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Abstract
Nanocomposites of LiZrO2 – LiBaZrO3: xMn2+ (x = 0–0.06 molar ratios) were prepared by the co-precipitation method. The prepared samples were characterized by X-ray diffraction (XRD), UV-visible spectrometry, photoluminescence (PL), scanning electron microscopy (SEM) and energy dispersive X-ray spectrometry (EDX) techniques. Analysis of XRD data shows two phases: the cubic phase of BaZrO3 perovskite and the tetragonal phase of ZrO2. Mn2+ ions were predominantly distributed in the tetrahedral sites and a few ions are situated at the octahedral sites of the composites. UV-visible spectroscopy of these samples presents two optical band energies, decreasing exponentially with increasing concentration of the Mn2+ ion. PL spectra of Mn2+-doped samples display three broad emission bands: a band centered at wavelength,λ =416 nm (blue), another with peak maximum atλ =527 nm (green) and the third (with relatively the lowest intensity) at about 600 nm (orange-red). The blue (λ =416 nm) band was dominant at low Mn2+ concentrations (x≤0.03) but the green band (λ =527 nm) became dominant at higher Mn2+ concentrations (x>0.03). The CIE coordinates revealed colour changes from blue to green at a concentration of 0.05 mole ratio and white light at 0.06 mole ratio. The phosphor presented in this work is a promising material for use in display devices such as flat panel displays, colour plasma displays, signage lights and backlights.
© 2020 Optical Society of America under the terms of the OSA Open Access Publishing Agreement
1. Introduction
Globally, energy crisis and environmental based issues have become the major problems in the 21st century. These problems are directly linked to high energy consumption from conventional lighting technologies, as they are highly inefficient; besides they pose a health risk to end users [1]. Significant reduction in energy wastage through efficient lighting technologies is crucial for human development and health [1–3]. White light-emitting diodes (W-LEDs) have emerged to replace conventional lighting sources, such as incandescent light bulbs and fluorescent tubes [4].
The focus of LED development in recent years has been to create efficient and inexpensive lighting phosphors for the production of white-LEDs for general illumination. Phosphor usually comprises of a host crystal material and one or more intentionally introduced impurities, called activators. The best choice of activators in phosphors is the ions with f-d or d-d electron configurations as they could emit visible and broad-band light under the influence of crystal field [5].
An example of transition elements (d – d) activators is the manganese ion (Mn2+). When this ion is incorporated into a host matrix, it emits broad band radiation with maximum emission peak depending on the valence state and site symmetry surrounding it. In the weak crystal field (tetrahedral), Mn2+ emits in the green-yellow spectrum, whereas in the strong field (octahedral), Mn2+ ion emits orange to red light [6,7]. Therefore, Mn2+ ion could be made to emit green, yellow, orange, red and even white light by carefully selecting a suitable host matrix. These multiple colour emissions have been reported for Mn2+ doped into ZnGa2O4, Zn3Ga2GeO10 and Zn1-xGa2-2xGexO4 host matrices comprising both tetrahedral and octahedral cation sites [8,9].
In this work, Mn2+ ion has been intentionally doped into a newly formulated host matrix, the LiZrO2-LiBaZrO3 nanocomposite. Zirconia (ZrO2) nanocrystals and BaZrO3 perovskite nanocrystals were selected to form the composite because both possess optical transparency in the visible region and have good thermal properties. Besides, ZrO2 is the 20th most abundant compound on the earth (more common than zinc, tin and mercury) [10]. Because of its relative abundance, phosphors made from zirconia would be cost-effective, thus display devices, signage lighting and backlight produced from these phosphors would become affordable. Perovskite, on the other hand, has shown remarkable enhanced emission properties when doped with rare-earth or transition ions. However, research reports on Mn2+-doped ZrO2 have never been shown to produce green emission but rather red and blue emissions [11–14]. Unfortunately, there is no report on Mn2+-doped BaZrO3 in the literature at least to our knowledge. Nevertheless, reports on Mn2+ doped into similar perovskites such as SrTiO3 and SrZO3 produced only blue and red luminescence [15,16] since manganese is generally known to substitute B-sites in ABO3 perovskites and at the +4 valence state [16] which produces essentially red emission.
The LiZrO2-LiBaZrO3 nanocomposite provides two cation sites for the Mn2+; a tetrahedral site in LiZrO2, and octahedral sites in the LiBaZrO3 perovskite, making it capable of multiple emissions from the single nanocomposite phosphor. Notwithstanding the probability of multiple emissions from the doped nanocomposites, the component phases may produce emissions that would complement the dopant ion emission. This will facilitate tuning of emissions from the doped nanocomposite by producing light colours that are not possible from a doped single-phase material that forms the composite.
2. Experiment
All reagents used for the synthesis of LiZrO2-LiBaZrO3: xMn nanocomposites (NCs) were of analytical grade and were used without further purification. These reagents include manganese (II) nitrate tetrahydrate (Mn(NO3)2.4H2O, 98.5%), barium nitrate (Ba(NO3)2, 99.0%), lithium acetate dihydrate (C2H3LiO2.2H2O, 99.999%), zirconium (IV) oxychloride (ZrOCl2.8H2O, 97%), ammonia solution (NH3) and absolute ethanol (C2H5OH, 99.7%).
