Unlike other 2D materials, MoTe2 allows easy phase modulation and efficient carrier type control in electrical transport. However, its unstable nature and low‐carrier mobility limit practical implementation in devices. Here, a deterministic method is proposed to improve the performance of MoTe2 devices by inducing local tensile strain through substrate engineering and encapsulation processes. The approach involves creating hole arrays in the substrate and using atomic layer deposition grown Al2O3 as an additional back‐gate dielectric layer on SiO2. The MoTe2 channel is passivated with a thick layer of Al2O3 post‐fabrication. This structure significantly improves hole and electron mobilities in MoTe2 field‐effect transistors (FETs), approaching theoretical limits. Hole mobility up to 130 cm−2 V−1 s−1 and electron mobility up to 160 cm−2 V−1 s−1 are achieved. Introducing local tensile strain through the hole array enhances electron mobility by up to 6 times compared to the unstrained devices. Remarkably, the devices exhibit metal–insulator transition in MoTe2 FETs, with a well‐defined critical point. This study presents a novel technique to enhance carrier mobility in MoTe2 FETs, offering promising prospects for improving 2D material performance in electronic applications.Molybdenum ditelluride (MoTe 2 device performance by inducing tensile strain through substrate engineering and encapsulation. This structure significantly improves hole and electron mobilities in MoTe2 field‐effect transistors, reaching up to 130 and 160 cm2 V s, respectively. Remarkably, the electron mobility increases by 6 times compared to the unstrained devices. The technique holds potential for 2D‐electronics.A deterministic method is presented to enhance MoTe A few‐layer 2H‐MoTe2 has an indirect bandgap of around 0.9 eV which makes it promising for applications in state‐of‐the‐art electronic and optoelectronic devices operating in the visible to near‐infrared range. The weak Fermi‐level pinning in MoTe2 and metal contacts allows for easy modulation of device polarity through contact engineering. Despite its potential, low carrier mobility and high contact resistance in MoTe2 field effect transistors (FETs), along with the material instability in atmospheric conditions, remain major obstacles to its integration into electronic devices. Theoretically, the room‐temperature hole and electron mobilities in MoTe2 devices can reach up to 200 cm−2 V−1 s−1, but in practice, the actual carrier mobilities of MoTe2 fall significantly below this theoretical range.MoTeications.2, 3 High dielectric constant (high‐κ) materials can also enhance the carrier mobility in 2D materials by strongly damping scattering from Coulombic impurities. Atomic layer deposition (ALD) grown Al2O3 (κ = 9.1) is one of the widely used dielectric materials for 2D transistors due to its role to reduce defects and suppress charge trapping effects at the interface. Additionally, ALD‐grown Al2O3 serves as a passivation layer for 2D transistors to improve thermal and chemical stability. Moreover, the metal‐oxide dopes the 2D materials with n‐type carriers and elevates the electron current level. While possessing all these advantages, another potential aspect of this ALD layer is its ability to introduce tensile strain to the 2D material when the substrate is engineered in such a way that the weight of the passivated ALD layer and the thermal expansion difference between Al2O3 and MoTe2 exert tensile strain on the material, thereby enhancing device performance.Several techniques have been developed to improve the performance of 2D‐material‐based devices, including chemical doping, inducing strain, and substrate engineering.15 Hig2 FETs. Our approach involves creating a hole array in the substrate, using ALD‐grown Al2O3 as an additional dielectric layer on SiO2, and employing Al2O3 as a passivation layer for the MoTe2 FET. Our proposed structure induces tensile strain on MoTe2, leading to significant enhancement of the carrier mobility and on‐off ratio. The improved performance of the MoTe2 devices can be attributed to the significant damping of the Coulomb and electron–phonon scattering as a result of the change in the band structure of MoTe2. Additionally, the stable and good device quality leads to the metal–insulator transition (MIT) in the strained MoTe2 devices during temperature‐dependent electrical measurements, characterized by a sharp MIT critical point. This observation highlights the potential of MoTe2 for the development of fast‐switching devices and memory elements, with applications in efficient sensors, photodetectors, and energy storage devices.In this work, we first demonstrate a simple technique to improve the performance of MoTe2Figure 2, where a 2D hexagonal lattice structure is formed by sandwiching one layer of Mo atoms between two layers of S atoms. Relatively, 2H‐MoTe2 is thermodynamically more stable compared to other phases. Figure 2 substrate using a few‐layer MoTe2 as the channel material. The room‐temperature Raman spectrum of pristine MoTe2 obtained using a 532 nm (≈2.33 eV) laser excitation is depicted in Figure 2 appear at 173, 233, and 288 cm−1 corresponding to out‐of‐plane A1g, in‐plane E2g1, and bulk in‐active B2g1 modes, respectively. The transfer characteristics curve of a ≈2 nm thick MoTe2 FET at room temperature is shown in Figure 2 V−1 s−1. It is worth noting that, we observe a very thin layer (≈2 nm) of MoTe2 behaving as p‐type or p‐dominant ambipolar with Ti/Au electrodes. However, as the thickness of MoTe2 increases, the n‐branch of the transfer curve becomes stronger, leading to n‐type MoTe2 FET.2O3 grown by ALD induces n‐doping in 2D materials when the metal oxide is used as a dielectric material and also as a passivation layer for 2D transistors. The effect of ALD doping is particularly noticeable in the electrical response of 2D FETs when compared with bare 2D FETs on Si/SiO2 substrate. Al2O3 dopes the material by injecting electrons in the channel, leading to a transition in carrier type from p‐type to p‐dominant, n‐dominant ambipolar, or even n‐type behavior. The µe in the devices also increases with an elevated electron current level. Meanwhile, Al2O3 contributes to a reduction in device hysteresis by minimizing carrier trapping and promoting a high‐quality interface. The high‐quality interface facilitated by Al2O3 may also enhance µh in the devices. Detailed discussion on the mechanism of doping by ALD‐grown Al2O3 is provided in the last section of this study.Usually, Al2 induced by hole‐array pattern on the Si/SiO2 substrate, while considering the ALD influences on the material. To achieve this, we fabricate hole arrays with various diameters on Si/SiO2, followed by deposition of 5 nm Al2O3 on the substrate. Subsequently, we transfer MoTe2 flakes onto the hole‐array substrate using the hot pick‐up technique. The adhesion between MoTe2 and ALD‐grown Al2O3 substrate relies on van der Waals bonding. The high‐quality and uniformity of layer‐by‐layer growth of Al2O3 in the ALD process enhances the adhesion. MoTe2 FETs with Ti/Au contacts are fabricated on the substrate, and an additional 50 nm layer of Al2O3 is deposited on the devices using a seeding layer to ensure uniform coverage on the MoTe2 because of the absence of dangling bonds on the surface of 2D materials . Figure 2O3. Two different thicknesses of MoTe2 flakes are selected to analyze the strain effects: a ≈2.7 nm flake thickness exhibits p‐dominant ambipolar behavior, as shown in Figure 5. The AFM measurement results of these samples are shown in Figure In this experiment, we investigate the role of local strain on MoTeechnique.22 TheCAl2O3 and CSiO2 are the capacitances of Al2O3 and SiO2, respectively, Vd represents the drain voltage, and dIddVg is the extrinsic transconductance of the devices. The thicknesses of the SiO2 and Al2O3 layers are 285 and 5 nm, respectively. The calculated µh of the device in Figure 2 V−1 s−1, while µe of the device in Figure 2 V−1 s−1. Note that the mentioned mobility values are underestimated due to the 2‐probe measurements, which include the effect of contact resistance. Accurate mobility numbers can be extracted from 4‐probe measurements.The 2‐probe field effect mobilities of these devices are calculated using the following equation2O3 deposition . The as‐transferred flakes of MoTe2 having a thickness of ≈5 nm, are suspended on the hole arrays experience relatively small strain. To enhance the localized tensile strain on the MoTe2 flakes, we grow a 50 nm layer of Al2O3 using the ALD process. As we know that Raman spectroscopy is a well‐established tool to investigate the stain in 2D materials, we use a 532 nm laser to characterize the nature and magnitude of strain in both strained and unstrained MoTe2. The distinct Raman peaks observed in both samples confirm the intact and high‐quality of the flakes. We notice that the Raman modes of strained MoTe2 are redshifted compared to the unstrained regions of the flake, as shown in Figure  Especially, in our samples, most strained samples exhibit a redshift of ≈2.98 and ≈2.22 cm−1 for prominent E2g1 and A1g modes, respectively. Notably, in the previous reports, the shift in the A1g mode under tensile strain was almost negligible. We notice a significant shift in the A1g mode of strained MoTe2. We assume that the shift is the result of using a few‐layer and unsupported MoTe2 or the effect of change in the dielectric environment in the hole‐array regions of our samples. Another major reason behind this larger shift in A1g could be the result of the high level of doping in MoTe2 induced by Al2O3. Figure 2 samples compared to the unstrained samples. It is evident that the shift of all Raman peaks is increasing with the increase in the hole diameter, with the highest shift observed when the diameter is 300 nm. This indicates that the 300 nm hole diameter exerts the strongest tensile strain on the MoTe2 flake around the peripheral region of the hole. The amount of strain approximated from the shift of E2g1 mode is ≈0.67%. The strain in the material decreases at 400 nm, which could be attributed to the increased flat or less strained region at the center of the hole.In order to get a deeper understanding, we create hole arrays in the substrate with different diameters, such as 100, 200, 300, and 400 nm. The target depth of the holes is 70 nm; however, the actual depth ranges from 40 to 55 nm after 5 nm Aln Figure 25 Esp2O3‐induced strain in the MoTe2 device performance. Local tensile strain resulting from our proposed structure enhances the electron current and mobility of MoTe2 FETs by up to ≈6 times compared to the unstrained flat sample. Figure e of the strained and unstrained samples. The optical image and height profile of the MoTe2 flake are presented in Figure 2. A schematic of a simplified band structure of MoTe2 before and after strain is presented in the inset of Figure 2. The bandgap EGT of MoTe2 decreases as the valley of the conduction band at the K point is lowered when tensile strain is applied to the few‐layer MoTe2 flake. As a result, the energy difference between K point and Q point (denoted as ΔEQK) increases, resulting in very low electron intervalley scattering by reducing the electron effective mass. This could be the major reason for the enhancement of µe in MoTe2 devices. In addition, the contact resistance of the devices decreases by reducing the Schottky barrier height in the devices when tensile strain is applied to 2D materials, also leading to the high µe. Among the 2D TMDCs, MoTe2 is more susceptible to defect formation and charge trapping. The application of tensile strain to MoTe2 could potentially enhance the carrier lifetime by mitigating the trapping of the charge carriers. This could also lead