximum at 7060.6 eV and a weak shoulder at 7057.9 eV. Similar to the ferrous case, upon resonant excitation in the pre‐edge region, no well‐defined Kβ′ feature is observed, distinctly demonstrating that the 1s3p RXES of the ferric species is not simply the equivalent of Kβ XES on a ferrous complex, despite their seemingly equivalent 1s23p53d6 FS. This indicates that one must consider the IS that are available to be populated by resonant excitation. In the section that follows, the detailed origins of the differences in Kβ XES and 1s3p RXES are evaluated.Figure 6 and d5 Kβ XES spectra . Atomic Russell Saunders terms for these processes are given in Table S4. For the high‐spin d6 ferrous case, the ground state term symbol is 5D. When a 1s electron is ionized, a 1s13d6 state is reached, giving rise to 4D and 6D configurations when alpha and beta 1s electrons are ionized, respectively. Following the Kβ emission process, a 3p hole is created giving rise to a 3p53d6 electron configuration. Coupling the 2P configuration of the 3p hole to the 5D ground state term symbol gives rise to 4,6P, D, and F final state multiplets (Table S4). The multiplet splitting is dominated by 3p–3d exchange, which separates the XES spectrum into a 4Γ (Kβ′) and 6Γ . In the d5 case, a similar picture may be derived, where the FS spectra are instead characterized by splitting of the 5Γ and 7Γ states. These simple pen and paper predictions are readily captured by atomic multiplet calculations  and fully consistent with previous assignments for Kβ XES spectra.13dn+1 IS  feature using multiplet simulations.This is supported by the multiplet simulation shown in Figure 5D term from the d6 configuration coupled with 2S of the 1s core hole results in a 6D IS. In the Kβ XES, 6F, 6D, and 6P FS multiplets may be accessed. In contrast to the ferrous case, there are no low lying multiplets that the 5D parent term can interact with, resulting in a more simplified FS multiplet structure relative to that of the ferrous. These trends are validated by multiplet calculations, Figure For the ferric case, the 2pd and G1,3pd) were systematically scaled from 100 % (e.g. the atomic limit) to 40 % . These simulations show that though the spectra are modulated by covalency, the general multiplet structure is maintained even for an unphysical 40 % reduction and the splitting between the 5G and 5F/5D terms is still large enough to be resolved experimentally. To further assess this assumption, data on highly covalent systems are presented in the next section.The contribution of covalency to 1s3p RXES spectra was investigated through computational studies in which the Slater Condon parameters (FRXES and ΔXES) generated by subtracting a ferric reference from ferrous species are given following the color of their ferrous parent (negative axis). The all‐ferric reference spectra are chosen to as [FeIIICl4]3− for the monomer . Additional RXES cuts can be found in the Supporting Information .The RXES spectra of all molecules are plotted in Figure r Figure  and L2Fes Figure  and C, tXES, and ΔRXES. Figure XES in comparison to ΔRXES. The changes are more pronounced for the diferrous dimer than the mixed valent dimer. In the diferrous case, resonant excitation into the pre‐edge gives rise to two features (at ∼7058.8 and ∼7060.5 eV), attributed to 5Γ derived from the 4F and 4P parent terms of the d7 configuration, as discussed above. In the mixed valent complex, two features are still present, however, the lower energy feature is reduced in intensity, consistent with the fact that only half the irons are in the ferrous form. This is also highlighted by the difference spectrum ΔRXES.Figure 3S4]2+ and [Fe4S4]2+ and the MoFe protein of nitrogenase were studied  and the existence of a low‐lying excited state with the same spin multiplicity with which mixing can occur via configuration interaction. This same phenomenon should occur for all TMs that reach an F term in the IS. Based on these observations, general groupings can be made for all first‐row TMs with high‐spin configurations:The results presented should be readily translatable to all high‐spin first‐row TMs of any d1,3 feature will be observed for all TMs in which an F term is its IS. E.g. d1, d2, d6 and d7 ground state configurations.A splitting of the 1s3p RXES Kβn≥5 TMs the 1s3p RXES at the pre‐edge, will be absent from the Kβ′ feature.For d1,3 feature in Kβ XES should be observed for TMs with F ground state terms, providing experimental resolution is sufficient .A splitting of the Kβn counts.The general pairings of resonant and non‐resonant XES multiplet behavior thus provide a means to distinguish electronic structural ambiguities for a range of 3dIn recent years, the literature has seen increasing applications of Kβ XES to a wide range of TM systems, including in situ and operando studies of catalysts,76, 77In our view, this work brings us one step closer to realizing the vision of Linus Pauling who said, “If scientific progress continues, the next generation may have a theory of valency that is sufficiently precise and powerful to permit chemistry to be classed along with physics as an exact science.” Although we are not there yet, we have added another tool to the chemists’ toolbox to bring the community one step closer to realizing this goal.The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


PLOS Biology revealed that Aβ from the liver can promote brain pathology, supporting that peripheral Aβ can contribute to neurodegeneration.Accumulation of amyloid beta (Aβ) in the brain in Alzheimer disease drives pathophysiology. A study in this issue of  Alzheimer disease is the most common type of dementia, and it is characterized by the presence of amyloid plaques in the brain that are comprised of aggregated amyloid beta (Aβ) peptides. The Aβ peptide is a fragment of the larger integral membrane protein called amyloid precursor protein (APP) that is expressed in neurons of the brain. When APP is cleaved by β and ɣ secretase enzymes, the resulting Aβ peptides form small soluble aggregates that induce toxicity in the brain and neuronal dysfunction followed by the deposition of Aβ peptides into insoluble amyloid plaques. Many studies have demonstrated that transgenic mice with expression of human APP carrying Alzheimer disease–causing mutations in the brain have high levels of toxic Aβ, neuronal pathophysiology, memory impairments, and plaque formation . The delHigher levels of Aβ peptide isoforms were found in the blood of individuals with pathological Aβ accumulation in brain and cognitive impairment . This coCommonly used transgenic mouse models with human APP transgenes carrying familial mutations that produce high Aβ levels in the central nervous system recapitulate important features of Alzheimer disease. Studies on these mice have long served as a foundation for the advancement of new therapies. The age-dependent pathological phenotypes reported in the HSHA mice created by Lam and colleagues include abnormal lipid accumulation in the brain, neurodegeneration, and the dysfunction of capillaries associated with lipofuscin aggregation. The HSHA mice also demonstrated impaired performance in the passive avoidance test of hippocampal-dependent memory at 12 months of age. Other human APP mouse models with Aβ produced in the brain can have more aggressive Alzheimer disease–related pathology and earlier onset cognitive impairments in a range of behavioral tests. It is important to consider that the extent of the pathological effect of peripheral Aβ on the brain may depend on a limited amount of Aβ that enters the brain and the region where it accumulates upon entry. Nevertheless, this new study supports that peripheral Aβ is sufficient to promote pathological disease–associated features in the brain and memory impairment.Could targeting peripheral Aβ be another strategy to prevent neurodegeneration? Pharmacological reduction of Aβ levels in the blood of wild-type mice by peripheral administration of a drug impermeable to the blood–brain barrier was linked to reduced Aβ levels in the brain . Dietary


Their moderate lithiation potential and reversible redox couples hold the great promise in the development of next‐generation lithium‐ion batteries (LIBs) that are characterized by high power density, long lifespan, and high safety. Despite these outstanding merits, there is still extensive advancement space for further enhancing their electrochemical kinetics. And the industrial feasibility of Wadsley–Roth crystallographic shear structure niobium‐based oxides as anode materials for LIBs requires more systematic research. In this review, recent progress in this field is summarized with the aim of realizing the practical applications of Wadsley–Roth phase anode materials in commercial LIBs. The review focuses on research toward the crystalline structure analyses, electrochemical reaction mechanisms, modification strategies, and full cell performance. In addition to highlighting the current research advances, the outlook and perspective on Wadsley–Roth anode materials is also concisely provided.Wadsley–Roth crystallographic shear structure niobium‐based oxides are of great interest in fast Li Recent progress of developing Wadsley‐Roth phase anode materials are summarized. The review focus on research toward the crystalline structure analyses, electrochemical reaction mechanisms, modification strategies, and full cell performance. With the aim of realizing the practical applications of Wadsley‐Roth phase anode materials in commercial LIBs, the outlook and perspective is also concisely provided. Goodenough, and M. Stanley Whittingham for the development of lithium‐ion batteries (LIBs).1 Sinc However, the implementation of LIBs in EVs to replace traditional