However, the slow reaction kinetics during discharge/charge seriously limits its development. Hence, a simple cation exchange strategy is developed to introduce Ru atoms onto a Co3O4 nanosheet array grown on carbon cloth (SA Ru‐Co3O4/CC) to prepare a single atom site catalyst (SASC) and successfully used in Li‐CO2 battery. Li‐CO2 batteries based on SA Ru‐Co3O4/CC cathode exhibit enhanced electrochemical performances including low overpotential, ultra high capacity, and long cycle life. Density functional theory calculations reveal that single atom Ru as the driving force center can significantly enhance the intrinsic affinity for key intermediates, thus enhancing the reaction kinetics of CO2 reduction reaction in Li‐CO2 batteries, and ultimately optimizing the growth pathway of discharge products. In addition, the Bader charge analysis indicates that Ru atoms as electron‐deficient centers can enhance the catalytic activity of SA Ru‐Co3O4/CC cathode for the CO2 evolution reaction. It is believed that this work has important implications for the development of new SASCs and the design of efficient catalyst for Li‐CO2 batteries.Li‐CO 3O4/CC is successfully applied to Li‐CO2 batteries. The doped single atom Ru sites have very high catalytic activity, which can not only improve the reaction kinetics but also optimize the growth pathway of the discharge products, thus comprehensively improving the electrochemical performance of the Li‐CO2 battery.A facile and eco‐friendly method to prepare single atom site catalysts is proposed, and the derived SA Ru‐Co In this case, attention is turning to technologies that can capture and use CO2, including electrochemical reduction, photoelectric chemical reduction, catalytic hydrogenation, catalytic reforming, metal‐CO2 battery, etc. Among them, Li‐CO2 battery is regarded as one of the main candidates for new energy storage and conversion system due to its high theoretical energy density (1876 Wh kg−1) and the ability to fix CO2. In addition, Li‐CO2 battery is considered to be a more promising priority in some high‐CO2 environments, such as underwater operations and Mars exploration. Though Li‐CO2 battery has been extensively studied and developed since it was first proposed in 2013, there are still many problems restricting its development, among which the main challenge is the hysteretic kinetics of the CO2 reduction reaction/CO2 evolution reaction (CRR/CER) process, which leads to unsatisfactory electrochemical performances including high overpotential, low capacity, pitiful reversible, and cycle stability of Li‐CO2 battery. Specifically, during the CRR process, the adsorption and activation of CO2 is difficult due to its thermodynamic stability. Moreover, the solid products (Li2CO3 and C) will occupy a lot of space and active sites of the cathode catalyst, which limits the CRR kinetics to a large extent. While during the CER process, Li2CO3 is a thermodynamically stable wide‐bandgap insulator, which requires a higher charging potential (>4.3 V vs Li+/Li) to decompose. At such a high potential, the electrolyte and commonly used carbon cathode are also prone to decomposition. To make matters worse, due to adverse contact between discharge products and the active site, Li2CO3 is difficult to be completely decomposed during charging. This results in the accumulation of discharge products on the surface of the cathode, which prevents the mass transfer process and passivates the catalyst, further leading to problems such as high overpotential, sudden failure, and poor cycle stability. Therefore, the development of cathode catalysts with suitable pore structure and high catalytic activity is still a pressing matter in the development of Li‐CO2 batteries. To this end, various materials including carbon based materials, transition‐metal oxides, noble metals, MOFs/COFs and their derivatives with rationally constructing hierarchical porous architectures, have been sprung up as efficient cathode catalysts to improve the performance of Li‐CO2 batteries. Actually, the implementation of these strategies can improve the performance of the battery to a certain extent. However, some unavoidable disadvantages, such as high cost, poor durability, relatively low capacity, poor cycle stability, and especially low utilization rate of active sites still exist. Hence, the development of cathode catalysts with high activity and durability under the premise of cost control is of great significance for Li‐CO2 batteries.Excessive use of fossil fuels has led to a rapid increase in COt urgent.2, 3 Compared with traditional nanoparticle catalyst, SASCs catalyst has many unique characteristics, such as high atomic utilization (AUE) (close to 100%) and maximizing accessible active sites, which can decrease the use of the metal and reduce the cost of the cathode. In addition, SASCs also have unique coordination and electron effects due to the interaction between the metal active sites and the supporter, which make SASCs performed extremely excellent in oxygen reduction reaction (ORR), oxygen evolution reaction (OER), CO2RR, hydrogen reduction reaction (HER), and other fields, and some progress has also been made in the application of SASCs in Li‐CO2 batteries. However, the reported SASCs supporters for Li‐CO2 batteries are carbon which is not stable in the battery system and the performance of battery needs to be further improved. In addition, understanding the deposition/decomposition process of the discharge products optimized by SASCs is an important role in determining the kinetics of the CRR and CER process; however, there are few studies in this aspect, especially in CER process. Most importantly, majority of reported SASCs synthesis processes tend to require harsh conditions and tedious process, which not only increase the preparation cost of SASCs but also bring great difficulties to their large‐scale synthesis and application. Therefore, it is of great significance to develop a simple and mild method to prepare SASCs for improving the electrochemical performance of Li‐CO2 batteries.Recently reported single‐atom site catalysts (SASCs) as a hot research direction in the catalytic field has aroused wide attention and is even regarded as the “holy grail” of catalysts.26 Com3O4 nanosheets surface supported on carbon cloth (marked as SA Ru‐Co3O4/CC) by combining template replication and cation exchange strategy, which is successfully used as a dual‐catalyst for Li‐CO2 batteries. It is worth mentioning that the synthesis method only uses water as the solvent, and calcination conditions are easy to reach (350 °C and air atmosphere), which greatly reduces the cost and difficulty of synthesis of SASCs. Experimental results show that Co3O4 nanosheets array with 3D architecture possess large specific surface area and suitable pores to facilitate the mass transfer of electrolyte and reactants, and provide enough space to store the discharge products (Li2CO3 and C). While atomically dispersed Ru species as active sites tailored the nucleation behavior of discharge products, thus optimizing their final distribution, which is favor of subsequent decomposition during charging process. Additionally, density functional theory (DFT) simulation demonstrated that the introduced Ru atoms as the driving force center can significantly enhance the intrinsic affinity of the key intermediates (Li2C2O4), which is beneficial to CRR process. Meanwhile, the calculation of Co/Ru electron charge shows that Ru atoms also act as an electron‐deficient center, which is beneficial to the CER process of the battery. Benefiting from these merits, Li‐CO2 batteries based on SA Ru‐Co3O4/CC cathode exhibit enhanced electrochemical performances, including low overpotential, high specific capacity, and long cycle stability. We believe that this work can not only guide the design of efficient Li‐CO2 batteries but also have significant implications for the design and synthesis of SASCs.In this study, we propose a facile and eco‐friendly method to load Ru atoms onto Co2Figure3O4/CC cathode catalyst. Firstly, to grow Co‐MOF homogenously, the substrate of carbon cloth (CC) was treated by Ar plasma, and after an ingenious subsequent ion exchange and template replication process, the aimed cathode catalyst was obtained. Observing the whole synthesis process, all the solvents (only water) used are eco‐friendly and the preparation process is mild and facile , and hence this method has the potential of large‐scale synthesis and application. The scanning electron microscope (SEM) images show the smooth carbon fiber with diameter of about 10 µm were woven with each other  spectra show that CoRu‐MOF and Co‐MOF have the same peak position, indicating that Ru3+ replaces Co2+, rather than forming new MOF or Ru ions adsorbed on the surface of Co‐MOF, as shown in Figure  it is supposed that this is due to the etching of MOF by H+ derived from the partial hydrolysis of Ru3+. To prove this, dilute hydrochloric acid was used to etch Co‐MOF/CC and a similar morphology to that of CoRu‐MOF was also found as shown in Figure 3O4, while the Ru species in CoRu‐MOF coordinates with O atoms in Co3O4, thus being anchored on the surface of Co3O4 nanosheet. The obtained SA Ru‐Co3O4 presented a slightly curled elliptical morphology  and (220) planes of cubic spinel Co3O4 (JCPDS card No. 42‐1467). To identify the distribution of Ru species on SA Ru‐Co3O4, elemental mapping was performed with energy‐dispersive X‐ray spectrometry (EDS) as shown in Figures 3O4. Figure 3O4/CC, and SA Ru‐Co3O4/CC, only the peaks of CC and Co3O4 were found on the XRD curves of SA Ru‐Co3O4/CC, without any peaks related to Ru species, and this result was confirmed by Raman spectra  images confirmed that the SA Ru‐Co species,37 and3O4/CC samples. The peak‐fitting analysis of Co 2p spectrum shows that there are two chemical states  analysis is used to preliminarily study the surface chemical status of SA Ru‐Cos Figure , corresp795.7 eV.37 The3O4, X‐ray absorption fine structure (XAFS) spectroscopies are conducted to characterize the electron state and coordination configurations of Ru in SA Ru‐Co3O4. Firstly, for identifying the 3D arrangement of Ru atoms at higher sensitivity, X‐ray absorption near‐edge structure (XANES) analysis is applied  of SA Ru‐Co3O4/CC in Figure 3O4/CC is even higher than that in RuO2 (+4). It should be taken into account that Ru atoms are mainly located on the surface of the precursor Co‐MOF nanosheets rather than inside, which means that Ru atoms may be fully oxidized to a higher valence state during the calcination process. And it is worth mentioning that a similar situation has appeared in some similar reports. Then, to verify the local environment of Ru species, the extended XAFS (EXAFS) spectra of Ru K‐edge of SA Ru‐Co3O4 were analyzed. The Fourier‐transformed EXAFS (FT‐EXAFS) curve of the Ru k‐edge of SA Ru‐Co3O4 shows a distinct peak at 1.45 Å , which should be assigned to the Ru—O bond in SA Ru‐Co3O4. Note that, there is a small peak at around 2.3 Å, which is slightly bigger than the peak of the Ru—Ru bond in the Ru foil. According to previous reports, this may be due to the Ru—Ru bond in the small Ru cluster, or the bond between Co and Ru , we will further explore the attribution of this peak through subsequent characterization methods. The coordination sphere of the atomically dispersed centers is further quantified by least‐squares EXAFS curvefitting analyses  simulations have also been conducted for interpreting a radial distance resolution in the K space, as shown in Figure 2, it is found that there are two obvious WT signals in Ru‐Co3O4, their positions are, respectively, around 3.9 and 6.8 Å−1. Among them, the WT intensity maximum near 3.9 Å−1 arising from the Ru—O coordination is well resolved at 1.0–2.0 Å for SA Ru‐Co3O4. While the signal near 6.8 Å−1 is different from the signal value of the maximum intensity of Ru foil (≈8.8 Å−1), then the WT signal should be assigned to the Ru—Co bond  or small Ru clusters. In addition, the detailed parameters of SA Ru‐Co3O4 measured by XAFS are shown in Table To further investigate the electron state and local structure of SA Ru‐Coed Figure2a. The Å Figure , located3O4, SA Ru‐Co3O4 samples were analyzed on the atomic scale by using atomic resolution aberration‐corrected high‐angle annular dark‐field scanning TEM (HAADF‐STEM). The results demonstrated that numerous bright spots (Ru atoms) are evenly distributed on the substrate; and some bright spots are close to each other, which means that small Ru clusters may also exist on SA Ru‐Co3O4 nanosheets, as shown in Figure 3O4 nanosheets  test results (3.16 wt%).To further verify the existence of atomically dispersed Ru atoms on SA Ru‐Cos Figure , and the2 batteries based on CC, Co3O4/CC, and SA Ru‐Co3O4/CC cathodes were assembled and their various electrochemical performances were tested. Figure−1 within 2.0–4.5 V. It is noted that battery with SA Ru‐Co3O4/CC cathode exhibits highest reduction onset potential and lowest oxidation onset potential, suggesting highest catalytic activity of SA Ru‐Co3O4/CC cathode compared to those of CC and Co3O4/CC cathodes. Figure 2 batteries with the different cathodes at a cut‐off capacity of 500 mAh g−1 with a constant current density of 100 mA g−1. Obviously, the battery with SA Ru‐Co3O4/CC cathode shows the much decreased overpotential (1.05 V) compared to batteries with CC (1.45 V) and Co3O4/CC (1.98 V) cathodes. And Li‐CO2 batteries with SA Ru‐Co3O4/CC cathode still maintain lowest overpotential even at high current density and large restrict capacity  kinetics; atomically dispersed Ru atoms as electron‐deficient centers promote the transfer of electrons in the oxidation (charge) process; and the deposition behavior and morphology of the discharge products were well tailored (vide infra), which is beneficial to the subsequent oxidation charging process. The tailored discharge products encouraged us to investigate the discharge capacity of batteries. Figure −1 and cut‐off voltage of 2.0 V. Unexpectedly, battery with SA Ru‐Co3O4/CC cathode exhibits a higher discharge capacity of 30 915 mAh g−1 than those with the CC (6468 mAh g−1), Co3O4/CC (19 228 mAh g−1) cathodes, even at a higher current density of 300 mA g−1, the discharge capacity of the SA Ru‐Co3O4/CC cathode can still reach 16 510 mAh g−1  was also prepared as shown in Figure 2‐Co3O4/CC cathode were tested as shown in Figure 2‐Co3O4/CC are improved compared with that of pristine Co3O4, the catalytic activity of RuO2‐Co3O4/CC is still lower than that of SA Ru‐Co3O4/CC, further indicating the advantage of single atom catalyst in the field of Li‐CO2 batteries.Li‐CO2 batteries performance. The discharge/charge profiles of batteries with different cathodes are shown in Figure −1 with a current density of 200 mA g−1. Obviously, compared with CC (41 cycles) and Co3O4/CC (97 cycles) cathodes, the cycling stability of SA Ru‐Co3O4/CC cathode (251 cycles) was significantly improved, which might be ascribed to the increased catalytic activity and optimized discharge products growth pathway caused by atomically dispersed Ru atoms (vide infra). Their corresponding time–voltage profiles are illustrated in Figure 3O4/CC cathode could sustain a stable discharge and charge potentials and operate more than 2000 h, and this is superior to that of battery with CC (328 h) and Co3O4/CC (776 h) cathodes, respectively. Li‐CO2 battery with SA Ru‐Co3O4/CC cathode can maintain comparable cycle stability even under deep charge/discharge cycles as shown in Figure 3O4/CC cathode catalyst, the performances of SA Ru‐Co3O4/CC cathode catalyst were compared with some previously reported Li‐CO2 battery cathode catalysts, as shown in Figure 2 batteries based on SA Ru‐Co3O4/CC cathodes is not shocking enough, but its properties are still better than that of most reported Ru‐based catalysts and SASCs in Li‐CO2 batteries system. Furthermore, considering that the load of precious metal Ru in SA Ru‐Co3O4/CC is much lower than other reported Ru‐based nanoparticles, such an overpotential is acceptable and competitive. In general, the electrochemical performances of Li‐CO2 battery with SA Ru‐Co3O4/CC cathode are superior to report in most literatures.Cyclic stability is an important index to evaluate Li‐COatalysts,48 as 2 battery, the evolution of the morphology and composition of the discharged and recharged cathodes was investigated. Under the current density of 200 mA g−1 and the cut‐off capacity of 800 mAh g−1, SEM technique was used to observe the pristine, discharged, and charged CC, Co3O4/CC and SA Ru‐Co3O4/CC cathodes as shown in Figure3O4/CC, and SA Ru‐Co3O4/CC cathodes present three different morphologies, which are granular, massive shape, and thin flaky shape, respectively, as shown in Figure 2 batteries. This phenomenon will be discussed in detail in the following sections by conjunction with the results of DFT calculations. After charging, there are still a small amount of small granular discharge products remaining on the surface of the CC cathode, while the discharge products on the Co3O4/CC and SA Ru‐Co3O4/CC cathodes are both decomposed  was used to monitor the consumption and production of CO2 under a current of 0.2 mA with a cut‐off capacity of 0.2 mA h, as shown in Figure To further investigate the operating mechanism of Li‐CO2 battery is a four‐electron transfer process, and the theoretical molar ratio of electron transfer to CO2 generation/consumption (e−/CO2) is about 1.33. In addition to the above paths, some other possible charging reaction paths were also proposed in recently reported literatures, such as Li2CO3 self‐decomposition and O2•− mediated Li2CO3 self‐decomposition according to Equations  and charge (1.39) process. Combined with the results of XRD pattern and Raman spectra, it is suggested that the discharge and charge reaction of Li‐CO2 battery based on SA Ru‐Co3O4/CC cathode may be in accordance with Equation . We supposed that this may be attributed to the following two reasons: During the charging process, CC and Co3O4/CC cathodes are difficult to promote the reaction between Li2CO3 and carbon species due to their insufficient catalytic activity, leading to a part of Li2CO3 is decomposed through O2•− mediated self‐decomposition pathway; and the higher charging potential leads to the decomposition of the electrolyte and cathode, resulting in a series of unknown parasitic reactions. This also explains why Li‐CO2 battery based on SA Ru‐Co3O4/CC cathode shows better cycle stability.However, based on the results of DEMS, no Od the Eq  was was quation  as shown in Figure3O4/CC cathode can promote the surface‐mediated reaction path of discharge products by enhancing the binding energy to the key reaction intermediate Li2C2O4. In addition, as encouraged by plentiful active sites triggered by atomically dispersed Ru atoms, homogeneously distributed discharge products were formed on the SA Ru‐Co3O4/CC cathode, which further verifies the SEM results. The homogeneously distributed discharge products increase the contact area with the cathode, which is conducive to the transfer of electrons and Li+ in the charging process. As a benefit, the high charge overpotential and parasitic reaction caused by high voltage can be greatly reduced. In addition, the uniform distribution of discharge products can alleviate the cathode volume effect to a large extent, which is beneficial to improve its stability. As for Li‐CO2 battery based on Co3O4/CC cathode, the growth of discharge products tended to be solvation‐mediated growth pathway according to the calculation results of adsorption energy, which eventually leads to the generation of massive shape discharge products.The morphology and distribution of discharge products depended on their growth pathway, which is important for the performance of the battery,23 whis Figure . Some ke2 batteries. However, the key factor affecting the performance of batteries, especially the round‐trip efficiency, lies more in their charging process, that is, the oxidation process of discharge products (CO2 evolution reaction). How does the catalyst essentially promote the oxidation process is the key to understanding the reaction mechanism of the battery and improving its performance. However, there has been little research in this area. It is well known that the doping of heterogeneous atoms will regulate the charge distribution on the surface of the supporter, and then change its catalytic properties. Inspired by this concept, The Bader charge of Co and Ru sites on the surface of Co3O4 and SA Ru‐Co3O4, respectively, was calculated as shown in Figure 3O4 has a more positive charge than that of octahedral Co site on Co3O4. According to previous reports, the doped metal atoms as electron‐deficient centers can improve the catalytic activity for the oxidation process, which is beneficial to reduce the charging overpotential of the battery. In conclusion, heterogeneous Ru atoms as the active centers not only enhance the CRR process of the battery, optimize the morphology and growth pathway of the discharge products, but also play a positive role in the battery charging process. All these results indicate that the synthesized SA Ru‐Co3O4 is an efficient dual catalyst for Li‐CO2 batteries.Generally, the adsorption/activation of reactants and intermediates as well as the regulation of the reaction path of the battery caused by catalyst can greatly improve the performances of Li‐CO process,57 whi2 batteries based on Co3O4/CC and SA Ru‐Co3O4/CC cathodes is shown in Figure 2 battery was changed to the surface‐adsorption growth pathway by enhancing the adsorption energy of the key intermediate Li2C2O4, and finally the thin and uniformly distributed discharge products were obtained. During the charge process, the thin and uniform discharge products are tightly combined with a large number of distributed Ru atoms active sites, making them easier to decompose. In sharp contrast, Li‐CO2 batteries based on Co3O4/CC cathodes generate large bulk products through solvation‐mediated pathway. In this case, the bulk discharge products have poor contact with the active sites, causing them to decompose at higher charging platforms. It should be noted that although the discharge product is solid phase and covers a certain amount of active sites, its influence on the cathode activity is very limited. The reason is that in the process of battery charge and discharge, the active site can not only catalyze the generation/decomposition of discharge products through direct contact but also promote the generation/decomposition of discharge products through the transfer of electrons and holes (hole polarons) with the help of defects in the discharge products, which has also been confirmed by previous reports.To show how different cathodes tailor the growth pathway of discharge products, a schematic diagram of the discharge process of Li‐CO reports.5933O4/CC was successfully applied to Li‐CO2 batteries as an efficient dual‐catalyst cathode. Experimental results show that the doped Ru atoms as drive force centers not only improves electrochemical performances of the battery but also tailors the morphology and distribution of discharge products. Critically, DFT simulations demonstrated that SA Ru active site could promote the adsorption of key intermediate (Li2C2O4), thus promoting the discharging process as well as optimizing the growth pathway of discharge products. Additionally, SA Ru active site as an electron‐deficient center also promotes oxidative decomposition of discharge products during charging, and this effectively reduces the charging overpotential. This work not only shares a simple and gentle preparation strategy for the synthesis of SASCs but also provides a new idea for the development of efficient cathode for Li‐CO2 batteries.In conclusion, a facile and eco‐friendly method to prepare SASCs was proposed, and the derived SA Ru‐CoThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


