Luminescence in Er³⁺ co-doped bismuth germinate glass–ceramics for blue and green emitting applications

Abstract

Glass samples with the composition 40 Bi₂O₃-60 GeO₂-x Er₂O₃ are prepared using the melt quenching procedure. Fourier-transform infrared spectroscopy (FTIR), differential thermal analysis (DTA), Ultraviolet/Visible/Near-infrared (UV/Vis–NIR) absorption, photoluminescence spectroscopy (PL), and X-ray diffraction (XRD) are used to characterize the glass and glass–ceramic materials. The glass transition and crystallization temperatures are determined. Crystallization kinetics was studied in non-isothermal conditions. The oscillator strengths and Judd–Ofelt parameters (Ω₂, Ω₄, Ω₆) for reported Er³⁺ absorption transitions are estimated, and they follow the trend Ω₂ > Ω₆ > Ω₄ for observed Er³⁺ absorption transitions. The iconicity of glass samples was unaffected by an increase in Er³⁺ ions. XRD and FTIR confirmed the formation of the Bi₄Ge₃O₁₂ phase after heat treatment of glass samples at the crystallization temperature. The CIE chromaticity diagram computes the CIE chromatic coordinates. The values for all glass and crystal samples are close to bright blue and green. Glass and glass ceramic samples are suitable for green and blue optoelectronics device applications.

1 Introduction

The glass crystal system offers two distinct advantages. It functions as a glass composition with various rare-earth or transition metal compositions. Glass samples containing rare-earth ions are of great interest due to their many practical and potential applications.

Due to the high conversion process of converting infrared rays into visible light and flat screens, which uses as glass lasers, optical fiber amplifiers, phosphors, electroluminescent devices, and memory devices [1, 2].

The application of bismuth and germanium in glass is significant because it uses as a layer for optical and optoelectronic devices, thermal and mechanical sensors, infrared transmitting windows, and active media for Raman optical fiber amplifiers. Due to its outstanding structural characteristics, this glass system uses in a wide range of fields [1], temperature sensors, 1.55 m amplifications, and planar waveguides [3]. The other method converts glass into a crystal phase through heat treatment. The crystal phase that results from this procedure has the same properties as a single crystal. It also has the advantages of being inexpensive and simple to make. Bismuth germinate oxide is a component of the glass–ceramic system.

The crystalline phase generates by heat treatment and is controlled by many parameters, such as the ratio of bismuth to germanium oxide, the melting temperature, and cooling rates [4].

Many researchers directed their attention to the preparation of bismuth germinate glasses and the Bi₄Ge₃O₁₂ phase. According to their physical properties, they can use optical waveguides in the IR (infrared) region, as well as active media for Raman optical fiber, scintillation counters, thermal and mechanical sensors, as well as other optical devices, such as optical fibers, optical switching, and optical memory [1, 5].

Chen et al. and Sanghi et al. [6, 7] have studied and compared the optical properties of different glass formations, such as bismuth glass, borate glass, and commercial laser glass doped with Er³⁺ ions. From this study, the researchers concluded that the spectral parameters of bismuth borate glass are better than those of other glasses.

Polosan et al. [8] investigated the melt-quenched glass-controlled crystallization of bismuth–germinate glass (40Bi₂O₃–60 GeO₂ mol %). And found that the formed phase affected by the annealing of the glass sample above the crystallization temperature (558 °C) formed the Bi₄Ge₃O₁₂ nano and traces of Bi₂GeO₅ crystallites. The number of nanocrystals grows as the annealing temperature rises, while their size remains constant in the 50 nm range.

Polosan et al. [9] investigated the production of Bi₄Ge₃O₁₂ (BGO) from a combination of oxides at 1323 K. The optical characteristics and dielectric functions of a sample made using BGO single crystals were compared.

