Deep sea equipment is an important guarantee for the exploration, development, and protection of deep-sea resources, including various types of deep-sea submersibles, landers, gliders, and their auxiliary equipment. In the deep-sea equipment above, the transparent protective cover for observation windows, cameras, and lighting must use high-strength transparent materials. The main transparent materials currently used are ultra-thick organic glass (observation window), inorganic glass, and sapphire (single crystal aluminum oxide). Their application in deep-sea equipment is possible due to the superior transparency of yttrium aluminum garnet (YAG) transparent ceramics like single crystals, as well as the characteristics of optical isotropy, high strength, high thermal conductivity, and stable physical and chemical properties. At present, there are many kinds of transparent ceramics, among which YAG transparent ceramics are the most representative. The existing researches on YAG transparent ceramics mainly focus on doping rare-earth elements to achieve optical functional applications, such as laser output, white LED luminescence, etc.. There are a few reports on the research of large-sized and complex shaped YAG transparent ceramics. In this paper, large-sized and complex shaped YAG transparent ceramics were prepared by a solid-state reaction sintering method, which could be applied to deep-sea equipment at 10 000 meter.

High-purity powders of α-Al₂O₃ (99.99% purity, Taimei Co., Japan), Y₂O₃ (self-synthesized by co-precipitation method) were used as starting materials. The starting powders were weighed according to a stoichiometric ratio of YAG, and then mixed with 0.5% tetraethyl orthosilicate (TEOS, 99.99%) and 0.05% MgO powder (analytical pure) as a sintering additive in ethanol by ball milling for 12 h. The mixtures were dried in an oven at 60 ℃ for 24 h and then sieved through 200-mesh screen. A portion of the dried powder was calcined at 900-1 500 ℃, and another part of the powder was pressed into circular discs with a diameter of 20 mm using a steel mold at 30 MPa, and then subjected to cold isostatic pressing at 200 MPa to obtain ceramic billets. The green body pellets were then sintered in a vacuum furnace with tungsten meshes as the heating elements under 10^-3 Pa at 1 200-1 800 ℃ for 2-20 h. After sintering in vacuum, the pellets were further annealed in air at 1 400 ℃ for 10 h. Finally, the both surfaces of the as-prepared ceramic samples were mirror-polished with different grade of diamond slurries.

The phase compositions of all the samples were qualitatively identified by a model D/Max-2550V X-ray diffractometer (Rigaku Co., Japan). The optical transmittances of all the samples of ceramics were measured by a model Carry 5000 UV-VIS-NIR spectrometer (Varian Co., USA). The surface microstructures of each group of samples after hot corrosion were characterized by a model JSM-6360 scanning electron microscope (JEOL Co., Japan).

YAG transparent ceramics were fabricated by a solid-phase reaction sintering method. The phase evolution and microstructural changes during sintering were investigated. At high temperatures, Al₂O₃ and Y₂O₃ react to form intermediate phases YAM and YAP, eventually forming YAG.

The effects of sintering temperature and holding time on the optical transmittance of YAG transparent ceramics were investigated. As the sintering temperature increases, the pores in YAG ceramics continue to discharge and the grains continue to grow. The optical transmittance firstly increases and then decreases, reaching its maximum value at 1 780 ℃. As the holding time prolongs, the transmittance of YAG ceramics gradually increases, but after more than 6 h, extending the holding time does not have a significant effect on improving the transmittance.

The YAG ceramic dome is formed using a specially designed rubber mold and a stainless steel hemispherical mold by a cold isostatically press. The YAG ceramic dome is then vacuum sintered in the optimized process. The optical window under 127 MPa was analyzed by finite element analysis (i.e., a software named ANSYS). The analysis results show that the maximum compressive stress of the 120 mm diameter dome is 773.7 MPa and the maximum tensile stress is 21.9 MPa, both of which are less than the maximum allowable stress of YAG transparent ceramic material (i.e., approximately 1 000 MPa). The YAG transparent ceramic dome and titanium alloy cylinder are sealed and packaged. The transparent ceramic dome is undamaged without cracking or leakage at 127 MPa for 1.5 h. The high-power deep-sea LED lighting and deep-sea camera, using YAG transparent ceramics as protective covers, are installed on a deep-sea equipment "Canghai" lander. In November 2020, the lander sank to the depths of 10 000 meters in the Mariana Trench (the deepest part of the oceans) for several times, and a large amount of valuable video data were recorded.

Highly transparent YAG transparent ceramics were fabricated in vacuum via solid-state reaction sintering. Under high temperature conditions, Al₂O₃ and Y₂O₃ reacted to form intermediate phases YAM and YAP in sequence, and finally formed YAG at 1 500 ℃. As the sintering temperature increased, the pores in YAG ceramics continued to discharge and the grains continued to grow. The optical transmittance firstly increased and then decreased, reaching its maximum value at 1 780 ℃. As the holding time prolonged, the transmittance of YAG ceramics gradually increased. The transmittance of YAG ceramic samples sintered in vacuum at 1 780 ℃ for 20 h was 84.7% at 1 100 nm and 82.8% at 400 nm, respectively. The YAG transparent ceramic dome was prepared via optimization. It could be applied to high-power LED lighting and camera protection covers in the deep-sea equipment at 10 000 meters.

This paper prepared a series of Eu³⁺ doped Gd₂Zr₂O₇ transparent ceramics by a vacuum sintering method, and charaterized the density and optical transmittance of cubic phase Gd₂Zr₂O₇. The effects of doping ions and their concentration on the crystal structure, optical transmittance, fluorescence performance, and scintillation performance were investigated. The results show that the prepared transparent ceramic samples have a good optical transmittance. The spectral characterization reveals that 20% Eu³⁺ doped Gd₂Zr₂O₇ has optimum fluorescence and irradiated luminescence intensity, and the most intense emission peak is located at 630 nm. The X-ray imaging shows that Gd_1.8Eu_0.2Zr₂O₇ transparent ceramics have an X-ray imaging resolution of up to 11 lp·mm^-1, having a potential for the imaging applications.

