J. Phys. Chem. C 2010, 114, 1337–1342
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Powder Synthesis of Y-r-SiAlON and Its Potential as a Phosphor Host Takayuki Suehiro,*,† Hiroaki Onuma,‡ Naoto Hirosaki,§ Rong-Jun Xie,§ Tsugio Sato,† and Akira Miyamoto‡ Institute of Multidisciplinary Research for AdVanced Materials, Tohoku UniVersity, Sendai 980-8577, Japan, Graduate School of Engineering, Tohoku UniVersity, Sendai 980-8579, Japan, and Nano Ceramics Center, National Institute for Materials Science, Tsukuba, Ibaraki 305-0044, Japan ReceiVed: October 9, 2009; ReVised Manuscript ReceiVed: NoVember 27, 2009
The Y-R-SiAlON fine powders have been successfully synthesized from the system Y2O3-Al2O3-SiO2, by using the gas-reduction-nitridation method. Completion of the nitridation was achieved at temperatures 2000 °C are attained within several seconds, which even enabled the preparation of generally metastable R′ phases such as those exclusively stabilized by Ce or Eu with an appreciable phase purity of 25-70%. The product powders were obtained by extensive mechanical grinding (24 h) of the crude reaction products, which is obviously disadvantageous from the viewpoint of optical applications. More recently, we have applied the gas-reduction-nitridation (GRN) method18-21 to the multicomponent oxide system CaO-Al2O3-SiO2, and successfully established a new powder-synthesis route for yellow-emitting Ca-R-SiAlON:Eu2+ phosphors.22-25 The synthesized powders possessed the superior characteristics for phosphor applications such as high phase purity, nonaggregated fine particle morphology, and very low impurity absorption throughout the visible spectral region. The better photolumi-
10.1021/jp9096748 2010 American Chemical Society Published on Web 12/23/2009
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Figure 1. FESEM micrograph of the raw Y-Si-Al-O powder.
nescent property compared to the conventional reaction-sintered powder was demonstrated,22 and the further improvement of the processing techniques attained the high external quantum efficiency of ∼44-55% under near-UV to blue light excitation.23 In the current work, we focused on the powder synthesis of Y-R-SiAlON using the GRN process, and explored its potential as a phosphor host material by Eu2+ doping. The photoluminescent properties were discussed and compared with the Castabilized system, with regard to the results of the first-principles calculations on their electronic structures. 2. Experimental Section 2.1. Powder Preparation by GRN. Homogeneous mixture of the system Y-Si-Al-O was prepared by a coprecipitation method using an amorphous SiO2 powder (SP-03B, Fuso Chemical Co., DBET ) 0.23 µm), for obtaining the target composition of Y0.60Si9.30Al2.70O0.90N15.10 (m ) 1.8, n ) 0.9). The resulting particle morphology is shown in Figure 1. The starting powder was contained in a high-purity alumina boat and set in a horizontal alumina tube furnace (inner diameter of 24 mm). The furnace was subsequently heated to the experimental reaction temperature (1400-1500 °C) at various heating rates in an NH3-1.5 vol % CH4 gas mixture, introduced at a constant flow rate of 1.3 L/min. After the predetermined reaction time, the sample was furnace cooled in an NH3 atmosphere. The extent of nitridation was estimated by measuring the fraction of observed and theoretical weight loss (∆Wobs/∆Wtheor), resulting from the reduction-nitridation reaction. The loss on ignition of the starting powder was predetermined and was taken into account. An additional postsynthesis heat treatment was conducted for the selected samples at 1600-1700 °C for 4 h in a N2 atmosphere. The Eu-doped sample with the nominal composition of Y0.60Eu0.05Si9.30Al2.70O0.80N15.20 was prepared by the postsynthesis activation (PSA) process23 conducted at 1700 °C for 4 h. 2.2. Characterization. Phase assemblage of the synthesized powders was analyzed by X-ray diffractometry (XRD) using Cu KR radiation (RINT2200, Rigaku). The mass fraction of R′ to R′ + β′ phase in the as-prepared samples was determined by a simplified formula.15,26,27 For the heat-treated samples, the multiphase Rietveld analysis was conducted to determine the mass fraction of each phase using the program RIETAN-2000.28 The actual solubility of the Y ion (x value) in the R-SiAlON matrix was determined by the refinement of the site occupancy factor, with the isotropic displacement parameter fixed to the literature value.2 The z value of the secondary β′ phase was estimated from the refined lattice constants.29 Specific surface area and the equivalent particle size of the samples were measured by the multipoint Brunauer-Emmett-Teller (BET) method (Autosorb, Quantachrome). Nitrogen, oxygen, and impurity carbon contents of the selected samples were analyzed by the selective hot-gas extraction method (TC-436, CS-444LS, LECO Co.). Electron micrographs of platinum-coated samples were obtained using a field-emission scanning electron micro-
