Structural and Luminescence Properties of Gd2O3:Eu3+ and ...

Journal of The Electrochemical Society, 152 共9兲 G707-G713 共2005兲

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0013-4651/2005/152共9兲/G707/7/$7.00 © The Electrochemical Society, Inc.

Structural and Luminescence Properties of Gd2O3:Eu3+ and Y3Al5O12:Ce3+ Phosphor Particles Synthesized via Aerosol Olivera Milosevic,a,z Lidija Mancic,a Maria Eugenia Rabanal,b Jose Manuel Torralba,b Bairui Yang,c,d and Peter Townsendc a

Institute of Technical Sciences of Serbian Academy of Sciences and Arts, 11000 Belgrade, Serbia and Montenegro, Yugoslavia University Carlos III, 28911 Leganes, Madrid, Spain c Scitec, University of Sussex, Brighton BN1 9QH, United Kingdom d Department of Physics, Beijing Normal University, Beijing 100875, China b

Opportunities for the synthesis of ultrafine spherical particles with uniformly distributed components and phases are of special importance when materials for photonic application are considered. In this study, the nanophase, spherical, luminescent polycrystalline Gd2O3:Eu and YAG:Ce particles were synthesized from aerosols of the corresponding nitrate solutions ultrasonically generated with frequencies of 1.7 MHz and 800 kHz, respectively. Detailed phase and structural analysis, compositional homogeneity, and particle morphology were determined in accordance to X-ray diffraction 共XRD兲, scanning electron microscopy 共SEM兲, and energy dispersive X-ray spectroscopy 共EDS兲. Quantitative SEM/EDS analysis indicated high material purity and compositional homogeneity. The phase development and structural changes imply nanocrystalline inner structure 共crystallites below 60 nm after thermal treating兲, which influences luminescence behavior. Luminescence measurements indicate that both the radioluminescence and thermoluminescence emission spectra of the Gd2O3:Eu are totally dominated by the line emission characteristic of the Eu, fully substituted onto host lattice sites. By contrast, the cerium dopants in YAG display only very weak Ce luminescence after annealing, but there are broad emission bands characteristic of host lattice defect sites. The spectra are clearly altered by heat treatments and the thermoluminescence indicates a range of independent defect emission sites. © 2005 The Electrochemical Society. 关DOI: 10.1149/1.1972319兴 All rights reserved. Manuscript submitted February 7, 2005; revised manuscript received April 28, 2005. Available electronically July 21, 2005.

Phosphors represent an inorganic crystalline material capable of emitting useful quantities of radiation in the visible or ultraviolet regions, or both, of the spectrum when bombarded with an external energy source such as an electron beam or photons.1,2 The origin of such behavior is the interaction between atomic states associated with the luminescent center and the host lattice material. Both rare earth 共Eu2+,3+, Ce3+, Tm3+, Tb3+, Nd3+兲 and transition metal 共Cr3+, Mn2+兲 ions are commonly used as luminescence centers.3,4 Due to their efficient luminescence properties, phosphor particles are utilized in cathode ray tube 共CRTs兲 screens for televisions and similar devices. More recently, phosphors powders have been used in a number of advanced displayed devices including flat panel displays 共FPD兲, liquid crystal displays 共LCDs兲, plasma displays 共PD兲, thick and thin film electroluminescent 共TFELs兲 displays, and field emission displays 共FED兲.1 The most important phosphor properties are the brightness they exhibit when excited, their spectral energy distribution, and the decay time.5 For many applications, the phosphor phase thermal stability is important.2 With regard to that, considerable interest in advanced phosphor materials today is reflected by progressive development of many convenient and reliable methods for producing them.6,7 In this context, the key for the development of highresolution displays which can achieve superior luminescence performance and lifetime requirements requires that several structural and morphological aspects for the constituent phosphor particles be fulfilled. These include a uniform distribution of the luminescent centers in the matrix of the host material, crystal structure formation having high crystallinity, a small particle size with a narrow particle size distribution, particles with a large surface area, spherical particle morphology, and the absence of agglomerates.8-10 It was reported that the luminescence and resolution exhibit improved behavior on the application abilities when submicronic, unaggregated, uniformly sized particles are considered in comparison with conventionally processed coarse-grained phosphors.8 Controlled surface chemistry is of considerable importance, especially in developments of waveguide lasers.11 Compared to other processing techniques, powder synthesis through aerosol routes enables the generation of fine, submicrometer