to high current and mobility in the devices.Next, we analyze the influence of the hole‐array pattern and top AlFigure 2.29 A s2O3 layer enhances the MoTe2 device performance significantly.Figure  and enhancing gate capacitance. An additional layer of Al2O3 on SiO2 has a similar influence on the device performance. Al2O3 provides strong gate controllability and reduces the Schottky barrier height (SBH) between metal and semiconductor, resulting in increased carrier mobility through the thermionic emission process.High‐κ dielectric materials improve the device performance by strongly reducing Coulomb scattering17 andFigure 2 transistor made with Ti/Au contacts at various temperatures. Remarkably, we observe MIT in MoTe2 devices. The drain current, measured as a function of gate voltage, shows a crossover at a specific gate voltage. This crossover point, also known as critical or MIT point is at ≈Vg = 62 V. Below the point, the MoTe2 device behaves as an insulator, as device conductivity increases with temperature increase. Meanwhile, above the MIT point, the conductivity decreases with increasing temperature, behaving as a metal. At the MIT point, the carrier density is estimated to be nD  ≈  3.74 × 10−12 cm−2, calculated using nD=Cox(Vg−Vth)q, where Cox is the overall capacitance of the dielectrics, Vth is the threshold voltage, and q is the charge of an electron. The dependency of the conductivity on the temperature is shown in the inset of Figure Ec. In the insulating regime, the Fermi energy EF moves away from the mobility edge which results in strong localization of the charge carriers. Therefore, the conductivity increases with the increase in the temperature due to thermal activation. Conversely, EF moves toward Ec in the metallic regimes, and charge carriers are weakly localized. As the temperature increases, phonon scattering becomes more dominant, leading to a decrease in the conductivity in this regime.For further analysis of the Schottky barrier, we investigate the transfer characteristics of the devices at different temperatures. 2 V−1 s−1 at 80 K, indicating strong screening of Coulomb scattering due to the low temperature and thin Al2O3 dielectric layer. The mobility of the transistor decreases with the temperature increase as the phonon scattering becomes dominant over Coulomb scattering. The hysteresis in the device increases with the temperature increase as well; however, the value of hysteresis is very small, even at room temperature compared to the FET made solely on flat Si/SiO2 substrate, as shown in Figure 2O3 than SiO2 in terms of low surface defects resulting from the layer‐by‐layer deposition of Al2O3 in the ALD process. This results in less carrier trapping during the conduction.The variation of mobility of the device with temperature is shown in Figure 2 devices, we can also analyze the quality of the contacts (Ti/Au) on the MoTe2 channel while considering the influence of top‐bottom Al2O3. The Arrhenius plot in Figure Id∝T32.exp[−qϕB/kBT], where ϕB represents the Schottky barrier height (SBH) and kB is the Boltzmann constant. The extracted SBH as a function of Vg is presented in Figure Vg of the SBH indicate thermionic emission of the carriers contributing to the current and leading to high mobility in the devices. This could be the result of the combined effect of the top Al2O3 used for inducing strain and the bottom Al2O3 used as an additional dielectric layer for MoTe2 FETs. Understanding the mechanism of doping in MoTe2 by ALD‐grown Al2O3 can provide further insight into the enhancement of device performance.From the temperature‐dependent transport measurements of MoTe2O3 is mainly attributed to the presence of oxygen vacancies in Al2O3. These oxygen vacancies create defect states within the bandgap of Al2O3 as shown in Figure  Depending on the electron occupation, each of the defect states has a charge and an energy level. At the interface of oxide‐MoTe2, if the defects occupy energy levels above the MoTe2 conduction band and donate their electrons to become positively charged, then MoTe2 becomes n‐doped, leading to high density and mobility of the electrons in the channel. The presence of Al2O3 both below and above MoTe2 enhances the doping effect, thus resulting in improved device performance.The n‐type doping in 2D materials by ALD‐grown Alin Al2O3.40 The32 FETs. The combination of the hole‐array pattern in the substrate and ALD‐grown Al2O3 layers as an additional dielectric and passivation layer enables us to achieve significant enhancements in both the hole and electron mobility of MoTe2 devices. The current and the electron mobility of our MoTe2 devices are increased by up to 6 times due to the application of tensile strain on MoTe2 by the proposed structure. The strain applied to MoTe2 changes the band structure of the material and reduces the electron‐phonon scattering in the samples, leading to high electron mobility in the devices. Additionally, we observe MIT in MoTe2 during analyzing the temperature‐dependent behavior of the devices. Overall, our proposed structure represents a promising approach for advancing the performance of MoTe2 devices for future electronics.In summary, we demonstrate a simple and effective method to modulate the carrier type and enhance the performance of few‐layer MoTe42 was cleaned with acetone and IPA. Patterns with different diameters of circles were created by electron beam lithography  after coating the substrate with PMMA A4. Then the substrate was subjected to a reactive ion etching (RIE) tool  for SiO2 etching. CHF3 and O2 were used as process gases and the rate was optimized to get 40 ± 5 nm min−1 etch rate. The etching was carried out for 2 min to get 70 nm of SiO2 etching from the PMMA opening. After etching the PMMA was dissolved in acetone.For making hole arrays, a p‐doped silicon substrate (0.001–0.005 Ω cm) covered with a 285 nm thick thermally grown SiO2O3 was grown on the hole‐array substrate using trimethylaluminium (TMA) and water at 200 °C, while the process parameters were optimized to deposit Al2O3 with 0.1 nm per cycle rate. Another deposition of Al2O3 was carried out after device fabrication. However, this time before growing the oxide, a 2 nm Al was deposited as a seed layer using e‐beam evaporation on the samples because MoTe2 has no dangling bond on the surface, therefore uniform growth of Al2O3 on 2D materials requires this seeding layer. After that, a 50 nm Al2O3 was grown using the same growth method mentioned before.The substrates with hole arrays were inserted in the ALD tool (Beneq TFS‐500). A 5 nm Al2 flakes were mechanically exfoliated from the commercially available MoTe2 crystal (2D Semiconductors) on a p‐doped silicon substrate (0.001–0.005 Ω cm) covered with a 285 nm thick thermally grown SiO2. A custom‐built 2D transfer setup was used to prepare the samples. After exfoliation, electron beam lithography  and metallization were carried out by using electron beam evaporator  under ≈10−7 torr chamber pressure to deposit a 5 nm Ti adhesion layer followed by 50 nm thick Au. The height of the flakes were determined using an atomic force microscope .The few‐layer MoTeThe room‐temperature Raman spectra were collected in back‐scattering geometry with a confocal micro‐Raman system . Samples were excited using a 532 nm laser with a spot size of less than 1 µm . Low laser power (<500 µW) was used to avoid laser‐induced damage in the samples.−1 heating rate and 100 °C min−1 cooling rate with < 0.1 °C temperature stability. All the measurements were repeated twice at each point to get confident data.All the electrical measurements were carried out with a custom‐built setup based on a Linkam LN600‐P probe station with environmental and temperature controller (T96‐S) using a source‐measure unit (Keithley 2400) and a Keithley 2700 multiplexing/voltage measurement unit. The temperature measurement system could achieve a 150 °C minThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


The properties of the bioactive solids are primarily dependent on the various types of noncovalent interactions such as hydrogen, halogen, chalcogen, pnictogen, tetrel, triel, π-stacking, cation-π and anion-π, and lone pair-π interactions . Among vSaeed et al. described the crystal structure analysis of N-((4-acetylphenyl) carbamothioyl) pivalamide, and the molecular conformation of this molecule is locked by an intramolecular N–H⋅⋅⋅O hydrogen bond that generatesa S (6) ring Saeed et al. The intermolecular N–H⋅⋅⋅O=C and C–H⋅⋅⋅π in addition to π-stacking interactions are formed in the solid state to stabilize the crystal structure. The molecular docking analysis of this compound against four biological targets, namely, α-amylase, urease, acetylcholinesterase (AChE), and butylcholinesterase (BChE), suggests that AChE and BChE have a relatively stronger binding to the compound than the other two protein targets. These findings are in good agreement with those from the in vitro experiment. Gunasekaran et al. identified potential bioactive phytochemicals from leaf extracts of Aegle marmelos. In particular, three phytochemicals, aegeline, marmin, and marminol, showed inhibition of sperm hyaluronidase activity, as evident from the in vitro assay Gunasekaran et al.. Further, in silico molecular docking analysis reveals the key active site residues are involved in the molecular recognition of aegeline, marmin and marmenol phytochemicals. These phytochemicals establish hydrogen bonds and hydrophobic interactions at the active site residues. The topological properties of the intermolecular interactions formed between the ligands and active site residues provided the strength of these interactions. Further, the noncovalent interaction (NCI) plot reveals the nature of interactions between ligand and enzymes.In this Research Topic, we collected four research articles which highlight the importance of various weak noncovalent interactions in different systems. Percino et al. synthesized two positional isomers of acrylonitrile derivatives, and X-ray diffraction analysis revealed that they exhibit similar crystal packing and hence similar optical properties Percino et al.. Hirshfeld surface analysis reveals the effect of positional isomers on the intermolecular interactions formed in the crystal structures. The characteristic intramolecular C–H⋅⋅⋅C(π) interaction is observed in these structure and helps to maintain the molecular conformation. Weak C–H⋅⋅⋅π and C–H⋅⋅⋅N interactions dominate the crystal packing. The energetics of molecular dimers held by various interactions and each of these interactions was further characterized by the QTAIM approach. This study clearly demonstrated the isomeric effect using X-ray structures, and various theoretical calculations help to design novel molecules with desired properties. Aguiar et al. demonstrated the effect of water molecules on paraquat salts. DFT calculations were carried out to obtain the chemical reactivity of the bipyridylium cation, and supramolecular arrangements were analyzed Aguiar et al.. The supramolecular assemblies are generated via various short intermolecular contacts such as O–H⋅⋅⋅O and C–H⋅⋅⋅O/Cl hydrogen bonds. It is interesting to note that the energy associated with C–H⋅⋅⋅Cl is remarkably different in hydrated and anhydrous systems of paraquat salts.In conclusion, the study of various weak molecular interactions in different molecular solids and bioactive compounds with their biological targets has gained great attention in recent years to help us understand the molecular behaviors in a solid state and in the biological system. This Research Topic clearly captures the importance of studying noncovalent interactions to better understand the system to design novel molecules with desired properties to some extent.


