Crystallization control is crucial during glass production. In glass crystallization theory, it has been hypothesized that there is a direct relationship between glass-crystal interfacial energy and nucleation rate. Since it is difficult to measure interfacial energy directly, classical nucleation theory is used to obtain it. However, the estimation process is complex; various measurement data such as the nucleation rate, viscosity, and Gibbs free energy barriers between the glass and crystal are required. Overall, estimating the glass-crystal interfacial energy is time-consuming. In this regard, the interfacial energy can be directly obtained using molecular dynamics (MD) simulations. In this study, we implemented an interface model to estimate the glass-crystal interfacial energy based on the theory of Tielemann et al., which was recently introduced to generate crystal orientation planes using minimum-energy cuts. To the best of our knowledge, this is the first work in which the crystal orientation plane determined based on a minimum-energy cutting process has been used to build a glass-crystal interface in a more realistic environment (i.e., hypothesizes that a minimum-energy structure may occur during crystal nucleation) and calculate the interfacial energy. To achieve this, we considered stoichiometric alkali (Li, Na, and K) disilicate (2SiO₂) glasses and crystals, as some early-stage experimental data have been reported, which are highly beneficial for validating the MD results. The effects of different potential models were also investigated and found that they had a significant impact on the reproduction of the experimental trend.

Simulations were performed using the LAMMPS package. The temperature and pressure were controlled with the Nose-Hoover thermostat and barostat, respectively. After setting up the glass-crystal interface using NPT, we ran MD simulations with NVT for another 200 ps to calculate the interfacial energy. The MD simulations were performed with a time step of 1 fs. Periodic boundary conditions were applied in all directions. Cutoff distances were applied according to references. We used three interatomic potential (SHIK, Du, and Pedone) models to estimate the interfacial energies of Li₂O-2SiO₂ glass-Li₂O-2SiO₂ (001), Na₂O-2SiO₂ glass-Na₂O-2SiO₂ (010), and K₂O-2SiO₂ glass-K₂O-2SiO₂ (001). All three of these potential functions are widely used. The glass structure was fabricated from the crystal structures by simulating a melt-quenching process. First, the Li₂O-2SiO₂ (001), Na₂O-2SiO₂ (010), and K₂O-2SiO₂ (001) crystal structures were prepared. Half the crystal was kept fixed, while the other half was melted at 3500 K with a canonical ensemble (NVT) for 300 ps and then quenched to 300 K at a cooling rate of 5 K/ps. After the melt-quenching process, the crystal part was unfixed and relaxed for 200 ps with a glassy structure using an isobaric-isothermal ensemble (NPT) at a temperature of 300 K. An interfacial model was developed by the theory of Tielemann et al, where the crystal plane was determined through the minimum energy cut in the crystal structure.

We first evaluated the glass density by calculating the local density profile along the z-direction of Li₂O-2SiO₂ glass-Li₂O-2SiO₂ (001), Na₂O-2SiO₂ glass-Na₂O-2SiO₂ (010), and K₂O-2SiO₂ glass-K₂O-2SiO₂ (001). The glass structures show density results comparable to the experimental data.

The calculated interfacial energy values using the SHIK potential show a similar experimental trend, while the Du and Pedone potentials are unable to reproduce the experimental trend. However, the potentials of Du and Pedone show better values of interfacial energy for the glass-crystal interface of Li₂O-2SiO₂ than the SHIK. Our MD results demonstrate that among the potential models, SHIK is a good candidate for calculating interfacial energy in terms of experimental reproduction.

We analyzed the interfacial coordination number (i.e., cation-anion) in the contact area between the glass and crystal surfaces. These coordination numbers are difficult to estimate experimentally. Typically, the coordination number is determined in the bulk region of a glass or crystal structure. We found that the interfacial coordination number varies significantly at the glass-crystal interface of Li₂O-2SiO₂, Na₂O-2SiO₂, and K₂O-2SiO₂. Among the interfacial systems, coordination number values of two, three, four, and five were observed, the most frequently observed value was one. The values were significantly lower at four and five for K₂O-2SiO₂ and Na₂O-2SiO₂, respectively. Our MD results demonstrated that the cation (Li, Na, and K)-anion environments were not the same in the interfacial domain. For further evaluation, we calculated the bond strength-coordination number.

