Introduction Gallium oxide crystal phase is widely used, and the metastable phase can be only obtained by means of epitaxy. Selecting suitable growth methods and improving the heteroepitaxy level are still the problems that the academic circles have been discussing. The mist chemical vapour deposition (Mist-CVD) method does not require a vacuum system, which is safe, economical, simple and controllable, demonstrating a potential application for metastable phase growth of gallium oxide. The level of heterogeneous epitaxy of various crystalline phases of gallium oxide still needs to be further improved. For instance, there is still a problem of mixing phases during the growth process, which leads to a difficulty of regulating the crystalline phases. In terms of the nucleation mechanism, there is a lack of process-specific research and explanatory notes. In this paper, the most suitable growth intervals of Ga2O3 crystalline phases on sapphire substrates were explored in a vertical mist chemical vapour deposition epitaxy furnace. The relationships among Ga2O3 crystal phase modulation, crystalline quality, growth rate and growth process parameters, as well as the optimisation of surface morphology by annealing were investigated respectively. The nucleation mechanism of the films and the precursor droplet motion law of Mist-CVD were analyzed, and the experimental results were discussed.Methods Heterogeneous epitaxial Ga2O3 films were prepared by the Mist-CVD method in a vertical mist chemical vapour deposition epitaxy furnace with a 2-inch commercial sapphire substrate. Gallium acetylacetonate (Ga(acac)3) aqueous solution was used as a precursor for Ga and O at 0.05 mol/L, and 1.5% HCl was used as a co-solvent. High-purity Ar (99.999%) was used as a carrier gas with Ar flow rate of 2 250 mL/min and growth pressure of 101.325 kPa, respectively, and O2 was used as a dilution gas. The growth parameters were adjusted in the growth temperature range from 600 ℃ to 750 ℃ and O2 flow rate from 225 mL/min to 600 mL/min for 1 h.A model D8-ADVANCE X-ray diffractometer (XRD) was used to determine the crystal structure. The elemental composition and content of the epitaxial layer surfaces were determined using a model ESCALAB-250 X-ray photoelectron spectrometer (XPS). In addition, the cross-sectional image and thickness of the Ga2O3 epitaxial layer were characterized by a model S-4800 scanning electron microscope (SEM). The AFM images and root-mean-square (RMS) surface roughness of Ga2O3 epitaxial layers were analyzed by a model ICON-type atomic force microscope (AFM). A LAMBDA-950 UV-Vis spectrophotometer (UV-Vis) was used to characterize the transmittance of the samples and calculate the forbidden band width.Results and discussion Based on the XRD analysis on the crystal structure of the epitaxial films of gallium oxide at 600-750 ℃, there exist α and ε mixed phases at 600-700 ℃, and the film diffraction peaks are weak with no obvious selective orientation. This is due to the higher lattice matching of the α phase to the c-plane sapphire substrate, which is more easily induced to grow on the sapphire substrate at low temperatures and stabilised under lattice mismatch strain. The ε-pure phase is grown at 750 ℃ with a selective orientation. The crystalline phase of gallium oxide at higher temperatures is not explored, and pure phase films of β-Ga2O3 are expected to be obtained at higher growth temperatures due to the limitations of the current epitaxial furnace equipment with a maximum set temperature of 750 ℃.ε-Ga2O3 pure phase is generated at different oxygen flow rates. However, the half-height width of the diffraction peaks of the films appears at a lower level of 0.141° and 0.150° for oxygen flow rates at 225-250 mL/min. When the oxygen flow rate exceeds 400 mL/min, the half-height widths of the samples become larger, i.e., 0.178°, 0.162° and 0.162°, respectively, and the crystalline quality of the films decreases, but the overall crystalline quality of ε-Ga2O3 is better. Based on the analysis on the cross-section of the film thickness by SEM, the growth rate of the film firstly increases and then decreases with the increase of the oxygen flow rate, reaching a growth rate of 3 μm/h at 250 mL/min and 400 mL/min. A higher growth rate can be obtained at 250-400 mL/min.For the ε-Ga2O3 samples grown at 750 ℃, the surface of the ε-Ga2O3 sample shows a cluster with coarser particles and a cloudy morphological features, and its surface roughness is 10.8 nm. From the 3D stereogram, the surface morphology shows multiple bulge-like or hill-like bumps, and the height of the surface bumps is between 50 nm and 100 nm, and its overall height is larger than the grain size of the film, indicating that the epitaxial film grows in accordance with the three-dimensional island-like growth mode. The β-Ga2O3 sample after annealing at 900 ℃ is similarly characterized by AFM. The surface roughness of the thin film of β-Ga2O3 samples decreases as RMS=5.42 nm, and the grain size becomes larger, ranging from 50 nm to 80 nm. At high temperatures, the height of the surface bumps is between 5 nm and 30 nm, which is smaller than the grain size, showing that the film densities increase at high temperatures.Conclusions Pure phase films of ε-Ga2O3 were obtained at 750 ℃ in a three-dimensional island growth pattern with a high surface roughness of 10.8 nm at low temperatures. The growth rate of the film of ε-Ga2O3 at an oxygen flow rate of 225 mL/min was 1 μm/h with a half-height width of 0.141°. The growth rate of the film increased to 3 μm/h with increasing the oxygen flow rate, and the film was still in the pure phase of ε-Ga2O3. However, the crystalline quality of ε-Ga2O3 was not improved, while rather deteriorated due to the high oxygen flow rate. β-Ga2O3 pure-phase thin films were obtained after annealing at 900 ℃, and the surface morphology was further optimized with a reduced roughness of only 5.42 nm. The pure crystalline phases of Ga2O3 with a considerable thickness could be prepared by a low-cost mist chemical vapor deposition method in a Mist-CVD equipment.

Introduction In the continuous casting of aluminum-killed steel, alumina (Al2O3) clogging frequently forms on the submerged entry nozzle (SEN) inner wall, causing SEN clogging. SEN clogging has several negative influences, i.e., flow field instability, reduced continuous casting efficiency, deteriorated billet quality, decreased quality of produced steel, and even interruption of continuous casting. The SEN clogging problem in the continuous casting of aluminum-killed steels has attracted much attention in metallurgy and refractory materials. Although many anti-clogging measures are proposed, the SEN clogging problem has not been solved yet. The main reason is that the mechanism of SEN clogging is still unclear. The clogging of SEN is mainly caused by the movement and adhesion of inclusions in molten steel to the SEN wall. It is thus necessary to clarify the mechanism of SEN clogging, especially force condition of inclusions in molten steel. Recent studies indicate that there is a correlation between the movement of oxide inclusions in molten metal and an applied electric field. The movement of inclusions under the action of electric field force is an important factor in SEN clogging. An effective way to solve SEN clogging problem is to analyze the effect of inclusions in refractory materials and molten steel from the perspective of electric field. In this paper, the charge of Al2O3 inclusions in SEN wall and molten steel during continuous casting was investigated to reveal the mechanism of SEN clogging and blockage, and to develop a technology for preventing alumina clogging in SEN. This work can provide a theoretical and technical support for solving the technical problems of SEN clogging in iron and steel industry.Methods The powerful mixers (Erich Co., Germany) were used for mixing. The blast drying oven was used to control the volatile matter of the material. The samples were prepared by four-column hydraulic press and isostatic press. The samples were heat treated under the condition of carbon embedded (i.e., 850 ℃ for 6 h). The wall charge of SEN in the continuous casting process was measured by a digital source meter. The charging properties of Al2O3 inclusions was determined based on the movement law of Al2O3 inclusions under the action of external electric field. The correctness of the experimental results was verified by applying a reverse electric field on the continuous casting site. The microstructure and morphology of the refractory were determined by a model JSM-700 F scanning electron microscope with a resolution of 3.0 nm, and the micro-area composition was analyzed by a model X-Max50 energy spectrum analyzer.Results and discussion The results of rotating flow experiment indicate that the SEN material is charged when rubbing with the molten steel. The experiments in the continuous casting site show that the SEN wall is negatively charged during the continuous casting process. There are three main reasons: 1) since the work function of Fe is 4.5 eV less than that of C is 5.0 eV, when the molten steel contacts with graphite in SEN, the electrons will migrate from the molten steel to SEN, making SEN negatively charged; 2) because the Fermi level of Fe (11.1 eV) is more than that of Al2O3 (8.7 eV), the electrons will migrate from the molten steel to SEN when the molten steel contacts with the corundum Al2O3 in SEN, making SEN negatively charged; 3) When the two conductors with different temperatures are in contact, the electrons will migrate from the conductor with a higher temperature to the conductor with a lower temperature. In the continuous casting process, the temperature of the SEN is less than that of molten steel, and the electrons will migrate from the molten steel to the SEN, eventually leading to the negative charge on the SEN wall.Al2O3 inclusions in molten steel move and adhere to the negative electrode under the action of electric field. This indicates that Al2O3 inclusions in molten steel are positively charged. The main reason is that in the continuous casting process, the molten steel is under high temperature and low oxygen conditions, and the oxygen vacancy in the defects of Al2O3 formed by the reaction of [Al] and [O] in the molten steel has the lowest energy.The effect of applied electric field on the SEN nodulation was carried out in the continuous casting site. The results show that the degree of nodulation in the inner cavity of SEN connected to the negative electrode of the power supply is higher than that of the reference SEN without electric field, while there is a slight nodulation in the inner wall of SEN connected to the positive electrode of the power supply. This further indicates that the SEN wall is negatively charged in the continuous casting process, and the Al2O3 inclusions in the molten steel are positively charged. SEN clogging can be effectively prevented via applying an external reverse electric field.Conclusions The clogging mechanism of SEN was analyzed based on the perspective of electrochemistry. In high-temperature simulation tests, actual working condition test and field verification, the charging characteristics of SEN wall surface, the charging behavior of Al2O3 inclusions in molten steel and their motion behavior under electric field were explored. The charging mechanism of alumina inclusions in molten steel was explained by double electric layer and friction charging theory, point defect theory in ion crystal and first-principle. The SEN was negatively charged in continuous casting, and the charge increased with the increase of casting speed. The Al2O3 inclusions in molten steel were positively charged, and the electrostatic force was an important driving force for the migration of Al2O3 inclusions to the SEN wall. The high-temperature simulation test and the actual working condition test both showed that the Al2O3 inclusions could move to the cathode and adhered to the Al2O3-C material, and the thickness of the attachment increased with the increase of the electric field strength.

Graphyne is a two-dimensional carbon allotropy that consists of different hybrid forms of carbon atoms in topological order. In different kinds of graphyne, carbon atoms undergo a rich bonding process through sp and sp2 hybridization, i.e., aromatic bonds, single bonds, and triple bonds, forming electron covalent frameworks and pores. Different bond lengths provide a higher structural flexibility and make coordination environmental regulation easier. For instance, the coexistence of sp and sp2 hybrid carbon atoms in graphdiyne (i.e., a typical and only artificially synthesized graphyne) makes the surface local electrons distributed unevenly, which makes it possible to design chemical reactions, site selective doping, and controllable atomic loading. The effective synthesis of graphdiyne (GDY) has thus brought a vitality to the research and development of carbon materials, providing a platform for carbon materials in various fields such as energy conversion and storage, optics, electronics, and magnetism, and having opportunities for the development of transformative materials. This unique alkyne bond rich structure provides almost infinite possibilities for the precise GDY tailoring towards different specific application scenarios. This review thus represented recent development on the GDY modifications and the corresponding functionalities of alkyne bonds.The precise modifications of graphdiyne (GDY) with the unique acetylene rich structure can be summarized as follows: 1) Compounding graphdiyne with nanoparticles (NPs), or the formation of metal graphdiyne bonding to enhance the charge transfer between materials. The structure rich in sp carbon can hybridize with other active components (such as molecules and NPs) to form a larger electron cloud density overlap, resulting in intense interactions. The composition with NPs can accelerate electron transfer and enhance the activity and stability for catalysis. GDY is easy to obtain and lose electrons, which is referred as an “electron sponge” property. A portion of GDY electrons can transfer with the hybrid across the interface. The electron cloud at a polarized region generates unique electron transfer enhancement characteristics.2) Utilizing acetylene bonds as the chemical reaction sites to achieve controlled doping of heteroatoms at desired sites. The sp and sp2 hybrid carbon atoms have a higher chemical activity, providing greater possibilities for the targeted introduction of heteroatoms. Especially, the triple bonds can create a form of nitrogen doping (i.e. sp-N) through pericyclic reaction and click reaction. Some work reveal a clear upward trend with the concentration of sp-N for ORR, OER and CO2RR reaction. Meanwhile, the doping of sp-N atoms with clear chemical sites paves a way for the development of high-performance carbon based catalytic materials via precise site modifications. Furthermore, optimized doping configuration and suitable spatial distance can enhance a catalytic activity when doping multiple elements. These findings clearly demonstrate that GDY can control the doping at desired sites for synergistic effects and catalytic performance optimization.3) The regulation of ions or atoms transportation or anchoring metal atoms. The two-dimensional plane topology of GDY is orderly distributed with molecular pores of specific sizes. The molecular pore and the surrounding acetylene rich structure endow it with a unique localized electron distribution. Therefore, regulating the size of stacking pores and the surrounding environment will affect the mass transfer process, thus applying graphene into more fields like seawater desalination and photothermal steaming evaporation. More importantly, the electrons from alkyne bond can interact with the metal empty orbits. The unique properties of GDY make it a suitable carrier for anchoring single metal atoms. The anchored metal atoms on the surface of GDY tends to be in zero valence, which in turn brings unique catalytic properties.Summary and prospects Graphene is a novel two-dimensional carbon material composed of sp and sp2 hybrid carbon atom topologies. This material has attracted widespread attention in various fields such as materials, chemistry, physics, information, biology, and the environment due to its high conjugation, abundant acetylene bonds, regular ordered pores, and adjustable electronic structure, which still has many challenges and opportunities in the future.1) Developing large-scale, high-quality, and low-cost synthesis technology can provide a solid material foundation for theoretical research and practical applications. The preparation methods for other types of graphene (such as GY, GY-3, and GY-4) are still in the exploratory stage. Obtaining novel graphene materials will expand the application field and clarify the structure-activity relationship.2) The alkyne bonds of graphene provide more designability. Designing chemical reactions based on alkyne bonds is of great significance for improving the performance of non-metallic catalytic materials. Synthesizing graphene and its derivatives with precise structures, controlling the degree of reaction and functionalization of alkyne bonds can effectively expand the scope and depth of “alkyne chemistry”.3) A more precise manipulation of the pore structure of graphene is needed. The precise regulation from molecular pores to nanopores will deepen the understanding of the atoms/ions transport and anchoring in GDY, and fully leverage its structural advantages. Accurately controlling the anchoring points, atomic numbers, and atomic types of metal elements is crucial for developing atomic catalytic materials, enhancing catalytic activity, selectivity and stability.The unique rich alkyne bond and pore structure of graphene provide infinite possibilities for the precise synthesis and controllable preparation. This review can provide a reference to understand the development of graphene materials for various applications.

Lithium-ion batteries (LIBs) dominate the field of energy storage due to their high specific energy density and long cycle life. However, the scarcity of lithium resources, potential safety issues, and high cost severely restrict their further large-scale energy storage applications. It is thus urgent to investigate other battery systems beyond LIBs. Alkali metal ions (Na+ and K+) with higher abundance and multivalent charge carriers (i.e., Zn2+, Mg2+, Al3+, etc.) have attracted much attention. Although sodium-ion batteries (SIBs) and potassium-ion batteries (KIBs) have the similar chemical properties to Li, they have lower energy density (i.e., 100-120 W·h/kg, 150-170 W/kg), toxic and flammable electrolytes, high operating costs, and safety hazards. Magnesium-ion batteries (MIBs) and aluminum-ion batteries (AIBs) involve multi-electron redox reactions. Although they can theoretically achieve greater specific capacity and energy density, the positive electrode materials available are just a few compounds, affecting their overall development. In addition, the passivation of the Mg anode greatly prevents a further transport of Mg2+. AIBs are in the primary development stage because of the formation of Al2O3 layers on the anode, leading to the corrosion of aluminum electrodes, decreased battery efficiency and low cycling stability.Among all available energy storage systems, aqueous zinc-ion batteries (AZIBs) are considered as a most promising large-scale energy storage due to their lower cost, higher cycling stability, aqueous electrolyte, and safe battery manufacturing process. AZIBs can directly use metallic zinc as a negative electrode material, and they also have unique advantages, such as high theoretical capacity (i.e., 820 mA·h/g and 5 855 mA·h/cm2), low redox potential (i.e., -0.76 V vs. SHE), abundant crustal reserves (i.e., 300 times greater than that of lithium), and good stability in the presence of oxygen and humid atmosphere. Multivalent AZIBs allow multiple electron transfers during the electrochemical reaction, providing an opportunity for achieving high energy and power density. In addition, compared to flammable organic electrolytes, the aqueous electrolytes typically used in AZIBs have lower cost, safer characteristics, and higher ionic conductivity.AZIBs have been developed since the invention of the Voltaic Pile (Zn-Ag) by Alessandro Volta. In 1870-1969, a large-scale production led to the invention of zinc-carbon batteries (i.e., Zn|NH4Cl|MnO2) and alkaline zinc-manganese batteries (i.e., Zn|KOH|MnO2), which are still used. In 1986, Yamamoto used weak acidic electrolytes instead of alkaline electrolytes to improve the reversibility of zinc-manganese batteries. In late 2011, Kang proposed a concept of “zinc-ion batteries” and confirmed the reversible insertion/extraction reaction of zinc ions in manganese dioxide in a zinc sulfate electrolyte system. Since then, various rechargeable battery systems based on near-neutral aqueous electrolytes (i.e., zinc-manganese, zinc-vanadium, zinc-cobalt, and zinc-iodine) have been developed. The existing work on AZIBs deals with aqueous solutions containing ZnSO4, ZnCl2, or Zn(CF3SO3)2 as an electrolyte with a high safety and a reversibility.There are numerous reports on the extensive research on AZIBs, i.e., zinc metal anode, cathode materials, electrolyte engineering, and potential applications. However, there exist still some challenges in the development of these batteries. For instance, the cathode materials suffer from the issues like structural instability, poor conductivity, and dissolution. The zinc anode has inevitable dendrite formation, hydrogen evolution reactions, surface corrosion, and passivation. The aqueous electrolyte has a low voltage window and corresponding parasitic reactions. In addition, most studies are also conducted under mild laboratory conditions without considering the entire battery system, including inactive components.Summary and prospects Rechargeable aqueous zinc-ion batteries based on neutral or weakly acidic electrolyte systems have been developed. However, there are still some challenges in the current battery systems such as low energy density and short cycle life. This review focused on the fundamental scientific issues associated with some aspects of AZIBs, and provided a comprehensive summary of the latest advancements in cathode materials, anode materials, electrolyte materials, and inactive component materials (such as separators, current collectors, and binders). The core issues and research strategies associated with each component were discussed., Some perspectives on the fundamental issues for high-performance water-based zinc-ion batteries were proposed based on the battery structure and electrochemical operating mechanisms.The existing work on AZIBs are mostly carried out in the laboratory, mainly single aspects of evaluation and lacking comprehensive assessments. From small button cells to large-scale pouch cells and prismatic cells, some issues and defects could be magnified, and the impact of side reactions on the electrochemical performance became more apparent. The future research and application of aqueous zinc-ion batteries could require a further exploration of their fundamental issues and multivariate optimization of battery performance.

