Organic compound formaldehyde (HCHO) is a toxic volatile gas with an irritating odour, which is mainly derived from building materials, furniture and various adhesive coatings. Living in an environment containing formaldehyde gas for a long time can cause a series of hazards to human health, such as headaches, allergies, nausea, and even genetic mutations and leukaemia. It is thus necessary to design a gas sensor with a superior sensing performance to detect formaldehyde in real time. ABO₃ type chalcogenide metal oxides have attracted much attention in sensor research, but their application is restricted due to the drawbacks such as poor sensitivity and high operating temperature. Many studies indicate that the modification of metal oxides with two-dimensional transition metal dihalides (2D TMDs) is one of the effective strategies to improve the gas sensing performance, and among the 2D TMDs materials, WS₂ has aroused much interest in sensor research because of its high specific surface area, good thermal stability and excellent electron mobility. The gas sensor performance of metal oxides can be improved via the combination of WS₂ to form a novel nanocomposite material. Hence, in this paper, a series of WS₂-CdSnO₃ composites with different mass ratios were prepared by an ultrasonic mixing method. The microstructure, crystal structure, and valence state structure of WS₂-CdSnO₃ composite material were characterized, and the gas sensing performance of the composite material was analyzed.

CdSnO₃ was prepared with Cd(NO₃)₂·4H₂O (AR, Shanghai Sarn Chemical Technology, China) and SnCl₄·5H₂O (AR) (Shanghai Aladdin Reagent Co., China). CdSnO₃ was ultrasonically compounded with WS₂ (99.9%) to obtain WS₂-CdSnO₃ composites. The physical phase composition, surface morphology, elemental composition, valence and pore size distribution and specific surface area of the composites were characterized by a model SmartLab SE X-ray diffractometer (XRD, Rigaku Co.,), scanning electron microscope (SEM, Tescan Mira Lms), a model K-Alpha X-ray photoelectron spectrometer (XPS, Thermo Scientific Co., USA), and a model ASAP 2460 fully automated specific surface and porosity analyser (BET, Micromeritics Co., USA).

For the determination of gas sensing performance of WS₂-CdSnO₃ composite material, WS₂-CdSnO₃ composite material was put into a mortar with 1-2 drops of terpineol, and thoroughly mixed and ground. The paste was evenly applied to the outer wall of the alumina ceramic tube with a small brush, and an Ni-Cr heating wire put through the ceramic tube and welded firmly to the hexagonal base. The operating temperature was indirectly controlled via adjusting the amount of power added to the two sections of the heating wire. The sensitivity to gases is expressed as a ratio of the stable resistance value of the gas sensing material in air to a stable resistance value in the gas tested. The gas sensing performance of the composites was investigated at a relative humidity of 55% (except for humidity tests).

From the SEM images, the microcube appeared is CdSnO₃, and the cubes structure is agglomerated on the surface of layered WS₂. Based on the EDS analysis, there are five elements Cd, Sn, O, W and S in the composites. It is indicated that WS₂-CdSnO₃ can be prepared. According to the XPS analysis, there are more oxygen vacancies in the sample WSC-2, which can provide more sites for oxygen adsorption, so that more target gases participate in the reaction on the surface of the material. The introduction of WS₂ increases the specific surface area of the sample WSC-2, and the material with a larger specific surface area has more gas adsorption active sites, thus improving the gas sensing performance of the sensor. The sample WSC-2 has the maximum response value of 15.7 for 0.01% formaldehyde at 140 ℃. To check the potential practical application of the prepared WSC-2 sensor, the selectivity and sensitivity of the sensor to different gases were investigated at 140 ℃. The results show that the sample WSC-2 has the optimum selectivity for 0.01% formaldehyde, and the response value is 4.2 times higher than that of CdSnO₃.