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to nanoscale powder, either single or complex, from a variety of precursor solutions.12-18 The processes involve formation of discrete droplets of precursor solution in the form of an aerosol, and the control over their thermally induced decomposition and phase transformation. Aerosols are most frequently formed ultrasonically, using a high-frequency 共100 kHz to 10 MHz兲 ultrasonic beam.19 Liquid atomization and aerosol formation occur for certain values of the acoustic wave amplitude, where the average droplet size depends on the solution properties 共viscosity, surface tension, concentration, density, etc.兲 as well as the ultrasound frequency. Gadolinium oxide becomes an effective crystal phosphor when it is activated by rare earths 共i.e., Eu3+ or Nd3+兲, and it exhibits both cathodoluminescence and laser action.20 Host gadolinia exhibits two polymorphic forms, low temperature 共cubic兲 and high temperature 共monoclinic兲.21 A reversible phase transformation occurs at 1523 K, influencing luminescence behavior.9,15 The garnet phase 共YAG兲 in the complex yttrium-aluminum oxide system is a suitable host material for full-color solid-state phosphors when doped with rare-earth ions.4,22 YAG belongs to the Ia3d space group with general formula C3A2D3O12, in which yttrium ions sit in dodecahedral C sites, while aluminum ions occupy both tetrahedral D site and octahedral A site.23 Yttrium cations can be partially substituted by rare-earth impurities. Several phases are evident in the Y2O3-Al2O3 system; these include Y3Al5O12 共YAG兲 with the garnet structure, YAlO3 共perovskite, YAP兲, and Y4Al2O9 共monoclinic, YAM兲.23 YAG is the phase typically produced by solid-state synthesis above 1873 K. Some authors have implied that the high heating rates and short residence times associated with spray pyrolysis cause either the formation of the kinetically stable YAM or YAP, rather than the YAG phase.24 It was shown that the aerosol synthesis process followed by additional thermal treatment may favor lower decomposition temperatures that allow particle spheroidization and maintain a uniform distribution of yttrium and aluminum, causing YAG phase formation.17,22 In continuation of our previous studies on both Gd2O3:Eu and YAG:Ce phosphor material synthesis via aerosols,15,17,18 the goal of this paper is to investigate the structural, morphological, and spectroscopic properties of the powders obtained under different temperature regime, residence time, and annealing temperatures.


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Table I. Synthesis and annealed conditions for Gd2O3:Eu and Y3Al5O12:Ce powder samples obtained via aerosol. Synthesis Residence timea or annealing Chemical Airflow rate in furnace/on temperature 共K兲 and time 共h兲 Sample composition 共dm3 /min兲 Tmax 共s兲

a

S1 S1-1 S1-2 S1-3

Gd2O3:Eu

1.5

84/31

973 1073-12 1173-12 1273-12

S2 S2-1

YAG:Ce

1.67

3

1173 1473-2

Calculated from the airflow rates and the geometry of the reactors.

