Synthesis and Characterization of Nb₂O₅:La³⁺,Eu³⁺ Phosphors Obtained by the Non-Hydrolytic Sol-Gel Process

Abstract

The phosphor-converted light-emitting diode technique is an important solid-state illumination strategy. Sulfide-based materials are the most often employed red phosphors, but they are chemically unstable and present lower efficiency in comparison to the blue and green phosphors. Therefore, it is important to find a new red phosphor source that can emit intense red light while absorbing light in the near ultraviolet (UV) spectral region. This paper describes the photoluminescence properties of Nb₂O₅:La³⁺,Eu³⁺ obtained by the non-hydrolytic sol-gel process. The X-ray results indicated that the thermal treatment allowed to obtain the different crystalline structures such as, the orthorhombic and monoclinic phases for the Nb₂O₅ and the orthorhombic phase for the La₂Nb₁₀O₂₈. This polymorphism was also confirmed by the Raman spectroscopy. The luminescence spectra revealed the existence of the Eu³⁺ ions in both crystalline phases for the samples annealed at higher temperature, depending of the excitation wavelength. The emission spectrum showed that increasing the annealing temperature promotes the narrowing of all intraconfigurational f-f transitions for the Eu³⁺ ions, due to the structural changes. In addition, all samples present good CIE (International Illumination Committee) chromaticity coordinates when excited in the UV (275 and 394 nm), blue (465 nm) and green (525 nm) radiation.

Keywords:
non-hydrolytic sol-gel route; niobia; red phosphor; luminescence

Introduction

Currently, solid-state white light emitters (LEDs) are considered the next generation of solid-state light sources. Because LEDs offer advantages such as longer lifetime and low energy consumption, they constitute potential replacements for incandescent and fluorescent light bulbs.^1,2 Bearing in mind that illumination consumes about 33% of all the generated energy, the development of energy-saving systems has become fundamental from a technological standpoint.³

Ultraviolet (UV) radiation provides high excitation energy, so a new strategy to obtain white light relies on the use of near UV radiation (ca. 400 nm) by LEDs coated with blue/green/red luminophores.⁴ An ultraviolet light emitting diode composed of indium gallium nitride (InGaN) or gallium nitride (GaN) present a high light emission efficiency in the light emission wavelengths from 370 nm to 410 nm, and particularly, has the highest light emission efficiency at wavelengths around 390 nm.⁵ Yttrium oxysulfide doped with Eu³⁺ (Y₂O₂S:Eu³⁺) is the red-light luminophore that is commercially available at present; however, it is less efficient than luminophores that emit blue and green light, has shorter lifetime under UV radiation, and is unstable because it releases hydrogen sulfide.^6-9 In this context, achieving novel and stable luminophores that can emit intense red light, exhibits excitation in the near UV region, and presents CIE (International Illumination Committee) chromaticity coordinates that meet the demands of the National Television Standard Committee (NTSC) is highly desirable.

To produce white light, the current phosphor materials based on near UV GaN-LEDs are BaMgAl₁₀O₁₇:Eu²⁺ (BAM) for blue, ZnS:(Cu⁺, Al³⁺) for green, and Y₂O₂S:Eu³⁺ for red. However, compared with the blue and green phosphors, the efficiency of Y₂O₂S:Eu³⁺ is about eight times lower and a mixture containing 10% blue, 10% green and 80% red is necessary to obtain appreciable color rendering.² Therefore, it is mandatory that scientists design new processes and search for novel materials to synthesize stable luminophores that can efficiently absorb near UV radiation while emitting red light. In this sense, niobium oxide doped with La³⁺ and Eu³⁺ has emerged as a promising material for multifunctional applications.

Nb₂O₅ is transparent over a wide range of wavelengths (0.35-9.0 µm), it has a wide band-gap (3.6 eV), it is stable under near UV radiation, it has a relatively low cut-off phonon energy (900 cm^-1), high refractive index (2.4), high permittivity (29 to 200 depending of the crystalline phase) and undergoes polymorphic transformations induced by treatment temperature.^10-13 All these properties enhance the optical features of rare earth ions, since efficient radiative emission output requires a relatively high refractive index and a low nonradioactive decay for excited states. The amorphous niobium oxide (V) begins to crystallize in a low-temperature form called the (T) form or γ-phase (orthorhombic, space group: Pbam) at about 500 ºC. Crystallization occurs more rapidly at higher temperatures until about 800 ºC, where upon a transition to a medium-temperature (M) form or β-phase (monoclinic, space group: I4/mmm). This transition continues more rapidly at higher temperatures and heating for 4 h at 1000 ºC brings about a complete conversion. At even higher temperatures, a third transformation to a high-temperature (H) form or β;-phase (monoclinic, space group: P2/m) is observed. These polymorphic transitions take place slowly, at temperatures which are not well-defined, and are irreversible.^14-16

Luminescent materials can absorb light from diverse sources and emit radiative electromagnetic energy in the visible, ultraviolet, and infrared regions.¹⁷ Among the various methodologies that are available to synthesize luminophores, the non-hydrolytic sol-gel process stands out as one of the most advantageous: it yields highly pure products (the metallic oxides originate in situ) with fewer pores; occurs at relatively low temperatures; allows for incorporation of organic species, to give more homogeneous hybrid materials; and is easier to reproduce.¹⁸

In the traditional sol-gel process, the water molecules within the Eu³⁺ ion coordination sphere contribute to a non-radiative decay, which has increased the interest in alternative low-temperature processes that take place in the absence of water.^19,20 The non-hydrolytic sol-gel route is a versatile way to prepare inorganic oxides¹⁸ during which non-hydrolytic condensation reactions furnish oxides and hybrid organic-inorganic materials; the oxo bonds originate from oxygen atoms of donors other than water.²⁰ This route provides strict control of stoichiometry, powder morphology, and phase purity; cations are distributed all over the polymeric structure. Thermal treatment at high temperatures (from 400 ºC) releases organic matter and produces well-ordered crystallites, which is important when one wishes to obtain high crystallinity and controlled distribution of the constituents within the crystalline network. The non-hydrolytic sol-gel process dismisses the need for solvents and may reduce or eliminate the formation of residual metal-OH groups.²¹

It is known that particle morphology and size affect the luminescent properties of Nb₂O₅ matrices; that is, the features of the Nb₂O₅ material depend on the methodology used to prepare the solid. Authors have employed countless methodologies to obtain Nb₂O₅ matrixes, like sol-gel,^10,11 Pechini,^22,23 chemical transport,¹⁴ hydrothermal route^24,25 and polymerized complexes.²⁶ Nevertheless, to date there are no reports on the use of the non-hydrolytic sol-gel technique to prepare these materials.

