Design of a Yellow-Emitting Phosphor with Enhanced Red Emission via Valence State-control for Warm White LEDs Application

Abstract

The phosphor-converted warm W-LED have being rapidly developed due to the stringent requirements of general illumination. Here, we utilized a strategy to synergistically enhance the red region and emission intensity of novel Eu-activated yellow-emitting LaSiO₂N phosphors. This was realized by predicting optimum crystal structure and governing the concentration of doping ions as well as preparation temperature. By using these straight-forward methods, we were able to vary the valence to enhance the red region and improve the quantum efficiency of LaSiO₂N phosphor. The warm W-LED lamp fabricated with this red region enhanced LaSiO₂N:Eu phosphor exhibited high CRI (Ra = 86), suitable CCT (5783 K) and CIE chromaticity (0.33, 0.36), indicating this synergistically enhanced strategy could be used for design of yellow-emitting phosphor materials to obtain warm W-LEDs.

Introduction

Due to the deficiency of red color, the cool and bluish-white light LEDs (typically a blue-light LED chip coupled with yellow-emitting YAG: Ce³⁺ phosphor) with high correlated color temperature (CCT) and low color rendering index (CRI) are gradually replaced by blue, green and red (RGB)-emitting phosphors with a UV/NUV chip^1,2. However, high cost and poor luminous efficiency of RGB-emitting phosphor became the main obstacles of their popularization because of self-adsorption occurring among these phosphor particles³. Currently, the white LEDs (W-LED) packaged by NUV chips (365–420 nm) with mixed blue and yellow-red emitting phosphors have attracted much attention because of their high CRI, tunable CCT and CIE chromaticity coordinates^{4,5,6,7,8,9,10,11}. Therefore, designing and developing tunable yellow-red emitting phosphors which can be effectively excited with NUV light are in great demand for W-LED industry.

Currently, La–Si–O–N system doped with Ce³⁺ ions has been widely reported as blue phosphors for the application in solid-state lighting, fluorescent lamps or plasma display panels (PDPs)^12,13. The emission properties of this La–Si–O–N system doped with Ce³⁺ or Eu²⁺ strongly depend on Si/La and N/O ratios, because the 5d electrons of Ce³⁺ and Eu²⁺ ions are unprotected and sensitive to the change of the strength of crystal field and covalency^14,15. According to the crystallographic examination for an equal amount of Ce³⁺ substitution, the degree of covalency increased in a sequence of La₅Si₃O₁₂N < La₄Si₂O₇N₂< La₂Si₆O₃N₈ < LaSiO₂N¹⁶. As supported by Dorenbos¹⁷, the emission position depended on nephelauxetic effect, crystal-field splitting (CFS) and Stokes shift. Herein, Eu²⁺ ions in LaSiO₂N should have stronger nephelauxetic effect due to its high covalency, compared with other La–Si–O–N system compounds. This effect would shift the centroid of the 5d band of Eu²⁺ ions to lower energy and result in the redshift of emission peak. Meanwhile, the higher formal charge of N³⁻ compared with O²⁻ makes the CFS become larger and the rigid lattice would lead to a smaller Stokes shift¹⁸. Thus, the LaSiO₂N doped with lanthanide, especially Eu²⁺, may emit long-wave bands. But until now, the Eu²⁺ photoluminescence in La–Si–O–N system has been rarely reported due to the charge mismatch of Eu²⁺ and La³⁺. However, in our recent experiment, we observed the Eu²⁺ photoluminescence in LaSiO₂N and this novel LaSiO₂N: Eu phosphor, as expected, exhibited broad emitting in yellow region. Further, we designed a strategy to cooperatively enhance the red region and emission intensity of this phosphor via altering the concentration of doping ions and preparation temperature for high CRI warm W-LED application. The red region enhancements were caused by varied valence which would affect the energy level of 5d band of Eu²⁺ and charge-transfer state of Eu³⁺, resulting in redshift and the increase of absorption efficiency as well as energy transfer from Eu²⁺ to Eu³⁺. Fabricated by using a blend of LaSiO₂N: Eu and commercial blue phosphors (BAM: Eu²⁺) with a 385 nm NUV LED chip, a warm white light with high CRI as well as suitable CCT and CIE chromaticity can be obtained.

