In-situ growth of Zn2SiO4@SiO2 ultra-low absorption glass-ceramic pigments, and preparation method and application thereof

Abstract

The application relates to in-situ growth Zn2SiO4@SiO2 ultra-low absorption glass ceramic pigments, a preparation method and application thereof. The in-situ growth Zn2SiO4@SiO2 low absorption glass ceramic pigment comprises an alpha-Zn2SiO4 crystal nucleus and a residual silicon-rich glass phase SiO2 shell layer coated on the crystal grain surface of the alpha-Zn2SiO4 crystal nucleus; wherein the alpha-Zn2SiO4 crystal nucleus and the residual silicon-rich glass phase SiO2 shell layer are connected by a continuous glass phase transition interface.

Description

Technical Field

[0001] This invention belongs to the field of spacecraft thermal control coating materials technology, specifically relating to an in-situ grown Zn2SiO4@SiO2 ultra-low absorption microcrystalline glass pigment, its preparation method and application. Background Technology

[0002] During spacecraft operation in orbit, solar irradiation is the primary external heat source for their surface thermal balance. Thermal control coatings maintain a low solar absorptivity (α). s α-thermal control coatings are used to regulate surface temperature and are a core component of thermal control systems. Paint-type thermal control coatings consist of white inorganic pigments and binders; the solar absorptivity of the pigments directly determines the coating's thermal control performance. Traditional white pigments are mainly metal oxides such as zinc oxide (ZnO) and zinc titanate (ZnTiO3). However, under combined proton, electron, and ultraviolet irradiation conditions, lattice defects such as oxygen vacancies and zinc interstices easily form in the ZnO lattice, leading to a significant increase in solar absorptivity—after 5000 ESH ultraviolet irradiation, the α-thermal control coating... s The increase typically exceeds 0.08, significantly deviating from thermal control design specifications. Therefore, developing novel white pigments that combine ultra-low initial solar absorptivity with excellent space irradiation stability is an urgent need in this field.

[0003] Zinc silicate (α-Zn₂SiO₄) has become a key focus for next-generation thermally controlled pigments due to its wide bandgap (approximately 5.5 eV). Its fundamental absorption edge is located at approximately 225 nm, far deeper than the main energy band of the solar spectrum. It exhibits no intrinsic absorption in the main energy band of the solar spectrum, with a solar absorptivity as low as 0.078 and a high ultraviolet reflectivity of approximately 95%. Its proton irradiation tolerance is superior to ZnO and ZnTiO₃. To further improve its irradiation stability, existing technologies employ a core-shell structure design—first synthesizing Zn₂SiO₄ powder, then coating its surface with an amorphous SiO₂ layer using a sol-gel method. However, the core-shell interface formed by this external coating method is a physically adsorbed weak interface. Under thermal cycling stress (-100–150 °C) and mechanical grinding (pigment pulverization process), the shell layer is extremely prone to peeling or cracking. Furthermore, the preparation cycle is long, with synthesis and coating performed in separate steps, resulting in high process complexity and difficulty in ensuring batch consistency.

[0004] Glass-ceramics are composite materials produced through the controlled crystallization of base glass, allowing for the in-situ precipitation of nano- to submicron-sized grains within the glass matrix. In the ZnO-Al₂O₃-SiO₂ system, glass exhibits a strong tendency for phase separation; by adjusting the composition and heat treatment regime, α-Zn₂SiO₄ nanocrystals can be precipitated. However, current research focuses primarily on the mechanical properties (such as hardness and fracture toughness) or optical transmittance of glass-ceramics, and has not yet utilized the component differentiation effect accompanying the in-situ crystallization process of glass-ceramics for constructing core-shell structured thermally controlled pigments.

[0005] In summary, the existing technology has the following main technical problems: (1) The shell of the existing core-shell structure Zn2SiO4@SiO2 is an external physical coating, and the bonding force at the core-shell interface is weak: The existing technology uses wet chemical methods such as sol-gel method to externally coat the surface of Zn2SiO4 particles with SiO2 shell, and the core and shell are physically adsorbed heterogeneous interfaces with obvious physical abrupt boundary; Under the typical thermal cycling conditions of spacecraft (-100~150 ℃), due to the difference in thermal expansion coefficient between Zn2SiO4 and SiO2 shell (approximately 1.5×10 ℃), the bonding force at the core-shell interface is weak. -6 / K), microcracks are easily generated at the interface, leading to failure of the core-shell structure; in addition, the external coating process is difficult to achieve complete and uniform coating of each grain by the shell layer, and there are still micron-level pores and uncoated areas at the core-shell interface. The coating process is sensitive to parameters such as TEOS concentration, pH value, and hydrolysis conditions. The uniformity of the shell thickness and the batch-to-batch differences are significant, making it difficult to meet the stringent requirements of spacecraft for pigment performance consistency; (2) The existing core-shell structure pigment preparation process is lengthy and does not utilize the in-situ crystallization mechanism of microcrystalline glass to construct the core-shell structure: The existing core-shell structure Zn2SiO4@SiO2 pigment preparation requires multiple processes such as "synthesis of Zn2SiO4 powder → wet chemical coating → drying → high temperature heat treatment". Synthesis and coating are carried out in steps, the process cycle is long, the controllability of the microstructure is limited, and the batch-to-batch consistency is difficult to guarantee. Therefore, existing technologies only focus on the crystallization phase transition of glass-ceramics and their mechanical / optical properties. They have not yet combined the component differentiation effect during the in-situ crystallization process with the core-shell structure design, nor have they recognized that this in-situ formed silicon-rich glass phase shell can replace the external coating layer, providing a continuous transition interface for chemical bonding of Zn2SiO4 grains, and fundamentally solving the problems of weak bonding force and incomplete coating of the external coating interface. Summary of the Invention

[0006] To address the aforementioned technical problems, the present invention aims to provide an in-situ grown Zn2SiO4@SiO2 ultra-low absorption microcrystalline glass pigment, its preparation method, and its application.

