Warm to neutral white light emissions from Dy3+–Eu3+ co-doped glass ceramics containing NaBSiO4 crystalline phase for W-LEDs applications

This paper reports on the structural, optical and luminescence studies of Dy3+–Eu3+ co-doped glass ceramics that are obtained via thermal treatment method. The initial confirmation on glass ceramics with the formation of NaBSiO4 crystalline phase was made through XRD study. The FTIR study showed the vibrations of network formers (B2O3, SiO2) and other functional groups. The quantity of non-bridging oxygens (NBOs) are increased in the glass ceramics with increase in annealing temperature. The absorption spectra (UV–visible-NIR) showed the possible transitions of Dy3+ and Eu3+ ions in the glass ceramics. A red-shift in the absorption band-edge and reduction in the optical band gap values were obtained for glass ceramics owing to their heating temperatures. Photoluminescence studies showed the excitations of Dy3+ and Eu3+ ions under 575 nm and 613 nm emission wavelengths. The obtained NaBSiO4 crystalline phase in the glass ceramics has enhanced the luminescence intensity, and lifetimes of Dy3+ and Eu3+ ions compared to unheated precursor glass when excited under 350 nm and 393 nm. Color chromaticity diagram and correlated color temperature (CCT) values showed a shift in the color of light from warm white of precursor glass to neutral white of glass ceramics. The overall results justify the possibility of considering the Dy3+-Eu3+ co-doped glass ceramics as efficient materials for solid-state lighting applications like W-LEDs.


Introduction
In today's fast growing world, the need for energy saving materials and devices are considered to be a primary source. Solid-state lighting devices such as white-light emitting diodes (W-LEDs) has recently fulfilled the human needs due to its lower power consumption typically 2-17 W of electricity which is equal to one-third of its counterparts such as incandescent/CFL bulbs, and also provides longer lifetime [1,2]. The conventionally available W-LEDs are obtained by one of the two methods i.e. (1) a blue LED chip (InGaN) combined with yellow Ce 3+ :Y 3 Al 5 O 12 phosphor and (2) combining three color LEDs under UV excitation. The former method resulted in poor color rendering index (CRI) and high correlated color temperature (CCT) because the 5d → 4f transition of Ce 3+ gives an intense blue emission which lacks the red spectral component needed for color mixing; whereas in the latter method different drive voltages are needed for the three different color emitting LEDs [3]. Thus, to compete with the commercially available phosphor converted W-LEDs, and also to substitute for the traditional lighting sources such as incandescent lamps, fluorescent lamps, and halogen lamps, rare earth doped solid materials like glasses are highly researched [4]. White light emission from glasses have always been regarded as a better choice due to their better thermal resistance, quick preparation, low cost-effective, easy dispersibility of trivalent rare earth ions. With a variety of applications including fiber amplifiers, solid-state lasers, telecommunications, optical displays and sensors [5], rare earth doped glasses deepen their roots in solid-state lighting (SSL) technology. Among the lanthanides, Dy 3+ is highly doped into the glasses because of their intense and reliable emissions from the 4 F 9/2 level to 6 H 15/2 (blue), 6 H 13/2 (yellow), and 6 H 11/2 (red) levels [6,7]. However, the red emission in Dy 3+ doped glasses are less intense which might be a disadvantage of obtaining a warm to neutral white light. Thus, by incorporating suitable lanthanide activators such as Eu 3+ , Sm 3+ or Pr 3+ the color of the white light can be tuned for desired applications [8,9]. Among the rare earths, Eu 3+ is considered as co-dopant for Dy 3+ due to its intense red emission at 613 nm ( 5 D 0 → 7 F 2 ). Altering the Eu 3+ percentage in the glasses with optimized Dy 3+ can tune the color quality of white light which is reported in one of our recently published papers [10]. However, the complete amorphous nature of the rare earth doped glasses may not enhance the emission as like crystalline phosphor and thus their practical applications are restricted. To obtain the crystalline property of phosphor with the limited growth of crystallite size, the possible method is to convert the glasses to glass ceramics through heat treatment process. Glass ceramics (GCs) will favour with the crystalline features of phosphor powders and the thermal stability of glasses. The rare earth ions can easily incorporate within the formed crystalline phase in the glass ceramics resulting in enhanced emission intensity. Some of the literature articles available on Dy 3+ -Eu 3+ co-doped glass ceramics are reported in the following references [11][12][13][14][15][16]. In this present work Dy 3+ (0.5 mol%) and Eu 3+ (1.0 mol%) co-doped glass is converted to glass ceramics under different heating temperatures. From the XRD studies, the crystalline phase formed in the glass ceramics is analysed which is referred to sodium boron silicate (NaBSiO 4 ), a mixed crystalline phase. The structural, optical and luminescence behaviour of the Dy 3+ -Eu 3+ co-doped glass ceramics with NaBSiO 4 crystalline phase is discussed in the paper. The color coordinates and the correlated color temperatures thus showed a transition from warm white light to neutral white light from the prepared glass ceramics.

