1. INTRODUCTION Rare-earth-doped transparent glasses and glass ceramics are important materials with many applications in photonics such as optical fibers, laser host materials, up-conversion lasers, amplifiers, and three-dimensional displays [1-3]. Although oxide glasses and glass ceramics are characterized by thermal, chemical, and mechanical durability, their high phonon energies (~1100 cm-1) and limited solubility of rare-earth ions have restricted their optical applications [4-8]. On the contrary, apart from the low phonon energies (~500 cm-1) of fluoride glasses, they suffer high devitrification rate during preparation as well as low chemical, mechanical, and thermal stability [9,10]. Since acceptable optical properties are not enough for an optical material, special efforts have been devoted to produce novel glasses based on both oxides and fluorides. Introduction of oxygen to fluoride glasses or addition of fluorine to oxide glasses increases their crystallization rates; however, it is revealed that oxygen can stabilize the amorphous state when added as specific oxides to fluoride melts . Accordingly, these types of glasses, called oxy-fluoride glasses, were introduced which received considerable attention since they benefited from advantages of both fluoride and oxide glasses [11,12]. These new materials were considered as excellent hosts for rare-earth ions; however, the structural defects caused by the amorphous structure can entrap electrons and holes, thus resulting in non-radiative processes . In order to solve this problem, oxy-fluoride glass ceramics were introduced and prepared by Wang and Ohwaki  for the first time. The precipitated crystalline phase in these glass ceramics was PbxCd1-xF2 nanocrystals, and the doped ions (Yb3+ and Er3+) were preferably segregated from the glassy matrix to the crystalline phase . Due to the environmental problems of PbF2 and CdF2, other oxy-fluoride systems were proposed . Among the several already existing systems, the species with MF2 (M= Ba, Ca, and Sr) nanocrystals have become of interest owing to their economical and non-toxic raw materials . Furthermore, trivalent rare-earth ions could be substituted for the divalent alkaline-earth cations and lead to suitable solubility of rare-earth ions in MF2 nanocrystals . BaF2 in its crystalline form exhibits lower phonon energy (~346 cm-1) than crystalline CaF2 (~466 cm-1), and its IR cutoff is placed at longer wavelengths. Consequently, the IR window provided by BaF2 is wider than CaF2 [18,19]. Nevertheless, contrary to the oxy-fluoride glasses and glass ceramics containing CaF2, there are only a few studies on optical properties of the aluminosilicate oxy-fluoride systems based on BaF2 [5,20-23]. Therefore, the present study aims to evaluate the optical properties of new oxy-fluoride glasses of SiO2-Al2O3-BaO-BaF2 system in the presence of different amounts of Bi2O3 additive. The main reason why this additive has been used is the important role of Bi2O3 in increasing the refractive index of the glasses [24,25]. To the best of the author’s knowledge, no report on its effects of the optical properties of oxy-fluoride glasses and glass ceramics has been found. In this regard, oxy-fluoride glasses with different contents of Bi2O3 were prepared using the melt-quenching method. Then, changes in the density, molar volumes, and microhardness of samples were studied. Structural changes and optical properties including transparency in UV-Vis-IR region, Fermi energy, direct and indirect band gap energies, Urbach tailing, and refractive index of glasses were also examined. 2. EXPERIMENTAL PROCEDURE 2.1. Materials, Sample Preparation, and Analyses Oxy-fluoride glasses with chemical compositions of 45SiO2-15Al2O3-25BaO-15BaF2-xBi2O3 (x=0, 1, 2.5, 4, and 6) (mole ratio) were prepared through the conventional melt-quenching method and they were nominated as GBi0, GBi1, GBi2.5, GBi4, and GBi6, respectively. The samples were obtained using high-purity materials such as Al2O3 (101077 Merck) and BaF2 (202746 Sigma-Aldrich). In order to reach high-purity SiO2 (approximately 99.5%), Hamedan silica was leached by HCl and calcined at 800 °C for 2 hours. BaCO3 (513779 Uni-Chem) was also used to supply BaO. In this respect, 30 g of batches were melted in covered alumina crucibles in an electric furnace at 1500 °C for one hour. The obtained melts were poured on stainless steel plates and then, they were pressed by another plate to produce disc-shaped