This review offers a prospective outlook on the development of oxidation-resistant coatings for C/C composites. Firstly, novel reinforcing phases, such as one-dimensional and two-dimensional materials, can be doped to design and optimize the synthesis of ultra-high-temperature oxidation-resistant ceramic materials via efficient artificial intelligence approaches. This involves more complex combinations of the existing oxidation-resistant coatings or mechanisms to actively explore new systems of oxidation-resistant coatings for C/C composites capable of withstanding at > 1700 ℃. Secondly, constructing a complete and reliable performance database for the existing C/C composite oxidation-resistant coatings can elucidate the related protective mechanisms, establish reasonable evaluation mechanisms, and explore key control methods, thereby providing design parameters and theoretical support for practical applications. Finally, a research should focus on the adaptability of oxidation-resistant coatings on C/C composite components of various sizes and configurations, optimizing preparation processes, revealing critical control factors, forming stable process control pathways, and developing relevant process standards.

Appropriate workability is a key to ensuring the smooth pouring, building of concrete, and quality of the project. However, in practical applications, the workability of concrete often undergoes unexpected drastic changes in the early mixing stage, such as rapid loss of flowability, false set and so on. The change of flowability of cement paste overtime in the early stage is mainly related to the early hydration of cement and the adsorption behavior of superplasticizer. Cement pastes with different mineral compositions can exhibit different change rules overtime. In the early stage of hydration, especially in the first hour, the hydration reaction of tricalcium aluminate (C₃A) mainly occurs. Since the reaction of C₃A with water is extremely rapid, a large amount of calcium sulfate is incorporated into cement, and gypsum dihydrate (G, CaSO₄·2H₂O) is usually the most important source of calcium sulfate. Although the existing studies focus on the influence of gypsum types on the hydration of C₃A, the fundamental reasons are still need to be explored.

In this paper, the effect of gypsum types was considered to elucidate the fundamental reasons behind the drastic changes in early flowability of fresh concrete. Five cements were prepared using clinker combined with different types of gypsum (i.e., G, gypsum dihydrate and H, hemihydrate gypsum). The study focused on the time-dependent changes in the flowability of pastes within the first hour, and examined the processes such as adsorption of superplasticizer, dissolution of mineral phase and precipitation of hydration products, thus analyzing the intrinsic mechanisms behind the time-dependent changes in flowability.

Based on the content of C₃A in the clinker, four types of cements with a single gypsum type were prepared via adding G and H, respectively, at two molar ratios of SO₄^2- to C₃A of 0.5 and 1.5, and one type of cement with mixed gypsum was prepared at a molar ratio of SO4^2- to C₃A of 1.5 and a molar ratio of G to H of 2. In the test, the water binder ratio was fixed at 0.3, and the amount of PC was adjusted to achieve the similar initial flowability (i.e., (250 ± 5) mm at 2 min). At 2, 5, 10, 20, 30 and 60 min after adding water, the paste was taken out for various tests (i.e., flow spread, analysis of interstitial liquid, adsorption and hydrated sample preparation).

The flow spread of paste was tested according to the Chinese standard GB/T 8077—2012. The adsorption behavior of PC and the concentrations of [Ca] and [S] in the interstitial liquid were obtained by TOC and inductively coupled plasma mass spectrometry (ICP-MS) on the interstitial liquid obtained by centrifugation. The hydration was stopped with cold isopropyl alcohol (5 ℃) to perform the microscopic analysis on the hydrated sample, including differential scanning calorimetry (DSC) and surface area measurement by BET based on nitrogen adsorption.

For five groups of cement pastes with the similar initial flowability, the addition of G leads to a rapid loss of flowability in the first 10 min, and increasing the amount of G can reduce the loss of flowability. The addition of H also leads to a rapid loss of flowability in the first 10 min, but then the flowability slowly increases. Moreover, the higher the amount of H added is, the more significant loss of flowability in the first 10 min and the greater the subsequent increase of flowability will be.

For cement with only G added, the precipitation of ettringite mainly occurs in the first hour, and AFt rapidly forms in the first 5 min and then increases slowly. However, for cement with H, two processes occur simultaneously, i.e., ettringite precipitation and H dissolution -G precipitation. H promotes the formation of ettringite, and the ettringite content in the paste is higher than that with only G added.