Experimental A Gd2O3:Eu source solution having the overall concentration of nitrates 0.1 mol/dm3 was prepared by dissolving, in distilled water, the appropriate amounts of Gd共NO3兲3·x6H2O and Eu共NO3兲3x6H2O in order to obtain the 0.90:0.10 Gd:Eu molar ratio. The solution was atomized ultrasonically 共ultrasonic atomizer with a resonant frequency of 1.7 MHz兲 and then introduced into a high-temperature 共973 K兲 tubular flow reactor using air as a carrier gas 共airflow rate was set at 1.5 dm3 /min, Table I兲. As-prepared powder samples, denoted as S1 in Table I, were additionally thermally treated in a chamber furnace at 1070-1273 K for 12 h. Composite particles based on YAG:Ce were prepared by combining 0.140 mol/dm3 water solutions of Y共NO3兲3·6H2O 共Aldrich, p.a., 99.9%兲 with 0.09 mol/dm3 Al共NO3兲3·9H2O 共Merck, p.a., 99.9%兲 in the desired 3:5 Y:Al molar ratio. Cerium was added in the form of cerium nitrate 共0.005 mol/dm3 Ce共NO3兲·36H2O, Aldrich, p.a., 99.9%兲 in order to achieve 1 at % of doping. The mixture was stirred with a magnetic stirrer at room temperature for approximately 20 min before being added to the ultrasonic atomization chamber. The precursors and precursors’ solution chemistry were carefully controlled by adopting and monitoring the solubility, concentration 共0.06 mol dm−3兲, pH 共2.7兲, viscosity 共1.0436 ± 0.0007 ⫻ 103 Pa s, MLW Viskosimeter B3兲, density 共1.0642 ± 0.0002 g cm−3, AP PAAR calculating density meter DMA55兲, and surface tension 共112.295 ± 0.001 mN m−1, digital tension meter K10T KRUSS兲. Measurements were done both at room 共293 K兲 and increased temperature 共323 K兲. Based on these values the average droplet size 共2.316 ␮m兲, the mean particle size 共745 nm兲 as well as the aerosol droplet number density 共⬍106 cm−3兲 were estimated in accordance with the previously proposed procedure.12 The details of the established synthesis conditions and calculated droplet/particle residence time are presented in Table I. The corresponding precursor solutions were atomized ultrasonically 共atomizing frequency, 800 kHz兲. The aerosol was introduced into a heated zone using air as a carrier gas. Airflow rate was set at 1.67 dm3 /min and the droplet or particle residence time was calculated as to be 3 s 共Table I兲. Aerosol decomposition was performed at 1173 K. As-prepared powders 共S2, Table I兲 were additionally thermally treated in a chamber furnace at 1473 K for 2 h. The experimental setup employed for the aerosol synthesis of both systems is presented elsewhere.15 Crystal phases of as-prepared and thermally treated samples were revealed by X-ray powder diffraction using Philips PW 1710 and 1877 Automated Powder Diffraction with Cu K␣ radiation. Step scanning was 0.02°/15 s. All peak positions were used for the determination of microstructural parameters. Structural refinements were carried out using the Rietveld-based program, Koalariet-Xfit.25 Compositional homogeneity and particle morphology were determined in accordance to scanning electron microscopy 共SEM, Hita-

Figure 1. SEM micrographs of as-prepared Gd2O3:Eu particles.

