NRP1) is a transmembrane co‐receptor involved in binding interactions with variety of ligands and receptors, including receptor tyrosine kinases. Expression of NRP1 in several cancers correlates with cancer stages and poor prognosis. Thus, NRP1 has been considered a therapeutic target and is the focus of multiple drug discovery initiatives. Vascular endothelial growth factor (VEGF) binds to the b1 domain of NRP1 through interactions between the C‐terminal arginine of VEGF and residues in the NRP1‐binding site including Tyr297, Tyr353, Asp320, Ser346 and Thr349. We obtained several complexes of the synthetic ligands and the NRP1‐b1 domain and used X‐ray crystallography and computational methods to analyse atomic details and hydration profile of this binding site. We observed side chain flexibility for Tyr297 and Asp320 in the six new high‐resolution crystal structures of arginine analogues bound to NRP1. In addition, we identified conserved water molecules in binding site regions which can be targeted for drug design. The computational prediction of the VEGF ligand‐binding site hydration map of NRP1 was in agreement with the experimentally derived, conserved hydration structure. Displacement of certain conserved water molecules by a ligand's functional groups may contribute to binding affinity, whilst other water molecules perform as protein–ligand bridges. Our report provides a comprehensive description of the binding site for the peptidic ligands’ C‐terminal arginines in the b1 domain of NRP1, highlights the importance of conserved structural waters in drug design and validates the utility of the computational hydration map prediction method in the context of neuropilin.Neuropilin‐1 (PDB with accession numbers PDB ID: 5IJR, 5IYY, 5JHK, 5J1X, 5JGQ, 5JGI.The structures were deposited to the  NRP1Neuropilin‐1NRP1‐b1B1 domain of Neuropilin‐1NRPsNeuropilin‐1 and Neuropilin‐2PDBprotein data bank165VEGF‐Avascular endothelial growth factor165 has been extensively studied, and is implicated in the processes of angiogenesis 165 binds to a well‐defined binding site on the b1 domain of NRP1 (NRP1‐b1) through its C‐terminal arginine residue. It is possible that other ligands containing a C‐terminal arginine will also bind in a similar fashion 165–NRP1 interaction 165Neuropilins (NRPs) are transmembrane receptors that take part in protein–protein interactions, and are essential for formation of the mammalian nervous and vascular systems We have conducted an extensive study of the C‐terminal arginine‐binding site in NRP1‐b1, focusing on protein side chain flexibility and conserved water molecules. Hydration of proteins is an important parameter to consider in drug design. Structural waters may play a key role in the function, shape and conformation of proteins, as well as in protein–ligand interactions KD) of the arginine analogues to immobilised NRP1‐b1. The binding data, together with computational methods and crystallographic information, was used to rationalise the side chain flexibility and hydration of the binding site in NRP1‐b1. The new crystal structures were analysed alongside publicly available NRP1‐b1 structures in the protein data bank (PDB). This comprehensive analysis and knowledge of the detailed architecture of the VEGF‐A165 binding site in NRP1‐b1 will facilitate future drug design efforts.We selected a number of arginine analogues with simple structural variations to perform this study, and obtained high‐resolution structures of six compounds bound to NRP1‐b1. Surface plasmon resonance (SPR) was used to determine the dissociation constant , 2 mouse  and 2 rat . These structures corresponded either to the isolated b1 domain or tandem b1b2 and a2b1b2 domains of NRP1. The structure of EG00229‐bound NRP1‐b1 (PDB code: 3I97) was selected for initial computational studies , which assesses the role of solvent in protein structures, using a three‐dimensional reference interaction site model (3D‐RISM). This method allows for the calculation of the solvation structure represented by the probability density function of finding interaction sites for the solvent molecules  at the specific point around the solute molecule  Nα‐substitution. Key interactions of EG00229 with NRP1‐b1 are established through hydrogen bonds between the guanidine and carboxylic acid groups of the ligand to the receptor. To gain more comprehensive understanding of the ligand‐binding site structure which would assist us in the design of NRP1 antagonists, we investigated interaction between NRP1‐b1 domain and multiple commercially available arginine analogues. The ligands selected for this study retain the arginine moiety as seen in EG00229, while providing structural diversity in the Nα‐site of the molecule  and seven commercially available Nα‐substituted‐l‐arginine analogues R2–R8 were selected to explore their binding affinities and the solvation profile of the NRP1‐b1 ligand‐binding site experimentally. In addition, underivatised l‐homoarginine (R9) was also included in the study  method was used to determine the binding affinities of the ligands to NRP1‐b1, and crystallisation trials were performed with the preformed complexes between NRP1‐b1 and the individual ligands.l‐Arginine R1 and l‐homoarginine R9 showed weak binding to NRP1‐b1, with dissociation constants (KDs) of 325 ± 5 and 637 ± 70 μm, respectively. These are the weakest binding ligands tested, with all other Nα‐substituted analogues exhibiting higher affinities and lower corresponding dissociation constants. The data suggest that the free amine is not favourable for binding, and that the longer alkane chain on l‐homoarginine might de‐stabilise key interactions with the receptor. Interestingly, a large Nα substituent, as seen in FMOC‐l‐arginine (R3) also resulted in a weakly binding ligand, with a KD of 201 ± 106 μm; however, this molecule's poor solubility might have affected the results. Tosyl Nα‐substituted l‐arginine (R2), acetyl‐l‐arginine (R4), benzoyl‐l‐arginine (R6) and benzoylglycyl‐l‐arginine (R8) exhibited a medium range of affinities with KDs between 70 and 22 μm  and Cbz‐l‐arginine (R7) exhibited highest affinities for NRP1‐b1 in the series, with KDs of 3 ± 1 and 17 ± 1 μm, respectively.All arginine analogues showed very fast rates of association and dissociation to immobilised NRP1‐b1, as determined by surface plasmon resonance experiments Fig. . Bindingrginine R, benzoyl165 for binding to NRP1‐b1. R5, the Arg analogue with the highest affinity in the SPR assay, exhibited clear sigmoidal curve in a ligand displacement essay with an IC50 of 2.6 ± 0.38 μm, which correlates extremely well with the SPR data  space group with a dimer in the asymmetric unit for all complexes except for that of R5 analogue which contained a tetramer in the asymmetric unit. The structures were refined using data at the resolutions higher than 2.1 Å (for R5‐bound NRP1 b1); full data and refinement statistics are available in Table Co‐crystallisation studies were attempted with all ligands selected and six complexes produced good‐quality crystals. All six complexes crystallised in the monoclinc P2 space gr165 binding site is located at the top of the β‐barrel core delineated by six juxtaposed loops connecting the β‐strands , we were able to identify the arginine analogue bound to each of the b1 molecules within the asymmetric unit of the respective unit cells. In the structure of the complexes with the ligands R4–R8, the side chain guanidine group forms bidentate hydrogen bonds with Asp320. The aliphatic part of arginine is located in the groove between Tyr297 and Tyr353, with a conformation favouring a π–π stacking between the arginine side chain and the phenyl rings. The carboxylate group anchors the small molecules specifically through hydrogen‐bond interactions such that one of the oxygen atoms binds to the hydroxyl of Ser346, while the other carboxylate oxygen forms hydrogen bonds with the hydroxyl oxygen atoms from Thr349, and Tyr353. These interactions are analogous to that observed for the C‐terminus of the peptide ligands 4RN5).The NRP1‐b1 domain belongs to the discoidin structural domain family that also includes homologous FV/VIII C domains. The fold in this type of domain is characterised by an eight‐stranded distorted jellyroll β‐barrel where a five‐stranded antiparallel β‐sheet packs against a three‐stranded antiparallel β‐sheet. The VEGF‐AR9 is the only analogue without the Nα‐substituent, and with an extra methylene group in the side chain when compared to the natural amino acid arginine. Longer aliphatic chain affected a shift in the positioning of the carboxyl group such that it binds differently to all other ligands in the ligand‐bound structures presented here. Compared to other ligands, the carboxylate of R9 is moved by more than 2 Å towards the protein core such that one of the oxygen atoms interacts directly with the hydroxyl from Ser346, whereas the other oxygen atom forms direct hydrogen bond with the nitrogen atom from the side chain of the residue Trp301. Furthermore, the guanidine group now forms single hydrogen bond with Asp320 compared to bidentate interaction observed in the structures of the b1 complexes with the other arginine analogues. Although the free amino group of R9 engages in a hydrogen bond with Glu348, the total number of direct hydrogen bonds to the protein chain is four compared to five observed with the other analogues. This shift in binding mode is accompanied by a significant decrease in binding affinity, as shown by SPR.Among the obtained crystal structures of the complexes, Overall, the protein components of the six arginine analogue‐bound structures show high structural similarity, with only discernible conformational differences evident in an area of the b1 domain engaged in forming the intermolecular interaction within the dimer/tetramer of the asymmetric unit in the bound structures and the side chains of residues Tyr297 and Asp320 in the ligand‐binding site. The two side‐chain differences reflect the structural differences of the bound arginine analogues. The differences are best depicted by the changes in the torsion angle of the Y297 side chain as shown in Fig. The high‐resolution X‐ray crystal structures showed additional density consistent with water molecules in the binding site. To conduct an analysis of the conserved water molecules