[0007] In a first aspect, the present invention provides an in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment, wherein the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment comprises: an α-Zn2SiO4 crystal nucleus, and a residual silicon-rich glass phase SiO2 shell coating the surface of the α-Zn2SiO4 crystal nucleus; wherein the α-Zn2SiO4 crystal nucleus and the silicon-rich glass phase SiO2 shell form a continuous glass phase transition interface.

[0008] Preferably, the α-Zn2SiO4 crystal nuclei have a particle size of 50-200 nm, and the residual silicon-rich glassy phase SiO2 shell has a thickness of 100-300 nm.

[0009] Preferably, based on the total mass of the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment as 100%, the mass percentage of α-Zn2SiO4 crystal nuclei is 30-50%, and the mass percentage of the residual silicon-rich glass phase SiO2 shell is 50-70%.

[0010] Secondly, the present invention provides a method for preparing the above-mentioned in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment, the preparation method comprising the following steps: (1) Mix the glass raw materials and melt them at high temperature, anneal them, and cool them to obtain the precursor glass; (2) The precursor glass is heated to undergo nucleation and crystallization treatments to obtain a microcrystalline glass block; (3) The microcrystalline glass block is crushed to obtain the in-situ grown Zn2SiO4@SiO2 low absorption microcrystalline glass pigment.

[0011] Preferably, in step (1), the glass raw material is composed of the following components by molar percentage: ZnO 34-40%, SiO2 42-50%, Al2O3 8-12%, K2O 4-8%, nucleating agent 0.2-0.5%, and the sum of the molar percentages of each component is 100%; Preferably, the nucleating agent includes at least one of ZrO2, SnO2, and Y2O3, and the molar ratio of ZnO to SiO2 is 0.72-0.85.

[0012] Preferably, in step (1), the high-temperature melting temperature is 1460-1500℃ and the time is 1-3 hours; Preferably, the annealing temperature is 500-550℃ and the time is 1-2 hours.

[0013] Preferably, in step (2), the temperature of the nucleation treatment is 30-60°C higher than the initial temperature Tg of the precursor glass during the glass transition process, the time is 1-2 hours, and the heating rate is 3-8°C / min; preferably, the Tg is 480-648°C.

[0014] Preferably, in step (2), the temperature of the crystallization treatment is 20-50°C lower than the peak temperature Tp of the precursor glass crystallization exothermic peak, the time is 2-8 hours, and the heating rate is 2-5°C / min; preferably, the Tp is 640-836°C.

[0015] Preferably, in step (3), the particle size D50 of the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment is <5 μm.

[0016] Thirdly, the present invention provides an application of the above-mentioned in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment in high-performance spacecraft white thermal control coatings.

[0017] Beneficial effects (1) Ultra-low solar absorptivity: The wide bandgap of Zn2SiO4 (approximately 5.5 eV) places the intrinsic absorption edge at approximately 225 nm, resulting in almost no interband absorption in the main energy band of the solar spectrum (200-2500 nm). The in-situ formed residual silicon-rich glassy phase shell completely covers the surface of α-Zn2SiO4 grains, effectively reducing the absorption and scattering losses of incident light due to surface defects. Based on the above mechanism, the solar absorptivity α of the pigment of this invention is extremely low. s It can be as low as 0.05, compared to existing externally coated core-shell pigments (α- s =0.09) and traditional ZnO pigments (α s Compared to (=0.15-0.20), the performance has achieved a qualitative leap; (2) Excellent resistance to ultraviolet radiation: This effect is achieved through the continuous glass phase transition interface and the thermal matching of the shell-nucleus. After 5000 ESH vacuum ultraviolet irradiation, the change in solar absorptivity Δα of the pigment of this invention is significantly reduced. s ≤+0.04, this effect stems from the suppression of irradiation defect formation by the in-situ core-shell structure. The residual silicon-rich glassy shell and the crystal nucleus form a continuous chemically bonded transition interface, without any externally added physical abrupt boundaries or micron-level pores. Irradiation-generated electron-hole pairs are difficult to capture at interface defects to form stable color centers. The shell itself is a natural continuation of the glass matrix, effectively isolating the crystal nucleus surface from external environmental factors. This contrasts with existing pure Zn2SiO4 pigments (Δα after proton irradiation). s =+0.134) and ZnO pigment (Δα after comprehensive irradiation) s Compared to (+0.08 or higher), the radiation resistance of this invention has significant advantages; (3) High core-shell interface bonding strength and excellent structural stability: This effect is achieved through a continuous glass phase transition interface. The core-shell structure of the present invention has a continuous glass phase transition between the shell and the crystal nucleus, and there is no physical abrupt boundary. The interface is spontaneously formed by in-situ component differentiation driven by thermodynamics. The shell and the crystal nucleus share part of the Si-O-Zn bonding network. After ball milling to D50 of about 2 μm, the core-shell structure can still be completely preserved, indicating that its interface resistance to mechanical peeling is significantly better than the physical adsorption type heterogeneous interface formed by external coating. The thermal expansion coefficients of this outer shell and the crystal nucleus are matched. Under the conditions of spacecraft on-orbit thermal cycling (-150~150 ℃), there will be no interface microcracks due to thermal mismatch. This feature provides a reliable guarantee for the structural integrity of the pigment in the coating preparation and long-term space service process. (4) Simple preparation process: The present invention adopts a one-step melting-quenching-controllable crystallization method. The core-shell structure is naturally formed during crystallization heat treatment, without the need for additional wet chemical coating steps. All raw materials are conventional industrial raw materials, and the preparation equipment is general equipment for the glass industry. The process flow only includes six core steps: "batching → melting → annealing → nucleation → crystallization → crushing". The preparation cycle is significantly shortened, the batch consistency is good, and the feasibility of mass production is high. Attached Figure Description

