The XRD patterns of the GAGG: xEu³⁺ transparent ceramics, shown in Figure 1A, confirm that all diffraction peaks can be well indexed to the standard Gd₃Al₃Ga₂O₁₂ (PDF#46-0448), with no detectable secondary phases. This result indicates that the incorporation of Eu³⁺ ions does not alter the garnet crystal structure of the host lattice. To further verify the phase purity and structural characteristics, Rietveld refinements were performed using GSAS-Ⅱ software^[26]. The corresponding refinement profiles are presented in Figure 1B-E. The obtained reliability factors (R_wp = 5.72%~8.47%) demonstrate satisfactory fitting quality. The refined structural parameters are summarized in Figure 1F and Table 1. As the Eu³⁺ concentration increases, both the lattice constant and unit-cell volume exhibit a slight increase, suggesting that Eu³⁺ ions are successfully incorporated into the lattice through substitution at the Gd³⁺ sites. The bulk densities (D) of GAGG: xEu³⁺ ceramics, measured by the Archimedes method, are approximately 6.4 g·cm^-3, corresponding to more than 98% of the theoretical density. These results indicate that highly densified ceramics were successfully obtained.

The schematic illustration of the GAGG crystal structure is presented in Figure 1G. GAGG crystallizes in a cubic garnet structure with space group O_h¹⁰-Ia3d. In this structure, Gd³⁺ ions occupy the dodecahedral [GdO₈] sites coordinated by eight O^2- ions, whereas Ga³⁺ and Al³⁺ preferentially occupy the octahedral [GaO₆] and tetrahedral [AlO₄] sites, respectively. However, partial cation disorder between Ga³⁺ and Al³⁺ ions may occur, resulting in the coexistence of [AlO₆] octahedra and [GaO₄] tetrahedra^[6]. Additionally, Gd³⁺ ions at dodecahedral sites can be readily substituted by rare earth ions with similar ionic radii, thereby facilitating the luminescence of diverse rare earth activators^[27,28]. The ionic radii of Eu³⁺ and Gd³⁺ are 0.0947 and 0.0938 nm, respectively. The slightly larger radius of Eu³⁺ facilitates its incorporation into the GAGG lattice via Gd³⁺, promoting efficient luminescence activation within the solid solution.

As shown in Figure 1H, the polished GAGG: xEu³⁺ transparent ceramics (thickness = 1 mm) exhibit excellent transparency without visible defects, confirming good optical quality. The prepared ceramics demonstrate optical transmittance values of approximately 70% within the 400-800 nm range. Among the investigated samples, the ceramic doped with the 2 at% Eu³⁺ exhibits the highest value of 77%, which is comparable to that reported for GAGG single crystals. Several sharp absorption peaks are observed in the visible region around 394, 466, and 527 nm, which correspond to Eu³⁺-related 4f-4f transitions^[29].

To further investigate the influence of Eu³⁺ doping on the microstructural evolution of GAGG ceramics, polished samples were thermally etched at 1,300 °C in air. The resulting SEM micrographs are shown in Figure 2. All samples exhibit dense microstructures with well-defined grain boundaries, and no obvious pores or secondary phases are observed. These results indicate a high degree of densification, which is consistent with the XRD and density measurements discussed above. To quantitatively evaluate the microstructural characteristics, grain-size distributions were analyzed, as in the inset of Figure 2. The average grain size shows a non-monotonic dependence on the Eu³⁺ concentration, increasing initially and then decreasing. The maximum average grain size of 8.32 μm occurs at 2 at% Eu³⁺, which also exhibits the highest optical transmittance. This enhancement is attributed to reduced light scattering at grain boundaries due to grain growth, effectively lowering optical losses and improving transparency in the ceramics.

The microstructure of the optimal GAGG: 2 at% Eu³⁺ ceramic was further characterized by TEM. The selected area electron diffraction (SAED) pattern shown in Figure 3A exhibits well-defined diffraction spots along the [211] zone axis, indicating the high crystallinity of the GAGG phase.