Results and Discussion

The phase purity of the as-prepared LaSiO₂N:Eu phosphors and LaSiO₂N host were substantiated by powder X-ray diffraction. As the Fig. 1 illustrated, all the diffraction peaks matched well with the standard pattern (JCPDS 71–1115) of LaSiO₂N, demonstrating that the doping of Eu ions and the increased preparation temperature did not significantly influence its crystal structure.

Figure 1

XRD patterns of as-synthesized LaSiO₂N:Eu phosphors, LaSiO₂N host and the standard pattern (JCPDS 71–1115) of LaSiO₂N.

Figure 2a presents the crystal structure, photoluminescence excitation (PLE) and photoluminescence (PL) spectra of the LaSiO₂N:0.01Eu phosphor prepared at 1500 °C. The LaSiO₂N crystallizes as a hexagonal structure whose space group is P-6c2 and lattice constants are a = 7.31 Å, c = 9.550 Å and V = 441.95 ų. LaSiO₂N has the α-wollastonite structure with N atoms present in the three-membered (Si₃O₆N₃) rings, occupying bridging sites between pairs of Si-centered tetrahedral and linking to two La atoms¹⁹. Hence, the N environment is more ionic and the lattice is more rigid than that of Si₃N₄ or Si₂N₂O. Based on the crystal structure, the LaSiO₂N:Eu phosphor could be predicted to emit broad and long-wave band and the following studies would verify this inference.

Figure 2

(a) Crystal structure, PLE (λ_em = 554 nm) and PL (λ_ex = 365 nm) spectra of the LaSiO₂N:0.01Eu phosphor (prepared at 1500 °C); (b) Eu concentration-dependent PL (λ_ex = 365 nm) spectra of LaSiO₂N:xEu (x = 0.01, 0.02, 0.04, 0.06, 0.08) phosphors (prepared at 1500 °C) and the corresponding schematic illustration.

As shown in Fig. 2a, The PLE spectrum of LaSiO₂N:0.01Eu monitored by 554 nm was composed of a broad excitation peaked at 365 nm ranging from 250–500 nm, which can be attributed to the 4f ⁷(⁸S7/2)–4f ⁶5d¹ transition of Eu²⁺ ions²⁰ and matched well with the emission of commercial N-UV chip (365–420 nm). The reduction of Eu³⁺ to Eu²⁺ in the trivalent La site can be explained with the charge compensating defect in the first anion (O²⁻ ion) coordination shell¹⁷. The PL spectrum of LaSiO₂N:0.01Eu under 365 nm light excitation exhibited a broad yellow band from 450 to 750 nm peaked at 554 nm with a full width at half-maximum (FWHM) of 115.87 nm which was larger than that of YAG phosphor (91.65 nm), indicating that the synthesized phosphor here is a suitable yellow-emitting phosphor candidate for W-LED application.

