Ettringite plays an important role in the early strength and volume stability of cement. In a high-concentration sodium sulfate solution, Na⁺ and SO₄^2- will be embedded between the structural layers of AFm to form a sodium-substituted AFm phase(U-phase). Therefore, the stability of ettringite and AFm under high temperatures and alkaline environments requires attention. Ettringite will not only transform into AFm but also form U-phase in a high-temperature and alkaline environment. However, the specific process of the mutual transformation of ettringite and AFm under the combined action of high temperature and alkalinity is still unclear. The phase transition of hydration products is bound to affect the performance of cement-based materials in high-temperature and high-alkalinity environments. Therefore, it is important and necessary to clarify the transformation process of Ettringite to AFm. Based on this, ettringite was first synthesized by adding a small amount of NaOH to the solution system, and the effect of NaOH on the synthesis and transformation of ettringite was explored, especially the effect on the evolution of the above-mentioned synthesized ettringite to AFm in an environment of 105 ℃. A detailed discussion was conducted, which provided new understanding of the stability of ettringite in high-alkali and high-temperature environments and also provided theoretical support for the performance degradation of cement-based materials under such conditions.

The raw materials Ca(OH)₂ and Al₂(SO₄)₃·18H₂O were used to synthesize the ettringite sample by solution method, and the alkaline environment was adjusted by NaOH. Samples were taken at high temperature at designed time intervals to explore the transformation process of ettringite and washed with deionized water and anhydrous ethanol in turn, dried in a vacuum drying oven at 40 ℃, and ground into powder. The phase composition of all samples was tested by the D8 ADVANCE X-ray diffractometer, and the structure of the samples was analyzed by TOPAS V6 software. ZEISS Gemini 360 FE-SEMs was used to observe the microscopic morphology of the samples. The thermal stability of the samples was tested by a NETZSCH STA449F3 synchronous thermal analyzer. Based on density functional theory (DFT), the Castep module in Materials Studio was used to optimize the crystal structure and calculate the electronic structure properties.

The addition of NaOH is beneficial to the formation of ettringite and has an inhibitory effect on the formation of dihydrate gypsum. Under certain alkalinity conditions, the formation of ettringite is favorable. With the increase of NaOH addition, the short rod-shaped ettringite decreases and tends to be fine needle-shaped, with a certain curvature parallel to the c-axis. The increase in OH^- ion concentration causes [Al(OH)₆]³ to form rapidly, resulting in an accelerated nucleation rate of ettringite and a smaller crystal size. In addition, the addition of NaOH reduces the crystal cell parameter c and the crystal cell volume of ettringite. The defect formation energy of Na⁺ replacing Ca1, Ca2, Al1 and Al2 sites in the ettringite crystal structure is Na@Al1>Na@Al2>Na@Ca1>Na@Ca2 from large to small, it can be concluded that Na@Ca2 is the most stable crystal configuration, which indicates that Na⁺ is more inclined to replace the Ca2 site. Compared with the N-0 sample without NaOH added, the decomposition temperature of the small-sized ettringite synthesized in the NaOH environment is lower, and the stability of the sample is relatively weak.

NaOH has a particularly significant effect on the transformation of ettringite to AFm. The intermediate phase U phase exists in the transformation of ettringite to AFm. The rod-shaped ettringite first split to varying degrees at both ends to form a flaky AFm phase. In addition, there are a small number of small needle-shaped rod-shaped ettringite crystals attached to the surface of AFm and transform from AFm into a thicker hexagonal flaky U phase., and finally form a stable AFm phase. Na⁺ and SO₄^2- are co-embedded between the layers of the AFm structure, resulting in a larger interlayer spacing of U-phase compared to AFm.

The ettringite samples synthesized by adding a small amount of NaOH at room temperature have fewer impurities than pure ones, the unit cell parameters are reduced, and Na ions easily replace some Ca2 sites in ettringite. With the increase in NaOH addition, the morphology of the synthesized ettringite changes from thick rods to curved needle rods, and the thermal stability of the ettringite decreases. The ettringite synthesized by adding a small amount of NaOH has an intermediate phase U-phase during the transformation to AFm under a high temperature environment, and all of them are transformed into the AFm phase after being kept at a high temperature of 105 ℃ for 24 h; however, the ettringite samples without NaOH almost do not undergo transformation. The transformation process of ettringite to AFm under NaOH and a high-temperature environment can be divided into three stages. Stage 1 (within 3 h): In this stage, most of the calcium sulfoxide directly transforms into the AFm phase. Stage 2 (4-6 h): Two reactions occur in the system. One is that Na⁺ and SO₄^2- ions react with AFm to form the intermediate phase U-phase with a larger interlayer spacing. At the same time, Ca²⁺ and SO₄^2- ions in the liquid phase react with AFm again to form calcium sulfoxide, which then decomposes into the intermediate phase U-phase; Stage 3 (8-24 h): The U-phase begins to gradually transform into the larger AFm stable phase.

In modern society, the electricity is the foundation of all intelligent facilities. To realize the intelligent development of cementitious materials, electrical conductivity is one of the important directions for the future research of engineering structures and building materials. Of which, ultra-high performance concrete (UHPC) has been widely used in large-span building structures, tunnel and bridge structures, and repair and reinforcement projects, etc. If the conductive properties of UHPC are realized, and its self-sensing characteristics are given to achieve the demand of intelligent monitoring, it is of great theoretical significance and practical value to ensure the service safety of the building structure. It has reached a consensus at present that UHPC is modified by adding conductive materials to reduce the resistivity. However, conductive materials are expensive and difficult to achieve large-scale engineering applications. The use of waste carbon fibers (WCFs) can not only meet the performance requirements, but also reduce the preparation cost, improve the utilization rate of resources, and avoid environmental pollution. But at this stage, there is a lack of research on the influence of WCFs on the electrical conductivity of UHPC, and the relationship between electrical conductivity and load-deflection needs to be explored. Therefore, this paper took the electrical conductivity and bending performance of UHPC as the research point, and analyzed the bending sensitivity and load-deflection relationship of C-UHPC under different WCFs content. And the variation law between stress, strain and resistance was explored.

