Contribution of oxygen-vacancy defect-types in enhanced CO2 sensing of nanoparticulate Zn-doped SnO2 films

Author’s Accepted Manuscript Contribution of oxygen-vacancy defect-types in enhanced CO2 sensing of nanoparticulate Zn-doped SnO2 films S. Deepa, K. Prasanna Kumari, Boben Thomas www.elsevier.com/locate/ceri

PII: DOI: Reference:

S0272-8842(17)32063-1 http://dx.doi.org/10.1016/j.ceramint.2017.09.134 CERI16305

To appear in: Ceramics International Received date: 2 July 2017 Revised date: 28 August 2017 Accepted date: 16 September 2017 Cite this article as: S. Deepa, K. Prasanna Kumari and Boben Thomas, Contribution of oxygen-vacancy defect-types in enhanced CO2 sensing of nanoparticulate Zn-doped SnO2 films, Ceramics International, http://dx.doi.org/10.1016/j.ceramint.2017.09.134 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Contribution of oxygen-vacancy defect-types in enhanced CO2 sensing of nanoparticulate Zn-doped SnO2 films Deepa S., PrasannaKumari K, Boben Thomas* Post Graduate and Research Department of Physics, Mar Athanasius College (Autonomous), Kothamangalam- 686666, Kerala, India

*

Corresponding author. Tel.: +91 9447329620. [email protected]

Abstract In spray deposited nanocrystalline SnO2 films, large number of defects involving vacancies and interstitials exist. Doping SnO2 lightly with zinc introduces new Sn2+ states and in fact more oxygen vacancies for charge compensation, which promotes the increase in concentration of chemisorbed oxygen. As the deposition temperature is higher, the XRD and HRTEM results exhibit a clear preferred (110) orientation. The Zn2+ doping in SnO2 introduces new surface and structural oxygen defects, as apparent from Raman and Photoluminescence analysis. The doping alters the polarization of atoms, as seen from the widening in Raman lines. Decreased UV emission peaks (~333 nm) can be attributed to enhanced crystallization and the incited oxygen vacancy (OV) defects. A dominant blue-green luminescence peak centered at 495 nm reveals the formation of deep trap, due to in-plane OVs. The red-shift in PL broad yellow-shoulder at 556 nm features an increase of deep oxygen defect energy levels occurring on the surface of the nanocrystal films. On doping, involvement of the defect composition and the contribution of different kinds of OVs play important roles in enhancing CO2 sensing response of 94.4% to 500ppm. In 2.4 wt% Zn doping at 310 °C, the in-plane OVs are found to dominate over bridging OVs.

Keywords: SnO2 nanoparticles; Gas sensors; Spray pyrolysis deposition; XPS; Raman spectra; Photoluminescence. 1. Introduction Semiconductor gas sensors function as a compact device to detect harmful, inflammable or toxic gases like CO2, LPG, H2 and CO. In that, metal oxide materials like SnO2 are used as gas sensors 1

because of their chemiresistance behavior [1-3]. Oxygen vacancies (OVs) and surface defects in oxide semiconductors have become an increasingly active area of research, connected with the broadening of awareness in gas response process of sensors [4]. Extensive research has been conducted in understanding the chemical and electronic mechanisms that control over SnO2 sensor performance, to extend the sensors for the detection of trace amounts of toxic or pollutant gases. In the case of thin or thick film material, the sensor performance depends on various factors such as doping, surface modification, exposed crystal planes, grain size, working temperature, and so on. For the last few years, a variety of different materials such as Cs, Mn, Cu, Co, Fe, Mg, Cd, Pb, and recently Zn also have been successfully doped in SnO2 films to improve the sensing parameters [5-10]. It has been evidenced that nanostructures provide an enhancement in both active surface area and analyte diffusion, so that the sensor exhibits better with quicker response and recovery properties. This is because the exclusive properties of these materials are completely different from their conventional polycrystalline counterparts, caused by their typical structure. In spray deposited nanocrystalline SnO2, large number of defects such as vacancies and interstitials exist in the material films. However, the involvement of the defect composition and the contribution of different kinds of OVs on gas sensing properties of SnO2 are much less investigated. It is well known that the conductivity of SnO2 depends on OVs that act as donors. On doping SnO2 lightly with Zn, the oxygen vacancies seem to be partially filled, thereby reducing anion OVs in SnO2, causing in an increase in electrical resistivity and activation energy for conduction. Investigations on the effect of doping, deposition temperature, grain size, defects and porosity offer important information about electrical conduction based on localized electric charge carriers in nanostructured materials. In this article, we report on the gas sensing performance of Zn-doped nanoparticulate SnO2 films by evaluating the effect of Zn in developing structural or oxygen-vacancy defects, to identify the oxygen-vacancy led sensing action. 2. Materials and methods Spray pyrolysis technique is used to prepare the pure and zinc doped SnO2 (Zn-SnO2) thin films. A solution of SnCl4.5H2O (Sigma-Aldrich 99.99%) with 20 ml of isopropanol (Merck GR grade) is prepared in 20 ml of distilled water. To this precursor solution, citric acid and then ethylene glycol (Merck GR grade) are added and stirred under closed condition, over night at 45 °C. The 2