Suehiro et al. scope (FESEM; JSM-6340F, JEOL). Diffuse reflectance of the undoped powders was measured with a spectrometer (Ubest V-560, JASCO), by referring to the Spectralon reflectance standard. Photoluminescent (PL) properties of the Eu-activated powders were evaluated at room temperature using a fluorescence spectrophotometer (F-4500, Hitachi). Quantum efficiencies and temperature-dependent PL properties were measured at the excitation wavelength of 450 nm using a multichannel spectrophotometer (MCPD-7000, Otsuka Electronics). 2.3. Electronic Structure Calculations. Computation of the electronic structures of Y- and Ca-R-SiAlON:Eu2+ was performed using the periodic quantum chemistry (QC) calculations based on the density functional theory (DFT). The DMol3 program implemented in the Materials Studio version 3.2 was employed. Geometry optimizations were performed at the local density approximation level employing the VWN functional.30 Only the atomic positions in the calculation model were optimized, while the cell sizes were not optimized. Electronic structure analysis and re-evaluations of total energy of optimized structures were performed at the generalized gradient approximation level employing the PBE functional.31 For all the calculations, the double numerical with polarization basis sets and effective core potentials were adopted and the k-point was set to 1 × 3 × 2. The models of Y- and Ca-R-SiAlON:Eu2+ with the compositions of Y3EuSi36Al12ON63 and Ca3EuSi36A112O4N60 were constructed using the structural parameters reported by Izumi et al.2 These models consisted of 1 × 2 × 2 superlattice of each hexagonal unit cell. Their pseudo-orthorhombic forms with the lattice vectors of a ) ah, b ) ah + 2bh, c )2ch were adopted for the calculations (Figure S1). The consequent cell dimensions were 7.83 × 13.56 × 11.42 Å3 and 7.84 × 13.58 × 11.41 Å3 for the Y- and the Ca-stabilized system, respectively. The compositions of the host lattices were set as close to the actual values as possible, with an identical concentration of Eu (25% with respect to the modifying cations). In the model of Ca-RSiAlON:Eu2+, the O atom was placed just above the vacant Ca site (2b) along [001], which was coordinated by the three Al atoms located at the closest neighboring 6c site (z ∼ 0.21 in the unit cell) forming an AlN3O trimer (Figure S1a). For Y-RSiAlON:Eu2+, the O atom was located in the 1 × 2 × 2 supercell constructed in the same manner, with the closest Eu-O distance of 4.59 Å (Figure S1b). 3. Results and Discussion 3.1. Powder Preparation of Y-r-SiAlON by GRN. Table 1 shows the characteristics of the as-prepared powders synthesized under the various reaction conditions. The samples showed rather complicated phase assemblages, composed mainly of β-SiAlON with smaller amounts of R-SiAlON, AlN, Y4Si2O7N2 (J-phase), and Y2Si4N6C,32,33 which was misregarded as “Y3Si6N11” in the earlier CRN experiments.16,34 With the processing temperatures of 1400-1450 °C (samples Y1 and Y2), formation of the highly reduced YN phase was detected, whereas completion of the nitridation was not attained at this low temperature region, resulting in an apparent carbon contamination with a soaking time longer than 2 h. The complete nitridation was achieved with the processing temperature of 1500 °C and the shorter reaction time of 1 h (sample Y3). The R′ phase content (with respect to R′ + β′) of the powders decreased with the increase of the processing temperature, ranging from ∼32% to 18%, which were much lower than the results observed for the Ca-stabilized system (∼90%).22 The sample Y4 obtained with a faster heating rate of 500 °C/h consisted predominantly
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TABLE 1: Characteristics of the As-Prepared Powders Synthesized under the Various Reaction Conditions sample Y1 Y2 Y3 Y4 a
reaction conditions 1400 1450 1500 1500
°C, °C, °C, °C,
2 2 1 1
h h h h
heating rate (°C/h)
phase assemblage
R′ /(R′ + β′)
∆Wobs/∆Wtheor
SBET (m2/g)
200 200 200 500
β′, R′ > AlN, Y2Si4N6C, YN β′ > R′ > AlN, YN, Y2Si4N6C, Ja β′ > R′, AlN, Y2Si4N6C, Ja β′ . Y2Si4N6C, R′, AlN, Ja
0.32 0.23 0.18
0.91 0.96 0.99 1.01
9.97 7.58 7.05 4.78
J: Y4Si2O7N2.
