A single and mixed-phases SnO2 (M-SnO2) nanostructures were synthesized by a simple spray pyrolysis method. The nanostructural crystallinity, surface morphology and optical evolution of Ba-doped tetragonal phase SnO2 with different Ba contents were studied by x-ray diffraction, atomic force microscopy, ultraviolet-visible spectroscopy and photoluminescence spectral measurements. The M-SnO2 with orthorhombic as well as tetragonal phases are formed in 6% Ba-doped SnO2 sample and it exhibits the highest average transmittance 86% with blue-shift of the optical band gap. The observed strong red emission at ∼ 615 nm might be encouraging for the implementation of red emission based on Ba-doped transparent conducting electrodes.
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Wide band gap semiconductor oxides, such as tin oxide (SnO2) has attracted extensive attention due to their potential applications in various state-of-the-art optoelectronic devices such as light-emitting diodes , solar cells , gas sensors [3,4] and photocatalysis . SnO2 is categorized by several promising properties for example low electrical resistivity, high optical transparency in the visible region and high chemical stability. These unique properties make it a suitable candidate for transparent conducting electrodes (TCEs) material [6–8]. Furthermore, it exhibits n-type conductivity with a direct wide band gap of 3.60 eV at 300 K [9,10]. The conductivity of SnO2 is due to the intrinsic defects such as oxygen vacancies and tin interstitial which act as donors in the host matrix [10,11]. It should be mentioned that the most common stable phase of SnO2 is rutile type tetragonal structure (T-SnO2) with a space-group symmetry of P42/mnm [12,13]. Along with the T-SnO2, it can also be synthesized in an orthorhombic phase SnO2 (O-SnO2) under high pressure and temperature [13,14].
In comparison to pure T-SnO2 and O-SnO2, the nanostructures with mixed-phases M-SnO2 (i.e., tetragonal and orthorhombic) show distinct physical and chemical properties . Thus, it is important to develop new synthesis approaches for the preparation of mixed-phases SnO2 at ambient pressure and low temperature. In this regard, several synthesis techniques such as pulsed laser deposition (PLD) , chemical precipitation method , sol-gel method , hybrid molecular beam epitaxy  have been used to prepare T-, O- and M-SnO2 nanostructures. Chen et al.  have fabricated O-SnO2 thin films at a pressure of 3×10−2 Pa and substrate temperature of 320°C using the PLD method. Kumar et al.  have synthesized mixed phases SnO2 and single-phase tetragonal structure. They showed that the coexistence of T- and O-SnO2 plays a crucial role in the crystallite size and morphology of SnO2 nanostructures. The M-SnO2 thin films were also synthesized by mist chemical vapour deposition and their electrical properties were investigated by Bae et al. . They showed that O-SnO2 acts as a metal-like transparent conductive oxide. Recently, Carvalho et al.  have investigated the M-SnO2 nanostructures under the doping process with a rare-earth cerium metal. They have obtained a mixture of O-SnO2 and T-SnO2 nanoparticles synthesized by sol-gel route at a temperature of 750 °C. From a different perspective, the nebulizer spray pyrolysis has several advantages including simplicity, low cost, high growth rate and easy to add doping materials .
The improvement of optoelectronic properties of T-SnO2 can be accomplished by the introduction of extrinsic defects through dopants, for instance, divalent or trivalent cations [19–22]. The extrinsic dopants such as indium (In) , antimony (Sb) , strontium (Sr) , neodymium (Nd) etc.  have been used to modulate the physical properties of SnO2. Among these dopants, indium-doped tin oxide is the most common transparent conducting material because of its balanced optical transmittance (T∼ 85–90%) in the ultraviolet-visible (Uv-Vis) region and excellent electrical conductivity (RSheet ∼ 10–15 Ω/sq) . However, indium is a rare-earth material and expensive . Ahmed et al.  synthesized alkaline Sr-doped SnO2 nanoparticles by sol-gel process. They have found a reduction of band gap from 4.14 to 3.71 eV as the Sr doping concentration is increased from 0 to 5%. He et al.  have investigated the effect of Mg2+ substitution on the electronic structure as well as optical properties of SnO2. In this respect, alkaline earth-metal barium (Ba) acts as a cationic dopant in the SnO2 lattice and it exists in nature abundantly. Besides, the incorporation of Ba into tetragonal SnO2 system may give rise to lattice distortion which perhaps produces stresses on SnO2.