The 2H‐VX2  materials have a nontrivial band gap in two spin channels is found and exhibit topologically protected corner states with spin‐polarization. These corner states only accommodate the quantized fractional charge (e/3). And the charge is bound at the corners of the nanodisk geometry 2H‐VX2  in real space. The corner states are robust against symmetry‐breaking perturbations, which makes them more easily detectable in experiments. Further, it is demonstrated that the SOTI properties of 2H‐VX2  materials can be maintained in the presence of spin‐orbit coupling and are stable against magnetization. Overall, the results reveal 2H‐VX2  as an ideal platform for the exploration of magnetic SOTI and suggest its great potential in experimental detection.The transition metal dichalcogenides, 2H‐VX 2  are identified as a two‐dimensional second‐order topological insulator with a ferromagnetic ground state by first‐principles calculation. They also show a topologically protected corner state with spin‐polarization. Remarkably, these corner states are robust against symmetry‐breaking perturbations, which makes them more easily detectable in experiments.2H‐VX Recently, the concept of second‐order topological insulator (SOTI) has attracted great interest. SOTI is featured by the gapless states on its (d‐2)‐dimensional boundary. For example, three‐dimensional (3D) SOTI has one‐dimensional (1D) hinge states in gapped surface states; And two‐dimensional (2D) SOTI has zero‐dimensional (0D) corner states in gapped edge states. In order to accurately evaluate whether a material is SOTI, some new bulk topological invariants with corresponding symmetries are used. So far, SOTIs have been predicted in some nonmagnetic and magnetic 3D materials, while 2D SOTIs have been found mainly in nonmagnetic materials.Topological insulator (TI) has been one of the most active fields in condensed matter physics since their discovery.7, 8 In addition to these properties, 2D TMDs have also been predicted to host a novel topological electronic structure. Recently, 0D corner states were discovered in the band gap of hexagonal TMDs (h‐TMDs) with a triangular prism structure. The robust (d‐2)‐dimensional boundary correspondence in h‐TMDs is protected by the rotation symmetry. Interestingly, based on the hidden breathing Kagome lattice (BKL) structure mimicked by the d‐orbitals from the transition metal atoms, this nontrivial topology in h‐TMDs can be explained by the higher‐order topological insulators concept in BKL. In Ref, it was also revealed that the band inversion of the d orbitals indicates a nontrivial Wannier center (WC), which is a signature of SOTI for the 2D h‐TMDs. And they used these theoretical analyses to confirm that some h‐TMDs materials, such as 2D MX2 , are indeed SOTIs. In comparison to graphdiyne and 2D hexagonal group‐V materials, 2D TMDs have a larger band gap (≈1.7 eV), facilitating separation of corner states from edge/bulk states in their electronic band structure. Notably, transition metal ions that have partially filled d shells are commonly associated with magnetism. Given the advantages of TMDs in nonmagnetic SOTI systems, magnetic TMDs may provide an ideal platform to study magnetic SOTI.In 2D materials, transition‐metal dichalcogenides (TMDs) are a class of materials that have garnered significant attention due to their exceptional performance in electronics, electrocatalysts, optical properties, and rich valley‐related physics.32 In  In particular, vanadium‐based dichalcogenides such as VX2  have been found to exhibit inherent ferromagnetism at room temperature and semiconducting properties in their 2H‐phase monolayer structure. In addition to the above characteristics, 2H‐VX2  also has a triangular prism structure, so 2H‐VX2  is expected to be a favorable candidate for magnetic SOTI. Moreover, the monolayer 2H‐VS2 has been successfully synthesized in the experiment, which enables the experimental detection of properties in the magnetic SOTI soon.Recent experimental and theoretical studies have suggested the presence of ferromagnetism in single‐layer TMDs.38 In 2  exhibit a ferromagnetic ground state and are magnetic 2D SOTIs. The nontrivial topology and 0D corner states in these materials are protected by the rotation symmetry of the lattice. Based on the first‐principles calculations, we confirm that VX2  have gapped bulk and edge states in two spin channels. Notably, the spin‐polarized corner states can be extracted in the gap. By symmetry analysis, we verify that the corner states carry a quantized fractional charge protected by C3 rotation symmetry. We also explore the stability of corner states and find that the SOTI phase is robust to symmetry‐breaking perturbations. The SOTI properties in VX2  can still be preserved even if holes are introduced to break C3 symmetry or spin‐orbit coupling (SOC) is included to break T symmetry. And the corner states accommodating a quantized fractional charge remain to be localized at the corner of the sample.In this work, we show that VX2  have a similar topology as well as the monolayer 2H‐VS2 has already been synthesized in experiment. So, as a typical example, we will focus on the results obtained for monolayer 2H‐VS2. The results for other materials are presented in Supporting Information.Since the electronic band structures for all the VX2 for first principles calculations by Vienna Ab initio Simulation Package (VASP). The exchange‐correlation potential is treated with the generalized gradient approximation (GGA) of Perdew‐Burke‐Ernzerhof function. Projector augmented wave potentials are provided by VASP and have been used on all elements. For the monolayer 2H‐VS2 crystal structure, a large vacuum space of 20 Å is established to avoid potential interactions between the layers. In our calculations, the cutoff energy is set as 500 eV. The Brillouin zone is sampled by a Monkhorst‐Pack k mesh with a size of 9 × 9 × 1. To optimize the crystal structure, the force and energy convergence criteria are set as 0.001 eV Å−1 and 10−6 eV, respectively. To investigate the dynamical stability of the monolayer 2H‐VS2, the phonon spectrum is calculated by the PHONOPY code. Since V is a transition metal element, the Coulomb interaction (U) between electrons is considered and GGA+U is used to calculate the band structure. In the main text, the calculation results are based on the effective Coulomb energy, Ueff, set as 2.5 eV. The conclusions in this work do not change over a large range of U values (0–4 eV) 40 for33.12  is depicted in Figure P6¯m2, and the corresponding point group is Dh3 generated by C3 and C2′. The primitive cell contains one V and two X atoms, and the lattice constant for the monolayer 2H‐VX2  is provided in Table 2 are consistent with the previous report.The lattice structure of 2H‐VX2’s stability can be assessed through its phonon spectrum and thermal stability. Figure 2. In addition, spin‐polarized ab initio molecular dynamics (AIMD) simulations are carried out in a 2 × 2 × 1 supercell to evaluate the thermal stability of the monolayer 2H‐VS2. Figure 2 is thermally stable. In addition, we have also investigated the stability of monolayer 2H‐VSe2 and 2H‐VTe2  is ≈ 292 K for the monolayer 2H‐VS2  materials). Before studying its electronic band structure, we also determine an easy plane of its ferromagnetic state by using the DFT+U+SOC calculations. Figure zox (zoy) and xoy planes. These results clearly indicate that the in‐plane (xy‐plane) magnetization dominates in monolayer 2H‐VS2.It is worth noticing that the monolayer 2H‐VS3.22 are given in Figure u) and spin‐down (Δd1) channels, respectively. Akin to monolayer MX2 , Bloch states near the band edges for the conduction and valence bands in the spin‐up channel are primarily contributed by the d‐orbitals of the transition metal atom . This is demonstrated in Figure dx2−y2,dxy) and dz2 for spin‐up states at K point. Following the work of Fang et al and Jung et al, if the model has C3, its polarization can be directly deduced from the eigenvalues of C3, giving rise to a nontrivial WC position, as shown in Figure d2 ≈ 0.74 eV) in the spin‐down channel is also nontrivial due to the presence of the d‐orbitals inversion.The electronic band structure and partial density of states (PDOS) of different spin channels for the monolayer 2H‐VS Se, Te),49 Blou and Δd1 can be attributed to the symmetry protection. It is necessary to point out that the spin‐up and spin‐down states are decoupled, which is a natural consequence of the absence of SOC in a FM state, such that the system preserves all the original symmetries in the absence of FM, including the time‐reversal symmetry. Consequently, the presence of C3 rotation symmetry allows for the existence of corner states with quantized fractional charge, which can be derived from the C3 symmetry eigenvalues of the occupied bands at high‐symmetry points (HSPs) in the Brillouin zone. We denote HSPs as Π(3), and the superscript refers to the threefold rotation symmetry. Thus, the eigenvalues of C3‐symmetry at HSPs can be expressed as[Πp(3)] can be defined based on the eigenvalues of C3 compared to a reference point Γ  =  , explicitly given by#Πp(3)(#Γp(3)) is the number of occupied bands with corresponding Πp(3)(Γp(3)). The topological indexes for both spin channels at HSPs have been listed in Table (Qcorner(3)) can be calculated using the following formula:n1aion stands for the ionic number at Wykoff position 1a, and ν is the number of occupied bands, e is the charge of a free electron. For the monolayer 2H‐VS2, we derive the topological index [Kp(3)]=, and Qcorner(3)=e/3 for both spin‐up and spin‐down states. The corner charges are also calculated for the other 2H‐VX2  materials, which are shown in Table In addition, the nontrivial topology for the band gaps Δ3.32 in the form of a triangle nanodisk that preserves the C3‐symmetry. The