2.1 Synthesis and characterization of the LiZrO2-LiBaZrO3: xMn2+ nanocomposite
Undoped LiZrO2-LiBaZrO3 nanocomposite was prepared using a co-precipitate method. All the chemicals were dissolved in ethanol/de-ionized water (DW) mixture in the ratio 3:7 ethanol:DW. A mixture of 0.1 M each of aqueous solutions of (Ba(NO3)2, C2H3LiO2.2H2O and ZrOCl2.8H2O, was stirred until a clear solution was obtained. While still stirring, 1.0 M aqueous solution of NH3 was added, dropwise, to the above mixture and the entire content was refluxed for 2 hours at a temperature of 80 °C to form a white precipitate. The white precipitate was repeatedly washed with deionized water and ethanol using a centrifuge until a neutral pH was achieved. The precipitate was dried for 12 hours at a temperature of 110 °C and allowed to cool to room temperature. Finally, the white (crystalline) solid sample crushed into a fine powder and then calcined at 800°C for 2 hours. The preparation of LiZrO2-LiBaZrO3: xMn2+ followed the same procedure used to prepare the undoped sample, except that predetermined molar ratios (x = 0.01, 0.02, 0.03, 0.04, 0.05 & 0.06) of Mn2+ from Mn(NO3)2.4H2O were added to the above mixture. The possible ionic reaction equation is presented in Eq. (1);
Different methods were employed to characterize the synthesized nanocomposites. These techniques include Zeiss Auriga field emission scanning electron microscope (FESEM) equipped with energy dispersive x-ray (EDX) for the determination of morphology and elemental compositions of samples, powder X-ray diffractograms were obtained using Bruker AXS D8 Advance powder x-ray diffractometer with Cu $K\alpha$ radiation, Perkin-ElmerLambda 1050 UV/Vis/NIR spectrophotometer equipped with an integrating sphere and an F-7000 spectrophotometer were employed for diffuse reflectance and photoluminescence measurements at room temperature, respectively.
3. Results and discussion
3.1 Morphology
Field emission scanning electron micrographs of LiZrO2-LiBaZrO3: xMn2+ nanocomposites (x = 0.04 & 0.06 mol ratios) presented in Figs. 1(a) and 1(b) show agglomerated particles with irregular size distributions. There is no clear distinction between the two images. The energy dispersive x-ray spectra in Figs. 1(c) and 1(d) show the presence of Ba, Zr, O, and Mn in weight percent. Li+ is not observed in any of the spectra because of its low atomic mass [17,18]. The EDX element mapping displayed in Figs. 1(e)–1(h) and Fig. 1(d) (inset) show a uniform distribution of host and dopant elements over grains and grain boundaries in the nanocomposite with no apparent enrichment of the doping element found at grain boundaries.
Fig. 1. FESEM micrographs of x = 0.04 (a) and x = 0.06 (b), and their corresponding EDX spectra (c) and (d), respectively of LiZrO2 -LiBaZrO3: xMn2+, and Elemental Mapping (e)-(h) of LiZrO2-LiBaZrO3:0.06Mn2+ nanocomposites.
3.3 Crystal structure and size
Figure 2(a) shows representative XRD patterns of LiZrO2-LiBaZrO3: xMn2+ nanocomposites (x = 0, 0.01 & 0.06 mol ratios). The presentation of only three patterns is for clarity of the peaks and because samples gave similar diffraction patterns. The XRD patterns possess two phases: the cubic phase of bulk BaZrO3 and the tetragonal phase of ZrO2. The peaks at $2\mathrm{\theta }$ values 37.47, 43.48, 53.90 and 71.5° correspond to the (111), (200), (211) and (310) reflections, respectively of mainly cubic BaZrO3 perovskite (JCPDS card No.06-0399); as well as the $2\mathrm{\theta }$ peaks values at 34.59, 35.363, 50.35, 51.43, 59.22, 60.23, 72.52 and 74.66° are assigned to the (002), (110), (112), (200), (103), (211), (004) and (220) planes, respectively mainly of bulk tetragonal- ZrO2 (JCPDS card No. 81-1545). . The refined lattice parameters presented in Fig. 2(b) for x = 0- 0.06, are also in agreement (qualitatively) with those reported in the above JCPDS files, except for some shift in the $2\mathrm{\theta }$ peak to lower angles as a result of substituting host cations with Mn2+ ions.
Fig. 2. (a) XRD patterns of LiZrO2-LiBaZrO3: xMn2+ nanocomposite and (b) lattice parameters at varying concentrations of Mn2+ ion.
