However, BaTaO2N still exhibits poor water‐splitting performance that is susceptible to its synthetic history, surface states, recombination process, and instability. This review provides a comprehensive summary of previous progress, current advances, existing challenges, and future perspectives of BaTaO2N for solar water splitting. A particular emphasis is given to highlighting the principles of photoelectrochemical (PEC) water splitting, classic and emerging photocatalysts for oxygen evolution reactions, and the crystal and electronic structures, dielectric, ferroelectric, and piezoelectric properties, synthesis routes, and thin‐film fabrication of BaTaO2N. Various strategies to achieve enhanced water‐splitting performance of BaTaO2N, such as reducing the surface and bulk defect density, engineering the crystal facets, tailoring the particle morphology, size, and porosity, cation doping, creating the solid solutions, forming the heterostructures and heterojunctions, designing the photoelectrochemical cells, and loading suitable cocatalysts are discussed. Also, the avenues for further investigation and the prospects of using BaTaO2N in solar water splitting are presented.Barium tantalum oxynitride (BaTaO 2N for solar water splitting. For enhancing the water‐splitting performance of BaTaO2N, various strategies, such as reducing the defect density, engineering the crystal facets, cation doping, forming the heterostructures/heterojunctions, designing the photoelectrochemical cells, and loading suitable cocatalysts are presented.This review provides a comprehensive summary of previous progress, current advances, existing challenges, and future perspectives of BaTaO However, the current global hydrogen production still heavily relies on fossil fuels (>95%), which is the cheapest option in most parts of the world. Solar‐driven water splitting is one of the most environmentally benign chemical processes to convert abundant solar energy into storable and transportable green hydrogen. Solar‐driven water splitting proceeds over cocatalyst‐assisted semiconductor according to three consecutive steps 1 and Hydrogen evolution reaction (HER):Oxygen evolution reaction (OER):2 and O2, several key criteria must be met simultaneously: i) sufficient voltage must be generated upon irradiation to split water, ii) the bulk band gap must be small enough to absorb a significant portion of the solar spectrum, iii) for the unbiased operation of a PEC cell, the band edge potentials at the surfaces must straddle the hydrogen and oxygen redox potentials (the conduction band minimum must be more negative than E0(H+/H2) = 0 V versus RHE at pH 0 and the valence band maximum must be more positive than E0(O2/H2O) = 1.23 V versus RHE at pH 0), iv) the system must be stable for a long period of reaction time, v) charge transfer from the semiconductor surface to the electrolyte must be facile to minimize energy losses, and vi) low‐cost and earth‐abundant elements must be used.Particularly, photoelectrochemical (PEC) water splitting is a powerful yet complex process, where the redox potentials for the decomposition of water determine the required band gap energy and band‐edge potentials for semiconductors to be used as photoanodes and photocathodes. However, to efficiently and sustainably split water into H The U.S. Department of Energy estimated the cost of green hydrogen produced by the PEC process to be US$5.7 kg−1 in 2020 (with a 20% STH efficiency) and to lower to US$2.1 kg−1 (with a 25% STH efficiency) in the more distant future. It has also been suggested that the STH efficiency of 25% and the photoelectrode lifetime of 10 years are required for the PEC systems to be economically consistent with fossil fuel‐based hydrogen production processes. Recently, monolithic systems and integrated PEC modules under light concentration reached STH efficiencies as high as 17.12% and 19%. However, the high capital and operational costs significantly hamper their economic viability. To date, no cost‐effective and highly efficient PEC systems have been developed to satisfy all those key criteria despite significant progress in this field. Thus, further studies are necessary to discover novel materials with unprecedented physicochemical and optoelectronic properties that can simultaneously meet those key criteria set for the design and application of the PEC system for green hydrogen production.In addition to these important requirements, the photoelectrode materials must exhibit at least >10% solar‐to‐hydrogen (STH) conversion efficiency for their commercial viability.11 The: i) water dissociation and formation of surface‐bonded OHads, ii) the further oxidation of OHads to Oads, and iii) the formation of OOHads, which is a precursor for O2 evolution:As one of the two key reactions of overall water splitting, the water oxidation reaction proceeds via the following general steps Figure 16]:16:P21/n is the most stable polymorph of WO3 at room temperature and has a perovskite‐like structure without an A‐site ion. Theoretical studies on the correlation between the crystal structure and the band gap of WO3 revealed that the bandgap energy values of different WO3 polymorphs decrease in the following order: monoclinic > orthorhombic > triclinic > tetragonal ≥ cubic. Also, the structural modification by introducing the nitride ions into the lattice of the monoclinic WO3 phase led to the reduction of the bandgap energy from 2.6 eV to 1.9 eV. Highly transparent nanoporous WO3 films fabricated by layer‐by‐layer deposition of a colloidal solution of tungstic acid and annealing exhibited a maximum IPCE of 75% and a photocurrent density of 2.4 mA cm−2 at 1.23 V versus RHE under simulated solar irradiation  offers a smaller bandgap energy and limits the formation of soluble tungstates due to the strong covalency in the metal oxo bonds, yielding higher photocurrent in neutral pH under visible light irradiation. The CuWO4 photoanodes exhibited photocurrent densities of ≈0.15–0.16 mA cm−2 and 0.62 mA cm−2 at the thermodynamic potential for water oxidation (1.23 V versus RHE) under simulated solar irradiation  has been intensively explored as a promising photoanode material for PEC water splitting due to its small bandgap energy, suitable valence band edge potential to thermodynamically drive the water oxidation reaction, excellent stability under alkaline conditions, earth abundance, and environmentally friendliness. With the reported bandgap energy of 2.0‐2.2 eV, α‐Fe2O3 has the potential to reach a theoretical maximum STH efficiency of 16.8% and a photocurrent density of 12.6 mA cm−2, which manifestly exceeds the STH benchmark efficiency of 10% required for practical applications. The α‐Fe2O3 nanowires‐based photoelectrode, which was fabricated by a chemical bath deposition method followed by hydrogen treatment and then loaded with ultrathin TiO2 overlayer and CoPi cocatalyst, exhibited a stable photocurrent density of ≈ 6 mA cm−2 at 1.23 V versus RHE over 100 h under AM 1.5G simulated sunlight and an IPCE value of ≈18% at 420 nm owing to its excellent light absorption and bulk charge separation capacity, the enhanced electrical conductivity, and the decreased surface recombination. A hematite nanosheets‐based electrode modified by Ag and CoPi nanoparticles showed an IPCE value of ≈40% at 420 nm and an STH efficiency of ≈0.55% due to the improved light harvesting and the facilitated charge transfer by Ag nanoparticles and the reduction of surface/bulk recombination and the stabilization of the photoelectrode surface by CoPi cocatalyst , short hole‐diffusion length (2‐4 nm), low charge carrier mobility (≈10–2 to ≈10–1 cm2 V−1 s−1), poor electric conductivity, and poor OER kinetics still hinder achieving the practical maximum STH efficiency of α‐Fe2O3.α‐Fet Figure .47]472O3 2O4 (space group Fd3¯m) also received much interest due to its narrow bandgap energy (Eg = 1.9 eV), high energy level conduction band minimum, outstanding photochemical stability, low cost, and magnetic recyclability. Interestingly, ZnFe2O4 with a relatively poor crystallinity but a higher spinel inversion degree (due to cation disorder) shows a superior efficiency in photo‐excited charge carrier separation and an improved majority charge carrier transport compared to ZnFe2O4 with higher crystallinity but a lower inversion degree. The optimization of these factors and the addition of a nickel‐iron cocatalyst overlayer resulted in a new benchmark photocurrent density of 1.0 mA cm−2 at 1.23 V versus RHE, and an IPCE value of about 25% was reached for ZnFe2O4 nanorod photoanode  is an n‐type semiconductor with a direct bandgap energy of 2.4 eV and a valence band edge potential of ≈2.4 V versus RHE, which is sufficiently positive than E0(O2/H2O) = 1.23 V versus RHE. Therefore, the monoclinic scheelite BiVO4 shows the highest photocatalytic water oxidation activity among the three polymorphs. A photocurrent density of 6.72 mA cm−2 at 1.23 V versus RHE and an IPCE of ≈90% were achieved by the combination of BiVO4 with more conductive WO3 nanorods in the form of core‐shell heterojunction  of 1.75% was found at a potential as low as 0.6 V versus RHE for nanoporous BiVO4 photoanodes. However, the water oxidation efficiency of BiVO4 is still limited by a high electron‐hole recombination rate, poor charge transport properties, a low carrier collection efficiency, inadequate water oxidation kinetics, and the bandgap energy limiting an AM 1.5G solar photocurrent density to 7.4 mA cm−2 and an STH efficiency to 9.1%.BiVOn Figure .56]564 cr2, WO3, CuWO4, Fe2O3, ZnFe2O4, and BiVO4 alone are unlikely to satisfy the criteria set for ensuring the practical application in solar to chemical energy conversion. This paved the way for the development of novel materials, including mixed‐anion compounds, for PEC applications.Despite their encouraging and relatively outstanding performance achieved so far, TiO3 Unlike single‐anion compounds, mixed‐anion compounds exhibit diverse structures, chemical and physical properties, and new functionalities because of different anionic characteristics, such as ionic radii, valency, electronegativity, and polarizability. Also, the combination of hetero‐anions can realize the structures of compounds that cannot be generally stabilized by homo‐anions. Especially, to develop photocatalytic materials that can efficiently function under visible light, anions less electronegative than oxygen can be simultaneously introduced. Having similar chemical, structural, and electronic properties, oxygen and nitrogen substitute each other in the anion site to form oxynitrides. In oxynitrides, the nitride anions (N3−) having atomic orbitals with potential energy higher than O 2p atomic orbitals of the oxide anions (O2−) shift the valence band maximum upward without affecting the conduction band minimum, leading to the increased covalency of metal‐anion bonds, the improved absorption property, and the decreased optical band gap 3 perovskites, BaTaO2N is regarded as one of the promising photocatalysts for solar water splitting due to its absorption of visible light up to 660 nm, small bandgap energy (Eg = 1.9 eV), good stability under light irradiation in concentrated alkaline solutions, and nontoxicity. Moreover, the conduction band minimum and valence band maximum of BaTaO2N are located at –0.4 V and 1.5 V versus NHE at pH 0, respectively, which should theoretically drive the water‐splitting reaction in the absence of an external bias. BaTaO2N can generate a photocurrent density of about 18 mA·cm−2 under AM 1.5G simulated sunlight based on the assumption of an IPCE efficiency of 100% at < 660 nm. Incidentally, the STH conversion efficiency of BaTaO2N is about 24%. Its large dielectric constant can promote a facile separation of photo‐excited charge carriers.As a typical representative of the 2N phase was first synthesized in 1986 by Marchand et al., and its centrosymmetric cubic crystal structure with space group Pm3¯m (no. 221) and a random O/N distribution were determined by a powder time‐of‐flight neutron diffraction method , Ta is at the cube center , and three anions (one nitrogen and two oxygen) are randomly distributed at the face centers. The Ba atom is 12‐fold surrounded by anions with a Ba– distance of 2.908 Å, and the Ta atom is octahedrally coordinated  with a Ta– distance of 2.056 Å. Although the conventional diffraction analyses using neutron and X‐ray radiations similarly confirmed an average cubic perovskite structure with Pm3¯m symmetry, the extended X‐ray absorption fine structure (EXAFS), zone‐axis electron diffraction (ED), pair‐distribution‐function (PDF), solid‐state magic‐angle spinning (MAS) NMR spectroscopy, constant‐wavelength neutron diffraction, and X‐ray absorption near edge structure (XANES) spectroscopy analyses revealed short‐range structural distortions around the Ta atoms and widely distributed Ta–O/N bond distances in BaTaO2N. First‐principles investigation conducted by Xu and Jiang found the existence of both short‐range order in the intra‐octahedron O/N occupation and randomness of inter‐octahedron O/N occupation, which is consistent with the cubic symmetry. Their study also revealed that the electronic properties of BaTaO2N do not strongly depend on the O/N configuration, whereas the dielectric properties indicate a stronger dependence on the O/N configuration. Since local structure relaxation is expected due to the different ionic radii and valency of the oxygen and nitrogen ions, the local symmetry of BaTaO2N was, however, not well defined. The DFT analyses suggested that the local symmetry in the crystal is lower, resulting in the different Ta–O and Ta–N bond lengths, and it predicted the possibility of either orthorhombic, tetragonal, or monoclinic symmetry without contradicting the macroscopic cubic description of the crystal structure of BaTaO2N with space group Pm3¯m and found BaTaO2N to be most stable in the Pmc21 (no. 26) structure type with an ordered anionic sublattice  analysis of the total neutron scattering reveals that although a long‐range anion order is not present in BaTaO2N, the TaO4N2 octahedra predominantly adopt a cis‐configuration with small Ta displacements toward the N atoms on a local scale. The cis‐chains propagate in all three dimensions in BaTaO2N, retaining the average cubic crystal symmetry. Nonetheless, the partial formation of a metastable trans‐type configuration by applying strain engineering can increase the N occupancy at axial sites. By applying an automated bonding analysis with Crystal Orbital Hamilton Populations , George et al. correlated the total energy of the systems with the strongest covalent interaction (Ta–N bond) of BaTaO2N and observed anti‐correlation with the bond energies of Ta–O, implying the importance of covalency of the Ta–N interactions in the anionic ordering of BaTaO2N. It was also found that the structure of BaTaO2N with a cis‐type local symmetry is much more stable than the corresponding trans‐type local symmetry.The BaTaOy Figure .2− anions, which are different from asymmetric cyanamide anions ([N = C (three bonds N]2−). These compounds can act as fluxes for the synthesis of perovskite oxynitrides. Kikkawa and co‐workers synthesized small cubic crystals (3.1 µm) of perovskite BaTaO2N via the formation of a Ruddlesden−Popper‐type oxynitride from the reaction between BaTaO2N powder and molten BaCN2 flux. On the other hand, BaCN2 was also used as a starting material along with Ta2O5 to synthesize highly dispersed BaTaO2N particles with a size of 100–850 nm at ≈900 °C for 5 h because of the partial dissolution of BaTaO2N in the BaCN2 melt. Further, to prevent the thermal decomposition and the partial nitrogen loss, spark plasma sintering with (100 MPa) and without pressure was applied to rapidly synthesize BaTaO2N powder for 3 and 10 min using BaCN2 and urea as additive and nitrogen sources, respectively. The BaTaO2N grains were bonded through dissolution and precipitation at their surfaces in the BaCN2 melt, and a minor Ba‐rich compound, possibly Ruddlesden‐Popper‐type Ba2TaO3N, was identified at the boundaries of the BaTaO2N grains. BaTaO2N synthesized by a pressureless spark plasma sintering technique exhibited a very high room‐temperature relative permittivity up to 9550 with a dielectric loss down to 0.001 at 100 Hz due to its O/N ordering in the cis configuration and high purity (97.8%). Gomathi et al. also employed urea as a nitrogen source to synthesize BaTaO2N nanoparticles with a size of about 60 nm at 1226 K for 2 h. NH3 is yielded upon urea decomposition at 523 K and can react with starting materials to form oxynitrides. By varying the amount of urea, the nitrogen stoichiometry can, therefore, be controlled since urea gives a nitrogen content close to the theoretical value. However, urea generally decomposes at a relatively low temperature that is far below the temperature required for the formation of perovskite oxynitrides and the ammonolysis of their oxide precursors. Therefore, it is still challenging to precisely control the synthesis process, purity, and stoichiometry of BaTaO2N. The presented synthesis approaches generally result in BaTaO2N powders, which need to be further deposited on substrates for PEC OER. The PEC‐OER performance depends on the thickness, morphology, crystallinity, grain boundaries, etc. of the BaTaO2N films.Alkaline earth metal cyanamides are compounds containing alkaline earth metals and nitrogen together as their constituents like oxynitride perovskites. They are expected to melt at temperatures lower than the temperature for the nitrogen release from perovskite oxynitrides. Since the partial decomposition of SrTaOcrystals.142 Ho62N was grown on a 100‐cut SrTiO3 single‐crystal substrate with a conducting buffer layer of SrRuO3 by a pulsed laser deposition technique at 760 °C in a mixed gas atmosphere of 100 mTorr N2/O2. The dielectric permittivity of the fabricated thin film was much different from the value reported for incompletely sintered ceramics of BaTaO2N with considerable porosity (45%). It was found that the epitaxial strain effect can cause a tetragonal distortion of the unit cell of perovskite BaTaO2N with a negligible volume change. The nanostructured thin films of BaTaO2N were formed on Ta substrates by the ammonolysis of hydrothermally grown Ba5Ta4O15 nanosheet layers at 1000 °C for 2 h. The fabricated BaTaO2N thin film deposited with a CoPi layer exhibited a photocurrent density of ≈0.75 mA cm−2 at 1.23 V versus RHE and generated O2 for 5 h with a Faradaic efficiency of > 90% without significant deactivation under AM 1.5G simulated sunlight due to the high crystallinity of the formed film  sputtering and heat treatment at 1173–1273 K for 0.5 to 1 h under N2 atmosphere. CoOx‐deposited BaTaO2N/Ta2N/Ta electrodes produced a photocurrent density of 4.6 mA cm−2 at 1.23 V versus RHE and exhibited a 9% IPCE at 600 nm during water oxidation under AM 1.5G simulated sunlight due to the remarkable conductivity of Ta2N, promoting an efficient transfer of electrons between BaTaO2N and Ti layers, and both H2 and O2 were produced with a Faradaic efficiency of almost 100%  method and electrophoretic deposition, respectively. Interestingly, they found that the particle‐based photoelectrodes could exhibit a higher photocurrent density due to the improved absorption properties and a larger electrochemical surface area  enables the probing of the k‐resolved electronic structure in the subsurface region despite atmospheric surface contamination and provides the opportunity to explore materials, surfaces, and interfaces as well as their reactivity and evolution. The anion arrangement in oxynitrides significantly influences their optical and electronic properties. However, it is difficult to assess anion arrangement in thin films. Yamamoto et al. demonstrated inverse photoelectron holography, which is an atomic‐resolution holography technique, to directly measure the local structure around atoms of light elements. The obtained holograms of anions in the SrTaO2N thin films showed a remarkable difference between O and N holograms, revealing a trans‐type anion ordering. The simulated hologram agrees well with the experimentally obtained data. To achieve high efficiency in BaTaO2N film, it is necessary to select an appropriate deposition technique and optimize the deposition parameters.Lippert and co‐workers151 coa Figure . The thi72N. Previously, the particle morphology, size, and porosity of BaTaO2N were modulated by the ammonolysis of the Ba5Ta4O15 precursor, synthesized by a flux method using BaCl2, KCl, RbCl, CsCl, KCl+BaCl2, and K2SO4 with different solute concentrations, with and without KCl flux. The flux method enabled to control the morphology, size, and dimension of Ba5Ta4O15 crystals. It was found that the ammonolysis of Ba5Ta4O15 with the KCl flux could lead to the formation of the dense crystals of BaTaO2N, whereas the ammonolysis of Ba5Ta4O15 without the KCl flux resulted in the porous structures of BaTaO2N  exhibited the highest photocurrent density of ≈3.11 mA cm−2 at 1.2 V versus RHE due to their average size of 1–2 µm and a porous surface. A further increase and a decrease in the particle size led to a reduction in the photocurrent density due to the existence of a long transfer distance to reach the surface for photo‐excited charge carriers and the surface degradation by excess high‐temperature nitridation, increasing the bulk and surface recombination rates, respectively. Within the first 3 hours of the photocatalytic reaction, BaTaO2N obtained by a flux‐inclusive ammonolysis of Ba5Ta4O15 crystals grown using BaCl2 at a 10 mol% solute concentration (sample B10K) showed a higher O2 evolution rate of 134.64 µmol h−1 (404 µmol) because of its larger particle size, predominantly exposed surfaces with a similar nature, high crystallinity, reduced surface defect density, and fewer grain boundaries  decreased within the first fragmentation passages due mainly to the presence of grain boundaries in the fragmented BaTaO2N particles, which became recombination hubs for photo‐excited electrons and holes. Also, the laser fragmentation resulted in the loss of N content and reduced crystallinity of BaTaO2N, which significantly affected the optical