solution thus prepared is used for spraying on to a preheated substrate at different temperature. Zinc (Zn) doping in SnO2 films is achieved by way of varying the amount of zinc acetate dihydrate in the spray solution. Crystallographic measurements are performed on Rigaku MiniFlex 600 X-ray diffractometer using Cu Kα radiation (λ = 0.15418 nm) with step size=0.5° and scanning rate of 0.0167°/sec. The changes in morphology and crystallinity of the samples are determined by Field Emission Scanning Electron Microscopy (FESEM) by Carl Zeiss SIGMA HV and Transmission Electron Microscopy (TEM) by JEOL JEM 2100. The X-ray photoelectron spectroscopy (XPS – Kratos Analytical – Axis Ultra) analyses are carried out with a monochromatic Al Kα source (hν = 1486.6 eV, 5 mA, 12 kV). The high-resolution narrow-scan spectra are recorded at pass energies of 10 eV (step 0.05 eV). Any charging shifts (charge neutralizer filament bias- 1.3V) produced by the samples are carefully removed using C 1s binding energy (BE) of the adventitious carbon line [10] at 284.66 eV. The error in all the measured BE values is within ±0.05 eV. The Ag 3d energy peak at 368.25 eV is used to calibrate the BE scale of the spectrometer. Instrument base pressure is 7 x 10–9 Torr. High-resolution spectra obtained are smoothed using Quadratic Savitzky-Golay method and subjected to non-linear least-squares curve fit by so-called Voigt profile, performed with Vision processing® software. The analysis of spectra from micro-Raman (WITec TS-150) and Photoluminescence (HORIBA Scientific Flurolog TCSPC) spectrometers are utilized to investigate the oxygen vacancy in the sensor samples. The films pyrolysed on a glass substrate are studied for gas sensing by loading the films in an in-house made set-up using quartz tubular furnace with probes attached [11] in Van der Pauw configuration. The samples are heated for sensing measurement by means of a PID control in the temperature range of 200–350 °C. The Keithley current-voltage source meter (model-2400) with LabVIEW-based user interface is employed for the resistance measurements. The test gas along with air flow is then monitored through gas flow meters and brought into the chamber in the desired concentration. The dynamic resistance is then measured by switching on and off the test gas. The CO2 sensor response (Sr%) is measured as the ratio of, change in resistance (ΔR) under air- CO2 mixture to that of the initial resistance in air (Ra) multiplied by 100, by Sr% = (ΔR/Ra)*100

(1)

3. Results and Discussion 3

3.1 Structural and Morphological studies The deposited films are visually uniform without any powdery feature or fogginess on the surface. The structural and morphological aspects of spray deposited Zinc doped tin oxide (SnO2) thin films, prepared from SnCl4.5H2O precursor have been studied in comparison to pristine films. 3.1.1 XRD Fig.1 (a) shows XRD patterns of pristine SnO2 (0% Zn dopant), 2.4 wt%, and 3.6 wt% Zn doped SnO2 prepared at 270°C (‘L’ - in the sense, lower temperature), whereas Fig.1 (b) presents the same compositions at 310°C (‘H’ - in the sense, higher temperature). The XRD pattern matches with the bulk JCPDS card No.21-1250 (cassiterite with P42/mnm space group), confirming that the material consist of tetragonal SnO2. No additional reflections related to impurity or dopant segregation are observed within the limit of instrumental sensitivity, which suggests that the materials are phase-pure and the feasible entry of zinc ions within the tin oxide lattice.

Fig. 1. XRD Spectra of Pristine, 2.4 wt.% and 3.6 wt.% Zn doped SnO2 thin films prepared at (a) 270°C and (b) 310°C. From the XRD spectrum, the average size of SnO2 crystallites are calculated from the (110) and (101) diffraction lines based on the Scherrer’s equation [12], which is comparable with the TEM results and given in Table 1. It is also observed that the XRD (110) peak intensities increase with deposition temperature, indicating their higher crystallinity. Depending on dopant amount in conjunction with formation temperature, Zn incorporation in the SnO2 film inhibits the growth of crystal grains within a range of concentration. On Zn doping, all the three lattice 4

parameters are indeed disturbed wherein, the lattice strain also is recorded a variation as evident from the Table1. When Zn2+ (ionic radius ~ 0.074 nm) replace Sn4+ (ionic radius ~ 0.071 nm) in SnO2 lattice, the difference in ionic radius provokes the lattice distortion which is likely to produce more oxygen vacancies for charge compensation. Except in the case of formation of composites or phase segregation, the doping with external atoms in a host material does not result in the emergence of new XRD peaks. However this leads to a gradual shift in the lattice parameter of the host material, as the dopant concentration is increased from the dilution limit. So upon doping, a shift arising from the reshuffle of the lattice parameters takes place to balance the change in the charge field. Fig. 2 shows a linear decrease in the ‘a/c’ lattice ratio of cassiterite films with increasing Zn2+ content at 310 °C, but for 270 °C of deposition, where it is lower than the bulk ratio 1.486 which is reaching a minimum around 2.4 wt% doping.

This lattice

deformation seems to be influenced by the distribution of local defects such as oxygen vacancies due to dopant incorporation, and thereby playing a crucial role in controlling the gas sensing properties of Zn doped SnO2 films.

