J. Phys. Chem. C 2008, 112, 18943–18947
18943
Structure and Optical Spectroscopy of Eu-Doped Glass Ceramics Containing GdF3 Nanocrystals Daqin Chen, Yuansheng Wang,* Yunlong Yu, and Ping Huang State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou, Fujian 350002, China ReceiVed: September 10, 2008; ReVised Manuscript ReceiVed: October 13, 2008
Transparent glass ceramics containing hexagonal GdF3 nanocrystals were fabricated by melt-quenching and subsequent heating of glass with composition of 44SiO2-28Al2O3-17NaF-11GdF3. The GdF3 nanocrystals with mean size of 6 nm were precipitated during quenching of the melt. Increasing the heating temperature of the quenched sample resulted in the further precipitation and growth of the GdF3 particles. The room temperature photoluminescence spectra of the glass ceramic excited at 398 nm show that, in addition to the characteristic sharp emissions of Eu3+ ions, an intense wide blue emission band ascribing to the Eu2+ 5d f 4f transition also appeared. The obvious Stark splitting of the emissions, the low forced electric dipole 5D0 f 7F2 transition, and the long decay lifetimes of Eu3+ evidenced the partition of Eu3+ ions into the GdF3 nanocrystals, while Eu2+ ions were found to reside in the glass matrix. The excitation and emission spectra demonstrated that the excitation energy of Gd3+ was transferred to both Eu2+ and Eu3+. 1. Introduction Rare earth (RE) ion doped luminescent materials have attracted a great deal of interest due to their potential applications in phosphor, X-ray imaging, scintillator, laser, and threedimensional solid-state display.1-4 Generally, the luminescent efficiency of these materials is limited by the dynamics of RE ions which depends on the interactions between RE ions and the host. Lanthanide trifluorides are very suitable hosts for doping RE ions because the lanthanide ions could be substituted easily by RE ions with the same valence, and more significantly, they have low phonon energy that makes it possible to reduce the nonradiative de-excitation probability of the luminescent RE ions by the multiphonon relaxation. However, the stability and machinability of fluorides still remain problems, which restrict their practical applications as the luminescent hosts. A good solution would be the RE ions doped transparent oxyfluoride glass ceramics (TGC) with low phonon energy fluoride nanocrystallites embedding among an oxide glassy matrix of excellent chemical and mechanical performances. Such a nanostructured composite is achieved by controlled crystallization of the precursor glass with appropriate chemical compositions, and the key factor for the efficient luminescence is the partition of the optically active RE ions into the precipitated fluoride nanocrystals.5-10 Due to the much smaller size of the precipitated crystals than the wavelength of the visible light, or the matching of the refractive indexes between nanocrystals and glassy host, the TGC exhibits high transparency. The preparation and luminescent properties of the REdoped LaF3 and YF3 nanocrystal embedded glass ceramics have been well studied.11-20 Unlike La3+ and Y3+ ions, Gd3+ ions could act as the sensitizer to transfer energy to the other RE ions to enhance their luminescence;21 for example, Gd3+ can absorb one vacuum ultraviolet photon excited in the 6GJ levels and relax through two-step energy transfer to Eu3+, yielding two visible photons with quantum efficiency approaching 200% * To whom correspondence should be addressed. E-mail: yswang@ fjirsm.ac.cn.
at room temperature.22,23 However, few investigations on optical behaviors of the glass ceramics containing RE-doped GdF3 nanocrystals were reported so far. In the present paper, the preparation, the selective partition of Eu ions and the optical spectroscopy of the Eu-doped TGC containing hexagonal GdF3 nanocrystals were systematically investigated. 2. Experimental Details The composition of the precursor sample was 44SiO228Al2O3-17NaF-11GdF3 (in mol%). The RE-doping was introduced by the addition of EuF3 with nominal concentration of 0.1%. The precursor sample was fabricated by melting a mixture of the reagent grade chemical compositions in a platinum crucible at 1400 °C for 30 min, and then pouring it into a 300 °C preheated copper mold to cool down to room temperature. To reduce the loss of the fluoride in the course of melting, a small amount of carbon powder was added to create a reductive atmosphere in the furnace. The obtained samples were then cut into 5 mm2 square coupons and heat-treated at 610, 630, 650, and 670 °C, chosen from the differential scanning calorimety (DSC) measurements, for 2 h respectively, to prepare the glass ceramics (donated as GC610, GC630, GC650, and GC670) through fluoride crystallization. DSC experiments of the as-quenched precursor sample were performed in air at a heating rate of 10 K/min in order to follow its thermal behavior. To identify the crystallization phase and determine the mean size of the crystallites, X-ray diffraction (XRD) analysis was carried out with a powder diffractometer (DMAX2500 RIGAKU) using Cu KR radiation (λ)0.154 nm). The microstructures of the samples were studied using a transmission electron microscope (TEM, JEM-2010) equipped with an energy dispersive X-ray (EDX) spectroscopy system. The emission, excitation spectra and the decay curves of the TGC were recorded on an Edinburgh Instruments FLS920 spectrofluoremeter equipped with both continuous (450 W) and pulsed xenon lamps. All the measurements were carried out at room temperature.