The high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) image acquired along the [211] direction is presented in Figure 3B. Clear lattice fringes can be observed, demonstrating the excellent crystallinity of the ceramic. The measured interplanar spacing is 0.87 nm, corresponding to the ($$ 00\bar{1} $$) plane, which agrees well with the theoretical value of 0.86 nm (PDF#46-0448). To evaluate the elemental distribution within the ceramic, energy dispersive spectroscopy (EDS) mapping was performed on a representative region, as shown in Figure 3C. The results indicate that Gd, Al, Ga, Eu, and O elements are uniformly distributed, with no other impurity elements detected. Additionally, no obvious elemental segregation or enrichment phenomenon is observed. Combining the TEM and EDS analysis results, it is confirmed that the prepared ceramics possess high uniformity and optical quality.

The PLE spectrum of the GAGG: 2 at% Eu³⁺ ceramic is shown in Figure 4A. Five distinct excitation bands are observed in 350-500 nm, which can be assigned to the following intra-4f transitions of Eu³⁺ ions: ⁷F₀ → ⁵D₄, ⁷F₀ → ⁵L₇, ⁷F₀ → ⁵L₆, ⁷F₀ → ⁵D₃, and ⁷F₀ → ⁵D₂ ^[30]. The PL spectra of GAGG: xEu³⁺ ceramics with different Eu³⁺ concentrations are presented in Figure 4B. A dominant orange-red peak is observed at 590 nm, corresponding to the ⁵D₀ → ⁷F₁ magnetic dipole transition of Eu³⁺ ions. Additional emission peaks located at 610 and 630 nm are attributed to the ⁵D₀ → ⁷F₂ electric dipole transition, while the weaker emission bands at 653, 695, and 710 nm originate from the ⁵D₀ → ⁷F₃ and ⁵D₀ → ⁷F₄ transitions, respectively^[31]. Based on the Eu³⁺ emission spectrum, a schematic energy-level diagram illustrating the photoluminescence process is proposed, as shown in Figure 4C.

Studies have shown that the luminescence behavior of Eu³⁺ ions is highly sensitive to their local crystal environment. In the PL spectrum, the ⁵D₀ → ⁷F₁ magnetic dipole transition remains dominant, indicating that Eu³⁺ ions occupy sites with relatively high symmetry in the host lattice^[32] The emission intensity initially increases with increasing Eu³⁺ concentration and reaches a maximum at 2 at%, beyond which concentration quenching becomes apparent. This quenching behavior is likely associated with reduced interionic distance between Eu³⁺ ions at higher doping concentrations, which facilitates energy transfer among activator ions and increases the probability of excitation energy being trapped by defect states or impurity levels. The trapped energy may subsequently be dissipated through non-radiative relaxation or transferred to impurity ions, thereby reducing the luminescence efficiency of Eu³⁺ ions^[33,34]. The non-radiative energy transfer between rare earth ions can occur through two primary mechanisms: exchange interaction and electric multipole interaction^[35]. According to the Blasse model^[36], the critical distance (R_c) for doped ions can be calculated using Equation (1):

(1) $$ R_{c}=2\left(\frac{3 V}{4 \pi X_{c} N}\right)^{\frac{1}{3}} $$

Where (X_c) is the critical doping concentration; (V) is the unit-cell volume and (N) represents the number of available substitutional sites per unit cell. For the GAGG ceramics, a = b = c = 12.214 Å, N = 8, and V = abc = 1,822.106 ų, the calculated (R_c) is 22.15 Å. When R_c < 5 Å, energy transfer is predominantly governed by exchange interaction. The calculated (R_c) is significantly larger than 5 Å, indicating that the concentration quenching behavior is primarily attributed to electric multipolar interaction rather than exchange interaction^[37]. To further clarify the electric multipolar interaction mechanism, Dexter’s theory^[38] was employed to analyze the relationship between the luminescence intensity (I) and the Eu³⁺ concentration (x), which can be expressed by Equation (2):

(2) $$ lg (I / x)=\mathrm{C}-(\theta / 3) lg x $$

where (I) represents the luminous intensity of the ⁵D₀ → ⁷F₁ transition, while (C) and (θ) are constants. The type of multipolar interaction can be identified according to the value of (θ), where θ = 6, 8, and 10 correspond to electric dipole-electric dipole (d-d), electric dipole-electric quadrupole (d-q), and electric quadrupole-electric quadrupole (q-q), respectively^[39]. The plot of lg(I/x) versus lg(x) is shown in Figure 4D, where the slope is equal to -θ/3. From this plot, the value of θ was calculated to be approximately 4.827, which is close to 6. This result indicates that the concentration quenching mechanism in GAGG: xEu³⁺ is governed by electric dipole-dipole interactions^[40].

