Layered rare-earth hydroxide and oxide nanoplates of the Y/Tb/Eu system: phase-controlled processing, structure characterization and color-tunable photoluminescence via selective excitation and efficient energy transfer

Abstract

Well-crystallized (Y_0.97−xTb_0.03Eu_x)₂(OH)₅NO₃·nH₂O (x = 0–0.03) layered rare-earth hydroxide (LRH) nanoflakes of a pure high-hydration phase have been produced by autoclaving from the nitrate/NH₄OH reaction system under the optimized conditions of 100 °C and pH ∼7.0. The flakes were then converted into (Y_0.97−xTb_0.03Eu_x)₂O₃ phosphor nanoplates with color-tunable photoluminescence. Detailed structural characterizations confirmed that LRH solid solutions contained NO₃⁻ anions intercalated between the layers. Characteristic Tb³⁺ and Eu³⁺ emissions were detected in the ternary LRHs by selectively exciting the two types of activators, and the energy transfer from Tb³⁺ to Eu³⁺ was observed. Annealing the LRHs at 1100 °C produced cubic-lattice (Y_0.97−xTb_0.03Eu_x)₂O₃ solid-solution nanoplates with exposed 222 facets. Multicolor, intensity-adjustable luminescence was attained by varying the excitation wavelength from ∼249 nm (the charge transfer excitation band of Eu³⁺) to 278 nm (the 4f⁸–4f⁷5d¹ transition of Tb³⁺). Unitizing the efficient Tb³⁺ to Eu³⁺ energy transfer, the emission color of (Y_0.97−xTb_0.03Eu_x)₂O₃ was tuned from approximately green to yellowish-orange by varying the Eu³⁺/Tb³⁺ ratio. At the optimal Eu³⁺ content of x = 0.01, the efficiency of energy transfer was ∼91% and the transfer mechanism was suggested to be electric multipole interactions. The phosphor nanoplates developed in this work may be incorporated in luminescent films and find various lighting and display applications.

1. Introduction

Rare-earth elements are used in phosphors for fluorescent lamps, white-light-emitting diodes, plasma display panels, flat-panel displays, field emission displays, and cathode-ray tubes. Current advances in high-definition display technologies require fine phosphor particles of well-defined morphology to improve the resolution by reducing pixel size, improve the overall luminescent performance and reduce the size of optoelectronic devices. As the traditional solid-state reaction route generally yields micron-sized phosphor particles with a wide size distribution and considerable aggregation [1], the past decade witnessed a boom in solution-based processing of phosphor particles with better-defined characteristics [2–7]. Y₂O₃ doped with Eu³⁺ and Tb³⁺ ions are among the best-known red-emitting and green-emitting phosphors, respectively, owing to their simple chemical composition, excellent luminescence efficiency, high color purity and long-term stability [8–10]. Tb³⁺ has also been used as a sensitizer to enhance Eu³⁺ emission in some phosphor systems such as Y₂O₂S:Eu³⁺ [11]. Among the low-dimensional Y₂O₃:RE phosphors (RE = rare-earth element), zero-dimensional (0D) nanoparticles [12–14] and 1D nanowires/nanotubes [15–19] have been extensively studied for their controlled synthesis, surface modification, size and shape dependent optical properties, and new applications in areas such as bio-analysis [20]. However, reports on 2D nanosheets were rare. In the solution-based processing methodology, oxide phosphors are mostly produced via controlled precursor synthesis followed by proper annealing. The lack of studies on 2D nanophosphors originated from the scarcity of precursor nanosheets or precursor compounds that could be delaminated into nanosheets. A breakthrough came in 2006 [21] when Gandara et al synthesized interlayer-anion exchangeable layered rare-earth hydroxides (LRHs) of the typical composition Ln₂(OH)₅A·nH₂O, where A is interlayer anion such as NO₃⁻, Cl⁻ or Br⁻; Ln is lanthanide (but may also be Y); and n = 1.5–1.8. Since then, considerable attention has been paid to LRHs owing to the interesting magnetic, catalytic and optical properties of the Ln elements. It was found that the LRH compounds tend to crystallize under roughly neutral pH, either through hydrothermal or refluxing synthesis, and the presence of mineralizers containing anion A helps stabilize the corresponding LRH phase [22–28]. Detailed crystal structure analysis by Sasaki et al [26] revealed two types of coordination environments for Ln³⁺, that is, eight-fold coordinated [Ln(OH)₇(H₂O)] (C₁ symmetry) and nine-fold coordinated [Ln(OH)₈(H₂O)] (C_4v symmetry) polyhedrons. In the LRH structure, each LnO₈ unit is linked to two other LnO₈ and four LnO₉ groups via edge sharing. The linked polyhedron units form a two-dimensional host layer parallel to the ab plane, with the free anions A located in the interlayer to support the layers and for charge balance. Photoluminescence studies on the Y/Eu binary [29] and Y/Gd/Eu ternary [30] LRH systems showed that the optical properties crucially depend on the local symmetry around Eu³⁺ ions. Luminescence systematically changes with the dehydration of LRH (decreasing n value), since the loss of crystal water shifts the Eu³⁺ symmetry from C_4v to C₁ [29]. It was also found that in the Y/Gd/Eu ternary system Gd³⁺ preferentially sensitizes the emission of Eu³⁺ ions possessing C₁ local symmetry [30].

