Amid the growing emphasis on energy conservation and environmental protection, designing phosphors with multifunctional applications has become an urgent challenge. The traditional approach of combining blue light-emitting diode (LED) chips with yellow phosphors to produce white light suffers from drawbacks such as low color rendering index (CRI) and high color temperature. To address the pressing demands of lighting and display technologies, developing high-efficiency red phosphors that can be excited by near-ultraviolet (near-UV) light is crucial for applications in white LEDs and flexible display technologies.

Building upon the garnet-like structure of Ca₃Zn₃(TeO₆)₂, a series of CaNaLaZn₃(TeO₆)₂:xEu³⁺ (CNLZT:xEu³⁺) red phosphors, capable of being effectively excited by near-UV light, were designed and synthesized using a chemical unit co-substitution strategy. In this approach, the [Na⁺-La³⁺] ion pair replaces the [Ca²⁺-Ca²⁺] ions. This co-substitution significantly enhances the luminescence properties of the CNLZT:0.15Eu³⁺ sample. White LED devices fabricated with this material exhibit excellent luminous performance and stability. Moreover, the CNLZT:0.15Eu³⁺/polydimethylsiloxane (PDMS) flexible film demonstrates promising luminance, making it well-suited for applications in flexible display technologies.

A series of CNLZT:xEu³⁺ red phosphors were synthesized using a high-temperature solid-state reaction method. Raw materials were accurately weighed according to stoichiometric ratios and thoroughly ground in an agate mortar for 25 min. The resulting mixture was loaded into alumina crucibles and sintered for 10 h in an air atmosphere at 1323 K. After cooling, the samples were reground to obtain fine powders.

The prepared CNLZT:0.15Eu³⁺ red phosphor was then mixed with BaMgAl₁₀O₁₇:Eu²⁺ (BAM:Eu²⁺, blue) and (Ba,Sr)₂SiO₄:Eu²⁺(BSS:Eu²⁺, green) in a specific mass ratio. This mixture, along with ZWL8820 organic silicone gel, was applied to a 395 nm near-UV chip and encapsulated to fabricate a white LED device. Additionally, a CNLZT:0.15Eu³⁺/PDMS flexible film was prepared by combining the CNLZT:0.15Eu³⁺ phosphor with PDMS in a predetermined mass ratio and cutting it to the desired size for subsequent testing.

X-ray diffraction (XRD) analysis and Rietveld structural refinements reveal that substituting [Ca²⁺-Ca²⁺] ion pairs with [Na⁺-La³⁺] does not alter the matrix structure. The Eu³⁺ replace La³⁺ in the matrix, causing lattice contraction and a slight shift in XRD diffraction peaks toward higher angles. The main absorption peak at 394 nm aligns well with the near-UV LED chip.

When x=0.15, the luminescence intensity of the optimal sample CNLZT:0.15Eu³⁺ reaches its maximum, 1.135 times higher than the commercial red phosphor Y₂O₃:Eu³⁺ and 1.908 times greater than the pre-substitution Ca₃Zn₃(TeO₆)₂:0.12Eu³⁺. The color purity is 99.78%, and the CIE color coordinates are (0.649 2, 0.350 4), close to the ideal red light point (0.670, 0.330), demonstrating excellent red luminescence. This confirms that the chemical unit co-substitution strategy effectively enhances luminescence properties.

At elevated temperatures, the CNLZT:0.15Eu³⁺ sample exhibits good color stability. The packaged white LED device has a correlated color temperature (CCT) of 6414 K and a CRI (R_a) of 87.3. Its CIE coordinates (0.314 9, 0.325 0) lie in the white light region. Even after 100 minutes of continuous operation at 50 mA, the white LED maintains stable performance and luminescence. Furthermore, the CNLZT:0.15Eu³⁺/PDMS film remains intact and efficiently luminescent after bending tests, demonstrating potential for flexible display applications.

A series of CNLZT:xEu³⁺ red phosphors, excitable by near-UV light, were synthesized using a high-temperature solid-state reaction method. By employing the chemical unit co-substitution strategy of [Na⁺-La³⁺] replacing [Ca²⁺-Ca²⁺], the optimized doping concentration of Eu³⁺ in the CNLZT matrix was effectively elevated to 15% (in mole). The luminous intensity of the CNLZT:0.15Eu³⁺ sample is 1.908 times that of the pre-substitution Ca₃Zn₃(TeO₆)₂:0.12Eu³⁺ sample and 1.135 times that of the commercial Y₂O₃:Eu³⁺.

The color purity of the CNLZT:0.15Eu³⁺ sample is 99.78%, and its color coordinates (0.649 2, 0.350 4) are close to the ideal red point (0.670, 0.330). The chemical unit co-substitution strategy effectively regulates the luminescent properties. Temperature-dependent fluorescence spectroscopy confirms that CNLZT:0.15Eu³⁺ possesses excellent color stability, with a chromaticity shift of only 0.008 52 at 420 K.

The packaged white LED achieves CIE coordinates of (0.314 9, 0.325 0) with a CRI (R_a) of 87.3, demonstrating precise color rendering and stable performance after 100 min of continuous operation at 50 mA. Moreover, the CNLZT:0.15Eu³⁺/PDMS flexible film shows promising luminance, making it suitable for applications in flexible display technologies.