The UHPC with conductive properties was prepared by WCFs and steel fiber (SF), and the final mix ratio was determined by combining theoretical analysis and pre-experimental research. The water-binder ratio of C-UHPC was 0.16, and the content of WCFs was 0%, 0.5%, 1.0%, 1.5%, and 2% (in volume), respectively. Due to some differences between WCFs and virgin carbon fibers (VCFs), a comparative study of the two fibers was carried out by SEM, Raman spectroscopy, and contact angle. The compressive strength tests of C-UHPC cured for 1, 3, 7 d, and 28 d were carried out. The conductivity of C-UHPC was tested by four-electrode method (28 d). The bending sensitivity of C-UHPC mainly studied the relationship between stress, strain, and resistance under bending load. During the test, the load was automatically recorded by the stress sensor of the test machine, the deflection was measured by the bottom displacement meter of the specimen, and the resistance was measured by the external resistance box. After loading, the change values of load, deflection, and resistance were collected synchronously. Finally, SEM was used to observe the microscopic morphology of fibers and fiber-matrix interface in C-UHPC.

Compared with VCFs, the surface of WCFs was smoother, but the wettability and graphitization degree were relatively poor. Compared with the reference group (WCFs content of 0%), the compressive strength of C-UHPC decreased after the addition of WCFs, which was attributed to the smaller bond force between the interface of WCFs and matrix. In addition, WCFs and SF were easily entangled with each other, and the positive synergism between fibers was weak. The resistivity of C-UHPC decreased with the increase of WCFs content, but when the content of WCFs was greater than or equal to 1%, C-UHPC reached the conductive threshold, that is, the effective lap probability between fibers satisfied the conductive path of electronic transmission. The main conductive pathways of WCFs in C-UHPC included overlapping, connecting cracks, and micropores. When C-UHPC was subjected to bending load, the specimen underwent uncracked, critical cracks, macroscopic cracks, and finally failure. Of which, the macro cracks were mainly axial cracks, which expanded from bottom to top.

A three-stage constitutive equation applicable to the C-UHPC bending load-deflection model was established with high prediction accuracy. When C-UHPC was subjected to bending load, the resistance changed, and this phenomenon indicated that C-UHPC had a bending sensitive property. Taking the peak load as the change node, the resistance curve of the pre-peak load decreased slightly, and the resistance curve of the post-peak load increased rapidly. This was because when C-UHPC was subjected to external loading, the internal microcracks and micropores were closed, resulting in a temporary increased and more complete conductive path, and the resistance decreased. After exceeding the peak load, the matrix gradually cracked, and the axial cracks expanded, interrupting most of the conductive paths, resulting in an increase in resistance. The stress sensitivity could be divided into two stages: first, the stress sensitivity decreased significantly first and then tends to be gentle with the increase of stress; second, the stress began to decrease after reaching the peak value, and the sensitivity increases. The Poisson's ratio was introduced to improve the formula for strain sensitivity, and the strain sensitivity perpendicular to the loading direction was obtained. The strain sensitivity decreased first and then increases with the increase of strain.

The addition of WCFs could reduce the compressive strength of C-UHPC, but improved its conductivity. Under bending stress, the load-deflection curve of C-UHPC could be divided into elastic stage, bending hardening stage, and failure stage. After mixing with WCFs, C-UHPC took the peak load as the change node, and the resistance curve of the pre-peak load basically did not fluctuate, and the resistance curve of the post-peak load increased rapidly. When WCFs content was 1%, C-UHPC reached the conductive threshold, and continuing to increase fiber content had no significant effect on the resistivity. The sensitivity of C-UHPC decreased first and then increases with the increase of stress and strain.

As a high-quality light source, the performance of phosphor-converted white light emitting diodes (pc-WLEDs) is closely related to the performance of the selected luminescent materials. The traditional method of combining blue LED chips and yellow phosphors to generate white light has problems of low color rendering index and high color temperature. To address the pressing requirements of lighting and display technologies, the development of novel red phosphors with high efficiency and excellent thermal stability is urgent. Perovskite tellurites have attracted extensive attention due to their elemental diversity and structural tunability. Based on the double perovskite Ca₃TeO₆, a series of NaLaCaTeO₆:xEu³⁺ phosphors that can be effectively excited by near-ultraviolet light have been designed and synthesized by elemental design. Through a co-substitution strategy of [Na⁺-La³⁺] substitution of [Ca²⁺-Ca²⁺] ion pair, the problem of low Eu³⁺ doping concentration in Ca₃TeO₆:Eu³⁺ is effectively improved by regulating the matrix structure. Meanwhile, the luminous intensity and thermal stability of the sample were also significantly improved.

Series of NLCT:xEu³⁺ (0.05 ≤ x ≤ 1.0) samples were successfully synthesized by a high-temperature solid-state reaction. The stoichiometric amounts of raw materials, including Na₂CO₃ (105.99 g/mol, 99.9%), CaCO₃ (100.09 g/mol, 99.9%), TeO₂ (159.6 g/mol, 99.99%), La₂O₃ (325.84 g/mol, 99.99%) and Eu₂O₃ (351.93 g/mol, 99.99%) purchased from Shanghai Aladdin Biochemical Technology Co., Ltd., were weighed and grounded into homogeneous powder for 30 min. Afterward, the obtained powder was put into crucibles and heated at 1473 K for 12 h in a muffle furnace. Eventually, the target substances were obtained after cooling down and grinding again. A red LED device was assembled using the NLCT:0.6Eu³⁺ phosphor with a 395 nm chip. The WLED device was fabricated by connecting tricolor phosphors including BaMgAl₁₀O₁₇:Eu²⁺ (BAM:Eu²⁺, blue), (Ba, Sr)₂SiO₄:Eu²⁺ (BSS:Eu²⁺, green), along with the prepared NLCT:0.6Eu³⁺ (red) with a 395 nm chip.

The results of XRD patterns showed that [Na⁺-La³⁺] substituted the [Ca²⁺-Ca²⁺] ion pair, which did not affect the structure of the sample. The diffraction peaks of the samples with x = 0.1-0.6 did not show obvious impurity peaks, indicating that the synthesized samples were pure phases. Eu³⁺ ions tended to occupy the site of La in the matrix rather than Te. DRS results also confirmed that Eu³⁺ ions were successfully introduced into the NLCT matrix. The dominant absorption peak at 395 nm matches well with InGaN-based LED, and the emission intensity reaches the maximum at x = 0.6. The primary mechanism among Eu³⁺ ions in NLCT:xEu³⁺ samples was proved to the dipole-dipole interaction.