High purity gadolinium oxide, europium oxide and zirconia (purity 4N) as raw materials were mixed according to a stoichiometric ratio (Gd³⁺,Eu³⁺)∶Zr⁴⁺ of 1∶1. The mixed raw materials were ground in anhydrous ethanol and polyethylene glycol (PEG) in a planetary mill with zirconia grinding chamber and ballsat a ball∶material mass ratio of 27∶1 for 72 h. Afterwards, the ground mixture slurry was washed with anhydrous ethanol after drying. Gd_2-xEu_xZr₂O₇ could be obtained via sintering in a box furnace at 1 400 ℃ according to the reaction: Gd_2-XEu_XZr₂O₇

The ceramic blank was formed via dry pressing and cold isostatic pressing, sintered in a vacuum tungsten furnace at 1 800 ℃ for 10 h, and then annealed in a box furnace in air at 1 250 ℃ for 6 h. Finally, a transparent ceramic with 1 mm thickness was obtained after double-sided polishing.

A ceramic powder prepared by ball grinding was sintered in a vacuum tungsten wire furnace to prepare transparent ceramics with a high light transmission rate. The XRD results show that when 20% Eu³⁺ is added to Gd₂Zr₂O₇, the crystal structure of Gd_1.8Eu_0.2Zr₂O₇ appears some superlattice diffraction peaks (331) of pyrochlore. The crystal structure changes from defective fluorite to pyrochlore. The optical transmittance spectra indicate that the optical transmittance decreases from 76% to 65% (@1 500 nm) when the doping content of Eu³⁺ increases from 0 to 28%. There exist 6 distinct excitation peaks between 230 nm and 500 nm. A wide excitation peak at 267 nm is attributed to the charge transfer band of O^2-—Eu³⁺. In addition, five excitation peaks at 362, 382, 393, 414 nm and 465 nm correspond to 4f-4f energy level transitions ⁷F₀→⁵D₄, ⁵G_J, ⁵L₆, 5D₃ and ⁵D₂ in Eu³⁺, respectively. Among six excitation peaks, a wide excitation peak at 267 nm has the most intense intensity. All the excitation peak intensities increase when Eu³⁺ concentration increases from 4% to 20%. When Eu³⁺ doping amount is 20%, the maximum excitation peak intensity appears. The further characterization of the scintillation properties of ceramic samples indicates that the X-ray irradiation intensity is a maximum value when the doping amount is 20%. At 630 nm, the radiation luminescence intensity value is 1.1×10⁴, and the material has an X-ray absorption capacity equivalent to that of BGO crystals, and also has a good anti-radiation damage ability. Finally, X-ray imaging tests on Gd_1.8Eu_0.2Zr₂O₇ ceramic samples of different materials show that the Gd_1.8Eu_0.2Zr₂O₇ ceramic sample can clearly distinguish metal and plastic materials, and has a limit spatial resolution of 11 1p·mm^-1, indicating that the modified material is a potential X-ray detection material.

A series of Eu³⁺ doped transparent ceramics were prepared by a vacuum sintering method. The crystal structure of Gd_2-xEu_xZr₂O₇ was gradually changed from defective fluorite to pyrochlore via adjusting the doping amount of Eu³⁺. At 630 nm, the Gd_2-xEu_xZr₂O₇ transparent ceramic sample showed the optimum fluorescence emission and X-ray irradiation intensity, and the material had a limit spatial resolution of 11 1p·mm^-1, which could clearly distinguish metal and plastic materials, confirming an application potential of the A₂B₂O₇ series of scintillating transparent ceramics X-ray imaging.

Lutetium aluminum garnet (Lu₃Al₅O₁₂, LuAG) is investigated as a classical scintillation material. LuAG can be fabricated in the form of single crystals, optical ceramics, thin films and other morphologies to meet different scintillation application demands. Cerium-doped LuAG (Ce:LuAG) ceramics can be used as a promising candidate for detection material in shashlik calorimeter for high-energy physics (HEP) experiments due to their high density, fast decay, efficient scintillation and relatively low cost. For HEP experiments with a high event rate, however, a slow scintillation component of Ce:LuAG ceramics leads to a pileup effect that needs to be diminished as much as possible. Divalent, optically inactive, Ca²⁺ co-doping is an effective strategy for suppressing slow scintillation component according to the defect engineering theory. In addition, the coprecipitated nano-powder can be also used as a raw material of ceramics, and improving the consistency of structure, composition and luminescence within the ceramics is also an important strategy to optimize the performance of Ce:LuAG scintillation ceramics. In this paper, Ce,Ca:LuAG ceramics with different Ce concentrations were prepared with the coprecipitation nano-powders. The effect of activator ion concentration on the decay behavior, light yield (LY) at different shaping times and afterglow intensity of Ce,Ca:LuAG ceramics were investigated.

A series of Ce,Ca:LuAG ceramics with the chemical composition of (Ce_xCa_{0.001 5}Lu_{0.998 5-x%})₃Al₅O₁₂; x=0.1, 0.2, 0.3, 0.4, 0.5 were prepared with coprecipitated nano-powders through vacuum pre-sintering and hot isostatic pressed (HIP) post-treatment. In the process of nano-powder preparation, high purity Lu₂O₃ (99.995%), CeO₂ (99.999%), and CaCO₃ powders were dissolved in nitric acid solution, and Al(NO₃)₃·9H₂O (AR) powders were dissolved in deionized water to obtain various nitrate solutions. The metal nitrate solutions were mixed and further diluted to a specific concentration. Ammonium hydrogen carbonate (AHC, 99%) was dissolved in de-ionized water as a precipitant solution, and ammonium sulfate (99%) was added into the precipitant solution as a dispersant. The mixed metal nitrate solution was dripped into the AHC solution under mild agitation at room temperature (RT). The resulting suspension was aged for 1 h and then washed with deionized water and ethanol. The precursor was dried, sieved and calcined to obtain Ce,Ca:LuAG nano-powders with a pure garnet phase. The powders were uniaxially dry-pressed into pellets at 20 MPa and then cold-isostatically pressed at 250 MPa. The green bodies were vacuum pre-sintered at 1 825 ℃ for 10 h, and then HIP in argon atmosphere under 176 MPa at 1 600 ℃ for 3 h. The fabricated Ce,Ca:LuAG ceramics were annealed in air at 1 450 ℃ for 10 h to remove oxygen vacancies introduced via sintering in a reducing atmosphere. The final Ce,Ca:LuAG ceramics were polished on the both sides to the thickness of 1 mm for the coming characterization.

