1 Introduction

Numerous studies on heavy metal-based oxide glasses containing PbO, Bi2O3, and other heavy metals have revealed substantial non-resonant optical infrared transmission up to about 7 μm [1,2,3]. Glass materials have various advantages like as effortlessly molded in multiple shapes, smoothly constructed and manufactured, transparent, to be thermally stable and chemically robust, and cover vast compositional range [4,5,6], as compared to other materials. Due to its higher absorption of dopants in an amorphous state, glassy materials provide new opportunities for the fabrication of compact, high-power devices [7]. Silicates are generally considered to be necessary glassy host materials for a variety of opto-dielectric applications. SiO2 is one of the most complex and diverse material families, occurring naturally as a composite of multiple minerals and a synthesized material (glass former) [7, 8]. The mechanical, chemical, optical, and luminous properties of silicate glasses are outstanding. Silicates are essential components in industrial and domestic glassware [9]. It is employed in structural materials, microelectronics (as an electrical insulator), thermal and dielectric sectors, and components. Nanoparticle-doped silicate glasses can be used for various biomedical applications [10]. High density, optical absorption, and tunable refractive index are the outstanding features of bismuth-based glasses [11]. Bismuth has a number of important applications in glass and glass ceramics, thermal and mechanical detectors, reflecting windows, and optical and optoelectronic devices [12, 13]. Bismuth oxide can be used to synthesize glass materials, but glass formation is not easy. However, other compounds, such as PbO and SiO2, can be used to grow glass materials by a simple synthesis pathway; these materials show structural units that are similar to BiO3 and BiO6. In BiO3, the pyramidal structural unit is represented as a glass former in which the bismuth ion is attached to three oxygen atoms and the outermost lone pair of electrons 6s2 is present at the top. BiO6 octahedral units act as modifiers and induce structural defects in the amorphous form [14,15,16]. It has been predicted that in glasses, bismuth oxide has dual role as a glass network and a modifier and this could be because of high polarizability and smaller ionic size of Bi3+. Bismuth can also exist in different ionic states, such as Bi+, Bi4+, and Bi5+, depending on its concentration and chemical composition [11, 17, 18]. When Bi2O3 is added to a silicate glass host, it transforms into a rigid, stretchy, non-corrosive, and thermally stable material. Because of their significant thermally stimulated luminescence, Bi2O3 dopant silicate glasses are commonly used as dosimeters for radiation therapy and protection. The addition of Bi2O3 to silicate glasses is suggested to enhance strength and provide gamma radiation protection [19]. PbO is another heavy metal oxide with a high atomic number, high refractive index, and low melting point [20]. Lead Oxide can be used as a filler/loading material in the matrix of various materials because it has an octahedral structure in its PbO6 form, which can be used to improve material qualities; however, it has a covalent bond structure in its PbO4 form, which can be used to grow glass materials [21,22,23]. Further, a high atomic weight of PbO and Bi2O3 in glass material expands the FTIR spectrum range and reveals different structural units [24]. Lithium ions, in addition to bismuth, have important applications due to their small size and ionic radii (≈0.76 Å), electro-positive nature, lightweight, ability to be used at high voltage, and high energy density [25, 26]. Lithium containing glasses can be utilized in solid-state lithium batteries and solid electrolytes because they have the highest ionic conductivity [27, 28]. Recently, Menazea et al. have reported the ac conductivity of lithium containing nanocomposite and found the suitability of material to be used in rechargeable battery applications [29]. Glass networks made of Bi2O3 and SiO2 have been employed recently and the properties of these composites have been improved by adding lithium ions into a mixed matrix of Bi2O3 and SiO2 [24, 30, 31]. All of the oxides mentioned above are commonly used to make high-resistance silicate-based compound glasses. Silicate-based glasses have a narrow cut-off wavelength and a large transmitting window, making them ideal for a variety of applications. Also, bismuth silicate glasses have essential applications, such as low-loss optical fibers, optical amplifiers, oscillators and IR transmitting materials [32, 33]. In the recent literature, there are several studies done on the glass systems Bi2O3·SiO2 [34], ZnO·Bi2O3·SiO2 [11], PbO·SiO2 [35], Bi2O3·TiO2·SiO2 [36], Li2O·Bi2O3·SiO2 [37], BaO·Bi2O3·SiO2 [38], Li2O·PbO·SiO2 [24], Fe2O3·Bi2O3·SiO2 [39], SiO2·B2O3·ZnO·Bi2O3 [19], Li2O·ZnO·Bi2O3·SiO2 [40], and Li2O·CdO·Bi2O3·SiO2 [41]. However, work on the physical, structural, and optical aspects of lithium lead bismuth silicate glass system has been not reported. Therefore, in the present research work, we synthesized 30Li2O·20PbO·xBi2O3·(50−x)SiO2 (where, x = 0 to 50 mol%) glasses and investigated the influence of bismuth oxide on the structural features of the samples by employing X-ray diffraction (XRD) and Fourier Transform Infrared (FTIR) Spectroscopy. In order to examine the role of different structural units spectra have been deconvoluted using origin software. A correlation between the physical and structural properties has been made. Further, for more insight optical properties viz., cut-off wavelength (λC), energy bandgap (Eg), Urbach energy (ΔE), theoretical optical basicity (Λth), oxide ions polarizability (α02−), refractive index (n), reflection loss (RL), molar refractivity (Rm), and metallization criterion (M), using the UV–VIS–NIR spectroscopy have also been analyzed.