Since the host layer of LRH is a close-packed (low energy) crystal plane, the LRH compounds tend to crystallize as nanoplates. The 2D morphologies of the pristine and the exfoliated LRHs are thus well suited for the construction of highly [001] oriented films with improved luminescence, since the [001] orientation maximizes the exposure of Ln activators to the excitation light [31–34]. Meanwhile, the [001] oriented LRH film would transform into a [111] oriented oxide film via quasitopotactic atomic arrangements under proper annealing [31–33], which should result in an enhanced exposure of the close-packed 222 facets of the oxide crystals and thus stronger luminescence [32–34]. Improved catalytic performance may similarly be expected from the maximized Ln exposure.

In view of the importance and wide application of Y₂O₃:Eu³⁺ red and Y₂O₃:Tb³⁺ green phosphors, Y/Tb binary and Y/Tb/Eu ternary LRHs were synthesized in this work for optical studies of the LRHs and their annealing-derived oxide nanoplates. With Eu³⁺ incorporation, color-tunable photoluminescence was achieved either by varying the excitation wavelength or through an efficient energy transfer from Tb³⁺ to Eu³⁺ ions. In the following sections, we report the phase purity controlled synthesis, detailed structure characterization and optical properties of the LRH compounds and their calcination-derived oxides.

2. Experimental procedures

2.1. Hydrothermal synthesis of LRH solid solutions

The starting rare-earth sources were Ln₂O₃ (Ln = Y, Tb and Eu; 99.99% purity; Conghua Jianfeng Rare Earth Co. Ltd, Conghua, China). Nitrate solutions of Ln³⁺ ions were prepared by dissolving the corresponding oxides with a proper amount of hot nitric acid. For each synthesis run, the total Ln³⁺ concentration was kept at 0.08 mol l⁻¹. Optimal Tb concentration in the Y₂O₃:Tb³⁺ green phosphor was previously established at 5 at.% [35, 36], and in this work the Tb/(Y + Tb + Eu) atomic ratio was fixed at 3 at.%. The Eu content in the (Y_0.97−xTb_0.03Eu_x) combination was varied in the range of x = 0–0.03 to reveal its effects on the optical properties. In a typical synthesis, a proper amount of ammonium hydroxide solution (25 wt%) was added to the mixed nitrate solution until pH ∼ 7, and the resultant suspension was then transferred into a Teflon-lined stainless steel autoclave after room-temperature homogenization for 30 min under magnetic stirring. The tightly sealed autoclave was placed in an electric oven preheated to a certain temperature in the range of 100–220 °C for a reaction period of 24 h. After natural cooling, the hydrothermal product was collected via centrifugation, washed with distilled water and ethanol, and finally dried in the air at 70 °C for 24 h to yield a white powder for characterization and further processing.

2.2. Characterization techniques

Crystalline phases were identified with an x-ray diffractometer (XRD, Model RINT 2200 V/PC, Rigaku, Tokyo, Japan) operated at 40 kV per 40 mA, using nickel-filtered Cu-Kα radiation and a scanning speed of 1° 2θ per minute. Fourier transform infrared spectroscopy (FTIR, Model 4200, JASCO, Tokyo) measurements were performed using KBr pellets. Morphology and microstructure were analyzed by transmission electron microscopy (TEM, 200 kV, Model JEM-2000FX, JEOL, Tokyo) and field emission scanning electron microscopy (FE-SEM, 10 kV, Model S-5000, Hitachi, Tokyo). Differential thermal analysis/thermogravimetry (DTA/TG, Model Thermo Plus TG8120, Rigaku, Tokyo) of the LRHs was carried out in stagnant air with a heating rate of 10 °C min⁻¹. Photoluminescence (PL), photoluminescence excitation (PLE) and fluorescence decay were analyzed at room temperature using an FP-6500 spectrofluorometer (JASCO, Tokyo) equipped with a 60 mm-diameter integrating sphere (Model ISF-513, JASCO) and a 150 W Xe lamp as the excitation source. Excitation and emission beams were monochromatized with a Rowland concave grating (1800 grooves mm⁻¹). Optical measurements were conducted at a scanning speed of 100 nm min⁻¹, wavelength accuracy of ±0.3 nm and slit widths of 5 nm both in excitation and emission. Spectral response of the system was corrected in the range of 220–850 nm with a Rhodamine-B solution (5.5 g l⁻¹ in ethylene glycol) and a standard light source unit (ECS-333, JASCO) as references.