electronic band structure obtained from TB calculations is consistent with the result derived from DFT calculations, as shown in Figure 2 has a zigzag metallic edge  model for the monolayer 2H‐VSC3‐symmetry, and then the corner states were recalculated. It was found that the corner states with the same energy and degeneracy still exist in the gap well preserved, as shown in Figure C3 symmetry still dominates. Upon checking the spatial positions of the corner states, we confirm that they are still localized at the original position, namely, the corners of the nanodisk. The robustness of these corner states would greatly facilitate their experimental detection in the future. One may notice that some additional 0D states are also appearing in the gap =e/3 is achieved as well, indicating a nontrivial topology. To check the corner states, a 14 × 14 supercell was constructed by Wannier Tools to form a parallelogram nanodisk. The wave function distribution is plotted in Figure C3 symmetry. The degeneracy associated with C3 symmetry was destroyed. Consequently, corner states appear randomly within the band gap. However, in real space, the corner states are still localized at the corners of the nanodisk, as shown in Figure C3 and T symmetries are broken. However, the band gap remains open during the rotation of the magnetization direction, as depicted in Figure Table For the monolayer 2H‐VS3.4 some sonic crystals and photonic crystals, and also electric circuits. However, it remains a challenge for the experimental observation of SOTIs in real materials due to the strict requirements of single crystal samples and scanning tunneling microscopy. Up to date, only a few 3D materials have been experimentally evidenced to show SOTI state with 1D hinge states, such as Bi, Bi4Br4, and Td‐WTe2. The detection of corner states has not yet been realized in experiment for 2D materials due to the lack of applicable materials. Fortunately, the high‐quality monolayer 2H‐VS2 has already been grown by a molten salt‐mediated precursor system and epitaxial mica growth platform. In addition, the semiconductor properties of monolayer 2H‐VS2 with a large global band gap have been experimentally confirmed, which greatly favor the experimental detection of corner states since these states are well separated from the global bands and the edge states  proposed in this work has already been synthesized by advanced experiments, and its magnetic ground state has also been confirmed. Thus, the SOTI band structure, the topological corner states, and the potential valley polarization effect proposed in the literature are potentially detected in monolayer 2H‐VS2.Quite recently, we come aware of a pioneer work that reported the magnetic SOTI state in two 2D materials including 2H‐RuClnus VSSe,65 wit42 as an example, we theoretically reveal the presence of magnetic SOTI states in a family of transition metal dichalcogenides 2H‐VX2 . The free‐standing monolayer 2H‐VS2 is dynamically stable and naturally shows in‐plane FM ordering. There are diverse size band gaps in different spin channels of the 2H‐VX2 , which is a magnetic semiconducting signature. Both band gaps are in topological non‐triviality with corner states possessing quantized fractional charge and ensured by the 3C rotation symmetry. These topological corner states are found to be robust against holes, SOC, and magnetization direction. In addition, the corner states in the 2H‐VX2  are fully spin‐polarized, which is different from those observed in nonmagnetic SOTIs reported previously. Our findings provide important insights into the high‐order topology in 2D magnetic systems.In conclusion, taking the monolayer 2H‐VSThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.


However, the detailed kinetics of the electrochemical processes in the cathodes is still unclear. In this work, the sodiation/desodiation process of Na3VP is first investigated using in situ high‐resolution transmission electron microscopy (HRTEM). The intermediate Na2V2(PO4)3 (Na2VP) phase with the P21/c space group, which would be inhibited by constant electron beam irradiation, is observed at the atomic scale. With the calculated volume change and the electrode–electrolyte interface after cycling, it can be concluded that the  Na2VP phase reduces the lattice mismatch between Na3VP and NaV2(PO4)3 (NaVP), preventing structural collapse. Based on the density functional theory calculation (DFT), the Na+ ion migrates more rapidly in the Na2VP structure, which facilitates the desodiation and sodiation processes. The formation of  Na2VP phase lowers the formation energy of NaVP. This study demonstrates the dynamic evolution of the Na3VP structure, paving the way for an in‐depth understanding of electrode materials for energy‐storage applications.Recently, all‐solid‐state sodium batteries (Na‐ASSBs) have received increased interest owing to their high safety and potential of high energy density. The potential of Na‐ASSBs based on sodium superionic conductor (NASICON)‐structured Na 3VP is demonstrated with in situ transmission electron microscopy (TEM). Outside the in situ area, the intermediary Na2VP phase with the space group P21/c is observed at atomic scale, which reduces the lattice mismatch between Na3VP and NaVP. The Na2VP phase lowers the formation energy of NaVP and possesses lower energy barrier for Na+ diffusion.The structural evolution of Na For the automotive sector, the development of LIBs with liquid electrolytes has stagnated owing to its flammable liquid electrolytes and insufficient energy density. The ever‐increasing demand for energy storage systems has invigorated the need to seek alternatives with higher performance, enhanced safety, and increased economic efficiency compared to conventional LIBs. In this respect, all‐solid‐state batteries (ASSBs) are regarded as the potential solution. Although the phenomenon occurring in a liquid electrolyte system and in a solid‐state electrolyte system is identical, ASSBs can eliminate the safety concerns of traditional batteries by replacing flammable liquid electrolytes with solid electrolytes. Furthermore, solid electrolytes are more chemically and electrochemically stable than liquid electrolytes. In addition, batteries with solid‐state design also show better adaptability to temperature changes, and have the potential of achieving both higher energy and power densities.Ever since Li The search for alternative high‐energy batteries using abundant and low‐cost materials is imperative. Among the possible candidates, sodium‐ion batteries (SIBs) possess advantages such as abundant sodium supply, high energy storage capacity, and high electrochemical activity, which have attracted the attention of researchers. Na3V2(PO4)3 (Na3VP) is a promising electrode material, which crystallizes in a NASICON structure. It shows very promising capability in Na‐ASSBs. Having the same structure as the highly conductive solid electrolyte Na3.4Zr2Si2.4P0.6O12 (NZSP), Na3VP showed excellent chemical compatibility with NZSP and great cycling stability when used as the positive electrode material.Li has been the preferred in battery chemistries in recent decades. However, the cost of future LIBs remains uncertain due to the increasing demand of LIBs and uneven geographical distribution of Li.7, 83VP is a 3D open framework composed of repeating lantern units, where two VO6 octahedra are joined together by three corner‐sharing PO4 tetrahedra. There are two different Na sites located in the framework: Na(1) and Na(2). The Na(1) sites are the octahedral sites between two VO6 octahedra, whereas the Na(2) sites are the tetrahedral sites between two adjacent PO4 tetrahedra along the c‐axis direction. The open diffusion paths and rigid structure in the polyanionic sublattice can be attributed to the strong covalent bonds between the oxygen atoms. The Na+ ions can reversibly diffuse through the interstitial sites without a significant volume change in the crystal structure, achieving high theoretical energy storage capacities and superior ionic mobility. Previous studies have demonstrated the high energy density, high rate, and long cycle life of Na3VP‐based sodium‐ion batteries, which are closely related to the Na3VP microstructure. However, the structural evolution of Na3VP during sodiation/desodiation remains ambiguous. Therefore, it is essential to investigate the Na+ ion insertion and extraction mechanism of Na3VP.The structure of Name change19 in  and in situ Raman methods. For example, the low strain of the Na3V2(PO4)2O2F lattice with a volumetric variation of only 2.56% during Na intercalation/extraction processes was disclosed by operando XRD. In situ Raman analysis revealed a highly reversible three‐phase transition as the sodium‐ion storage mechanism of Na2FeFe(CN)6 during sodiation/desodiation processes. Among various in situ techniques, in situ TEM possesses the capability of resolving the microstructural evolution and a providing diffraction information of the electrode materials. It can create a variety of external fields, including electric and thermal fields, and can observe the dynamic structure evolutions at the same time. Based on the advantages mentioned above, in situ transmission electron microscopy (TEM) has also been utilized to observe various reaction phenomena, such as catalytic phenomena, electromigration, and electrochemical reaction of batteries. In this work, an all‐solid‐state battery cell was prepared and the phase transition of Na3VP under charging/discharging was investigated using operando X‐ray diffractometer (XRD) analysis. Moreover, the structural evolution of Na3VP during sodiation/desodiation process was observed via in situ high‐resolution transmission electron microscopy (HRTEM) for the first time. The selected area electron diffraction (SAED) images of Na3VP demonstrated the evolution of lattice