In this article, we direct our efforts toward studying glass samples prepared in two steps: firstly, study the glass structure and determine the physical, thermal, optical, and emission properties, and determine the effect of adding Er₂O₃. Secondly, investigate the production and characteristics of the Bi₄Ge₃O₁₂ phase.

2 Experimental work

Glass samples having a composition of 40 Bi₂O₃-60 GeO₂-x Er₂O₃ mol%, created with x=0, 0.5, 1.0, 1.5, and 2 mol%. The samples were placed in an oven at 1100 °C for one h., a molten poured between two copper plates in the air. XRD system from Philips Analytical (type PW3710) used a Cu tube anode with wavelengths of K_α1 =1.54060 and K_α2 =1.54439 to examine the nature of all samples. The starting angle (2θ) was 10°, and the ending was 70°. The FTIR JASCO, 6300 infrared spectrometer using to measure the absorption spectra in the range of 400–4000 cm^-1 of all samples using the KBr pellets technique. The prepared samples were analyzed using differential thermal analysis (DTA) at temperatures ranging from 30 to 1200 °C, with heating rates of 5, 10, 20, and 30 °C/min, and Al₂O₃ powder as a reference material (Shimadzu DTA-50 analyzer) measured. A computerized recording spectrophotometer uses to measure the optical transmission spectra in the region of 190 to 2500 nm (JASCO, V-570). The emission measurements were calculated by using the (JASCO–FP-6300) Spectrofluorometer in the (200–800 nm) wavelength range.

3 Results and discussion

3.1 XRD

The state of the samples is amorphous or crystalline examined using XRD equipment. Fig. 1 reveals that the XRD pattern of the materials under analysis is devoid of identifiable peaks.

Fig. 1

XRD for glass samples prepared

Figure 1 indicates that the samples are amorphous.

3.2 FTIR

FTIR is a guide instrument that detects any structural changes formed in the vibration groups of the glass network. Fig. 2 shows the glass free of Er, an obvious band for Bi and Ge beginning at 475 cm⁻¹, representing Bi-O vibration in BiO₆ [10]. The band at 540 cm⁻¹ represents the distorted BiO₆ [1, 11]. The bands around 560–580 cm⁻¹ and 1045 cm⁻¹ represent the stretching vibration of GeO₄ overlapped with the BiO₆ vibration band [1, 10]. The two bands at 730 cm⁻¹ and 770 cm⁻¹ represent the stretching of GeO₄ [1]. The band at 860 cm⁻¹ represents BiO₆ [1] and the asymmetric stretching of GeO₄ [10]. The band at 1035 cm⁻¹ indicates the stretching band of Ge–O in GeO₄ [1]. The band at 1100 cm⁻¹ represents the Bi-O-Bi bond or Bi-O-Ge [1]. The first addition of Er₂O₃ begins to appear in several bands as the band at 495 cm⁻¹ represents the BiO₆ octahedral units [7, 12] or the bending vibration of BiO₃ units [1]. A little band at 600 Cm⁻¹ represents the Bi-O- stretching vibration in BiO6 [1]. The band at 443 cm⁻¹ was attributed to symmetry stretching of bridging oxygen (BO) in 6-membered GeO₄ rings Ge–O–Ge [13], and the continuous addition of Er broadened the band and could include bands. The appearance of a new band at 520 cm⁻¹ represents the defect mode of GeO₄ [13], while the band at 630 Cm⁻¹ represents the stretching of BiO₆. The bands at 710–750 cm⁻¹ represent GeO₄ [1]. The band at 845 cm⁻¹ corresponds to the symmetric stretching of BiO₃ [1, 7], while the band ranging from 750 to 900 cm⁻¹ corresponds to the Bi-O stretching vibration in BiO₃ units [11] and the asymmetric stretching of bridging oxygen in Ge–O-Ge [13].