10.1021/jp808061x CCC: $40.75 2008 American Chemical Society Published on Web 11/08/2008
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Figure 1. DSC trace of the as-quenched precursor sample recorded at a heating rate of 10 K/min.
Figure 2. XRD patterns of the Eu-doped samples heat-treated at various temperatures; the inset shows the photograph of glass ceramic sample heat-treated at 670 °C for 2 h.
3. Results and discussion The DSC trace of the as-quenched precursor sample is presented in Figure 1. Two exothermic peaks are observed locating around 610 and 780 °C respectively. XRD analysis demonstrated that the former peak is ascribed to the precipitation of GdF3 and the latter the bulk crystallization of the glass matrix. The large difference up to 170 °C between the crystallization temperatures for the GdF3 phase and the bulk matrix provides the possibility to control the precipitation and growth of the GdF3 crystallites among the glass matrix. The XRD patterns of the as-quenched precursor sample and the glass ceramics heat-treated at various temperatures are shown in Figure 2. For the precursor sample, besides the amorphous diffuse humps, several diffraction peaks corresponding to the hexagonal GdF3 also appear, indicating the precipitation of GdF3 during the glass-melt quenching. Since the XRD data for the hexagonal GdF3 have not been reported, we indexed this phase according to the isostructural LaF3 (JCPDS No. 84-0942). With the increase of heating temperature, the GdF3 diffraction peaks tend to be sharper. Based on the peak widths, the mean crystallite size was evaluated by the Scherrer equation
D)
0.9λ β cos θ
(1)
where λ is the X-ray wavelength, β is the full width at halfmaximum of the peak, and θ is the diffraction angle. The determined mean crystallite sizes are 5, 10, 12, 16, and 18 nm for the precursor and the GC610, GC630, GC650, and GC670 samples, respectively. The photograph of the investigated sample obtained by heating the glass at 670 °C for 2 h, presented in the inset of Figure 2, verifies the high transparency of the glass ceramic. High-resolution TEM images and the corresponding selected area electron diffraction pattern of the precursor and the GC670 samples are presented in Figure 3, which demonstrate 3-6 nm GdF3 nanocrystals already precipitated in the glass
Figure 3. TEM images of the precursor (a) and the GC670 (b) samples; the inset of (b) is the selected area electron diffraction pattern.
matrix of the as-quenched precursor sample, and they grew up to 20-30 nm when the sample was subsequently heated at 670 °C, in consistent with the XRD result. The room temperature excitation and emission spectra of the as-quenched precursor and the GC670 samples are presented in Figure 4. There are two sets of excitation signals appearing in the excitation spectra shown in Figure 4a for the 592 nm (Eu3+ 5D0 f 7F1) emission: one consists of several characteristic Eu3+ excitation peaks from the 7F0 ground-state to the indexed excited states, and the other consists of two excitation bands at 274 and 310 nm respectively, originated from the Gd3+ 8S7/2 f 6I and 8S 6 J 7/2 f PJ transitions, which demonstrates the existence of energy transfer from Gd3+ to Eu3+. In addition, a broad excitation band peaking at 235 nm, corresponding to the O2-Eu3+ charge transfer (CT) excitation, is also observed in both excitation spectra. It is interesting that the CT band of the GC670 is much weaker than that of the as-quenched precursor sample.
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Figure 5. Room temperature excitation (λem ) 450 nm) (a) and emission (λex ) 335 nm) (b) spectra of the 670GC sample.
Figure 4. Room temperature excitation (λem ) 592 nm) (a), and emission (λex ) 398 nm) (b) spectra of the as-quenched precursor and the GC670 samples. The symbols J-J′ represent the transitions from the excited level 5DJ to the lower levels 7FJ′ of the Eu3+ ions. The inset shows the dependence of intensity ratio β on heating temperature.