A series of WS₂-CdSnO₃ composites with different ratios were prepared via ultrasonic dispersion mixing. WS₂-CdSnO₃ had a large specific surface area, which was beneficial for gas adsorption and diffusion. The gas sensing performance of CdSnO₃ with different ratios of WS₂-CdSnO₃ composites were investigated, and the results showed that 3.5% WS₂-CdSnO₃ had the optimum selectivity for formaldehyde and the superior gas sensing performance at 140 ℃, and the sensitivity for 0.01% formaldehyde was 4.2 times higher than that of CdSnO₃ material, and the detection limit was 0.0001%. Therefore, the 3.5% WS₂-CdSnO₃ composite had potential applications in indoor formaldehyde detection. This study indicated that the modification of metal oxides by two-dimensional TMD could be one of the effective strategies to improve the gas sensing performance.

Glass scintillators are commonly utilized for the detection of high-energy rays in various applications such as nuclear medicine imaging, industrial non-destructive testing, high-energy nuclear physics, safety inspection, and environmental monitoring. The scintillation glass with a high density is critical for improving the absorption and interception capabilities of rays and particles, thereby favoring the enhancement of the detector efficiency. This review represented recent development on high-density scintillation glass doped with rare-earth ions like Ce³⁺, Tb³⁺, and Eu³⁺. The critical parameters such as density, decay time, light output, and radiation resistance affecting the performance of the scintillation glass were discussed.

Overall, the research delves into the scintillation mechanism of scintillation glass and presents an overview of the preparation process for high-density variants. The research progress on rare-earth ion-doped high-density scintillation glasses containing Ce³⁺, Tb³⁺, and Eu³⁺ ions was summarized, highlighting characteristics like density, light yield efficiency (LY), and decay time (τ). Some factors affecting these properties (i.e., substrate selection, energy transfer among ions, and reduction atmosphere) were also discussed.

Ce³⁺-doped glass exhibits a nanosecond-level decay time. At the present, the density of Ce³⁺ doped scintillation glass can reach up to 6 g/cm³, but the light yield of this component is significantly low. The luminescence efficiency of the glass can be improved via introducing Gd₂O₃, which enhances the energy transfer between Gd³⁺ and Ce³⁺, while increasing the glass density. Also, incorporating fluoride into the oxide glass can lower the phonon energy, decrease the chances of non-radiative transition, and enhance the light yield. Ce³⁺ is prone to oxidation to non-luminescent Ce⁴⁺ during high-temperature melting. This oxidation process can be effectively controlled by using a reducing atmosphere (CO, H₂) or adding reducing agents such as Si₃N₄, SiC, AlN, etc.

The density of Tb³⁺-doped scintillation glass can reach 7.15 g/cm³, with a decay time in the millisecond range. Enhancing the luminescence efficiency of Tb³⁺ is achievable through the energy transfer between Gd³⁺-Tb³⁺, Ce³⁺-Tb³⁺, Dy³⁺-Tb³⁺, while the introduction of fluoride and reducing agents can effectively boost the light yield. Eu³⁺-doped scintillation glass can achieve a density of 6.60 g/cm³, and the addition of Gd³⁺ and Tb³⁺ can effectively enhance the emission intensity of Eu³⁺ ions.

High density scintillating glasses doped with other rare-earth ions, such as Pr³⁺, Dy³⁺, Sm³⁺, and Er³⁺, can achieve a high density and exhibit distinct scintillation properties because of their unique energy level structures. However, the light yield of these glasses is too low to be used in practical applications.

A high-density and high-yield scintillation glass is needed. Despite the efforts of scientific research personnel, producing flashing glass with both a high density and a high yield remains a challenge. The main reason is that in high-density glass, although the introduction of heavy metal ions increases the density of the glass, it may lead to an increase in non-radiative transition channels, hindering the transfer of energy from the substrate to the luminescent center. Also, an increase in density can enhance the absorption and scattering of light, thereby reducing the effective amount of excitation light reaching the luminescent center. In addition, high density glass matrix may cause the local environment around rare-earth ions to become complex, affecting their energy level structure. To prepare scintillation glass that can meet both high-density and high light yield, several aspects can be considered. Firstly, the use of a reducing atmosphere such as H₂ or CO and the addition of reducing agents like carbon powder can effectively control the valence of luminescent ions (i.e., Ce³⁺, Tb³⁺) during the melting process, preventing the formation of high valence ions. Using fluoride instead of partial oxides can reduce the melting point of the material and improve the light output via altering the band gap structure of the glass. The smaller volume of fluoride ions simultaneously contributes to the increased density of the glass. The utilization of energy transfer between ions can significantly enhance the light yield of glass. Some impurities and defects in glass can be minimized via utilizing high-purity raw materials and optimizing process steps such as melting and annealing. This can enhance the transparency and uniformity of the matrix, leading to a reduction in non-radiative transitions and an improvement in luminescence performance. It is possible to induce the formation of microcrystalline structures in glass via controlling the cooling rate or subsequent heat treatment. This optimization can enhance the energy transfer path, minimize scattering losses, and ultimately improve the luminescence efficiency.