Thus, the present study employed the non-hydrolytic sol-gel route to prepare Nb₂O₅:La³⁺,Eu³⁺ matrices at Nb₂O₅:La³⁺_(1.0%x%)Eu³⁺_(x%) stoichiometric ratios, where x = 0.1 and 0.3 mol%, to obtain red phosphors with a wide excitation range and adequate CIE chromaticity coordinates. The samples underwent thermal treatment at 550, 750, 900, and 1200 ºC, for 4 h, followed by investigation of their luminescent properties.

Experimental

All the solvents and reagents were of commercial grades (Merck and Sigma-Aldrich) unless otherwise stated. NbCl₅ was donated by Companhia Brasileira de Metalurgia e Mineração (CBMM).

Preparation of Nb₂O₅:La³⁺_(1%-x%)Eu³⁺_(x%)

The preparation of the gel was carried out in oven-dried glassware. The material was synthesized via modification of the method described by Acosta et al.²⁷ Niobium pentachloride (NbCl₅, 5.55 mmol) was previously suspended in 5 mL of dry dichloromethane and 50 mL of ethanol; it was then reacted in the presence of lanthanum chloride (LaCl₃, 5.55 × 10^-2 mmol) under reflux at 110 ºC, under argon atmosphere. For the samples doped with Eu³⁺ ions, an amount of europium chloride (EuCl₃, 5.55 × 10 ^-3 or 1.66 × 10 ^-2 mmol) was added in the precursor solution. The condenser was placed in a thermostatic bath at −5 ºC. The gel was formed after 4 h of reaction; after 0.5 h, a solid material started to precipitate. After reflux, the mixture was cooled and aged overnight in the mother liquor, at room temperature. The solvent was removed under vacuum, and the solid was dried at 100 ºC. The powders were annealed at 550, 750, 900, and 1200 ºC for 4 h.

Characterization

X-ray diffraction (XRD) was performed at room temperature using a Rigaku Geigerflex D/max-c diffractometer with monochromated CuKα radiation (λ = 1.54 Å). Diffractograms were recorded in the 2θ range from 10 to 80º at a resolution of 0.05º. Raman spectra of the materials were obtained in the Raman spectrometer Horiba Jobin Yvon model LabRAM HR 800, operating with a He-Ne laser at 632.81 nm equipped with a charged couple device (CCD) camera model DU420A-OE-325 from the same manufacturer. The spectrum was obtained in the region from 100 to 4000 cm^-1 with a hole of 100 µm and length of 5 to 30 cycles. Photoluminescence data were obtained at room temperature, under continuous Xe lamp (450 W) excitation in a Horiba Jobin Yvon Fluorolog-3 spectrofluorimeter equipped with an excitation and emission double monochromator and a photomultiplier R 928 Hammatsu. The emission was collected at 90º from the excitation beam. The slits were placed at 2.0 and 0.5 nm for excitation and emission, respectively; the bandpass was 0.5 nm, and the integration time was 0.5 ms. G1227 emission filters were employed (transmittance 100% for λ > 450 nm). The decay curves were acquired using a phosphorimeter accessory equipped with a Xe Lamp (5 J per pulse). Particle morphology was investigated by scanning electron microscopy (SEM) on a Shimadzu microscope model VEGA 3 SB.

Results and Discussion

Figure 1 shows the XRD patterns of the samples sintered at different temperatures. Samples without thermal treatment do not display any peaks. Addition of increasing, but small amounts of the dopant Eu³⁺ does not shift the diffraction peaks, because the unit cell varies within the detection limit of the instrument. Samples sintered at 550 ºC present peaks indexed to the (T) phase of Nb₂O₅ with an orthorhombic structure (JCPDS 30-873) and the Nb₃O₇Cl orthorhombic structure (JCPDS 73-295). Samples treated at 750 ºC contain Nb₂O₅ polycrystalline phases indexed to the Nb₂O₅ orthorhombic phase (JCPDS 30-873) and tetragonal structure (JCPDS 18-911). The crystalline (M) phase of the Nb₂O₅ monoclinic structure (JCPDS 15-166) and the La₂Nb₁₀O₂₈ orthorhombic phase (JCPDS 20-547) appear after sintering at 900 ºC. After the thermal treatment at 1200 ºC, a totally crystalline (H) phase indexed to the Nb₂O₅ monoclinic structure (JCPDS 37-1468) and the La₂Nb₁₀O₂₈ orthorhombic phase (JCPDS 20-547) were identified.

Figure 1
X-ray diffractograms of samples treated at different temperatures: (a) Nb₂O₅:La_(1.0%); (b) Nb₂O₅:La_(0.9%)Eu_(0.1%); and (c) Nb₂O₅:La_(0.7%)Eu_(0.3%).

Figure 2 illustrates the Raman spectra of the samples after calcination at different temperatures. Untreated samples display broad bands in the region of the 60-400 and 760-1100 cm^-1 related to the precursor (Nb-Cl) and to the initial linkages between niobium and oxygen.^11,28 Confirming the data observed in X-ray diffraction, the results obtained by Raman spectroscopy confirmed the polymorphism of niobia, which has many Nb-O polyhedra, being NbO₆, octahedrons, the most common presenting in the T, M and H phases in different degree of distortion. In the T phase, broad bands are always present due to the large number of hexa or hepta coordination of Nb per unit cell, which increased the distortion and the number of different sites occupied by niobium atoms. The M and H phases are very similar and less distorted.^10,11

Figure 2
Raman spectra of samples treated at different temperatures: (a) Nb₂O₅:La_(1.0%); (b) Nb₂O₅:La_(0.9%)Eu_(0.1%); and (c) Nb₂O₅:La_(0.7%)Eu_(0.3%).