The bond strength-coordination number has been reported for chalcogenide glasses. In the present work, the bond strength-coordination number is introduced for glass-crystal interface systems. The calculated results of bond strength-coordination number show a decreasing trend in the order Li > Na > K, which is similar to interfacial energy. Overall, analysis revealed that the alkali (Li, Na, and K)—O bond plays a crucial role in the interfacial strength.

We developed an interface model using molecular dynamics simulations based on the minimum energy cut of crystals with glassy structures. The minimum energy cut of the crystal was determined by applying Tielemann theory. The modeled interfaces were between glass and crystals in stoichiometric alkali (Li, Na, and K) disilicates. The interface model reproduced the experimental interfacial energy trend of Li > Na > K; the interatomic potential model was found to have a significant effect on reproduction. The MD results indicate that the SHIK-based MD model could reproduce the experimental results better than the other potential models. In addition, the interfacial coordination number was also calculated, which varies depending on the alkalinity of Li, Na, and K. Using the interfacial coordination number, the bond strength-coordination number was estimated, and found that the interfacial energy trend Li > Na > K is related to the bond strength-coordination number.

Scintillators are functional materials that immediately emit low-energy photons after absorbing high-energy ionizing radiations. Glasses have advantages as a material form for scintillator use, such as low manufacturing costs, ease of formability, high optical transparency, and high compositional tunability. A ⁶Li-glass scintillator doped with Ce is a commercial glass scintillator, and it has been used for thermal neutron detections owing to ⁶Li(n,α)³H neutron capture reactions. In general, the ³He counter is used in practice for neutron detections; however, the depletion of the ³He supply has driven vigorous research and development of alternative materials. ¹⁰B-contained glasses have an advantage of their large thermal neutron capture cross section (3840 barn). In this study, we focused on MgO-P₂O₅-B₂O₃ glass systems. They are composed of light elements; hence, neutron detection signals can be easily distinguished from noise due to X- and γ-rays. Furthermore, alkali-earth-embedded P₂O₅-B₂O₃ glasses realize high optical transparency with good chemical durability. As luminescence centers, Eu was selected. Eu exists in two different states: Eu²⁺ in reduction states and Eu³⁺ in oxidation states. The former shows broad emission bands with fast decay times of ~μs, whereas the latter shows sharp emission bands with slow decay times of ~ms. In general, Eu³⁺ is governed in the glasses prepared in an air atmosphere; however, alkali-earth-embedded P₂O₅-B₂O₃ glasses can exhibit Eu²⁺ owing to the localized reducing atmosphere generated by NH₃ derived from the raw material (NH₄H₂PO₄). Here, Eu-doped MgO-P₂O₅-B₂O₃ glasses with various concentrations of Eu were synthesized by the conventional melt quenching technique in air, and their optical, photoluminescence (PL), and scintillation properties were examined.

Undoped and 0.1%, 0.3%, 1.0%, and 2.0% (in mole fraction) Eu-doped 25MgO-30P₂O₅-45B₂O₃ (MPB) glasses were synthesized by the melt quenching method. Raw materials were homogeneously mixed and transferred into an alumina crucible. The powders were melted at 1100 ℃ for 1 h with an electrical furnace in air. The melt was flowed onto a preheated stainless-steel plate to quench, and pressed into batches. After that, the glasses were annealed at 400 ℃ for 1 h to remove thermal and mechanical strains. The surfaces of the glasses were polished for the following optical and scintillation measurements. The glass transition temperature (T_g) of the undoped sample was measured with a TG-DTA system. The powder X-ray diffraction (XRD) patterns were measured using a diffractometer. Diffuse transmission spectra were measured using a spectrophotometer (Shimadzu, SolidSpec-3700). PL excitation and emission spectra, PL quantum yields (QYs), and PL decay curves were measured using a Quantaurus-QY and Quantaurus-τ. Scintillation spectra, scintillation decay curves, and afterglow curves under X-ray irradiations, and pulse height spectra of ²⁴¹Amα-rays and ¹³⁷Cs γ-rays were measured with our original setups.