glasses with the thickness of 3-5 mm. To relieve internal stresses, the shaped glasses were annealed at 500 °C for one hour and cooled to room temperature with a controlled cooling rate. To investigate the crystallization behavior and determine the crystallization temperature of samples, Differential Scanning Calorimetry (DSC) was performed (NETZSCH STA 449 F3) at the heating rate of 10 °C/min. In addition, X-Ray Diffraction (XRD) patterns of glasses and crystallized samples were recorded (Philips Xpert MMD system) to identify the amorphous nature of glasses and precipitated crystalline phases in glass ceramics. Vickers microhardness of the glasses was obtained using HV-1000Z Technologies PACE instruments under the load of 1 N for 15 s. The FTIR spectra were recorded by FTIR Tensor 27 Brucker to assess the structural changes. Optical transmittance spectra of bulk glasses in the UV-Vis-IR range of wavelengths were achieved using UV-Vis-NIR Shimadzu 3100 and FTIR Shimadzu 8400S. 2.2. Calculation of Density and Molar Volume Density (d) of a glass is calculated using Equation (1) with considering its weights in air (W1) and water (W2): "d= " "W" _"1" /〖"W1-W" 〗_"2" (1) Obviously, the relationship between molar volume (Vm) and density is defined by Equation (2), as shown in the following: "V" _"m" "=" ∑▒〖"(" "Mi" /"d" ")" 〗 (2) where Mi is the molar mass of component “i” in glass and is equal to that in Equation (3): M_i=C_i A_i (3) where Ci and Ai are the molar concentration and molecular weight of component “i”, respectively. 2.3. UV-VIS Spectra and Optical Constants Measurements 2.3.1. Fermi Energy, Band Gap, and Urbach Tailing For transparent glasses, Fermi energy level (EF) is determined using the Fermi-Dirac distribution function (Equation (4)): K(λ)=1/(1+exp((E_F-E)/(k_B T))) (4) In Equation (4), EF and E stand for Fermi energy and energy of the probing photon, respectively. In addition, kB and K(λ) are Boltzmann constant and extinction coefficient, respectively. This equation can be written as follows: "k" _"B" "T Ln" ("1" /"K" "-1" )"=" "E" _"F" "-E" (5) Moreover, K is calculated as follows: "K =" "αλ" /"4π" (6) where α is the absorption coefficient obtained from UV-Vis spectra. Therefore, plotting K(λ) vs. incident photon energy (hν) and linear fitting of Equation (5) to the linear part of these plots make the calculation of EF possible [26,27]. According to the model proposed by Tauc and Davis-Mott, light absorption by an amorphous material depends on its optical band gap (Eg) and energy of incident photon (hν) [28,29]. This behavior is represented in Equation (7), as shown in the following: "(αhν)=" "β" ^"2" 〖"(hν-" "E" _"g" ")" 〗^"n" (7) where β is a constant and n is an index that exhibits the type of optical transition that takes the values of 2, 3, 1/2, and 1/3 for indirect allowed, indirect forbidden, direct allowed, and direct forbidden transitions, respectively [30-33]. Here, Tauc plots ((αhν)1/n vs. hν plots) are employed, and the linear part of these curves is taken into account to compute the band gap energies. In other words, the band gap energy of a glassy material was calculated using the intercept of the linear part of Tau plot divided by its slope. The disordered structure of amorphous materials is the tailing of electrons density of states into the band gap. The energy of these tails is known as Urbach energy (EU). Equation (8) shows the relationship between the absorption coefficient and EU: α=βexp(hν/E_U ) (8) On the basis of UV-Vis spectra, Ln(α) against hν diagrams can be drawn and EU can be estimated by least square fitting of Equation (8) to these diagrams [34-37]. 2.3.2. Calculation of Refractive Index The refractive indices of the samples in the UV-Vis region of wavelengths were measured using Fresnel equations (Equations (9) and (10)), reflectance, and transmittance spectra: R=(〖(N_t-N_i)/(N_t+N_i ))〗^2 (9) "T=(" 〖〖"2N" 〗_"t" /("N" _"t" "+" "N" _"i" ) ")" 〗^"2" (10) where Nt and Ni are the complex refractive indices of the glass and air, respectively. Moreover, Nt is defined as Equation (11): N= n-iK (11) where K is the extinctions coefficient and n the refractive index. The Reflectance and transmittance spectra of a sample were considered, and a system of two equations and two unknowns was solved by Macleod media, hence the formation of a curve of refractive index vs. wavelength [38,39]. 