3. Results and discussion

3.1. Phase-controlled processing, structure characterization and optical properties of the Y/Tb/Eu LRHs

The LRH phase was widely observed to crystallize under approximately neutral pH [22–30]. When keeping solution pH at ∼7.0, we found that processing temperature significantly affects the phase composition of the hydrothermal product, as shown in figure 1 for the example of Y_0.965Tb_0.03Eu_0.005 system. The sample synthesized at 100 °C exhibits a series of strong 00l and sharp 220 diffractions, suggesting that a Ln₂(OH)₅NO₃·nH₂O LRH compound with highly ordered host layers has been formed [21–34, 37–40]. The samples processed at higher temperatures up to 200 °C similarly show the 220 diffraction of LRH, but splitting of the 002 and 004 diffractions was clearly observed. Lattice spacing calculations (table S1; see the supplementary material available from stacks.iop.org/STAM/14/015006/mmedia) confirmed that the two additional peaks belong to oscillation diffractions in each case. The above results thus imply that these higher-temperature samples contain two phases with different interlayer distances. Dehydration was believed to be responsible for the appearance of the LRH phase with a shorter interlayer distance. Sasaki and co-workers [26–28] revealed that, in the LRH structure, water molecules are directly coordinated to the Ln atoms to form coordination polyhedrons, instead of being intercalated between the layers, and the loss of hydration water shortens the interlayer distance. This explanation was confirmed by water content analysis of the samples (figure S1). The LRH phase dehydrates up to ∼260 °C [29, 30], and the average water contents derived from the TG curves at 260 °C indeed tend to decrease toward a higher synthesis temperature (table S1). The two phases are thereafter referred to high hydration (HH, larger interlayer spacing) and low hydration (LH) phases. The LH content can be evaluated from the intensities of the 002 diffractions of the HH (I_HH002) and LH (I_LH002) phases using the ratio I_LH002/(I_HH002 + I_LH002). Results show that the samples synthesized at 130, 150, 180 and 200 °C have ∼40.2, 31.2, 45.8 and 49.2 wt% of the LH phase, respectively. Though the LH fraction tends to increase for a higher processing temperature, attempts to obtain a pure LH phase failed, as already at 220 °C the product changed into an anhydrous Ln₄O(OH)₉NO₃ compound (JCPDS file no. 79-1352, figure 1(f)).

Figure 1.

Powder XRD patterns of the hydrothermal products fabricated at (a) 100, (b) 130, (c) 150, (d) 180, (e) 200 and (f) 220 °C. All samples have the same cation combination of Y_0.965Tb_0.03Eu_0.005; HH and LH represent high-hydration and low-hydration phases, respectively.

FTIR spectra (figure S2) of the 100 °C sample show a single and sharp absorption band centered at ∼1387 cm⁻¹, characteristic of free nitrate ions (the v₃ vibration mode) in the interlayer of LRHs [21–28, 37, 41, 42]. The additional absorption at ∼3600 cm⁻¹ is indicative of the hydroxide anion (OH⁻) [38, 39], while the 3380 and 1640 cm⁻¹ peaks can be assigned to the O–H stretching vibrations (ν₁ and ν₃) and the H–O–H bending mode (ν₂) of the water molecules, respectively [38, 39]. This analysis reveals that an LRH compound has been crystallized with free NO₃⁻ anions intercalated between the layers, in agreement with the XRD results. Although a shortened interlayer spacing might lead to the coordination of NO₃⁻ anions to Ln in the LH phase, FTIR spectra of the samples processed up to 200 °C exhibit the single absorption of free NO₃⁻ anion at ∼1387 cm⁻¹, and are clearly different from the spectrum of the Ln₄O(OH)₉NO₃ compound obtained at 220 °C (spectrum d in figure S2). Since the nitrate anions are indirectly linked to Ln via the hydrogen bonding between OH⁻ and NO₃⁻ anions in Ln₄O(OH)₉NO₃ [37], two peaks appear at ∼1370 and 1410 cm⁻¹, corresponding to the ν₃ vibration of NO₃⁻ anion and the ν₄ asymmetric stretching of the O–NO₂ group, respectively [37, 41, 42].

FE-SEM observation indicated that the product obtained at 100 °C crystallizes as hexagonal nanoflakes with lateral sizes of up to ∼500 nm and thicknesses of up to 80 nm (figure S3). Raising the hydrothermal temperature yields larger crystallites, and at 150 °C some crystallites of up to ∼2 μm in lateral size and over 150 nm in thickness were produced. For the sake of phase purity, however, 100 °C was chosen as the hydrothermal temperature for the further synthesis.

