Extend AbstractIntroductionThermoelectric materials, capable of directly converting heat to electricity, have a significant potential in sustainable power generation and thermoelectric cooling. Research in this area focuses on optimizing multiple interdependent and competing thermoelectric parameters to maximize the dimensionless figure-of-merit (ZT). Achieving an optimal ZT value requires materials with a large Seebeck coefficient, an excellent electrical conductivity, and a low thermal conductivity. However, these transport properties are interdependent and often in competition, necessitating a balance between the related parameters. A critical aspect of this optimization involves the trade-off between the Seebeck coefficient and electrical conductivity, and both of which are strongly affected by carrier concentration. Lower carrier concentrations generally enhance the Seebeck coefficient, but reduce the conductivity, whereas higher carrier concentrations favor the conductivity at the cost of a lower Seebeck coefficient. This competition also extends to the effective mass, where an increased effective mass can boost the Seebeck coefficient, but often reduce the conductivity due to the decreased carrier mobility. Effective thermoelectric optimization thus requires balancing enhanced effective mass with a high mobility. Furthermore, the coupling between electronic and lattice thermal conductivities significantly impacts the overall performance, positioning a thermal transport as a pivotal element in this optimization. To improve the ZT at low carrier concentrations, achieving electron-phonon decoupling while enhancing carrier mobility are an effective strategy. Based on some simplified thermoelectric models, this work demonstrated that electron-phonon decoupling could improve the thermoelectric performance at near room temperature under the condition of enhanced carrier mobility. In the narrow region of a low carrier concentration where the Seebeck coefficient could change minimally, the effect of the ratio of electronic conductivity to lattice thermal conductivity on the enhancement of ZT was evaluated. Finally, we emphasized the important role of electron-phonon decoupling in enhancing carrier mobility to optimize the thermoelectric performance.MethodsIn this work, a simplified theoretical model was proposed to evaluate a relationship between the ZT enhancement and the ratio of electronic conductivity to lattice thermal conductivity (i.e., the degree of electron-phonon decoupling). A simplified phonon-electron decoupling parameter model was designed via assuming that the Seebeck coefficient and carrier concentration could remain constant. In this model, the increase in electronic thermal conductivity was considered to be consistent with the increase in mobility. The lattice thermal conductivity was primarily affected by some complex factors such as crystal structure, defects, and anisotropy, and was therefore treated as a constant term without detailed consideration. In addition, the single-band Kane model was also utilized under approximate conditions, effectively simulating the impact of electron-phonon decoupling on the thermoelectric performance at near room temperature when enhancing the carrier mobility.Results and discussionElectronic conductivity and lattice thermal conductivity are two important thermal conductivity parameters in semiconductors, and their difference lies in the mechanism and main carrier of heat conduction. The enhancement of carrier mobility can significantly affect the electronic thermal conductivity without changing the lattice thermal conductivity. Therefore, the improvement of mobility can achieve the enhancement of ZT while unchanging the lattice thermal conductivity, and the improvement effect becomes dominant as the carrier mobility increases when using the electron-phonon decoupling parameter model. In addition, for systems with a lower lattice thermal conductivity, the electronic thermal conductivity can be a low level in order to achieve optimal optimization when using optimization strategies to increase carrier mobility. This indicates that the ratio of electronic conductivity to lattice thermal conductivity (i.e., the degree of electron-phonon decoupling) must be concerned when optimizing ZT by increasing carrier mobility in different systems because this ratio directly determines the extent of ZT improvement. Since the ratio of electronic conductivity to lattice thermal conductivity is usually limited in the range of [0.1, 10] when the ZT value reaches its peak as a function of carrier concentration, this range ensures that the material can achieve the optimal performance. Also, the ratio of electronic conductivity to lattice thermal conductivity increases with increasing temperature. This indicates that optimizing carrier mobility can have a more significant effect on the improvement of near-room temperature thermoelectric performance. Since the carrier concentration is directly proportional to the electronic thermal conductivity, and the lattice thermal conductivity does not depend on the carrier concentration, there is a positive correlation between the carrier concentration and the ratio of electronic conductivity to lattice thermal conductivity. This means that the electronic thermal conductivity increases quickly as the carrier concentration increases, resulting in an increase in the ratio. Therefore, the decrease in carrier concentration can significantly enhance the degree of electron-phonon decoupling as the carrier mobility increases. The thermoelectric performance can be effectively improved via regulating carrier concentration and optimizing carrier mobility, providing an important theoretical basis for the design and optimization of novel thermoelectric materials. In practical situations, the ratio of electronic conductivity to lattice thermal conductivity is usually limited to the interval [0.1, 10] when the ZT reaches an optimum value as a function of carrier concentration. Therefore, the improvement range of ZT value can also form different upper and lower limits, thereby restricting the improvement of thermoelectric performance. Although the improvement of carrier mobility can effectively improve the thermoelectric performance, there are differences in the electron-phonon coupling strength of different thermoelectric materials, which can make it difficult to improve the ZT to the theoretical optimal level. The single-band Kane model used is combined with some assumptions to simulate the electron-phonon decoupling discipline that is closer to the actual situation. For instance, the peak of ZT moves in the direction of lower carrier concentration, and achieves a larger increase in amplitude, thus forming a potential optimization area due to the increase in carrier mobility and the enhancement of electroacoustic decoupling (i.e., the deformation potential decreases and the B parameter increases).ConclusionsThis study systematically explored the discipline of electron-phonon decoupling at near room temperature when enhancing the carrier mobility based on the simplified electron-phonon decoupling model and single-band Kane model. The results showed that the ratio of electronic conductivity to lattice thermal conductivity could be concerned when optimizing the thermoelectric performance by increasing the carrier mobility, that is, the electron-phonon decoupling strength. The increase in this ratio with increasing temperature further verified the importance of optimizing the carrier mobility for improving the room-temperature thermoelectric performance. Meanwhile, the enhanced ZT of some thermoelectric materials was evaluated by the model, and the increase was far from the theoretical expectation. For the electron-phonon coupling, this synergistically promoted the shift of the ZT peak toward lower carrier concentrations and achieved significant improvements due to the increase in mobility and electron-phonon decoupling (i.e., the reduction of deformation potential and the improvement of B parameter), revealing a potential optimization range for the near room temperature thermoelectric performance. In summary, how to effectively manipulate electron-phonon decoupling while improving mobility to optimize the near room temperature thermoelectric performance could become a challenge in the field of future thermoelectric cooling.

IntroductionCaTiO3-based compounds emerge as a promising thermoelectric material due to their environmentally benign, thermally stable, and cost-efficient merits. Nonetheless, pristine CaTiO3 manifests inherently inferior electronic transport properties. In this paper, Ca1–xDyxTi0.95Nb0.05O3 (x=0–0.15) bulk ceramics were prepared via solid-phase sintering combined with hot-pressing sintering, and the composition, microstructure, and thermoelectric properties were analyzed. The results show that Dy and Nb doping can significantly increase the carrier concentration and effectively improve the electronic transport properties. Also, the lattice thermal conductivity is drastically reduced due to the introduction of large mass-field strain and stress-field strain, thus scattering high-frequency phonons. Ca0.85Dy0.15Ti0.95Nb0.05O3 bulk ceramic has a zT maximum of 0.29 at 1073 K, showing that the CaTiO3-based materials have a promising prospect for thermoelectric applications.MethodsCa1–xDyxTi0.95Nb0.05O3 (x=0–0.15) bulk ceramics were synthesized by solid-phase sintering and hot pressing. First, powders of CaCO3, Nb2O5, Dy2O3, and TiO2 were mixed and ground in a planetary ball mill for 12 h, then dried and cold-pressed. The pre-burned samples were ground for 6 h and the dried powders were cold-pressed again. The pressed samples were then placed in a graphite mold with a diameter of 13 mm for hot pressing at 1573 K for 1.5 h to obtain dense disks with a thickness of 2 mm.The phase composition of the bulk samples was performed by a model EMPYREAN X-ray diffractometer (XRD, PANalytical Co., the Netherlands). The elemental distribution of the samples was measured by a model JXA-8530F electron probe X-ray micro-analyzer (EPMA, JEOL Co., Japan). The Seebeck coefficient and resistivity of the samples were measured simultaneously by a model LSR-3 device (Linseis Co., Germany). The samples used for the test were long strips with the sizes of approximately 11.0 mm × 2.0 mm × 2.5 mm. The carrier concentration and Hall mobility of the system were measured by a model 8400 Hall measurement system (Lake Shore Co., Ltd., USA). The thermal conductivity of the system was measured by a model LFA-457 laser flash thermal conductivity meter (Netzsch Co., Germany), and the specific heat capacity of the bulk samples was calculated according to the Debye-Dulong law.Results and discussionThe polycrystalline Ca1–xDyxTi0.95Nb0.05O3 (x=0–0.15) samples with a single phase were prepared. The main phase of the synthesized samples is an orthorhombic structure, with a space group of Pnma. Dy3+ and Nb5+ occupy Ca2+ and Ti4+ sites in the matrix, respectively, causing the lattice expansion and introducing the donor impurity levels, transforming CaTiO3 from an insulating wide-bandgap semiconductor to a good electrical conductivity thermoelectric material. The actual amounts of elements Dy and Nb in the bulk samples are basically consistent with the nominal doping concentrations, indicating that elements Dy and Nb are incorporated into the CaTiO3 matrix effectively. The electrical conductivity decreases with increasing temperature, while the absolute value of the Seebeck coefficient increases with increasing temperature, showing typical characteristics of degenerate semiconductor electron transport. At high temperatures, the lattice vibrations are violent, and the acoustic wave scattering dominates the carrier transport, with other scattering mechanisms having little effect on the electronic transport properties. As Dy concentration increases, the effective mass of the density of states decreases. When x=0.15, the maximum zT value can reach 0.29 at 1073 K, which is comparable to the others reported CaTiO3 based thermoelectric performances in the literature.ConclusionsPolycrystalline Ca1–xDyxTi0.95Nb0.05O3 (x=0–0.15) samples with a single phase were prepared by solid-phase reaction combined with HP sintering. Dy and Nb doping could significantly increase the carrier concentration and effectively improve the electronic transport performance. Meanwhile, the high-frequency phonons were scattered due to the large mass difference and covalent radius difference between Dy3+ and Ca2+, resulting in a substantial reduction in the lattice thermal conductivity of the system. Ca0.85Dy0.15Ti0.95Nb0.05O3 achieved the maximum zT value of 0.29 at 1073 K due to the simultaneous optimization of electrical and thermal properties, indicating that CaTiO3 oxide thermoelectric materials could have a promising prospect for thermoelectric applications.

IntroductionThermoelectric materials have a promising prospect as they can directly convert thermal energy into electrical energy. Some thermoelectric materials are discovered in recent years, especially bismuth telluride (Bi2Te3)-based materials that are capable of large-scale commercialization. At present, the average zT value and conversion efficiency of Bi2Te3-based materials can be further enhaced. The low-valent cation doping is achieved via doping to optimize the electrical conductivity, while introducing defects as phonon scattering centers to effectively reduce the lattice thermal conductivity. This strategy is verified to be the most effective optimization method. In this paper, Na2S was selected as a p-type dopant to dope Bi0.42Sb1.58Te3 (BST) alloys, the conductivity was optimized via replacing the cation position of the BST matrix, and improving the solid solubility of Na and strengthening its mechanical properties by element S. The lattice distortion caused by Na2S doping in the BST alloys and the nanopore structure generated by the volatile element Te in the matrix were investigated.MethodsThe original high-purity Bi powder (99.99%, in mass, the same below), Te powder (99.99%), Sb (99.99%), and Na2S powder (Aladdin Co., China) were precisely weighed in an Ar atmosphere glove box based on their nominal composition (Bi0.42Sb1.58Te3 + x% Na2S, where x = 0, 0.2, 0.5, and 1.0). The weighed mixed powder was then placed in a quartz tube, evacuated to a vacuum degree of 10–4 Pa, and sealed. The quartz tubes were pre-plated with carbon to avoid the possibility of Na corrosion. The mixed powder in the sealed quartz tube was heated in a vertical resistance furnace for melting at 1125 K for 12 h and kept for 16 h. The ingots obtained after melting were ground by a model QM-3SP2 planetary ball mill at 425 r/min for 6 h. Finally, the powder was sintered at 698 K and 50 MPa by spark plasma sintering to prepare the bulk samples.The phase structure of the samples was analyzed by X-ray diffraction (XRD, Rigaku Co., Japan) with Cu K radiation ( = 1.540 6 A) in a diffraction angle range of 20°–60° with a step size of 0.02° (5 (°)/min). The microstructure of the samples was examined by scanning electron microscopy (SEM, ZEISS Co., Germany). The thermoelectric performance of the samples was analyzed via measuring their Seebeck coefficient, electrical conductivity, and power factor in a model ZEM-3M10 Seebeck coefficient/electric resistivity measuring system (Ulvac-Riko Co., Japan) under a thin helium atmosphere. The thermal diffusivity of the samples was measured by a model LFA457 laser flash instrument (NETZSCH Co., Germany), and the thermal conductivity was calculated based on tot=DCp, where D is the thermal diffusivity, Cp is the specific heat capacity deduced via the Dulong-Petit limit, and is the density of samples. The density was determined according to the Archimedes principle. The carrier concentration and mobility of the samples were measured at 295 K under an applied magnetic field of 1.5 T and an electrical current of 30 mA in a model PPMS-9T physical properties measurement system (Quantum Design Inc., Japan).Results and discussionWhen Na2S is used as a dopant, the electrical conductivity of 0.5% Na2S-doped BST sample increases from 598 S/cm of the pure sample to 749 S/cm at 300 K, and the lattice thermal conductivity of the sample decreases to 0.55 W/(m·K) at 300 K. The peak zT of 0.5% Na2S-doped BST sample reaches 1.3 at 300 K due to the increase of power factor and the decrease of thermal conductivity, which is 49.4% higher than that of the undoped sample. The thermoelectric conversion efficiency of the single-arm device reaches 3% as T=275 K due to its excellent thermoelectric figure of merit. In addition, the average hardness of the doped samples also increases to 1.09 GPa.ConclusionsIn this work, p-type BST thermoelectric materials were prepared by a solid-state method and a spark plasma sintering technology. Utilizing Na2S as a dopant achieved a low-valent substitution at cation sites, and introduced holes, thus enhancing the carrier concentration and optimizing the electrical conductivity of BST matrix thermoelectric materials. The peak zT of 0.5% Na2S-doped BST sample reached 1.3 at 300 K, which was 49.4% higher than that of the undoped sample. Furthermore, a large number of point defects enhanced phonon scattering and reduced the thermal conductivity of the material. In addition, element S could also enhance the solid solubility of Na in the BST matrix and the solid solution strengthening. The average hardness of the doped samples increased to 1.09 GPa.

IntroductionPyroelectric materials are widely used in the field of infrared detection. With the development of miniaturization and integration of devices, it is particularly important to develop lead-free pyroelectric sensitive elements with superior pyroelectric properties and high depolarization temperature compatible with reflow soldering process. In this paper, different mass fractions of Mn ions were doped into 0.96NaNbO3–0.04BaTiO3 matrix materials via a component gradient design. The effect of doping amount on the microstructure, ferroelectric properties, dielectric and pyroelectric properties of 0.96NaNbO3–0.04BaTiO3 ceramics was systematically investigated. The results show that when the doping amount of Mn is 0.4%, the ceramics have the optimum pyroelectric properties, which has a practical prospect of uncooled infrared detection.MethodsAll the components of sodium niobate ceramics were prepared by a solid-phase sintering method. The series of components were weighed according to the stoichiometric ratio, and the raw materials with anhydrous ethanol were ground in a stirred bead mill with zirconia beads at 400 r/min for 6 h. After being dried and shaped, the ground material was was pre-sintered at 1000 ℃ for 4 h, and then subjected to milling and drying. Green wafers with a diameter of 12 mm were prepared by pressure cold isostatic pressing at 200 MPa for 5 min after adding 10% PVA. The wafers were degummed at 600 °C for 2 h to remove organic matter and other impurities. To prevent the volatilization of sodium at a high temperature, the padding with the same composition was used for burial treatment, and the padding was kept at the corresponding temperature for 3 h and then cooled naturally. The sintered ceramic wafers were ground to fine particles with the size of 300 m, coated with silver electrodes on both sides, and polarized in silicone oil under an electric field of 8-10 kV/mm at 80 ℃ for 15 min. The electrical properties were tested after 24 h.Results and discussionThe microstructural morphology, dielectric properties, temperature stability, and pyroelectric properties of 0.96NN–0.04BT ceramics doped with 0–0.6% Mn are investigated. The results show that doping Mn ions significantly improves the sintering characteristics of 0.96NN–0.04BT ceramics, enhances the density, and the relative densities of all components reach above 90%. The NN–BT ceramic grain size increases with the increase of the doping amount. The substitution position of Mn ions causes changes in the electrical properties of the ceramic as the Mn doping content increases gradually. The concentration of oxygen vacancies caused by Na+ evaporation decreases when Mn ions substitute for A-site vacancies, which is beneficial to reducing the ceramic dielectric constant and dielectric loss, thus significantly improving the detection sensitivity factor. However, as the Mn doping content further increases, the imbalance of charges is prone to attract oxygen vacancies to form defect dipoles due to the heterovalent substitution of Mn ions on the B-site, thereby producing domain wall pinning effects. Also, the defect dipoles cause the increasing coercive field of the doped material and make it difficult to fully polarize, resulting in the material that is unable to fully utilize its pyroelectric properties.ConclusionsThe ceramic could have more suitable characteristics for pyroelectric requirements via appropriately adjusting the Mn doping concentration. The ceramic with Mn ion doping mass fraction of 0.4% had the optimal pyroelectric performance (i.e., the pyroelectric coefficient of 2.11×10–8 C·cm–2·K–1, and the Fv value of 3.74×10–2 m2·C–1), which was greater than 1.3 times and 1.8 times of the un-doped component. The coercive field reduced from 4.04 kV/mm of the un-doped component to 3.17 kV/mm, which was decreased by 21.5%. The results of the thermal stability test proved that the depolarization temperature could be maintained at 310 ℃, which had a practical prospect for non-cooled infrared detection.