Fig. 2. Lattice constant ratio a/c (in left axis) together with Crystal cell volume (in right axis) against Zn doping quantity in wt%. In films prepared at 270°C and 310°C, the cell volume calculated from XRD is obviously higher than the bulk value 71.566 Å3, for all the doping concentrations but for undoped films deposited at 270°C, where it remains in a narrowly lower value. The acceptor doping with (Zn2+) brings about charge compensation by forming oxygen vacancies in the lattice which creates electric dipoles. These dipoles of Zn2+ ion have effective negative charge and oxygen vacancies have positive charge. At the Sn4+ site, zinc sits in a divalent oxidation state forming Zn2+, causing 5

a swelling of the interlayer space between adjacent Sn–Sn chains due to the slightly bigger ionic size of Zn2+ than Sn4+. This accounts for the enhancement in unit cell size estimated in the XRD. A. K. Singh et al gave evidence that tin interstitials (Sni), oxygen interstitials (Oi) or doubly charged oxygen vacancies (VO2+) increase the size of the lattice, while neutral oxygen vacancies (VO0) induces a reduction in lattice size [13]. This is because the nearest-neighbor Sn atoms equilibrium Sn-O bond length increase outward by 5.6% and 10% for VO+ and VO2+, respectively, while relax inward by 2.5% in the VO0. It is observed that for the sensor response maximum (see section 3.3.1), the cell volume and dislocation density to be an optimum lower value. It has also noticed that for a nanocrystal size of nearly 11 nm, optimum number of oxygen vacancy formation takes place as a result of incorporation of the zinc in 2+ states, to record a peak sensor response. Table 1 Sensor response along with structural parameters (lattice constant, grain size and strain) of undoped and Zn doped SnO2 films. The quantities given in parentheses are the determined standard deviations, which derive clearly from the least-squares calculations. Sensor Crystall Δa (Å) Sample Δc (Å) a/c value Cell Volume Strain Stac Dislocatio determine determined Description Respon ite Size king n Density se (%) from (Å)3 2 d Fault (lines /m ) XRD (nm)

undoped at 270°C 2.4 wt.% Zn at 270°C 3.6 wt.% Zn at 270°C Undoped at 310°C 2.4 wt.% Zn at 310°C 3.6 wt.% Zn at 310°C

43 (0.2) 65 (0.3) 76 (0.3) 39 (0.8) 94.4 (0.04) 65 (0.34)

5.73 (1.3) 8.58 (0.6) 8.18 (0.9) 12.04 (0.8) 10.9 (0.8) 10.4 (0.8)

(Standard=

1.486)

–0.05

–0.003

1.4719

–0.0078

+0.0615

1.4557

–0.0029

+0.021

1.4756

+0.0735

–0.035

1.526

+0.0375

–0.0034

1.4996

–0.011

+0.0547

1.4577

(Standard =71.566

71.1774 (0.0230) 73.2771 (0.0393) 72.5209 (0.0311) 72.1592 (0.0408) 72.1938 (0.0160) 72.9916 (0.0293)

0.0224 ± 0.02 0.0149 ± 0.02 0.0159 ± 0.03 0.0118± 0.03 0.0125± 0.03 0.0150± 0.03

0.012 79 0.008 66 0.009 21 0.006 90 0.007 2 0.008 54

3.8876 E16 1.8499 E16 2.0755E16 1.1564E16 1.2795E16 2.0253E16

The dislocation densities of the films are provided by the Williamson and Smallman’s relation [14]. δ = n/D2

(1)

where δ is dislocation density defined as the length of dislocation lines per unit volume, n is a factor equal to unity at the minimum dislocation density and D is the crystallite size. For doped

6

films with higher response, it is found that the dislocation density is generally in an intermediate range (1.25 x1016 lines m–2). 3.1.2 Texture The grain growth process is directly involved with the reduction of surface energy. Leite et al [15] have determined the specific surface energy, and found that the surface energy of the (001) surface is 1.533 times higher than that of the (110) surface, and that of the (101) surface is 1.192 times than that of the (110) surface. The texture determination shows that at a temperature of 270°C, the grains predominantly grow in preferred (110) orientation which increases further at 310°C. But (101) orientations are also visible irregularly, as the surface demand only slightly higher energy than the (110) direction, to grow in. As the Zn content in the film improves, the XRD pattern exhibits a strong preferred (110) orientation as seen from Table 2. Table 2 Texture coefficients of different planes in doped along with undoped SnO2 films, deposited at temperature 270°C and 310°C. Number in the bracket is the standard deviation in the evaluated data. Substrate Temperature undoped at 270°C

110 1.299(5)

TC of prominent planes 101 200 211 0.854(7) 0.976(1) 0.869(5)

2.4 wt.% Zn at 270°C

1.431(1)

0.569(9)

1.399(7)

0.599(1)

3.6 wt.% Zn at 270°C

1.935(7)

0.485(7)

0.986(2)

0.592(2)

Undoped at 310°C

1.992(1)

0.238(9)

0.980(6)

0.788(3)

2.4 wt.% Zn at 310°C

2.148(1)

0.315(6)

0.849(2)

0.686(9)

3.6 wt.% Zn at 310°C

2.721(5)

0.154(5)

0.641(6)

0.482(3)