2 Experimental details

2.1 Synthesis of glasses

The chemicals required Li2CO3, PbCO3, Bi2O3, and SiO2 in the synthesis of desired samples were purchased from high media chemicals with 99.5% purity and analytical-grade mark. The glass samples were made using the melt-quenching technique with the composition 30Li2O·20PbO·xBi2O3·(50−x)SiO2 (where x = 0, 10, 20, 30, 40, and 50 mol%). The prepared mixture of chemicals was placed in a porcelain crucible and then heated in an electric muffle furnace at a temperature of 1150 °C. As obtained melt was occasionally stirred in between during the entire procedure of 30 min and we obtained homogeneous mixture. Finally, the resultant mixtures were splashed onto a stainless steel plate and pressed immediately with another stainless steel plate. The obtained samples were annealed below the glass transition temperature (Tg) for 3 h to remove thermal stress occurred during quenching. The prepared samples were labeled as S0, S1, S2, S3, S4, and S5.

2.2 Characterization techniques

The X-ray diffractograms of powdered samples were recorded at room temperature using a Rigaku Miniflex-II X-ray diffractometer. The glass transition temperature was recorded using a Differential Scanning Calorimeter (Model Mettler Toledo Q20) maintained at a temperature of 10 °C/min. For exploring the molecular bonding properties, Fourier transformation infrared spectroscopy was employed through PerkinElmer (BX-II) spectrometer at room temperature. For the FTIR measurements, acceptable glass powder samples were mixed with KBr in the proportion of 1:20. After preparing the thin pellets, the infrared spectra were recorded quickly to avoid moisture. The optical absorption spectra of the polished glass samples were recorded in the wavelength range of 190–3300 nm using a Shimadzu spectrometer (UV-3600 Plus).

3 Results and discussion

3.1 Structural and molecular bonding measurements

Figure 1a represents the XRD spectra of present glass samples (S0–S5). The XRD patterns of prepared samples were recorded between the angles 10 to 80°. Because of the highly short-range disordering of atoms in glasses, the XRD pattern of all the samples has a large hump around ~ 27° which confirms that the glass samples are amorphous in nature. Also, the variation is observed in the broad hump when loading takes place which confirms the fact of interaction between bismuth and silicate particles.

Fig. 1
figure 1

a XRD patterns of different 30Li2O·20PbO·xBi2O3·(50 − x)·SiO2 glass compositions. b FTIR spectra of different 30Li2O·20PbO·xBi2O3 (50 − x)·SiO2 glasses at room temperature