Figure 2 shows XRD patterns of the powders synthesized at 100 °C, with varying x in the Y_0.97−xTb_0.03Eu_x combination. The strong 00l and 220 diffractions imply that well-crystallized LRHs have been produced. Splitting of the 00l diffractions was no longer observed in any sample. Closer observations revealed that the 00l diffractions slightly shift to the high-angle side and the 220 diffraction to the low-angle side with increasing Eu³⁺ concentration. As the 00l diffractions reflect the stacking of host layers along the c-axis of the crystal structure while the 220 diffraction originates within the host layers, the above results thus suggest that the interlayer distance shrinks while the ab plane expands at a higher Eu³⁺ content. The 002 basal spacing and 220 lattice spacing calculated from the XRD patterns are shown in table S2 (see the supplementary material available from stacks.iop.org/STAM/14/015006/mmedia). The shrinking interlayer distance is again caused by lowered hydration, as seen from the n value (table S2) derived via TG analysis (figure S4). Sasaki and co-workers [26–28] reported that, for the nitrate family of LRH, the hydration number n in Ln₂(OH)₅NO₃·nH₂O decreases for a bigger Ln³⁺ ion, and this accounts for the successively lower hydration with increasing Eu³⁺ incorporation. We have reported similar phenomena in the Y/Eu binary [29] and Y/Gd/Eu ternary [30] nitrate-LRH systems. The expansion of the ab plane (host layer) results from the incorporation of relatively large Eu³⁺ ions (for eight-fold coordination,r_Y³⁺ = 0.1019 , r_Tb³⁺ = 0.1040 and r_Eu³⁺ = 0.1066 nm; for nine-fold coordination, r_Y³⁺ = 0.1075 , r_Tb³⁺ = 0.1095 and r_Eu³⁺ = 0.1120 nm) [43]. These results demonstrate that well-crystallized Y/Tb/Eu LRHs have been produced in the solid solution form.

Figure 2.

Powder XRD patterns of the Y_0.97−xTb_0.03Eu_x LRH samples synthesized at 100 °C.

A low-magnification TEM image (figure 3(a)) shows the typical flake-like nanocrystals of the Y_0.965Tb_0.03Eu_0.005 LRH sample. The hexagonal nanoplates exhibit a relatively wide size distribution and the biggest has a lateral size of ∼500 nm. Platelets as thin as ∼15–20 nm are seen in the micrograph. FE-SEM observations indicate that Eu content does not significantly affect the crystal morphology (figure S5). High-resolution TEM (HR-TEM) analysis clearly resolves lattice fringes of the LRH crystal, and the spacing of 0.312 nm may well correspond to the 220 plane of the host layer (figure 3(b)). Selected area electron diffraction (SAED) yielded well-defined diffraction spots (figure 3(c)), indicating that the object under observation is a well-crystallized single crystal. Indexing the SAED pattern (figure 3(c)) is consistent with the ab plane of the LRH crystal.

Figure 3.

Low-magnification TEM (a) and HR-TEM (b) images and SAED pattern (c) of the Y_0.965Tb_0.03Eu_0.005 LRH.

To elucidate the optical properties of the Y/Tb/Eu ternary system, two binary LRHs of the Y_0.995Eu_0.005 and Y_0.97Tb_0.03 combinations were synthesized and their PL and PLE spectra are shown in figure S6 (see the supplementary material available from stacks.iop.org/STAM/14/015006/mmedia). The PLE spectrum of the Y_0.995Eu_0.005 LRH obtained by monitoring the red emission of Eu³⁺ ions at 615 nm is composed of a relatively strong charge transfer (CT) excitation band at ∼256 nm and a series of sharp lines in the range 350–550 nm that are ascribed to the intra-4f⁶ electronic transitions of Eu³⁺, as labeled in the figure (figure S6(a)). The CT band is assigned to electron transfer from the 2p orbital of O²⁻ to the 4f orbital of Eu³⁺. Such a strong band was not observed before for the Y/Eu and Gd/Eu binary systems [29, 32] and its origin is still unclear. The PLE of the Y_0.97Tb_0.03 LRH, recorded by monitoring the green emission of Tb³⁺ at 545 nm, exhibits a less-defined broad 4f–5d transition band in the 230–295 nm region [44], along with peaks at 295–500 nm that are attributed to the intra-4f⁸ transitions between the ⁷F₆ and ⁵H_7–4, ⁵D_1,0, ⁵L_10–7,⁵G_6–2, and ⁵D_2–4 levels of Tb³⁺ (figure S6(a)) [45]. The Y_0.965Tb_0.03Eu_0.005 ternary LRH shows PLE spectra (figure 4(a)) almost identical to those of the Y_0.995Eu_0.005 and Y_0.97Tb_0.03 binary LRHs for the 615 and 545 nm emissions, respectively, implying that the characteristic emissions of Eu³⁺ and Tb³⁺ can both be obtained by selective excitation.

Figure 4.

(a) Luminescence excitation spectra of the Y_0.965Tb_0.03Eu_0.005 LRH obtained by monitoring the Eu³⁺ emission at 615 nm (red line) and Tb³⁺ emission at 545 nm (green line). (b) Emission spectra of the Y_0.965Tb_0.03Eu_0.005 LRH obtained under the different excitation wavelengths indicated in the figure.