3.1.3 FESEM The surface topography of the films is examined in fairly high field depth, by scanning electron microscopy (SEM). The films are polycrystalline in nature and the agglomeration depends on the temperature of deposition and doping. Fig 3 (a & b) show the FESEM image of undoped films prepared at 270°C and 310°C, respectively. The images shown Fig 3 (c & d) display 2.4wt.% Zn doped, whereas Fig 3 (e & f) display 3.6wt% Zn doped SnO2 thin film prepared at 270°C and 310°C, respectively. In all the films, except 3.6wt% Zn at 310°C, the grains are spheroid in shape and are uniformly distributed over a wide region. On close examination, it can be evinced that the grains are composed of agglomerates formed by a superposition of nanoscaled particles. The voids are also meagerly visible. It is observed that at a deposition temperature of 270°C, the 7

average particle size is about 5- 9 nm whereas 10-12 nm at 310°C. Formation of smaller crystallites in lower temperature deposited samples is due to the enhancement in the density of nucleation centres [16]. The micrographs of 2.4 wt.% doped films formed at a higher temperature of 310°C show distinct variation from pristine films. The micrograph of 3.6 wt.% doped samples at 310°C show considerably agglomerated crystalline particles with elongated grain shape having an average diameter of about 10.4 nm, which is slightly less than that of pristine films.

8

Fig. 3. FESEM picture of undoped SnO2 film at (a) 270°C and (b) 310°C; 2.4wt.% Zn-doped SnO2 film at (c) 270°C, and (d) 310°C, and of 3.6wt.% Zn-doped SnO2 film at (e) 270°C, and (f) 310°C. 3.1.4 HRTEM In the High Resolution Transmission Electron Micrograph (HRTEM) images, nanoparticles are clearly visible without any indication of segregation of dopant ions. The images of 2.4 wt.% Zn doped SnO2 films at 270°C (Fig. 4a) shows roughly spherical agglomerated nanoparticles in an amorphous background with an average diameter of about 8.5 nm. The particle size is slightly larger (~ 10 nm) in 310°C deposited films (Fig. 4b) in contrast with films at 270°C, but less than undoped films (~ 12.04 nm, from XRD) at same temperature. Zn is looked upon as a dopant which has an ability of suppressing the surface diffusion process to serve as a structure-directing agent or acts as a growth inhibitor to a certain level [17]. This observation is already reported in literature for Nd-doped SnO2 [18] and Fe-doped SnO2 [19] and is usually ascribed to a reduced surface energy and thus a decreased particle growth [20]. An HRTEM together with SAED study, that has been carried out is consistent with the phase analysis.

(a)

9

(b) Fig. 4. HRTEM micrograph of 2.4 wt.% Zn doped SnO2 thin film deposited at ( a) 270°C and (b) 310°C, along with the corresponding SAED pattern. The lattice fringes are observed in 2.4 wt.% Zn doped films at 270°C, using the high resolution imaging mode. The interlayer spacing is approximately 0.333 nm, corresponds to tetragonal rutile (1 1 0) crystal plane and for the same doping concentration at 310°C, the lattice spacing is about 0.328 nm (left panes in Fig.4) corresponding to the same plane. However, lattice spacing 0.267 nm matching to (1 0 1) crystal planes are also infrequently visible. In these doped films, the texture and particle size results obtained in TEM, agree well with the x-ray diffraction results. Rings observed in the selected area electron diffraction pattern (right panes in Fig.4) demonstrate explicitly, the randomly oriented polycrystalline SnO2 nanoparticulate films. The Electron Diffraction positions corresponding to lattice planes obtained by SAED are consistent with the results of the XRD study. 3.2 Oxygen stoichiometry by XPS

10

Fig. 5. XPS O 1s deconvoluted spectra of (S) SnO2 powder, (a & b) Pristine SnO2 thin films prepared at 270°C & 310°C respectively , (c & d) 2.4 wt.% Zn doped SnO2 thin films prepared at 270°C & 310°C respectively, and (e & f) 3.6 wt.% Zn doped SnO2 thin films prepared at 270°C & 310°C respectively. The binding energy in XPS is determined by the chemical environment of atoms at the surface and close to the surface (~100-150 Å). Hence it is possible to study the oxidation state of Sn by means of the oxygen (1s) spectra (Fig. 5), and the influence of Zn dopant in the lattice. In SnO2 commercial powder (Fig. 5 (S)), the foremost O 1s peak occurs at about 530.6 eV. In general, the 11