To elucidate this crystal structure hypothesis, these XRD data were refined by the Rietveld method using the Fullprof software [19]. The analysis was applied to doped (x = 0.01–0.06) and undoped (x = 0) samples. Calculated XRD patterns were obtained using a $Fm\bar{3}m$ space group for c-BaZrO3 cubic and a $P{4_2}/nmc$ space group for the t-ZrO2 tetragonal phase of NCs. The Li+/Mn2+ ions were considered in six-fold and eight-fold oxygen coordination during the Rietveld fit, for cubic and tetragonal symmetries respectively. Figure 3 shows the profile fitting of the XRD patterns with lower (x=0, x=0.01) and higher (x=0.05, x=0.06) Mn2+ ion concentrations annealed at 800°C. From x = 0.01 to x=0.05 compositions, the crystal structure of Li+/Mn2+ co-doped c-BaZrO3 and Li+/Mn2+ co-doped t-ZrO2 is stable, after these composites a secondary phase of LiMnO4 is induced [* in the XRD pattern Fig. 3(d)]. In this case, the Mn2+/Li+ ions in the 1a site position for cubic symmetry could reach a maximum site concentration in x=0.05, where the ions cannot be moved through a diffusion path, therefore these ions in the 1asite position stop the diffusion of more ions as a consequence of the reduction in its mobility capacity inducing thus the formation of a small amount of LiMnO4. The quantitative analysis and crystal structure results of c-BaZrO3 and t-ZrO2 phases for x = 0, 0.01, 0.05, 0.06 annealed at 800°C are presented in Table 1. From the results and their goodness of fit, it is clear that the Mn2+ ion reaches a maximum concentration into the 6-fold and 8-fold sites occupancy of cubic and tetragonal symmetries. From Rietveld results a maximum of x=0.05 Mn2+ions occupying both sites could be permitted. Although in general, the unit cell volume increases with Mn2+ions concentration for both cubic and tetragonal phases, there is a slight volume decrease in the 0.05 mole ratio composition compared to x=0.6. This could be due to the orthorhombic-LiMnO4 secondary phase formation in the x=0.06 composition. Here, an organized arrangement of Li+/Mn2+ ions into the BaZrO3 cubic and ZrO2 tetragonal phases occurs; where, during the crystallization of LiMnO4 structure there is a competition in the nucleation and growth of these phases, inducing thus, a not systematic change in the lattice parameters for x=0.6 composition, hence, increasing their unit cell volume. The above results are confirmed by the quantitative phase analysis, where the tetragonal phase concentration increases from 65% (x=0) to 75% (x=0.05), after these compositions, the tetragonal phase concentration is reduced. Tetragonal ZrO2 has a face-centered cell, where Zr4+ ions occupy sites with D2d lower symmetry or D4h symmetry [20]. The structure has a combination of two distorted tetrahedral with zirconium surrounded by 8-oxygen atoms (coordination number, CN = 8); a flattened tetrahedron and an elongated tetrahedron rotated 90° relative to the former [21], with an average Zr – O bond length of 2.22${\pm}$0.02$\AA\; $[22]. On the other hand, Li-doped BaZrO3 form an ABO3 cubic perovskite structure with Ba2+ at each of the 12 –coordinated cavities of the cuboctahedra framework [23–25]. The Zr4+ occupies the center of the octahedral interstices formed by 6-oxygen ions seated at the face center. In this cubic perovskite structure, both the Ba2+ and Zr2+ sites have the Oh local point group symmetry, while the oxygen atom at the face centers has a local D4h group point symmetry [26,27].
Fig. 3. Rietveld refinement of LiZrO2-LiBaZrO3: xMn2+ compositions (x = 0 (a), x=0.01 (b), x=0.05 (c), x=0.06 (d) mol ratio); the experimental, calculated and difference profiles are shown.
Table 1. Crystal structure results and quantitative phase analysis by Rietveld methoda
Since the synthesized LiZrO2-LiBaZrO3 is a composite of tetragonal Li-doped ZrO2 and cubic Li-doped BaZrO3 perovskite, there is the need to determine the distribution of Mn2+ in the two phases. To achieve this, we first consider the site occupancy of Mn2+ in the LiZrO2-LiBaZrO3nanocomposite using the radius percent difference (Rr) between the dopant cation (Mn2+) and host cations (Li+, Ba2+, $\textrm{Zr}_t^{4 + }$ and $\textrm{Zr}_c^{4 + }$) [28]:
where CN is the coordination number, Rh(CN) is the radius of the host cation, and fRd(CN) is the radius of the dopant ion, $\textrm{Zr}_t^{4 + }$ and $\textrm{Zr}_c^{4 + }$ are zirconium ions in LiZrO2 (CN = 8) and cubic LiBaZrO3 (CN = 6) phases, respectively. The acceptable percentage difference in ionic radii between dopant ions and host (substituted) ions must not exceed 30% [28]. For the LiZrO2 phase, the sizes of Mn2+ and $\textrm{Zr}_t^{4 + }$ in 8-coordinations are 0.96Å and 0.84Å, respectively. The ionic radius percentage difference for Mn2+ substituting $\textrm{Zr}_t^{4 + }$ in LiZrO2 is -14.3%. The incorporation of a larger cation at a small cation site is expected to increase the cell constant values, $a$ and c (enlargement of unit cell volume) [25,29]. The values for cell volumes of tetragonal LiZrO2 show increment for all Mn2+ doped nanocomposites (Fig. 4) when compared to the undoped nanocomposite, suggesting that some Mn2+ ions are residing on the $\textrm{Zr}_t^{4 + }$ sites.Fig. 4. (a) Cell volume, and (b) structural occupancy & average crystallite size of LiZrO2-LiBaZrO3: xMn2+ nanocomposites.
