Intense photoluminescence was observed from mixed-valence Eu-doped nanocrystalline BaAl2Si2O8/LaBO3 glass ceramics. For preparation in air, the ratio between Eu3+ and Eu2+ luminescence depends on synthesis temperature. XRD, TEM and IR absorption spectra were employed to characterize the crystallization process and structural properties of the precursor glass and corresponding glass ceramics. When annealed at 950 °C, the material exhibited photoluminescence more than ten times stronger than was found in its glassy state. Spectroscopic data indicate that during such a heat treatment, even in air, a significant part of the Eu3+ ions is reduced to Eu2+. Lifetime of the 5D0 state of Eu3+ increases with increasing heat treatment temperature. Eu3+ species are largely incorporated on La3+ sites in LaBO3 crystallites whereas Eu2+ locates on Ba2+ sites in the hexacelsian phase. A mechanism for the internal reduction of Eu3+ to Eu2+ is proposed. Spectroscopic properties of the material suggest application in additive luminescent light generation.
© 2010 OSA
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 . 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 . 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 . 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 . 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 , alkaline earth borate  or infiltrated nanoporous silica , 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 , BaMgSiO4 , BaAl2O4 , and Li2SrSiO4 . 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.
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 , 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 (, JCPDS card no. 00-073-1149) due to the incorporation of impurities such as, eventually, Eu-ions . Both phases can readily be observed by TEM (Fig. 1, right).
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 ), due to the impact of the Ba2+ network modifier ion deconvoluted into two peaks at 441 cm−1 and 481 cm−1 , 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 , the Si-O-Si asymmetric stretching vibration at 1052 cm−1 , 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 . 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.
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→7FJ (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.
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 . 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.
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 5DJ(1-3) states by multiphonon interaction. In the spectral range of 550 to 570 nm, photoemission (from 5DJ(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 → 7F1, 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 . This is further confirmed by the transformation of monoclinic LaBO3 into an orthorhombic polymorph, LaxEu1-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) .
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 7FJ (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.
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.
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 .
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 (La1-x, Eux)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.
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.
The authors would like to acknowledge the funding of the Deutsche Forschungsgemeinschaft (DFG) through the Cluster of Excellence Engineering of Advanced Materials.
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