Figure 2. XRD patterns of the powders synthesized at 1200-1500 °C without isothermal holding. The value in parentheses denotes the nitridation extent of each sample.
Figure 3. FESEM micrographs of the powders synthesized at (a) 1200, (b) 1300, (c) 1400, and (d) 1500 °C without isothermal holding.
of β-SiAlON, and possessed a markedly lower specific surface area. This result indicates that the abundant presence of oxygenrich liquid phases during heating promotes the formation of β′ phase and significant particle coarsening, in conformity with the results observed for the Ca system.22 Figure 2 shows the XRD patterns of the samples synthesized at 1200-1500 °C with the heating rate of 200 °C/h without isothermal holding (0 h soaking), which represents the phase evolutions during the reduction-nitridation. The corresponding particle morphologies are also shown in Figure 3. At 1200 °C, the only crystalline phase detected was Y3Al5O12 on the particle surface, while a large part of the core SiO2 remained an amorphous state. The core SiO2 particles reacted partly with surface Al2O3 layers to form β- and O-SiAlON by further heating to 1300 °C, at which the nitridation proceeded through a solid-state reaction, as indicated by no significant changes in the particle morphology (Figure 3b). The extent of nitridation further increased to 64% at 1400 °C, i.e., at above the lowest
ternary eutectic temperature (∼1371 °C35) of the Y2O3Al2O3-SiO2 system, through the partial contribution of the dissolution-precipitation process to form β-SiAlON, due to the persistence of Y3Al5O12 phase. This is evident from the coagulated particle morphologies shown in Figure 3c,d. Upon reaching the soaking temperature of 1500 °C, Y3Al5O12 was decomposed/nitrided into Y4Si2O7N2, Y2Si4N6C, and AlN phases, while a minor amount was consumed to form R-SiAlON by reacting with β′ phase. The system was almost fully nitrided (∼96%) at this stage, and thus provides less possibility of liquid phase formation resulting in only a slight increase in R′ content after the 1 h soaking (sample Y3). While the nitridation kinetics was found to be comparable to the Ca system under the identical processing conditions,22 the formation of pure R′ phase could not be attained in the current system, due to the aforementioned trapping of the modifying cation by the unfavorably stable intermediate phases. 3.2. Effects of Postsynthesis Heat Treatment. The postsynthesis heat treatment was conducted for the nitrided sample Y3, and the resulting powder characteristics are summarized in Table 2. The R′-phase content significantly increased to 83.4% with the heat treatment at 1600 °C, while appreciable amounts of undissolved β-SiAlON (z ) 0.21) and AlN phases still remained (sample YR1). The transformation into the R′-phase further proceeded at the higher temperature of 1700 °C, at which the synthesis of almost phase-pure (98.3%) R-SiAlON powder containing trace amounts of AlN and Y2Si3O3N4 (M-phase) was achieved (sample YR2). To promote the formation of R′-phase through the expected dissolution-precipitation, an addition of YF3 flux (melting point of 1152 °C36) to the reaction system was also attempted (sample YR3), which effectively lowered the temperature required for attaining an equilibrium, resulting in a comparably high phase purity of 96.8% at 1600 °C. The XRD patterns of the near single-phase samples YR2 and YR3 are shown in Figure 4, and their particle morphologies are shown in Figure 5. The sharp diffraction profiles indicate the high crystallinity and compositional homogeneity of the heat-treated powders. The x values determined by the Rietveld refinement was found to be ∼90% of the nominal composition (x ) 0.60) in both cases, suggesting the presence of a very minor amount of Y-containing noncrystalline phases. The heat-treated powders consisted of uniform discrete primary particles of ∼1-2 µm, consistent with the particle sizes determined by the surface area measurement (Table 2). Table 3 summarizes the analyzed nitrogen, oxygen, and impurity carbon contents of the heattreated powders, together with those of the as-prepared sample Y3. The as-prepared powder contained excess oxygen and an appreciable amount (0.56 wt %) of residual carbon, which were reduced via the partial CRN reaction under the heat treatment in N2, resulting in the final compositions with the almost ideal anion stoichiometries. The presence of a minor amount of residual carbon might also prevent the sintering of primary particles, as was indicated by our previous work.23 3.3. Optical and Photoluminescent Properties. The PL properties of the Eu-activated Y-R-SiAlON has been briefly
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TABLE 2: Characteristics of the Powders after the Postsynthesis Heat Treatment sample
heat treatment temperature (°C)
phase assemblage
x
SBET (m2/g)
DBET (µm)
YR1 YR2 YR3c
1600 1700 1600
R′ (83.4%), β′ (10.8%), AlN (4.6%), Ja (1.2%) R′ (98.3%), AlN (1.2%), Mb (0.5%) R′ (96.8%), Mb (2.5%), AlN (0.7%)
0.482(3) 0.535(2) 0.542(2)
1.77 1.32 1.37
1.02 1.37 1.31
a
J: Y4Si2O7N2. b M: Y2Si3O3N4. c With 2.5 wt % (0.10 mol %) YF3 added.