The formation of mixed-phases coexisting SnO2 nanoparticles and their gas sensing performance have been investigated to date, however, the structural, optical and photoluminescence properties of the mixed phases SnO2 are still not completely explored [4,13,14,21]. In this work, the nebulizer spray pyrolysis method is used to synthesize the Ba-doped SnO2 nanostructures. This work also aims to provide a comprehensive investigation of the nanostructural, optical and photoluminescence properties of Ba doped T-SnO2 as well as M-SnO2 structures.
2. Experimental methods
2.1 Synthesis of undoped and Ba-doped SnO2 nanostructures
Sn1-xBaxO2 thin films of different composition with x = 0.00, 0.03 and 0.06 were synthesized on the glass substrates at a constant temperature of 300°C by a nebulizer spray pyrolysis method. SnCl2 · 2H2O (purity 99.99%, Sigma-Aldrich) was used as host precursor and BaCl2 · 2H2O (purity of 98%, Sigma-Aldrich) was used as a doping material. Typically, 2.226 g of SnCl2 · 2H2O was dissolved in distilled water to prepare 0.1 M concentration. Then, a few drops of HCl was added to obtain a clear and homogenous solution. For Ba doping, 1.221 g (0.1 M) of BaCl2 · 2H2O was dissolved in water and stirred it for 1 h. The dopant precursor was then added drop wise to the host precursor until total volume becomes 100 ml. The visual appearance of the host solution and dopant solution is presented in Fig. 1. Ba-doped SnO2 thin films with different Ba concentrations of 0, 3 and 6 mol. % were prepared by changing the dopant solution in the mixture. The solution was then sprayed through a nozzle using compressed air at 4.15×105 Pa. The normalized distance between the substrate and spray nozzle was ∼ 3 cm. Finally, the prepared samples were annealed at 400°C in the open air for 3 hours.
2.2 Materials characterization
Structural phase and crystalline quality of the films were analyzed by x-ray diffractometer (XRD) using D2 PHASER (BRUKER) system. The pattern was recorded using CuKα radiation (λ = 0.1541 nm) and the scanning range was between 10° and 80° in the 2θ scale with a scanning step of 0.02°. Crystallite size was calculated from the XRD data by using the Debye–Scherrer (D–S) formula and Williamson–Hall (W–H) method. The morphological analysis of Ba-doped SnO2 thin films was carried out using atomic force microscopy (AFM) (Nanosurf Flex AFM Axiom with C3000 controller) and scanning electron microscopy (SEM) (Philips XL30 FEG). The optical transmittance spectra were measured using an ultraviolet-visible spectrophotometer (J-A Woolson Jewel, M-2000) in the wavelength range of 300–900 nm. To estimate the film thickness, the spectroscopic ellipsometry method was used that has been explained in our previous work . The thickness of the all films was maintained 120 nm (± 15 nm) by controlling the deposition rate of ∼ 10 nm/min. The room temperature photoluminescence (PL) spectra were measured using a Hitachi fluorescence spectrophotometer (Model no F-4600).