The durability of cement-based materials is significantly influenced by the transport processes of water, gases such as CO₂, and ions like Cl^-. Water transport, particularly capillary water absorption, is crucial for durability performance. In most engineering practices, cement-based materials are non-saturated, making capillary water absorption the primary mode of water transport. This process is driven by capillary pressure arising from the meniscus within pores and is governed by the extended Darcy's law or the Richard's equation. In one-dimensional capillary absorption scenarios, theoretically there exists a linear relationship between the cumulative absorbed volume per unit area and the square root of absorption time, which is known as the square root of time linear law. The slope obtained from linear fitting is defined as the capillary sorptivity, which serves as a quantitative indicator of durability. However, long-term capillary water absorption often deviates from this linear law due to specific physicochemical interactions between water and cement hydrates, especially calcium silicate hydrate (C-S-H) gel. The widespread use of supplementary cementitious materials such as fly ash, slag, and silica fume alters the composition and micro-structure of C-S-H gel through secondary hydration reactions, significantly impacting the properties of cement-based materials. Although the capillary sorptivity has been well recognized as a durability indicator, the anomalous phenomena observed during long-term capillary water absorption are frequently overlooked, leading to incomplete research and inconclusive findings. This study focuses on the long-term capillary absorption processes of water and IPA into cement mortars with representative supplementary cementitious materials, including fly ash, slag, and silica fume. The aim is to elucidate the influence mechanisms of supplementary cementitious materials on the capillary absorption properties of cement mortars.

The experimental section focused on investigating the impact of supplementary cementitious materials on the capillary absorption behavior of cement mortars. Samples were prepared with sands, white cement and high contents of fly ash, slag, and silica fume. These specimens were subjected to both standard curing at 20 ℃ and accelerated curing at 60 ℃ to assess the influence of curing temperature. Testing methods encompassed capillary absorption tests using IPA and water, along with low-field magnetic resonance relaxation technique to analyze pore structure evolution. The capillary sorptivity were determined by fitting the experimental data to a square root of time linear law, and deviations from this law were analyzed to characterize water sensitivity. Data analysis involved examining changes in capillary sorptivity, deviation times, and pore structure characteristics to elucidate the mechanisms by which supplementary cementitious materials affect the capillary absorption properties of cement mortar.

The incorporation of fly ash, slag, and silica fume significantly alters the capillary sorptivity of both IPA and water into cement mortar. Specifically, all three supplementary cementitious materials reduce the capillary sorptivity of IPA, with fly ash exhibiting the most pronounced effect. During the initial stage of capillary absorption of water, slag increases the sorptivity, while silica fume and fly ash decrease it. In the later stage, high silica fume content notably enhances the sorptivity, whereas high slag and fly ash contents decrease it. The water sensitivity of cement-based materials is also affected. High slag and fly ash increases the water sensitivity of cement mortars, leading to earlier deviation from the initial linear law and higher degrees of deviation in later stages. These findings indicate that supplementary cementitious materials modify the composition and microstructure of C-S-H gel, thereby affecting the water sensitivity and durability of cement-based materials.

The addition of three supplementary cementitious materials will reduce the capillary sorptivity of isopropanol, and the reduction effect of fly ash is most significant. In the initial stage of capillary absorption of water, slag significantly increases the initial sorptivity of mortar, while silica fume and fly ash show inhibiting effects. In the later stage, silica fume significantly increases the secondary sorpvitity of water into mortar, while slag and fly ash both decrease their secondary sorptivity. The addition of slag and fly ash will significantly increase the difference between the initial and secondary capillary sorptivity. High content of silica fume can reduce the water sensitivity of cement-based materials, while high content of slag and fly ash can improve the water sensitivity, so as to advance the deviation of capillary absorption of water from the square root of time linear law and to increase the degree of deviation. The improvement of water sensitivity is the most significant for the mortar with high slag substitution. Curing at high temperature slightly reduces the water sensitivity of mortar with high slag content, but almost has no effect on the water sensitivity of mortars with high silica fume and fly ash contents.

Ceria (CeO₂) is an important rare earth metal oxide abundant in nature. Due to the easy conversion of the oxidation valence between cerium ions, CeO₂ has a high oxygen storage capacity. The band gap value of ceria is similar to that of other traditional semiconductor photocatalyst of TiO₂, so it has the potential to be a suitable and efficient photocatalyst. Specifically, ion doping can introduce new energy levels into the band structure of ceria, reduce its band gap width, and enable more visible light to be absorbed, thereby ultimately improving the photocatalytic efficiency. Secondly, when other cations enter the CeO₂ lattice, they form oxygen vacancies (oxygen defects) under structural deformation and charge compensation mechanisms. The oxygen vacancy formed in this process can act as an effective trapping point for photogenerated electrons, effectively reduce the recombination rate of electrons and holes, and significantly improve the photocatalytic efficiency of the material. In addition, ion doping can also change the electronic structure of cerium oxide, making it have better conductivity and charge transport performance, which is conducive to the separation and transport of photogenic carriers. Generally speaking, it is very difficult to introduce ions into the CeO₂ lattice. The author focuses on designing the synthesis system of CeO₂ photocatalyst doped with Sm ions and systematically studying and summarizing the reaction ratio relationship of raw materials, so as to elucidate the influence of ion doping on photocatalytic performance from the perspective of photocatalytic mechanism.

Firstly, 5.88 g C₆H₅Na₃O₇·2H₂O and 60 mL deionized water were stirred magnetically together at room temperature for 30 min. Then, 2.40 g urea was added and mixed for 30 min continually (solution 1). At the same time, 1.00 g Ce(NO₃)₃·6H₂O and different molar ratios of Sm(NO₃)₃·6H₂O were dissolved in 20 mL deionized water and stirred for 30 min, respectively (solution 2). Next, solution 2 was added into solution 1 and continue to stir for 30 min until the solution shows a light-yellow color. Then, the mixed reaction liquid was transferred to a 100 mL hydrothermal reactor, and cooled naturally after the hydrothermal reaction at 120 ℃ for 36 h. After used deionized water and ethanol to clean at least three times, and finally dried for 12 h at 70 ℃ and calcined at 500 ℃ for 4 h to obtain the target product. At the same time, CeO₂ product without adding Sm(NO₃)₃·6H₂O precursor were used as a comparison sample.