Preparation of glass ceramics
With respect to our previous reported work [10] the glass composition is taken as 20S iO 2 -18. 5B 2 O 3 -10 Al 2 O 3 -10 ZnO -30 NaF -10 ZnF 2 -0. 5Dy 2 O 3 -1.0Eu 2 O 3 . The raw materials in the glass composition were grinded well, melted at 1320 °C for 2 h, and quenched on a pre-heated brass plate (350 °C) to obtain the solid glass. The quenched glass pieces were annealed at the quenched temperature 350 °C for nearly 2 h to maintain a good transparency of the glasses. Later, the glass pieces were polished well, and they were heated above their glass transition temperature ( T g ) (obtained through DTA plot given in Fig. 1). In this work four selective temperatures were chosen at 470 °C, 510 °C, 550 °C, and 590 °C to heat the glasses for a duration of 3 h. After heat-treatment the base glasses are converted to glass ceramics which is quantified through the XRD studies (Fig. 2).

Characterization and Instrumentations
The differential thermal analysis was performed using using STA7200 HITACHI thermal analyser system operated in the nitrogen atmosphere presence at a heating rate of 10 °C/ min. X-ray diffraction studies were carried out using Rigaku MiniFlex 600 X-Ray Diffractometer with Cu-Kα radiation source of wavelength 1.5418 Å in the range 10-70° with a step size of 0.02° and scan rate of 1 degree per minute. Fourier transform infrared spectroscopy studies were performed with IR Affinity-1 Shimadzu FTIR Spectrophotometer of high resolution 0.5 cm −1 . Field emission scanning electron microscopy studies were recorded on the surface of the glass ceramic using CARL Zeiss FESEM 03-81 instrument. Ultra-violet-Visible-NIR absorption studied were done using Perkin Elmer Lambda 750 s Spectrophotometer. Photoluminescence measurements were carried out with JASCO FP-8500 spectrofluorometer. Decay analysis is recorded  Figure 1 shows the differential thermal analysis (DTA) plot of precursor glass (DE-PG). The first characterization temperature called the glass transition temperature (T g ) is marked at 434 °C with a small dip at the curve due to the phase transition from a glassy state into a viscous or rubbery state. The second temperature at 562 °C is the onset crystallization (T x ) temperature such that the viscous glass start to form crystal growth. Followed by this is the peak crystallization (T c ) temperature at 593 °C where the growth of crystallites can be more and fast. The final temperature marked in the plot is the melting temperature (T m ) at 743 °C characterized by an endothermic peak similar to glass transition such that at this point the crystalline or semi-crystalline region of the glass change to a solid amorphous phase. From the reported temperatures, the thermal stability of the precursor glass is given as

Differential thermal analysis
The ΔT value is 128 °C for the precursor glass. The thermal stability above 100 °C tells that the glass is stable enough to prevent from fast crystallization [17,18] and it also implies that the glass can be used for photonics and optoelectronic applications. From the aforementioned temperatures, it is noted that the T g value is obtained at 434 °C, thus heat treating the precursor glass above this temperature can induce growth of tiny crystallites in the glass domain.
Therefore, four temperatures such as 470 °C, 510 °C, 550 °C, and 590 °C were selected to convert the precursor glass to glass ceramics.