in the binding site of NRP1‐b1, we considered all new structures reported here together with the already publicly available structures. Only structures of human NRP1 with a resolution higher than 2.1 Å were selected for analysis. Table –5, on the molecular surface of the NRP1‐b1 ligand‐binding site were identified as containing the conserved water molecules  indicate that these are stable waters, and that their replacement would be favourable.The polypeptide chains were aligned to allow the visualisation of the crystallised water molecules in the binding site of NRP1‐b1. A water molecule was considered ‘conserved’ when observed in at least four out of ten superimposed molecules from the crystal structures (see methods for details). In total, five localised areas, labelled 1les Fig. A. It wasles Fig. . Free‐enles Fig. . In the −1) suggests that this cluster could also be target for displacement by a ligand's functional group. In fact, in the crystal structure of EG0229 (PDB code: 3I97), this site is occupied by a sulphonamide group. The conserved water molecule in site 4 is engaged in a hydrogen‐bond interaction with the nitrogen atom on the Trp301. This is observed in six structures. Additionally, this molecule is also seen engaging in hydrogen bonds with E348. In the structures with ligands R4–R8, the water molecule in site 4 additionally forms a hydrogen bond with the oxygen atom of the arginine analogue carboxyl group. However, in the structure of R9‐bound NRP1, the conserved water molecule in site 4 has been replaced by an oxygen atom of the R9 carboxylate. As mentioned above, R9 shows very weak binding affinity to the NRP1‐b1 domain, which can be interpreted by the loss of favourable direct hydrogen bonds of R9 to the polypeptide chain together with a disruption of a water network caused by the displacement of the water molecule in site 4. This carboxyl/oxygen–water–Trp301/nitrogen interaction might be a key network in the binding of ligands to NRP1‐b1. We suggest this water molecule, which corresponds to the highest free‐energy site , should be included in docking studies for drug design, and used as a bridge between the ligand and the protein instead of attempting to displace it or replace it with functional groups. The last conserved water molecule observed was in site 5, found in only four structures, and located near the protein surface, with a possible hydrogen bond to Tyr353. Since it is superficial water, we would not consider it significant for drug design.Sites 1 and 2 are occupied by water molecules in 9 out of 10 analysed structures. These are deep areas in the binding site, which could be classified as sub‐pockets, and could be targeted for replacement in drug design with the aim of enhancing affinity. The water molecule in site 3 is present in five structures. This molecule establishes a hydrogen bond with the side chain of Ser298 and is present in both apo‐ and ligand‐bound structures. Calculated free energy of site 3  contained carbamate group.We studied a known binding site in the NRP1‐b1 domain. To disrupt ligand binding to this domain, a protein–protein interaction, with a potent small molecule, it was important to perform an in‐depth study of the architecture and the solvation map of the target site. NRP1 shows a well‐defined arginine‐binding site where the guanidine part can establish a bidentate hydrogen‐bond interaction with the side chain of Asp320, and the C‐terminal carboxyl group can establish hydrogen bonds with the side chain oxygen atoms from Thr349 and Ser346. Although these interactions are conserved when small molecules containing a We obtained the crystal structures of six arginine analogues bound to NRP1‐b1 and analysed these alongside other publicly available structures in the PDB. The binding site of NRP1‐b1 shows little main chain and/or side chain flexibility, with only Tyr297 and Asp320 showing significant variations, which are dependent on the ligand bound. Tyr297 has been previously shown as required for VEGF binding, with NRP1Y297A/Y297A homozygous knock‐in mice exhibiting reduced VEGF‐induced angiogenesis and tumourigenesis R1, R2, R8 , R3, R7, R9 (AK Scientific), R4, R5, R6 , and used without further purification. All experiments were conducted using commercial reagents and standard procedures, unless stated otherwise.Arginine analogues were obtained from commercial sources: Molecular modelling was performed using the software package MOE .3I97. Water and co‐crystallised ligands were removed from the structure. An arginine residue was modelled from the coordinates of EG00229  and used to define the boundaries of the binding site. Hydrogen atoms were added to the structure by selecting the function ‘add hydrogen’ in MOE before the calculations were performed. The binding site for solvent prediction was defined by all atoms included in a 10 Å proximity from the ligand, extended by residues of those atoms.The structure used for the prediction of binding site hydration was PDB code: Binding site hydration prediction was carried out using the ‘Solvent Analysis’ function in MOE. The application uses the three‐dimensional reference interaction site model (3D‐RISM) method 3I97 (NRP1‐b1) as a template. Only PDBs with a resolution better than 2.1 Å and sequence identity higher than 90% were retained for analysis. This resulted in three PDBs, which are 4RN5, 2QQI and 1KEX. In cases where the X‐ray crystal structure had more than one chain, both chains were analysed. Combined with the six structures obtained in‐house, a total of 9 PDBs or 10 chains were analysed. Among these structures, 2QQI and 1KEX are in the apo form.The SAS tool (Sequence annotated by Structure) The protein binding site was defined by a 5 Å proximity of all the ligand atoms present. Only the water molecules observed in the binding site were analysed. A water molecule is considered ‘conserved’ when observed in the same position of at least four superimposed crystal structures.Surface plasmon resonance experiments were performed using a Biacore 4000 instrument at a constant temperature of 25 °C. Sensor chips, buffer stock solutions and immobilisation reagents were purchased from GE Healthcare.N‐hydroxysuccinimide for 10 min. NRP1‐b1 at a concentration of 20 μg·mL−1 in 10 mm sodium acetate pH 5 was injected onto the surface for 10 and 5 min in spots 1 and 2, respectively, to generate surfaces with high and low density. The immobilisation levels ranged from 3487 to 3652 resonance units (RU) on spot 1 and from 1275 to 1756 RU on spot 2. Spot 3 was left unmodified and used as a reference.PBS containing 0.05% surfactant P20 was used as the running buffer during immobilisation. NRP1‐b1 was immobilised onto a CM5 chip using random amine coupling. The four flow cells were treated in the same way to optimise throughput. In summary, immobilisation spots 1 and 2 were activated with the coupling reagents, 1‐ethyl‐3‐(3‐dimethylaminopropyl)carbodiimide and m. Dilution series were then optimised for three compounds to obtain a more accurate measurements and dissociation constants. R1 and R9 were tested at 12–1500 μm, and R5 was tested at 0.16–20 μm. Binding responses to high‐ and low‐density surfaces were processed independently and the average ± SD is presented. KDs reported are derived from steady‐state binding responses assuming a 1 : 1 interaction and therefore correspond to the equilibrium binding affinity of the compounds.PBS containing 0.05% surfactant P20 was used as the running buffer and sample dilution buffer throughout these experiments. Dose‐responses were obtained using a two‐fold sample dilution, generating eight data points for the concentration range, and using an injection time of 60 s. Surface regeneration between injections was not necessary, but a wash step with 1M NaCl was included after injection of the highest concentration sample for each compound. All compounds were initially tested at a wide dilution range from 0.6 to 300 μE. coli strain Rosetta‐gami2‐(DE3)pLysS (Novagen) and purified as previously described –10 mg·mL−1 and 10 μL was mixed with 1 μL ligand at 10–100 mm in water. The complex was crystallised in a 1 : 1 volume mixture with 10–30% w/v PEG 3350, and 0.2 m ammonium chloride at 16 °C using hanging drop vapour diffusion method and micro‐seeding with 1/10 dilution of seeds of apo‐NRP1‐b1 crystals. Seeds were prepared according to Hampton Research's seeding kit. Crystals appeared within 2–3 days with needle morphology. A single crystal was transferred to a solution containing the crystallisation condition plus 20 % v/v ethylene glycol and was afterwards flash‐frozen in liquid nitrogen.NRP1‐b1 was expressed in 1KEX) as the search model. The refinement was carried out by Refmac5 X‐ray diffraction data were collected at 100 K on beamline I04 at the Diamond Light Source, Didcot UK. All data sets were processed with xia2‐3d automated software system −1 overnight at 4 °C. On the following day, the plates were treated with blocking buffer (PBS containing 1% BSA) and washed three times with wash buffer (PBS containing 0.1% Tween‐20). The various concentrations of compounds diluted in PBS containing 1% DMSO were added, followed by addition of 0.25 nm bt‐VEGF‐A165. After 2 h of incubation at room temperature, the plates were washed three times with wash buffer. The bt‐VEGF‐A165 bound to NRP1‐b1 was detected by streptavidin‐horseradish peroxidase conjugates and the enzyme substrate and measured using a Tecan Genios plate reader at 450 nm absorbance with a reference wavelength at 595 nm. Non‐specific binding was determined in the absence of NRP1‐b1‐coated wells of the plates and subtracted from the readings obtained in the presence of NRP1‐b1. GraphPad Prism was used to carry out a non‐linear regression analysis and to generate inhibition curve and IC50 value. IC50 values for the particular ligand were obtained based on the three separate experiments.The assay was run as previously described DS, IZ, PF and SD were consultants for Magnus Life Science.CF, RR and TY carried out the X‐ray crystallography; AT assisted with protein preparation. EC did the computational work. JH and DS selected the compounds for analysis. AOL carried out VEGF displacement assays. FM performed the SPR experiments, analysed the results and wrote the paper. DS, IZ, PF and SD supervised the research team and edited the paper.