The influence of different doping concentrations on the photocatalytic performance of ceria was studied in detail, and the optimal doping ratio of Ce and Sm was determined to be 1.00∶0.15 (0.15-SC). SEM images indicated that Sm-doped material still maintained a broom-like morphology, but the surface smoothness, average diameter and length of the nanorods showed slight changes. From XRD analysis, all the samples exhibited sharp diffraction peaks, indicating high crystallinity, and no impurity peaks are detected, meaning that the doping process does not affect the crystal purity of the final product. UV-Vis absorption spectrum shows that the bandgap value (E_g) of these Sm-doped samples is significantly reduced compared to that of pure CeO₂, and the E_g of the 0.15-SC sample is the smallest at only 2.88 eV. The factors that lead to the reduction of the band gap can be attributed to two points: First, the nanoscale CeO₂ composed grains can induce the quantum limiting effect, which leads to the spectral redshift; Second, the decrease of particle size will increase the Ce³⁺ content at the edge of CeO₂. With the decrease of Ce³⁺ content, the UV-Vis spectra will also show a redshift phenomenon.

Compared with other Sm-CeO₂ samples, pure phase CeO₂ showed the strongest characteristic peak strength and the lowest peak strength when the doping ratio was 1.00∶0.15, indicating that Sm-CeO₂ samples had the best effect on promoting the separation of photogenerated holes and electron pairs and photocatalytic performance at this ratio. The environmental pollutant bisphenol A (BPA) was used as a photocatalytic degradation model, and it was found that the efficiency of all the Sm-doped CeO₂ samples was improved, and the degradation efficiency of the 0.15-SC sample was the highest, reaching 99.83%, which was 3.52-times superior than that of the pure CeO₂ (28.31%). The capture experiment of the active substances can be judged that both hydroxyl radicals (·OH) and holes (h⁺) played very important roles in the photocatalytic degradation of BPA on the 0.15-SC sample. And after four performance cycling experiments, the degradation rate of the 0.15-SC sample was still maintained 89.53%, indicating that it has good recyclability.

In summary, Sm-doped ceria material exhibits high photocatalytic purification activity for environmental pollutants, fully proving that ion's doping can act as an effective method to efficiently improve photocatalytic performance, and can further provide effective ideas for designing ion-doping systems for similar metal oxide materials in the future.

Carbon emissions in the construction industry are increasing now. Against the backdrop of environmental factors and low-carbon economy, the research on reducing carbon emission in cement industry has garnered widespread attention. During the cement production procedure, the decomposition of carbonates during the calcination of raw materials is the main source of CO₂ emissions. Therefore, reducing the calcium content in cement clinker is expexted to significantly reduce the carbon emissions, thereby achieving energy conservation and emission reduction in the industry. Unlike the C₃S mineral in ordinary Portland cement clinker, CS, C₃S₂, and γ-C₂S are three types of low calcium silicate minerals, but they are all non hydraulic minerals that cannot obtain a certain mechanical strength through traditional cement hydration. Former research revealed that the carbonation activity is higher than the hydration activity of above three low calcium minerals, and cement products can be prepared through carbonation curing.

To solve the problems of high energy consumption and high CO₂ emissions in the production of ordinary Portland cement, low calcium silicate minerals (CS, C₂S, and C₃S₂) with carbonation activity are used to replace C₃S minerals. The strategy could not only reduce carbon emissions in cement production, but also enhance utilization of carbon. Research has found that incorporating β-C₂S into the low calcium system of γ-C₂S can significantly enhance the mechanical properties of the system. In most reported research work, the single minerals are often blended after firing, and then the carbonization properties of the mixed minerals are studied. How to burn the C₂S minerals with the optimal crystal ratio in one step and to achieve the symbiosis of γ-C₂S and β-C₂S in the system remains a challenge. This work focuses on C₂S minerals. By changing the calcination temperature, C₂S minerals with different crystal ratios are treated in one step, and the relative content of γ-C₂S and β-C₂S in low calcium cement clinker is regulated. The symbiotic mechanism of the two and the influence of different mineral contents of β-C₂S and γ-C₂S on the carbonization performance of the material are studied. The carbonization products and mechanisms are explored, providing a theoretical basis for the application of low calcium fixed carbon cement in carbonization and new ideas for the carbonization research of C₂S minerals. which has certain guiding significance for energy conservation and emission reduction in the cement industry.

This work used limestone and sandstone to obtain low calcium cement clinker with different proportions of γ-C₂S and β-C₂S by adjusting the calcination temperature. The carbonization and hardening mechanism of low calcium cement clinker was analyzed using thermogravimetric analysis, X-ray diffraction, pH value, conductivity and other testing methods. The results showed that the molar ratio of γ-C₂S/β-C₂S was between 0.19 and 2.63. As the ratio of γ-C₂S/β-C₂S decreased, the compressive strength of low calcium cement clinker after 24 hours of carbonization gradually increased. When the calcination temperature was 1340 ℃ and the ratio of γ-C₂S/β-C₂S was 0.19, the compressive strength of the system reached 163.32 MPa. The change in CO₂ absorption of the sample is related to the carbonation activity of calcium silicate. The early carbonization reaction of the experimental group, mainly composed of γ-C₂S minerals, is severe. There is a phenomenon of incomplete reaction in the test block. The experimental group mainly composed of β-C₂S minerals has a longer duration of exothermic reaction, which is more conducive to the progress of carbonization reaction. Among them, the heat released by γ-C₂S during the carbonization process promotes the easier dissolution of Ca²⁺ from β-C₂S, facilitating the carbonization reaction of β-C₂S. The addition of a small amount of γ-C₂S (i.e. 8.13%) has a positive effect on the carbonization degree and strength of β-C₂S. In addition, the dissolution of β-C₂S leads to a gradual increase in the pH value of the suspension, changing the liquid phase environment and accelerating the dissolution rate of γ-C₂S in water under alkaline conditions. The excellent compressive strength of the sample is attributed to its high degree of carbonization and dense microstructure. The higher the content of β-C₂S minerals (γ-C₂S/β-C₂S ratio=0.19), the better crystallized calcite type calcium carbonate can be observed after carbonization, presenting a stacked and dense morphology. The samples mainly composed of γ-C₂S minerals have weak particle bonding and loose structure after carbonization, showing amorphous calcium carbonate with many pores.