3. Results and discussion
3.1 Structural characterization
Figure 2(a) shows the XRD patterns of the undoped SnO2 and Ba-doped SnO2 thin films annealed at 400°C in the open air. With the increase of Ba doping content, the diffraction peaks become sharper, suggesting a better crystallinity of 6% Ba-doped SnO2 sample. All the XRD peaks of undoped SnO2 and 3% Ba-doped SnO2 samples can be indexed to the rutile type tetragonal phase T-SnO2 (ICDD card No. 41-1445). No diffraction peaks corresponding to the formation of other crystalline impurities are detected in the 3% Ba-doped SnO2 sample, indicating that 3% Ba doping does not affect the T-SnO2 structure. However, 6% Ba-doped SnO2 shows a small number of distinct peaks along with the majority of the peaks those corresponds to T-SnO2. The presence of extra diffraction peaks at 2θ = 29.50°, 35.89° and 48.44° (indicated by squares) corresponding to (101), (002) and (200) planes, respectively, can be assigned to the orthorhombic (ICDD 78-1063) crystal structure. The above results suggest that both O-SnO2 and T-SnO2 phases coexist in the 6% Ba doped-SnO2 film. In the previous study, Zhang et al.  have investigated the generation of the O-SnO2 at ambient conditions for the samples doped with 5 mol.% of Mn, prepared by the co-precipitation method.
Figure 2(b) exhibits the shifting of diffraction peaks with the increase of Ba content. The peak position of the planes (110), (101) and (211) for the 3% Ba sample moves slightly towards lower 2θ angles, indicating effective Ba incorporation into the SnO2 lattice . This shift can be attributed to the doping induced lattice deformation and disorder in the SnO2 crystal . Also, the shifting of peak position may be correlated to change in lattice parameters [11,30]. The estimated lattice parameters values are listed in Table 1. For the 6% Ba doped SnO2 sample, the shifting of diffraction peak towards higher Bragg angles may result from the decrease in lattice parameters. The microstrain (ɛ) of undoped and Ba-doped SnO2 samples were estimated by using the W–H equation [31, 32].
The W–H method is generally considered as an accurate method to determine the microstructural parameters. According to the W–H method, the broadening of diffraction peaks is the result of the combined effect of crystallite size and microstrain where the strain is considered to be uniform in all crystallographic directions. The W–H method can be expressed by the following equation:
For 6% Ba, it can also be seen from Table 1 that the crystallite size of T-SnO2 is greater than that of O-SnO2 nanostructured. On the contrary, the value of strain for the O-SnO2 is found to larger compared to the strain of T-SnO2. The similar observations were reported by Carvalho et al.  for the rare earth Ce-doped SnO2 nanoparticles synthesized by sol-gel route. The rise in lattice strain for 3% Ba suggests that majority of the Ba2+ ions replaces Sn4+ ions and this replacement of ions produces defects in the SnO2 lattice [37,38].
Dislocations are regarded as the most significant crystal defects which could be a result of the effect of the internal stresses. The dislocation density (δ) can be determined using the Williamson–Smallman equation [31,39], which is followed as δ = 1/D2, where D is the value of crystallite size. The estimated dislocation density for all the samples is displayed in Table 1. The dislocation density of T-SnO2:Ba 3% is found to increase remarkably. This result can be attributed to the deterioration of the crystallization process. Hence, it may be noted that 3% Ba reduces the crystallite size and increases the microstrain and dislocation density as well as the number of grain boundaries which may be due to the increment in point defects such as oxygen vacancies [25,34].
3.2 Preferred orientation and texture coefficient
The texture coefficient (TC) was estimated to quantify the preferred crystallographic orientations of the material. The texture coefficient can be calculated according to Harris’s method :
3.3 Morphological studies
Figures 4(a-c) show the AFM images (10 × 10 µm2) of the Ba-doped SnO2 thin films grown with different Ba contents. For the SnO2 sample, the grains are almost ellipsoid-like, roughly uniform in shape and have sizes of ∼ 150 nm, as depicted in Fig. 4(a). The surface morphology of SnO2 samples is dependent on Ba doping content. For the Ba-doped SnO2 samples, spherical-like grains are distributed over the surface of the films. As shown in Fig. 4(c), the spherical grains of M-SnO2:Ba 6% sample are more distinct and uniformly distributed over the surface. Nevertheless, the average grain size increases in the range of ∼ 180-200 nm. This change in grain size may be due to the enhancement of grain growth resulting from the reduction of stress and distortion . The grain size of the AFM measurement is much higher than the crystallite size obtained from XRD analysis. This result likely suggests that the grains are not single crystallite; nonetheless, they can be composed of many crystallites . In AFM measurement, grain sizes are usually determined by the distances between the grain boundaries, while XRD generally measures the crystallite sizes which are the coherent diffracting crystalline domains. Surface morphology has been further studied by using SEM technique.