spacing, verifying the Na+ ions insertion/extraction during sodiation/desodiation process. Density functional theory (DFT) computations were employed to determine the formation energies of the cathode and the energy barrier for Na+ ions diffusing in the cathode structure. In‐depth information about the mechanism of the phase transition in the Na3VP–NaV2(PO4)3 (NaVP) system will be discussed in future research. This study unveils the structural evolution of Na3VP during the sodiation/desodiation process, providing fundamental information on Na3VP and laying the foundation for further research and future applications of Na‐ASSBs.There have been growing interests in developing in situ techniques for battery studies, such as operando X‐ray diffraction (XRD) technique23 and23VP/NZSP sample. The XRD spectra of the sample showed that the diffraction peaks of Na3VP and NZSP were both indexed to the NASICON structure with the R3¯c space group 3 (8.68%), Na4MnCr(PO4)3 (7.7%), and Na3.8MnV0.8Zr0.2(PO4)3 (4.9%). The result is against previous study since the larger volume change usually leads to poor cyclability.Figure   are closed to the diffraction peaks (14.76° and 36.05°) appearing in the mid‐charge. It is suggested that the new reflections are attributed to the new intermediate Na2VP phase.To investigate the structural evolution of Na° Figure . During 0.8266 Å.41 Bas3VP cathode is investigated via in situ HRTEM. Figure 2¯6) and (13¯5) planes of the Na3VP crystal. After the desodiation process, the lattice spacings narrowed to 1.835 and 1.84 Å, respectively, which correspond to the same lattice planes of NaVP with space group R3¯c   and (13¯5) upon desodiation. This result indicates that the lattice contracted upon the extraction of Na+ ions. Figure 2V2(PO4)3 (Na2VP), with a different crystal structure from Na3VP and NaVP, which was observed outside the in situ observation area  of Na3VP, Na2VP, and NaVP are 241.277, 227.159, and 220.105 Å3, respectively. The calculated volume changes from Na3VP to Na2VP and from Na2VP to NaVP were 5.9% and 3.1%, respectively, while the volume change in the phase transition between Na3VP and NaVP was 8.8%.To further understand the Na‐ion insertion/extraction mechanism, the desodiation process of the Naoup R3¯c 42 3]3− framework. Therefore, the total volume changes from Na3VP to Na2VP and from Na2VP to NaVP (9.0%) is slightly larger than that from Na3VP to NaVP (8.8%). It is suggested that the gradual volume change, that is, from Na3VP to Na2VP and from Na2VP to NaVP, could reduce the lattice mismatch between Na3VP and NaVP, which will be further validated below. EDS mapping analysis of the desodiated sample showed that the Na signal of the cathode weakened after desodiation, indicating the extraction of Na+ ions 3]3− framework  and performance degradation. In Figure 2VP phase was observed, which also implies structural evolution during the desodiation process. However, the high stability and cycling ability of Na3VP has been demonstrated previously, which seems to contradict the studies mentioned above unless the occurring lattice distortions are reversible. Hence, an investigation of the phase transition between Na2VP and Na3VP is essential. Figure2VP to Na3VP, with the sample sodiated to 4 V. Initially, the lattice spacings were 2.542 and 2.992 Å, corresponding to the (602¯) and (32¯2¯) lattice planes of the Na2VP phase. During sodiation, the dihedral angle gradually changed from 55.2° to 57.8°. The lattice spacings were 2.519 and 2.788 Å after sodiation for the (300) and (116¯) lattice planes of the Na3VP phase, respectively (the sodiation process is shown in Movie 2¯) ((300)) and (32¯2¯) ((116¯)) upon sodiation. The EDS mapping of the sodiated sample showed that the Na signal became stronger  becomes larger, causing the structural degradation of cathode material. To avoid this situation, the in situ area with constant electron beam irradiation (to avoid the formation of Na2VP) was observed under relative low magnification during applying bias. After reaching 4 V, the interface was further observed with higher magnification. After three cycles, the interface outside the in situ area (with Na2VP) remained continue   shown in Figure 3VP and NSZP still exhibited excellent contact. It is suggested that the optimization of the in situ device mitigated the thermal effect on the sample, improving its rechargeability. EELS was utilized to investigate the valence states of V. Figure 2,3 spectra of pristine, desodiated, and sodiated Na3VP. The L‐edge shifted toward higher energies due to Na+ ion extraction. The valence state of V changed from +3 to +4, which corresponds to the desodiation from Na3VP to NaVP. The L‐edge shifted toward lower energies when the Na+ ions were inserted, suggesting the recurrence of Na3VP.In addition to the reversibility of the sodiation/desodiation process as confirmed via the in situ TEM and SAED patterns, rechargeability is also an important quality for a secondary battery. For the TEM sample fabricated by FIB, the sample has a lamellar shape and a thickness of ≈100 nm. The lamellar shape has a much smaller cross‐sectional area compared with the bulk Na‐ASSB, resulting in a higher current density. Moreover, contacts of Pt wires and the sample  site and the other two are located at the Na(2) sites. The Na(2) ions are considered to be the main participants in the transportation of Na+ ions because of their longer Na–O distance than that of the Na(1) site. Figure+ ions were extracted. The lantern unit consists of two VO6 octahedra sharing corners with three PO4 tetrahedra. The yellow and blue spheres represent Na+ ions at the Na(1) and Na(2) sites, respectively. In the rhombohedral Na3VP, both Na(1) and Na(2) sites are partially occupied  sites, resulting in less occupancy of this position. The extraction of the Na(2) ions reduces the rhombohedral symmetry due to a distortion of the lattice, transforming the rhombohedral structure into a monoclinic structure  sites, the monoclinic structure then transforms back to the rhombohedral phase , while Na2VP and NaVP manifest lower formation energies of −2.264 and −2.280 eV per atom, respectively. This indicates that the Na3VP cathode prefers to transform to the Na2VP or NaVP structures. Moreover, the desodiation energy from Na3VP to Na2VP is much lower than that from Na3VP to NaVP, demonstrating that the Na3VP structure is more likely to transform to Na2VP structure during the desodiation process . For NaVP, where all Na ions only occupy Na(1) sites, the extraction of Na ions may affect the structural integrity. Therefore, Na migration between Na(1) sites at NaVP is hindered by a higher migration barrier than Na3VP. This suggests that the Na+ ion migrates more rapidly in the Na2VP structure compared to Na3VP and NaVP, thereby decreasing the resistance of the sample  remain the same in spite of the state of the aperture. Although the electron irradiation of TEM may cause irreversible damage to the sample, it can provide a feasible stimulus that mimics the effect of electrochemical cycling. Huang et al. reported that electron beams could reclaim Li metal from the degradation product and serve as a local Li source, which triggers the lithiation of NiFe2O4/carbon nanotubes. LiNi0.4Mn0.4Co0.18Ti0.02O2 particles, repeated electron beam irradiation induced a phase transition from an R3¯m layered structure to an Fm3¯m rock‐salt structure, which is attributed to the stoichiometric lithium and oxygen removal from R3¯m 3a and 6c sites, respectively. This implies that the electron beam can be used as a high‐energy particle flow to change the microstructure of sensitive materials. In Figure 3VP particle was exposed to a high‐energy electron beam throughout the entire desodiation process. Na+ ions could be easily extracted from Na3VP owing to the interaction between electrons and atoms. Therefore, the Na3VP phase could transform to the NaVP phase directly without the formation of the intermediary Na2VP phase  sites in the rhombohedral structure. The Na2VP phase lowers the formation energy of NaVP and possesses lower energy barrier for Na+ diffusion. This study allows a better understand of the behavior of fast‐rate electrode materials that undergoes phase transition, providing insight into the future design of Na3VP.In conclusion, a Na4−1 kg per batch) using a solution‐assisted solid‐state reaction method (SA‐SSR). NaNO3 (VWR), ZrO(NO3)2 (Aldrich), Si(OCH2CH3)4 (Merck), and NH4H2PO4 (Merck) were used as starting materials. All the chemicals applied in this study were analytically pure. Stoichiometric amounts of NaNO3 and ZrO(NO3)2 were dissolved in deionized water. A stoichiometric amount of Si(OCH2CH3)4 was also added to the solution while stirring. When Si(OCH2CH3)4 was hydrolyzed, the stoichiometric amount of NH4H2PO4 was added to the system during stirring. The homogeneous aqueous solution immediately showed formation of complex zirconium oxyphosphate compounds. The whole mixture was dried at 85 °C. The dried powder was then calcined at 800 °C for 3 h. After calcination, a white powder was obtained. The calcined powder was then milled in ethanol with zirconia balls on a milling bench for 48 h, and dried at 70 °C for 12 h.NZSP powder was synthesized on a laboratory scale (10 g(SA‐SSR).55 NaNThe NZSP powder was put into a cylindrical pressing mold (13 mm in diameter) and pressed with a uniaxial pressure of about 100 MPa at room temperature. The pressed pellets were sintered between 1260 and 1300 °C for 5 h. The obtained pellets had a diameter of about 10 mm, and thickness of 1–2 mm. The relative density of the sintered pellets was >95%.3VP precursor solution was prepared by mixing ethanolamine, de‐ionized H2O, NaH2PO4, and NH4VO3 (weight ratios of 1:2:0.71:0.46). A cotton swab was dipped into the Na3VP precursor solution and then swept over the surface of the sintered NZSP pellet, leaving