Therefore, it can be concluded that the synergistic carbonization of γ-C₂S and β-C₂S mainly exists in two processes: the dissolution process before CO₂ is introduced, and the carbonization reaction process after CO₂ is introduced. Due to the action of water, β-C₂S preferentially dissolves Ca²⁺, changing the liquid-phase environment around γ-C₂S particles and promoting Ca²⁺ dissolution. After the introduction of CO₂, the carbonization reaction of γ-C₂S becomes more intense, releasing a large amount of heat that accelerates the dissolution of β-C₂S minerals and promotes their carbonization reaction. The final carbonized product is mainly composed of stacked calcite, with a small amount of high polymer silica gel interspersed to bond the calcite. Therefore, the high content of β-C₂S carbonized product shows a more tightly aggregation of particles and thus better performance.

Cement-based materials are widely used because of their relatively low cost and strong mechanical properties. However, during hydration hardening and service, such materials are often subjected to multiple effects such as plastic shrinkage, self-shrinkage, drying shrinkage, freezing and thawing cycles, and external loading, which in turn induces the formation of internal microcracks. The emergence and development of internal microcracks in cement-based materials can easily become a penetration channel for erosion media such as sulfate and chloride salts, which adversely affects the stability, durability and safety of the structure. Therefore, the characterization and evaluation of internal cracks in cement-based materials have received extensive attention from scholars at home and abroad. Traditional methods still face certain challenges in measuring the width of internal microcracks in samples quickly and accurately. This paper proposes a micro-crack testing method based on the directional transport of the inert gas argon, which can rapidly and accurately determine the micro-crack width by analyzing the diffusion behavior of the gas in the cracks and the cement matrix through the comprehensive consideration of the sample size and the pressure gradient at both ends.

The cement paste cracking samples with specific crack widths were preset by the insert method, and the preset crack widths were 0.03, 0.05, 0.10, 0.15 mm and 0.20 mm. A cylinder mold with a diameter of Φ50 mm×100 mm was adopted. Through the insert method, steel sheets with different thicknesses were placed in the center of the mold in advance to control the formation of penetrating cracks with different standard widths.

The gas transport system was mainly composed of sample chamber, gas transport control system and computer data acquisition system. During the test, the sample needed to be placed on a closed sample table and connected to the upper and lower outlets through a thermoplastic tube. The sample room maintained an oil bath environment to ensure sealing. The gas mass flowmeter was adjusted to the required flow rate, so that the gas passed through the device and recorded the pressure gradient changes at different gas flow. The crack width was calculated by Darcy's law.

When the crack width is in the range of 0.03-0.15 mm, the gas transport method can measure the width value more smoothly. When the crack width is 0.1 mm, the gas transport test results are closest to the actual crack width, and the error is less than 5%. When the crack width reaches 0.15 mm, the gas transport test can still maintain a high measurement accuracy, and the error is less than 10%. When the crack width reaches 0.2 mm, the error is more than 110% because the pressure gradient at both ends is not obvious. When the gas flow is 10-50 mL/min, with the increase of crack size, the increase of gas flow helps to reduce the error. Combined with the crack width and gas flow, it can be concluded that when the crack width is about 0.03 mm, the gas flow of 50 mL/min should be adopted; when the crack width is about 0.05 mm, the gas flow of 30 mL/min should be adopted. When the crack width is 0.10-0.15 mm, the gas flow of 20 mL/min should be adopted. In addition, there is a certain complementarity between the gas transport method and the optical microscope test in measuring the crack width. The optical microscope is suitable for the measurement of the width value of the larger crack surface, and the gas transport test can obtain the statistical average value of the penetrating crack inside the sample.

In the self-healing test, with the increase of gas flow (≥20 mL/min), the equivalent crack width obtained by the test is basically in a stable state. This is because the high gas flow is more conducive to the passage of argon through the gap between SAP hydrogel particles and the formation of channels, thereby further increasing the gas fluidity in the sample and improving the gas transport effect. By comparing the crack width before and after self-healing, it can be found that the healing effect of SAP hydrogel significantly reduces the crack equivalent width, and the healing rate calculated according to the crack equivalent width.

At the gas flow of 20-50 mL/min, the error is less than 10%, and the test error gradually decreases with the increase of crack width. The gas transport test can accurately measure the crack volume equivalent width of the crack in a short time (400 s), and has good rapidity and reliability. When the crack width reaches 0.2 mm, the detection accuracy is significantly reduced because the pressure difference gradient at both ends is not obvious, and the error is more than 110%. When the crack width is 0.03-0.15 mm, the fluctuation range of the test results increases with the increase of the crack size, and the appropriate increase of the gas flow is helpful to reduce the test error. Therefore, the gas transport method can better quantify the microcracks within 0.15 mm. The suitable gas flow range for crack self-healing in gas transport test is recommended to be 20-50 mL/min, and the average self-healing rate of 24 h is 39.1%, indicating that SAP hydrogel can effectively seal the microcracks inside cement-based materials in a short time, reduce the equivalent crack width, and play a good self-healing sealing effect.

Magnesium potassium phosphate (MKP) cements are normally composed of dead-burnt magnesia and monopotassium phosphate (KH₂PO₄), which react in water and yield the main hydrate of K-struvite (MgKPO₄·6H₂O). MKP cements have been frequently used as rapid repair materials in the field of civil engineering and as solidification/stabilization agents in the field of waste managements, due to the unique properties, such as fast setting, high early strength, low drying shrinkage, strong bonding to old Portland cement concrete substrates, low pH value, and so on. Boron compounds are the de facto standard retarders of MKP cements. Normally the increase of boron compound dosage could prolong MKP cement setting, but result in slower strength development, especially at early ages. Therefore, in order to obtain sufficient setting time and early strength at the same time, research attempts on seeking new retarders have been carrying out in recent years. Aluminum salts, Al(NO₃)₃ and Al₂(SO₄)₃, were among those reported chemicals, showing promising retardation effect, especially at low M/P molar ratios. However, influence of aluminum salts on properties and hydration of MKP cements are yet to be explored, as several important aspects remain unclear, such as their impact differences on MKP cement early and late strengths, roles of Al³⁺ and of different anions on governing the hydration. Therefore, in this work influence of three common aluminum salts, AlCl₃, Al(NO₃)₃ and Al₂(SO₄)₃, on properties and hydration of MKP cements at magnesium-to-phosphate (M/P) molar ratio of 4 and water-to-cement (w/c) ratios of 0.27 and 5.00 was studied.