The Si₃N₄ ceramic is an ideal substrate material for high power device, and its high thermal conductivity is essential. However, there is still a significant gap between the actual and theoretical values of the thermal conductivity of Si₃N₄ ceramics due to the influence of lattice oxygen impurity and microstructure composition. The approaches to obtain Si₃N₄ ceramics with a high thermal conductivity are to decrease the lattice oxygen content, reduce the grain boundary phase and obtain the double peak structure composed of fine grains and a high aspect ratio β-Si₃N₄ grains through improving the purity of raw materials, selecting appropriate sintering aids and forming methods and optimizing the sintering system.

To obtain high thermal conductive Si₃N₄ ceramics, a research on the raw material powder with a high purity and a low oxygen content should be strengthened. The high quality equipment for Si₃N₄ ceramics production should be developed through controlling the sintering and processing technology. The practical application research should be promoted continuously and the completed application closed loop should be formed among powder production companies, product manufacturing enterprises and users. The timely feedback and response about the problems discovered during this process should be provided quickly to improve the performance of high thermal conductive Si₃N₄ ceramic products.

These researches involve 1) the introduction of organic dispersants to increase the dispersion of C and Al₂O₃ particles, and the formation of a uniform binding interface between particles, 2) the introduction of a binder (i.e., an organic carbon source) on the uniform coating of Al₂O₃ particles and pyrolysis to form a uniform combination of cracked carbon and Al₂O₃ particles, 3) the decomposition of organic matter, combustion synthesis and sol-gel and foam-injection coagulation to obtain precursors with a porous structure and a high specific surface area and the uniform mixing of aluminium and carbon sources in a micron scale, thus reducing the conditions for complete reaction and synthesizes high-quality nanopowders of AlN and 4) increasing the substitution (flow) efficiency of the reaction atmosphere, reducing the partial pressure of CO and increasing the partial pressure of N₂ to promote the positive synthesis reaction. According to the existing research, the future industrial development can be based on the combination of key process technology and key equipment. The continuous precursor manufacturing and continuous high-temperature synthesis equipment, especially continuous high-temperature synthesis equipment to eliminate differences in temperature zones and in airflow fields, need to be developed for the enhanced performance and efficiency of AlN powder synthesized by the CRN method.

This review summarizes recent research progress on the preparation of aerogels by chemical vapor deposition. The template-free chemical vapor deposition (CVD) method of the aerogel preparation can produce carbon material aerogels and multifunctional silicon carbide nanowire aerogels with superior performances for several hours. The template-assisted CVD method can customize the microstructure of the aerogel according to the template in order to achieve some special functions (i.e., superelasticity, negative Poisson's ratio, etc.). The chemical vapor deposition process for the preparation of aerogel can completely avoid the influence of surface tension during the drying process and the preparation process is more efficient due to the absence of liquid phase, compared with the traditional aerogel liquid-phase preparation process. The existing research in the related field is not sufficient due to the presence of some problems. Some possible aspects of subsequent research include:

1) To further simplify the process. The template-assisted chemical vapor deposition method in the chemical vapor deposition preparation process for aerogels usually requires the specialized preparation of templates with aerogel structures, which increases the complexity of the process and is not conducive to large-scale production. Industrially produced fumed silica or fumed alumina has a structure similar to that of an aerogel, and the use of fumed silica or fumed alumina as a template to grow the desired material on its nano-skeleton is expected to reduce the cost of preparation of partial oxide aerogels.