a thin and homogenous Na3VP precursor solution layer on the NZSP pellet. The pellets were then heated to 740 °C in Ar‐2% H2 for 4 h to form the Na3VP phase on the NZSP surface.The NaFor the operando XRD analysis, Au was then sputtered on the positive electrodes to serve as current collector. The cell was assembled inside Ar‐filled glove box, subjected into X‐ray examination during charging/discharging at a rate of 0.4 C. Bruker D2 Phaser and BioLogic SP‐150 Potentiostat are used as X‐ray diffraction instrument and a modulo‐battery‐mode electrochemical tester, respectively.3VP/NZSP was prepared using a high‐resolution dual‐beam focused ion beam (FIB) system (FEI Versa 3D). Because of the unevenness of the Na3VP/NZSP sample  pattern of the pristine sample was measured using a Brucker D2 Phaser with Cu Kα radiation. After the samples were sodiated/desodiated to 4 V, the valence states of V in Na with the projector augmented wave (PAW) method based on the density functional theory (DFT). The Perdew–Burke–Ernzerhof (PBE) functional within the generalized‐gradient approximation (GGA) was used to describe the exchange‐correlation interaction for the structure optimization. The cut‐off energy was set to 600 eV for all the calculations. The first Brillouin zone was sampled with 3 × 3 × 1, 2 × 3 × 3, and 3 × 3 × 1 k‐points grids for Na3VP, Na2VP, and NaVP crystals, respectively. The energy convergence criterion was set as 1.0 × 10−5 for all the calculations. The optimized lattice parameters  are good consistent with the experimental values.First‐principles calculations were conducted by using the Vienna ab initio simulation package (VASP)57 witEf) of the NaVP crystal is expressed by Equation (ENax+yV2(PO4)3 and N are the total energy and the total number of atoms of the NaVP crystals per formula units ; in and iµ are the number of atoms and the chemical potential of the species i, respectively. The chemical potentials of Na, V, and P were conducted using the total energy per atom in the bulk system, and that of O was obtained using O2 molecule.The formation energy (Equation (2)Ef=ED) of the NaVP crystal during the desodiation process is expressed by Equation (ENax+yV2(PO4)3 is the total energy of the Nax+yVP crystals per formula units The desodiation energy (Equation (3)ED= and the diffusion barriers related to the energy difference between the initial‐state (IS) and the TS. All atoms were allowed to relax until the force components were less than 0.05 eV Å−1.The transition‐state (TS) was determined using the climbing image nudged elastic band (CI‐NEB) calculation through the VTST code62 andThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.Supplemental Movie 1Click here for additional data file.Supplemental Movie 2Click here for additional data file.


Thus the ionic transport and diffusion properties of the N–H co-catalyst are fundamentally important to understanding and developing such syntheses. Here we show hydride ion conduction in these materials. Two distinct calcium nitride-hydride Ca2NH phases, prepared via different synthetic paths are found to show dramatically different properties. One phase (β) shows fast hydride ionic conduction properties (0.08 S/cm at 600 °C), on a par with the best binary ionic hydrides and 10 times higher than CaH2, whilst the other (α) is 100 times less conductive. An in situ combined analysis techniques reveals that the effective β-phase conducts ions via a vacancy-mediated phenomenon in which the charge carrier concentration is dependent on the ion concentration in the secondary site and by extension the vacancy concentration in the main site.Recently nitrogen-hydrogen compounds have successfully been applied as co-catalysts for mild conditions ammonia synthesis. Ca 2NH) are explored and connected to their respective ionic conductivity. These materials have previously been investigated as a co-catalysts/support material for ammonia synthesis.The detailed structures of two calcium nitride-hydride phases (Ca The long standing issue for MCAS to overcome relates to a scaling relation, due to the saturation of the transition metal  with adatoms of H2, N2, and NHx9. For MCAS, lowering of the activation energy (Ea) of adsorption for reactants (N2 and H2) also causes an increase in the adsorption energy (ΔE) of intermediate products (NHx); the result being that the catalyst’s surface is quickly saturated and the overall reaction rate drops dramatically. Recent research has shown that catalytic supports including alkali and alkaline earth metal hydrides (H−), nitrides (N3−), imides (NH2−), amides (NH2−), and nitride-hydrides (N3−-H−) can significantly promote the synthesis of NH3(g)8.Recently, nitrogen–hydrogen (N–H) compounds have attracted interest as potential co-catalysts in mild conditions ammonia synthesis 2 Ea and NHx ΔE as they have a positive reaction order with respect to H2(g). The mechanism that underlies this change is widely believed to result from the transfer of H-adatoms from the TM surface to the N–H compounds7; in some cases N-adatoms are additionally transferred10. This removal of the H and N adatoms frees the catalytic surface for the reaction of NHx species into in NH3(g). Thus, good transport properties, specifically, ionic conductivity of the support materials are crucially important for the overall heterogeneous system. Recently, several high performing hydride ion conductors (on the order of 10−2 S cm−1 or better) have been reported in the literature13. Both barium hydride and the oxygen doped lanthanum hydride have been reported to show good MCAS activity14, while similar oxyhydrides and oxynitride-hydrides to Ba-Li based oxyhydride of Takeiri et al. have also shown good activity16. This paper reports on a fast H- ionic conductor: β-Ca2NH, as well as establishing a nomenclature to distinguish between species of calcium nitride-hydride.These N–H compounds are proposed to do this by breaking the scaling relationship between H2NH was first reported by Brice et al. as forming in the Fd-3m space group with Ca2+ ions (32e) forming a slightly distorted face-centred cubic close pack arrangement, and the nitride and hydride species occupying the octahedral positions 17. A 25% intrinsic vacancy concentration due to H- ions moving to the 96 g tetrahedral position was observed. More recently, Verbraeken et al. suggested that the structure was better fitted using the R-3m space group, which differs from the Fd-3m in the ordering of the octahedral positions18. Additionally, they found a 13% intrinsic vacancy concentration at ambient temperature. The R-3m space group is also observed for the barium and strontium analogues20. Finally, Verbraeken et al. noted the presence of a secondary quasi-imide phase that constituted approximately 10% of their system. This phase disappeared upon heating. Sichla et al. and Chemnitzer et al. observed previously that the nitride-hydrides of Sr form mixed phases containing nitride, hydride, and imide anions21. Additionally, these types of systems have been observed and studied in lithium NH compounds23. These disordered phases comes in many forms depending on the synthesis route. A paper by Weidner et al. gives a considered discussion around these quasi-imide phases24. Importantly, the disruption in site occupancy in these mixed phases posed a significant challenge for the application of these materials in reversible solid-state hydrogen storage24. Makepeace et al. suggest that a compositional compromise must be struck in cycling lithium imide to lithium amide during ammonia synthesis as the prior has significantly better electrical performance than the latter23.CaWe are aware of four approaches to produce calcium nitride-hydride and have summarized these in Supplementary Table 2 and H2 gases were chosen: a hydride25 (CaH2) and a nitride (α-Ca3N2)27. These compounds were heated to 600 °C under flowing 5% H2/D2 and N2 in Ar respectively. The properties of the resulting phases were analysed from data collected from in situ neutron powder diffraction (NPD) collected on Polaris Diffractometer at ISIS28 and impedance spectroscopy . Details of these analyses are available in the Supplementary Methods. Next, the CaH2(D2) system was exposed to N2 gas, while the α-Ca3N2 system was exposed to H2/D2. The result in both systems was the emergence of a Ca2NH phase.Two precursors that had been previously shown to react with N3N2 . This space group assignment has a lattice parameter that is half the size for the analogous structure using the Fd-3m assignment. The best fit was produced by placing the Ca2+ ions in 4a position and N3− and D− (deuteride) ions in the 4b position. Furthermore, a 20% concentration of imide (NH2-) protons were placed in the 24e position which correspond to proton positions in the published calcium imide structure30. Thus, the reaction of α-Ca3N2 with H2(g) at 600 °C produces a mixed/disordered phase which we label α-Ca2NH. We also note that similar to other systems containing secondary anion species23, the system exhibits pronounced asymmetric peaks upon cooling  as a result of a phase equilibrium between the produced Ca2ND phase and the N-doped CaD2 precursor  Å vs 10.23062(24) Å for the α- and β-phases, respectively. Previous research has shown that solid solutions of NH compounds have larger lattice parameters than the pure compounds in agreement with the findings here32.The Ca17. Full refinement yielded an occupancy of the D1 octahedral site of 0.822(11) as opposed to the value of 0.75 reported by Brice et al., with the remainder being located in a tetrahedral site  giving a stoichiometric compound (Ca2NH0.997(34)). One major difference between the fit in this paper and that of Brice et al. is the position of the D2 site. It was found that refinement of the 96g position proposed by Brice et al. was unstable and required that the atomic displacement parameter (ADP), occupancy, and