Fig. 2

FTIR spectra of prepared glass sample

As for undoped and Er-doped Bi₂O₃-GeO₂ glass samples, there are broadband located between 850 cm⁻¹ and 1200 cm⁻¹, which should be corresponding to the vibrations of 850 to 930 cm⁻¹, 1016 to 1031 cm⁻¹ and 1090 to 1110 cm⁻¹, respectively. However, the line shape of the vibrations becomes broad compared to that of the 2Bi₂O₃-3GeO₂ glass sample, which suggests that the structures of 2Bi₂O₃-₃GeO₂ and Er-doped Bi₂O₃-GeO₂ are slightly different. With the deconvolution the broad bands consist of three bands at of 850 to 930 cm⁻¹, 1016 to 1031 cm⁻¹ and 1090 to 1110 cm⁻¹. Fig. 3 shows the effect of Er₂O₃ addition on these bands. From Fig. 3, observed the Bi–O–Ge linkage (1090 to 1110 cm⁻¹) decreases and the vibration of Bi–O bonds in BiO6 units and Ge–O–Ge asymmetric stretching of GeO4 increases as Er₂O₃ equal and greater than 1 mol %, while GeO4 vibration groups (1016 to 1031) constant for all samples. It is clear the structure variation between samples.

Fig. 3

Effect of Er₂O₃ content on vibration groups resulting from broadband analysis from 850 to 1200 cm⁻¹

3.3 Optical absorption

Figure 4 depicts the absorption spectra of all glass samples. The absorption spectrum of the glass sample free from Er³⁺ ions doesn’t have any transition band.

Fig. 4

the optical absorption of glass sampled doped with different concentrations of Er₂O₃

The first addition of Er₂O₃ in bismuth germanium glasses causes a range of Er³⁺ absorption transitions, which translate from the ⁴I_15/2 ground state to various excited states due to the f—f transition. The resolved bands in the 400–2000 nm range arise from the ⁴I_15/2 ground state to several excited states [7, 14,15,16,17] as ⁴F_7/2, ²H_11/2, ⁴S_3/2, ⁴F_9/2, ⁴I_9/2, ⁴I_11/2, and ⁴I_13/2, with center wavelengths of 486, 522, 542, 650, 796, 980, and 1524 nm, respectively.

3.4 Judd–Ofelt parameters

The intensity of the optical absorption line of Ln³⁺ ions is used in the modified Judd–Ofelt theory to determine the laser and optical characteristics of ions in crystals and glass [18]. The Judd–Ofelt hypothesis aids in determining the oscillator strength empirically. According to the following relationship [19]:

$${\text{f}}_{{{\mathbf{exp}}}} \,\, = \,\,{ 4}.{32} \times {1}0^{{ - {\mathbf{9}}}} \int {\varepsilon \left( \nu \right)d\nu }$$

(1)

ε (ν) is the molar absorbance of the peaks at a wave number (cm⁻¹).

The experimental oscillator strength of an electronic transition contains the electric and magnetic dipole strengths.

The magnetic dipole's intensity is determined by the properties of the rare-earth ion, not by the host. Magnetic dipole transitions have a lower probability of occurring than electric dipole transitions. The experimental oscillator line strength determines using [2, 17, 20, 21]:

$$fcal = \frac{{8\pi^{2} mc(n^{2} + 2)^{2} }}{3h\lambda (2J + 1)9n}\sum\nolimits_{(t = 2,4,6)} {\Omega_{t} \left( {\psi J\left\| {U^{\lambda } } \right\|\varphi^{^{\prime}} J^{^{\prime}} } \right)}^{2}$$

(2)

where λ is the wavelength (nm) of each transition from the ground state (J) to the excited state (J'), c is the velocity of light in a vacuum, m is an electron's rest mass, ${\boldsymbol{\Omega }}_{{\varvec{t}}}$ is the Judd–Ofelt parameters, n is the refractive index, h is plank's constant, and ${{\varvec{U}}}^{{\varvec{\lambda}}}$ is a doubly reduced matrix element.