The samples calcined at 550 and 750 ºC presented bands at around 130, 230 and 700 cm^-1 related to the optical vibration of Nb₂O₅, to the bending modes of the Nb−O−Nb linkage (ν₄) and to the vibrational mode related to the Nb−O symmetric stretching (ν₁). These vibrations are characteristic of the (T)-phase of Nb₂O₅ as observed in the X-ray results and they are in agreement to the literature. A broad band between 560 and 790 cm^-1 is related to the vibrations of niobium polyhedrons in the (T)-phase of the Nb₂O₅. In addition, there is a weak shoulder around 890 cm^-1, which corresponds to the symmetric stretching mode of surface Nb=O. For the samples annealed at higher temperature, an intense peak at 992 cm^-1 was assigned to the Nb−O symmetric stretching (ν₁) of the NbO₆ polyhedra, characteristic of the highly-ordered (M)- or (H)-phase.^11,29 There are low intensity bands at 895 and 847 cm^-1, which correspond to the symmetric stretching mode of surface Nb=O and to the vibrational mode (ν₁) of the NbO₄ tetrahedron, once that one of the 28 Nb atoms in each unit cell is present in a tetrahedral site, which occurs at some block junctions. Several bands were also observed between 620 and 680 cm^-1, which are related to the transverse optic modes in different types of octahedral arrangements.^11,15,29

Low intensity bands at 350, 470 and 548 cm^-1 assigned to the mode (ν₅) are related to the block structures and the band at around 347 cm^-1 is ascribed to the corner-shared octahedra. The octahedral distortion provides the appearance of broad multiplet bands at 260-290 cm^-1 assigned to the infrared-active modes ν₃ and ν₄; the bands at 204 and 265 cm^-1 are characteristic of ν₆ modes and exist as a broad multiplet, also because of octahedral distortion.^11,29 As observed for the (T)-phase, the samples calcined at 900 and 1200 ºC (M- and H-phase) also presented a band at around 120 cm^-1 but in this case, it probably arises from a more ordered structure. Finally, bands at 102 and 200 cm^-1 were assigned to the La-O vibrational mode,^30,31 due to the lanthanum niobate phase also detected by the X-ray diffraction.

SEM images (Figure 3) reveal some interesting morphologies. Non-calcined samples and samples sintered at 550 ºC (Figures 4a_1,2, 4b_1,2 and 4c_1,2) present porous surface; surface particles are irregular. Samples treated at 750 ºC (Figures 3a₃, 3b₃ and 3c₃) correspond to regular pellets with typical crystalline grain measuring around 0.2-0.7 µm. Higher sintering temperature increases the crystalline grain size; indeed, the samples calcined at 900 and 1200 ºC (Figures 3a_4,5, 3b_4,5 and 3c_4,5) have crystals with particle size around 0.7-5.0 µm; the pellets are arranged as sheets with a thickness estimated of around 60 nm (Figure 4). This is in agreement with the values obtained by the Scherrer analysis,³² which provided crystallite size of around 50 nm to the diffraction plane h k l (3 0 1) at 2θ = 17.25º for the samples annealed at 1200 ºC.

Figure 3
Scanning electron micrographs of samples (a) Nb₂O₅:La_(1.0%); (b) Nb₂O₅:La_(0.9%)Eu_(0.1%); and (c) Nb₂O₅:La_(0.7%)Eu_(0.3%). The indexes 1-5 indicate the thermal treatment temperature of 100, 500, 750, 900, and 1200 ºC, respectively.

Figure 4
Scanning electron micrographs of samples: (a) Nb₂O₅:La_(1.0%); (b) Nb₂O₅:La_(0.9%)Eu_(0.1%); and (c) Nb₂O₅:La_(0.7%)Eu_(0.3%) treated at 1200 ºC.

Figures 5 and 6 depict the luminescence spectra of the samples Nb₂O₅:La_(0.9%)Eu_(0.1%) and Nb₂O₅:La_(0.7%)Eu_(0.3%), respectively, calcined at different temperatures. All the samples display a broad band at 275 nm in the excitation spectra (λ_em: 614 nm, Figures 5a and 6a), irrespective of the thermal treatment temperature. This band was more intense for the Nb₂O₅:La_(0.7%)Eu_(0.3%) calcined at 750 and 900 ºC (Figure 5a). Many literature papers describing niobium oxide doped with lanthanides have reported that the charge transfer (CT) band emerges in the excitation spectrum. Zhang et al.³³ detected the CT band at 280 nm for the Li₅(La_2.7Eu_0.3)Nb₂O₁₂ oxide obtained by solid-state reaction; the most intense band of the ƒ → ƒ transition appeared at 463 nm and corresponded to the ⁷F₀ → ⁵D₂ transition. The (La_0.99Eui_0.01)NbO₄:Eu³⁺ matrix prepared by spray pyrolysis method display a large CT band between 240 and 310 nm, ascribed to CT band.³⁴ The YNbO₄-EuNbO₄ system produced via hydrothermal route also presented the CT band between 240 and 270 nm.³⁵ On the other hand, the nanocrystalline powder Ln³⁺ doped Nb₂O₅ obtained by the Pechini methodology did not have a CT band.³⁶ In our work, we hypothesize that this band is related to the La₂Nb₁₀O₂₈ micro-crystalline phase; which was assign to the charge CT band resulting from the 2p orbital of oxygen to empty 4f orbitals of Eu³⁺ (Eu³⁺ → O^2-), and to charge transfer from the ligand oxygen to the empty 4f orbital of the central niobium atom of the group NbO₄^3- (Nb⁵⁺ → O^2-).^5,34,37 This fact is in agreement with the results obtained by the X-ray diffraction, that reveled the presence of the La₂Nb₁₀O₂₈ as a secondary phase in the niobia host. It is important to note that in several crystals, such as lithium niobate, strontium barium niobate, and barium sodium niobate, the location of the Ln³⁺ in the lattice is not obvious, as the trivalent lanthanide ions cannot easily replace the constitutional cations (Li⁺, Na⁺, Sr²⁺, Ba²⁺, Nb⁵⁺) due to clear size and/or charge mismatches. This location has been addressed by several studies where optical and/or structural techniques have been employed.²² Nevertheless, it is still a matter of debate whether the Ln³⁺ can substitute for the smaller and higher charged Nb⁵⁺ cation in a crystalline lattice. Falcomer et al.²² reported that pentavalent niobium ions can be substituted by trivalent lanthanide ions in crystalline niobates, but this substitution is accompanied by a strong disorder around the Ln³⁺ ions.