For cement with only G added, the precipitation of ettringite leads to a large increase in the specific surface area of the system, which in turn causes the loss of paste flowability over time. For cement with H added, the precipitation of ettringite and large-sized G leads to a rapid loss of the early flowability of the paste, but the gradually decreasing sulfate ion concentration in the interstitial liquid is conducive to the adsorption of PC on the surface of the mineral phases, which in turn leads to a reverse increase in the subsequent flowability.

The dissolution of mineral phases and the precipitation/growth of hydration products were key factors causing the drastic changes in early fluidity of the paste. For cement pastes containing dihydrate gypsum, the formation of ettringite led to a significant increase in the specific surface area of the samples, causing a time-dependent loss in fluidity. In contrast, for cement paste containing hemihydrate gypsum, in addition to ettringite, the precipitation of large-sized dihydrate gypsum accelerated the early loss of flowability. However, the gradually decreasing concentration of SO₄^2- was conducive to the adsorption of superplasticizer, thus leading to an increase in flowability of paste.

In coastal engineering, the durability of reinforced concrete structures is often compromised due to the corrosion of steel reinforcements caused by chloride ingress. This issue significantly affects the safety and longevity of these structures. In concrete materials, the interfacial transition zone (ITZ) between the cement matrix and aggregates is a critical point, affecting the pathways and rates of chloride ion diffusion. Although previous studies demonstrate that the incorporation of nanomaterials can enhance the durability of cement-based materials, the uneven dispersion and agglomeration of these nanomaterials restrict their effectiveness. In this paper, a modification method was proposed via coating graphene oxide (GO) onto aminated sand particles to improve the dispersion of GO within cement-based composites, thereby enhancing their durability. This approach could have a promising potential for extending the service life of coastal engineering structures.

In this study, a GO dispersion was firstly prepared and dispersed by ultrasound. Subsequently, a specific chemical treatment method was employed to coat GO onto standard sand surfaces, creating aminated functionalized sand (GO-MAs). The preparation process of GO-MAs included sand pretreatment, amination, and chemical bonding with GO. In the experiments, cement-based composites with different GO contents (i.e., 0.1%, 0.3%, 0.5% by mass) were prepared with high-performance polycarboxylate superplasticizers to enhance GO dispersion. The chloride ion resistance of these composites was evaluated by rapid chloride migration (RCM) tests and natural chloride diffusion tests. In addition, the mechanisms of GO-coated sand enhancing the microstructural properties of cement-based materials, particularly the ITZ, were analyzed by mercury intrusion porosimetry (MIP), scanning electron microscopy (SEM), energy-dispersive X-ray spectroscopy (EDS), X-ray diffraction (XRD), and thermogravimetric analysis (TG/DTG).

The experimental results indicate that the incorporation of GO significantly reduces the chloride ion penetration depth and migration coefficient in cement-based materials. At GO content of 0.3%, the chloride ion penetration depth and diffusion coefficient when using the method of using GO-coated sand can be reduced by 24.8% and 18.5%, respectively. The MIP tests further reveal that the size of harmful pores (i.e., greater than 50 nm) is decreased by 36.8%, indicating that GO-coated sand effectively improves the pore structure of cement-based materials, reduces porosity, and enhances material density. The microstructural analysis results reveal that GO-coated sand primarily enhances the performance of cement-based composites via optimizing the ITZ. The coupling agent forms stable chemical bonds with the sand surface, while the covalent bonding of amide bonds between GO and the coupling agent improves the dispersion of GO and promotes ITZ hydration, reducing defects and cracks, thereby enhancing resistance to chloride ion transport. The microstructural morphology analysis shows that the samples without GO have more and larger pores, while the samples with GO-coated sand exhibit significantly fewer and denser pores. The EDS analysis indicates that GO-coated sand promotes the enrichment of hydration products at the ITZ interface of cement-based materials, enhancing the ITZ density. Based on the XRD and TG/DTG analysis, GO-coated sand promotes the hydration reaction of cement, thus generating more hydration products, especially C-S-H and Ca(OH)₂ crystals. The formation and arrangement of these hydration products play a crucial role in maintaining a high alkaline environment within cement matrix and protecting a passivation film on the steel surface. The qualitative comparative analysis of the crystallographic orientation index of Ca(OH)₂ crystals indicates that GO-coated sand reduces the number of Ca(OH)₂ crystals perpendicular to the interface in the ITZ, optimizing the directional arrangement of Ca(OH)₂ crystals in the ITZ, which is beneficial for strengthening the weak interface and enhancing its density.