The appearances of undoped and Eu-doped MPB glasses were transparent and included some bubbles. Under ultraviolet light at 360 nm, Eu-doped samples showed bluish-red light. Some parts of the glasses were crushed into powder and used for the XRD measurements. Precipitations of crystalline phases were not observed in the XRD patterns; hence, the prepared samples formed glass phases with no periodical structures. T_g of the undoped glass was estimated to be 535 ℃. The transmittances of all the glasses were 70%-90% at 400-700 nm. Both the undoped and Eu-doped samples showed absorption peaks at 200-250 nm; hence, this can be due to the hosts. In addition, absorption peaks emerged at 250-400 nm in Eu-doped samples. The absorption bands at 250-280 nm and 280-400 nm originated from the 4f⁷-4f⁶5d¹ (T_2g and E_g) transitions of Eu²⁺, respectively. An absorption peak due to the ⁷F₁-⁵D₃ transitions of Eu³⁺ was confirmed at 415 nm in the spectra of 2% Eu. A broad emission band due to the electronic transitions of Eu²⁺ was observed at 400-600 nm under excitation bands at 250-410 nm in the Eu-doped samples. Generally, emissions due to Eu²⁺ are difficult to observe in the glasses prepared in air because of the oxidizing atmosphere. The emissions appeared owing to the localized reduction atmosphere, which derived from NH₄H₂PO₄ and (NH₄)₂O·5B₂O₃·8H₂O. When excitation and monitored wavelengths were respectively set to 340 nm and 420 nm, the PL decay curves were fitted by a sum of two exponential functions. The PL decay time constants of both the fast (0.2-0.5 μs) and slow (0.7-1.2 μs) components were similar to those of other Eu-doped phosphors; hence, they are reasonable as 4f⁶5d¹(T_2g)-4f⁷ and 4f⁶5d¹(E_g)-4f⁷ transitions of Eu²⁺. Eu-doped samples showed emission peaks at 350-500 nm under X-ray irradiations. These emission wavelengths were consistent with those of the PL spectra. Afterglow levels (AL) of undoped and 0.1%, 0.3%, 1.0%, and 2.0% Eu-doped glasses at 20 ms passed after X-ray irradiation of 2 ms were respectively estimated to be 1700×10^-6, 4500×10^-6, 3400×10^-6, 800×10^-6, and 160×10^-6. Pulse height spectra of ²⁴¹Am α-rays (5.5 MeV) were measured using the prepared MPB glasses. Eu-doped glasses showed clear full energy absorption peaks. Light yields (LY) of 0.3%, 1.0%, and 2.0% Eu-doped MPB glasses were respectively calculated to be 70, 150 photons/5.5 MeV, and 40 photons/5.5 MeV.

Eu-doped MPB glasses were synthesized by the conventional melt quenching method. All the samples showed halo peaks with no periodic patterns in XRD measurements. The transmittances were 70%-90%, and absorptions due to electronic transitions of Eu²⁺ and Eu³⁺ were observed. All the samples showed luminescence, which originated from the 4f⁶5d¹-4f⁷ transitions of Eu²⁺. PL QYs of 0.1%, 0.3%, 1.0%, and 2.0% Eu-doped samples were respectively calculated to be 46.7%, 37.2%, 20.4%, and 9.3% when monitored at 350-560 nm under excitation at 320 nm. ALs at 20 ms passed after X-ray irradiations of 2 ms were obtained to be 1700×10^-6, 4500×10^-6, 3400×10^-6, 800×10^-6, and 160×10^-6 in undoped, 0.1%, 0. 3%, 1.0%, and 2.0% Eu-doped samples, respectively. Pulse height spectra of ²⁴¹Am α-rays were measured using the prepared samples. LYs of 0.3%, 1.0%, and 2.0% Eu-doped samples were respectively determined to be 70, 150 photons/5.5 MeV, and 40 photons/5.5 MeV.

Solid oxide fuel cells (SOFCs) are high-efficient solid-state energy conversion devices. However, all-ceramic self-supported SOFCs face several challenges such as high brittleness, difficulty in mechanical processing, poor thermal shock resistance, and limited weldability, which result in high manufacturing costs and restrict the applications in mobile power systems. In contrast, metal-supported SOFCs (MS-SOFCs) with metal materials as the external structural support, exhibit remarkable mechanical strength, low cost, and rapid start-up capability, making it highly promising for mobile applications. The anode is a critical component of MS-SOFCs, serving as the site where fuel oxidation occurs to generate electrons. Its microstructure significantly influences the density and effectiveness of the triple-phase boundaries (TPB), where the gas phase, the ionic phase, and the electronic phase intersect. The TPB density largely determines the polarization resistance, with its low-frequency component being inversely related to anode gas diffusion. A common strategy to enhance gas transport is the incorporation of pore-formers, such as graphite, into the anode raw materials. Most studies focus on the type, particle size, and content of pore-formers, which directly affect the pores number, size, and distribution. In this work, atmospheric plasma spraying (APS) was employed to fabricate three types of anodes and corresponding cells. APS reduces thermal input to the metal substrate, effectively preventing oxidation, deformation, and elemental interdiffusion between the metal support and the anode at high temperatures. This study systematically investigates the influence of graphite incorporation methods on the microstructural evolution of the anode and the resulting cell performance, providing important theoretical insights into the operational mechanisms of SOFC anodes.