properties and photoelectrochemical performance. In general, the surface area becomes greater when the particle size is minimized. The reduced particle size may give a quantum size effect, resulting in a broader band gap. Therefore, a high degree of crystallinity is often preferred over a large surface area for uphill water splitting reaction because of the recombination of photo‐excited charge carriers in the defects.According to Anbalagan and Thomas,154 the Figure . Despite2N crystals with well‐developed {111} facets and co‐exposed {100} and {110} facets could exhibit a significantly enhanced photocatalytic activity for H2 evolution in comparison to the BaTaO2N crystals with only {100} facets. Hence, a number of theoretical studies have been conducted to understand the impact of the surface local structures on various properties of BaTaO2N. The influence of different surface terminations on the electronic, optical, and photocatalytic properties of trans‐ and cis‐BaTaO2N was studied using DFT calculations by Zhou et al. The calculated work functions showed that the Ba‐terminated surfaces have smaller work functions than the Ta‐terminated ones. This indicates that the former is more favorable to promoting the separation of photo‐excited charge carriers. The hydrogen evolution reaction proceeds more easily on the surfaces terminated with Ta, O, and N atoms, while the oxygen evolution reaction proceeds more readily on the surfaces terminated with N atoms than on the surfaces terminated with O atoms. In another work, they also found that the water dissociation prefers to occur on the BaNbO2N surface terminated with N atoms than on the surface terminated with O atoms. Dissociative water adsorption on surfaces with the (100) and (001) terminations is thermodynamically favorable. Therefore, the exposed O atoms (cis‐(100)‐Nb2O3N, trans‐(100)‐NbO2, and trans‐(001)‐BaO) are the most favorable sites for HER, while cis‐(001)‐BaO is the most suitable for OER.The water‐splitting performance of photocatalysts is sensitive to the surface local structures. Recent studies on the effect of different exposed surfaces found that the BaTaO1} facets134 an2 evolution and HO*, O*, and HOO* for O2 evolution) on the predominant (110) surface of BaTaO2N (BTON) terminated with TaO6, TaN6, and TaO4N2 octahedra /(rBTON)] versus ‐(EXBTON−EBTON)) plot. The experimental photocatalytic reaction rates were satisfactorily described using the adsorption energies of intermediates (H* for HER and HO* and O* for OER) estimated by MD calculations, rationalizing the effects of aliovalent cation doping  and surface chemical termination  of BaTaO2N. A satisfactory match for all dopants was obtained for the BaTaO2N surface terminated with the TaO6 octahedra. Nevertheless, the BaTaO2N:Zr surface terminated with the TaO4N2 octahedra and the BaTaO2N:Mg surface terminated with the TaN6 and TaO4N2 octahedra showed trend deviations possibly due to simplifications in the MD simulation.Molecular dynamics (MD) simulation was involved to gain insights into the respective effect of doping with various aliovalent cations on the adsorption energy of water molecules and formed intermediates  facet TaO2N‐terminated surface under (photo)electrochemical conditions  or acceptors  and change the carrier concentration.Doping is one of the effective strategies to modulate the electronic structure, optical properties, electrical conductivity, charge density, charge mobility, and charge separation and transfer, leading to enhanced water‐splitting efficiency.157 Ea8.1A‐site of perovskite BaTaO2N, monovalent alkali‐metal cations are prone to sublimation during high‐temperature synthesis and ammonolysis under an NH3 flow. Nevertheless, small amounts of lithium, sodium, and potassium were found in BaTaO2N synthesized by the ammonolysis of LiBa4Ta3O12 and a flux method using the KCl and RbCl fluxes, respectively. In the time course of the Z‐scheme overall water splitting reaction over BaTaO2N synthesized by the ammonolysis of LiBa4Ta3O12, the evolution rates of H2 and O2 in the first hour were 3.1 and 1.55 µmol, respectively, confirming the overall water splitting with the stoichiometric H2/O2 ratio of 2:1 due to the efficient transfer of photo‐excited charge carriers to the surface  under visible light irradiation, resulting in an apparent quantum yield for the H2 evolution reaction of ≈1% at 420 nm, and during Z‐scheme overall water splitting, a mixture of Na‐containing Pt/BaTaO2N and H+–Cs+‐modified PtOx/WO3 evolved stoichiometric amounts of H2 and O2 from an aqueous NaI solution under visible light up to 12 h (Zn1/4Ta3/4)O3 improved the OER activity of BaTaO2N (699 µmol h−1), with an apparent quantum yield of 11.9% at 420 nm  in BaTaO2N by  or  modified its electronic structure. The compositional variation (BaLi0.2Ta0.8O2.8N0.2 and BaNa0.2Ta0.8O2.8N0.2) could control the lattice iconicity, optical band gap, and color, which also play an important role in solar water splitting  of BaTaO2N crystals grown by a flux method. The Mg doping along with Cr2O3/(Na)Rh and IrO2 co‐catalyst loading led to the one‐step excitation overall water splitting with an apparent quantum yield of 0.08% at 420 nm and an STH conversion efficiency of 4 × 10−4% for BaTaO2N synthesized by an RbCl flux‐assisted ammonolysis  doping and strain engineering on the OER for differently terminated surfaces of BaTaO2N. The smallest theoretical overpotential (0.53 V) was obtained for Ca2+‐doped TaON‐terminated BaTaO2N (001) clean surface, with 4% tensile uniaxial strain in the [010] direction  by Johnston et al. The Ba1–xSrxTaO2N with cubic Pm3¯m structure at x = 0–0.2 is consistent with a 3D distribution of disordered cis‐chains, while neutron occupancies reveal that a long‐range 2D confinement of anion chain layers is present across the tetragonal P4/mmm symmetry at x = 0.4–1. The local structures of ATaO2N  were investigated by Page and co‐workers using a combination of experimental and theoretical approaches, including neutron total scattering, density functional theory (DFT), and ab initio molecular dynamics (AIMD) simulations. Their study shows that the local cis ordering and Ta off‐centering can play decreasing roles in overall lattice stability, overshadowed by the stabilizing effects of octahedral tilting, by replacing Ba2+ with Sr2+ or Ca2+ in BaTaO2N. The findings indicate that the anion order may have a larger influence on local dipoles and dipole ordering in such perovskite systems. Zn2+ with an effective ionic radius of 0.68 Å, which is close to that of Ta (0.64 Å), was substituted in the crystal lattice site of Ta in BaTaO2N by the synthesis and ammonolysis of the Ba(Zn1/3Ta2/3)O3 precursor. The reduction and sublimation of Zn2+ during the ammonolysis process and the effective charge balance by Zn2+ resulted in an extremely low concentration of Ta4+‐associated defects, efficient charge separation, and improved photocatalytic HER rate of about 65 µmol within 5 h  doping was found to be useful for shifting the respective conduction band potential (ECB) of BaTaO2N to more negative values. The photocatalytic HER rate increased significantly from 6.59 µmol H2 h−1 (pristine BaTaO2N) to 21.92, 10.27, and 12.12 µmol H2 h−1 by doping 5% Al3+, Ga3+, and Sc3+, respectively, revealing the advantageous effect of trivalent dopants. In contrast, the photocatalytic OER rate decreased from 316.3 µmol (pristine BaTaO2N) to 252.5 and 188.0 µmol when 5% Ga3+ and Al3+ were substituted in BaTaO2N, whereas a 5% Sc3+ dopant increased the photocatalytic OER rate within 5 h. Interestingly, 1% La3+ doping resulted in the HER rate that is three times lower than that of pristine BaTaO2N.Trivalent cation O6 (space group P21/n). The impedance spectroscopy analysis revealed an interesting capacitor geometry in BaSc0.05Ta0.95O2.1N0.9 in which the semiconducting oxynitride grains are separated by the insulating secondary phases. Although the Rietveld refinement implies the presence of secondary phases like BaSc2O4, Sc2O3, or ScN, no impurity phases were detected in the synchrotron XRPD measurements. The diffusion of Sc‐rich phase out of bulk to form amorphous shells surrounding crystalline BaTaO2N grains was proposed as a plausible model for such behavior.Based on synchrotron X‐ray powder diffraction analyses, Kim and Woodward173 fo3+–Mg2+ dual substitution (5%) was applied to engineer structural defects and to modulate optoelectronic and surface properties and photocatalytic activity of BaTaO2N. The optical absorption edge of BaTaO2N shifted to shorter wavelengths after (co)substitution of Al3+ and/or Mg2+ for Ta5+, leading to the increase in the optical bandgap energy. This effect was more pronounced in the BaTaO2N samples with higher levels of Mg2+ dopant because a large number of O2− were substituted for N3− to compensate charge balance. The initial photocatalytic reaction rates for O2 and H2 evolution confirmed the enhancement of the photocatalytic performance of BaTaO2N due to the dual substitution of Al3+ and Mg2+ for Ta5+  with an apparent quantum efficiency of 0.18% at 420 nm, while BaTaO2N modified with 3.47% Al3+ + 1.52% Mg2+ produced the highest quantity of H2 (18.94 µmol h−1) with an apparent quantum efficiency of 0.64% at 420 nm. The enhancement achieved by partial Al3+–Mg2+ dual substitution is related to the changes in the defect density, dynamics of charge carriers, electronic band structure, improvement in water and methanol adsorption, and a favorable shift in the band energy levels with respect to water reduction and oxidation potentials.Partial Al BaTaO2N.174 Th8.44+ and Ti4+) doping lowered the photoelectrochemical performance of BaTaO2N photoanodes due to the decreased donor density and reduced titanium species (Ti3+), which generated anion defects simultaneously, increasing the donor density but facilitated the recombination rate between photo‐excited electrons and holes through the redox cycle. The stabilization of the +4 oxidation state of titanium can be achieved by the inductive effect of doped rare‐earth elements that share some electrons with the closest TM–O,N bond to enhance its covalency. In contrast, 5% Zr‐doped BaTaO2N showed a high photocatalytic activity in both O2 (446.8 µmol) and H2 (80.4 µmol) evolution half‐reactions within 5 h due to the altered potentials of the valence and conduction bands by Zr4+ doping. Introducing Zr4+ did not change the ECB value compared to that of pristine BaTaO2N presumably due to the compensation given by the change of valence band (EVB) to more positive values. This is possible if it is considered that the incorporation of a dopant with a larger ionic radius (Zr4+) leads to a greater substitution of O2− for N3− to compensate ionic charge imbalance, justifying the change in the EVB in the Zr‐doped BaTaO2N. Recently, Li et al. found that the incorporation of 1% Zr in BaTaO2N could enhance the photocatalytic H2 evolution activity to the greatest extent (≈500 µmol h−1) because of the extension of the lifetime of electrons and the promotion of electron injection into the Na‐Pt cocatalyst, while the higher doping levels reduced the activity. The Z‐scheme overall