Figure 6. UV-vis diffuse-reflectance spectra of the synthesized Yand Ca-R-SiAlON powders.
Figure 4. X-ray Rietveld refinement patterns of the heat-treated powders: (a) sample YR2, (b) sample YR3.
Figure 5. FESEM micrographs of the heat-treated powders: (a) sample YR2, (b) sample YR3.
TABLE 3: Nitrogen, Oxygen, and Carbon Contents of the As-Prepared and Heat-Treated Powdersa
a
sample
N
O
C
Y3 YR2 YR3
14.50(4) 15.10(4) 15.11(4)
1.50(1) 0.90(1) 0.89(1)
0.28(1) 0.04(1) 0.04(1)
In atomic units.
reported in our previous works37,38 for the reaction-sintered powders. In the current work, the powder with a composition Y0.60Eu0.05Si9.30Al2.70O0.80N15.20 was prepared from the GRNderived sample Y3 via the PSA process23 conducted at 1700 °C for 4 h. The Eu-activated powder thus-obtained possessed the phase purity comparable to the undoped sample, containing only 1.4% of Y2Si3O3N4 and 1.7% of AlN as secondary phases (Figure S2). The average (Y, Eu)-(O, N) bond distance
Figure 7. PL excitation and emission spectra of the synthesized Yand Ca-R-SiAlON:Eu2+ powders.
estimated from the Rietveld refinement was 2.604(4) Å, which showed an insignificant difference from that in the undoped sample (2.602(4) Å) and the Ca-stabilized sample having the similar composition of m ) 1.6 (2.611(9) Å, sample C8R in ref 22). Figure 6 shows the UV-vis diffuse-reflectance spectrum (DRS) of the undoped sample YR2, together with the data for the aforementioned Ca-stabilized sample for comparison. The Y-R-SiAlON powder synthesized by the current process exhibited the comparably high reflectance of ∼88-91% over the whole visible spectral range, as expected from the sufficiently low impurity carbon content attained after the final heat treatment (Table 3). The onset values of the absorption edge derived from the Kubelka-Munk transformed spectra were estimated to be 3.76 and 4.43 eV for the Y- and the Ca-stabilized samples, respectively, indicating the marked narrowing of the band gap (Eg) in the current system. Figure 7 shows the PL excitation and emission spectra of the synthesized Y-R-SiAlON:Eu2+ powder, wherein the data for Ca0.80Eu0.05Si9.60Al2.40O0.70N15.30 prepared under the identical PSA conditions23 is also shown. Both the excitation and emission spectra of Y-R-SiAlON:Eu2+ resemble closely those of the Castabilized counterpart implying that the almost equivalent coordination environment around the Eu2+ ions leads to essentially similar emission properties. The detailed PL properties derived from the observed spectra as well as the results of the quantum efficiency measurement for both systems are summarized in Table 4. The excitation spectrum of the Y system
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TABLE 4: Photoluminescent and Optical Properties of the Y- and Ca-r-SiAlON:Eu2+ Powders Prepared by the GRN quantum efficiencies
CIE coordination
M
Abs
QEint
QEext
x
y
λd (nm)
λpeak (nm)
CFS (eV)
Ecent (eV)
SS (eV)
Eg (eV)
Y Ca
0.749 0.684
0.172 0.568
0.129 0.388
0.566 0.528
0.425 0.466
589.5 582.8
610.6 591.2
1.31 1.01
3.42 3.57
0.73 0.96
3.76 4.43
consisted of two distinct bands peaking at ∼305 and ∼450 nm, which revealed the stronger crystal field splitting (CFS) and the lowered energy centroid (Ecent) of the 5d level of Eu2+ being 1.3 and 3.4 eV, respectively. The emission spectrum showed an appreciably red-shifted peak wavelength of 611 nm, compared to 591 nm in the Ca system. The corresponding Commission Internationale de l’Eclairage (CIE) 1931 chromaticity coordinates were x ) 0.566 and y ) 0.425 with the dominant wavelength (λd) of 590 nm, giving an orangish yellow emission with high color purity. The absorption efficiency (Abs), corresponding to the allowed 4f-5d transition in Eu2+ was even higher in the Y-stabilized system, whereas the internal quantum efficiency (QEint) was found to be a much lower value of 17.2%, resulting in the external quantum efficiency (QEext) being about one-third of that in the Ca system. The significant luminescent quenching observed for the Y system is