Structural evolution induced preferential occupancy of designated cation sites by Eu2+ in M5(Si3O9)2 (M = Sr, Ba, Y, Mn) phosphors

In this paper, we present new insight into a changing Eu2+ crystallographic site preference in Eu-doped M5(Si3O9)2 (M = Sr, Ba, Y, Mn), which is related to the structural variation induced by M cation substitutions. The effect of the local structural geometry on the luminescence properties of Eu2+ is revealed. By substitution of Ba2+ for Sr2+, the lattice expansion is restricted to specific cation sites, resulting in the abrupt blue shifted emission of Eu2+ ions. The abnormal blue shift on replacing Sr2+ with Mn2+ is attributed to the preferential 6-fold coordination for Mn2+ that moves the Eu2+ ions to other sites. The results elucidate the mechanisms of emission band adjustment by local site coordination change and it can be potentially extended to crystals which properties are sensitive to local lattice variations.

Rare-earth-activated silicates are widely used as phosphor materials for white light-emitting diode (WLED) illumination because of their high quantum efficiency and low cost. [1][2][3] However, unsatisfactory thermal stability and spectral position of the excitation/emission bands limit the applications of these materials. To overcome these disadvantages, different cation substitutions were used to adjust the spectroscopic parameters of silicate phosphors, and numerous corresponding mechanisms were proposed to elucidate the relationship of the structural and luminescence properties. These mechanisms include the "cation size mismatch", "neighboring cation substitution", "nanosegregation and neighbor cation control" and "chemical unit co-substitution" effects. [4][5][6][7][8][9][10][11][12] However, these mechanisms cannot ultimately solve all problems, and many questions remain unanswered. To date, the cation substitution effect, which can systematically tune spectral position and thermal quenching by changing the coordination environment of the selected cation sites and controlling the preferential activator ion occupancies, has rarely been considered. In the present study, we present a new insight into the structural variation induced site-preferential occupancy of Eu 2+ in M 5 (Si 3 O 9 ) 2 (M ¼ Sr, Ba, Y, Mn) phosphors through sizemismatched cation substitutions applicable to spectral positions and thermal stability tuning.
In this study, the Sr 2.97Àx Ba x Eu 0.03 Y 2 (Si 3 O 9 ) 2 (Ba series) and Sr 2.97Ày Mn y Eu 0.03 Y 2 (Si 3 O 9 ) 2 (Mn series) solid solutions (0 # x # 1.59, 0 # y # 0.63) were successfully prepared. The doping concentration of Eu 2+ was controlled at the level of 1 at% of Sr 2+ in the Sr 3 Y 2 (Si 3 O 9 ) 3 host. The synthesis route and characterization description can be found in ESI. † Firstly, the phase composition and structural properties of the samples were identied by XRD analysis. For Ba series, all diffraction peaks of the compounds with x ¼ 0, 0.06 and 0.09 were indexed in the monoclinic cell (C2/c) of Sr 3 Y 2 (Si 3 O 9 ) 3 , as plotted in Fig. S1 (ESI †) and Fig. 1a. 13 A similar C2/c cell was obtained for the composition range of 0.18 < x < 1.59. However, a noticeable difference was observed in the diffraction patterns in comparison with those recorded from the low-doped samples (Fig. S1 (ESI †)), implying a possible phase transition induced by the substitution of Ba 2+ for Sr 2+ . Accordingly, these phases can be reasonably differentiated as Phase 1 (0 # x # 0.09) and Phase 2 (0.18 < x # 1.59). Samples with x ¼ 0.12, 0.15, 0.18 represent the mixture of these two phases. Notably, the pure phase state of the Ba series is destroyed at x > 1.59, and many impurities appeared, including, besides (Sr,Ba) 3 Y 2 (Si 3 O 9 ) 2 , the known silicates BaSi 2 O 5 , Ba 5 Si 8 O 21 and BaY 2 Si 3 O 10 ( Fig. S1 (ESI †)). Thus, the structural analysis was mainly focused on the evolution from Phase 1 to Phase 2.
The evident chemical shi from $604.2 to $650.4 cm À1 at x $ 0.15 was observed for the representative band from the Raman spectra of the Ba series, verifying the phase transition emergence, as shown in Fig. S2 (ESI †). To determine the effect of this phase transition on the coordination environment of M (Sr/Ba/Y/Eu) ions, Rietveld renement was performed for the Ba series. The crystal structure of Sr 3 Y 2 (Si 3 O 9 ) 2 was taken as a starting point, 13 The renement parameters and cell parameters a, b, c, and V are presented in Table S1, Fig. S3-S17 (ESI †), and Fig. 1a. The chemical composition of the compounds was close to the nominal compositions, as obtained by the inductively coupled plasma element analysis (Table S2, (ESI †)). Thus, the total site occupancies were constrained during renement in accordance to the designed compositions. As shown in Fig. 1a, the linear increase of cell parameters and volume on the Ba concentration increase at x < 0.09 and x > 0.18 in accordance with Vegard's rule proves the solid solution formation at Phase 1 and 2 ranges. The cell parameter jump appeared at x $ 0.15. The a and c values markedly decreased at x $ 0.15, whereas b increases abruptly.