Similarly, the ionic radius percentage difference for Mn2+ (CN = 12, RMn = 1.27Å) [30] ions substituting Ba2+ (CN = 12, ${R_{Ba}}$ = 1.61Å) and $\textrm{Zr}_c^{4 + }$ (CN = 6, RZr = 0.83Å) sites in the LiBaZrO3 phase are 21.6% and -15.3%, respectively. These results indicate that Mn2+ will preferentially occupy the $\textrm{Zr}_c^{4 + }$ site in the octahedron than the Ba2+ site in the cuboctahedron of the perovskite structure. According to Pradhan and Roy, compositions with stoichiometry that are likely to produce perovskite with B-site vacancies would rather crystallize as a mixture of perovskite and none-perovskite phases [31]. Though, some reports have suggested that cations with different valence charges could substitute the B-cations in the octahedral with the evolution of vacancies and without necessarily producing a second phase [25,29,32]. Nevertheless, this may be a reason (in addition to the synthesis technique and temperature) for composite formation in this work. Manganese has oxidation states which include Mn2+ and Mn4+ which might isovalently occupy position A and/or B in the perovskite structure [30]. X-ray diffraction pattern at higher Mn2+ ions concentration (x = 0.06) shows the phase belonging to LiMnO2, suggesting phase segregation and oxidation of Mn2+ to Mn4+ (Fig. 3). This assertion is consistent with the report by Berlin et al. [13] that increasing Mn2+ concentration increases the oxidation of Mn2+ to higher oxidation states (Mn3+ and Mn4+).However, for this work and for reasons that may become apparent in photoluminescence analysis, we consider only the possibility of Mn2+ions occupying either or both of the two sites (Ba2+ and $\textrm{Zr}_c^{4 + }$ sites) provided by the LiBaZrO3 perovskite structure. On the other hand, the substitution of Li+ at both Ba2+ and Zr4+ sites may be less likely because of the requirement for large charge compensation compared to Mn2+ ion substitution. The question therefore is, will Li+ occupy interstitial sites in the LiBaZrO3 structure? It is reported that the high packing density in perovskite does not support the presence of interstitial ions [31], however, substitutional concentrations of A-site vacancies can be accommodated. Therefore, Li+ ions will preferentially occupy grain boundaries as well as the surface in both phases than in the bulk because of its low solubility limit in small size polycrystals. However, some small amount of Li+ ions may be doped into the bulk lattice, where they serve as flux. On the other hand, a comprehensive explanation on Li+-doped ZrO2 nanocrystals is given in the reported by Liu et al. [26].
Perovskites have a unique structural characteristic that enables them to accommodate smaller ions at the A-site. This is termed ‘Goldschmidt’s tolerance factor’ t, which also determine the stability of the perovskite and is dependent on the relative ionic radii of A-site cation, (${r_A}$), B-site cation (${r_B}$) and the oxygen ion (${r_0}$), expressed as [33];
Therefore, the tolerance factor can be regarded as a measure of the degree of distortion of perovskite from the ideal cubic structure with $0.9 < t < 1.0$[34]. The distortion to cubic structure occurs if $0.71 < t < 0.9$[34], that is, when ${r_A}$ is smaller or ${r_B}$ is larger than the ideal ion size. Smaller ${r_A}$ values mean t is lesser than 1 ($0.71 < t < 0.9$), causing the BO6 octahedron to tilt to fill the space provided by A-site cation. Tilting of the octahedral could also reduce the coordination number of A-cation to as low as 8 [25,35], consequently reducing the symmetry of the crystal structure. The tolerance factor of ABO3 perovskite with Ba2+ substituted at A-site and $\textrm{Zr}_c^{4 + }$ at B-site (i.e. BaZrO3) is 0.999, allowing Ba2+ to be seated stably at A-site and forming an undistorted cubic structure. However, when Mn2+ (CN = 12, ${\textrm{r}_{\textrm{Mn}}}$ = 1.27Å) substitutes Ba2+ at A-site, t = 0.890, with the result that Mn2+ could create more distortion in the perovskite by leaving more space at A-site. However, the distortion caused by the substitution of smaller size Mn2+ ions at A -sites are partially compensated for by the substitution of a larger size Mn2+ ion at the $\textrm{Zr}_c^{4 + }$-site. A lesser tilting and rotation of ZrO6 octahedral takes place to effectively accommodate the smaller ions at A-sites. At the same time, the coordination number and symmetry at A-site may be slightly lowered. For this reason, the 2θ peak shift is largely determined by the relative substitution of Mn2+ at $\textrm{Zr}_t^{4 + }$ sites in the LiZrO2 phase. This may be the reason why the 2θ peak shift is toward smaller angles for all Mn2+ concentrations in the nanocomposite. So far, we have been able to establish that Mn2+ would preferentially substitute Zr4+ in LiZrO2 structure and Ba2+ in the LiBaZrO3 perovskite structure of LiZrO2 - LiBaZrO3 nanocomposite. The question is how are Mn2+ ions distributed in the two composite structures?To answer the question posed in the preceding paragraph we consider the change in unit cell volume and percentage occupancy as a function of Mn2+ concentration in the nanocomposite. Figure 4(a) shows larger cell volumes for Mn2+-doped samples in comparison to the undoped sample. This increase in cell volume is consistent with the reported volume increase when a larger ion substitutes a smaller ion in the lattice [25]. The non-monotonic increases in cell volume may be as a result of some of the Mn2+ and Li+ ions entering Ba2+ sites. Also presented in Fig. 4 is the percentage phase (tetragonal LiZrO2 and cubic LiBaZrO3) occupancy in the LiZrO2-LiBaZrO3 nanocomposite [Fig. 4(b)]. The percentage occupancy of each phase was calculated from values of unit cell volume using the (200) and (112)/(200) reflections inherent to the tetragonal LiZrO2 and cubic LiBaZrO3 phases, respectively of the nanocomposites. It can be seen from Fig. 4(b) that the values of phase occupancy are almost equal for the two phases, except at x = 0.02 and x = 0.03 mole ratios. This result suggests that Mn2+ ions are uniformly distributed in the two phases especially at low and high Mn2+ concentrations. The result is in agreement with the EDX mapping shown in Fig. 1(h), which indicates a uniform distribution of the Mn2+ ions in the sample. The result is equally in good agreement with the nearly equal values of ionic radii percentage differences obtained for both phases.