The fit quality between the experimental and estimated oscillator strengths is determined by the root mean square deviation (δ_rms). The following relationship is used to calculate δ_rms:

$$\delta_{rms} \,\, = \,\,\left[ {\frac{{\sum (fexp - fcal)^{2} }}{N - 3}} \right]^{\frac{1}{2}}$$

(3)

The total number of energy levels is N. In Table 1, f_exp, f_calc, and δ_rms are calculated and tabulated. The δ_rms value indicates the fitting between experimental f_exp and calculated f_calc strength. Furthermore, using Judd–Ofelt’s theory, intensity parameters JO Ωλ (2, 4, and 6) for all samples are calculated and presented in Table 1.

Table 1 The Fexp, Fcal, δ, and Judd Ofelt parameter (Ω2, Ω4, Ω6) calculation

For all glass samples, the values of Judd Ofelt's parameters have a Ω₂ > Ω₆ > Ω₄ trend. The higher the value of the Ω₂ intensity parameter, the more asymmetric and covalent the rare-earth ion sites [22]. The value of Ω_λ showed the glass composition containing 1 mol% Er₂O₃ has the largest Ω₂, Ω₄, and Ω₆. The spectroscopy quality factor (χ) is connected to rigidity, laser determination, and amplifier application [16] and is equivalent to the ratio between Ω₄ and Ω₆ [22].

The resulting value was less than 1, indicating a hard glass [7, 23] used in the application of a fiber amplifier with high emission intensity (²H_11/2 → ⁴I_15/2) [12, 15, 24] in comparison to the literature [12, 25,26,27,28]. The glass composition doped with rare earth causes a wavelength shift, named the nephelauxetic effect. The nephelauxetic effect parameter was determined using the relation [23]:

$$\,\beta = \nu_{{\mathbf{c}}} /\nu_{{\mathbf{a}}}$$

(4)

where ν_c is the wave number of the Er³⁺ ion in the host and ν_a is the wave number of the Er³⁺ ion in aqua ions.

The bonding parameter is calculated using the formula δ = (1-β')/β' *100 where β' is the average of nephelauxetic. The bonding parameter determines the type of bonding between rare-earth ions and their environment. ν_a, ν_c, β, β', and δ are calculated and tabulated in Table 2.

Table 2 The values of νa, νc, β, β’, and δ bonding

The band position appears in the same place in Fig. 3, and the negative sign of the value denotes the ionic link between Er₂O₃ ions and the surrounding host.

The value of band gap energy E_opt (energy excites the electron from the top of the valance band to the bottom of the conduction band) indicates the bridging and non-bridging oxygen formed in the glass structure.

The optical band gap determines by drawing the relationship between the absorption coefficient and (αhv)² [Tauc plots] with hv, as shown in Fig. 5. Table 3 summarizes and tabulates the results gathered.

Fig. 5

The transmission of glass composition 40Bi₂O₃-60GeO₂-xEr₂O

Table 3 The values of Eopt, Tg, Tx, Tc, Tm, S, HR, n, and Ec of crystallization

Table 3 shows that the sample without Er has a higher E_opt value. The value of E_opt decreases with the increase in Er³⁺, which indicates an increase in NBOs (non-bridging oxygen) [23, 30].

As shown in Fig. 2, the glass cuts off the UV and a small strip of the visible region, indicating the possibility of using glass samples as a filter [31].

3.5 Density and molar volume

The glass system’s density ρ and molar volume V_m are vital in determining a deep sight at any structural alterations in the glass network formation, density, and molar volume are estimated using the relationships [32].

$$\rho\,\,\,\, = \,\,\,\,\frac{0.865Wa}{{(Wa - Wb)}}\,\,g\,cm\,^{ - 3}$$

(5)

where W_a signifies the sample's weight in the air, W_b denotes the sample's weight in the immersion fluid, and 0.865 is the immersion fluid's density (toluene).

$$V_{m} \,\, = \,\,\frac{{\sum X_{i} M_{i} }}{\rho }cm^{3} /mol$$

(6)

where ${X}_{{\varvec{i}}}$ and ${M}_{{\varvec{i}}}$ are the concentration and molecular weight of each oxide.