[0018] Figure 1 X-ray diffraction patterns of the pigments prepared in Examples 1 and 2; Figure 2 The image shown is a DSC diagram of the precursor glass in Example 1. Figure 3 The ultraviolet-visible-near-infrared diffuse reflectance spectrum of the pigment prepared in Example 1; Figure 4 The changes in solar absorptivity of the pigments prepared in Example 1, Comparative Examples 1 and 2 before and after ultraviolet irradiation; Figure 5 This is a SEM image of the pigment prepared in Example 1. Detailed Implementation

[0019] The present invention will be further illustrated by the following embodiments. It should be understood that the following embodiments are for illustrative purposes only and are not intended to limit the present invention.

[0020] First, this invention provides an in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment. The in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment may include: an α-Zn2SiO4 crystal nucleus, and a residual silicon-rich glassy phase SiO2 shell coating the surface of the α-Zn2SiO4 crystal nucleus; wherein, the α-Zn2SiO4 crystal nucleus and the silicon-rich glassy phase SiO2 shell form a continuous glassy phase transition interface.

[0021] In some embodiments, the grain size of the α-Zn₂SiO₄ nuclei can be 50-200 nm. This invention limits the grain size to 50-200 nm, which is an optimal window based on the dual constraints of scattering optics and core-shell structure integrity: when the grains are too small, scattering enters the Rayleigh region, and the scattering efficiency decreases proportionally to the sixth power of the grain size, resulting in a significant increase in solar absorptivity. Furthermore, insufficient crystallization leads to an excessively thin or discontinuous residual silicon-rich glassy shell, preventing the formation of a complete in-situ core-shell protective structure. When the grains are too large, the Mie scattering efficiency plateaus and no longer increases, while the number of scattering centers per unit volume decreases. Simultaneously, excessive crystal growth causes the shell to be stretched and thinned, or even fractured, and the exposed grain surface becomes a precursor site for color center formation under ultraviolet irradiation. Within this 50-200 nm range, the grain size is within the optimal Mie scattering range to achieve an ultra-low solar absorptivity, while the shell remains intact and continuous to provide reliable ultraviolet irradiation protection.

[0022] In some embodiments, the thickness of the residual silicon-rich glass phase SiO2 shell is 100-300 nm.

[0023] When the shell layer is too thin, it cannot effectively isolate the Zn₂SiO₄ crystal nucleus surface from photochemically reactive species such as O₂ and H₂O in the external environment. Ultraviolet photons can penetrate the shell layer to reach the crystal nucleus surface, inducing the formation of color centers at surface defect sites, thus reducing the resistance to ultraviolet radiation. When the shell layer is too thick, although the shell layer itself is transparent amorphous SiO₂, the excessive thickness of the shell layer will dilute the volume fraction of high-refractive-index Zn₂SiO₄ scattering nuclei in the filler, reducing the number of effective scattering centers per unit volume, resulting in a decrease in solar reflectivity and an increase in solar absorptivity. The thickness range of 100-300 nm is the optimal window based on the above balance between shielding effect and scattering efficiency.

[0024] It should be noted that the continuous glass phase transition interface is a compositionally gradual transition region from the Zn2SiO4 crystal nucleus (with a fixed composition of 2ZnO·SiO2) to the residual silicon-rich glass phase shell (SiO2 content > 80 wt%), without any physical abrupt boundary. From the crystal nucleus side to the shell side, the ZnO concentration continuously decreases while the SiO2 concentration continuously increases. The interface region is a Zn-O-Si bonded network, and the shell and crystal nucleus are a chemically bonded continuous glass phase transition body, fundamentally different from a physically adsorbed heterogeneous interface with an externally coated core-shell structure.

[0025] In some embodiments, based on the total mass of the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment as 100%, the mass percentage of α-Zn2SiO4 crystal nuclei is 30-50%, and the mass percentage of the residual silicon-rich glass phase SiO2 shell is 50-70%.

[0026] When the proportion of crystal nuclei is too high, the crystallinity is too high, resulting in an excessively thin or discontinuous residual silica-rich glassy shell that cannot completely encapsulate the crystal nuclei. This leads to a decrease in UV radiation resistance, and some crystal grains may come into direct contact under high crystallinity, losing the "inherent isolation" advantage against agglomeration. Conversely, when the proportion of crystal nuclei is too low, the crystallinity is too low, meaning insufficient scattering core content in the filler, a reduced number of effective scattering centers per unit volume, decreased solar reflectivity, and increased solar absorptivity. At the same time, a high residual glassy phase content and an excessively thick shell dilute the volume fraction of high-refractive-index scattering cores, which is also detrimental to achieving ultra-low solar absorptivity. A mass ratio of 30-50% crystal nuclei and 50-70% shell represents the optimal balance between scattering efficiency and radiation protection: it provides sufficient high-refractive-index scattering cores to achieve ultra-low solar absorptivity while ensuring that each crystal grain is completely and continuously encapsulated by a silica-rich glassy shell to provide reliable UV radiation protection.

