1. Introduction

In the past decades, trivalent and divalent rare earth ions doped luminescent materials have attracted great attention due to their promising applications in white light-emitting-diodes (W-LED), temperature sensor, laser, scintillator, and so forth [1–7]. However, the luminescence efficiency of rare earth ions has a great relationship with the phonon energy of the host material [8]. The lower phonon energy of the host is, the higher luminescence efficiency could be. Compared with oxide glass matrix, fluoride host has lower phonon energy, but its poor chemical and mechanical stabilities restrict its practical applications [9]. Therefore, much attention has been devoted to transparent glass ceramics (GCs) which are embedded with fluoride nanocrystals in the oxide glass matrix. The oxyfluoride GC has comparatively low phonon energy because it contains fluoride nanocrystals, and also has high chemical and mechanical stability due to the oxide glass matrix. Consequently, rare earth ions doped transparent GCs containing various kinds of fluoride nanocrystals have been investigated. Compared with CaF₂ [10], BaF₂ [11–13], SrF₂ [14,15], PbF₂ [16], and other difluorides, lanthanide trifluorides nanocrystals are much suitable for rare earth ions doping because the same valence of lanthanide trifluorides can be easily substituted by active rare earth ions [8,17–19].

Rare earth ions doped scintillating materials have been extensively studied due to their applications in high-energy physics, nuclear radiation detection, medicine imaging and other X-ray detection fields [20–23]. Compared with single crystal and ceramic scintillator, rare earth ions (especially Ce³⁺ and Tb³⁺) doped glass scintillating materials have been widely investigated owing to their low cost, easy preparation, and easy to be made into various shapes [22,24,25]. However, the low luminescent efficiency of glass scintillators due to the amorphous structure of the glass hinders their practical applications. Therefore, it is important to improve their luminescence efficiency. Thus, rare earth ions doped scintillating GCs have become promising candidates [20,23,25–28] due to the advantages of glass scintillating materials compared with scintillating crystal, as well as the superiority of high luminescence efficiency compared with scintillating glass.

As one of the most widely studied rare earth ion, Eu³⁺ was often studied for its intense red emission (⁵D₀→⁷F_J transition) and probe function for the rare earth ion site structure (intensity ratio of ⁵D₀→⁷F₂ and ⁵D₀→⁷F₁ transitions) [6,7,17,29]. Moreover, it also could exhibit intense red emission around 610 nm under the excitation of X-ray. Thus Eu³⁺ doped glass have been researched for its potential application as scintillating materials in the recent years. Most of those works are focused on the borosilicate and borogermanate glass system [30,34,35]. However, due to the amorphous structure of those glasses, the luminescence efficiency of them are not so high. Replacing glasses by glass ceramics may be a effective way to improve the luminescence efficiency [26]. But there is rare research on Eu³⁺ doped GC as a scintillating material.

In the present work, Eu³⁺ doped transparent germanate GCs containing LaF₃ nanocrystals were successfully prepared. The incorporation of Eu³⁺ into LaF₃ was evidenced by XRD, PL spectra and decay time measurements. The photoluminescence and the scintillating properties of Eu³⁺ doped GCs were investigated by PL, PLE and XEL spectra.

2. Experimental

Eu³⁺ doped germanate glasses with nominal composition of 58GeO₂-8Al₂O₃-10Na₂O-10LiF-(14-x)LaF₃-xEuF₃ (x = 0.1, 1, 2, 3, 4 and 5, in mol%) were prepared by a conventional melt quenching technique. About 20 g raw materials of GeO₂ (99.999%), Al₂O₃ (99.99%), Na₂CO₃ (99.99%), LiF (99.99%), LaF₃ (99.99%) and EuF₃ (99.99%) were mixed homogeneously in an agate mortar for each batch, and then melt in a covered alumina crucible at 1470 °C for 45 min in an electric furnace in air. The melt was poured onto a preheated stainless steel plate to obtain the bulk glass. In order to release the inner stress, the obtained bulk glass was annealed at 500 °C for 3 h, then cooled to room temperature naturally. Transparent glass ceramics were formed after heat treatment on 3 pieces of 4 mol% Eu³⁺ doped germanate glass (the precursor glass, denoted PG) at 620 °C for 3 h, 640 °C for 3 h and 640 °C for 6 h, which were named as GC620°C-3h, GC640°C-3h and GC640°C-6h, respectively. All of the PG and GC samples were cut and polished for further optical measurements with the size of 10 mm × 10 mm × 2 mm.