The luminous intensity of the prepared NLCT:0.6Eu³⁺ red phosphor is about 1.98 and 3.97 times than that of CT:0.2Eu³⁺ and Y₂O₃:Eu³⁺, confirming the feasibility of the co-substitution strategy. The internal quantum yield and color purity of NLCT:0.6Eu³⁺ are 67.6% and 99.87%, respectively. The CIE color coordinates of this sample (0.658 1, 0.341 0) are closer to the ideal red point (0.670, 0.330) than that of the purchased commercial red phosphor Y₂O₃:Eu³⁺. Compared with CT:0.2Eu³⁺ (about 77.5%), the thermal stability of NLCT:0.6Eu³⁺ is significantly improved. When the temperature rises to 420 K, the emission intensity at 616 nm remains 91.7% of that at 300 K. Furthermore, the sample exhibits excellent color stability by the CIE coordinates shifting only slightly from (0.658 2, 0.340 8) to (0.646 6, 0.351 9) with increasing temperature. The CCT and R_a of the prepared WLED device are 6215 K and 80.6, respectively, and the CIE coordinates (0.317 6, 0.334 0) are located in the white light region.

A series of innovative NaLaCaTeO₆:xEu³⁺ phosphors were synthesized using a high temperature solid-state reaction method. Under 395 nm excitation, the NaLaCaTeO₆:Eu³⁺ samples exhibit a strong red emission at 616 nm (⁵D₀→⁷F₂) of Eu³⁺ ions. Through a co-substitution strategy where [Na⁺-La³⁺] replaced the [Ca²⁺-Ca²⁺] ion pair, the optimized doping concentration of the NaLaCaTeO₆:xEu³⁺ is elevated to 60% (molar fractior), leading to exceptional color purity and high internal quantum efficiency reaching 99.87% and 67.60%, respectively. Furthermore, compared to the Ca₃TeO₆:0.2Eu³⁺ sample pre-substitution, the luminous intensity and thermal stability of NaLaCaTeO₆:0.6Eu³⁺ phosphor witness a notable increase. Even at 420 K, the emission intensity remains at 91.7% compared to 300 K. The resultant WLED displays a color rendering index R_a of 80.6 and CIE chromaticity coordinates of (0.317 6, 0.334 0), underscoring the broad application potential of NaLaCaTeO₆:0.6Eu³⁺ in solid-state lighting.

Alkali-activated fly ash cementitious materials are characterized by several advantageous properties, including corrosion resistance, high-temperature resistance, and the ability to solidify heavy metals. Additionally, these materials demonstrate a high utilization rate of fly ash and are produced through an environmentally friendly process. Despite these benefits, a significant limitation is their slow solidification and hardening at room temperature, which restricts their practical engineering applications. Consequently, alkali-seed composite activated solutions were used to prepare alkali-seed composite excitation fly ash (ASAFA) specimens, and their properties and structural evolution were explored. The results indicated that N-A-S-H gel and zeolite (sodalite and chabazite-Na) were the main components of alkali-seed composite activated solution, and their proportion played a crucial role in determining the excitation effect. Through adjustments in the activity of raw materials, reaction temperature and duration, it is possible to increase the total content and proportion of N-A-S-H gel and zeolite in the seed, thereby encouraging accelerated and uniform growth of the reaction products within the ASAFA system, and reducing the proportion of harmful pores in the hardened structure. As a result, the setting and hardening process and compressive strength of ASAFA was promoted.

The study utilized first-class fly ash (FA1) with a specific surface area of 430 m²/kg as the primary raw material. FA1 was subjected to sorting-roller grinding to achieve specific surface areas of 715 m²/kg and 1014 m²/kg, resulting in the production of FA2 and FA3, respectively. The activity of the three types of fly ash was assessed in accordance with GB/T 1596-2005 "Fly Ash Used in Cement and Concrete." The results indicated that the mechanical activation from the grinding process significantly enhanced the activity of FA2 and FA3 at all ages compared to FA1, with activity values at 28 days reaching 92% and 114%, respectively. An 8 mol/L NaOH solution was employed as the alkali activator for the synthesis process.

The alkali-seed composite activated solution was prepared by treating fly ash with the NaOH solution at elevated temperatures. During the preparation, four technological parameters were systematically adjusted: ① Type of fly ash used; ② Mass ratio of fly ash to NaOH solution; ③ Synthesis temperature; And ④ synthesis time. The ASAFA specimen was created using the alkali-seed composite activated solution in conjunction with FA1. For comparative purposes, a reference specimen (GB) was prepared using a mass ratio of 0.5:1.0 of NaOH solution and FA1. Cubic specimens measuring 40 mm × 40 mm × 40 mm were fabricated and cured at a temperature of (20 ± 3) ℃ and 60% relative humidity. Following the initial curing period, the specimens were demoulded and subjected to additional curing for 7 and 28 days prior to the compressive strength testing.

In the preparation of the alkali-seed composite activated solution, a significant amount of sodium hydroxide solution is utilized, with a mass ratio of 100:4 to fly ash. The simultaneous application of heating and stirring notably accelerates the geopolymerization process. After a synthesis duration of 0.5 h, a compound activator is produced, characterized by a zeolite phase that includes sodalite and sodium-chabazite, exhibiting a markedly higher content than that observed in conventional polymerization methods. As the synthesis time is extended to 1 hour, the sodalite content within the composite activator continues to rise, with the zeolite phase predominantly existing at the nanometer scale. However, upon further extending the synthesis time to 2 hours, a slight decrease in sodalite content is observed, accompanied by an increase in the N-A-S-H gel content. This phenomenon may be attributed to the interaction between the zeolite phase and the strong alkali present in the sodium hydroxide solution, which facilitates the dissolution of aluminum ions from the zeolite phase into the alkali solution. Consequently, the N-A-S-H gel is generated through a polymerization reaction, a process further evidenced by the blurring of grain corners of the zeolite phase. Overall, in the sample designated as AD0.5, a greater coexistence of N-A-S-H gels and zeolite is observed. In contrast, AD1 contains only a limited number of unreacted fly ash particles, while the N-A-S-H gel undergoes a secondary reaction, resulting in the formation of smaller zeolite grains. In AD2, the content of zeolite particles decreases slightly, yet the structure becomes denser, and the particle size increases, alongside a rise in the N-A-S-H gel phase content. The seeds present in the alkali-seed composite activator primarily consist of N-A-S-H gel and zeolite crystal phases. These nano-products serve as nucleation sites that facilitate the geopolymerization of fly ash. Additionally, they fill the spaces between fly ash particles, promoting uniform growth of the induced polymerization products throughout the system. Notably, an increased content of crystal seeds in the composite activator, particularly a higher proportion of small zeolite crystal phases, significantly enhances their positive impact on the advancement of geopolymerization.