Results showed that the W–1vol.%Zr/2vol.%Sc2O3 composites exhibited optimal performance with the best relative density of up to 98.93% and high Vickers microhardness of approximately 583 Hv. The thermal conductivity of W–Zr/Sc2O3 composites decreased initially and then increased as the Zr content increased. The moderate Zr alloying element could combine well with Sc2O3 particles and W grains and form a solid solution. However, excess Zr element leads to agglomeration in the grain boundaries. W–1vol.%Zr/2vol.%Sc2O3 composite had a good deuterium irradiation resistance very closing to pure tungsten compared with the other Zr element contents of composites. Under 500 K, D2 retention and release of them were similar to those of commercial tungsten, even lower between 400 K to 450 K. Pre-irradiation with 5 keV-He+ ions to a fluence of 1 × 1021 He+/m2 resulted in an increase in deuterium retention (deuterium was implanted after He+ irradiation), thereby shifting the desorption peak to a high temperature from 550 K to 650 K for the W–1vol.%Zr/2vol.%Sc2O3 composite.Dense W and W–Zr composites reinforced with Sc Nuclear fusion energy is considered to be the principal way to effectively solve the future energy problem as a clean and infinite energy resource. And it is being developed internationally via the International Thermonuclear Experimental Reactor (ITER) Project, which aims to demonstrate the extended burn of deuterium-tritium (D–T) plasma in a fusion reaction134672, La2O3, CeO2, Y2O3, and TiC) can effectively inhibit recrystallization and grain growth as well as improve high-temperature strength and creep resistance by hindering grain boundary (GB) sliding9101112151617Many studies have shown that several disperse second-phases particles  into the first wall results in irradiation defects upon fusion reactor operation1922232O3 particles dopant can not only refine the grains and increase the tungsten alloy density but can also improve the strength of the samples2O3 powder was prepared in the current study through powder metallurgy and then consolidated through spark plasma sintering (SPS), which can produce fine-grain and high-density materials at a relatively low temperature34A previous study showed that Sc2/Sc2O3 powder  were set to investigate the effects of different contents of Zr alloying element on the microstructure and properties of the composites. An omnibearing ball milling process was utilized to mix the powder in stoichiometric proportions uniformly and to diminish the size of particles under a purified N2 atmosphere for 40 h, which seemed to be an appropriate time for both mixing and refining powders. A ball-to-powder weight ratio of 10:1 and a rotation speed of 400 rpm were adopted. Powder collection was conducted in a vacuum glove box with ultimate vacuum of 10 Pa to prevent powders from being exposed to the air as much as possible. The W–ZrH2/Sc2O3 powder was consolidated through SPS  at 1700 °C under 58 MPa pressure for 3 min in Ar + 3% H2 atmosphere. Then it was loaded in an electrically and thermally conductive graphite die with a diameter of approximately 20 mm. DC current was applied in pulses together with uniaxial high mechanical pressure during the process. The temperature and pressure profiles of the sintering program in this study are illustrated in 2 to decompose completely into Zr and H2. Afterward, the powder samples were maintained at 1350 °C  for 10 min to allow Zr and Sc2O3 to disperse uniformly. Finally, the specimens were heated at 1700 °C for 3 min and then cooled down. The heating and cooling rates in the sintering process were both 100 °C/min. The pressure was increased from 15 Mpa to 58 Mpa whereas the temperature was increased from 700 °C to 1350 °C; both were maintained afterward until the sintering process was over. The samples were approximately 20 mm in diameter and 3 mm in thickness.Three different contents of W–ZrH2O3 were 19.35 g/cm3, 6.49 g/cm3, and 3.86 g/cm3, respectively. The polished sintered samples were subjected to Vickers microhardness testing with MH–3L by measuring from the center to the sample edges with a loading weight of 300 g held for 15 s .The SPS-sintered samples were machined through wire-electrode cutting. The density of the sintered samples was measured with Archimedes method. The relative densities were calculated from the volume fraction. The theoretical densities of W, Zr, and ScA thermal conductivity test was performed with a laser flash thermal analyzer ; the disk samples had 6 mm diameter and 2 mm thickness. Thermal conductivity (λ) was calculated from the thermal diffusivity (α), density (ρ), and specific heat (Cp) through the expression λ = αCpρ. The specific heat capacity was determined with the theoretical rule of mixtures according to the following formula:W, VZr, and VSc2O3 are the volume fractions of W, Zr, and Sc2O3, respectively, whereas CPW, CPZr, and CPSc2O3 are the specific heat values of W, Zr, and Sc2O3, respectively. The data of CPW, CPZr, and CPSc2O3 at different temperatures were obtained from ref. V2O3 composites were mechanically polished until the surface was mirror-like. These composites were irradiated with D2+–only with 5 keV in extremely high vacuum and the irradiation dose was 1 × 1020 ions/m2, according to the previous experimental results. W–1vol.%Zr/2vol.%Sc2O3 composites were also exposed to He+ ions with a fluence of 1 × 1021 He+/m2 and subsequent D2 ions with a fluence of 1 × 1020 D2+/m2. After these ion implantations, TDS measurements were performed from room temperature (RT) to 900 K via infrared irradiation with a heating rate of 1 K/s to investigate the D retention behavior. In all cases, the samples were stored in air between the end of exposure and the beginning of the TDS measurement372O3 composite prepared with ion-thinning technology because of its highest second phase content which was easier to observe. Fractured surface was obtained by breaking up the samples at RT artificially. The grain sizes were measured from the SEM images of the fractured surface morphology of the SPS-sintered samples.The deuterium inventory in the samples was examined through nuclear reaction analysis (NRA) and thermal desorption spectroscopy (TDS). The W–Zr/Sc2O3, ZrH2, and the as–prepared powders were observed through FE–SEM equipped with EDS. The different powders shown in 2O3, ZrH2, and as–prepared powders were 3.954 μm, 16.699 μm, 2.396 μm, and 1.472 μm, respectively, as shown in Before sintering, W, Sc2O3 composites that were all higher than 95%, particularly the W–1vol.%Zr/2vol.%Sc2O3 samples which exhibited a relative density of 98.93%, as listed in 2O3 composites seems to apparently decreased with the addition of Zr. The microhardness of samples was ruleless; it decreased after the first increase with the addition of Zr. The highest average hardness is 614.9 Hv, which corresponded to the W–3vol.%Zr/2vol.%Sc2O3 samples. The lowest microhardness of 456.1 Hv, which corresponded to the W–5vol.%Zr/2vol.%Sc2O3 samples, may be attributed to the low relative density and excess Zr, which led to the agglomerations at GBs. The average grain size of the composites was obtained according to the fracture morphologies of the sintered samples. The results demonstrated that the grain size of the samples with different Zr content was inconspicuous at approximately 1–2 μm. Therefore, the grain size of the samples containing much Zr element for oxidation could not be observed with ease.The relative density, average grain size, and Vickers microhardness of pure W and the samples are shown in 2O3, W–3vol.%Zr/2vol.%Sc2O3, and W–5vol.%Zr/2vol.%Sc2O3 samples with HA probe and the energy spectrum diagrams of the second-phase particles in the W–1vol.%Zr/2vol.%Sc2O3 samples. These images confirm that the doped phase could distribute in the samples uniformly. Energy dispersive X-ray spectroscopy was also measured to analyze the black dots qualitatively. The corresponding energy spectrum diagrams of the pane region e and f in 2O3 composite, except Sc2O3 particles. This observation reveals that W, Sc2O3, and Zr can combine well with one another to form a solid solution. Since the content of Zr in W–1vol.%Zr/2vol.%Sc2O3 sample is less than W–3vol.%Zr/2vol.%Sc2O3 and W–5vol.%Zr/2vol.%Sc2O3 samples, there are still many Sc2O3 particles exist alone with large particle size. With the increasing of Zr, it can bond well with Sc2O3 particles and W, form a solid solution and make the size of second phase tinier. However, in the W–5vol.%Zr/2vol.%Sc2O3 sample, the content of Zr element is excess, which leads to the agglomerations in the GBs.2O3 particles hindered the growth of tungsten grains. The moderate Zr purified the GBs while strengthening the tungsten grains and increasing the interface bonding. Their combination improved the strength and toughness of tungsten to some extent. First-principles calculations also indicated that the cohesion effect of transition metals, like Zr, Hf, and Re on W GBs could significantly strength the GBsThe SEM images of the fracture surface of the sintered composites at RT are presented in 2O3 dispersed phase and Zr alloying element, as shown in 2O3 particle and particle A. The interplanar spacing of each phase, as measured from the high resolution TEM (HRTEM) images, corresponded well with the calculated lattice spacing of W (220) and Sc2O3 (400), (0–62), and (23–1), as shown in 2O3 particle was still present even though there were excess Zr in the W–5vol.%Zr/2vol.%Sc2O3 composite. Particle A  in the GBs and could capture these impurities and react to form compounds. Since the compound was ZrO2, the impurities were mostly oxygen. Zr alloying element could reduce the O content in the alloy, enhance the cohesion of the GBs, and improve the distribution of elements in the alloy, thereby improving alloy mechanical properties. Other compounds, such ZrC or ZrN were also present theoretically. Some tiny particles existed in the W grains  in the GBs to form compounds to a certain degree. The formed oxides and carbides strengthened the GBs and hindered thermal conduction, which were the roles of Sc2O3. It corresponds well with those from the TEM analysis  in the GBs and could capture these impurities and react to form compounds, which could purify the GBs, enhance the cohesion of the GBs, and reduce the vacancies of material itself. The solid solution combined with W, Zr, and Sc2O3 can also enhance the cohesion of particles with tungsten and reduce the vacancies. On the other hand, the relative density of 98.93% of W–1vol.%Zr/2vol.%Sc2O3 composites also signifies less defects comparing with the 98.6% of W–2vol.%Sc2O3 composites. Large D2 peaks were observed clearly for W–3vol.%Zr/2vol.%Sc2O3 and W–5vol.%Zr/2vol.%Sc2O3 composites at a temperature ranging from 400 K to 550 K and 450 K to 700 K, respectively. With the relative density decreased, defects of materials themselves increased served as trapping sites for deuterium, which could generate to defect clusters with temperature increasing and result in the peak of desorption rate shifted to the high temperature side and became higher. Excess Zr element would agglomerate in the GBs according to the SEM analysis, which also leading to the increase of defect clusters such as vacancies and reducing the resistance to radiation. This finding suggests that excessive Zr content is not beneficial for the D2 ion irradiation resistance of W alloys. It reveals that moderate Zr can purify the GBs, enhance the cohesion of the GBs, and combine with W and Sc2O3, therefore reduce the vacancies of material itself, which improves the irradiation property, thereby drastically reduces the peak of desorption rate from W–2vol.%Sc2O3 samples.The D2+ ions at RT was measured via TDS to investigate the effects of helium irradiation on deuterium retention properties in W–1vol.%Zr/2vol.%Sc2O3. The samples were pretreated with He+ ions with a fluence of 1 × 1021 He+/m2 and subsequent D2+ ions with a fluence of 1 × 1020 D2+/m2. 2 thermal desorption spectra for W–1vol.%Zr/2vol.%Sc2O3 composite exposed to D2+ only and He+–D2+. The significant difference in the retention properties between the D2+ only and He+–D2+ samples was caused by differences in the microstructures of the samples. Only small dislocation loops were formed in the sample irradiated with D+ ions, whereas both dislocation loops and bubbles were formed in the sample irradiated with He+ ions32392+ only irradiation, the W–1vol.%Zr/2vol.%Sc2O3 composites treated with both He+ and D2+ ion irradiation had higher desorption rate from 550 K to 900 K. The D2 desorption peak of He+–D2+ samples shifted to a higher temperature from 550 K to 650 K compared with D2+–only samples. These results demonstrate that pre-implantation of energetic He+ ions at RT induced a significant increase in D retention in W. Energetic He+ ions can cause structural damage directly through elastic collisions with target atoms2O3 samples and Japanese commercial tungsten irradiated by D2+ only and He+–D2+ are also shown in The deuterium from the samples irradiated with a sequential irradiation of 5 keV-D2O3 composites were sintered through SPS. After sintering, their relative density reached more than 95% and even 98.93%, which was the highest value for the W–1vol.%Zr/Sc2O3 composite. The microhardness of W–3vol.%Zr/Sc2O3 was 614.9 Hv, which was greater than that of the other two groups of samples. The grain sizes of composites were generally fine (approximately 1–2 μm) and had a slight difference. The thermal conductivity of the samples decreased as the temperature increased from 300 K to 1100 K. With increasing Zr content, the thermal conductivities of samples initially decreased and then increased. The results revealed that the Zr alloying element can combine well with Sc2O3 particles and W grains, thereby forming a solid solution. However, excessive Zr element led to agglomerations in the GBs.W–Zr/Sc2O3 composites possessed optimal mechanical and irradiation properties. The D2 retention and release of W–1vol.%Zr/2vol.%Sc2O3 composites were similar to those of Japanese commercial tungsten exhibiting satisfactory irradiation resistance compared with the other samples. Comparing with the different phenomenon of W–2vol.%Sc2O3 and W–1vol.%Zr/2vol.%Sc2O3 composites on deuterium retention properties, GBs seem to be the dominant factor. Moderate Zr element can purify the GBs, enhance the cohesion of the GBs, and combine with W and Sc2O3, therefore reduce the defects of material itself such as vacancies, which improves the irradiation property. However, pre-irradiation with 5 keV-He+ ions to a fluence of 1 × 1021 He+/m2 resulted in an increase in deuterium retention (the deuterium was implanted after He+ irradiation). Thus, the desorption peak shifted to a high temperature. This finding demonstrates that pre-irradiation with He+ ions can promote D retention by affecting surface topography and by adding dislocation loops and bubbles as trapping sites for deuterium.The W–1vol.%Zr/2vol.%ScHow to cite this article: Chen, H. et al. Effects of zirconium element on the microstructure and deuterium retention of W–Zr/Sc2O3 composites. Sci. Rep.6, 32678; doi: 10.1038/srep32678 (2016).