Lithium-sulfur (Li-S) battery is considered as a next generation energy storage system with a great application prospect because of its high theoretical energy density (i.e., 2 600 W·h/kg), low cost and environmental friendliness. However, the shuttling of lithium polysulfides (LiPSs) is considered as one of the main bottlenecks hindering the commercial application of Li-S batteries. Conventional physical confinement and chemisorption are proved to be important strategies to solve a problem of shuttle effect, but they are still “passive” solutions and cannot solve the problem from the source. The shuttle effect is due to the slow transformation of the liquid intermediate LiPSs (i.e., Li2Sn, 4≤n≤8) to the solid product Li2S2/Li2S, which accumulates continuously in the electrolyte and diffuses to the negative electrode driven in the concentration gradient and electric field. The introduction of high-efficiency catalysts can reduce the energy barrier of sulfur conversion, accelerate a “solid-liquid-solid” conversion process, improve the utilization of sulfur, and inhibit the shuttle effect, playing a key role in improving the capacity and cycling performance of the battery. Therefore, a further understanding the sulfur conversion process in Li-S battery chemistry and putting forward the fundamental method to promote the rapid conversion of LiPSs and the uniform deposition of solid products from the perspective of electrocatalytic reaction kinetics can have important theoretical value and practical significance for promoting the practical development of Li-S battery.In the chemical industry, catalysts can reduce the activation energy of chemical reactions and increase the reaction rate. The desulfurization technology in conventional sulfur chemistry gives inspirations for the catalysis in Li-S battery. The key to this technology is that the C—S bond in petroleum distillate is broken by multi-step catalytic hydrogenation, and sulfur is gradually removed in the form of H2S. The charge-discharge process of Li-S battery involves the conversion of solid phase S8, liquid phase polysulfide (i.e., Li2Sn, 4≤n≤8) and solid phase Li2S2/Li2S. The reaction process is similar to the industrial step-by-step desulfurization process.In this review, from the perspective of battery chemistry for Li-S batteries, the origin of the catalysis for Li-S batteries, catalytic mechanism and research progress of catalysts were summarized, and the activity descriptors guiding the rational design of catalysts for Li-S batteries were proposed. Finally, the future development of Li-S batteries was prospected. The idea of industrial desulfurization was introduced into Li-S battery to realize the combination of conventional chemical engineering and Li-S battery chemistry. The “lithium-sulfur catalysis” strategy is demonstrated as a “proactive” solution to the problem of shuttle effect of Li-S batteries, accelerating the transformation of polysulfides and reducing their accumulation in electrolyte, which is expected to fundamentally inhibit the shuttle effect. Since 2014, transition metal oxides, sulfides, nitrides and carbides have been developed as effective catalysts for Li-S battery. In addition, heterostructure catalysts with different functions have been also developed with multi-functions. With the comprehensive understanding of the sulfur redox process, selective catalysis, bidirectional catalysis and homogeneous catalysis have been proposed. Recent development on quantum chemistry and AI technology provides a novel research method for reasonable prediction and analysis of the properties and reaction process of catalysts. Several key activity descriptors, such as the lattice matching, orbital hybridization and p charge density are developed to guide the rational design of catalysts. However, there are some differences in the performance of different types of catalysts, the relevant internal reasons still need to be further revealed, and the universality of descriptors also need to be further explored.Summary and prospects “Lithium-sulfur catalysis” is considered as an important strategy to solve the bottleneck of practical application of Li-S batteries. Although a variety of catalytic material systems are developed, the catalytic mechanism still remains unclear, and advanced theories need to guide the rational design of high-efficiency catalysts. Therefore, it is extremely urgent to develop a novel lithium-sulfur battery system and explore an effective research paradigm of Li-S catalysis for the conversion reaction process of Li-S battery. Future research directions focus on three aspects: 1) The failure mechanism and stability of catalysts should be investigated; 2) The structure-activity relationship between the electronic structure and catalytic performance of the catalyst should be deeply explored by AI technology; and 3) The conversion mechanism of solid-state Li-S battery should be clarified, and a novel practical solid-state Li-S battery system should be constructed for some future extreme applications.

Electrochemical energy storage (EES) is an essential technology in modern lifestyle. As two typical EES devices, batteries and supercapacitors have their own advantages and disadvantages. Pseudocapacitive materials have attracted recent attention. Such materials store charges via battery-like redox reactions, but have supercapacitor-like electrochemical behaviors. Therefore, pseudocapacitive materials have a considerable potential to simultaneously achieve both high energy and high power density as well as excellent stability, merging a unbridgeable gap between batteries and supercapacitors. However, the development of pseudocapacitive materials still encounter some serious challenges such as unclear energy storage mechanism, lack of alternative pseudocapacitive materials and vague future development directions. This review represented recent development on pseudocapacitive materials and briefly discussed their energy storage mechanism and the imperfections of the existing issues. Furthermore, we proposed a concept of bulk pseudocapacitance and realized a goal for the future development of pseudocapacitive materials. In addition, we also discussed some key factors and strategies for achieving such a bulk pseudocapacitance.The energy storage mechanisms of pseudocapacitive materials were discussed. The first pseudocapacitive material, i.e., RuO2 thin film, was reported by Trasatti et al. in 1971. The thin film electrode exhibited a supercapacitor-like cyclic voltammetry curve but differed from the electric double layer in faradic origin. Although the electrochemical performance of RuO2 electrodes were enhanced via fabricating various nanostructures and introducing structural water, the energy storage mechanism of this electrode was not fully understood, and limited to reversible redox reaction at the surface, the content of structural water in crystal, fast proton transport and crystallinity of electrode, etc.. MnO2 is another typical pseudocapacitive electrode, which was firstly reported by Lee and Goodenough in 1999. Simon and Gogotsi reported the pseudocapacitive response of this electrode to successive multiple surface redox reactions. Recently, Augustyn et al. indicated the pseudocapacitive response to the existence of nanoconfined interlayer structural water, mediating an interaction between charge carrier and electrode. Dunn et al. investigated the pseudocapacitive response for T-Nb2O5 electrode and proposed a model of pseudocapacitance, i.e., intercalation pseudocapacitance. Mxenes is one of pseudocapacitive materials with a unique structure that is responsible for its pseudocapacitive behaviors. Although it is still controversial whether it is a pseudocapacitive electrode, Co(OH)2 demonstrates a coupled performance of batteries and supercapacitors. The energy storage mechanism of this electrode was investigated by in-situ X-ray adsorption spectroscopy. It was found that the high specific capacitance originates from a battery-mimic bulk reaction, while the high stability performance roots in the crystal structure stability in charge-discharge process. Furthermore, we highlighted that the Co(OH)2 electrode blurs the distinction between supercapacitors and batteries, which is in accordance with the continuous transition from electric double layer to battery proposed by Fleischmann and co-workers. In future, more intensive studies should be conducted on the energy storage mechanism of pseudocapacitive materials, especially on the unique characteristics of charge carrier and electron transportation, charge transfer, interaction between charge carrier and electrode in an atom-level.The energy density of pseudocapacitive materials is still far below than the actual demand, even the batteries. To address this issue, a concept of bulk pseudocapacitance is proposed. In fact, intercalation pseudocapacitance depicted by Dunn et al. is a typical example of bulk pseudocapacitance. Bulk pseudocapacitance refers to the pseudocapacitive response that happens throughout the bulk of the electrode material rather than the surface. Meanwhile, the proposal of bulk pseudocapacitance can provoke a research on boosting the pseudocapacitive response of battery materials to alleviate the lack of the alternative pseudocapacitive materials. Reducing the size of active materials to nano-scale, i.e., nanostructured material, as one of the strategies, can be used to achieve the bulk pseudocapacitance. For instance, LiCoO2 nanocrystalline with the sizes of ?11 nm exhibits a pseudocapacitive response. In addition to nanostructured materials, a strategy named diffusionless-like conversion reaction to achieve pseudocapacitive response in conversion reaction was proposed. Based on the conversion reaction between Fe(OH)2 and δ-FeOOH, a battery-like conversion reaction achieves a high specific capacitance, and a diffusionless-like transformation between charge and discharge phases without a massive atom movement presents high rate and stability performance. For battery materials, their pseudocapacitive response can be enhanced via manipulating the charge carrier transport, electron transfer, and designing the interaction between charge carrier and electrode in atom-scale. In addition, such a provoking pseudocapacitive response in battery materials extends the scope of bulk pseudocapacitive materials. Li and co-workers enhanced the Li+ transport by pre-intercalation of NH4+ into interplay of Mo2CTx. The mechanism of FeHCF storing NH4+ is clarified, indicating that the weak Fe—N interaction mitigates the volumetric expansion induced by NH4+ insertion. These findings demonstrate that modulating on charge carrier transport and interaction between charge carrier and electrode can improve the pseudocapacitive response of battery materials. Summary and prospects Future development of pseudocapacitance relies on a deep understanding of the energy storage mechanism of pseudocapacitive materials in an atomic level, i.e., the charge carrier and electron transport, charge transfer, and the interaction between charge carrier and electrode. In addition, exploiting and designing novel pseudocapacitive materials as well as boosting pseudocapacitive response of existing battery materials are crucial for the application of pseudocapacitive materials.

As the world grapples with a persistent issue of fossil fuel overconsumption, two prominent challenges have emerged on the global stage, i.e., energy crises and environmental pollution. These hurdles cast a shadow over humanity’s pursuit of sustainable development, underscoring an urgent need for renewable and clean energy sources. Among the array of technologies aiming to tackle these challenges, semiconductor-based photocatalysis technology shines as a promising avenue. This technology holds a key to harnessing solar energy and converting it into chemical energy, offering a promising application potential. However, the widespread adoption of this technology is hindered due to the inefficiency of single-component photocatalysts. This inefficiency is fundamentally since photogenerated electron-hole pairs within these single-component photocatalysts recombine readily.Some heterojunction photocatalysts are developed to address this limitation. Constructing such photocatalysts is considered as a pivotal approach to preventing the recombination of photogenerated electron-hole pairs, thus enhancing the photocatalytic performance. Recent emergence of S-scheme heterojunction photocatalysts, featuring reduction photocatalysts (RP) and oxidation photocatalysts (OP), is an effective strategy for averting the recombination of these electron-hole pairs.The conduction band (CB) and Fermi level (Ef) positions of RP surpass those of OP. When OP intimately contacts RP, electrons from RP migrate to OP through the interface until equilibrium is reached. This leads to an upward bending of the energy band near the RP interface due to electron depletion. Conversely, energy bands near the OP interface bend downward due to electron accumulation. This creates a built-in electric field (IEF) at the interface, pointing from RP to OP. Furthermore, the Ef of RP gradually decreases from RP's bulk to the interface, while the Ef of OP gradually increases from OP's bulk to their interface, until they meet at the same point. Under irradiation, electrons in OP and RP are excited from their valence band (VB) to CB, respectively. The built-in electric field at the interface propels the transfer of photogenerated electrons in the CB of OP to the VB of RP. Simultaneously, Coulombic repulsion, band bending, and the built-in electric field inhibit the transfer of electrons from the CB of RP to the CB of OP (i.e., hole transfer from VB of OP to VB of RP). In this scenario, the original high reduction ability of photogenerated electrons in RP remains, and the original high oxidation ability of photogenerated holes in OP preserves. Consequently, this heterojunction ensures an efficient charge separation and augments the redox capabilities of charge carriers, ultimately enhancing a photocatalytic performance. From a macroscopic perspective, the electron transfer is akin to ascending a “staircase”. Accurately characterizing electron transfer at the interface of S-scheme heterojunctions is vital for understanding photocatalytic mechanisms and for providing experimental and theoretical guidance for the preparation of high-efficiency photocatalysts.Recent methods are developed to investigate the charge transfer behavior in S-scheme heterojunctions. These strategies are categorized into direct and indirect verification of S-scheme heterojunctions. Direct methods include in-situ irradiation X-ray photoelectron spectroscopy (ISI-XPS), zeta potential measurements, and femtosecond transient absorption spectroscopy (fs-TAS). ISI-XPS and surface potential measurements directly evaluate electron accumulation or depletion by assessing relative energy shifts and electrostatic surface charge distribution. In contrast, fs-TAS tracks ultrafast electron transfer processes at the RP/OP interface. Indirect methods encompass work function measurements, electron paramagnetic resonance techniques (EPR), selective deposition of metal nanoparticles, and photocatalytic reactions. These indirect or complementary methods provide the validation of charge transfer behavior within S-scheme heterojunctions.Summary and prospects This review represented various methods used to explore the electron transfer mechanism within S-scheme heterojunctions. Each method was introduced with a fundamental explanation of the mechanism, followed by practical application examples. Some challenges persisted in elucidating the electron transfer mechanism within S-scheme heterojunctions. Despite the progress in characterization techniques for electronic transfer mechanisms, understanding the S-type heterojunctions and optimizing the photocatalytic systems need some techniques with atomic-scale resolution to analyze the dynamics of photogenerated charge carriers from a microscopic perspective. Such techniques, like ISI-XPS, zeta potential measurements and fs-TAS, should be complemented by advanced methods for in-situ characterization, providing research avenues for understanding the photocatalytic mechanisms of S-scheme heterojunctions in solar energy conversion.

With the rapid development of cutting-edge technology fields such as aerospace and defense, high-temperature structural materials are needed for key hot-end components. High-entropy alloy and ceramic eutectic composites are a kind of eutectic composites composed of high-entropy alloy and ceramic, having high density and high-temperature properties. The eutectic composites exhibit superior high-temperature strength and room-temperature plasticity via combining the performance advantages of high-entropy alloys, ceramics, and eutectic composites, which have attracted recent attention.High-temperature high-entropy alloy and ceramic eutectic composites are mainly composed of refractory metal elements and high-melting non-metallic elements such as C, Si and B, which have a high density up to 13.31 g/cm3. The phase composition and eutectic structure can be effectively predicted by a phase diagram software named CALPHAD, while the predicted difference in eutectic points originates from the limited simulation databases. The eutectic composites are composed of refractory body-centered cubic (BCC) high-entropy alloy and ceramic phase with a eutectic structure, in which the two phases exhibit a semi-coherent interface with a low interface energy and a high stability. After annealed at 1 573 K for 168 h or deformed at 1 473 K, the eutectic composite maintains a phase and structural stability. The high-temperature high-entropy alloy and ceramic eutectic composites have superior high-temperature strength and room-temperature plasticity via combining the synergistic effects from strong bonding interface, fine grain strengthening, and second phase strengthening, which are better than the reported high-temperature high-entropy alloys and their ceramic composites. In addition, the high-temperature high-entropy alloy and ceramic eutectic composites derived from the optimization of alloying elements also possess high-temperature oxidation resistance, wear resistance and corrosion resistance.Summary and prospects The high-temperature eutectic composites composed of high-entropy alloy and ceramic have superior eutectic structural stability and performance such as high-temperature strength, which are expected to be applied in the high-temperature fields. However, the corresponding research on these eutectic composites is still in infancy, a relationship among composition, structure and performance is relatively unclear. In the future, the development of these eutectic composites is mainly carried out from five aspects, i.e., theoretical design and computational simulation, synthesis and preparation of new materials, multi-scale structure design, high-temperature oxidation resistance of enhanced materials, and improvement of mechanical properties of materials.