The glass transition temperature (Tg) and the crystallization temperature (Tx) of the glass were characterized by differential thermal analysis (DTA) on powdered samples using a NetzschSTA9/C differential scanning calorimeter at a heating rate of 10 K/min. The X-ray diffraction (XRD) patterns were carried out on powdered samples using an X-ray diffraction (Bruker D2 PHASER) with Cu-Kα radiation and 2θ data were collected between 10° and 80° with a 0.1° step size. Transmission electron microscope (TEM) and high-resolution TEM (HRTEM) were used in order to study the microstructure of GCs by a transmission electron microscope (FEI TF20). The transmittance spectra were recorded on a spectrophotometer (Shimadzu UV-3600) in the range of 300-750 nm. The photoluminescence (PL) spectra, photoluminescence excitation (PLE) spectra and fluorescence decay were performed by a fluorescence spectrophotometer (Jobin-Yvon Fluorolog3) equipped with Xe lamp and a Spectra LED as external excitation light source. X-ray excitation spectra (XEL) were recorded by a spectrometer (Ocean Optical QE65000) with a X-ray tube (Copper target, 80 kV, 1.5 mA). All of the structural and optical measurements were performed at room temperature.

3. Results and discussions

The DTA curve of powdered 4 mol% Eu³⁺ doped oxyfluoride germanate glass sample is shown in Fig. 1. The glass transition temperature (T_g) is 528 °C and two exothermic peak with the value at 645 °C (T_x1) and 805 °C (T_x2) are observed. XRD analysis shows the first peak is due to the precipitation of LaF₃ nanocrystal and the latter peak is due to the glass crystallization. The thermal stability factor [ΔT = (T_x-T_g)] of this glass is 117 °C, which indicates the oxyfluoride germanate glass has a good thermal property [1]. Based on these results, two temperatures, 620 °C and 640 °C, near the LaF₃ crystallization temperature were chosen as the crystallization temperature to form transparent glass ceramics.

 

Fig. 1 DTA curve of 4 mol% Eu³⁺ doped PG.

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The XRD patterns of PG, GC620°C-3h, GC640°C-3h and GC640°C-6h are displayed in Fig. 2(a). PG sample shows amorphous structure characterized by the broad hump without any crystalline diffraction peaks of its XRD pattern. On the contrary, after heating the PG at 620 °C and 640 °C for 3h or 6h, sharp diffraction peaks occur in GC samples, and can be easily assigned to the hexagonal LaF₃ (JCPDS No.32-0483) without any other phase. The intensities of these peaks become stronger with the increasing of heat treatment temperature and heat treatment time, which indicates the gradual growth of nanocrystals. The mean size of LaF₃ nanocrystal can be calculated from the peak widths by Scherrer equation [16,36]:

 

Fig. 2 (a) XRD patterns; (b) partial enlargement of XRD patterns in the range of 27° to 29°; (c) experimental and calculated XRD patterns of the Rietveld refinement of GC640°C-6h; (d) Transmittance spectra; (e) Photographs of PG, GC620°C-3h, GC640°C-3h and GC640°C-6h (All samples are 2 mm thick).

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Herein, D is the crystal size, K = 0.89 is a dimensionless shape factor, λ = 0.15406 nm is the wavelength of X-ray, β is the full-width at half maximum (FWHM) of the diffraction peak (derived from XRD pattern) and θ is the angle of diffraction peak, respectively. The calculated sizes of LaF₃ nanocrystals are 24.4 nm, 29.3 nm and 32.1 nm for GC620°C-3h, GC640°C-3h and GC640°C-6h, respectively. The results demonstrate that the mean size of LaF₃ nanocrystal can be easily tuned by changing the conditions of heat treatment. As shown in Fig. 2(b), the diffraction peak of (111) shifts to larger Bragg angle (other diffraction peaks have the same tendency to change), which is due to the incorporation of Eu³⁺ (with ionic radius 0.095 nm) into LaF₃ nanocrystals by substituting La³⁺ (with ionic radius 1.032 nm).

To get the detailed crystal structure information on the obtained GC sample, the experimental, calculated, different XRD profiles and Bragg positions for the Rietveld refinement of GC640°C-6h are presented in Fig. 2(c). According to the refinement data, the GC640°C-6h crystallizes in a hexagonal LaF₃ phase with a P-3c1 space group. Additionally, the lattice parameters were determined to be a = b = 7.19 Å, c = 7.35 Å, α = β = 90°, γ = 120°, V = 328.8 ų, Z = 6, and the refinement finally converged to Rwp = 5.53%, Rp = 4.30% and GOF = 2.18. The results reveal that the actual structure agrees well with the initial structure.

Figure 2(d) shows the transmittance spectra of PG, GC620°C-3h, GC640°C-3h, and GC640°C-6h in the range of 300 nm to 750 nm. It is obviously noted that all the PG and GC samples have a transparency over 71% in the range of 450 nm to 700 nm. In order to display the transparency of the GCs intuitively, the digital photographs of PG and GC samples are shown in Fig. 2(e). Owing to the nanosized LaF₃ crystal is much smaller than the wavelength of visible light, so the obtained GC samples still maintain a high transparency. In addition, the high transparency of GC samples indicates the successful prevention of the collapse of the GeO₂ glassy network after heat treatment and demonstrates the promising prospect of optical application [9].