In these 5–3 silicides Nb atoms can be substituted by Ti atoms. The type of stable Nb5Si3 depends on temperature and concentration of Ti addition and is important for the stability and properties of the alloys. The effect of increasing concentration of Ti on the transition temperature between the polymorphs has not been studied. In this work first-principles calculations were used to predict the stability and physical properties of the various Nb5Si3 silicides alloyed with Ti. Temperature-dependent enthalpies of formation were computed, and the transition temperature between the low (α) and high (β) temperature polymorphs of Nb5Si3 was found to decrease significantly with increasing Ti content. The γNb5Si3 was found to be stable only at high Ti concentrations, above approximately 50 at. % Ti. Calculation of physical properties and the Cauchy pressures, Pugh’s index of ductility and Poisson ratio showed that as the Ti content increased, the bulk moduli of all silicides decreased, while the shear and elastic moduli and the Debye temperature increased for the αNb5Si3 and γNb5Si3 and decreased for βNb5Si3. With the addition of Ti the αNb5Si3 and γNb5Si3 became less ductile, whereas the βNb5Si3 became more ductile. When Ti was added in the αNb5Si3 and βNb5Si3 the linear thermal expansion coefficients of the silicides decreased, but the anisotropy of coefficient of thermal expansion did not change significantly.Nb-silicide based alloys could be used at T > 1423 K in future aero-engines. Titanium is an important additive to these new alloys where it improves oxidation, fracture toughness and reduces density. The microstructures of the new alloys consist of an Nb solid solution, and silicides and other intermetallics can be present. Three Nb


Interference with this will impair relaxation. The factors responsible for regulation of diastolic [Ca2+]i, in particular the relative roles of the sarcoplasmic reticulum (SR) and surface membrane, are unclear. We investigated the effects on diastolic [Ca2+]i that result from the changes of Ca cycling known to occur in heart failure. Experiments were performed using Fluo‐3 in voltage clamped rat ventricular myocytes. Increasing stimulation frequency increased diastolic [Ca2+]i. This increase of [Ca2+]i was larger when SR function was impaired either by making the ryanodine receptor leaky (with caffeine or ryanodine) or by decreasing sarco/endoplasmic reticulum Ca‐ATPase activity with thapsigargin. The increase of diastolic [Ca2+]i produced by interfering with the SR was accompanied by a decrease of the amplitude of the systolic Ca transient, such that there was no change of time‐averaged [Ca2+]i. Time‐averaged [Ca2+]i was increased by β‐adrenergic stimulation with isoprenaline and increased in a saturating manner with increased stimulation frequency; average [Ca2+]i was a linear function of Ca entry per unit time. Diastolic and time‐averaged [Ca2+]i were decreased by decreasing the L‐type Ca current (with 50 μm cadmium chloride). We conclude that diastolic [Ca2+]i is controlled by the balance between Ca entry and efflux during systole. Furthermore, manoeuvres that decrease the amplitude of the Ca transient (without decreasing Ca influx) will therefore increase diastolic [Ca2+]i. This identifies a novel mechanism by which changes of the amplitude of the systolic Ca transient control diastolic [Ca2+]i.The intracellular Ca concentration ([Ca 2+]i) must increase during systole to activate contraction and then fall, during diastole, to allow the myofilaments to relax and the heart to refill with blood.For the heart to function as a pump, intracellular calcium concentration ([Ca2+]i in rat ventricular myocytes.The present study investigates the control of diastolic [Ca2+]i is increased by manoeuvres that decrease sarcoplasmic reticulum function. This is accompanied by a decrease of systolic [Ca2+]i such that the time‐averaged [Ca2+]i remains constant.We show that diastolic [Ca2+]i is controlled by the balance between Ca2+ entry and Ca2+ efflux during systole.We report that diastolic [Ca2+]i.The results of the present study identify a novel mechanism by which changes of the amplitude of the systolic Ca transient control diastolic [Ca Although the mechanisms that control the amplitude of the systolic rise of [Ca2+]i are well understood    Act, 1986 and Directive 2010/63/EU of the European Parliament. The experiments were approved by the University of Manchester Ethical Review Board. Male Wistar rats (weighing ∼200–250 g) were killed by stunning and cervical dislocation. Single ventricular myocytes were isolated by digestion with collagenase and protease as described previously (Eisner m) 135 NaCl, 11.1 glucose, 1 CaCl2, 10 Hepes, 1 MgCl2 and 4 KCL. Then, 4‐aminopyridine (5 mm) and BaCl2 (0.1 mm) were added to inhibit K+ currents and the solution was titrated to pH 7.4 using NaOH. Probenecid (2 mm) was added to reduce loss of indicator from the myocytes. Micropipettes (< 5 MΩ) were filled with a solution consisting of (in mm): 125 KCH3O3S, 12 KCl, 10 NaCl, 10 Hepes, 5 MgCl2 and 0.1 EGTA; titrated to pH 7.2 with KOH; and a final concentration of amphotericin B of 240 g ml−1. Cells were voltage clamped with the perforated patch clamp technique using the discontinuous switch clamp mode (frequency 1–2 kHz and gain 0.3–0.7 nA mV−1) of an Axoclamp 2A voltage clamp amplifier . Cells were voltage clamped and stimulated at a range of frequencies (0.2 to 3 Hz) with a 40 mV, 100 ms duration pulse from a holding potential of −40 mV. All experiments were performed at room temperature.Isolated myocytes were superfused with a solution (control) consisting of (in mm for 10 min) and allowed to de‐esterify before use. An aliquot was then placed in a superfusion chamber mounted on the stage of an inverted fluorescence microscope. To measure changes of [Ca2+]i, at the end of each experiment, the maximum fluorescence (Fmax) was measured by damaging the cell with the patch pipette. The dissociation constant of Fluo‐3 (Kd) was taken as 864 nm , R&D Systems  or Fisher Scientific . Caffeine was added as required. Ryanodine and thapsigargin were both stored as 1 mF test. P < 0.05 was considered statistically significant.Data are reported as the mean ± SEM where applicable for descriptive analysis. Statistical comparisons were made using two‐way ANOVA. Regression lines were compared with an A show the effect of periods of stimulation at 2 Hz. Under basal conditions, the diastolic level of [Ca2+]i during the 2 Hz stimulation was only slightly greater than that during rest. The subsequent addition of caffeine (1 mm) to increase RyR leak decreased the amplitude of the Ca transient and modestly elevated diastolic [Ca2+]i. The effects of β‐adrenergic stimulation with isoprenaline (ISO) (1 μm) were then investigated on these phenomena. In agreement with previous work, ISO increased the amplitude of the systolic Ca transient  show values normalized to those obtained at a slow stimulation rate (0.5 Hz) in the absence of both caffeine and ISO. It is clear that the effects of ISO on both systolic and diastolic [Ca2+]i are greatly affected by the presence of caffeine. For example, in the absence of caffeine, ISO greatly increases the amplitude of the Ca transient , whereas there is no effect on diastolic [Ca2+]i . By contrast, in the presence of caffeine, ISO had a much smaller effect on the amplitude of the Ca transient  but markedly increased diastolic [Ca2+]i .The first series of experiments examined the effect of increased RyR leak on the response to an increased stimulation rate. The grey traces in Fig. caffeine  mm to incaffeine  mm to incaffeine  mm to in2+]i at the same time as decreasing the systolic rise of [Ca2+]i. Accordingly, we next investigated the effect of caffeine on time‐averaged [Ca2+]i as shown by the blue trace in Fig. A. It is clear that average [Ca2+]i is: (i) increased by stimulation; (ii) increased by ISO; and (iii) unaffected by caffeine in both control and ISO. These observations are confirmed by the mean data of Fig. D, which show that the average [Ca2+]i does not significantly change upon addition of caffeine . A similar finding is seen in the presence of ISO . Finally, these data also show that ISO increases average [Ca2+]i in both the presence and absence of caffeine (P < 0.001).The above data show that caffeine increases diastolic [Ca2+]i were more prominent in the presence of ISO, all subsequent experiments were performed in the presence of ISO. The experiment shown in Fig. A shows the data obtained in the presence and absence of caffeine. As the frequency of stimulation increased, the Ca transient amplitude decreased slightly and diastolic [Ca2+]i increased. The decrease in Ca transient amplitude was accompanied by (and is presumably at least in part caused by) a decrease in the amplitude of the L‐type Ca current  . Average [Ca2+]i did, however, increase in a saturating manner with an increasing frequency of stimulation. Figure B sheds light on this saturation of average Ca. The Ca influx via the L‐type Ca current on each pulse decreased with an increasing stimulation rate  as a result of increasing inactivation  was a saturating function of frequency and therefore paralleled the frequency dependence of average [Ca2+]i. Data obtained in the presence of caffeine  were identical to those in its absence . The correlation between Ca entry per unit time and average [Ca2+]i is emphasized by Fig. D, which shows a clear linear relationship between these two parameters that is not statistically different  in the absence and presence of caffeine.The above observations of the frequency dependence of average Ca+i are emA, the application of thapsigargin decreased systolic [Ca2+]i and increased diastolic [Ca2+]i during stimulation at 0.5 Hz. On increasing the stimulation rate to 2 Hz, there was a more marked increase of diastolic and decrease of systolic [Ca2+]i. When stimulation was stopped, [Ca2+]i declined to a level similar to the original diastolic one. The mean data of Fig. B show the effects of thapsigargin at 0.5 and 2 Hz. . Thapsigargin decreased the amplitude and increased diastolic [Ca2+]i at the same time as having no effect on average [Ca2+]i.The purpose of the experiments shown in Figs A, the application of ryanodine decreased systolic [Ca2+]i and increased diastolic [Ca2+]i. The dark line shows that average [Ca2+]i remained constant during this period. When stimulation was stopped, [Ca2+]i declined to a level similar to the original diastolic one. Subsequent stimulation at 2 Hz increased average [Ca2+]i. Ryanodine was reapplied to increase its effect. When stimulation was recommenced at 0.5 Hz, the Ca transient was considerably smaller than that observed previously at this frequency. The average [Ca2+]i was, however, almost identical. Increasing the stimulation rate to 2 Hz