The average crystallite sizes of all samples were determined by considering the four most prominent (and symmetric) diffraction peaks and using Debye Scherer’s formula [36]:
where Dhkl is the crystallite size, $\lambda $ is the X-ray wavelength of Cu_K$\alpha $ radiation (0.15406 nm), K is the shape factor taken as 0.89 for spherical particles, ${\beta _{hkl}}$ is the experimental full-width at half maximum (FWHM in radians), and ${\theta _{hkl}}$ is the Bragg’s angle. The average crystallite size was obtained by taking the average crystallite sizes for the four peaks of each sample. The average crystallite size of undoped LiZrO2-BaZrO3 nanocomposite is 15 nm, and those of the doped samples shown in Fig. 4(b) did not display any special variation with the addition of more Mn2+ ions. It should be mentioned that pure (single-phase) metastable tetragonal-ZrO2 nanocrystals without stabilizing agents have been obtained previously by other authors [37–39]; however, the domain size of these nanocrystals should be $\le$10 nm [11,40,41] or < 14 nm [39]. From the result shown in Fig. 4(b), it is clear that the x = 0 and x = 0.06 compositions are within this threshold values, while x = 0.01 to 0.05 compositions are above this threshold values but less than 20 nm which is the threshold for tetragonal to monoclinic phase transformation [42,43]. Therefore the tetragonal symmetry obtained in this work for LiZrO2could be due to the Li+ ion for the x = 0 composition [44], and Li+/Mn2+ ions for the other Mn2+-doped compositions [37,45,46], where these ions act as stabilization agents.3.3 Diffuse reflectance spectroscopy
Diffuse reflectance measurement carried out on LiZrO2-BaZrO3: xMn2+ (x = 0, 0.01, 0.02, 0.03, 0.04, 0.05 and 0.06 molar ratios) gave the spectra shown in Fig. 5. The undoped sample displayed three absorption bands at 227 nm, 308 nm and 480 nm. For Mn2+ doped samples, these bands became broader and an additional band is situated at 511 nm. The band at 308 nm is the broadest and is red shifted to 321 nm for Mn2+-doped samples. The energy band gap of the nanocomposites was estimated from the combined Kubelka-Munk and Tauc’s-Wood relations [47–49]:
where hν is the photon energy, A is a constant related to the properties of the valence and conduction bands, and the exponent ‘n’ is ½, 2, 3/2 and 3 for allowed direct, allowed indirect, forbidden direct and forbidden indirect transitions, respectively [47,49], ${E_g}$ is the energy band gap, $F(R )$ is Kubelka-Munk (SKM) remission function which is defined as [47,48]: where R, K and S are the reflectance, absorption and the scattering coefficient, respectively. From the method described by Sangiorgi et al.[50], the electronic transition type, n was determined for selected Mn2+ concentrations; x = 0, 0.01, 0.02 and 0.06 to be 1.7, 1.3, 1.0 and 0.5, respectively. The results suggest that the indirect electronic transition translates to a direct transition after doping with Mn2+ ions. The indirect to direct band gap translation has been previously reported theoretically for BaZrO3 doped with Cd2+ by [51] and experimentally for doped SnS2 by [52]. Rizwan and co-workers [51] attributed the band gap translation to the Ba-6p and shifting of O-2p states. On the other hand, Sainbileg and Hayashi [52] ascribed the energy band transition from indirect to direct to the hybridization of d-orbitals of Ni with the 3p orbitals of S atom of SnS2 in the conduction band at M-point. Similarly, the hybridization of Mn d-orbitals with the 2p orbitals of oxygen together with the induced localized states from complex clusters may lead to O-2p shifting, hence the energy band gap translation from indirect to direct. Another reason for the band gap translation is that, since the composite in this work is made up of two structures having dissimilar electronic transition types, it is obvious that the dominant structure which in this case is the direct band gap ZrO2 (from Rietveld refinement data) will influence the nature of transition in the nanocomposite, especially when collaborated by other defect energy states. Therefore, for all doped samples, the direct transition is employed to evaluate the energy bands.Fig. 5. Normalized Diffuse Reflectance spectra of LiZrO2-LiBaZrO3: xMn2+ (x = 0 -0.06 mol ratio) nanocomposite.
Figures 6(a)-(c) show plots energy band for LiZrO2-LiBaZrO3: xMn2+ (x = 0, 0.01 and 0.02) nanocomposites. The band gap energy values were obtained from the points on the hν axis, where the extrapolated linear fit of the ${[F(R)h\nu ]^n}$ verses h$\textrm{v}$ plots (Fig. 6) gave ${[F(R)h\nu ]^n}$ = 0. The undoped sample has single band gap energy at 5.07 eV [Fig. 6(a)] for direct transition (LiZrO2), while two energy bands are observed for the indirect transition which is often related to BaZrO3. For the indirect transition, energy bands at 2.79 and 4.80 eV were obtained [Fig. 6(b)]. Yuan et al. [53] and Moreira et al. [54] had reported band gap energy for BaZO3 at 4.80 eV and 4.89 eV, respectively. Figure 6(c) is the energy plot for the Mn2+ doped sample (x = 0.02), displaying a single energy band. Figure 6(d) presents energy band variation as a function of Mn2+ ion concentrations in LiZrO2-BaZrO3 nanocomposites. The band gap energy display an exponential decrease with Mn2+ ion concentrations [Fig. 6(d)], similar to the report for Mn-doped ZnO [55]. This exponential decrease of Eg is attributed to an increase in lattice defects, which in turn increases localized states in the band gap region of disordered LiZrO2-BaZrO3 nanocomposite [56–58]. The obtained energy bands for Mn2+-doped samples 2.98 eV (416 nm), 2.77 eV (447 nm), 3.03 eV (409 nm), 2.83 eV (438 nm), 2.51 eV (498 nm) and 2.85 eV (435 nm) for x = 0.01, 0.02, 0.03, 0.04, 0.05 and 0.06, respectively correspond to emission transition states associated with oxygen vacancy defects in the nanocomposite as can be seen in the PL results (Fig. 7). The narrowing of the band gap in disordered LiZrO2-BaZrO3 nanocomposite is consistent with the report by Cavalcante et al. [59] for BaZrO3 and Berlin et al. [13] for ZrO2 that the optical band gap narrows down as the structural order decreases, and oxygen vacancies (Vo) increase [13,60]. Also, the incorporation of Li+ at any of the available sites will increase oxygen vacancies (Vo) and localized states in the band gap of the composite and further narrowing the energy band gap [61,62].