XRD analysis
The XRD profiles for Dy 3+ -Eu 3+ co-doped glasses heattreated at different temperatures are provided in Fig. 2. The base glass looks completely amorphous with no peaks but only broad hump in the range 20°-40°. But the heat treatment to the base glasses at different temperatures resulted in formation of crystal growth. The obtained crystalline phase belongs to sodium boron silicate (NaBSiO 4 ), a hexagonal crystal system associated to the JCPDS file No: 77-0718. The peaks indexed with the obtained crystal phase is seen at 32.26°, 38.72°, 46.69°, 58.23° corresponding to (112), (300), (221), and (223) planes. There are no impurity or secondary formation as observed from the XRD spectra but a slight improvement in the peak intensity is seen which could be a signature of higher growth of crystallites. The lattice constants for NaBSiO 4 are given as a = 8.035 Å and c = 7.703 Å and the space group lattice is P6 3 (173). Due to very low intense peaks observed the crystallite sizes were not calculated for the glass ceramics.

FTIR analysis
The FTIR absorbance spectra of the glass ceramics is given in Fig. 3a. The major modes of vibrations such as (i) B-O-B, Si-O-B, AlO 4 , (600-800 cm −1 ) (b) BO 4 , SiO 2 (800-1200 cm −1 ) and (c) BO 3 (1200-1600 cm −1 ) are marked in the spectra. The vibration of hydrogen bonds and hydroxyl (O-H) groups are also marked in the range 2300-2500 cm −1 and 3200-3600 cm −1 [19]. To evaluate the number of underlying peaks within the range  Fig. 3b. The respective band assignments with their peak centres and relative areas are given in Table 1 [17,[19][20][21][22][23][24]. The glass network consists of tetragonally coordinated network units ( BO 4 ) and trigonally coordinated network units (BO 3 ) that are associated to the network formers such as B 2 O 3 and SiO 2 . When Dy 3+ and Eu 3+ dopants are doped into the glass network, they break the network bonds by replacing the cationic positions like B and Si. In such a case more of BO 4 gets converted to BO 3 . Similarly, when a glass is converted to glass ceramics through heat-treatment a new crystalline structure is formed which ruptures the existing BO 4 units and convert them even more to BO 3 units. Therefore, the percentage of BO 3 increases in the glass ceramics (or) we can say that non-bridging oxygens ( BO 3 ) are formed more whereas there are fewer bridging oxygens ( BO 4 ). These bridging (BOs) and non-bridging oxygens (NBOs) are denoted by the symbols N 4 and N 3 and they are calculated using the areas under the deconvoluted peaks as reported in our previous article [10,[24][25][26]. Table 2 shows the amount of bridging and nonbridging oxygens present in the precursor glass and glass ceramics. The amount of NBOs are increased with heat treatment and growth of NaBSiO 4 crystal phase in glass  ceramics. This signifies that the defects in glass network were improved when the glasses were subjected to higher heating temperatures.

FESEM studies
The FESEM image of the Dy 3+ -Eu 3+ co-doped precursor glass in Fig. 4 looked plain, non-agglomerated and nonporous, and the corresponding EDAX spectrum showed the presence of all the raw elements in the glass network.
The FESEM images of glass ceramics at different heating temperatures are displayed in Fig. 5a-d. With increasing heating temperatures, the images exhibited evident agglomeration changes with a non-uniform structure. The structure of NaBSiO 4 obtained for the heat-treated glasses is similar to those images reported by Aran Rafferty et al. [27].
In Fig. 7a, it is noted that the absorbance band-edges of the glass ceramics showed a red-shift on increasing the heating temperatures. This shift in band-edges could be attributed to the formation of new energy levels or defect states near the ultraviolet region with the growth of NaBSiO 4 crystals. The optical band gap energies are determined through the Tauc's plots as shown in Fig. 7b [30]. The band gap values tend to decrease for the glass ceramics compared to that of precursor glass due to the additional energy levels formed between valence band and conduction band with formation of NaBSiO 4 crystals. The observed values are given in Table 3. Also, the results match with the FTIR studies such that an increased trend in non-bridging oxygens for the glass ceramics showed a decreased trend in the band gap values.
Under both 350 nm and 393 nm excitations, the emission intensity of glass ceramic is initially dropped for lower annealing temperature i.e., at 470 °C (DE-GC1) due to the very low crystal growth in the glass system. Such a drop in emission intensity subjected to thermal treatment is known Fig. 7 a Absorbance band-edges and b indirect optical band gap plot of Dy 3+ -Eu 3+ co-doped glass ceramics  as thermal quenching [32]. However, the glass ceramics obtained at a temperature above 470 °C showed consecutive increase in their emission intensity and achieved a maximum for DE-GC4 glass ceramic. In this case, the energy transfer between the rare earths is minimized or not noticed as the concentration of both Dy 3+ and Eu 3+ are maintained constant for all the glass ceramics provided only their heating temperatures are varied. It is obvious that the intensity of emission spectra was enhanced after crystallization. It is also understood that crystals have a lower phonon energy compared to that of rare earth doped glass. If the phonon energy is less in the glass matrix, then the non-radiative transition rate will be minimized. Therefore, the enhanced luminescence of the glass ceramics can be attributed to the incorporation of Dy 3+ and Eu 3+ ions into the NaBSiO 4 crystal sites which lowers the non-radiative energy transfer between Dy 3+ and Eu 3+ .