Photoluminescence spectra of the Y_0.995Eu_0.005 and Y_0.97Tb_0.03 binary LRHs are shown in figure S6(b). Under 394 nm excitation (the ⁷F₀ → ⁵L₆ transition of Eu³⁺), the Y_0.995Eu_0.005 sample exhibits characteristic Eu³⁺ emissions at ∼593 (⁵D₀ → ⁷F₁), 615 (⁵D₀ → ⁷F₂) and 699 nm (⁵D₀ → ⁷F₄). Exciting the Y_0.97Tb_0.03 sample at 368 nm (the ⁷F₆ → ⁵D₂ transition of Tb³⁺, the strongest on the PLE curve) yields a strong green band at ∼545 nm (⁵D₄ → ⁷F₅ transition), weaker blue bands at ∼450 (⁵D₃ → ⁷F₅) and 492 nm (⁵D₄ → ⁷F₆), and minor red bands at 587 (⁵D₄ → ⁷F₄) and 628 nm (⁵D₄ → ⁷F₃) [45]. These typical Tb³⁺ emissions are also observed in Tb³⁺-doped Y₂O₃ [35, 36], Lu₂O₃ [46, 47] and yttrium aluminum garnet (YAG) [48, 49] green phosphors. Tunable emission (figure 4(b)) was achieved with the Y_0.965Tb_0.03Eu_0.005 ternary LRH by varying the excitation wavelength from 260 (the CT band of Eu³⁺) to 394 nm (the ⁷F₀ → ⁵L₆ transition of Eu³⁺), showing the advantages of Tb³⁺ and Eu³⁺ codoping. Under 260 nm excitation, the PL spectrum is dominated by Eu³⁺ emissions, along with weaker ones from Tb³⁺. The Tb³⁺ emission is seen because the excitation wavelength resides on the left slope of the 4f–5d transition band of Tb³⁺ (figure 4(a)). Upon excitation at 284 nm (the 4f⁸–4f⁷5d¹ transition of Tb³⁺), the sample exhibits strong Tb³⁺ and weak Eu³⁺ signals. Eu³⁺ emission appears since this excitation wavelength is located on the right-hand tail of the CT band of Eu³⁺ (figure 4(a)). Only Tb³⁺ and Eu³⁺ peaks were recorded under 368 and 394 nm excitations, respectively, in accordance with the excitation spectra shown in figure 4(a).

The effects of Eu content on photoluminescence of the Y_0.97−xTb_0.03Eu_x ternary LRHs are shown in figure 5 for the two typical excitation wavelengths of 260 (figure 5(a)) and 284 nm (figure 5(b)). Irrespective of the x value, strong Eu³⁺/weak Tb³⁺ and strong Tb³⁺/weak Eu³⁺ peak combinations were recorded under 260 and 284 nm excitations, respectively, conforming to the results of figure 4(a). The Tb³⁺ emission intensity decreases while that of Eu³⁺ increased with Eu³⁺ content (figure 5(c)), implying the occurrence of energy transfer from Tb³⁺ to Eu³⁺. The energy transfer rate increased with Eu³⁺ content, thereby affecting the emission intensities of both the activators. Figure 5(c) suggests that the energy transfer is more efficient under 284 nm excitation that corresponds to the 4f⁸–4f⁷5d¹ transition of Tb³⁺. Tunable emission colors can thus be obtained with the Tb³⁺ and Eu³⁺ codoped LRHs.

Figure 5.

Emission spectra of the Y_0.97−xTb_0.03Eu_x LRHs under 260 nm (a) and 284 nm (b) excitation. Panel (c) shows relative intensities of the 615 nm Eu³⁺ and 545 nm Tb³⁺ emissions, as a function of the Eu³⁺ content (the x value).

3.2. Characterization and optical properties of the (Y_0.97−xTb_0.03Eu_x)₂O₃ solid solutions

Tb³⁺ ions readily oxidize at elevated temperatures to form Tb⁴⁺, and thus flowing NH₃ (300 ml min⁻¹) was used as a protective/reducing atmosphere upon LRH calcination. The Tb³⁺-bearing oxides processed up to 900 °C were light brown, suggesting the presence of Tb⁴⁺, while those calcined at 1100 °C were white. XRD analysis (figure 6) showed that the 1100 °C samples had high crystallinity and exhibited almost all the diffractions corresponding to the cubic Y₂O₃ (space group: , JCPDS no. 89-5591). The d spacing and lattice constant calculated from the 222 diffraction become steadily larger with increasing Eu³⁺ addition (table S3 see the supplementary material available from stacks.iop.org/STAM/14/015006/mmedia), in accordance with the observed gradual shift of the diffraction peak towards the low angles. This result confirms the successful incorporation of larger Eu³⁺ and Tb³⁺ ions into the Y₂O₃ lattice to form solid solutions.

Figure 6.

Powder XRD patterns of the (Y_0.97−xTb_0.03Eu_x)₂O₃ samples calcined from their corresponding LRHs at 1100 °C for 4 h under flowing NH₃.

FE-SEM images revealed that the oxides largely retain the thin-plate-like shapes of the original LRH crystals, though some plates have partially collapsed due to the strain arising from thermal decomposition (figures 7(a)–(e)). HR-TEM analysis of the lateral surface of an isolated nanoplate clearly resolved lattice fringes with a spacing of ∼0.307 nm that corresponds to the 222 crystal plane of the oxide (figure 7(f)). This complies with the findings of Sasaki and co-workers [32], which showed that the host layer (the ab plane) of LRH would quasitopotactically transform into the 222 plane of the oxide owing to the similar atomic configurations of the two planes.

Figure 7.