Respiratory motion models establish a correspondence between respiratory‐correlated (RC) 4‐dimensional (4D) imaging and respiratory surrogates, to estimate time‐resolved (TR) 3D breathing motion. To evaluate the performance of motion models on real patient data, a validation framework based on magnetic resonance imaging (MRI) is proposed, entailing the use of RC 4DMRI to build the model, and on both (i) TR 2D cine‐MRI and (ii) additional 4DMRI data for testing intra‐/inter‐fraction breathing motion variability.Repeated MRI data were acquired in 7 patients with abdominal lesions. The considered model relied on deformable image registration (DIR) for building the model and compensating for inter‐fraction baseline variations. Both 2D and 3D validation were performed, by comparing model estimations with the ground truth 2D cine‐MRI and 4DMRI respiratory phases, respectively.3), with higher values in the presence of large inter‐fraction motion . In the 2D validation, the median estimation error on anatomical landmarks’ position resulted below 4 mm in every scenario, whereas in the 3D validation it was 1.33 mm and 4.21 mm when testing intra‐ and inter‐fraction motion, respectively. The range of motion described in the cine‐MRI was comparable to the motion of the building 4DMRI, being always above the estimation error. Overall, the model performance was dependent on DIR error, presenting reduced accuracy when inter‐fraction baseline variations occurred.The median DIR error was comparable to the voxel size (1.33 × 1.33 × 5 mmResults suggest the potential of the proposed framework in evaluating global motion models for organ motion management in MRI‐guided radiotherapy. The management of respiratory motion in external beam radiotherapy is of primary importance for tumour coverage and sparing of the surrounding healthy tissues.Recently, magnetic resonance imaging (MRI) has emerged as an ideal image modality for treatment guidance, with in‐room MRI‐linacs clinically availableTo estimate 3D TR information of tumour and surrounding organs, respiratory motion modelling techniques have been explored from the 2D or 1D information available.In this study, we propose an MRI‐based validation framework for the assessment of global motion models performance in MRI‐guided treatments of patients with abdominal lesions. The framework relies on repeated RC 4DMRI and TR 2D cine‐MRI being able to account for breathing motion variabilities. For our analyses, we consider a global motion model already adopted in the literature,The proposed framework entails two acquisition sessions (S1 and S2) on each patient, performed on different days with the same setup in order to sample inter‐fraction breathing variations. In each session, two RC 4DMRI (R1 and R2) were acquired to capture intra‐fraction breathing pattern variability. In both sessions, also a TR 2D cine‐MRI scan was performed to sample cycle‐to‐cycle breathing variations.The respiratory motion model is built on the first‐acquired RC 4DMRI R1) and tested on successively acquired data sets, considered as ground truth for model evaluation. This means that, for each testing data set, the corresponding surrogate is used to feed the motion model, and the resulting model estimates are compared against the ground truth data. Only a 2D comparison between the ground truth 2D cine‐MRI and the corresponding slice of the model estimate is viable , whereas a 3D comparison is possible between the ground truth respiratory phases of the testing RC 4DMRI and the model estimates . The evaluated scenarios are described in Table  and testData from 5 pancreas  and 2 liver patients  were acquired at the National Centre for Oncological Hadrontherapy https://plastimatch.org/), and its error was evaluatedThe considered respiratory motion modelThe DIR error affecting the motion model performance was evaluated in a patient‐specific fashion.For the 3D validation, the 3D MRIs estimated by the model were compared against the corresponding volumes of the (ground truth) testing 4DMRI. A global evaluation was performed by automatically extracting corresponding points in the estimated and the ground truth volumes by means of the SIFTthe distance between landmarks in corresponding slices, quantifying the estimation error (cine‐error);the maximum displacement of cine‐MRI landmarks with respect to the end‐exhale MRI (which was used as reference for model building), thus evaluating respiratory motion during cine‐MRI acquisition (cine‐motion);the landmarks range of motion in the building 4DMRI, for comparison with cine‐MRI motion (4DMRI‐motion).For cine‐MRI, only a 2D evaluation of the model performance was possible on coronal and sagittal planes. Specifically, the slices of the estimated 3D MRI (referred to as estimated slice) corresponding to the (ground truth) coronal/sagittal cine‐MRI frame were considered. Anatomical landmarks manually identified on the first cine‐MRI frames (Appendix S3) were tracked through template matching on all cine‐MRI frames and on the corresponding estimated slices, as well as on the end‐exhale and end‐inhale MRI slices of the building 4DMRI. In the 2D validation, the SIFT was not used due to poor performance on the estimated coronal slice. Then, the model performance was investigated by considering:Overall, the median (interquartile range) error was 1.33 (2.97) mm, 1.33 (2.66) mm and 2.97 (4.41) mm for the 4DMRI S1, 4DMRI S2 and INTER DIR, respectively. For 4DMRI S1 and 4DMRI S2 DIR, the median error was below 1.88 mm in all patients Fig. , whereasThe error on corresponding points between the estimated and ground truth volumes, as well as the errors in terms of COM, reflected the trend observed on DIR error. Specifically, the error on corresponding points resulted in median of 1.33 mm over all cases when testing intra‐fraction motion variations (3D intra‐S1 and 3D intra‐S2), with a tumour COM error in median of 1.81 mm, whereas slightly higher errors  were quantified in the presence of inter‐fraction variations (3D inter‐R1 and 3D inter‐R2). Nevertheless, for all the investigated scenarios, the estimation error was below the range of motion to be estimated by the model , as clear from Fig. The proposed validation framework confirms that the global motion model performance is dependent on DIR accuracy, which therefore requires a patient‐specific evaluation,One limitation of the study consists in the restricted field of view of the 4DMRI (12.5 cm in the right‐left direction), which was due to imaging constraints and the trade‐off between temporal and spatial resolution.In conclusion, the acquisition of 2D cine‐MRI and repeated 4DMRI on patients with abdominal tumours allowed us to provide the groundwork for the definition of a validation framework that is put forward to create the required imaging data set for the evaluation of global motion models in MRI‐guided treatments.Appendix S1. Respiratory motion model schematic.Appendix S2. Phase and amplitude estimation.Appendix S3. Manually identified landmarks.Appendix S4. Deformable Image Registration validation.Appendix S5. Qualitative results.Appendix S6. Jacobian analysis of model output.Click here for additional data file.


However, the cathode structure collapse caused by H2O co‐insertion in aqueous solution dramatically deteriorates the electrochemical performance and hampers the operation reliability of V2O5–Zn batteries. The real‐time phase tracking and the density functional theory (DFT) calculation prove the high energy barrier that inhibits the Zn2+ diffusion into the bulk V2O5, instead the ZnCl2 “water‐in‐salt electrolyte” (WiSE) can enable the dominant proton insertion with negligible lattice strain or particle fragment. Thus, ZnCl2 WiSE enables the enhanced reversibility and extended shelf life of the V2O5–Zn battery upon the high temperature storage. The improved electrochemical performance also benefits by the inhibition of vanadium cation dissolution, enlarged voltage window, as well as the suppression of the Zn dendrite protrusion. This study comprehensively elucidates the pivotal role of a concentrated ZnCl2 electrolyte to stabilize the aqueous batteries at both the static storage and dynamic operation scenarios.Aqueous V 2 “water‐in‐salt” electrolyte in V2O5—Zn battery can effectively enhance the electrochemical performance no matter in dynamic or static conditions. This improvement is ascribed to the inhibition of cathode dissolution, extended voltage window, and the suppression of the Zn dendrite protrusion. Moreover, it is proved that proton inertion contributes most of the capacity, benefiting the particle integrity.The application of ZnCl Meanwhile, zinc metal possesses a low equilibrium redox potential (−0.76 V vs standard hydrogen electrode) and high hydrogen evolution overpotential, balancing the appropriate energy density with aqueous electrolyte compatibility. Despite these merits, the AZBs suffer from cathode structural collapse and electrolyte decomposition, especially at the slow cycling rates. Besides, the self‐discharge issue of AZBs has yet to be explored, posing the limited calendar life for the practical deployment of the AZBs. V2O5 and its derivative are the most studied cathode materials for the AZBs because of the natural abundance and rich oxidation states of vanadium. Noteworthy, the charge charrier for V2O5 in typical AZBs is the hydrated Zn ions other than Zn ions, due to the high charge density of Zn ion and the charge shielding effect of water molecular. The water and Zn ion co‐insertion would cause dramatic lattice variation up to 300% and deteriorate the structural stability of the strained V2O5 upon the long‐term cycling.With the uptake of intermittent renewable generation (namely wind and solar power), the grid‐scale energy storage and behind‐the‐meter storage markets rapidly grow with the more demanding requirements for the greener, cleaner, and more cost‐efficient chemistry formats.3, 4 Additionally, the cation/ molecular pre‐intercalation methods, including Li+, Na+, K+, Mg2+, Cu2+, Zn2+, Ni2+, Mn2+, H2O, and polyaniline, were proposed to buffer the lattice expansion of the layered galleries and alleviate the internal stress. For instance, Yang et al. reported a cotton‐like Li+ pre‐intercalated V2O5 material, demonstrating an enlarged layered spacing (13.77 Å) and facile Zn2+ diffusion (up to 3.37 × 10−8 cm2 s−1). Liu and his co‐workers synthesized polyaniline intercalated V2O5 with constant interlayer distance upon the intercalation and extraction of the Zn2+. Even though the nano‐engineering and pre‐intercalation strategies achieved great success in terms of capacity and cycle stability at high rate, the complex pre‐treatment procedures hinder the practical application of the as‐synthesis materials, while the pre‐intercalation processing would also sacrifice the rated gravimetric/volumetric energy density of the AZBs. At the high voltage status, moreover, the large surface area of the nanostructure would cause serious interfacial side reactions upon the electrolyte contact. Thus, the innovative design of the practical AZB should put more emphasis on the component compatibility of the AZB system, besides the design ingenuity of the electrodes only.Pioneer studies mainly focused on the nano‐engineering techniques, that is, the doping process or interfacial modification for structural stabilization.18 Add A composite electrolyte comprising 20 m LiTFSI and 1 m Zn(TFSI)2 was also employed to stabilize the reversibility of Zn batteries (Coulombic efficiency (CE) of 99.5% for 200 cycles at 1 mA cm−2). However, the heavy use of costly salts goes against the cost benefits of the aqueous electrolyte. Recently, Zhang et.al. used a new type of WiSE based on the low‐cost ZnCl2 to suppress the zinc dendrites. Jin