Decay analysis of Dy 3+ -Eu 3+ co-doped glass ceramics
By providing a suitable excitation and emission wavelengths to the intermediate levels of Dy 3+ and Eu 3+ , the decay times of these rare earths in the glass ceramics can be determined. In the case of Dy 3+ , 4 F 9/2 is the intermediate level from which the intense emission falls to the lower level 6 H 13/2 (575 nm). Thus, the lifetime of Dy 3+ ions in 4 F 9/2 level are measured by providing the excitation and emission source of 350 nm and 575 nm. The decay curves are shown in Fig. 10a which exhibited a bi-exponential nature. The lifetime values are evaluated through exponential fitting of the decay curves to Exp-Dec2 Fit as shown for DE-GC4 glass ceramic in Fig. 10b. The biexponential decay fitting equation is given as [33] where A 1 and A 2 are the constants, t 1 and t 2 are the luminescence decay times. From the inset table given in Fig. 10b two lifetime values are obtained and thus the average lifetime is determined via the equation given as Similarly, the lifetimes of Eu 3+ ions lying in the intermediate level 5 D 0 are obtained by performing the decay measurements at an excitation and emission wavelengths at 350 nm and 613 nm as shown in Fig. 11a. Here all the curves showed the bi-exponential behaviour and the lifetimes of Eu 3+ are obtained as similar to the above case. The bi-exponential behaviour with a fitting plot shown for DE-GC4 glass ceramic in Fig. 11b. Table 4 .

Color coordinates and correlated color temperatures
The CIE chromaticity plot for the glass ceramics are given in Fig. 12a, b. Under 350 nm, the color coordinates of precursor glass (DE-PG) is located a little way from the white light zone, whereas for glass ceramics with higher heating temperatures, the color coordinates lie much closer to the white light. The correlated color temperatures (CCT) under 350 nm are given in Table 5 [34]. With increasing the annealing temperatures, the white light moved from a   Table 6 [11,[35][36][37][38]. The results suggest that the present work showed better neutral white light emission making the glass ceramic favourable for W-LEDs applications. On exciting the glass ceramics at 393 nm, a red emission is observed whose color coordinates and CCT values are also given in Table 5. The corresponding CIE chromaticity plot given in Fig. 12b also showed the tuneable red emissions from the glass ceramics. Therefore, the prepared co-doped glass ceramics with the NaBSiO 4 crystalline phase can be used as a suitable material for neutral white-LEDs as well as red-LEDs application under ultraviolet excitations.

Conclusion
Glass ceramics co-doped with Dy 3+ and Eu 3+ ions were synthesized using thermal treatment method at different temperatures. The formation of NaBSiO 4 crystalline phase is confirmed via XRD studies. The FTIR spectra showed the different vibrational groups and increased quantity of non-bridging oxygens in the glass ceramics. The absorption

Declarations
Conflict of interest All authors certify that they have no known conflicts of interest that could have appeared to influence the present work.
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