FE-SEM (a)–(e) and HR-TEM (f) images of the (Y_0.97−xTb_0.03Eu_x)₂O₃ solid solutions calcined under flowing NH₃ at 1100 °C for 4 h. Panels (a)–(e) correspond to x =0, 0.005, 0.01, 0.02 and 0.03, respectively; the same sample was used for panels (b) and (f).

Figure 8 shows excitation spectra of the Y_0.965Tb_0.03Eu_0.005 ternary and the Y_0.995Eu_0.005 and Y_0.97Tb_0.03 binary solid solution oxides obtained by monitoring the red emission of Eu³⁺ at 613 nm and the green emission of Tb³⁺ at 543 nm. The Y_0.995Eu_0.005 sample (spectrum c) shows a sharp and strong CT band at ∼250 nm, together with a weak host excitation band at ∼212 nm, as widely observed for Eu³⁺-doped Y₂O₃ red phosphors [5]. The Y_0.97Tb_0.03 sample exhibits a broad and strong excitation band in the range of ∼225–320 nm, corresponding to the well-documented 4f⁸–4f⁷5d¹ Tb³⁺ transition (spectrum d) [35]. The doublet at ∼278 and 303 nm arises from the low-spin and high-spin inter-configurational f–d transitions of Tb³⁺ electrons, respectively [26, 50]. Unlike their LRH counterparts, the oxides show negligible excitations at longer wavelengths, indicating that the most efficient Eu³⁺ and Tb³⁺ emissions can be obtained through exciting the CT and 4f⁸–4f⁷5d¹ transition bands, respectively. Monitoring the 543 nm green emission of Tb³⁺ in the Y_0.965Tb_0.03Eu_0.005 ternary sample yielded a PLE spectrum (curve (b)) almost identical to that of the Y_0.97Tb_0.03 (curve (d)). Monitoring the 613 nm Eu³⁺ emission of the same sample, however, produced both the f–d transition band of Tb³⁺ and the CT band of Eu³⁺ in the PLE spectrum (curve (a)), with the former even significantly stronger than the latter. This provides persuasive evidence of an efficient energy transfer from Tb³⁺ to Eu³⁺, and predicts that stronger Eu³⁺ luminescence can be achieved by exciting the Tb³⁺ ions at 278 or 303 nm rather than through the CT band at ∼250 nm.

Figure 8.

Luminescence excitation spectra of the Y_0.965Tb_0.03Eu_0.005 (a, b), Y_0.995Eu_0.005 (c) and Y_0.97Tb_0.03 (d) oxides recorded by monitoring the 613 nm Eu³⁺ emission (a, c) and 543 nm Tb³⁺ emission (b, d).

Figure 9 shows the effects of excitation wavelength (249–278 nm) on the luminescence of Y_0.965Tb_0.03Eu_0.005 oxide. The PL spectra exhibit similar features irrespective of the excitation wavelength and are composed of two major emission bands at ∼595 (⁵D₀ → ⁷F₁) and 615 nm (⁵D₀ → ⁷F₂) for Eu³⁺ and at ∼485 (⁵D₄ → ⁷F₄) and 543 nm (⁵D₄ → ⁷F₃) for Tb³⁺. The Tb³⁺ and Eu³⁺ activators can thus be simultaneously excited by UV light. The emission intensity of both Tb³⁺ and Eu³⁺ increases with increasing excitation wavelength. The increase for Tb³⁺ is readily understandable, since the excitation becomes closer to the center of the strong f–d transition band of Tb³⁺ (figures 8(a) and (d)). The enhancement of Eu³⁺ emission, despite moving the excitation away from the center of the CT band, is clearly due to the energy transfer from Tb³⁺ to Eu³⁺. Plotting the intensity ratio of the green to red emission (I₅₄₃/I₆₁₅) against the excitation wavelength shows that, though the red emission is dominant in each case, green emission gradually gains intensity at longer excitation wavelengths up to 278 nm (figure 9, inset). Color-tunable emission can thus be obtained via selective excitation of the Y_0.965Tb_0.03Eu_0.005 sample. When excited at 249, 260, 265, 270 and 278 nm, the sample has CIE chromaticity coordinates of (0.48,0.39), (0.47,0.42), (0.46,0.42), (0.46,0.43) and (0.46,0.43), respectively. All these emission colors fall into the yellow to greenish-yellow region of the CIE chromaticity diagram but with different proportions of green.

Figure 9.

Emission spectra of the Y_0.965Tb_0.03Eu_0.005 solid solution oxide under the various excitation wavelengths indicated in the figure. Inset shows the intensity ratio of the green to red emission (I₅₄₃/I₆₁₃) as a function of excitation wavelength.