Figure 5
Luminescence spectra of the sample Nb₂O₅:La_(0.9%)Eu_(0.1%) calcined at different temperatures: (a) excitation spectra (λ_em: 614 nm); (b) emission spectra (λ_exc: 275 nm); (c) emission spectra (λ_exc: 394 nm); (d) emission spectra (λ_exc: 465 nm) and (e) emission spectra (λ_exc: 525 nm).

Figure 6
Luminescence spectra of the sample Nb₂O₅:La_(0.7%)Eu_(0.3%) calcined at different temperatures: (a) excitation spectra (λ_em: 614 nm), (b) emission spectra (λ_exc: 275 nm), (c) emission spectra (λ_exc: 394 nm), (d) emission spectra (λ_exc: 465 nm) and (e) emission spectra (λ_exc: 525 nm).

In the excitation spectrum, a transition is observed only if this level is efficient in populating the emitting level and thus in generating luminescence. If an energy level is absent it means that this level is not efficient in absorbing the excitation light and/or is not able in populating the emitting level.³⁸ Transitions within the ⁷F ground term are only observed for Eu³⁺ ions doped in inorganic matrices with low phonon energies, since these transitions are otherwise masked by the much stronger overtones and combination bands of the vibrations of the host matrix or ligands.³⁸

The sharp lines in the 380-600 nm range correspond to Eu³⁺ intraconfigurational 4f-4f transitions in the host lattices; the main excitation bands at 394, 465, 525 and 532 nm refer to the Eu³⁺transitions: ⁷F₀ → ⁵L₆, ⁵D₂, ⁵D₁ and ⁷F₁ → ⁵D₁, respectively. As the ⁷F₀ → ⁵D₁ is a magnetic dipole transition, its intensity does not change. However, the increase of the annealing temperature promotes an increase of crystallinity (from orthorhombic to monoclinic) changing the local site symmetries for the Eu³⁺ ions and that can provide some differences in the intensities of the others electric dipole transitions. Higher annealing temperatures promote a decrease of the relative intensity of the ⁷F₀ → ⁵L₆ transition and favors the ⁷F₁ → ⁵D₁ transitions.

The emission spectrum for both samples show different profiles depending of excitation wavelengths and annealing temperatures. The samples excited at CT band (275 nm) exhibit broad bands for all intraconfigurational f-f transitions of Eu³⁺ ions, as presented in Figures 5b and 6b. Similar profiles were observed for the samples annealed at 750 ºC and 900 ºC, which were different in comparison with the samples annealed at 550 ºC and 1200 ºC. These three different behaviors are explained by the different crystalline phase. All samples calcined at 550 ºC present inhomogeneous broad bands for the ⁵D₀ → ⁷F_J (J = 0, 1, 2, 3 and 4) due to the presence of many sites of symmetry, which are different to each other. As mentioned above, it is characteristic of (T) phase of niobia, that presents in its unit cell, eight of the Nb ions are present in distorted octahedra, while another eight Nb ions occupy pentagonal bipyramids. The remaining 0.8 Nb ion per unit cell is located in interstitial 9-coordinated sites in the unit cell.¹⁵ The samples annealed at 750, 900 and 1200 ºC presented a narrowing in the emission bands in comparison with the samples calcined at 550 ºC. For the samples calcined at higher temperatures, the (M) and (H) phase of the Nb₂O₅, have a shear structure consisting of blocks of NbO₆ octahedra (3 × 4 and 3 × 5) that share corners with octahedra in their own block and edges with octahedra in other blocks. One of the 28 Nb atoms in each unit cell is present in a tetrahedral site, which occurs at some block junctions.¹⁵ In addition, as the ⁵D₀ → ⁷F₄ transition is also sensitive to the local symmetry, the structure changes can be observed by the variation of its intensity. The emission spectra (Figures 5a and 6a) excited at 275 nm for these samples present similar profiles to those already published in the literature,^1,33,34,36 which describe the preparation of niobates matrices doped with the Eu³⁺ ions, suggesting that under CT band excitation, the majority emitting ions are located in the lanthanum niobate phase.

The presence of the ⁵D₀ → ⁷F₀ transition is an indication that the Eu³⁺ ion occupies a site with C_nv, C_n or C_s symmetry.^38,39 In addition, this transition is also useful to confirming the presence of non-equivalent sites in a host structure, because due to the non-degeneracy of the ⁷F₀ and ⁵D₀ levels, it is expected only one peak for a single site or species. The observation of more than one peak indicates that more than one site or species is present. Slight differences between two sites can provide small energy differences between the two different peaks in the ⁵D₀ → ⁷F₀ region, and the presence of more than one site for the Eu³⁺ ions is verified by the asymmetric shape of this transition or as a shoulder.³⁸ Although this transition is strictly forbidden according to the standard Judd-Ofelt theory, it occurs due to J-mixing. J is the total angular quantum number and it indicates the relative orientation of the spin and the orbital momenta. The J quantum numbers are well defined in the free Eu³⁺ ion, but J-mixing is due to the crystal-field perturbation which causes a mixing of the wave functions of terms with different J values.³⁸

The emission spectra excited directly in the Eu³⁺ ions transitions (Figures 5c-e and 6c-e) prompts the appearance of bands characteristic of the transition from the excited state ⁵D₀ to the fundamental level ⁷F_J (J = 0, 1, 2, 3, and 4). The emission profiles for the samples excited at 394, 465 and 525 nm presented great differences in comparison to those excited at 275 nm, suggesting the existence of different local sites symmetry for the Eu³⁺ ions. As the lanthanum ions are present in low concentration in comparison with the niobium ions, the majority of Eu³⁺ ions are most probably located in the niobia phase. For these samples, it was observed a narrowing of the bands in function of the thermal treatment temperature. The samples calcined at 900 and 1200 ºC presented two bands in the ⁵D₀→ ⁷F₀ transition region at 577.5 and 582.6 nm, evidencing that at least two active sites exist. This fact corroborates with the X-ray diffraction results, which show distinct phases for the calcined samples. An increase of the ⁵D₀ → ⁷F₀ transition’s intensity was observed for the samples calcined at higher temperatures due to an increase system’s cristallinity. Similar results were published for Sr₂TiO₄:Eu³⁺, Ln₂O₂SO₄ (Ln = La, Gd, Y), Gd₂O₃ and Y₂O₃ materials.^38,40-42 In these cases, the high intensity of the ⁵D₀ → ⁷F₀ transition was ascribed to ordered crystal structure, which leads to large linear terms in the crystal-field potential.³⁸