water splitting (ZOWS) system based on Cr2O3/Na–Pt/BaTaO2N:Zr0.01 as the HEP, CoOx/Au/BiVO4 as the OEP, and [Fe(CN)6]3−/4− as redox ions evolved H2 and O2 under visible light up to 520 nm, and its STH conversion efficiency reached 0.022% and the apparent quantum yields at 420 nm and 520 nm under monochromatic light were 1.5% and 0.2%, respectively  was found to increase the donor density effectively, enhancing the photoelectrochemical performance. In contrast, the decreased photoelectrochemical performance in W‐doped BaTaO2N was noted due to the facilitated recombination rate through the redox cycle between W4+ and W6+ species.Hexavalent Mo and W species with the [Kr] 4dn Figure .4− was employed as the electron donor. The Z‐scheme system based on Cr2O3/Na–Pt/BaTaO2N:Zr0.01 as the H2‐evoling photocatalyst, CoOx/Au/BiVO4 as the O2‐evolving photocatalyst, and [Fe(CN)6]3−/4− as redox ions evolved H2 and O2 under visible light up to 520 nm, with the STH conversion efficiency of 0.022%. Compared with an individual loading, the combination of pre‐loading of CoOx on the BaTaO2N particles and post‐loading of RhOx was found to be highly efficient in improving both the photocurrent efficiency and stability under visible light irradiation and shifted the onset potential for water oxidation negatively (≈300 mV)  nanoparticles and BaCO3, was achieved for the first time by modification with Rh (or Ru), Cr2O3, and IrO2 as cocatalysts. It was found that the photocatalytic activity of pristine BaTaO2N was influenced by the concentrations of the chosen cocatalysts, and the optimal amounts of Rh (or Ru), Cr2O3, and IrO2 were determined to be 4 wt% (2 wt%), 1 wt%, and 0.3 wt%, respectively. The apparent quantum yield and STH conversion efficiency of Rh/Cr2O3/IrO2‐loaded BaTaO2N reached 0.1% at 400 nm and 5 × 10−4%, respectively. Dual cocatalysts loaded by photodeposition can promote simultaneous oxidation and reduction reactions of photocatalysts. However, it is still challenging to photodeposit oxygen evolution cocatalysts because of weak driving forces for oxidation reaction. Recently, Kobayashi et al. successfully loaded the FeOx cocatalyst onto a Mg‐doped BaTaO2N photocatalyst by an oxidative photodeposition technique, and BaTaO2N coloaded with Pt and FeOx exhibited an apparent quantum yield of 1.2% at 420 nm during the oxygen evolution reaction. It is noteworthy that this photodeposition does not require a further heat treatment, which is advantageous for the design of novel cocatalyst‐loaded photocatalysts that are prone to thermal decomposition at elevated temperature.Cocatalysts are often employed to activate and boost the photocatalytic and photoelectrochemical performance of semiconductor‐based photocatalysts because cocatalysts, their interfaces with semiconductors, dispersion, concentration, and size play an important role in improving the charge separation and transfer, and extending the lifetimes of photo‐excited states.207 PtR Figure . Using ap Figure .[67]E Figure . The Na–N Figure .[211132N for the water‐splitting reaction is presented in the previous sections with respect to its synthesis methods, film fabrication, doping, solid solution, heterostructures, heterojunctions, and cocatalysts. This analysis also includes the effect of morphology, particle size, porosity, and surface chemistry. As a final objective, most of the studies seek to efficiently generate H2 from water splitting using BaTaO2N under simulated AM 1.5G solar illumination. It is emphasized that all these strategies are based on the requirement that the valence band potential of semiconductor must be more positive than that required for the oxidation of water, while the conduction band potential must be more negative than that required for the reduction of water. The generated H2 flux, apparent quantum yield (AQY), and rate constants are often used to compare the photocatalytic performance. Quantitative approaches toward the evaluation of the photocatalytic water splitting were previously recommended by Takanabe and Kisch and Bahnemann.The photocatalytic and photoelectrochemical performance of BaTaO Takanabe93 and2N‐based photocatalysis, the photocatalytic performance has been reported to be closely dependent on the synthesis method, indicating a number of general aspects that allow exploiting the water decomposition reaction using BaTaO2N‐based materials. Some of these general aspects include synthesis methods, anion doping, and cation doping to improve overall photocatalytic performance, combination with other materials for solid‐state Z‐scheme photocatalytic water splitting, and the use of co‐catalysts to promote OER  and HER . Table 2N‐based photocatalysts.In the realm of BaTaO methods,171 an2N is an n‐type semiconductor, thin films of BaTaO2N can be fabricated on conductive substrates and used as photoanodes for photoelectrochemical characterization. Under simulated AM 1.5G sunlight, a photocurrent density of ≈18 mA cm−2 and a solar‐to‐hydrogen (STH) conversion efficiency of about 24% can be theoretically achieved for BaTaO2N. However, the experimentally obtained results are much lower than the predicted values due to unsuitable surface states, photocorrosion, and various types of defects formed during the synthesis process. Bandgap engineering through mono and dual substitution, solid solutions, defect density control, thin film fabrication, and the optimization of exposed surfaces, particle morphology, porosity, and size have been explored to enhance the photoelectrochemical performance of BaTaO2N.The photoelectrochemical studies provide important information to the understanding of the performance of the photocatalysts in the water‐splitting reaction. Since BaTaO8 mA cm−267 and2N to ≈0.03 mA cm−2, >1.2 mA cm−2, 2.05 mA cm−2, 4.2 mA cm−2, ≈4.5 mA cm−2, and 6.5 mA cm−2 at 1.2 V versus RHE  limit the PEC response and can be improved by applying state‐of‐the‐art approaches.The photoelectrochemical studies have enabled the resolution of trends in visible‐light‐assisted water oxidation based on BaTaOhenomena.124 At142N has emerged as a promising photocatalyst for solar water splitting due to its capability to absorb visible light up to 660 nm, suitable band‐edge potentials for overall water splitting in the absence of an external bias, and theoretical solar‐to‐hydrogen energy conversion efficiency of ≈24% under AM 1.5G simulated sunlight. Over the past three decades, the crystal structure, anion ordering, electronic structure, dielectric, ferroelectric, and piezoelectric properties of BaTaO2N were investigated. The apparent quantum yield, photocurrent density, incident photon‐to‐current conversion efficiency, and solar‐to‐hydrogen conversion efficiency of BaTaO2N have been progressively improved. However, its solar‐to‐hydrogen conversion efficiency lags behind the benchmark efficiency required for practical application. This review presented various strategies applied for achieving enhanced solar water‐splitting efficiency. Namely, high‐ and low‐temperature synthesis techniques, advanced synthesis approaches, film fabrication, surface and bulk defect density controlling, crystal facet engineering, and particle morphology, size, and porosity tailoring were overviewed. Also, the impacts of cation doping , creating the solid solutions, forming the heterostructures and heterojunctions, designing the photoelectrochemical cells, and loading suitable cocatalysts were discussed. Other members of the Ba‐Ta‐O‐N system were highlighted for solar water splitting.In this review, the basic principles of photoelectrochemical water splitting and the merits and demerits of oxide‐based photocatalysts researched for photoelectrochemical oxygen evolution reactions were introduced. As one of the most intensively explored representatives of the transition metal (oxy)nitrides family, perovskite BaTaO2N for solar water splitting. These challenges encompass the precise control of various factors, such as homogeneity , crystallinity, particle morphology/size, and defect density. Furthermore, there is a pressing need to identify a suitable deposition technique for BaTaO2N particles, which can preserve their physicochemical and optoelectronic properties. The ability to maintain a long‐term water‐splitting efficiency relies on a comprehensive understanding of surface properties and stability during photocatalytic and photoelectrochemical reactions. Scaling up the photoelectrochemical cells and photocatalytic reactors beyond the laboratory scale is yet another facet of the challenge.Despite significant progress made in the past decades Figure , several2N is known to exhibit defects that have a substantial impact on its properties and water‐splitting performance. Hence, it is imperative to gain a deep understanding of the defect chemistry, including anion and cation vacancies, within BaTaO2N. This is crucial for the development of high‐efficiency BaTaO2N for future solar water‐splitting applications.Similar to other transition metal (oxy)nitrides, perovskite BaTaO2N as a promising material for solar water splitting. Also, in‐depth photoelectrochemical studies can provide valuable information on the interplay between electron transfer dynamics and recombination phenomena. Machine learning and advanced synthesis techniques have enormous potential for the discovery of new photocatalytic materials within the Ba‐Ta‐O‐N system. The optimization of photocatalyst‐electrocatalyst systems is also of utmost importance. Ongoing scientific efforts to integrate transition metal borides, carbides, nitrides, and chalcogenides in water oxidation as electrocatalysts/cocatalysts are very promising.Addressing these challenges with a rigorous approach involves understanding the intricate structure‐property relationships and unraveling the nature of the photocatalyst‐cocatalyst interface at the atomic level. Furthermore, the continued exploration of the underlying kinetics and mechanisms, using operando/in‐situ characterization along with theoretical studies, may pave the way for further development of BaTaO2N‐based photoelectrochemical systems, it is crucial to directly observe Fermi level pinning or unpinning, mitigate photocorrosion and recombination processes, and minimize energy losses during PEC reactor design, construction, and operation.To achieve high solar‐to‐hydrogen (STH) conversion efficiency, the challenges related to the fabrication of semi‐transparent photoanodes need to be addressed for the development of tandem systems. To overcome these limitations and scale up BaTaOThe authors declare no conflict of interest.


In the original publication , there wThere were two figure numbers that were writing mistakes in Sections 2.3 and 2.4. The authors have replaced “there was no alteration in the NaHS treatment group ” with “there was no alteration in the NaHS treatment group ” in Section 2.3. And the authors replaced “Figures 3J and 4H” with “Figures 3I and 4H” in Section 2.4.The authors apologize for any inconvenience caused and state that the scientific conclusions are unaffected. This correction was approved by the Academic Editor. The original article has been updated.