The raw materials for the preparation of MKP cements were industrial-grade dead-burnt magnesia and KH₂PO₄. The chemicals in analytical grade were borax (Na₂B₄O₇·10H₂O), Al₂(SO₄)₃·16H₂O, Al(NO₃)₃·9H₂O, and AlCl₃·6H₂O. The MKP cement pastes and suspensions were prepared at same M/P molar ratio of 4, and at w/c ratios of 0.27 and 5.00. The aluminum salts were dissolved in the portioned mixing water before use, and the Al concentrations for the pastes and the suspensions were 600 mmol/L and 32.4 mmol/L, respectively. For comparison purpose, borax was used as reference, of which the dosages were 5.15% and 6.08%, by weight of cement, equal to the dosages of Al₂(SO₄)₃·16H₂O and Al(NO₃)₃·9H₂O, respectively. Vicat needle test was used to determine the paste final setting time. Flexural and compressive strengths of the pastes with the dimensions of 25 mm × 25 mm ×125 mm were measured after 1, 3 h, 1, 7, 28, 90 d and 400 d. Additional strength measurements were carried out after 6 h and 12 h for the paste with Al(NO₃)₃, due to the low strengths within the first 3 h. Hydration heat of the pastes up to 24 h was measured using an isothermal calorimeter (TAM) at 20 ℃, using internal mixing method. Moreover, hydrate assemblages of the pastes after 28 d and 400 d were analyzed using XRD, TGA and FTIR methods. To track the aqueous evolution, pH and electrical conductivity of the MKP cement suspensions were continuously monitored up to 24 h at 25 ℃. Also the aqueous compositions were determined using ion chromatography and inductively coupled plasma after 15, 33 min, 3 h, 1 d and 7 d.

Compared to the reference paste, the additions of Al₂(SO₄)₃, Al(NO₃)₃ and AlCl₃ could extend the setting time by around 2.5, 3.1 times and 5.1 times, respectively, corresponding to around 25%, 45% and 80% increases as compared to borax at 6.08%. It indicates better retardation effect of the aluminum salts and the anions taking effect in the order of Cl^-, NO₃^-, and SO₄^2-. In contrast to Al(NO₃)₃ and Al₂(SO₄)₃, AlCl₃ leads to clear sample expansion and cracks, suggesting its unsuitable use in MKP cement-based materials. Moreover, compared to the reference paste, both Al(NO₃)₃ and Al₂(SO₄)₃ lower the early strengths. The strength weakening effect of Al₂(SO₄)₃ mainly occur within the first 3 h, and of Al(NO₃)₃ is more pronounced and lasts longer up to around 12 h. However, thereafter, the strengths of the pastes with Al(NO₃)₃ and Al₂(SO₄)₃ gain fast, leading consequently to similar or even slightly higher strengths at late ages compared to the reference paste.

Al(NO₃)₃ shows better effects on slowing down hydration heat release and on decreasing cumulative heat up to 24 h, compared to Al₂(SO₄)₃, in agreement with the setting time results. Both Al(NO₃)₃ and Al₂(SO₄)₃ do not change the main hydrate, which remains K-struvite, but slightly decrease the formation contents. Compared to Al₂(SO₄)₃, Al(NO₃)₃ lowers more the system pH value and extends hydration stages to later reaction times, consistent with the results of setting time and calorimetry, revealing stronger effect of NO₃^- than SO₄^2- under same Al³⁺ concentration. The aqueous composition results suggest well that Al³⁺ is precipitated rapidly at the beginning of reaction, leading to certain amorphous Al-containing phases. At low w/c of 0.27 the anions of NO₃^- and SO₄^2- are precipitated as crystalline KNO₃ and K₂SO₄, filling micro-pores in hardened paste matrices, which indicates attention required on matrix volume stability as high dosages of these aluminum salts are used.

The main findings of this work are concluded as follows. All the investigated aluminum salts (AlCl₃, Al(NO₃)₃ and Al₂(SO₄)₃) could retard MKP cement setting. Furthermore, they show better retardation effect than borax at the same dosage. In contrast to Al(NO₃)₃ and Al₂(SO₄)₃, AlCl₃ could lead to sample expansion and cracks, indicating its unsuitable use in MKP cement-based materials. Moreover, both Al₂(SO₄)₃ and Al(NO₃)₃ could lower the cement early strengths, which occurs within the first 3 h and 12 h, respectively, consistent with their retardation effects. But they show no adverse impacts on the paste strengths at late ages, compared to the reference paste. Compared to Al₂(SO₄)₃, Al(NO₃)₃ is more effective on reducing system pH, extending hydration stage, slowing down hydration heat release, and on inhibiting slightly K-struvite precipitation, which contribute to better retardation effect. In addition to K-struvite, small contents of amorphous Al-containing hydrates and of crystalline KNO₃ and K₂SO₄ are precipitated in the hydrated matrices, affecting the paste properties.

Large amounts of CO₂ emissions have caused escalating challenges of global warming risk. The construction industry significantly contributes to global CO₂ emissions, prompting the need for sustainable building materials. Carbon capture, utilization, and storage (CCUS) technologies provide a promising solution by reducing emissions and integrating carbon storage capabilities into materials. To this end, foam concrete, a lightweight porous material (normally 500-1600 kg/m³), has gained increasing attention for its significant advantages in energy-efficient and low emission buildings. Traditional foam concrete is lightweight with good insulation properties but has low strength and limited environmental benefits. In contrast, CO₂ foam concrete offers the potential of reducing emissions during production and CO₂ sequestration, making it a more environmentally friendly alternative.