also reflected in the temperature-dependent PL properties shown in Figure 8. The PL intensity in the Y system decreased steeply to 90% by heating from room temperature to 50 °C, and showed further decrease to 50% at 150 °C, at which the Ca system still retained the high relative intensity of 88%. The PL intensity fell off to 19% by further heating to 300 °C, while it recovered completely the initial value upon cooling, showing the high thermal stability of the host lattice itself. By considering the high lattice rigidity (i.e., low thermal expansion39,40) and the consequent smaller Stokes shift (SS) in the current system, the observed thermal quenching might mainly result from the ionization of excited 5d electron of Eu2+ into the conduction band, caused by the remarkable decrease of the Eg value. The occurrence of photoionization is also suggested by the decreased emission efficiency under the excitation below ∼350 nm region, as indicated by the PL excitation spectrum shown in Figure 7. The above results conclusively showed the practical limitation of the use of pure Y-R-SiAlON:Eu2+ for the white LED applications, whereas (Y, Ca)-R-SiAlON:Eu2+37,38 substituted with an appropriate amount of Y exhibited an effectively redshifted emission and moderate thermal stability (typical relative PL intensity of ∼60-80% at 150 °C), which may enable it to be used as a supplementary phosphor for the multichromatic white LEDs possessing an improved color rendering property. 3.4. Electronic Structure Calculations. The electronic structure computation based on the DFT was performed to
clarify the observed correlation between the modifying cations and the PL properties in the R-SiAlON system. Figure 9 shows the calculated partial density of states (PDOS) for the aforementioned two R-SiAlON models, Y3EuSi36Al12ON63 and Ca3EuSi36A112O4N60. As indicated by Figure 9a,b, the component of the conduction band (CB) was significantly different between the Y- and the Ca-stabilized system. The Si 3p and Al 3p orbitals were dominant for both systems, while the Y 4d orbitals showed great contributions to the CB in the Y system. The energetic position of the lowest 5d state of Eu2+ was also different. The lowest Eu 5d state was located around the bottom of the CB in the Ca system, whereas in the Y system, it was located within the CB. The present electronic structure analysis enables one to discuss the PL properties of Y- and Ca-R-SiAlON:Eu2+ in the following way. In the case of the Y-stabilized system, unoccupied states consisting of Y 4d orbitals exist, which extend to a lower energy level compared to the Eu 5d states (Figure 9a). This means that the excitation energy might transfer nonradiatively from the Eu 5d states to the lower energy Y 4d states in the Y-stabilized system. In contrast, the lowest Eu 5d state in the Ca-stabilized system was located just below the CB (Figure 9b), indicating that the radiative electron transitions between the Eu 5d and 4f orbitals might occur with a higher probability. Therefore, the results of the present calculation revealed that the existence of Y 4d states in the energy region lower than the Eu 5d states might be the main reason for the decrease of the QEint value observed for the Y-stabilized system. The host gap energy was also analyzed to reveal the difference of the Eg between the two systems. The host gap energy was taken as the energy difference between the top of the valence band consisting mainly of N 2p orbitals, and the lowest unoccupied molecular orbitals. The host gap energies were calculated to be 3.66 eV for the Y-stabilized system and 4.75 eV for the Ca system, showing close agreements with the values determined by the DRS (Table 4). These results clearly showed that the contribution of the unoccupied Y 4d states lying beneath the Eu 5d states caused the narrowing of the host band gap and, thus, easier photoionization of the Eu 5d electrons into the CB, which resulted in the significant thermal quenching observed for the Y-stabilized system.
