1. Introduction

Over the last few decades, optical properties of rare-earth (RE) ions doped into various host materials have been studied extensively, mainly with respect to potential application in solid state lighting, displays, lasers, optical amplifiers, sensors, etc [1]. Of the various RE species, europium is traditionally occupying a dominant role. The electronic configuration of Eu is [Xe]4f⁷5d⁰6s². In oxide matrices, it is usually incorporated as Eu³⁺ or Eu²⁺. Depending on valence state, its red (Eu³⁺) or bluish (Eu²⁺) luminescence is widely known and made use of in various applications. In the case of Eu³⁺, photoluminescence typically originates from f-f transitions which are practically independent of ligand field strength. The corresponding emission spectrum comprises relatively narrow bands, in glasses usually dominated by the ⁵D₀ → ⁷F₂ transition at a wavelength of about 612 nm. On the other hand, for Eu²⁺, photoluminescence evolves from f-d relaxation. In this case, corresponding emission bands are typically rather broad and strongly field-dependent (e. g. [2–4]).

Doping Eu³⁺ into glass ceramics may be motivated by various aspects. Most importantly, compared to conventional solid-state reactions, the glass ceramic route enables to benefit from all the advantages of glass processing. That is, glasses with very broad compositional flexibility, very high homogeneity and well-controlled dopant concentration may be prepared (and recycled) by conventional melting, and can, in principal, be processed almost universally [5]. In a second step of thermal annealing, one or more crystalline phases are then precipitated from the glass, preferably by internal nucleation without affecting macroscopic geometry of the as-processed body. The structure of the precipitated crystallite species is typically controlled by composition and annealing procedure. Depending on the availability of sufficiently large (or small) lattice sites, the dopant will at least partially enter (or not enter) the crystalline phase [6]. As a consequence, the spectroscopic properties of the material change. For instance, if the dopant does not enter the crystalline phase, it will ultimately be enriched in the residual glass phase and, e.g., concentration quenching may be observed. Alternatively, emission intensity may be increased as a result of multiple scattering at the newly developed glass-crystal interfaces [7]. Multiple scattering may also occur in the contrary case, if the dopant enters the crystalline phase. As a result of the changing ligand field, distinct changes may occur in the emission spectrum as well as in the lifetime of the excited states. Internal QE may be affected by a change in the phonon-electron interaction and, hence, the probability of non-radiative transfer. Additionally, a change in the absorption cross section may further lead to increasing emission intensity. Often, it is thus highly desirable to, via developing glass ceramics, combine advantages of glasses and polycrystalline materials. From a standpoint of optical application, it is usually even more desirable to prepare glass ceramics with high crystallite number density and low crystallite size. Such nanocrystalline glass ceramics, suitable as host material of Eu-dopants, can be obtained only from a rather limited number of chemical systems [8–11].

Usually, in order to stabilize Eu²⁺, synthesis of the doped material requires a reducing atmosphere (e.g. H₂ or CO). However, it has been shown that in relatively acidic glasses such as Al₂O₃-SiO₂ [12], alkaline earth borate [13] or infiltrated nanoporous silica [14], the redox equilibrium can be pushed towards Eu²⁺ even for melting in oxidizing atmosphere. Similarly, various polycrystalline materials, typically prepared via solid state reaction, are now known where Eu²⁺ can be stabilized even when synthesizing in air, e.g. alkaline earth aluminosilicates [15], BaMgSiO₄ [16], BaAl₂O₄ [17], and Li₂SrSiO₄ [18]. However, little attention has been paid in this respect to Eu-doped glass ceramics. The present work therefore focuses on a novel mixed-valence Eu-doped nanocrystalline glass ceramic material in which, after synthesis in air, the luminescence ratio of Eu²⁺/Eu³⁺ can be controlled by the temperature of annealing and the degree of crystallization.