The method of using GO-coated sand effectively could enhance the resistance of cement-based materials to chloride ion erosion, and the optimal effect occurred at a GO content of 0.3%. This modification method optimized the ITZ through a "dual-bond effect," reducing porosity and increasing material density, thereby enhancing the barrier against chloride ions. This research could provide an effective technical approach to improving the durability of cement-based materials and offers important scientific evidence for the protection of concrete structures in coastal engineering. A future research needed to explore the long-term performance of GO-coated sand under different environmental conditions and its application effects in practical engineering, thus providing comprehensive theoretical support and technical guidance for the modification and engineering application of cement-based materials.

Retarders are necessary in construction projects such as long-distance transportation, high-temperature construction in summer and large-volume concrete construction. However, in practical applications, abnormal setting phenomena such as rapid or non-setting of concrete often occur due to the possibly incorrect or excessive mixing, which can affect the progress of the project and cause economic losses. It is thus of great significance for accurately regulating the setting and hardening properties of cement-based materials to clarify the cement hydration process and its mechanism under the action of retarders. Citric acid (CA) and its salts (mainly sodium citrate (SC) and potassium citrate (PC)) are commonly used as retarders. However, there are still some controversies about their effect and mechanism in the existing studies. In this paper, CA, SC and PC with different dosages and addition methods were used to clarify the effect and mechanism of these retarders on cement hydration and solve the problem of abnormal setting during the use of such retarders. Their effects on the setting time of cement paste, hydration heat, and adsorption capacity under the action of different sulfate carriers (anhydrite (AH) and hemihydrate (HH)) were investigated. In addition, the possible mechanism was also analyzed.

A fixed water-to-cement mass ratio of 0.4 was used. The dosage of citric acid and its salt was calculated based on a mass ratio of citrate to cement, and the dosage of citrate ion was 0.045%, 0.090%, 0.180% and 0.360% (mass fraction), respectively. Two addition methods were used, i.e., W addition method and S addition method. For W addition method, the retarders were first dissolved in water and then mixed with cement. For S addition method, the retarders were first mixed with cement as a solid powder, and then the mixture was added to water. Two kinds of cement with different types of sulfate carriers were prepared, i.e., AH-contained cement prepared with anhydrite and HH-contained cement prepared with hemihydrate.

The setting time was measured according to the Chinese standards GB/T 1346—2011. The hydration of cement paste at (25.00±0.02) ℃ was measured by a model TAM Air 8-channel isothermal calorimeter (Newcastle, DE, USA). The adsorption capacity of different retarders after hydrating for 5 min was measured by a model TOC-L total organic carbon analyzer (Shimadzu Co., Japan). The change in conductivity when AH and HH dissolved in CA, SC and PC solutions with a citrate concentration of 0.047 6 mol/L was determined by a model FE38 conductivity meter (Mettler Toledo Co., Switzerland). In addition, the ionic concentrations of calcium and sulfate during the dissolving process were measured by a model 5800 inductively coupled plasma emission spectrometer (Agilent Co., USA).

The effect of retarders on the cement hydration process can be affected by different addition methods, dosages, cations of retarders and different sulfate types in cement. For cement with different sulfate carriers, the setting time of HH-contained cement is longer than that of AH-contained cement regardless of the dosage of retarder and the addition method used. The abnormal rapid setting phenomenon of retarder can occur in the AH-contained cement rather than in HH-contained cement. Regardless of the type of retarders used, the initial dissolution capacity of HH is greater than that of AH. For different addition methods, the retarder with S addition method shows a greater retarding ability. The effect of citric acid and citrates indicates that the retarding ability of PC is greater than that of SC.

The results of hydration heat show that a phenomenon of abnormal rapid setting in AH-contained cement can be caused due to the rapid hydration of the aluminate-containing phase. The results of conductivity and ionic concentration indicate that HH quickly dissolves and releases sufficient sulfate ions, thus inhibiting a rapid reaction of C₃A and avoiding the occurrence of rapid setting. Moreover, the ability of PC to promote the dissolution of sulfate carrier is greater than that of SC possibly due to its greater complexing ability with calcium ions. Therefore, the complexation can play an important role in determining the effect of retarders.

The type of sulfate carrier in cement could change the effect of citric acid/citrate. For AH-contained cement, there existed a phenomenon of first retarding and then rapid setting with the increase of retarder dosage. Such a phenomenon was avoided in HH-contained cement. The addition method of citric acid/citrate greatly changed its effect. The S addition method of retarders enhanced its retarding ability. The adsorption and complexation of citric acid/citrate with calcium ions both played an important role in determining the cement hydration process.