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As shown in Fig. 1b, molecular bonding measurements were performed using FTIR in the wavenumber range 400–2000 cm−1 at room temperature. This region of mid-infrared spectrum is the fingerprint region, where IR vibrations are most active. The FTIR spectra of present samples contain a number of fundamental bonds that confirm the desired samples growth. Due to the abundance of the heavy metal bismuth oxide (Bi2O3) and modifier cations (Li2O, PbO), the position of absorption bands in the spectra of these glasses is similar to the usual ranges of lithium zinc bismuth silicate and borate glasses [40, 41]. Two absorption bands are identified in the wavelength range 750–1250 cm−1 centered at 981 cm−1 and 400–600 cm−1 centered at 498 cm−1, respectively. From Fig. 1b, it is also observed that in the presence of a modifier, when the SiO2 is replaced by an unusual glass former (Bi2O3) the intensity as well as the position of these bands changes. FTIR spectra in band range from 400 to 600 cm−1 infer that the band presents around wavenumber ~ 425 cm−1 as a result of symmetric oxygen bending-rock mode (R) BO’s bonding and around ~ 450 cm−1 ascribed to the vibration of Pb–O in the PbO4 structural unit [42, 43]. Furthermore, the Li–O–Li and Bi–O bending and stretching vibrations of bonds can be linked to the bands located between 400 and 600 cm−1 [44, 45]. Stalin et al. have assigned the band in this region to Li+ and Bi–O–Bi linkage in BiO6 octahedral unit [46]. Similarly, Kaur et al. have also linked the bands in region < 650 cm−1 to vibrations of the BiO6 and in region 420–460 cm−1 to lithium cation vibration [47]. A sharp peak at 624 cm−1 is also observed behind this range, which shifts toward lower wavenumber when Bi2O3 percentage in glass composition increases and owing to Bi–O stretching vibrations in the BiO6 octahedral unit [48]. A combined broad valley peak is observed in the wavenumber range 750–1250 cm−1. To identify all present peaks in this region, this wavenumber range was deconvoluted into five components using Lorentzian and Gaussian curve fitting for all the samples. It is observed that the area under all of the bands change when the doping concentration of component increases, as shown in Fig. 2a–f. The calculated parameters such as peak position (Xc), amplitudes (A), the full width at half maxima (W), and the corresponding IR band assignments are tabulated in Tables 1 and 2, respectively. Figure 2a shows the deconvoluted FTIR spectra of pure (x = 0) samples, which consists of four peaks, whereas deconvoluted spectra of composite (10 ≤ x ≤ 50) samples contain five peaks indicating the development of a new peak when loading concentration is introduced in the matrix of the pure sample. In the deconvoluted spectra, the band at around ~ 868 cm−1 is attributed with the vibration of PbO6 structural units and Bi–O stretching vibrations of Bi–O bonds in BiO3 units [37, 45, 49], while band ~ 944 cm−1 is associated with O–Si–O bonds in [SiO4]4− units (Q0) stretch asymmetrically without bridging oxygen ions per silicon [35, 50]. Similarly, bands at 1029 cm−1 and 1104 cm−1 can be assigned to asymmetrical stretching vibrations of O–Si–O bonds in SiO4 tetrahedral units in a pure sample. Peaks situated at around ~ 969 cm−1 may be assigned to the combined vibrations of PbO6 structural units and asymmetric stretching vibrations of [SiO4]3− units (Q1; with one bridging oxygen ions per silicon) [51, 52]. Moreover, another band at around ~ 1035 cm−1 is due to the asymmetric stretching vibrations of [SiO4]2− units (Q2; with two bridging oxygen ions per silicon). However, another band observed at around ~ 1055 cm−1 is ascribed to the asymmetric stretching vibrations of [SiO4]− units (Q3; with three bridging oxygen ions per silicon). Also, the band situated at ~ 1105 cm−1 can be assigned to the linkage vibrations of Bi(3)–O–Bi(6) associated with NBOs and can be connected to the asymmetric stretching vibrations of [SiO4] tetrahedral units (Q4; with four bridging oxygen ions per silicon) [52,53,54].

Fig. 2
figure 2

a–f Deconvoluted FTIR spectra of 30Li2O·20PbO·xBi2O3·(50 − x)·SiO2 glasses

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Table 1 Peak position (Xc), Amplitude (A), and full width at half maxima (W) of deconvoluted peaks of FTIR spectra of different compositions of 30Li2O·20PbO·xBi2O3·(50 − x)·SiO2 glass system
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Table 2 Data of FTIR spectra of 30Li2O·20PbO·xBi2O3·(50−x)SiO2 glasses (band position are in cm−1)
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3.2 Physical properties

3.2.1 Density measurements

The density of the prepared samples was determined using the Archimedes principle at room temperature. Xylene was employed as the buoyant fluid in this technique. The adequate molar volume (Vm) for all the samples was estimated using the formula given below [11]

$$V_{{\text{m}}} = M/D,$$
(1)

where M is the molar mass, and D is the density of glass samples (D for xylene is 0.861 g/cm3). The expressive crystalline volume (VC) of samples was computed using the expression [11]

$${V}_{\text{C}}=\sum {x}_{\text{i}}{V}_{\text{i}},$$
(2)

Here xi is the molar fraction and Vi is the crystalline molar volume for each constituents. For Li2O, PbO, Bi2O3, and SiO2 crystalline molar volume is 14.84, 23.42, 52.36, and 22.68 cm3, respectively [22].