The alkali-seed composite activated solution primarily consists of N-A-S-H gel and zeolite phases, with their content and proportions being influenced by several factors, including the activity and dosage of fly ash, as well as the synthesis temperature and duration. By optimizing these parameters, it is possible to achieve the final setting of ASAFA pastes within as little as 2 h. Furthermore, these optimized specimens exhibit significant compressive strength, reaching 4.5 MPa at 7 d and 13 MPa at 28 d. In contrast, alkali-activated fly ash paste demonstrates a considerably longer final setting time of nearly 100 h, along with a much lower compressive strength of only 0.2 MPa at 7 d. This comparison underscores the advantages of using alkali-seed compounds in enhancing the performance of ASAFA pastes.

Supersulfated cement (SSC) is a low-carbon material made from slag, gypsum, and a small amount of alkali activator. However, the poor carbonation resistance limits its large-scale application. In this study, activated carbon with high CO₂ adsorption capacity, as well as activated carbon modified by calcium hydroxide (CH), were doped into SSC (0.5% and 1.0%, in mass) to investigate the effects on the mechanical properties, carbonation depth, phase composition and microscopic morphology before and after the carbonation of SSC. This research aims at providing a effective methond to increase the carbonation resistance of SSC.

P·I 42.5 Portland cement (conforming to GB 175—2023), S95 powder, gypsum and coconut shell-based activated carbon particles (Jiangsu Hartel Carbon Materials Technology Co., Ltd.) were used. The activated carbon powder particles were ball-milled and passed through a 150 μm sieve. The activated carbon particles were immersed in saturated CH solution for 48 h and dried to obtain CH-modified activated carbon. The water-cement ratio of both paste and mortar specimens was 0.4. The dosage of activated carbon was 0, 0.5%, and 1.0% of the total mass, respectively. The compressive strength of mortar specimens (40 mm×40 mm×160 mm) were tested after curing for standard cured for 3, 28 d and after carbonation for 1, 3, 7, 14 d. The hydration degree of slag in SSC before carbonation was analyzed by EDTA method. The phase composition of hardened SSC before and after carbonation were measured by combination of a thermal gravimetric analyzer (TGA) and X-ray diffraction (XRD), The pore structure and micro-morphology hardened SSC before and after carbonation was tested by low-field nuclear magnetic resonance (LF-NMR, field emission scanning electron microscopy (ZEISS GeminiSEM 560), respectively.

Activated carbon significantly improved the carbonation resistance and mechanical properties of SSC. Similar to nanomaterials such as carbon nanotubes, the active functional groups on the surface of activated carbon can provide nucleation sites for the growth of C-(A)-S-H gels, a phenomenon that has been confirmed by SEM results. Combined with the analysis of TG and XRD, it can be seen that this nucleation effect increases the number of C-(A)-S-H gels in the SSC. Meanwhile, according to the LF-NMR, the activated carbon would decrease the porosity and critical pore size of the SSC, thus enhance the compressive strength before and after carbonation.

After carbonation, the activated carbon group generated more calcite than the control group. This indicates that activated carbon promotes the formation of C-(A)-S-H gel and possibly increases the proportion of Ca²⁺ in the gel. This is related to the strong adsorption capacity of activated carbon for Ca²⁺. Therefore, the mechanism by which activated carbon enhances the carbonation resistance of SSC is as follows: Activated carbon provides nucleation sites for C-(A)-S-H gel, promotes the formation of C-(A)-S-H gel, densifies the pore structure, thereby improving the mechanical properties and carbonation resistance of SSC. Furthermore, the clusters formed by activated carbon and C-(A)-S-H gel contribute to improve the distribution uniformity of the gel within SSC, further enhancing structural stability.

Furthermore, activated carbon has a high CO₂ adsorption capacity. After the carbonation decomposition of the C-(A)-S-H gel encapsulated on the surface of activated carbon, the exposed activated carbon can absorb the invading CO₂. Saturated activated carbon might act as a 'barrier' to CO₂, preventing CO₂ penetration and further enhancing carbonation resistance.

The carbonation resistance effect of CH-modified activated carbon is significantly lower than that of activated carbon. The CH distributed in the activated carbon could have dissolved into the liquid phase of the SSC and mainly act as an alkali activator. This leads to an excessive amount of alkali activator in the SSC system, thus reducing the C-(A)-S-H gel content, increasing the AFt content, and making the overall structure looser and more susceptible to carbonation. Therefore, activated carbon is effective in enhancing SSC's carbonation resistance and can also be used to load other functional components.

This study demonstrates that activated carbon could significantly enhance the performance of SSC, particularly the mechanical properties and carbonation resistance. The results show that activated carbon can provide more nucleation sites for the generation of C-(A)-S-H gels, enhance the Ca/Si of C-(A)-S-H gels, increase the amount of C-(A)-S-H gel, densify the pore structure, and improve the compressive strength by 10.6% at the dosage of 0.5%. However, the enhancement effect diminishes at higher activated carbon content.

In contrast, CH-modified activated carbon would lead to increase in the carbonation depth and porosity of SSC, significantly reducing its mechanical properties and carbonation resistance. This could be due to CH-modified activated carbon fails to effectively adsorb and immobilize CO₂, instead of raising the alkali activator content in SSC. The excess alkali activator results in a decrease in C-(A)-S-H gel content, an increase in AFt content, and a more porous and easily carbonated structure.

This paper demonstrates the significant practical value of activated carbon as a low-cost material. It not only greatly enhances the mechanical properties and carbonation resistance of SSC, but also offers greater functionalization potential due to its porous structure. By loading other components into the pores of activated carbon, activated carbon is expected to provide an effective means for the functionalization of SSCs and even other materials. The multifunctionality of activated carbon not only opens up new paths for optimizing the properties of materials, but also offers the possibility in a wider range of construction materials in the future.