However, the conventional spot mode analysis is restricted by depth effects at the spatial resolution < 0.5–1 μm. Thus sufficient signal amount cannot be achieved due to limited analytical depths, resulting in low analytical precision (1.5‰). Here we report a new method that simultaneously improves spatial resolution and precision of sulfur isotopic analysis based on the NanoSIMS imaging mode. This method uses a long acquisition time  for each analytical area to obtain sufficient signal amount, rastered with the Cs+ primary beam of ∼100 nm in diameter. Due to the high acquisition time, primary ion beam (FCP) intensity drifting and quasi-simultaneous arrival (QSA) significantly affects the sulfur isotopic measurement of secondary ion images. Therefore, the interpolation correction was used to eliminate the effect of FCP intensity variation, and the coefficients for the QSA correction were determined with sulfide isotopic standards. Then, the sulfur isotopic composition was acquired by the segmentation and calculation of the calibrated isotopic images. The optimal spatial resolution of ∼ 100 nm (Sampling volume of 5 nm × 1.5 μm2) for sulfur isotopic analysis can be implemented with an analytical precision of ∼1‰ (1SD). Our study demonstrates that imaging analysis is superior to spot-mode analysis in irregular analytical areas where relatively high spatial resolution and precision are required and may be widely applied to other isotopic analyses.NanoSIMS has been widely used for  The main reason is that when small probe sizes are used to achieve a high spatial resolution, the ion current of the primary beam and the intensity of the secondary beam are both reduced. Generally, the acquisition time is lengthened to obtain a sufficient signal for statistical purposes. However, rastering on a small area with a long acquisition time can cause noticeable depth effects, which result in changes in the angle at which the second ion is sputtering. The Instrumental Mass Fractionation (IMF) will change accordingly, resulting in a decrease in analytical accuracy . To redu2) for S isotopic analysis can be implemented with an analytical precision of ∼1‰ (1SD). This method is particularly suitable for sulfur isotopic analysis of submicron-scale to nanoscale oscillatory zoning or complex core-rim structure.An important function of NanoSIMS is its ion image mode. This mode is primarily used to observe the isotopic or elemental distributions using small probe sizes (down to 100 nm). The secondary ion image is acquired by rastering the primary beam over a larger area, thereby eliminating the crater effect. The isotopic composition can be determined by segmenting and calculating the secondary ion images. We report here a new method based on the NanoSIMS imaging mode that improves both spatial resolution and precision of sulfur isotopic analyses. The optimal spatial resolution of ∼ 100 nm (Sampling volume of 5 nm × 1.5 μm34S/32S) ratios are reported as δ34S in the standard per mil notation (‰):34S/32S value of 044162. Three natural pyrite standard samples have been used in our method, including PY-CS01 (δ34S = 4.6‰ ± 0.1‰ 1SD), PY 1117 (δ34S = 0.3‰ ± 0.1‰)and PY-SRZK (δ34S = 3.6‰ ± 0.05‰). Pyrite standard Py-1117 was used as the reference material, and the sulfur isotopic measurement was conducted on the standards of PY-CS01 and PY-SRZK used as the test samples. The bulk sulfur isotopic compositions of the pure mineral separations were determined by a conventional method (gas-source mass spectrometry). All the standards were embedded in epoxy resin and prepared as a polished disk with a diameter of 1 inch. Then coated with carbon and mounted in the sample holder.The measured . Ions image mode is used to acquire the ions image of 32S and 34S by raster the primary beam counting with electron multiplier (EMs). To optimize the primary beam setting with the spatial resolution and the signal amount for statistics, the Cs + beam of ∼ 1 pA and a diameter of∼ 100 nm was used (32S and 34S are ∼3–4 × 105 and ∼1.2–1.6 × 104 counts per second (CPS), respectively. These counting conditions can meet the needs of a sufficient signal amount of 34S and reduce the EM aging effect of 32S. Details on the aging effect and the signal amount for statistics are discussed in later sections. In order to eliminate interference from 32SH2−, 33SH1− on 34S−, the entrance slit (30 μm), apertures slit (200 μm) and exit silt (90 μm) were used to achieve a mass resolution of ∼7,000. Multi EM detectors were set to acquire the ions image of 32S, 34S, 60Ni, 80Se, 63Cu32S, and 75As32S. Among them, 60Ni-, 80Se-and 75As32S-are zoning related elements. 63Cu32S is used to eliminate some areas where chalcopyrite and pyrite are accompanied. In addition, the secondary electrons are also detected to represent the surface appearance characteristics. The NMR Probe is used to stabilize the magnetic field during the multi-collection.The was used . In this2. Before collecting the ions image, each area was presputted with a primary beam of ∼1 nA for∼ 30 s to 1 min to remove the carbon film on the surface and implant ions to obtain a stable yield. To compensate for the changes of sample height of different analytical location, EOS centering and SIB (secondary ion beam) centering were applied before each measurement. It is necessary to clarify that if the standard sample and unknown sample are in different holder, the EOS voltage are required to be adjusted to be consistent by changing the Z-axis. In addition, automatic peak centering was employed before acquiring the ions images. Under the primary beam condition of ∼1 pA with 20 × 20 μm2 raster size, the total counting time takes at least 2 h to acquire the analysis precession of 1‰ at the spatial resolution of 100 nm (discussed in later sections). The count rate of 32S is ∼ 300000–400000 cps. In order to reduce the impact of image drift, each area has been raster in 30 frames, with 256 × 256 pixels and 5 ms/pixel dwell time. The total duration is 9,900 s for each area and the sputter depth is ∼5 nm calculated using the sputter rate (0.2 nm*μm2 pA−1 s−1) report in the previous study . According to Poisson’s statistical theory, the analytical precision of d in the .MAX is located in the 240–290 mv interval, which can eliminate aging effects to obtain a high precision (0.5‰) isotope analysis. These two methods are combined in our study. The standards are used to monitor the drift, and the PHDMAX of EMs is maintained within a required range. We have tested the age effect on each analytical section. Each area has been raster in 30 frames, with 256 × 256 pixels and 5 ms/pixel dwell time. The total duration is 9,900 s for each area. As shown in Compared with the several minutes of analysis time used in spot mode, the longer detected time of each analysis was employed with electron multiplier. The aging effect of the EM is required to be considered. The aging effect of EM will occur when used to receive higher intensity secondary ions (more than 400000 cps), lowering the gain of the electron multipliers (the number of secondary electrons per secondary ions). It causes the isotope ratio drift of more than 1‰–2‰ per hour. In order to eliminate the aging effects on the isotopic analytical precision, two methods are mainly used for correction: 1) after analyzing ∼10 unknown samples, measure the references standards and correct the ratio. It is more commonly used and suitable for spot analysis with a shorter analysis time for each measurement. 2) Increasing the high voltage of EMs, the peak height distribution (PHD)34S = 0.3‰ ± 0.1‰) and Py-CS01(δ34S = 4.6‰ ± 0.1‰). In order to determine the coefficient β of various K values, different widths of aperture slits  have been used. The ion image acquisition time of each condition varies from 10 to 30 min, depending on the ion intensity. The sulfur isotopic data of QSA correction and K values measured using various aperture slit settings are given in 34S and Kcor before and after QSA correction is shown in 34Sexp exhibits a linear relationship to Kcor, which is related to the QSA effect. The QSA coefficient β (slope) of ∼ 0.98 is determined in our method. The slope is consistent in these two pyrite standards with different S isotopic compositions. The gap between the fitting line intercept indicates the difference in the compositions of the S isotope. Therefore, the QSA correction formula is determined:Quasi-simultaneous arrivals (QSA) effect may induce a specific type of bias. It is a challenge in high-precision analysis using EMs. For the QSA effect correction of S isotope analysis, Slodzian calculated Kcor (the average number of secondary ions ejected per primary ion) with an iterative procedure through probability distribution :Kcor=Ke34Scor measured using different silts is not shown as consistent. The instrumental mass fractionation (IMF) may vary with the used AS width, and the mass resolution is inconsistent. Therefore, IMF correction of the corresponding instrument settings, in which AS1 is used in this method, should be carried out after QSA correction. Moreover, the plot of raw δ34S of each frame on PY-1117 was measured using different aperture slits before and after QSA correction is given in 34Sexp presented a trend of less instrumental mass fractionation with the aperture slit width decreasing. Indeed, with the decrease of the K value, the influence of the QSA effect gradually becomes smaller.Note that if the instrument switches to the other analysis method and returns to the S isotope measurement, the β value and the calibration curve must be re-established. In addition, as shown in 2, with other instrumental settings remaining the same. As shown in Compared with the QSA coefficient β measured by previous studies in the range of 0.7–0.8, the value determined in our experiment is relatively higher. It is maybe due to the IMF difference between QSA correction in spot analysis mode and QSA correction in image analysis mode. QSA in spot analysis mode was also tested in our experiment. The raster size was decreased to 2 × 2 μm+ primary ion beam intensity is mainly caused by the variation of the evaporation of Cs₂CO3 in the ion source and the change of ionization of Cs+, which is mostly monotonic change within a few hours. Consequently, the primary ion beam intensity, which is called FCP (Faraday Cup Primary) on the NanoSIMS, was measured before (FCPbefore) and after (FCPafter) the multilayer isotope image analysis. Then the FCPbefore and the FCPafter were interpolated into each layer of images to determine the FCP(i) of frame i  is the FCP of the frame i after correction. Accordingly, the primary ion beam bombarded on the sample of each frame, which was used to calculate the K value, was determined by the relationship between the primary ion beam FCPi)( and the bombarded ion beam on the sample surface (FCO in the NanoSIMS). Then, the QSA correction was performed using the corresponding K value and β (before) and after (FCPafter) in the ions image acquisition, is applied to calculation the K value. The δ34S values were calculated after correction of the QSA effect using this K value; 2) The interpolated FCP using our method  to 0.8‰. Hence, it is necessary to perform primary beam current (FCP) correction on the results of each frame in analyzing high-precision stable isotopes in image mode. Then, the QSA correction can be performed to obtain more precise results.Where ue and β . To illur method  was apply change . The dec34S/32S ratio with the reference values measured using gas source isotope ratio mass spectrometry (GS-IRMS). Similarly, the IMF calibration in image mode was determined by the standards:34Smeasure and δ34Strue are the measured and recommended values of δ34S, respectively. Then, isotope image analysis was carried out on the sample to be tested, and IMF corrected the S isotope ratio to obtain the true value. In this study, the δ34S result of unknown sample was reported with associated 1SE uncertainty (err bar). And the uncertainty of the each ROI on the unknown sample (SEsample)is estimated as the square sum of the reproducibility of δ34S measurements on the corresponding reference pyrite of PY 1,117 (SD standard), the internal precision of each ROI on the sample (SE ROI) and the uncertainty of the reference values of the standards 1,117 (SD ref.):The matrix effect means instrumental mass fractionation varies with different matrixes. The standard–sample–standard bracket method was commonly used to calibrate the matrix effects. The IMF calibration should be performed using the same instrumental settings. In spot mode analysis, it was estimated by comparing the measured (1) Instrumental setting: According to the requirement of stable isotope image analysis, the instrument is set, including primary beam configuration, mass spectrometer configuration, and multi-collection configuration.