coordinates be fixed. Using the 48f position instead allowed for the full refinement of the position, although the low occupancy and high multiplicity create larger than normal errors. A comparison between the fit results of the two phases is shown in Table 2NH from CaH2 required temperatures above 400 °C. Additionally, heating β-Ca2NH to 800 °C resulted in the appearance of shoulder peaks  data collected during the reaction of CaDsee Fig. . The ris2 instead of D2) found that the upon reaction of α-Ca3N2 with hydrogen, the measured conductivity plummeted by 4 orders of magnitude  at 1.0 ppm and imide at (NH2-) at −2.7 ppm. The amide appears more prevalent in the spectrum due to the fact that its peak is quite sharp compared to the imide which is not only positionally distributed around the bonded N atom, but also randomly distributed throughout the host lattice. The final peak at −1.3 ppm is thought to be associated with OH- groups formed during loading and unloading the sample into the NMR instrument.To further explore the structural difference between the two phases, NMR and Raman spectra were collected at room temperature. Figure −1)7. This peak is also shifted to lower wavenumbers as compared to the β- (328 cm−1) suggesting that the configuration of the α-phase is lower energy. This result matches well with published results on disordered lattices36. The α-phase also shows the characteristic broad peaks between 100 and 1000 cm−1 of the imide species, as well as the sharp peak at 3125 cm−122.Raman data also show broadening in the α-phase spectrum for the peaks associated the nitride-hydride species  at 600 °C. NPD patterns were collected before and after each dose, while EIS and quasi-elastic neutron scattering (QENS) data were collected over the course of several hours. Further details and example data and fits are available in Supplementary Figs. N followed by a numeral representing the dose number, while the number after the hyphen is the dataset for that dose.The β-phase conductivity was explored further on Osiris spectrometer at ISIS2NH as measured by NPD as well as a jump in the measured ionic conductivity as measured by EIS (2NH phase is again correlated with the rise in - ions. The QENS data were modelled by the Chudley-Elliott jump diffusion model (CEM)38 which produces two important physical characteristics of the diffusion process: the jump length (l), and the residence time (τ). In order to understand the mechanism of diffusion it is important to contrast the extracted l with site to site distances. Figure 2NH phase refinement of the Osiris diffraction patterns as well as the l extracted from the CEM. The results show that the l values are close to the octahedral site to site distances. Therefore, conductivity values were calculated for a H1-H1 pathway using the Einstein Diffusion equation in conjunction with the Nernst-Einstein equation:n is 2x dimensionality of the diffusion process , i (atoms cm−3), and i. The number densities were calculated from the NPD structure refinements using the following equation:i, and −1 for 3D diffusion, and ~0.7 S cm−1 for 2D) than the measured −1). However, it was noticed that if the −1). Since, the 2NH structure which refines as overall stoichiometric. Thus, a modified Nernst-Einstein equation that takes into account the fraction of charge carrier sites that have neighbouring vacancies  Fig. . Thus, teffective charge carrier concentration:-) from - . A similar result is seen in the difference in conductivity between lithium imide  and lithium amide  where the former showed good conductivity properties due to the creation of 2 Frenkel pair defects40, similar to what we see here.There are two equivalent ways to think about the effect of this term. First, one can consider an 2 shows that the 2NH system is higher than that measured for BaH229; while the overall conductivity at 600 °C for the two systems in significantly higher for BaH2—the difference arising from the significantly higher effective charge carrier concentrations in the BaH2 due to the presence of a concerted migration mechanism. Table Comparing the results of this combined in situ QENS and EIS experiment with previously published research on BaH- ion conductor β-Ca2NH and also shown that the closely related α-phase (containing disorder due to the presence of secondary anionic species) has poor ionic conductivity potentially as a result of the lack of intrinsic vacancies caused by the presence of the imide species in the phase; this result has important implications for MCAS in that the ionic conductivity of the support material may play a fundamental role in promoting the TM catalyst via the removal of adatoms from the surface. This is evinced by the Mars-van Krevelen mechanism observed by Kitano et al. and by the chemical looping of BaH2 systems by Gao et al.7, both implying significant contributions to the overall process from the bulk of the N–H support materials.We have presented the HThis paper has provided direct evidence for a vacancy-mediated conductivity pathway by using combined analysis of simultaneously collected QENS and EIS data. Finally, this paper has shown that secondary characteristic techniques such as NMR and Raman are important to verifying and understanding the crystal structure of N–H materials.All materials were handled under Ar atmosphere, as the samples are air-sensitive. Under these conditions, the phases are stable.2 was synthesized from pure calcium shot in a sealed reactor at 800 °C for 6 h under flowing 5% D2 in Ar at 100 cm−3 min−1. 25 mm pellets were pressed of CaD2 and α-Ca3N2 (purchased from Alfa Aesar >99%) and sintered at 800 °C for 2 h under flowing 5% D2 / N2 respectively in Ar at 100 cm−3 min−1. The sintered pellets were then painted with palladium ink (C2031105P2) from Gwent Group and the electrodes were dried inside an oven in a glovebox at 100 °C for 10 min. A microscopic camera was used to photograph the electrodes along with a calibration slide. These pellets were loaded into a specialized in situ cell (the St Andrews in situ cell) at ISIS inside a glovebox and then placed inside the RAL4 furnace loaded on Polaris28. All diffraction data structures were refined using GSAS41.CaD2 was purchased from Fluka. A 25 mm pellet of CaH2 was prepared in identical manner as the NPD experiment (with the substitution of H2 for D2) except the pellet was extremely thin to allow for the 90% transmission rule . The St Andrews in situ rig was used again but this time in conjunction with the RAL3 furnace and the Osiris spectrometer at ISIS37. Simultaneous EIS, diffraction, and QENS data were collected. The sample was oriented at a 45° angle to the incoming neutron beam to maximize the exposed sample area. The sample was heated to 600 °C under flowing Ar and H2 (95 and 5 cm−3 min−1 respectively). A Solartron 1280Z was used to collect EIS data with a 40 mV perturbation voltage between 20 kHz and 1 Hz.CaHα-Ca2NH and β- were synthesized at St Andrews at 600 °C with exposure to 5/5/90 cm−3 min−1 N2/ H2/ Ar for the nitride and hydride precursors. These materials were loaded into air-tight sample holders inside a glovebox and then transported to the NMR and Raman facilities. The NMR spectra were collected on a 9.4 T spectrometer with a magic angle spin of 14 kHz. The Raman spectra were collected using Renishaw inVia Raman Microscope with a 532 nm laser at 5% power collected over 10 s from 100 cm−1 to 3500 cm−1.Supplementary InformationPeer Review File


Interestingly, only the larger Fen– (n ≥ 15) clusters support the observation of O2-intake, while the smaller clusters Fen– (n = 7-14) are nearly nonreactive. What is more interesting is that Fe17O10– shows up with prominent abundance in the mass spectra indicative of its distinct inertness. In combination with DFT calculations, we unveil the stability of Fe17O10– within an interesting acordion-like structure and elucidate the spin accommodation in such a strongly ferromagnetic iron cluster oxide.Isolated clusters are ideal systems for tailoring molecule-based magnets and investigating the evolution of magnetic order from microscopic to macroscopic regime. We have prepared pure Fe n- (n = 7-31) and observe the stable, strongly ferromagnetic iron oxide cluster Fe17O10- featuring an accordion-like structure.Iron oxide nanoclusters have high potential for high-density energy storage and spintronics microdevices, but systems that are both stable and display strong ferromagnetism remain challenging to obtain. Here, the authors study the gas phase collision reactions of the iron clusters Fe Sub-nano clusters of iron oxides could exhibit altered properties in comparison with larger nanoparticles, enabling highly tuneable magnetism and chemical activity, and thus are ideal candidates for high-density storage or spintronics microdevices when embedded in semiconductors. On the premise and aim of a full understanding of the spin-exchange interactions and structure-property relationship8, ongoing continuous efforts have also been devoted to understanding the gas-phase reactivity of such clusters9, as well as fundamental mechanisms and potential applications relating to metalloenzymes and metallic corrosion and catalysis11.Magnetic iron oxide nanoparticles are well-known for their significance in targeted drug delivery13. Typically, FeO is antiferromagnetic at low temperatures but with a high-spin 5Δ ground state based on the DFT-GGA calculations14; Fe2O3 is antiferromagnetic but Fe3O4 is known to be ferrimagnetic. Meanwhile, oxo, peroxo, and superoxo isomers15, dioxygen, and oxygen complexes Fe(O2)n (refs. 18) and oxygen-rich clusters22, as well as sequential stoichiometries of (FeO)m (ref. 23), FemOm+1,2 (refs. 28), (Fe2O3)n clusters31 and (Fe2O3)kFeO+ (ref. 32) clusters have been extensively studied35. Studies of such unique stoichiometries not only elucidate varying electronic and geometric