chi, S-800 and JSM 5300-JEOL兲 and energy dispersive X-ray spectroscopy 共EDX, HORIBA, QX-2000 and Philips XL, Filament W兲. Qualitative and semiquantitative sample analyses were done in accordance to spot and square 共222 ⫻ 322 nm兲 analysis. Luminescence data were recorded with a wavelengthmultiplexed pair of spectrometers. One spectrometer has a grating blazed for the blue region, which activates a UV/blue sensitive photon imaging tube. The second spectrometer is optimized for the green/red region. Spectra are typically recorded every second across the wavelength range 220 to 800 nm. Excitation here was made with a dc 25 kV Cu X-ray unit; heating cycles during radioluminescence 共RL兲 measurements were recorded at a heating, or cooling, rate of 10 K per min over the temperature range from 25 K to room temperature. Thermoluminescence 共TL兲 heating was at the same rate following irradiation at 25 K. Following the data collection the signals were processed to account for the spectral sensitivity dependence of the two spectrometers and their photocathode detectors. Accurate intensity matching at the crossover between spectrometers was not perfect for the type of sample used here; this results in an artifact of a small intensity step near 450 nm in much of the YAG:Ce data. Results and Discussion Gd2O3:Eu-based phosphor particles.— SEM micrographs of asprepared and thermally treated Gd2O3:Eu powder samples are presented in Fig. 1 and 2, respectively. The figures imply formation of individual nonaggregated spherical particles with smooth particle surfaces and filled morphology for the as-prepared particles. Rough particle surfaces are evident after annealing due to the additional promotion of the crystallization process. A narrow particle size distribution is determined for all powdered samples. Qualitative and semiquantitative EDS mapping proved uniform distribution of the constitutive elements inside the particles: gadolinium 共Gd L␣1, M␣兲, europium 共Eu L␣1, M␣兲, oxygen 共O K␣兲, and, for asprepared samples, nitrogen 共N, K␣兲. EDS spot analysis performed in randomly chosen particles in samples from series S1 imply a slightly changeable content of activator ion in different-sized particles, but the uniformity of the presence of the activator ranged from 2.1 to 4.4 atom %, as determined for all samples. A typical XRD powder pattern after annealing together with the resulting difference curve is presented in Fig. 3; this implies the presence of the cubic Gd2O3 phase 共JCPDS file card 43-1014兲. Besides Gd2O3, another cubic phase is evident for as-prepared powders 共12 wt % in S1兲.26 The latter phase is structurally similar to the Gd2Te6O15 共SSI兲共JCPDS file card 37-1400兲, and corresponds to the intermediate solid solution with disordered CaF2 type structure in which both Eu2+ ions can exchange Te4+, and Eu3+ ions can substitute for Gd3+.15 Extensive ion exchange is possible because of a


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Figure 3. Typical XRD pattern of thermally treated 共a兲 Gd2O3:Eu and 共b兲 YAG:Ce 共S2-1兲 powder samples.

order to provide some of the main crystallographic data. The structural changes expressed through average crystallite size and lattice microstrain, presented in Table II, imply a gradual increase of crys-

Table II. Crystal phases and structural parameters (KoalarietXFit) (Ref. 25) for as-prepared and annealed Gd2O3:Eu and Y3Al5O12:Ce powder samples obtained via aerosol.

Sample S1

Figure 2. SEM micrographs of thermally treated Gd2O3:Eu particles: 共a兲 S1-1; 共b兲; S1-2; and 共c兲 S1-3.

close match in their ionic radii: values are for Eu2+: 0.112 nm and Te4+: 0.111 nm; Eu3+: 0.095 nm and Gd3+: 0.094 nm 共all having octahedral coordination兲. The corresponding peak ratio of SSI 共111兲共2␪ ⬇ 27.5°兲-to-cubic gadolinium oxide 共222兲共2␪ ⬇ 28.5°兲-phase, ISSI 共111兲 /I共222兲 is 0.33 for the as-prepared powder samples S1. It seems that this intermediate phase evolution is additionally influenced by the synthesis conditions that enable partially Eu3+ ion reduction to Eu2+, due to the oxygen depletion at the particle surface caused by adsorbed nitrogen.15 The intermediate solid solution phase 共SSI兲 disappeared after annealing above 673 K.15 Welldeveloped peaks of cubic gadolinia are evident in all samples after annealing. Based on these spectra, peak fittings have been done in

S1-1 S1-2 S1-3

S2-1

a

b

Phase 共wt %兲

Unit cell parameters 共Å兲

Gd2O3 SSI,a 12 Gd2O3 Gd2O3 Gd2O3

Cub. Ia3, 10.8040 Cub. Cub. Ia3, 10.8212 Cub. Ia3, 10.8244 Cub. Ia3, -

Y3Al5O12 Cub. Ia3d C,b 71.27 12.0318 ± 0.00004 Y3Al5O12 Cub. Ia3d ICb, 22.94 12.0414 ± 0.00014 CeO2 Cub. Fm-3m C, 5.08 5.4128 ± 0.0004 CeO2 Cub. Fm-3m IC, 0.71 5.4139 ± 0.0016