IntroductionSrO(SrTiO3)2 (Hereinafter referred to as Sr3Ti2O7) is a layered perovskite-like structural material, which can be considered as a SrTiO3 superlattice interspersed with SrO layers. Compared to SrTiO3, Sr3Ti2O7 exhibits higher asymmetry and more significant dipole moment changes, thus providing a broader space for infrared radiation performance regulation. Also, the interface between SrO layer and SrTiO3 layer enhances phonon scattering, significantly reducing the thermal conductivity of electron-doped Sr3Ti2O7, compared to cubic SrTiO3. For the high-temperature stability and environmental friendliness of Sr3Ti2O7, there is a motivation to explore the relationship between its thermoelectric transport properties and infrared radiation properties to develop novel multifunctional materials. However, little researches on the intrinsic relationship between these two properties in Sr3Ti2O7 have been reported. Exploring the relationship between thermoelectric and infrared radiation properties and further studying the combination of thermoelectric materials with infrared materials can enhance the conversion efficiency and functional applications of thermoelectric materials, driving technological development and innovation, and open new avenues for applications in environmental monitoring, energy management, and efficient detection across multiple fields.MethodsSrCO3, TiO2, Nb2O3, and Gd2O3 were mixed as raw materials with an appropriate amount of alcohol according to a specific molar ratio. The mixed powder was ground in a ball mill for 1.5 h. Afterwards，the prepared raw powder in a mold was pressed by an electric press at 3.0 MPa for 2.0 min, and then demold. The pressed pellets were tightly wrapped with a plastic wrap and pressed by cold isostatic pressing at 39.5-44.5 MPa for 8.0 min. Afterwards, the sample wrapped in carbon paper was placed in an alumina crucible. The sample in carbon powder was heated in a muffle furnace at 1200 °C for 5.0 h. The sintered pellets were ground into a powder. The powder was pressed in an electric press at 5.0 MPa for 1.5 min, and then by cold isostatic pressing again at 39.5–44.5 MPa for 8.0 min. The pellets wrapped in carbon paper were buried in carbon powder and heated in the muffle furnace at 1500 °C for 10.0 h to obtain the final sample.Results and discussionThe XRD patterns of (Sr0.8Gd0.2)3(Ti1–xNbx)2O7 (0.01≤x≤0.05) show that the samples are similar to the standard material of Sr3Ti2O7, with a small amount of Gd2O3 phase. The most intense peak shifts between the planes (105) and (110), indicating that Nb doping affects the phase composition. As Nb doping increases, the XRD peak positions between 31° and 33° shift to lower angles due to the lattice distortion caused by Nb5+ with the radius of 0.690 substituting for Ti4+ with the radius of 0.605 . The SEM image reveals that a small amount of Nb doping (1%, in mole) improves the density, but higher doping levels increase the porosity and reduce the density.The electrical conductivity of the sample decreases with increasing temperature. The maximum electrical conductivity for the sample doped with 1% (in mole) Nb can be obtained, which can be enhanced by 503% at 340 K and 95% at 973 K, compared to undoped samples. The Seebeck coefficient of Gd20Nb1 significantly reduces, while the power factors of Gd20Nb1 and Gd20Nb2 are similar at high temperatures. The thermal diffusivity, lattice thermal conductivity, and total thermal conductivity all decrease with increasing temperature. The sample Gd20Nb1 exhibits the maximum thermal conductivity, while the sample Gd20Nb5 reaches the minimum value of 1.4 W·m–1·K–1 at 973 K. The zT value is maximum at 0.19 for the sample doped with 2% (in mole) Nb, which is primarily affected by the power factor. A lower thermal diffusivity of the sample Gd20Nb2 resulted in a lower thermal conductivity, compared to the sample Gd20Nb1, achieving the maximum zT value.The Nb-doped samples show an intense infrared absorption in the range of 8–14 m, as the infrared emissivity increases with temperature due to the enhanced carrier scattering. At 973 K, the infrared emissivity of the sample Gd20Nb1 reaches 95% in the range of 8–14 m, indicating potential applications in high-temperature protective coatings. This demonstrates a promising potential for simultaneous tuning of thermoelectric and infrared properties, thus offering some possibilities for Sr3Ti2O7-based materials in radiation heat shielding and other fields.ConclusionsGd/Nb co-doped Sr3Ti2O7 ceramics were synthesized by a solid-state method. The influence of Nb doping concentration on the thermoelectric transport and infrared radiation properties as well as the interrelationship between these properties were analyzed. The results showed that a small amount of Nb doping significantly enhanced the electrical conductivity of the material, and the zT value reached an optimal 0.19 at 973 K when the Nb doping concentration reached 2% (in mole). Furthermore, the Nb-doped samples exhibited an intense absorption in the 8–14 m range, and their infrared emissivity increased with the increase of the power factor and zT value. At 973 K, the sample Gd20Nb1 achieved the maximum average infrared emissivity of ~95% in the 8–14 m atmospheric window. The sample Gd20Nb5 had the minimum thermal conductivity of ~1.4 W·m–1·K–1. The results revealed the the synergistic effects of Sr3Ti2O7-based ceramics on the thermoelectric and infrared radiation properties, demonstrating their application potential in radiant heat protection and high infrared emissivity materials. This study could provide important references for the further optimization and design of novel thermoelectric and infrared materials.

IntroductionWith the rapid development of portable and wearable microelectronics, flexible thermoelectric materials have attracted much attention due to their ability to enable self-powered devices via utilizing the temperature difference between the skin and the environment. However, the scarcity of tellurium resources and its high toxicity pose some challenges to the commercial application of tellurium-based thermoelectric materials. It is thus critical to develop high-performance, low-cost, and non-toxic alternatives for flexible thermoelectric films. Copper(I)selenide (i.e., Cu2Se), as a representative liquid-like thermoelectric material, has attracted recent attention. Cu2Se is an intrinsic p-type semiconductor composed of abundant, low-cost, and non-toxic elements, thus offering significant advantages in the field of sustainable energy. The existing studies report the fabrication of high-performance Cu2Se thin films using conventional preparation techniques. However, these processes face significant limitations in pattern refinement and miniaturization, hindering their application in the fabrication of micro-thermoelectric devices. Inkjet printing technology, as a low-cost, efficient, and high-precision digital printing process, offers a promising solution to these challenges. This technique enables the preparation of complex patterns without masks or templates and is suitable for a variety of substrate materials. Therefore, this study was to prepare Cu2Se nanopowders and explore their dispersion properties in different solvents to develop Cu2Se inks suitable for inkjet printing technology, thereby offering insights into the preparation of flexible thermoelectric films and their application in next-generation wearable devices.MethodsCu2Se nanoparticles were synthesized by the following methods. In method 1 (i.e., hydrothermal method), 5 mmol copper chloride, 2.5 mmol selenium powder, 0.05 mol sodium hydroxide were mixed with 50 mL deionized water, and continuously stirred at room temperature for 20 min, and then a certain amount of hydrazine hydrate solution was added, and the mixture was transferred to 100 mL PPL–stainless steel autoclave. Subsequently, the sealed autoclave reactor was heated at 180 °C for 20 h and then naturally cooled to room temperature. The precipitate at the bottom of the reactor was centrifuged at 10 000 r/min and washed with ethanol and distilled water for several times to obtain Cu2Se powder. In method 2 (i.e., hydrothermal method), 1 mmol of selenium dioxide and 2 mmol of copper acetate were mixed with 65 mL of deionized water, and then stirred at room temperature until there was no sediment. Afterwards, hydrazine hydrate was slowly added and stirred for 10 min. The mixture was transferred to 100 mL PPL–stainless steel autoclave, and then the sealed autoclave reactor was heated at 180 ℃ for 24 h. After cooling to room temperature, the resulting sediment was collected and centrifugally cleaned. The content of hydrazine hydrate was changed by the control variable method in the synthesis process to investigate its effect on the morphology. In method 3 (i.e., wet synthesis), 1.54 g polyvinylpyrrolidone (PVP) was dispersed in 77 mL deionized water, followed by 1.92 mmol selenium dioxide (9.6 mL) and 11.52 mmol ascorbic acid (28.8 mL) solutions and stirred for 15 min. Afterwards, 3.84 mmol cupric sulfate pentahydrate (9.6 mL) and 15.36 mmol -ascorbic acid (38.4 mL) solution were added successively. The solution was magnetically stirred at room temperature for 16 h. After washing with deionized water for several times, Cu2Se core–shell nanoparticles were collected by centrifugation at 11 000 r/min.For Cu2Se ink configuration, the prepared Cu2Se powder was dispersed in deionized water and ethanol by an ultrasonic dispersion method. After ultrasonic treatment for 1 h, the dispersion effect was analyzed.Results and discussionCu2Se nanoparticles are prepared by the three methods, and the crystal structure of the synthesized products is verified by the XRD patterns. The morphology and size of Cu2Se particles prepared by different methods are different. In method 1, Cu2Se powder prepared under alkaline conditions has a flake structure with a particle size of more than 10 m, which does not meet the size requirements of inkjet printing. In method 2, under the influence of different contents of hydrazine hydrate, the synthesized powders show a large irregular flake structure (~500 nm) and small granular structure, and the particle size does not decrease significantly. The particle size of the powder is reduced when PVP is introduced into the subsequent synthesis process, indicating that the particles do not continue to grow into sheets during the synthesis process due to the coating of PVP, but remain small and irregular nanoparticles, and the particles are still agglomerated, which can cause a nozzle clogging in the subsequent printing process. In method 3, Cu2Se particles synthesized by the wet method are spherical with a particle size of approximately 40 nm, and uniform in sizes without particle agglomeration, which can meet the requirements of material size for inkjet printing.Cu2Se particles synthesized by the wet method and dispersed under ultrasound in ethanol show a well-dispersion, and the particles in ink do not settle significantly after placing for a period of time, showing a good stability and suitable for inkjet printing.ConclusionsCu2Se particles were prepared by the three methods for Cu2Se ink suitable for inkjet printing, respectively, and nano-sized particles of Cu2Se with uniform particle size and non-agglomeration were prepared by the wet method. The effective dispersion and stability of Cu2Se particles were mainly attributed to the coating of PVP on the surface of the particles, improving the solubility of the particles in ethanol, effectively preventing the particle aggregation. This study could provide a foundation for the subsequent development of high-performance Cu2Se thermoelectric thin films and their microdevices.

IntroductionHigh-capacity anode materials are essential for high-energy sodium-ion batteries. Sn anodes with their high theoretical capacity can face significant volume expansion (i.e., up to 420%) during sodiation, thus affecting material stability. Previous strategies for long-cycling micron-sized Sn are low loadings and inadequate for commercial needs. Enhancing cyclic stability under a high mass loading is thus crucial for high-performance sodium-ion batteries. In this paper, we simply mixed Sn with HC to form an HC "fence" that could prevent Sn particle reagglomeration and allow self-evolution. Adding 50% (in mass) HC boosted Sn/HC cyclic stability without shedding or cracking after 100 cycles at 70% (in mass) HC. The optimized Sn30HC70 delivered 1.8 mA·h/cm2 at 200 mA/g and retained 93.3% capacity after 200 cycles at 500 mA/g. The Sn30HC70||Na3V2(PO4)3 (NVP)full cell had 186 W·h/kg energy density at 0.5 C and retained 92.11% capacity at 6 C.MethodsA slurry was prepared via mixing Sn/HC, conductive agent (SP), and sodium alginate (SA) in deionized water in a mass ratio of 8:1:1. The mixture was coated onto carbon-coated aluminum foil, dried under vacuum at 80 ℃ for 12 h, and then cut into electrode disks with the diameter of 11 mm. The electrode disks with different SnxHC(100–x) ratios were obtained, i.e., HC, Sn20HC80, Sn30HC70, Sn40HC60, and Sn50HC50, via adjusting Sn content in the mixed anode (x=0%, 20%, 30%, 40%, and 50%). In an argon-filled glovebox (the volume fractio of H2O, O2 ≤ 1×10–6), 2325-type coin-cell half-cells were assembled with sodium as a counter/reference electrode, Celgard2500 as a separator, and 1 mol/L NaPF6 in Bis(2-methoxy ethyl)ether as an electrolyte. The full cells of (Sn30HC70||NVP and Sn||NVP) were also constructed with Sn30HC70/Sn as an anode and NVP as a cathode. The calculations of capacity and energy density excluded the masses of binder and conductive agent.Results and discussionAt a loading of 5.5 mg/cm2, Sn experiences a rapid capacity decay, losing up to 60% after 200 cycles. However, blending 50% (in mass) HC enhances Sn/HC cyclic stability under the same conditions. The SEM images of the sample after 100 cycles show the best-preserved electrode surface morphology with 70% (in mass) of HC, absent of active material shedding or surface cracking. This indicates that 70% (in mass) of HC achieves a structural stability for Sn anode, exhibiting a promising potential for developing long-term cyclic stability. At a current density of 500 mA/g, this electrode retains 93.3% of its initial capacity after 200 cycles.The results show that the capacity contribution of HC within the electrode material increases as the HC content increases, via comparing the voltage-capacity profiles of Sn/HC anodes with different HC proportions. The charging platforms of Sn at 0.2 V and 0.31 V merge into a platform, potentially due to the reduced transport resistance facilitated by HC. This can be further clarified via comparing the electrochemical impedance spectroscopy (EIS) data of Sn30HC70 and pure Sn, where the intrinsic impedance and diffusion impedance of Sn/HC both are lower than those of pure Sn.The ex-situ X-ray Diffraction (XRD) patterns indicate the phase transitions of Sn/HC at different potentials. A peak at 2 of 35.7° appears and disappears during the charge-discharge process, indicating the high reversibility of Sn material during cycling.The Sn30HC70||NVP full cell validates Sn/HC practicality in sodium-ion batteries. achieving 186 W·h/kg energy density at 0.5 C and retains 92.11% at 6 CConclusionsMixing HC with Sn effectively could enhance the cyclic stability of micron-sized Sn. This was attributed to the HC fence structure around Sn particles. The optimized Sn30HC70 anode had an areal capacity of 1.8 m·h/cm2 at a current density of 200 mA/g and retained 93.3% after 200 cycles at a current density of 500 mA/g. The full cell assembled with NVP maintained a capacity retention of 92.11% after 100 cycles at 0.5 C, further validating the effectiveness of the HC "fence" structure design in improving the stability of Sn anodes under high-mass loading conditions.

IntroductionWith the development of science and technology, precision positioning technology becomes one of the key technologies driving the prosperity and development of modern science and advanced industrial technology in the 21st century. It is widely used in high-tech fields such as precision machining and aerospace. Among the core components of precision positioning technology, the micro-displacement actuator with a high positioning accuracy plays a critical role. However, conventional micro-displacement actuators suffer from several drawbacks, including long transmission chain, complex structure and low precision, which pose some challenges to precision positioning system. Electrostrictive effect has attracted considerable attention due to its advantageous properties, such as the absence of hysteresis and independence from the direction of the electric field. The relaxor ferroelectrics, which exhibit minimal hysteresis are commonly studied for their electrostrictive behavior. Although extensive research has been conducted on the large electrostrictive effect of materials, there is still no analytic expression to describe electrostrictive strains. In this paper, the electrostrictive strain of high-energy electron irradiated P(VDF-TrFE) 68/32 relaxor ferroelectric copolymers was measured as a function of applied electric field using a laser-assisted micro-displacement measurement setup. An analytic expression of electrostrictive strain as a function of applied electric field was derived based on the thermodynamic phenomenological theory.MethodsAn electrostrictive effect test platform was built, and the electrostrictive effect of high-energy electron irradiated P(VDF-TrFE) 68/32 relaxor ferroelectric copolymers was examined. The tested results were analyzed. The phenomenological theory was used to derive the analytic expression of electrostrictive strain as a function of applied electric field for relaxor ferroelectrics.Results and discussionThe results reveal that the electrostrictive strain has a quadratic relationship with the electric field at lower electric fields. At higher electric fields, the relationship transitions to a power of 2/3, and at even higher electric fields, it further shifts to a power of 2/5. The electrostrictive strain as a function of electric field is analyzed for relaxor ferroelectric ceramics and polymers using this analytical expression. The fitting results confirmed the validity of this relationship across a wide range of electric fields. The strain and crossover electric field of the material can be designed in terms of the analytical expression of strain as a function of electric field.ConclusionsThe high-energy electron irradiated P(VDF-TrFE) relaxor ferroelectric polymers were prepared, and their electrostrictive strain as a function of electric field was determined. In addition, an analytical expression was also derived based on the thermodynamic phenomenological theory, which was used to fit the electrostrictive strain as a function of electric field for irradiated copolymers. The electrostrictive strain had a quadratic relationship with the electric field at lower electric fields. At higher electric fields, the relationship transitioned to a power of 2/3, and at even higher electric fields, it further shifted to a power of 2/5. These relationships were consistent with the experimental results. Furthermore, the expression was used to fit the electrostrictive strains versus electric field for relaxor ferroelectric ceramics and polymers, and it was procured that the expression could be applicable. Moreover, the transition electric field of the electrostrictive strain versus electric field was proportional to 3/2，and inversely proportional to 1/2.