Figure 6 depicts the connection between density and molar volume as Er concentration rises. It detected a reversal of the relationship, the density values decreasing up to 1 mol percent Er₂O₃ and then increasing in the other manner with M_V values. The density changes attributed to the addition of Er³⁺ is due to the breakdown of bonds, and the creation of NBOs, attributed to increased volume as Er ions enter the interstitial site.

Fig. 6

The density and molar volume calculation of glass samples

Another option is that the RE enters the interstitial site generated by the polymerization of glass [20, 33], which is responsible for the compactness of glass, increasing the molecular weight (MW) of Er₂O₃.

3.6 Excitation and emission of glass

Figure 7 shows the excitation spectra of Er-doped glass samples at 550 nm emission. Peaks at 317, 368, 381, 410, 453, and 490 nm describe the transition [34, 35] from the ground state ⁴I_15/2 to the excited states ²P_3/2, ⁴G_9/2, ⁴G_11/2, ²G_9/2, ⁴F_5/2, and ⁴F_7/2. The high intensity of two-level transitions at ⁴G_9/2 and ⁴F_7/2 choose to be the emission energy of the title glass shown in Fig. 8.

Fig. 7

The excitation spectra of Er-doped glass samples at 550nm emission

Fig. 8

The emission spectra of Er-doped glass samples excited at A) 380nm and B) 490nm

It observed the difference in emission energy resulting from different energy peak transitions. The resulting emission spectra showed three peaks at 530, 550, and 639 nm, corresponding to Er³⁺ [12, 23, 25] transitions ²H_11/2 → ⁴I_15/2 (green emission), ⁴S_3/2 → ⁴I_15/2 (green emission), and ⁴F_9/2 → ⁴I_15/2 (red emission).

As seen in Fig. 8, emission peaks have two physical qualities. The splitting of the peak results from the stark splitting. The intensity of peak emission increases up to 1 mol% Er₂O₃ and then decreases. This decreases due to the presence of the quenching process caused by non-radiative resonant cross-relaxation at the ⁴S_3/2 level [12, 23, 36]. The elimination of red emissions is due to this change [25]. The peak is responsible for the blue emission in the 474–490 region assigned to the ⁴F_3/2, _5/2 → ⁴I_15/2 transition in another exciting energy at 380 nm [34, 37].

The color coordinates of the CIE chromaticity diagram use to check the variation in light color according to the emission spectra.

As shown in Fig. 9, the CIE chromaticity diagram could be affected in prepared glass mainly by two parameters, the change of Er concentration and the change of exciting wavelength. The values of chromaticity coordinate listed in Table 4 indicate the samples excited with λ = 489 nm, used in the emission of green laser [36]. And the emission ranged from white to green light when excited with high energy at 380 nm.

Fig. 9

The CIE chromaticity coordinates for BiGeEr glasses under excitation wavelength (a) 489 nm and (b) 380 nm

Table 4 Chromaticity coordinated of glass, glass ceramic samples and decay lifetime

Pandey et al. [37] reported that white light forms from the combination of three colors: blue, green, and red. The sample contains 0.5 mol% close to the expected value of the white light coordinate (0.33, 0.33).