2) To develop a chemical vapor deposition preparation process for aerogels with milder conditions. At present, the reaction temperature is mostly more than 1 000 ℃ when using chemical vapor deposition to prepare aerogels. The reaction conditions become milder and the requirements of equipment reduce through the introduction of catalysts and other methods, which is an important research aspect for the future development of chemical vapor deposition process of aerogel.

3) To broaden the scope of application of the material type of aerogel chemical vapor deposition process. From the existing reports, the chemical vapor deposition process can be used for the preparation of some aerogel types (i.e., mainly carbon materials, carbides, borides and other aerogels). The reaction mechanism of the aerogel chemical vapor deposition preparation process needs to be further investigated. The chemical vapor deposition process can be used to prepare oxide aerogels (i.e., silica and alumina) via utilizing the advantages of the chemical vapor deposition process (i.e., short cycle time and no drying).

Moderate heat Portland cement (MHC) enhances early concrete strength, reduces hydration heat, and increases crack resistance, making it widely used for dam construction. Most dams located underwater therefore having to face to leaching issues. Prolonged concrete exposure to water leads to the dissolution of Ca(OH)₂ from hydration products, thus decreasing internal alkalinity so as to cause the breakdown of ettringite, C-(A)-S-H, and clinker. The phenomena result in deterioration of mechanical properties and impermeability. Current research on the leaching of cementitious material mainly focuses on ordinary Portland cement (OPC) and its blends with mineral admixtures, while little attention is paid to MHC and moderate heat Portland cement-fly ash (MHC-FA) systems. The latter two MHCs are commonly employed in dam construction. In this work, leaching tests were conducted on MHC pastes and MHC-FA pastes. Comparison with OPC pastes system was further carried out, focusing on the evolution of leaching depth, deterioration of compressive strength, and the impact on phase composition and C-(A)-S-H structure.

Specimens with a water-to-cement ratio of 0.5 were prepared from MHC pastes (M0), MHC with 30% FA pastes (M30), and OPC pastes (P0). The size of the formed specimens was 40 mm× 40 mm×40 mm. After curing in a standard concrete curing room for two years, five surfaces of the specimens were sealed with epoxy resin, leaving one surface exposed for accelerated leaching in a 5 mol/L NH₄Cl solution. The specimens were taken out from the solution after 7, 14, 28 d, and 56 d, following a test on compressive strength and leaching depth. The leaching depth is determined by spraying a 1% phenolphthalein solution in alcohol on the cut-open specimens. Samples were taken from areas 0-3 mm and 3-6 mm from the leaching exposure surface and pristine surface at 56 d. The thickness of 0-3 mm is labeled as the external layer of the leaching zone, and the thickness of 3-6 mm is labeled as the internal layer of the leaching zone. Characterization techniques included X-ray diffraction (XRD), Thermogravimetric Analysis (TGA), and Solid-State ²⁹Si Nuclear Magnetic Resonance (²⁹Si NMR) spectroscopy. In addition, the porosity of samples was evaluated using the hydrostatic weighing method.