3. RESULTS AND DISCUSSION 3.1. Density, Molar Volume, and Microhardness Table 1 presents the values of densities and molar volumes of glasses. Incorporation of one mole ratio of Bi2O3 decreased the density from 3.84 to 3.74 (g/cm3); however, higher amounts of Bi2O3 increased the density again. In fact, changes in molar mass and molar volume generate variations in density values. As a result of the enhancement of molar mass caused by increasing the Bi2O3 content, an increase in the density is required. In contrast, in case Bi2O3 plays the role of network modifying, Vm must increase as a consequence of the emergence of more Non-Bridging Oxygens (NBOs) and the density decreases. The calculated value of Vm (Table 1) increase upon adding Bi2O3 and approving the network modifying role of this oxide. The effect of Vm outweighs the molar mass in the case of sample GBi1; for other samples, the opposite holds. Table 1 lists the results from microhardness measurements. Hardness is usually influenced by the introduction of glass network modifiers, which is provoked with the creation of more NBOs and breakup of the glass network . The decreasing trend of the microhardness of the samples with higher mole ratios of Bi2O3 is in accordance with the above-mentioned statement that proves it to some extent. TABLE 1. Some physical properties of glasses with different amounts of Bi2O3 Sample Code d (g/cm3) Vm (cm3) Microhardness Hv (MPa) GBi0 3.84 27.69 720.21 GBi1 3.74 28.41 695.00 GBi2.5 3.77 30.42 677.30 GBi4 3.86 31.05 642.50 GBi6 3.98 31.60 613.10 3.2. Structural Studies As mentioned in the previous section, Bi2O3 acts as a network modifier and affects the glass structure by creating NBOs. To evaluate this claim, FTIR spectra should be studied (Figure 1). All of the samples exhibit three absorption bands at ~440, ~680, and ~970-980 cm-1 which are related to rocking, symmetric, and asymmetric stretching vibrations of Si-O-Si bonds, respectively . The wide band with the highest intensity is composed of three over-lapped peaks corresponding to different vibrational modes of Si-O bonds in all silicate units, i.e., Qn (n=1, 2, 3). The band at 1080-1100 cm-1 is generated by stretching vibrations of Si-O bonds with single NBOs (Q3), and the other one placed at 970-980 cm-1 is attributed to the stretching vibration of bonds with two NBOs (Q2). Finally, the peak at 900-930 cm-1, which is not distinguishable, is created by the stretching vibrations in silicate units with three NBOs (Q1). Further, Q1 represents the stretching vibrations of Si-O-Al in aluminosilicate glasses [42,43]. As demonstrated in Figure 1, the position of this broad band shifts to lower wavenumbers from GBi0 to GBi6 and its maximum value is observed at wavenumbers near Q2 and Q1 species. Therefore, the number of silicate units with more NBOs grew with the addition of Bi2O3. The other band at ~587 cm-1 is also related to Si-O-Al asymmetric stretching vibrations , and the changes of this band confirm the increase in this type of bonds for samples with Bi2O3. Moreover, the broad peak of asymmetric vibrations of Si-O-Si is intensified more than the symmetric vibrations band. All these changes are indicative of the higher numbers of NBOs and disorderliness in the presence of Bi2O3 . There are three weak peaks at 1460, 1640, and 1741 cm-1 resulting from the vibrations of Al-F bonds . In oxy-fluoride glasses, it is preferred that F- ions bond to Al3+ cations instead of Si4+ to decrease the fluorine loss as SiF4 . Of note, in the spectra of glasses with the exception of GBi6 sample, another weak peak is observed at 534 cm-1, which is probably related to Bi-O bonds. Sample GBi6 lacks this bond mainly because the Bi introduced by six mole ratios of Bi2O3 forms Bi˚ particles (colloidal Bi) instead of Bi-O in this glass, which will be discussed in detail in the next section . Figure 1. FTIR spectra of glasses containing different amounts of Bi2O3 3.3. UV-VIS-IR Transmittance Spectra and Evaluation of Optical Properties UV-Vis transmittance spectra are plotted in Figure 2. Although there is not any significant difference between the transmittance of samples, their absorption edge has a red-shift to longer wavelengths. Moreover, absorption edges are not sharply defined due to the amorphous nature of the glasses. These absorption edges can be assumed as Urbach fundamental edges [48,49]. A broad absorption peak at ~400-600 nm emerged only for sample GBi6, and it is usually considered as evidence of Bi˚ particles in glass. In other words, the surface Plasmon resonance of Bi˚ particles