Figure 2b depicts the Eu concentration dependent PL spectra of LaSiO₂N:xEu (x = 0.01, 0.02, 0.04, 0.06, 0.08) phosphors with a 365 nm excitation. It is noteworthy that the shoulder peaks at 596 nm, 614 nm and 660 nm arose increasingly along with the increased Eu concentration. These shoulder peaks is reasonable to attribute to the unreduced Eu³⁺ ions and can be assigned to 4f-4f transitions of Eu³⁺ (⁵D₀-⁷F_J (J = 1, 2 and 3))²¹. As the doping concentration of Eu increased to 6 mol%, the emission intensity reached the maximum and then declined dramatically with a further increase of concentration. Generally, the declined intensity with increased concentration is caused by the concentration quenching effect²². Such effect is mainly caused by the energy consumption via energy transfer from one activator to another²³. When the concentration of Eu increased gradually, the interatomic distance between the two Eu ions reduced and the energy transfer rate between Eu²⁺-Eu²⁺ as well as the probability of energy transfer to luminescent killer sites increased²⁴. Simultaneously, the interaction was more intensive with the reduction of interatomic distance. As a result, the 5d band of Eu²⁺ ion decreased and led to the redshift of emission peak²⁵. As depicted in Fig. 2b, an obvious redshift of emission peak occurred, indicating the intensive interaction between the identical activators²⁶. However, as shown in Fig. 3, the decay curves of LaSiO₂N:xEu (x = 0.01–0.08) monitored at 565 nm obviously consist of two lifetimes and all decay curves could be well fitted via the second-order exponential equation²⁷:

Figure 3

The decay curves of LaSiO₂N:xEu (x = 0.01, 0.02, 0.04, 0.06, 0.08).

where I means the luminescence intensity; A₁ and A₂ are constants; t is time; and τ₁ and τ₂ are the lifetimes for the exponential components. As previous researches report that there is only one single La site can be substituted by Eu ion^14,16. Thus, the existence of two lifetimes may due to the two kinds of decay forms for Eu:²⁺ one is the process of electrons from the excited stated to ground state; the other is the energy transfer process between Eu²⁺ and Eu³⁺ because of the coexisting of Eu²⁺ and Eu³⁺. The average lifetime τ* could be reckoned according to the following equation:

based on eqs 1 and 2, the τ* can be estimated to 2.92, 2.27, 1.87, 1.58 and 1.17 μs for LaSiO₂N:xEu with x = 0.01, 0.02, 0.04, 0.06 and 0.08, respectively. The donor decay times decreased as Eu concentrations increased, indicating the existence of energy transfer processes^28,29. Therefore, the phenomenon of declined intensity in LaSiO₂N:Eu might both result from the energy transfer between Eu²⁺-Eu²⁺ and Eu²⁺-Eu³⁺.

The excitation spectra of LaSiO₂N:0.06Eu were shown in Fig. 4a. Although the excitation band of Eu³⁺ were covered by that of Eu²⁺, the charge-transfer state (CTS) could be identified from the excitation spectra of LaSiO₂N:0.06Eu under different emission features at 565 nm, 596 nm, 615 nm and 660 nm. The relative intensity of shoulder peak at around 300 nm increased from 565 nm to 660 nm excitation, suggesting it played a key role in generation of the Eu³⁺ emission. Theoretically, charge-transfer state (CTS) or the energy transfer from host lattice (HL) to Eu³⁺ can generate this shoulder peak²⁵. For detailed investigation of the source of the shoulder peak, the diffuse reflectance spectra (DRS) of LaSiO₂N host and LaSiO₂N:0.06Eu were presented in Fig. 4b. The LaSiO₂N host showed an energy absorption in the short-wavelength UV region (peaked at 225 nm) and a high reflection ranging from 275 to 800 nm. The band gap was estimated to be about 5.08 eV (244 nm) based on the Kubelka-Munk function³⁰. When Eu ions were introduced into the LaSiO₂N host, two broad absorption bands were observed in the 275–350 nm and 350–500 nm, demonstrating that the shoulder peak unlikely originated from HL. The emission spectra of LaSiO₂N:0.06Eu under different excitations were also illustrated in Fig. 4a. Each emission spectrum was consist of luminescence characteristics of both Eu²⁺ and Eu³⁺ ions and the relative intensity of Eu³⁺ reached the maximum at 300 nm, further proving the shoulder peak at 300 nm is assigned from CTS. It suggests that the coexistence of Eu³⁺ and Eu²⁺ can be realized to enhance the red region, while the emission efficiency of Eu³⁺ is insufficient to generate a marked effect.