Thermal expansion analysis is a most widely used method for determining the firing temperature of ancient ceramics, which infers the original firing temperature of ancient ceramics via analyzing the inflection point of the thermal expansion curve. However, for the pottery at an original firing temperature that is lower than the formation temperature of glass phase, the inflection point on the thermal expansion curve is actually the formation temperature of glass phase rather than original firing temperature. To solve a problem of inaccurate determination of low-temperature pottery firing temperature by a thermal expansion method, clay samples fired at different temperatures were prepared with clay as a raw material, and analyzed by X-ray diffraction (XRD), synchronous comprehensive thermal analyzer (TG-DTA), and scanning electron microscopy (SEM). In addition, the first derivative curve corresponding to the top rod thermal expansion curve of clay samples fired at different temperatures fitted to the dehydroxylation expansion peak of kaolinite, and the relationship curve between the dehydroxylation expansion peak of kaolinite and the original firing temperature was also obtained.

A clay as a raw material was processed by grinding, screening, drying, pressing, and sintering at different temperatures. Afterwards, the cylindrical samples with a diameter of 5 mm and a length of 25 mm were prepared with clay by polishing. The thermal behavior of clay raw materials was determined by thermogravimetry-differential thermal analysis (TG-DTA). The chemical composition, phase composition, and microstructure of the raw materials and the samples fired at different temperatures were characterized by X-ray fluorescence spectroscopy (XRF), X-ray diffraction (XRD), and scanning electron microscopy (SEM). The linear expansion or contraction process of the samples during the refiring process at different firing temperatures was tested by a thermal expansion instrument.

The XRF spectra show that the chemical composition of raw materials and samples at different firing temperatures is the similar. The XRD patterns and TG-DTA analysis of raw materials indicate that the dehydroxylation of kaolinite occurs at 300-600 ℃. Based on the XRD patterns, the phases in clay samples fired at different temperatures are consistent with those in clay raw materials. At < 900 ℃, the characteristic diffraction peak intensity of kaolinite in clay samples fired at different temperatures changes with the increase of firing temperature. The SEM images show that the typical layered structure of kaolinite in clay raw material begins to lose its structural hydroxyl groups and gradually transforms into metakaolinite, and this transformation gradually deepens as the firing temperature increases. The results of the thermal expansion test of clay samples fired at different temperatures indicate that there is a linear relationship between the dehydroxylation expansion peak of kaolinite on the first derivative curve of thermal expansion and its original firing temperature. This linear relationship can be expressed by a formula, i.e., A= -2.278 481 7×10^-6T+3.337 870 9×10^-3. To verify the accuracy of the formula, the thermal expansion test of clay samples is conducted using a push rod, and the calculated firing temperature is not significantly different from its original firing temperature, which are 41℃ and 35 ℃, respectively. The push rod thermal expansion method can be thus used to determine the original firing temperature of low-temperature pottery made from clay.

Before the glass phase formation temperature, kaolinite in the sample could lose the hydroxyl group in its structure and gradually transform into metakaolinite as the firing temperature increases, thus causing the expansion of the thermal expansion curve. The intensity of the thermal expansion peak of kaolinite dehydroxylation gradually decreased as the firing temperature increased. The first-order derivatives of the thermal expansion curves of the samples fired at different temperatures were taken, fiting the kaolinite dehydroxylation peak. This linear relationship with the firing temperature could be expressed by a formula, i.e., A= -2.278 481 7×10^-6T+3.337 870 9×10^-3. The clay samples fired at 550 ℃ and 650 ℃ were used for the thermal expansion curves, respectively. The clay samples fired at 550 ℃ and 650 ℃ were used to verify the reasonableness of the fitting formula. It was indicated that the firing temperature of low-temperature pottery made of clay could be determined by the top bar thermal expansion method.

To address these challenges, some solutions are explored. On one hand, the molecular structure of PCE is improved to enhance its stability and adaptability in complex environments. On the other hand, some concepts from nanotechnology and smart materials are proposed to develop self-healing PCE products. These innovative studies can improve the overall performance of concrete and pave some effective pathways for the green development of the construction industry. Polycarboxylate superplasticizers play a vital role in driving the concrete industry towards energy saving and emission reduction. A future research will focus on enhancing PCE performance, cost reduction, and expanding its application fields. With the advancement of technology and the guidance of innovative thinking, we can believe that PCE will play an even greater role in promoting sustainable development in the construction industry.