internal combustion engine technology still faces great challenges in terms of energy density and charging time. For instance, the driving mileage of Tesla Model 3 is about 500 km with a fully charging time of more than 1 h by using a specific quick charger according to its owner's manual, while the conventional petrol‐driven cars can be refueled within 5 min with a longer driving mileage of above 750 km. Therefore, developing hybrid electric vehicles (HEVs) may be a compromising solution, which is more acceptable and economically more practical. Generally, high‐power batteries are required for HEVs to continually capture/release energy during braking, starting, and accelerating the vehicle in a short period of time. Unfortunately, the commonly used graphite anodes fail to meet the requirements for high‐power LIBs in HEVs owing to their insufficient electrochemical kinetics stemming from the sluggish Li‐ion diffusion rate. Recently, Wadsley–Roth crystallographic shear structure niobium‐based oxides including Nb12O29, TiNb2O7, Ti2Nb10O29, TiNb24O62, Nb16W5O55, Nb18W8O69, PNb9O25, etc., have been extensively investigated as high‐rate anode materials for LIBs. Their unique features of 3D open crystalline structure and reversible redox couples hold the promise of developing next‐generation LIBs characterized with high power density and long lifespan.In recent years, the rapid development of electric vehicles (EVs) has motivated the technical innovation of LIB industry.3 Howe Moli Energy's first attempt to commercialize rechargeable Li‐based batteries with metallic Li as the anode failed due to serious safety hazards resulting from the uncontrolled growth of Li dendrites during cycling, forcing them to withdraw the product from the market. One of the Nobel Laureates mentioned above, Akira Yoshino proposed replacing reactive Li in the anode with carbon‐based materials, which later became the foundation of commercially viable LIBs. In a fully charged LiMO2//graphite battery, the lithiated graphite anode (LiC6) possesses similar chemical reactivity to metallic Li. Thus, the solid electrolyte interface (SEI) play an important role in stabilizing the electrode/electrolyte interphase and constraining the continuous reaction between LiC6 and electrolyte while using graphite as the anode. However, the decomposition temperature of SEI was identified to be as low as 60 °C, which would trigger chain‐reactions between LiC6 and electrolyte, leading to the thermal runaway potentially. Moreover, lithium plating issue of the graphite anode during high‐rate cycling is also considered as an important incentive for easier thermal runaway or even the internal short circuit of batteries, due to its relatively low lithiation potential (<0.2 V vs Li+/Li). However, Li‐ion insertion into Wadsley–Roth crystallographic shear structure niobium‐based oxides mainly occurs above 1.0 V (vs Li+/Li), where the formation of SEI film is not necessary. Therefore, Wadsley–Roth phase anode materials can afford great advantages over the conventional graphite anode in terms of rate capability and safety performance, which should be promising candidates for high‐power LIB anodes in HEVs. Despite the outstanding characteristics of Wadsley–Roth crystallographic shear structure niobium‐based oxides, there is still extensive advancement space for further enhancing their electrochemical kinetics. And the industrial feasibility of Wadsley–Roth crystallographic shear structure niobium‐based oxides as anode materials for LIBs requires more systematic research.Besides, safety feature has always been a key design consideration for the practical application of LIBs, particularly high‐power ones.13 MolThis article is mainly concerned about how to realize the practical applications of Wadsley–Roth phase anode materials for high‐rate and high‐safety LIBs. The rest of the manuscript is organized as follows. Section 2Figure Eg of the highest occupied molecular orbital (HOMO) and lowest unoccupied molecular orbital (LUMO) of the electrolyte is accepted as its “stability window” (1.0–4.1 V vs Li+/Li). If the electrochemical potential at the anode (μA) is above the LUMO energy level, the electrolyte will be reduced. Similarly, if the electrochemical potential at the cathode (μC) is below the HOMO energy level, the electrolyte will be oxidized. Therefore, to ensure that the batteries are thermodynamically stable, μA and μC must be within the stability window of the electrolyte. A passivating SEI film can provide an extended window (≈5.0 V) through kinetic protection, allowing the graphite electrode to be used as the anode. However, at high charge rates, especially under extreme conditions, polarizations  on the graphite electrode can increase the propensity of the SEI to break down and drive the anode potential to below the Li+/Li equilibrium potential. These negative effects lead to the formation of dendritic Li, which penetrates the separator and results in the short‐circuiting of the cells. To overcome the disadvantages of graphite anodes and realize a better rate capability and safety characteristic, the development of high‐performance anode candidates is essential.μA > 1.0 V vs Li+/Li), to eliminate electrolyte decomposition and side reactions at a fundamental level. Thus, Li4Ti5O12, TiO2, and other Ti‐based oxides based on the redox couple Ti4+/Ti3+ are emerging as new anode materials for high‐performance LIBs. Among these, Li4Ti5O12 has been investigated extensively for the small volume change that occurs during the electrochemical Li intercalation/deintercalation process (0.2%) and the safe lithiation voltage (≈1.55 V vs Li+/Li). However, the relatively low theoretical specific capacity (175 mAh g−1) and the inherently low electronic conductivity restrict its rate performance. On the other hand, recently reported Wadsley–Roth phase compounds in the Nb2O5‐TiO2 and Nb2O5‐WO3 phase diagrams are some of the most promising families of anode materials for high‐safety LIBs, on the account of their open host structure for rapid Li+ diffusion and uninterrupted band nature for fast electron transport. Moreover, the two‐electron transfer per Nb atom (Nb5+/Nb4+ and Nb4+/Nb3+ redox couples) contributes to a high theoretical capacity and appropriate lithiation potentials in the range of 1.0–2.0 V (vs Li+/Li). Table n), calcuated density, maximum theroretical gravimetric (volumetic) capacity for related Wadsley–Roth phase niobium‐based oxides in this review. It is obvious that the theoretical gravimetric capacities of Wadsley–Roth phase anode materials are much higher than that of Li4Ti5O12 (175 mAh g−1) due to the multi‐electrons transfer characteristic. Importantly, the theoretical volumetric capacity of Li4Ti5O12 is calculated to be only 592.2 mAh cm−3 based on the calculated density of 3.384 g cm−3. However, the theoretical volumetric capacities of these Wadsley–Roth phase compounds are above 1600 mAh cm−3. Therefore, Wadsley–Roth phase anode materials can be powerful substitutes for Li4Ti5O12 in high‐power LIB markets.The key for developing high‐safety anode materials is identifying ideal redox couples with moderate lithiation potential ∞ blocks, which are connected by edge‐sharing type only ∞ ReO3‐type blocks, which are joined within the same a–b plane by corn‐sharing tetrahedra and to blocks offset along the c‐axis by edge‐sharing crystallographic shear planes  Abundant lithium sites in the Wadsley–Roth phases provide considerable active sites for Li+ intercalation and ensure a relatively high theoretical specific capacity.Kocer et al. demonstrated that there are three types of lithium sites in shear structures by density functional theory (DFT) calculations: vertical window positions, horizontal window positions, and pocket positions Figure .28]28[2ii)16W5O55 through bond valence energy landscape calculations  maps and possible Li pathways in the Wadsley–Roth‐phase Nbs Figure .10]1016W5iii)6 octahedron are partially disorderly distributed in the block. Partial cation order in Nb14W3O44 was observed by Cheetham and Allen through neutron diffraction characterizations  belongs to the Wadsley–Roth family of crystallographic shear structures consisting of (3 × 4)1 and (3 × 5)∞ ReO3‐type blocks, denoted as [(3 × 4)1 + (3 × 5)∞]  within the ReO3 blocks and LiO5 (five‐coordinated) square‐pyramidal sites on the block edges intercalation. The evolution of the lattice during the Li+ intercalation process can also be visualized by operando transmission electron microscopy images  analysis, as shown in Figure  H‐Nb2O5 is lithiated via two routes during the Li+ intercalation/extraction processes, in the voltage range of 1–3 V (vs Li+/Li): a two‐phase reaction corresponding to the flat regions of the charge/discharge profiles; and a solid–solution reaction corresponding to the slope regions. During the lithiation process, a slight right shift is observed in the crystalline planes parallel to the b‐axis, and an obvious left shift is observed in the crystalline planes of (110), (−311), (−511), (−215), and (−616). The expansion of the structure along the b‐axis and the minor crystal contraction along both the a and c planes are corroborated by the results.Monoclinic H‐Nb] Figure 3a. The b5 Figure  shows thmmodated.30 Hows Figure .