Figure 8. Temperature dependence of the emission intensity for Yand Ca-R-SiAlON:Eu2+ powders, under the excitation at 450 nm.
Figure 9. Calculated PDOS for (a) Y3EuSi36A112ON63 and (b) Ca3EuSi36A112O4N60.
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4. Conclusions The reduction-nitridation synthesis of Y-R-SiAlON fine powders from the multicomponent oxide system Y2O3-Al2O3-SiO2 was performed using the GRN process. Formation of the pure R′ phase could not be attained via the one-step GRN process due to the persistence of the Y-containing intermediate phases, while the postsynthesis heat treatment under the moderate processing conditions enabled the formation of nearly single-phase Y-R-SiAlON powders possessing the favorable characteristics for the phosphor applications: high phase purity of g97%, uniform discrete primary particles of ∼1-2 µm, and high reflectance of ∼90% throughout the visible region. The synthesized Y-R-SiAlON powder activated with Eu2+ exhibited a saturated orangish yellow emission (CIE chromaticity of x ) 0.566, y ) 0.425) under the near-UV to blue light excitation, whereas its internal quantum efficiency and quenching temperature were significantly lower compared to the Castabilized counterpart, suggesting the limited applicability to the white LEDs. QC calculations based on the DFT clearly indicated that the large contribution of the Y 4d states to the conduction band, which lies in more stable energy level compared to the Ca 3d states, may lead to the lower emission efficiency and thermal stability in Y-R-SiAlON:Eu2+. Acknowledgment. T.S. is grateful to Dr. M. Mitomo, formerly of National Institute for Materials Science, for many helpful discussions. H.O. acknowledges Dr. A. Suzuki and Professors H. Tsuboi, N. Hatakeyama, A. Endou, H. Takaba, and M. Kubo of Tohoku University for their continuous support. Supporting Information Available: Structural models for the DFT calculations and X-ray Rietveld refinement pattern of the Eu-activated sample. This material is available free of charge via the Internet at http://pubs.acs.org. References and Notes (1) Hampshire, S.; Park, H. K.; Thompson, D. P.; Jack, K. H. Nature 1978, 274, 880. (2) Izumi, F.; Mitomo, M.; Bando, Y. J. Mater. Sci. 1984, 19, 3115. (3) Mitomo, M.; Tanaka, H.; Muramatsu, K.; Ii, N.; Fujii, Y. J. Mater. Sci. 1980, 15, 2661. (4) Cao, G. Z.; Metselaar, R. Chem. Mater. 1991, 3, 242. (5) Ekstro¨m, T.; Nygren, M. J. Am. Ceram. Soc. 1992, 75, 259. (6) Chen, I.-W.; Rosenflanz, A. Nature 1997, 389, 701. (7) Xie, R.-J.; Mitomo, M.; Uheda, K.; Xu, F.-F.; Akimune, Y. J. Am. Ceram. Soc. 2002, 85, 1229. (8) van Krevel, J. W. H.; van Rutten, J. W. T.; Mandal, H.; Hintzen, H. T.; Metselaar, R. J. Solid State Chem. 2002, 165, 19.