Chloride induced corrosion damage of the embedded reinforcement is one of primary reasons for the deterioration of concrete structures in civil engineering. Admixing organic corrosion inhibitors are considered as an effective way to halt the corrosion damage of the embedded reinforcement for its convenient use and high cost-efficiency. However, the corrosion inhibition efficiency of conventional corrosion inhibitors exists due to the limited amount of adsorption functional groups in molecular structures. It is thus essential to develop novel corrosion inhibitors with a great adsorption capacity to satisfy the efficient corrosion protection of reinforcement concrete in harsh marine conditions. As typical crystalline hybrid materials, metal-organic frameworks (MOFs) are self-assembled via coordination between metal nodes and organic ligands. Zeolitic imidazolate frameworks (ZIF-8, a representative MOF material) exhibit many heteroatoms and π-electron system, thus enabling its interaction and high corrosion inhibition potential for reinforced concrete. In this study, ZIF-8 corrosion inhibitors were synthesized and its inhibition effect for reinforced mortar was investigated. In addition, the characteristics of the reinforcement/mortar interface were also evaluated by image analysis based on scanning electron microspcopy (SEM) images and chloride titration.

In the experiments, 0.60 g Zn(NO₃)₂·6H₂O was dissolved in 4.5 mL deionized water. 11.50 g 2-methylimidazole was dissolved in 40.0 mL deionized water. The zinc nitrate solution above was then added into 2-MeIm solution in dropwise under stirring. After stirring for 24 h, white ZIF-8 corrosion inhibitors were collected via centrifugation at 10 000 r/min and then cleaned with deionized water for three times. The water to cement (w/c) ratio was 0.5 and the sand to cement (s/c) ratio was 3. The dosage of ZIF-8 corrosion inhibitors was 0% (blank), 0.2%, 0.4% and 0.8% per dry cement weight in mortar specimen, respectively. The reinforcement electrode was centrally embedded in cylindrical mortar specimen (Φ50×100 mm). The reinforced mortar specimens were then half-immersed in 3.5% (mass fraction) NaCl solution (i.e., 25±1 ℃ and 60±5% relative humidity). The chemical composition, thermal stability, morphology, and specific surface area of ZIF-8 corrosion inhibitors were determined by Fourier transform infrared spectroscopy, thermal gravimetric analysis (TG), field emission scanning electron microscopy (FE-SEM) and N₂ physisorption, respectively. Electrochemical workstation was used to evaluate the electrochemical behavior of the reinforced mortar during 210 d immersion with three electrode systems. The morphologies and chemical elements of the reinforcement were characterized by FE-SEM combined with energy dispersive spectroscopy (EDS). The compositions of product film on the reinforcement surface were evaluated by X-ray photoelectron spectroscopy. The total porosity at the reinforcement/mortar interface was obtained by image analysis based on the FE-SEM images acquired in backscattered electrons (BSE) mode. Chloride content in mortar matrix within 2 mm from the embedded reinforcement surface was measured by a potentiometric titrator.

ZIF-8 is prepared by a facile and eco-friendly solvent method. The negative OCP shift of the embedded reinforcement is dramatically prolonged by the admixed ZIF-8 corrosion inhibitors. The more pronounced beneficial effect is relevant to a higher dosage of ZIF-8 corrosion inhibitors. Corrosion initiation time for the embedded reinforcement in mortar specimens is delayed by 8.3%-43.3% at 0.2%-0.8% (mass fraction) ZIF-8 corrosion inhibitors. The corrosion damage of the reinforcement embedded in mortar specimens with ZIF-8 corrosion inhibitors merely occurs after 180 d immersion. ZIF-8 corrosion inhibitors can fill in the pore space in cement matrix due to small particle size (i.e., 75 nm), and the bulk mortar resistance of the reinforced mortar are continuously reduced with immersion age. Compared to the blank specimen, a higher resistance of the formed product layer and a lower capacitance of the formed product layer are relevant to the specimens with ZIF-8 corrosion inhibitors at the same immersion stage. Meanwhile, a higher charge transfer resistance and a lower electric double-layer capacitance are also related to the specimens with ZIF-8 corrosion inhibitors at the same immersion stage. Besides, a lower anodic current density appears for the embedded reinforcement in a mortar with ZIF-8 corrosion inhibitors, and this beneficial effect is more pronounced for a higher dosage of ZIF-8 corrosion inhibitors. The inhibition efficiency is 62.1% in a reinforced mortar in the presence of 0.8% (mass fraction) ZIF-8 corrosion inhibitors after 180 d immersion. The reduced dimensions and increased packing density of the corrosion products occur for the specimen with a higher dosage of ZIF-8 corrosion inhibitors. The chemisorption of ZIF-8 corrosion inhibitors is confirmed by the formation of coordination bond between nitrogen atoms in ZIF-8 corrosion inhibitors and iron atoms on the reinforcement surface. Furthermore, the enhanced peak intensity of Fe₂O₃ and FeOOH, together with the reduced peak intensity of FeO appear for the specimen with 0.8% ZIF-8 corrosion inhibitors. The porosity of cement matrix at the reinforcement/mortar interface is reduced by 9.8 % and the chloride content in the vicinity of the embedded reinforcement is efficiently reduced by 28.6% with the admixed ZIF-8 corrosion inhibitors in mortar specimen.