The XRD patterns of all the ceramic samples are in reasonable agreement with the standard XRD patterns of LuAG (PDF 73-1368), indicating that the main phase of the prepared ceramics is a LuAG garnet phase. The lattice parameters of Ce,Ca:LuAG ceramics with different Ce concentrations increase gradually with increasing doping concentration. The presence of Ce³⁺ can act as a sintering aid to accelerate the migration of grain boundaries during sintering of garnet ceramics and increase the average grain size. The absorption spectra of Ce,Ca:LuAG ceramics indicate that the amplitudes of 4f→5d₁ transition at 447 nm and the 4f→5d₂ transition at 345 nm of Ce³⁺ increase with the starting content of Ce³⁺ that evidences its entering into the garnet lattice. An intense absorption below 350 nm is due to the charge transfer absorption band induced by Ce⁴⁺ in the garnet lattice. The photoluminescence decay curves of the 5d₁→4f emission of Ce³⁺ (510 nm) under the 452 nm excitation show that the fast component intensity I₁ of all Ce,Ca:LuAG ceramics is > 95%. The main decay time shows an acceleration at the Ce concentration of > 0.3% because the formation of closely spaced (Ce⁴⁺—Ca²⁺) pairs can increase with the increase of Ce concentration. The scintillation decay under γ-excitation slightly slows down with increasing Ce concentration, since the proportion of induced Ce⁴⁺ decreases with increasing total Ce concentration, and the competition in accelerating scintillation process becomes progressively weaker. The X-ray excited luminescence (XEL) intensity of the Ce,Ca:LuAG ceramics firstly increases and then decreases with the Ce concentration, reaching the maximum value as x=0.3. The integral scintillation efficiency of Ce,Ca:LuAG ceramics is 146%-190% of a BGO crystal. The variation of the LY and the integral scintillation efficiency obtained from XEL have a similar pattern, i.e., increasing and then decreasing with the increase of Ce concentration. The optimum energy resolution (ER) is obtained as x=0.3. The normalized afterglow intensity after continuous X-ray irradiation monotonically decreases with the Ce concentration due to the more intense competition ability of cerium centers for charge carrier capture.

Ce,Ca:LuAG scintillation ceramics with different Ce concentrations were prepared with coprecipitation nano-powders via vacuum pre-sintering and HIP post-treatment. A small amount of Al-rich secondary phase existed in the ceramics due to the loss of lutetium during coprecipitation and washing, which could be avoided by component design in the future ceramic preparation. At the Ce concentration of 0.3%, Ce,Ca:LuAG ceramic had the optimum XEL integral efficiency, LY and ER. This study demonstrated that Ce doping concentration had a certain effect on the PL decay, scintillation decay and afterglow of Ce,Ca:LuAG ceramics, providing a guidance for the subsequent selection of doping concentration in practical applications.

The lattice environment of low phonon energy of strontium fluoride can suppress the energy loss caused by non-radiation process of trivalent lanthanide ions, so it is an ideal matrix material for active-ions doping. The fabrication and spectroscopic properties of SrF₂ transparent ceramics doped with trivalent lanthanide ions have attracted recent attention. Trivalent lanthanide ions tend to form clusters in SrF₂ crystal lattice, and the distance between ions shortens, resulting in energy exchange process and concentration quenching within these clusters, especially for rare-earth ions with a rich energy level structure. The concentration quenching effect leads to a reduction in luminescence intensity, and their applications are limited to a certain extent. Co-doping buffer ions into Pr:SrF₂ system is an effective solution to alleviate the concentration quenching of Pr³⁺ luminescence intensity. In this paper, Y³⁺ ions were selected as buffer ions. SrF₂ raw powder was synthesized by a chemical precipitation method, and Pr³⁺ and Y³⁺ ions co-doped SrF₂ transparent ceramic were fabricated in vacuum by a hot pressing (HP) sintering technology. The effect of Y³⁺ doping levels on the microstructure and luminescence spectroscopic characteristics of Pr, Y:SrF₂ transparent ceramics was discussed to regulate the luminescence properties of Pr³⁺ in SrF₂ transparent ceramics.

SrF₂ particles were synthesized by a chemical precipitation method with commercial strontium nitrate and potassium fluoride reagents. Cationic (Sr²⁺) and anion (F^-) solutions were prepared, and then separated via high-speed centrifugation after washing for three times. The as-synthesized SrF₂ particles were mixed with commercial available PrF₃ and YF₃ particles, where the doping level of PrF₃ was fixed at 3.0% and the level of YF₃ was varied from 0 to 10.0%. Pr, Y:SrF₂ transparent ceramics were fabricated in vacuum by a hot pressing sintering technology. The sintering temperature was 1 000 ℃, the pressure was 30 MPa and the holding time was 120 min. After the heat preservation, the ceramic samples were cooled to room temperature in the furnace, and the ceramics were polished on the both sides.

The transmittance and absorption spectra of the ceramics were determined by a UV-vis-NIR spectrophotometer, and the phase composition of the transparent ceramics was characterized by a X-ray diffractometer. The polished transparent ceramics were immersed in 6 mol/L hydrochloric acid solution and corroded at room temperature for 10 min. The microstructure of transparent ceramics was measured by a phenom Prox model scanning electron microscope. The luminescence spectra and fluorescence lifetime of transparent ceramics were measured by a fluorescence spectrofluorometer.

The diameter of Pr, Y:SrF₂ transparent ceramics is16 mm and the thickness is 2 mm. The lattice parameters of ceramics are gradually decreased after Sr²⁺ are replaced by the smaller Y³⁺. Pr, Y:SrF₂ ceramics co-doped with 0-5.0% YF₃ have the similar transmittance, and the maximum transmittance of the ceramics is close to 90%. The ceramic co-doped with 5.0% YF₃ has the maximum transmittance, and the transmittance at 400 nm wavelength is 90.1%, which is similar to the theoretical value. The transmittance of Pr, Y:SrF₂ ceramics decreases significantly with the increase of YF₃ co-doping level to 10.0%.