TEM bright field image and HRTEM image of GC640°C-6h sample are presented in Fig. 3(a) and (b). As shown in Fig. 3(a), nearly spherical nanoparticles (with dark appearance) were distributed among glass matrix. The LaF₃ crystals scatter more electrons due to larger density than glass matrix and therefore appear dark [9]. The size of these nanocrystals as observed in TEM image is ~30 nm, which matches well with the values evaluated from XRD pattern by Scherrer equation. The small crystal size of near 30 nm avoids light scattering and result in the high transparency as show in Fig. 2(e). As verified by Fig. 3(b), the distance of lattice fringes is about 0.317 nm, which is in good agreement with the (111) plane in hexagonal LaF₃ crystal. These results are the further evidences for the precipitation of hexagonal LaF₃ in the glass matrix.

 

Fig. 3 TEM image (a) and HRTEM image (b) of GC640°C-6h.

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Figure 4(a) depicts the emission spectra of precursor germanate glasses with different concentration of Eu³⁺ under the excitation wavelength of 393 nm. As shown in the figure, the emission spectra are composed of four dominant emissions with peaks at 582 nm, 595 nm, 616 nm, 654 nm and 701 nm in the range of 575 nm to 750 nm, which can be readily assigned to ⁵D₀→⁷F₀, ⁷F₁, ⁷F₂, ⁷F₃ and ⁷F₄ transitions of Eu³⁺ ion, respectively. The emission at 616 nm (⁵D₀→⁷F₂ transition) is the predominant and has the strongest intensity among them. Moreover, the emission intensity increases monotonously with the increasing of Eu³⁺ concentration until it reaches 4 mol%, then decreases due to the concentration quenching. The results indicate the optimum concentration of Eu³⁺ is 4 mol%. Thus, herein we choose the glass which is doped with 4 mol% Eu³⁺ for further heat treatment to obtain transparent GCs.

 

Fig. 4 (a) Emission spectra of different Eu³⁺ doped PGs under excitation at 393 nm; (b) Emission spectra of 0.1% mol Eu³⁺ doped PG and GC620°C-3h under excitation at 393 nm.

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To reveal the surrounding structural environment of Eu³⁺ in germanate glass and GCs, the emission spectra of PG and GC doped with 0.1 mol% Eu³⁺ were recorded under the excitation wavelength of 393 nm, as exhibited in Fig. 4(b). Clearly, the emission spectra of the PG and the GC samples are quite different. The emission intensities of ⁵D_{1, 2, 3} to ⁷F_{J (J = 0, 1, 2, 3, 4)} transitions of Eu³⁺ in the GC sample are much stronger than that in the PG sample. For Eu³⁺ ions, the higher excited states ⁵D_{1, 2, 3} will be quenched through the multi-phonon relaxation owing to the high phonon energy of host matrix [31]. So, the reason for the difference in the emission spectra of PG and GC samples can be attributed to the incorporation of Eu³⁺ ions into LaF₃ nanocrystals of lower phonon energy (~300 cm⁻¹) after heat treatment.

The spectra of ⁵D₀ level is also an effective tool for detecting the location of Eu³⁺ in glass ceramic [31]. Due to the intensity of the electric dipolar ⁵D₀→⁷F₂ is depended on the Eu³⁺ ligand field while the magnetic dipolar ⁵D₀→⁷F₁ is not, therefore the ratio (R) of the peak intensities of ⁵D₀→⁷F₂ and ⁵D₀→⁷F₁ transitions is an indicator of the structural environment change of Eu³⁺. Herein, the R of the PG sample is 2.79 while GC sample is 0.93. What is more, it can be seen clearly that Stark split occurs at ⁵D₀→⁷F₂ and ⁵D₀→⁷F₄ transition in GC sample. And a weak emission peak appears at 749 nm in the GC sample, which is due to the ⁵D₀→⁷F₅ of the Eu³⁺ ion [32,33]. All these results confirm that after the crystallization of PG, Eu³⁺ entered into the LaF₃ nanocrystal phase.

Figure 5(a) shows the excitation spectra of 4 mol% Eu³⁺ doped PG and GC samples measured by monitoring 617 nm emission (⁵D₀→⁷F₂ transition of Eu³⁺). The excitation spectra of PG and GC samples consist of a number sharp peaks corresponding to ⁷F₀→⁵D₄ (361 nm), ⁷F₀→⁵L₇ (381 nm), ⁷F₀→⁵L₆ (393 nm), ⁷F₀→⁵D₃ (414 nm), ⁷F₀→⁵D₂ (463 nm), and ⁷F₀→⁵D₁ (531 nm) transitions of Eu³⁺ ions, respectively. In addition, the intensity of each excitation peak increases with the increasing of heating temperature and heating time, which is similar to the emission spectra.