Unexpectedly, V noticeably decreases at x $ 0.15, as the Ba concentration increase should imply a cell volume increase. This situation is unusual and it may lead to an unexpected change in the coordination environment of Eu 2+ .
According to the renement results (  (Fig. 1a). In addition, the dependence of occupancies in the three sites was investigated. It is clearly shown in Fig. 1c (line 2) that only M2 sites were occupied by Ba 2+ ions aer the phase transition, whereas M1 and M3 were preferential sites for Y 3+ ions. To emphasize the main features of phase transition associated with three M (Sr/Ba/Y/Eu) sites, the model of structural transformation was built, as shown in Fig. 2a (Fig. 1b). The preferential occupation of M3O 6 octahedron by smaller Y 3+ ions results in a shrinkage of this polyhedron.
Monocapped trigonal prism M2O 7 is markedly shrank along the c-axis direction from 2.67Å to 2.42Å (Fig. 1b). All these shrinkages in the M sites nally lead to the cell volume decrease and a shrinkage between two SiO 4 polyhedra. Although the   shrinkage occurs along the c-axis of the M2O 7 polyhedron, the nal d(M2-O) increases and is, particularly, increasing abruptly during phase transition. This abnormal phenomenon originates from the expansion along the a-and b-axes. This situation also causes the expansion between two SiO 4 polyhedra along the a-axis direction. Generally, with the Ba 2+ content increase from the phase transition point, only M2O 7 polyhedra expand, whereas M1O 8 and M3O 8 polyhedra shrink regardless of a nal increase in the cell volume. Therefore, Ba 2+ ions preferably occupy the M2 site, whereas the M1 and M3 sites are preferential for Y 3+ ions.
The 5d-4f transition of Eu 2+ is sensitive to the structural variation of the host lattice that even a slight change in the local coordination environment around Eu 2+ can cause a big effect for its luminescence. [15][16][17][18][19][20] The photoluminescence excitation (PLE) and emission (PL) spectra of the Ba series are shown in Fig. 3a and b, respectively. For x ¼ 0, the PLE consists of a broad band from 250 to 430 nm with the maximum at 365 nm. Under 365 nm UV, a bluish-green emission is given with CIE color coordinates (0.168, 0.258). The PL spectrum covers a broad range from 425 to 575 nm, centered at 474 nm. The asymmetric emission results from the three Sr 2+ sites available for Eu 2+ , as shown by the three Gaussian tting peaks at 21 978 cm À1 (455 nm), 20 833 cm À1 (480 nm), and 19 231 cm À1 (520 nm) at M1, M2 and M3 sites (Fig. 3b). With the Ba 2+ -doping in Phase 1, a slight blue shi emission from 474 nm (x ¼ 0) to 469 nm (x ¼ 0.09) was observed, as shown in Table S3 and Fig. S18 (ESI †). This shi is attributed to random occupation of larger Ba 2+ over Sr 2+ sites in Phase 1 and the resulting cell enlargement. Thus, the crystal eld splitting (CFS) of Eu 2+ 5d energy levels in these enlarged sites weakened, resulting in blue shi emission. During the phase transition, the large blue shi emission from 468 to 438 nm occurred. The result implied that the lattice environment around Eu 2+ became looser, and thus the average Eu-O bond length (d) increased. Generally, the crystal eld strength is proportional to 1/d 5 Therefore, Eu 2+ ions in the looser sites with the longer bond length will possess a higher energy emission and it will generate a blue shi of the emission band.
In view of the size-difference between Ba 2+ ions and Sr 2+  with x. Given x > 1.59, the stable structure was destroyed, and numerous foreign phases appeared. Hence, certain low-energy emission peaks emerged, generating a broadened emission and red shi, as shown in Fig. S19 (ESI †). Generally, a large spectral blue shi in the Ba series was induced by the change in the crystal eld environment at specic cation sites because of the phase transition. Thus, it offers a novel and efficient route to tune the optical properties of phosphor materials.
On the basis of the same principle, a spectral red shi can be expected in the Mn series because of the enlarged CFS of Eu 2+ by the substitution of smaller Mn 2+ for larger Sr 2+ . Actually, unusual blue shis in the Mn series spectra were observed with the Mn 2+ content (y) increase, as shown in Fig. 4a and b. To reveal the mechanism of the luminescence blue shi, the phase purity and structure variation in the Mn series were rst analyzed by XRD and Rietveld renement. Evidently, the Mn series samples crystallize in the monoclinic Sr 3 Y 2 (Si 3 O 9 ) 2 phase, space group C2/c, and the XRD diffraction peaks continuously shi to larger angles with the y increase, as shown in Fig. S20 (ESI †). This is in agreement with Vegard rule, revealing the solid solution formation. 4 For the Mn series, the renement results are presented in Table S4 and Fig. S21-S26 (ESI †). The results verify the generation of host-type structures due to Mn/Eu doping. The cell volume decrease with the y increase reects the formation of Mn series solid solutions (Fig. S27 (ESI †)).