is the reason for this absorption peak . Figure 2. UV-Vis spectra of glasses containing different amounts of Bi2O3 Owing to the strong UV absorption rather than visible wavelengths, the extinction coefficient of samples follows Fermi-Dirac distribution function , and K(λ) vs. hν plots (Figure 3) are used to calculate Fermi energies, the results of which are demonstrated in Table 2. For this purpose, the linear part of these plots steeping to lower energies was taken into consideration and the formula of the best fitted line to this part was obtained through linear regression analysis. Then, this formula was substituted into Equation (5) and the Fermi energy at each wavelength was calculated. Higher EF values of glasses compared to kBT indicate that they are insulators and their insulating behavior varies very slightly to the semiconducting behavior upon adding Bi2O3. Figure 3. Extinction coefficient vs. energy plots of glasses containing different amounts of Bi2O3 Tauc plots of glasses are depicted in Figure 4. The direct and indirect optical band gap energies can be obtained by determining the extrapolation formula of the linear part of these plots and dividing the intercept of this line by its slope. Table 2 includes the values of optical band gaps obtained from these plots. Formation of dangling bonds like NBOs with higher Bi2O3 contents reduces the band gap energy by developing the localized states within the band gap and putting the valence and conduction bands closer [50,51]. TABLE 2. Optical properties of oxy-fluoride glasses containing different amounts of Bi2O3 Energy (eV) Sample Code EF Eg (indirect) Eg (direct) EU GBi0 4.307 3.481 3.950 0.170 GBi1 4.108 3.222 3.752 0.193 GBi2.5 3.805 3.007 3.495 0.204 GBi4 3.759 2.970 3.411 0.209 GBi6 3.649 2.958 3.380 0.212 As stated earlier, in amorphous materials, Urbach tails or the tails of the density of states in forbidden gap spark off the short-range order of these materials. This is the reason why Urbach energy is regarded as a degree of crystallinity and orderliness . To calculate Urbach energies, the slope of the linear part of Ln(α) vs. hν plots (illustrated in Figure 5) is reverted, the results of which are listed in Table 2. In the case of adding Bi2O3, Urbach energy increases and this increment is associated with the higher degree of disorderliness caused by increasing the NBOs. (a) (b) Figure 4. Tauc plots for calculation of (a) direct and (b) indirect band gaps Figure 5. Ln(α) vs. energy diagrams for determination of Urbach energy Figure 6 illustrates how Bi2O3 can affect the IR cut-off position and UV-Vis-IR transmitting window. All of these glasses are transparent in UV-Vis (300-1100 nm) region with transmittance of approximately 90%. In the IR region, the transparencies are acceptable and the IR cut-offs are placed at ~4.7-5 μm. Accordingly, from the IR transmittance point of view, the present glasses can compete with some commercial IR glasses like Schott IRG2 and Schott IR 11 . Figure 6. UV-Vis-IR transmittance spectra of glasses containing different amounts of Bi2O3 Generally, light absorption in the IR region is different from that in the UV-Vis region of spectrum. Most optical absorptions in the IR region in glasses result from vibrational transitions. The frequency (ν) of a vibrational absorption is given as follows: ν=(1/2π)√(F/μ) (12) where F is the force constant for bond energy and μ the effective mass [53,54]. According to Section 3.1., the molar mass and number of NBOs increase by introducing Bi2O3. Generation of NBOs indicates the F parameter and absorption in the IR region and shifts the IR cut-off to longer wavelengths . However, in case the amount of Bi2O3 exceeds four mole ratios, transmittance descents and IR cut-off demonstrate a blue-shift again, mainly due to the light scattering by Bi˚ particles in GBi6. The values of the refractive index of glasses as a function of wavelength are presented in Figure 7. In case the quantity of Bi2O3 increases, the higher refractive index ensues. In fact, the refractive index of a glass is determined by the interaction of light with electrons of constituent atoms of the glass. In addition, upon increasing the electron density or polarizability of ions, the refractive index would increase. Furthermore, NBOs are more polarizable than bridging oxygens with the ability to increase the refractive index. The polarizability of cations plays a significant role . Of note, the main reasons for such increase