[0027] In some embodiments, the solar absorptivity α of the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment is... s It can be ≤0.08, with α being preferred. s ≤0.06, the change in solar absorptivity Δα after 5000 ESH vacuum ultraviolet irradiation s ≤+0.05, preferably Δα s ≤+0.02.

[0028] The following is an exemplary description of the preparation method of the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment. The preparation method may include the following steps: (1) Mix the glass raw materials and melt them at high temperature, anneal them, and cool them to obtain the precursor glass; (2) The precursor glass is heated to undergo nucleation and crystallization treatments to obtain a microcrystalline glass block; (3) The microcrystalline glass block is crushed to obtain the in-situ grown Zn2SiO4@SiO2 low absorption microcrystalline glass pigment.

[0029] In some embodiments, in step (1), the glass raw material may be composed of the following components by molar percentage: ZnO 34-40%, SiO2 42-50%, Al2O3 8-12%, K2O 4-8%, nucleating agent 0.2-0.5%, and the sum of the molar percentages of each component is 100%; preferably, the nucleating agent may include at least one of ZrO2, SnO2, and Y2O3, and the molar ratio of ZnO to SiO2 may be 0.72-0.85.

[0030] This invention limits the composition of glass raw materials to the above-mentioned range, based on the synergistic optimization of the triple constraints of target crystal phase purity, core-shell structure integrity, and glass-forming ability. When the ZnO content is too low, Zn... 2+ Insufficient supply prevents the full analysis of the target crystalline phase; excessively high content results in insufficient residual glass phase, incomplete shell, and reduced glass-forming ability. Insufficient SiO2 content leads to insufficient glass network formation, poor glass-forming stability, and insufficient shell thickness; excessively high content results in a low ZnO / SiO2 ratio and insufficient crystallinity. Insufficient Al2O3 content leads to insufficient glass chemical stability and mechanical strength; excessively high content leads to insufficient Al2O3 content. 3+ With Zn 2+ Competition generates ZnAl2O4 impurity phases; when the K2O content is too low, the fluxing effect is insufficient and phase separation is incomplete; when it is too high, the network deagglomerates excessively, which may lead to the precipitation of impurity phases such as KAlSi2O6 or glass devitrification; when the nucleating agent content is too low, there are insufficient nucleation sites; when it is too high, it may introduce impurity phases or affect optical properties; when the ZnO / SiO2 molar ratio is too low, the crystallinity is insufficient; when it is too high, the residual glass phase is insufficient, resulting in discontinuous shell layers, and ZnO impurity phases may precipitate. The above ranges together ensure that α-Zn2SiO4 is the only main crystalline phase and is covered by a complete and continuous in-situ core-shell structure.

[0031] In some implementations, in step (1), the mixing method can be ball milling at a speed of 200-500 r / min for 2-4 hours.

[0032] In some embodiments, in step (1), the temperature of the high-temperature melting can be 1460-1500°C, and the time can be 1-3 hours.

[0033] When the melting temperature is too low, the glass batch melts insufficiently, resulting in high melt viscosity and inadequate component diffusion. This leads to uneven composition and unmelted particles, causing uneven nucleation and an incomplete core-shell structure during subsequent crystallization. When the melting temperature is too high, the volatilization of ZnO and K2O intensifies, causing the actual glass composition to deviate from the designed ratio, affecting the purity of the crystalline phase and the thickness of the shell. Insufficient holding time results in incomplete melt clarification, leaving bubbles and streaks, and poor glass uniformity. Excessive holding time similarly exacerbates ZnO and K2O volatilization and intensifies crucible erosion, potentially introducing Al2O3 contamination. Holding at 1460-1500℃ for 1-3 hours provides an optimized window that balances sufficient melting and homogenization with inhibiting component volatilization.

[0034] In some implementations, in step (1), the annealing temperature can be 500-550°C and the time can be 1-2 hours.

[0035] When the annealing temperature is too low, atomic migration is insufficient, and the internal thermal stress of the glass cannot be fully released, making it prone to uneven crystallization or even cracking during subsequent crystallization. When the annealing temperature is too high, approaching the glass transition temperature, uncontrollable nucleation or phase separation may occur prematurely, interfering with subsequent controllable crystallization. A short holding time may result in insufficient stress relief; conversely, a long holding time, especially at temperatures close to the glass transition temperature (Tg), also carries the risk of uncontrollable crystallization. Holding at 500-550℃ for 1-2 hours is the optimal window to balance sufficient stress relief with preventing premature crystallization.

[0036] In some embodiments, in step (2), the temperature of the nucleation treatment can be 30-60°C higher than the initial temperature Tg of the precursor glass during the glass transition process (determined by differential scanning calorimetry DSC test), the time can be 1-2 hours, and the heating rate can be 3-8°C / min; preferably, the Tg is 480-648°C.