Crystallite size Microstrain 共nm兲共%兲 16.2 60.7 54.4 -

1.183 0.233 0.214 -

197.5 ± 6.5

0.476

-

1.446

41.9 ± 3.9

0.241

-

0.559

Intermediate solid solution: a phase structure similar to the Gd2Te6O15 共file card JCPDS 37-1400兲. C-crystalline; IC-intercrystalline.

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Figure 5. The thermoluminescence of Gd2O3:Eu at a heating rate of 10 K per min.

Figure 4. Radioluminescence of Gd2O3:Eu with 共a兲 temperature and 共b兲 an example of the spectrum taken at 135 K.

tallinity as well as a decrease of microstrain with increasing annealing temperature. The main feature of as-prepared particles is the presence of the small primary crystallites. The value of 16.2 nm, determined for sample S1, confirms that although SEM reveals visually uniform particle surface, particles are characterized with composite structure and the presence of subgrains. Promoted crystallite growth is evident in thermally treated samples. Particle surface roughness, evident in Fig. 2c, results from the thermally induced crystallite growth, collision, and aggregation into clusters. The estimated values of the cell parameters for the cubic gadolinium oxide phase, presented in Table II, imply the tendency of crystal cell increase with increase of thermal treatment temperature. Gradual increase of crystal cell could be associated with the Eu3+ ion incorporation into gadolinium oxide matrix, because it is a thermally induced diffusion process. The substitution of Gd3+ lattice sites with europium ion in gadolinium oxide matrices is additionally confirmed by TL and RL measurements. The shape of the emission band for as-prepared powder samples is probably related to the intermediate solid solution phase 共SSI兲, as revealed by XRD. This is in agreement with the previously reported results.15 Figure 4a indicates that the radioluminescence spectra, throughout the temperature range 25 to 300 K, contains the same components with relatively minor changes in the ratio of the various europium lines. Spectra taken from the isometric data maps comprise the line component shown in Fig. 4b. The same line pattern occurs in all the samples, independent of the thermal processing treatments, although the intensity of the radioluminescence signal increases as the result of annealing above 1073 K. Following the 1173 and 1273 K anneals, there is a minor enhancement of the main line relative to the others between 25 to ⬃100 K. This might be linked in part to a contribution to the RL from a small thermoluminescence term. Intentional measurement of thermoluminescence following X-ray excitation at 25 K gives an emission pattern, which is totally dominated by the Eu rare earth dopants. Figure 5 shows that there are no indications of other rare-earth lines or any of the broader emission bands consistent with luminescent defect sites of the host