resulted in a smaller Ca transient than that seen at 2 Hz before. Once again, the average Ca was unaffected. These data were interrupted by exposures to caffeine plus 2,3‐butanedione monoxime to release Ca from the SR and thereby obtain an estimate of the fall of SR Ca content. The mean data of Fig. B confirm that ryanodine decreases the amplitude of the Ca transient  and increases diastolic [Ca2+]i , whereas there is no effect on average [Ca2+]i .In the experiment shown in Fig. 2+]i depends on a balance between Ca influx and efflux. If this is the case, diastolic [Ca2+]i would be expected to be decreased by reducing Ca influx. In the experiment shown in Fig. A, caffeine had been added when stimulating at either 0.5 or 3 Hz. In agreement with the results reported above, the rise of diastolic [Ca2+]i was greater at the higher frequency. Addition of cadmium chloride (50 μm) to decrease the L‐type Ca current was observed to decrease diastolic [Ca2+]i at both stimulation rates . On average, cadmium decreased the L‐type Ca current to 33 ± 5% (n = 5 cells). Figure C shows the mean data obtained from five cells (at 0.5 Hz). One‐way ANOVA showed that cadmium decreased average and diastolic [Ca2+]i, as well as the amplitude of the Ca transient . Specifically, cadmium decreased the average [Ca2+]i to 69.6 ± 1.7%.The above experiments suggest that the level of diastolic [Ca2+]i in rat ventricular myocytes. In agreement with previous work, we find that increasing stimulation frequency increases diastolic [Ca2+]i  this is accompanied by a decrease of systolic [Ca2+]i, such that the time‐averaged [Ca2+]i remains constant; (iii) time‐averaged [Ca2+]i is increased by β‐adrenergic stimulation and is a saturating function of frequency reflecting a frequency‐dependent decrease of Ca influx per beat; and (iv) diastolic [Ca2+]i is regulated by the difference between the Ca influx and the systolic efflux on each beat.The present study investigates the control of diastolic [Ca2+]i . This may be because inhibition of SERCA greatly slows the decay of the Ca transient. At low rates, the increased duration of the Ca transient compensates for the decreased amplitude in maintaining efflux such that diastolic [Ca2+]i does not increase. This is impossible at high rates because the decay of the transient is interrupted by the next stimulus.Previous work has shown that ryanodine and thapsigargin increase diastolic [Ca+]i Figs , 2 and 4gin Fig. B. This 2+]i such that average [Ca2+]i is maintained constant when SR function is altered? We suggest that the answer resides in the mechanisms that maintain Ca flux balance. In the steady‐state, during each cycle of stimulation, the influx mediated by the Ca current must be precisely equal to the Ca efflux, largely via sodium‐calcium exchange (NCX) . It is also possible that an increase of Ca leak from the SR, possibly via a Ca/calmodulin‐dependent kinase II (CaMKII) mechanism  increased Ca influx per unit time ; (ii) decreased systolic efflux; and (iii) decreased diastolic time for efflux to occur in. Ca flux balance can therefore only be established with an increase of diastolic [Ca2+]i. If SR function is depressed, increased frequency produces a larger decrease of systolic [Ca2+]i and therefore a larger rise of diastolic [Ca2+]i would be expected. In agreement with these predictions, we find that increasing frequency increases diastolic [Ca2+]i and this increase is potentiated by increasing SR leak or decreasing SERCA activity. Average [Ca2+]i is a saturating function of frequency  presumably because the Ca entry per unit time also saturates with frequency as a result of increased inactivation of the L‐type Ca current , . It should be noted that, if the L‐type Ca current did not decrease at higher frequencies, the rise of diastolic [Ca2+]i would be even greater.An increase of stimulation frequency will increase Ca influx per unit time, thereby requiring an increase of average Ca+i to maiD. It is clear that, even with zero influx through the L‐type Ca current, [Ca2+]i has a finite value. This has been accounted for by a background Ca entry  may have on NCX activity. [Na+]i will be increased by an increase of stimulation rate and as a consequence of the increased NCX activity in response to the Ca loading produced by β‐adrenergic stimulation. It will be decreased as a consequence of phosphorylation of phospholemman and stimulation of the sodium pump  a component activated by the diastolic level of [Ca2+]i and (ii) an additional component that occurs during the systolic Ca transient .


Using the CALYPSO structure searching method, multicenter bonding H3+ ions were stabilized under high pressure in a hydrogen rich H–Cl system. nCl, n = 2–7) compounds. In comparison to the group 1 and 2 hydrides, the structural modification associated with increasing pressure and hydrogen concentration is much less dramatic. The polymeric HCl chains already present in the low temperature phase under ambient pressure persist in all the high pressure structures. No transfer of electrons from the chlorine atoms into the interstitial sites is found. This indicates the chemical bonding at high pressure in group 17 elements is fundamentally different from the alkali and alkaline elements. It is found that almost perfectly triangular H3+ ions can be stabilized in the crystalline structure of H5Cl.The particle-swarm optimization method has been used to predict the stable high pressure structures up to 300 GPa of hydrogen-rich group 17 chlorine (H Although, the predicted hydrides have yet to be confirmed by experiment, the theoretical studies have revealed a myriad of novel structures for the high pressure hydrides and enriched the understanding of the nature of chemical bonding in the rarely explored high pressure regime.The suggestion that the hydrides of main group elements at high pressure may be superconductors with high critical temperature has stimulated recent interest in the search for the possible existence of hydrogen-rich alloys by theoretical and experimental means.6 Most of the structures and structural trends can be explained from the simple concept of electron transfer from the metal to the hydrogen due to the large electronegativity differences between the alkali and alkaline elements and hydrogen molecules. The predicted compounds display rich H species distinguished from the well-known H-ion in hydrides at traditional stoichiometric ratios. Perhaps one of the most exciting predictions is the emergence of symmetric and linear H3– at high pressures as observed in dense CsH3 and BaH6.8 Moreover, the formation pressures of these compounds of just a few tens of GPa are accessible by experiments. In comparison, the bonding pattern is quite different for group 14 and transition elements. For example, a Van der Waals solid with such molecular H2 units was found experimentally in SiH4 at low pressure.9 At higher pressure, the atoms of group 14 elements tend to aggregate to form a 2D layered structure decorated with molecular like H2 species as predicted for SiH4 and SnH4.11 In comparison, the high pressure chemistry of hydrogen with electron-rich group 17 halogens has not been investigated. In this paper, we present results on a study of the crystal structures and phase stabilities of hydrogen-rich HCl–H2 system. A major finding is the stabilization of cationic (H3)+ (H2) species in H5Cl. The geometry of H3+ becomes almost an equilateral triangle under very high pressure.The hydrides of group 1 and 2 elements formed from the reaction of the metal and hydrogen molecules have been the most studied.3+ in the solid state is new and significant. The isolated molecule is important in various branches of science, such as physics, chemistry and astronomy. For example, it is known that the H3+ ion with a triangular configuration is stable in the interstellar medium thanks to the low temperature and low density of the interstellar space and the H3 molecule is commonly formed from the neutralization reaction of H3+ and an electron, and rather evanescent as a result of the repulsive nature of its ground state.12 Furthermore, the multicenter (H3)+ (H2) bond in group 17 Cl compounds signifies a deviation in the nature of chemical bonding from the charge transfer interactions in group 1 and 2 hydrides and covalent bonding in group 14 hydrides.The observation of a triatomic hydrogen cation HnCl system was based on global minimization of free energy surfaces using ab initio total energy calculations and the particle–swarm-optimization scheme as implemented in the CALYPSO  code.14 The performance and reliability of this method has been demonstrated on many known systems. An example is the success on the prediction17 of an insulating orthorhombic  structure of Li and, more recently, two low-pressure monoclinic structures of Bi2 Te3. In both cases, the predicted structures were later confirmed by experiments.19 Structural searching was performed at 100, 200, and 300 GPa with a simulation cell consisting of 1–4 formula units. Ab initio electronic structure calculations and structural relaxations were carried out using density functional theory with the Perdew–Burke–Ernzerhof (PBE) exchange-correlation20 implemented in the Vienna ab initio Simulation Package (VASP) code.21 The predicted stable structures were carefully optimized at a high level of accuracy. A plane wave energy cutoff of 1000 eV was employed. Large Monkhorst-Pack k point sampling grids22 were used to ensure that all the enthalpy calculations were well converged to an accuracy of 1 meV per atom. The atomic charges were obtained from Bader topological analysis25 with very large grids to ensure sufficient accuracy. We also performed additional calculations employing the DF-2 van der Waals (vdW) functional26 to validate the results, in particular for the low pressure structures. Phonons were calculated with the supercell method27 implemented in the PHONOPY program.28 In essence, from finite displacements, the Hellmann–Feynman atomic forces computed at the optimized supercell by the VASP code were input to the PHONOPY code to construct the dynamical matrix. Diagonalization of the dynamical matrix gives the normal modes and their frequencies. Converged results were obtained with the use of a 2 × 2 × 1 supercell and 4 × 4 × 6 k-meshes for the Cc structure, and a 2 × 2 × 1 supercell and 4 × 4 × 6 k-meshes for the C2/c structure.The search for stable high pressure structures of the HnCl polymers, the relative energetics of the H–Cl system with H-rich stoichiometry from 100 to 300 GPa are summarized in the convex hull plot shown in nCl structure with solid HCl and H2 at the selected pressures. Since hydrogen has a small atomic mass, the zero point energy (ZPE) may be important. To investigate the vibrational effects on the phase stability, ZPEs for H2Cl and H5Cl were estimated at 100–300 GPa from the corresponding phonon spectra using the quasi-harmonic approximation.29 It is found that the ZPEs are quite small and the inclusion of ZPEs in the phase diagram only resulted in a slight shift in the formation pressures but the stability of both phases remains unaltered. Structures lying on the convex hull are thermodynamically stable or metastable and, in principle, can be synthesized. 