Introduction During the long-term service of the tunnel in high-geothermal environment, one side of the lining structure is in contact with the high-temperature rock wall, and another side is the normal temperature opening area of the tunnel. The temperature field is formed inside the lining concrete. The durability deterioration of lining concrete is the result of the coupling effect of temperature field and sulfate attack. The existence of temperature field under the coupling effect accelerates the diffusion of SO42- from the outside to the inside of concrete as well as the chemical reaction and damage degree of attack. The deterioration mechanism of sulfate attack in high-geothermal environment is different from that in normal temperature environment. However, an indoor simulation test in high-geothermal environment adopts a hot sulfate solution soaking method. Although this method can achieve the acceleration of high temperature, it cannot simulate the effect of temperature field. The concrete corner is subjected to two-way high temperature under the method of hot sulfate solution immersion, resulting in more serious corner deterioration. Therefore, the scientific basis and rationality of the hot sulfate solution soaking method are still debatable. In addition, the pore structure characteristics of concrete are a key to affecting sulfate attack. The existing research on the pore structure of sulfate attack concrete mostly adopts the index of pore size distribution and its corresponding pore content, which cannot fully reflect the mechanism of sulfate attack. Fractal theory is proved to quantitatively characterize the complexity and irregularity of the pore structure of cement-based materials, and can also link the pore structure characteristics of cement-based materials with macroscopic properties. Therefore, this paper proposed a durability simulation system for the coupling of one-dimensional temperature field and sulfate attack. A fractal dimension method was used to investigate the pore structure characteristics of lining concrete under the coupling of temperature field and sulfate attack. Based on the layered damage theory, the relationship between the pore structure characteristics and the average compressive strength of the damaged layer, and the overall concrete strength under the coupling effect were analyzed.Methods A Portland cement 42.5 (GB175—007 standard) and a fly ash with grade II (GB/T 1596—2017 standard) were used. The fineness modulus of the river sand was 2.7, and the coarse aggregate was gravel with a gradation of 5-25 mm. A water reducing agent was polycarboxylate as a superplasticizer, and the water reducing rate was 25%. According to the actual engineering test, a mix ratio is cement of 320 kg/m3, fly ash of 80 kg/m3, water-binder ratio of 0.38, as well as sand and gravel of 740 kg/m3 and 1 108 kg/m3. After de-moulding, the specimens were placed and cured in a curing box for 28 d at different curing temperatures (i.e., 40, 60 ℃ and 80 ℃, respectively). After reaching the curing age, the durability test of specimens was carried out in a self-designed high geothermal environment attack test device. The wet sand containing SO42- solution was used to simulate the attack environment. The wet sand and heating system were placed on the top of the test block. The coupling effect of single-sided attack and one-dimensional temperature field from top to bottom was realized. Water was added every 12 h to keep the humidity of the wet sand constant. The wet sand changed every 7 d to keep the salt concentration constant. The upper part of the wet sand was sealed with a plastic film to prevent water evaporation.Results and discussion The change of compressive strength of concrete under the coupling effect of temperature field at 40 ℃ and 60 ℃ and sulfate attack are divided into the early growth and the subsequent decline. The coupling effect is 30.0% and 23.6% lower than that of single factor. However, the compressive strength of the coupling effect of temperature field at 80 ℃ decreases with a maximum decrease of 22.5%. The temperature field at 40-80 ℃ accelerates the attack chemical reaction, but a greater initial damage of concrete under the temperature field at 80 ℃ occurs. The most probable pore size and pore size distribution of concrete increase with the increase of temperature field. The large pores and transition pores in concrete are dominant, which are 40.38%-46.72 % and 31.09%-38.51%. Under the coupling effect of temperature field at 60 ℃, the content of gel pores and capillary pores decreases, while the content of macropores increases with the increase of attack depth. The surface of transition pores and capillary pores of concrete under coupling action has fractal characteristics, while the surface of gel pores and macropores does not have fractal characteristics. The pore tortuosity also has fractal characteristics in different pore size ranges. The Dt of gel pores increases with the increase of temperature field. The Dt of macropores decreases with the increase of temperature field, but the Dt of transition pores, capillary pores and macropores does not change with the depth. The influence of the surface roughness of the transition pore and the large pore on the average compressive strength of the damaged layer concrete is dominant. The surface roughness of the gel pore has a less influence. On the contrary, the complexity of gel pore structure has a dominant effect on the average compressive strength of damaged layer concrete. The influence of fractal dimension of pore structure on the strength of integral concrete is not consistent. There is a negative linear correlation between Ds of macropores and the compressive strength of integral concrete. However, the influence of Ds of capillary pores, transition pores and gel pores on the compressive strength of integral concrete is dominant, which is different from the results of previous studies. However, the Dv of macropores, capillary pores and transition pores is negatively linearly correlated with the compressive strength of the whole concrete.Conclusions The overall compressive strength of concrete under the coupling effect of temperature field at 40-80 ℃ was 22.5%-30.0% lower than that of single factor. The proportion of gel pores in concrete was decreased by 43.17%, and the capillary pores were increased by 573.10% as the temperature field was increased from 40 ℃ to 80 ℃. The surface of transition pore became more rougher, and the spatial structure of gel pore and transition pore was more complex than that of capillary pore and macropore under the coupling action. The tortuosity of gel pore, transition pore, and capillary pore increased but the macropore decreased. The surface roughness of transition pores and macropores, and the structure complexity of gel pore had a dominant influence on the concrete damaged layer average compressive strength. There was a negative linear correlation between the fractal dimension of pore volume of macropores/capillary pores/transition pores and the overall concrete compressive strength.

Introduction In the marine environment, reinforced concrete components are extremely vulnerable to corrosion by harmful ions in seawater, resulting in a serious damage. The impressed current cathodic protection method is one of the effective methods for corrosion protection of reinforced concrete structures. Conventional anode materials such as stainless-steel mesh, magnesium mesh and zinc alloy materials can effectively prevent the corrosion of passivated steel bars, but metal anode materials have some problems such as difficult laying operations, high-cost, susceptibility to corrosion and poor long-term durability. Thus, non-metallic composites, especially conductive cementitious materials can be used as promising anode materials. Conductive cementitious materials are simple to prepare, low-cost, adaptable to different concrete surfaces, and have a possibility of being reused. In this paper, a high-strength, low-permeability cement-based anode material was proposed due to its surface protection and conductive anode functions. Two kinds of cement-based anode materials were prepared in the laboratory, which were high-strength conductive mortar materials mixed with carbon fibers (CF) and carbon nanofibers (CNF) as well as conductive mortar materials embedded with carbon fiber net cloth, respectively. The mechanical properties, resistivity, and impermeability of the two anode materials were tested, and the evolution of the properties under the ICCP accelerated electric field was investigated.Methods The cementitious material of the mortar matrix was composed of P.O 42.5 cement and silica fume, the particle sizes for two types of quartz sand were 0.11-0.21 mm and 0.21-0.38 mm, a high-performance polycarboxylate superplasticizer was used, and a water-binder ratio is 0.26. In the carbon fiber conductive mortar, carbon fibers (in volume fractions of 0%, 1.0%, 1.5%, 2.0%) and carbon nanofibers (in volume fractions of 0%, 0.6%, 0.9%, 1.2%) were mixed. The single layer carbon fiber net cloth with a mesh size of 5 mm and was placed vertically on the surface of the specimen. The anode materials were cured at 90 ℃ for 72 h.The anode materials were subjected to the flexural and compressive strength tests (GB/T 17671—2021). Electrical resistivity tests (two-electrode method) were conducted on anode materials cured for different durations. Also, gas permeability tests and one-sided water absorption tests (EN 13057) were carried out to assess the permeability resistance of the anode materials. The performance degradation of the anode materials under the influence of electrical current was tested through cathodic protection acceleration experiments. The designed current density for external electrical current was 20 mA/m2. Acceleration tests were conducted by increasing the current density to 100 times (i.e., 2 000 mA/m2) and 150 times (i.e., 3 000 mA/m2). The corrosion status of the rebar was determined by a model ASTM C876 standard test method to assess the effectiveness of the anode materials in the cathodic protection test. The microstructure of the anode materials after ICCP acceleration test at an acceleration voltage of 30.0 kV was characterized by a model VEGA3 TESCAN scanning electron microscope.Results and discussion For carbon fiber conductive mortar, the mixing of CF with CNF improves the mechanical properties and conductive properties of the carbon fiber conductive mortar matrix. The compressive strength and flexural strength of mortar with 1.0% CF + 0.6% CNF are increased by 28% and 66%, respectively. There is a binding effect between CF and matrix, which can significantly improve the bending strength and increase the toughness of the material, CNF has a nucleation effect and a filling effect, which refines the pore structure inside the material, and the bridging effect of CNF improves the mechanical properties of the material. The resistivity of the mortar fiber-doped with 2.0% CF + 1.2% CNF is 160 Ω·cm, which is as low as nearly one ten-thousandth of that of the mortar fiber-undoped, this is a result of the formation of a conductive network inside the mortar specimen formed via the interconnection of CF and CNF.For carbon fiber conductive mortar with embedded carbon fiber net cloth, the introduction of carbon fiber net cloth improves the mechanical properties of the material. This is due to the fixed distribution of a layer of carbon fiber net cloth in the matrix, which constrains the cracking process of the material. Carbon fiber mesh cloth is woven from carbon fibers. The current conduction path between the fibers is more complete. The introduction of carbon fiber net cloth, CF, and CNF improves the conductivity of the composite mortar materials.Carbon fiber conductive mortar and carbon fiber conductive mortar with embedded carbon fiber net cloth both have superior impermeability properties. Intrinsic gas permeability Kv and capillary water absorption rate S are as low as one-tenth of ordinary concrete, indicating that the prepared anode material has a superior physical protection function.The cathodic protection accelerated test proves the effectiveness of the anode material. The mechanical properties and electrical conductivity of the anode material after energized slightly deteriorates. The compressive and flexural strengths are decreased by only 14.5% and 12.1%, and the resistivity is increased by 18.1%. The anode impermeability of carbon fiber conductive mortar decreases. The impermeability of a conductive mortar embedded with carbon fiber mesh cloth and carbon fiber as an anode material does not change significantly because carbon fiber mesh cloth has a main conductive effect to protect the integrity of the substrate.Conclusions The cement-based anode materials both had the superior mechanical properties, i.e., the compressive and flexural strength of 103.0 MPa and 14.6 MPa, respectively, the resistivity of 160 Ω·cm, the gas permeability and capillary water absorption rate of only 1/10 of the ordinary concrete. The concrete structure had a positive physical protection effect. After the ICCP acceleration, the mechanical properties and electrical conductivity of the two materials only produced a slight deterioration phenomenon, in which the anode material with carbon fiber net cloth performed more consistently. It could be suitable for cathodic protection of concrete with applied current. The cement-based anode materials can be also used for electrochemical chloride extraction of concrete components under high chloride salt environment conditions instead of the conventional anode materials.

Introduction The existing alkali-activated repair materials cannot adapt to the corrosion repair of concrete in marine engineering with high humidity and high salt environments. In this paper, a polyacrylic acid lotion (PAA) modified alkali-activated slag repair material with high strength, low shrinkage and high durability was developed by adding PAA to regulate the flexibility of the consolidation body, and its performance formation mechanism was analyzed. The results indicate that the mechanical properties of alkali-activated slag repair materials are positively correlated with the content of sodium silicate added when the modulus of sodium silicate is greater than 1.2. The optimal activating ratio for slag repair materials is 1.2 modulus and 15% content of sodium silicate. PAA filling in the gaps of amorphous cementitious substances in alkali-activated slag repair materials significantly enhances its mechanical properties, bonding properties, volume stability, and sulphate resistance.Methods A powder used was cement (i.e., P·II 52.5 grade Portland cement) mixed with slag (i.e., a granulated blast furnace slag) with a hydraulic coefficient of 2.18, an activity coefficient of 0.25, an alkalinity coefficient of 1.11, and a quality coefficient of 2.07. Alkali activator used was sodium silicate (with a modulus of n=3.3 and a Baume degree of 40) and sodium hydroxide(granular sodium hydroxide with a purity of ≥98%). The organic lotion was a polyacrylic lotion with a molecular weight of 3 000-5 000 and a solid content of ≥30%. The aggregate was river sand with the maximum particle size of of 2.5 mm, the bulk density of 1.6 g/cm3, and the apparent density of 2.5 g/cm3. Water used was tap water.Sodium hydroxide weighed was fully dissolved in sodium silicate under stirring magnetically. The solution was colourless and transparent, and there were no suspended sodium hydroxide particles in the solution. Subsequently, a beaker with the solution was covered with a thin film to prevent the evaporation of water due to the dissolution and heat release of sodium hydroxide. The solution was cooled to room temperature before use. The prepared solution was poured into a mixing pot, and mixed with the powder under stirring. After curing until the specified age, the relevant tests were conducted.The physical and mechanical properties of alkali activated slag repair materials were determined in accordance with the provisions of GB/T50081—2019 “Standard Test Methods for Physical and Mechanical Properties of Concrete”. The bond strength was measured via bond bending tests. According to GB/T 50082—2009 “Standard Test Methods for Long term Performance and Durability of Ordinary Concrete”, the shrinkage rate of alkali activated slag repair materials was measured by a contact method. According to GB/T 50082—2009 “Standard Test Methods for Long term Performance and Durability of Ordinary Concrete”, the sulphate attack resistance of the alkali activated slag repair material was determined, and the corrosion resistance coefficient (K) was used as an indicator to measure the sulphate attack resistance of the sample. The phase composition of alkali activated slag repair materials was analyzed by a model D8ADVANCE X-ray diffractometer. The main functional group composition and chemical bond types in alkali activated slag repair materials were characterized by a model TENSOR27 Fourier infrared spectrometer. The microstructure of alkali activated slag repair materials was characterized by a model ZEISS Gemini 500 scanning electron microscope.Results and discussion Increasing the modulus of sodium silicate can improve the strength of alkali activated slag repair slurry. Adding 15% sodium silicate with a modulus of 1.2 is an optimal excitation approach for slag repair materials. When the content of PAA is 1%, the flexural strength and compressive strength of PAA modified alkali activated slag repair slurry are increased by 7% and 14%, respectively, compared to alkali activated slag repair slurry, reaching 9.90 MPa and 57.87 MPa. High-temperature curing at 50 ℃ further improves the strength of alkali activated slag repair mortar, and the strength improvement of PAA modified alkali activated slag repair mortar is greater than that of unmodified mortar. After PAA modification, the bonding strength of alkali activated slag repair mortar is improved, and its resistance to sulphate attack is enhanced. The bonding strength of PAA modified alkali activated slag repair mortar is greater than 2 MPa, and the corrosion resistance coefficient is greater than 0.85, meeting the requirements of JC/T 2381—2016 for the bonding strength and corrosion resistance coefficient of repair mortar. The shrinkage rate of alkali activated slag repair mortar at all ages increases with the increase of sodium silicate modulus. The PAA modification further reduces the shrinkage rate of alkali activated slag repair mortar. When the PAA content is 1%, the shrinkage rate of alkali activated slag repair mortar is reduced by 54%. Conclusions The early flexural strength development of alkali activated slag repair slurry was slower, while the early compressive strength development was faster. Adding a small amount of (10%) sodium silicate could stimulate the activity of the slag, slightly improving the strength of the slurry, and significantly increasing the strength at 15% and 20% dosages. The addition of PAA did not change the product composition in the alkali activated slag repair material, but filled in the gap of amorphous cementitious material. Before PAA modification, the cross-section of the alkali activated slag repair material was smooth and flat. After PAA modification, the cracks on the both sides of the alkali activated slag repair material were connected by fibrous filaments, further enhancing the strength of the material.

Introduction Steel corrosion is particularly severe in marine environments with a high concentration of chloride ions, which causes huge economic losses and also hinders engineering construction in coastal areas. Coating treatment is one of the common methods to protect steel from corrosion. Conventional thicker organic coatings are prone to local breakage, reduced mechanical interlocking force between steel and concrete, and being easy to deteriorate. Since the anions and cations of layered double hydroxides (LDH) both are tunable, this designability gives them a variety of potentials for metal corrosion protection, such as physical barrier, chlorine fixation, corrosion inhibition and chemical self-repair. In this paper, Mg/Al-NO3?-LDH and Zn/Al-NO3?-LDH films were grown in-situ on the surface of steel Q235 by an electrodeposition-hydrothermal method, and their adaptability in a simulated high-alkaline concrete environment and corrosion inhibition performance in NaCl solution were investigated.Methods The in-situ growth of LDH film on steel surface includes electrodeposition and hydrothermal treatment. In the electrodeposition, a three-electrode system was used, with a steel sheet as a working electrode, an Ag/AgCl electrode as a reference electrode, and a platinum sheet as a counter electrode. The electrolyte solution consisted of 50 mL of 0.045 mol/L Mg(NO3)2·6H2O and 0.015 mol/L Al(NO3)3·9H2O, and the steel sheet was immersed in the electrolyte solution to stabilize it for 500 s, and then a potential of ?1.5 V was applied for 300 s. The hydrothermal treatment used a hydrothermal solution consisting of 50 mL of 0.06 mol/L Mg(NO3)2·6H2O and 0.02 mol/L Al(NO3)3·9H2O. NH3·H2O was added dropwise to adjust the pH value of the solution to 9. Subsequently, the electrodeposited steel sheet was vertically placed into a Teflon?-lined autoclave containing the hydrothermal solution, and reacted in a homogeneous reactor at 90 ℃ for 18 h. After the reaction was completed and cooled down to room temperature, the sheet with the film was removed and named M1. Similarly, the obtained sample was named Z1 when replacing Mg(NO3)2·6H2O with Zn(NO3)2·6H2O. To examine the adaptability of the two samples in a high-alkaline concrete environment, the samples M1 and Z1 were immersed in saturated Ca(OH)2 solution for 24 h and then taken out, resulting in samples named M2 and Z2, respectively.The crystal structures of the different samples were characterized using a model D8 Advance X-ray diffractometer (Bruker Co., Germany, λ=0.154 06 nm). Sample’s functional groups and chemical bonds were detected using a model Spotlight 200 Fourier transform infrared spectrometer (PerkinElmer Co., USA). The surface morphology and chemical composition of the samples were determined by a model APREO S high-resolution scanning electron microscope with an X-ray energy spectrometer (Thermo scientific Co., The Netherlands). The impedance and corrosion resistance of the four samples in 3.5% (mass fraction) NaCl solution were analyzed via electrochemical impedance spectroscopy (EIS) and potentiodynamic polarization in a ParStat 4000 electrochemical workstation (AMETEK Co., USA).Results and discussion The electrodeposition and hydrothermal method was employed to in-situ grow a dense and uniform Mg/Al-NO3?-LDH film and a lamellar porous Zn/Al-NO3?-LDH film on the surface of steel Q235. After immersing these two LDH films in a saturated Ca(OH)2 solution for 1 d, the lamellar structure of Mg/Al-LDH remains unchanged, and the interlayer anions are partially exchanged from NO3? to OH? due to the slightly higher affinity of OH? in the solution. Also, the lamellar structure of Zn/Al-LDH undergoes a chemical reaction due to the presence of a large amount of Ca2+ and OH?, resulting in the generated portion of Ca/Al-NO3?/OH?-LDH and amorphous hydroxides remaining on the carbon steel surface. The results above indicate that Zn/Al-LDH is difficult to be stabilized in highly alkaline concrete environments since its constituents, i.e., Zn(OH)2 and Al(OH)3, both are amphoteric hydroxides. In contrast, Mg(OH)2 is an alkaline hydroxide, allowing Mg/Al-LDH to maintain a structural stability in high alkaline environments.The EIS results show that in a 3.5% (in mass fraction) NaCl solution, the impedance of the Zn/Al-NO3?-LDH film is only 3.18 Ω/cm2, whereas the Mg/Al-NO3?-LDH film exhibits a greater impedance of 2 722 Ω/cm2, which is attributed to the dense structure of the Mg/Al-NO3?-LDH film. The results of dynamic potentiodynamic polarization tests indicate that the corrosion protection efficiency of the Zn/Al-NO3?-LDH film and the Mg/Al-NO3?-LDH film for carbon steel is 63.86% and 97.43%, respectively. Clearly, the Mg/Al-NO3?-LDH film is more effective in protecting carbon steel against chloride ion erosion due to its superior physical barrier and the double protection effect of "fixing chlorine".Conclusions Zn/Al-NO3?-LDH and Mg/Al-NO3?-LDH films with a high crystallinity were grown in-situ on the surface of steel Q235 by an electrodeposition-hydrothermal method for corrosion protection. The microscopy showed that the Zn/Al-NO3?-LDH films were lamellar and porous, and the Mg/Al-NO3?-LDH films were dense and uniform. The electrochemical results indicated that in 3.5% (in mass fraction) NaCl solution, the corrosion protection efficiency of Zn/Al-NO3?-LDH film on carbon steel substrate was 63.86%, while the Mg/Al-NO3?-LDH film had a higher impedance and a lower corrosion current density, and its corrosion protection efficiency was 97.43%, effectively protecting the carbon steel against chloride ion erosion. In a high-alkaline saturated Ca(OH)2 solution, the Zn/Al-NO3?-LDH film exhibited an instability, leading to the formation of partially Ca/Al-NO3?/OH?-LDH and amorphous hydroxides. In contrast, the Mg/Al-NO3?-LDH film remained stable, and the anion exchange resulted in the formation of Mg/Al-NO3?/OH?-LDH film.