The corrosion initiation time of the embedded reinforcement in mortar was efficiently prolonged by 8.3%-43.3% in the presence of 0.2%-0.8% ZIF-8 corrosion inhibitor. ZIF-8 corrosion inhibitors could reduce the corrosion current density and increase the charge transfer resistance and polarization resistance of the embedded reinforcement after corrosion initiation. The inhibition efficiency of 0.8% ZIF-8 corrosion inhibitors was maintained at 62.1% after 180 d immersion in 3.5% NaCl solutions. ZIF-8 corrosion inhibitors exhibited an effective adsorption on the reinforcement surface, thus reducing chloride content on the reinforcement surface. Also, more highly protective products (FeOOH and Fe₂O₃) could form on the reinforcement surface. ZIF-8 corrosion inhibitors refined the pore structure at the reinforcement/mortar interface and reduced chloride content at the vicinity of the reinforcement. For specimen with 0.8% ZIF-8 corrosion inhibitors, the porosity of cement matrix at the reinforcement/mortar interface was reduced by 9.8%, and the chloride content of cement matrix in the vicinity of reinforcement was declined by 28.6%.

Thermal barrier coatings (TBCs) are capable of effectively isolating high-temperature gases, protecting the turbine blade matrix, reducing blade surface temperatures, and consequently enhancing aero engine thrust. TBCs with a thickness ranging from 100 to 500 μm in conjunction with internal turbine blade cooling can decrease the surface temperature of the superalloy by 100-300 ℃. In recent decades, yttrium-stabilized zirconia (YSZ) containing 7%-8% (by mass) is widely utilized as a material for thermal barrier coatings (TBCs) due to its superior properties. With the increasing demands for thrust-to-weight ratio, weight reduction, and stability in aviation turbine engines, the surface temperature of blade is expected to surpass 1300 ℃. When the operating temperature exceeds 1200 ℃, the YSZ coating undergoes the phase transformation and sintering. For cooling, it transitions from the tetragonal phase to the monoclinic phase, leading to coating volume expansion. Moreover, calcium-magnesium-alumino-silicate (CMAS) compounds derived from dust, sand, and volcanic ash can melt on the TBC surface, causing a severe erosion and ultimately leading to coating failure. The existing CMAS erosion becomes a challenge for TBCs. It is thus necessary for the evolving requirements of future turbine engine development to develop novel thermal barrier coating materials. Recent studies focus on new TBCs materials, including rare-earth zirconates, phosphates, and hafniates, which exhibit an enhanced CMAS erosion resistance, compared to YSZ. In this paper, the CMAS erosion resistance of rare-earth zirconates was investigated.

Y₂Zr₂O₇ (H1), (Y_0.5Gd_0.5)₂Zr₂O₇ (H2), (Y_0.33Gd_0.33Er_0.33)₂Zr₂O₇ (H3), (Y_0.25Gd_0.25Er_0.25Yb_0.25)₂Zr₂O₇ (H4), (Y_0.2Gd_0.2Er_0.2 Yb_0.2Lu_0.2)₂Zr₂O₇ (H5) were synthesized by a solid-state synthesis method at 1650 ℃ for 10 h. The glass phase was prepared according to the classical composition of 33CaO-9MgO-13Al₂O₃-45SiO₂. In the CMAS erosion resistance test, CMAS powder was uniformly applied to ceramic discs at a density of 20 mg/cm². The discs were then placed in a crucible and maintained in an air atmosphere at 1300 ℃ for 5 h to conduct the erosion experiment. Subsequently, the samples were sectioned along the centerline, and the cross-sections were ground and polished to achieve a surface roughness of 1 μm, facilitating the observation of the morphology and thickness of the eroded layer. To analyze the reaction products of different zirconate ceramics during CMAS erosion, CMAS powder was mixed with five types of zirconate ceramic powders at a mass ratio of 1∶1 and heated in an air atmosphere at 1300 ℃ for 5 h.