The effects of Eu content (the x value) on photoluminescence of the Y_0.97−xTb_0.03Eu_x oxides were investigated by exciting the 4f⁸–4f⁷5d¹ transition of Tb³⁺ at 278 nm, and the results are shown in figure 10(a) with the Y_0.995Eu_0.005 spectrum included for comparison. The Y_0.97Tb_0.03 sample (x = 0) exhibits the typical emission of Tb³⁺ in Y₂O₃ and is characterized by strong bands at ∼543 nm (⁵D₄ → ⁷F₃ transition, green emission) and ∼485 nm (⁵D₄ → ⁷F₄ transition, blue emission). The Y_0.995Eu_0.005 binary sample only produces weak Eu³⁺ emissions because the excitation wavelength of 278 nm falls on the tail of the Eu³⁺ CT band (figure 8, curve (c)), as mentioned above. In the presence of only 0.5 at.% Eu³⁺ (Y_0.965Tb_0.03Eu_0.005), the intensity of the 543 nm green emission of Tb³⁺ was drastically reduced to ∼21% of the Eu-free sample. Meanwhile the 613 nm Eu³⁺ emission increased ∼seven-fold compared to the Y_0.995Eu_0.005 oxide, indicating a very efficient energy transfer from Tb³⁺ to Eu³⁺. The Tb³⁺ signals gradually weakened with increasing x whereas the Eu³⁺ emission showed a maximum at x = 0.01.

Figure 10.

(a) Emission spectra of the Y_0.97−xTb_0.03Eu_x ternary oxides under 278 nm excitation for x = 0, 0.005, 0.01, 0.02 and 0.03 (curves (a)–(e), respectively). Curve (f) corresponds to the Y_0.995Eu_0.005 oxide and is included for comparison. The right part shows luminescence of the x = 0 and 0.10 samples under 254 nm excitation from a hand-held UV lamp. (b) Intensities of the 543 and 613 nm emissions normalized to 10, and the I₆₁₃/I₅₄₃ intensity ratio versus Eu content (the x value).

Figure 10(b) shows the relative intensities of luminescence and the red-to-green intensity ratio (I₆₁₃/I₅₄₃) for the Y_0.97−xTb_0.03Eu_x series. The maximum I₆₁₃/I₅₄₃ value of ∼8.6 is observed for x = 0.02. The Tb³⁺ to Eu³⁺ energy transfer and the varied emission intensities of both the activators conform well to the PLE spectra obtained by monitoring the 613 and 545 nm emissions (figure S7 see the supplementary material available from stacks.iop.org/STAM/14/015006/mmedia). The optimal activator concentration in Y₂O₃ is ∼5 at.% for either Tb³⁺ or Eu³⁺ [35, 36], and this value would also hold for the Tb³⁺/Eu³⁺ codoped system. At the maximum Eu³⁺ emission intensity, the total activator concentration C_Tb+Eu of 4 at.% is indeed close to 5 at.%. Overdoping leads to the concentration quenching of luminescence via two mechanisms: (i) shortened separation among the neighboring activators, if a uniform dopant distribution was assumed, and (ii) localized dopant distribution (dopant clustering) owing to the lattice strain induced by surplus incorporation of larger Tb³⁺ and especially Eu³⁺ ions. Concentration quenching of luminescence was not observed for the LRHs in this and previous studies [28, 29, 32] because in the LRH structure the activators reside at the centers of the [Ln(OH)₇(H₂O)] and [Ln(OH)₈(H₂O)] polyhedrons and are well separated from each other. The maximum I₆₁₃/I₅₄₃ ratio at x = 0.02 (C_Tb+Eu = 5 at.%) observed in figure 10(b) might result from an interplay between the efficiency of Tb³⁺ to Eu³⁺ energy transfer and the relative extent of concentration quenching between Tb³⁺ and Eu³⁺ emissions. In the absence of concentration quenching, the efficiency of Tb³⁺ to Eu³ energy transfer (η _ET ) can be calculated from the fluorescence intensity using the formula [51–53]

where I_S and I_S0 are the integrated intensities of Tb³⁺ emission in the presence and in the absence of Eu³⁺, respectively. Using the experiment results of figure 10(a) and equation (1) the η _ET values were estimated at ∼86% for x = 0.005 and 91% for x = 0.01. These high efficiencies of energy transfer primarily originate from the significant spectral overlap between the ⁵D₄ → ⁷F_j emission of Tb³⁺ and the ⁷F_0,1 → ⁵D_0,1,2 absorption of Eu³⁺ [54]. The Tb³⁺ to Eu³⁺ energy transfer may occur via exchange interaction and electric multipole interaction; it depends on the average distance (R) between the Tb³⁺ donor and Eu³⁺ acceptor ions. Exchange interaction generally requires an overlap of the donor and acceptor orbitals and an R value of less than 0.3–0.4 nm; otherwise, the electric multipole interaction may dominate [55]. The average separation R can be estimated from the following equation proposed by Blasse and Bril [56, 57]

where C_Tb+Eu is the total concentration of Tb³⁺ and Eu³⁺ ions, N is the number of sites that the rare-earth ions can occupy per unit cell and V is the cell volume. Cubic Y₂O₃ has 80 atoms per unit cell, among which 32 are cations, and therefore N = 32 [5, 58]. With the cell volume shown in table S3, calculations according to equation (2) yield R = 1.328, 1.262, 1.209, 1.125 and 1.061 nm for C_Tb+Eu = 0.03 , 0.035, 0.04, 0.05 and 0.06, respectively. All these values are significantly larger than the 0.3–0.4 nm required for the exchange interaction mechanism, implying that the Tb³⁺ to Eu³⁺ energy transfer largely takes place via electric multipole interactions.