Introduction Using as supplementary cementitious materials (SCM) is one of the most important research topics in the resource utilization field of steel slag. However, the iron oxide content in steel slag is greater than that of other cementitious materials (i.e., clinker, granulated blast furnace slag, and fly ash). The solid solution formed by FeO and CaO-MnO-MgO (RO phase), hematite, wustite, Ca2Fe2O5(C2F) and other Fe-bearing phases have little contribution to the activity of ground steel slag. Some researchers use reductants such as H2, CO, anthracite, silicon carbide, and carbon powder to reduce iron oxides from molten steel slag into ferroalloys. Aluminosilicate melt is quenched and resultant residual used as SCM. These methods have a challenge of high cost of reducing agents and difficulty in industrial utilization. This paper used residual carbon in coal gangue as a reductant to reduce iron oxides of steel slag. The remaining aluminosilicate was quenched to prepare slag with a high hydration activity, which could achieve the synergistic disposal of two types of industrial solid waste.Methods Based on the chemical composition of steel slag and coal gangue, as well as the fixed carbon content of coal gangue, the amount of coal gangue required for the sufficient reduction of iron oxide in the two raw materials was determined. After mixing ground coal gangue and ground steel slag thoroughly, the mixture was heated to 1 200, 1 300, 1 350, 1 400, 1 450 ℃ and 1 500 ℃ at 40 ℃/min in an argon atmosphere (3 L/min), and then quenched with water after holding for 30 min. The resultant strip off iron particles, and the remaining aluminum silicate is called water quenched residue (WQR). The iron recovery rate was determined according to the chemical composition of the WQR. The mineral phases were determined by an X-ray diffractometer (XRD, Rigaku Co., Japan), and the glass content in water quenched residue was calculated using the Rietveld quantitative analysis method.The paste and mortar samples were prepared by adding 30% (in mass fraction) WQR into the reference cement. The cross-sectional morphology was determined by a model GeminiSEM 500 scanning electron microscope (Zeiss Co., Germany). The elements distribution of the paste samples was analyzed. The heat release curve of hydration for the first 72 h was recorded with a model TAM-Air-8microcalorimeter (TA Instruments Co., USA). The thermal behavior of the pure slurry after hydration for 3 d, 7 d and 28 d were analyzed by a model TGA/DSC1/1600 thermogravimeter (TG, Mettler Toledo Co., USA). The mortar specimens with a dimension of 40 mm×40 mm×160 mm were used for testing flexural and compressive strength after standard curing to the specific age.Results and discussion The reduction temperature is an important factor affecting the iron recovery rate in steel slag and the vitrification of WQR. When the reduction temperature exceeds 1 300 ℃, the reduction rate significantly increases. When the temperature exceeds 1 350 ℃, the Fe2O3 content decreases from 20% in the steel slag raw material to less than 2% in the WQR. The main mineral phases in the reduced residue at 1 100 ℃ are merwinite and gehlenite. When the temperature rises to 1 350 ℃, the diffraction peaks of gehlenite disappear. At 1 450 ℃, the diffraction peaks of merwinite almost disappear. When the reduction temperature is less than 1 300 ℃, the glass content in WQR is less than 60%. When the temperature reaches 1 350 ℃, the glass increases rapidly, approaching 80%. The glass exceeds 90% with further increasing the reduction temperature to 1 450 ℃. The fusion temperature is an important factor for the degree of vitrification of WQR.The hydration heat release curve of cement paste shows that the second exothermic peak of sample prepared with WQR generated at 1 500 ℃ is similar to the peak of the reference cement, but the exothermic peak attributed to the formation of the second ettringite decreases. The second exothermic peak of cement with WQR at 1 400 ℃ is 1 mw/g lower than that of WQR cement at 1 500 ℃. A lower heat release rate is due to the presence of crystalline phases in the WQR at 1 400 ℃. After 72-h hydration, a higher cumulative heat of the paste prepared with WQR at a higher melting temperature occurs. The cumulative heat of WQR cement at 1 400 ℃ and 1 500 ℃ is 71.65% and 85.64% of the reference cement, respectively. This is due to the high reduction temperature leading to a high degree of vitrification, thereby improving the hydration activity of the WQR.The XRD and TG results indicate that the Ca(OH)2 content in the paste gradually decreases with the increase of steel slag reduction temperature and the hydration time of WQR cement. This is because the glass content and pozzolanic activity of WQR increase with the increase of reduction temperature. Consequently, the Ca(OH)2 content of WQR cement decreases.The strength results of the mortar show that the 3-d flexural strength of cement with 30% WQR is almost equivalent to that of the reference cement. The growth rate of flexural strength of WQR cement is higher than that of reference cement with curing time, and its strength reaches 130% of reference cement strength at 28 d. The compressive strength of WQR cement at 3 d is 60% of the reference cement. The compressive strength development rate of WQR cement is higher than that of reference cement. After 28-d hydration, the compressive strength of cement prepared with WQR at 1 400 ℃ reaches 90% of the reference cement strength.Conclusions The recovery rate of Fe could reach 63%~94% by using carbon in coal gangue as a reducing agent to reduce iron oxide of steel slag. The reduction rate increased as the reduction temperature increased. When the reduction temperature was below 1 300 ℃, the mineral phases of the WQR mainly consisted of merwinite and gehlenite. The crystalline phase gradually decreased as the temperature increased. When the reduction temperature was >1 450 ℃, the glass content maintained at 91%. The glass phase content and the pozzolanic activity in WQR increased as the reduction temperature increased. The cement prepared with WQR had good mechanical properties. After 28-d hydration, the flexural strength and compressive strength of WQR cement reached 130% and 90% of the reference cement, respectively.

Introduction Conventional concrete is a porous heterogeneous quasi-brittle material, which has some problems such as poor toughness, low flexural tensile strength and poor durability, and cannot meet the needs of modern construction for high strength, high crack resistance and high durability of concrete materials. The existing methods used to toughen concrete are fiber reinforcement and matrix toughening. Among them, fiber is an extrinsic toughening. The toughness of cement matrix can be improved with polymers and nanomaterials. In addition, compared with mixing polymers directly, the modified method of in-situ polymerization of organic monomers can improve the flexural strength of cement pastes. However, the effect and mechanism of this method on toughening concrete are still unclear. Therefore, the effect of in-situ polymerization of acrylamide (AM) monomer and wollastonite (CS) whisker on the mechanical properties, hydration heat release, phase composition and microstructure of concrete were investigated to provide an idea for toughening concrete matrix.Methods Raw materials were P·Ⅱ 52.5 Portland cement, quartz sand with the fineness modulus of 2.6-2.9 and basalt with the particle sizes of 5-16 mm. The composite modified materials used were all adscititious, in which an organic monomer was acrylamide (AM), an initiator was ammonium persulphate (APS), and a crosslinking agent was N,N'-methylenebisacrylamide (MBA). The inorganic modified material was wollastonite (CS) whisker, and its micromorphology was dense and short rod fibrous. The water-cement ratios (W/C) of concrete and cement paste were 0.47 and 0.40, respectively. The dosages of CS were 1% (in mass fraction), 2% and 4% by the mass of cement. The dosages of AM were 3% and 5% by the mass of cement, the dosages of APS were 3% and 5% by the mass of AM, and the dosage of MBA was 0.1% by the mass of AM.The compressive strength, flexural stress-strain curve and fracture energy of concrete specimens (100 mm×100 mm×100 mm and 100 mm×100 mm×400 mm) were tested after curing under standard curing condition (20 ℃, 95% RH) for 3, 7 d and 28 d. The hydration exothermic curve of the cement paste in the early 72 h was recorded at 20 ℃ by a model TAM-Air micro-calorimeter (TA Instruments Co., USA). After the cement paste was cured for 3, 7 d and 28 d, the phase composition was qualitatively and quantitatively analyzed by a model D8 ADVANCE X-ray diffractometer ( Bruker Co., Germany). The hydration degree of cement was determined by a model TG-209F3 thermal gravimetric analyzer (Netzsch Co., Germany), and the functional groups were characterized by a model Nicolet IS10 Fouri infrared spectrometer (Thermo Scientific Co., USA). The microstructure was observed by a model Quanta 3D FEG scanning electron microscope (Fei Co., USA).Results and discussion The effect of single-doped CS on the flexural strength of concrete is not significant, but it can improve the compressive strength to a certain extent at 3 d and 7 d in the early hydration stage. Adding AM has a negative effect on the bending resistance of concrete at each stage. The addition of APS and MBA can make AM crosslinked and polymerized in cement hydrated matrix, which can significantly improve the peak flexural stress and fracture energy. However, the film forming characteristics of AM polymerization lead to a poor compatibility with cement matrix, directly affecting the hydration process of cement and reducing the compressive strength. Under the synergistic effect of CS and AM polymerization induced by APS, the maximum flexural stress and fracture energy of concrete at 28 d are increased by 31.5% and 91.4%, respectively, compared with control group. Note that the compressive strength of concrete at 28 d remains unchanged.Hydration heat results show that the main exothermic processes of AM monomer polymerization and cement hydration are not synchronized. AM polymerization can immediately release a lot of heat and delay the hydration reaction of cement, while the addition of CS has little effect on the hydration process of cement. Neither CS nor AM polymerization can change the types of cement hydration products, but CS can participate in cement hydration reaction and increase calcium hydroxide (CH) content. Under the action of initiator, AM monomer polymerization can promote ettringite (AFt) formation and reduce CH content. The addition of CS can fill the void in the hydrated paste, and promote the hydration reaction of cement, thus reducing the porosity of concrete. However, AM monomer polymerization can form an organic plasticizing zone, strengthen the bonding force between hydration products such as CH, and interact with needle-shaped AFt, effectively forming an organic-inorganic interpenetration network structure.Conclusions AM in-situ polymerization and CS could improve the toughness of concrete. The 28-d flexural stress peak and fracture energy were increased by 31.5% and 91.4%, respectively, and the compressive strength remained unchanged. CS whiskers could promote cement hydration, exert filling effect and reduce the porosity. AM polymerization could build a plasticized zone in cement pastes and weaken stress concentration in concrete. The synergistic effect enhanced the bonding force between the hydration products, formed an organic-inorganic interpenetrating network structure, optimized the structure of concrete across scales, and provided an effective technology for toughening concrete matrix.

Introduction Although different experiments on the co-corrosion of concrete vulcanization and carbonation are performed, the data are scattered and are not systematic. There are little reports on the mechanism of the performance degradation of cement-based materials due to the coupled effects of high temperature, high humidity, SO2 and CO2. It is thus of a great theoretical and practical significance to scientifically evaluate the durability of industrial building concrete structures and investigate the durability damage mechanism and performance degradation of industrial building concrete structures.In this paper, the neutralization process of cement-based materials under the coupling action of high temperature, high humidity, SO2 and CO2 was investigated with the three levels of cement paste, mortar and concrete. The appearance, neutralization depth, and quality change rules of cement-based materials with different water-cement ratios were analyzed. The porosity and pore solution of cement-based materials were determined by nuclear magnetic resonance (NMR) and electrochemical impedance (EIS). In addition, the neutralization mechanism of cement-based materials was discussed, and a prediction model for the neutralization depth of cement-based materials was also established.Methods The cement used was P.O 42.5. The fine aggregate was a natural river sand with a fineness modulus of 2.34, and the coarse aggregate was a granite crushed stone with a continuous gradation of 5-25 mm. Three types of samples were cement paste, mortar, and concrete. In the test, a prism-shaped sample of 100 mm×100 mm×400 mm was used to test the neutralization depth of the sample, and a cubic sample of 100 mm×100 mm×100 mm was used for quality and microscopic tests, with 3 samples in each group. The tests began after 28 d of standard curing of the samples.A concentration ratio of CO2 and SO2 in the workshop was approximately 20:1, and the temperature and humidity were at 50 ℃ and 70%, respectively. The rapid carbonation test method based on the Chinese National Standard GB/T50082—2009 was used. The ambient temperature of the vulcanization test box was 50 ℃, the relative humidity was 70%, and the SO2 concentration was 0.9% (the maximum concentration of SO2 in this test box was 0.9% due to the limit of the SO2 sensor). For the carbonization box, the ambient temperature was 50 ℃, the relative humidity was 70%, and the CO2 concentration was 20%. The samples were covered with epoxy, leaving only two sides. The sample was placed on a bracket in the vulcanization box, and the distance between each sample was not less than 50 mm. The sample was vulcanized for 1 d and carbonized for 1 d as one cycle. The samples were tested when the cycles reached 0, 2, 4, 7 and 14. The appearance, quality and neutralization depth of cement-based materials were examined. The pore characteristics and pore solution changes in the samples before and after corrosion cycles were analyzed by NMR and EIS.Results and discussion Based on the analysis of cement slurry, mortar, and concrete samples, the neutral depth of cement pure pulp is the largest, and the neutral depth after 14 cycles is approximately 35 mm. The neutral depth of the mortar sample after 14 cycles is approximately 13 mm. The overall neutralization depth of the concrete is the smallest, and the difference is not large. The neutral depths after 14 cycles are 6.42, 6.95 mm and 8.19 mm, respectively. The depth of the neutral depth of the ingredients on the cement base is obvious. The aggregate can reduce gas transmission in the sample. The maximum neutral depth of mortar and concrete is decreased by 59.85% and 77.52%, respectively, compared with cement paste.The pore rate for paste is increased by approximately 3%. For mortar, the pores increase slowly. After 7 cycles, the pore rate is increased rapidly by approximately 8%. For concrete, the pores decrease first and then increase slowly by approximately 1%. For different materials under the common action of carbonization and vulcanization, the pore rate of cement base materials is generally increased.The initial proportion of harmful pores in concrete is larger than that in paste and mortar. However, the increase in harmful pores in mortar and cement paste is significantly larger than that in concrete as corrosion cycle increases. The harmful pores of concrete increase slowly, and the harmful pores decrease significantly after the 14th cycle. Concrete harmful pores are increased by approximately 0.4%. The harmful pores of the mortar are increased by approximately 0.8%. The harmful pores of the cement paste are increased by approximately 0.6%.Conclusions In the early stage of the corrosion cycle, the concentration of CO2 was relatively high, and the reaction was mainly carbonization. In the middle of the corrosion cycle, CaCO3 filled in the pores of the sample so that the reaction slowed down and the mass increased slowly. SO2 could react with calcium-containing substances, resulting in expansion cracks. SO2 could also lead to decalcification of CSH, further increasing the porosity and accelerating the transmission of gas inside the sample. In the later stage, the carbonization reaction speed accelerated, and the sample mass began to increase again. The neutralization depth of different samples changed approximately in the same manner, i.e., firstly increasing rapidly and then increasing slowly as the number of corrosion cycle increases. The porosity changes in the sample at all levels were similar, and the less harmful pores gradually increased and decreased after 14 cycles. For harmful holes and more harmful holes, the cement paste and mortar samples increased gradually, while the concrete samples firstly increased and then decreased after 14 cycles.