Figure 3 illustrates that the density and the molar volume of the studied glasses show an increasing trend with the rise in concentration of Bi2O3. The values of density increase from 3.54 to 5.39 (g/cm3) with the concentration of Bi2O3 (Table 3). Molar volume and crystalline volume show similar variations. The large molecular mass of bismuth (465.98 a.m.u) in comparison to silicate (60.08 a.m.u) entailed this increase in density and molar volume. Similar trends have been observed by Meena for bismuth containing glasses [9]. Also, there is no much variation in the density values after x = 30 shows some sort of structural changes occurring at this compositions. Further, for all glass compositions, Vm values exceed VC values, indicating the existence of excess structural volume, i.e., the potential of glass formation rather than crystallization.

Fig. 3
figure 3

Compositional dependency of density, molar volume, and crystalline volume for all the samples of 30Li2O·20PbO·xBi2O3·(50 − x)·SiO2 glasses

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Table 3 Physical and optical parameters of 30Li2O·20PbO·xBi2O3·(50 − x)SiO2 glasses
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3.2.2 Differential scanning calorimetry (DSC)

The DSC thermographs of prepared glass samples are shown in Fig. 4. From these curves glass transition temperature for each composition was determined and values are listed in Table 3. The intensity of the chemical bond and the density of crosslinks in the glass matrix are hypothesized to affect the glass transition temperature [55]. The glass transition temperature increases with the bismuth concentration at the expense of silicate, as shown in Table 3. From density measurements, it can be predicted that as Bi2O3 content rise structure becomes more denser causing an increase in Tg. A slight decrease in the value of glass transition temperature for the S4 sample is due to the effect of Bi2O3 as a modifier usually weakening the glass structure, also confirmed by the FTIR spectra.

Fig. 4
figure 4

DSC thermographs for different compositions of 30Li2O·20PbO·xBi2O3·(50 − x)·SiO2 glasses

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3.3 UV–VIS–NIR spectroscopy

3.3.1 Optical absorption spectra

In the present study, the optical bandgap and absorption coefficient, α(ν) of glass samples were measured by studying the optical absorption spectra. The Beer–Lambert law is used to compute the absorption coefficient α(ν),

$$\alpha (\nu )= \frac{A}{t},$$
(3)

where A is the absorbance and t is the thickness of glass sample. A relationship between absorption coefficient α(ν) and the role of photon energy (hν) gives the direct and indirect optical transitions as well as optical bandgap energy and is known as Tauc’s relation [56].

$${\alpha \left(v\right)=\frac{B\left(hv-\text{Eg}\right)}{hv}}^{n},$$
(4)