[31]s Figure . The lat2O5, there are also other polymorphs of Nb2O5 such as T‐Nb2O5, TT‐Nb2O5, and B‐Nb2O5, which do not belong to Wadsley–Roth phases. As shown in Figure 2O5 is similar to that of tungsten bronzes but composed of highly distorted NbO6 octahedral and NbO7 pentagonal bipyramidal rather than regular octahedra. TT‐Nb2O5 can be described as a disordering modification of T‐Nb2O5. The B‐Nb2O5 phase displays a TiO2(B)‐like structure. Among these Nb2O5 polymorphs, T‐Nb2O5 usually exhibits superior rate capability, owing to the fast Li+ intercalation pseudocapacitance and high Li+ diffusion rate in its room‐and‐pillar layered structure.In addition to H‐Nbctahedra.33 TT‐3.22O7 with a space group of C2/m is built from (3 × 3)∞ blocks along the b‐axis  and X‐ray absorption fine structure results show simultaneous and equivalent changes in the charge compensation of Nb and Ti with the increase of the degree of lithiation in LixTiNb2O7 . However, larger energy barriers (0.71–1.00 eV) have been found in the crossblock diffusion, indicating the kinetic hindering of Li+ diffusion perpendicular to the tunnels near the crystallographic shear planes.As shown in analyses.39 They Figure –f. The as Figure . The phe+ diffusion was observed within the tunnels along the b‐axis and between the neighboring channels ∞ blocks composed of corner‐shared NbO6 octahedra that infinitely extend along the b‐axis. A more exact expression of Nb12O29 is (Nb4+)2(Nb5+)10O29, with an average oxidation of 4.83, which means that there are two Nb4+(4d1) per 12 Nb sites. The resulting two unpaired electrons per block and shorter Nb–Nb distances endow Nb12O29 a metallic conductor with high electronic conductivity. According to different stacking forms, two similar Nb12O29 polymorphs of monoclinic A2/m and orthorhombic Cmcm structures are present . Fu et al. prepared monoclinic CrNb11O29 through a thermal method and subsequent low‐temperature calcination. The introduction of conductive Cr3+ (t32ge0g) ions with three free electrons in the 3d orbitals not only improves the conductivity, but also increases the unit‐cell volume, leading to a relatively high reversible capacity (401 mAh g−1). The shift and recovery of characteristic peaks of CrNb11O29 in in situ XRD measurements indicate that the lithiation–delithiation processes in CrNb11O29 are based on a representative intercalation reaction mechanism O6 octahedra. The volume change during the Li+ insertion/extraction is ≈6.5%. The volume expansion is anisotropic and mainly attributed to the expansion of the crystal along the b‐axis (+6.8%). In sharp contrast, the alteration of c is only +0.4%, and a even decreases by −1.1%. Zheng et al. also investigated the structural evolution of FeNb11O29 using in situ XRD characterization  can be classified into four stages corresponding to two different lithiation mechanisms: I) The initial sloping region between 3.0 and 1.7 V is attributed to solid–solution reaction; II) between 1.70 and 1.58 V, a two‐phase transformation was observed during the formation of Li3.29FeNb11O29, induced by Li+ insertion into the 4i(1) and 4i(2) cavity sites; III) the continuous insertion of Li+ into the 4i(3) and 4i(4) cavity sites during the evolution of Li3.29FeNb11O29 into Li8.34FeNb11O2; IV) another solid–solution reaction region assigned to the last sloping region in the voltage range of 1.58–1.0 V resulting from the insertion of Li+ into the 4i(5) and 8j cavity sites during the evolution of Li8.34FeNb11O2 into Li17.4FeNb11O2. Through the Rietveld refinement of XRD data, the lattice volume expansion was estimated to be 6.92%, which is consistent with that estimated via neutron diffraction.Another alternative Nbctahedra.43 Then Figure .[44]2Nb10O29 is also a promising Nb12O29 analogy, and its theoretical capacity can go as high as 396 mAh g−1, based on the 24‐electron transfer. Similar to TiNb2O7, the partial cation order was also observed in Ti2Nb10O29 evidenced by power neutron diffractions, in which the Ti occupancy in the center octahedral sites of the block is 4.5%, while that in the corner octahedral sites is 64.5%. The lithiation mechanism of Ti2Nb10O29 was studied using in situ high‐energy synchrotron X‐ray diffraction . Generally, the transformation from M1 to M4 is irreversible, resulting from the rearrangement of the Ti2Nb10O29 host to tolerate the insertion of Li+. During the delithiation process, the growth of the lattice parameter b is continuously inhibited, while a/c expands slightly, indicating that the Li+ diffusion and accumulation may have occurred along the b‐axis.Tiis 64.5%.35 Then Figure .[45]3.42O5‐WO3 binary‐phase diagram was first studied by Roth and Wadsley in 1965. Recently, several niobium tungsten oxides, such as Nb18W8O66, Nb16W5O55, Nb14W3O44, and Nb12WO33 have been extensively investigated as high‐safety anode materials for LIBs due to their unique inherent framework structures. In this section, we summarize some pioneering and representative studies about the lithiation mechanisms of the Wadsley–Roth phase niobium tungsten oxides.The existence of stable niobium tungsten oxides with the Wadsley–Roth phases in the Nb in 1965.46 Rec3.4.112WO33 with the space group of C2 belongs to Wadsley–Roth phases with (3 × 4)∞ ReO3 structure blocks. Unlike the structures of TiNb2O7 and Nb12O29 analogies in which the metal atoms are highly disordered, the crystalline structure of Nb12WO33 is highly ordered because all W atoms in Nb12WO33 are located at the tetrahedral positions at the junctions, with the remaining Nb atoms located at the octahedral sites  and LiO4 sites, respectively  and horizontal “window” positions (denoted as yellow circles), depending on the direction of the window, with the planar block as reference.Monoclinic Nbs Figure 6a. KoçerNb12WO33.47 Sim5 sites  exhibit the lowest site energy, which is energetically more favorable. For the horizontal “window” positions , the Li atom may be located at the top of the plane created by four adjacent oxygen atoms with slightly higher site energies. The remaining vertical “window” positions may be too large for establishing the coordination between the Li and O atoms, and therefore, show the highest site energies. Generally, from the perspective of thermodynamics, the order in which lithium is inserted into the active sites of Nb12WO33 may be LiO5 sites, horizontal “window” LiO4 sites, and vertical “window” LiO4 sites.The simulation result of a single Li atom in different active sites shows that Li atoms located at LiO12WO33 during Li+ insertion can also be computed ) and bond distance. Compared with pristine Nb12WO33, the block height of Li13Nb12WO33 increases from 0.381 to 0.409 nm, whereas the Nb–Nb distance along the shear plane decreases from 0.337 to 0.294 nm. The Li+ accumulation along the b‐axis and the decrease of bond distance along the a and c axes not only lead to the anisotropic evolution of the lattice parameters, but also improve the crystalline symmetry. This improvement is confirmed by the most conspicuous change of Δ(θoct) occurs between x = 5 and x = 11, which exactly corresponds to the region of intense shrinkage observed for lattice parameters a and c.A map of the structural evolution of Nbd Figure . During D Figure . However+ migration path is necessary for understanding the lithiation mechanism, the minimum energy paths and activation barriers (NEB) for Li+ motion in Nb12WO33 are summarized in Figure b‐axis) and ii) diffusion parallel to the block plane . The diffusion parallel to the block plane can be further divided into two types: intrablock and crossblock, respectively. The crossblock diffusion, either by migrating through cavity IV next to the W tetrahedral (H↔H′↔I) or by crossing the shear plane (A↔B), presents the highest energy barriers .The determination of the Lia–c plane, the energy barrier needed for Li+ to hop over from LiO5 pocket sites at the block edges into the block center (B↔C and H↔F′ Ea = 0.44 eV) is relatively higher than that between the cavities within the same block . All Li+ hops along the b‐axis exhibit exceedingly low energy barriers . Based on the above analysis, the Li+ diffusion features in Nb12WO33 can be summarized as follows: i) Li+ hopping between cavities perpendicular to the block plane has the lowest energy barrier and cover distances of 0.019 nm, which is half that in case of the lattice parameter b, revealing the rapid long‐range Li+ conduction along the b‐axis; ii) the Li+ hopping between the two tunnels in the center is much more facile than that between the two tunnels close to the edges; iii) the energetics of Li+ diffusion across the crystallographic shear planes of the contiguous blocks are greatly impeded.For the intrablock diffusion along the 3.4.214W3O44, (4 × 4)∞ blocks on the same plane are connected by tetrahedrally coordinated W atoms. In separated planes, the blocks are joined along the c‐axis by crystallographic shear edge‐sharing. Unlike the ordered Nb11WO33 blocks, the Nb14W3O44 blocks are composed of residual WO4 and all NbO4 octahedra. The statistically distributed Nb and W atoms probably cause disorder among the cations. Recent studies have reported that the distributions of Nb and W atoms in the blocks were not entirely random, and that W atoms may be inclined to reside in block‐central sites within the block.In tetragonal Nb‐sharing.49 Unl14W3O44 with high crystallinity and studied its structural change and redox behavior during cycling using in situ synchrotron high‐energy XRD and in situ XANES  between 3.0 and 1.7 V (vs Li+/Li). In region I2, which corresponds to the voltage range of 1.7–1.3 V, c continues to expand, implying that the phenomenon of Li+ diffusion may be occurring mainly along the c‐axis. The contraction of a can be observed in this region as well, which can be attributed to the electrostatic attraction between the Li and O atoms. Subsequently, a reversal is observed in a as it increases again till the cut‐off voltage of 1.0 V in region I3 is reached, which may be related to the deep reduction of Nb4+/Nb3+ and W5+/W4+ redox couples verified by in situ XANES. O2p may compensate for the additional charges of Nb3+ and W4+ via the highly hybridized Nb‐O and W‐O routes. Thus, the interaction between the O and Li atoms weakened with further lithiation.In a previous study, we synthesized micro‐sized NbS Figure 7a–c.1118W8s Figure . The Li+3.52O5, TiNb2O7, Nb10O29 and its analogies, and Nb‐W oxides, there are also some other Nb‐based oxides with the related Wadsley–Roth phase. TiNb24O62 also belongs to the Wadsley–Roth phase, which is consisted of 3 × 4 blocks split into pairs, described as (3 × 4)2  peak of TiNb24O62 disappears and a new phase is formed at ≈44.13°, which can be assigned to the lithiated TiNb24O62 (Li15.7TiNb24O62) phase. And the occurrence of phase transformation during cycling was also verified by the galvanostatic intermittent titration technique measurement of TiNb24O62 by Griffith et al. The subsequent intercalation of Li+ ions further causes the shift of this peak to a lower θ. Finally, TiNb24O62 is fully lithiated to Li28TiNb24O62 with a volume expansion of 17.5%. Upon charging, all the reflection peaks including the (020) peak of TiNb24O62 recover to the initial positions. It should be noticed that the recent study about the entropy stabilization of TiO2‐Nb2O5 Wadsley–Roth shear phases indicated that TiNb24O62 may convert into energetically more stable phases during repeated cycling in batteries, which will affect the long‐term cycling stability.In addition to the extensively studied H‐Nb2 Figure 9a.