Introduction The orientation of steel fibres has an impact on the mechanical properties of ultra-high performance concrete (UHPC). It is difficult for the conventional experimental methods to maintain the same fibre orientation and distribution characteristics even under identical conditions, and there is a lack of mesoscale model that can simultaneously capture the complex mechanisms of crack evolution and fibre-mortar interaction in UHPC. This impedes the quantitative analysis and simulation validation. This paper thus conducted direct tensile tests of UHPC in an external magnetic field to achieve a certain fibre orientation, and the distribution density of fibre inclination angles was obtained via image analysis. The influence of fibre orientation on the fracture processes and stress-strain curves was investigated by a digital image correlation (DIC) technique. Also, a 3D microscale numerical model of UHPC was established to explicitly consider fibre orientation and distribution, single fibre pullout force-slip relations based on inclination angles, and discrete cohesive cracking of mortar. The model was validated by the experiments, the complex 3D fracture process of UHPC was quantified, and the influences of mesoscale factors (i.e., fibre distribution and orientation) on the strain hardening characteristics, crack evolution and final morphology, and fibre bridging effects were investigated.Methods In the experiments, steel fibre orientation was controlled through an external magnetic field, which was created by direct current in a coil loop. The dog-bone specimens were vibrated in the magnetic field during the sample preparation. As a result, the fibres were subjected to the electromagnetic force and underwent a directional distribution, i.e., along the magnetic field direction. Three fibre orientations were a) parallel to the tensile direction, b) perpendicular to the tensile direction, and c) no magnetic field applied. This resulted in three groups of specimens termed aligned-UHPC (AUHPC), perpendicular-UHPC (PUHPC), and UHPC, respectively. A tensile testing device for dog-bone specimens was used to investigate the fibre orientation-dependent mechanical properties of UHPC for subsequent validation of simulations. The loading process was recorded by a high-speed camera, providing data for the fracture process analysis based on the DIC technique. In addition, the cut-off sections were analyzed on specimens AUHPC, UHPC, and PUHPC to obtain the orientation distribution characteristics of fibres for later simulations.In the mesoscale simulation, cohesive elements were pre-inserted in the mortar to simulate the energy dissipation in the fracture process zone and the opening/closing of discrete cracks. Different fibre orientation distributions in specimens AUHPC, UHPC, and PUHPC were generated by a random algorithm. The fibre-mortar interaction was equivalently simulated through fibre constitutive laws, which were curve-fitted from the pullout force-slip relations extracted from single fibre pullout tests for various inclination angles. The mesoscale model was validated via single fibre pullout tests and direct tensile tests. In addition, the mesoscale influencing factors on the complex multiple cracking and fibre bridging mechanisms were also quantified.Results and discussion The experimental results reveal that the specimen AUHPC exhibits the most pronounced strain-hardening behavior and maximum peak stress, which are 2.3 times and 1.4 times greater than those of specimens PUHPC and UHPC, respectively. The first cracking stress of specimen AUHPC is also the maximum, while there is no obvious change in the elastic modulus. The increase in peak stress can be explained by the fracture process using DIC: when fibres are more aligned with the tensile direction, there are more microcracks in the specimen. This is because when the fibre orientation follows the tensile direction, the fibre bridging effect is the more intense, thus having the maximum crack resistance. When the fibre orientation is perpendicular to the tensile direction, the fibre bridging effect is the minium. The final crack patterns of specimens AUHPC and UHPC are complex and tortuous, while the crack paths in specimen PUHPC are smoother and mostly perpendicular to the tensile direction. The distinct crack patterns also indicate the differences in the stress-strain curves. Moreover, from the cur-off sections of specimen AUHPC determined by optical microscopy, the fibres after the test have obvious pullout scratching on the fibre surfaces with spalled mortar clinging to them. In addition, there is no obvious necking at the fibre ends, indicating that the fibre failure is pullout rather than yield fracture. The mesoscale mechanisms (i.e., interfacial failure, mortar detachment, and fibre pullout) indicate a much energy dissipation, which improves the material ductility.The mesoscale simulations predict the fracture patterns and stress-strain curves, which are in reasonable agreement with the experimental results. Based on the analysis of fibre stress evolution, this work can quantify the process of microcrack initiation, multiple cracking, strain hardening, localized major cracks or macrocracks, and gradual closure of other microcracks with increasing the displacement. The stresses of fibres far away from the localization zones decrease, while the stresses of fibres crossing the major cracks continue to increase (i.e., 2 000 MPa) due to fibre bridging. Afterwards, the specimen gradually enters the softening stage until complete failure. The specimen AUHPC exhibits three macrocracks, while specimens UHPC and PUHPC have only one macrocrack each. Besides, the quantities of microcracks in specimen AUHPC are much larger than those in specimen PUHPC, and the fibre average stress level in specimen AUHPC is much higher. This is because when the fibre is aligned with the tensile direction, more stresses can be transferred into the mortar, causing re-cracking, which is a repeated process, leading to multiple cracking in specimen AUHPC. As a result, the stresses of bridging fibres consistently increase. On the contrary, the efficiency of stress transfer for the fibres perpendicular to the tensile direction is the lowest, with the minimum fibre average stress of 94 MPa in specimen PUHPC, while AUHPC and UHPC have 328 MPa and 188 MPa, respectively. It is evident that the fibre orientation has an impact on the fibre utilization level. The simulated stress-strain curves demonstrate that specimen AUHPC has the most pronounced strain hardening and the maximum ductility, while specimen PUHPC is relatively brittle and shows no strain-hardening characteristics. This indicates that aligning the fibres with the tensile direction is an effective way to enhance the strength and ductility of specimen UHPC.Conclusions When the fibre orientation was consistent with the principal stress direction, the fibre had a bridging effect on cracks. Therefore, specimen AUHPC had more microcracks, more tortuous paths, higher peak stress, more pronounced strain hardening, and more energy dissipation. specimen PUHPC had only one single, much smoother crack and showed no strain hardening. In addition, specimen AUHPC had the maximum fibre utilization level with the maximum fibre average stress, thus having the most intense bridging effect. The fibre average stress in specimen PUHPC was the minimum with insufficient ductility and significant brittleness. Therefore, optimizing the fibre orientation could be an effective way to improve the strength and ductility of specimen UHPC. Although the fibre-mortar interfaces were not simulated directly, the proposed mesoscale model with rigid embedment of fibres and equivalent fibre stress-strain constitutive relation could prove computationally efficient in simulating the bridging effects from massive random fibres. Therefore, the model could be used for the optimization of material parameters such as fibre dimensions, shape, distribution and volume fraction.

Introduction As a non-homogeneous material, concrete has a large number of microcracks and even macro-defects on its surface and inside. Cracks in concrete dams can impair the integrity of the dam, change its stress state, and shorten its service life. In practical engineering, most concrete structures are subjected to dynamic loads such as earthquakes and dynamic water pressure. Dynamic loading has uncertainty and complexity, which has an important impact on the load-carrying capacity and durability of concrete structures. Research on the crack resistance of concrete is conducive to delaying and mitigating the cracking of concrete as well as taking measures to stop the cracking of cracked concrete, thus improving the durability and safety of concrete structures.Dam concrete is usually fully-graded concrete, subject to various conditions usually uses wet-screening concrete to carry out experimental research, but its test results do not fully reflect the mechanical properties of fully-graded concrete due to the sieve out of the large aggregate. It is thus necessary to carry out the fracture test of fully-graded concrete. In addition, many concrete high dams in China are constructed in the southwest and northwest earthquake areas, and the existing research on the fracture under dynamic loading is seriously lack. This paper was to carry out tests during rapid loading to obtain the dynamic fracture parameters of fully-graded concrete.Methods Fully-graded concrete specimens were prepared using the same mix ratio as the Baihetan Arch Dam. The 15MN dynamic and static material was used and an axial tensile fracture test was carried out. The specimens had a diameter of 450 mm and a height of 450 mm. The central part of the concrete specimen was cut by a cutting machine to form a ring prefabricated crack with a depth of 45 mm and a ratio of the length of the prefabricated crack to the length of the ligament of 0.2.To compare the fracture properties of fully-graded concrete under static and dynamic loading, a strain rate of 10-6 s-1 was set as the quasi-static condition. According to the range of seismic characteristic strain rate, a strain rate of 10-3 s-1 was set as the dynamic working condition. The test was controlled via displacement loading, and the change process of load and loading point displacement was recorded during the test. At the cracks, one trans-seam displacement extensometer was arranged at every 90° degree along the ring, with a total of four extensometers at a distance of 50 mm, and the direction of the displacement extensometers was parallel to the length of the specimen, which was used to determine the change curve of the crack opening displacement with the load in the loading process. Laser displacement gauges were arranged adjacent to each of the displacement gauges at a distance of 350 mm to measure the overall deformation of the specimen. The fracture toughness was calculated via the experimental and numerical data, where the numerical method was an extended finite element method, and the eccentricity phenomenon that occurred during the test was also taken into account.Results and discussion All the specimens are fractured along the prefabricated crack surface. For the specimens under quasi-static loading, the fracture surface is rough and exposed more coarse aggregate, and the interface between aggregate and matrix stripping mostly appears on the fracture surface, and the fracture of aggregate is rare. In the specimens under dynamic loading, the fracture surface is slightly flat, and in addition to the interface between the aggregate and matrix stripping, a number of coarse aggregates are pulled off in the section. During the slow increase in load, the cracks tend to expand along the weak surface, i.e., the aggregate-matrix interface. Under dynamic loading, the cracks do not have time to develop along the weakest part but expand along the shortest path, leading to damage to the specimen.The increase in the strength of concrete specimens under dynamic loading is because the process of extension along the prefabricated cracks runs directly through the coarse aggregate, without having time to pass through the weak part of the cement mortar-coarse aggregate bond. In essence, the inertial forces and inhomogeneity of the concrete material itself lead to an increase in the strength of the specimens under dynamic loading. Under quasi-static loading, the rate of strain energy release and aggregation in concrete specimens is slower, and crack expansion occurs along the weak part of the material. In contrast, under dynamic loading, the rate of strain energy release and aggregation in concrete specimens accelerates and cracks may propagate through the stronger parts of the material, leading to an increase in strength.Conclusions The peak load, energy absorption capacity, fracture energy, and fracture toughness of fully-graded concrete specimens under dynamic loading were increased by 58.51%, 145.75%, 124.48% and 47.71%, respectively, compared to static loading, while the characteristic length was reduced by 11.67%. The numerical simulation calculated the fracture toughness taking into account the eccentricity phenomenon during the test, which was increased by 71.98% for quasi-static loading and by 60.35% for dynamic loading, compared to the uniform loading method without taking into account the eccentricity phenomenon.

Introduction Carbon dioxide emissions generated during the production of cement-based materials are one of the important sources of global greenhouse gas emissions. Various technological innovations and improvement measures are introduced to reduce carbon emissions in cement industry. Among these, CO2 can be fixed in concrete to form a stable calcium carbonate by solidifying CO2 or accelerating carbonization, thereby reducing the release of CO2 into the atmosphere. This carbonization process can effectively reduce the carbon dioxide produced by concrete and improve the comprehensive performance of concrete to a certain extent. However, the carbonization process is complicated due to the diverse components of concrete materials. The existing studies indicate that the components such as calcium silicate hydrate, ettringite, and calcium hydroxide in the concrete can participate in the carbonization reaction, and different components have different effects on the formation process of calcium carbonate. Therefore, this paper investigated the nucleation and growth process of calcium carbonate in confined space of different components by a molecular dynamics simulation method to reveal the formation mechanism of calcium carbonate in confined space. In addition, the structural characteristics of each component in concrete and the mechanism of their interaction with carbon dioxide were also analyzed.Methods The adsorption models of calcium carbonate in C-S-H, AFt, CH and SiO2 confined space were established, respectively, to investigate the cluster formation process at the interface between calcium carbonate and different components of concrete. For C-S-H, the model includes C-S-H matrix and aqueous solution containing calcium carbonate. The C-S-H substrate with a Ca/Si ratio of 1.7 was used as a substrate model, and the supercell was cut along the crystal plane parallel to (001) to form the surface of the C-S-H substrate. The overall model size is 43.20 ?×45.04 ?×80.00 ?, and the pore size of the substrate is 4 nm. A total of 60 CO32-, 60 Ca2+ and 2 000 water molecules were randomly placed in the pores, and the ion concentration in the solution was 1.4 mol/L. The other three matrices have the same size as the C-S-H matrix, and the concentration of Ca2+ and CO32- in the pores is the same as that of the above C-S-H system. Also, four matrix models with confined space of 3 nm and 5 nm were established, respectively, to explore the influence of pore size on calcium carbonate. Before the formal simulation, the conjugate gradient method was used to minimize the energy, and then the isothermal isobaric ensemble (NPT) was used in the simulation process. The temperature was set to 300 K, the running time was 10 ns, and the time step was set to 1 fs. The clay force field (ClayFF) was used to simulate the C-S-H, AFt, CH, and SiO2 matrix, respectively, and the carbonate solution used a separate force field parameter. The application of the mixed force field to different components in the model could more accurately simulate the dynamic characteristics and interactions between the components.Results and discussion In the case of bonding and local structure, the C-S-H, AFt and CH matrix form a stable and strong bonding structure between the interface region and the ions in the solution. Especially in the C-S-H system, Cac and CO32- in the pores form more bonds with the ions at the interface of the matrix and have a stronger interaction. In contrast, the adsorption capacity of the ions in the pores is obviously weak, and only a small amount of ions are adsorbed due to the lack of calcium ions on the surface of the SiO2 matrix.The nucleation mechanism of calcium carbonate clusters is more consistent with the pre-nucleation theory. The whole simulation process can be divided into three stages, i.e., rapid agglomeration, cluster growth and cluster densification. The stronger interaction between the ions at the interface of the C-S-H matrix and the free ions in the solution lead to the aggregation of calcium carbonate clusters at the interface, providing more nucleation sites for the further nucleation and growth of calcium carbonate. Therefore, calcium carbonate has the optimum nucleation and agglomeration effect in the C-S-H matrix. Also, the pore size of the confined space has a great influence on the formation process of calcium carbonate. When the pore size is 3 nm, the surface of the C-S-H, CH and AFt matrix has a stronger attraction to the free ions in the solution, which is conducive to promoting the rapid nucleation and growth of calcium carbonate. When the pore size increases to 5 nm, the nucleation and growth process of calcium carbonate clusters slows down due to the weakened effect of the matrix interface on the ions in the solution.Conclusions The ions at the interface of the matrix had a great influence on the formation of calcium carbonate clusters. For different components of concrete, calcium carbonate and matrix surface ions attracted each other to form ionic bonds, which were then adsorbed at the interface for nucleation and growth. It was easier to bond with carbonate and calcium ions in the pores because C-S-H had more free hydroxyl groups and calcium ions at the interface, providing more nucleation sites for the agglomeration and growth of calcium carbonate at the interface. Also, the attraction of the substrate surface to the free ions in the solution was stronger when the pore size was 3 nm promoting the nucleation and growth of calcium carbonate. The nucleation mechanism of calcium carbonate clusters in the matrix pores did not follow the classical nucleation theory, but calcium ions and carbonate ions combined in the form of ionic bonds to form smaller clusters and then aggregated to grow and nucleate. In this process, it was unnecessary to overcome a large nucleation energy barrier, which was consistent with the pre-nucleation mechanism.