3.7 Thermal properties

DTA curve plays a vital role in the two-point explanation. It explains the glass structure through the T_g value, which is affected by the coordination number, the density of the crosslink, and the bond. strength [11, 38, 39]. On the other hand, it could explain thermal stability. DTA curve characterizes by an endothermic appearance related to annealing, softening, nucleation, and reversible phase transformations determined through T_g and T_m. The exothermic peak represents crystallization and phase change, as defined by the crystallization temperature [40, 41]. The obtained values tabulated in Table 3 show higher values of T_g compared to glass formed as bismuth boro-tellurite [42], TeO₂-Ga₂O₃-GeO₂ system [43], boron bismuth/fluorooxy glass [44], calcium bismuth silicate glass–ceramics [45], Ga-Ge-Sb–Se glass composition [46]. The Tg values increase with the increase of Er concentration as the increase in the compactness of the glass is formed by the presence of Bi in two coordination numbers, BiO₃ and BiO₆ [47], which agrees with the molar volume and FTIR result. Fig. 10 indicates two peaks of different intensities, representing the presence of two crystalline phases, and with the addition of Er₂O₃, the second peak disappears. It interprets as the formation of a single homogenous phase. The thermal stability characterizes by three parameters T_g, the onset of crystallization T_x, and ΔT = T_p-T_g. The higher T_g value could indicate glass has a higher damage threshold due to the transmitted high-power pump, and the higher thermal stability (ΔT is larger than 100 °C) [42, 48] indicate a wide range of working temperatures for the benefit of fiber drawing [49, 50].

Fig. 10

DTA curves of glass samples.

Another parameter is the Hrupy parameter HR = (T_p-T_g)/(T_m-T_p) [38], which comes from a relationship that shows the glass's stability and anti-crystallization ability to form [51, 52].

Relatively such a high value of activation energy for the crystallization, resistance of nucleation and high thermal stability for glass samples under study indicates the glass promises for high-powered optoelectronic device applications [53].

The crystallization mechanism of glass is studied using the kinetic parameters obtained from DTA data. Fig. 11 shows the DTA curve for the glass samples at different rates at 5, 10, 20, and 30 °C/min. Fig. 11 shows that the peak of crystallization temperature becomes sharp and shifts to a higher temperature with the increase in the heating rate [54]. Figs. 12, 13.

Fig. 11

Gass samples of different rates ——5, ——10, ——20, and —— 30

Fig. 12

The relationship between ln (α /Tp) and 1000/Tp

Fig. 13

The relation between ln[-ln(1-x)] and lnα

The activation energy of crystallization can calculate using various approaches [52, 54,55,56]. The activation energy of crystallization glass samples under investigation is determined using the Augis and Bennett method.

$${\text{Ln }}(\alpha /T_{{\varvec{P}}} )\,\, = \,\, - \frac{{E_{{\varvec{C}}} }}{{RT_{{{\varvec{p}} }} }} + Ln k_{0}$$

(7)

where α is the heating rate, Tp is the crystallization temperature, R is the gas constant, and ko is a constant.

The relationship between ln (α /Tp) and 1000/Tp (Fig. 12)  gives a straight line, and from the slope, the value of the activation energy is determined and tabulated in Table 3.

Ozawa equation used to calculate Avrami coefficient (r) according to [54]:

$$Ln \, \left[ { - Ln\left( {1 - x} \right)} \right] \, = \, - r \, ln \, \alpha \, + \, const.$$

(8)

X is the crystallization fraction calculated using the equation X = (A_T / A), and A_T is the partial area at a given temperature. A is the total area of exothermic crystallization [57].

r value determines from the slope of the straight line of the relation between ln[-ln(1-x)] and lnα  (Fig.13). The r values of the prepared glass samples are 1.9, 1.56, 1.56, 1.89, and 1.54. This value indicates that surface crystallization occurred during one-dimensional growth [54, 58].

3.8 XRD of heat-treated samples

Young [59] investigated the heat-treated glass sample (40Bi₂O₃-60GeO₂ mol %) at 521 and 584 °C. The formation of two phases, Bi₂GeO₅ and Bi₄Ge₃O₁₂, can be observed. They also formed glass samples with different concentrations of Bi₂O₃ and GeO₂ containing B₂O₃ or Sm₂O₃ mol %, and by thermally treating them at 630 °C, the crystalline phase of the Bi₄Ge₃O₁₂ cube formed.