and co‐workers developed a low‐cost 19 m (17 m NaClO4 + 2 m NaOTF) WISE for aqueous sodium ion batteries, in which NaClO4 was mainly used to reduce water activity as well as a small amount of NaOTF was applied to assist formation of robust SEI, enabling a wide electrochemical stability window of 1.6–4.4 V (vs Na+/Na). Despite the enhanced CE values achieved in these concentrated electrolytes, there remains a paucity for the spontaneous interfacial instability or electrochemical‐driven redox reaction in the aqueous batteries. It is noteworthy that the influence of electrolyte concentration on reaction mechanism of V2O5 cathode in AZBs has not been studied even though the water plays a pivotal role in the charge storage process.The high concentration of the “water‐in‐salt” (WiSE) electrolyte was recently proposed to optimize the cyclability of AZBs. Wang and coworkers reported the 21 m LiTFSI (lithium bis(trifluoromethanesulfonyl)imide) salt could extend the electrochemical window of the aqueous electrolytes to ≈3.0 V with the aid of solid electrolyte interface (SEI) formation.33 A c2O/Zn2+ co‐insertion into the V2O5 cathode on the micro‐ and electrode level. Correspondingly, we employed a high concentration of ZnCl2 WiSE (30 m) to eliminate the water co‐insertion process. The real‐time phase tracking and DFT calculations of the V2O5 cathode validate the negligible lattice variation accompanied by H+ intercalation/de‐intercalation upon the cycling in the WiSE. Moreover, the prototype realizes robust capacity retention at the low current rate (9.5% capacity loss after 300 cycles at 50 mA g−1) as well as self‐discharge suppression (74.66% capacity retention after 300 h at 55 °C). Based on these multiscale structural and compositional characterizations, we propose a plausible deterioration model of the vanadium‐based cathode in the dilute electrolyte: severe internal lattice stress of the V2O5 cathode that derives from the aqueous solvent media would induce the particle pulverization and cation dissolution, which further exacerbates the parasitic interfacial reaction in turn. On the contrary, the ZnCl2 WiSE could significantly immobilize the free water and stabilize the interfacial electrochemistry in the AZB, at both the high‐temperature storage and dynamic operation scenarios.Herein, we systematically investigated the detrimental impact of the H22O5 particles are ≈4–5 µm in size with irregular interparticle pores distributed on the surface. In addition, the energy‐dispersive X‐ray spectroscopy (EDS) result of the top‐view electrode reveals the uniform dispersion of the graphite and V2O5 particles on the scale of micrometer. The as‐prepared V2O5 cathode was directly assembled as a reference to the zinc foil to illuminate the impact of various electrolytes on cycle performance. The commonly employed electrolyte (1 m ZnSO4) realized a maximum capacity of 302 mA h g−1 for AZBs  and phase‐pure V2O5 (PDF NO.75‐457), demonstrating that the ball‐milling process does not induce any undesired phase transition . Obviously, a new peak that appeared at the discharged state was indexed to the (001) peak of ZnxV2O5·nH2O, suggesting the co‐intercalation of the Zn2+ as well as the accompanied water molecules insertion into the V2O5 structure. Pioneer works also reported that the water molecular would intercalate into the layered spacing of V2O5 and induce the dramatic lattice variation. In fact, the V2O5 cathode in dilute electrolyte would experience a serious lattice swelling from 4.4 to 13.9 Å (319%). Post‐mortem morphological analysis was carried out to probe the electrode evolution upon such mechanical stress. Before cycling, the cathode displays randomly distributed V2O5 particles. Then, the particles gradually evolve into the interconnected porous structure in the dilute electrolyte after 200 cycles. Noteworthy, this microstructure evolution is accompanied by serious active material loss, evidenced by the gradually decreased V:C ratio from 5.20 for the pristine electrode to 0.25 for the cathode cycled in 1 m ZnSO4  to probe the relationship between electrochemical reaction and water content. The application of WiSE effectively improved the cycle stability of V2O5 cathode: the maximum capacity was evaluated 341.0 mA h g−1 and only 9.5% capacity decay can be observed after 300 cycles at 50 mA g−1  peak of V+ and Zn2+ diffusion in V2O5 were also calculated by simulating the structural evolution in the absence of water. The migration pathways for H+ and Zn2+ in V2O5 are shown in Figure + and Zn2+ in V2O5 are 0.10 and 6.57 eV. The diffusion barrier of V2O5 can even block the Zn2+ intercalation in WiSE. According to the aforementioned experimental results, Figure 2O5 cathode in dilute electrolyte would experience a serious lattice swelling from 4.4 to 13.9 Å (319%) while the application of WiSE effectively suppresses the lattice expansion due to the small radius of the proton. Based on the plausible proton insertion mechanism, the large volume change and high internal stress of the V2O5 cathode could be remarkably avoided in the WiSE.To fully understand the competitive relationship between the charge carriers, the dynamic kinetics of H2O5 cathode and electrolyte in detail. As shown in Figure −1 upon 300 cycles, while the cells containing 1 m ZnSO4 piled up the irreversible capacity consumption to 235.6 mA h g−1, consistentwith the excellent cycle stability. The self‐discharge rate is a critical indicator to assess the rechargeable battery in the practical application scenarios, especially upon the operation at elevated temperatures. Therefore, the above various cells were charged to 1.6 V and kept idling at 55 °C to evaluate the self‐discharge rate. Subsequently, all the cells were discharged and charged again at 20 mA g−1. As compared in Figure 2+ could be extracted in the subsequent charge process, suggesting that reversible side reactions had been reduced and no irreversible structural damage occurred for the cathode.Besides the cycle performance and bulk structure evolution of cathode, we also evaluated the compatibility of V−1 . We believe that the large surface area and good conductivity of activated carbon can reduce the electrolyte oxidation, and we find that the Cl2 evolution occurred above 2.0 V versus Zn/Zn2+. Thus, it should be no Cl2 evolution considering our onset potential is only 1.6 V versus Zn/Zn2+.Additionally, the capability of various electrolytes against oxidation reaction at high voltage was measured by linear sweep voltammetry (LSV) with titanium foil as the working electrode in three‐electrode cells at a scan rate of 0.1 mV s Herein, the AZBs were shelved in various electrolytes at 55 °C for 10 days. Figure m ZnSO4 electrolyte, which subsequently induced the rapid capacity loss both upon the galvanostatic cycling and the high temperature storage. In sharp contrast, the concentrated ZnCl2 electrolyte would inhibit the vanadium dissolution due to the fewer free solvent molecules that can coordinate with the metal cations. In this regard, the usage of WiSE indeed practically obstructs the unwanted side reactions on the cathode/electrolyte interface both at the static and dynamic conditions.The concentrated electrolyte was proposed to suppress the metal ion dissolution of the layered cathode.35 Her2O5 cathode in WiSE, cyclic voltammetry (CV) curves at different scan rates were tested to quantitatively identify the capacitive contribution as shown in Figure m ZnSO4). Upon the scan rate increase, the voltage hysteresis between the reduction peaks and the oxidation peaks gradually enlarged. According to previous reports, the capacitive effect of the battery system can be calculated by the relation i = bav. If b reaches 0.5, a faradic intercalation controls the charge storage process while b equals 1.0 for the pseudo‐capacitance process. The b values are calculated as 0.84, 0.72, 0.82, and 0.70 for peaks 1–4, respectively, suggesting the capacity of V2O5 particles is coherently dominated by the capacitive and diffusion behaviors. At a scan rate of 0.2 mV s−1, capacitive behavior contributes 53.46% of the total capacity and the capacitive contribution increases with scan rate. Therefore, the capacitance is a major contribution to the total capacity, especially at high charge/discharge rates. The galvanostatic intermittent titration technique (GITT) was also employed to investigate the ion diffusion rate. Despite the multivalence of Zn2+, the Zn2+ diffusion coefficient (≈10−13 cm2 s−1) in 1 m ZnSO4 is measured as one order magnitude higher than the H+ diffusion coefficient in WiSE (≈10−14 cm2 s−1). We ascribe this lower ion diffusion rate to the low charge charrier (H+) concentration in the WiSE, which in return guarantees the structural robustness of the V2O5 cathode at the low current density.To further understand the energy storage behavior of Vm ZnSO4).15 Upo After 200 cycles, the zinc foil in 1 m ZnSO4 electrolyte shows the micro‐slice morphology due to the uncontrolled zinc plating 2, cannot convert to the original MnO2 with tunnel structure. Thus, we believe that the V2O5–Zn battery with WiSE, which shows continuous layered structure change in cathode, is an alternative choice to be a large‐scale, low‐cost, and rechargeable energy supply.The electrochemical reactions that occurred inside the Vcapacity.48 How32 WiSE through the real‐time phase tracking, both at the dynamic cycling and static storage scenarios. The experimental and theoretical analysis validated the enhanced mechanism of WiSE, namely the charge storage of the V2O5 cathode based on the H+ insertion/extraction process, which significantly inhibited the internal mechanical stress of the cathode structure as well as the Zn dendrite deposition. The interfacial stabilization of the AZB configuration thus echoes with the impressive cyclability (300 cycles at a low current density of 50 mA g−1) and the mitigated self‐discharge rate (74.66% capacity maintained after 300 h rest at the 55 °C). This work systematically provides a new insight into the AZBs integration, namely the component compatibility of the ZnCl2 WiSE, vanadium‐based cathode and the Zn foil. Additionally, a long‐standing puzzle for the structural evolution of the V2O5 cathode is elucidated, which highlights the crucial ZnCl2 WiSE in designing the highly efficient AZBs devices, both at the static storage and dynamic cycling conditions.In conclusion, we probed the reversibility and stability origin of the AZBs with the ZnCl42O5 cathode material was prepared by a facile ball‐milling process. Briefly, V2O5 (99%) and graphite (99%) were mixed under ball‐milling with a predetermined weight ratio of 4:1 for 2 h at a speed of 250 r min−1. The ZnCl2 is easy to absorb the moisture in air, so the chemical should be stored and weighted in glove box filled with Ar. 30 m ZnCl2 aqueous solutions were prepared by dissolving 20.44 g zinc salts in 5 mL water solvent at 60 °C under stirring for 12 h. The as‐prepared samples also needed to be sealed for storage.All the chemicals were purchased from Sigma–Aldrich and used as received without further purification. Vα1 radiation with the wavelength (λ) of 0.70930 Å, operating at 50 kV and 40 mA. The crystal structure and lattice parameters of the obtained powder were analyzed by the same equipment. The particle size and morphology of the samples were examined by field emission scanning electron microscopy (Quanta 600 FEG) at 15 kV. Transmission electron microscopy (TEM) images were captured at 200 kV on an FEI Talos F200X with energy‐dispersive X‐ray spectroscopy (EDS). The compositional ratio of V in the cycled electrolyte was analyzed by inductively coupled plasma (ICP) analysis, which was performed on Agilent 5110.Operando XRD analysis was carried out to study the phase transition during the discharge/charge process; the tests were performed in a transmission mode X‐ray diffractometer (STADIP STOE) with a