Introduction Cementitious materials are widely used in construction due to their wide range of raw material sources and easy construction. However, the characteristics of brittleness and easy cracking restrict their application in ductile structures. Fiber toughening is an effective approach to enhance the toughness of cementitious materials. The damage of cementitious materials is a gradual, multi-scale occurrence of the process. The first micrometer-scale cracks emerge, gradually expand and become larger with the increase of the load, forming the microscopic cracks, following convergence, developing macroscopic cracks, and ultimately leading to the material failure. Therefore, the toughening and crack-resisting effect of single-scale fibers are difficult to meet the multi-scale cracking process of cementitious composites. To correspond to this process, the hybrid fiber is an effective method, and the synergistic effect between different types of fibers can achieve the effect of “1+1>2”. There exist a few studies on the static flexural performance of hybrid fiber-toughing cementitious composites (HFTCC), but there is paucity of studies on the dynamic flexural behavior of HFTCC. To understand the hybrid effect of steel fiber and polyethylene fiber (PE fiber) on the static and dynamic flexural properties of cementitious composites, the static and dynamic flexural performances of HFTCC were investigated via three point bending test and drop weight impact test. In addition, the hybrid effect of steel fiber/PE fiber and the corresponding toughening mechanism were also analyzed.Methods The three-point bending tests were carried out based on a reference (GB/T 17671—2021). An electro-hydraulic servo universal testing machine (INSTRON 1342) was used to carry out the three-point bending test, at a loading rate of 0.2 mm/min. Load transducers and displacement sensor were used to record the load and mid-span deflection (δ) changes during the test, and the loading was stopped when the δ value reached 5 mm.The impact flexural properties were determined by a model INSTRON-CEAST 9340 drop weight impact testing machine . The required velocities were obtained by releasing the impactor from a predetermined height. The mass of the impact hammer head used in this test was 4.08 kg, the set impact speeds were 3 m/s and 4 m/s, and the impact force was collected by a high-speed data acquisition system (within 22 kN) at a data acquisition frequency of 0.5 MHz.Results and discussion The failure pattern of specimens indicates that fibers can effectively improve the dynamic flexural properties of specimens due to the fiber bridging and crack-blocking effect. Compared with PE fibers, steel fibers have a higher stiffness and a better ductility, which can resist the deformation of the specimen and limit the degree of opening of the main crack, thus improving the flexural properties of the specimen.Comprehensive analysis of load-displacement curve of HFTCC specimens shows that the static flexural performance of HFTCC is positively correlated to the PE fiber content, and the dynamic flexural performance demonstrates a positive correlation with the steel fiber content. The specimen with 1.5% (in volume fraction) PE fiber and 0.5% steel fiber exhibits the optimal static flexural peak stress, while the specimen with 1.5% steel fiber and 0.5% PE fiber presents the optimal dynamic flexural energy absorption. Under static flexure, the peak strength and energy absorption of HFTCC are increased by 26.5%-31.7% and 14.8%-56.8%, respectively, compared to the specimen with 2.0% steel fiber content.The dynamic increase factor (DIF) and energy absorption of HFTCC exhibit a significant strain rate effect. Under dynamic loading, the toughening effect of PE fibers is weaker than that of steel fibers, and the incorporation of a large number of PE fibers is not conducive to the enhancement of the dynamic energy absorption capacity of the specimen, which is related to the change in the failure mode of PE fibers. A small amount of PE fibers instead of steel fibers plays a toughening effect, which can exert the enhancement of the specimen dynamic energy absorption capacity.Synergistic effect analysis reveals that the steel fiber plays the dual effect of enhancement and toughening on the higher rate of dynamic flexure when the steel-PE fiber system with a high dosage of steel fiber due to the high tensile strength, stiffness, and ductility of the steel fiber, effectively enhancing the energy absorption capacity of the HFTCC under the lager deformation.The schematic diagram of HFTCC physical model was established via the analysis of SEM images. The steel-PE fibers effectively prevent the development of micro-cracks and the formation of macro-cracks, and the steel-PE fibers have a positive effect on the improvement of the mechanical properties and impact resistance of the cementitious materials during the static and dynamic loading process. Also, the micro-morphological analysis of HFTCC indicates that the steel-PE fibers form a three-dimensional mesh structure to have the effect of fiber toughening and crack-resisting on the matrix.Conclusions 1) The steel-PE hybrid fibers improved the static and dynamic flexural properties of HFTCC, in which the 0.5% (in volume fraction) steel-1.5% PE hybrid fiber HFTCC exhibited the optimal static peak bending and tensile stresses, while the 1.5% steel-0.5% PE hybrid fiber HFTCC had the optimal energy absorption under the dynamic loads. The DIF and energy absorption of HFTCC had significant strain rate effects.2) The steel-PE fibers in HFTCC showed a synergistic effect. For the peak load, static flexure mainly showed a positive synergistic effect, and the contribution of PE fibers to the peak stress of HFTCC was weaker than that of steel fibers under the dynamic impact loads. The specimen with 0.5% steel-1.5% PE hybrid fiber under the static flexure effectively improved the energy absorption of HFTCC, with the maximum hybrid coefficient S of 1.239, while the impact dynamic loading under the impact dynamic loading of the specimen with 1.5% steel-0.5% PE hybrid fiber effectively enhanced the energy absorption of HFTCC, with the maximum S of 1.086.3) The toughening mechanism of HFTCC mainly originated from the synergistic crack-blocking and toughening effect between steel and PE fibers. Especially, the micro-fine steel fibers and PE fibers could fully utilized their respective mechanical properties at their appropriate dosages, thus effectively improving the static and dynamic flexural performances of HFTCC.

Introduction Alkali-activated fly ash (AAFA) can reduce the carbon emissions from the production of ordinary Portland cement (OPC) by >80%. The utilization of AAFA and the research on the electrochemical behavior of steel in AAFA have thus attracted much attention. Compared with OPC, a higher resistance to chloride-induced corrosion can be achieved in AAFA due to its high alkalinity and zeolite-like adsorption layer. However, the corrosion of steel in AAFA cannot be completely inhibited in harsh environments contaminated with chlorides.Carboxylate groups-based corrosion inhibitors (such as sodium citrate) improve the pore structure and compressive strength of concrete, and have a high inhibition efficiency for steel. Especially at a high chloride concentration, the pitting potential (Epit) of steel treated with carboxylate groups-based corrosion inhibitors is greater than that treated with amine salt and amino acid. Therefore, sodium citrate can be considered as a promising organic corrosion inhibitor to enhance the corrosion resistance of steel in AAFA. However, the inhibition mechanism of sodium citrate for steel in AAFA is still unclear. In this paper, the effect of sodium citrate on the corrosion behavior of reinforcing steel in extracted AAFA solution under chloride attack was investigated by electrochemical measurements and surface characterization techniques, and the inhibition mechanism of sodium citrate was revealed.Methods In this study, the preparation of AAFA solution was carried out at room temperature ((25±2) ℃). Firstly, a low calcium fly ash was mixed with an alkali-activated solution (80 g/L NaOH solution) at a mass ratio of 1:1, then the mixture was stirred in a closed vessel for 24 h, and finally it was filtered to remove the residues. The concentration of sodium citrate added into the AAFA solution was 100 mmol/L.Conventional ribbed carbon steel (HRB400E) with a diameter of 16 mm was used. The electrochemical measurements, namely electrochemical impedance spectroscopy (EIS) and the cyclic potentiodynamic polarization (CPP), were performed in a classic three-electrode electrochemical system, i.e., the working electrode (HRB400E steel), the reference electrode (saturated calomel electrode, SCE) and the counter electrode (platinum electrode). The surface morphology and chemical composition of steel passive films were characterized by atomic force microscopy (AFM) and X-ray photoelectron spectroscopy (XPS), respectively. After chloride attack, the corrosion morphology and corrosion products of steel were analyzed by scanning electron microscopy (SEM) with energy dispersive X-ray spectroscopy (EDS) in a model FEI Quanta 3D FEG environmental scanning electron microscope.Results and discussion In the absence of protective passive film, the capacitive loop of the non-passivated steel (the Blank specimen) decreases in AAFA solution containing chlorides, thereby gradually reducing the corrosion resistance of steel. However, the capacitive loop and impedance modulus (|Z|) of steel treated with citrate (the Cit specimen) increase. Citrate can effectively improve the corrosion resistance of the non-passivated steel in AAFA solution containing chlorides.Moreover, the EIS results of pre-passivated steel in chlorides-free AAFA solution show that the capacitive loop of steel treated with citrate (the P-Cit specimen) is somehow smaller than that of the blank group (the P-Blank specimen), indicating that citrate is unfavorable to the formation of steel passive films in AAFA solution. However, in the presence of chlorides, an enhanced corrosion resistance can be obtained for the P-Cit specimen, which is the same as that of the non-passivated steel.According to the results of AFM and XPS measurements, citrate can modify the surface morphology and chemical composition of steel passive film in AAFA solution. There are a large number of fine particles on the steel surface in AAFA solution, and the surface roughness of steel is less and more uniform due to the adsorption effect of citrate, compared with that without citrate. For steel with citrate, the intensities of C—C/C—H and C==O peaks increase and the oxidation of Fe may be suppressed to a great extent. Furthermore, O2?/OH? ratio of steel passive film treated with citrate slightly decreases, and the corresponding content of bound water in steel passive film decreases.Clearly, citrate ions can combine with iron ions to form an adsorption film on the steel surface, therefore suppressing the formation of passive film to some extent. However, when steel is attacked by chlorides, the adsorption film of citrate can prevent chloride ions from reaching steel surface. In addition, the negatively charged carboxylic acid group is also dominant in the competition with chloride ions, effectively inhibiting the depolarization of chloride ions, and finally improving the corrosion resistance of steel in AAFA solution.Conclusions Although sodium citrate had a negative effect on the formation of a stable steel passive film, it was suitable for improving the corrosion resistance of steel in AAFA solution. Citrate ions could be adsorbed on the steel surface to form an adsorption film, thereby acting an effective physical barrier. Also, citrate ions played a dominant role in the competitive adsorption with chloride ions on the steel surface. The corrosion resistance of steel in AAFA solution contaminated with chlorides was improved by the addition of sodium citrate.

Introduction Phosphates are one of commonly used inorganic retarding agents in Portland cement, including monophosphates (i.e., sodium phosphate) and polyphosphate compounds (i.e., sodium pyrophosphate, sodium tripolyphosphate, and sodium hexametaphosphate). Different phosphate salts exhibit distinct retarding mechanisms. Among these, sodium phosphate is extensively investigated as an alkaline metal monophosphate-type retarder. However, there exist varying interpretations of the retarding action of sodium phosphate in different studies, potentially impacting its practical application, and necessitating a clear elucidation of its retarding mechanism. Ammonium dihydrogen phosphate as an acid phosphate can also serve as a retarder and exhibits a more enduring retarding effect rather than the commonly used sodium phosphate. Compared to PO43-, the volatility of NH4+ or its potential to mitigate the influence of alkaline metal ions on the cement hydration and the larger ionization constant of H2PO4- can have a quicker reaction with alkaline cement particles, affecting the hydration process. Some work indicate that ammonium dihydrogen phosphate has a favorable retarding effect on cementitious slurries, but its specific retarding mechanism remains less explored. It is thus necessary for the analysis of their impact on the kinetics of Portland cement hydration to conduct a comparative study between ammonium dihydrogen phosphate and sodium phosphate.Methods Cement used was conformed to the GB 8076—2008 standard with a strength grade of 42.5. Ammonium dihydrogen phosphate and sodium phosphate were of analytical reagent grade (AR). The dosage of ammonium dihydrogen phosphate and sodium phosphate was 0.5% (in mass fraction), with a fixed water-to-cement ratio of 0.4. Cement samples for microstructural tests were prepared in 60 mL plastic bottles. The cement paste was prepared via dissolving ammonium dihydrogen phosphate and sodium phosphate, and mixed with cement. The well-mixed cement paste was placed in the plastic bottles and cured under natural conditions at (20±2) ℃ until the specified age.Cement setting time and paste flowability were determined according to GB/T 1346—2011 and GB/T 8077—2000 standards, respectively. The hydration exothermic curve of the cement paste in the early 72 h at 20 ℃ was recorded by a model TAM isothermal calorimeter. The zeta potential within the first 10 minutes of cement paste hydration was measured by a model DT-1200 electroacoustic spectrometer. The element concentration of the pore solution was measured by a model PS-6 inductively coupled plasma optical emission spectrometer (Baird Co., USA). The thermogravimetric analysis was conducted by a model STA449C synchronous thermal analyzer (NETZSCH Co., Germany). The X-ray diffraction (XRD) patterns of the samples were obtained by a model Advance D8 fully automated powder diffractometer (Bruker Co., Ltd., Switzerland). The surface morphology and elemental identification of the samples were observed by a model Nova NanoSEM230 scanning electron microscope.Results and discussion The incorporation of sodium phosphate and ammonium dihydrogen phosphate results in a corresponding extension of both initial and final setting time of the cement paste. Ammonium dihydrogen phosphate with an equal mass exhibits longer initial and final setting time, compared to sodium phosphate. Compared to the pure cement group, the addition of sodium phosphate and ammonium dihydrogen phosphate improves the flowability of the cement paste and delays the time-dependent loss of flowability, showing a more pronounced retarding effect.Hydration heat, thermogravimetric and XRD results indicate that sodium phosphate and ammonium dihydrogen phosphate both can inhibit the dissolution rates of C3S and C3A, thus retarding the formation of hydration products. This leads to a reduction in the rate of heat evolution during Portland cement hydration, an extension of the induction period of hydration, a decrease in the total heat evolved during the first three days of hydration, and an overall inhibition of Portland cement hydration. However, an inhibitory effect of ammonium dihydrogen phosphate is dominant.Based on the analysis of zeta potential and SEM results, after the addition of sodium phosphate and ammonium dihydrogen phosphate in cement, a coating layer is formed on the surface of the cement particles. Sodium phosphate precipitates as calcium phosphate on the surface of Portland cement particles, while ammonium dihydrogen phosphate precipitates as hydroxyapatite. Ammonium dihydrogen phosphate exhibits a greater retarding effect rather than sodium phosphate due to its higher adsorption capacity of hydroxyapatite for inorganic ions, compared to calcium phosphate.The early pore solution analysis indicates that ammonium dihydrogen phosphate can be adsorbed more rapidly on Portland cement particles, leading to the rapid precipitation of solid-phase products. In addition, no nitrogen (N) element appears when analyzing the early pore solution of the cement paste with ammonium dihydrogen phosphate, indicating that NH4+ rapidly escape into air as ammonia when adding to the Portland cement paste. This escape of NH4+ effectively prevents the influence of alkali metal ions in sodium phosphate on the early hydration of Portland cement.Conclusions Ammonium dihydrogen phosphate and sodium phosphate both were capable of prolonging the setting time of Portland cement and retarding the time-dependent loss of paste flowability with ammonium dihydrogen phosphate exhibiting a more pronounced retarding effect. These agents could inhibit the dissolution rates of C3S and C3A, thereby reducing the rate of heat evolution during Portland cement hydration, extending the induction period of hydration, lowering the total heat evolved during the first three days of hydration, and overall suppressing Portland cement hydration. However, the inhibitory effect of ammonium dihydrogen phosphate was dominant. Compared to sodium phosphate, ammonium dihydrogen phosphate could be more rapidly adsorbed on Portland cement particles, leading to the rapid precipitation of solid-phase products, and thereby inhibiting cement hydration. Sodium phosphate precipitated as calcium phosphate on the surface of Portland cement particles, suppressing cement hydration, while ammonium dihydrogen phosphate precipitated as hydroxyapatite on the surface of Portland cement particles, similarly inhibiting cement hydration.

Introduction Calcium sulphoaluminate cement has a potential application in sub-zero temperature environments due to its rapid hydration, rapid setting, fast strength development, and high early strength. In cold construction sites, such as polar regions, the materials and the concrete both are exposed to sub-zero temperatures for extended periods. It is thus necessary to investigate the formation, condensation, hardening and microstructures of calcium sulphoaluminate cement at sub-zero temperatures. Cold concreting method widely recognized involves directly blending raw materials on-site at the same temperature as the ambient environment, even at sub-zero temperatures, follows via pouring and curing fresh concrete without supplemental heating. Choosing a suitable mixing solution is crucial for the effective cold concreting technology. The mixing solution must prevent freezing of both the solution and the cement pastes at sub-zero temperatures, initiating the cement hydration reaction and enabling continuous progression. Also, the properties of the mixing solution (i.e., pH value and ion concentration) affect the dissolution and hydration of clinker minerals, leading to consequential changes in cement characteristics. In this paper, the performance evolution laws and mechanisms of calcium sulphoaluminate cement mixed with different mixing solutions at sub-zero temperatures were investigated.Methods Deionized water was used as a solvent. Six inorganic analytical pure chemical reagents and four organic analytical pure chemical reagents were used as solutes. Each of these reagents was dissolved individually in deionized water at room temperature. The solution was stirred until in a homogeneous state and left standing for 24 h. To ensure the temperature of the materials as the environmental temperature, the calcium sulphoaluminate cement and the mixing solution both were placed in a sub-zero temperature experimental system for a storage period of at least 24 h. All processes, i.e., sample shaping and curing, were carried out in the sub-zero temperature experimental system.At ?10 ℃, the mixing solution and cement were poured sequentially into a mixing bowl. They were stirred at a low speed for 120 s, paused for 15 s, and then stirred at a high speed for 120 s. The resultant paste was poured into the mold. After pouring, the vibrating table was used for 60 compactions. The samples were then placed on a rack for cured without being covered in the sub-zero temperature environment. Once the samples were cured to the specified age, a portion of them was taken out from the sub-zero temperature experimental system for the measurement of compressive strength. Another portion was crushed and sampled on-site. These samples were placed in anhydrous ethanol to terminate hydration at room temperature for at least 7 d. Afterwards, they were dried in vacuum at 40 ℃ and ?0.08 MPa for no less than 24 h. A portion of the samples dried was used for the microscopic observation, while another portion was ground and further analyzed for the phase composition.Results and discussion CaCl2/Ca(NO3)2 solution is capable of accelerating hydration, resulting in an increased hydration temperature and a greater early strength. The peak temperature reaches up to 23.6 ℃ and 20.4 ℃, and the compressive strength at 12 h reaches 21.5 MPa and 16.5 MPa, respectively. Furthermore, the compressive strengths at 28 d are 101.8 MPa and 19.6 MPa, respectively. K2CO3 solution generates an increased hydration temperature and results in a high early strength, with a peak temperature of 7.7 ℃ and a compressive strength of 13.9 MPa at 12 h. However, the compressive strength only increases to 19.3 MPa after a 28-d period. MgCl2, NaCl and NaNO2 solutions exhibit a limited effect on enhancing hydration, presenting a low heat evolution and an early strength at their peak temperature of <5 ℃ and the compressive strength of <5 MPa after 12 h. Nonetheless, after a 28-d period, their compressive strength increases to 74.0, 42.7 MPa and 39.2 MPa, respectively. After 12-h hydration, the main products formed from CaCl2, Ca(NO3)2, MgCl2, NaCl and NaNO2 in calcium sulphoaluminate cement are AFt and AFm phases, having a needle-bar and flake morphology. Also, the main products formed from K2CO3 solution-mixed samples are CaCO3 and K2SO4 phases, having a cube and ellipsoid morphology.CH3OH, CH3CH2OH, (CH2OH)2 and C3H8O3 solutions have a slight promotion effect on the hydration, resulting in a negligible heat evolution (peak temperature of < -5 ℃), minimal strength development at 12 h, and extremely low compressive strength for the samples after 28 d. However, the compressive strength of C3H8O3 solution-mixed samples reaches 23.5 MPa after 28 d. For the samples mixed with four organic solutions, no hydration products such as AFt appear after 12-h hydration.Conclusions CaCl2 and Ca(NO3)2 solutions with the same ions (Ca2+) as the cement minerals promoted cement hydration, resulting in an increased temperature and a greater early strength. MgCl2, NaCl and NaNO2 solutions had negligible effects on the hydration, leading to a decreased temperature and a weaker early strength. CO32? in K2CO3 solution reacted with Ca2+ produced due to the dissolution of anhydrite to create CaCO3 precipitates. This reaction increased the freezing point of the solution and reduced its frost resistance, causing the sample to freeze and preventing the effective cement hydration reactions.The hydroxyl group of alcohol molecule in CH3OH, CH3CH2OH, (CH2OH)2 and C3H8O3 solutions was adsorbed by Ca2+ on the surface of cement particles to form an adsorption film. This film delayed hydration. In addition, the interaction between water and cement particles also impeded when the alcohol molecule contained a methyl group, exacerbating a retarding effect.