In recent years, the high-voltage and high-power multi-layer ceramic capacitors (MLCCs) have attracted extensive attention with the development of new energy vehicles, 5G communication and high-power pulse technologies. At present, the materials used in the MLCC inner electrode are mainly precious metal (Ag/Pd or Au) and base metal (Ni or Cu). Base metal materials have more cost advantages compared with the precious metal materials. In order to prevent the oxidation of the base metal Ni or Cu at high temperatures, it is necessary to sintering in a low oxygen or reducing atmosphere. In order to obtain MLCC products with excellent performance, it is necessary to strictly control the matching of the sintering atmosphere between the inner electrode and the ceramic body during the sintering process. The crystal structure and electrical properties of the ceramic will be significantly reduced in the reducing atmosphere or low oxygen atmosphere environment, mainly because the precipitation of metal lead will destroy the lattice structure, or produce lattice defects. Therefore, it is of great significance to study the properties of PbLaZrSnTiO₃-based (PLZST-based) antiferroelectric ceramics in low oxygen sintering atmosphere. In this paper, PLZST-based antiferroelectric ceramics were sintered in four kinds of low oxygen sintering atmosphere. The crystal structure, microstructure and electrical properties of PLZST-based antiferroelectric ceramics were studied. Moreover, the energy storage properties of all samples were studied.

Pb_0.95La_0.02Sr_0.02(Zr_0.50Sn_0.40Ti_0.10)O₃ (PLSZST) antiferroelectric ceramics doped with were prepared using the traditional solid-state method. High quality raw materials, namely Pb₃O₄(95%, Sinopharm), La₂O₃(99.99%, Acros), ZrO₂(99%, Aladdin), SnO₂(99.8%, Sinopharm), TiO₂ (99.8%, Sinopharm) and SrCO₃(99%, Sinopharm), were weighted 500 g according to stoichiometric and ball milling for 12 h with ethyl alcohol as a solvent and dried at 70 ℃. These raw materials were calcined at 850 ℃ for 2 h in a closed crucible. Then, 0.6% BaCO₃-B₂O₃-SiO₂-K₂CO₃(BBSK, in mass) was added to the mixed powder according to the solid weight ratio as the calcination powder. After ball milling for 12 h, the powder was dried. Next, the powder was pressed into pellets at 20 MPa with 5% Polyvinyl Alcohol (PVA, in mass) as the binding agent 2 mm thick and 10 mm in diameter. Finally, the green ceramics with BBSK addition were sintered at 1020 ℃ and low oxygen sintering atmosphere (5.3×10^-10, 2.1×10^-9, 2.1×10^-8, 1.3×10^-7 atm) and air for 2 h in closed crucibles filled with the calcination powder.

In order to study the effect of sintering atmosphere on the crystal structure of PLSZST-0.6% BBSK (in mass) antiferroelectric material, XRD were conducted as shown in Fig. 1. It can be seen that all samples have perovskite structure, and a small amount of second phase is generated, which may be pyrocholate phase produced by Pb volatilization at high temperature. From the splitting peak near 44.2°, it can be seen that the intensity and Angle of the two splitting peaks do not change under the condition of low oxygen atmosphere, which indicates that the tetragonal structure of the ceramic body does not change. The peaks of (200) and (002) move to the lower Angle at lower oxygen content, indicating that the low oxygen atmosphere leads to the lattice distortion of the material. Then, the microstructure was studied. As can be seen in the surface micrograph, there are no metal lead or other substances precipitated from the ceramic body as observed in those samples treated at low oxygen sintering atmosphere. It can be seen from the f dielectric temperature spectra that all samples show the same variation characteristics. With the increase of temperature, the dielectric constant first increases to the maximum value and then slowly decreases. This dielectric anomaly peak is caused by the transformation of tetragonal antiferroelectric phase into paraelectric phase. With the increase of frequency, the dielectric constant and the peak value of dielectric anomaly do not change significantly, which indicates that the ceramics are a diffuse phase transition, and does not change with the change of oxygen content, which indicates that the oxygen content does not change the phase transition characteristics of antiferroelectric. With the decrease of oxygen content, the breakdown field strength and maximum polarization strength of the ceramic decrease, resulting in the gradual decrease of the saturated energy storage density and recoverable energy storage density of the ceramic body. The ceramic sintered at 5.3×10^-10 atm has a recoverable energy storage density of 2.76 J·cm^-3 which is lower than that of the ceramic sintered at air (3.56 J·cm^-3), showing a discharge energy density of 2.21 J·cm^-3.

Multi-layer ceramic capacitors with base metal inner electrodes have been widely concerned because of its lower cost advantages. In order to prevent the oxidation of the base metal inner electrodes (Ni or Cu) at high temperatures, it needs to be sintered in low oxygen or reducing atmosphere. The effect of low oxygen atmosphere on the electrical properties of PLSZST antiferroelectric ceramics was studied. With the decrease of oxygen content, the insulation characteristics and energy storage characteristics of antiferroelectric ceramics were reduced. Under the oxygen content of 5.3×10^-10 atm, the ceramic has a recoverable energy storage density of 2.76 J·cm^-3, which is lower than that of sintered sample under air (3.56 J·cm^-3), has a discharge energy storage density of 2.21 J·cm^-3, and has a discharge current of 28.6 A. This work provides understanding of anti-reduction characteristics of PLZST-based antiferroelectric ceramics, and further promotes the development and production of PLZST-based antiferroelectric multilayer ceramic capacitors with internal electrodes in base metals.

Salt crystallization plays an important role in the physical weathering process of materials, and is generally considered to be one of the main problems affecting the durability of materials, and most of the weathering damage of porous materials is related to salt crystallization. Since the permeation and crystallization of salt solutions are unavoidable. Therefore, it is of great value to explore the crystallization process and mechanism of salt in pores for the study of material durability in the fields of construction engineering, historical building protection engineering, traffic engineering, underground and geotechnical engineering, etc, and is of great significance for improving the disaster prevention and mitigation capabilities of historical buildings and infrastructure, and reasonably predicting the service life of structures. At present, the supersaturated crystallization pressure hypothesis is generally accepted as the cause of salt crystallization failure of porous materials, which holds that the crystallization pressure exerted by crystal on the pore surface is the main cause of the failure of cement-based materials. However, existing study has found that the actual crystallization pressure produced by NaCl evaporation crystallization is about 0.46 MPa, which is much lower than the tensile strength of concrete, and the concrete material still subjected to damage. Therefore, in this paper, the finite element method is used to study the crystallization process of NaCl in the micro pore of hardened cement paste, and the mechanism of crystallization pressure during the crystallization process is revealed. In addition, the effects of pore and crystal geometry, crystallization pressure and filling degree on the maximum tensile stress inside the hardened cement paste are also discussed.