2. Experimental

Samples of nominal composition (mol.%) 33.3SiO₂-10Al₂O₃-16.7B₂O₃-35BaO-5La₂O₃-0.2Eu₂O₃ (SABBL) were prepared by conventional melting and quenching from a 100 g batch of analytical grade reagents SiO₂, Al₂O₃, H₃BO₃, BaCO₃ La₂O₃ and Eu₂O₃. Melting was performed in alumina crucibles at 1600 °C for 2 hours. Glass slabs were obtained after pouring the melts into preheated (500 °C) graphite moulds and annealing for 2 h. From these slabs, disks of 10x10x1 mm³ were cut and polished (SiC/water). Subsequently, samples were heat-treated on an alumina substrate at temperatures between 800°C and 950 °C (step size 50 °C) for 2 hours in ambient atmosphere. During this procedure, glasses were transformed into translucent glass ceramics. Structural characterization was performed by X-ray diffractometry (XRD, Siemens Kristalloflex D500, Bragg-Brentano, 30 kV/30 mA, Cu Kα) and infrared (IR) absorption spectroscopy (FTIR spectroscopy, Pekin-Elmer 1600) in the wave number range of 400 - 2000 cm⁻¹ with a resolution of 2 cm⁻¹. Photoluminescence was studied with a high-resolution spectrofluorometer and by single photon counting (Horiba Jobin Yvon Fluorolog-3), using a static Xe lamp (450 W) and a Xe flashlamp (75 W) as excitation sources, respectively. All spectroscopic analyses were performed at room temperature. Transmission electron microscopy (TEM) was performed on a Philips CM30 at 300kV. For that, samples were cut in slices, polished, dimpled ans subsequently ion-thinned and coated with a thin carbon layer.

3. Results and discussion

3.1 Structural characterization

XRD patterns of Eu-doped SABBL samples are shown in Fig. 1 . The as-melted specimen does not exhibit any discrete diffraction peaks, confirming its amorphous nature. Annealing resulted in the gradual precipitation of hexacelsian [19], BaAl₂Si₂O₈, (JCPDS card no. 00-012-0725), first observed after heat-treating for 2 h at 850 °C. Subsequently, after treatment at 950 °C, an orthorhombic high-temperature polymorph of lanthanum orthoborate can be observed (JCPDS card no. 00-012-0762). The latter is assumed to be a result of lattice distortion in monoclinic LaBO₃ ([20], JCPDS card no. 00-073-1149) due to the incorporation of impurities such as, eventually, Eu-ions [21]. Both phases can readily be observed by TEM (Fig. 1, right).

 

Fig. 1 XRD pattern of as-made Eu-doped SABBL and corresponding samples after annealing at various temperatures (left). Right: Bright field image showing diffraction contrast of needle-shaped hexacelsian and globular LaBO₃ crystallites. Inset: Corresponding diffraction pattern, resulting from contributions of several crystals with different orientation.

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Corresponding FTIR absorption spectra are presented in Fig. 2 . Spectra were normalized to the intensity of the peak at around 960 cm⁻¹. In Fig. 2(a), the deconvolution of the spectrum of the as-made glass into eight individual peaks is shown. These are attributed to Si-O-Si asymmetric bending of SiO₄ tetrahedra (~450 cm⁻¹ [22]), due to the impact of the Ba²⁺ network modifier ion deconvoluted into two peaks at 441 cm⁻¹ and 481 cm⁻¹ [12], the B-O-B bending vibration in trigonal BO₃, overlapped by vibrations of the AlO₆-group at 702 cm⁻¹ [23,24], B-O-B stretching in tetrahedral BO₄, overlapped by vibrations of Si-O^- non-bridging oxygen at ~950 cm⁻¹ [13], the Si-O-Si asymmetric stretching vibration at 1052 cm⁻¹ [25], asymmetric stretching of the B-O^- non-bridging oxygen at 1233 cm⁻¹ and 1395 cm⁻¹ [13,26] and boroxol rings at 1306 cm⁻¹ [27]. Upon transition to the glass ceramic [Fig. 2(b)], distinct sharpening and splitting of these bands can be observed. I. e., splitting of the Si-O-Si asymmetric bending increases. At the same time, the B-O-B bending vibration of trigonal BO₃ shifts to higher energy and decreases in intensity, and a new band evolves at ~650 cm⁻¹. Also the broad resonance at around 960 cm⁻¹ sharpens, shifts towards higher energy and, upon crystallization, deconvolutes further into at least three distinct bands. The intensity ratio of the bands at ~960 cm⁻¹ and ~700 cm⁻¹ increases with increasing crystallization progress, indicating a transition in boron coordination from trigonal to tetrahedral. Consequently, also the resonances of the asymmetric stretching vibrations of non-bridging oxygen sharpen, and the band centered at ~1233 cm⁻¹ (B-O^- of BO₄) increases in intensity.