Figure 4

(a) PLE and PL spectra of the LaSiO₂N:0.06Eu phosphor monitored by different wavelengths of emission and excitation, respectively; (b) Diffuse reflection spectra of LaSiO₂N host and LaSiO₂N:0.06Eu (prepared at 1500 °C and 1550 °C). The inset shows the absorption spectrum of LaSiO₂N host calculated using the Kubelka-Munk equation.

For further improving the red region of LaSiO₂N:Eu phosphor, the preparation temperature was changed and all temperatures were controlled in the range from 1500 °C to 1550 °C to insure the obtained samples are single phase. Figure 5 shows the room temperature PLE (λ_em = 565 nm, 596 nm, 617 nm and 660 nm) of the LaSiO₂N:0.06Eu phosphors prepared at 1550 °C and PL (λ_ex = 365 nm) spectra of the LaSiO₂N:0.06Eu phosphors as a function of the preparation temperature (1500 °C, 1525 °C and 1550 °C). The enlargement of shoulder peak (325 nm) from Eu²⁺ emission feature excitation to that of Eu³⁺ also proved the shoulder peak derived from CTS. It is interesting that not only all the emission intensities were enhanced through the increase of the preparation temperature, but also the relative intensity of Eu³⁺ characteristic peak was enhanced compared with that of Eu²⁺. One of reason for the increase of holistic emission intensity may be due to grain growth and the increased degree of crystallization at a higher preparation temperature^31,32. It can be substantiated by the micro-morphology of the crystalline LaSiO₂N:0.06Eu phosphors which were observed via SEM, TEM and XRD Refinement. As depicted in Fig. 6, the particles prepared at 1500 °C (Fig. 6a,b) had irregular morphology with the diameters ranging from 0.3 to 0.8 μm. When the preparation temperature increased, the edges and corners of irregular particles became clear and a dramatic increase in particle sizes was observed (Fig. 6c,d). In addition, a mass of primary crystals reunited to be particles and the obvious “sintering necks” between primary crystals suggested that grain growth occurred during the process of synthesis. The typical TEM images were illustrated in Figure S1 to further prove the grain growth. After Jade software refined, the relative crystallinity of LaSiO₂N:0.06Eu prepared at 1500 °C and 1550 °C were estimated about 82.38% and 89.91%, respectively, demonstrating the increased degree of crystallization at a higher preparation temperature.

Figure 5

The schematic illustration of the mechanism of red enhanced Yellow-emitting LaSiO₂N:Eu phosphor.

PLE (λ_em = 565 nm, 596 nm, 617 nm, 660 nm ) spectra of the LaSiO₂N:0.06Eu phosphors (prepared at 1550 °C) and PL (λ_ex = 365 nm) spectra of the LaSiO₂N:0.06Eu phosphors as a function of the preparation temperature (1500 °C, 1525 °C and 1550 °C).

Figure 6

SEM images of the LaSiO₂N:0.06Eu phosphors prepared at (a,b) 1500 °C and (c,d) 1550 °C.

To compare the ratio changing of Eu²⁺ and Eu³⁺, the high-resolution XPS spectra at the Eu 3d of LaSiO₂N:0.06Eu phosphors prepared at 1500 °C, 1525 °C and 1550 °C were detected, respectively. As exhibited in Fig. 7, two peaks were found at around 1128 eV and 1135 eV and the shapes as well as binding energies of the Eu3d signals in LaSiO₂N:0.06Eu agreed well with the signals of Eu²⁺ 3d_5/2 and Eu³⁺ 3d_5/2, respectively, indicating the existence of Eu²⁺ and Eu³⁺ ions³³. Additionally, the relative intensity of Eu³⁺ 3d_5/2 signals was gradually decreased with the increased preparation temperature, revealing the promotion of reduction process of Eu³⁺. This may due to the amount of thermal defects increased with the increasing preparation temperatures^34,35, which could charge compensate the Eu²⁺ in the La³⁺ site and improved the reduction from Eu³⁺ to Eu^2+ 17. Thus, the increase of Eu²⁺ concentration may be a reason for the enhanced emission intensity of Eu²⁺. However, decreased ratio of Eu³⁺ concentration compared with Eu²⁺ was inconsistent with the enhanced emission intensity of Eu³⁺. Hence, other assistance might be involved to contribute the characteristic emissions of Eu³⁺ in LaSiO₂N.