Porous 430L stainless steel substrates were used as supports. Three different NiO-GDC (Gd_0.2Ce_0.8O_1.9) anode powders were prepared. C1 is the baseline without a pore-former. C2 contains 40% (in volume fraction) graphite, which is mixed by spray granulation process to produce composite particles. C3 is made by mechanically mixing 40% of the same graphite with C1 powder. All powders are spherical with good fluidity. Anode layers were deposited via APS. Subsequently, the C2 and C3 anodes were heat-treated in air to remove the graphite pore-former at 750 ℃ for 2 h. A ScSZ (Sc₂O₃-ZrO₂) electrolyte and an LSCF (La_0.6Sr_0.4Co_0.2Fe_0.8O_3-δ) cathode were subsequently deposited by APS to build single cells. The microstructure and porosity of the anodes were characterized using scanning electron microscopy (SEM) and image analysis software. The surface roughness was measured by profilometer. The electrochemical performance, including the open-circuit voltage (OCV), current-voltage-power (I-V-P) and electrochemical impedance spectroscopy (EIS), were evaluated in the range of 600-750 ℃ using humidified H₂ as fuel and air as oxidant. The equivalent circuit fitting of EIS data is carried out to quantitatively analyze the contribution of charge transfer, surface adsorption/dissociation and gas diffusion in polarization impedance.

The incorporation of graphite pore-former significantly modified the pore size distribution and total porosity within the anodes. The measured porosity of C1, C2, and C3 was 26%±2%, 37%±3.1%, and 42%±2.3%, respectively. The C1 anode featured a relatively dense structure with uniformly distributed pores, which primarily consisted of submicron cracks and fine pores originating from the thermal stress inherent to the APS process. In contrast, the C2 anode showed a notable increase in both the number and size of pores, which were homogeneously dispersed without significant agglomeration. The C3 anode, however, contained a substantial amount of large pores, mostly ~5 μm in diameter, attributed to the agglomeration of graphite particles during mechanical mixing, resulting in coarse and irregular pore structures after heat treatment. Furthermore, the addition of graphite modified the thermal response characteristics of the agglomerated powder during the spraying process, promoting the formation of a more uniformly melted microstructure in the anode layer. The average surface roughness (R_a) values for C1, C2, and C3 were 6.64, 7.06 nm, and 7.66 μm, respectively, indicating that graphite addition increased anode surface roughness. This phenomenon is due to the thermal decomposition of graphite during the plasma spraying process, where high temperatures cause partial oxidation of graphite in the open atmosphere, thereby generating CO₂ gas. The release of this gas from the incompletely solidified anode surface etches irregular pits and protrusions, ultimately leading to increased surface roughness. The cell without graphite pore-former (C1) consistently demonstrated the highest OCV and maximum power density, reaching 1.0 V and 957 mW·cm^-2 at 750 ℃, respectively. EIS analysis revealed that the anodes with graphite pore-former (C2 and C3) exhibited improved charge transfer capability, thereby reducing the high-frequency polarization resistance. Despite this, the overall output performance of C2 and C3 did not show effective enhancement, which is attributed to their increased ohmic resistance (R_o) and lower OCV. The elevated surface roughness and inherent porosity in the C2 and C3 anodes adversely affected the quality of the subsequently sprayed electrolyte layer, introducing microcracks and gas permeation pathways. This resulted in increased R_o and reduced OCV, ultimately weakening the benefits gained from the reduced polarization resistance.

The method of graphite pore-former addition significantly affects the anode microstructure and overall cell performance. Compared with mechanical mixing, spray granulation produces a superior and uniform pore structure. However, contrary to conventional expectation, the introduction of graphite pore-forming agent into the APS anode reduces overall cell performance due to induced electrolyte defects, which elevated R_o and lowered OCV. Future optimization should focus on strategies to reduce anode surface roughness and refine pore structure without affecting the quality of electrolyte deposition. It is anticipated that this will further enhance the output performance of MS-SOFCs.