[9D Figure .[50]9O25 with the space group of I4/m belongs to Wadsley–Roth phases with (3 × 3)∞ ReO3 structure blocks. As shown in Figure 9O25 is built up from nine NbO6 octahedra corner‐share to form a 3 × 3 block in the ab plane, and these blocks are held together by corner‐sharing PO4 tetrahedra. It is interesting that the P atoms in PO4 tetrahedra can also be substituted by redox‐active V to obtain VNb9O25 with a similar crystal structure. The lithiation mechanism of PNb9O25 is a combination of three solid–solution regions  and two biphasic regions (two plateaus around 1.62 and 1.11 V), determined by in situ XRD  the transfer of electrons and migration of de‐solvated Li+ across the interface; and iii) diffusion of Li+ into the bulk structure of the anodes. Because of the absence of an SEI film on the surface of the Wadsley–Roth phase anodes, the de‐solvation active energies of Li+ should be essentially determined by the nature of the electrolyte solvent and Li salts. Therefore, the enhancement electrochemical kinetics of Wadsley–Roth phase anode materials should mainly focus on the manipulation of (ii) and (iii). We analyzed the corresponding works published from 2017 to the most recent ones and present the following conclusions.Although Wadsley–Roth crystallographic shear structure anode materials have shown great potential for high‐rate Lis Figure 10: i) L4.1+ diffusion in a host material (process iii) is not only determined by the intrinsic Li+ diffusion energy barriers but also by the diffusion length in the material, which can be represented as: τ=L2DLi . Therefore, the reduction of particle size L should reduce the diffusion time energetically up to a certain extent. Furthermore, reduced particle size can also help alleviate the strain arising from volume variations, because the toughness and adhesion effects within the grain boundaries may increase at the nanoscale, which is favorable to the stabilization of the crystalline host during charge–discharge.To enhance the high‐rate cycling performance of Wadsley–Roth phase anode materials, researchers have explored methods such as reducing the size of the anode and constructing the anode with a porous structure. Li14W3O44 nanowires with a diameter of ≈400 nm by using the electrospinning method with stoichiometric amounts of ammonium metatungstate hydrate, niobium oxalate, nitric acid, and poly(vinyl pyrrolidone) (PVP) dissolved in a mixture of ethanol and deionized water as the electrospinning solution. The relatively high apparent Li+ diffusion coefficient of Nb14W3O44 nanowires (8.02 × 10−13 cm2 s−1), which is comparable with that of Li4Ti5O12, contributed to the high initial charge/discharge capacities of 215.9 and 264.7 mAh g−1, respectively. However, the cycling current density was limited to only 0.8 A g−1. AlNb11O29, TiNb24O62, Nb12WO33, Ti2Nb10O29, and TiNb2O7 nanowires have also been synthesized by the electrospinning method, and their improved electrochemical performance compared with that of bulk materials has clearly demonstrated the influence of the nanowire architecture on the Li+ diffusion rate.The design of 1D nanostructure is a commonly investigated route to engineer Wadsley–Roth phase anode materials. One such design method, the electrospinning method, is used for the fabrication of nanofibers, nanowires, and nanotubes. Yan et al. fabricated Nbsolution.57 The A representative sample is the 1D TiNb2O7 nanofibers synthesized by the one‐step dual‐nozzle coaxial co‐electrospinning method. After calcination, the inner PVP and polyacrylonitrile polymer were removed, and hollow nanofibers with a wall thickness and outer thickness of ≈600 and 200 nm, respectively, were formed by connecting the nanosized TiNb2O7 particles along the orientation of the fiber. The hollow TiNb2O7 nanofibers exhibited smaller redox polarizations and better cycling stabilities than the TiNb2O7 nanofibers.By further modifying the electrospinning parameters, hollow nanofibers and hollow nanotubes with larger surface areas can be also obtained.60 A r Lou et al. synthesized porous Ti2Nb10O29 nanospheres with a diameter of ≈300–500 nm through a facile one‐pot solvothermal method. The unique porous hierarchical architecture of the anode demonstrably enhanced the pseudocapacitive‐based charge storage, and a high capacity of 141.5 mAh g−1 can be retained after 1000 cycles at a high current density of 10 C. Ti2Nb10O29 with a crosslinked hierarchical porous structure was reportedly obtained using a versatile bioinspired solvothermal‐template method that uses biomass‐derived carbon as the sacrificial templates. A considerable capacity of 192 mAh g−1 can be maintained after 300 cycles at 1 C in the solid electrolyte.In addition to the aforementioned 1D nanostructures, porous structures provide another alternative for the structural modification of Wadsley–Roth phase anode materials without sacrificing much volumetric energy density.61 Lou4.2 graphene, carbon nanotubes (CNTs), and carbon nanofibers (CNFs), to form a hybrid structure has been proved to be an effective strategy for overcoming the intrinsically poor electronic conductivity of these materials. In particular, carbon modification and structural engineering approaches can be adopted simultaneously to combine the benefits of the synergetic optimization of processes (ii) and (iii) for fast charge storage.The integration of Wadsley–Roth phase anode materials with carbon‐based material, such as amorphous carbon,64 gra2O7@S‐C core/shell arrays. In this architecture, TiNb2O7 nanoparticles are sandwiched between the VG skeleton and an S‐C shell transformed from an electrodeposited poly layer, ensuring the continuous and percolating electrons/ion transport network. Moreover, the increased active electron density of the S‐doped carbon shell benefits the adsorption performance of Li+ ions, which is expected to improve the electrochemical performance further. Consequently, the VG/TiNb2O7@S‐C electrode exhibits a capacity of 193 mAh g−1 with a capacity retention of 78% after 5000 cycles at 10 C.Shen et al. used solvothermal electrodeposition (ST‐ED), to construct binder‐free vertical graphene (VG)/TiNbl arrays.68 In 2O5 (MSC‐Nb2O5) to obtain N‐C@MSC‐Nb2O5. The N‐doped carbon shell not only promotes the homogeneity of electron transport in the particles, but also eliminates the spatiotemporal asynchronization of the lithiation/delithiation processes caused by the local inhomogeneity of Li+ ions. The improved electrochemical reaction kinetics and reversibility endow the N‐C@MSC‐Nb2O5 with better cycling stability (170 mAh g−1 after 1000 cycles at 2 A g−1) than that of MSC‐Nb2O5 without carbon modification (only 65 mAh g−1). 1D CNFs or CNTs can also be used to form a conductive network for anode modification. Tang et al. combined a combination of electrophoretic deposition and solvothermal methods to grow hierarchical CNFs@Ti2Nb10O29 arrays on carbon cloth (CNFs@Ti2Nb10O29/CC). Ti2Nb10O29 particles with a size of 20–50 nm were tightly anchored to the intertwined and interconnected CNFs to form a whole conductive network. Moreover, the Li+ diffusion coefficients of CNFs@Ti2Nb10O29/CC can reach as high as 2.59 × 10−11 cm2 s−1, which is approximately two orders of magnitude larger than that of bulk Ti2Nb10O29 (1.45 × 10−12 cm2 s−1).In the work of Song et al., an amorphous N‐doped carbon film is covered on the surface of micrometer‐sized single‐crystal H‐NbSC‐Nb2O5.32 The4.3+ diffusion rate and electronic conductivity. For Wadsley–Roth phase anode materials, the main doping sites are cation sites. According to Yang et al., and more recently Deng et al., Cr3+ doping by substituting Ti4+ in the crystalline structure of Ti2Nb10O29 has been reported to increase the conductivity to 2.80 × 10−4 S cm−1 by forming of impurity bands around the Fermi level, which is an astounding six orders of magnitude higher than that of an undoped one (≈1.61 × 10−10 S cm−1). A serendipitous advantage of this method is that Cr3+ doping increases the unit‐cell volume and decreases the particle size. The as‐prepared Cr0.6Ti0.8Nb10.6O29 sample, among different doping amounts of Cr3+ showed the best cycling stability with a capacity loss of only 0.01% per cycle over 500 cycles at 10 C. In addition, Cr3+ doping can be also introduced into the crystalline structures of FeNb11O29 and Nb12O29. Moreover, other dopants of Cu2+, Mn2+, and V5+ have been investigated and shown to be advantageous in improving the