 

Fig. 2 FTIR absorption spectra of SABBL samples. In (A) the deconvolution of the spectrum of the as-made glass is shown. (B) shows spectra of samples after annealing at different temperatures.

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3.2 Photoluminescence from Eu³⁺ centers

As noted before, it is well know that in a solid matrix, photoluminescence from Eu²⁺ is dominated by the 4f⁶5d → 4f⁷ transition, while emission from Eu³⁺ ions results from ⁵D₀→⁷F_J (J = 0-4). Both species can thus be unambiguously distinguished by luminescence spectroscopy.

Room temperature excitation spectra of Eu³⁺ (monitoring the 612 nm emission) in Eu-doped SABBL glass and glass ceramics are shown in Fig. 3(a) . Broad excitation peaks correspond to Eu³⁺ transitions from the ground state (⁷F₀) to the indicated excited levels. The most intense excitation peak corresponds to the ⁷F₀→⁵L₆ transition at 393 nm and was used in the following for recording Eu³⁺ emission spectra [Fig. 3(b), relaxation from ⁵D₀ to the indicated states]. As visible from Fig. 3(b), upon heat-treatment, the luminescence properties of Eu³⁺ ions undergo significant changes. That is, firstly, upon crystallization, intensity of all excitation and emission peaks increases notably. In principle, this may be a result of either multiple scattering, increasing absorption cross section or increasing quantum efficiency, eventually caused by the incorporation of Eu³⁺ in a crystalline environment. In a more detailed consideration, it further becomes visible that particularly the emission bands of ⁵D₀ → ⁷F₁ (587 nm), ⁵D₀ → ⁷F₃ (648 nm) and ⁵D₀ → ⁷F₄ at (702 nm) become sharpen and that, relative to the overall increase, their intensity increase is stronger (i.e., the intensity ratio of these peaks to the 612 nm - peak increases). This change is most prominent when the annealing temperature reaches 950 °C and is taken as a clear indicator that at this temperature, Eu³⁺ species are incorporated into one of the crystalline phases.

 

Fig. 3 Eu³⁺-excitation (A) and emission (B) spectra of as-melted glass and SABBL samples after annealing at various temperatures. Inset of (B): Zoom at the spectral region of 500-570 nm.

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The ⁵D₀→⁷F₂ transition is electric-dipole allowed and, hence, the intensity of the corresponding photoemission depends strongly on the symmetry of the Eu³⁺ environment. Contrary, ⁵D₀→⁷F₁ is magnetic-dipole allowed and is independent of local symmetry [9]. As a result, the ratio R of the emission intensities of ⁵D₀→⁷F₂ and ⁵D₀→⁷F₁ can be taken as a probe of the ligand asymmetry in the vicinity of Eu³⁺ ions. High values of R indicate low ligand symmetry and high bond covalency. For the present case, R is plotted as a function of the annealing temperature in Fig. 4 . Its value decreases with increasing annealing temperature from about 3.8 to 1.30. Hence, with higher heat treatment temperature, Eu³⁺ locates in an increasingly symmetric environment of less covalent character.

 

Fig. 4 Asymmetry ratio of the emission intensities of Eu³⁺ transitions of ⁵D₀→⁷F₂ and ⁵D₀→⁷F₁ for as-melted as well as annealed (2h) SABBL samples for different annealing temperatures.

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For the glass ceramic, the emission bands of ⁵D₀ → ⁷F₁ and ⁵D₀ → ⁷F₂ exhibit strong Stark splitting into two or three, respectively, distinct bands. For the as-melted glass, splitting of the ⁵D₀ → ⁷F₁ transition into three peaks indicates the low symmetry of the Eu³⁺ environment. Upon crystallization, these split bands become sharper and gradually overlap. This is a further indicator of the increasing symmetry of the Eu³⁺ environment. A similar conclusion can be drawn from the evolution of the split emission bands of ⁵D₀ → ⁷F₃ and ⁵D₀ → ⁷F₄.