[0037] Nucleation is the first stage of the two-step heat treatment process for glass-ceramics. Its purpose is to induce local relaxation of the glass network structure within a temperature range slightly above the glass transition temperature (Tg), allowing atoms to begin to possess limited migration capabilities. During this process, a large number of uniformly dispersed crystal nuclei are formed within the glass matrix, providing high-density preferential nucleation sites for the subsequent crystallization stage. The nucleation stage determines the number density and size uniformity of the subsequent crystals and is a crucial step in ensuring the uniform precipitation of α-Zn₂SiO₄ grains and the spontaneous formation of a complete core-shell structure.

[0038] In some cases, excessively low nucleation temperatures result in a rigid glass network with insufficient atomic migration and weak nucleation driving force, making it impossible to form a sufficiently dense nucleus population within the preset time. Excessively high nucleation temperatures lead to an excessively rapid nucleation rate, resulting in too many nuclei forming instantaneously, limiting grain growth space, and ultimately causing smaller grain sizes. This leads to entry into the Rayleigh scattering region, decreased scattering efficiency, and increased solar absorptivity. Insufficient nucleation time results in an inadequate and unevenly distributed number of nuclei, leaving some areas without nuclei for subsequent crystallization, leading to a wide grain size distribution and poor core-shell structure uniformity. Excessively long time results in excessively high nucleus density, limiting crystal growth space and ultimately causing smaller grain sizes. Conversely, excessively slow heating rates lead to prolonged stays in the low-temperature region, potentially causing uncontrollable slow nucleation before reaching the target nucleation temperature, affecting the consistency of nucleus density. Excessively rapid heating rates result in significant temperature overshoot, potentially causing the actual nucleation temperature to exceed the preset range, deviating from the optimal nucleation window, and leading to uneven nucleus density and distribution.

[0039] In some embodiments, in step (2), the temperature of the crystallization treatment can be 20-50°C lower than the peak temperature Tp of the precursor glass crystallization exothermic peak (determined by differential scanning calorimetry DSC test), the time can be 2-8 hours, and the heating rate can be 2-5°C / min; preferably, the Tp is 640-836°C.

[0040] Crystallization is the second stage of the two-step heat treatment process for glass-ceramics. Its purpose is to allow the numerous crystal nuclei formed during the nucleation stage to grow synchronously at a controlled rate within a temperature range slightly below the peak exothermic temperature Tp, precipitating the target crystalline phase α-Zn₂SiO₄. During crystal growth, the crystal nuclei consume Zn from the surrounding glass matrix. 2+ This process allows SiO2 to accumulate automatically in the crystal front region, spontaneously forming a residual silicon-rich glassy phase shell that tightly encapsulates the grain surface. The crystallization stage directly determines the final size and morphology of the grains, as well as the integrity and shell thickness of the in-situ core-shell structure.

[0041] In some cases, excessively low crystallization temperatures result in slow atomic diffusion rates, insufficient driving force for crystal growth, smaller grain sizes, and entry into the Rayleigh scattering region, leading to decreased scattering efficiency and increased solar absorptivity. Conversely, excessively high crystallization temperatures, approaching or exceeding the crystallization peak temperature, result in excessively rapid crystal growth, larger grain sizes, and overstretching and thinning of the shell, potentially causing fractures and compromising the integrity of the core-shell structure. Excessively high temperatures can also lead to the precipitation of ZnO or ZnAl2O4 impurity phases, reducing the purity of the main crystalline phase and optical properties. Insufficient crystallization times result in incomplete crystal growth, low crystallinity, and smaller grain sizes. Excessively long crystallization times may lead to abnormal grain growth or Ostwald ripening, causing small grains to dissolve and large grains to grow further, resulting in a wider size distribution. If the heating rate is too slow, the transition time from the nucleation temperature to the crystallization temperature is too long, causing some crystal nuclei to grow slowly and prematurely in the low-temperature region, resulting in uneven grain size and poor core-shell structure uniformity. If the heating rate is too fast, there will be significant temperature overshoot, and the actual crystallization temperature may exceed the preset window, leading to uncontrolled crystal growth rate, larger and unevenly distributed grain size. This temperature window is based on the dual optimization of crystal growth kinetics and controllability of core-shell structure: within this range, the crystal grows at an appropriate rate, and the grain size can be controlled within the optimal Mie scattering range of 50-200 nm, while Zn... 2+ The full diffusion allows the component differentiation effect at the crystal front to be fully utilized, forming a complete and continuous residual silicon-rich glass phase shell, which takes into account both ultra-low solar absorptivity and excellent radiation resistance.

[0042] In some embodiments, in step (3), the particle size D50 of the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment can be <5 μm.

[0043] The microcrystalline glass pigment provided by this invention is prepared by controlled crystallization of precursor glass. The main crystalline phase α-Zn₂SiO₄ precipitates in situ within the glass matrix, while a residual silicon-rich glassy phase shell spontaneously forms around the grains. A continuous glassy phase transition interface exists between the shell and the crystal nucleus. When α-Zn₂SiO₄ precipitates from the silicon-rich precursor glass, the grains consume ZnO, causing SiO₂ to accumulate automatically in the crystal front region, forming a residual silicon-rich glassy phase. This provides the thermodynamic and kinetic conditions for the spontaneous formation of a core-shell structure. This in-situ generated residual silicon-rich glassy phase can spontaneously form a shell on the grain surface, and the shell and crystal nucleus form a chemically bonded continuous transition interface, fundamentally avoiding the physical abrupt boundary problems associated with external coatings. This unique "in-situ growth core-shell structure" provides a novel technical approach for developing new thermally controlled pigments that possess both ultra-low solar absorptivity and excellent radiation resistance.