All the five rare-earth zirconate ceramics exhibit a single, defective fluorite structure. After five rare-earth zirconate ceramics are eroded at 1300 ℃ for 5 h, the CMAS erosion depths are ranked as H2 > H3 > H1 > H4 > H5. Despite the fact that the CMAS corrosion resistance of rare-earth zirconate ceramics does not improve with increasing entropy, the CMAS corrosion resistance of high-entropy rare-earth zirconate ceramics is superior to that of low- and medium-entropy ceramics. After heating at 1300 ℃ for 5 h, the erosion depth of H5 is only 22 μm. The mechanism of resistance to CMAS erosion in rare-earth ceramics can be elucidated through the "dissolution-reprecipitation" mechanism. Upon interaction between rare-earth zirconate ceramics and CMAS, RE³⁺ and Zr⁴⁺ dissolve from the ceramic matrix into the CMAS melt. RE³⁺ ions with larger ionic radii (such as Gd³⁺ and Y³⁺) react with Ca²⁺ and Si⁴⁺ in the CMAS to form RE-apatite. Conversely, RE³⁺ ions with smaller ionic radii (such as Er³⁺, Yb³⁺, and Lu³⁺) preferentially incorporate into ZrO₂ lattice, forming a fluorite-structured RE-ZrO. When the precipitation rates of RE-apatite and RE-ZrO₂ exceed the erosion rate of CMAS, a dense protective layer forms, effectively preventing further CMAS erosion. The average rare-earth ion radius and optical basicity both affect the corrosion resistance of five rare-earth zirconate ceramics to CMAS.

Y₂Zr₂O₇, (Y_0.5Gd_0.5)₂Zr₂O₇, (Y_0.33Gd_0.33Er_0.33)₂Zr₂O₇, (Y_0.25Gd_0.25Er_0.25Yb_0.25)₂Zr₂O₇, and (Y_0.2Gd_0.2Er_0.2Yb_0.2Lu_0.2)₂Zr₂O₇ were synthesized by a high-temperature solid-state reaction method. These five rare-earth zirconate ceramics exhibited a defective fluorite crystal structure. The analysis of the mechanism of resistance to CMAS erosion in these ceramics revealed that the erosion layer consisted of two distinct layers, i.e., an infiltration layer and a reprecipitation layer. The reprecipitated layer primarily comprised near-spherical fluorite-structured RE-ZrO₂ and RE-apatite crystalline pillars, forming a dense barrier that prevented further CMAS penetration into the substrate. The impact of configurational entropy on the CMAS corrosion resistance of these five rare-earth zirconate ceramics was investigated. The CMAS corrosion depths were found to be H2 > H3 > H1 > H4 > H5, and H5 exhibited the optimum CMAS corrosion resistance, showing a corrosion depth of only 22 μm after heating at 1300 ℃ for 5 h. The findings indicated that while the resistance to CMAS erosion in rare-earth zirconate ceramics did not strictly increase with increasing configurational entropy, high-entropy rare-earth zirconate ceramics demonstrated a superior resistance. compared to low- and medium-entropy counterparts. In addition, the CMAS corrosion resistance of rare-earth zirconate ceramics was influenced by both configurational entropy and the radius of rare earth ions. A concept of optical basicity was introduced to evaluate the corrosion resistance of five rare-earth zirconate ceramics against CMAS. Under identical conditions, a larger optical basicity difference (ΔΛ) between the rare earth zirconate ceramics and CMAS resulted in a higher reaction rate and deeper erosion by CMAS. The optical basicity could serve as a valuable reference for selecting new thermal barrier coating materials.