(2) The QSA correction coefficient: Isotope imaging was performed on the isotope standard samples with aperture slits of different sizes to obtain the QSA correction slope for isotope image analysis.(3) The Instrumental Mass Fractionation (IMF) correction: On the sulfur isotopic standard sample, which matches the matrix of the sample to be tested, the instrumental mass fractionation is calculated according to the measured isotope ratio and the recommended value of the standard sample.2, a total of 30 frames and each layer acquisition time of ∼300 s. The standard–sample–standard bracket method to calibrate the matrix effects and QSA effects. Specifically, the isotope imaging on the standard was performed before the ion image acquirement on each area of unknown sample.(4) Ions image acquisition on the unknown sample: acquire the S isotope ions image on the unknown samples with the same analytical conditions as those on the standard samples, including a raster size of 20 × 20 μm32S and 34S in the ROI area was calculated and output from Image J as the raw data of the isotope for processed. The raw data were first corrected for dead time effect. After that, the measured 34S/32S were corrected for FCP, QSA and Matrix effect.(5) Data process of ion image on unknown sample: The 30 frames of ion images were automatically aligned and added. Then the Region of Interesting (ROI)s were selected depended the Experimental requirements. The total counts of In summary, the sulfur isotopic analysis method based on the NanoSIMS imaging mode was established.2 (∼78 nm per pixel) to 40 × 40 pixels2 to 200 × 200 pixels2 (32S and 34S of a single frame is (34Si/32Si)ROI, and each ROI has 30 frames ratio data. The relative internal precision calculation formula is:Compared with the detected using Faraday Cup, the internal precision of the isotope measurement employing EMs is limited to shot noise (Poisson error) and independent of thermal/JN noise . Therefo pixels2 . In spot34Si/32Si)ROI is the standard deviation of the ratio of 30 frames in a certain ROI, and MEAN (34Si/32Si)ROI is the average value of the ratio of 30 frames in a certain ROI. The Poisson error (%) is calculated following the formula in the 32S are more than 20 times 34S, the statistical error of 32S can be ignored. The error is estimated as 34S/32S vs. the ROI with the different size used is given in 34S count rate and with the ROI size increasing. The internal and Poisson errors are also relatively consistent in value. In addition, the dispersion of the ratio data decreases with the ROI size increase  and the higher the internal precision can be obtained, while external reproducibility is affected by more factors, such as signal statistics, sample topography, and surface potential.34S total counts and the external reproducibility of the pyrite standard of Py-1117. In the ions image with the size of 20 × 20 μm2 (256 × 256 piexl2), 100 ROIs with a size of 20 × 20 piexl2 (∼1.5 × 1.5 μm2) are randomly selected. In order to change the signal statistics of 34S total counts of the ROIs, different cumulative frames were adopted, which are 20, 40, 80, 100, 120, and 150 frames. The standard deviation of isotopic data of each group with the signal statics of 34S total counts is given in 34S total counts, the external reproducibility of ROI to ROI is improved. However, when it is close to 1‰, the external reproducibility improvement is limited when the total counts are increased. The main reason is that the influence of signal statistics on external reproducibility can be neglected compared to other factors.We determined the relationship between the signal statistics of 2 compared with the rectangular of 1 × 4 μm2 has the same sampling area and count statistics. However, its lateral spatial resolution has been improved from 2 μm to 1 μm. We have measured the S isotope values of ROIs with the same area size of 2.25 μm2 but different shapes on PY 1117 pyrite. The ROIs are shown in 2 in the center of the image. For the second group, 62 square ROIs of 1.5 μm × 1.5 μm regions were randomly selected without overlapping within the same region.Using ions image mode for sulfur isotope analysis, its analytical spatial resolution depends entirely on the size and the shape of ROI. The shape of ROI can be arbitrarily selected, which is the advantage of ions image analysis. For example, the square ROI with the size of 2 × 2 μm34S on PY-1117 with the rectangular and square ROIs are shown in 34S of 36.41 ± 1.1(1SD). It is 36.49 ± 1.1(1SD) with the 62 square ROIs of 1.5 μm × 1.5 μm. The results show that the different ROI shapes do not affect the isotope ratio and external reproducibility of ROI to ROI when the ROI area is the same. This result demonstrates the advantages of ions image analysis, optimizing ROI to improve spatial resolution for sample analytical requirements and ensuring higher precision by obtaining more count statistics.The result of RAW δ34S/32S IMF measured on Py-1117 is 1.0006 with the reproducibly of 1‰ (1SD), which is used to correct the measured δ34S of the test samples. The corrected results of δ34S on PY-SRZK with the ROIs of 1.5 μm2 and the PY-CS01 with the ROIs of 2.5 μm2 are shown in 34S result was reported with associated 1SE uncertainty , and twenty measurements on PY-CS01 gave an average δ34S of 4.3 ± 1.2 (1SD). The results were consistent with the recommended values of 3.6‰ ± 0.05‰ and 4.6‰ ± 0.1‰ within analytical uncertainty. The IMF was consistently shown between the two samples, with an uncertainty of 1.1‰ (1SD) obtained for the sixty measurements. In addition, the selected ROI areas were 1.5 μm2 and 2.5 μm2. Therefore, considering the primary beam spot size of ∼100 nm, the spatial resolution can reach 100 nm × 15 μm or 150 nm × 10 μm with an analytical accuracy of ∼1‰.Forty measurements on PY-SRZK gave an average δin-situ analysis using NanoSIMS imaging. A focused Gaussian Cs + beam of 1 pA was utilized to achieve a lateral resolution of 100 nm for S isotopic analysis. The acquisition time of 3 h for each analytical area is applied to obtain sufficient counts. The sulfur isotopic composition was acquired by the segmentation and calculation of the secondary ion images. The deadtime, FCP, QSA, and matrix effects have been corrected to improve the analytical accuracy to ∼1‰. This method is suitable for sulfur isotopic analysis of submicron-scale to nanoscale oscillatory zoning or complex core-rim structure. More importantly, this study demonstrates that the imaging mode of NanoSIMS is a powerful tool for high spatial resolution isotope analysis, providing a strategy that can be widely applied to improve both precision and spatial resolution for the analysis of other isotopes.This study developed a new method that improves the spatial resolution and precision of sulfur isotopic


In the published article, there was an error in the legend for Supplementary Material. Supplementary Figure S1 was displayed as “CCK-8 to determine the optimal drug concentration of mefunidone for 24 h for HK-2 cell viability”; Supplementary Figure S2 was displayed as “Mefunidone alleviated kidney fibrosis and inhibited EMT in IRI-induced CKD.” The correct Supplementary Figure S1 is “Mefunidone alleviated kidney fibrosis and inhibited EMT in IRI-induced CKD”; The correct Supplementary Figure S2 is “CCK-8 to determine the optimal drug concentration of mefunidone for 24 h for HK-2 cell viability.”Furthermore, there was an error in the The authors apologize for this error and state that this does not change the scientific conclusions of the article in any way. The original article has been updated.


Conspicuous ferroelectric properties including record‐high remanent polarization (2rP ≈ 60 µC cm−2), wake‐up‐free operation, and high compatibility with advanced semiconductor technology nodes, are achieved in the sub‐6 nm thin film. The prominent antiferroelectric to ferroelectric phase transformation is ascribed to the in‐plane tensile stress introduced into ZrO2 by the HfO2 seeding layer. Based on the high‐resolution and high‐contrast images of surface grains extracted precisely by helium ion microscopy, the evolution of the MPB between tetragonal, orthorhombic, and monoclinic phases with grain size is demonstrated for the first time. The result indicates that a decrease in the average grain size drives the crystallization from the tetragonal to polar orthorhombic phases.Atomic layer engineering is investigated to tailor the morphotropic phase boundary (MPB) between antiferroelectric, ferroelectric, and paraelectric phases. By increasing the HfO 2 seeding layer, triggering the morphotropic phase boundary effect correlated with the dramatic transition from antiferroelectric to ferroelectric phase in ZrO2. Furthermore, clear images of surface grains taken by helium ion microscopy are employed to characterize the transformation of crystallographic structures for the first time.Significant ferroelectricity in sub‐6 nm thin films is boosted by atomic layer engineering on the HfO In comparison with conventional perovskite materials, the fluorite HfO2/ZrO2‐based materials have favorable advantages including higher remanent polarization even with a thickness down to sub‐10 nm, larger bandgap, environmental friendliness, and high compatibility with advanced semiconductor technology. The distinguished superiority makes the HfO2/ZrO2‐based thin films very attractive in many fields such as high‐k gate dielectrics, non‐volatile memory devices, energy‐storage devices, and so on.The morphotropic phase boundary (MPB) between the antiferroelectric (AFE) and ferroelectric (FE) phases is a decisive factor to tailor the crystalline structures, which has caught widespread spotlight and interest recently.1 The  In addition, concerning the integration with the back‐end‐of‐line (BEOL), the processing temperatures for fabricating the FE devices should be restricted below 400 °C to reduce the impact of high thermal budget. However, in terms of various dopants such as Zr, Si, Al, Y, and La in well‐known high‐performance HfO2‐based FE thin films, the post‐annealing temperatures generally exceed 550 °C, even up to 800 °C, as the film thickness is lower than 6 nm, which is incompatible with the BEOL processes. The thickness scaling of the HfO2‐based thin films to the sub‐6 nm regime also gives rise to significant degradation of the ferroelectricity, which is ascribed to the increasing proportion of the dead layer (non‐polar phases) in the film. As a consequence, the improvement of crystallization into the polar o‐phase is regarded as a challenge to the FE thin films with a thickness scaling below 6 nm.With the evolution of 3D stacked memory structures, the continuous thickness scaling of FE thin films is urgently demanded to further increase the memory density.7 In a2 thin films. The relatively low crystallization temperature gives ZrO2 an edge over the HfO2‐based materials because of its more compatibility with BEOL. Therefore, it is of great significance to explore the pathway to induce the phase transformation from the non‐polar t‐phase to the polar o‐phase in ZrO2 thin films. It has been pointed out that the presence of MPB causes the transition between the AFE and FE phases by adjusting the composition, crystallization temperature, applied electric field, etc. In this research, a novel method is proposed for achieving the AFE to FE phase transformation by manipulating the MPB through monolayer engineering on a nucleation seeding layer. The various types of seeding layers such as Al2O3, TiO2, ZrO2, and HfO2 have been intensively investigated to improve the FE performance of nanoscale Hf0.5Zr0.5O2 (HZO) thin films. The