lower binding energy components of the O 1s spectra are associated with oxygen atoms in the oxide crystal, whereas the higher binding energy component represents the oxygen ions in the oxygen-deficient regions or at the adsorbed sites. So any oxygen peak below 530.5 eV shows up from lattice-oxygen below the surface, while peaks at 532.7.eV appear from O22− ions [21], and peaks near 534.5 eV is due to surface O2– ions or OH groups. In pristine films at 270°C, O− ions component in oxygen-deficient region appear prominently around 531.7 eV (Fig. 5a) along with the surface oxygen peaks at 532.5 eV and lattice-oxygen peaks with less prominence. With increase of the temperature of depositions to 310°C, the less coordinated peaks of oxygen at 530.4 eV gain importance with a minor contribution from O− ions in oxygen-deficient region around 531.7. On 2.4 wt% Zn doping, the intensity of slightly oxygen-deficient or less coordinated peaks of oxygen around 530 eV increases (Fig. 5c), together with contributions from O− ions in oxygen-deficient region come out around 531.1 eV and surface oxygen peaks at 532 eV. As the deposition temperature is increased to 310 °C, the peak shifts indicate the conversion of the O groups to a new chemical species plus the enhancement of O− ions around 531.6. As doping increases to 3.6 wt% Zn, lattice-oxygen deficient peak at 529.94 eV is dominant (Fig. 5f) with the suppression of O− ions in oxygendeficient region. Accordingly, in undoped films when the temperature increases from 270 °C to 310 °C, the XPS peak intensity corresponding to O− ions in oxygen-deficient region decreases to form a slightly oxygen deficient, primarily non-stoichiometric SnO2 film. With 2.4 wt% Zn doping, the lattice-oxygen deficiency crests at 310 °C, with the weakening of other peaks having higher binding energy. But in 3.6 wt% Zn doped films lattice-oxygen deficiency steadily intensifies with increase of deposition temperature from 270°C to 310°C. In brief, all the samples invariably show surface oxygen and oxygen-deficient lattice energies including 2.4 wt% Zn doped at 310°C, where the significant part of the lattice atoms are oxygen-deficient with a strong participation from O– ions in oxygen-deficient region. It is to be noted that in doped samples a peak in the region of 529.5 eV routinely appears and the sensing maximum depends on its optimum intensity relative to adsorbed oxygen peaks. In samples deposited at 270°C, the surface adsorption oxygen energies are more fabulous whereas in higher temperature depositions, additional oxygen-deficient lattice energies develop depending on doping level. Accordingly, a moderate oxygen-deficient non-stoichiometric compound formation is predominantly favoured for enhanced gas sensing by means of instigating improved surface oxygen adsorption. 12

3.3 Sensor Response to CO2 It has been established that the oxygen molecules are absorbed onto the film surface on deposition, by capturing free electrons from the Zn doped n-type SnO2-x according to O2(g) + e– (CB)→ O2– (ad), which decreases the carrier density in the films and hence the films show a higher resistance. On heating, the surface oxygen molecules get enough thermal energy to desorb and the sample resistance decreases. When exposed to 500 ppm of CO2 gas at a favorable operating temperature, the gas molecules react with the captured oxygen species on the surface of the sensor. Fig.6 (a &b) shows the dynamic sensor response of undoped and Zn doped SnO2 thin films prepared at 270°C and 310°C, and it is evident that the latter shows better response compared to that prepared at 270°C.

Fig. 6. The dynamic response in 500 ppm of CO2 for undoped and Zn doped SnO2 thin films prepared at (a) 270°C and (b) 310°C. The transient CO2 response of the undoped and doped films is measured at temperatures 200°C, 250°C, 300°C and 350°C, which is shown Fig 7. As the operating temperature increases, the sensor response increases to a remarkable level. At higher operating temperature, the gas molecule gets enough thermal energy to interact with the surface adsorbed oxygen species [22]. The highest sensor response of 94.4% is obtained at 350 °C for 2.4 wt.% Zn doped SnO2 sample prepared at 310°C and the response time is also good. On lowering the operating temperature, the sensor response gradually decreases. For the undoped one, the maximum sensor response is 39% at 350°C for the sample prepared at 310°C.

13

Fig. 7. Transient response in 500 ppm of CO2 versus time, at different operating temperature for 2.4 wt.% Zn doped SnO2 film prepared at 310°C. It is clear from the Fig.7 that, the CO2 ‘response times’ goes on varying with rise in operating temperature. The 90% response and recovery levels attained for 2.4% H films at 300 °C are 55 and 82 seconds, but at 350 °C they are 90 and 40 seconds, respectively. The slow response and recovery times for CO2 are notable features of the nanoparticulate Zn doped SnO2 thin film sensor. The extra amount of Sn2+ states as a result of Zn2+ doping together with lattice oxygen deficiency and the slow diffusion of charge carriers may be responsible for sluggish response and recovery features. 3.3.1 Response dependence on Temperature and Doping concentration The variation in CO2 response with temperature in the range 200°C to 350°C for undoped and doping concentrations of 2.4 & 3.6 wt.% Zn is shown in Fig. 8 (a & b). Regarding the influence of operating temperatures lesser than 250 °C on the response, it is found to decrease drastically for all the samples.

Fig. 8. Variation of response in 500 ppm of CO2 against operating temperature, for undoped, 2.4 wt.% and 3.6 wt.% Zn doped samples prepared at (a) 270°C (b) 310°C.