The objective of this research was to develop a high-strength CO₂ foam concrete (HSCFC) that incorporates carbon capture and storage (CCS). By enhancing foam stability and optimizing the concrete mixture, this study firstly developed a new HSCFC product with assessing the CO₂ sequestration efficiency and micro/macro performance of the devloped material, thereby verifying its potential as both a sustainable building material and a solution for carbon reduction.

The design concept of HSCFC was based on integration of a CO₂ foam precursor into a high-strength cement-based paste, utilizing particle packing theory to optimize strength and density. The CO₂ foam was produced using a amphiphilic nano-silica modified sodium dodecyl sulfate foaming agent, enhancing the stability of the foam and preventing collapse of foam concrete. The HSCFC was prepared by mixing the CO₂ foam with a high-strength cement slurry, followed by curing and a series of micro and macro testing to evaluate HSCFC performance, including compressive strength, concrete density, pore structure and microstructures.

1) The CO₂ foam exhibited a well-distributed pore size of 50-100 μm, which grew over time due to natural drainage and coalescence processes. The stability of the foam was attributed to the surfactant and nanoparticle interactions, which prevent CO₂ diffusion and enhance liquid film strength. The overal performance of CO₂ foam met the requirments of the Chinese standard for the concrete foam, which thus could be used to fabricate the foam concrete products. 2) The rheological behavior of HSCFC followed the Bingham model, with an increase in yield stress as CO₂ foam volume increases, up to 110% compared to the control group. However, excessive foam content led to a reduction in yield stress due to the weakening of the matrix from enlarged pores. The plastic viscosity was inversely proportional to foam content, enhancing workability by reducing internal friction. 3) The increase in CO₂ foam content led to the decrease in compressive strength of HSCFC. The compressive strength of HSCFC was found to be significantly higher than conventional foamed concretes, which is more than double that of typical foam concretes of similar density. Microhardness tests revealed values above 90 HV, close to ultra-high-performance concrete, despite the presence of foam. 4) X-ray computed tomography (XCT) analyses indicated that the pore structure of HSCFC became more complex as foam content increased, with the peak value of the pore size enlarging from 162.8 μm to over 1100 μm. This indicated the importance of balancing the compressive strength and density for the HSCFC. Additionally, using CO₂ foam improved the cement hydration degree and led to the in-situ growth of calcium carbonate (CaCO₃) on pore walls, enhancing pore wall strength. However, excessive foam amounts resulted in non-uniform pore distribution, leading to reduced material strength. 5) The thermal conductivity of HSCFC was lower than conventional concretes due to its porous structures. This property made the material suitable for energy-efficient buildings, where insulation is critical for reducing energy consumption. Therefore, the use of CO₂ foam could realize the synergy of active and passive carbon reductions by carbon mineralization and decreasing engergy uses respectively.

This research demonstrated the successful design and development of a high-strength CO₂ foam concrete system with superior mechanical properties, enhanced carbonation, and thermal insulation capabilities. The use of CO₂ foam not only provided an effective CCUS solution but also significantly improved the performance. The synergistic effects of foam stability, internal carbonation, and matrix densification showed a promising pathway for the sustainable development of carbon-neutral building materials. Future research could focus on improving the overall performance of HSCFC and investigating its long-term service values. This material presents a novel approach for CO₂ capture and utilization, alongside advancements in foam concrete technology. The developed materials can be further optimized to improve carbon sequestration efficiency, facilitating their use in innovative structures like energy-efficient and floating buildings.

Due to the distinctive climatic characteristics of the Qinghai-Tibet Plateau, such as an extremely frigid climate and significant diurnal temperature fluctuations, concrete structures in this region necessitate substantial freeze-thaw durability. Presently, one of the most efficacious approaches to enhance the freeze-thaw durability of concrete involves incorporating tiny air bubbles into the mixture through the use of air-entraining agents (AEAs). Nevertheless, researchers have discovered that the low atmospheric pressure (LAP) prevalent on the plateau may result in inadequate air content within air-entrained concrete (AEC). Only a limited number of research findings have confirmed that LAP deteriorates the air-void structure of AEC, with relatively few studies investigating its impact on freeze-thaw durability of AEC. Furthermore, there is limited research on the impact of LAP on the distribution of air-void diameters in AEC. Therefore, the air content, air-void structure parameters and freeze-thaw durability of the AEC prepared under normal atmospheric pressure (NAP) and LAP were examined in this study. The primary focus is to investigate the impact of LAP on the distribution of air-voids with varying diameters, thus to further explore how these influences degrade both the air-voids structure parameters and freeze-thaw resistance durability of AEC.

Ordinary Portland cement was used in the experiment. The coarse aggregate was crushed stone and the fine aggregate was river sand. In the experiment, two types of AEAs and a high-efficiency superplasticizer were used to produce the AEC. The experiments were conducted in laboratories located in Lhasa, China (atmospheric pressure values of 64 kPa) and Beijing, China (101 kPa). Three air content design levels (D) were used in this study, namely, 3%, 5%, and 7%. Air content tests, air-void analyzing tests and freeze-thaw durability tests were conducted.