electrochemical kinetics of Wadsley–Roth phase anodic materials.Ion doping can optimize the electrochemical performance of battery materials by changing the lattice constants and band gap to improve the Lig et al.,70 Cr34.4 Oxygen vacancy, as a type of intrinsic defect, can be infused to modify the crystalline structure and electronic properties of a Wadsley–Roth phase anodic material and substantially improve the Li+ storage performance.Recently, defect engineering has emerged as a powerful technology to impart features such as more active sites, reduced stress, and electrostatic repulsion between interactive atoms, and rapid ion diffusion in insertion‐type electrode materials.73 Oxy Deng et al. proposed the assembly of Ti2Nb10O29 on highly conductive bacterial cellulose carbon (BCC). Oxygen vacancies are generated in the final BCC/Ti2Nb10O29−x after annealing the compound at 500 °C for 1 h in a H2 atmosphere. The H2 reduces Ti4+ and Nb5+ to lower valence states and adjusts the surface chemical nature, this has been confirmed using X‐ray photoelectron spectroscopy, Electron paramagnetic resonance, and XANES spectra. DFT calculations show that the band gap of O2‐vacancy‐modulated Ti2Nb10O29−x is 1.0 eV, which is lower than that of pristine Ti2Nb10O29 (≈1.81 eV). This demonstrates the improvement of electronic conductivity by the enhancement of electron density. The resultant BCC/Ti2Nb10O29−x electrode delivers a high capacity of 253 mAh g−1 at 10 C after 500 cycles. In another study, Deng et al. developed a dual‐functional, low‐temperature carbon‐coating method that decomposes acetylene (C2H2 = 2C + H2) to simultaneously produce an oxygen vacancy in Ti2Nb10O29. In addition to the enhanced electronic conductivity, the Li+ diffusion energy barrier of defective Ti2Nb10O29−x (0.71 eV) is also lower than that of perfect Ti2Nb10O29 (0.91 eV), indicating the easier Li+ migration. The regulation of oxygen defects might even weaken the lattice dynamical anisotropy during the cycling and lead to better structural stability, as demonstrated by the in situ synchrotron high‐energy XRD.Among the many methods to introduce oxygen defects, high‐temperature treatment in a reductive or inert atmosphere is the most commonly used method.73 Den2O7 using the supercritical ethanol treatment at 250 °C for 1 h. The relative content of oxygen vacancies introduced by this strategy is nearly twice of that prepared by Ar atmosphere annealing, leading to the enhanced pseudo‐capacitance effect. Accordingly, the as‐prepared TiNb2O7−x exhibits a high initial discharge capacity of 231 mAh g−1 at 0.1 A g−1 due to the optimized active sites.The use of reductive chemical agents can also cause defects on the surface of metal oxides. Zhang et al. synthesized the oxygen vacancy‐rich TiNb for 1 h.75 The The existence of extended planar intergrowth defects (Wadsley defects) has been proved to be reasonable in Wadsley–Roth phase materials. However, point defects such as oxygen vacancies in Wadsley–Roth crystallographic shear structures are largely absent due to the oxygen‐vacancy‐compensating nature of Wadsley defects. Therefore, the practical function of the oxygen vacancy modification strategy should be further examined carefully.Although many previous studies have demonstrated that oxygen vacancies can improve the electrochemical performance of Wadsley–Roth phase anode materials, there is also disagreement in the literature. A very recent perspective was put forth by Griffith et al. showing that the oxygen defect theory has more counter‐evidence than supporting‐evidence and may be oversimplified.76 The4.5Table Through the above summarized modification strategies of structural engineering, carbon modification, ion doping, and oxygen vacancies, the electrochemical performance of Wadsley–Roth phase anode materials has been significantly enhanced. However, developing large‐scale synthesis methods also should be carefully considered for realizing their practical application in LIBs. As shown in 2O5  is higher than TiO2 . However, the synthesis of Li4Ti5O12 requires the addition of lithium compounds such as Li2CO3 . Thus, the cost of raw materials for producing Li4Ti5O4 may be close to that for producing Nb‐based oxides. The development of LIB technology is always accompanied by trade‐offs; there is no known perfect electrode material. Considering that Wadsley–Roth phase anode materials possess higher volumetric energy densities than Li4TiO5, the slightly increased material cost is accepted. Because the volumetric energy density is a crucial indicator, which is even more important than the gravimetric energy density for most applications.From the perspective of scalable production, the cost and abundance of niobium should be also considered. The price of Nb Therefore, the reserves of niobium can entirely meet the demand for producing potential Wadsley–Roth phase Nb‐based oxides.According to the U.S. Geological Survey, “the estimated global reserves and resources of niobium are large and appear more than sufficient to meet global demand for the foreseeable future, possibly the next 500 years.”79 The54Ti5O12 anodes (such as the “SCiB” launched by Toshiba) have exhibited outstanding rate performance and improved safety properties. Wadsley–Roth phase anode materials not only have lithiation potentials similar to those of Li4Ti5O12, but also have a higher theoretical specific capacity. Therefore, exploring high‐power and high‐safety lithium‐ion full cell systems with Wadsley–Roth phase anodes is feasible.To evaluate the feasibility and commercial perspective of Wadsley–Roth phase anode materials, the electrochemical performance of full cells with such anodes must be tested. As mentioned in the introduction, conventional graphite anodes suffer from poor rate capability and potential Li‐plating risks, and hence, cannot satisfy the increasing requirements for high‐output LIBs. Recently developed LIBs with Lioperties.84 Wad2, LiNixMnyCozO2, LiFePO4, LiMn2O4, and LiNi0.5Mn1.5O4, which show varying electrochemical characteristics. In an ideal full cell design, the capacity of the positive electrode  × mp ) must be matched with that of the negative electrode  × mn ), and the capacity of the cell (Qcell) follows the law: 1/Qcell = 1/Qp + 1/Qn. The value of the average working voltage (cellU) can be further used to calculate the energy density of the cell (cellE), i.e., cell = Qcell×UcellE. In practical LIBs, the N/P (Qn/Qp) ratio should be critically investigated to obtain a balance between cycling stability, safety feature, and cost consideration. In this section, we detail the electrochemical performance and matching principles of Wadsley–Roth phase anode materials‐based full cell assemblies using different cathodes.The major commercial cathode materials for LIBs are LiCoO5.15.1.12 is a first‐generation commercialized cathode material, and is still widely used in portable devices because of its high tap density (3.6–4.2 g cm−3). Its operating voltage is ≈3.9 V (vs Li+/Li), with an available capacity of ≈138 mAh g−1 (corresponding to the extraction of 0.5 Li+). Li et al. studied the electrochemical performance of a full cell constructed using ZrNb14O37 nanowires and commercial LiCoO2. The as‐fabricated ZrNb14O37//LiCoO2 cell had an average voltage output of ≈2.17 V, and a reversible capacity of 84 mAh g−1 . The cell retains 81.3% of capacity after 100 charge/discharge cycles at 0.1 A g−1.LiCoOl LiCoO2.82 The1−x−yCoyMnxO2) consist of closely packed oxygen atoms with alternating layers of transition metals and Li in the interstitial sites. The main advantages of LiNi1−x−yCoyMnxO2 are its increased capacity, low toxicity, and reduced cost due to the replacement of expensive Co with inexpensive Ni and Mn. The available capacity of LiNi1−x−yCoyMnxO2 varies from 155 to 220 mAh g−1 according to the composition. Takami and co‐workers, of the Toshiba Corporation, fabricated large‐size TiNb2O7@C//LiNi0.6Co0.2Mn0.2O2 batteries with a capacity of 49 Ah. The cell exhibited a high volumetric energy density of 350 Wh L−1, ultralong cycle life (capacity retention after 7000 cycles at 1 C is as high as 86%), and excellent fast‐charging characteristics (the time to reach 90% SoC is less than 6 min at 10 C), which should make a viable prototype for commercial EV applications.Layered Ni–Co–Mn oxides  is slightly higher than that of LiCoO2, but its available capacity is much lower (≈110 mAh g−1). Zhu et al. synthesized porous Nb12MoO33 microspheres with diameters in the range of 1.5–3.0 µm and assembled as an anode into a full battery with a commercial LiMn2O4 cathode. The Nb12MoO33//LiMn2O4 cell exhibits a reversible capacity  of 234 mAh g−1 and a relatively high working voltage of ≈2.58 V at 0.1 C. In addition, a high capacity retention of 93.9% was achieved for the cell up to 1000 cycles at 5 C. Interestingly, Lakhnot et al. proposed a new type of aqueous LIB with niobium tungsten oxides as the anode, LiMn2O4 as the cathode, and 21 m LiTFSI (lithium bis(trifluoromethane sulfonyl)imide) aqueous solution as the water‐in‐salt electrolyte. Both the Nb16W5O55//LiMn2O4 and Nb18W8O93//LiMn2O4 cells can be cycled stably in the voltage window of 1.0–2.5 V. After 100 cycles at 1 C, the gravimetric capacity of Nb18W8O93 (Cn) in the aqueous lithium‐ion full battery remains at 69 mAh g−1, showing a capacity retention of 88%, which is better than that of Nb16W5O55 (48 mAh g−1 with a capacity retention of 85%). Notably, the volumetric energy density of the Nb18W8O93 electrode can reach as high as 200 Ah L−1 due to the high tap density (2.21–2.97 g mL−1), which is superior to that of the majority anode materials in aqueous systems.The low‐cost and eco‐friendly LiMno LiCoO2.87 The5.24 is a cathode material that crystallizes into an