Fig. 6. Band gap energies for (a) undopeddirect transition, (b) undoped indirect transition, (c) Mn2+ doped (direct transition) and (d) energy bandsvariation of LiZrO2-LiBaZrO3: xMn2+ nanocomposites.
Fig. 7. Photoluminescence (a) excitation spectra (x = 0 & 0.01) and (b)-(h) emission spectra (x = 0 to 0.06) of LiZrO2-BaZrO3:xMn2+ nanocomposites monitored at 527 nm emission and 240 nm excitation, respectively.
3.4 Photoluminescence excitation and emission spectroscopy
Figure 7(a) shows the excitation spectra of undoped and 0.01Mn2+-doped LiZrO2-LiBaZrO3 nanocomposites, obtained by monitoring emission at 414 nm and 527 nm, respectively. The undoped nanocomposite shows three broad excitation peaks at 204 nm, 242 nm and 272 nm. On the other hand, Mn2+ doped nanocomposites display two energy peaks at 204 and 242 nm; the peak at 204 nm is associated with the band edge excitation of LiZrO2, the peak at 242 nm is assigned to the tetrahedron-localized Zr-O charge transfer band and the peak at 272 nm is related to the octahedron-localized Zr-O charge transfer band of LiBaZrO3 perovskite [63–66].
Figure 7(b) shows the emission spectrum of undoped LiZrO2-LiBaZrO3 nanocomposite with a wide asymmetric band in the range 350 to 550 nm. This band is deconvoluted into Gaussian component peaks at 404 nm (P1), 417 nm (P2), 438 nm (P3) and 472 nm (P4) which are attributed to defect state transitions (DSTs) in LiZrO2 and LiBaZrO3 structures of the nanocomposite [67]. Figures 7(c)–7(h) show photoluminescence emission spectra of Mn2+-doped LiZrO2-LiBaZrO3 nanocomposites at varying Mn2+concentrations. The PL spectra of all doped samples are similar in shape that is, a broad band in the range 400 nm to 700 nm, comprising of two distinct peaks at about 416 nm and 527 nm. The PL spectra of x = 0.01 and 0.04 Mn2+-doped samples were also deconvoluted into four Gaussian components that peaked at about 411, 449, 527 and 600 nm [Figs. 7(c) and 7(e)]. The blue peaks at 411 nm (D1) and 449 nm (D2) (which are wavelengths red-shifted) belong to both the LiZrO2 and LiBaZrO3 of the host nanocomposite [11,13]. The peaks at 527 (D3: green) and 600 nm (D4: orange-red) are mainly attributed to the 4T1 – 6A1 transition of Mn2+ ion in the host matrix [68]. These four bands (D1, D2, D3 & D4) are also present in the spectra of other Mn2+ doped samples (x = 0.02, 0.03, 0.05 & 0.06), though, with varying intensities. There were no characteristic sharp red emission bands of Mn4+ ions suggesting the LiMnO2 phase did not contribute to the overall luminescence of the nanocomposite. The broad orange-red emission band at 600 nm is often associated with Mn2+ ions at the octahedral site with Oh coordination environment [69,70]. All PL spectra were obtained by using a 240 nm excitation wavelength source.
Though the 4T1 – 6A1 transition of Mn2+ ion is partially spin forbidden, perturbation of the host crystal field caused by the distorted lattice makes this transition to be partially electric dipole allowed [27]. The energy difference between the first excited-state 4T1 and the ground-state 6A1 reduces as the crystal field increases [7]. Thus, the emission from the 4T1 – 6A1 transition wavelength is red-shifted with increasing crystal field [6,71]. The crystal field is inversely proportional to the average distance between anions and cations displaced by the Mn2+ [72]. The average bond distances of Zrc - O, Zrt - O and Ba – O are 2.099Å, 2.22Å and 2.969Å, respectively [21,73]. The Zrc - O bond distance for the octahedral coordination is shortest, which suggests that the crystal field effect at the $\textrm{Zr}_c^{4 + }$ site is stronger than at $\textrm{Zr}_t^{4 + }$ and Ba2+. Therefore, the green emission at 527 nm originates from Mn2+ situated in a weak crystal field (tetrahedral coordination) [7,9], and the orange-red emission at about 600 nm is due to Mn2+ seating at a site with a stronger crystal field (octahedral coordination) [7,9,74]. From the photoluminescence spectra shown in Figs. 7(c)–7(g), the nanocomposite phosphors emit predominantly green light, confirming that Mn2+ ions are indeed mostly positioned at $\textrm{Zr}_t^{4 + }$ site in the LiZrO2 phase of the nanocomposite. However, a fraction of the Mn2+issituated at the octahedral sites of the LiBaZrO3 of the nanocomposite in accordance with the XRD result. As Mn2+ concentration increases, the PL intensity of the blue emission band becomes weaker while that of the green emission gets stronger and reaches maximum intensity at x = 0.05 (see Fig. 8). Beyond x = 0.05, the intensity of the green emission band decreased because of concentration quenching. The increasing intensity of the green emission band with decreasing host band emission (blue emission band) demonstrates energy transfer from host to Mn2+ ion. These results are consistent with the tetragonal/cubic phase concentration in the composites obtained by the Rietveld method, wherein a high tetragonal phase concentration (75%), the green emission gets stronger and it reduces to blue emission as the cubic phase decreases as a concrescence of the Mn2+ concentration.