Compared to the AEC-N (AEC prepared under NAP), the LAP resulted in a significant increase in the loss of air content of fresh concrete (A_f), with a greater magnitude observed as D increased. The average and maximum losses in air content of hardened concrete (A_h) for AEC-L (AEC prepared under LAP) were approximately 0.5% and 0.7% higher than those for AEC-N, respectively. Furthermore, due to decreased atmospheric pressure, microvoids within AEC experienced a more substantial decrease with increasing D. Notably, air-voids ranging from 0-100 μm exhibited the most pronounced influence. However, increasing D still effectively enhanced the proportion of microvoids within AEC. The proportion of A₁₀₀₀ (the air content in voids with diameters no greater than 1000 μm, the same below) in A_h showed an overall increasing trend followed by a subsequent decreasing trend as A_h increased, both under LAP and NAP conditions. However, when comparing AEC-N to AEC-L with similar D, the former demonstrated significantly higher proportions of A₁₀₀₀. With similar A_h, the reduction in atmospheric pressure led to a significant decrease in the proportion of microvoids (especially A_(0,100]), a pronounced increase in the proportion of trapped air-voids proportion (A_(1000,4000]), and small changes in the proportion of mesovoids (A_(300,500]) and macrovoids (A_(300,500]) within the AEC. The variation in the proportion of microvoids and trapped air-voids significantly influenced the spacing factor (). Specifically, an increase in the air content of microvoids (particularly A_(0,100]), leads to a decreased in the of concrete. Conversely, an increase in the air contentof trapped air-voids resulted in an increase in the . However, the proportion of air content in the mesovoids and macrovoids was most weakly correlated with the .

Concrete subjected to a plateau environment may experience more severe freeze-thaw conditions. when the water-cement ratio was 0.44, the required A₁₀₀₀ under LAP and NAP was about 3.90% and 2.92%, respectively, for achieving a frost resistance durability factor (F_D) of 100%.

The analysis presented in this study demonstrates that the maximum allowable F_D in concrete under LAP is around 3.93%, whereas under NAP it is 4.73%. Clearly, despite the deterioration of air void structures caused by LAP, as long as the air content introduced by AEAs in AEC reaches a certain threshold level, the frost resistance durability requirements can still be met. It is suggested that while meeting the mechanical performance requirements of AEC design, there should be an increase in the maximum air content level of fresh concrete by approximately 2.0% compared to plain areas. This adjustment aims to ensure that the frost resistance performance of concrete meets the design requirements.

The LAP caused a more pronounced loss in the air content of hardened AEC-L, and the loss increased as the D increased. The average and maximum loss was about 0.5% and 0.7% higher than those under AEC-N, respectively. Compared to the AEC-N with the same D, a decrease in atmospheric pressure resulted in a reduction in the air content of the microvoids within the AEC-L. This effect was particularly pronounced for air-voids with diameters not exceeding 100 μm, and consequently leading to the deterioration of the air-void structures and a reduction in frost resistance performance of the AEC-L. The variation in the proportion of air content in microvoids and trapped air-voids has a significant impact on the of the AEC. The decreases with an increase in the proportion of air content of microvoids, particularly A_(0,100] and A_(100,200], while it increased with an increase in the proportion of air content of trapped air-voids. The increase in the D of the concrete effectively enhanced the air content proportion of microvoids, thereby enhancing the frost resistance durability of concrete. It is recommended that when designing AEC for plateau areas with severe freeze-thaw environments, the D of fresh concrete should be increased by a maximum of approximately 2.0% compared to plain areas.

Ultra-high performance concrete is an advanced cement-based material with high strength and high durability. Due to the low water-binder ratio and high binder consumption of UHPC, the early shrinkage of UHPC during the setting and hardening process is larger than that of conventional concrete and high-strength concrete, which easily leads to the cracking of engineering structures and affects the performance of structures. Therefore, it is of great significance to study the early shrinkage characteristics of UHPC for the optimization of the material and the prediction of early cracking. In this paper, the effects of binder-sand ratio, water-binder ratio, different fiber types and contents and curing environment on the early shrinkage performance of UHPC were investigated, and the early shrinkage model of UHPC was established by combining BPNN and WOA-BPNN neural network.

In this study, a total of seven groups of specimens were set up, considering four control factors, i.e., the cement-sand ratio, water-binder ratio, curing environment and polypropylene fiber volume content. Each group consisted of four specimens with size of 25 mm×25 mm×280 mm (including three specimens for drying-shrinkage tests and one specimen for autogenous-shrinkage tests). The specimens were cured in a curing box with a temperature of (25±2) ℃ and a relative humidity of (98%±2%). At the same time, a control group was set up, which was cured in a natural environment with a temperature of (28±5) ℃ and a relative humidity of (60%±15%). The shrinkage deformation of the specimens was measured by a specific length meter at 0, 1, 2, 3, 4, 6, 8, 10, 12 h and 1.0, 1.5, 2.0, 2.5, 3.0, 4.0, 5.0, 6.0, 7.0, 9.0, 11.0, 14.0, 18.0, 21.0, 25.0, 28.0 d. Combined with BPNN and WOA-BPNN machine learning models, the measured data were trained with small samples, and finally a neural network model that can be used to predict the early drying and autogenous shrinkage performance of UHPC was obtained.

The early drying-shrinkage and autogenous-shrinkage of UHPC increased with an increase of the cement-sand ratio, and the MIC values of drying shrinkage and cement-sand ratio also increased. The early drying-shrinkage and autogenous-shrinkage of specimen with a cement-sand ratio of 1.2 increased by 108% and 60%, respectively, compared with specimen with a cement-sand ratio of 0.8. The early drying-shrinkage and autogenous-shrinkage of UHPC increased with a decrease of the water-binder ratio. The appropriate amount of polypropylene fiber and steel fiber mixture would limit the early shrinkage of UHPC. With an increase of polypropylene fiber content, the inhibitory effect on both the early drying-shrinkage and autogenous-shrinkage grew firstly and then weakened. The specimens with fiber content of 0.10 % exhibit the best inhibitory effect. The drying-shrinkage expressed a great correlation with the curing method, and the MIC value was 0.56. The correlation between the autogenous-shrinkage and curing method is small, and the MIC value is 0.27. The drying-shrinkage rate and self-shrinkage rate under dry curing conditions were 2.5 times and 1.2 times that under standard curing conditions, respectively.

The two machine learning (ML) algorithms were used in the shrinkage prediction and exhibit good accuracy. The WOA-BPNN algorithm expressed delightful predicted accuracy, whose R², RMSE and MAE for drying-shrinkage model are 0.959, 0.050 and 0.040, respectively, and R², RMSE and MAE for autogenous-shrinkage model are 0.896, 0.076 and 0.053, respectively. The predicted results indicated that the whale optimization algorithm could improve the ML model effectively.