After YF₃ co-doping, the absorption capacity of the ceramic is enhanced, and the strongest absorption capacity is located at 443 nm. The ceramic has a dense microstructure without residual pores or impurity phase. The lattice distortion caused by the radius difference between Y³⁺ and Sr²⁺ promote the diffusion rate of ions at a high temperature and increase the average grain size of ceramics. ³P₂ states of Pr³⁺ are directly populated from ³H₄ ground states under light excitation at 443 nm, and then fall back to adjacent ³P₁, ³P₀ and ¹D₂ states through non-radiative relaxation, multiphonon relaxation and cross relaxation processes. These radiation transitions from generated high energy states to lower states forming different spectral bands. Y³⁺ break the cluster structure of Pr³⁺ , releasing more luminous centers, and reducing the energy transfer process within Pr³⁺, thus improving the luminous intensity. Y³⁺ co-doping causes the change of the local coordination environment around Pr³⁺, varying the transition probability between different energy states. The increase of the transition probability corresponds to the increase of the emission intensity. The proportion of emission band intensity to the total luminous intensity of Pr, Y:SrF₂ ceramics varies with the co-doping levels of YF₃, and the proportion of orange luminescence band intensity increases in the spectrum. The co-doping of Y³⁺ ions leads to serious distortion of SrF₂ lattice, decreases the symmetry of matrix structure, and accelerates the energy transition rate. The fluorescence lifetime of Pr³⁺ energy transition decreases with the increase of Y³⁺ co-doping level.

SrF₂ particles were synthesized by a chemical precipitation method, and Pr, Y:SrF₂ transparent ceramics were fabricated by a hot pressing technique. The ceramic co-doped with 5.0% YF₃ had the mximum transmittance, and the transmittance at 400 nm was 90.1%. As the co-doping level of YF₃ increased to 10.0%, the transmittance of Pr, Y:SrF₂ transparent ceramics decreased, and the average grain size increased to 200.1 μm. The co-doping of Y³⁺ ions into Pr, Y:SrF₂ transparent ceramics improved the absorption capacity of the ceramics. The absorption cross-sections of Pr:SrF₂ and Pr, Y:SrF₂ transparent ceramics at 443 nm were 3.59×10^-21 cm² and 1.13×10^-20 cm², respectively. The luminescence spectral properties of Pr, Y:SrF₂ transparent ceramics were regulated via changing the clusters and local coordination structures of Pr³⁺ in the matrix, and the luminescence of the ceramics gradually changed from reddish to orange-yellow. The lifetime of ³P₀→³H₄, ¹D₂→³H₄ and ³P₀→³F₂ transition decreased from 114.2, 100.4 μs and 91.8 μs to 28.6, 24.4 μs and 30.3 μs, respectively.

The quality of the powder is crucial for the performance of transparent yttria ceramics. The powder characteristics, such as size, shape, distribution, and agglomeration, directly affect the densification and microstructure of the ceramics. Co-precipitation is a common method for preparing yttria powder, which has the advantages of simplicity, low cost, and scalability. However, it also has some drawbacks, such as uneven particle size distribution, severe agglomeration, and complicated post-treatment processes. The agglomeration of the powder not only reduces its activity, but also impairs its sinterability. Therefore, how to effectively disperse the powder is one of the key techniques for improving the quality of transparent yttria ceramics. Adding dispersant is a common method for dispersing the powder, among which PEG (polyethylene glycol) is a widely used dispersant, which can form strong hydrogen bonds with the surface of hydroxide colloids, and enhance the stability of the colloids by steric hindrance effect. Moreover, since PEG is an organic substance, it can be completely removed during the high-temperature calcination process, without causing any negative impact on the properties of the ceramics. In this study, we use PEG as the dispersant, and investigate its influence on the optical properties of yttria nanocrystalline powder and its sintered ceramics.

Zirconium oxynitrate and yttrium nitrate hexahydrate were dissolved in anhydrous ethanol to prepare dilute solutions as reaction materials. Two salt solutions were mixed according to a stoichiometric ratio of (Y_0.97Zr_0.03)₂O₃. Ammonium hydroxide was dissolved in 100 mL of anhydrous ethanol as a precipitant, and to investigate the effect of the addition amount of polyethylene glycol 4000 as a dispersant on the dispersion of samples, and PEG4000 with different molar amounts of 0, 0.1% (mole fraction), 0.3%, 0.6%, 0.9% of Y³⁺ content was added to the precipitant. The precipitant solution was gradually added to the reactants at 4 mL/min under vigorous stirring until the pH reached 9. The resulting precipitates were aged at room temperature for 4 h, then washed for 4 times with deionized water to remove impurity ions, and collected by centrifugation. The washed precipitates were freeze-dried for 10 h to obtain well-dispersed precursor powders. The precursor powders were calcined in a muffle furnace at 400, 600 ℃ and 850 ℃ for 4 h, respectively. The calcined yttria powders were wrapped with tantalum foil and then loaded into a customized graphite mold. The graphite mold was then placed in a vacuum hot-pressing furnace and sintered under vacuum. Another ceramic sample without tantalum foil shielding was also sintered using the same process for comparison. After the hot-pressing step, the Ta foil was removed, and the samples were annealed at 900 ℃ in air for 2 h and then ground and polished for the coming characterizations.

The precursor was completely decomposed and crystallized into well-crystallized yttria powder with little organic and impurity residues after calcination at 850 ℃, as revealed by TG-DSC, FT-IR and XRD analyses. The average particle size of the yttria powder prepared with different dispersant concentrations was in the range of 20-30 nm, indicating that the addition of PEG had little effect on the grain growth. However, the addition of PEG changed the agglomeration state of the particles in the powder. At low PEG addition, the surface charge of the powder was unbalanced, and the particles tended to agglomerate due to van der Waals force or electrostatic force, because PEG could not form a uniform and thick adsorption layer on the particle surface. This agglomeration reduced the energy demand for high-temperature sintering of the nanopowder, but also resulted in the microstructural inhomogeneity of the green body. At high PEG addition, the particles were effectively isolated from the surface charge by the uniform and thick adsorption layer formed by PEG on the particle surface, which suppressed the agglomeration effect of the particles, and also provided steric hindrance and solvation effects, preventing the particles from contacting and bonding.