14

It can be observed that the sensing performance is clearly superior in Zn doped samples at operating temperatures higher than 300°C. Further, it is to be noted that the sensing performance is remarkable for 2.4 wt.% Zn doped sample prepared at 310°C. The incorporation of zinc ion in lower concentration in SnO2 films has the mostly shown increase of resistance [2325]. This is because the oxygen deficiency at the intra-grain SnO2 will be partially quenched by Zn2+ addition, modulated by chemisorbed oxygen at the SnO2 surface. This promotes the appearance of a potential barrier at the interface with the regular SnO2 lattice [26]. Accordingly, it is deduced that the increase of concentration of chemisorbed oxygen at the SnO2 surface in the range of higher temperatures [27]. Thus the formation of defects caused by lower percentage doping exerts strong influence on gas sensing properties. Consequently the surface oxygen defect is an important factor to be examined while analysing the sensor response. As the formation temperature is high, the SnO2 (110) surface will be losing its lattice oxygen’s (bridging, in-plane or sub-bridging positions). The contribution of in-plane to surface bridging oxygen atoms is a factor, which is to be established by PL and Raman measurements. 3.3.2 Selectivity of the sensor The selectivity of Zn-doped SnO2 sensor has been explored by exposing it to both reducing and oxidizing gases, comprising of LPG (liquid petroleum gas), CH4 (methane), NH4 (ammonia), and Cl2 (chlorine), in a range of operating temperatures. The responses are evaluated for deposition temperatures of 270°C and 310°C with a concentration of 500 ppm of gases. Fig. 9 depicts the selectivity of CO2 gas in comparison with other gases tested in the temperature range 200 to 350°C. It is clear from histogram that the sensor presents a positive resistance response to gases CO2, Cl2, NH3 while a negative trend to reducing gases LPG and CH4. It is also noted that for all the sensor films, the response of CO2 as a rule goes on increasing with rise in operating temperature. On exposure to oxidizing gas Cl2, the formed oxygen ions on the film surface at elevated temperature begin to react with chlorine, resulting in a decrease in concentration of electrons and an increase of the sensor film resistance. The films exhibit marginal selectivity to chlorine along with highest response of 72% in 500 ppm at 350 °C, with fast response time of 40 s, and short recovery time of 32 s. The sensors show improved response but a weak selectivity to chlorine gas below 350 °C temperature. In the case of NH3 it is curious to note that the film resistance increases just like in an oxidizing gas, though it is categorized as a reducing gas. Haizhou Ren et al have reported that 15

when the SnO2 surface is positively charged, the attracted NH3 molecules get their nitrogen atom adsorbed on the surface [28]. They can react with chemisorbed O2− ions on the sensor surface to form NOx, causing to behave as an oxidizing material to show an increase of sensor resistance. The sensor shows a weak selectivity amongst NH3 and CO2 gases under a sensing temperature of 350 °C.

Fig.9. Histogram shows the selectivity of Zn-SnO2 in 500 ppm concentration for different test gases. sample DT in parenthesis denotes the Zn-doped sensor deposition temperature. The magnitude of the highest transient response of about 63.2% and 48.2% in LPG and CH4 respectively, is markedly lower than that of CO2 and the resistance changes are also in opposite magnitude. Thus, the sensor based on 2.4 wt% Zn doped SnO2 films at 310 °C shows an excellent selectivity for CO2 against LPG or CH4. This selectivity arises from the fact that the Zn-doped SnO2 contains more Sn2+ that can be oxidized to Sn4+ by an oxidizing gas like CO2, to give a better response than reducing gases. The 90% response and recovery times recorded as 7 and 18 seconds for LPG and 25 and 18 seconds for CH4 , respectively are obviously eye-catching (for plots of dynamic responses and discussion, see supplementary material). Consistent with the results, the Zn-doped SnO2 sensor can selectively detect CO2 in a reducing gas ambience clearly, but only with some interference in an atmosphere of oxidising gas. The wide variations in the response and selectivity behaviour of these films are due to the deposition parameters in addition to the operating temperature. 3.4 Raman analysis

16

(a)

(b)

(c)

(d)

(e)

(f)

Fig.10.Room-temperature deconvoluted Raman spectra of undoped (a & b), 2.4 wt% (c & d) and 3.6wt% (e & f), Zn doped SnO2 nanoparticulate films deposited at 270° and 310°C respectively. 17

A detailed study of the Raman and Photoluminescence (PL) behavior would be helpful to elucidate the type of oxygen vacancies imparted by Zn2+ ions in SnO2 host lattice, from the perspective of enhanced gas sensing. The Raman active modes of SnO2 occur because of the polarization taking place as a result of the displacements of Sn and oxygen in the corresponding sub-lattices in SnO2. As the Raman mode is sensitively dependent on the surface disorder (the defects and structure in the surface regions), identifying these modes is equally well significant in resolving gas sensing action. Out of the three fundamental Raman peaks of rutile SnO2, the most intense peak observing at 635 cm‒ 1 is ascribed to the A1g mode, related to the expansion and contraction vibrations of Sn–O bonds. This mode is found to shift to 630 in 2.4% L and 627 cm‒ 1 in 3.6% L doped samples [Fig.10], whereas the peak appears at 635/634 cm‒ 1 in doped samples deposited at higher temperatures. In undoped samples deposited at temperatures 270 °C and 310 °C, the mode emerges at 623 and 625 cm–1 respectively. It is obvious that the intensity of A1g mode is strengthened along with the width narrowed when the Zn doping amount is 2.4 wt% at deposition temperature of 310 °C. It is reported that bridging OVs (OB) causes a downward shift in position of A1g mode at 635 cm–1 [29]. Further it is conveyed from theoretical calculation that in nanocrystals (NCs), the decline of OB causes to shift from 617.6 to 632.3 cm–1 [30]. We consider that the blue-shift of the A1g mode in the doped SnO2 sample must be related to the substitution of Sn by Zn ions in the lattice. Another characteristic mode B2g, arises from the asymmetric stretching of Sn-O bonds which is expected at 778 cm−1, red shifts to 768 cm−1 in 2.4% H, providing evidence of decrease of bridging OVs [31] in that set of samples. Samples deposited in other conditions by and large show the absence of this line, excluding 3.6% H. Depending on temperature of deposition, medium Raman bands at around 695/672 and 337 cm‒ 1 appear principally in pristine samples which seem to correspond to IR-active A2u LO and Eu (3) TO modes respectively, thought to come from surface modes of SnO2 as a result of OVs. The former band identified in pristine samples may caused by dominance of OB in a previous work [32]. In higher temperature doped samples, band at around 310/304 cm−1 characterizes the oxide of Zn in a nano- background or NCs. [33] and it is evident that the mode wavenumber decreases as the Zn content increases. Another fundamental Raman active mode, Eg, is observed in the range 464-492 cm−1, in which two oxygen atoms vibrate opposite to each other in the direction of the c-axis and this mode is highly sensitive to oxygen vacancies than the other modes. The variation in the position 18