The micro pore model and NaCl crystal model of hardened cement paste are established respectively, and the NaCl crystal model is horizontally put into the micro pore model of hardened cement paste to obtain the NaCl crystallization process model in the micro pore of hardened cement paste. In order to accurately distinguish the crystallization process of NaCl crystal, a two-dimensional spatial filling degree is introduced. The model of NaCl crystallization process in micro pore of hardened cement paste is established in three steps. The first step is to determine the geometric properties of the crystal. The size change of the crystals in the NaCl crystallization process model is controlled by the filling degree, and the crystals without pore wall confinement are obtained, and the crystals in the pore confined state are obtained by Boolean operation. The second step is to split the model. The NaCl crystallization process model is divided into two parts: the hardened cement paste micro pore model and the NaCl crystal model. The third step is to apply crystallization pressure. The stress distribution diagram of hardened cement paste and NaCl crystal is obtained by applying crystallization pressure to the contact parts between the crystal and the pore wall in the micro pore model and the NaCl crystal model. Results and discussion During the crystallization process of NaCl in the micro pore of the hardened cement paste, the tensile stress region is distributed within 2 μm of the pore radius, and the maximum tensile stress in the crystallization process is about 3 times of the crystallization pressure, which appears at the apex of the long axis of the pore, while the NaCl crystal is mainly compressed and the tensile stress is small. NaCl crystal do not fail before hardened cement paste. The pore axial ratio, filling degree and crystallization pressure are the main factors influencing the maximum tensile stress in the cement-based material during the crystallization process. The maximum tensile stress of hardened cement paste increases with the increase of pore axial ratio, filling degree and crystallization pressure during the crystallization process, while the crystal axial ratio has no obvious effect on the maximum tensile stress.

There are free crystallization and crystallization restricted stages in the crystallization process of salt in micro pore. The crystallization pressure generated by the crystallization restricted stage is a major factor causing damage to cement-based materials, and the tensile stress concentration caused by crystallization pressure is the direct cause of the damage of hardened cement paste. The crystallization pressure leads to the tensile stress concentration in the hardened cement paste, and the maximum tensile stress is about 3 times of the crystallization pressure, which appears at the apex of the long axis end of the pore. NaCl crystals will not be damaged before hardened cement paste. In addition, the pore axial ratio, filling degree and crystallization pressure are the main factors affecting the maximum tensile stress in the hardened cement paste during the crystallization process.

CO₂ curing technology has been extensively studied due to its dual benefits of carbon sequestration and enhancement in the properties of cementitious materials. During CO₂ curing, the carbonatable binders react with CO₂ and form carbonation products, such as calcium carbonate and silica gel, leading to matrix densification and rapid strength development. The development of this technology enables the use of low lime calcium silicate minerals, such as wollastonite (CS), rankinite (C₃S₂) and γ-dicalcium silicate (γ-C₂S) with very low hydration reactivity to produce low CO₂ footprint binders through CO₂ curing. The carbonation activity, carbonation products and mechanical properties of the carbonated matrix of calcium silicate minerals have been widely studied. However, the relationship between the mechanical properties and microstructure of carbonated calcium silicate minerals remains unclear, which potentially limits the application of CO₂ cured cement and concrete as well as low-calcium cementitious materials. In this work, the microstructure and mechanical properties of CO₂ cured Portland cement, β-C₂S, γ-C₂S, C₃S₂ and CS were investigated, and the correspondence between the structure and properties of CO₂ cured calcium silicate minerals and the carbonation process was analyzed.

The raw materials used in this study include Portland cement (PC, PI 42.5), β-dicalcium silicate (β-C₂S), γ-dicalcium silicate (γ-C₂S), rankinite (C₃S₂) and wollastonite (CS). Pastes were prepared with a water-to-binder ratio of 0.18 for microstructural and phase analysis, while mortars were prepared with a constant water-to-binder ratio of 0.25 and sand-to-binder ratio of 2 for mechanical and porosity analysis. Specimens were prepared by compaction molding, with a pressure of approximately 10 MPa for pastes and 25 MPa for mortars. After molding and pre-conditioning, the compact specimens were placed in a pressure chamber. The curing chamber was vacuumed to a pressure of around -0.1 MPa, and maintained for 3 min. After that, CO₂ gas, with a purity of 99%, was injected and maintained at 0.2 MPa at (20 ± 2) ℃ and RH of (60% ±5%) for 3 d.

The carbonation degree of calcium carbonate minerals was determined by a thermal gravimetric analyzer. The phase composition of CO₂ cured calcium silicate minerals was analyzed by XRD analysis, and FTIR spectra were obtained with a Thermo-Scientific IS10 FTIR instrument. Moreover, the microstructure and morphology of paste specimens were examined using Phenom LE SEM. The compressive strength of cylindrical mortar specimens (Φ = 25 mm and h = 25 mm) was tested after CO₂ curing. Meanwhile, the pore structure was tested by a MAG-MED proton nuclear magnetic resonance spectroscopy.

The carbonation degree of γ-C₂S, C₃S₂ and CS exceeded 60%, followed by 50.3% for β-C₂S, while that of PC was only 29.7%. The lowest carbonation degree of PC was because cement particles release more Ca ions than other minerals during early CO₂ curing, rapidly forming a carbonate shell that blocks further carbonation inside the compacts. The carbonation products of calcium silicate minerals were mainly calcium carbonate and silica gel. Calcite was the main crystalline calcium carbonate, along with minor amounts of aragonite and vaterite. The silica gel phases in CO₂ cured γ-C₂S, C₃S₂ and CS showed higher polymerization degree and exhibited a clear demarcation from calcium carbonate, while the silica gel phases in the outer layer of CO₂ cured PC particles were intermixed with calcium carbonate. This was related to the high hydration reactivity of PC, which could react with water and form a certain amount of hydration products (i.e. C-S-H and portlandite) in the pores. The carbonation of porous C-S-H results in the intermixing of calcium carbonate with silica gel. Additionally, the synergy of cement hydration and carbonation also facilitates the leaching of Ca²⁺, thus leading to an evaluated content of amorphous phases in CO₂ cured PC.

After 3 days of CO₂ curing, the compressive strength of β-C₂S and γ-C₂S mortars exceeded 50 MPa, followed by PC and C₃S₂, while the strength of CO₂ cured CS was only 11.6 MPa. The compressive strength of CO₂ cured calcium silicate minerals showed a linear relationship with CO₂ uptake and porosity, increasing with a decrease in porosity and an increase in CO₂ uptake. With the increase of Ca/Si, the CO₂ uptake of CS, C₃S₂ and γ-C₂S increased, the porosity decreased, and thus the compressive strength increased. However, PC and β-C₂S were subjected to both hydration and carbonation, which promote the formation of amorphous phases, leading to lower porosity and higher compressive strength. Moreover, the crystal size of calcite, the content of amorphous calcium carbonate and the interfacial properties of calcium carbonate and silica gel can also impact the compressive strength and microstructure evolution.