Yttria transparent ceramics were obtained after sintering the yttria powder. The use of tantalum foil effectively shielded the carbon contamination. The sintering temperature of 1 500 ℃ was the optimal condition for the sample performance. The density and transmittance of the samples decreased when the PEG addition was low. The transmittance of the samples increased slightly when the PEG addition increased to 0.9% (the content of PEG in ethanol). This was because the addition of a small amount of PEG caused the agglomeration of yttria powder, which affected the forming, sintering and densification processes of the ceramics. With the increase of PEG content, the agglomeration degree of yttria powder decreased, and the ceramic grain size also gradually decreased. When the transmittance of the samples reached the maximum value, the ceramic grain size was 1-2 μm. However, the transmittance of all samples in the visible light range was still not high, which might be related to the oxygen vacancies and carbon impurities in the samples.

Yttria nanopowders were synthesized by a precipitation method in ethanol solvent with ammonia as a precipitant, PEG as a dispersant, and 3% ZrO₂ as a sintering aid. The optimal calcination temperature of the precursor was determined to be 850 ℃ by thermal analysis, at which a high purity cubic yttria phase could be obtained. Y₂O₃ ceramics were prepared by a vacuum hot-pressing sintering technique at 1 450-1 600 ℃ under 30 MPa, and the samples were wrapped with tantalum foil to prevent carbon contamination. The results showed that 1 500 ℃ was the optimal sintering temperature, at which Y₂O₃ ceramic had the maximum optical transparency, with a linear transmittance of 48.4% at 1 100 nm wavelength for samples with the thickness of 2 mm. In addition, the PEG addition also affected the microstructure and optical properties of Y₂O₃ ceramic, and when the PEG content was 0.9%, Y₂O₃ ceramic had uniform and fine grains, which was beneficial to improving the transparency. This work could provide a reference for the effective preparation of a high-performance Y₂O₃ ceramic.

Tremendous efforts are devoted to improving the luminous performance of white light-emitting diodes (WLEDs) as so-called fourth generation lighting source. The most common approach for manufacturing WLEDs is to cover a blue LED chip with a yellow-emitting Y₃Al₅O₁₂:Ce³⁺ (YAG:Ce³⁺) phosphor. This construction has an advantage of high lumen efficiency, but an insufficient red light component in the spectrum leads to a low color rendering index and a high color temperature. An effective solution is to enhance the red emission of YAG:Ce³⁺ via introducing red-emitting ions such as Pr³⁺, Eu³⁺, Sm³⁺, Mn²⁺, Mn⁴⁺ and Cr³⁺. In Y₃Al₅O₁₂ and many other garnet-structured aluminates, Mn²⁺, Mn⁴⁺ and Cr³⁺ are suffered from serious luminescent thermal quenching, while Pr³⁺ and Sm³⁺ are severely limited by concentration quenching. In comparison, the luminescence of Eu³⁺ in YAG is highly anticipated because it has an ion radius similar to Y³⁺, which is beneficial to increasing doping concentration and reducing luminescence quenching. In this paper, a series of Ce³⁺-Eu³⁺ co-doped YAG transparent ceramics with different Eu concentrations were fabricated, and the effect of Eu³⁺-doping on the structure, optical transmittance and photoluminescence properties was also evaluated.

Ce_0.01Y_2.99-xEu_xAl₅O₁₂ (0 < x < 0.75, Eu contents of 0-25%) fluorescent ceramics were synthesized in vacuum via high-temperature solid-state reaction. High purity Y₂O₃, α-Al₂O₃, Eu₂O₃, and CeO₂ powders were used as starting materials. Oleic acid and tetraethoxysilane were used as a molding agent and a sintering aid, respectively. These materials were weighed in stoichiometric ratios, and ground in ethanol for 20 h. The obtained slurry was dried, ground and sieved through a 100 mesh sieve. The resulting powders were made into 20 mm discs under uniaxial pressure of 5-10 MPa, annealed in oxygen atmosphere at 800 ℃ for 10 h, and cold isostatically pressed under 200 MPa. Finally, the ceramics was mechanically thinned and polished to the thickness of 1mm for the subsequent structural and performance characterizations.

For YAG:Ce³⁺/Eu³⁺ ceramics sintered at 1 600 ℃, a garnet phase appears when Eu content is in the range of 0-25%. However, for the ceramics sintered at 1700 ℃, the grains grow as Eu content (x) increases, and eventually become over-sintered and textured when x ≥ 15%. YAG:Ce³⁺/Eu³⁺ ceramics show some characteristic absorption/emission bands for both Ce³⁺ and Eu³⁺. For Ce³⁺, the PLE spectrum involves two broad bands at 320-370 nm and 400-520 nm, and the PL spectrum spans from 440 nm to 650 nm and is dominated by yellow-green emission. Unlike a broad band absorption/emission of Ce³⁺, the PLE/PL spectra of Eu³⁺ are composed of a series of lines originating form f-f transitions. There are some intense excitation lines around the maximum absorption locating at 395 nm, constructing a quasi-continuous broad band. It is therefore inferred that YAG:Ce³⁺/Eu³⁺ ceramics can be efficiently pumped by NUV LED chips. The PL spectrum of Eu³⁺ ions covers the spectral range from orange to deep red. The orange emission from ⁵D₀→⁷F₁ transition is more intense than the red emission from ⁵D₀→⁷F₂ transition because Eu³⁺ ions occupy the lattice sites with an inversion symmetry.