[0044] In summary, this invention has the following outstanding improvements over the prior art: 1. Breaking through the traditional coating process to achieve an in-situ core-shell structure with continuous glass phase transition inside the glass matrix: This invention abandons the two-step route of "first synthesizing Zn2SiO4 powder and then adding an outer SiO2 shell" in the prior art. For the first time, it applies the component differentiation effect in the in-situ crystallization process of microcrystalline glass to the construction of core-shell structured thermally controlled pigments. Through an integrated process of melting-quenching-controllable crystallization, α-Zn2SiO4 crystals are directly precipitated in-situ from the glass matrix. During crystal growth, Zn... 2+ A large portion of the Zn is consumed in the formation and growth of the crystal nucleus, and the glass matrix surrounding the crystal is affected by the Zn. 2+The extracted residual silicon-rich glass phase automatically forms, tightly encapsulating the surface of α-Zn2SiO4 grains, spontaneously forming a natural core-shell structure. The shell layer and the crystal nucleus form a continuous glass phase transition interface with chemical bonding, without any physical abrupt boundary. This is the fundamental difference between in-situ grown core-shell structures and externally coated core-shell structures. Thus, compared to existing external coating schemes, the in-situ core-shell structure of this invention exhibits stronger bonding between the shell layer and the crystal nucleus, and withstands thermal cycling in spacecraft (-150~150°C). Under ℃) and mechanical grinding stress, the shell will not peel off or crack. The shell completely covers each grain, without micron-level pores and uncovered areas, which fundamentally solves the problems of weak bonding force of the external coating interface, incomplete coating and poor batch consistency; (2) The formation of the core-shell structure and the precipitation of the crystal phase are combined in the same heat treatment process and completed simultaneously. The process is simple and controllable, and the microstructure is highly adjustable: The preparation of existing core-shell structure Zn2SiO4@SiO2 pigments requires the process of “synthesizing Zn2SiO4 powder → wetting”. The traditional process involves multiple steps such as coating, drying, and high-temperature heat treatment, with synthesis and coating performed separately, resulting in a long process cycle. Furthermore, the shell thickness and uniformity depend on sensitive parameters such as TEOS concentration and pH value. This invention combines the formation of the core-shell structure with the precipitation of the α-Zn₂SiO₄ crystal phase in a single, controllable crystallization heat treatment process. The construction of the core-shell structure requires no additional coating steps, and the shell thickness and grain size can be determined by the precursor glass composition (ZnO / SiO₂ molar ratio) and the heat treatment regime (nucleation temperature, crystal growth rate, and grain size). (3) The addition of an amorphous SiO2 layer is replaced by a silicon-rich glass phase shell. The thermal expansion coefficients of the shell and the crystal nucleus are matched, and the shell itself is a continuation of the glass matrix, with extremely low interface stress. In the existing addition coating scheme, there is a difference in the thermal expansion coefficients between the SiO2 shell and the Zn2SiO4 crystal nucleus. Under thermal cycling, the interface stress is extremely low. Microcracks are easily generated. In this invention, the residual silicon-rich glass phase shell is a natural continuation of the precursor glass matrix after crystallization. It is a continuous transition interface with chemical bonding between the shell and the crystal nucleus. There is no heterogeneous material interface. The difference in thermal expansion coefficient between the shell and the crystal nucleus is much smaller than the difference between the added SiO2 layer and Zn2SiO4. The interfacial thermal stress is extremely low. Thus, under the thermal cycling conditions experienced by the spacecraft in orbit, the in-situ core-shell structure of this invention will not generate interfacial microcracks due to thermal mismatch. The long-term structural stability is significantly better than that of the external coating scheme.

[0045] Therefore, the core-shell structure provided by this invention suppresses the formation and accumulation of irradiation-induced color centers at the intrinsic material level, solving the technical bottleneck of severe irradiation degradation in existing Zn2SiO4 pigments, and is suitable for high-performance white thermal control coatings for spacecraft. In the aforementioned applications, the microcrystalline glass pigment can be used as a carrier with inorganic binders or organic resins to prepare white inorganic pigments for low solar absorptivity thermal control coatings.

[0046] The following examples further illustrate the present invention in detail. It should also be understood that the following examples are only for further explanation of the present invention and should not be construed as limiting the scope of protection of the present invention. Any non-essential improvements and adjustments made by those skilled in the art based on the above description of the present invention are within the scope of protection of the present invention. The specific process parameters, etc., in the following examples are merely examples within a suitable range; that is, those skilled in the art can make appropriate selections within the range based on the description herein, and are not intended to be limited to the specific values ​​in the examples below. Unless otherwise specified, the technical means used in the examples are conventional means well known to those skilled in the art.

[0047] Unless otherwise specified, the raw material information is as follows: zinc oxide (ZnO, purity ≥99.9%), silicon dioxide (SiO2, purity ≥99.9%), aluminum hydroxide (Al(OH)3, purity ≥99.9%), anhydrous potassium carbonate (K2CO3, purity ≥99.9%), zirconium oxide (ZrO2, purity ≥99.9%, nucleating agent), and tin oxide (SnO2, purity ≥99.9%, nucleating agent).