IntroductionFrequency-dependent dielectric response is one of the important properties of ferroelectrics, reflecting the polarization response to high-frequency electric fields. Polarizations are closely related to ferroelectric domain structures, such as single domain, which represents the region with homogeneous polarizations direction. Ferroelectrics usually possess complex domain structures with domain walls (DWs) separating adjacent homogeneously polarized domains. DWs have attracted much attention during the past two decades due to their properties and potential for device designing. The related issues include DW motion, nonvolatile memory, topological defects, enhanced susceptibility, enhanced quality factor, low dielectric loss, and others. (Ba0.8,Sr0.2)TiO3 (BST0.8) is a ferroelectric usually with multi-domain structures. Previous work identified two typical types of domain walls (DWs), i.e., 90° DWs and 180° DWs. The enhancement of dielectric response in systems with 90° DWs is now well understood, and the behavior of dielectric response in multi-domain systems with 180° DWs remains unclear. Therefore, gaining insights into how 180° DWs affect the dielectric response can clarify the effects in multidomain systems.MethodsWe performed molecular dynamics simulations using the ALFE-H code with the first-principles-based effective Hamiltonian method to study the BST0.8 system. All the calculations were performed in the NPT ensemble using the Evans-Hoover thermostat, and periodic boundary condition (PBC) along all three directions. To simulate the substrate, a uniform biaxial strain was fixed to the 1.55% in-plane strain. To analyze the multi-domain with different DWs, the simulations started with a self-constructed initial multi-domain polarization configuration. Subsequent 50 ps MD simulation was performed under chosen strains for structural relaxation. In the initial configuration, the magnitude of non-zero components of soft mode on each site was set to 0.1 , atomic occupations (alloying) were randomized, and unless otherwise specified, all other mode variables were set to zero. The trajectory of local mode averaged over the supercell during MD simulations was extracted to calculate the dielectric response. The 8 ns MD simulations were performed to obtain an autocorrelation function for any time t ranging from 0 to 1 ns by one step of 10 fs. The fast Fourier transformation (FFT) was performed to calculate the dielectric response. Then two uncoupled damped harmonic oscillators (DHOs) were used to fit the data of dielectric response.Results and discussionThe dielectric response of single domain at 300 K with the different electric fields along [110] from 0 to 2 MV/cm was computed. The computational results can be well fitted with the model of two uncoupled DHOs. The real and imaginary parts of the predicted dielectric response at each chosen electric field both exhibit two peaks. As the electric field increases, the low-frequency mode with 300 GHz at zero field in the system gradually disappears, and a high-frequency mode of larger than 8 THz appears when electric field is larger than 1 MV/cm. The high frequencies modes of 3 THz at zero filed and 8 THz under 1 MV/cm shift towards higher frequencies as the electric field increases. In other words, the present simulations reveal that it is possible to manipulate the frequency of peaks in dielectric response via changing the magnitude of the external electric field.The dielectric responses of multi-domain with 90° DWs and 180° DWs are further analyzed. According to the experimental PFM results, the multi-domain structures are simulated and the dielectric response through MD simulations is calculated. The analysis of the dielectric response of single domain structure and multi-domain structures shows that the single domain structures exhibit high-frequency peaks at > 300 GHz, whereas the multi-domain structures exhibit low-frequency peaks at 8 GHz and 120 GHz for 180° DWs system and at 10 GHz for 90° DWs system, revealing that there exists a low-frequency mode related to collective oscillation of DWs in multi-domain structures. In addition, the frequencies of peaks in multi-domain with DWs are in a gigahertz range, whereas the single domain structure exhibits peaks in a terahertz range. The contribution of DWs to the dielectric response primarily arises from the timescale of DWs motion, such as sliding or breathing, which differs significantly from the high-frequency vibrations of optical phonon modes. The vibrational frequency of DWs is much lower, with relaxation times in the order of nanoseconds, resulting in a response frequency in GHz range, which is far below the terahertz range of optical phonon modes. Therefore, DWs oscillations dominate the dielectric response at a low frequency. Moreover, multi-domain structure with 180° DWs exhibits a unique low frequency mode at 120 GHz, which is significantly different from single domain and 90° DWs system. In other words, multi-domain structures with 180° DWs and 90° DWs exhibit different dielectric responses. There exists a common low-frequency mode related to the oscillations of DWs in BST0.8.ConclusionsIt was possible to manipulate the frequency of peaks in dielectric response of single domain through changing the magnitude of the external electric field. Domain walls oscillations dominated the dielectric response in a low frequency gigahertz range, whereas the single domain structures exhibited resonant peaks in a terahertz range. Moreover, multi-domain structures with different domain walls in BST0.8 had different dielectric responses, but the both have a same low-frequency mode at 10 GHz related to the domain walls oscillations. The results of this study indicated the dielectric response behaviors of ferroelectrics induced in an external electric field and internal multi-domain configurations and provided the potential mechanisms and guidance for optimizing application performance.

IntroductionMulti-layer ceramic capacitor (MLCC) is a highly promising dielectric capacitor due to its high capacitance, favorable dielectric temperature stability, and high breakdown strength. The conventional MLCC uses expensive silver and palladium metals as inner electrode materials. In recent years, base metals such as nickel or copper are employed as inner electrodes to reduce the related costs. However, these metals must be co-fired with the dielectric layer in a reducing atmosphere. BaTiO3 ceramics are the principal dielectric materials used in the manufacture of MLCC. The sintering of BaTiO3 ceramics in a reducing atmosphere generates a significant number of oxygen vacancies and free electrons, which significantly impair the insulation performance and reduce the insulating properties. Also, the pronounced dielectric peak of BaTiO3 ceramics occurs at the elevated temperature end (i.e., 125℃), leading to an inadequate temperature stability.Recent research focus on the relaxation properties and dielectric properties of BaTiO3-Bi(Me)O3(Me = Fe3+, Co3+, Y3+, etc.) solid solutions. The co-substitution of the A and B sites via introducing cations with different ion radius and valence states is shown to disrupt the long-range ferroelectric ordering, induce diffusive phase transitions, and broaden the dielectric temperature profile. Furthermore, Zn2+, Mg2+, Mn4+, Co3+, and other rare-earth element ions are introduced into BaTiO3-based ceramics as host-accepted dopants with the objective of forming defective dipoles with oxygen vacancies, which can bind the free charge movement. Among these, multivalent ions can bind electrons, thereby modifying the valence and regulating the electron concentration within the ceramics. This reduces the likelihood of conversion of Ti4+ to Ti3+.In this paper, 0.95BaTiO3-0.05BiCoO3 dielectric ceramics were used as a matrix. A small amount of MnO2 was selected for doping and modification, and subsequently sintered in a reducing atmosphere. The structure, micro-morphology, dielectric, and insulating properties were characterized, and the mechanism of anti-reduction was investigated.MethodsIn the preparation of the ceramics with a composition of 0.95BaTiO3–0.05BiCoO3–x% (in mass) MnO2 (x=0.1, 0.2, 0.3, 0.4, and 0.5), the materials were sintered at 1275 ℃ for 2 h under a reducing atmosphere (99.5%N2+0.5%H2, in volume fraction), forming dense ceramic samples. Subsequently, the ceramic samples were thinned and polished, and silver paste electrodes were applied to the both sides to measure their electrical properties.The physical structure of the samples was analyzed using a model X'Pert Pro X-ray diffractometer (XRD, PANalytical Co., the Netherlands). The microscopic morphology of the sample section was determined by a model 450 FEG field emission scanning electron microscope (FE-SEM, FEI Quanta Co., USA). The dielectric properties of the samples were measured by a model PolyK PK-CPT1705 low temperature dielectric tester. The resistivity of the ceramic samples was tested by a model Concept 80 low frequency module (Novocontrol Co., Germany). The ferroelectric properties of the samples were analyzed by a model PK-CPE1801 ferroelectric polarisation return and dielectric breakdown test system (PolyK Co., USA). The elemental valence analysis was conducted by a model 250Xi X-ray photoelectron spectrometer (XPS, EscaLab Co., USA). The thermally stimulated depolarization currents (TSDC) of the samples were carried out by a model Concept 80 TSDC (Novocontrol Co., Germany).Results and discussionAll the ceramic samples exhibit a pseudo-cubic perovskite structure. The crystal surface spacing of the ceramics increases as the amount of MnO2 doping increases. This indicates that Mn4+ is reduced to Mn3+ or Mn2+ during sintering in a reducing atmosphere. The radius of Mn3+ or Mn2+ is larger than that of Ti4+ (i.e., 0.060 5 nm, CN=6) ionic radius, leading to the crystal lattice distortion and increased cell volume. The cross section of the ceramics reveals fine grains (i.e., 0.24–0.32 m), clear grain boundaries, and a close arrangement between grains. The change in valence of Mn effectively inhibits the concentration of free electrons. When x = 0.5, the insulation resistivity increases to 1.84×1013 ·cm, the breakdown voltage reaches 210 kV/cm, thus improving the anti-reduction performance of ceramics. The introduction of Mn reduces the dielectric loss of the ceramics and improves their dielectric temperature stability. The optimum dielectric properties are achieved when x = 0.5 (i.e., a dielectric constant of 2354, a dielectric loss of 0.0075, and C/C(25℃)≤±15% at –55 to 154 ℃). The XPS spectra reveal the presence of both divalent and trivalent Co ions in the ceramics. The result of TSDC test shows a TSDC peak associated with trapped charge, which is related to the change in valence states of the variable valency acceptor ions Co and Mn. Two TSDC peaks are also related to defect dipoles, which are formed by Co and Mn with oxygen vacancies in the ceramic sample.ConclusionsThe effect of MnO2 doping on the phase structure, microstructure, and electrical properties of 0.95BT–0.05BC ceramics was investigated. The results demonstrated that Mn4+ ions underwent a conversion to Mn3+ or Mn2+ ions during sintering process in a reducing atmosphere, having the larger ionic radius. This transformation led to an expansion of the cell volume. Mn doping significantly impacted the broadening of the Curie peak and enhanced the dielectric temperature stability of 0.95BT–0.05BC ceramics. Specifically, as x = 0.5, the dielectric constant at 25 ℃ was 2354, and the dielectric loss was 0.007 5. Furthermore, at –55 to 154 ℃, C/C(25℃)≤ ±15%. The resistivity at room temperature increased to 1.84×1013 ·cm, and the breakdown voltage increased to 210 kV/cm. The XPS spectra revealed the presence of both divalent and trivalent Co ions in the ceramic, indicating that Co ions could capture free electrons and reduce their valence state. The result of TSDC test further demonstrated the existence of two types of defective dipoles, which were associated with the acceptor substitution of Co and Mn. The combination of these two factors reduced the carrier concentration in the ceramics, thereby hindering the migration of oxygen vacancies and free electrons and improving the anti-reduction performance.

IntroductionCement industry continues to face significant environmental challenges due to its substantial CO2 emissions, primarily stemming from carbonate decomposition in raw materials and fossil fuel combustion during production. Carbonatable binders and carbonation curing technologies have a promising potential for CO2 reduction in cement production, and the existing methods involve the calcination of calcareous and siliceous raw materials to produce low-lime silicates for their emission reduction limits. Natural wollastonite, primarily composed of calcium silicate (CS), has an attractive alternative due to its abundant reserves and ease of extraction. Utilizing natural wollastonite as a carbonatable binder can eliminate both fossil fuel combustion and the need for calcareous raw materials, thereby avoiding the carbonate decomposition entirely. However, its application is currently limited due to its low carbonation reactivity. This study was to investigate the enhancement of carbonation reactivity of natural wollastonite via optimizing calcination activation to develop a high-performance carbonatable binder.MethodsNatural wollastonite was calcinated under various calcination activation condictions (i.e., temperature, time, and cooling regimes). The activated wollastonite powder was mixed with water and compressed at 4 MPa to form the samples with different two dimensions (i.e., 20 mm × 20 mm × 20 mm and 20 mm × 80 mm × 20 mm). These samples underwent carbonation curing for 24 h before characterization. The mechanical properties were evaluated through compressive and flexural strength tests. The phase composition was characterized by X-ray diffraction (XRD) on both the wollastonite powder and carbonated samples before and after calcination activation. The structure of the samples was determined by Fourier-transform infrared spectroscopy (FT-IR) to analyze both the calcination activation effects and the resulting carbonation products. The microstructural evolution was determined by scanning electron microscopy (SEM). CO2 uptakes were analyzed by thermogravimetric analysis (TGA).Results and discussionThis study reveals that calcination activation significantly enhances the carbonation hardening performance of natural wollastonite via increasing carbonation reactivity. Optimal conditions are identified at a calcination temperature of 1200 ℃, effectively transforming Para-CS and -CS into metastable -CS. The effective calcination time for the complete transformation is 60 min, while preserving the material valuable fibrous structure. Water quenching is found to be the most effective cooling method, retaining higher concentrations of crystal defects in the metastable -CS. The transformation from a chain structure to a ternary ring structure at 1200 ℃ results in an increased structural distortion, thus leading to an enhanced carbonation reactivity of -CS. The carbonation analysis reveals the extensive formation of carbonation products, primarily calcium carbonate and highly polymerized silica gel, resulting in a dense microstructure. Note that the preservation of the fibrous structure through calcination activation allows unreacted wollastonite to serve as a fibrous reinforcement within the carbonation hardening system, contributing to enhanced mechanical properties.ConclusionsThe results showed that calcination temperature exhibited a dominant effect, followed by calcination time and cooling regime. The optimal calcination temperature of 1200 ℃ led to the improvements in compressive strength, flexural strength, and CO2 uptakes of the carbonated samples. The calcination time was critical in ensuring a complete phase transformation, while maintaining structural integrity, and a rapid cooling could preserve crystal defects that enhanced carbonation reactivity. The resulting carbonatable binder developed a carbonation hardening system comprising calcium carbonate, highly polymerized silica gel, and unreacted wollastonite fibers that could serve as an effective reinforcement within the carbonation hardening system. These findings could provide a promising pathway for the large-scale utilization of natural wollastonite in sustainable construction materials, having a potential for reducing the environmental impact of cement industry.

IntroductionThermal barrier coatings (TBCs), primarily composed of yttria-stabilized zirconia (YSZ), are critical materials for protecting the hot section components of gas turbine engines due to their high fracture toughness, melting point, and thermal stability. However, increasing operational temperatures has some challenges, particularly from molten CMAS (CaO–MgO–Al2O3–SiO2) deposits that can penetrate TBCs, reduce their strain tolerance, and accelerate failure. The performance of YSZ deteriorates upon interaction with CMAS, leading to phase transformation and coating detachment. Research efforts focus on modifying YSZ with TiO2, Cr2O3, or Al2O3 based on sacrificial protection strategies, enhancing resistance to CMAS via promoting the crystallization of compounds such as gehlenite (CaAl2Si2O8). This study was to investigate the effect of Ti on the reactivity of sacrificial protective Al2O3 coatings with CMAS. Al and Ti layers were deposited on YSZ by magnetron sputtering and heat treatment to form mixed oxide layers to improve the protective efficacy against CMAS attack.MethodsIn this study, graphite was utilized as a substrate for YSZ coatings. After degreasing and cleaning, graphite substrates were roughened with alumina sand (WA22). A YSZ powder (KF231) was sprayed onto the substrates in Ar-H2 as a plasma gasby a model Multicoat APS system (Oerlikon Metco Co., Switzerland) to deposit a YSZ coating with the thickness of 1500 m. The graphite substrate was removed by heat treatment at 750 ℃ for 3 h. Subsequently, Al and Ti films were deposited on the YSZ surface via magnetron sputtering, and followed via annealing at 1000 ℃ for 2 h to form TiO2–Al2O3 modified YSZ (TA–YSZ) coatings. An additional control group of Al2O3-modified YSZ (A–YSZ) without Ti deposition was also prepared.CMAS powders with a composition of 33%CaO–9%MgO–13%AlO1.5–45%SiO2 (in mole fraction) were synthesized via melt quenching. The melting and crystallization temperatures of CMAS and its mixtures with Al2O3 and TiO2 were characterized by differential scanning calorimetry (DSC). The thermal corrosion test was performed by painting the coated surface with a CMAS slurry and exposing it to 1300 ℃ for 5 min or 1250 ℃ for 2 h and 10 h (after 2 h heat treatment, the temperature was lowered in a furnace and then raised to 1250 ℃ for 2 h, and this was repeated for 5 times). The phase structure was determined by X-ray diffraction (XRD, Rigaku Smart Lab., Japan). The surface and cross-sectional microstructures of the corroded samples were analyzed by a model Phenom ProX field emission scanning electron microscope (SEM, Thermo Scientific, Netherlands) equipped with an energy-dispersive X-ray spectroscopy (EDS) system. The corroded samples were embedded in epoxy resin, cut and diamond-polished before performing the SEM analysis of the coating cross-section. The phase structure of YSZ was determined by a model InVia-Reflex Fourier-transform Raman spectroscopy (Renishaw Co., UK).Results and discussionAl and Ti films are sequentially deposited on the surface of YSZ coatings via magnetron sputtering. Al film exhibits a typical physical deposition structure. After heat treatment, the Ti element shows a dotted band above Al layer and forms an oxide layer. The results of DSC analysis reveal that the crystallization and melting characteristic temperatures of the three powder samples (i.e., CMAS, CMAS+Al2O3, and CMAS+Al2O3+TiO2) shift towards lower temperatures in such an order. This indicates that the addition of Al2O3 and TiO2 effectively reduces the crystallization temperature of CMAS and affects its melting characteristics.TA–YSZ and A–YSZ coatings coated with CMAS and subjected to treatments at 1300 ℃ for 5 min or 1250 ℃ for up to 2 h and 10 h show that Ti promotes reactions between CMAS and Al2O3 to form high-melting-point crystalline phases such as gehlenite (CaAl2Si2O8) and spinel (MgAl2O4). The introduction of Ti significantly enhances the interfacial reaction kinetics between the coating and the CMAS melt. Ti facilitates the formation of high-melting-point crystalline phases. Also Ti accelerates the crystallization process of corrosion products, forming a dense reaction barrier layer that effectively consumes molten CMAS and inhibits its penetration into the YSZ substrate. Moreover, the time and spatial distribution characteristics of the reaction products can occur during corrosion testing due to the different reaction kinetics characteristics of high-melting-point crystalline phases. Under the sacrificial protection mechanism of the surface coating, no phase transformation appears in the YSZ substrate.ConclusionsYSZ coatings prepared by the atmospheric plasma spraying method were subjected to magnetron sputtering to deposit Al and TiAl films on their surface, which were then in-situ oxidized to form Al2O3 and TiO2–Al2O3 (TA–YSZ) layers. The synergistic effect of Ti and Al on the enhancement of the CMAS corrosion resistance of YSZ coatings was investigated. The results showed that Ti significantly accelerated the reaction between CMAS and Al2O3, leading to the formation of high-melting-point products such as gehlenite and spinel, and promoted the crystallization of corrosion products. These actions effectively hindered the penetration of molten CMAS into the YSZ coating interior.