The PL spectra of YAG:Ce³⁺/Eu³⁺ ceramics vary with excitation wavelength due to the difference in absorption spectra of Ce³⁺ and Eu³⁺. Therefore, the effect of excitation wavelength on the PL spectra of YAG:Ce³⁺/Eu³⁺ ceramics with various Eu³⁺ concentrations were investigated. Under the excitation at 363 nm and 466 nm, the emission bands/peaks of both Ce³⁺ and Eu³⁺ appear. The PL spectrum excited at 442 nm contains almost only a broadband emission of Ce³⁺ ions, as Eu³⁺ ions hardly absorb photons near this wavelength. In contrast, the PL spectrum under the excitation at 395 nm is dominated by Eu³⁺ emission due to a weak absorption of Ce³⁺. Clearly, the PL spectrum and color coordinates of YAG:Ce³⁺/Eu³⁺ exhibit an excitation wavelength dependence, which has a potential application. In addition, under excitation at 442 or 466 nm, the emission of Eu³⁺ is weak, while that of Ce³⁺ rapidly decays with increasing Eu content. As a result, the emission intensity of YAG: Ce³⁺/Eu³⁺ ceramics decreases rapidly with the increase of Eu content when excited by blue light. However, when excited by UV light (at 363 nm), the attenuation of Ce³⁺ emission is compensated by the enhancement of Eu³⁺ emission to some extent, and the overall photoluminescence intensity can be maintained at a high level. Especially, the samples with Eu contents of 1%-3% can emit an intense orange-yellow light, thus increasing the red component of YAG:Ce³⁺ effectively. This kind of phosphor has a certain practicability in WLEDs. Also, under the excitation at 395 nm, the PL spectrum originates mainly from the f-f transitions of Eu³⁺, and the quenching concentration of Eu³⁺ in YAG:Ce³⁺ ceramics is 9%, which is much higher than that of Pr³⁺ (i.e., 0.8%) and Sm³⁺ (i.e., 3%). For rare-earth ions with forbidden f-f transition, a high quenching concentration is crucial for achieving a high fluorescence emission intensity and meeting the inevitable requirements of LED applications.

Thermal quenching is a key factor in evaluating luminescent materials. For the YAG:Ce³⁺/Eu³⁺ ceramics with Eu contents of 1% and 9%, the profile of PL spectrum does not change with temperature, leading to a good temperature stability of color coordinates. The relative PL intensity at 575 K remains 83% and 61% of that at room temperature (i.e., 300 K) for 1% and 9% Eu³⁺-doped samples, respectively, indicating that the photoluminescence thermal quenching in YAG:Ce³⁺/Eu³⁺ ceramics is rather weak. Furthermore, the relative PL intensity (I_T/I₀) and the ratio of emission arising from ⁵D₀ to different ⁷F_J levels, including (⁵D₀→⁷F₁)/(⁵D₀→⁷F₂)、(⁵D₀→⁷F₄)/(⁵D₀→⁷F₂) and (⁵D₀→⁷F_1,2)/(⁵D₀→⁷F₄), vary linearly with temperature. This feature can be used in the field of fluorescent thermometers, which is one of the potential applications of YAG:Ce³⁺/Eu³⁺ transparent ceramics.

In YAG: Ce³⁺/Eu³⁺ transparent ceramics, Eu³⁺ could be excited by UV (at 363 nm), NUV (at 380-405 nm), and blue (at 466 nm) light, emitting a series of lines in the spectra range from 580 nm to 750 nm, i.e., from orange to deep red. These PL lines had an important application in lighting and display. Eu³⁺ showed the maximum emission when its concentration was between 5% and 9%. The high quenching concentration had a advantage in enhancing the red component in the emission spectrum of YAG:Ce³⁺. Furthermore, the red component in the spectrum could be regulated via adjusting the doping concentration of Eu³⁺. However, the increase in Eu content led to a decrease in Ce³⁺ emission intensity. The similar phenomena appeared in various dual/multi activators co-doped systems. The solution of this common problem was crucial for the application of this kind of luminescent materials in WLEDs. A highlight of YAG:Ce³⁺/Eu³⁺ transparent ceramics was that they could absorb NUV light from 380 nm to 405 nm and emit an orange-to-red light with a high quantum efficiency and a high thermal stability. Therefore, they could be used as red-emitting phosphors excited by NUV LED chips. Besides, the ratios of emission arising from ⁵D₀ to different ⁷F_J levels varied linearly with temperature, having a promising potential application in the field of fluorescent thermometers.

Transparent aluminum oxynitride (AlON) ceramics have superior optical transparency properties, high strength and hardness, having great potential applications as window, dome, and other optical components requiring high strength combined with optical transparency. Recent efforts are made to fabricate transparent AlON ceramics, i.e., synthesis of high purity AlON powder, optimization of the particle size distribution of AlON powder, and doping with various additives in different amounts. In these works, sintering additives, such as Y₂O₃, Y₂O₃/La₂O₃, Y₂O₃/La₂O₃/MgO, etc., are commonly used to accelerate densification process of AlON. However, a long dwelling duration of ≥6 h is generally required to eliminate pores to obtain a high transmittance in pressureless sintering. Although the dwelling duration of pressureless sintering AlON ceramics can be shortened to 2.5 h, it is crucial to carefully match the doping amount of Y₂O₃ with the particle size of AlON powder. The excessive Y₂O₃ doping can result in a significant decomposition of AlON into α-Al₂O₃ and AlN, and the massive formation of α-Al₂O₃ can lead to particle aggregation/coarsening, and even separation between components during the early stage of sintering, which in turn retards the subsequent densification process. Studies on using CaCO₃ as an additive to fabricate AlON ceramics indicate that it should be a promising way to fabricate highly transparent AlON ceramics via suppressing AlON decomposition during heating with sintering additive. In this paper, La₂O₃ was respectively doped to AlON powders with median particle size (D₅₀) of 1.1 μm and 2.0 μm to prepared transparent AlON ceramics via pressureless sintering (PS) at 1 880 ℃ for 2.5 h. The effect of La₂O₃ doping amount on the transmittance of AlON ceramics was investigated, and the phase composition, microstructure and densification process of samples during heating were monitored to investigate the fast densification mechanism of AlON.

Pure AlON powder and La₂O₃ powder were used as starting materials. The AlON powders doped with 0-0.50% La₂O₃ were ground in absolute ethyl alcohol in a grinding mill with Si₃N₄ balls, and AlON powder with D₅₀ of 1.1 μm and 2.0 μm were obtained, referred to as P1.1 and P2.0, respectively. The ground powders were loaded into a pellet of 13 mm in diameter at 50 MPa and then cold isostatically pressed at 120 MPa to prepare green bodies. The green bodies were pressureless sintered in N₂ in a graphite furnace at a heating rate of 20 ℃/min. One group of samples were sintered at 1 880 ℃ for 2.5 h to prepare transparent AlON ceramics, and another group samples were respectively heated to 1 400-1 900 ℃ (no dwelling) to investigate the phase transformation, microstructure evolution and densification process during heating. The transparent AlON ceramics were ground and mirror polished on the both sides to a thickness of 2 mm for the optical transmittance measurement. The polished samples were hot etched at 1 640 ℃ for 40 min to determine their microstructure.