[0048] Example 1 The method for preparing in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment provided in this embodiment includes the following steps: (1) Weigh the raw materials zinc oxide (ZnO), silicon dioxide (SiO2), aluminum hydroxide (Al(OH)3), anhydrous potassium carbonate (K2CO3), and zirconium oxide (ZrO2) according to the molar percentage of ZnO 37.5%, SiO2 47.5%, Al2O3 9.5%, K2O 5.0%, ZrO2 0.5%, and ZnO / SiO2 molar ratio of 0.789. Add all raw materials to a zirconium oxide ball mill jar with a ball-to-material ratio of 3:1 and dry ball mill at 250 rpm for 4 hours. Then pass the mixture through a 100-mesh sieve. Place the mixture into a corundum crucible (pre-fired at 1100℃ for 2 hours to remove gas) and place it in a high-temperature heating furnace. Melt the mixture according to the following process: room temperature → 300℃ (5℃ / min, holding for 30 min) → 900℃ (8℃ / min, holding for 60 min). (min) → 1480℃ (10℃ / min, hold for 2.5 hours, stir with a quartz rod every 30 minutes), after the melt is cooled to 1380℃, it is poured into a graphite mold preheated to 300℃, and then quickly transferred to a muffle furnace at 520℃ for annealing for 2 hours. It is then slowly cooled to room temperature with the furnace at a rate of 2℃ / min to obtain the precursor glass. (2) The precursor glass powder was subjected to DSC analysis at 10℃ / min (room temperature → 1200℃, air atmosphere). The DSC curve showed Tg = 640℃ and Tp = 836℃. The nucleation temperature was set to Tg + 55℃ = 695℃ and the crystallization temperature was set to Tp - 36℃ = 810℃. Nucleation stage: the temperature was increased to 695℃ at 5℃ / min and held for 1.5 hours to promote uniform nucleation. Crystallization stage: the temperature was increased to 810℃ at 3℃ / min and held for 4 hours to allow the crystal nuclei to grow fully. The powder was then slowly cooled to room temperature in the furnace to obtain a microcrystalline glass block. (3) The crystallized block was coarsely crushed to <5 mm, and wet-milled in a zirconium ball mill jar for 6 hours (350 rpm, ball-to-material ratio 5:1) with anhydrous ethanol as medium, and dried at 80°C for 12 hours to obtain the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment.

[0049] Example 2 The preparation method of the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment provided in this embodiment is the same as in Example 1, with the main difference being: In step (1), the nucleating agent is SnO2; In step (2), the DSC test showed Tg=632℃, Tp=830℃, the nucleation temperature was set to 685℃, and the crystallization temperature was set to 800℃.

[0050] Example 3 The preparation method of the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment provided in this embodiment is the same as in Example 1, with the main difference being: In step (1), the ZnO / SiO2 molar ratio of the precursor glass is adjusted from 0.789 to 0.828, and the composition is adjusted to: ZnO3 8.5%, SiO2 46.5%, Al2O3 9.5%, K2O 5.0%, ZrO2 0.5%; In step (2), the DSC test showed Tg=648℃, Tp=821℃, the nucleation temperature was set to 705℃, and the crystallization temperature was set to 790℃.

[0051] Comparative Example 1 The pigment preparation method provided in this comparative example includes the following steps: The raw materials were weighed with ZnO and SiO2 in a molar ratio of 2:1, and ball-milled with anhydrous ethanol for 4 hours. After drying, they were calcined in solid phase at 1050℃ for 3 hours and ball-milled to a D50 of about 2μm.

[0052] Comparative Example 2 The preparation method of the pigment provided in this comparative example is the same as that in Example 1, with the main difference being: In step (1), the ZnO / SiO2 molar ratio of the precursor glass is adjusted from 0.789 to 1.0 (Zn2SiO4 stoichiometry), and the composition is adjusted to: ZnO 44.5%, SiO2 44.5%, Al2O3 6.0%, K2O 4.5%, ZrO2 0.5%.

[0053] Table 1 below compares the performance data of the pigments prepared in Examples 1-3 and Comparative Examples 1-2: .

[0054] Figure 1 The figures show the X-ray diffraction patterns of the pigments prepared in Examples 1 and 2. As can be seen from the figures, all diffraction peaks of the pigments obtained in Examples 1 (ZrO2 nucleating agent) and 2 (SnO2 nucleating agent) correspond one-to-one with the α-Zn2SiO4 standard card (JCPDS 37-1485), with no ZnO, ZnAl2O4, or other impurity phase diffraction peaks. This indicates that both nucleating agents can induce the precipitation of pure α-Zn2SiO4 with good crystallinity.

[0055] Figure 2The image shows the DSC curve of the precursor glass in Example 1. As can be seen from the curve, an endothermic step appears near 640℃, corresponding to the glass transition temperature Tg, indicating that the glass network begins to transition from rigid to viscoelastic at this temperature. A sharp exothermic peak appears near 836℃, corresponding to the crystallization peak temperature Tp of α-Zn₂SiO₄. The peak shape is symmetrical and there are no impurity peaks, indicating that the crystallization process is a single phase transition without the formation of intermediate or impurity phases. Based on the DSC data, the nucleation temperature was set to Tg + 55℃ = 695℃, and the crystallization temperature was set to Tp - 30℃ = 810℃, corresponding to the optimal temperature windows for sufficient relaxation of the glass network to promote uniform nucleation and synchronous growth of crystals at a controlled rate, respectively.