IntroductionA photocatalytic material of Ag10Si4O13 is widely concerned due to its band structure suitable for visible light response and built-in electric field promoting photogenerated carrier separation. However, in the photocatalysis process, Ag+ in Ag10Si4O13 is easily reduced by photoelectrons and accumulates Ag elements, affecting the cycle stability of the photocatalytic system. Graphene oxide can effectively transfer photogenerated electrons, and amorphous SiO2 can inhibit ion diffusion leaching, which are expected to slow down the photoreduction of Ag10Si4O13. In this paper, x% SiO2/Ag10Si4O13/GO (SAG) composites with different amounts of amorphous SiO2 were prepared via adding graphene oxide (GO) and silica sol in the process of Ag10Si4O13 preparation by a sol-gel method. The effect of amorphous SiO2 addition on the crystal structure, morphology and photoelectrochemical properties was investigated. Methylene blue (MB) was used as a target. The results show that a small amount of amorphous SiO2 can greatly improve the photoreduction resistance of Ag10Si4O13 photocatalysts, and broaden the photoresponse range and increase the adsorption capacity of the material for MB. However, the addition of amorphous SiO2 can simultaneously reduce the migration rate of photogenerated carriers and accelerate the recombination of photogenerated electron-hole pairs. Therefore, the photocatalytic degradation efficiency of the composite system firstly increases and then decreases with the increase of SiO2 addition. The degradation rate of SiO2/Ag10Si4O13/GO (SAG-1) samples can reach 99% under visible light for 40 min, showing a good cyclic stability. After 5 cycles, the crystal structure of the materials becomes stable and the degradation rate can still reach 97%.MethodsIn pre-hydrolysis of ethyl orthosilicate: 0.47 g of citric acid was added to 26 mL of anhydrous ethanol under magnetic stirring until complete dissolution, and then 5 mL of ethyl orthosilicate and 1.65 mL of distilled water were added in dropwise to the solution and the solution was stirred for 1 h and sealed for 1 week aging to obtain the silica sol. In preparation of Ag10Si4O13/GO, 2.1 mL of silica sol and GO sol containing 1%(in mass) Ag10Si4O13 were added to 10 mL of distilled water under ultrasound and magnetic stirring. 1.6 mL of 0.38 g/mL AgNO3 solution was added in dropwise to the mixed solution above under stirring for 1 h and aging for 24 h to obtain the precursor sol. The precursor sol was dried, ground and heat-treated at 400 °C for 5 h to obtain a brick-red Ag10Si4O13/GO powder=(i.e., AG). In preparation of SiO2/Ag10Si4O13/GO, x%(in mass) Ag10Si4O13 (x=1, 2, 3, 4) of silica sol was added, stirred continuously for 1 h and aged for 24 h to obtain x% SiO2/Ag10Si4O13/GO precursor sol. The subsequent steps were the same as for AG to obtain x% SiO2/Ag10Si4O13/GO powders (i.e., SAG-1, SAG-2, SAG-3, and SAG-4), respectively.The crystal structure of the materials was characterized by a model D8/axs X-ray diffractometer (XRD, Bruker Co., Germany). The microscopic morphology of the materials was observed by a model JSM-6701F transmission electron microscope (TEM, Electro-Optics Co., Japan). The chemical state was analyzed by a model PHI5702 X-ray photoelectron spectrometer (XPS, Physical Electronics Inc., USA). The photoresponsive properties were determined by a model U-3900H diffuse reflectance spectrometer (UV-Vis, Hitachi High-Technologies Co., Japan). The fluorescence emission spectra of the materials were determined by a model F97 fluorescence spectrophotometer (FSP, Shanghai Prismatic Technology Co., Ltd., China). The impedance, transient photocurrent, and Mott-Schottky curves were determined by a model CHI660D electrochemical workstation (Shanghai Tatsuwa Instrumentation Co., Ltd., China).The sample of 0.1 g was weighed and stirred in 100 mL of 20 mg/L MB solution, which was protected from light. For every 10 min, 3 mL of the sample was removed and the supernatant was centrifuged to determine the concentration of MB by a UV spectrophotometer. The adsorption performance was evaluated by a ratio of the concentration of the solution after adsorption, ct, to the initial concentration of MB, c0, at time t (ct/c0). After dark adsorption equilibrium, a 300 W xenon lamp was used as a light source, a filter (>420 nm) was added to simulate visible light, the distance from the light source to the liquid surface was fixed at 15 cm, the sample of 3 mL was taken for every 5 min, and the supernatant was centrifuged to determine the concentration of MB by a UV spectrophotometer. The photocatalytic performance was evaluated by a ratio of the concentration of the degraded solution to the initial concentration of MB, c0, at the time of t' (ct/c0).Five cycling experiments were performed on AG and SAG-1 to evaluate the stability of the photocatalytic properties of the materials. The stability of the crystal structure of the materials was evaluated based on the XRD patterns after each cycle.Results and discussionThe XRD patterns show that the introduction of SiO2 has no effect on the crystal structure of Ag10Si4O13. The TEM images of SAG-1 indicate the presence of graphene oxide, Ag10Si4O13 and SiO2. The UV-Vis diffuse reflectance spectra show that the edge of the absorption band of SAG-1 extended to 660 nm, which can be caused by the defects formed during the preparation of the material, but the forbidden bandwidth changes slightly, indicating that the protective layer of SiO2 does not cause any significant changes in the band gap. The PL spectra show that the fluorescence intensity of the material increases gradually with the increase in the content of amorphous SiO2, showing that the photogenerated carrier complexation rate increases with the increase of SiO2 incorporation. The photocatalytic performance of Ag10Si4O13 decreases significantly from 99% to about 90% after 5 cycles, while the photocatalytic performance of SAG-1 does not change after 5 cycles and still reaches 97%. The results show that the introduction of a small amount of SiO2 can slow down the reduction production of Ag+ during the photocatalytic degradation of Ag10Si4O13, and the crystal structure remains stable after some cycles, which greatly improves the stability of the material structure.ConclusionsSiO2/Ag10Si4O13/GO ternary composites with different SiO2 composite amounts were synthesized by a sol-gel method with GO and excess silica sol in an one-pot method. The complexing agent citric acid was used as a "bridge" to allow Ag10Si4O13 nanoparticles to grow uniformly on GO, and the excess silica sol was attached to the precursor surface, forming an amorphous SiO2 protective layer after heat treatment. The introduction of the amorphous SiO2 protective layer did not affect the crystal structure of Ag10Si4O13, which could extend the visible-light response range of the composite photocatalyst and significantly improve the adsorption performance for cationic dyes. However, SiO2 could reduce the migration rate of photogenerated carriers and accelerate the assembly of photogenerated electron-hole pairs, affecting the photocatalytic performance of the material. An appropriate amount of amorphous SiO2 protective layer could improve the photoreduction resistance and cycling stability of Ag10Si4O13 photocatalysts, while maintaining the original photocatalytic activity. The degradation rate of the SAG-1 sample was still 97% after 5 cycles, and the crystal structure remained stable after 5 cycles (I(13ˉ2)/I(111)=1.16). This work could provide an reference for solving the problem that the materials of Ag-based efficient photocatalysts could be susceptible to photocorrosion during the photocatalytic process.

The exponential explosive growth in transistor density enhances chip performance, while introducing significant thermal management challenges based on the Moore law. Effective thermal management is critical to maintaining chip performance and preventing damage in high-power-density systems. Conventional passive cooling, such as heat sinks and thermal interface materials, struggle to address the transient and localized heat flux in modern chips. It is thus necessary to develop innovative cooling solutions beyond conventional techniques. Thermoelectric cooling (TEC), based on the Peltier effect, stands out due to its precise temperature control, rapid thermal response, and high reliability. Furthermore, its adaptability to complex and uneven thermal profiles renders it particularly effective in managing localized hotspots in high-performance chips. However, realizing the full potential of TEC requires the development on material design, device integration as well as system-level optimization.This review represents the foundational principles of thermoelectrics (TEs) and analyzes the theoretical formulations pertaining to maximum heat dissipation power and cooling coefficient. It highlights that chip-level TEC differs from conventional goals, prioritizing a high the dimensionless figure of merit and a low thermal conductivity. Instead, chip applications require a balance between the relatively high thermal conductivity and elevated power factor to facilitate efficient heat dissipation. This review evaluates mainstream TE materials, including bismuth telluride-based materials, selenide based-materials, and magnesium based-alloys, alongside promising emerging materials such as Heusler alloys and magnon-drag metals. Some fabrication strategies, including nanostructure design, doping, and interface engineering, are emphasized.Except for materials, device design is critical for chip thermal management. TEC encounters commercialization challenges due to the need for ultra-thin (i.e., <50 μm) film structures directly attached to chip hotspots. Optimizing the geometric dimensions of TE legs (i.e., n-type and p-type material width ratios) and modular layouts enables efficient localized cooling, while minimizing interfacial thermal resistance and electrical losses. Devices must also handle high heat flux densities (>1000 W/cm2) and provide millisecond-scale thermal responses. Two primary micro-TEC architectures including out-of-plane and in-plane TE devices are examined. Note that multi-stage module configurations can significantly improve cooling performance via reducing cold-side temperatures and enhancing heat extraction efficiency.For the encapsulation architectures, two-dimensional (2D) and three-dimensional (3D) architectures with system-level optimization of thin-film TECs are explored. Integrating TEC with liquid cooling or heat pipes address the constraints of single method, thus providing effective solutions for both hotspot cooling and uniform heat dissipation. For instance, TE modules can decrease the cold-end temperature in liquid cooling systems or provide localized cooling in heat pipe setups, thereby ensuring optimal thermal distribution. As chip designs increasingly trend towards higher integration and miniaturization, the scalability and compactness of TE modules become critical. Advances in nanotechnology, 3D integration, and composite materials propel the development of ultra-thin, high-power-density TE modules that seamlessly integrate into chip architectures. Innovations in cost-effective manufacturing and material reliability are also essential for advancing commercialization and long-term sustainability.Summary and ProspectsTEC technology offers precise temperature control, addressing challenges like localized hotspots and complex heat flows in high-power-density chips, which is a key technology for sustainable and high-efficiency thermal management. However, the existing researches still face significant challenges in achieving the ultra-thin designs and high-efficiency required for modern chip architectures. The ability to manage transient high heat flux while maintaining scalability and cost-efficiency remains a concern. Future efforts should focus on the balance of cost, durability, and TE properties of materials through nanoengineering, doping, and interface engineering to enhance their performance. In terms of device design, optimizing multi-stage architectures and achieving miniaturization alongside intelligent functionalities will enable more efficient heat dissipation. These improvements are indispensable for catering to the increasing demands of high-power chips. At the system level, the integration of TEC with liquid cooling or heat pipes will enhance overall thermal management efficiently. The incorporation of TE generation provides new opportunities for constructing sustainable and energy-efficient cooling systems. With continued innovation in materials preparation, device design, and system integration, TEC may pave a way for efficient, reliable, and scalable solutions for the next generation of electronic systems.

Flexible thermoelectric devices with the advantages of deformability, small size, portability, stability and reliability are widely studied, which can realize the combination of complex heat sources and maintain good thermoelectric conversion under dynamic deformation. Compared with conventional block thermoelectric devices composing of block thermoelectric materials and rigid components, flexible thermoelectric devices are more flexible in adapting to the geometry of the heat source and can collect more energy. Therefore, the development of flexible thermoelectric devices as a research direction with a great application potential in the field of thermoelectric technology has attracted much attention. Flexible thermoelectric devices can effectively adapt to the complex surface shape of the heat source in the real scene, as well as the special conditions caused by dynamic deformation, and meet the unique requirements of flexible electronic devices. As a result, flexible thermoelectric devices have a great potential for some applications, such as versatility, power for small portable devices and sensor applications.In this review, the performance evaluation indexes of flexible thermoelectric devices are introduced, and the key performance, interface factors and the important influence of flexible stability indexes on flexible thermoelectric devices are summarized. This review represents the existing relevant research and summarizes the progress in three areas, i.e., flexible substrate devices, low-dimensional flexible devices, and fiber thermoelectric devices. Flexible substrate device is a typical form with the superior output performance. In contrast to the application limitations imposed by the immutability of bulk thermoelectric materials, the flexible substrate device is initially flexible through a flexible design that combines flexible components with rigid thermoelectric legs. The flexibility and wearability of substrate devices are significantly enhanced via optimizing the adaptability of flexible components and employing systematic designs. Also, some studies focus on special application scenarios, and deeply explore the flexible endowing methods for such devices to achieve optimization and innovation of preparation methods. The advent of low-dimensional flexible devices effectively overcomes the inherent limitations of flexible substrates in improving flexibility, which is reflected in improving the applicability of flexible devices in complex, small and dynamic environments. The different application for different heat flow directions are clarified through the in-depth investigation of the design of external and internal heat flow transfer. In the case of thin films for out-of-plane heat flow transfer, although it is difficult to maintain the temperature difference in the film thickness direction, the films are thin and easy to process, thus having an advantage in small and micro-integrated applications. In the case of thin films for in-plane heat flow transfer, the direction of heat flow is designed to propagate in the in-plane direction, and the temperature difference can be maintained in a large lateral range for efficient conversion of heat energy. In the domain of fiber thermoelectric devices, there exist applications that depend directly on structural flexibility. These applications involve the coating or integration of thermoelectric materials directly onto a fiber substrate with highly deformable properties. There are examples of thermoelectric materials prepared in fiber form and optimized for thermoelectric efficiency by virtue of a unique textile architecture that combines the advantages of material flexibility and structural flexibility. These flexible thermoelectric materials retain a high degree of flexibility and permeability, and have the advantage of manufacturing cost-effective and reliable high-volume products through industrial processes at room temperature. This property makes the material show a great application potential in flexible electronic devices.This review systematically summarizes the related research progress and points out the synergistic enhancement of flexibility, stability, and output performance during the preparation and design optimization process. This provides a useful reference and guidance for the preparation and design optimization of flexible thermoelectric devices.Summary and prospectsAlthough flexible thermoelectric technology has made remarkable progress in waste heat power generation, wearable devices and smart textiles, having a great application potential, there are still a series of problems in practical application. Therefore, a more systematic and comprehensive study is needed. The output of flexible substrate thermoelectric devices depends on the heat harvesting area and temperature gradient, and can produce power densities ranging from a few microwatts per square centimeter to a few milliwatts per square centimeter. However, the interface thermal resistance of low-dimensional flexible devices will lead to a low conversion efficiency, and its preparation technology is still limited to a large-scale application. Major challenges for fiber thermoelectric devices include building larger temperature gradients in the direction of heat flow, designing scalable architectures, improving mechanical stability, and enhancing comfort. In addition, the development of advanced preparation technologies such as magnetron sputtering and pulsed laser deposition will also achieve a large-scale production. Also, the combination of boost and energy storage technologies, such as efficient DC-DC converters and new energy storage systems, will effectively solve the problem of low output voltage of thermoelectric devices, so that they can directly drive low-power sensors or communication modules. Finally, with the further development of interface stability and flexible design of flexible substrate thermoelectric devices, flexible thermoelectric devices will be more widely used in wearable technology, smart textiles and some related fields.