lattice. The pattern of rare-earth impurities efficiently controlling the luminescence decay pathways is typical of many low-temperature TL studies, for example with the rare-earth halide LaF3,27,28 Nd:YAG,3 or from Bi4Ge3O12.29 In those cases there were still residual signals from the host lattice, such as low-temperature TL peaks from excited exciton decay or other intrinsic defects. Such signals were normally suppressed only with high dopant levels. The current behavior of the Gd2O3 therefore suggests that the Eu ions have efficiently occupied lattice sites, and so dominate the decay routes to give solely Eu emission spectra. The main TL peaks are noted near 63 and 127 K. Y 3Al5O12:Ce based phosphor particles.— SEM reveals asprepared particles are unagglomerated and spherical, having smooth particle surfaces and sizes ranging from 300 to 800 nm 共Fig. 6兲. After annealing, the particle size and morphology remain mostly unchanged because the majority of particles persists in its unagglomerated form 共Fig. 7兲. In just a few cases a group of particles appeared in which particle bonding and neck formation is evident. The roughness of the particle surface indicates thermally induced “primary particle” crystallization and growth 共Fig. 7c兲. EDS area analysis proved there is a uniform distribution of the constitutive elements inside the examined particles: aluminum 共Al,K␣兲, yttrium 共Y,L␣兲, cerium 共Ce,M␣兲, and oxygen 共O,K␣兲. The results of the semiquantitative analysis and particle profiling indicate similar composition of the activator inside the different particles 共0.4-0.6 atom %兲. XRD pattern of the as-prepared powder samples indicates poorly crystallized phases: Y2O3 共JCPDS file card 43-1036兲, orthorhombic Y4Al2O9 共JCPDS file card 14-0475兲, and cubic Y3Al5O12 共JCPDS file card 33-0040兲. Peaks at 1.7864, 1.5881, and 1.5106 nm might be connected with 共024兲, 共211兲, and 共018兲共hkl兲 planes of Al2O3 共JCPDS file card 31-0026兲, respectively.17 Table II summarizes the results of the peak fitting done by Koalariet-XFit after annealing at 1473 K together with the weight content 共wt %兲 of certain phases and their structural parameters. As evident, cubic YAG, Y3Al5O12 共JCPDS file card 33-0040兲 is the prevailing phase, although there is some content of cubic CeO2 共JCPDS file card 34-0394兲共Fig. 3b兲. Note that satisfactory peak fitting is obtained even after proposing that the phases obtained are twin domains, where one domain relates solely to the crystalline component and the other to the intercrystalline component that has a high defect content. This is presumably caused by the nanophase particle internal structure. A similar approach has been applied previously to a single-phase metal system comprised of different-sized fractions of particles.30 The higher value of YAG lattice parameters 共Table II兲 compared with the theoretical value 关12.0089共3兲 Å兴 is possibly the consequence of substituting cerium 共CeY兲 for Y in the YAG matrix, because, although there is close ionic radius match within ⬃1.1% 共0.101 nm Ce3+ and 0.09 nm for Y3+兲, the distortions should be considered in terms of the volume mismatch, which is nearer to 3%. However, the associ-


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Figure 6. SEM micrographs of as-prepared YAG:Ce precursor powders 共S2兲.

ated strains from this substitution are partially suppressed with the formation of CeO2, and this route needs to be optimized in further research. Our previous results imply both the decrease in the concentration and increase in the lattice parameter of cubic CeO2 关theoretical value: 5.411 34 共12兲 Å兴 with increasing temperature,17 suggesting the tendency for Ce4+ to be reduced to the larger Ce3+ ion to stabilize the fluorite structure.31 This is probably followed by the diffusion of Ce3+ into the garnet structure, which together with increasing the YAG crystallinity with temperature, influences luminescence behavior. The Ce emissions located in the 480-560 nm spectral range with the broad band centered near 525 nm have been tentatively attributed to the Ce3+ intershell transition 共5d → 4f兲 in the YAG lattice.5,32 The results obtained are in agreement with the ultrasonically synthesized 共 f = 1.75 ⫻ 106 s−1兲 Y:Al = 3:5 system containing 0.5-3 at % Ce.4 However, the shape of the luminescence spectrum obtained, and the peak shifting in comparison to a single crystal,5 are presumably associated with the polycrystalline nanostructure, associated with a high content of interfaces. By contrast with Gd2O3, the emission spectra of cerium-doped YAG show broadband emission spectra with negligible evidence for narrow line emission signals, except after high-temperature annealing. The change might thus be evidence for both diffusion and changes in the charge state. Further, the RL spectra change steadily with heating over the 25 to 300 K range. There is also a very pronounced spectral change cause by annealing the material at 1473 K,

Figure 7. SEM micrographs of thermally treated YAG:Ce powders 共S2-1兲.

and a dramatic reduction of the TL signal in the annealed powder. These features are exemplified by the following figures. Figure 8 contrasts the RL signals taken during cooling of the YAG:Ce. The intensity units are nominally arbitrary for the two samples, but the signals are somewhat stronger for the annealed sample. For the asprepared sample there are at least two broad emission bands which increase during cooling, and the band near 600 nm passes through an intensity maximum near 140 K 共Fig. 8a兲. By contrast, the annealed sample has the main red emission bands of almost constant

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Figure 10. The TL of the as-prepared YAG:Ce. No data are shown for the annealed samples as the signals were extremely weak.