The oxide nanoplates exhibit vivid luminescence colors under 254 nm UV excitation, as shown in figure S8 for all the compositions and in the right part of figure 10(a) for two typical ones. Color-tunable photoluminescence was thus achieved with the Y_0.97−xTb_0.03Eu_x phosphors by varying the Eu content, and the CIE chromaticity coordinates (x,y) of the emissions were (0.34,0.53) for x = 0, (0.46,0.43) for x = 0.005, (0.48,0.41) for x = 0.01, (0.50,0.39) for x = 0.02, and (0.49,0.39) for x = 0.03, roughly corresponding to yellowish-green, yellow, yellowish-orange, yellowish-orange and yellowish-white colors, respectively.

Fluorescence kinetics were studied for the ⁵D₀ → ⁷F₂ emission of Eu³⁺ at 613 nm and the ⁵D₄ → ⁷F₅ transition of Tb³⁺ at 543 nm. The decay curves (figure S9) can be fitted with the single-exponential function of I = A exp(-t/τ_R) + B, where τ_R is the fluorescence lifetime, t is delay time, I is relative intensity and A and B are constants, and the derived fluorescence lifetimes are summarized in table 1. The lifetime of Eu³⁺ emission slowly yet continuously decreases with increasing Eu content, which might be understood from the following two aspects. Raising the Eu concentration shortens the separation among the luminescent centers, which increases the probability of non-radiative energy transfer among them via electric multipole interactions. It also increases the probability of radiative (resonant) energy transfer among the Eu³⁺ ions followed by non-radiative transfer to the surface sites. Both mechanisms shorten the fluorescence lifetime. The fluorescence lifetime of Tb³⁺ emission rapidly decreases from 3.03 ms for x = 0 to 2.10 ms for x = 0.005 and then to 1.89 ms for x = 0.01 due to the energy transfer from Tb³⁺ to Eu³⁺, which is frequently observed for donor/acceptor codoped systems [54, 59]. Beside the Tb³⁺ to Eu³⁺ energy transfer, other mechanisms such as defect clustering due to lattice strain (mechanism (2) mentioned above for Eu³⁺) may contribute to the lifetime shortening at larger x values.

Table 1.

Fluorescence lifetimes of the 543 nm Tb³⁺ and 613 nm Eu³⁺ emissions for the (Y_0.97−xTb_0.03Eu_x)₂O₃ nanoplates, τ_R (Tb³⁺) and τ_R (Eu³⁺), respectively.

x	0	0.005	0.01	0.02	0.03
τR (Tb3+) ms	3.03±0.01	2.10±0.01	1.89±0.01	1.47±0.01	1.44±0.01
τR (Eu3+) ms	–	2.56±0.01	2.33±0.01	2.27±0.01	2.04±0.01

4. Conclusions

Multicolor phosphor nanoplates of the Y/Tb/Eu system were produced by optimized hydrothermal synthesis of (Y_0.97−xTb_0.03Eu_x)₂(OH)₅NO₃·nH₂O LRHs and (Y_0.97−xTb_0.03Eu_x)₂O₃ solid solutions (x = 0–0.03). Detailed characterization by XRD, FTIR, FE-SEM, TEM and optical spectroscopies yielded the following conclusions:

• 1.Phase-pure LRH can be obtained at pH ∼7.0 and 100 °C, whereas mixtures of high-hydration and low-hydration phases were produced at higher temperatures up to 200 °C.
• 2.Irrespective of the hydration extent, the LRHs are solid solutions with free NO₃⁻ anions intercalated between the layers. The incorporation of relatively large Eu³⁺ and Tb³⁺ ions expands the ab planes while reducing the hydration degree and the interlayer distance.
• 3.Characteristic Tb³⁺ and Eu³⁺ emission lines are observed from ternary LRHs when selectively exciting the two types of activators, and the occurrence of energy transfer from Tb³⁺ to Eu³⁺ was confirmed.
• 4.The (Y_0.97−xTb_0.03Eu_x)₂O₃ solid-solution nanoplates annealed from the LRHs at 1100 °C possess exposed 222 facets. The color and intensity of luminescence can be tuned by varying the excitation wavelength from ∼249 nm (the charge transfer excitation band of Eu³⁺) to 278 nm (the 4f⁸–4f⁷5d¹ transition of Tb³⁺).
• 5.Efficient Tb³⁺ to Eu³⁺ energy transfer was observed for the (Y_0.97−xTb_0.03Eu_x)₂O₃ phosphor, allowing one to tune the emission color from approximately green to yellowish-orange by varying the Eu³⁺/Tb³⁺ atomic ratio.
• 6.At the optimal Eu³⁺ content of x = 0.01 for (Y_0.97−xTb_0.03Eu_x)₂O₃, the efficiency of energy transfer was ∼91%, and the transfer was attributed to electric multipole interactions.