The Leaching depth of M0 develops slower than OPC, while the M30 is the fastest. Initially, the compressive strength of M0 is higher than P0, but adding 30% fly ash to M0 lowers its compressive strength. Before leaching, the compressive strength of M0 is higher than that of P0, and the M30 has the lowest compressive strength. The compressive strength of the specimens decreases as the leaching time increases. Among them, the compressive strength of M0 is always the highest, and M30 is always the lowest. Compared with OPC, MHC has higher initial compressive strength and a lower strength decline rate, benefiting its long-term corrosion resistance. Although M30 has the lowest initial intensity, its intensity decreases more slowly. Additionally, before leaching, M30 has the highest porosity, P0 is in the middle, and M0 the lowest. Leaching increases the porosity of all samples evenly. Despite M30 has the highest initial porosity, the increasing rate of porosity is the lowest. Similar phenomena are observed for M0 and P0, while P0 shows the highest increasing rate of porosity. There is no significant difference in hydration products among the P0, M0 and M30 in the unleaching regions. The prominent XRD diffraction peaks of the sample before dissolution include Ca(OH)₂, sjoegrenite, ettringite, CaCO₃ and gypsum, and C₃S, C₂S and C₄AF. M30 also contains quartz peaks. The long-time curing leads to carbonation in the external layers. Thus, after leaching, a significant CaCO₃ peak appears. For the leached external layers of P0 and M0, the main phases include CaCO₃, sjoegrenite, ettringite and gypsum, as well as C₃S, C₂S and C₄AF. All ettringite dissolves in the leached external layers of M30, showing weakened peaks of C₃S and C₂S, yet ettringite remains in the internal layer. TGA reveals the Ca(OH)₂ contents of 26.0% for P0, 22.1% for M0, and 5.3% for M30, showing that P0 has the strongest leaching buffering effect. Based on the former result that M0 has the best leaching resistance, the experiments reveales no guaranteed relationship between strong buffering and optimal leaching resistance. This can be understood as that Ca(OH)₂ is the most soluble, and is therefore easier to create more pathways for leaching, leading to significant microstructural damage. The ²⁹Si NMR results reveal that a significant increase in MCL is observed in all the samples due to the decalcification of C-(A)-S-H. Both the internal and external layers of M30 exhibit the highest growth proportions, accompanied by the formation of Si-Al gel. Although the MCL of M0 exceedes that of P0 in both the internal and external layers after leaching, the growth of silicate chain is still lower than that of P0. However, M30 experiences the lowest compressive strength reduction rate, suggesting that microstructural changes from soluble substance leaching may show greater impact on declining mechanical properties than silica gel formation due to C-(A)-S-H decalcification. During the leaching, the stability of AlVI is greater than AlIV.

The main conclusions are summarized as follows: Compared to OPC pastes, MHC pastes experience a slower leaching depth development, leading to enhanced residual compressive strength, decreased pore degradation, mitigated declines in compressive strength and changes in C-(A)-S-H gel structure. Before leaching, OPC pastes has the highest Ca(OH)₂ content, MHC pastes the middle and MHC-FA pastes the lowest. Leaching leads to a significant increase in the MCL of C-(A)-S-H in the paste. After leaching, both the internal and external layers MHC-FA pastes exhibit the largest increase in MCL, accompanied by the production of silica-alumina gel. The average increase of silicate chain length is observed greater in C-(A)-S-H of OPC pastes than that of MHC pastes. During leaching, the stability of AlVI exceeds that of AlIV.

De-icing salt has been a common way to melt snow accumulated on concrete structures. But this often leads to the deterioration of concrete structures. Therefore, frost resistance durability of concrete subjected to de-icing salt is crucial. Conventionally, concrete is allowed to freely deform during single-side freezing and thawing (SSFT) tests. However, a concrete specimen is only equivalent to a unit of the concrete component in actual, and its deformation subjected to SSFT is restrained by the surrounding concrete, leading to alterations in salt frost damage. The differences in the evolution of concrete's frost damage between the laboratory and actual engineering are highlighted. The different internal damage of concrete subjected to freezing and thawing with pure water and de-icing salt is usually related to the water migration and phase transformation inside the concrete. However, the only difference between restrained concrete and unrestrained concrete is the restraint effect. The significant difference in internal damage of the two concretes subjected to SSFT cannot be explained by classical theory. Additionally, the influence of restraint effect on water migration in concrete also needs to be clarified. Thus, it is crucial to study the effect of restraint on the internal damage of concrete subjected to SSFT, which will provide guidance for the mix design of high frost resistant concrete in cold regions. In this paper, restraint ring and anchors were designed to confine the deformation of concrete to simulate the serving condition during SSFT tests. The relative dynamic elastic modulus and water absorption in volume of restrained and unrestrained concrete with a water cement ratio of 0.60 were tested after every 4 SSFT cycles. The temperature and humidity sensors were used to monitor the evolution of the relative humidity (U_IRH) inside the concrete during the SSFT test. The influence and mechanism of restraint effect on the internal damage of concrete subjected to SSFT were clarified from the perspective of moisture migration.