seeding layer can facilitate the crystallization of the polar o‐phase, reduce the leakage current, adjust the dielectric properties, and improve the cycling reliability of HZO thin films. Although the studies have been conducted so far, there is still much room for significant improvement in ferroelectricity. Critical issues need to be addressed, including the thickness scaling, the reduction of process temperature, and the need for the wake‐up operation. As opposed to the comprehensive research on HZO thin films, the literature about the impact of the seeding layer on nanoscale ZrO2 thin films is scarce. In order to achieve pronounced FE properties in sub‐6 nm thin films, the atomic layer engineering on the HfO2 nucleation seeding layer is used to tailor the MPB for the AFE to FE phase transition in the overlying ZrO2 layer in this paper.One of the promising strategies to boost the ferroelectricity in sub‐6 nm thin films is the MPB‐correlated AFE to FE phase transition. It has been well recognized that the t‐phase is the most stable crystallographic structure in nanoscale ZrOin films.12 The The wake‐up effect has been recognized to be relevant with the conversion from the non‐polar crystalline phases to the polar o‐phase or the depinning of FE domains in the HZO thin films. The charged defects such as oxygen vacancies are redistributed more uniformly after the implementation of the cycling bipolar electric field , which leads to a homogeneous internal electric field for the reduction in the energy barrier of the transition from the non‐polar to the o‐phases. The formation of oxygen vacancies may be suppressed by the enhancement of crystal quality since high crystallinity can effectively confine the flexibility of lattice structures and restrain atomic relaxation. The monolayer engineering on the HfO2 seeding layer in this work might increase the degree of crystallinity, which is highly favorable for a decrease in the content of oxygen vacancies in the ZrO2 layer to alleviate the wake‐up effect, leading to the wake‐up‐free operation in various FE devices including ferroelectric field‐effect transistors, ferroelectric random access memories, and ferroelectric tunnel junctions.On the other hand, the prerequisite wake‐up process to achieve the optimal ferroelectricity in HZO devices is considered a practical challenge and a deterioration in the device reliability.15 The2 is investigated in detail by atomic layer engineering on the HfO2 seeding layer. The monolayer tailoring of materials is realized by the self‐limiting chemical reactions at the surface in the atomic layer deposition (ALD) process. In addition, the comprehensive electrical and material characterizations are conducted carefully to explore the MPB‐related phase transition, including the AFE/FE hysteresis loops, dielectric constant, high‐resolution transmission electron microscopy (HRTEM), nano beam diffraction (NBD), and X‐ray photoelectron spectroscopy (XPS). The out‐of‐plane, grazing incident, and in‐plane high‐power X‐ray diffraction (XRD) measurements and the sin2ψ method in the grazing incident configuration are also utilized to probe the residual strains in ZrO2 induced by the HfO2 seeding layer, revealing the correlation of MPB with the in‐plane tensile stress. Especially, the impacts of the grain size and its distribution on the AFE, FE, and PE phase transformation are analyzed using helium ion microscopy (HIM) for the first time to capture high‐resolution and high‐quality surface images. Conventionally, scanning electron microscopy (SEM) is a widely used method to acquire plane‐view images. However, it is inappropriate for SEM to characterize non‐conductive thin films due to the accumulation of charges at the surface. Since the charging effect causes blurred surface images of non‐conductive thin films, coating a heavy metal thin film such as gold or platinum is necessary to mitigate the surface charging in SEM. Nevertheless, the heavy metal coating leads to a distortion of the surface morphology or results in a loss of subtle surface features. Benefiting from the larger mass of He ions in HIM than that of electrons used in SEM, the scattered He ions keep almost the same direction when penetrating into the sample. Therefore, the interaction area on the sample surface is more convergent in HIM than that in SEM, giving rise to a much smaller spot size, higher resolution, and greater depth of focus in HIM. Besides, the issue of the charging effect could be avoided because of the deeper penetration depth of He ions and the neutralization by the electron flood gun equipped in HIM. As a consequence, the surface charging effect can be considerably alleviated in non‐conductive samples without any heavy metal coating in HIM, and so real images of tiny surface structures can be acquired. Moreover, a higher yield of secondary electrons further enhances the contrast of HIM images, which is conducive to clearly defining grain sizes and boundaries on the sample surface.In this paper, the MPB phenomenon for the transformation between the AFE, FE, and paraelectric (PE) crystalline phases in nanoscale ZrO2 seeding layer in this study, the dramatic AFE to FE phase transition in ZrO2 is achieved at a low processing temperature of only 400 °C, which leads to the realization of robust ferroelectricity with a record‐high remanent polarization (2rP ≈ 60 µC cm−2) in sub‐6 nm FE thin films, along with being free of the wake‐up operation. The excellent FE performance and the low thermal budget render the HZ (ZrO2 on the HfO2 seeding layer) structure indispensable and highly anticipated in nonvolatile memory devices in advanced semiconductor technology nodes.Based on the precise atomic layer engineering on the HfO22.1Figure 2 layer, and the monolayer engineering on the HfO2 seeding layer was realized by the accurate thickness control of atomic resolution to modulate the crystalline phases in the ZrO2 layer as the schematic illustration in Figure 2 seeding layer was prepared with 0, 2, 4, 6, 8, 10, and 14 ALD cycles in this paper, and so the stacking structures are denoted as H0Z, H2Z, H4Z, H6Z, H8Z, H10Z, and H14Z, respectively. The impact of the monolayer engineering by ALD to induce the AFE, FE, and PE phase transformation is illustrated in terms of crystal structures coupled with the Landau‐free energy landscapes in Figure AFE, FE, and dielectric properties of the HZ structure were examined by metal‐insulator‐metal (MIM) devices using ruthenium (Ru) as the electrodes, which are schematically plotted in 2.2P–V) and current density versus voltage (J–V) characteristics of all the MIM devices are shown in Figure P–V and J–V curves reveal the AFE feature of the H0Z, H2Z, and H4Z samples. The observed remanent polarization in the pinched P–V curves suggests the coexistence of the AFE and FE crystalline phases, that is, the presence of MPB, in the thin films. It is noticeable that the remanent polarization increases from the H0Z to H4Z samples, implicating that an increase in the thickness of the HfO2 seeding layer (THfO2) results in an increase of the FE phase in the HZ structures.The polarization versus voltage  Thus it can be concluded that the optimal and robust ferroelectricity is achieved in the H6Z sample, of which the double remanent polarization (2rP) is as high as ≈60 µC cm−2. It is worth noting that the 2rP ≈ 60 µC cm−2 is the record high in nanoscale FE thin films of a thickness below 6 nm. The result demonstrates the significant impact on the AFE and FE phase transformation by the monolayer layer engineering on the seeding layer. The polarity asymmetry of the P–V curves is discussed in Section The sharp change in the electrical properties of the HZ thin films can be observed in Figure 2.3Figure 4Z and H6Z samples, and the clear lattice fringes reveal a high degree of crystallinity in the films. It can be seen that the thicknesses of the H4Z and H6Z layers are ≈5.6 and 5.9 nm, respectively, which are in good agreement with the estimation based on the growth rate per ALD cycle of ZrO2 and HfO2. The NBD patterns of the H4Z and H6Z thin films are shown in Figure 2/nmcP4 space group along the [1¯11] zone axis, suggesting the presence of the non‐polar t‐phase in the H4Z layer. Whereas, the NBD spots of the H6Z thin film correspond well to the 1Pca2 space group along the [01¯1] zone axis in Figure 2 seeding layer triggers the t‐ to o‐phase transition in the overlying ZrO2 layer.°, and almost parallel to the sample surface can be probed by different configurations of the incident X‐ray and detector, as illustrated in Figure 0Z sample, the XRD peak emerges at the 2θ ≈ 30.5°, which is related to the t(101)/o(111) diffraction peak of ZrO2. The increase in THfO2 leads to a shift of the diffraction peak toward the higher 2θ direction for the H2Z, H4Z, and H6Z samples. Despite that, the XRD peaks of the t‐ and o‐phases are difficult to separate at 2θ ≈ 30.5°; however, it has been indicated that the diffraction angle of the o‐phase is smaller than that of the t‐phase. Therefore, the shift in the XRD peak toward the higher 2θ direction is unexpected for the H6Z sample because of the higher degree of crystallinity into the o‐phase as revealed by its outstanding ferroelectricity as shown in Figure THfO2.The XRD patterns were measured with the out‐of‐plane, grazing incident, and in‐plane scanning modes, from which specific crystallographic planes with the normals perpendicular, inclined at an angle of ≈14.8 of ZrO2.26 The6Z sample, the out‐of‐plane and in‐plane XRD patterns of the H4Z and H6Z samples were detected, which are shown in Figure 4Z sample. However, there is an obvious deviation between the out‐of‐plane, grazing incident, and in‐plane diffraction peaks in the H6Z sample as shown in Figure °) to the sample surface, whereas the diffraction peak detected from the in‐plane XRD measurement is attributed to the t(101)/o(111) plane almost perpendicular to the film surface. As compared with the out‐of‐plane and grazing incident XRD peak in Figure θ of the in‐plane XRD peak indicates a larger interplanar distance (d‐spacing) of the t(101)/o(111) plane in the in‐plane direction, thus revealing the presence of the in‐plane tensile stress in the ZrO2 layer of the H6Z sample. The result can be understood by the surface energy effect as predicted by the thermodynamical theoretical studies, which indicates that the ≈1 nm HfO2 thin film is prone to form the polar o‐phase. Since the lattice constant of the o‐phase in HfO2 is larger than that of the t‐phase in ZrO2, in‐plane tensile stress is introduced into the overlying ZrO2 layer. The residual tensile stress in the in‐plane direction is favorable for the boost of the phase transformation from the t‐ into o‐phases as reported in the literature, and accordingly the ZrO2 layer experiences the transition into the polar o‐phase. Therefore, the XRD patterns  plane of the monoclinic (m‐) phase (space group: 1/cP2)) emerge in H8Z and then dominate in the H10Z and H14Z samples as shown in Figure 2 thickness facilitates the formation of the non‐polar m‐phase. Actually, Y. Wang et al. have reported the observation of the m‐phase in HfO2 with a thickness of 2.48 nm, which is extremely close to the THfO2 (≈2.3 nm) of the H14Z sample. Therefore, the substantial degradation of 2rP in the H8Z, H10Z, and H14Z samples can be ascribed to the suppression of the crystallization to the o‐phase in the overlying ZrO2 layer.As the ALD cycles for preparing the HfOp: P21/c)26) emTHfO2 on the crystalline phase in the overlying ZrO2 layer, a model based on the phase transition temperatures of HfO2 and ZrO2 is proposed as illustrated in Figure 2 and ZrO2, the t‐/c‐phase is initially formed by the annealing treatment, followed by a transition from the t‐/c‐ to the m‐phases during the cooling process. However, it has also been reported that the metastable o‐phase may be present as the samples are cooled down. The emergence of the o‐phase is attributed to a lower kinetic energy barrier between the t‐/c‐ and o‐phases than that between the t‐/c‐ and m‐phases, which can be deduced from the effects of dopants, interface strain, surface energy, film thickness, and so on. Compared to HfO2, the transition temperature between the t‐/c‐ and m‐phases is much lower in