structures of iron oxides36, but also provide fundamental information on the kinetics and thermodynamics of coordination and catalytic reactions43. Notably, iron atoms in most of these oxide clusters, except a few (FeO)n clusters23, are separated by oxygen atoms but without direct Fe–Fe bonds, resulting in a tendency of antiferromagnetic rather than ferromagnetic properties.Iron compounds could exhibit a variety of magnetic properties, with appropriate nuclearity and topology to function as molecular magnets45. Reasonable research interest has been devoted to polynuclear clusters50, especially the magic-number iron clusters and iron oxide clusters such as a ring cluster Fe10 (ref. 53), cubic Fe13O8 and cage Fe12O12 clusters61. However, the antiferromagnetic states were often found to be much more stable than the ferromagnetic counterparts62. There seems an incompatible contradiction for iron oxide clusters to bear high stability and strong ferromagnetism.It has been recognized that the magnetic moments of metal clusters depend on geometry and electron localizationn clusters with oxygen were studied. It was found that Co13O8 dominates the mass distribution in sufficient collisional reactions with O2. The distinctive stability of neutral Co13O8 was demonstrated as a class of metalloxocubes63. The prominent mass abundances of Ni13O8+ and Fe13O8+ were also verified in such a flow tube reactor65. Notably, the cubic M13O8+  clusters prefer high-spin states, and the spin multiplicity increases from Ni (3d84s2) to Co (3d74s2) and Fe (3d64s2), shedding light on the importance of spin accommodation. Likely due to the strong magnetic effect, there were challenges to observe and monitor the reactions of anionic iron clusters.Recently utilizing the customized deep-ultraviolet laser ionization mass spectrometry technique, the reactions of neutral Con– clusters, which enables us to meticulously study their reactions and probe the likely magic-number clusters. Here we observe the gas-phase reactions of the iron clusters Fen– (n = 7–31) with oxygen and find a stable, strongly ferromagnetic iron oxide cluster Fe17O10– featuring an accordion-like structure.Until very recently we have made a breakthrough in the preparation of pure Fen– (n = 7-31) clusters in the absence and presence of different doses of oxygen reactants (10% in He), corresponding to the on-time of the pulse valve at 185 μs and 200 μs respectively. These bare Fen– clusters display a nearly Gaussian distribution of their mass abundances centered at n = 16; however, the oxidation products were observed first for the larger Fen– (n = 15–27) clusters, and the formed FenO2– clusters also display a Gaussian distribution centered at n = 20  clusters benefit the O2 adsorption and thermal balance by sufficient collisions with helium; meanwhile, the larger Fen–·O2 clusters could allow for more flexible vibrational structure relaxation to disperse the energy gain of O2-adsorption. In comparison, the smaller Fen– clusters (n ≤ 14) may be subject to a smaller vibrational density-of-states (DOS) and shorter lifetime of the vibrational excitation states thus allowing for fragmentation of the likely formed Fen–·O2 intermediates. With a further increased dose of the O2 reactant  do react with a low rate. Griffin and Armentrout69 reported a study on the reactions of Fen+ (n = 2–18) with O2 by guided ion-beam apparatus and showed that the dioxide cluster cations are dominant products for the larger ones. The size-dependent cross-section for dehydrogenation of ethylene on the cationic Fen+ (n = 2–28) clusters was also studied by Ichihashi et al.70, showing that the dehydrogenation reaction increases rapidly above a certain cluster size. It is inferred that the dramatic size effect of iron clusters is associated with the Fe-Fe bonding and variable 3d electrons and complex spin-exchange interactions.This observation of size dependence is, to some extent, consistent with the previous studies on gas-phase reactivity and interactions of cations and neutral iron clusters. For example, early studies by Whetten et al.17O10–, we have performed a global structure search by USPEX employing the VASP software package. The typical structure isomers and electronic spin-state isomers of the Fe17O10– are shown in Fig. 17O10– (C2v) has 10 μ3-O atoms capped at the 10 hollow sites of Fe17 which, however, exhibits a different geometry compared with the C3v structure of nascent Fe17–  correspond to magic numbers in the transition-metal oxide clusters48.To understand the prominent mass abundance of Fe17– Fig. . The sub17O10– finds an extremely high magnetic moment . Considering that the Fe atom has multiple spin-unpaired electrons and the ground state of O2 is in a triplet state, we have compared the energy of Fe17O10– at the GGA + U method with a variety of magnetic moments being considered (from 4 to 56 μB), as shown in Fig. 17O10– corresponds to the accordion-like C2v structure with a magnetic moment of 56 μB. Notably, the spin magnetic moment of Fe17– cluster is estimated to be 54 μB , indicative of spin-excitation to form the strongly ferromagnetic Fe17O10– . Within this high-spin state, the accordion-like C2v structure allows for a resonating structure isomer of 0.07 eV higher in energy m clusters. When connecting the convex polygon with all the FenOm– clusters being involved, Fe17O10– and Fe13O8– are both located in the lowest position  observed in the mass spectrometry experiments, we further analyzed the reaction processes of Fe17– and Fe13– with oxygen by DFT calculations based on the Gaussian software package. As shown in Fig. 17– in the side-on or end-on orientation, the O2 molecule dissociates spontaneously giving rise to two μ3-O atoms capped on the neighboring hollow sites of the Fe17– cluster. We also tested Fe17O8– + O2  giving rise to likely O–O dissociation at a small energy barrier  clusters, as shown in Fig. 17O10– to Fe17O12– (−2.23 eV) is much smaller than that for each step of Fe17O2– → Fe17O4– → Fe17O6– → Fe17O8– → Fe17O10–, indicating that the formation of Fe17O10– is significantly faster than its subsequent conversion to Fe17O12–. This could be the reason why Fe17O10– instead of Fe13O8– is dominated in the reactions of Fen– clusters with oxygen. In addition, the DFT calculation results show that Fe11– also allows for oxygen addition and O–O dissociation  pertaining to strong ferromagnetism; in comparison, the cubic Fe13O8– optimized by the same GGA + U approach finds a ground state of ferrimagnetic characteristics (10 μB). To fully understand the novelty behind the two clusters, Fig. 73. It is shown that the NPA charge distributions display similar negative values on all the μ3-O atoms  for the DOSs of Fe17O10– dominate in the energy range of −2 ~ −7.5 eV, while the dominant minority-spin bands (beta electrons) in the range of 0 ~ 5.0 eV. The strong spin polarization in Fe17O10– in contrast to Fe13O8– is also depicted by electron localization function (ELF)74 analysis, as shown in Fig. Also, the NICS values on the quasi-square surfaces of the two clusters are comparable with each other Fig.  vs. 5h; ces Fig.  vs. 5e. ity Fig. ; howeverter Fig. , which a17O10–. Global-minima structure search by USPEX interfaced with the VASP software package has determined the ground state of Fe17O10– corresponding to a high-spin ferromagnetic accordion-like structure. DFT calculations of thermodynamics and dynamics rationalized its reasonable stability and the feasible reaction dynamics for Fe17– to react and form dissociative oxygen addition successfully. The observation of an absence of Fe13O8– is consistent with the fact that small Fen– (n = 7–14) clusters are less reactive with oxygen to form stable oxides. We highlight the strong ferromagnetism of Fe17O10– for sub-nano materials with well-defined components and geometric/electronic structure. Such a strongly ferromagnetic stable cluster could be an ideal candidate for high-density storage or spintronics microdevices.In summary, utilizing a customized mass spectrometer we studied the reactions of iron clusters with oxygen and observed the prominent inertness of FeL = 60 mm), and a home-built laser vaporization (LaVa) source. In this experiment, the Fen– clusters were prepared by ablating a clean and rotating iron disk (99.95%) with a pulsed 532 nm laser (Nd:YAG) at a repetition rate of 10 Hz in the LaVa source. In a nozzle (L = 35 mm), the clusters were cooled down during a supersonic expansion process by the pulsed He carrier gas  which was controlled by a pulsed valve . For reactions between the Fen– clusters and O2, the reactant gas of 10% O2 in He (1.0 atm) was injected into the flow tube reactor (at room temperature) and the pressure of the carrier gas inside this tube was kept at ~35 Pa. More detailed descriptions can be found elsewhere , it consists of a customized vacuum system, a mini flow tube reactor , Fe2O3–, Fe13O8–, and Fe17O10– clusters and all possible spin states are optimized by applying the Gaussian 09 software package, with the Becke’s exchange and Perdew−Wang’s correlation functionals (BPW91) and the 6-311 G(d) basis set for all elements . Vibrational frequency calculations and zero-point vibrational corrections were performed for all of the optimization and energy calculations. The transition states (TS) were checked to ensure that there is only one imaginary frequency, and the intrinsic reaction coordinate (IRC) scan was employed to ensure a connection with both intermediates in the reaction pathways. Gibbs free energy were calculated at a temperature of 298 K. The structures and NPA population are plotted by the software package of VESTA and visual molecular dynamics (VMD). More detailed descriptions of the computational methods can be found elsewhere (Supplementary SThe ab initio evolutionary algorithm USPEXm– n = 3-0, m ≤ n,Peer Review FileSupporting Information