P·I 42.5 cement complying GB 8076—2024, natural river sand with a fineness modulus of 2.68, and graded stone by uniformly mixing the crushed stone with 5-10 mm and 10-20 mm in a mass ratio of 4∶6 were used to prepare fresh concrete with water to cement ratios of 0.60. The slump and air content of fresh concrete are 120 mm and 1.4%, respectively, tested according to GB/T 50080—2016. The 28 d cube compressive strength of the concrete specimen is 28.3 MPa, tested according to GB/T 50081—2019.

Two groups of moulds were used to cast concrete specimens. For unrestrained specimens, cylindrical plastic moulds with a diameter of 100 mm and a height of 70 mm were used. For restraint specimen, a restraint device was designed, which is used as the mould to cast a restrained concrete specimen meanwhile used as the restraining device to limit concrete deformation in SSFT tests. The device consists of a restraint ring, 24 restraint anchors, a counter-force ring, and a flange. All parts of the device are made of 304 stainless steel.

The curing and pre saturation of concrete specimens were carried out according to the salt freezing method in CIF test and GB/T 50082—2009. The CDF/CIF TESTER produced by Schleibinger Geräte was used to provide SSFT cycles. After every 4 SSFT cycles, the relative dynamic elastic modulus (r_n) and water absorption in volume (V_n) of the concrete specimens were tested. The evolution of internal relative humidity (U_IRH) of 28 SSFT cycles was monitored, and the pore size ratios of restrained and unrestrained concrete during the first cooling cycle were calculated.

The results show that after every one freeze-thaw cycle, the U_IRH of both restrained and unrestrained concrete increased, resulting in hysteresis of the U_IRH-temperature curves, which reflects the process of water migration inside concrete subjected to SSFT cycles. The hysteresis of the U_IRH-temperature curves of unrestrained concrete is more significant than that of restrained concrete. After 28 SSFT cycles, the U_IRH increment, r_n loss and V_n of restrained concrete are 30%, 32% and 28% less than those of unrestrained concrete, respectively.

The calculated results based on the U_IRH-temperature curves show that the radius of pores (r_pr) in restrained and unrestrained concrete increases with SSFT cycles. After 28 SSFT cycles, the r_pr of restrained concrete is 58% less than that of unrestrained concrete. A model for the internal water migration of concrete subjected to SSFT is established to analyze the influence of restraint on the water migration. The ratio of pore radius of the restrained and unrestrained concrete before freezing point during the cooling process in first SSFT cycle was calculated. The results indicate that under the influence of restraint, the water migration inside the concrete driven by SSFT is effectively weakened.

The main conclusions of this paper are summarized as following. The conventional SSFT test of unrestrained concrete overestimates the internal damage. The restraint effect weakens the water migration before the freezing point of concrete during the cooling process, which is the basis for reducing internal damage induced by SSFT cycles. Moreover, the restraint effect also reduces the pore size expansion of concrete subjected to SSFT cycles, which is manifested macroscopically as less loss of relative dynamic elastic modulus and less increase of water absorption in volume.

In cold regions, concrete is widely used as a material in hydraulic structures. Its durability is crucial for ensuring the stability of concrete structures over service life. As a porous medium, concrete undergoes freeze-thaw cycles where the water in the pores freezes and expands, causing water migration and pore pressure. This phenomenon results in internal structural damage and thus degradation of concrete performance. Therefore, it is important to analyze the relationship between temperature, humidity, and pore pressure of concrete during the freeze-thaw cycles. Scholars have proposed various mechanism theories, and a general knowledge is given that the thermo-hydro coupled behavior of concrete imposes the major machinal loads inducing the freeze-thaw damage. But the dominant mechanism is still not clear. The hydrostatic pressure theory and the crystallization pressure theory explain the freeze-thaw damage of concrete from the perspective of the effective stress principle of fluids. The theory of poroelasticity explains the freeze-thaw damage of concrete from the perspective of internal pressure and deformation generated by the pores inside the concrete during the formation and thawing of ice crystals. Hence, it is believed that an additional attention should be paid to the combined mechanical effect of solid phase expansion and the load contribution of the fluid phase. By conducting a detailed analysis of pore pressure and ice expansion deformation during the freeze-thaw process, the dominant mechanism of freeze-thaw damage can be further analyzed.