ZrO2. Therefore, it can be expected that an increase in THfO2 leads to an increase in the phase transition temperature in the HZ thin film, as depicted in Figure 2. Therefore, the content of the o‐ or m‐phase increases as the THfO2 increases in the HZ thin film.To further understand the impact of the 2.4d‐spacing of specific crystallographic planes with different orientations. As shown schematically in Figure 1, S2, S3) on the sample surface and  upon the crystal plane, are employed to describe the experimental configuration. The ψ (psi) is the angle between the normals of the sample surface (S3) and the crystal plane (C3), and the χ (chi) is the angle of tilting the sample normal along the S1 axis. One of the approaches to extract the residual stress is known as the “sin2ψ method,” which employs the standard Bragg–Brentano configuration  with the variation in the ψ angle. However, the sin2ψ method is not appropriate for stress analysis of thin films because significant penetration of the X‐ray irradiation into the substrate greatly reduces the diffraction intensity from the thin film. An alternative method that is advantageous for thin films is the grazing incident configuration with the asymmetric diffraction geometry (ω≠θ and ψ≠χ) as shown in Figure  where the relationship between the ψ and χ angles is given by:The more rigorous analysis of residual stress in materials involves using XRD to measure the ntations.39 As d‐spacing of specific diffraction planes will vary as measured at different ψ angles. Due to the effect of Poisson's ratio, when a thin film is subjected to an in‐plane tensile stress, a corresponding compressive stress will be generated in the out‐of‐plane direction. Figure χ = 0°, 9°, 18°, 27°, 36°, and 45° using the grazing incident configuration, with φ = 0° and 180° at each χ angle (φ (phi) is the azimuth angle of rotating around the normal of the sample surface as plotted in Figure χ is greater than 45°. The d‐spacing at different sin2ψ is derived from the 2θ position of the XRD peak. It can be seen that the slopes of the regression lines are 0.098, 0.103, and 0.113, respectively, for the H0Z, H2Z, and H4Z samples. The slight increase in the slope reveals a gentle rise in the in‐plane tensile stress as the number of monolayers in the HfO2 seeding layer increases from 0 to 4. It is noteworthy that there is a steep increase in the slope to 0.214, accompanied by a maximum coefficient of determination (R2) of 0.875, for the H6Z sample. This indicates the existence of substantial and homogeneous in‐plane tensile stress, contributing to significant ferroelectricity of the H6Z device. Compared to the H6Z sample, the slope (0.152) and the R2 value drop in the H8Z sample, manifesting a release of residual stress that is preferable for the emergence of the m‐phase as widely reported in the literature. As the THfO2 further increases, the m‐phase becomes dominant in the H10Z and H14Z samples. Consequently, both the slopes of the regression lines and the R2 values abruptly decline, demonstrating that the in‐plane tensile stress almost vanishes in the thin films.If residual stress is present in a thin film, the 's ratio,43 whe2.5Figure r–Vε) curves of all the devices, which were extracted from the capacitance versus voltage (C–V) measurement. The evolution of the dielectric constant with the number of monolayers in the HfO2 seeding layer is plotted in Figure rε of the t‐, o‐, and m‐phases in ZrO2 are ≈46.6, ≈25, and ≈19.7 respectively. Two pairs of butterfly‐like peaks (the AFE feature) are observed in the r–Vε characteristics of the H0Z, H2Z, and H4Z samples, and their rε (taken at 2.5 V) are in the range of ≈28–30, which suggests the presence of the mixed state of the t‐ and o‐phases. When considering the rε at 0 V of the H4Z thin film, it is noteworthy that the MPB effect enhances the rε up to ≈38 and so a low equivalent oxide thickness (EOT) of only ≈0.54 nm is achieved, which is attractive for high‐k capacitors. With an increase of THfO2 over 4 ALD cycles, the r–Vε curves of the H6Z and H8Z samples convert to clear butterfly‐like double peaks , and their rε decreases to about 26 at 2.5 V. The results are highly consistent with the predominant o‐phase in the H6Z and H8Z layers. The rε taken at 2.5 V of the H10Z and H14Z samples further decreases to ≈20–21, which can be ascribed to the dominant m‐phase in the thin films. The dielectric properties of all the samples agree well with the electrical and crystalline phase analyses as revealed by the P–V, J–V, NBD, and XRD measurements.rε and low EOT of the H4Z device, further discussion on the nucleation of the HfO2 seeding layer is essential. By examining the P–V curves of the H0Z, H2Z, and H4Z devices in Figure rP of the H0Z, H2Z, and H4Z samples increases with the THfO2, suggesting that the HfO2 seeding layer may serve as the nucleation sites for the formation of the polar o‐phase in the overlying ZrO2 layer. However, the HfO2 seeding layer is too thin in the H2Z and H4Z samples, which results in inhomogeneous nucleation sites and so is insufficient to trigger the overwhelming t‐ to o‐phase transition. This can also be confirmed by Figure 2 seeding layer in the H2Z and H4Z samples induces the obvious MPB effect, which leads to a considerable increase in rε at 0 V. The dramatic phase transformation to the prominent ferroelectricity in the H6Z sample proposes that the 6‐cycle HfO2 seeding layer contributes to uniform nucleation sites for polar o‐phase formation.To gain a deep insight into the origin of MPB which leads to the high 2.62 seeding layer on the ZrO2 layer, the XPS measurement was carried out to analyze the O 1s spectra, which can give information on oxygen vacancies in the thin film. Figure 0Z and H6Z samples, respectively. The spectra can be deconvoluted into two main peaks: the one at a lower binding energy of ≈530.5 eV originates from the lattice oxygen, and the other at a higher binding energy of ≈532 eV is correlated with the oxygen vacancies. Accordingly, the contents of oxygen vacancies in the films can be estimated by the areas under the fitting curves in the O 1s spectra. One can see that the oxygen vacancies account for ≈12.6% and ≈4.5% of the O 1s spectra in the H0Z and H6Z samples, respectively, which indicates that the oxygen vacancies are substantially suppressed by inserting the HfO2 seeding layer. Since the displacement of atoms tends not to take place in the crystalline region, oxygen vacancies are not prone to present in the area with higher crystallinity. Hence the lower amount of oxygen vacancies in the H6Z sample can be attributed to the superior crystallinity as revealed by its higher XRD peak intensity than that of the H0Z samples. The O 1s XPS spectra of all the samples and the impact of oxygen vacancies on the crystalline phases are discussed in Section To further explore the effect of the underlying HfOacancies.47 AccrP as a function of wake‐up cycles of the H6Z device, revealing that the 2rP nearly keeps constant up to 105 wake‐up cycles. The P–V curves of the pristine state and that treated with 105 wake‐up cycles are shown in Figure P–V loops indicate the wake‐up‐free characteristic of the H6Z sample. It has been widely reported that the local electric field induced by oxygen vacancies will impede the switching of FE dipoles, resulting in weak ferroelectricity at the pristine state. The redistribution of oxygen vacancies by the cycling electric field during the wake‐up process alleviates the local built‐in electric field in the FE layer, which is the so‐called wake‐up effect for the enhancement of FE properties. As shown in the XPS measurement  HfO2 seeding layer provides the o‐phase for the nucleation of overlying ZrO2, which contributes to the smaller average grain size and the more uniform grain size distribution of the ZrO2 layer in the H6Z sample , H2Z (12.04 nm), and H4Z (11.34 nm) samples are smaller than 30 nm, the crystallization to the stable t‐phase can be expected as predicted by the thermodynamic theory. As shown in Figure 2 seeding layer leads to an increase in the surface energy, which gives the driving force to overcome the energy barrier for the transformation to metastable o‐phase from the t‐phase in ZrO2. This can be appropriately responsible for the correlation between the pronounced ferroelectricity, dominant o‐phase, and smaller average grain size of the H6Z sample. On the contrary, the HfO2 seeding layer may be too thin to form a continuous nucleation layer in the H2Z and H4Z samples. The random and non‐uniform nucleation sites bring about a larger variation of the grain size  of the H6Z sample, which is highly in conformity with the predominance of the m‐phase in the H14Z sample  can be expressed as follows:E is Young's modulus, ν is Poisson's ratio, γs is the surface energy of the film, γgb is the grain boundary energy, and D is the grain size. The equation reveals that a decrease in the grain size gives rise to an increase of the in‐plane tensile stress. Therefore, the minimum grain size in the H6Z thin film is also an indispensable factor responsible for significant changes in both the residual stress and the electrical properties of the HZ samples.Another noteworthy perspective is the correlation between the grain size and the in‐plane tensile stress. The 6‐cycle HfO follows:50‐hafnium 2]4), tetrakis‐(dimethylamido)‐zirconium 2]4), and O2 plasma, respectively. The deposition temperature of HfO2 and ZrO2 was 300 °C and the growth rates per cycle of both HfO2 and ZrO2 were ≈0.16 nm, which were estimated from the cross‐sectional HRTEM images as shown in Figure 2 ambient for 60 s. Then, the 75 µm × 75 µm square Ru top electrodes were precisely defined on the HZ thin films by the photolithography and lift‐off processes. The layer structure of the HZ sample is shown in the low‐magnification cross‐sectional HRTEM image in Section As plotted schematically in Figure P–V, J–V, and C–V curves. The P–V hysteresis loops and J–V curves were obtained from the bipolar triangular voltage waveform at 2 kHz. The C–V characteristics were measured with a small AC signal of 30 mV at 1 MHz superimposed on a DC bias sweep. The wake‐up process was performed by the cycling bipolar triangular electric field at 4.7 MV cm−1 with a frequency of 2 kHz.A semiconductor parameter analyzer (Keithley 4200‐SCS) was employed to characterize the electrical properties including the The cross‐section samples were prepared by the focus ion beam (FEI Helios). Then the HRTEM  was utilized to take the cross‐sectional images to extract the thickness of HZ layers and the diffraction patterns from the nano electron beam.αK X‐ray source at λ = 0.154 nm (Rigaku TTRAX3) to evaluate the crystalline structures and crystallinity in the nanoscale HZ thin films. The sin2ψ method in the grazing incident configuration was conducted by the XRD system (Bruker D8 DISCOVER with GADDS) equipped with the Cu αK X‐ray source (λ = 0.154 nm) at 1.6 kW for the residual stress analysis.Out‐of‐plane, grazing incident, and in‐plane XRD patterns were acquired by the high‐power X‐ray diffractometer with the Cu αK line at 1486.6 eV. The XPS spectra were then deconvoluted using the Gaussian‐Lorentzian peaks by the Multipak software .The ratio of oxygen vacancies in the HZ thin films was obtained according to the XPS (PHI VersaProbe scanning microprobe) with monochromatic Al The plane‐view images with clear crystalline grains and grain boundaries were taken by helium ion microscopy (Zeiss Orion NanoFab). On the other hand, SEM (JeoL JSM‐7800F Prime) was also used to take the surface image for comparison with the HIM images.The authors declare no conflict of interest.M.‐J.C. conceived the idea and coordinated the project. C.‐H.C. and S.‐H.Y. proposed the conceptualization. C.‐H.C. conducted methodology, investigation including electrical properties, XRD analysis, and writing—original draft. C.‐H.C., C.‐Y.C., and J.‐J.S. were responsible for the XPS measurement. C.‐H.C. and T.‐Y.W. contributed to the HIM characterization. M.‐J.C. was in charge of resources, writing—review and editing, supervision, project administration, and funding acquisition. C.‐H.C., Y.‐S.J., and M.‐J.C. were responsible for review questions.Supporting InformationClick here for additional data file.