The ⁵D₀ → ⁷F₁ transition directly reflects the crystal-field splitting of the ⁷F₁ level. In a high symmetry such as, cubic or icosahedral crystal-fields, the ⁷F₁ level is not split. In hexagonal or tetragonal crystal-fields, it is split into a non-degenerate and a twofold degenerate crystal-field level and in orthorhombic or lower symmetries, the total removal of crystal field degeneracies results in three sublevels for ⁷F₁.³⁸ The samples presented an increase of splitting for the ⁵D₀ → ⁷F₁, indicating an increase of the local symmetry for the Eu³⁺ ions.

The peak at 582.5 nm could in principle be assigned to the ⁵D₀ → ⁷F₁ transition or in our case, to the ⁵D₀ →⁷F₀ transition. According to Antic-Fidancev, a correlation exists between the position of the ⁵D₀ level of Eu³⁺ and the position of the barycenter of the ⁷F₁ manifold.⁴³ If the ⁵D₀ level is at 17300 cm^-1 (578.0 nm), the barycenter of the ⁷F₁ manifold is expected at about 400 cm^-1, which is not in agreement with our experimental values (17316 cm^-1 and 240 cm^-1, respectively).

The most intense line appears at 614 nm and corresponds to the forced Eu^{3+ 5}D₀ → ⁷F₂ electric transition, indicating the absence of inversion symmetry in the sites occupied by Eu³⁺ ions.⁵ All samples presented a narrowing and an unfolding of the ⁵D₀ → ⁷F₂ transition with an increase of the annealed temperature. That fact is related to the changing of structure crystal as indicated by the results discussed heretofore. Only Nb₂O₅:La_(0.7%)Eu_(0.3%) treated at 1200 ºC and excited at 275 nm presented unfolding of this transition around 619 nm, probably as a consequence of higher occupation of different active sites by the emitting ion, due to its larger concentration in the matrix.

The ratio between the integrated intensities of these two transitions, I_0-2/I_0-1, functions as a probe of the local cation surroundings. The ratio values decrease as a function of the thermal treatment temperature for both Nb₂O₅:La_(0.9%)Eu_(0.1%) and Nb₂O₅:La_(0.7%)Eu_(0.3%). For an example, the values found for the Nb₂O₅:La_(0.9%)Eu_(0.1%) sample excited at 465 nm were 6.18, 4.39, 2.64 and 2.02 for calcination temperatures of 550, 750, 900 and 1200 ºC, respectively; and for the Nb₂O₅:La_(0.7%)Eu_(0.3%) sample excited at 465 nm, the values ratio were 5.63, 4.72, 2.46 and 2.08 for calcination temperatures of 550, 750, 900, and 1200 ºC, respectively. The ratio between the intensities of the ⁵D₀→⁷F₂ and ⁵D₀ → ⁷F₁ transitions varies as a function of the thermal treatment temperature, showing that Eu³⁺ exists in a distorted (or asymmetric) cationic environment in the samples calcined at 550 and 750 ºC. At higher treatment temperatures, the Eu³⁺ ions are preferentially located close to more crystalline and symmetric regions.

The ⁵D₀ → ⁷F₃ transition presented a similar behavior to the ⁵D₀ → ⁷F₀ transition, with an increase of intensity in function of the thermal treatment for all samples. In general the ⁵D₀ → ⁷F₃ transition is very weak, because it is forbidden according to the Judd-Ofelt theory, and as the ⁵D₀ → ⁷F₀ transition, the ⁵D₀ → ⁷F₃ transition can only gain intensity via J-mixing.³⁸ The increase of its intensity for the higher tempetarure annealed samples indicates an increase of J-mixing and crystal-field perturbation, caused by the structural changes.

Other indicative of structural changes can be observed by the variation of the ⁵D₀ → ⁷F₄ transition. The intensity of this transition is determined not only by symmetry factors, but also by the chemical composition of the host matrix.³⁸ As can be observed in the emission spectrum, the ⁵D₀→⁷F₄ transition presented different profiles with an increase of the annealing temperature and the europium concentration. Comparing the samples calcined at different temperatures, it was observed an increase of intensity, a narrowing and an higher unfolding for this transition, caused by the structural changes in the matrix. These structural changes allowed to obtain a deep-red emitting phosphor (ca. 700 nm). This fact is important because if a larger color gamut is required, e.g., for high-quality TV, monitors for design applications, etc., the primaries, i.e., the light transmitted by the blue, green, and red sections of the color filter, must be more separated from each other.⁴⁴ In addition, slight changes on the intensity and split of this transition were also observed with an increase of europoium concentration, as showed in the emission spectrum for the samples annealed at 900 and 1200 ºC in Figures 5b-e and 6b-e.

Figure 7 brings the decay curves for the emission from the ⁵D₀ states for both samples, treated at different temperatures. Table 1 lists the excited state lifetimes. The decay does not follow a simple exponential law, confirming the distribution of symmetry sites for the Eu³⁺ ion along the sample and providing a mean lifetime (τ_exp). In general, the Eu³⁺ ion lifetime increases in function of the thermal treatment temperature, which causes an increase on the structural crystallinity of the host as also confirmed by the X-rays diffraction and by the narrowing of the intraconfigurational f-f transtitions of the Eu³⁺ ions in the emission spectrum.

Figure 7
Decay curves for the Eu³⁺ ion in the matrixes (a) Nb₂O₅:La_(0.9%) Eu_(0.1%) and (b) Nb₂O₅:La_(0.7%)Eu_(0.3%) treated at different temperatures.