Magnesium silicate hydrate (M-S-H) cement is a low-carbon cementitious material. This paper discusses the mechanism of dipotassium phosphate (DKP) accelerating the hydration of hydrated magnesium silicate cement, and confirms that the introduction of DKP accelerated the formation of M-S-H gel due to the alkali induction effect and K-struvite seed crystal induction effect. This improves the compressive strength and subsequent strength of hydrated magnesium silicate cement, which can increase 3.25, 13.78, 22.23, 34.58 MPa and 37.96 MPa at the age of 1, 3, 7, 14 d and 28 d, respectively. Meanwhile, different additives can induce the formation of different M-S-H crystal phases. The acceleration of the hydration process by these additives is confirmed through X-ray diffraction (XRD), ²⁹Si nuclear magnetic resonance spectroscopy (NMR), scanning electron microscopy (SEM) and density functional theory calculations. M-S-H is the main source of strength for hydrated magnesium silicate cement, and this paper also emphasizes that suitable additives can induce a faster formation of M-S-H structures with stronger cementitious properties, thereby improving the hydration rate and early compressive strength of magnesium silicate hydrate cements, and further expanding the application range of magnesium silicate hydrate cements.

MgO: SF mass ratio for preparing MSHC samples was 40 ∶ 60 (i.e., the amount of substance was 1∶1), the water/cement (w/c) ratio was 0.45, the SHMP content was 1% (mass fraction) of the total solid mass, and the samples were all cured at 25 ℃ and 95% RH for 28 d. When the samples were cured at different ages (i.e., 1, 3, 7, 14 d and 28 d), they were broken into small pieces with the particle sizes of 1-3 cm and soaked in isopropyl alcohol to stop their further hydration. The broken samples were dried in a vacuum drying oven at 40 ℃ until there was no more mass loss.

The crystal phases of the samples were determined by a model Empyrean Rietveld X-ray diffractometer (XRD, Malvern Panalytical Co., The Netherlands). The morphology of the samples was analyzed by a model JSM-IT300 scanning electron microscope (SEM, JEOL Ltd., Japan). The nuclear magnetic resonance map of the samples was analyzed by a model Advance III 400 nuclear magnetic resonance spectroscope (NMR, Bruker Co., Germany).

This work was also analyzed based on density functional theory calculations (DFT). All the models were geometrically optimized and subsequent calculations were performed by a software named Vienna Ab-initio Simulation Package (VASP). The projection enhanced plane wave method (PAW) was used in the calculation, and the cutoff energy was set to 500 eV, which was obtained after the convergence test. The exchange-correlation potential was approximated by Perdew-Burke-Ernzerhof (PBE) under generalized gradient approximation (GGA). A 15×15×9 K-point grid was used to sample the first Brillouin region for the stress-strain calculation. Young's modulus and Poisson's ratio of each M-S-H crystal phase were calculated via vaspkit.

The DKP increases the compressive strength of MSHC by 3.25, 13.78, 22.23, 34.58 MPa and 37.96 MPa at 1, 3, 7, 14 d and 28 d, respectively. This mechanism is since the DKP induces the formation of more suitable M-S-H crystal phase with a higher mechanical strength, and the DKP accelerates the hydration rate of M-S-H, especially the early hydration rate. The SEM images confirm the acceleration effect of the DKP on the hydration. The results of Rietveld XRD and DFT calculation indicate that the DKP can induce the formation of M-S-H crystal structure with a greater gellability. The M-S-H crystal phase-D induced by SHMP is of the lowest Young's modulus and the highest Poisson's ratio. Also, its M-S-H crystal phase generation is rather small, compared with other groups, resulting in its lowest theoretical compressive strength. The M-S-H principal crystal phase induced by DKP is a superhydrous phase B, which is of the highest Young's modulus and the lower Poisson's ratio, leading to the highest theoretical compressive strength. This indicates that the DKP induces the formation of M-S-H crystalline phase with a greater compressive strength. The main crystal phase of M-S-H induced by KOH and K-struvite are talc and anthophyllite, respectively, which are of higher Young's modulus and lower Poisson's ratio, thus improving the compressive strength.

The admixture DKP could improve the compressive strength of MSHC via inducing the MSHC system to generate more suitable and higher mechanical strength M-S-H crystal phase, accelerating the hydration process of MSHC, and increasing the compressive strength of MSHC samples (i.e., 3.25, 13.78, 22.23, 34.58 MPa and 37.96 MPa) at different ages (i.e., 1, 3, 7, 14 d and 28 d). The results by Rietveld X-ray diffraction refinement showed that different admixtures induced the hydration of MSHC to form different M-S-H crystal phases, which could affect the cement strength. The DKP induced M-S-H to form a superhydrous phase B, which was a M-S-H crystal phase with a high Young's modulus and a low Poisson's ratio. This was one of the mechanisms due to the DKP enhancing MSHC. This wor'k indicated that suitable admixtures could induce the formation of M-S-H crystal structure with a greater cementitious property, thereby improving the hydration rate and early strength of MSHC.