The spatial discretization method of flow lattice network (FLN) considering the concrete mesoscopic heterogeneity is adopted to describe the heat transfer and mass transport behavior. Within the framework of FLN, a thermo-hydro coupled model for concrete during the freeze-thaw cycles is proposed and its mechanical mechanism are discussed based on the poroelasticity and effective stress principle. The proposed model implemented into UEL subroutine and solved by standard/implicit solver in ABAQUS. The proposed model is used to simulate the freeze-thaw process of concrete under uniform and non-uniform temperature change.

Firstly, the numerical simulations of concrete under freeze-thaw conditions are conducted based on existing concrete ice content and the corresponding deformation experiments. As the temperature decreases, the relative ice content predicted by the model begins to increase at 0 ℃, and the ice crystallization first increases and then decreases. At -55 ℃, the relative ice content reaches 0.42. Subsequently, due to the consideration of freeze-thaw hysteresis, the relative ice content does not decrease with the increase in temperature, but starts to gradually decrease only when the temperature rises to the melting temperature (-15 ℃). The ice completely thawed at 0 ℃, which is in good agreement with the experimental results. In the simulation with constant and temperature-dependent thermal expansion coefficient, the total deformation and deformation trend calculated by the freezing stage model are in good agreement with the freeze-thaw deformation test data. However, during the thawing stage, due to the damage (cracking strain) in the experimental, the model calculation results do not match the experimental results. The strain of ε_Σm and ε_Δρ show expansion, and ε_th shows contraction deformation. ε_Δρ dominates the over deformation. Additionally, through the calculation of relative errors in deformation, it can be observed that the calculation results with constant a_l and a_c keep relatively consistent with the experimental results. During the cycling process, the peak pore pressure of 5 MPa during the freezing stage of concrete appears at 4.4 hours, and the peak pore suction pressure of 9 MPa during the thawing stage appears at 21.5 h. In the evolution of frost heave strain, the peak strain appears at 13 h.

In the simulation of non-uniform temperature distribution, the variation depth of relative ice content gradually expands towards the temperature invariant region as the number of freeze-thaw cycles increases. The peak value also increases with the pore pressure expands towards the rear side. In addition, the peak deformation occurs alternately with the peak pore pressure. During the process of water freezing and thawing, the mechanical loads of the ice-water mixture in pores on the solid skeleton presents two stages. The first stage is dominated by fluid pressure and is computed by the principle of generalized effective stress. The second stage is ice crystallization induced expansion computed by the poroelasticity.

A thermo-hydro coupled model for concrete during the freeze-thaw cycles is proposed within the framework of FLN. The numerical results of relative ice content and the corresponding deformation are in good agreement with the experimental results, which verifies the accuracy of the proposed model. The consideration of the temperature-dependent mathematical pore distribution in the calculation of ice content can simulate the hysteresis of ice content during freezing and thawing. It is numerically found that the pore pressure and frost heave strain play the rotatory dominant role during the freeze-thaw cycle. One stage is dominated by fluid pressure and another stage is dominated by ice crystallization induced expansion computed. The effective stress caused by the pore fluid and the calculation of frost heave strain are comprehensively considered.

The wide range and strip-like distribution of high-speed rail crossings, as well as the exposed service characteristics, result in ballastless track concrete not only directly bearing train loads, but also facing freeze-thaw damage in severely cold or cold areas such as high-latitude cold regions and high-altitude mountainous areas. In the service process, the coupling beween environmental effects and train loads has significantly affected the ballastless track concrete at the micro- and macro-levels. However, current standards do not consider the effects of high-speed train fatigue loads and freeze-thaw cycles on the service performance of concrete. During the operation of high-speed trains, the flexural fatigue load frequency that causes damage to the concrete exceeds 20 Hz, which is a significant difference from the test system commonly used in current fatigue tests.