Figure 7

High-resolution XPS spectra at Eu 3d position of LaSiO₂N:0.06Eu phosphors prepared at (a) 1500 °C, (b) 1525 °C and (c) 1550 °C.

As illustrated in Fig. 8, the PLE (λ_em = 596 nm) spectra of the LaSiO₂N:0.06Eu phosphors prepared at 1500 °C, 1525 °C and 1550 °C were deconvoluted into three Gaussian components, respectively. The relative intensity of CTS (fit peak 3) was enhanced, indicating that the charge transfer from the O²⁻ to Eu³⁺ was enhanced and more efficient with the increasing preparation temperatures. This phenomenon can be explained based on the increase of oxygen vacancies (V_o) with the increase of nonequivalent substitution of La³⁺ by Eu²⁺ in the host²⁵. These V_o might act as sensitizers for the energy transfer from host to Eu³⁺ ion due to the strong mixing of charge transfer states^36,37,38,39. Thus, the relative emission of Eu³⁺ was enhanced and the red region of this yellow emitting phosphor was promoted. Meanwhile, an obviously redshift of CTS occurred as the preparation temperature increased, which can be attributed to the increase of Eu³⁺-O²⁻ bond length. Since the Eu²⁺ ions have smaller electronegativity than that of Eu³⁺ ions, the Eu–O bond strength became weaker with increasing concentration of Eu²⁺, resulting in weakening of bond strength. Hence the Eu³⁺-O^2- bond length became longer. As reported by Lin et al.⁴⁰, the longer the Eu-O bond, the shorter the energy difference between the 4f and O 2p electrons and the lower energy position of the CTB.

Figure 8

Normalized PLE spectra (λ_em = 596 nm) for LaSiO₂N:0.06Eu samples (1500 °C, 1525°C and 1550 °C) and normalized PL spectra (λ_ex = 365 nm) for LaSiO₂N:0.06Eu sample.

The inset shows the magnified Fit Peaks 1 of LaSiO₂N:0.06Eu samples.

Additionally, the coordination number of Eu²⁺ was reduced because the V_o increased with an increase of the Eu²⁺ content in the host, which caused the formation of a centroid of the 5d state at a lower level. Thus, as shown in Fig. 8, the 5d excitation band (fit peak 2 and 3) red shifted and the overlapping between PLE of 596 nm and PL enlarged, indicating that the energy transfer ratio between Eu²⁺ and the ⁵D₀ level of Eu³⁺ was enhanced. The energy transfer mode between Eu²⁺ and Eu³⁺ in LaSiO₂N: Eu was proposed in Fig. 9. Under the NUV light excitation, Eu²⁺ was excited from the ground state 4f ⁷(⁸S7/2) to the excited state 4f ⁶5d¹. Partial energy relaxed to the ground state through the inherent transition of Eu²⁺, generating a yellow light emission; the rest of energy transferred to the nearest level ⁵D₀ of Eu³⁺ and then 597 nm, 617 nm, 660 nm and 707 nm emissions appeared by a transition to the ⁷F_j (J = 1, 2, 3 and 4) ground state. With the increase of the preparation temperature, the depressed 5d level showed more overlapping with the ⁵D₀ of Eu³⁺, resulting in the enhancement of the energy transfer between the Eu²⁺ and Eu^3+ 25. The increased emission peak position from 550 nm at 1500 °C to 565 nm at 1550 °C and FWHM of emission peak from 116 nm at 1500 °C to 120 nm at 1550 °C also verified the redshift of 5d level^25,26. The decay curves of LaSiO₂N:0.06Eu prepared at different temperatures monitored at 565 nm were depicted in Fig. 10. All decay curves also could be well fitted via the second-order exponential equation and the lifetimes were estimated 1.58, 1.23 and 0.96 μs for LaSiO₂N:0.06Eu prepared at 1500 °C, 1525 °C and 1550 °C, respectively, demonstrating the existence of energy transfer. Consequently, via the increase of preparation temperature, the holistic emission intensity was enhanced owe to the increased crystallinity and reduction process. The relative emission intensity of Eu³⁺ was also increased due to the enhancement of the energy transfer. As a result, the red region of yellow emitting LaSiO₂N: Eu phosphor was successfully enhanced.