Figures 5(a-c) show SEM micrographs of Ba-doped SnO2 samples with a variation of doping concentrations. The micrographs exhibit a smooth surface morphology with appearing few voids. The surface morphology of the Ba-doped SnO2 samples is somewhat different from the surface of undoped SnO2 regarding the size and shape of the grains.
3.4 UV-Vis spectroscopy
The optical transmittance spectra of the undoped and Ba-doped SnO2 samples are illustrated in Fig. 6(a). In the visible region, the undoped SnO2 sample and the film deposited at 3% Ba show an average visible transmittance (AVT) of ∼77% and ∼ 62%, respectively. On the other hand, M-SnO2:Ba 6% sample exhibits the highest AVT of 86%. This enhancement of transmittance might be due to grain size growth as observed in the XRD and AFM analysis, resulting in less scattering effects from the grain boundaries [30,41].
As shown in Fig. 6(a), the undoped and 3% Ba-doped SnO2 have a peak at ∼720 nm. This peak can be associated with structural defects absorption. It is also observed that the transmittance threshold moves towards higher wavelengths region (red shift) at 3% Ba while the transmittance threshold for the 6% Ba-doped SnO2 is found to blue shift or increase in band gap. Similar red shift in absorption edges of tetragonal phase SnO2 with Ba doping synthesized by spray pyrolysis method was also reported by Bannur et al. . The previous investigations showed that the optical characteristics near the absorption edge are sturdily affected by the quality of films, which combined several factors such as chemical composition, crystallite size, lattice strain and the number of structural defects [44, 45]. Red shift in the absorption edge with a decrease in particles size arises due to the defects. Blue shift in the absorption edge might be due to a decrease in lattice strain of the M-SnO2:Ba 6% sample .
At the near band-edge UV region (310-335 nm), all the samples exhibit almost zero transmittance, which is the result of strong absorbance in this region. The optical band gap of the samples was estimated from the absorption coefficient (α) using Tauc relation for the direct type transition :6(b-d) show the Tauc plot of (αhν)2 versus hν for the SnO2:Ba samples with various Ba doping concentrations. For the direct allowed transition, extrapolating the linear portion of this plot to the zero absorption provides the value of band gap. The optical band gap is estimated to be 3.85 eV, 3.79 eV and 3.94 eV for the undoped T-SnO2, T-SnO2:Ba 3% and M-SnO2:Ba 6%, respectively. The estimated band gap of the sample at 6% Ba exceeds the reported values of doped SnO2 nanostructures [29,38]. We believe the substitution of 6% Ba improves the crystallinity of the SnO2 nanostructures, declines the fraction of grain boundaries as well as causes a decrease of defect states near the band gap, and consequently increases the band gap. These tunable optical properties are particularly important for applications in solar cells for which SnO2 is used as a TCE.
The Urbach energy (EU) is associated with the width of the tail of localized defect states in the band gap and it can be determined by the following equation:7(a). Below the optical band gap region, the inverse of the slope of the straight line provides the value of the Eu and the variation of Eg and Eu as a function of Ba content are shown in Fig. 7(b). It can be observed that the T-SnO2:Ba 3% sample decreases the band gap, thus there is an increase in Eu. This result can be associated with the increase in the defect centers in the sample which is roughly further verified by the PL spectra.