Fig. 8. Relative emission intensity of blue (host emission) band and green (Mn2+) band in LiZrO2- LiBaZrO3: xMn2+ nanocomposites.
3.5 Band structure and energy transfer mechanism in LiZrO2-LiBaZrO3:xMn2+ nanocomposites
The schematic energy band structure is shown in Fig. 9 for LiZrO2 and LiBaZrO3, which form the nanocomposite. The electronic band structure of LiZrO2 has a valence band which is mainly composed of occupied 2p energy state of O atom and the conduction band is constituted 4d energy state of Zr atom [11]. The blue emission originating from the LiZrO2is from defect centres associated with oxygen vacancies [11,13]. When an electron is captured on the surface of the oxygen vacancy, F centres are generated at the surface. Meanwhile, Zr3+ ions are formed (creating states below the conduction band edge) when Zr4+ ion adjacent to oxygen vacancy captures an electron. The transition from the F centres to the valence band produces the blue emission bands. The substitution of Mn2+ into the LiZrO2 lattice increases the number of oxygen vacancies. Similar to the band structure of LiZrO2, the electronic band structure of LiBaZrO3 has a conduction band which is a mixture of Ba (5d) and Zr (4d) states and a valence band which is mainly composed of 2p states. The blue emission from the host nanocomposite can also be attributed to transitions involving intermediary energy states within the energy band gap of LiBaZrO3. These localized states which are associated with the formation of zirconium clusters due to oxygen vacancies as well as Zr – O bond breaking in the BaZrO3 lattice are non-homogenously distributed in the band gap in a way that several photons can excite electrons [62,75]. According to the wide-band model and in analogous to the complex cluster formation in SrZO3[58], the [BaO11.${{\textrm{V}^{\prime\prime}}_\textrm{o}}$] and [BaO11.${\textrm{V}}_\textrm{o}^{ \bullet }$] complex clusters in LiBaZrO3 are linked to shallow defects in the bandgap from where the blue emissions emerged. On the other hand, [ZrO5.${{\textrm{V}^{\prime\prime}}_\textrm{o}}$] and [ZrO5.${\textrm{V}}_\textrm{o}^{ \bullet }$] complex clusters are ascribed to defects deeply inserted in the band gap from where the green-yellow-orange-red emissions emerged [76,77]. The energy transfer (ET) mechanism for both LiZrO2 and LiBaZrO3 in the Mn2+-doped LiZrO2-LiBaZrO3 nanocomposite are depicted in Fig. 9. The defect levels were estimated from the deconvoluted bands of PLE and emission spectra as follows; the excitation bands at 204 nm, 242, 272, 285 and 304 nm, and emission bands at 404 nm, 417 nm, 438 nm and 472 nm of the host, respectively. Similarly, the transition levels of Mn2+ were estimated using the same method. The emission process begins with electrons being excited from the ground state into the conduction band or defect state by absorbing UV radiation at 205 nm or 242/272/285/305 nm, respectively. In the case of ZrO2, the electrons then relax nonradiatively to a metastable state associated with oxygen vacancy, from where radiative transitions to the valence band produce violet-blue emissions at 417 and 438 nm. In the case of BaZrO3, after band-to-band excitation or defect related excitation, electrons relax nonradiatively and are trapped in localized (deep- or shallow- trap) states in the lattice from where electron-hole recombination takes place at the valence band by several paths producing blue-band emission with a peak at 477 nm. The broad blue-emission band suggests an emission mechanism characterized by the participation of several energy levels. Since the emission from BaZrO3 is essentially blue light, it is an indication that the transitions involved are the shallow traps states. Meanwhile, in both cases, energy is transferred from the defect levels of the host to 4A1, 4E (G) and 4T2 (G) levels of Mn2+ (in the tetrahedral LiZrO2 and octahedral LiBaZrO3 phases) because of the comparable values of the energy levels. The emission wavelength from the host at 417 nm, 438 nm and 477 nm overlap the 6A1 – 4T2(4D), 6A1 – 4A1(4G) and 6A1 – 4A1(4E) excitation transitions reported for Mn2+[71]. Then electrons in the Mn2+ ion from the ground state are excited by the energy received from the defect excited states to the 4A1, 4E (G) and 4T2 excited states of Mn2+ ions in the tetrahedral LiZrO2 phase. From the excited states of Mn2+, the free electrons relax nonradiatively via intermediate energy levels to the lowest excited state 4T1 (G), followed by a radiative transition to the ground state 6A1 (6S), giving rise to a green emission band at 527 nm. The weak red emission band also from the 4T1 (G) → 6A1 (6S) transition of Mn2+ ion follows the same process as described for the green emission, except that the process, in this case, involves the 4T1 emitting state of Mn2+ ions situated in the octahedral sites of LiBaZrO3. This emitting state is lower than that of tetrahedrally coordinated Mn2+ in the ZrO2 phase of the nanocomposite. The decrease in the intensity of the blue band may be due to a reduction in the shallow traps (emergence of deep traps as seen in the red emission of low intensity) as a result of an increase Mn2+ concentration or energy transfer from the trap state emission to the dopant ion.