The carbonation products of calcium silicate minerals were mainly calcium carbonate and silica gel. The carbonation degree of non-hydraulic γ-C₂S, C₃S₂ and CS was relatively higher, followed by β-C₂S, while PC demonstrated the lowest carbonation degree. The silica gel phases in CO₂ cured γ-C₂S, C₃S₂ and CS showed higher polymerization degree and exhibited a clear demarcation from calcium carbonate, while the silica gel phases in the outer layer of CO₂ cured Portland cement particles were intermixed with calcium carbonate. The compressive strength of CO₂ cured calcium silicate minerals showed a linear relationship with CO₂ uptake and porosity. With the increase of Ca/Si ratio, the CO₂ uptake of CS, C₃S₂ and γ-C₂S increased, the porosity decreased, and thus the compressive strength increased. Portland cement and β-C₂S were subjected to both hydration and carbonation, leading to lower porosity and higher compressive strength. The smaller size of calcite crystals and the higher content of amorphous calcium carbonate also contribute to an increased mechanical property of CO₂ cured Portland cement.

The pore structure of cement-based materials significantly influences their compressive strength. Studies have shown that the nanoscale pore structure significantly changes with variations in water content. Traditional pore measurement methods like mercury intrusion porosimetry and gas adsorption require the drying of specimens, potentially altering or even destroying the pore structure. The accuracy and representativeness of the compressive strength-pore structure relationship model established on this basis are therefore questionable. To accurately describe the important correlation between compressive strength and pore structure of cement-based materials, low-field magnetic resonance relaxation technique is utilized to carry out in-situ, nondestructive testing of saturated cement mortar doped with air-entraining agent. This is then combined with compressive strength measurement results to validate and refine the model describing the relationship between compressive strength and pore structure characteristics.

(1) Specimen preparation To avoid the ambiguous effects of ferromagnetic substances when using by low-field nuclear magnetic resonance (LF-NMR) technique, white Portland cement with low Fe₂O₃ content was used to prepare cement mortars with water-to-cement ratio of 0.4. Cement to sand ratio is 1∶2. And air-entraining agents (SJ-2 type) were incorporated. Chemical composition of white cement were 0.50% Fe₂O₃, 64.60% CaO, 21.71% SiO₂, 4.60% Al₂O₃, 2.80% SO₃, 2.43% MgO, 0.48% R₂O. The mass ratio of air-entraining agents to cement was 0%, 0.05%, 0.10%, and 0.15%. Considering different mass ratios, they are named as WP-BLK, WP-SJ05, WP-WP-SJ10, and WP-SJ15, respectively.

According to the designed proportions, the white cement, sand, air-entraining agents and water were mixed and casted into prisms of size 20 mm×20 mm×50 mm. After demounting at 24 h, they were all cured in saturated lime water at (20±2) ℃ for 180 d. After reaching the specified age, ten specimens were taken from each group of mortar and put into the high-pressure water saturation equipment for 3 d to ensure that the specimens were completely saturated and set aside.

(2) Pore structure test The pore structure of saturated mortars were tested by LF-NMR technology. Utilising a 2 MHz NMR analyser manufactured by Limecho Ltd., China, transverse relaxation test was performed through the Carr-Purcell-Meiboom-Gill (CPMG) pulse sequence at controlled room temperature about 20 ℃. In detail, the echo time τ_e (s) was adopted as short as 60 μs to detect the water confined in nanoscale pores of fast relaxation. Preliminary experimental tests had revealed that echo time τ_e ≤100 μs was short enough to make the contribution of diffusion relaxation negligible due to the controlled low concentration of Fe₂O₃ in white cement pastes. Moreover, A long sequence of N = 40 000-60 000 echoes generated at time τ= nτ_e (n=1, 2, …, N) (s) was recorded to capture the water stored in large capillary pores and even air voids with long relaxation time. 256-1024 scans with enough long repetition time 5-15 s were averaged. Except for the echo time, the rest of the parameters were flexibly adjusted according to water content of specimen. One can see detailed steps for testing pore structure by low-field magnetic resonance techniques in the relevant literature. (3) compressive strength test The compressive strength test of specimens were carried out by YAW-300 microcomputer automatic cement folding testing machine, which was produced by Jinan Hengruijin Testing Machine Co., Ltd., with a maximum capacity of 100 kN. The size of the compression surface of the specimen is 20 mm×50 mm, and the loading rate is fixed at 0.2 kN/s.

LF-NMR technology accurately captures the pore structure characteristics of saturated mortar. Empirical models using total porosity or graded porosity as variables are generally effective, but the empirical parameters obtained from model fitting are not stable enough, limiting the model's applicability. An improved empirical model, based on the pore size distribution curve and assuming a linear relationship between the weights of different pores on compressive strength and their logarithmic pore sizes, offers a concise form, significantly improves fitting accuracy, and accurately reflects the differences within the specimen groups. Furthermore, building on Griffith's fracture theory, the theoretical model, which uses total porosity and logarithmic mean radius as independent variables, exhibits complex expressions and low predictive accuracy. If pores with a radius above 10 nm are regarded as defects, their influence can be explained by Griffith's fracture theory. Meanwhile, pores with a radius of 10 nm or less are regarded as internal pores of the gel, and its effect on the strength of the matrix is approximated using the model description proposed by Powers. The resulting modified Griffith fracture theory model is very high in prediction accuracy, stable, and interpretable, accurately reflecting how minor differences in pore structure lead to small changes in compressive strength.

The incorporation of air-entraining agents mainly increases the capillary pores above 10 nm in radius but also decreases the matrix pores below 10 nm in radius. LF-NMR technology can obtain the full-scale pore structure of cement mortar, aiding in the description of the fundamental correlation between pore structure characteristics and compressive strength.