Figure 4b exhibits the emission spectra with 398 nm excitation for the precursor and the GC670 samples. The transitions from the excited 5DJ levels to the lower 7FJ′ levels of Eu3+ are observed. The inhomogeneous broadening of the spectrum expected for the glass appears in the precursor sample. However, the spectrum changes remarkably, i.e., the emission bands become structured and narrowed, suggesting the presence of Eu3+ in crystalline environment, after heat treatment at 670 °C. The intensity ratio (β) of the 5D0 f 7F2 and 5D0 f 7F1 transitions is determined by the symmetry of the crystal sites in which Eu3+ ions are located. The intensity of the magnetic dipolar 5D f 7F transition does not depend on the ligand field of Eu3+, 0 1 while the electric dipolar 5D0 f 7F2 one is known to be forbidden in the centrosymmetric environment.24 Therefore, it can be concluded that the decreasing in intensity for the 5D0 f 7F transition is related to an increase in the symmetry of the 2 ligand field for the Eu3+ ions, due to further precipitation and growth of the hexagonal GdF3 crystallites verified by XRD and TEM. As presented in the inset of Figure 4b, theβ value decreases rapidly from 1.6 to 0.9 after heating at 610 °C, and then keeps almost a constant with the increase of heating temperature from 610 to 670 °C, indicating the crystal field environments of Eu3+ are almost unchanged for the samples heat treated at temperatures above 610 °C. Besides the sharp emissions of Eu3+, a very intense broadband emission centered at 450 nm is also recorded when the samples are excited at 398 nm, as shown in Figure 4b, which is believed originated from the Eu2+ 5d f 4f transition. A similar broadband emission was also reported in the Eu2+ doped transparent glass ceramics containing CaF2 nanocrystals.25,26 The characteristic excitation and emission spectra of Eu2+ ions in the GC670 sample are illustrated in Figure 5. The excitation
Figure 6. Fluorescence decay curves of the Eu3+ 5D0 level for the as-quenched precursor and the GC670 samples.
band with a maximum at 335 nm corresponding to the Eu2+ 4f f 5d transition is recorded when monitored at 450 nm, while a very intense blue emission band centered at 420 nm corresponding to the Eu2+ 5d f 4f emission appears when excited at 335 nm. The luminescence decay curve for Eu2+ exhibits a single exponential decay behavior with a lifetime of 3 us, in good agreement with the result previously reported.27 Fu et al.25,26 found that the Eu2+ emission in the glass ceramic containing CaF2 nanocrystals was much stronger than that in the precursor glass, due to the partition of Eu2+ ions into the CaF2 crystalline phase. However, for the presently investigated samples, the emission intensity and lifetime of Eu2+ are not dependent on the heating temperature, revealing that no Eu2+ ions entered into the GdF3 nanocrystals. This result is reasonable because the radius of Eu2+ (0.139 nm) is far larger than that of Gd3+ (0.120 nm), and the charge compensations are needed for the substitution of Gd3+ by Eu2+ in the crystal lattice. The experimental decay lifetimes for the Eu3+ excited states (5D0, 5D1, 5D2, and 5D3) provide important information to evaluate the environments of the Eu3+ ions. Figure 6 shows the fluorescence decay curves of the 5D0 level for the asquenched precursor and the GC670 samples. The decay curve for the GC670 sample can be approximately fitted by a single exponential function with a lifetime of 9.5 ms, while that for the precursor sample can be fitted by a two-exponential function with one lifetime of 9.2 ms and the other of 4.1 ms, suggesting the presence of two Eu3+ sites, i.e., one in the glass matrix (with shorter lifetime) and the other in the crystalline environment (with longer lifetime). Because of the nonexponential nature of the decay curve for the precursor sample, the experimental decay
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Figure 8. Dependence of the intensity ratio of 5D0 f 7F2/5D0 f 7F1 (β) on the delayed time for the precursor and the GC670 samples.
Figure 7. Room temperature time-resolved luminescence spectra of the precursor (a) and the GC670 (b) samples with various delayed time.