where B is a constant and known as the band tailing parameter. The energy of photon is hν that is incident on glass materials and n is depending on the type of glass transition. The value of n = 2, 1/2, 3, and 1/3 depends upon the electronic transition of the absorption factor. For n = 2 and 3, transitions are considered as indirectly allowed and indirect forbidden, whereas for n = 1/2 and 1/3 transitions are directly allowed and direct forbidden, respectively [57]. Indirect optically allowed transition is possible in solid amorphous glass material. There are three regions found in the absorption coefficient. The first region is the high absorption region which is known as the “Tauc region” and depicted in Fig. 5. The energy gap of optical band was estimated from the linear section of the curve toward the energy axis at \((\alpha h\nu {)}^{1/2}=0\) and obtained values are listed in Table 3. It is clear from Fig. 6, that with an increase in the concentration of Bi2O3 in the glass system, both the optical absorption spectra and the cut-off wavelength are observed to be red shifted. Similarly, the bandgap energy is seen to decrease from S0 (3.11 eV) to S5 (1.92 eV). This trend shows that when the concentration of bismuth ion rises (0 to 50 mol%), the non-bridging oxygen ion (NBO) rises, lowering the bandgap energy. Large difference in Vm and VC values with increase in bismuth concentration also supports this assumption. The spectral investigations also revealed that NBOs are associated with asymmetrical stretching vibration in the SiO4 tetrahedral unit. Because when a metal–oxygen bond breaks, bond energy is released. The highest energy state of the valence band model consists of O(2p) orbitals, whereas the lowest energy level is made up of the conduction band, which is made up of M(nS) orbitals. The non-bridging oxygen atoms have higher energy than bonding orbital. As a result, the increase in non-bridging oxygen (NBOs) concentration results in an increased VBM, lowering the optical bandgap energy [58]. Initially, polarization due to Bi3+ ions in the glass structure is modest at low content of Bi2O3 resulting in a small shift in the bandgap energy Eg. However, after adding Bi2O3, substantial polarization occurs owing to Bi3+ ions, leading to a large decrease in the bandgap energy Eg (Table 3). Thus, Bi2O3 enters the interstitial sites as a modifier upto 30 mol% that produces maximum NBOs. Bismuth ionic bonds are developed with NBOs at the place of covalent bonds. When the bismuth concentration is increased at 40 to 50 mol%, it enters in the network as a former. These results are also supported by the structural variation depending upon the composition in the present glasses.

Fig. 5
figure 5

Tauc’s plot for all the compositions of 30Li2O·20PbO·xBi2O3·(50 − x)·SiO2 glasses for n = 2

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Fig. 6
figure 6

Optical absorbance spectra as a function of wavelength for all compositions of 30Li2O·20PbO·xBi2O3·(50 − x)·SiO2 glasses

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The second region, known as the “Urbach region,” appears as a result of structural disorientation of the materials. It depends on various factors, like thermal vibration in the lattice, temperature, statics, and induced disorder, and photon energy. The slope of the linear region curve drawn between ln(α) and hν, as illustrated in Fig. 7, was used to calculate the Urbach energy. The relationship can be expressed as,

Fig. 7
figure 7

Urbach’s Plot for all the compositions of 30Li2O·20PbO·xBi2O3·(50 − x)·SiO2 glasses

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$$\text{ln }\alpha (\nu ) = \frac{h\nu }{\Delta E} +\text{ constant}.$$
(5)

The obtained values of ΔE are presented in Table 3 and it was concluded that the value of ΔE decreases as the concentration of Bi2O3 content increases and the minimum value of Urbach energy is observed for S5. It may be due to the decrease in broadening that is correlated with the static disorder [59, 60]. The value of ΔE is consistent with the width of the band tail of the electron state. Due to phonon-assisted indirect electronic transition between localized states, the small value of ΔE generates the exponential tail [61]. Similarly, the significant value of ΔE reveals that defects are maximum and reduce the long-range order. In addition, as compared to other compositions, low values of Urbach energy at a high value of Bi2O3 suggest the possibility of long-range order locally developing the defect concentration [62].

The third region in UV spectra arises due to weak absorption. The values of refractive index increases and becomes maximum (Fig. 8) at a concentration of 50 mol% due to the maximum concentration of both modifier and former oxides of lead and bismuth. Also, a decrease in bandgap energy which causes an increase in refractive index is due to electronic band structure, as illustrated in Eq. (6)

Fig. 8
figure 8

The variation of bandgap energy and refractive index with different concentrations of Bi2O3

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$$\left(\frac{{n}^{2}-1}{{n}^{2}+2}\right)=1-\sqrt{\frac{{E}_{g}}{20}}.$$
(6)

3.3.2 Optical parameters

The value of theoretical optical basicity \({(\Lambda }_{\text{th}})\) for the prepared glasses has been calculated by acid–base properties that represent in terms of electron density carries by oxygen using the expression (7) [63, 64].