Rechargeable aluminum (Al) metal batteries (RAMBs) are considered as one of the most attractive alternatives to the existing battery technologies because of their low cost, high safety, high abundance of Al (the most earth-abundant metal element, over 3 700 times that of lithium), and the well-established Al industry. Furthermore, Al anode has a high theoretical gravimetric capacity of 2 981 mA·h··g&#x2212;1 and the highest theoretical volumetric capacity of 8056 mA·h·cm&#x2212;3, as well as a relatively low redox potential of &#x2212;1.66 V versus the standard hydrogen electrode, thus enabling RAMBs to have a potentially high energy density. The importance of reversible Al plating and stripping is particularly pronounced for RAMBs. However, it is still non-trivial to achieve high Al plating/stripping reversibility and superior lifespan, regardless of whether in aqueous, ionic liquid (IL) or organic electrolytes. Therefore, the reversibility and development of Al anodes await a breakthrough in addressing significant challenges, i.e., passivation, corrosion, and dendritic growth. Unfortunately, there is still a lack of systematic discussions on the strategies for achieving highly reversible and long-life Al anodes.This review discusses some challenges facing Al anodes in current electrolyte systems, analyzes failure mechanisms, and presents some advanced solutions. Specifically, in aqueous electrolytes, the passivation of Al by the formation of a thin yet dense Al2O3 layer on the Al surface hinders reversible stripping and plating. Moreover, the thermodynamic instability of Al in aqueous electrolytes induces and exacerbates the corrosion and hydrogen evolution side reactions during resting and cycling. To solve these problems, the structural optimization of Al anode and the development of high-performance modification layers can effectively alleviate side reactions and enhance cycling stability. Also, regulating electrolyte components and concentrations can facilitate the formation of a more efficient solid electrolyte interphase (SEI), thus preventing the performance degradation caused by hydrogen evolution and dendrite growth. Room-temperature ILs with high concentrations of AlCl&#x2083; salt can promote reversible Al plating and stripping under ambient conditions. However, their notorious corrosion nature causes continuous consumption of Al, and exacerbates the uneven nucleation and growth during deposition, thus resulting in Al dendrites. In addition to conventional strategies (i.e., anode structure modification, host design, and interfacial coating), Al alloying and control of preferential crystal orientation are proved as effective approaches to enhance the corrosion resistance of Al and reduce nucleation barriers. These approaches promote uniform deposition and mitigate dendrite growth. The incorporation of functional additives, such as leveling agents, into the electrolyte also offers an effective approach. These additives stabilize Al morphology and suppress corrosion. In the case of organic electrolytes, the high-volumetric-charge density and intense electron-withdrawing nature of Al3+ result in slow de-solvation kinetics and severe interface passivation, which represent major bottlenecks for Al anodes. To address these challenges, the introduction of chloride ions into organic electrolytes can enhance the reversibility of Al. Chloride ions partially disrupt the passivation layer and, more importantly, facilitate the formation of Al2Cl7&#x2212; active anions, thereby improving the de-solvation kinetics of Al3+ in organic electrolytes.Summary and ProspectsRAMBs are regarded as one of the most promising alternatives to the existing battery technologies for large-scale energy storage due to their low cost, high capacity, superior safety, and abundant availability of aluminum. However, RAMBs still encounter significant challenges, with the lack of highly reversible, long-lifetime Al anodes being a critical bottleneck. Although considerable progress has been made in this field, it remains in its early stages and requires further comprehensive investigation. A future research can focus on the following aspects, 1) Revealing the in-depth failure mechanisms and enhancing the stability of Al metal anodes. Whether through anode modification or electrolyte optimization, the underlying mechanisms must be thoroughly investigated, with particular attention to the failure processes and stability of the anode interface. In-situ characterization techniques are highly preferred to identify the fundamental causes of interface failure. Also, there is a need to refine some methods for evaluating anode stability and to establish standardized performance parameters and evaluation criteria. 2) Exploring high-performance, non-corrosive, and cost-effective RAMBs systems. The performance of RAMBs is highly depended on the properties of the electrolyte and its compatibility with Al anodes. To address the corrosion challenges posed by chloride-based ILs on both Al and other metallic parts, the design of non-corrosive electrolytes is essential for practical applications. However, employing such non-corrosive electrolytes necessitates overcoming an issue of natural Al2O3 passivation layer on Al anode. 3) Developing high-capacity and stable cathode materials. The extremely high charge density of Al3+ results in an intense Coulomb interaction with intercalation-type cathode materials, which slows down the intercalation/deintercalation kinetics of Al3+ and leads to an irreversible structural damage to the cathode. For conventional conversion-type cathodes, their intense bonding with Al3+ severely limits the reversibility of redox reactions. It is thus critical for advancing RAMBs to develop novel cathode materials that can efficiently utilize the three-electron-transfer mechanism of Al3+ and exhibit fast kinetics.In this review, we emphasize the critical importance of the fundamental failure mechanisms of Al anodes and clarify Al3+ solvation and de-solvation behaviors and breakthroughs in the design of pivotal materials such as non-corrosive and high-performance electrolytes and high-capacity cathodes. This review provides valuable insights for the development of next-generation RAMBs and contributes to accelerating their practical application.

Lithium-ion batteries are extensively utilized in portable electronic devices, electric vehicles, and large-scale energy storage due to their high energy density, long cycle life, and various other advantages. However, the limited natural abundance of lithium resources and their uneven geographical distribution imped the further development of the lithium-ion battery industry. Sodium-ion batteries have an application potential in large-scale energy storage due to their advantages such as abundant sodium resources, low cost, and compatibility with existing lithium-ion battery production lines. Replacing the conventional electrolyte with a solid electrolyte possessing flame retardant properties can effectively address the issues of thermal runaway and explosion associated with sodium-ion batteries. All solid-state sodium batteries (ASSB) offer benefits, including high energy density, enhanced safety, and low cost, aligning with the development goals of energy storage. As a critical component of ASSBs, the electrochemical properties of solid electrolytes play a pivotal role in determining their performance. Composite solid electrolytes (CSE), characterized by a good flexibility, high interfacial compatibility, and ease of processing, are considered as the most promising solid electrolytes for future large-scale commercial applications.This review summarizes recent research progress on organic-inorganic composite solid electrolytes (CSE) for ASSB and further analyzes the ion transport mechanisms within CSE. CSE are primarily composed of an organic polymer matrix, inorganic fillers, and sodium salts in specific ratios. The polar groups (i.e., &#x2212;S&#x2212;, C≡N, &#x2212;O&#x2212;, C=O) present in the polymer form group-ion complexes with the dissolved sodium salt. Within the amorphous regions, individual segments of the polymer chains exhibit relative freedom to rotate and bend, facilitating the movement of group-ion complexes within a limited spatial domain. As the chain segments reposition themselves appropriately, the group-ion complexes on these segments begin to segregate, allowing ions to interact with the functional groups of adjacent chain segments, thereby forming new group-ion complexes. This process is reiterated to facilitate ion transport. It is widely accepted that Na+ conduction predominantly occurs in the amorphous phase regions of the polymers above the glass transition temperature. In addition to these amorphous regions, certain crystalline domains exist within the polymers. In inorganic solid electrolytes, ionic hopping migration serves as the primary ion transport mechanism, which are largely affected by defects within the crystal lattice. The complexity of the interfacial region is further compounded in the presence of inorganic fillers that possess highly reactive surface defects, which readily interact with the polymer matrix. Two main explanations for ion transport in the interfacial region are proposed, i.e., 1) the interaction between the functional groups on the surfaces of the inorganic fillers and the polymer matrix, as well as the sodium salt, which weakens the interaction between the polymer matrix and Na+, resulting in a higher concentration of free Na+ on the filler surfaces and the formation of ion-transport channels, and 2) the space charge layer effect arises from the disparity in Na+ concentration between the inorganic filler and the polymer matrix, as well as the electrochemical potential difference. This results in the spontaneous formation of a Na+-rich space charge region, which serves as an efficient channel for Na+ transport. Solid electrolytes, including oxides and sulfides, are commonly utilized as active fillers, wherein the Na+ transport mechanisms encompass both vacancy and interstitial mechanisms.Based on their ionic conductivity, fillers are classified into inert fillers and active fillers. Inert fillers include Al2O3, ZnO, TiO2, SiO2, Y2O3, ZrO2, MgO, and BaTiO3, while active fillers encompass NASICON-type, calcite-type, and sulfide solid electrolytes. Fillers are further categorized by their shape and dimension into zero-dimensional (0D), one-dimensional (1D), two-dimensional (2D), and three-dimensional (3D) structures. Modifying the chemical properties of the filler surface enhances compatibility and interfacial bonding between the filler and the polymer matrix and allows for the modulation of ionic transport pathways within the electrolyte, significantly impacting overall electrochemical performance. Graft copolymers serve as interfacial modifiers, interacting with both the polymer and the filler to markedly improve interfacial bonding and enhance ionic conductivity. This strategy also mitigates the aggregation of fillers, which arises from the disparity in surface energy between the inorganic filler and the polymer matrix. Furthermore, the quantity of inorganic filler added plays a crucial role in ionic transport within CSE, making it essential to clarify the relationship between the amount of inorganic filler and ionic conductivity to effectively regulate ionic transport.Summary and prospectsComposite solid electrolytes (CSE), which consist of a polymer matrix combined with inorganic fillers, leverage the advantages of both solid polymer electrolytes (SPE) and inorganic solid electrolytes, offering promising prospects for practical applications. However, the development and implementation of CSE are still in their early stages, revealing a significant gap between their existing performance and application requirements, necessitating further enhancements. To address this, CSE with specific morphologies and microstructures must be designed to improve the connectivity of ion migration channels, facilitating efficient ion transport. An effective interfacial conductive network can be established via carefully controlling the ratio, size, dispersion, and other characteristics of the inorganic fillers, significantly enhancing ionic conductivity. The strategic application of simulation and characterization techniques to elucidate the ion transport mechanisms within CSE is crucial for advancing the performance of these materials. Lastly, to lower the costs associated with ASSB, the fabrication process of CSE should be streamlined to be simple, time-efficient, and material-saving, while also being compatible with the existing battery production lines to enhance the practicality.

As a fundamental physical parameter reflecting the microscopic motion within matter, the measurement of temperature holds a critical significance in industrial production, biomedical fields, and aerospace engineering. Conventional contact-based thermometry, requiring direct interaction with measured objects, demonstrates limitations in extreme environments due to the susceptibility to external interference and inherent measurement perturbations, thus failing to meet modern technological demands for high sensitivity, non-invasiveness, and rapid response. These challenges propel the development of non-contact thermometric technologies featuring enhanced measurement sensitivity, robust anti-interference capability, non-invasive characteristics, and fast response. Fluorescence-based temperature sensing has emerged as an innovative solution via leveraging the distinctive energy level configurations of rare-earth ions and their temperature-dependent luminescent properties. This technique has attracted substantial research attention due to its rapid response kinetics, exceptional sensitivity, and remarkable adaptability to harsh operational environments.Fluorescence thermometry based on the correlation between fluorescence intensity and temperature represents one of the earliest developed methods in temperature measurement technology. This approach primarily falls into two categories, i.e., single-energy-level fluorescence intensity thermometry and fluorescence intensity ratio (FIR) thermometry. The single-energy-level method, as the most straightforward technique, determines temperature via monitoring the intensity variation of a specific emission peak with temperature changes, directly demonstrating the relationship between fluorescence intensity and temperature. However, its applications reduce in recent years due to the inherent limitations such as susceptibility to fluorescence loss during detection and strong dependence on excitation light intensity, hindering precise control in measurement processes.In contrast, FIR thermometry has attracted much attention due to its insensitivity to external disturbances, high measurement accuracy in complex environments, and excellent reproducibility. This technique is further divided into thermally coupled energy level FIR thermometry and non-thermally coupled energy level FIR thermometry. The former operates based on the principle that particles in thermally coupled energy levels (denoted as Level 1 and Level 2) reach a thermal equilibrium within a short timeframe, where the population redistribution between these levels induces measurable changes in fluorescence intensity ratio. While a larger energy level spacing generally enhances temperature sensing performance, the inherent limitation of thermal coupling energy gaps poses some challenges for improving relative sensitivity. To address this constraint, non-thermally coupled FIR thermometry through co-doping dual luminescent centers is porposed. This strategy enhances relative sensitivity and measurement reliability via leveraging the distinct temperature-dependent luminescence characteristics of two different centers. The intensity ratio between dual centers provides more accurate temperature determination through mutual calibration and effectively compensates for measurement errors caused by external factors. Common dual-center configurations include rare-earth ion pairs and rare-earth/transition metal ion combinations, showing significant improvements in relative sensitivity for advanced thermometric applications.The application of rare-earth-doped fluorescence intensity-based thermometry focuses on two primary domains, i.e., temperature measurement within the physiological range and in ultra-high-temperature environments. Distinct requirements for the upper temperature limit and sensitivity arise across these applications. For instance, in physiological temperature monitoring, slight temperature variations can lead to significant biological effects, necessitating exceptionally high relative sensitivity in low-temperature regimes to ensure sufficient measurement accuracy. Conversely, in high-temperature scenarios, the primary objective is to achieve an upper measurement limit compatible with extreme thermal conditions. Consequently, the design and development of novel materials should be closely aligned with specific application contexts, emphasizing performance optimization tailored to operational demands, such as tunable temperature thresholds, enhanced thermal stability, and environment-specific signal responsivity. This approach ensures that material systems balance sensitivity, durability, and temperature range adaptability for targeted technological implementations.Summary and ProspectsThe development of fluorescence intensity-based thermometry faces two major challenges, i.e., enhancement of temperature measurement performance and expansion of application domains. Regarding performance improvement, the existing fluorescence intensity thermometric methods primarily focus on fluorescence intensity ratio (FIR) techniques involving thermally coupled and non-thermally coupled energy levels. However, several critical issues have emerged.Firstly, rare-earth ions and certain transition metal ions exhibit significant thermal quenching effects at elevated temperatures, thus leading to substantial attenuation of fluorescence signals in high-temperature environments. Secondly, in thermally coupled energy level-based FIR thermometry, despite the abundant energy levels of rare-earth ions, some studies on potential thermally coupled energy levels for specific rare-earth ions remain insufficient. Moreover, systematic investigations into performance variations among different thermally coupled energy level pairs for FIR thermometry are notably lack. Thirdly, the fundamental mechanisms underlying non-thermally coupled FIR thermometry, particularly the operational principles of dual-luminescent-center systems, have yet to establish universally accepted theoretical explanations. The energy transfer processes and interaction mechanisms between different luminescent centers’ energy levels require a further elucidation. In terms of luminescent center selection, more combinations such as rare-earth–rare-earth ion pairs and rare-earth–transition metal ion systems warrant a comprehensive exploration. Lastly, although numerous temperature-sensitive materials demonstrate either high upper measurement limits or superior relative sensitivity, materials simultaneously with a high relative sensitivity in a broad temperature range remain scarce. Therefore, the design and development of novel temperature-sensitive materials will constitute a crucial research frontier.Concerning application expansion, fluorescence intensity thermometry has attracted considerable attention due to its unique characteristics. The existing thermometric materials exhibit diverse advantages, each possesses inherent limitations. A critical challenge lies in rationally designing application scenarios based on material properties to maximize their functional advantages for practical implementations. Current research predominantly focuses on material development itself, with insufficient exploration of potential application fields. Note that raw materials cannot be directly employed in industrial production but require integration into temperature sensor architectures coupled with complete testing systems. However, the research framework spanning from material development to practical applications, encompassing complete testing systems and material optimization, remains underdeveloped. Furthermore, although applications have extended to biomedical and aerospace fields, investigations under extreme conditions such as ultra-low temperatures and highly corrosive environments remain inadequate.It is anticipated that these challenges above will be progressively addressed with sustained research efforts. The development of advanced temperature-sensitive materials with enhanced performance and their subsequent integration into industrial applications are expected to drive significant advancements in fluorescence intensity-based thermometry technology.

IntroductionRoom-temperature sodium-sulfur (RT Na-S) batteries, as one of effective candidates for next-generation high-energy-density battery systems, have the advantages of high theoretical energy density (i.e., 1274 W·h·kg–1), high elemental abundance (i.e., S and Na) and low cost. However, the practical application of RT Na-S batteries is restricted due to the poor electronic conductivity of sulfur, sluggish reaction kinetics, and sodium polysulfides (NaPSs) shuttle effect. The related studies are performed on cathode materials of RT Na-S batteries. Among various materials, two-dimensional layered materials have attracted extensive attention due to their unique structures and physicochemical properties, showing a substantial promise for applications. Compared to conventional experimental research methods, first-principles computational techniques can assist in designing novel high-performance electrode materials in atomic and electronic scales. This paper was thus to investigate g-C3N4 as a catalyst for RT Na-S batteries based on the first-principles calculation methods. In addition, the chemical interactions between g-C3N4 and NaPSs, electronic structure, and reaction energy barriers were also analyzed.MethodsDensity functional theory (DFT) calculations were carried out by a software named Vienna ab initio simulation package (VASP). The exchange-correlation energy was described using the Perdew-Burke-Ernzerhof (PBE) functional within the framework of the generalized gradient approximation (GGA). A plane-wave cutoff energy of 500 eV was chosen. A vacuum layer larger than 20 was used in the calculations to avoid the interlayer interactions. The convergence of energy and force criteria on the atoms were set to be 10–5 eV and 0.02 eV·–1, respectively. The DFT-D3 method was used to calculate the long-range van der Waals interactions. The 3×3×1 and 4×4×1 Monkhorst-Pack K-points were set in the first Brillouin zone for geometric optimization and calculation of density of states (DOS), respectively. The adsorption energy of NaPSs adsorbed on g-C3N4 monolayer was calculated byEads=Eg-C3N4/NaPSs&#x2212;Eg-C3N4&#x2212;ENaPSs(1)where Eg-C3N4/NaPSs,Eg-C3N4&#x00A0;and&#x00A0;Eg-C3N4 are the calculated total energies of g-C3N4 monolayer with adsorption of NaPSs, g-C3N4 before adsorption, and isolated NaPSs. The differential charge density was calculated by=g-C3N4/NaPSs&#x2212;g-C3N4&#x2212;NaPSs(2)where g-C3N4/NaPSs，g-C3N4&#x00A0;and&#x00A0;NaPSs are the calculated electron densities of g-C3N4 monolayer with adsorption of NaPSs, g-C3N4 monolayer, and isolated NaPSs.Results and discussionThe g-C3N4 studied belongs to a hexagonal crystal system, where C and N atoms are bonded in sp2 hybridized configuration, forming a conjugated -electron structure. The density of states (DOS) of g-C3N4 monolayer shows semiconductor characteristics with a band gap of 1.12 eV. The adsorption energy is a fundamental criterion for assessing whether a material can anchor polysulfides. The adsorption energies of NaPSs and S8 on the g-C3N4 surface were calculated, and all the values range from –1.0 eV to –5.0 eV, indicating that g-C3N4 can be ideal candidates for anchoring NaPSs. The analysis of DOS and differential charge density of these adsorption systems shows that electron transfer occurs through Na-S and Na-N bonds and the band gap decreases, compared to the pristine g-C3N4 (except S8 adsorption system), facilitating electron transfer and providing electrons for the redox processes of NaPSs. The Gibbs free energy calculation for the entire discharge process reveals that the energy barrier of the rate-determining step is only 0.70 eV. These results emphasize a pivotal role played by the g-C3N4 in accelerating the conversion of NaPSs.ConclusionsBased on first-principles calculations, we systematically investigated the anchoring and catalytic behavior of g-C3N4 toward NaPSs. The results showed that g-C3N4 could have a great potential as a sulfur host and catalytic material for RT Na-S batteries. g-C3N4 had an anchoring effect on NaPSs, contributing to improved battery cycle life. g-C3N4 could effectively capture NaPSs from the electrolytes, suppressing the shuttle effect. And g-C3N4 accelerated the conversion kinetics of NaPSs, enhancing sulfur utilization. All these findings could underscore an immense potential of g-C3N4 in the design of RT Na-S battery cathodes.