3.5 Photoluminescence properties of the SnO2 and Ba-doped SnO2 samples
Photoluminescence spectroscopy is an important tool to determine the defect states in SnO2. Figure 8 shows the room temperature PL spectra of the SnO2:Ba thin films in the wavelength range of 450–750 nm with an excitation wavelength of 300 nm. There are three distinct emission bands in the PL spectra: a blue emission at ∼ 470 nm (2.64 eV) nm, a high-intensity red emission at ∼ 615 nm (2.02 eV) and a relatively weak green emission at ∼523 nm (2.37 eV). The red and green emission peaks were also investigated by Ling et al.  for the SnO2 nanoparticles synthesized via vapour-phase transport method. Since all emission peaks at 2.64 eV, 2.02 eV and 2.37 eV are much lower than the band gap of SnO2, the visible emission cannot be attributed to the direct recombination of a conduction electron in the Sn 4p band and a hole in the O 2p valence band . However, these emissions are usually assigned to the crystal defects or surface defect including oxygen vacancies, tin vacancies and tin interstitials in the samples . Mishra et al.  have also reported that the visible emissions of SnO2 nanoparticles in the range of ∼ 400-500 nm occur due to the formation of doubly ionized oxygen vacancies (Vo++). It is likely that the photogenerated holes in the valence band are first trapped at the oxygen vacancy sites. These defect sites trap holes then recombine with electrons in (Vo+ centers to form Vo++ centers . Hence, the direct recombination does not produce PL emission.
The observed blue emission at ∼ 470 nm can be attributed to the recombination of a Vo++ center with an electron from the conduction band . The red and green emission band is commonly observed in the PL spectra and they may coexist in the sample. A strong emission band around at 615 nm can be associated with the deep trapped states forming the defect energy levels inside the band gap [1,49]. It can also be seen from Fig. 8 that the intensity of PL emission and peak broadening is found to be maximum for the 3% Ba doped SnO2 sample. These results may be attributed to the increase in defects such as oxygen vacancies [52, 53].
Previous articles reported that the substitution of divalent/trivalent ions in the SnO2 lattice increases the number of oxygen vacancies . As mentioned in XRD analysis, the 3% Ba-doped SnO2 sample decreases the crystallite size of SnO2 sample as well as increases the strain and dislocation density, which may be due to the increment of point defects. Thus, a significant amount of oxygen vacancies defects can be formed on the material surface during the growth process. These oxygen vacancies interact with interfacial Sn vacancies, producing oxygen vacancies related trapped states within the band gap; this result gives rise to the PL intensity [52–54].
In summary, the nanostructures of the Ba-doped tetragonal phase SnO2 were synthesized by a nebulizer spray pyrolysis method. The prepared samples were annealed at 400°C in air. The XRD results reveal that the undoped and 3% Ba-doped SnO2 thin films are the rutile type tetragonal phase SnO2. However, the obvious reflections peaks in the XRD pattern suggests the presence of orthorhombic phase SnO2 in 6% Ba-doped SnO2 sample. AFM micrograph shows the ellipsoid shaped morphology for the undoped sample, it varies with increasing the doping content. The optical and photoluminescence properties of the undoped and Ba-doped SnO2 samples exhibit a correlation with the strain and crystallite size. The mixed-phases SnO2 sample shows the highest crystallinity with a preferential orientation along the <110> direction. Also, it exhibits the highest AVT of 86% with a wide band gap of 3.94 eV. The Urbach energy results hint the existence of oxygen vacancies defects in the system, which is also roughly verified by using PL spectra.
University Grants Commission of Bangladesh; Rajshahi University of Engineering and Technology research system (Project No. DRE-6-RUET-258-7).
The authors acknowledge the Department of Materials Science and Engineering, City University of Hong Kong, Hong Kong SAR, P. R. China for providing the lab facilities (X-ray diffraction and spectroscopy ellipsometry).
There are no conflicts of interest to declare.
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