Introduction Tensile softening (σ-w) curve is a critical parameter necessary to model the non-linear fracture behavior of concrete structures by the finite element method. The uniaxial tensile test is the most direct approach to determine the softening curve, but it is a challenge due to its insufficient stiffness of the loading machine, eccentricity, asymmetric fracture modes, and multiple cracking. Alternatively, a softening curve can be derived via minimizing the deviation between the predicted and experimental results based on the load-crack mouth open displacement (F-lCMOD) curve measured by the bending or wedge-splitting tests. Inverse analysis methods are commonly used to derive σ-w curve include the J-integral, poly-linear, and global optimization. Although the J-integral method is computationally least demanding, it requires conducting the tests on two specimens with different notch sizes for deriving each σ-w curve, yielding a lower precision and a higher dispersion. For the given test data, the poly-linear method can produce a unique solution without any pre-assumptions regarding the softening curve’s shape. However, it is sensitive to arbitrarily small measurement errors, yielding a softening curve with severe oscillation, and accumulating all previous errors in the current analysis step. The global optimization technique is not easy to converge when there are many fitting variables, and it needs to pre-assume the shape of the softening curve. There is no guarantee that the outputs of this technique will be globally optimal. Therefore, this paper proposed a self-modifying inverse analysis method to determine the softening curve of concrete after combining global best-fitting with an automatic correction technique. Methods The self-modifying inverse analysis method belongs to the global optimization technique with more additional constraints. This method consists of two modules, i.e., global best fitting and parameter self-modifying. The former mainly focuses on the optimum fitting of the F-lCMOD curve to minimize the cumulative deviation between the simulated and measured values. For this purpose, a computational program was developed a software named MATLAB. In this program, the ability to define search boundaries for each control variable increases the robustness of the fitting algorithm and lessens the likelihood of finding a pseudo-optimal solution. Despite, the optimal solution may be still the local minimum in some situations, so it is also necessary to check whether the local error requirement is satisfied at each discrete point and to find out the maximum load deviation as the second criterion for evaluating the prediction accuracy. Thus, the role of the second module is to automatically modify the optimal value of the fitting parameter when the second criterion is not met. After providing the corrected parameter values (including the boundaries) to the first module, the global optimization procedure is executed again until the load tolerance error is satisfied. The current model cannot accurately simulate the crack propagation process if the load condition cannot be satisfied after the maximum number of iterations. It is required to add a new line segment to the current softening model and conduct the global fitting procedure again with the model containing additional variables until the load criterion is satisfied or the maximum number of line segments is reached. In this algorithm, it is possible to output a precise global optimal solution even when the initial estimate of fitting parameters differs from the ideal solution or the initial softening model is relatively simple.Results and discussion For the sample, in all regions of the F-lCMOD curve, the simulated responses match well with the experimental results, and the maximum relative error between the calculated and measured loads is less than 5%. This indicates that the developed method can accurately simulate the crack propagation process of concrete, and this method is highly versatile, applying to different specimen forms and various concrete materials. The fracture energy determined by the F-lCMOD curve is in a reasonable agreement with that by the σ-w curve, and their deviation does not exceed 2%, further confirming the rationality of the recommended technique. Concrete exposed at different humidities exhibits a similar shape to the softening curve, consisting of a fast descending, a slow descending, and a tail. The tensile strength (ft) increases with decreasing age while critical crack width (wc), finally tending to their stabilizing values. For a given age, ft and wc decrease with reducing ambient humidity, and this decrease is more significant in the later stages, which is due to the lack of sufficient water in the pores for hydration reaction, the non-uniform shrinkage deformation between aggregate and mortar, and the disturbance of local structure of C-S-H. Although there are multiple “optimal” solutions at different initial fitting parameters, the ideal σ-w curve is rather narrow, implying that the algorithm converges to the similar solutions. The determined softening curve is independent of the initial estimate, and the algorithm is robust. In the case of different initial models, the optimal σ-w curve is similar with only some minor differences in the tails. This demonstrates that the softening model’s initial form has little effect on the optimization outputs, and trustworthy optimal solutions can be obtained even by a simple beginning model.Conclusions The self-modifying inverse analysis method could consider the effect of local response on the softening curve, allow the definition of searching boundaries of the fitting parameters, and eliminate the need to predefine the final shape of the softening curve, reducing the probability of obtaining a "pseudo-optimal solution" and weakening the dependence of the optimization output on the initial guesses. In all the region of the F-lCMOD curve, the calculated load was correlated well with the experimental result, and the fracture energy determined by the F-lCMOD curve differed by no more than 2% one by the σ-w curve. The proposed method was highly robust and versatile. It could accurately determine the σ-w curve of various concrete materials with different specimen configurations, and the resultant optimal solution was independent of the initial estimate of fitting parameters and the initial shape of the softening model. Concrete exposed at different humidities exhibits a similar shape to the softening curve. The ft increased with decreasing age while critical crack width wc, finally tending to their stabilizing values. ft and wc decreased as the ambient humidity decreased, and this reduction was more pronounced in the later age.

Introduction Application of silica nanoparticles in concrete has attracted much attention. Its unique filling characteristics and high volcanic ash characteristics can improve the pore structure of concrete and promote hydration reaction, having a positive effect on the mechanical properties of concrete. After concrete is subjected at a high temperature, the pore pressure inside the structure, the deterioration of the transition zone of the slurry aggregate interface and the change of phase are all reasons for the decline of the mechanical properties of concrete at a high temperature. Silica nanoparticles have a filling effect, increasing the compactness of the material, while silica nanoparticles can undergo a volcanic ash reaction to generate more hydrated calcium silicate (C-S-H) gel, improving the high-temperature resistance of concrete. The inside of alkali slag mortar is alkaline environment and its hydration reaction is completely different from concrete, and the influence of silica nanoparticles on the high-temperature resistance of alkali slag cementitious material is different from that of ordinary concrete. Therefore, silica nanoparticles can improve the high-temperature mechanical properties of cement-based materials, but the influence of silica nanoparticles on the high-temperature resistance of seawater sea sand alkali slag mortar is more complicated. This paper was to investigate the effect of silica nanoparticles on the high-temperature resistance of seawater sea sand alkali slag mortar.Methods Simulated seawater and sea sand with a fineness modulus of 2.3 were used as mixing water and fine aggregates. Liquid sodium silicate and solid sodium hydroxide were formulated as alkali initiators in a molar ratio of 1.7 for SiO2 and Na2O. Silica nanoparticles with the sizes of 20-50 nm and a specific surface area of 670 m2/g as an admixture were supplied by Taihong Shengda New Materials Co., Ltd., China. Four sets of alkali slag mortar with different concentrations of silica nanoparticles (i.e., 0%, 1%, 2% and 3%) were produced. After mold removal, it was placed in a curing room at (20±2) ℃ and a relative humidity of > 90% for 27 d, and then dried naturally for 1 d. The changes of mass and length in specimens were recorded by scales and vernier calipers before and after high-temperature heating. The flexural and compressive strengths were determined on seawater sea sand alkali slag mortar specimens (40 mm×40 mm×160 mm) at different high temperatures (i.e., 200, 400, 600 ℃ and 800 ℃).Results and discussion At room temperature, silica nanoparticles have little effect on the compressive strength of seawater sea sand alkali slag mortar, and the compressive strength of the specimen is only increased by 1.5% and 2.3% at the concentrations of silica nanoparticles of 2% and 3%, respectively. Also, silica nanoparticles can reduce the pH value of alkali initiator, which is unstable in salt solutions where Mg2+ is > 0.1% (in mass fraction). Monomeric silicon ([SiOn(OH)4-n]n-) is formed when silica nanoparticles are added into the base initiator, improving the SiO2/Na2O ratio of the solution. At 200 ℃, compared with ordinary concrete, the addition of silica nanoparticles can increase the compressive strength of seawater sea sand alkali slag mortar. The compressive strength of the specimen with silica nanoparticles content of 1%-3% is 103%-123% of the normal temperature state. This may be related to the superior dispersion of silica nanoparticles in alkaline environments. OH- can convert the silicon hydroxyl group on the surface of silica nanoparticles to Si-O- groups, so that the surface of silica nanoparticle generates a negative charge and increases the dispersion. The deformation results at a high temperature show that the residual deformation of the specimen can be effectively alleviated when the silica nanoparticles content is 3%. Combined with the analysis of the results of mass loss, the addition of silica nanoparticles can reduce the content of evaporable water in the specimen. This decreases the pressure generated by the evaporation of partially pore water, thereby reducing the partial crack generation.Conclusions The addition of silica nanoparticles improved the compressive strength of seawater sea sand alkali slag mortar at 200 ℃. This could be related to the dispersion of silica nanoparticles in alkaline environments rather than in neutral environments. Silica nanoparticles could reduce the evaporable water content inside seawater sea sand alkali slag mortar. At the silica nanoparticles content of 3%, the deformation and cracks of the specimen at a high temperature were effectively inhibited.

Introduction Cerium is the most abundant rare-earth element and is widely used in industry as catalyst, magnet, phosphor, polishing powder and ceramic colorant. Cerium occurs naturally in the form of Ce3+. Ce4+ is easily extracted and separated from cerium fluocarbonate. Therefore, Ce3+ is often converted to Ce4+ during the reprocessing stage in mineral extraction and metallurgical engineering. Ce4+ is likely to accumulate in the environment, food chain and organisms due to its extensive development and application, causing serious damage to the environment and biological systems. It is thus necessary to detect Ce4+. AA can be also called vitamin C as one of the most important vitamins to maintain the normal activities of the human body. Previous studies indicate that the level of AA in the human body is closely related to many diseases, lack of AA may suffer from the following diseases like mouth ulcer, hair loss, Alzheimer's disease, scurvy and so on. A simple and healthy way to consume AA is to eat more fruits and vegetables that contain AA. AA concentration is an important index to determine fruit freshness and nutritional quality. It is thus of great significance to check the AA of fruits and vegetables correctly and reliably to prevent diseases and evaluate the freshness of vegetables. The development of a simple, reliable and highly selective AA assay is crucial for analytical applications and clinical disease diagnosis, becoming a key topic in current chemistry research. In the past few decades, a series of methods for the detection of Ce4+ or AA are reported, and a fluorescence method has attracted recent attention due to its advantages such as simple instrument, low cost and low sample consumption. However, such fluorescence analysis methods still have some technical challenges and inevitable shortcomings. For instance, the detection results with single emission fluorescence method are often interfered by some factors unrelated to the detection environment, such as background signal, fluorescence substance concentration, and surrounding environment. To solve this problem, a fluorescence analysis strategy, namely ratio fluorescence analysis, was proposed. The ratio fluorescence analysis method realized the self-calibration of the analysis signal through the double fluorescence signal, easily obtaining the more accurate detection results rather than the fluorescence method with single emission signal. It provides an effective way for the analysis and detection of target objects.Methods According to the method described in the literature, 0.25 g thiourea and 0.25 g citric acid were dissolved in 20 mL ultra-pure water, treated with ultrasound for 30 min, then placed at 80 ℃ for 8 h, centrifuged to collect the supernatant, and dialyzed twice with 1 000 Da dialysis bag. The solution obtained was N, S-CQDs solution.BSA-AuNCs were prepared according to the method reported in the literature. After 50 mg/mL BSA solution and 10 mmol/L HAuCl4 solution were mixed under stirring for 2 min, 1 mL NaOH (1 mol/L) was added drop by drop and stirred at 100 ℃ for 7 min. After dialysis in ultra-pure water with 1 000 Da bag for 12 h, BSA-AUNCS solution was obtained.BSA-AuNCs and N, S-CQDs were mixed at a volume ratio of 1:2 and stored at 4 ℃ for the coming use. 5 mL N, S-CQDs@BSA-AuNCs as a working solution was mixed with Ce4+ solution at different concentrations. Deionized water was added into the mixture for a fixed volume to 10 mL, and stabilized for a period of time. Afterwards, the fluorescence test was performed at the excitation wavelength of 345 nm. 5 mL N, S-CQDs@BSA-AuNCs solution was firstly mixed with 100 μL 1 mmol/L Ce4+, and then mixed with AA solution at different concentrations at a constant volume of 10 mL under stirring for 2 min. The fluorescence test was performed at the excitation wavelength of 345 nm after stability for a period of time.Results and discussion In this work, nitrogen and sulfur doped carbon quantum dots (N, S-CQDs) and gold nanoclusters (BSA-AuNCs) are combined to form a N, S-CQDs@BSA-AuNCs fluorescence sensor for the continuous determination of Ce4+ ion and AA. Bsa-auncs were synthesized with bovine serum protein (BSA) as a stabilizer. The BSA-AuNCs have superior fluorescence properties, good biocompatibility, high quantum yield, excellent colloidal stability and ultra-small size. It is indicated that BSA-AuNCs selectively recognizes Ce4+ ions based on a fluorescence enhancement mechanism, and an "open" NIR fluorescence sensor platform for detecting Ce4+ is established. The Ce4+ mediated BSA-AuNCs "off" NIR fluorescence sensing strategy is developed to detect AA via adding AA to the NIR fluorescence sensor platform of N, S-CQDs@BSA-AuNCs system.Conclusions The N, S-CQDs@BSA-AuNCs ratio fluorescence sensor was constructed by N, S-CQDs and BSA-AuNCs for the continuous determination of Ce4+ and AA. When Ce4+ was added to the system, BSA-AuNCs were oxidized by Ce4+ and aggregated into coarser particles with an enhanced fluorescence. Therefore, an "open" NIR fluorescence sensing platform for Ce4+ detection was constructed. The fluorescence intensity ratio of IF429/IF692 showed a linear detection relationship with the concentration of Ce4+, and the detection limit was 0.12 μmol/L. Subsequently, a Ce4+ mediated NIR fluorescence sensing strategy based on N, S-CQDs@BSA-AuNCs "off" was developed to detect AA, with a detection limit of 0.19 μmol/L via adding the NIR fluorescence intensity of AA quenching N and S-CQDs@BSA-AuNCs/Ce4+ system. The ratio fluorescence sensor was applied to the detection of Ce4+ concentration in actual tap water samples, AA concentration in commercial peach juice and vitamin C tablets.

Introduction Hydrated salt phase change energy storage materials enable sustainable storage and release of thermal energy through the phase transition of crystalline water. This low-cost, environmentally friendly material has a high heat storage density. However, the significant heat loss and liquid phase leakage after multiple cycles become persistent challenges, restricting the large-scale application of the original hydrated salt phase change material. Consequently, the preparation of shape-stabilized composite phase change materials by infiltrating hydrated salt phase change materials into porous support structures is one of the investigated methods to address the encapsulation challenges of hydrated salt phase change materials. In this paper, carbon aerogels with a high adsorption capacity were prepared with melamine foam as a precursor. The carbon aerogels were used to encapsulate hydrated salt phase change materials. Methods Carbon aerogel was prepared via multistage temperature carbonization with melamine foam as a precursor. A composite eutectic salt material (EHS) of Na2SO4·10H2O-Na2HPO4·12H2O with 0.2% glucose carbon nanoparticles was added as the phase change matrix. Shape-stabilized melamine carbon aerogel/eutectic salt composite phase change materials were prepared via vacuum melt impregnation.The apparent morphology of the prepared aerogels was determined by scanning electron microscopy (SEM). The composition of the aerogels was analyzed by X-ray diffractometry (XRD), and the interactions between the MF and the eutectic hydrate salts in the composite phase change materials were investigated by the Fourier infrared spectrometry (IR). The thermal conductivities of the samples at room temperature were determined by the Hot Disk thermal constant analysis. Also, the latent heat of the phase change material was analyzed by differential scanning calorimetry (DSC), and the thermal stabilities of the samples were assessed by the simultaneous thermal analysis.Results and discussion Melamine foam before carbonization has large pore and smooth surface. The size of pore decreases with increasing carbonization temperature. Meanwhile, the elemental of carbon aerogels mainly includes C, N and O. The surface functional group disappears as the carbonization temperature rises. The XRD patterns and IR spectra of the carbon aerogel reveal the presence of numerous defects in conjunction with the carbon material, indicating a high degree of graphitization. Moreover, the graphitization degree shows a gradual increase with increasing carbonization temperature. The specific surface areas of carbon aerogels MF, MF6, MF7 and MF8 measured by BET method are 1.4, 107.9, 485.1 m2·g-1 and 313.4 m2·g-1, respectively. The DFT analysis of the pore size distribution shows that carbon aerogel MF has a number of pore sizes larger than 50 nm, and the volume of carbon aerogel MF7 decreases significantly. Also, a large number of micropores and a part of the mesoporous structure appear. The destruction of melamine formaldehyde structure at high temperatures leads to the reduction of the original micropores and the generation of new pores, which is consistent with the SEM results.The loadings of carbon aerogels MF and MF7 reach 116 and 122 times after 5 000 solid-liquid cycles, and the loss of carbon aerogel MF7 is less, indicating its better cycling stability and high potential for application. The initial eutectic salts lose completely at 75 ℃. The eutectic salts after encapsulation by carbon aerogels has a better thermal stability. All of them lose water completely at 145 ℃. The thermal conductivities of composite phase change materials PCM6, PCM7 and PCM8 prepared are 1.11, 1.24 W·m-1·K-1 and 1.37 W·m-1·K-1, respectively. The phase transition latent heat value of PCM is the maximum, but after multiple solid-liquid cycles, its enthalpy of melting is decreased by 15.0% and crystallization enthalpy is decreased by 16.3%, leading to a poor cycling stability. Also, the enthalpy of melting of PCM7 is 278.3 J·g-1 and the enthalpy of crystallization is 227.0 J·g-1, which are only decreased by 4.5% and 2.6% after solid-liquid cycling, resepctively. PCM7 has a superior cycling stability.Conclusions Carbon aerogel of a three-dimensional network structure with high loading, high elasticity and good ductility was prepared via multistage carbonization with melamine foam as a precursor, and carbon aerogel eutectic salt composite phase change materials were prepared.The carbonization of carbon aerogel MF at different temperatures had an effect on the surface elements, functional groups, morphology, and the degree of graphitization of the pore structure. The carbon aerogel prepared at 700 ℃ had an advantage over the encapsulated eutectic salt phase change materials.Carbon aerogel MF had a high loading for EHS, but its leakage prevention, thermal conductivity and cycling stability were poor. Carbon aerogel MF7 had a better pore structure and a needle fiber surface, making it maintain 122 times of its own mass adsorption after 5 000 solid-liquid cycles.The enthalpies of melting and crystallization of PCM7 were 278.3 J·g-1 and 227.0 J·g-1, and the enthalpies were decreased by 4.5% and 2.6%, respectively, after 5 000 solid-liquid cycles. PCM7 had a good cycling stability and a great application potential.