 

Fig. 5 (a) Excitation spectra monitored at 617 nm emission; (b) Emission spectra under excitation at 393nm; (c) Fluorescence decay curves for ⁵D₀→⁷F₀ (617 nm) transition under excitation at 393 nm; (d) XEL spectra of PG, GC620°C-3h, GC640°C-3h, GC640°C-6h and BGO.

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The emission spectra of 4 mol% Eu³⁺ doped PG and GC samples under the excitation wavelength of 393 nm light are displayed in Fig. 5(b). Compared with the PG sample, the intensities of emission peaks of GC samples are increased after crystallization. And the intensity ratios (R) of the peak intensities of ⁵D₀→⁷F₂ and ⁵D₀→⁷F₁ transitions is about 2.68, 1.19, 1.12 and 1.07 for PG, GC620°C-3h, GC640°C-3h and GC640°C-6h, respectively. Impressively, the emission of the ⁵D_{1, 2, 3}→⁷F_{J (J = 1,2,3,4,5)} and ⁵D₀→⁷F₅ transition of Eu³⁺ disappeared in these GC samples which is due to the concentration quenching of Eu³⁺ in LaF₃ nanocrystals [6]. However, despite reaching the quenching concentration, Stark split still occurs at the emission peaks of 617 nm (⁵D₀→⁷F₂) and 701 nm (⁵D₀→⁷F₄).

The fluorescence decay curves for 617 nm (⁵D₀→⁷F₂) of PG and GCs samples under the excitation wavelength of 393 nm are shown in Fig. 5(c). All the curves are well fitted to a double-exponential decay function [14],

where I stand for the intensity of luminescence, A₁ and A₂ are fitting constants, t is the time, τ₁ and τ₂ are fast and slow lifetimes components, respectively. According to the above parameters, the average lifetime (τ*) of Eu³⁺ can be evaluated by:

According to Eq. (3), the average lifetime could be calculated to be 2.14, 2.59, 2.73 and 2.77 ms of PG, GC620°C-3h, GC640°C-3h and GC640°C-6h, respectively. It is obvious that the lifetime of GCs have a slight increase (from 2.14 to 2.77 ms) after crystallization of PG sample. In GC sample, Eu³⁺ entered into LaF₃ crystalline phase with a low phonon energy environment, so reduce the non-radiative relaxation and enhance the lifetime.

With the aim of exploring the potential application of these PG and GC samples in X-ray detection, the XEL spectra of the prepared samples and BGO scintillating crystal were recorded under the excitation of X-ray (80 kV, 1.5 mA) and exhibited in Fig. 5(d). The emission bands with peaks centered at 589, 614, 651 and 700 nm can be readily assigned to ⁵D₀→⁷F₁, ⁷F₂, ⁷F₃, ⁷F₄ transitions of Eu³⁺ ion, respectively. Compared with the emission spectra that under ultraviolet light (393 nm) excitation, the emission peaks of XEL spectra have a slight move to the short wavelength, which is attributed to the different excitation mechanism between ultraviolet and X-ray. Under X-ray excitation, the energy of X-ray has a direct interaction with electrons and holes in the matrix material, producing secondary electrons. Then, the energy of secondary electrons is transferred to the luminescent centers (Eu³⁺). Finally, the excited luminescent centers relax and emit the characteristic scintillating emission. On the contrary, Eu³⁺ were excited directly by ultraviolet light [34,35]. It is worthy to note that, the integrated XEL intensities (the integral of the emission peaks) of PG, GC620°C-3h, GC640°C-3h and GC640°C-6h samples in the range of 300 to 800 nm are 12%, 13%, 18% and 20% of that of the BGO crystal, respectively. It is clearly that the emission intensity is enhanced after crystallization. And the integrated XEL intensities are much higher than that of the reported in the literature [35].

4 Conclusions

In summary, Eu³⁺ doped transparent GC containing LaF₃ nanocrystals have been successfully prepared by melt-quenching technique with subsequent heat treatment. The XRD and TEM results show that hexagonal LaF₃ were precipitated in glass matrix. The transmittance spectra show that all the GC samples still remain a high transparency over 71% in Vis light region. Characteristic emission spectra of Eu³⁺ in GC show that Eu³⁺ were incorporated into LaF₃ nanocrystals and the emission intensity enhanced after crystallization. Additionally, the integrated XEL intensity of GC640°C-6h is about 20% of that of the BGO crystal scintillator. The results indicate that Eu³⁺ doped transparent germanate GC containing LaF₃ nanocrystals could be a promising scintillating material for X-ray detection.