To evaluate the luminescence stability of GAGG: xEu³⁺ ceramics, temperature-dependent PL spectra of the GAGG: 2 at% Eu³⁺ sample were recorded over the temperature range of 300~550 K, as shown in Figure 5A. The luminescence intensity of Eu³⁺ gradually decreases with increasing temperature and retains 50.28% of its initial intensity when the temperature reaches 550 K. The decrease in luminescence intensity with increasing temperature is attributed to thermal quenching. As the temperature rises, electrons in the excited states of the luminescence centers can be thermally activated to higher vibrational levels and subsequently return to the ground state through non-radiative relaxation processes, leading to a reduction in luminescence intensity^[41]. The thermal quenching is closely related to the strength of phonon-electron coupling^[42]. The activation energy (ΔE_a) for thermal quenching was calculated by fitting the temperature-dependent luminescence intensity using the Arrhenius^[43] Equation (3):

(3) $$ I(T)=I_{0} /\left[1+B exp \left(-\Delta E_{a} / {k} T\right)\right] $$

where (I₀) is the initial luminescence intensity, I(T) is the luminescence intensity at a given temperature; (B) is a constant and (k) is the Boltzmann constant. The fitted curve for ΔE_a across the temperature range of 300-550 K is shown in the inset of Figure 5A. The fitting process yields ΔE_a = 0.13 eV. A higher ΔE_a indicates enhanced stability, as it corresponds to a higher non-radiative energy barrier that must be overcome during the transition from the excited state to the ground state. This elevated barrier effectively suppresses nonradiative relaxation, thereby confirming the excellent thermal stability of the GAGG: 2 at% Eu³⁺ ceramic^[44].

To further examine the influence of defect states on the thermal stability of radioluminescence in ceramics, TSL measurements were performed on the 2 at% Eu³⁺ ceramic, and the corresponding spectra are presented in Figure 5B. Unlike the sharp Eu³⁺ emission lines observed in the PL spectra, the TSL spectrum exhibits a broad red emission band in the 650-750 nm range, indicating that TSL process does not simply follow the PL emission mechanism. During TSL measurements, thermally released charge carriers undergo recombination through defect-related or activator-related radiative pathways. Therefore, the broad emission band centered at approximately 700 nm is likely associated with defect-mediated recombination processes within the GAGG lattice, involving oxygen vacancies, antisite defects, and local lattice distortions, together with a possible contribution from the weak Eu^{3+ 5}D₀ → ⁷F₄ transition. Continuous trap distributions are observed over the temperature range of 340~550 K. This is likely due to the stronger lattice vibrations at higher temperatures, which cause the release of electrons (or holes) from traps, leading to light emission. Their release difficulty depends on both the depth of the trap and the temperature of the material. Shallow trap levels can release charge carriers at relatively low temperatures, whereas deep trap levels require higher temperatures for carrier release. The relationship between the temperature and the trap depth (E) can be calculated using Equation (4)^[45]:

(4) $$ E=\left(0.94 \ln \frac{s}{\beta}+4.12\right) \times {k} T m $$

where (Tm) is the temperature corresponding to the TSL peak, (s) is the frequency factor, (β) is the heating rate, and (k) is the Boltzmann constant. In this measurement, (β) was fixed at 1 °C/s. Considering that the frequency factor (s) typically falls within the range of 10¹⁰~10¹⁴, a characteristic value of 10¹² was adopted to simplify the calculation while maintaining reasonable accuracy. The calculated trap depths of the GAGG: 2 at% Eu³⁺ ceramic range from 1.62 eV to 1.85 eV, indicating the presence of medium and deep trap levels. Generally, shallow traps can release charge carriers at relatively low temperatures and may participate in rapid carrier recombination processes. However, an excess concentration of shallow traps can increase the probability of non-radiative recombination, thereby reducing luminescence efficiency. In contrast, deeper traps require higher thermal energy for carrier release and can temporarily store charge carriers during excitation. These traps are beneficial for regulating carrier capture and delayed release processes, thereby influencing the thermal stability and radioluminescence performance of the material^[46].