olivine structure. Its theoretical capacity based on one‐electron transfer is 170 mAh g−1, with a stable potential plateau of 3.5 V (vs Li+/Li). The differential scanning calorimetry measurements of overcharged layered, spinel, and olivine cathodes demonstrate the superior thermal stability of LiFePO4. Therefore, matching the LiFePO4 cathode with Wadsley–Roth phase anodes may prevent the risk of thermal runaway by a substantial margin. Deng et al. demonstrated a full cell based on the VG/Ti2Nb10O29/HMB  core/shell sandwich arrays anode and a LiFePO4 cathode. The full cell exhibited a high discharge capacity (Cn) of 314 mAh g−1 at 0.5 C, and about 82% capacity retention after 200 cycles at 5 C.LiFePO Li+/Li).89 The+ storage performance of NC‐TiNb12O29@TiC‐C anodes combined with LiFePO4 cathodes (TiNb12O29 nanoparticles are sandwiched between an N‐doped C shell and a TiC/C nanowire backbone). The operating voltage of the TiNb12O29@TiC‐C//LiFePO4 cell is close to about 1.7 V. It retains a capacity of 125 mAh g−1 (Cn) at a high current density of 10 C after 5000 cycles. Based on the total mass of the anode and the cathode, the cell delivers energy densities of 177 and 85 Wh kg−1 at power densities of 195 and 2000 W kg−1, respectively. The same group of authors also reported the use of Ti2Nb10O29@TiC/C‐NC (N‐doped Ti2Nb10O29 is loaded into a TiC/C‐NC core‐branch skeleton) as the anode. In addition to the Ti2Nb10O29‐based anodes, TiNb2O7 anodes compatible with the LiFePO4 cathode have also been investigated. Apart from the aforementioned Ti2Nb10O29‐based anodes, TiNb2O7 anodes have also been investigated to match with LiFePO4 cathode. The as‐fabricated VG/TiNb2O7@S‐C//LiFePO4 full battery with an operating voltage of ≈1.65 V can maintain a capacity (Cn) of 230 mAh g−1 after 1000 cycles at 10 C, corresponding to a capacity retention of 95.5%. In a previous study, we proposed a Nb14W3O44//LiFePO4 full cell configuration, and studied its electrochemical performance at an evaluated temperature of 50 °C. Under such an extreme condition, the discharge capacity of the Nb14W3O44//LiFePO4 full cell retained 100.5 mAh g−1 after 1000 cycles at 10 C. A high energy density of 143.8 Wh kg−1 is achieved in this cell configuration.Yao et al. studied the Liackbone).81 The5.30.5Mn1.5O4 is a promising high‐voltage cathode material with a flat plateau of ≈4.7 V (vs Li+/Li) that can deliver a large capacity of 147 mAh g−1 based on the two‐electron reaction of Ni2+/Ni4+ redox couples, in which the residual Mn4+ ions are left electrochemically inert. Benefiting from the high discharge potential, the relatively high working voltage and power density are predictable while matching with Wadsley–Roth phase anode materials to construct full cells. Lin et al. fabricated an AlNb11O29//LiNi0.5Mn1.5O4 full cell that exhibited a high operating voltage of ≈3.05 V and a reversible capacity (Cn) of 136 mAh g−1, with a high capacity retention of 93.2% after 100 cycles at 1 C. They also comprehensively evaluated the electrochemical performance of Zn2Nb34O87 nanofibers//LiNi0.5Mn1.5O4 full cell. It exhibited a large capacity (Cn) of 166 mAh g−1 at 1 C, and retained 92.3% of the initial capacity after 100 cycles. Even better cycling stability with only 0.0035% decay per cycle was obtained at 5 C for up to 1000 cycles. Han and Goodenough reported a TiNb2O7 anode combined with a LiNi0.5Mn1.5O4 cathode to assemble a 3‐V class TNO//LiNi0.5Mn1.5O4 full cell and studied the capacity matchup principles. The results indicated that the cycling stability of the anode‐limited battery is obviously better than that of the cathode‐limited one. Liang et al. fabricated a 18 650‐type full cell with a capacity of 1.2 Ah, by using a Mg‐doped LiNi0.5Mn1.5O4 cathode coupled with a TiNb2O7 anode. A flat operating voltage at ≈3.0 V and a minor polarization effect were observed in the TiNb2O7//Mg0.1‐LiNi0.5Mn1.5O4 cell. A capacity retention of ≈96.3% was achieved after 68 cycles at 0.1 C, which corresponds to a gravimetric energy density (Ecell) of 91.4 Wh kg−1.LiNily inert.91 BenTable + ions, which will reduce the crystal stability of the host structures. 2) The TiNb2O7//LiNi0.6Co0.2Mn0.2O2 full cell shows the longest lifespan among all these battery systems. 3) The TiNb2O7//LiNi0.5Mn1.5O4 full cell exhibits the highest operating voltage and energy density. From the perspective of energy utilization efficiency, LiNi0.5Mn1.5O4 may be a more appropriate cathode to match with Wadsley–Roth phase anodes. However, the conventional carbonate‐based electrolyte is easy to decompose and many adverse side reactions may occur between the cathode/electrolyte interface due to the high working voltage of LiNi0.5Mn1.5O4, which will exert negative influence on the cycling stability and safety performance.6The unique features of Wadsley–Roth phase niobium‐based oxides, including 3D open crystalline structure, moderate lithiation potential, and reversible redox couples, can usher in a new era of LIBs with high power density, long lifespan, and high safety. Despite this, there is still extensive advancement space for further enhancing their electrochemical kinetics. In this paper, we systematically analyzed structural features and electrochemical reaction mechanisms of various representative Wadsley–Roth phase anode materials and then explored the improvement strategies. We also summarized the electrochemical properties of a series of full cells that demonstrate the compatibility of Wadsley–Roth phase anode materials with commercial cathode materials to fabricate LIBs applicable for commercial use.+ storage mechanism, modification of electrode materials, full‐cell‐assembly parameters, safety considerations, and industrial feasibility should be considered in future studies.Yet, several significant aspects, such as LiListed below are the recommendations following the findings of this study:i)Researchers must comprehensively study the lithiation mechanism and electrode–electrolyte interface behavior of Wadsley–Roth phase anode materials using various in situ or operando techniques and theoretical methods. The relationship between the electrochemical performance, intrinsic structure, and physicochemical properties of the anode materials should be established to obtain a guideline for exploiting more precise and valid electrode‐modification strategies.ii)−2, which means that the loading mass of Wadsley–Roth phase anode materials in the electrode should be more than 10 mg cm−2. Most of the existing material modification strategies for improving the rate performance and cycling stability are based on results obtained at much lower mass loading levels (≈1–2 mg cm−2). Whether or not these modification methods are effective at higher loading mass conditions remains unanswered. Besides, the initial Coulombic efficiency is also an important indicator for commercial LIBs, and it should be more than 90% for full cells to realize high energy density. However, this modification approach often causes serious side reactions at the interface, resulting in much lower initial Coulombic efficiency. Therefore, Wadsley–Roth phase anode materials must be optimized to achieve high power density and balance the gravimetric/volumetric energy densities at a high loading mass condition.For commercial applications, the areal specific capacity of anodes must be no less than 3.0 mAh cmiii)The reduction of the electrolyte by the anode and thermal decomposition of the anode at elevated temperatures are the possible reactions that trigger the thermal runaway of batteries. It is plausible that the liberation heat related to the Wadsley–Roth phase anode materials may be apparently suppressed owing to their relatively high lithiation potentials, but the detailed thermal reaction mechanisms between Wadsley–Roth phase anode materials and electrolytes still need to be critically evaluated.iv)There are still no criteria for assembling and testing Wadsley–Roth phase anodes‐based LIB full cells. Parameters, such as cell design (thickness and porosity of electrodes), electrode capacity balance (anode‐limited or cathode‐limited), battery capacity, and abuse tolerance (high‐ or low‐temperature cycling stability), lack any organized criteria. Moreover, the reliable safety features of Wadsley–Roth phase anode‐based LIB full cells must be systematically evaluated using characterizations such as accelerating rate calorimetry, nail penetration test, and overcharge test.v)+.The use of high‐safety Wadsley–Roth phase anode materials in solid‐state electrolytes may be the ultimate route to address the safety issues pertaining to the carbonate‐solvent‐based LIBs. In addition, these anode materials can also be applied in well‐designed, noninflammable, aqueous electrolytes due to the moderate lithiation potentials that prevent the reduction of HThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


To this end, we determine all duality threshold parameters from a Bayesian fit for the first time. Our results, obtained at small momentum transfer q2, are extrapolated to large q2 where they agree with precise lattice QCD results. We find that a modification to the commonly used BCL parametrization is crucial to interpolate the scalar form factor between the two q2 regions. We provide numerical results for the form factor parameters — including their covariance — based on simultaneous fit of all three form factors to both the sum rule and lattice QCD results. Our predictions for the form factors agree well with measurements of the q2 spectrum of the semileptonic decay Vub| = (3.77 ± 0.15) · 10−3, which is in agreement with the most recent inclusive determination at the 1 σ level.We revisit light-cone sum rules with pion distribution amplitudes to determine the full set of local The online version contains supplementary material available at 10.1007/JHEP07(2021)036. ESM 1(TGZ 3 kb)