of the observed broad peak for the Eg mode for different temperature depositions and doping concentrations establishes this point. We have found that, as the concentration of the oxygen vacancies get modified, the main Raman peak is broadened and shifted away from 476 cm−1 [34]. It is interesting to note that when the mode is shrinked along with intensity diminished, the gas sensing is remarkably good. At lower temperature deposited undoped films, an infrared active transverse optical mode A2u TO, appears prominently at around 502 cm–1. The oxygen absorption in SnO2 surface changes the vibrations of the O–Sn–O chain which produces shift in the Raman lines. Deconvoluted Raman bands (Fig. 9) of doped samples show scattering modes in the range 556-586 cm–1. It is established in literature that the surface related mode appears at 556 cm–1, identified as optically inactive band A2g. The band is related to the small size effect as stated by Matossi force constant model [35] with large amount of surface disorder and defects which disrupt the translational symmetry of the ideal crystal structure. As a result, the selection rules get modified to transform the infrared active modes or silent modes to become Raman active. This band is not observable in pristine SnO2, shifts to higher wavenumber (blueshift) with the increase in ‘O’ (oxygen) vacancy. Liu et al [36] has reported that the in-plane OVs (OP) are in fact responsible for the Raman mode at ~574 cm−1. In 2.4% H films, the band occupies the range 586 cm–1 in the deconvoluted plots is certainly due to more number of in plane oxygen vacancies. Raman active mode observed below 300 cm–1 at 298-283 cm–1 is assigned to OVs in nano-structure surface and in the range of 223-216 cm–1 can be attributed to the substoichiometric Sn3O4 phases [37] present in samples. The nanocrystal ZnO scattering lines typically appear at 197 cm–1. Hence the lines between 197-192 cm–1 in doped samples uphold the doping realization. Besides, ZnO itself has a Raman scattering mode at 304 cm–1 [33]. The area of the mode is found to increase as the Zn content increases, validates the Zn incorporation into the SnO2 lattice and the Raman results corroborate the XRD and XPS findings. Novel peaks appear in the Raman spectra wavenumber range 128-168 cm−1 of films, identified as the partially oxidized stoichiometric phase Sn2O3/ Sn3O4 [38]. Furthermore, broad bands deconvoluted into peaks viewed in the span 102-117 cm–1 can be assigned to SnO phase centering at 115 cm−1 (B1g). An intense shoulder peak in the lower wavenumber region (75 to 92 cm−1) of the spectra decreases, when the doping concentration and deposition temperature increase. This is attributable to characteristic peak of Sn3O4 at 72/90 cm−1. In short, the broad

19

hump in the wavenumber range 76-168 cm−1 corresponds to the sub-stoichiometric SnOx phases (1
20

Fig. 11A. Photoluminescence spectra (room-temperature ) of (a & d) pristine, (b & e) 2.8wt% and (c & f) 3.6wt%, Zn doped SnO2 nano-particulate films deposited at 270° and 310°C respectively.

Fig. 11B. Photoluminescence spectra from 580-700 nm (room-temperature ) with excitation 370 nm, of (a & d) pristine, (b & e) 2.4 wt% and (c & f) 3.6wt%, Zn doped SnO2 nano-particulate films, deposited at 270° and 310°C respectively. From the Raman spectra it is evident that defects are present in the samples. To understand the features of these defects more, the emission PL spectra have been studied. PL is a very sensitive technique in terms of crystal defects, surface effects, energy bands and exciton 21

fine structure in nanocrystals. Fig. 11A shows the PL spectra of the films recorded with excitation wavelength 290 nm (corresponding to photon energy 4.28 eV). The luminescence peaks originate mainly from the energy states present in the band gap due to defects like tin interstitials, dangling bonds and oxygen vacancies [41]. It is worth noting from the spectra that the emission spectral intensity of Zn-doped SnO2 nanoparticulate films at deposition temperature 270°C decreases as the doping concentration increases, whereas the intensity increases up to 2.4 wt. % doping at 310°C and then decreases. In contrast, for the higher temperature depositions the green region displays a relatively higher loss of intensity. This decreased intensity cannot be solely ascribed to the change in the particle size as, the same sized particle show differing PL intensity and have different concentration of oxygen vacancies, to show dissimilar gas sensing response depending on deposition temperature. Another important aspect observed is that, there is decrease in emission spectral intensity with the increasing dopant concentration in doped samples. This behavior is attributed to the modification in the crystal field or quenching of oxygen defect states through interaction of the dopant atoms with the support oxide, which is exhibited by the reduction of peak intensity. The samples show photoluminescence at 333, 412, 496, 536, 557 nm. Considering the emission spectra reported by other researchers [42–46], the emissions of SnO2 at room temperature between 350 and 580 nm have usually been accredited to deep or shallow energy levels of oxygen defects or vacancies, where the electron in an oxygen vacancy recombines with a hole to emit a photon. The emission peaks, lesser than the band gap energies cannot be assigned to direct recombination of a conduction electron in Sn 4d band and a hole in O 2p valence band, but labeled as a localized states [47]. A strong near-UV emission band at 333 nm (3.72 eV) remaining invariable in its position (do not record a shift) with respect to doping or higher temperature of deposition, is habitually designated to the band-edge/ near band-edge emission in the SnO2 films containing nanoparticles. A shoulder perceptible near 374 nm (3.32 eV) is attributed to a recombination of near band edge free excitons emission. This strong UV band is characteristic of the near-band-edge emission originating from the recombination of free excitons. Increased UV emission bands can also be attributed to the decrease of the defects [48] and increase of ZnO in crystallites. The UV emission intensity increases with respect to Zn content up to 2.4 wt% and it decreases for 3.6 wt% doping. This is may be due to the superior crystalline nanoparticle SnO2 films formed at