[0056] Figure 3 The image shows the UV-Vis-NIR diffuse reflectance spectrum of the pigment prepared in Example 1. As can be seen from the figure, the pigment exhibits extremely high diffuse reflectance across the entire main energy band of the solar spectrum (250-2500 nm), with a reflectance exceeding 90% in the UV band (250-380 nm). Based on the ASTM E490 standard solar spectrum weighted integral, the solar absorptivity α... s It is only 0.05.

[0057] Figure 4 The graph shows the change in solar absorptivity of the pigments prepared in Examples 1, 1, and 2 before and after ultraviolet irradiation. As can be seen from the graph, the α-absorbance ratio of the pigments prepared in Example 1 before irradiation... s It was only 0.05 at first, but rose to 0.09 after irradiation, Δα s The value was +0.04, indicating the smallest degradation; Comparative Example 1 (solid phase method) α before irradiation s It was 0.09 at first, and rose to 0.16 after irradiation, Δα s The degradation was as high as +0.07, approximately 1.8 times that of Example 1. This is because the Zn2SiO4 particles obtained by the solid-state method have exposed surfaces without a core-shell protective layer, and surface defect sites continuously generate color centers as precursor sites under ultraviolet irradiation; Comparative Example 2 (ZnO / SiO2=1.0) before irradiation α s It was 0.10, and after irradiation it rose to 0.18, Δα s The degradation was as high as +0.08, indicating the most severe degradation. In addition to the lack of a complete core-shell structure, the excess ZnO existed as an independent impurity phase. The irradiation degradation of ZnO itself (zinc vacancy ionization produces a 425 nm absorption band) further aggravated the performance deterioration. This comparison directly proves that the in-situ core-shell structure under silicon-rich composition (ZnO / SiO2=0.72-0.85) is the key to endowing Zn2SiO4 pigments with ultra-low solar absorptivity and excellent radiation resistance.

[0058] Figure 5The image shows a SEM image of the pigment prepared in Example 1. As can be seen from the image, a large number of uniformly dispersed near-spherical grains are visible in the cross-section of the particles, with a grain size of about 80-120 nm. The grains are isolated by the residual glass phase, and there is no direct contact or sintering of the grains, which indirectly confirms the core-shell structure characteristics of the grains being completely covered by the glass phase.

[0059] Although the present invention has been described in detail through the preferred embodiments above, it should be understood that the above description should not be considered as a limitation of the present invention. Various modifications and substitutions to the present invention will be apparent to those skilled in the art after reading the above description. Therefore, the scope of protection of the present invention should be defined by the appended claims.

Claims

1. An in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment, characterized in that, The in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment comprises: α-Zn2SiO4 crystal nuclei, and a residual silicon-rich glass phase SiO2 shell coating the surface of the α-Zn2SiO4 crystal nuclei; wherein, there is a continuous glass phase transition interface between the α-Zn2SiO4 crystal nuclei and the silicon-rich glass phase SiO2 shell.

2. The in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment according to claim 1, characterized in that, The α-Zn2SiO4 crystal nuclei have a particle size of 50-200 nm, and the residual silicon-rich glassy phase SiO2 shell has a thickness of 100-300 nm.

3. The in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment according to claim 1 or 2, characterized in that, Based on the total mass of the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment as 100%, the mass ratio of α-Zn2SiO4 crystal nuclei is 30-50%, and the mass ratio of the residual silicon-rich glass phase SiO2 shell is 50-70%.

4. A method for preparing an in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment according to any one of claims 1-3, characterized in that, The preparation method includes the following steps: (1) Mix the glass raw materials and melt them at high temperature, anneal them, and cool them to obtain the precursor glass; (2) The precursor glass is heated to undergo nucleation and crystallization treatments to obtain a microcrystalline glass block; (3) The microcrystalline glass block is crushed to obtain the in-situ grown Zn2SiO4@SiO2 low absorption microcrystalline glass pigment.

5. The preparation method according to claim 4, characterized in that, In step (1), the glass raw material is composed of the following components by molar percentage: ZnO 34-40%, SiO2 42-50%, Al2O3 8-12%, K2O 4-8%, nucleating agent 0.2-0.5%, and the sum of the molar percentages of each component is 100%. Preferably, the nucleating agent includes at least one of ZrO2, SnO2, and Y2O3, and the molar ratio of ZnO to SiO2 is 0.72-0.

85.

6. The preparation method according to claim 4 or 5, characterized in that, In step (1), the high-temperature melting temperature is 1460-1500℃ and the time is 1-3 hours; Preferably, the annealing temperature is 500-550℃ and the time is 1-2 hours.

7. The preparation method according to any one of claims 4-6, characterized in that, In step (2), the temperature of the nucleation treatment is 30-60°C higher than the initial temperature Tg of the precursor glass during the glass transition process, the time is 1-2 hours, and the heating rate is 3-8°C / min; preferably, the Tg is 480-648°C.

8. The preparation method according to any one of claims 4-7, characterized in that, In step (2), the temperature of the crystallization treatment is 20-50°C lower than the peak temperature Tp of the exothermic peak of the precursor glass crystallization, the time is 2-8 hours, and the heating rate is 2-5°C / min; preferably, the Tp is 640-836°C.

9. The preparation method according to any one of claims 4-8, characterized in that, In step (3), the particle size D50 of the in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment is <5 μm.

10. The application of an in-situ grown Zn2SiO4@SiO2 low-absorption microcrystalline glass pigment according to any one of claims 1-3 in a high-performance white thermal control coating for spacecraft.

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