In this study, hypercrosslinked polymer (HCP) photocatalysts have been developed for the selective reduction of CO2 to CO, combining excellent CO2 sorption capacities, good general stabilities, and low production costs. HCPs are active photocatalysts in the visible light range, significantly outperforming the benchmark material, TiO2 P25, using only sacrificial H2O. It is hypothesized that superior H2O adsorption capacities facilitate access to photoactive sites, improving photocatalytic conversion rates when compared to sacrificial H2. These polymers are an intriguing set of organic photocatalysts, displaying no long‐range order or extended π‐conjugation. The as‐synthesized networks are the sole photocatalytic component, requiring no added cocatalyst doping or photosensitizer, representing a highly versatile and exciting platform for solar‐energy conversion.The design of robust, high‐performance photocatalysts is key for the success of solar fuel production by CO This is how it fuels: Hypercrosslinked polymers are an intriguing set of organic photocatalysts. They seemingly display no extended π‐conjugation and are able to photoreduce gaseous CO2 to CO using H2O under visible light. This conversion can be done by a variety of processes, including photochemical , thermochemical, and electrochemical reactions. However, overcoming the high thermodynamic and kinetic barriers to conversion is challenging, and so a catalyst is required to improve energy efficiency and, ultimately, render these processes viable.[The ever‐increasing global energy demand requires a significant overhaul of current production processes if humanity is to address climate change. Carbon management and renewable energy must play a key role in our energy outlook, challenging researchers to reshape our energy portfolio. Research efforts are focused on the development of efficient carbon capture, utilization, and storage (CCUS) technologies, as well as the improvement of methods to harness renewable energy.s viable.2, 32, CdS, ZnO, WO3, Ru‐, Re‐ and Pd‐based complexes have received much attention as photocatalysts, owing to their ability to generate charge carriers under light irradiation.[Herein, we focus on a photochemical route to solar fuel production, namely photocatalysis, whose main advantage lies in the simplicity of its implementation. To date, ‘traditional’ semiconductors, such as metal oxides or sulfides, and transition metal complexes, such as TiOadiation.6, 7 H2 photoconversion to CO using visible light.2 to CO with 92 % selectivity, using visible light and sacrificial H2O.2 photoreduction to CO with a 97.8 % selectivity.2 photoreduction rates in the presence of acetonitrile and sacrificial triethanolamine.2 photoreduction in the presence of water.[2 photoreduction, the synthesis of these photocatalysts generally required the use of rare‐earth metals, or specifically polymerizable monomeric units, presenting implementation barriers due to relatively high‐costs and poor sustainability.The development of new classes of photoactive materials, including inorganic–organic hybrids, such as metal–organic frameworks (MOFs), or organic‐based materials, such as porous organic polymers, have emerged as promising alternatives to traditional photocatalysts.11, 122‐graphene composite for the photoreduction of CO2, with the HCP component aiding CO2 adsorption and diffusion.Hypercrosslinked polymers (HCPs) represent a class of materials with excellent tunability and relatively low cost. HCPs are densely crosslinked amorphous networks that are produced by using simple Friedel–Crafts chemistry. Nonfunctional aromatic compounds  can be ‘knitted’ together using an external crosslinker, requiring only iron(III) chloride as catalyst.2 to CO. Photocatalytic conversion was achieved using only visible light in the presence of sacrificial H2O, without additional sacrificial agents or cocatalysts, significantly outperforming TiO2 P25 (7.5‐fold improvement). The influence of the reducing agent was investigated . Surprisingly, employing sacrificial H2O led to significant improvements in photoconversion rates. We hypothesize that the preferential adsorption of H2O concentrates the sacrificial agent at the HCP's surface, driving photocatalytic performance. Owing to their lack of requirement for precious‐metal catalysts, as well as their easily scaled chemistry, HCPs present an exciting platform for the further design and discovery of high‐performance organic photocatalysts.Herein, we present HCPs as a new class of photocatalyst capable of selectively reducing CO2 adsorption selectivity through the inclusion of amino groups. Photographic images of the networks  reveal a color gradient from dark to light brown from HCP‐1 to HCP‐3. We successfully incorporated the aromatic monomers into the networks, as confirmed using a variety of characterization techniques. Solid‐state cross‐polarization magic angle spinning 13C NMR  of all polymers showed signals at around 139 and 129 ppm corresponding to quaternary aromatic carbons (CAr) and aromatic CAr−H.We synthesized three HCPs of varied chemical structure – HCP‐1, HCP‐2, and HCP‐3 – by a Friedel–Crafts alkylation reaction using external crosslinkers to ‘knit’ together aromatic monomers. A general reaction scheme and representative HCP structures are shown in Figure 3 and CARC−O, respectively. The signal at 74 ppm is absent in HCP‐3, but a new signal emerges at around 102 ppm, suggesting this signal may be shifted downfield in this case. For HCP‐3, a signal at 166 ppm is assigned to triazine carbons and in HCP‐2 the C‐NH2 signal is expected at around 146 ppm and is therefore likely eclipsed by the substituted aromatic peak. Fourier‐transfer infrared spectroscopy (FTIR) showed distinct stretches for primary amines in HCP‐2 and HCP‐3, and the triazine tertiary amines in HCP‐3 . HCP‐2 and HCP‐3 contain 1.3 and 7.1 at% N, respectively, as determined by elemental analysis (Table S1), which may explain the lack of a signal in ssNMR for C‐NH2 in HCP‐2. These values correspond to final polymer compositions of roughly 1 : 4 and 1 : 3 monomer to crosslinker ratios for HCP‐2 and HCP‐3, respectively. Scanning electron micrographs revealed HCP‐1 and HCP‐3 as agglomerated spherical particles, whereas HCP‐2 shows a more fibrous structure . Thermogravimetric analysis demonstrated the high thermal stability of all HCPs, with decomposition temperatures of >300 °C in both N2 and air atmospheres . Char yields in N2 at 900 °C were >60 % in all HCPs, while complete degradation was observed in air at >550 °C. Powder X‐ray diffraction confirmed the amorphous nature of all HCP networks .Signals at around 36 ppm are assigned to newly formed methylene bridges, whereas weak signals at around 54 and 74 ppm correspond to unreacted methyl ethers O−CH2 sorption measurements at −196 °C to assess the porous nature of the networks. HCP‐1, HCP‐2, and HCP‐3 exhibited BET surface areas of 951, 311, and 357 m2/g, respectively , dropping to 0.13 and 0.16 cm3/g for HCP‐2 and HCP‐3, respectively, reflecting trends seen in the polymer's BET surface areas. Figure 2 uptake ability of HCPs, we collected adsorption isotherms at 25 °C up to 1 bar . Although the HCPs followed the expected trend, whereby higher surface area polymers adsorbed more CO2, the CO2 capacities did not reflect the large differences in surface areas (Table S2). The presence of the amino groups in HCP‐2 and HCP‐3, which are known to impart CO2 selectivity to hypercrosslinked polymers,2.[We used Nlymers,2.29, 302 uptake in polar HCPs, due to competitive adsorption.2O and CO2 co‐adsorption, since we used sacrificial H2O vapor in CO2 photoreduction (see below). We exposed the samples to humid air (>99 % humidity) for at least 48 h before collecting CO2 adsorption isotherms at 25 °C up to 1 bar. Crucially, samples were not degassed prior to CO2 adsorption measurements . The CO2 sorption capacities at 1 bar are given in Figure 2 capacity in comparison to its dry equivalent, while the amine‐containing HCP‐2 and HCP‐3 showed negligible difference in CO2 uptake between the wet and dry networks. These uptake capacities demonstrate that CO2 can still adsorb in the presence of water. To further investigate the impact of water, we measured the catalytic ability of “wet” HCPs  in a pure CO2 stream. The HCPs still displayed CO2 photoreduction activity (Table S3), albeit decreased, which indicates that CO2 could still adsorb onto the materials and react. We note that during the ‘normal’ photoreduction experiments, we did not expose HCPs to water vapor prior to CO2 reduction, but rather a stream of CO2 containing water vapor.The presence of adsorbed water was shown to impede CO2 photoreduction, we investigated their optoelectronic properties using UV/Vis diffuse reflectance spectroscopy (UV/Vis DRS). The UV/Vis spectra show all HCPs absorb light in both the UV and visible range, with estimated optical gaps of 3.56, 3.54 and 3.19 eV for HCP‐1, HCP‐2 and HCP‐3, respectively  and the VB onset, also known as VB offset . Placing the VB offset on the absolute energy scale requires knowledge of the position of the Fermi level, which was determined by measuring the secondary electron cut‐off through XPS work function measurements . Finally, we identified the CB position using the UV/Vis spectra absorption onset. For all HCPs, the CB is located above the reduction potential of CO2/CO and the VB below the oxidation potential of both H+/H2 and O2/H2O. Hence, the band diagrams point towards a sufficient thermodynamic driving force to enable the reduction of CO2 to CO using either gaseous H2 or H2O as a sacrificial agent. The Fermi level of all HCPs lies closer to the CB than the VB, suggesting electrons are the majority charge carriers, which is desirable for CO2 photoreduction. Overall, the light absorption properties of HCPs, as well as their CO2 adsorption capacities, make them a priori attractive candidates for CO2 photoreduction.The time at which the photoluminescence signal has decayed to half its initial amplitude  was 1.9 ns, 2.2 ns, and 3.2 ns for HCP‐1, HCP‐2, and HCP‐3, respectively, which demonstrates that HCP‐3 has a substantially longer excited state lifetime than HCP‐1 or HCP‐2. We estimated the HCP's valence band (VB) and conduction band (CB) positions by complementing our UV/Vis DRS data with XPS measurements . No cocatalyst or photosensitizer was required, but some residual iron is present from the HCP synthesis and may play a role in the overall photocatalytic activity (Table S1) as shown with Pd for other organic materials in the context of H2 evolution.