2 complexes, H2Cl and H5Cl are stable at 100 GPa. At this pressure H2Cl has the most negative enthalpy of formation. With increasing pressure, the stability of H5Cl relative to H2Cl increases and becomes the most stable phase at 300 GPa.Before embarking on a detailed discussion of the structures of the predicted high pressure H26 The results show that the differences between calculations with and without vdW corrections on the formation enthalpies of the structures considered in 2Cl increased from 21.2 to 21.3 GPa, while for H5Cl it increased from 50 to 60 GPa. Otherwise, the energetic order remains the same.vdW effects may play an important role in the stabilization of a molecular solid. We have thus performed additional calculations on the H–Cl system with the vdW-DF2 method.nCl (n = 2–7). The starting point is the crystal structure of HCl under ambient pressure. At low temperature, X-ray and neutron diffraction show HCl crystallized in an orthorhombic structure (Bb21m).30 In the crystal, HCl molecules are linked via the H atoms forming zigzag chains running parallel to the crystallographic b axis. The nearest neighbour Cl–H and second nearest neighbour Cl···H distances are 1.25 Å and 2.44 Å, respectively. The Cl···Cl separation is 3.88 Å and the H–Cl···H valence angle is 93.6°. The predicted crystalline phase of H2Cl at 100 GPa has a C2/c space group and the structure is shown in 2 molecules. In this case, the H in the HCl chain is midway between the two Cl atoms with an H–Cl distance of 1.45 Å. The H–Cl–H angle has opened to 97.9° and the Cl···Cl separation is shortened to 2.90 Å. The H2 units in the structure all have a H–H distance of 0.74 Å, which is almost identical to that of the isolated molecule. The Bader charges for the H in the chain and Cl atoms are +0.44 and –0.35, respectively and 0.0 for the H atoms in the H2 units. The closest contact between a Cl atom and the H2 molecule is 1.98 Å. The crystal structure of H2Cl at 300 GPa differs little from that at 100 GPa. The H atoms in the H–Cl chains are still situated at the middle of the two neighbouring Cl atoms with a H–Cl distance of 1.35 Å. The H–Cl–H angle is 95.5° and the shortest separation between two Cl atoms has reduced further to 2.69 Å. The H–H bond length in the H2 unit is 0.73 Å. The closest H2···Cl distance is. 1.70 Å. Compression has a significant effect on the interatomic distances of the H–Cl chains but does not alter fundamentally the underlying bonding pattern. A longer H–Cl distance in the chain suggests increased ionicity of the Cl–H bonds.Now we examine the development of the high pressure crystal structures in H3Cl and H4Cl are only metastable, it is instructive to examine the evolution of the crystal structure with increasing H2 concentration. The structures of H3Cl at 100 and 300 GPa are shown in 2Cl, the H atom is equidistant from the two nearest Cl atom with H–Cl bond distances of 1.44 Å at 100 GPa and 1.43 Å at 300 GPa. The most significant difference between the low and high pressure structures is that the Cl–H–Cl angle is almost linear at 100 GPa but bends to 135° at 300 GPa. At 300 GPa, the closest contact between the H2 and the H in the chain is 1.27 Å. However, in both cases, the H–H distance of the H2 molecule remains 0.73 Å. The structure of H4Cl at 100 GPa differs dramatically from all the structures within this series of compounds. Instead of H–Cl chains, the structure is composed of isolated HCl and H2 molecules. The H–Cl distance is 1.38 Å and the H–H bond length is 0.74 Å. For comparison, the H–Cl bond of a free molecule is 1.276 Å. Therefore, the distance in the solid state at 100 GPa is slightly longer. The structure of H4Cl at 300 GPa again is different from that at 100 GPa. The basic building units are isolated Cl atoms, H2 molecules with H–H distance of 0.74 Å and a novel 2-D layer of slightly puckered fused hexagonal rings formed from 3 HCl units with additional H atoms attached to the Cl atoms. Each H atom in the ring is bonded to three Cl atoms. In addition, each Cl is bonded to an extra H atom which is not coordinated to other species in the crystal. The H–Cl distances in the fused ring are 1.59 Å and the terminal H–Cl is substantially shorter at 1.49 Å. Interestingly, the terminal H–Cl–H (ring) angles are 77° and the in-plane H–Cl–H and Cl–H–Cl angles are between 114–115°.Although H5Cl at 100 GPa. Although chains formed from Cl and H atoms are still clearly visible, the detailed construction of the chain is very different. In H5Cl, instead of placing one H atom midway between the two nearest neighbour Cl atoms, it is replaced by an H3 unit. The H3 is a distorted isosceles triangle and can be described as a loosely bound unit of an H atom and an elongated H2 with an apical angle of 63.8°. The apical H atom is linked to the two nearest Cl atoms in the chain with H–Cl distances of 1.47 Å. The distances from the apical H atom to the two H forming the H2 are 1.01 Å and 0.97 Å, respectively and the intermolecular H–H distance is 0.81 Å. The remaining H2 units in the structure have H–H distances of 0.74 Å. Moreover, the shortest distance from these H2 to the H3 is 1.36 Å and, therefore, may be considered as non-interacting molecules.An interesting structure was observed in H5Cl to 300 GPa does not change the space group symmetry. The chain pattern with interpose H3 units is still maintained, but the local H···H interactions have changed dramatically. The H3 unit now approaches an equilateral triangle. The H–H lengths are 0.87, 0.87 and 0.88 Å with bond angles 59.7, 59.7 and 60.5°. The Cl–H distance has elongated from 1.47 Å at 100 GPa to 1.60 Å! The large lengthening of the Cl–H clearly suggests a substantial change in the Cl–H bonds. More significantly, the isolated H2 molecules are now pushed towards the H3 units and interact with one of the H atoms forming almost two H···H bonds at 1.15 Å. Concomitantly, the distance in H2 is lengthened to 0.76 Å. The Bader charges for the H atom in the H3 and H2 units and for the Cl atom are +0.16 and +0.014 and –0.48 respectively. In comparison to H2Cl the ionicity on both the H and Cl atoms have increased substantially. The plot of the electron localization function (ELF) shown in 3 ring and in the H2 molecule (ELF over 0.8). Weak pairing is also observed between one of the H in the H3 ring with the two H atoms of H2.Compression of H3 unit to an isolated trihydrogen cation H3+.It is tempting to relate the positively charged H3+ has a perfect triangular structure with H–H bond distance of 0.90 Å. For comparison, at 300 GPa the average intramolecular H–H distance in the H3 unit in H5Cl is 0.87 Å and a total charge of +0.48 (3 × 0.16). In addition, the H–H distance of 0.76 Å in the H2 unit is only slightly perturbed from the isolated molecule. Therefore, it is not unreasonable to suggest that the high pressure H5Cl structure is composed of H3+ stabilized in the solid state through primarily ionic interactions with the Cl atoms and secondary weak interactions with a pair of H2 molecules. To investigate further the properties of the H3+ unit in H5Cl, the phonon densities of states calculated at 300 GPa are shown in –1 are dominated by the Cl atomic motions. The phonon modes in the region from 1800 to 3400 cm–1 can be assigned to H3+ molecular vibrations and then near 4000 cm–1 to the H2 molecule. The H3+ bend vibrations are split into two bands centered at 1800 cm–1 and 2000 cm–1. The peak at 3350 cm–1 is attributed to the stretching vibration. In comparison, the fundamental frequencies for isolated H3+ are ν (stretch) = 3220 × 48 cm–1 and ν (bend) = 2545 × 99 cm–1, the latter is a degenerated mode.31 The vibration frequency is 4350 cm–1 for a free H2 molecule. The main mechanism for the synthesis of H3+ by experiment is via the chemical reaction:32 H+ + H2 → H3+. The concentration of H2+ is the limiting factor on the rate of this reaction. H2+ can only be produced in interstellar space by the ionization of H2 by a cosmic ray. In H5Cl, the electronegative chlorine atoms acquire electrons from the hydrogen. From previous studies, it is known that high pressure facilitates the transfer of electrons from the more electropositive element into the interstitials of a crystal, forming electrodes. In this case, these electrons, being transferred to the chlorine, are originating from the hydrogen. The formation of multicenter H3···H2 clusters helps to delocalize the positive charge and stabilizes the system. This explains the unexpected formation of H3+-like units in the crystal structure.HThe nature of the bonding in the group 17 hydrides at high pressure is different from group 1 and 2 and group 14 hydrides. One observed a gradual shift in the chemical interaction from electron transfer in electropositive group 1 and 3 compounds to covalent bonding group 14 and finally the ionic bonding in group 17 elements. Although this study was focused on Cl, we anticipate a similar bonding mechanism is applicable to other halogen hydrides.2 system at high pressure using the PSO algorithm in combination with ab initio density functional based electronic calculations. Between 100–300 GPa, two stable phases with the H5Cl and H2Cl stoichiometries were found. The basic structure of the high pressure phases is similar to the low temperature ambient pressure structure of HCl. H2Cl consists of zigzag H–Cl chains and non-interacting H2 molecules. The most usual and informative finding is that while the chain like structure is preserved in H5Cl, the H atoms connecting the Cl in the chains are replaced by units consisting of weakly interacting H2···H3. The H3+ is positively charged and stabilized from the formation of multicenter bonds. The similarity in the local structure, vibrational frequencies and electronic charge compel us to relate the unit to the isolated H3+ molecule. It is also found that the effect of pressure on the electronic structure of group 17 hydrides is very different from the more electropositive group 1 and 2 elements: the electron-rich Cl atom and anion do not transfer their electrons into interstitial space of the crystal under very high compression. In fact electrons are removed from the H atoms leading to the formation of cationic clusters that benefit from multicenter bonding.We have investigated the phase stability of the HCl–H