Figure 9

Schematic of energy transfer in LaSiO₂N:0.06Eu.

Figure 10

The decay curves of LaSiO₂N:0.06Eu prepared at 1500 °C, 1525 °C and 1550 °C, respectively.

The quantum efficiency of phosphors, which is an important technological parameter for practical application, were also been compared. The internal quantum efficiencies (IQE) of LaSiO₂N:0.06Eu prepared at different temperature were measured and calculated by the following equations⁴¹:

where L_S represents the luminescence emission spectrum of the sample; E_R is the spectrum of the excitation light from the empty integrated sphere (without the sample); E_S means the excitation spectrum for exciting the sample. As given in Figure S2, the IQE of the LaSiO₂N:0.06Eu prepared at 1500 °C, 1525 °C and 1550 °C were estimated to be about 3.48%, 12.48% and 20.01%, respectively, under 365 nm excitation. The enhanced IQE matched well with the variation trend of emission intensity. Although the IQE of LaSiO₂N:0.06Eu is lower than that of commercial YAG (61%)⁴², it can be further improved by optimization of the preparation conditions, because the QE depends closely on the prepared conditions, crystalline defects, particle size and morphology of the phosphor^43,44.

In order to well understand the luminescent performance of this phosphor, the temperature-dependent luminescent properties of LaSiO₂N: 0.06Eu phosphor was measured during the temperature ranges of 25–200 °C. As presented in Fig. 11a, the emission intensity decreased with increasing temperature. At 100 °C and 150 °C, the PL intensity quenched to 62.8% and 42.3%. The thermal quenching effect of this phosphor is more intense than the commercial YAG:Ce^3+ 45, while similar to La–Si–O–N system phosphors¹². The thermal quenching can be explained by the model of thermal excitation of the 5d electron to conduction band states^46,47. When Eu²⁺ substituted in a trivalent site, the Eu²⁺ had the trend to be ionized to Eu³⁺, which reduced the activation energy and lead to stronger thermal quenching. To determine the activation energy for thermal quenching, the Arrhenius equation was used to estimate the thermal quenching⁴⁸:

Figure 11

(a) Temperature-dependent PL spectra (λ_em = 365 nm) and (b) the plot of ln (I0/I) vs. 1/T of LaSiO₂N:0.06Eu phosphors. The inset shows the emission intensity versus temperature.

where I₀ and I mean the luminescence intensity of LaSiO₂N:Eu at room temperature and a given temperature, respectively; A is a constant; k presents the Boltzmann constant (8.617 × 10⁻⁵ eV K⁻¹). Figure 11b plots ln[(I0/I)-1] variation dependence of 1/(kT) and the ΔE was calculated to be about 0.24 eV.