[2 capture and CO2 conversion, avoiding limitations owing to poor CO2 solubility in water . In each experimental set, we compared HCP performance to that of the benchmark TiO2 P25. For all HCPs, after 3 h of irradiation the primary carbonaceous product observed was CO, with a gaseous product selectivity of up to 93 % and 98 % using sacrificial H2 or H2O, respectively . Quantum efficiencies were calculated for HCP‐3 in both UV/Vis and visible light alone and compared to a selection of competitive organic photocatalysts for gaseous CO2 reduction (Table S5 and S6). We note that this is intended to provide some perspective rather than a direct comparison. To investigate the production of less volatile compounds remaining on the polymer surface, we washed HCP‐3 with water post‐irradiation and analyzed the resulting solution using HPLC. Low concentrations of methanol and formic acid were also detected , the buildup of which at the surface could lead to decreased catalyst activity with time. Interestingly, when using H2O as sacrificial agent, we did not detect O2. Instead, we observed H2O2 after photocatalysis . The formation of H2O2 may originate from two possible routes. H2O might be oxidized first to O2 and then reduced to H2O2 on the surface of the catalyst, resulting in a lack of detectable O2. This route has been reported for a different study.2O might be directly oxidized to H2O2.After establishing their COvolution.32, 332 P25 under UV/Vis light and was up to 7.5 times better when irradiated with only visible light , likely increasing its H2O adsorption ability. Adding to this effect, the high porosity and particularly microporosity will also contribute to the water uptake of the HCPs, thereby explaining the trend in water sorption. A higher concentration of sacrificial agent at the HCP's surface should favor high CO2 photoreduction rates, explaining the increase of photoactivity when using H2O as reducing agent.Interestingly, COs Figure . From th2O adsorption on the photoactivity, we decreased the humidity inside the photoreactor by decreasing the temperature of the H2O vapor saturator to around 1 °C. At 1 °C, the partial pressure of H2O at saturation is reduced (6.6 mbar), thereby dramatically reducing the concentration of H2O introduced into the system. After sealing the reactor, we conducted CO2 photoreduction measurements at 25 °C, under atmospheric pressure. The decreased humidity of the stream resulted in a 38 % decrease in activity for HCP‐3 . Indeed, H2O uptake in this network is reduced to around 0.4 mmol/g at 6.6 mbar, as estimated from the pure water sorption isotherm . Recyclability tests were also performed. After 5 cycles the photoactivity decreased by 9 % Figure f, which 3 aliphatic crosslinker and the amount of partial crosslinked C−O slightly decreased after irradiation . Elemental analysis showed minimal change in bulk polymer composition after the irradiation process (Table S1). We also collected ssNMR spectra for all networks after the photoreduction process , which also showed no major structural change to bulk composition. We then probed the structural and chemical stability of HCP‐3 further by collecting N2, CO2, and H2O sorption isotherms and FTIR spectra, all of which showed no significant alteration, with the exception of a change in the shape of the N2 isotherm . The slight degradation of the polymer seen in XPS measurements post‐irradiation may arise from the formation of H2O2 which decomposes into hydroxyl radicals under UV/Vis irradiation and leads to structural changes.To assess the possible degradation, we collected XPS spectra of all polymers before and after irradiation. Deconvolution of C1s, N1s and O1s spectra for all HCPs showed that the amount of sp2 conversion over HCPs photocatalysts, we conducted a series of control experiments, in an inert atmosphere (N2/H2 or N2/H2O), without catalyst, without light, and with isotopic labeling of 13CO2 (details of control experiments are given in Table S3). In the absence of CO2, the activity decreased by 77 % to 88 % depending on the atmosphere (N2/H2O vs N2/H2). We attribute the trace CO detected under inert atmosphere to the degradation of some carbon sp3 aliphatic crosslinker and residual oxygen‐containing functional groups of the HCP crosslinkers. To verify the photocatalytic production of CO from CO2, we conducted isotopic labeled 13CO2 tests using H2O as a sacrificial agent under UV/Visible light irradiation. We observe a 13CO peak (m/z=29) after light irradiation, confirming the ability of HCPs to photoconvert CO2 to CO . No 13CH4 was observed, which is likely due to the good selectivity of the HCP toward CO  for the first time as a photocatalytic platform for CO2 P25 , benzene, dimethoxymethane, 2,4‐diamino‐6‐phenyl‐1,3,5‐triazine, aniline, 1,2‐dichloroethane, N,N‐dimethylformamide and iron(III) chloride were all purchased from Sigma‐Aldrich. 4,4′‐bis(methoxymethyl)biphenyl was purchased from Tokyo Chemical Industry (TCI) UK and methanol (reagent grade) was purchased from VWR. All reagents were used as received. All gases were purchased from BOC.TiO2, before the solution was purged with N2 for at least a further 30 min. After purging, iron(III) chloride  was quickly added to the solution and the mixture was heated to 80 °C for a further 24 h, during which the reaction was kept under an inert atmosphere. The resulting solid was allowed to cool before it was filtrated using a Büchner funnel and washed with methanol until the filtrate was almost colorless. The polymer was then further washed by Soxhlet extraction in methanol for 24 h. Finally, the polymer was dried in a vacuum oven overnight at 70 °C.HCP‐1: Anhydrous 1,2‐dichloroethane (20 mL) was added to benzene  and dimethoxymethane  under NHCP‐2: The overall procedure remained the same as HCP‐1 with 1,2‐dichloroethane (20 mL) added to aniline  and 4,4′‐bis(methoxymethyl)biphenyl  before iron(III) chloride  was added.HCP‐3: Again, the overall procedure remained the same as HCP‐1 with 1,2‐dichloroethane (20 mL) added to 2,4‐diamino‐6‐phenyl‐1,3,5‐triazine  and 4,4′‐bis(methoxymethyl)biphenyl  before iron(III) chloride  was added. In addition to methanol, HCP‐3 was also washed in chloroform by Soxhlet extraction, to ensure the removal of unreacted 2,4‐diamino‐6‐phenyl‐1,3,5‐triazine. Yields for HCP‐1, HCP‐2 and HCP‐3 were 84 %, 66 % and 91 %, respectively, based on hypothetical 100 % polycondensation.All samples were fully characterized and details about the characterization techniques can be found in the Supporting Information.2 photocatalytic measurements . Tests were conducted at ambient temperature. The photocatalysts were deposited on a stainless‐steel plate with a fixed area of 9.6 cm2. To do so, 10–15 mg of the ground photocatalyst was dispersed in DI water (1.2 mL), sonicated for 30 sec and drop cast onto the sample holder. Research grade (99.999 %) CO2 and H2  were flowed at controlled rates using mass flow controllers . For experiments using H2 as sacrificial agent, the photoreactor (35 cm3) was vacuumed and replenished with a gas mixture of CO2 and H2 (1.5 vol/vol ratio) six times. The same gas mixture of CO2 and H2 was subsequently passed over the catalyst bed in the photoreactor for 15 residence times before it was sealed at 1.25 bar and irradiated for 3 h. For experiments using water as sacrificial agent, H2O vapor was generated from a saturator at room temperature (25 °C), unless stated otherwise, providing a gas mixture of CO2 and H2O (1.4 wt% H2O). The photoreactor was vacuumed and replenished with a gas mixture of CO2 and H2O at least six times prior to measurement. A xenon arc lamp  equipped with a water filter was used as the irradiation source. The distance from the lamp to the sample was 9.5 cm with a broadband irradiance at the sample surface of 1830 W m−2. For visible light measurements, a long‐pass UV filter  was used.A gas/solid photoreactor was assembled to conduct CO13CO2 . The reactor was flushed with He for 1 h before injecting 15 mL of 13CO2. After irradiation, the evolved gases were analyzed by a mass spectrometer (Shimadzu MS) equipped with a Q‐bond and a MolSieve column with gas sampling valves connected directly to the photoreactor. The photocatalytic CO2 reduction tests were repeated 3 times for each material under the same reaction conditions.Evolved gases were analyzed by a gas chromatograph (GC) with gas sampling valves connected directly to the photoreactor. The GC  was equipped with HayeSep  and molecular sieve  packed columns in series, a thermal conductivity detector (TCD) and a flame ionization detector (FID). For recyclability tests, the aforementioned process was repeated after each 3 h irradiation cycle without opening the photoreactor. In addition, isotopic tracing experiments were performed with The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.