Suehiro et al. (9) Xie, R.-J.; Hirosaki, N.; Mitomo, M.; Yamamoto, Y.; Suehiro, T.; Ohashi, N. J. Am. Ceram. Soc. 2004, 87, 1368. (10) Xie, R.-J.; Hirosaki, N.; Sakuma, K.; Yamamoto, Y.; Mitomo, M. Appl. Phys. Lett. 2004, 84, 5404. (11) Xie, R.-J.; Hirosaki, N.; Mitomo, M.; Yamamoto, Y.; Suehiro, T.; Sakuma, K. J. Phys. Chem. B 2004, 108, 12027. (12) Xie, R.-J.; Hirosaki, N.; Mitomo, M.; Uheda, K.; Suehiro, T.; Xu, X.; Yamamoto, Y.; Sekiguchi, T. J. Phys. Chem. B 2005, 109, 9490. (13) Sakuma, K.; Omichi, K.; Kimura, N.; Ohashi, M.; Tanaka, D.; Hirosaki, N.; Yamamoto, Y.; Xie, R.-J.; Suehiro, T. Opt. Lett. 2004, 29, 2001. (14) Mitomo, M.; Takeuchi, M.; Ohmasa, M. Ceram. Int. 1988, 14, 43. (15) van Rutten, J. W. T.; Terpstra, R. A.; van der Heijde, J. C. T.; Hintzen, H. T.; Metselaar, R. J. Eur. Ceram. Soc. 1995, 15, 599. (16) Ekstro¨m, T.; Shen, Z.-J.; MacKenzie, K. J. D.; Brown, I. W. M.; White, G. V. J. Mater. Chem. 1998, 8, 977. (17) Jiang, J.-X.; Wang, P.-L.; He, W.-B.; Chen, W.-W.; Zhuang, H.R.; Cheng, Y.-B.; Yan, D.-S. J. Am. Ceram. Soc. 2004, 87, 703. (18) Suehiro, T.; Tatami, J.; Meguro, T.; Matsuo, S.; Komeya, K. J. Eur. Ceram. Soc. 2002, 22, 521. (19) Suehiro, T.; Tatami, J.; Meguro, T.; Komeya, K.; Matsuo, S. J. Am. Ceram. Soc. 2002, 85, 715. (20) Suehiro, T.; Hirosaki, N.; Terao, R.; Tatami, J.; Meguro, T.; Komeya, K. J. Am. Ceram. Soc. 2003, 86, 1046. (21) Suehiro, T.; Hirosaki, N.; Komeya, K. Nanotechnology 2003, 14, 487. (22) Suehiro, T.; Hirosaki, N.; Xie, R.-J.; Mitomo, M. Chem. Mater. 2005, 17, 308. (23) Suehiro, T.; Hirosaki, N.; Xie, R.-J.; Sakuma, K.; Mitomo, M.; Ibukiyama, M.; Yamada, S. Appl. Phys. Lett. 2008, 92, 191904. (24) Li, H.-L.; Hirosaki, N.; Xie, R.-J.; Suehiro, T.; Mitomo, M. Sci. Technol. AdV. Mater. 2007, 8, 601. (25) Li, H.-L.; Xie, R.-J.; Hirosaki, N.; Suehiro, T.; Yajima, Y. J. Electrochem. Soc. 2008, 155, J175. (26) Grand, G.; Demit, J.; Ruste, J.; Torre, J. P. J. Mater. Sci. 1979, 14, 1749. (27) Haviar, M. J. Am. Ceram. Soc. 1994, 77, 2425. (28) Izumi, F.; Ikeda, T. Mater. Sci. Forum 2000, 321-3, 198. (29) Ekstro¨m, T.; Ka¨ll, P. O.; Nygren, M.; Olsson, P. O. J. Mater. Sci. 1989, 24, 1853. (30) Vosko, S. H.; Wilk, L.; Nusair, M. Can. J. Phys. 1980, 58, 1200. (31) Perdew, J. P.; Burke, K.; Ernzerhof, M. Phys. ReV. Lett. 1996, 77, 3865. (32) Liddell, K.; Thompson, D. P. J. Mater. Chem. 2001, 11, 507. (33) Ho¨ppe, H. A.; Kotzyba, G.; Po¨ttgen, R.; Schnick, W. J. Mater. Chem. 2001, 11, 3300. (34) Ekstro¨m, T. C.; MacKenzie, K. J. D.; Ryan, M. J.; Brown, I. W. M.; White, G. V. J. Mater. Chem. 1997, 7, 505. (35) Kolitsch, U.; Seifert, H. J.; Ludwig, T.; Aldinger, F. J. Mater. Res. 1999, 14, 447. (36) Sobolev, B. P.; Fedorov, P. P.; Seiranyan, K. B.; Tkachenko, N. L. J. Solid State Chem. 1976, 17, 201. (37) Sakuma, K.; Hirosaki, N.; Xie, R.-J.; Yamamoto, Y.; Suehiro, T. Phys. Status Solidi C 2006, 8, 2701. (38) Sakuma, K.; Hirosaki, N.; Xie, R.-J.; Yamamoto, Y.; Suehiro, T. Mater. Lett. 2007, 61, 547. (39) Mitomo, M.; Izumi, F.; Greil, P.; Petzow, G. Ceram. Bull. 1984, 63, 730. (40) Okamoto, T.; Ukyo, Y. J. Mater. Sci. 1997, 32, 2473.
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