After a 2.5 h dwell at 1 880 ℃, all the La₂O₃ doped AlON ceramics prepared (i.e., P1.1 and P2.0) show higher transmittance than their undoped counterparts, and the transmittances at 3 850 nm enhance from 17% and 0% to 80% and 84%, respectively. Moreover, the doping of La₂O₃ broadens the transmission wavelength range of the fabricated samples at 2 500-6 000 nm for P1.1 and the optical range at 220-6 000 nm for P2.0. Doping amount of La₂O₃ has a similar effect on the transmittance of AlON ceramics prepared by P1.1 and P2.0. When the doping amount of La₂O₃ ≤0.15% in P1.1 and ≤0.40% in P2.0, increasing the doping amount of La₂O₃ contributes to the enhanced transmittance of the AlON ceramics. Doping 0.15%-0.25% and 0.40% La₂O₃ into P1.1 and P2.0, respectively, results in the transmittances of AlON ceramics up to 80%-84% at 3 850 nm. The high transmittance of fabricated AlON ceramics is mainly due to thehigh relative density. For the two highly transparent AlON ceramics prepared by P1.1 doped with 0.20% La₂O₃ and P2.0 doped with 0.40% La₂O₃, their average grain sizes are 97.3 μm and 60.2 μm, and their Vickers hardness values are (16.00±0.31) GPa and (16.36±0.31) GPa.

For the samples prepared by P1.1 doped with 0.20% La₂O₃ and by P2.0 doped with 0.40% La₂O₃, although AlON undergoes decomposition and reformation during heating. α-Al₂O₃ content in the samples does not exceed 25.1%. Moreover, a lower amount of α-Al₂O₃ appears in the sample P2.0, compared to the sample P1.1 at 1 400-1 700 ℃. Furthermore, compared to the samples without additive doping, the two samples above exhibit lower α-Al₂O₃ contents at ≤1 600 ℃. It is indicated that the decomposition of AlON during heating is slightly suppressed by doping La₂O₃ as an additive. No obvious particle aggregation/coarsening or no composition decomposition occur on the fracture surface of the samples due to the reduced α-Al₂O₃ content in the samples. It is beneficial for the subsequently densification process.

AlON ceramics with a high infrared transmittance of 80%-84% at 3 850 nm were fabricated via pressureless sintering at 1 880 ℃ for 2.5 h after doping 0.15%-0.25% and 0.40% La₂O₃ to AlON powders with D₅₀ of 1.1 μm and 2.0 μm, respectively. During heating, La₂O₃ could suppress the decomposition of AlON and lead to a lower amount of α-Al₂O₃ being decomposed from AlON. This could reduce the particle aggregation/coarsening and even the separation of composition at the early stage of sintering. The obtained exceptional microstructure was beneficial for the subsequent fast densification. Moreover, La₂O₃ could provide sufficient sintering kinetics to effectively eliminate pores at the later stage of sintering, thus promoting the densification of AlON.

Transparent ceramics are polycrystalline inorganic nonmetallic materials that are transparent in the visible range and have some superior characteristics (i.e., high melting point, high strength, high insulation, corrosion resistance, high temperature resistance, and good light transmission). Diamond as a hardest substance in nature has an extremely high thermal conductivity, a wide spectral transmission range and a good chemical stability as an ideal transparent ceramic material in harsh and extreme environments. However, the single crystal of diamond is brittle and easily broken along the cleavage plane (111), greatly restricting the application of single crystal diamond. Nanocrystalline diamond (NPD) is a polymer of small diamond particles in nano-scale, and the grains are directly bonded by diamond to form a compact structure. The macroscopic properties of nanocrystalline diamond show an isotropy, and there is no directional and disfoliating plane due to the disordered accumulation between particles, having the better mechanical properties, compared to single crystal diamond. The nano-polycrystalline diamond can be prepared, but the preparation of transparent nano-polycrystalline diamond and its application in the field of light transmission are rarely reported. Little work on the transparent mechanism and defect formation mechanism of nano-polycrystalline diamond have been done yet. In this paper, a nanocrystalline diamond was prepared by a high-temperature and high-pressure method. In addition, the mechanical properties were also investigated.

A high purity graphite powder (99.999 9% in mass fraction) as a precursor was prepressed into a cylindrical shape (Φ=2.5 mm, h = 2 mm), wrapped with aluminum oxide, and assembled. The pressure was uniformly increased to 15 GPa for 8 h. Afterwards, a rhenium (Re) heater was used for heating, and the temperature was uniformly increased to 1 800, 2 300 ℃, and 2 600 ℃ at 200 ℃/min. After holding for while, the temperature was slowly reduced to room temperature at 200 ℃/min, and the pressure was slowly relieved to normal pressure after 15 h. The polished sample was determined by an optical microscope. The phase composition of the sample was analyzed by a model Rigaku FR-X X-ray diffractometer (target Mo, wavelength λ = 0.709 3 Å). The Raman spectra were determined by a model Mono Vista CRS+ Raman spectrometer at an excitation wavelength of 532 nm. The surface morphology of the samples was analyzed by a model Hitachi SU-70 scanning electron microscope. The transmission samples were prepared by focusing ion beam (FIB), and the electron diffraction (SAED) and high-resolution transmission electron microscopic images of the samples were characterized by a model JEM-2200FS transmission electron microscope (HRTEM).

A high-temperature and high-pressure method is an effective way to prepare high performance transparent ceramics. Graphite as a raw material can be transformed into a transparent nano-polycrystalline diamond at 15 GPa and 2 600 ℃ for 1 min. Also, the temperature gradient of the sample cavity leads to a non-uniform transparency of the sample, and the increase of temperature is conducive to the increase of the transparency of the nanocrystalline diamond. The results of morphology analysis show that the sample has a layer structure of martensitic phase transformation and a uniform particle structure of diffusion phase transformation. The layer thickness and grain size are approximately 100 nm. High pressure and stress cause a partial dislocation within the sample, improving the mechanical strength of the sample. Nano-polycrystalline diamond as the maximum hardness transparent ceramic is expected to be widely used in the field of special optical windows under extreme conditions.