To reflect the actual service conditions of ballastless track more realistically in cold regions, a coupled test system of 'freeze-thaw cycles and 20 Hz fatigue loads' was designed. The effects of freeze-thaw cycles on the performance of concrete were analyzed from the perspectives of mechanical properties, fatigue life, and stiffness. The freeze-thaw cycles were divided into freezing and thawing stages, and their damage characteristics and energy transfer laws were analyzed. A life prediction model considering the coupling of freeze-thaw and fatigue was established from a probabilistic perspective to predict the service life of ballastless track manufactured sand concrete in cold regions.

The study used C60 manufactured sand concrete as the research object, with specimen dimensions of 100 mm (width) ×100 mm (height) × 400 mm (length). A SUNS-890 electro-hydraulic servo static and dynamic universal testing machine was used for fatigue tests. The stress level (maximum stress f_max/flexural strength f_f) was fixed at 0.6, stress ratio (minimum stress f_min/peak stress f_max) was 0.1, loading frequency was 20 Hz, and displacement data was simultaneously recorded during loading. After being saturated, the concrete underwent a series of freeze-thaw cycles, including 0, 150, 300, 450 cycles, and 600 cycles, before undergoing fatigue testing. To analyze the evolutionary of concrete performance during different stages, the freeze-thaw cycles were divided into two stages: freezing (-20 ℃) and thawing (0 ℃ and 20 ℃). After saturation, the specimens were placed at -20, 0 ℃, and 20 ℃ to conduct bending fatigue tests. Simultaneously, in order to analyze the temperature rise caused by fatigue loads at a constant temperature, a temperature measurement point was taken every 25 mm along the side of the concrete during the loading process, and a thermometer was used to measure its temperature.

After 300 freeze-thaw cycles, the ballastless track concrete exhibits accelerated brittle failure characteristics. Both flexural strength and fatigue performance show a trend of accelerated decline with freeze-thaw cycles exceeding 300 times, resulting in a 46.3% decrease in fatigue life and a 12.9% increase in stiffness degradation after 600 freeze-thaw cycles. The internal and surface defects caused by freeze-thaw cycles serve as initiation points for fatigue damage propagation, with the increased actual stress level making the concrete more prone to fatigue fracture. Defect connectivity induced by the cycles becomes the main factor contributing to the deterioration of concrete performance. From an energy perspective, the loading process involves both potential and internal energy, with concrete dissipating energy mainly through deformation energy generated by deformation. The decrease in maximum strain and increase in residual strain caused by freeze-thaw cycles result in reduced deformability of the concrete, thereby negatively affecting its fatigue performance. Ice within the concrete pores during the freezing stage has a reinforcing effect, leading to an improvement in fatigue performance. Additionally, ice helps alleviate internal friction and accelerates internal energy dissipation, thereby mitigating fatigue damage during the freezing stage. By introducing a fatigue damage factor, a life analysis model for concrete under freeze-thaw + fatigue service conditions are established. Predictions of concrete life under different environments are made based on freeze-thaw environmental criteria, revealing that C60 manufactured sand concrete can meet the freeze-thaw design requirements of cold region at level D1 and level D2, but considerations must be made for the intensified damage from fatigue loads.

A test regime of 'freeze-thaw cycle + high-frequency fatigue' was established to investigate the damage mechanism of concrete under the coupled action of freeze-thaw and fatigue from the perspectives of damage progression and energy transfer. It was found that high-speed train fatigue loads exacerbated the freeze-thaw damage of ballast-free track concrete, with defect connectivity caused by freeze-thaw cycles being the main reason for the decline in concrete fatigue performance. Freeze-thaw cycles lead to a decrease in maximum strain and an increase in residual strain, resulting in reduced deformability of the concrete. Compared to fatigue performance at room temperature, ice within the concrete pores during the freezing stage has a reinforcing effect, and ice in the frozen state can accelerate internal energy dissipation, thereby alleviating fatigue damage. A life analysis model for concrete under 'freeze-thaw + fatigue' service conditions, considering a fatigue damage factor, was successfully established to predict the service life of ballast-free track C60 manufactured sand concrete.