The extensive use of fossil fuel resources has a serious impact on the environment, and hence energy conservation and CO₂ emission reduction have become the focuses of current development. Solid oxide fuel cells (SOFCs) have attracted the research attention because of the outstanding energy conversion efficiency, wide range of fuel choices, no pollution, and high quality waste. H₂ is the ideal fuel for SOFCs, but its low volumetric energy density and the requirement of high pressure to be liquefied, making it difficult to store and transport. At ambient temperature and pressure, methanol is liquid and its volumetric energy density is much higher than that of hydrogen, which makes methanol easy to storage and transport as a good hydrogen carrier. Furthermore, methanol has been quite mature production process. The industrial methanol production was hydrogenation of CO and CO₂. SOFCs can generate power with direct internal methanol reforming. However, the main problem of direct methanol fed SOFCs is carbon deposition. The formation of carbon deposition not only reduces the cell performance, but affects the anode structure as well. In this study, the anode surface of flat-tube anode supported SOFCs was modified by Barium. The cell performance, durability, and methanol conversion of the unmodified and modified SOFCs were compared and evaluated under methanol atmosphere with a steam/carbon ratio of 0.75.

The anode support was composed of nickel oxide with 3%Y₂O₃-stabilized ZrO₂ (NiO-3YSZ, in mole), and the anode functional layer was made up of nickel oxide with 8%Y₂O₃-stabilized ZrO₂ (NiO-8YSZ). The electrolyte, barrier layer, and cathode layer were 8YSZ, Gd_0.1Ce_0.9O_2-δ (GDC), and La_0.6Sr_0.4Co_0.2Fe_0.8O_3-δ-GDC (LSCF-GDC), respectively. The active cathode area was 60 cm². The cell is denoted as NiO-3YSZ|NiO-8YSZ|8YSZ|GDC|LSCF-GDC. Using 15.68 g Ba(NO₃)₂ powder dissolved in 200 mL deionized water, being stirred at 70 ℃ for 2 h, to obtain 0.3 mol/L barium nitrate precursor solution for wet impregnation. After sealing one end of the Ni/YSZ cell with wax, the prepared barium nitrate precursor solution (70 ℃) was added into the fuel channels of through the sampler, and the other end of the cell was sealed with wax. The cell was kept in an oven at 50 ℃ for 24 hours to ensure that the Ba(NO₃)₂ solution was fully immersed in the anode channels. Then, the cell was dried at 90 ℃ for 1 h to remove the wax, and the cell obtained as above preparation processes was recorded as Ba-Ni/YSZ cell. After the cell was reduced at 750 ℃ by H₂, 0.60 L/min (SLM) H₂ was used as fuel for the initial cell performance test. When methanol was used as fuel for the test, the methanol solution with a steam/carbon ratio of 0.75 was added to the anode channels, and the temperature of the water steam generator was set at 130 ℃. At this time, the corresponding methanol flow rate was 0.30 g/min, and 5 L/min air was passed into the cathode. Fuel cell testing equipment (SOFC-100W, Wuxi Lead Equipment Co., LTD.) and electro-chemical workstation (VMP3B-20, France BioLogic Co., LTD.) were used to test the electro-chemical performance and durability of the cell under different operating condition. The frequencies of electrochemical impedance spectroscopy (EIS) tests were ranging from 20 mHz to 30 kHz. The AC amplitude of EIS testing was 10 mV.

The scanning electrical microscopy (SEM) with energy dispersive spectrometer analysis proved that Ba element was effectively loaded on the surfaces of Ni/YSZ anode fuel channels. Under H₂ fuel, the current voltage (I-V) characteristics of Ni/YSZ and Ba-Ni/YSZ cells at 750 ℃ indicate that the open-circuit voltage (OCV) of Ni/YSZ and Ba-Ni/YSZ cells were 1.12 V and 1.07 V, respectively, which revealed that the cell air tightness was good. At 750 ℃, the maximum power density (P_max) of Ni/YSZ and Ba-Ni/YSZ cells were 555.66 mW/cm² and 507.03 mW/cm², respectively, indicating that Ba element had no obvious influence on the cell output power of flat-tube SOFC fueled by H₂. The electro-chemical performance of Ni/YSZ and Ba-Ni/YSZ cells in methanol atmosphere [r(S/C) = 0.75] showed that the power density of Ni/YSZ and Ba-Ni/YSZ cells at 750 ℃ and 0.8V were 326.72mW/cm² and 452.87mW/cm², respectively. It can be seen that the power density of Ba-modified cell was significantly improved when fueled by methanol. Under OCV condition, the methanol conversion rates of Ni/YSZ and Ba-Ni/YSZ cell were 89.40% and 92.70%, respectively, because Ba on the surfaces of the fuel channels promoted methanol conversion within the cell. The hydrophilicity of the Ba-modified cell was enhanced, and the water cracking into OH^- groups that were adsorbed on the Ni surface was promoted. The durability of Ni/YSZ and Ba-Ni/YSZ cells was tested at 750 ℃ under methanol [r(S/C) = 0.75] and 200 mA/cm². The results showed that the electro-chemical performance of Ni/YSZ cell deteriorated sharply shortly after 3 h to 4 h, while the voltage of Ba-Ni/YSZ cell kept stable. After running for more than 500 h, the OCV value of the Ba-Ni/YSZ cell remained stable, indicating that the air tightness of cell was good, and the voltage was 0.89 V with a deterioration rate of 0.01 (%)/h. The gas composition analysis of the Ba-Ni/YSZ cell during the durability test showed that the conversion rate of methanol in the modified cell during the test was stable, and the conversion rate was larger than 90%. According to the SEM images of the Ba-Ni/YSZ cell before and after the durability test, the average Ni particle size of anode functional layer of the Ba-modified cell lowered from 0.97 μm to 0.92 μm. This deduced that the Ni particle lost in the anode function layer of the Ba-Ni/YSZ cell was obvious after the durability test, leading to the performance deterioration of the cell. Raman spectra showed that Ni/YSZ cell has detectable carbon deposition after the durability test, while Ba-Ni/YSZ has no carbon deposition. This showed that Ba content on the surface of the anode support of flat-tube SOFC could effectively inhibit carbon deposition.

The performance and durability of the flat-tube Ba-Ni/YSZ cell were investigated in this study. The Ba-Ni/YSZ cell was prepared by wet impregnation. The power density of the Ba-Ni/YSZ cell with methanol under r(S/C)=0.75 at 750 ℃ was 452.87 mW·cm^-2 at 0.8 V, larger than that of Ni/YSZ cell. The durability of Ba-Ni/YSZ with methanol under r(S/C)=0.75 at 750 ℃ reached 500 h, while the Ni/YSZ cell revealed a sudden performance degradation. The Ni particle agglomeration was detected for the Ba-Ni/YSZ cell after the durability test, causing the performance degradation during the test. No carbon was observed on the surface of the Ba-Ni/YSZ after the durability test.