The data indicate that the key to nanobrush formation is controlling the atomic surface transport processes to balance defect aggregation and surface‐smoothing processes. High‐resolution scanning transmission electron microscopy data reveal that defect‐mediated aggregation is the key to TiO2 nanobrush formation. The large concentration of defects present at the intersection of domain boundaries promotes aggregation of PLD growth species, resulting in the growth of the single‐crystalline nanobrush architecture. This study proposes a model for the relationship between defect creation and growth mode in nonequilibrium environments, which enables application of this growth method to novel nanostructure design in a broad range of materials.This study demonstrates that precise control of nonequilibrium growth conditions during pulsed laser deposition (PLD) can be exploited to produce single‐crystalline anatase TiO The ideal 2D growth mode can be readily interrupted by growth instabilities, which cause growth front roughness. When the growth front roughness is extremely enhanced due to limited surface diffusion, a columnar structure will be developed. Kinetic control between growth instabilities and surface equilibration is therefore critical to break down the 2D epitaxy, increasing the possibility of constructing 1D nanostructures. Recent reports of growing highly crystalline 1D nanostructures that exhibit good ionic conductivity in yttria‐stabilized zirconia (YSZ), enhanced low‐field magnetoresistance in La2 because 1D TiO2 tubes,2, anatase is the most widely explored phase for practical applications.2, obtained by an analysis of surface energies, is dominated (>94%) by low energy {101} facets.2 nanostructures with a large percentage of reactive {001} facets is highly desired. Solvothermal synthesis using hydrofluoric acid (HF) as a morphology controlling agent is currently the most effective and popular method for stabilizing the {001} facets of anatase TiO2.2 nanostructures with a high concentration of the {001} facets using physical vapor deposition method.In this report, we focus on 1D nanostructured TiO2 1D “nanobrush” architectures fabricated on (001) SrTiO3 (STO) substrates within a small window of kinetic growth parameters. The 1D nanobrushes are epitaxially stabilized on a continuous 2D anatase underlayer and feature a large fraction of exposed {001} facets. Based on a systematic study of the microstructural evolution of a TiO2 nanobrush film, we describe the growth morphology in terms of defect‐mediated aggregation, stressing the important role of the misfit defects in the heteroepitaxial growth of 1D nanobrush architectures that are generated by operating PLD in the defect producing modality.In this study, we report single‐crystalline anatase TiO2Figure2, revealing a spatially separated roof‐like surface morphology  consisting of triangular lamellae that are stacked along the <110> directions. Film densification occurs at a higher substrate temperature of 600 °C or a lower p(O2) of 50 mTorr. The breakdown of the compact 2D epitaxy that can be seen from the sample grown at 500 °C in 100 mTorr is attributed to the reduced adatom mobility and resultant aggregation that prevails when surface diffusion is suppressed. However, further reduction of the temperature or an increase in oxygen pressure resulted in an undesired cauliflower‐like surface morphology.The plan view scanning electron microscopy (SEM) images of films in Figure2 film despite its high porosity observed in the SEM image shown in Figure 2 101 and STO 101 reflections at ψ = 68.32° and 45°, respectively, confirm the epitaxy of a nanostructured TiO2 film on STO with a clear fourfold symmetry as shown in Figure The morphology of the porous film with a repeated roof‐like surface structure suggested that the film was obviously not continuous. Porosity often originates from the breakdown of 2D epitaxy and results in the formation of polycrystalline or even amorphous phases. Surprisingly, neither impurity phases nor different orientations were detected in X‐ray diffraction (XRD) scans  image of this film in Figure FigureHigh‐resolution STEM using high‐angle annular dark field (HAADF) imaging of a single nanobristle and its branches along the STO [010] projection direction was performed to reveal the surface morphology of the bristles as shown in To better elucidate the structure of each bristle, a series of atomic‐resolution STEM images were taken from the marked regions in Figure 2, the angle between the (001) and (101) planes is 68.3°. If we consider a branch terminated by equal contributions of (001) and (101) surface as observed and take into account the 18° tilt of the (001) plane with respect to the STO(001) surface plane  STO substrates by PLD. The STO substrates were chemically etched by buffered HF for 30 s and then thermally treated at 1100 °C in air for 1.5 h. TiO2 nanobrush films with a thickness of 1.5 µm were deposited with a KrF excimer laser using a TiO2 polycrystalline target. In order to check the viability of growing single crystalline nanoarchitectures by PLD, various growth conditions were explored. In order to control the surface diffusion of adatoms, the substrate temperature (T = 400–600 °C) and oxygen background pressure (p(O2) = 50–200 mTorr) were systematically varied. We grew TiO2 films with a fixed number of laser pulses (50 000 pulses), and the typical deposition rate was 0.3 Å per laser pulse. For structural characterization, XRD was carried out using a high‐resolution X‐ray diffractometer with Cu Kα1 radiation. The plan view morphology of samples was characterized by SEM. High‐angle annular dark‐field (HAADF) imaging was carried out in Na ion UltraSTEM200 operated at 200 kV. The microscope is equipped with a cold field‐emission gun and an aberration corrector for sub‐angstrom resolution. Inner/outer angles of 70/240 mrad were used for HAADF imaging. The convergence semi‐angle for electron probe was set to 30 mrad.Anatase‐phase TiOSupplementaryClick here for additional data file.


Beverages play an important role in the diets of adolescents because they help to maintain hydration and can provide important nutrients, such as calcium, vitamin D, and vitamin C ” and “During the past 7 days, how many times did you drink 100% fruit juices such as orange juice, apple juice, or grape juice? ” Response options were “I did not drink soda or pop during the past 7 days” or “I did not drink 100% fruit juice during the past 7 days,” “1 to 3 times during the past 7 days,” “4 to 6 times during the past 7 days,” “1 time per day,” “2 times per day,” “3 times per day,” or “4 or more times per day.” Students who selected “1 time per day,” “2 times per day,” “3 times per day,” or “4 or more times per day” were categorized as daily soda or juice drinkers; all other students were categorized as non-daily soda or juice drinkers. Daily milk consumption was assessed with the question “During the past 7 days, how many glasses of milk did you drink? ” Response options were “I did not drink milk during the past 7 days,” “1 to 3 glasses during the past 7 days,” “4 to 6 glasses during the past 7 days,” “1 glass per day,” “2 glasses per day,” “3 glasses per day,” or “4 or more glasses per day.” Students who selected “1 glass per day,” “2 glasses per day,” “3 glasses per day,” or “4 or more glasses per day” were categorized as daily milk drinkers; all other students were categorized as non-daily milk drinkers.Data from each survey were weighted to provide national estimates. Statistical software was used to account for the complex survey design of the YRBS. Prevalence estimates were computed overall and by school grade , sex , and race/ethnicity . Other and multiple racial/ethnic subgroups were excluded from the race/ethnicity subgroup analysis because the numbers were too small for meaningful analysis. Research indicates that income plays a role in the dietary choices of adults; however, little research has been done on the impact of socioeconomic factors on adolescents’ beverage choices  of studeDuring 2007–2015, daily soda consumption decreased from 33.8% to 20.4%. During 2007–2011, daily milk and juice consumption did not change; however, during 2011–2015 daily milk and juice consumption decreased from 44.4% to 37.5% and from 28.2% to 21.6%, respectively .Among students in grade 9, daily soda consumption decreased significantly during 2007–2011 and then further decreased significantly during 2011–2015 . Among sBeverages contribute approximately 20% of calories to the diets of children and adolescents and can contain important nutrients (,Findings from this report and other recent studies indicate that adolescents are consuming less soda (Recent analysis of national data also indicates a decrease in juice consumption (***The findings in this report are subject to at least three limitations. First, data are self-reported and might be subject to reporting and social desirability bias. A recent study showed that YRBS beverage questions underestimated the prevalence of daily non-diet soda intake but overestimated prevalence of daily milk and 100% juice intake compared with a 24-hour dietary recall interview (Multiple measures are needed to address adolescents’ beverage consumption and should reach settings where adolescents spend their time, such as homes, schools, and the community at large. Parents can influence the home nutrition environment through their food purchases (Community-based strategies also should be considered. For example, health care providers can screen and counsel patients and their families on decreasing SSB intake, and organizations can implement social marketing campaigns to promote consumption of healthier beverages. Although the results of this report indicate a decline in soda consumption, there is a continued need to help adolescents shift beverage consumption patterns to more healthful options.Beverages contribute approximately 20% of calories to the diets of children and adolescents and can contain important nutrients, but beverages can also contribute to excess consumption of added sugars and calories. Previous research has indicated that daily consumption of milk, juice, and non-diet soda has been decreasing over time, but little is known about trends among subgroups of youth.During 2007–2015, daily soda consumption among U.S. high school students decreased significantly from 33.8% to 20.4%. During 2007–2011, daily milk and juice consumption did not significantly change, and then during 2011–2015 daily milk and juice consumption decreased significantly from 44.4% to 37.5% and from 28.2% to 21.6%, respectively.Although the significant downward trends in daily soda consumption suggest that interventions encouraging reduced consumption of soda are working, overall prevalence of daily soda consumption remains high. Policy and educational approaches should continue to promote healthier beverage options in place of sugar-sweetened beverages.