The main conclusions of this paper are given as follows: 1) The fine aggregate has an inhibitory effect on the early drying-shrinkage and autogenous-shrinkage of UHPC. When the cement-sand ratio decreases from 1.2 to 0.8, the early drying-shrinkage and autogenous-shrinkage increase by 108% and 60%, respectively. 2) The early drying-shrinkage and autogenous-shrinkage of UHPC increase with a decrease of water-binder ratio, and the decrease of water-binder ratio leads to the advance of UHPC self-drying phenomenon. 3) The research found that the polypropylene fiber volume content of 0.10% exhibit the best inhibitory effect on the UHPC shrinkage. 4) The curing method had a great influence on its early drying-shrinkage and autogenous-shrinkage. The drying-shrinkage rate of specimens under dry curing condition is 2.5 times that under standard curing condition, and the autogenous-shrinkage rate of specimens under dry curing condition is 1.2 times that under standard curing condition. The early drying-shrinkage and autogenous-shrinkage predicted results of UHPC based on the WOA-BPNN neural network show nice accuracy and robustness compared to that of BPNN.

Freeze-thaw damage is a significant factor affecting the durability of concrete in cold regions. The research on the mechanisms of freeze-thaw damage in concrete remains inconclusive. Currently, the primary methods for investigating freeze-thaw issues in concrete include experimental research and numerical simulation studies. The research focus can be categorized into two areas: macro structural performance degradation and micro pore structure analysis. However, most current value methods are limited to the micro level when simulating the freeze-thaw behavior of concrete, and there is a lack of research on the macroscopic mechanical properties of concrete following freeze-thaw cycles. At the same time, grid dependence arises when addressing crack problems. Consequently, this work proposes a mesoscopic-scale peridynamic freeze-thaw model. The relationship between pore frost heaving force and frost heaving force state is established. The pore frost heaving force is incorporated into the peridynamic motion equation as a force state, enabling the simulation of cracks and the prediction of residual bearing capacity in freeze-thaw damaged concrete.

The commercial concrete purchased by Qingdao Yehua Jianzhong Construction Engineering Co., Ltd. is used. The concrete size is 100 mm × 100 mm × 100 mm, the concrete strength is C40, and the water-cement ratio (W/Cs) is 0.51. The freeze-thaw cycle test was carried out by TDR-15D concrete rapid freeze-thaw test machine. The number of freeze-thaw cycles was set to 0, 30, 60, 90 times and 120 times, and the temperature range was (-18±2)-(18±2) ℃. Three concretes were set up for repeated tests under each cycle. Before the start of the test, the concrete was first immersed in water for 4 d to ensure that the water fully penetrated the concrete. Then the freeze-thaw test machine was started, and the freeze-thaw cycle test was carried out according to the preset cycle times and temperature. After the test, the concrete test block was taken out to dry the surface scum and water, numbered and weighed, and the dynamic elastic modulus was detected. The uniaxial compression test was carried out on the MTS Exceed E64.305 electro-hydraulic servo universal testing machine.

The open porosity measured by the test is 1.68%, and the pore frost heaving force calculated by the finite element method is 58.3 MPa. After obtaining the relevant parameters, the freeze-thaw simulation is realized by matlab programming based on the peridynamics theory.

The results of the freeze-thaw tests and uniaxial compression tests were compared with the simulation outcomes. The diagram illustrating simulated freeze-thaw damage depicts the phenomenon of surface spalling and internal crack propagation, which aligns with the test results. The uniaxial compression stress-strain curves exhibit a similar trend, with a maximum difference in peak stress of 4.7%. After undergoing freeze-thaw cycles, the elastic modulus of the test group was initially low but subsequently increased. This is due to the cracks being filled with ice and broken concrete chips, resulting in a loose state of the concrete, which leads to a low elastic modulus during the initial stages of loading. As the load increases, the stiffness of the concrete also increases once the crack has closed.

The open porosity and pore frost heaving forces of concrete are critical factors influencing its freeze-thaw performance. The freeze-thaw damage and residual bearing capacity of concrete with open porosities of 1.0%, 3.0%, 5.0%, 7.0% and 9.0%, as well as pore frost heaving forces of 40, 50, 60, 70 MPa and 80 MPa, are discussed in detail. The influence of open porosity and pore frost heaving forces on the freeze-thaw performance of concrete follows a similar pattern. When the two values are low, there are no freeze-thaw cracks or minor freeze-thaw cracks in the concrete. When the two values are large, the freeze-thaw cracks cover the whole concrete. The rate of decrease in the residual bearing capacity of concrete initially increases and then subsequently decreases. This is because when the freeze-thaw damage is small, each new damage may become the starting point of the freeze-thaw crack, so the damage will promote the new damage. The damage caused by freezing and thawing in the later stages reaches saturation, resulting in a limited number of cracks and overall damage in the concrete. The concrete has been damaged and should be regarded as having completely lost its load-bearing capacity.

The main conclusions of this paper are summarized as following. The peridynamic freeze-thaw model closely resembles the test used to simulate surface spalling and crack propagation in concrete subjected to freeze-thaw cycles. The simulation results of uniaxial compression for freeze-thaw concrete align closely with the experimental findings. The maximum difference in stress values is 4.7%, while the maximum difference in residual bearing capacity is 3.7%. The rate of freeze-thaw damage in concrete increases gradually during the early stages and then slows down in the later stages. The freeze-thaw damage of concrete is characterized by surface damage spalling and internal crack staggered expansion. Through regression analysis, the linear relationship between freeze-thaw damage and residual bearing capacity is established, with a coefficient of determination reaching 0.966. Compared to the mechanical model calculations, the regression curve is utilized to predict the residual bearing capacity based on freeze-thaw damage, resulting in an efficiency increase of approximately 360 times.