in the refractive index include the high polarizability of Bi3+ cations and creation of more NBOs. Figure 7. Refractive index curves of glasses containing different amounts of Bi2O3 3.4. Crystallization Behavior and Feasibility Study of Transparent Glass-Ceramic Preparation As reported by other researchers, there are usually two exothermic peaks in the DSC thermo-grams of oxy-fluoride glasses based on BaF2 [55-57]. In the DSC results of the present glasses (Figure 8), two exothermic peaks are observed. The first peak at lower temperatures is ascribed to the crystallization of BaF2, and the second one at higher temperatures is related to the crystallization of glassy matrix . According to this figure, the shape of the first peak changes with higher amounts of Bi2O3. The crystallization peak of BaF2 moves to lower temperatures in the presence of Bi2O3, followed by increase in its content since Bi2O3 increases the number of NBOs and enhances crystallization. Bocker and Russel  proposed a self-organizing model for nano-crystallization of BaF2 from oxy-fluoride glasses where a highly viscous layer enriched by network former cations was formed that acted a diffusion barrier hindering the crystal growth. Figure 8. DSC curves of glasses containing various amounts of Bi2O3 In case the number of NBOs in oxy-fluoride glasses increased, the residual glassy matrix became less viscous and ions diffused the barrier easier. Therefore, NBOs facilitate the crystallization of BaF2 and reduce its crystallization temperature . In the XRD patterns of as-made glasses (Figure 9(a)), no diffraction peak is observed since samples are amorphous, hence no unwanted crystallization. Glasses were heat-treated at their first peak temperature for two hours, the XRD patterns of which are depicted in Figure 9(b). Despite the DSC results and our expectations, BaF2 crystals were not precipitated in samples. Heat-treating at higher temperatures could not solve the problem and apparently, the heat-treated samples lost their transparency. Finally, in glasses crystallized at their second peak temperature, BaF2 was obtained in company with BaAl2Si2O8 (Figure 9(c)). Hence, it is assumed that crystallization is a surface crystallization process. GBi4 was chosen to justify this possibility, and a fine sample of it was prepared to compare its DSC results with those of the coarse glass sample (Figure 10). The crystallization peak of BaF2 for the fine sample moved to lower temperature and got sharper, which proved that the crystallization process began from the surface and made it impossible to prepare transparent glass ceramics from the glasses under study. Moreover, hanging the crystallization circumstances has not shown any positive effect on the crystallization of BaF2 in these glasses and preparation of transparent oxy-fluoride glasses with BaF2 nanocrystals . (a) (b) (c) Figure 9. XRD patterns of (a) as-made glasses and glasses heat-treated at (b) first peak and (c) second peak temperature of DSC results Figure 10. DSC curves of fine and coarse samples of GBi4 glass 4. CONCLUSIONS Oxy-fluoride glasses containing different amounts of Bi2O3 were prepared using the melt-quenching method. All of the samples were amorphous without any unwanted crystallization as demonstrated by XRD patterns. Bi2O3 played the role of network modifier, and the value of Vm increased from 27.69 to 31.60 cm3 for samples GBi0 to GBi6. FTIR and UV-Vis spectra proved the presence of Bi˚ particles in the sample with six mole ratios of Bi2O3. By adding Bi2O3, Fermi energy level and band gap energies decreased and the insulating behavior of glasses was mitigated. Urbach energy of sample with more contents of Bi2O3 increased from 0.170 to 0.212 eV due to the formation of more NBOs and increment of disorderliness. Bi2O3 up to four mole ratios increased transmittance and shifted the cut-off to longer wavelengths in IR region. In the case of the glass GBi6, refractive index increased to 1.7 due to the higher polarizablility of Bi3+ ions and enhancement of NBOs. The first exothermic peak in DSC curves related to the crystallization of BaF2 was moved to lower temperatures due to the creation of more NBOs in the presence of Bi2O3. Moreover, transparent glass ceramics were not obtained because of the surface crystallization process. ACKNOWLEDGEMENTS We would like to show our gratitude to Ali Rahimian and Laleh Farahinia, our colleagues who provided insight and experties that assisted this research.