Graphite powder was selected as a precursor material and pre-pressed into a 2.5 mm diameter and 2 mm high cylinder. The Kawai 6-8 press device (1 000 t) was used to synthesize a nano-polycrystalline diamond at 15 GPa and 1 800-2 600 ℃. The results showed that the partially transparent nano-polycrystalline diamond could be synthesized under the conversion boundary conditions of synthetic transparent diamond, and the complete transformation from graphite to diamond was completed at 15 GPa and 2 600 ℃, and the completely transparent nano-polycrystalline diamond was synthesized. The results by XRD and SEM analysis indicated that the prepared samples were a pure cubic phase, there was no preferred orientation, and the grain size was approximately 100 nm. Nano-polycrystalline diamond as the hardest transparent ceramic could be used to manufacture high-performance transparent ceramics, which have a promising application in military, industrial and other fields.

Gadolinium oxysulfide (Gd₂O₂S, GOS) scintillation ceramics as luminescent materials are widely used in nuclear medical imaging and security inspections due to their high light output, low afterglow, and strong X-ray stopping capability. However, the synthesis process of GOS powder is intricately complex, having some challenges for achieving accurate component ratios and controlling the second-phase impurities, thus leading to difficulties in producing high-quality GOS scintillation ceramics. In this paper, GOS precursor powders were firstly synthesized by a hot water bath reduction method with sub-micron and nano-scale gadolinium oxide (Gd₂O₃) as raw materials, and Gd₂O₂S:Pr scintillation ceramics were subsequently prepared by a two-step sintering method (i.e., atmospheric pressure pre-sintering and HIP post-treatment). In addition, the microstructure and optical properties of the powder and ceramics were also characterized.

Gd₂O₂S:Pr phosphors were synthesized by a hot water bath reduction method. Gd₂O₃ (99.999%, Ganzhou Berier New Materials Co., Ltd., China), Pr₆O₁₁ (99.995%, Ganzhou Berier New Materials Co., Ltd., China), and concentrated H₂SO₄ (AR, Shanghai Macklin Biochemical Co., Ltd., China) were used as high-quality raw materials. Gd₂O₃ and Pr₆O₁₁ were firstly weighed according to (Gd_1-xPr_x)₂O₂S (x=0.001-0.010), and then mixed with a diluted sulfuric acid solution at a molar ratio of Gd₂O₃:H₂SO₄ of 1∶1. The suspension was stirred and uniformed by a magnetic stirrer. Also, the suspension was heated in a hot water bath at 90 ℃ for 150 min. After the reaction was completed, the suspension was cooled to room temperature, and then the suspension was filtered, washed/dried, ground and sieved. Finally, the ground powders were calcined in a tube furnace, The resulting mixture was subsequently calcined in a tube furnace under a 20% hydrogen-nitrogen mixture reduction atmosphere at 750 ℃ for 6 h to yield the gadolinium oxide sulfide powders.

The Gd₂O₂S:Pr powders obtained were formed in a 15 mm cylindrical mold under uniaxial dry pressing at 10 MPa and then treated by cold isostatic pressing (CIP) at 200 MPa. Subsequently, the green bodies were pre-sintered in a flowing hydrogen-nitrogen atmosphere at 1 350 ℃ for 2 h, and HIP post-treatment was performed in argon at 1 600 ℃ and 180 MPa for 2 h. The sintered ceramics were polished and annealed for coming characterizations.

The microstructure of Gd₂O₂S:Pr powders prepared with submicron-sized particles of Gd₂O₃ and nano-sized particles of Gd₂O₃ was examined, respectively. The results reveal differences in particle size and morphology between the two powders. Gd₂O₂S:Pr powder derived from submicron-sized Gd₂O₃ powder has irregular columnar particles and sheet-like structures, whereas that prepared with nano-sized Gd₂O₃ powder is coralloid particles. The fluorescence emission spectrum of the ceramics is consistent with that of the powders although the excitation peak partially obscures by the absorption peak. High-temperature sintering gradually eliminates the differences in particle size and leads to the formation of dense sintered bodies in the ceramics. Consequently, the discrepancy in fluorescence intensity caused by differences in particle size is reduced. Specifically, ceramics prepared with Gd₂O₃ nano-powder display a low fluorescence intensity, while those produced with Gd₂O₃ submicron-sized powder exhibites a high fluorescence intensity. To further investigate the factors influencing the optical properties of the ceramics, the doping concentration of Pr³⁺ was explored. Fluorescence and XEL spectroscopy tests were conducted on Gd₂O₂S:xPr (x=0.1%, 0.2%, 0.4%, 0.6%, 0.8%, and 1.0%) ceramics prepared using Gd₂O₃ submicron-sized powder with different Pr³⁺ doping concentrations. The experimental results reveal that the fluorescence and XEL spectral intensities reach their maximum values when the Pr³⁺ doping concentration is 0.2%. The SEM images demonstrate that the ceramics prepared using submicron-sized Gd₂O₃ powder have a dense structure with well-defined grain boundaries and uniform grain sizes. The EDS element distribution results confirm the effective doping of Pr element into the host lattice. Gd₂O₂S:Pr ceramic prepared exhibits a high transmittance.

The utilization of Gd₂O₃ submicron-sized powder resulted in Gd₂O₂S:Pr powder and the scintillation ceramics displayed a higher fluorescence intensity, compared to that synthesized with Gd₂O₃ nano-sized powder. Furthermore, Gd₂O₂S scintillation ceramics pre-sintered in a weak reducing atmosphere and subsequently sintered in a hot isostatic pressing exhibited improved density and optical properties. The maximum fluorescence intensity of the scintillation ceramics was achieved at a Pr³⁺ doping concentration x of 0.2%. In addition, the scintillation ceramics achieved a transmittance of 31% at a wavelength of 513 nm. This study developed a simple and cost-effective ceramic preparation process without any byproduct pollution, thereby offering a promising potential for large-scale production of high-performance Gd₂O₂S scintillation ceramics.