position‐sensitive detector and Mo KN‐methyl‐2‐pyrrolidone (NMP). Then, the homogenous slurry was spread onto a titanium foil and dried at 120 °C for 12 h in a vacuum oven. The final product was punched into discs (Φ = 12 mm) as the working electrode. The electrochemical performance was tested via a pouch cell where the zinc foil and Whatman GF/D glass fiber were used as a counter electrode and separator, respectively. The galvanostatic charge/discharge measurements were implemented on a Neware‐Battery Testing System 4 battery tester at ambient temperature within the potential range of 0.4–1.6 V. Linear sweep voltammetry (LSV) and cyclic voltammetry (CV) tests were carried out on a Gamry interface 1000 electrochemical workstation. LSV of electrolytes were also measured using a on a three‐electrode device at a scan rate of 0.1 mV s−1; the work electrode and counter electrode were both Ti foil, and Ag/AgCl electrode worked as the reference electrode. Ionic conductivity was measured at room temperature (298 K) by electrochemical impedance spectroscopy (EIS) with Gamry interface 1000 electrochemical workstation. The galvanostatic intermittent titration technique (GITT) measurements of the electrodes were conducted using a Neware‐Battery Testing System 4 battery tester at room temperature after initial three cycles. The test procedure was as follows: a 10 min galvanostatic discharge pulse (20 mA g−1) was applied to the cells, followed by 30 min of relaxation time without any current being passed through the cell. The cycle, consisting of a charge pulse and a relaxation period, was applied to the cell until its voltage reached 0.4 V versus Zn/Zn2+. Then, a charge pulse  was applied to the cells, followed by a 30 min relaxation period without any current flowing through the cell. The cycle was applied to the cells until its voltage increased to 1.6 V versus Zn/Zn2+.The cathode slurry was prepared by mixing the active materials (70%), acetylene black (20%), and polyvinylidene fluoride (PVDF) binder (10%) in The chemical diffusion coefficient can be obtained as:τ is the constant current pulse duration (10 min); nm and Vm are the moles (mol) of V2O5 and molar volume (cm3 mol−1), respectively; S is the electrode–electrolyte interface area ; ∆Es and ∆Eτ are the change in the steady state voltage and overall cell voltage after the application of a current pulse in a single step GITT experiment, eliminating the iR drop, respectively.Where  The 2 × 1 × 1 supercell of V2O5 (010) with space group Pmm was used. The cutoff energy for all calculations was 400 eV. The Brillouin zone integration was accomplished with 2 × 2 × 8 Monkhorst–Pack k‐point mesh. The structures were relaxed until the energy difference and the residual force on the atoms reached 1 × 10−4 and 0.05 eV Å−1. During the optimization of the structures, the DFT + U with U = 3.25eV was considered. DFT‐D2 was used for the van der Waals' interactions between layers. To model ionic diffusion in the selected host structure, the climbing‐image nudged elastic band (CI‐NEB) method was employed to couple with density functional theory (DFT). For each CI‐NEB calculation, six images were interpolated between the initial and final structures. DFT + U was not used for the simulation of diffusion. The convergence threshold of the total energy was set to 1 × 10−4 eV, and a tolerance of 0.1 eV Å−1 for the forces used in the CI‐NEB procedure.All calculations were performed using Perdew–Burke–Ernzerhof (PBE) generalized gradient approximation (GGA) exchange–correlation implemented in the Vienna ab initio simulation package (VASP).51 TheThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


On cooling below 43 K the Mn2+ centres in MnSrTa2O7 adopt a canted antiferromagnetic state, with a small spontaneous magnetic moment. On further cooling to 38 K there is a further transition in which the size of the ferromagnetic moment increases coincident with a decrease in magnitude of the polar distortion, consistent with a coupling between the two polarisations.Preparing materials which simultaneously exhibit spontaneous magnetic and electrical polarisations is challenging as the electronic features which are typically used to stabilise each of these two polarisations in materials are contradictory. Here we show that by performing low-temperature cation-exchange reactions on a hybrid improper ferroelectric material, Li 2O7 by performing low-temperature cation-exchange reactions on Li2SrTa2O7.Fabricating materials with simultaneously spontaneous magnetic and electrical polarisations is challenging due to contradictory electronic features. Here, the authors report a synthesis path toward a perovskite MnSrTa However, the preparation of magnetoelectric multiferroic materials is hampered by the contrasting electronic structures typically required for ferroelectric and ferromagnetic behaviour5. Specifically, the noncentrosymmetric, polar crystal structures, which are a prerequisite for ferroelectric materials7, are typically stabilised by second-order Jahn–Teller distortions8 of either d0 transition metals11 or ns2 post-transition metals15, and these closed-shell electronic configurations are not consistent with magnetism.Magnetoelectric multiferroic materials20. In this mechanism, two nonpolar structural distortions, typically the low-energy tilts of BO6 octahedral units in A3B2O7 and A′AB2O7 layered perovskite phases, combine to break the inversion symmetry of the host framework and couple to a third polar-distortion mode, leading to a spontaneous electrical polarisation28. A key feature of the trilinear coupling between the three distortion modes is that the polar-distortion mode is not the primary order parameter of the ferroelectric-phase transition of this class of material and is unstable in the absence of the two nonpolar distortions, hence the improper label.The “hybrid improper” (HIP) mechanism offers an alternative method for inducing the noncentrosymmetric polar crystal structures necessary for ferroelectric behaviour0 or ns2 diamagnetic ions, it should be easier, in principle, to combine magnetic behaviour with this class of ferroelectric. Indeed, one of the first materials predicted to exhibit HIP ferroelectric behaviour, Ca3Mn2O716, adopts a canted antiferromagnetic state (weak ferromagnet) below 115 K29, in which magnetoelectric coupling has been observed30, in line with theory predictions16.As the HIP mechanism does not rely on the presence of dB-site cations are needed to stabilise the highly distorted perovskite frameworks required for the HIP mechanism to function27. When combined with the high B-site charges needed in these frameworks, this size requirement rules out the inclusion of the vast majority of paramagnetic transition-metal cations, and to date, the only candidate magnetoelectric multiferroic materials based on the HIP mechanism are Ca3Mn2O730 and [Ca0.69Sr0.46Tb1.85Fe2O7]0.85[Ca3Ti2O7]0.1531, the latter of which uses the introduction of some large diamagnetic Ti4+ ions onto the perovskite B site, via alloying the magnetic iron phase with diamagnetic Ca3Ti2O7, to help stabilise the required distorted perovskite framework.However, further investigation of the Ruddlesden–Popper and Dion–Jacobson phases that exhibit HIP ferroelectric behaviour reveals that large A′AB2O7 Dion-Jacobson phases to tune the structural distortions of this family of materials to induce HIP ferroelectric behaviour35. Here we describe how this cation-exchange chemistry can be used to substitute paramagnetic cations onto the A-sites of HIP ferroelectric Ruddlesden–Popper oxides to yield an additional class of magnetoelectric multiferroic materials.Recently we have been using the facile cation-exchange chemistry of the 2SrTa2O7 was prepared by a ceramic synthesis route from a combination of Li2CO3, SrCO3, and Ta2O5. Neutron powder diffraction (NPD) data indicate that Li2SrTa2O7 adopts a polar a–a–c+/a–a–c+ distorted, n = 2 Ruddlesden–Popper structure, described in space group A21am, in which the lithium cations reside in pseudo tetrahedral coordination sites between the SrTa2O7 perovskite double sheets  by the application of the X3– and X2+ distortion modes that tilt the TaO6 octahedra  collected from the cation-exchanged material at room temperature were initially indexed using a modified unit cell based on that of Li2SrTa2O7 and fit using a structural model based on Li2SrTa2O7 (space group A21am), but with the Li+ cations replaced by a 50% occupancy of Mn2+ ions, as shown in Fig. 2, Supplementary Fig. MnSrTaA21am(0β0)000 with a propagation vector q =  compared with the previously refined commensurate A21am symmetry unit cell—a situation that is clearer in a further NPD data set collected at 200 K using the WISH diffractometer.Close inspection of the fit of this commensurate model to the NPD data revealed a series of weak diffraction peaks that could not be indexed by this model and that could not be assigned to impurity phases Fig. . However2O7 framework led to the structural modulation. As shown in Fig. 2+ cations can adopt a chequerboard-ordered arrangement within the tetrahedral coordination sites, analogous to the arrangement of Na+ cations in NaNdNb2O727, resulting in an array of apex-linked MnO4 units, which maximises the Mn–Mn separation. However, there are two distinct ways of ordering the Mn cations (A and B), which are simply related by switching the occupied and unoccupied tetrahedral sites.A series of test structural models were constructed and refined against the data to establish which components of the MnSrTa2+ occupancy, which corresponds to an incommensurate intergrowth of these two ordering patterns, as shown approximately in Fig. Supplementary Information39. In addition, there is a coupled modulation of the “equatorial” O4 oxide ions that are adjacent to the tetrahedral coordination sites  cm3 K mol−1 consistent with the value expected for S = 5/2, Mn2+ cations. On cooling below 43 K, the ZFC and FC data diverge sharply, with field-dependent data exhibiting weak hysteresis at 5 K.Zero-field-cooled (ZFC) and field-cooled (FC) magnetisation data collected from MnSrTa2O7 Fig.  can be fk = , of the incommensurate crystallographic unit cell. The best fit to the magnetic diffraction data was achieved using a model described in magnetic space group A21a’m’(0β0)000, shown in Fig. B at 1.5 K) are aligned parallel to the z-axis and ordered in an approximate G-type antiferromagnetic structure that is then perturbed by the modulated Mn cation order. During the refinement of the magnetic model, the crystallographic model was also simultaneously refined, and the structural modulation vector q =  was observed to be unchanged.NPD data collected at 1.5 K Fig.  show a sA21a’m’(0β0)000) allows a ferromagnetic component parallel to the x axis, consistent with this observation. Furthermore, coupling between the antiferromagnetic order (which transforms as the mX3+ irreducible representation) and the X2+ structural distortion stabilises a weak ferromagnetic moment (which transforms as the mΓ5+ irreducible representation) via a trilinear invariant, as described in the 2O7 is required to have a weak ferromagnetic component by symmetry. However, the small size of this ferromagnetic component means that it could not be observed in the NPD data.Magnetisation data indicate that there is a small ferromagnetic component (canting) to the magnetic structure (0.04 μB per Mn at 5 K). The magnetic space group (TN = 43.5(1) K, as shown in Fig. xy plane above TN. Indeed, the value of the exponential term obtained from the power-law fit (β = 0.27(1)) is in excellent agreement with other systems that show strong 2D magnetic interactions, but weak 3D coupling40.The temperature dependence of the ordered moment obtained from NPD data can be fitted by a power law, TN = 43 K  are nonlinear and exhibit weak hysteresis, consistent with a canted antiferromagnetic state below