Figure 8. 共a兲 Radioluminescence taken during cooling for the as-prepared YAG:Ce; 共b兲 and the more intense signals seen for YAG:Ce after heating at 1473 K.

intensity throughout the cooling range. There is also evidence in Fig. 8b of a small, wide-line feature near 300 nm, which develops at the lower temperatures. This is a typical emission wavelength for Ce.

Figure 9. 共a兲 Temperature dependence of the RL for the original YAG:Ce and 共b兲 comparison of emission spectra between as-prepared and annealed YAG:Ce.

The spectral slices from the isometric images indicate more clearly for the as-prepared material that the emission spectra are temperature sensitive, and Fig. 9 contrasts spectra at 30, 120, and 275 K. While some component features exist over the range, it is clear that their relative intensities differ with temperature. A far more dramatic difference in emission spectra is seen on comparing the 30 K signals for the as-prepared and high-temperature annealed samples, Fig. 9b. Annealing reduces the long-wavelength signals and greatly enhances those near 550 nm. The defect pattern revealed by thermoluminescence is quite different for the two types of sample. Whereas the as-prepared material has relatively strong TL signals which depend on wavelength, the annealed samples gave almost negligible TL signals, although there was evidence for weak emission in the red near 160 K. The details of the TL for the original powder are first shown in an isometric plot, Fig. 10, with TL curves as a function of wavelength in Fig. 11a,

Figure 11. 共a兲 TL as a function of emission wavelength for as-prepared YAG:Ce; and 共b兲 TL spectra at three temperatures.

Journal of The Electrochemical Society, 152 共9兲 G707-G713 共2005兲 and the spectral changes recorded in Fig. 11b. The interpretation of the luminescence signals from the YAG:Ce powder is that the cerium ions are not readily accommodated on the YAG lattice sites, and well-separated cerium ions on such sites only occur after a high-temperature anneal. Even then the cerium ions do not offer a favorable excitation decay path; instead, there are broad emission bands typical of the host lattice defect sites and a very weak luminescence via the cerium ion pathway near 300 nm. The complexity of the temperature dependence of the TL as a function of wavelength also suggests that a variety of independent defect sites contributes to the TL. Comparison with the TL data of Zych et al.33 is in broad agreement in that they also record a strong TL signal near 130 K; however, their data include many other narrow TL peaks in the low-temperature range. They also indicate numerous anomalies in the temperature dependence of their luminescence lifetime data. Reference to the YAG:Nd results3,32 might suggest that some of the anomalous behavior might be linked to trapped impurity nanoparticles, such as CO2 or water ice; therefore, the more numerous glow peaks that they report may merely indicate that their crystals are less pure than in the present experiment. Conclusion Spherical polycrystalline Gd2O3:Eu and YAG:Ce particles have been synthesized from aerosols. They differ in terms of their responses to thermal treatments, and the Eu dopants produce strong luminescence entirely via the emission lines of Eu, whereas the Ce-doped material is characterized by broadband luminescence spectra. The data indicate that this is an effective route for the production of phosphor material. Acknowledgments The authors are grateful to the Ministry of Science and Technology of the Republic of Serbia for financial support, as well as NEDO International Joint Research Grant Program 01MB7, Japan. References 1. P. D. Rack and P. H. Holloway, Mater. Sci. Eng., 4, 171 共1998兲. 2. R. C. Ropp, Luminescence and the Solid State, Elsevier Science Publishers B. V., New York 共1991兲. 3. M. Maghrabi, P. D. Townsend, and G. Vazquez, J. Phys.: Condens. Matter, 13,

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