The energy transfer efficiency (η) from the defect states of the host nanocomposite to Mn2+ was estimated from the expression [78];
where ${I_s}$ and ${I_{so}}$ arethe intensities of host defects in the presence and absence of Mn2+, respectively. The calculated average ET efficiency for LiZrO2-BaZrO3 nanocomposites doped with Mn2+ is 25%. The lowest ET efficiency value, 8% was obtained for low Mn2+ concentration due to large separation between the ions and the effective defect states responsible for the energy transfer process. The largest energy transfer efficiency is obtained at the highest Mn2+ ion concentration (x = 0.06). Considering the energy transfer process from host to Mn2+ to be via multipolar interaction, then Dexter’s energy transfer relation can be applied to know the dominant multipolar processes responsible for the ET in this nanocomposite [79]; where ${I_s}$ and ${I_{so}}$ are defined for Eq. (7), C is Mn2+ ion concentration, and n = 6, 8, and 10 correspondings to dipole-dipole, dipole - quadrupole, and quadrupole - quadrupole interaction, respectively. The plots of $({{{I_{so}}} / {{I_s})}}$ verses ${C^{n/3}}$ are shown in Figs. 10(a)–10(c). The best linear fit is obtained for n = 6, indicating that the energy transfer process from the host to Mn2+ in the nanocomposite follows a nonradiative electric dipole-dipole interaction.Fig. 10. Dependence of ${{{I_{so}}} / {{I_s}}}$ on host defects emission on (a) ${C_{M{n^{2 + }}}}^{6/3}$, (b) ${C_{M{n^{2 + }}}}^{8/3}$ and (c) ${C_{M{n^{2 + }}}}^{10/3}$
On the other hand, the quenching of Mn2+ emission at high concentration beyond x = 0.05 is a result of a large population of Mn2+ ions which shortens Mn2+ to Mn2+ average distance and enhances energy transmission between them. There are two mechanisms responsible for energy transfer among the Mn2+ ions, viz; exchange interaction and multipolar interaction. Therefore, to identify the energy transfer process responsible for the concentration quenching of Mn2+ emission, the critical distance (Rc) for ET between Mn2+ ions was determined using the Blasse formula [80–82];
where V is the unit cell volume taken from Fig. 4(a) with average values 73.64 Å for tetragonal LiZrO2 and 68.33 Å for cubic LiBaZrO3 phases, Xc = 0.05 is the critical concentration of the activator ion (Mn2+) and N = 4 is the number of Mn2+ ions per unit cell. The obtained values; 8.9 Å and 8.7 Å for tetragonal and cubic phases, respectively are larger than 5 Å for the case of exchange interaction [79], suggesting that radiation quenching for LiZrO2-BaZrO3: Mn2+ nanocomposites phosphors are dominated by the multipolar interaction process.3.4 Commission Internationale de l’éclairage (CIE)
The CIE coordinates and chromaticity diagram of undoped and LiZrO2 - LiBaZrO3: xMn2+ nanocomposites obtained from the PL emission values are presented in Table 2 and Fig. 11. The CIE coordinates of the undoped phosphor (x = 0.1607, y = 0.0669) is clearly in the blue region. Figure 11 shows that the colour emitted by LiZrO2 -LiBaZrO3: xMn2+ nanocomposite is tunable from blue to green and then near-white light by only changing Mn2+ content. It is evident that the more Mn2+ is added to the host matrix, the more intense the green emission is, except at x = 0.06 mol ratio, where the CIE coordinates (x=0.26086, y=0.36230) appears to be in the cool white region as shown in Fig. 11.
Fig. 11. CIE chromaticity coordinate of LiZrO2-BaZrO3:xMn2+ [x = 0.00(0), 0.01(1), 0.02(2), 0.03(3), 0.04(4), 0.05(5) & 0.06(6)] nanocomposites.
Table 2. CIE coordinates for different Mn2+ concentration
4. Conclusion
LiZrO2-BaZrO3:xMn2+ nanocomposites were synthesized using the co-precipitation method. The XRD results show a structural match with the cubic phase of bulk BaZrO3 and the tetragonal phase of ZrO2. The nanocomposites are capable of been excited by UV-light with maximum wavelengths at 240 nm and 270 nm. PL spectra display broad emission bands in the blue and green regions. The blue emission band originates from the host matrix, while the green emission band is due to the activator (Mn2+) situated at tetrahedral coordinate sites in the nanocomposite. The emission intensity of the green band increases with increasing Mn2+ concentrations. The emission colour of the nanocomposite can be tuned from blue (due to host matrix) to green by only changing the concentration ratio of Mn2+ from 0 to 0.05 mol ratio and then white light at 0.06 mole ratio. Therefore, the nanocomposite presented in this work is a promising material for use as a multicolour phosphor in field-emission displays.
Disclosures
The authors declare no conflicts of interest.
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