IntroductionPorous silicon-based anode materials with their high theoretical specific capacity (i.e., approximately 4200 mA·h·g&#x207B;1) and unique porous structure have a promising application potential in the field of lithium-ion batteries. Their porous structure provides effective buffering space for the significant volume change (i.e., 300%) of silicon during charging and discharging, effectively mitigating material pulverization and maintaining the integrity of the electrode structure, thereby enhancing battery cycle stability and Coulombic efficiency, and greatly increases the utilization rate and conductivity of active materials, thus enhancing the energy density of lithium-ion batteries. However, despite the enormous application prospects of porous silicon-based anode materials, their high preparation costs, low initial Coulombic efficiency, and capacity fade after long-term cycling remain some critical factors restricting their commercialization.To address these issues, recent studies mainly focus on nanostructure design, composite material development, electrolyte optimization, and large-scale preparation techniques. The emergence of silicon-carbon composite materials, which embed silicon nanoparticles into a porous carbon matrix, can construct a good conductive network, effectively improving material conductivity, and further enhance the buffering capacity for silicon volume changes, significantly improving battery cycle stability and rate performance. In this paper, a porous silicon-carbon layered silicon-based composite was prepared as an anode material. In addition, the electrochemical performance of this material was also investigated.MethodIn the preparation process, silicon nanopowder was first uniformly dispersed with a mixed solvent of deionized water and anhydrous ethanol, and then sonicated under ultrasound to ensure uniform dispersion of silicon nanopowder. Subsequently, PVP (polyvinylpyrrolidone) and glucose/sucrose were added to the dispersed suspension as a pore-forming agent and two carbon sources, and stirred at room temperature to form a uniform precursor sol. The two precursor sols, i.e., Si/PVP/glucose and Si/PVP/sucrose, were placed into the ink cartridges of a microelectronic printer, and the gels were uniformly printed onto copper foil by a 3D printing technology. Finally, the carbon-containing substances in the electrode were completely reduced to amorphous carbon after vacuum drying and carbonization treatment. Note that different carbon sources (i.e., glucose and sucrose) released different amounts of gas during the carbonization process, leading to subtle changes in the electrode structure and subsequently affecting its electrochemical performance.The prepared electrode materials were analyzed by scanning electron microscopy (SEM), transmission electron microscopy (TEM) and other characterization techniques.Results and discussionThe results show that larger pores with a uniform pore distribution exist in Si@PVP/glucose structure, and the electrode exhibits minimal cracking after charging and discharging cycles, demonstrating the superior mechanical stability and resistance to expansive internal stresses. In contrast, the pore distribution of Si@PVP/sucrose is relatively uneven, and the electrode exhibits greater cracking after cycling, leading to a decrease in cycle stability and electrochemical performance. These results demonstrate that the Si@PVP/glucose structure electrode exhibits a greater ability to withstand expansive internal stresses and has a higher cycle retention rate. The electrochemical performance tests further validate the superiority of the Si@PVP/glucose composite material. After 100 cycles of charging and discharging, the specific capacity of the Si@PVP/glucose electrode reaches 1095 mA·h·g&#x207B;1, with a reversible capacity retention rate of 97.4%, while the specific capacity of the Si@PVP/sucrose electrode is only 970 mA·h·g&#x207B;1, with a reversible capacity retention rate of 91.0%. It is indicated that the Si@PVP/glucose composite material has significant advantages in optimizing the performance of porous silicon-based anode materials, particularly in improving cycle stability and rate performance.ConclusionsThis study prepared a porous silicon-carbon layered silicon-based composite (i.e., Si@PVP/glucose) as an anode material, and the performance was optimized. The results showed that the Si@PVP/glucose electrode material exhibited a greater ability to withstand expansive internal stresses and had a higher cycle retention rate. The electrochemical performance tests further validated the superiority of the Si@PVP/glucose material.It is indicated that this porous silicon-based anode material could have a promising application potential in the field of energy storage and conversion for high-energy-density and long-life batteries.

IntroductionThe inverted inorganic perovskite solar cells (IPSCs) with a bandgap of 1.7 eV are prospective candidates for the next generation of photovoltaic cells due to their elemental composition and unparalleled light and thermal stability. However, the higher defect state density and energy level mismatch at the interface result in an inferior photovoltaic performance to that of organic-inorganic hybrid perovskite solar cells (HPSCs) with the same bandgap. In this paper, tetrabutylammonium iodide (TBAI) was used to treat the surface of the optical absorption layer of the inverted device. The interface engineering technique can effectively reduce the density of defect states in the inverted device, suppress non-radiative recombination, and decrease the carrier accumulation at the interface. Finally, the photoelectric conversion efficiency (J) of IPSCs prepared in air is increased from 17.10% to 19.76%, and the PCE is increased by 15.56%. In addition, the stability of the device is also considerably improved.MethodThe NiOx solution was spun on the surface of FTO glass and annealed at 130 ℃ for 20 min. After cooling to room temperature, the P3CT-N solution was spun coated on the surface of the annealed NiOx and annealed in air at 130 ℃ for 10 min. Afterwards, the prepared perovskite precursor liquid was spun coated on the surface of P3CT-N preheated at 70 ℃，and annealing on a hot plate at 190 ℃ for 5–8 min. The annealed films were transferred to a N2 atmosphere glove box for dynamic spin coating of TBAI (IPA) solution, and then annealed at 80 ℃ for 5 min. MgF2 was thermally evaporated on the surface of TBAI by an evaporation equipment. PC61BM (CB) was spun as the electron transport layer, and BCP (IPA) was spun as the hole barrier layer. Finally, the evaporation device vaporized Ag with the thickness of 80.0 nm on the surface of the BCP as a metal electrode.Results and discussionThe SEM images show that the particles of TBAI treated films are fuller and the holes significantly reduce. These holes are the center of non-radiative recombination, seriously restricting the improvement of photovoltaic performance of devices. The results of UV-absorption test indicate that TBAI post-treatment does not change the bandgap of inorganic perovskite CsPbI2.85Br0.15. The results of TRPL test show that the TBAI-treated films have a longer carrier lifetime, indicating that the non-radiative recombination is inhibited. Based on the results of the Mott-Schottky test, the TBAI-treated device demonstrates a built-in electric field of 1.196 V, which is much higher than that of the reference group (i.e., 1.038 V), and a larger built-in electric field means a more intense carrier separation drive, corresponding to a higher VOC. The steady-state power output (SPO) efficiency of the TBAI treated device is tested for 300 s, and the average SPO efficiency of the device is 19.36% for over 400 s. A pure hole device is constructed to analyze the defect state density of the two films. The VTFL values of the control and TBAI treated devices are 1 V and 0.95 V, respectively, and the corresponding defect state densities are 1.68×1016 and 1.56×1016 cm–3, respectively. The TBAI post-treatment reduces the defect state density of the films, inhibiting non-radiative recombination. The humidity stability of the two films is monitored under specific conditions. The decomposition the films in the control group appears when the films in the control group are in humid air for 72 h, while that in the experimental group almost constancy.ConclusionsThe defect state density at the interface of IPSCs was high, and the non-radiation recombination loss was serious, thus restricting the further improvement of the photovoltaic performance of the device. In this paper, TBAI was used to post-treated the surface of CsPbI2.85Br0.15, which improved the crystal quality of the film, effectively reduced the density of defect states, inhibited non-radiative recombination, and reduced unnecessary carrier accumulation at the interface. IPSCs prepared in air treated by this interface engineering achieved a champion PCE of 19.76% with a VOC of 1.200 V. This work could provide a method for preparing highly efficient and stable IPSCs in air, thus promoting the practical process.

IntroductionDye-sensitized solar cells (DSSC), as the third generation of solar cells, have the advantages of low cost, environmental protection, pollution-free, and abundant raw materials. Conventional liquid-state electrolyte of DSSC restricts an absorption capacity of visible light and is corrosive, which is considered as an obstacle to the long-term use of DSSC. Solid-state electrolyte has attracted much attention because it overcomes the shortcomings of liquid-state electrolyte (i.e., volatile organic solvent, poor stability and easy leakage). TiO2 is often used as a photoanode material for DSSC, and its conduction band can match the energy level of common dyes to accelerate electron transfer. However. TiO2 has a low electron mobility, a low conductivity, a long-term ultraviolet irradiation and the existence of oxygen vacancy, which hinders the effective current output and reduces the photoelectric conversion efficiency. ZnO material can replace TiO2 as a photoanode because of its high electron mobility and transparency, and it is also a wide-band gap semiconductor material. To improve the photoelectric conversion efficiency of DSSC and solve the problem of liquid-state electrolyte leakage, ZnO solid-state dye-sensitized solar cells were simulated by a software named SCAPS-1D, and the photoelectric performance of solid-state DSCC was analyzed.MethodsBased on FTO/ZnO/N719 / PEDOT: PSS/Au structure modeling by the software SCAPS-1D, a solid-state DSSC model in each layer material interface lose band gap, dielectric constant, mobility and other parameters, and inside layer and layer between the defects was set up under AM1.5G 1 standard solar lighting at an incident light power of 100 mW/cm2. The material defect was set as a neutral defect to ensure that the defect could generate SRH recombination without generating space charge because the charge recombination mode of DSSC was mainly SRH. The Poisson equation, carrier continuity equation and carrier drift diffusion equation were numerically solved by the software of SCAPS-1D. The influences of material thickness, operating temperature, series parallel resistance and back contact metal type on the DSSC photoelectric parameters were simulated, and the internal mechanism of the influences of various factors on the DSSC was analyzed.Results and discussionThe open circuit voltage, short circuit current density and photoelectric conversion efficiency increase significantly, and the thickness continues to increase to 1000 nm, as well as the photoelectric parameters increase slowly when the thickness of the dye layer increases from 100 nm to 500 nm. The maximum photoelectric conversion efficiency can be obtained at the thickness of electron transport layer of 20 nm. The photoelectric conversion efficiency increases with the increase of hole transport layer thickness. The efficiency decreases with the increase of operating temperature, and the maximum photoelectric conversion efficiency is 15.17% at 300 K. The open circuit voltage and short circuit current density change little, and the filling factor and photoelectric conversion efficiency decrease when DSSC series resistance increases. The short circuit current density basically unchanges, and other parameters improve when the shunt resistance increases. The effect of changing the type of back contact metal on the photoelectric performance of DSSC is analyzed. The open circuit voltage, short circuit current density and photoelectric conversion efficiency all increase when the back contact metal work function value increases, and the maximum photoelectric conversion efficiency can be obtained when Au is used as the back contact metal.ConclusionsThe optimal thickness of dye layer was 500 nm, and the maximum photoelectric conversion efficiency could be obtained at the thickness of electron transport layer of 20 nm, and the photoelectric conversion efficiency increased with the thickness of hole transport layer. The maximum photoelectric conversion efficiency was 15.17% at the operating temperature of 300 K. The series parallel resistance had an important effect on the photoelectric performance of DSSC. The smaller the series resistance was, the larger the parallel resistance could be, for the improvement of the photoelectric conversion efficiency of DSSC. The electrical conversion efficiency increased when the back contact metal work function value increased, and the maximum photoelectric conversion efficiency could be obtained as Au was used as the back contact metal.

IntroductionX-ray multilayer enables an efficient reflection of X-rays in specific wavelength, serving as an important optics for obtaining monochromatic X-ray. The period thickness of multilayers typically falls within a nano-scale, and it is generally composed of two kinds of thin films alternately due to the short wavelength of X-ray. The multilayer can realize a high reflection of specific wavelength X-ray based on the Bragg diffraction, yielding a pure monochromatic X-ray. Compared to crystal monochromators, the period of multilayer can be adjusted readily, and the multilayer can be suitable for X-rays with different energies. The energy bandwidth of reflected X-ray obtained by the multilayer is 1–2 orders of magnitude broader than that of the crystals, thus providing a higher photon flux. At present, X-ray multilayer is widely used in synchrotron radiation for applications such as X-ray imaging, and X-ray small-angle scattering. According to the working principle of X-ray mutlilayer, the period of multilayer is only a few nanometers for the efficient X-ray reflection at a larger grazing angle due to the short wavelength of X-ray. Therefore, the coating method with a high precision thickness control is essential to prepare X-ray multilayer. Magnetron sputtering is the primary preparation method for X-ray multilayers, although achieving a high precision thickness control remains a challenge. Atomic layer deposition (ALD) method can realize uniform growth of single atomic layer thickness film on substrate and the film thickness control is extremely accurate because of the special film growth process. The ALD method has unique advantages in the preparation of small-period multilayer. It is identified that HfO2/Al2O3 is an effective material pair for X-ray multilayer at a wavelength of 0.154 nm and the X-ray reflectivity is acceptable.MethodsTo further investigate the properties of HfO2/Al2O3 X-ray multilayer prepared by the ALD and explore its application, HfO2/Al2O3 X-ray multilayer with a period of 3.8 nm and a period number of 60 was prepared by the ALD method. The surface 3D profile of the multilayer was obtained by a model ICON2-SYS atomic force microscope (AFM, Bruker Co., Germany), and the surface roughness of the multilayer was analyzed by a software named NanoScope Analysis. The microstructure of multilayer was determined by a model Talos F200s G2 transmission electron microscope (TEM, Thermo Fisher Scientific Co., USA) with a software named Digital Micrograph. The X-ray reflectivity of the multilayer was measured by a model D8 Discover X-ray diffractometer (XRD, Bruker Co., Germany) at X-ray wavelength of 0.154 nm. A software named IMD was used to fit the measured reflectance data. The fitting was based on a two-layer model, and the structure parameters of multilayer was obtained, including period and duty cycle. The multilayer was also analyzed by a modle AL-Y3500 diffractometer (Dandong Aolong Ray Instrument Group Co., Ltd., China). Before the installation of multilayer, the X-rays emitted from the source were directly illuminated on the sample by a collimator. After the installation of multilayer, the X-rays emitted from the ray source were firstly reflected by the multilayer and then irradiated on the sample. The diffraction patterns before and after installing multilayer were analyzed. This work consisted of three main parts. Firstly, a software named Fluent was used to simulate the gas flow distribution according to the structure of ALD equipment chamber. The results show that the gas flow rate is uniform, indicating that the chamber is suitable for preparing a film with uniform thickness. Secondly, HfO2/Al2O3 multilayer is prepared by the ALD method, and the surface roughness, period thickness uniformity and X-ray reflectivity are characterized. Finally, the multilayer is installed on the X-ray diffractometer to test the diffraction patterns of a silicon sample. The results show that the quality of the diffraction patterns is improved after the installation of the multilayer.Results and discussionThe characterization results show that the surface root-mean-square roughness of the multilayer is 0.77 nm, the maximum deviation of the period is 0.11 nm and the maximum X-ray (i.e., 0.154 nm) reflectivity of the multilayer is approximately 45%. The multilayer is in an amorphous state, and the interface between two layers is clear and sharp. The XRD pattern of Si powder sample is obtained without and with the HfO2/Al2O3 multilayer in X-ray diffractometer. The results indicate that the background intensity of the pattern with multilayer reduces, and no diffraction peaks of other X-ray appear. The reason is that the X-rays irradiated on the sample are composed of high-intensity Cu characteristic X-rays, other stray energy X-rays and low-intensity bremsstrahring radiation when the multilayer is not installed. After the multilayer is installed, the pure Cu k characteristic X-rays are irradiated on the sample and thus there is basically no interference peak.ConclusionsThis work investigated the characteristics of HfO2/Al2O3 multilayer prepared by the ALD method and explored its application in X-ray diffractometers. The results showed that HfO2 and Al2O3 layers were amorphous with a maximum X-ray (0.154 nm) reflectivity of approximately 45%. The quality of the diffraction patterns after the installation of the multilayer was improved, demonstrating a potential application of the HfO2/Al2O3 multilayer.