The current study comprises adding different amounts of Er₂O₃ mol % to a 40 Bi₂O₃-60 GeO₂ mol % glass sample manufactured and heat-treated at 650 °C. The development of crystals causes bulk samples to become opaque. According to the XRD study, the only Bi₄Ge₃O₁₂ phase developed is isolated tetrahedral GeO₄, and distorted BiO₆ units make up the BGO crystal structure [59, 60]. Fig. 14 shows the XRD patterns at room temperature for materials generated by heat treatments at 650 °C for 5 h. The Bi₄Ge₃O₁₂ crystalline phase (ICDD: 01–074-6243) ascribed to the XRD peaks seen in samples free of Er³⁺ ions and containing Er₂O₃.

Fig. 14

XRD of heat-treated glass samples at 650 °C

The size of particles computed using the Scherrer equation [9] was between 55 and 73 nm. Another approach considers all diffraction peaks, although the findings are relatively similar.

3.9 FTIR of heat-treated samples

FTIR measurements demonstrate any structural changes when the glass samples are converted to glass–ceramic using heat treatment (Fig. 15). As well as compare the results of heat treatment on the glass structure.

Fig. 15

FTIR of heat-treated glass samples at 650 °C

The range of 490–540 cm⁻¹ indicates a significant quantity of crystallization (Bi₄Ge₃O₁₂ phase). [60], which is consistent with the XRD results.

The band at 540 to 830 cm⁻¹ becomes wide and more intense than observed in the glass samples, representing the overlapping structure of the BiO₆ and GeO₄ compositions.

Heat treatment separates the band created in glass samples ranging from 830 to 1100 cm⁻¹ into two bands, broadening and intensifying it. An increase in BiO₆ groups causes a sharp band at 870 cm⁻¹. Bi-O-Bi bonds or the interaction between Bi-O-Ge [1] cause the band at 1110 cm⁻¹.

The formation of symmetric vibrations of hydrogen bond (OH) groups is due to the addition of a modifier or change of the bismuth from the former to the modifier in the glass network, which breaks the bridging oxygen and leads to the formation of NBO atoms. The formation of NBO leads to give vacant oxygen atoms and these oxygen atoms will take part in the formation of OH groups. And the high concentration of OH groups in glass can increase the crack growth, impurities and trapping and leads to a decrease in the overall strength and luminous intensity of the glass. The OH content present in a glass can be evaluated using the following formula [61, 62]:

$$\alpha_{OH} \,\, = \,\,\,\frac{{{\text{ln }}(\frac{{\tau_{o} }}{{\tau_{D} }} )}}{L}$$

(9)

Here, ${\tau }_{o}$ is the maximum transmission, ${\tau }_{D}$ is the transmission at 3000 cm⁻¹ and L is the thickness of the glass sample. The OH content can be calculated according to the following equation:

$${\text{OH content }}\left( {{\text{ppm}}} \right) \, \, = \,{ 3}0{\text{ x}}\,\, \alpha_{OH}$$

(10)

The estimated OH content for the glass samples under study is found between 94 to 102 ppm, which is relatively low when compared with the other glasses such as LBTAF (120 ppm), LHG-8 (773 ppm), GeS2 (175.5 ppm) and KBS (133 ppm) [63,64,65,66].

Hence present systems of glasses are quite suitable for visible photonic applications.

3.10 Excitation and emission of heat-treated sample

Due to the importance of Bi₄Ge₃O₁₂ phase formation in glass ceramics, it uses in laser host crystal and scintillation detector applications [67, 68].

This phase could exhibit green emission [69]. Fig. 16 displays the excitation curve for the glass–ceramic sample doped with 2 mol % Er.