To evaluate the high-energy X-ray detection capability of GAGG: xEu³⁺ transparent ceramics, as shown in Figure 6A. The GAGG: Eu ceramics exhibit X-ray absorption capability comparable to that of commercial scintillators. The scintillation properties were further characterized by X-ray-excited emission spectroscopy, and the corresponding spectra for different Eu³⁺ concentrations are presented in Figure 6B. The radioluminescence spectra are generally consistent with the PL spectra. Among the prepared GAGG: xEu³⁺ ceramics, the highest luminescence intensity is observed at 3 at%. When the Eu³⁺ concentration exceeds this value, concentration quenching becomes apparent. The quenching phenomenon observed under X-ray excitation differs from that observed in the PL spectra. The discrepancy in the optimal Eu³⁺ concentration between PL and X-ray-excited luminescence can be attributed to the different excitation and energy-transfer pathways. Under UV excitation, Eu³⁺ ions are mainly excited through localized 4f-4f transitions, and the emission intensity is strongly influenced by concentration-dependent energy migration among Eu³⁺ ions. In contrast, under X-ray excitation, the scintillation process involves X-ray absorption by the host lattice, generation of secondary electrons and holes, carrier transport, defect trapping and release, and subsequent energy transfer to Eu³⁺ emission centers. Therefore, the scintillation performance is jointly influenced by Eu³⁺ concentration, defect states, and optical transparency^[39]. The CIE 1931 chromaticity coordinates derived from X-ray-excited emission spectra are shown in Figure 6C. All GAGG: xEu³⁺ ceramics exhibit red emission in the visible region, indicating their potential for optoelectronic applications.

In addition, the dependence of the radioluminescence intensity of the GAGG: 2 at% Eu³⁺ ceramic on X-ray tube power was investigated, as shown in Figure 6D. The results show that the luminescence intensity increases steadily with increasing X-ray power. This enhancement is attributed to the increased flux of incident X-ray photons at higher power levels. These photons interact with the scintillator through processes such as the photoelectric effect and Compton scattering, generating a larger number of secondary electrons. The generated secondary electrons subsequently transfer energy to the luminescent centers, resulting in enhanced visible light emission. These findings suggest that the prepared ceramic scintillator holds promise for applications in X-ray detection.

To explore the potential of the prepared GAGG: xEu³⁺ transparent ceramics as scintillation layers for X-ray imaging, an X-ray imaging system was employed, as illustrated in Figure 7A. In this system, X-rays generated by an X-ray tube penetrate the object and subsequently irradiate the ceramic scintillator, where the incident X-rays are converted into visible photons. The emitted visible light is then captured by a digital camera and processed to generate the corresponding image. In the experiment, the GAGG: 2 at% Eu³⁺ transparent ceramic scintillator was selected for imaging, as shown in Figure 7B. A capsule containing a miniature spring was used as the imaging target [Figure 7C]. Under a tube voltage of 50 kV, the X-ray intensity was varied by adjusting the tube current from 20 μA to 80 μA. As shown in Figure 7D, both the image brightness and clarity improved with increasing current. At lower currents of 20 μA, the X-ray intensity was insufficient, resulting in a blurred image of the spring. As the current increased, the image quality gradually improved. At 80 μA, the scintillator exhibited the highest imaging brightness, and the internal structure of spring could be clearly distinguished. These preliminary imaging results demonstrate the feasibility of using GAGG: Eu³⁺ transparent ceramics as scintillation converters for X-ray imaging.