TABLE 1: Experimental Lifetimes of the Eu3+ 5DJ (J ) 0-3) Levels for the As-Quenched Precursor and the GC670 Samples excited levels of Eu3+ 5
D0 D1 5 D2 5 D3 5
τexp (ms) as-quenched
τexp (ms) GC670
6.3 2.9 1.5 0.6
9.5 3.9 2.6 0.9
time was calculated by τexp ) ∫ I(t) dt/Ip, where I(t) represents the function of the luminescence intensity on the time t and Ip is the peak intensity in the decay curve. The obtained lifetime of 6.3 ms is shorter than that of the GC670 sample, owing to the partition of more Eu3+ ions into the low phonon energy GdF3 nanocystals after heating. Similar results are also observed for the other excited states of Eu3+, as demonstrated in Table 1. To verify the results of the fluorescence decay, the timeresolved emission spectra of the precursor and the GC670 samples were measured with 398 nm excitation, as presented in Figure 7. The emission intensity originated from 5DJ (J ) 0-3 respectively) decreases when the delayed time is set increasing gradually from 0.1 to 12 ms. Due to the short lifetimes of Eu3+ ions resided in the glass matrix, the decay of these Eu3+ ions is faster than that of others partitioned in the GdF3 nanocrystals. Therefore, an obvious change in the intensity ratio of 5D0 f 7F2/5D0 f 7F1 occurs for the precursor sample, i.e., the intensity ratio decreases from 1.60 to 0.75 with delayed time increasing from 0.1 to 12 ms, as shown in Figure 8. Combined with the XRD and TEM results, it is concluded that some of the Eu3+ ions were incorporated into the GdF3
Figure 9. Room temperature emission spectra of the precursor and the GC670 samples with the delayed time of 0.1 ms, under 274 nm excitation (Gd3+:8S7/2 f 6IJ); the inset presents the energy level diagram of Gd3+, Eu2+ and Eu3+ ions in the glass ceramics, showing the possible energy transfer processes from Gd3+ to Eu2+ and Eu3+; the solid arrows denote the absorption and emission transitions, the curved arrows denote the energy transfers, and the dotted arrows denote the nonradiative transitions.
nanocrystals during quenching of the melt. As for the GC670 sample, the variation of the intensity ratio with the increase of the delayed time is not severe since most of the Eu3+ ions have been partitioned into the grown GdF3 crystallites after heating, inconsistent with the fluorescence decay measurement stated above. As demonstrated by the excitation spectra of Eu3+, the energy transfer from Gd3+ to Eu3+ has been verified. Figure 9 shows the room temperature emission spectra of the as-quenched precursor and the GC670 samples with delayed time of 0.1 ms, under 274 nm excitation of the Gd3+:8S7/2 f 6IJ transition. The 0.1 ms delayed time was chosen in order to eliminate the influence of the self-excited Eu2+ emission. Both spectra exhibit the characteristic emissions of Eu2+ and Eu3+ ions, indicating the existence of the energy transfer from Gd3+ to Eu2+ and Eu3+, which is depicted in the inset of Figure 9. With 274 nm excitation, Gd3+ ions were excited to 6IJ states and then decayed nonradiatively to the 6PJ states of lower energy. Because the Gd3+:6PJ, Eu2+:5d and Eu3+:5HJ states are energetically close to each other, the energy transfer from Gd3+ to Eu2+ or Eu3+ can easily proceed. Consequently, the Eu2+ ions are excited to the 5d state from which the blue emission occurs; and the Eu3+ ions are excited to the 5HJ states and then decay nonradiatively to the 5DJ states from which the sharp 4f f 4f emissions arise.
Eu-Doped Glass Ceramics 4. Conclusions The GdF3 nanocrystals embedded transparent glass ceramics were successfully prepared by melt-quenching and subsequent heating. Doping Eu into the nanocomposite resulted in the characteristic emissions of both Eu3+:4f f 4f and Eu2+:5d f 4f transitions. Evidently, the Eu3+ ions mainly partitioned into the GdF3 nanocrystals while the Eu2+ ions resided in the glass matrix for the glass ceramics heat treated at 610∼670 °C. The energy transfers from Gd3+ to Eu2+ and Eu3+ were verified. Acknowledgment. This work was supported by NSFC (50672098), the Science & Technology Projects of Fujian (2007HZ0002-2, 2008F3114, and 2006L2005), the projects of CAS (KJCX2-YW-M05) and FJIRSM (SZD07004, 2006K02), the Knowledge Innovation Program of CAS, and the State Key Laboratory of Structural Chemistry (20080039). References and Notes (1) Auzel, F. Chem. ReV 2004, 104, 139. (2) Wang, X.; Zhuang, J.; Peng, Q.; Li, Y. D. Inorg. Chem. 2006, 45, 6661. (3) Yan, R. X.; Li, Y. D. AdV. Funct. Mater. 2005, 15, 763. (4) Zeng, J. H.; Su, J.; Li, Z. H.; Yan, R. X.; Li, Y. D. AdV. Mater. 2005, 17, 2119. (5) Wang, Y. H.; Ohwaki, J. Appl. Phys. Lett. 1993, 63, 3268. (6) Fan, X.; Wang, J.; Qiao, X.; Adam, J. L.; Zhang, X. J. Phys. Chem. B 2006, 110, 5950. (7) Dantelle, G.; Mortier, M.; Vivien, D.; Patriarche, G. Chem. Mater. 2005, 17, 2216. (8) Driesen, K.; Tikhomirov, V. K.; Go¨rller-Walrand, G.; Rodriguez, V. D.; Seddon, A. B. Appl. Phys. Lett. 2006, 88, 073111.
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