Fig. 16

The Excitation spectra of glass–ceramic doped with 2 mol% Er measured at an emission wavelength of 550 nm

Figure 16 shows the same exciting peaks in the glass sample curve (Fig. 7) in addition to the band at 522 nm. The intense emission obtains by exciting the sample with the energy corresponding to the prominent excitation band. Furthermore, glass–ceramic excited at 380 nm and 522 nm emit green light, whereas glass–ceramic excited at 380 nm emit blue light as shown in Fig. 17A and 17B.

Fig. 17

The emission spectra of glass–ceramic samples at two excitation wavelengths λ_exc at A = 380 nm and B = 522 nm

The CIE chromaticity program shows the emission color. The x–y color parameters calculate and included in Table 4. From Table 4, the emission difference is due to the difference in the emitted wavelengths. Excitation at 522 nm coordinates the green emission chromaticity. The apparent excitation at 380 nm coordinates the blue-green emission chromaticity. Figs. 18A and18B.

Fig. 18

The CIE chromaticity coordinates for glass–ceramic samples under excitation wavelength (A) 380 nm and (B) 522 nm

3.11 lifetime decay

The lifetime is an important factor for potential laser materials. The fluorescence decay curve of Er³⁺ ion excited at 380 nm and monitored at 522 nm at room temperature of the present samples was measured and is shown in Fig. 19. The decay curves have deviated from single exponential, the decay profile is found to be biexponential for all samples except the sample containing 1.5 mol% Er₂O₃ follow single exponential decay. The bi-exponential nature may be because of the reason that the energy transfer from excited donor ion to unexcited acceptor ion transpires. The intensity of luminescence can be articulated as follows [66]:

$$I = I_{o } + A_{1 } \exp \left( { - \frac{t}{{\tau_{1} }}} \right) + A_{2 } \exp \left( { - \frac{t}{{\tau_{2} }}} \right)$$

(11)

where, I₀ is the intensity at t = 0, and I is intensity at time t, A₁ and A₂. are the amplitudes of decay constants and τ₁ and τ₂ are the luminescence lifetimes for fast and slow channels of decay, respectively. The average values of decay time are calculated by using the formula given as [66]:

Fig. 19

The decay curve of Er³⁺ doped bismuth germinate glass–ceramics at room temperature

$$\tau_{avg } \,\, = \,\,\frac{{\left( {A_{1 } \tau_{1}^{2} \,\, + \,\, A_{2} \tau_{2}^{2} } \right)}}{{\left( {A_{1} \tau_{1} \,\, + \,\,A_{2} \tau_{2} } \right)}}$$

(12)

The average lifetime values for Er³⁺ ion-doped glasses are inset in Table 4 and are shown in Fig. 18. It can be seen the average lifetime decreases with an increase in Er³⁺ ion concentration signifying the energy transfer between Er³⁺— Er³⁺ ions [71–73].

4 Conclusion

Bismuth germanium glass systems doped with Er³⁺ ions were studied using different techniques and concluded that:

• 1. The pyramidal [BiO₃] octahedral [BiO₆] units and GeO₄ tetrahedral units are the glass network of samples under study.

• 2. The Ω₂ and Ω₄/Ω₆ values were higher for the sample containing one mol% Er³⁺, which was attributed to the covalent nature and suggested it as a better glass for laser material applications.

• 3. The XRD spectra show the Bi₄Ge₃O₁₂ crystalline phase (50–73 nm size nano-crystals) participating in the glass samples.

• 4. The values of the chromaticity coordinates refer to the samples excited at 489 nm used for green laser emission, while excited at 380 nm used for emission ranging from white light to green emission.

• 5. The luminescence quenching of > 1 mol of Er³⁺ ion concentration occurs under both excitation in the prepared glass–ceramic samples due to the formation of a non-radiative cross-relaxation.

• 6. Based on the structure study, OH content, thermal analysis, emission lifetime and CIE coordinates, it is concluded that the Er3 + ions doped Bismuth germanium glasses are suitable for green and blue optoelectronics device applications.