To further assess the potential application of the LaSiO₂N: Eu phosphors, the yellow emitting LaSiO₂N:0.01Eu phosphor prepared at 1500 °C and LaSiO₂N:0.06Eu phosphors prepared at 1550 °C were mixed with blue-emitting BAM: Eu²⁺ phosphor, respectively and then the mixtures were severally combined with a 385 nm NUV chip to fabricate white LED lamps. The electroluminescent (EL) spectra of these lamps driven by 30 mA current were depicted in Fig. 12. The CIE color coordinates, CCT and CRI of the fabricated W-LED lamp with LaSiO₂N:0.01Eu phosphor prepared at 1500 °C were determined to be (0.32, 0.38), 5959 and 76, respectively (Figure S3 and Fig. 12a). Via utilizing the redshift, varied valence and efficient energy transfer, the fabricated W-LED lamp with the red region enhanced yellow-emitting LaSiO₂N:0.06Eu phosphor displayed an entire white spectrum with a CIE color coordinates of (0.33, 0.36), a CCT of 5783 K and a CRI of 87 (Figure S3 and Fig. 12b). Compared with the W-LED lamp using commercial YAG:Ce phosphor in previous study (CIE = 0.302, 0.315; CCT = 7272 K, Ra = 78.38)³, the CCT of as-fabricated LEDs was relative low and the Ra was high, suggesting that the high CRI warm W-LEDs could be easily obtained by altering the concentration of doping ions and the preparation temperature.

Figure 12

Electroluminescence spectra of W-LED lamps fabricated using a 385 nm NUV chip in combination with the mixed phosphors consisting of blue-emitting BAM:Eu²⁺ phosphor and yellow-emitting (a) LaSiO₂N:0.01Eu phosphor prepared at 1500 °C, (b) LaSiO₂N:0.06Eu phosphor prepared at 1550 °C. Insets show the digital images of the LED package with an input of 30 mA current.

Conclusions

In summary, a novel Eu-activated LaSiO₂N yellow-emitting phosphor has been synthesized and evaluated for the application in W-LEDs. With the aid of crystal structure, valence-varied, redshift and energy transfer, the red region enhanced yellow-emitting LaSiO₂N: Eu phosphor has been designed and realized by controlling the concentration of doping ions and the preparation temperature. A high CRI and warm W-LED lamp was obtained in combination with this phosphor, proving that the red region enhanced LaSiO₂N: Eu phosphor has a great potential for the application in W-LEDs. More importantly, this study would provide a new strategy for designing Eu-activated yellow-emitting phosphors by synergistically enhancing the red region and emission intensity to adjust the CCT and CRI for warm W-LEDs application without reducing the other luminescence properties.

Methods

Materials and Synthesis

The LaSiO₂N:Eu was synthesized from stoichiometric mixtures of La₂O₃ (analytical reagent (A. R.)), α-Si3N4 (A. R.) and Eu₂O₃ (A. R.). The ground powders were placed in alumina crucibles and fired for 6 h in a reducing atmosphere (10% H₂ + 90% N₂) at 1500 °C, 1525 °C and 1550 °C, respectively. Then, the precursor was reground and heated again at same condition. Thereafter, the samples were cooled down to room temperature naturally and powdered for subsequent analysis.

Materials Characterization

Powder X-ray diffraction on a D8 Advance diffractometer (Germany) with graphite-monochromatized Cu Kα radiation (λ = 0.154 06 nm) was recorded for the structure of all samples. Photoluminescence spectra were collected using a Hitachi F-4600 fluorescence spectrophotometer (Japan) equipped with a 150 W Xe lamp as the excitation source. Diffuse reflection spectra were recorded using a Shimadzu UV-3600 UV−vis−NIR spectrophotometer attached with an integrating sphere. BaSO4 was used as a reference for 100% reflectance. The morphology was observed using scanning electron microscopy (SEM; JSM-6460LV, JEOL, Japan). X-ray photoelectron spectroscopy (XPS) measurements were performed in a PHI 5300 ESCA system using an Al Ka X-ray source with constant pass energy of 55.00 eV. The charge effect referred to the C1s signal (284.6 eV). The room-temperature luminescence decay curves were obtained from a spectrofluorometer (Horiba, Jobin Yvon TBXPS) using a tunable pulse laser radiation (nano-LED) as the excitation. Quantum efficiency was measured by a fluoromax-4 spectrofluorometer (Horiba, Jobin Yvon) with an integral sphere at room temperature.