As shown in Fig. 4a, at y < 0.45, the PLE spectra of the Mn series are similar to those of the Ba series, except for the maximum at 355 nm. The PL spectra consist of two broad bands in the range of 400-650 nm, which are centered at 472 nm and 545 nm, respectively. The rst band is attributed to the 5d-4f transition of Eu 2+ , whereas the second band is related to the 4 T 1g ( 4 G)-6 A 1g ( 6 S) transition of Mn 2+ . Notably, the emission positions of the Eu 2+ and Mn 2+ bands remain persistent up to Mn 2+ doping content y $ 0.45. However, the emission intensity of the Mn 2+ band gradually increased with the y increase, whereas the Eu 2+ emission intensity rst increases and then decreases. Thus, a tunable single-composition white light  This implies that Mn 2+ ions are more easily coordinated with six oxygen atoms, following the priority of MnO 6 > MnO 8 > MnO 7 . Therefore, a possible mechanism is proposed for the abnormal blue shi of Eu 2+ emission (Fig. 4d). At low Mn 2+ doping levels (y < 0.45), Eu 2+ and Mn 2+ ions randomly enter M1-M3 sites. At y > 0.45, the d(M3-O) is lower than 2.33Å, which is smaller than the sum of Eu 2+ and O 2À ion radii [r(Eu 2+ ) + [6] r(O 2À ) ¼ 1.17Å + 1.4Å ¼ 2.57Å], but larger than the sum of Mn 2+ and O 2À ion radii [r(Mn 2+ ) + [6] r(O 2À ) ¼ 0.67Å + 1.4Å ¼ 2.07Å] [6] . Therefore, Mn 2+ ions preferentially occupy the M3 sites, driving Eu 2+ ions to the M2 and M1 sites. As mentioned in the previous section, the Eu-O bond length signicantly affects the crystal eld strength (D q ), that is, D q is proportional to 1/d 5 , and the looser site accommodating Eu 2+ ions should correspond to a higherenergy (shorter wavelength) emission peak, while a lowerenergy (longer wavelength) emission appears. 21 Therefore, at y ¼ 0.63, the Eu 2+ ions mainly stay in the M1 sites, as conrmed To ensure high efficiency for the phosphor-converted WLED devices, a comprehensive understanding of the thermal quenching of phosphors is necessary. [22][23][24][25] The relative emission intensity (I T /I 0 ) for the Ba and Mn series from room temperature to 573 K are shown in Fig. 5a and b. The emission intensity decreases with the environmental temperature increase for all the samples because of the thermal quenching effect. However, the overall trend of thermal stability improvement by doping is evident in the Ba series (x ¼ 0-1.59) and the Mn series (y ¼ 0-0.63), with respect to the parent Sr 2.97 Eu 0.03 Y 2 (Si 3 O 9 ) 2 sample. In particular, the Ba series shows a clear thermal stability increase at the phase transition. The results can be governed by the increasing quenching activation barriers in both series (insert in Fig. 5b). Generally, excellent thermal stability could be expected in certain phosphors with high covalency, high rigidity and a small Stokes shi. The blue shis in the Ba and Mn series cause the Stokes shi decrease, indicating a thermal energy increase (E a ). The E a values of the Ba and Mn series were calculated by relation I T /I 0 ¼ [1 + D exp(ÀE a /kT)] À1 , where I T (intensity at T), I 0 (intensity at T ¼ 0), D, and activation energy E a are rened variables. 10 Clearly, the E a values of the Ba and Mn series gradually increase with x and y (insert in Fig. 5b), suggesting that the probability of nonradiative transition is weakened. Therefore, the thermal stability of the phosphors gradually increased on doping. A clear thermal quenching decrease appeared at the phase transition point, as shown by the dashed circle in Fig. 5a. This nding is also consistent with the E a turning point during the phase transition (the insert in Fig. 5b). Thus, a jump change of thermal stability is possible in phosphor materials at phase transition. It is noted that the Mn series have a lower enhancement for the thermal stability than the Ba series, which can be attributed to the different thermal decay behavior between Eu 2+ and Mn 2+ .  To sum it up, the lattice-site control effect for Eu 2+ ions discovered in M 5 (Si 3 O 9 ) 2 (M ¼ Sr, Ba, Y, Mn) crystals can efficiently tune the photoluminescence and thermal quenching properties via the cation-substitution approach. This effect provides a new insight into the structural variations of the single-cation sites that are induced by phase transition and sitepreferential occupancy, which are driven via size-mismatched cation substitution and can tune the luminescent properties of phosphor materials. This work reveals the mechanisms of optical adjustment by the coordination environment changing at specic sites. In particular, the abrupt structural change owing to phase transition offers unexpected and large-scale changes in optical properties. This effect can be extended to tune other properties, including the electric and magnetic properties that are sensitive to the structural variation at local sites. [26][27][28]