Improved light harvesting efficiency and charge transporting performance in IOT‐ETL based PSCs yield high power conversion efficiency of 13.11%. Moreover, the TiO2 nanoparticles, TiO2 mesoporous structure, 1D TiO2 and ZnO have similar electron‐injecting mechanism and have been widely applied as mesoporous scaffold layers.2 or ZnO is utilized as both compact layer and mesoporous scaffold layer, these two layers could be named as electron transport layer (ETL). Moreover, enhanced light harvesting efficiency (LHE) is necessary for a good‐performance solar cell, while the LHE of conventional mesoporous structure TiO2 ETLs based perovskite solar cells has large potential to be improved.The structures of perovskite solar cells can be widely divided into two classes: mesoporous structure based perovskite solar cell and planar heterojunction based perovskite solar cell.2 electron transport layer (IOT‐ETL) by a simple polystyrene‐assistant method for the first time, which combines the functions of the compact layer and mesoporous scaffold layer in perovskite solar cells. It demonstrates that the porous structure in IOT‐ETL could enhance the devices' LHE obviously. Moreover, the bottom of IOT‐ETL plays the same role as the traditional compact layer, transporting electrons and inhibiting the recombination of electrons and holes. As a result, the IOT‐ETL based perovskite solar cell yields a best PCE of 13.11%, higher than that of the conventional P25 mesoporous layer based perovskite solar cells (11.00%). To the best of our knowledge, this is the first study utilizing multifunctional inverse opal‐like TiO2 as ETL to obtain highly efficient perovskite solar cells with PCE >10%.Here, we prepare a multifunctional inverse opal‐like TiOFigure4)2 aqueous solution and Ti(SO4)2 in PS emulsion) were spin‐coated on the etched FTO substrate respectively. After sintered at 773 K for 30 min, the template was removed and the porous structure was obtained. It is worth noting that a bottom of the IOT‐ETL film was fabricated by spin‐coating the Ti(SO4)2 aqueous solution. This bottom is significant for the improvement of the devices' overall PCE and will be discussed in the following text. CH3NH3PbI3 was prepared by a typical two‐step method, in which PbI2 was deposited on the IOT‐ETL films first followed by spin‐coating the CH3NH3I solution. After the spiro‐MeOTAD was deposited on the CH3NH3PbI3 film, Ag was thermally evaporated to form the back contact.The fabrication procedure of IOT‐ETL based perovskite solar cells is shown in Figure2, all diffraction peaks match well with those of the anatase TiO2 (JCPDS No. 21–1272) . The inset image in Figure 3NH3PbI3 cuboid above the IOT‐ETL film is between 100–150 nm, which is similar to the diameter of the circular pore. The larger pore size than that of conventional TiO2 mesoporous layer is beneficial for the filling of perovskite absorber.2 wall with a wall thickness of 6 nm. Moreover, conventional P25 mesoporous film with around 150 nm thickness was prepared by spin‐coating self‐made P25 paste. The light transmittance spectra of both IOT‐ETL, P25 mesoporous film and bare FTO were measured. As shown in Figure 2 diffraction peaks can be detected by XRD pattern , which indicates pores structure in IOT‐ETL film is suitable for the infiltration of CH3NH3I. We also investigated the morphology of the bottom in IOT‐ETL film obtained from the first spin‐coating of Ti(SO4)2 aqueous solution. Figure S3, Supporting Information, shows that the bottom is a dense and uniform film, which fully covers the whole FTO substrate, similar to the TiO2 compact film in previous reports.As shown in J)–voltage (V) characteristics of the pervoskite solar cells (4000 rpm 30 s for IOT‐ETL film preparation) with reverse scan under a simulated AM 1.5 G solar irradiation and in the dark are shown in FigureJsc), the open‐circuit potential (Voc) and the fill factor (FF) are 22 mA cm−2, 940 mV and 0.62, respectively, yielding a PCE of 12.78%. It is well known that the thickness of ETL films is closely related to the PCE of solar cells.−1 demonstrates a PCE value of 13.11% with 21.93 mA cm−2, 972.8 mV, and 0.61 for Jsc, Voc, and FF, respectively , where R and A represent reflectance and absorbance, respectively.Jsc and improve the PCE of pervoskite solar cells. J–V curves of two kinds of ETL films based devices are shown in Figure Jsc from 19.9 mA cm−2 to 21.93 mA cm−2 leads to high‐performance of IOT‐ETL film based pervoskite solar cells, which confirms our inference. The external quantum efficiency (EQE) spectra of the best performing IOT‐ETL film based perovskite solar cell is close to full‐conversion as shown in Figure S6, Supporting Information. The integrated Jsc calculated from the EQE data is 20.59 mA cm−2, which agrees well with the measured Jsc.Typical UV–vis absorption spectrum of CH4)2 aqueous solution could prevent the direct contact between pervoskite and FTO glass and was critical to the improvement of the devices' PCE. FigureJ–V curves for perovskite solar cells based on different IOT‐ETL films. The bottomless IOT‐ETL film based solar cell exhibits a Jsc of 20.70 mA cm−2, a Voc of 840 mV and a FF of 0.48, yielding a PCE of 8.28%. Obvious decrease of Voc and FF is responsible for the low PCE, which attributes to the increased charge recombination without the bottom.3NH3PbI3 and the bottom of IOT‐ETL and the electron–hole pair recombination is restricted.J–V characteristics under dark conditions for perovskite solar cells based on different IOT‐ETL films . Might due to the charging‐discharging processed in perovskite or IOT‐ETL film, the lowest currents do not happen at zero bias.Voc  Nanospheres Emulsion: Monodisperse PS nanospheres emulsion with the nanosphere size of 100 nm was prepared through a modified dispersion polymerization approach.3NH3I)Synthesis of Methylammonium Iodide  conducting glass substrate  was patterned by etching with Zn powder and 2 m HCl diluted in DI water. The etched FTO substrate was then cleaned in an ultrasonic bath containing acetone for 20 min, ethanol for 20 min, rinsed with DI water and dried with clean dry air. The IOT‐ETL with 100 nm pore size was prepared by a simple two‐step spin coating titanium precursor solution (solution A and solution B) followed by sintering. Solution A was prepared by maxing 85 mg Ti(SO4)2  with 1 mL DI water and solution B was made by mixing 85 mg Ti(SO4)2 and 1 mL the obtained PS emulsion followed by stiring and sonication. Solution A was spin‐coated on the FTO substrate at 4000 rpm for 30 s, which was heated at 398 K for 5 min. After cooling to room temperature, the solution B was spin‐coated on the obtained substrate at various spin speed for 30 s. Finally, the formed films were annealed at 773 K for 30 min with a heat ramp of 2 K min−1 to form inverse opal‐like TiO2 film. The IOT‐ETL film was immersed in 0.02 m aqueous TiCl4  solution at 333 K for 40 min. After rinsing with DI water and drying, the film was heated at 773 K for 30 min. The bottomless IOT‐ETL film was prepared by the same method without the spin coating of solution A. Compact TiO2 blocking layer in conventional P25 mesoporous ETL was spin‐coated on the FTO substrate at 2000 rpm for 20 s using 0.15 m titanium diisopropoxide bis(acetylacetonate)  in 1‐butanol  solution, followed by heated at 398 K for 5 min. After cooling to room temperature, the diluted P25 paste was spin‐coated on the compact TiO2 blocking layer at 5000 rpm for 30 s, where the pristine P25 paste was prepared by previous report−1). The same thermal and TiCl4 treatment for IOT‐ETL film was further applied to P25 mesoporous ETL.Fabrication of Hybrid Perovskite Solar Cells: The CH3NH3PbI3 was deposited on the ETL film using a two‐step method.m PbI2 solution was prepared by dissolving 462 mg PbI2  in 1 mL of N,N‐dimethylformamide , stirring at 343 K for 12 h. The IOT‐ETL film was infiltrated with PbI2 by spin‐coating the 1 m PbI2 solution at 3000 rpm for 5 s and 6000 rpm for 5 s followed byheated at 373 K for 5 min. After cooling to room temperature, 100 μL of 6 mg mL−1 CH3NH3I solution in 2‐propanol  was loaded on the PbI2‐coated film for 40 s, which was spun at 4000 rpm for 20 s and annealed at 373 K for 5 min. The HTM was then deposited by spin coating 2,2′,7,7′‐tetrakis‐9,9‐spirobifluorene (spiro‐MeOTAD) solution at 4000 rpm for 30 s. The spiro‐MeOTAD solution was prepared by dissolving 72.3 mg of spiro‐MeOTAD in 1 mL of chlorobenzene, to which 28.8 μL of 4‐tert‐butyl pyridine and 17.5 μL of lithium bis(trifluoromethanesulfonyl)imide (Li‐TFSI) solution  were added. Finally, 110 nm Ag was thermally evaporated on the spiro‐MeOTAD‐coated film to form the back contact. The obtained solar cells were left in a desiccator for 72 h before tested. All the device fabrication processes were carried out under controlled atmospheric conditions and a humidity of <1%.Characterizations: The morphology and structure of the IOT‐ETL films were characterized by field emission scanning electron microscopy  and high‐resolution transmission electron microscopy . The XRD spectra of the prepared IOT‐ETL films were measured using powder X‐ray diffraction . The absorption spectra of the perovskite coated IOT‐ETL films were measured by using a Cary 500 UV–vis‐NIR Spectrophotometer. The photoluminescence measurement was acquired at room temperature using a UV–vis‐NIR fluorescene spectrophotometer (Fluorolog‐3‐P) with an excitation wavelength of 600 nm. The solar cells were illuminated using a solar light simulator  and the power of the simulated light was calibrated to 100 mW cm−2 using a Newport Oriel PV reference cell system (model 91150 V). The J–V curves of solar cells were measured using a Keithley 2400 digital sourcemeter. Devices were masked with a metal aperture to define the active area of 0.0625 cm2. The EQE was measured using a Newport‐74125 system (Newport Instruments).As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.