Table 1
Judd-Ofelt intensity parameters (Ω₂, Ω₄), probability of spontaneous Einstein emission (A_RAD), radiative and experimental lifetimes (τ_RAD, τ_EXP), and quantum efficiency (q) of the calcined Nb₂O₅:La_(0.9%)Eu_(0.1%) and Nb2O5:La_(0.7%)Eu_(0.3%) samples

Sample	Temperature / °C	Ω2 / (10-20 cm2)	Ω4 / (10-20 cm2)	ARAD / ms-1	τRAD / ms	τEXP / ms	q
Nb2O5:La(0.9%)Eu(0.1%)	550	6.2 ± 0.3	5.7 ± 0.3	0.54 ± 0.03	1,85 ± 0.09	0.66 ± 0.07	0.36 ± 0.05
	750	4.4 ± 0.2	4.7 ± 0.2	0.42 ± 0.02	2.38 ± 0.1	0.75 ± 0.08	0.32 ± 0.05
	900	2.6 ± 0.1	5.8 ± 0.3	0.39 ± 0.02	2.59 ± 0.1	0.94 ± 0.09	0.36 ± 0.05
	1200	2.0 ± 0.1	6.1 ± 0.3	0.38 ± 0.02	2.64 ± 0.1	0.93 ± 0.09	0.35 ± 0.05
Nb2O5:La(0.7%)Eu(0.3%)	550	5.6 ± 0.3	4.8 ±0.2	0.49 ± 0.02	2.05 ± 0.1	0.70 ± 0.07	0.34 ± 0.05
	750	4.7 ± 0.2	6.0 ± 0.3	0.48 ± 0.02	2.08 ± 0.1	0.71 ± 0.07	0.34 ± 0.05
	900	2.4 ± 0.1	6.8 ± 0.3	0.41 ± 0.02	2.41 ± 0.1	0.96 ± 0.10	0.40 ± 0.06
	1200	2.1 ± 0.1	5.3 ± 0.3	0.37 ± 0.02	2.71 ± 0.1	0.97 ± 0.10	0.36 ± 0.05

The emission spectra allowed us to calculate the spontaneous emission coefficient (A_RAD), the radiative lifetime (τ_RAD), the quantum efficiency (q), and the Judd-Ofelt (Ω₂, Ω₄) parameters (Table 1). The ⁵D₀ → ⁷F₁ transition is purely magnetically dipolar, and its radiative rate does not depend on the local field imposed by the environment. The relation A_0-λ = A_0-1(S_0-λ/S_0-1)(σ_λ/σ₁), where S_0-λ is the area under the curve related to the ⁵D₀→⁷F_λ transition obtained from the spectral data, and σ_λ is the energy barycenter of the 0-λ transition, furnishes the A_0-λ values. Here, we assumed that A_0-1= 50 s^-1, and that ⁵D₀→⁷F_5,6 have negligible intensities. Although the calculations of radiative rates were based in lanthanide complexes in solutions, they have been used for lanthanides doped in solid state matrices as a form to evaluate its luminescent properties.^45-48 The details of these calculation were given by Werts et al.⁴⁹ The mean lifetime was calculated using the equation <τ> = ∑α_i.τ_i/∑α_i; α and τ represent the pre-exponential factor or amplitude and lifetime of the decay components, respectively.⁵⁰

All the samples have presented similar q values between 30 and 40%. Rising calcination temperatures reduce A_RAD and Ω₂. Reisfeld and Jørgensen⁵¹ consider that Ω₂ is a useful parameter, because it is sensitive to the local symmetry of the ligand field and bond covalency. The value of Ω₂ increases as the local symmetry of the ligand field decreases and the bond covalency increases. The Ω₄ and Ω₆ are related to the viscosity and rigidity of the host medium in which the ions are situated, but there is no theoretical prediction for this sensibility to macroscopic properties.^51,52

The increase of the annealed temperatures promoted a decrease of the Ω₂ values for the samples, indicating an increase of the local symmetry for the Eu³⁺ ions.⁵³ This fact corroborates with the results obtained by the X-ray diffraction that shows a higher crystallinity phases for the samples calcined at higher temperatures. In addition, these structural changes were also observed in the emission spectra of the Eu³⁺ ions in different annealed samples. However, the low Ω₂ and Ω₄ values obtained suggest that a weakly polarizable and rigid chemical environment surrounds Eu³⁺.⁵³

The stronger transition at 614 nm favors saturated CIE chromaticity. We generated chromaticity coordinates for the Nb₂O₅:La_(0.9%)Eu_(0.1%) and Nb₂O₅:La_(0.7%)Eu_(0.3%) samples using the software Spectra Lux 2.0⁵⁴ and the respective emission spectra recorded at room temperature (Tables 2 and 3). The CIE chromaticity coordinates of all the samples excited at 465 nm lie above the NTSC standard values (x = 0.670 and y = 0.330).⁵ Moreover, the full width at half maximum (FWHM) values diminish with increasing treatment temperature: Nb₂O₅:La_(0.9%)Eu_(0.1%) calcined at 1200 ºC affords the lowest value among the investigated samples, 3.5 nm.

Table 2
Chromaticity coordinates and FWHM of the Eu^{3+ 5}D₀ → ⁷F₂ transition in calcined Nb₂O₅:La_(0.9%)Eu_(0.1%) and Nb₂O₅:La_(0.7%)Eu_(0.3%) samples excited at 275 and 394 nm

Sample	Temperature / ºC	λexc:a 275 nm			λexc:a 394 nm
		CIE		FWHMb / nm	CIE		FWHMb / nm
		x	y		x	y
Nb2O5:La(0.90)Eu (0.1)	550	0.666	0.333	12.0	0.676	0.322	12.4
	750	0.660	0.338	5.2	0.662	0.336	6.3
	900	0.669	0.329	6.0	0.670	0.329	5.7
	1200	0.658	0.341	9.0	0.680	0.318	4.4
Nb2O5:La(0.7)Eu(0.3)	550	0.681	0.318	12.0	0.683	0.316	13.8
	750	0.657	0.342	6.2	0.678	0.321	7.7
	900	0.651	0.347	5.0	0.671	0.327	4.8
	1200	0.649	0.350	9.0	0.676	0.322	5.2

• ^a
  Excitation spectra;
• ^b
  full width at half maximum.