TN. Equivalent data collected at 35 K and 30 K (below TA) show a significant increase in both remnant magnetisation and coercive field consistent with a sharp increase in the ferromagnetic component (increased canting) of the magnetic state.As noted above, the ZFC and FC magnetisation data diverge at 3 K Fig. . Howeverata Fig.  collecteTA are also observed in the crystallographic data, with both the b and c lattice parameters exhibiting sharp minima at this temperature  in a paramagnetic state. On cooling below T ~ 100 K, 2-dimensional magnetic correlations build up, before the system undergoes a transition to 3-dimensional magnetic order at TN = 43 K. On cooling to TA = 38 K, there is a further transition in which the magnitude of the ferromagnetic moment of the phase increases, the X3− distortion increases, and the Γ5– polar mode decreases.Taking all these observations together, we observe that at 300 K, MnSrTa2O7  is directly coupled to the Γ5– (HIP) polarisation through a quadrilinear invariant involving the mΓ5+, Γ5–, mX3+ and the X3− octahedral tilting. Data in Figs. 5–, mX3+, and X3- distortions undergoe a step change at TA, coincident with a sharp increase in the weak ferromagnetic moment (mΓ5+) of MnSrTa2O7, consistent with magnetoelectric coupling.A symmetry analysis of MnSrTa2O7 see  shows th3- or the X2+ octahedral tilting due to the HIP mechanism. If the X2+ octahedral tilting changes sign, this will also require the reversal of the antiferromagnetic mX3+ mode or reversal of the weak ferromagnetic moment (mΓ5+), due to the trilinear invariant coupling of these distortions to the X2+ tilting. This gives a clear path for the electrical control of the weak ferromagnetic moment via the X2+ tilting. The reverse mechanism, the magnetic control of the electrical polarisation, is less likely since the application of a magnetic field to switch the weak ferromagnetic moment will presumably result in the change of sign of the mX3+ antiferromagnetic mode and not in a change of the X2+ tilting.This quadrilinear coupling also gives an indication that it is possible to reverse both improper ferroic orders, leaving the energy of the system invariant. The actual switching mechanism is more complex and can include different paths, depending on the distortions involved. Indeed, the reversal of the improper ferroelectric polarisation, achievable by the application of an external electric field, requires the reversal of either the X3Mn2O716. However, the appearance of a weak ferromagnetic moment at TN prior to the onset/enhancement of magnetoelectric coupling at TA does not appear to have an analogue in the observed behaviour of Ca3Mn2O7. Accurate theoretical investigations and switching experiments will be needed to identify the details of the coupling mechanism.The symmetry analysis and couplings detailed above are analogous to those described by Benedek and Fennie for Ca3O4 exhibits a magnetic ordering transition (TN ~ 41−43 K)42 at a temperature similar to the magnetic events observed for MnSrTa2O7. However, as shown above, the only features observed in the magnetisation data collected from MnSrTa2O7 are observed at TN and TA, which can be correlated to features in the NPD data collected from MnSrTa2O7  of MnSrTa2O7. Thus, we can conclude that even if small quantities of Mn3O4 are present in the sample of MnSrTa2O7, any contribution they may make to the magnetisation data does not affect the veracity of the analysis presented above. We emphasise that there is no sign of any binary manganese oxide phases in any of the diffraction data collected from any of the samples prepared.It should be noted that Mn+ A-cations of the HIP ferroelectric phase Li2SrTa2O7 with paramagnetic Mn2+ cations allows the preparation of a coupled magnetoelectric material, MnSrTa2O7. Given that this type of cation-exchange chemistry should be broadly applicable, this suggests that a wide variety of M2+ transition-metal cations can be substituted into polar Li2AB2O7 host phases, allowing the preparation of a large number of additional, potentially magnetoelectric, materials.In conclusion, a directed cation-exchange reaction, which replaces the diamagnetic Li2SrTa2O7 were prepared by combining suitable stoichiometric ratios of SrCO3 (99.994%) and Ta2O5  and a 10% stoichiometric excess of Li2CO3 (99.998%) to compensate for metal loss due to volatility at high temperature. This mixture was then heated at 600 °C in air for 12 h, reground, and pressed into pellets. The pellets were then heated to 1250 °C at a heating rate of 5 °C/min and kept at 1250 °C for 12 h, followed by cooling at 5 °C/min to room temperature. Samples were then reground, pressed into pellets, and heated at 1250 °C for 3 h, before being quenched to room temperature. Polycrystalline samples of MnSrTa2O7 were prepared by reacting Li2SrTa2O7 with 5-mole equivalents of anhydrous MnCl2 (98%). The mixture was ground together in an agate pestle and mortar in an argon-filled glovebox, loaded in a Pyrex tube, and heated for four days at 375 °C under flowing argon. The mixture was then washed with distilled water to remove the remaining chlorides, and dried for 12 h at 140 °C in air.Polycrystalline samples of Liα1 radiation). High-resolution synchrotron X-ray powder diffraction (SXRD) data were collected using instrument I11 at Diamond Light Source Ltd. using Si-calibrated X-rays with an approximate wavelength of 0.825 Å, from samples sealed in 0.3 mm-diameter borosilicate glass capillaries. Low-temperature X-ray powder diffraction data were collected using Rigaku Smartlab diffractometer fitted with a Ge crystal monochromator  and an Oxford Cryosystems Phenix cryostat. Neutron powder diffraction (NPD) data were collected on the HRPD and WISH diffractometers at the ISIS neutron source, with the samples contained in indium-sealed vanadium cans. Rietveld profile refinements were performed using JANA200643 and the GSAS suite of programs44. Symmetry analysis and distortion-model quantification was performed using the ISODISTORT software46.Reaction progress and the final sample purity were assessed using X-ray powder diffraction data collected using a PANalytical X’pert diffractometer incorporating an X’celerator position-sensitive detector  and field-cooled (FC) data were collected in an applied field of 100 Oe. Powder second harmonic generation measurements were performed on a Kurtz-NLO system using a Nd:YAG laser with a wavelength of 1064 nm. The SHG signal was recorded and compared with a standard sample of Supplementary Information


Moreover, molecular oxygen is employed as the terminal oxidant in this transformation. A library of 36 compounds as a new class of BF2 complexes with remarkable photophysical properties is delivered in good to excellent yields, showing a substituent‐dependency on the photophysical properties, derived from the π–π* character of the photoexcited state. In addition, aggregation‐induced emission (AIE) is observed and quantified for the brightest exemplars. The excited state properties are fully investigated in solids and in THF/H2O mixtures. Hence, a new series of photofunctional materials with variable photophysical properties is reported, with potential applications for sensing, bioimaging, and optoelectronics.The development of organic materials with high solid‐state luminescence efficiency is highly desirable because of their fundamental importance and applicability in optoelectronics. Herein, a rapid construction of novel BF 2 complexes with N,O‐bidentate ligands through a copper‐catalyzed cascade C—H activation/acyloxylation and difluoroboronation, which avoids the need for pre‐composing the N,O‐bidentate ligands. A library of 36 compounds as a new class of AIE‐active BF2 complexes with remarkable photophysical properties is delivered in good to excellent yields.A novel strategy is developed for the rapid construction of aggregation‐induced emission (AIE) active BF Despite significant developments, the majority of organic BF2 complexes exhibit strong luminescence in dilute solutions. Their emission in aggregates or in the solid state are usually spoiled by self‐absorption and dissipative intermolecular interactions, referred to as aggregation‐caused quenching (ACQ) effects. As a result, the application of organic BF2 complexes has been confined to the fields of fluorescence sensing and bioimaging for a long time. Due to the high demand for organic luminogens with strong solid‐state emission as organic optoelectronic materials for applications including organic lasers, organic light‐emitting diodes (OLEDs), and organic light‐emitting transistors (OLETs), the development of organic BF2 complexes with high solid‐state luminescence efficiency is highly desirable.Since the seminal discovery of organic difluoroboron  mainly focus on linking AIE‐active units such as tetraphenylethylene (TPE) and triphenylamine (TPA) to the chromophoric core  describes the behavior of organic molecules that exhibit dim or no emission in dilute solutions but significantly enhanced luminescence in aggregates or in the solid state, which has been widely accepted to be an efficient strategy exploiting the reduction of rotovibrational relaxation pathways.2 complexes are prepared by the reaction of a prefabricated organic bidentate ligand precursors with boron trifluoride–diethyl etherate in the presence of a base —H bond and poor tolerance of sensitive functional groups such as acyl, cyano, methylthio, and alkenyl. To the best of our knowledge, the transition‐metal catalyzed C—H bond activation/acyloxylation and difluoroboronation remain an unknown transformation.Generally, organic BFe Scheme .[2a] T4, which provides a rapid and efficient method to access a new library of organic BF2 complexes with N,O‐bidentate ligands (Scheme 4)2•6H2O (50 mol%) as a catalyst and BF2 source to transform a broad range of 2‐vinylpyridine derivatives into the corresponding BF2 complexes with N,O‐bidentate ligands in excellent yields; 2) molecular oxygen (precisely: 3O2) is employed as the terminal oxidant (electron acceptor) in this transformation, which eliminates by‐product formation of stoichiometric silver salts and thereby enhances the practicality of this protocol; 3) most notably, the photophysical properties of these obtained BF2 complexes were fully investigated and quantified, both at room temperature and 77K. In addition, the AIE nature of selected BF2 complexes was studied and rationalized through the character of the excited states. Finally, these remarkable results were further evaluated and interpreted by a combination of experimental and theoretical calculations.Herein, we describe a full study of the construction of solid‐state emissive luminogens via copper‐catalyzed C—H activation without using AgBFs Scheme . Notewor22.1E)‐2‐pyridine (1a) as a model substrate under a previously established reaction system for 2‐phenylpyridine derivatives, which was composed of Cu(OAc)2 (20 mol%), AgBF4 (1.5 equiv.), and PivOH (1.5 equiv.) in toluene (0.2 m) under air atmosphere at 140 °C for 20 h 2•6H2O (50 mol%) and AcOH (3.0 equiv.) in toluene under air atmosphere at 140 °C for 20 h as the optimum system to deliver the desired BF2 complex 2a in 95% isolated yield (Scheme 4)2•6H2O services as a catalyst and BF2 source, and the O2 in air is employed as the terminal oxidant for this transformation. A summary of the optimization results is shown in Table 2a was reduced to 34% when running the reaction under argon atmosphere, suggesting that the O2 in air plays the role of the oxidant (Table  (3) compound 2a could be obtained in 83% yield when reducing the loading of Cu(BF4)2•6H2O to 20 mol% and adding 1.0 equiv. of additional NaBF4 as BF2 source (Table 2a was further reduced to 72% when replacing Cu(BF4)2•6H2O (20 mol%) by Cu(OAc)2 (20 mol%) as catalyst  as the oxidant ‐2‐pyridine (1a) undergoes C—H activation to form intermediate IM1, oxidation to form Cu(III) intermediate