The morbidity and mortality of patients with chronic wounds increase with the aggravation of social aging. The sharp rise in the burden of the medical system leads to an increased demand for high clinical efficacy, developed therapies, and innovative products. Bioactive glass (BGs) is well considered as a promising clinical material in biomedical fields, especially in wound healing. BGs stimulate the production of growth factors due to its unique controlled dynamic ion release properties, enhancing cell proliferation and regulating the gene expression of related cells. Particularly, the bioglass composed of borate and borosilicate exhibits a controllable degradation rate, which can continuously release functional ions and regulate the wound microenvironment for various pathological stages of wound healing to induce angiogenesis and accelerate the steady recovery of pathological wound tissue. In over 50 years of clinical and fundamental research, the efficacy of bone grafts in repairing bone tissue has been conclusively demonstrated, and biomaterial scientists have analyzed in depth the molecular biological mechanisms underlying the interaction between BGs and bone. Emerging applications of BGs for soft tissue repair still require a further research. This review represented a relationship between the structure and properties of boric acid/borosilicate bioactive glass, focusing on the clinical application research progress and future prospects for the development of bioglass wound dressings.Complicated factors make chronic wounds difficult to manage. The excessive inflammation is a key factor in wound pathogenesis. To treat chronic wounds, it is essential to comprehend how to prevent the development of microorganisms at the wound site, reduce ROS production and tissue inflammation, induce cells to release more growth factors, alter the cell vitality, chemotaxis and mobility, and initiate new blood vessel growth. The development of BGs reveals its positive role in the treatment, restoration, and regeneration of hard and soft tissues in the human body. Compared to conventional inactive biomaterials, its clinical application significance and potential are substantial. Borosilicate bioactive glass (BBGs) with a unique glass network structure has a faster degradation rate, compared to silicate bioactive glass, making it suitable for soft tissue repair. BBGs with specific functional properties can be produced via adjusting the proportion of biologically active elements (i.e., Na, K, Ag, Au, B, Ca, Cu, Co, Ga, Mg, Sr and Zn) doped into the glass network structure according to the therapeutic needs of target tissues or organs. BBGs with different ion compositions can continuously release functional ions and regulate the wound microenvironment for different pathological stages of wound healing, thereby inducing angiogenesis and accelerating the steady-state recovery of pathological wound tissue. In addition to the superior clinical therapeutic effect by using fiber as clinical application dressing, BBGs can also combine with polymer to form functional hydrogel dressing, and can prepare multi-functional composite materials by electrospinning technology or 3D tissue engineering printing technology to meet the personalized needs of different stages of wound repair. The majority of chronic wound patients are affected with the intensification of global aging. Boric acid/borosilicate bioactive glass wound repair materials are low-cost, easy to store, and have significant social and economic benefits.Summary and prospects Although clinical and basic application studies have confirmed the potential of BBGs in wound healing applications, the structural composition of BBGs, especially the types and concentrations of therapeutic ions released, as well as the regularity and molecular biological mechanisms of pathological wound healing, still need a further exploration. The risk of soft tissue and organ toxicity, and even systemic toxicity caused by BBGs, is needed to be evaluated. Also, it is necessary to effectively evaluate and ensure the biosafety of BBGs dissolved ions, and to construct a more advanced in-vitro simulated wound dynamic microenvironment. In addition, combining BBGs as additives with other treatment strategies (such as stem cell therapy or growth factors or drugs) or technologies (such as negative pressure drainage) could be a possible synergistic strategy for promoting wound healing and repair. In summary, the existing researches indicate that the full potential of BGs in medicine is not fully developed, and the related market is expected to further grow in the future.

All inorganic perovskite nanocrystals of CsPbX3 (X=Cl, Br, I) have superior properties such as high fluorescence quantum yield, narrow half-peak width and easy adjustment of emission wavelength, etc., which have wide application prospects in light-emitting diodes, solar cells and photodetectors. However, CsPbX3 nanocrystals have a poor stability under light, moisture and heating conditions due to their ionic properties, restricting their practical application. Metal ion doping can affect the electronic band structure of CsPbX3 nanocrystals, thus effectively improving their photoluminescence properties. The doped metal ions can be divided into main group metal ion, rare-earth metal ion and transition metal ion. Transition metal ion has a lower price, and it can efficiently improve the stability of perovskite nanocrystals. This review represented recent work on the effect of transition metal ions (i.e., Mn2+, Ni2+, Cu2+, Zn2+, etc.) doping on the properties of perovskite nanocrystals. In addition, the challenges in the development of doped all-inorganic perovskite nanocrystals were analyzed, and the future development direction was also prospected.Transition metal ion doping CsPbX3 nanocrystals (NCs) are explored extensively. Mn2+ doping can improve the formation energy of CsPbX3 NCs with a high stability. Mn2+ doping exists in wide-band-gap perovskite hosts where the excitation energy is transferred to Mn d-state, resulting in short-range tunable yellow-orange d-d emissions. Low concentration Mn2+ doping is beneficial to enhancing the exciton emission, and the Mn2+-doping sample has a long PL lifetime. Ni2+ doping can eliminate the structural defects of CsPbX3 NCs, resulting in the improvement of the lattice order. Compared with the undoped sample, doping Ni2+ also induces the formation energy of CsPbX3 NCs and triggers the phase transition of CsPbBr3 from orthorhombic to cubic, which is attributed to the lattice strain due to the size mismatch of Ni2+ and Pb2+. Ni2+ doping reduces the non-radiative recombination efficiency and increases the fluorescence intensity of the CsPbX3 NCs. Cu2+ doping sample maintains a cubic crystal structure of the initial phase. When Cu2+ with the smaller radius (i.e., 0.72 ?) replaces Pb2+ ion with the larger radius (i.e., 1.19 ?), the shrinkage of lattice contraction regulates the tolerance factor of the [PbBr6]4- octahedron, and increases the lattice formation energy, resulting in high stability and fluorescence intensity of NCs. Compared with Pb2+ (i.e., 1.19 ?), Zn2+ has smaller ion radius (i.e., 0.74 ?) and a lower toxicity. The introduction of Zn2+ can effectively passivate the surface and internal defects of CsPbX3 NCs. Doping Co2+ can increase the formation energy, improve the short-range order of the lattice, and enhance the different monochromatic band edge emission without introducing new composite channels. Doping Cr3+ can regulate the structure and the optical properties of NCs. Doping Fe2+/Fe3+ has some advantages like low cost, high electrical conductivity and magnetic properties. Moreover, an appropriate amount of Fe2+ doping improves the size uniformity of CsPbX3 NCs, as well as increases the PLQY and the average PL lifetime due to the decrease of the defect state and nonradiation recombination of NCs. Furthermore, the application of transition metal ion doped CsPbX3 NCs in white light-emitting diode (WLED), solar cells and laser devices was outlined. The employment of doping transition metal ions into perovskites can improve the optical properties and stability, which energetically facilitate their applications. Finally, the main future research directions of transition metal ion doped CsPbX3 NCs were concluded.Summary and prospects Transition metal doping can promote the radiative recombination of excitons, and the defect formation energies of VCs, VBr and VPb are effectively improved, resulting in an improvement of the luminescence performance of CsPbX3 NCs. However, some problems need to be solved in the future. The PL linewidth of nanocrystals for the application of bioluminescence imaging and LED fields needs to be reduced. The external quantum efficiency (EQE) and luminous efficacy (LE) of transition metal doped CsPbX3 NCs still need to be improved, compared to that of Cd-based QDs. Finally, there are still potential risks associated with the reaction, and the PLQY still needs to be further improved. The investigation of transition metal ion doping all-inorganic CsPbX3 NCs has a profound impact on the scientific research and commercial applications of fluorescent materials.

Foam concrete can be used as a non-load-bearing structure material due to its low density, high porosity and heat and fire resistance, and foam concrete consists of a foaming agent, water, cement, sand and admixture. Air bubbles generate the pores in foam concrete, and the hydration products of the cement, sand, and admixture form a supporting structure. Foam concrete often involves foam collapse, low product strength and susceptibility to cracking during its preparation and practical use. This review was to reveal the foam property of foam concrete and the influence of the matrix materials on the structure and performance. This review also summarized the types and amounts of foaming agents, foam stabilizers, matrix materials, fibers, and additives, and clarified the correlation mechanism between macro- parameters and performance/microstructure.The property of foam concrete is closely related to its foam performance and matrix materials, and foaming agents and foam stabilizers have a dominant impact on the foam performance. At present, foaming agents mainly include physical foaming agents and chemical foaming agents. Physical foaming agents refer to surfactants or active substances with a high activity, forming bubbles via reducing the surface tension of the liquid and producing a double electrical layer on the surface of the liquid film. Physical foaming agents mainly include rosin resin, ionic, protein and composite. Chemical foaming agents produce bubbles via chemical reactions by adding chemical reagents (i.e., H2O2, Zn and NaHCO3, etc.). Rosin resin foaming agents are the earliest and most commonly used foaming agents, but they have problems such as a low foaming ratio and a poor foam stability. Ionic foaming agents increase the bubble foaming ratio. Protein-based foaming agents have higher strength and stability, but have shortages such as higher foaming costs. Compound foaming agents are the fourth-generation foaming agents with synergistic effects and take into account ionic foaming agents. The foaming agent has some advantages like high foaming multiple and good stability of protein foaming agent.As the single-use of foaming agents produces a poor stability of foam and results in a low strength of concrete, the foam stability is used to promote the performance of foam concrete by the foam stabilizing agent with its thickening and increase the viscosity of the liquid film and block drainage. Foam stabilizers can be divided into three categories, i.e., synergistic foam stabilizers, tackifier foam stabilizers, and nano-solid foam stabilizers. The mechanism of action of the synergistic foam stabilizer is to enhance the adsorption of surface molecules, increase the strength of the liquid film, and reduce air permeability, thereby improving the stability of the foam. Viscosity-increasing foam stabilizers are used to increase the liquid phase viscosity of foam to reduce the bubble drainage rate, extend the half-life, and improve foam stability. Based on the adsorption of nanoparticles at the gas-liquid interface, nano-particles have a good foam stabilizing effect. Nano particles transform the original "gas-liquid" foam into a "solid-liquid-gas" three-phase foam, hindering bubble drainage, thereby improving foam stability and refining bubble size.The matrix mainly composed of cement, mineral admixtures and alkali-activated cementitious materials/geopolymer materials is an important factor affecting the performance of foam concrete. Comprehensive consideration of the cement grade and type, the type and amount of mineral admixtures, and the ratio of activators and solid waste materials of alkali activating materials and geopolymer materials is a key to preparing a matrix with a high strength. In addition, the water-cement ratio (W/C), fiber dosage and admixtures are also a key to improving the performance of foam concrete.Summary and prospects The superior performance of foam concrete has promising applications in the fields of building energy saving, roadbed and mine filling. Although the effects of foaming agents and foam stabilizers on the bubble performance are investigated, the bubble performance by different matrix materials and slurries still needs to be further improved, achieving a synergistic enhancement of matrix and bubbles. The large-volume utilization of solid waste matrix materials in foam concrete is a key to realizing the promotion of foam concrete considering the wide range and low price of solid waste materials. The durability of foam concrete still needs to be systematically discussed. In addition, it is also necessary to conduct in-depth exploration of the green multi-purpose application of foam concrete to provide theoretical and experimental supports for the high-quality development of building materials.

Alkali-activated materials (AAMs) are one of green cementitious materials in building materials industry and beneficial to the goals of carbon peaking and carbon neutrality. Compared with ordinary Portland cement (PC) based materials, however, the raw material composition, reaction products and pore solution composition of AAMs are complex and thus their reaction mechanisms and performance evolutions still need to be further clarified. Thermodynamic modelling is an effective method in analyzing AAMs. It can predict the phase assemblage and pore solution composition based on the raw material composition and given reaction conditions, which is of great significance to profoundly investigate the reaction mechanisms and performance evolutions of AAMs. The existing thermodynamic modelling is increasingly applied in AAMs and the related results are achieved. However, the corresponding comprehensive review on the state-of-art in thermodynamic modelling of AAMs is lack. A clear and systematic knowledge of the principles, thermodynamic databases, methods, challenges and gaps remains implicit for thermodynamic modelling of AAMs. In this context, this review summarized recent progress on thermodynamic modelling of AAMs, pointed out the deficiency gaps of current thermodynamic modelling research work and put forward the relevant prospects. This review could provide a theoretical guidance for thermodynamic modelling of AAMs.Chemical reactions in AAMs follow the laws of thermodynamics. There exists two thermodynamic equilibriums in AAMs, i.e., one is between the precursor and aqueous solution and another is between the reaction products and aqueous solution. Thermodynamic modelling can be performed to predict the phase assemblage and pore solution composition of AAMs by assuming the thermodynamic equilobriume. The accuracy and reliability of results by thermodynamic modelling largely depend on the quality of thermodynamic database that consist of solubility products (Ksp), heat capacity (), entropy (), Gibbs free energy(), enthalpy () and molar volume () for all solid, liquid and gas phases involved in the system. The thermodynamic database of AAMs is usually established based on the thermodynamic database of PC via introducing the unique reaction products of AAMs. The unique reaction products and their thermodynamic parameters are available for alkali-activated high-Ca and alkali-activated low-Ca systems.Thermodynamic modelling of alkali-activated slag was initially conducted via the thermodynamic database of PC. Although the modelling results can predict the phase composition evolution, it still needs the corresponding experimental measurements to calibrate. with the established CNASH~~ss model for describing C-(N-)A-S-H gel, thermodynamic modelling is increasingly used to investigate the phase assemblage evolution of alkali-activated slag cements. Besides the phase evolution, thermodynamic modelling is also applied to predict the phase diagram, providing a theoretical basis for the refined design of chemical properties of alkali-activated slag cement. In recent years, thermodynamic modelling tends to be used to investigate the durability of alkali-activated slag cements under single factor action such as carbonation, chloride attack and sulfate attack, as well as under multi-factors action, i.e., the combined attack by chloride and sulfate salts. Thermodynamic modelling is also applied to predict the phase assemblage evolution of alkali-activated low- and medium-Ca systems. However, it is less applied to those for alkali-activated high-Ca system. This is mainly due to the less developed thermodynamic database for alkali-activated low- and medium-Ca systems.In addition, thermodynamic modelling is also coupled with other simulation techniques to numerically analyze AAMs. For instance, a novel numerical model GeoMicro3D was proposed by coupling thermodynamic modelling and lattice Boltzmann method to simulate the reaction process and microstructure formation of alkali-activated slag cement, clarifing the interaction mechanisms between chemical reaction, multi-ions transport and microstructure formation. However, the numerical studies by coupling thermodynamic modelling and other simulation techniques are still limited for AAMs when compared to those for PC based materials.Summary and prospects Thermodynamic modelling has a robustness in studying the phase evolution and durability performance of AAMs induced by chemical reactions. Firstly, thermodynamic modelling can predict the reaction products assemblage and pore solution composition of AAMs. Secondly, thermodynamic modelling can calculate the phase evolution of AAMs under the action of aggressive media, and then study the deteriation mechanism of AAMs. Finally, thermodynamic modelling can be combined with other numerical simulation techniques to investigate AAMs. At present, however, there are still some issues that need to be further studied as follows:1) The incomplete thermodynamic database for alkali-activated low-Ca system is an important reason for the limited thermodynamic modelling studies on alkali activated low and medium calcium systems. It is expected that a thermodynamic model describing the N-A-S-H gel can be established by ab-initio calculations and molecular dynamics simulations with the development of atomic- and molecular-scale simulation techniques.2) It is generally assumed that the amorphous phases in precursors are dissolved synchronously in current thermodynamic modelling of AAMs. However, the heterogeneous distribution of composition and structure of precursor makes this assumption in doubt. The non-uniformity of the dissolution of amorphous phases in precursor is an issue to be further considered in future thermodynamic modelling studies.3) The phase evolution of AAMs is actually a process coupling thermodynamics and kinetics. However, most of the thermodynamic modelling studies only focus on the phase assemblage in the equilibrium state, ignoring the kinetic issues before reaching the equilibrium. Considering the kinetic parameters (i.e., dissolution rate and reaction rate, etc.) in thermodynamic modelling should be a focus of current and future thermodynamic modelling studies.4) The phase evolution, microstructure damage and ions transport are three inter-dependent aspects for studying the durability performance of AAMs. However, the current thermodynamic modelling studies mainly focus on the phase evolution under the chemical attacks, while ignoring the interaction between the phase evolution, microstructure damage and ions transport. In future studies, it is necessary to consider the interaction and establish a chemical-damage-transport model to numerically analyze the durability performance of AAMs.