Table 3
Chromaticity coordinates, quantum efficiency (q) and FWHM of the Eu^{3+ 5}D₀ → ⁷F₂ transition in calcined Nb₂O_5:La_(0.9%)Eu_(0.1%) and Nb₂O₅:La_(0.7%) Eu_(0.3%) samples excited at 465 and 525 nm

Sample	Temperature / ºC	λexc.: 465 nm			λexc.: 525 nm
		CIE		FWHMa / nm	CIE		FWHMb / nm
		x	y		x	y
Nb2O5:La(0.9%)Eu(0.1%)	550	0.680	0.319	10.5	0.671	0.328	13.5
	750	0.681	0.317	7.0	0.663	0.335	7.7
	900	0.678	0.321	5.0	0.662	0.337	6.0
	1200	0.682	0.317	3.5	0.660	0.339	3.5
Nb2O5:La(0.7%)Eu(0.3%)	550	0.685	0.313	12.5	0.683	0.316	13.5
	750	0.687	0.312	11.3	0.689	0.309	9.1
	900	0.683	0.316	5.0	0.682	0.316	5.0
	1200	0.682	0.317	4.0	0.682	0.316	4.5

• ^a
  Full width at half maximum.

Conclusions

In this work, we prepared Nb₂O₅: La³⁺, Eu³⁺ phosphors by the non-hydrolytic sol-gel process. The SEM images revealed that particles resembled pellets arranged into sheets with thickness of ca. 60 nm and the samples annealed above 900 ºC elicited a totally crystalline monoclinic phase. On the contrary of Y₂O₂S:Eu³⁺ which efficiently absorbs light at around 370 nm^55,56 the luminescence results shown that all samples presented a wide excitation reange, including the light emission region of a light emitting diode emitting UV (275, 394 nm), the blue (465 nm) and green (525 nm) light. The CIE chromaticity coordinates were above the NTSC standard values. Hence, these two phosphors are promising red components for LED applications.

• FAPESP has sponsored the publication of this article.

Acknowledgments

The authors would like to deeply express our most sincere thanks to Jeannette Dexpert-Ghys (CEMES) for her profitable assistance and cooperation. The agencies CNPq, and CAPES (Brazilian research funding agencies), grant 2011/51858-0 Brazil-France cooperation program FAPESP-CNRS; 2011/09823-4 and 2012/11673-3 E.J.N.; 2011/15757-4 L.A.R.; São Paulo Research Foundation (FAPESP). The authors thank the Laboratory of Photonic Materials (Prof S. J. L. Ribeiro) for the use of the RAMAN spectrometer and CBMM (Companhia Brasileira de Metalurgia e Mineração) for donating NbCl₅.