IM2, and reductive elimination to form intermediate IM3, successively. Then, intermediate IM3 can react with Cu(BF4)2 through anion exchange to afford IM4, which decomposes to give intermediate IM5 along with CuF due to its instability at high temperature. Subsequently, intermediate IM5 can deliver the desired product 2a by releasing a molecule of MeCOF, which is presumably quenched by copper salts in the reaction system. Finally, the generated Cu(I) species is re‐oxidized to the Cu(II) species by O2 to close the catalytic cycle.Based on the above observations and our previous work,2.3Table 2 complexes with N,O‐bidentate ligands were obtained as a library of 36 compounds via copper‐catalyzed C—H activation, which features a broad substrate scope and a high tolerance level for sensitive functional groups. (E)‐2‐pyridines (1) bearing electron donating or electron withdrawing substituents on the ortho‐, meta‐, or para‐position of the phenyl rings smoothly underwent this transformation, producing the corresponding products in good to excellent yields (Table 2a–2w). Satisfactorily, a variation of the substitution pattern on the pyridine moiety was also tolerated in this reaction (Table r–w). Moreover, the pyridyl unit could be replaced with isoquinolinyl unit as chelating group without any decrease in reactivity (Table x). Other N‐heterocycle units such as pyrimidinyl, pyrazolyl, imidazolyl, thiazolyl, and oxazolyl could not be applicable in this protocol. Furthermore, various (hetero)aryls could be installed on the 4‐position of the pyridine ring, delivering the corresponding BF2 complexes in excellent yields (Table 2y–2ai). A variety of functional groups, especially sensitive functional groups such as TMS, MeO, MeS, CHO, Ac, CO2Me, NO2, and even CN were tolerated well.With the optimized reaction conditions in hand, the scope of this reaction was examined. As summarized in 3a–3h). Due to its nonplanar characteristics, excellent electron‐donating properties, and hole‐transporting capabilities, the triphenylamine (TPA) unit and its derivatuves have been widely used in organic optoelectronic materials. To our delight, the TPA‐containing pyridines could smoothly undergo this protocol to form the desired BF2 complex 3h in 72% yield. It needs to be noted that the triphenylamine unit could not be tolerated in our previous method. Therefore, compared with our previous method, the current protocol exhibited broader substrate applicability and could be used as a general method for the synthesis of organic difluoroboron complexes with N,O‐bidentate ligands.To further demonstrate the efficiency of this protocol, a series of 2‐arylpyridine derivatives were employed as reactants in this transformation, delivering the corresponding products in good yields  and emission maxima of all the obtained BF2 complexes were measured in the solid state, and the results are summarized in Table ΦF can reach up to 50%. These variations depend on the π‐donating ability of the substituents, which most likely destabilize the highest occupied molecular orbital (HOMO) while diminishing the HOMO–LUMO gap and the excited state energy (vide infra).As documented in the bibliographic literature,2a, 2k, 2l, 2s, 2y, and 2aj, as they display the highest ΦF of each family of compounds. For these complexes, time‐resolved photoluminescence decays and the resulting excited state lifetimes (τ) at room temperature (RT) were measured. Moreover, ΦF and τ were also evaluated in diluted frozen 2‐Me‐THF glassy matrices at 77 K , only very weak fluorescence intensities can be traced, whereas an intense solid‐state photoluminescence can be detected Table . This reF Figure . In thisΦF is accompanied by prolonged excited state lifetimes, pointing to reduced rotovibrational rate constants upon rigidification in frozen glassy matrices, solids, and aggregates.In general, the enhancement of the photoluminescence 2.52 complexes, X‐ray diffractometric analysis of single crystals from two selected compounds, namely 2a and 2ai, was performed. The conformation and packing patterns of 2a and 2ai were determined from crystalline phases, which were grown by using ethyl acetate as a solvent through slow evaporation at room temperature. As shown in Figure 2a and 2ai display twisted conformations and the phenyl rings  are not coplanar. The dihedral angles between rings B and C, rings B and D, rings E and F, rings G and H, and rings G and I are 66.8°, 26.1°, 36.3°, 88.3°, and 59.4° respectively. This twist enables the inhibition of intermolecular π–π stacking and thereby reduces the exciton quenching in the solid state, which would otherwise induce aggregation‐caused quenching (ACQ). As anticipated, no face‐to‐face intermolecular π–π stacking was found but multiple intermolecular interactions including C—H⋯π and C—H⋯F interactions were observed within the packing structures of 2a and 2ai  computations were performed for a series of representative compounds (see Supporting information for the details). Geometry optimizations revealed that – in excellent agreement with the X‐ray crystallographic analysis – both phenyl rings are not coplanar with the central π system, but adopt a highly twisted conformation. As a result of this minimal conjugation, frontier molecular orbitals are mainly located on the central bicyclic core. Both the π‐type HOMO and the π*‐type LUMO are delocalized over this heterocyclic core, with only minor contributions from the aryl substituents. Time‐dependent DFT computations revealed that the experimentally observed absorption and emission bands can be assigned to spin‐allowed transitions between the ground state and the lowest excited singlet, where the latter can be described as a local π→π* monoelectronic excited configuration accessible via photoexcitation. This assignment applied to all investigated examples with different electronic properties, and the obtained trends in TD‐DFT excitation energies nicely correlate with the experimentally observed emission wavelengths.To further investigate the transitions involved in the absorption and emission of these AIE‐active BF32 complexes with N,O‐bidentate ligands through a copper‐catalyzed cascade C—H activation/acetoxylation and difluoroboronation, which avoids the need for pre‐composing the N,O‐bidentate ligands and features a broad substrate scope and a high tolerance level for sensitive functional groups. This straightforward synthetic strategy enabled us to obtain a library of 36 compounds as a new class of BF2 complexes with N,O‐bidentate ligands in good to excellent yields. The photophysical properties of these compounds were studied in different conditions, such as solid state, frozen glassy matrices, in dilute solution, and under conditions that induce aggregation with concomitant AIE properties. We observed that depending on the substituent pattern, and consequently the π‐electron‐donating ability of the substituents, different emission maxima and quantum yields are obtained, depending principally on the π–π* character of the excited state configuration. X‐ray diffractometry and crystallographic analysis of selected BF2 complexes demonstrated that the weak intermolecular C—H⋯π and C—H⋯F interactions, which could rigidify their highly twisted conformation and impede intramolecular rotation, are responsible for intense luminescence in the solid state. Overall, a new series of photofunctional materials with variable photophysical properties is reported, with potential applications for sensing, bioimaging, and optoelectronics.In summary, we have demonstrated a novel strategy for the rapid construction of aggregation‐induced emission (AIE) active BFThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


Recently, first insights into the regulation and the role of the RNA‐binding protein IMP2 in macrophage activation have been published by Wang et al. This study addresses differences in the regulation of IMP2 between the human and murine system. While the expression of IMP2 in anti‐inflammatory macrophages is synchronous in mice and men, IMP2 expression is regulated differently in inflammatory macrophages. Recently, first insights into the regulation and the role of the RNA‐binding protein IMP2 in macrophage activation have been published by Wang et al. This study addresses differences in the regulation of IMP2 between the human and murine system. While the expression of IMP2 in anti‐inflammatory macrophages is synchronous in mice and men, IMP2 expression is regulated differently in inflammatory macrophages. Advanced Science, Wang et al. for the first time report a role of the insulin‐like growth factor 2 (IGF2) mRNA‐binding protein 2 (IGF2BP2/IMP2/VICKZ2) in macrophage activation.With recently published data in tivation.1 However, beyond the knowledge on its promotion of inflammatory processes in fatty liver disease and autoantibody‐induced glomerulonephritis, little is known about the role of IGF2BP2 in inflammatory conditions.As a highly conserved RNA binding protein, IGF2BP2 plays an essential role in the translation, stabilization, localization, modification, and processing of various mRNA targets and, accordingly, influences physiological and pathophysiological processes in the context of metabolism and malignancy.3, 4Igf2bp2 leads to an enhanced inflammatory M1 phenotype of murine macrophages by stabilizing Tsc1 and Pparg.Wang et al. suggest that a loss of nd Pparg.1IGF2BP2 in polarized human monocyte‐derived macrophages (HMDMs) and our analysis of IGF2BP2 expression levels in various inflammatory conditions in the human system. These data suggest a more complex view of IGF2BP2 in inflammation and, to some extent, contradict the conclusions of Wang et al.Here, we would like to report our data on the regulation of IGF2BP2 mRNA expression  for up to 24 h, we found a significant reduction in on Figure1A, whiln Figure. SimilarIGF2BP2 reduction in HMDMs was accompanied by the induction of the inflammatory cytokines TNF and IL6 and a reduction of the anti‐inflammatory cytokine TGFB , as demonstrated by Wang et al.At early activation time points, B Figure . Importan Figure .7] 7 IGF2TSC1, which was suggested by Wang et al. to facilitate IGF2BP2 action, correlated with IGF2BP2 and was distinctly reduced upon LPS treatment  activation. The presence of acute inflammation was confirmed by an elevated TNF expression of the specific samples. Our analyses demonstrate a reduced IGF2BP2 levels in all models of acute inflammation, although statistical significance was not reached in all of them  macrophages act as effector cells in the Th1 response and are usually generated in vitro by the stimulation with LPS and IFNrization.10IGF2BP2 expression revealed that human M1 macrophages display no alterations in expression, while M2 polarized human macrophages show a significant induction on gene and protein level , Crohn's diseases, type 2 diabetes, and rheumatoid arthritis (Figure The anti‐inflammatory M2 phenotype plays a role in chronic inflammation by supporting tissue remodeling and repair.12 In s Figure .IGF2BP2 is differentially regulated in acute inflammation in humans versus mice. Mouse macrophages in general show a limited responsiveness toward LPS, which has been suggested to be due to a rapid induction of negative feedback regulators. Since IGF2BP2 is consistently induced in both murine and human M2 macrophages, which are desensitized toward activation, the induction of IGF2BP2 in murine macrophages by LPS might contribute to the attenuated responsiveness of murine macrophages.In conclusion, our data suggest that gulators.13 SinIGF2BP2 expression in chronic inflammatory conditions between mouse and man. Future studies are required to gain insight into the role of IGF2BP2 at different stages of inflammation.In contrast to LPS treatment, there is a high consistency of The authors declare no conflict of interest.Supporting InformationClick here for additional data file.