IntroductionThe treatment of dye wastewater remains a formidable challenge within the realm of industrial wastewater management. The existing methods for treating dye wastewater include adsorption, electrochemical oxidation, photocatalysis, biodegradation, and membrane separation. Among these, membrane separation is particularly advantageous due to its energy efficiency, high efficacy, and ease of process control. Separation membranes can be categorized into two primary types, i.e., organic and inorganic membranes. Inorganic membranes have attracted significant attention due to their high mechanical strength, controllable size distribution, resistance to high temperature and pressure, and chemical stablility. Despite these benefits, the advancement of conventional inorganic ceramic membranes is hindered by their complex preparation processes and high energy demands. This limitation underscores a urgent need for the development of simple and cost-effective inorganic membranes. A promising innovation in this field is a geopolymer-zeolite composite membrane, which is synthesized through in-situ hydrothermal conversion following the formation of reactive silica-aluminum materials under alkali or acid excitation. Coal gangue (CG), as a solid waste from coal mining primarily composed of SiO2 and Al2O3, can be used as a raw material for the production of geopolymer-zeolite composite membrane. In this paper, A cost-effective geopolymer-zeolite composite membrane was prepared by a polymerization-hydrothermal method with CG as the main raw material for the separation of RhB in water. The objective was to propose a novel utilization strategy for CG and to lay a foundation for the development of membrane separation technology.MethodsCoal gangue (CG) (Zhunger, Inner Mongolia, China) was calcinated at 800 ℃ for 2 h, resulting in the formation of calcined coal gangue (CCG). The CCG was then uniformly mixed with other additives in a mass ratio of CCG:NaOH:Na2SiO3: H2O (100:10:40:45). This mixture was poured into polytetrafluoroethylene moulds with the diameters of 50 mm and thicknesses of 5 mm. The samples were cured at 70 ℃ for 24 h to produce coal gangue-geopolymer membrane (CCG-GM). Afterwards, the CCG-GM was hydrothermally treated in a 100 mL 1 mol/L NaOH solution at 140 ℃ for 12 h, resulting in a coal gangue-geopolymer-zeolite composite membrane (CCG-GZCM).The prepared CCG-GZCM was incorporated into a custom-built terminal filtration system, and its performance was evaluated via measuring the pure water flux, membrane flux, and RhB removal rate under varying conditions (i.e., system pressure, initial RhB solution concentrations, and pH values of the RhB solution). In addition, the circular utilization of CCG-GZCM was also examined to assess its reusability. The selectivity of CCG-GZCM for the removal contaminants (i.e., RhB, methylene blue (MLB), methyl violet (MV), methyl blue (MB), methyl orange (MO), and eosin Y (EY)) was determined.Results and discussionThe XRD patterns reveals that CG primarily consists of quartz and kaolinite. Upon calcination at 800 ℃ for 2 h, kaolinite in CG transforms into metakaolinite, thus froming CCG. For alkali activation, a broad diffraction peak of CCG initially appears at 15°–25°, and then shifts to 20°–40°, indicating the conversion of CCG into a geopolymer, thus forming CCG-GM. Subsequently, CCG-GZCM, with a main phase composition of NaP1 zeolite, is obtained via in-situ hydrothermal transformation of CCG-GM. The FTIR spectra of CCG-GZCM confirm the presence of NaP1 zeolite absorption bands. The XPS spectra further identify Na, Si, O, and Al as the predominant elements on the surfaces of both CCG-GM and CCG-GZCM. Note that CCG-GZCM exhibits a larger specific surface area and pore volume, compared to CCG-GM, albeit with a smaller average pore size. The TG-DSC spectra indicate that the weight loss of CCG-GZCM primarily occurs at 30–253 ℃, indicating its good thermal stability. In addition, the compressive strength of CCG-GM also increases from 18.84 MPa to 47.01 MPa (CCG-GZCM) after hydrothermal transformation.CCG-GZCM exhibits varying removal efficiencies for different dye types. At an initial concentration of 10 mg/L and a pH value of 7, the pure water flux and membrane flux to RhB solution of CCG-GM and CCG-GZCM increase linearly with the increase of system pressure, while the RhB removal rate decreaseslinearly. At a system pressure of -0.08 MPa and a pH value of 7, the RhB removal rate by CCG-GZCM gradually decreases with the increase of initial concentration within 60 min. Conversely, at the same system pressure and an initial RhB concentration of 10 mg/L, the RhB removal rate by CCG-GZCM gradually increases with increasing pH value of the RhB solution within 60 min, and the maximum is 99.11%. However, when the system pressure maintains at –0.08 MPa, at an initial RhB concentration of 10 mg/L and an initial pH value of 11, the RhB removal rate by CCG-GZCM decreases with an increasing number of cycles, reducing to 94.25% after 6 cycles. In addition, CCG-GZCM demonstrates different removal rates within 30 min for different dye types. It achieves a high removal rate of exceeding 99% for cationic dyes including RhB, MLB, and MV. In contrast, the removal rates for anionic dyes, including MB, MO, and EY, are 70.00%, 25.82%, and 42.69%, respectively.ConclusionsThe pore structure of geopolymer membrane was modified and converted it into CCG-GZCM with larger specific surface area, higher pore volume and smaller average pore size by a hydrothermal method. The CCG-GZCM was characterized by an amorphous geopolymer matrix with NaP1 zeolite and a minor presence of NaA zeolite on the surface. At a system pressure of –0.08 MPa, an initial RhB concentration of 10 mg/L, and a pH value of 11, the CCG-GZCM achieved a remarkable RhB removal rate of 99.11% for 60 min. In addition, the CCG-GZCM also demonstrated a high circular utilization efficiency for RhB removal, maintaining a removal rate of exceeding 99% for cationic dyes for 30 min. The CCG-GZCM could be used as a promising approach for the effective utilization of coal gangue, having a significant potential for application in water pollutant treatment.

Extend AbstractIntroductionThe solid solution structures of potassium tantalum niobate (i.e., KTa1–xNbxO3, KTN) crystals are investigated, which are formed by the solid solution of potassium tantalate (KTaO3) and potassium niobate (KNbO3). The KTN crystals exhibit superior piezoelectric, ferroelectric and electro-optic properties. However, a difficulty in preparing large-sized, high-quality KTN crystals restricts their practical applications. Some studies indicate that the Curie temperature of potassium tantalum niobate single crystals with a composition of KTa0.63Nb0.37O3 is near room temperature. Currently, single crystals of solid solutions are mostly grown by a melt method. The seedless solid-phase crystal growth technique is a method for growing single crystals in the solid-phase without the need for artificially introduced seed crystals. This method offers some advantages such as low cost, simplicity in process, and short growth cycles, providing a novel approach to crystal growth and demonstrating an application potential. This paper was to investigate the growth of potassium tantalum niobate single crystals by a seedless solid-phase technique, to expand the application scope of crystal growth and obtain an effective pathway for preparing potassium tantalum niobate single crystals. In addition, the mechanism of the transformation from KTN ceramic preforming to single crystals and delving into the this transformation was also discussed.MethodsHigh-purity K2CO3 (99.0%, in mass, the same below), Nb2O5 (99.5%), Li2CO3 (98.0%), Bi2O3 (99.0%), and Ta2O5 (99.5%) were selected as raw materials, and 0.3% (in mole fraction) of LiBiO3 was used as a flux. In the preparation by a seedless solid-phase method, the raw materials with anhydrous ethanol were placed in a nylon jar with zirconia balls as grinding media. The mixture was then ground in a ball mill for 24 h. After drying, dried powdered material was pre-fired at 750 ℃ for 10 h. After pre-firing, the powdered material was further ground for 24 h. The dried powdered material was sieved, and then pressed into a circular shaped preform with a diameter of 20 mm and a thickness of approximately 2–3 mm at 100 MPa. The circular green preform was sintered in a box-type resistance furnace at 1168–1198 ℃ for 21 h to ensure that the green preform could be fully sintered and converted into the desired crystal structure. After sintering, the furnace temperature was slowly reduced to room temperature at a rate of 0.5 ℃/min. The crystal structure of the samples was analyzed using X-ray diffraction (XRD). The microstructure of the single crystals and ceramics was determined by scanning electron microscopy (SEM), and the elemental composition of the samples was characterized by energy-dispersive spectroscopy (EDS).Results and discussionThe KTN single crystals grow in a layered mode, with a close contact between the crystals and the polycrystalline regions, featuring flat interfaces and regular geometric shapes. In the growth process, solute molecules or ions arrange in layered stacks on the crystal surface under the action of driving forces. There are differences in growth rates due to the different driving forces required for growth in various directions, leading to the formation of stepped structures on the crystal surface. The single crystals gradually engulf and integrate with the dispersed micro-grains in the ceramic region to achieve their growth. However, crystal surface energy exists during the actual growth process. The crystal planes with a higher surface energy have more intense attraction and binding capabilities for atoms, resulting in faster growth rates. In the growth of KTN crystals, the surface energy and growth rates of the crystal planes (100) and (110) jointly determine the main morphology of the crystals. The growth direction of the crystals follows a directional pattern of layered arrangement, and the crystals follow a two-dimensional layered accumulation growth mechanism in the radial dimension. The growth of KTN strip-shaped crystals is achieved through two-dimensional layered accumulation. In the solid-phase growth of KTN crystals, the crystals grown by two-dimensional layered accumulation create a "shell–core" structure with different coexistence interface geometries. The large KTN crystals grown by two-dimensional layered accumulation are enclosed within the "shell" of the fine grain region, resulting in their "shell–core" coexistence. The XRD patterns of the natural surface of the KTN bulk sample indicate a preferred orientation of the crystal plane (210). The addition of K, Bi, and Li elements promotes the formation of the liquid phase, which facilitates the atom migration and thereby drives the continuous growth of the crystals.ConclusionsThe strip-shaped KTN single crystals were prepared as LiBiO3 was used as a flux, and the single crystallization of ceramic preforms in this system via the seed-free solid-state crystal growth was explored. The obtained KTN single crystals grew in a two-dimensional layered stacking manner and possessed a cubic perovskite structure. A "shell-core" structure was formed within the strip-shaped crystals, consisting of outer layers that were short and fine, and inner layers that were long and coarse. The exposed surface of the crystals was the crystal plane (210). Based on the theory of two-dimensional crystal growth, liquid-phase-assisted mechanisms, and the anisotropy of interfacial energy and interfacial migration, the mechanism of seed-free solid-state crystal growth of KTN crystals was analyzed. This mechanism was the promotion of ordered expansion of the crystals on a two-dimensional plane through the auxiliary action of the liquid phase. Also, the differences in interfacial energy and the anisotropic nature of their migration without the need for seed crystal guidance could intervene the direction and morphology of crystal growth, thereby achieving the seedless solid-phase growth of KTN crystals.

IntroductionWith the development of aero-engine towards large thrust, high thrust-to-weight ratio, high thermal efficiency, low fuel consumption and long life, the turbine inlet temperature is bound to continuously increase. Thermal barrier coating (TBC) with high thermal insulation and long life becomes popular. Multi-rare-earth doped zirconia (especially Gd2O3-Yb2O3-Y2O3 co-doped ZrO2, GdYb-YSZ) as one of the most potential candidate materials for ultra-high temperature TBCs has a better high-temperature phase stability and a lower thermal conductivity rather than YSZ. It is easy to generate thermal stress during thermal cycling due to GdYb-YSZ with a low fracture toughness, resulting in a poor thermal cycling performance of GdYb-YSZ/NiCrAlY TBC. The design of double ceramic layer (DCL) TBC can alleviate the thermal stress mismatch and effectively solve this problem. In the DLC TBCs system, the thickness ratio between ceramic layers can seriously affect the thermal insulation effect and thermal cycle life of the coating. However, little work on the effect of thickness ratio of GdYb-YSZ/8YSZ ceramic layers on the mechanical properties and thermal cycle life of GdYb-YSZ/8YSZ DLC TBCs system has been reported yet.In this paper, four kinds of GdYb-YSZ/8YSZ DLC TBCs systems with different ceramic layer thickness ratios were prepared. The microstructure, mechanical properties and water quenching-thermal shock cycle properties of GdYb-YSZ/8YSZ DLC TBCs systems with different ceramic layers thickness ratios were investigated. The failure behavior and failure mechanism of GdYb-YSZ/8YSZ DLC TBCs at 1150 ℃were also discussed.MethodsGdYb-YSZ/8YSZ DLC TBCs systems with different ceramic layer thickness ratios (i.e., GdYb-YSZ:8YSZ=5:1, 2:1, 1:1, and 1:2) were prepared with GH4169 nickel-based superalloy as a matrix, NiCoCrAlY as a bond coating, 8YSZ as the first ceramic layer, and GdYb-YSZ as the second ceramic layer., and then four TBC systems on the surface of GH4169 nickel-based superalloy were prepared by a model UniCoatProTM atmospheric plasma spraying system (Oerlikon Metco Inc., USA). Subsequently, the microstructure and phase composition of four kinds of GdYb-YSZ/8YSZ DLC TBCs systems were characterized. The bonding strength and water quenching-thermal shock cycle behavior at 1150 ℃ were tested, and the failure mechanism was analyzed. To analyze the water quenching-thermal shock cycle behavior at 1150 ℃, the coated sample was kept in a high-temperature furnace at 1150 ℃ for 5 min and then quickly cooled in distilled water to room temperature. After drying, the sample surface was recorded as a cycle experiment. Repeating the steps above until the cracking or peeling area of the coating surface of greater than 5% was considered as the TBC failure.Results and discussionMost of the GdYb-YSZ and 8YSZ spraying powders were melted in a high-temperature plasma flame flow, and interacted with the matrix to form the smooth structures via high speed jetting. A small number of unmelted and partially melted particles can form a rough microstructure during deposition, accompanied by micro-cracks and pores. The interfaces of each layer are closely bonded and clearly demarcated in GdYb-YSZ/8YSZ/NiCoCrAlY/substrate TBC system. The average bonding strength of the four GdYb-YSZ/8YSZ thermal barrier coating systems all is greater than 40 MPa. The macro-image of fracture after tensile test shows that most of the fracture occurs at the interface between 8YSZ and NiCoCrAlY. The average porosity of GdYb-YSZ/8YSZ TBC systems with different ceramic layer thickness ratios (i.e., 5:1, 2:1, 1:1, and 1:2) calculated by a software named Image J is 3.39%, 2.65%, 4.0%, and 2.64%, respectively. The GdYb-YSZ ceramic layer is composed of 98% (mole fraction) c-ZrO2 phase and 2% m-ZrO2 phase, which is basically the same as the phase composition in the spraying powder.Four kinds of TBCs with different thickness ratios are intact after 80 quenching-thermal shock cycles at 1150 °C. GdYb-YSZ/8YSZ TBC systems with four ceramic layer thickness ratios (5:1, 2:1, 1:1, and 1:2) begin to peel off at the edge of the coating after 85, 165, 120, and 195 water quenching-thermal shock cycles, respectively. The water quench-thermal shock cycle resistance of the GdYb-YSZ/8YSZ TBC system becomes better as the thickness of the GdYb-YSZ ceramic layer decreases and the thickness of the 8YSZ ceramic layer increases at a constant total thickness of the ceramic layer. The GdYb-YSZ/8YSZ TBC systems with different ceramic layer thickness ratios (i.e., 5:1, 2:1, 1:1, and 1:2) have a quenching-thermal shock cycle life of 125, 195, 200, and 205 times at 1150 ℃, respectively. The GdYb-YSZ/8YSZ TBC systems with different ceramic layer thickness ratios are mainly characterized via the peeling of ceramic layer during the water-quenching-thermal shock cycle at 1150 ℃. The thermal stress is one of the main factors leading to the failure of TBCs. The residual stress of 8YSZ ceramic layer with different layer thickness ratios of 5:1, 2:1, 1:1 and 1:2 is 0.01, 0.08, 0.12 GPa and 0.15 GPa, respectively, after the water quenching-thermal shock cycle failure at 1150 ℃. The stress increases gradually with the increase of the thickness of the ceramic interlayer 8YSZ. The residual stress (i.e., compressive stress) in the TGO layer in the coating system with different layer thickness ratios is 0.51, 0.29, 0.72 GPa and 0.62 GPa, respectively. The compressive stress of the TGO layer will cause the 8YSZ/TGO interface to buckle, and the tensile stress will be generated at the crest of the wave, thus promoting the generation of cracks. In the process of water quenching-thermal shock cycle of GdYb-YSZ/8YSZ TBCs, some micro-cracks are easy to form on one side of the 8YSZ ceramic layer near the 8YSZ/TGO interface under the combined action of residual stress in the ceramic layer and TGO layer. The microcracks can further spread into TGO layer as the TGO layer increases, seriously weakening the interface bonding force between ceramic layer and adhesive layer, and finally leading to the failure of coating spalling.ConclusionsGdYb-YSZ/8YSZ double ceramic layer thermal barrier coatings (i.e., DCL TBC) with different layer thickness ratios (i.e., GdYb-YSZ/8YSZ = 5:1, 2:1, 1:1, and 1:2) were prepared by an atmospheric plasma spraying technology. GdYb-YSZ ceramic top coated with four coatings were composed of c-ZrO2 phase and a small amount of m-ZrO2. The bonding strength of the four coatings was greater than 40 MPa, and the average bonding strength of the coatings was the maximum when the thickness of GdYb-YSZ/8YSZ was 5:1. The thermal shock cycle resistance of GdYb-YSZ/8YSZ TBC system became better as the thickness of the GdYb-YSZ ceramic coat decreased and the thickness of the 8YSZ ceramic coat increased. At the thickness ratio of GdYb-YSZ/8YSZ of 1:2, the water quenching-thermal shock of the TBC was up to 205 cycles at 1150 ℃, having the optimum thermal shock resistance. In addition, GdYb-YSZ/8YSZ TBCs with different ceramic layer thickness ratios did not undergo a phase transformation under the water-quenching-thermal shock cycle at 1150 ℃. The failure mechanism of GdYb-YSZ/8YSZ TBCs with different ceramic layer thickness ratios under the quench-thermal shock cycle at 1150 ℃ was as follows, i.e., in the process of water quenching-thermal shock cycle at 1150 ℃, the residual stress in 8YSZ coating increased with the increase of the thickness of 8YSZ ceramic layer, and micro-cracks were easy to form on one side of 8YSZ ceramic layer at near 8YSZ/TGO interface. In addition, the formation of TGO in the process of thermal cycling was accompanied by the generation of compressive stress. TGO further increased, and the micro-cracks in the coating expanded into TGO with the increase of the number of thermal cycles, reducing the binding force of the coating, and eventually causing the coating to flake and fail.