$${\Lambda }_{\text{th}} = {\Lambda }_{{\text{Li}}_{2\text{O}}}.{\text{X}}_{{\text{Li}}_{2}\text{O}}+{\Lambda }_{\text{PbO}}.{X}_{\text{PbO}}+{\Lambda }_{{\text{Bi}}_{2}{\text{O}}_{3}}.{X}_{{\text{Bi}}_{2}{\text{O}}_{3}}+{\Lambda }_{{\text{SiO}}_{2}}.{X}_{{\text{SiO}}_{2}},$$
(7)

where \({X}_{{\text{Li}}_{2}\text{O}}\),\({X}_{\text{PbO}}\),\({X}_{{\text{Bi}}_{2}{\text{O}}_{3}}\), and \({X}_{{\text{SiO}}_{2}}\) are equivalent fraction of different oxides and \({\Lambda }_{{\text{Li}}_{2\text{O}}},{\Lambda }_{\text{PbO}},{\Lambda }_{{\text{Bi}}_{2}{\text{O}}_{3}},\text{and }{\Lambda }_{{\text{SiO}}_{2}}\) are their optical basicities. The optical basicity values \({\Lambda }_{{\text{Li}}_{2\text{O}}}\) = 0.87, \({\Lambda }_{\text{PbO}}\) = 1.18, and \({\Lambda }_{{\text{Bi}}_{2}{\text{O}}_{3}}\) = 1.19,\({\Lambda }_{{\text{SiO}}_{2}}\)= 0.50 are taken from the literature [65].

Simultaneously, it gives us a relationship between oxide ion polarizability (\({\alpha }_{0}^{2-}\)) and the optical basicity of the oxide medium,

$${\Lambda }_{\text{th}}=1.67\left(1-\frac{1}{{\alpha }_{0}^{2-}}\right).$$
(8)

The value of theoretical optical basicity increases with the concentration of Bi2O3 content, as seen in Table 3. The polarizability of the Bi3+ cation is high, and it has a lone pair in the outermost shell. As a result, when the concentration of Bi2O3 increases, the NBOs and theoretical optical basicity increase. From Table 3 it can be seen that oxide ion polarizability also increases. The electron donor ability of the oxide ions is thought to be much more vital in these samples. The molar refractivity (Rm) is related to the polarizability of constituent ions of the glass and can be computed from the Eg values, using the relation given as [66],

$${R}_{\text{m}}={V}_{\text{m}}\left[1-\sqrt{\frac{{E}_{\text{g}}}{20}}\right]= \left(\frac{{n}^{2}-1}{{n}^{2}+2}\right)\left(\frac{{\text{cm}}^{3}}{\text{mol}}\right).$$
(9)

The value of Rm increases with the increasing concentration of bismuth from 0 to 50 mol% rapidly and the reverse trend is observed in bandgap energy. The molar polarizability (\({\alpha }_{\text{m}}\)) is also related to molar refractivity as given by the relation [67],

$${\alpha }_{\text{m}}=\left(\frac{3}{4\pi {N}_{A}}\right){R}_{\text{m}}.$$
(10)

The reflection loss is also calculated by the given equation [62],

$${R}_{\text{L}}={\left(\frac{n-1}{n+1}\right)}^{2}.$$
(11)

The metallization criterion (M) of oxide based on its bandgap energy is given as [68],

$$M=1-\frac{{R}_{\text{m}}}{{V}_{\text{m}}}.$$
(12)

Metallization criterion provides us the information on nature of the material. When the ratio of Rm/Vm ≥ 1, the material is metallic in nature, and when the value of Rm/Vm < 1 material is of non-metallic nature. Table 3 shows that the reported values of M are less than 1, indicating that our samples are non-metallic in nature and may possess non-linear optical properties.

4 Conclusion

A study on the effect of substitution of SiO2 by Bi2O3 on the physical and structural properties of Li2O·PbO·Bi2O3·SiO2 glasses has been carried out. The diffused XRD patterns ~ 27° confirm the amorphous nature of the as-prepared glass samples. The density, molar volume, and crystalline volume were increased with concentration of bismuth oxide. FTIR structural analysis reveals that Bi2O3 acts as network former and modifier and can exist in the structural units as BiO3 and BiO6. The indirect-allowed optical transition is possible in glass sample. For the present studied glass composition cut-off wavelength increases from 358 to 539 nm and bandgap energy decreases from 3.11 to 1.92 eV due to increase in bismuth concentration that increase the number of non-bridging oxygen ions which result in decrease of bandgap energy. The Urbach formula is used to determine Urbach energy. Smaller value ΔE at high Bi2O3 content shows the possibility of long-range order arising locally as the defect concentration grows. The values of metallization criterion for all the samples are less than 1 (0.395–0.310), which indicate that studied glasses can be explored for non-linear optical applications.