References

• ¹
  Huang, J.; Zhou, L.; Liang, Z.; Gong, F.; Han, J.; Wang, R.; J. Rare Earths2010 , 28, 356.
• ²
  Neeraj, S.; Kijima, N.; Cheetham, A. K.; Chem. Phys. Lett 2004 , 387, 2.
• ³
  Thejo Kalyani, N.; Dhoble, S. J.; Renewable Sustainable Energy Rev 2012 , 16, 2696.
• ⁴
  Wu, Z. C.; Shi, J. X.; Wang, J.; Gong, M. L.; Su, Q.; J. Solid State Chem 2006 , 179, 2356.
• ⁵
  Zhou, L.; Huang, J.; Gong, F.; Lan, Y.; Tong, Z.; Sun, J.; J. Alloys Compd 2010 , 495, 268.
• ⁶
  Ferrari, J. L.; Cebim, M. A.; Pires, A. M.; dos Santos, M. A. C.; Davolos, M. R.; J. Solid State Chem 2010 , 183, 2110.
• ⁷
  Kodaira, C. A.; Stefani, R.; Maia, A. S.; Felinto, M. C. F. C.; Brito, H. F.; J. Lumin 2007 , 127, 616.
• ⁸
  Pires, A. M.; Serra, O. A.; Davolos, M. R.; J. Alloys Compd. 2004, 374, 181.
• ⁹
  Kutty, T. R. N.; Nag, A.; J. Mater. Chem 2003 , 13, 2271.
• ¹⁰
  Aquino, F. T.; Ferrari, J. L.; Ribeiro, S. J. L.; Ferrier, A.; Goldner, P.; Gonçalves, R. R.; Opt. Mater. (Amsterdam, Neth.)2013 , 35, 387.
• ¹¹
  Aquino, F. T.; Pereira, R. R.; Ferrari, J. L.; Ribeiro, S. J. L.; Ferrier, A.; Goldner, P.; Gonçalves, R. R.; Mater. Chem. Phys 2014 , 147, 751.
• ¹²
  Palatnikov, M.; Shcherbina, O.; Sidorov, N.; Skab, I.; Bormanis, K.; Ukr. J. Phys. Opt.2012 , 13, 207.
• ¹³
  Rooksby, H. P.; White, E. A. D.; Acta Crystallogr 1963 , 16, 888.
• ¹⁴
  Emmenegger, F. P.; Robinson, M. L. A.; J. Phys. Chem. Solids1968 , 29, 1673.
• ¹⁵
  Nowak, I.; Ziolek, M.; Chem. Rev 1999 , 99, 3603.
• ¹⁶
  Blanquart, T.; Niinistö, J.; Heikkilä, M.; Sajavaara, T.; Kukli, K.; Puukilainen, E.; Xu, C.; Hunks, W.; Ritala, M.; Leskelä, M.; Chem. Mater 2012 , 24, 975.
• ¹⁷
  Blasse, G.; Grabmaier, B. C; Luminescent Materials; Springer: Berlin, 1994.
• ¹⁸
  Nassar, E. J.; Avila, L. R.; Pereira, P. F. S.; de Lima, O. J.; Rocha, L. A.; Mello, C.; Ciuffi, K. J.; Carlos, L. D.; Quim. Nova2005 , 28, 238.
• ¹⁹
  Ricci, G. P.; Rocha, Z. N.; Nakagaki, S.; Castro, K. A. D. F.; Crotti, A. E. M.; Calefi, P. S.; Nassar, E. J.; Ciuffi, K.; Appl. Catal., A2010 , 389, 147.
• ²⁰
  Mutin, P. H.; Vioux, A.; Chem. Mater 2009 , 21, 582.
• ²¹
  Sommerdijk, N.; Wright, J. D.; J. Sol-Gel Sci. Technol.1998 , 13, 565.
• ²²
  Falcomer, D.; Speghini, A.; Ibba, G.; Enzo, S.; Cannas, C.; Musinu, A.; Bettinelli, M.; J. Nanomater 2007 , 2007, DOI http://dx.doi.org/10.1155/2007/94975.
  » http://dx.doi.org/10.1155/2007/94975
• ²³
  Graça, M. P. F.; Meireles, A.; Nico, C.; Valente, M. A.; J. Alloys Compd 2013 , 553, 177.
• ²⁴
  Luo, H.; Wei, M.; Wei, K.; Mater. Chem. Phys.2010 , 120, 6.
• ²⁵
  Wei, M.; Wei, K.; Ichihara, M.; Zhou, H.; Electrochem. Commun.2008 , 10, 980.
• ²⁶
  Bueno, P. R.; Faria, R. C.; Bulhões, L. O. S.; Solid State Ionics2005 , 176, 1175.
• ²⁷
  Acosta, S.; Corriu, R. J. P.; Leclercq, D.; Lefèvre, P.; Mutin, P. H.; Vioux, A.; J. Non-Cryst. Solids1994 , 170, 234.
• ²⁸
  Konings, R. J. M.; Booij, A. S.; Vib. Spectrosc.1994 , 6, 345.
• ²⁹
  McConnell, A. A.; Anderson, J. S.; Rao, C. N. R.; Spectrochim. Acta, Part A1976 , 32, 1067.
• ³⁰
  Boldish, S. I.; White, W. B.; Spectrochim. Acta, Part A1979 , 35, 1235.
• ³¹
  Weckhuysen, B. M.; Phys. Chem. Chem. Phys 2003 , 5, 4351.
• ³²
  Burton, A. W.; Ong, K.; Rea, T.; Chan, I. Y.; Microporous Mesoporous Mater 2009 , 117, 75.
• ³³
  Zhang, X.; Meng, F.; Wei, H.; Seo, H. J.; Ceram. Int 2013 , 39, 4063.
• ³⁴
  Freiria, G. S.; Nassar, E. J.; Verelst, M.; Rocha, L. A.; J. Lumin.2015 , DOI dx.doi.org/10.1016/j.jlumin.2015.06.022.
  » dx.doi.org/10.1016/j.jlumin.2015.06.022
• ³⁵
  Hirano, M.; Dozono, H.; Mater. Res. Bull 2014 , 50, 213.
• ³⁶
  Hsiao, Y. J.; Fang, T. H.; Chang, Y. S.; Chang, Y. H.; Liu, C. H.; Ji, L. W.; Jywe, W. Y.; J. Lumin 2007 , 126, 866.
• ³⁷
  Mahesh, S. K.; Rao, P. P.; Francis, T. L.; Reshmi, V. R.; Koshy, P.; Mater. Lett.2014 , 120, 115.
• ³⁸
  Binnemans, K.; Coord. Chem. Rev 2015 , 295, 1.
• ³⁹
  Rocha, L. A.; Avila, L. R.; Caetano, B. L.; Molina, E. F.; Sacco, H. C.; Ciuffi, K. J.; Calefi, P. S.; Nassar, E. J.; Mater. Res 2005 , 8, 361.
• ⁴⁰
  Blasse, G.; Chem. Phys. Lett.1975 , 33, 616.
• ⁴¹
  Antic-Fidancev, E.; Hölsä, J.; Lastusaari, M.; J. Alloys Compd 2002 , 341, 82.
• ⁴²
  Porcher, P.; Svoronos, D. R.; Leskelä, M.; Hölsä, J.; J. Solid State Chem 1983 , 46, 101.
• ⁴³
  Antic-Fidancev, E.; J. Alloy Compd 2000 , 300-301, 2.
• ⁴⁴
  Jüstel, T.; Möller, S. ; Winkler, H. ; Adam, W. ; Ullmann's Encyclopedia of Industrial Chemistry, Wiley-VCH Verlag GmbH: Weinheim, 2012. DOI 10.1002/14356007.a15_519.pub2.
  » https://doi.org/10.1002/14356007.a15_519.pub2
• ⁴⁵
  Rocha, L. A.; Ciuffi, K. J.; Sacco, H. C.; Nassar, E. J.; Mater. Chem. Phys.2004 , 85, 245.
• ⁴⁶
  Yan, B.; Li, Q. P.; Microporous Mesoporous Mater 2014 , 196, 284.
• ⁴⁷
  Shao, Y.-F.; Yan, B.; Microporous Mesoporous Mater 2014 , 193, 85.
• ⁴⁸
  Gibelli, E.; Kai, J.; Teotonio, E.; Felinto, M.; Brito, H.; J. Photochem. Photobiol., A2013 , 251, 154.
• ⁴⁹
  Werts, M. H. V.; Jukes, R. T. F.; Verhoeven, J. W.; Phys. Chem. Chem. Phys 2002 , 4, 1542.
• ⁵⁰
  Basu, B. B. J.; Vasantharajan, N.; J. Lumin 2008 , 128, 1701.
• ⁵¹
  Reisfeld, R.; Jørgensen, C. K. In Handbook on the Physics and Chemistry of Rare Earth; Gschneidner, K. A., Eyring, L., eds.; North Holland: Amsterdam, 1984.
• ⁵²
  Barve, R. A.; Suriyamurthy, N.; Panigrahi, B. S.; Venkatraman, B.; Phys. B (Amsterdam, Neth.)2015 , 475, 156.
• ⁵³
  Santos, J. G.; Dutra, J. D. L.; Alves, S.; de Sá, G. F.; da Costa, N. B.; Freire, R. O.; J. Braz. Chem. Soc 2013 , 24, 236.
• ⁵⁴
  Santa-Cruz, P. A.; Teles, F. S.; Spectra Lux Software; Ponto Quantico Nanodispositivos: Recife, 2003.
• ⁵⁵
  Fu, Z.; Geng, Y.; Chen, H.; Zhou, S.; Yang, H. K.; Jeong, J. H.; Opt. Mater. (Amsterdam, Neth.)2008 , 31, 58.
• ⁵⁶
  Thirumalai, J.; Jagannathan, R.; Trivedi, D. C.; J. Lumin 2007 , 126, 353.

Publication Dates

• Publication in this collection
  Nov 2015

History

• Received
  09 June 2015
• Accepted
  18 Sept 2015