This reaction can treat wastewater at a decomposition rate above 0.077 min−1, even when interacting with organic pollutants at concentrations up to 1500 ppm. Theoretical calculations by Dmol3 simulation elucidates that the reactions proceed spontaneously, and the kinetic constant (kobs) for this redox reaction with 10 ppm RhB dye is 0.53 min−1, which is 65 times faster than the titanium dioxide photocatalytic wastewater treatment. More importantly, the residual waste solution can be further utilized as a precursor to reconstruct the MoSe2 nanoflowers. To demonstrate the effectiveness and reusability, the treated effluent is directly used as the sole source of irrigated water for plants with no adverse effect. This method offers an eco‐friendly and more accessible way to treat industrial wastewater with zero‐discharge.Water pollution is one of the leading causes of death and disease worldwide, yet mitigating it remains a challenge. This paper presents an efficient new strategy for the processing of wastewater utilizing an accessible redox reaction with MoSe 2 nanoflowers, which shows a strong oxidizing ability and permits the decomposition of dye molecules in dark environments without the need of an external power source.In this work, an efficient new strategy is presented for the processing of wastewater utilizing an accessible redox reaction with MoSe Water can be polluted in a wide variety of ways from many different pollutants, but one of the most impactful is wastewater produced by industrial contamination, which contains a wide range of different pollutants at different concentrations and combinations. To clean the water these pollutants need to be processed at a high level of safety and effectiveness, but also at a relatively low cost in order to maintain accessibility to the public. This can be done by using methods that do not require a high amount of technical knowledge or labor to perform, which can reduce labor costs. Currently, the price of wastewater processing is extremely high, and is projected to reach nearly 456 billion dollars worldwide in 2026 due to population growth and freshwater shortages. Thus, developing an efficient, low‐cost, and scalable technology to process wastewater is an urgent issue to tackle.The developmental level of a country is directly linked its water accessibility, as it affects not only the health of their citizens but also the health of their ecosystems. In fact, water pollution is one of the leading causes of death and disease worldwide1, 2 precipitation, carbon adsorption, flotation, ion exchange, membrane filtration, electrochemical methods, bioremediation, and advanced oxidation process technologies. Conventional chemical and physical water treatment are economically advantageous, technologically simple, efficient, and offer a wide range of commercial applications. However, the high construction and repair‐maintenance costs, high energy consumption, and sludge production are major disadvantages to using these processes to treat polluted water. Biological processes involve the use of microbes and are thus eco‐friendly, cost‐effective, and have diverse amount of processing options, but maintaining a well‐controlled environment for the microbe is difficult limiting microbial degradation efficiency. As a result, many alternate methodologies have been developed to increase the efficiency of the treatment process such as a fenton‐like system, electro‐fenton, photocatalytic processes, and piezocatalysis. However, other issues such as sludge production, electrolyte addition, short‐lived reactive species, and catalyst deactivation limit their commercialization. The limitations with current treatment methods highlight the need to create a sustainable and accessible wastewater treatment technology that can treat large amounts of water with ultrahigh concentrations of pollutants by integrating innovative resource recovery technologies into treatment‐process designs.Current wastewater treatment methods include a multitude of physical, chemical, and biological processes including solvent extraction,8 prec2 NFs) and hydrogen peroxide (H2O2) solution to form an oxidized solution. This oxidized solution exhibits excellent degradation activity and can treat ultra‐high concentrations (1500 ppm) of Rhodamine B (RhB) organic pollutants, a commonly used dye in industrial processes. Moreover, this method can lead to 100% wastewater reusability, as the treated effluent can be used as a watering source for plants with no notable negative outcomes on the plants’ growth. Furthermore, the treated effluent can be employed as a precursor to re‐synthesize MoSe2 NFs by adding hydrazine through the hydrothermal process, achieving the aim of zero discharge in a wastewater treatment process. Calculation of all Gibbs free energy change (∆G) reactions are shown in this work, and the results demonstrate the entire degradation process is a spontaneous chain reaction that can be performed in dark conditions without light irradiation.Herein, we developed a new degradation method using molybdenum diselenide nanoflowers  The solution, which now has the potential to oxidize wastewater, forms a thermodynamic nonequilibrium state of H2SeO4 and H2SeO3 (the MS oxidized solution). 3) The ultrahigh concentration of dye molecules (>1500 ppm) is efficiently decomposed by the MS oxidized solution. At the same time, H2SeO4 is reduced to H2SeO3, which is further oxidized by H2O2 back into H2SeO4, allowing for continuous decomposition of organic components and demonstrating the ultrahigh degradation capabilities of the MS oxidized solution. 4) The MoSe2 NFs are re‐synthesized through the hydrothermal process by adding hydrazine into the treated solution. A comparison of the dye removal concentrations between our zero‐discharge system and many recent reports in premier publications is shown in Figure The decomposition of dye molecules in the zero‐discharge system can be described in four steps, as illustrated in n Figure 522SeO4) solution  . In addition, the intensity in the ESI‐MS spectra of both H2SeO3 and H2SeO4 were identical, indicating a state of thermodynamic equilibrium in the solution. In comparison, the MS oxidized solution contained the same molecular ion region lines as selenic acid, but the intensity of H2SeO4 was significantly higher than H2SeO3, indicating an extremely non‐equilibrium thermodynamic state (lower spectrum). The dramatically increased formation of H2SeO4 in MS oxidized solution is further demonstrated by a comparison of the ESI‐MS spectra of the commercial selenic acid solution and MoSe2 NFs mixed with various concentrations of H2O2 from 0 to 90%  was used to investigate the deprotonated molecules in the MS oxidized solution, using a control sample prepared with a commercial selenic acid (H) Figure . In the 2SeO3 in the presence of high concentrations of H2O2 reacted to form H2SeO4, enhancing the concentration of H2SeO4 and forming the MS oxidized solution with H2O, as shown in Equation  to form H2SeO3 with by‐products of CO2, NH3, Cl2, and H2O, as given by Equation , which exhibit abundantly flower‐like morphologies with 3D hierarchical architecture nanosheets highly dispersed around the edge sites. After decomposing the dye molecules, the treated solution was further used as a precursor for preparing the MoSe2 NFs precursor by adding additional hydrazine (MoSe2‐B), whose reconstruction process is shown in Figure 2SeO4, H2SeO3, and MoO3 to Se and Mo monomers, which adjusted the pH value to 7. Next, the solution was transferred to an autoclave for the hydrothermal process (220 °C for 24 h). Finally, after centrifugation and drying, the reconstructed MoSe2‐B can be obtained. The morphologies of MoSe2‐B are very similar to MoSe2‐A, as revealed through the SEM image of MoSe2‐B , as shown in Figure 2‐A, the MoSe2‐B was slightly shifted toward the low‐angle region, suggesting that residual strain was created in MoSe2‐B. However, both MoSe2‐A and ‐B show a significant exposure of active edge sites, which provide a high specific surface area to facilitate rapid reaction with H2O2 t, where kobs (min−1) is the pseudo‐first‐order rate constant for decolorization of the dye solution by the MS oxidized solution and yields a straight line with relatively high regression coefficients (R2 > 0.95), as shown in Figure kobs) for different dye concentrations of 10, 500, 1000, 1500 ppm yielded values of 0.53, 0.19, 0.14, 0.077 min−1, respectively  photocatalysis. Moreover, the kobs using the MS oxidized solution with 1500 ppm dye solution was 0.077 min−1, which is still superior to using the conventional photocatalytic process for a 10 ppm dye solution. Figure 2O2). These results support the conclusion that the dye degradation ability results from the MS oxidized solution, rather than the high concentration H2O2solution. The versatility of our process is further demonstrated in Figure kobs in 30 min for the different food colorings and industrial molecules can be found in Figure −1).The efficiency of the system for wastewater treatment can be evaluated by its degradation activity. As shown in y Figure . The hig2.3Figure 2O2 through the syringe filter with a constant flow rate (2 mL min−1) controlled by a linear motor. As shown in Figure 2 syringe filter maintained a 100% dye removal ratio. The complete degradation of life video is shown in Video 2 NFs by filling it with carbon felt and then the MoSe2 NFs were uniformly grown on its surface, which has an area of 1.8 cm2 , which is necessary to evaluate for industrial applications, reached a steady state of almost zero after 120 min as illustrated by the degradation curve in Figure To simulate the decomposition of the organic pollutants systemically, a degradation filter was fabricated using a syringe, as shown in the schematic diagram of 2 Figure . The SEM2.4Figure −1 and C—C bonding at 1238 cm−1. The leaf mesophyll cells are the main contributors to photosynthetic activity, so the contamination of the mesophyll cells strongly affected the plant's growth. In contrast, the spectrum of leaves irrigated with treated effluent showed no difference to regular leaves. Furthermore, as shown in their optical microscope (OM) images and inset images in Figure To achieve 100% reuse, the treated effluent could be used to grow plants. Here, cabbage seedlings were continuously irrigated with either the treated or untreated effluent to observe the growth of the seedlings, as shown in the schematic diagram in 32 NFs were added with H2O2 solution to generate a solution that contains a thermodynamic nonequilibrium state of H2SeO4 and H2SeO3, which possesses a strongly oxidizing ability for the decomposition of ultrahigh concentrations of up to 1500 ppm various dye molecules, including both industrial dyes and food coloring, with a kobs of 0.077 min−1. After decomposing the dye solution, the treated effluent can be used as a source of water for plant irrigation without any adverse effects on the plants. Alternatively, the treated effluent containing H2SeO4 and H2SeO3 can be used as a precursor for the reconstruction of the MoSe2 NFs by adding hydrazine, which achieves an entirely zero discharge system by reusing 100% of the wastewater by‐products. The fastest noted reaction constant (kobs) is 0.53 min−1, acting on a solution with a dye concentration of 10 ppm, which is 65 times faster than titanium dioxide photocatalysis. The entirely spontaneous degradation process occurs in dark conditions without external energy, which is a significant breakthrough for ultra‐high concentration wastewater treatment and a promising method for industrial applications.The simple process of processing wastewater with MS oxidized solution resulted in a high‐efficiency, scalable, and cyclic process for treating ultra‐high concentrations of diversifying organic pollutants. The MoSe42 NFs were prepared using the hydrothermal method. 0.206 g of Na2MoO4·H2O and the 0.222 g of SeO2 powder were dissolved in 58 mL of deionized water. The solution was stirred for 20 min to disperse completely, followed by the slow addition of 2 mL of hydrazine into the solution. After that, the solution was transferred into a teflon‐lined stainless‐steel autoclave and maintained at 220 °C for 24 h. After the heating process, the autoclave was cooled naturally down to room temperature to obtain the MoSe2 NFs.The MoSel method.57 0.2α radiation and a cold‐field‐emission scanning electron microscope (Hitachi SU‐8010) were used to characterize the crystal phase structure and surface morphology of MoSe2 NFs. Raman spectroscopy was used to identify the composition of the MS oxidized solution. High‐performance liquid chromatography  with a C18 inversed‐phase column was employed for detecting the content ratio of H2SeO4 and H2SeO3, and the mobile phase was prepared by ammonium acetate and methanol  mixing solution.An X‐ray diffractometer (Bruker D2) with monochromatic Cu K2O2 solution up to 90%. To analyze the degradation activity, the 10 mg MoSe2 NFs were added into the 10 mL RhB and 10 mL 90% H2O2 mixing solution. The UV–Visible spectrophotometer (Hitachi U‐3900) was then employed to measure the decomposition ratio of different dye concentration. Carbon felt  was used as a substrate together with MoSe2 NFs precursor and transferred to an autoclave with PTFE lining, and maintained at 220 °C for 24 h. After that, the MoSe2 NFs were grown on the surface of carbon felt.A rotary evaporator was used to prepare a high concentration H2 NFs‐B.After the dye was decomposed by the MS oxidized solution (as referred to as treated solution), hydrazine, which acts as the reducing agent, was added into the treated solution to adjust the pH value up to seven. The solution was transferred to an autoclave with PTFE lining and maintained at 220 °C for 24 h to obtain the MoSeTwo germinate cabbages were irrigated in either treated or untreated solutions. The burette's tip was fixed 2 cm from the seedling to observe growth status, which was recorded entirely with a time‐lapse camera.G) of degradation reactions was calculated by the first‐principles simulation DMol3 module of Materials Studio, which indicated that the degradation process was dominated by a spontaneous reaction. Generalized gradient approximation (GGA) with Perdew, Burke, and Ernzerhof exchange‐correlation functionals was employed. The ΔG of the reactions was calculated by the density functional theory based on the DMol3 module using Material Studio. First, the molecular structure of MoSe2, H2O2, H2SeO4, H2SeO3, RhB, MoO3, H2, O2, H2O, CO2, NH3, and Cl2 were built‐in from the database. Next, the GGA with BLYP (Becke‐Lee‐Yang‐Parr) exchange‐correlation functionals was employed for geometry optimization and frequency analysis. After that, the total energy , entropy (S), heat capacity (Cv), enthalpy (H), and free energy (G) at room temperatures (298.15 K) for all reactants and products from G1, ΔG2, and ΔG3) were calculated as shown in Note The Gibbs free energy change (Δnctionals59 wasThe authors declare no conflict of interest.J.M.W. and J.C. supervised the project. J.T.L. and J.M.W. conceived the idea. J.T.L., J.M.W., and J.C. designed the experiments. J.T. conducted the experiments and theoretical calculation using Materials Studio. J.T.L., J.C., and J.M.W. analyzed the experimental data and working principles, drew the figures, and prepared the manuscript. S.M. and S.S. reviewed and made technical comments on the manuscript.Supporting InformationClick here for additional data file.Supplementary VideoS1Click here for additional data file.Supplementary VideoS2Click here for additional data file.Supplementary VideoS3Click here for additional data file.