IntroductionThe Aldol condensation reaction, as a pivotal organic reaction in forming C-C bonds, plays a crucial role in various fields such as pharmaceutical synthesis, total synthesis of natural products, and the industrial production of perfumes and dyes. Conventional Aldol condensation reactions often rely on liquid alkaline catalysts, which suffer from some issues such as difficulty in recovery, complex waste liquid treatment, strong corrosion to equipment, and severe environmental pollution, thereby restricting their application in industrial production. Therefore, the development of efficient, recyclable, and environmentally friendly solid alkaline catalysts becomes a key to promoting the greening and efficiency of the Aldol condensation reactions. Layered double hydroxides (LDHs), with their unique layered structure and adjustable composition have a great potential in catalysis. Recent research indicates that functional modifications, such as ion exchange, intercalation, and doping, can further enhance their catalytic performance. Among them, sulfate-intercalated magnesium aluminum LDH (MgAl-SO42&#x207B;-LDH) as a special LDH material, through the introduction of sulfate ions, can enhance the structural stability of LDH and introduce additional acidic sites. This makes MgAl-SO42&#x207B;-LDH effectively activate aldehyde or ketone molecules during the catalysis of Aldol condensation reactions, thereby promoting the progress of the condensation reaction. In addition, the morphology of the catalyst has a significant impact on its catalytic performance. The flower-like spherical structure, as a special nanostructure, has attracted much attention in the field of catalysis due to its high specific surface area, abundant pore structure, and superior mass transfer performance. Combining the flower-like spherical structure with MgAl- SO42&#x207B;-LDH can be expected to further improve the catalytic efficiency of the catalyst.MethodsThe flower-like MgAl-SO42&#x207B;-LDH catalyst was prepared in the molar ratios of n(Mg) : n(Al)=2 : 1 and n(urea) : n(NO3–)=1 : 1. In the preparation, magnesium nitrate hexahydrate (Mg(NO3)2·6H2O), aluminum nitrate nonahydrate (Al(NO3)3·9H2O) and urea were mixed and dissolved in deionized water under stirring. The mixed solution was then transferred to a hydrothermal reaction vessel and reacted at 120 ℃ for 6 h. The resulting product was washed with deionized water for several times and then dried thoroughly in an oven. The dried material was calcined in a muffle furnace at 500 ℃ for 8 h to obtain the magnesium-aluminum composite oxide (LDO). Subsequently, a series of ammonium persulfate solutions ((NH4)2S2O8) with different concentrations (i.e., 0, 0.05, 0.10, 0.15 mol/L, and 0.20 mol/L) were prepared. The reaction was proceeded in a nitrogen atmosphere at room temperature for a certain period, and the solvent was removed through filtration and drying, yielding the S2O82– intercalated magnesium-aluminum hydrotalcite compounds (i.e., S2O82–/LDH). Finally, the prepared S2O82–/LDH was calcined in a muffle furnace at 550 ℃ for 8 h to obtain the modified magnesium-aluminum composite oxide. The modified magnesium-aluminum composite oxide was uniformly dispersed in deionized water, and the hydration-reduction process was repeated when sulfate ions were introduced to replace persulfate ions. A final catalyst MgAl-SO42&#x207B;-LDH was obtained after completing the hydration-reduction and subsequent treatment.Results and DiscussionMgAl- SO42&#x207B;-LDH catalyst prepared has a flower-like spherical structure assembled from two-dimensional MgAl-SO42&#x207B;-LDH sheets with a uniform morphology and a diameter of 2 m. The SEM images reveal that, in contrast to the irregular shape of traditional MgAl-LDH catalysts, the MgAl-SO42&#x207B;-LDH sample after sulfate ion intercalation modification exhibits a flower-like spherical structure. This interconnected sheet structure divides the catalyst surface into numerous small chambers, which provides abundant reaction sites for the catalytic reaction and promotes the mass transfer process between reactants and products, thereby contributing to an enhanced catalytic performance. The results of EDS analysis confirm that sulfate ions are introduced into the catalyst structure and distributed uniformly.The XRD patterns indicate that the crystallinity of MgAl-SO42&#x207B;-LDH catalyst is inferior to that of the MgAl-LDH, exposing more active sites due to the intercalation of sulfate ions. The FTIR spectra reveal that MgAl-SO42&#x207B;-LDH catalyst contains a rich number of Brnsted acid sites, enhancing the activation of acetone and p-nitrobenzaldehyde, thus effectively improving catalytic performance. The optimized MgAl-SO42&#x207B;-LDH catalyst has a p-nitrobenzaldehyde conversion rate of 70.5%, and after one regeneration the catalytic conversion rate remains 64.03%, demonstrating a good reusability. The FTIR spectra of the reaction products indicate a mixture of 4-(4-nitrophenyl)-3-hydroxy-2-butanone and 4-(4-nitrophenyl)-3-penten-2-one.ConclusionsThe study prepared a flower-like sulfate-intercalated MgAl-SO42&#x207B;-LDH catalyst with a uniform microstructure and a diameter of 2 m, assembled from two-dimensional MgAl-SO42&#x207B;-LDH sheet structures. The catalyst showed an improved surface activity due to the intercalation of sulfate ions, thus increasing the number of active sites. The FTIR spectra confirmed the abundance of the Brnsted acid sites in MgAl-SO42&#x207B;-LDH catalyst, which significantly enhanced the activation of acetone and p-nitrobenzaldehyde, effectively improving catalytic performance. The optimized catalyst showed a high conversion rate of 70.5% for the aldol condensation of acetone with p-nitrobenzaldehyde, and after regeneration it maintained a conversion rate of 64.03%, indicating a good reusability. The reaction products consisted of a mixture of 4-(4-nitrophenyl)-3-hydroxy-2-butanone and 4-(4-nitrophenyl)-3-penten-2-one, indicating a possible mechanism involving the Brnsted acid sites on MgAl-SO42&#x207B;-LDH catalyst facilitating the nucleophilic addition reaction between the activated enol of acetone and the protonated aldehyde group of p-nitrobenzaldehyde.

High-energy Lithium-based battery technology is a crucial support for achieving the national "dual-carbon" goals. In the technology, silicon (Si)-based anodes are considered as one of the most promising choices for next-generation high-energy Li-ion batteries (LIBs) due to their ultra-high specific capacity, which is 10 times greater than that of graphite (i.e., silicon 4200 mA·h/g). When a "Si-majority" anode (mass fraction of silicon greater than 50% in the anode) works with a high-nickel LiNixMnyCozO2 (x+y+z=1, x≥0.6) cathode, the cell-level energy density can increase to over 400 Wh/kg. However, silicon usually experiences significant volume changes (i.e., over 300%) during lithiation and delithiation, leading to structural degradation, pulverization, unstable electrode-electrolyte interfaces, and gas generation, which pose substantial challenges to the cycling life of silicon-based anode batteries. Also, there is still a lack of clarifying key scientific issues such as interfacial evolution and failure mechanisms under practical conditions. Recent research on silicon-based anodes can be summarized as two main directions.1) Structure design and composites: Since nano-sized Si as an anode for LIBs was proposed, most of studies focus on the design of various nanostructured Si materials to improve their cycling stability, such as nanowires, hollow, porous, and core-shell Si anodes. However, nano-sized Si anodes suffer from a low initial Coulombic efficiency, a low packing density, and a high cost. To address these issues, Si-graphite composite materials are invented to buffer the volume expansion of Si via embedding those into graphite layers, thus improving the cycling stability and become a recent hot topic.2) Electrolyte engineering: For Si-majority or pure-Si anodes, neither structural design nor graphite confinement can fully alleviate the issues caused by volume expansion, such as the cracking of the solid electrolyte interphase (SEI) and pulverization. Another strategy based on electrolyte engineering is thus proposed to form stable SEI with high ionic conductivity and favorable mechanical properties on the Si surface. Carbonate-based electrolytes, known for their strong oxidation resistance, are highly compatible with high-voltage cathodes (>4 V vs. Li/Li+). For Si-based anodes, fluorinated ethylene carbonate (FEC) is considered as the most effective electrolyte component. However, FEC suffers from its poor thermal stability, leading to poor high-temperature cycling life and serious gassing at high temperatures. Ether-based electrolytes, featuring their strong reductive resistance, are extensively studied for Si-based anodes. For instance, some studies report that the rationally designed ether-based electrolytes can enable a high specific capacity of ~2400 mA&#x0387;h/g with stable cycling for over 200 cycles. These findings demonstrate that the interface stability of Si-majority/pure-Si anodes can be improved through electrolyte design and optimization, thus enhancing their cycling life.From the perspective of actual battery operating conditions, a research on electrolytes still faces some challenges. For silicon anodes with FEC-based carbonate electrolytes, the cycling performance is less than ideal, particularly due to the poor high-temperature stability of FEC and the issues of gas generation, which remain difficult to resolve. Ether-based electrolytes can improve the cycling stability of silicon-based anodes, and exhibit a poor oxidation resistance. This poses challenges such as oxidative decomposition and high-temperature gas generation, especially when paired with high-voltage cathodes and operating in high-temperature environments. It is thus essential for the practical application of such electrolytes to overcome these issues. An analysis spanning from fundamental mechanisms to practical implementation identifies the key points for the design and development of electrolytes tailored for silicon-based anodes:Summary and prospectsThis review represents the research progress on Si-based anodes from material structure to electrolyte design, and highlights the need to clarify the degradation mechanisms of Si-based anodes, develop novel characterization methods, and design new molecular structures to improve interfacial stability. Some efforts in future research should be directed towards the following aspects, i.e., 1) further elucidating the degradation mechanisms of Si-based anodes from the chemomechanical viewpoint. The Si-based anodes work under cyclic stress and chemical corrosion, which are highly intertwined. This coupling effect is highly analogous to the "stress corrosion cracking" well known in metals. Clarifying and decoupling these two failure mechanisms can guide the design of both electrodes and electrolytes, 2) developing novel interdisciplinary characterization and research methods. These methods should employ both in-situ and ex-situ techniques in microscale to investigate crack propagation and structural evolution of electrodes and interfaces under various chemical environments, 3) evaluating novel electrolytes under realistic conditions. It is essential to assess the performance of batteries, including cycling and gas generation, under realistic conditions such as high-loading electrodes, lean electrolytes, large-format cells and extreme temperatures to analyze the battery degradation mechanisms, 4) electrolyte design for Si–C anodes must consider the compatibility of both Si and Graphite. Since the requirements for electrolytes differ between graphite and silicon, a design needs to balance the compatibility of different solvents and additives with the both materials, and 5) developing novel additives and solvent molecules. Designing new molecules to replace FEC and address the gassing issue at high temperatures is quite meaningful for practical applications. The design of these molecules should ensure superior oxidation resistance, compatibility with high-voltage cathodes, stability with trace amounts of water in the electrolyte, and good chemical stability at high temperatures.

Carbon aerogels can be used as exceptional carrier materials in many applications like energy storage, wastewater treatment, and electronics due to their large specific surface area and high porosity. However, some challenges such as high production costs, complex manufacturing processes and the use of toxic precursors hinder their applications. Lignin as a naturally abundant, environmentally friendly, and cost-effective aromatic polymer presents a promising solution, which can be used as a precursor for carbon aerogels. Recent studies focus on harnessing the unique chemical structure and reactive groups of lignin to develop high-performance and lignin-based carbon aerogels.However, the widespread adoption of carbon aerogels is hindered by several challenges. The high costs for conventional preparation methods restrict their large-scale production mainly due to expensive precursors and complex manufacturing processes. Also, controlling the pore structure and morphology during synthesis remains a significant hurdle, affecting the consistency of their performance. To address these issues, some researchers turn their attention to lignin as a natural and abundant polymer. Lignin is cost-effective and environmentally friendly and features a complex chemical structure with numerous reactive functional groups, making it an ideal precursor for carbon aerogels.Significant progress has been achieved in the preparation techniques for lignin-based carbon aerogels. In the gelation process, some research efforts focus on optimizing reaction conditions to improve the degree of cross-linking in lignin. Researchers can tailor gel structures that serve as a foundation for the subsequent formation of the aerogel network via adjusting the types and quantities of initiators and cross-linkers. In the drying phase, multiple methods are explored. Freeze-drying is widely used due to its ability to largely preserve the original pore structure of gel, yielding carbon aerogels with a high porosity and a uniform pore distribution. Supercritical drying also offers some benefits in maintaining nanostructure though it requires specialized equipment and expensive solvents. Meanwhile, ambient drying is simple and great cost-effective ratio, which still has some challenges like uneven drying and structural shrinkage that need to be addressed. For carbonization, the temperature, rate and duration of this process significantly affect the final properties of the carbon aerogel. High-temperature carbonization can enhance carbon content and specific surface area. However, it also leads to pore collapse. A judicious combination of these parameters can yield carbon aerogels with optimized pore structures and improved mechanical properties.In applications, lignin-based carbon aerogels have superior performance. In energy storage, they can enhance the capacitance and cycling stability of supercapacitors. Their extensive specific surface area provides more active sites for charge storage, and the porous structure facilitates rapid ion diffusion. As catalysts or catalyst supports, the high surface area and porous nature of lignin-based carbon aerogels effectively disperse active components, thereby improving catalytic activity and selectivity. In gas storage and separation, their well-defined pore structures enable the selective adsorption of specific gas molecules. In wastewater treatment, they efficiently remove various pollutants, including heavy metals, dyes, and organic contaminants, based on the physical and chemical adsorption mechanisms.This review represents the preparation and application of lignin-based carbon aerogels. The gelation, drying, and carbonization processes are critical steps that determine the final properties of the carbon aerogel. Freeze-drying is a preferred drying method, and some efforts in this method should direct towards further reducing its energy consumption and costs. Precise control of the carbonization process is essential for achieving carbon aerogels with the desired pore structures and properties.Summary and ProspectsThe gelation of lignin often requires considerable time, which hampers production efficiency. Also, the carbonization process is complex and difficult to control with precision, leading to inconsistencies in product quality. A future research should focus on modifying lignin to enhance its reactivity, thereby reducing gelation time. The development of more advanced composite templates for carbonization should facilitate a better control for pore structure. In addition, investigating novel methods to introduce functional groups into lignin-based carbon aerogels can broaden their applications in emerging fields like flexible electronics and biosensors.

Inorganic silicate coatings, as a type of inorganic paint, exhibit the superior long-term anti-corrosion performance, heat resistance, and weather resistance. Compared to the conventional organic coatings, inorganic silicate coatings are more environmentally friendly, aligning with the development of eco-friendly coatings. However, pure inorganic silicate coatings face some issues such as high brittleness, susceptibility to cracking, and poor water resistance, thus restricting their practical application. Extensive research has been conducted on the modification of these coatings to enhance the mechanical properties and anti-corrosion performance of inorganic silicate coatings. It is essential to review the current research on the modification of inorganic silicate coatings to accelerate progress in both research and practical applications in engineering. This review represents the modification of inorganic silicate coatings with various materials affecting the overall performance of the coatings in varying degrees. Based on a spatial multi-scale approach, the relevant modifying materials are categorized into three scales, i.e., macro-scale, meso-scale, and micro-scale.In the macro-scale, the modifying materials include inorganic substances such as silica sol, aluminum phosphate (AlPO4), and calcium hydrogen phosphate (CHP), while organic materials consist of organo-silicon emulsions, silane coupling agents, acrylic resins, and epoxy resins. Silica sol significantly enhances the corrosion resistance of coatings via increasing their degree of crosslinking. Phosphate compounds promote the polymerization of silanol groups by providing H&#x207A; ions, thereby accelerating the curing of the coatings and improving their water resistance. Organic modifiers increase the internal crosslinking density and flexibility of the coatings, while also imparting hydrophobic and self-cleaning properties, by introducing functional groups such as hydroxyl and carboxyl groups into the silicate curing process.In the meso-scale, the modifications primarily utilize micron-sized fillers, i.e., zinc-type fillers, layered structure fillers, and other metallic and metal oxide fillers. Zinc-type fillers encompass zinc powder, zinc oxide, and zinc silicate. Layered structure fillers include mica powder, two-dimensional transition metal carbides (Ti3C2Tx), and layered double hydroxides (LDHs), and other metallic and metal oxide fillers comprise aluminum powder, zinc-aluminum alloy powder, mica iron oxide, titanium (Ti), and titanium oxide (TiO2). Metal and metal oxide modified fillers enhance the anti-corrosion performance of inorganic silicate coatings through a threefold mechanism of filling, physical shielding, and electrochemical protection, which delays and prevents the penetration of corrosive media into the substrate surface. In addition, these fillers can also chemically bond with silicate matrix, thereby enhancing the adhesion and chemical stability of the coating. Layered fillers, such as mica powder, Ti3C2Tx, and layered double hydroxides (LDHs), provide the superior physical shielding and create a "maze effect" within the coating, which extends and convolutes the diffusion pathways of corrosive media like chloride ions and water molecules. This ultimately slows down the corrosion reactions of the substrate and improves the anti-corrosion performance of the coating.In the micro-scale, the materials for the modification are mainly nano-scale fillers, such as nano-silica (SiO2) and graphene-based nano-scale fillers. This review discusses the impact of modification materials in different scales on the performance of inorganic silicate coatings and elucidates their modification mechanisms. This review addresses the existing shortcomings of modification materials in various scales in current research and outlines the future development trends. Nano-SiO&#x2082; enhances the bonding strength and anti-corrosion performance of inorganic silicate coatings via participating in the construction of the silicate network, thereby increasing the internal connectivity of the coating. Meanwhile, nano graphene-based fillers can improve the weather resistance and anti-corrosion performance of the inorganic silicate coatings due to their exceptional chemical stability, conductivity, and barrier properties.Summary and ProspectsIn the macro-scale modification, constructing an environmentally friendly coating system to reduce the use of organic solvents is a future development direction. However, inorganic modifiers have some limitations in enhancing the coating performance. The use of organic modifiers can lead to the emission of volatile organic compounds (VOCs), which negatively impacts the environment. It is thus important for the dual goals of optimizing coating performance and environmental protection to develop eco-friendly inorganic silicate coatings that balance the use of inorganic and organic modifiers. The performance of the prepared inorganic silicate anti-corrosion coatings is relatively singular. A future research can explore multifunctional composite anti-corrosion coatings that can be used under various environmental conditions, such as super-hydrophobicity, self-cleaning, self-healing, oxidation resistance, and wear resistance. In the micro-scale, the modification of inorganic silicate coatings with nanoparticles may reduce economic viability due to the high cost of nano-materials. A future research should focus on the development of more cost-effective methods for synthesizing nanoparticles, such as microwave-assisted synthesis, solvothermal methods, and sol-gel techniques. In addition, introducing specific functional groups on the surface of nanoparticles or combining them with other materials to impart the coatings with multifunctional properties like self-healing, super-hydrophobicity, antibacterial activity, and flame retardancy can be an important development direction. In the molecular/atomic scale, studies on the modification of inorganic silicate coatings primarily focus on experimental approaches, failing to delve into the molecular or atomic mechanisms. This results in an incomplete understanding of the coatings performance. A future research should focus on molecular design and optimization, as well as molecular dynamics simulations, to clarify the modification mechanisms for inorganic silicate coatings.