Typically, photoluminescence from Eu³⁺ ions doped into glasses always results from ⁵D₀, regardless of which was the highest excitation band. The reason for this is the occurrence of fast non-radiative relaxation of the ⁵D_J(1-3) states by multiphonon interaction. In the spectral range of 550 to 570 nm, photoemission (from ⁵D_J(1-3)) is therefore very improbable. However, as shown in the inset of Fig. 3(b), for the crystallized samples, photoluminescence from higher lying excited levels ⁵D₁ → ⁷F_{1, 2, 3, 4}, peaking at 525, 535, 551, 560 nm can indeed be observed and the intensity of these bands increases with increasing crystallization. As will be discussed in the following, the reason for this is the lower maximum phonon energy in the vicinity of Eu³⁺ species embedded in the crystalline phase, resulting in less-likely non-radiative energy transfer (see also Fig. 2).

Considering charge and ionic radius, it appears highly probable that the incorporation of Eu³⁺ occurs via substitution of La³⁺ in LaBO₃ [28]. This is further confirmed by the transformation of monoclinic LaBO₃ into an orthorhombic polymorph, La_xEu_1-xBO₃, which appears, as already discussed, related to the presence of a stabilizing impurity species (i.e., in the present case, Eu³⁺). In the consequence, photoluminescence from Eu³⁺ centers is strongly enhanced. Dynamic emission data for the ⁵D₀→⁷F₂ and ⁵D₀→⁷F₁ transitions of Eu³⁺, respectively, are shown in Fig. 5(a) and 5(b). For all samples, the observed decay curves follow a single exponential equation. Corresponding emission lifetimes were found to increase from 1.81 to 2.56 ms (⁵D₀→⁷F₂) and 1.95 to 2.63 ms (⁵D₀→⁷F₁), respectively, from glass to glass ceramic. Longer emission lifetimes as observed for the crystallized specimens indicate lower probability of non-radiative energy transfer. This may originate from the lower maximum phonon energy of XBO₃ which lies within the range of 1296 cm⁻¹ (X = La) and 1040 cm⁻¹ (X = Eu) [28].

 

Fig. 5 Luminescence decay curve of Eu³⁺ emission resulting from ⁵D₀→⁷F₂ (A) and ⁵D₀→⁷F₁ (B) in as-melted Eu-doped SABBL and after annealing at various temperatures. Inset: Exponential plot of the same data, shown for clarity.

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3.3 Photoluminescence from Eu²⁺ centers

Monitoring photoemission at 450 nm, a different set of excitation spectra of Eu-doped SABBL glass and glass ceramics is shown in Fig. 6(a) . The strong excitation band peaking at 350 nm is assigned to the 4f⁷→4f⁶5d¹ transition of Eu²⁺. Corresponding emission spectra (excited at 350 nm) are shown in Fig. 6(b). Emission bands in the spectral range of 570 to 670 nm clearly belong to the already-discussed transitions from ⁵D₀ to ⁷F_J (J = 0, 1, 2, 3 and 4) of Eu³⁺. However, in samples which were annealed at 800 °C or higher, additional broad emission bands appear at around 400 and 510 nm. These are readily assigned to the transition 4f⁶5d→4f⁷ of Eu²⁺ ions (noteworthy, blank reference samples of SABBL glass and glass ceramics did not exhibit any photoluminescence in the considered spectral range). Interestingly, the quantity of divalent europium ions appears to be affected by the annealing temperature: For the temperature regime of 800-900 °C, Eu²⁺-related emission intensity only slightly increases with increasing annealing temperature (whereby the as-melted glass sample did not shown any photoluminescence in this spectral range). When treated at 950 °C for 2 h, a sudden and strong jump is observed in the emission intensity from Eu²⁺ centers. Although while at the same time, photoluminescence intensity from Eu³⁺ centers, as well, increases, the ratio of Eu³⁺/Eu²⁺-related emission intensity strongly decreases with increasing annealing temperature. This is interpreted as a redox transition from Eu³⁺ to Eu²⁺, caused by the annealing process. Hence, the quantity of Eu²⁺ emission centers increases with increasing annealing temperature.

 

Fig. 6 Eu²⁺-excitation (A) and emission (B) spectra of as-melted glass and SABBL samples after annealing at various temperatures. Inset of (B): Zoom at the spectral region of 375-675 nm.

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The decay curve of the Eu²⁺-related photoemission from a crystallized sample of SABBL (950 °C, 2h), monitored at 450 nm, is shown in Fig. 7 . An effective non-exponential emission lifetime of 29.3 μs is obtained from these data. The decay curve can best be fit by a second-order exponential decay function, resulting in two distinct emission lifetimes, i.e. 17.6 and 34.5 µs. This is taken as further evidence for the present of at least two types of Eu²⁺-emission sites.

 

Fig. 7 Luminescence decay curve of Eu²⁺ emission from SABBL after annealing for 2 h at 950 °C. Inset: Exponential plot of the same data, shown for clarity.

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In the following, a possible mechanism for the internal reduction of Eu³⁺ to Eu²⁺ is proposed on the basis of charge compensation [4,18–20]. Firstly, it is assumed that Eu²⁺ is incorporated on Ba²⁺ sites in the BaAl₂Si₂O₈ hexacelsian phase. While in principle, considering ionic radius alone, Eu²⁺ could readily be incorporated on the relatively large interstitials as well as on either of the cationic sites, this assumption is based on charge and coordination equivalence. In order to incorporate Eu-species in this lattice, for reasons of charge balance, two Eu³⁺ ions are needed to substitute for three Ba²⁺ ions, resulting in the formation of one barium vacancy. This vacancy then acts as the donor of electrons to Eu³⁺. The whole process is presented in the following equations:

Only a comparably small number of Eu³⁺ species enters the hexacelsian phase where it is reduced to Eu²⁺. There, it may be present on one of the two differently coordinated (sixfold and eigthfold, respectively) Ba²⁺ sites, what gives rise to the occurrence of the distinct emission bands [Fig. 6(b)] and the fact that the decay curve (Fig. 7) does not follow a single exponential equation. Compared with the Ba²⁺ site in BaAl₂Si₂O₈, the La³⁺ site in LaBO₃ represents an highly suitable host for Eu³⁺ dopants [28].

Hence, the evolution of luminescence properties with increasing degree of crystallization can be understood as follows: For annealing at ~850 °C, the onset of hexacelsian precipitation can be observed (Fig. 1, left). In this temperature range, however, crystallization kinetics appear rather slow. Crystallization is promoted by increasing the annealing temperature. During hexacelsian precipitation, the composition of the residual glass phase shifts towards higher molar content of B₂O₃ and La₂O until at ~950 °C, crystalline LaBO₃ precipitates as the secondary crystalline phase. In accordance to this assumption, TEM indicates precipitation of LaBO₃ primarily at the glass- BaAl₂Si₂O₈ interface (Fig. 1, right). During this process, Eu³⁺ species are first partly incorporated into the hexacelsian phase, where they are immediately reduced to Eu²⁺, occupying one of the two available distinct Ba²⁺-sites. With the subsequent precipitation of LaBO₃, another part of the Eu³⁺ ions is incorporated on La³⁺ sites, leading to the formation of orthorhombic (La_1-x, Eu_x)BO₃ [21,28]. Consequently, photoluminescence occurs simultaneously for Eu²⁺ and Eu³⁺ species. The whole process is overlapped by additional changes in the extent of multiple scattering, especially with the occurrence of the secondary crystalline phase at 950 °C.

4. Conclusions

In summary, the photoluminescence properties of Eu-doped SABBL glasses and glass ceramics were studied. XRD, FTIR and TEM analyses indicate the formation of hexacelsian BaAl₂Si₂O₈ and monoclinic LaBO₃ after annealing at ≥ 850 °C and ≥ 950 °C, respectively. While no Eu²⁺ species are present in the as-melted glass, upon annealing, Eu³⁺ ions are partly incorporated into the hexacelsian phase. There, they are immediately reduced to Eu²⁺. The reduction process may be understood on the basis of a charge compensation model. In the hexacelsian environment, Eu²⁺ ions occupy the two distinct Ba²⁺ sites. LaBO₃ precipitates as secondary crystalline phase and provides a further host for Eu³⁺ species. Their incorporation leads to the stabilization of an orthorhombic polymorph of LaBO₃, and simultaneous photoemission from Eu²⁺ as well as Eu³⁺ centers can be observed. At the same time, Eu³⁺- related photoemission strongly intensifies. The lifetime of the excited state of ⁵D₀ increases as a result of decreasing phonon energy. Spectroscopic properties of the material suggest application in additive luminescent light generation.