22

lower doping concentration. Also the PL spectra show the blue shift in the band edge emission, on increasing the dopant concentration. The appearance of the violet broad sharp luminescence bands at 412 nm (~ 3.01 eV) in the sample is attributed to surface oxygen deficiencies of shallow VO0 states. In the 0% L and 2.4 wt.% L, the NPs have a particle size of ~5 nm, showing more intense luminescence at 411 nm in comparison with the other samples. The position of blue-green luminescence (BGL) at 496 nm (~2.5 eV) decreases to 495 nm, on Zn doping of 2.4 wt% and 3.6 wt%, in both the deposition temperatures of 270 and 310°C. The emission about 496 nm related to deep level emissions (DLE) is due to electron transition, mediated by singly/doubly ionized oxygen vacancies concentration. DFT modeling using generalized gradient approximation (GGA) predicted luminescence bands around 486 nm (2.55 eV in materials with band gap of 3.6 eV) which was attributed to the in-plane ‘O’ vacancies [49]. S.-T. Jean et al have shown in Sb-additivated SnO2 nanostructures that the exhibited peak at 2.55 eV is associated with in-plane surface oxygen species [50]. Thus the peak at 495 nm is attributable to formation of deep trap due to a singly charged OP. A remarkable observation made during the measurements is that as deposition temperature increases, up to a Zn doping level of 2.4 wt% in SnO2, there is a continued strengthening of the blue-green band (496 nm), which is commonly attributed to oxygen vacancies, habitually occurring on the surface of the nanocrystals. This observation is an apparent evidence that the oxidation state of the tin ions changes from Sn4+ to Sn2+ as the Zn-content increases in low concentration. The photoluminescence yellow broad shoulder at 556 nm (2.23 eV) of the prepared SnO2 films is found to be slightly red-shifted with both doping concentration (up to 2.4 wt%) and surface area ratio, evidencing an increase in CO2 response, which is attributed to an increase in deep oxygen defect energy levels [51]. This broad shoulder band is regarded as a result of OB habitually occurring on the surface of the nanocrystal films [44]. The PL emission from bluish-violet 2.90 eV (427 nm) to yellow 2.29 (541 nm) is usually regarded as a result of OP in literature [30, 52]. In the case of 3.6% L and 2.4% H, the NPs have a size of ~ 10-11 nm, the OP related luminescence is strongly observed at 536 nm in supplement to the Raman mode at 573 cm–1, which shows a congruent strong scattering in nanocrystals. Hence the peak at 536.6 nm (2.31 eV) can be looked upon as of OP vacancies. In contrast, one may conclude that the PL position and the line shapes of nanocrystalline Zn doped SnO2 films may slightly differ from those of microcrystalline ones, depending on crystallite size. 23

The PL emission with an excitation 370 nm (Fig. 11B) shows a major peak at 609/610 nm (2.03 eV), which is considered to stem from bridging oxygen vacancies. It is to be noted that, the ratio of in-plane to bridging ‘O’ vacancies is found to peak nearly at 33 nm sized particles. It is crucial to note that when the PL peak area ratio of 496 nm to 609 nm, corresponding to inplane to bridging ‘O’ vacancies respectively is nearly 2 with a particle size (~10 nm), then the CO2 response is found to be higher. This is dissimilar with the observation made by Bonu et al [53] that the SnO2 nanostructures showed decreased response around 9 nm, while 4 nm and 25 nm NPs showed enhanced CH4 sensing. Zn2+ doping in SnO2 happens to introduce new Sn2+ states at Sn4+ sites, and in fact more oxygen vacancies. It is plausible that the crystallites in the samples are apparently by the coexistence of SnO2 and SnO crystalline phases. Oxygen’s in the environment of six-fold coordinated Sn are commonly bridging oxygens and that lying on the surfaces is referred as inplane oxygen. Loss of oxygen from the inside leads to a lower coordination number of five and four of Sn, which maintains charge neutrality due to the multivalence of Sn i.e. Sn4+ and Sn2+. Lowering the coordination number of Sn from six to four denotes the absence of bridging oxygen, evident in (110) surface. But in the case of (110) surface, additional in-plane oxygen atoms have to be removed to obtain a surface layer that shows only Sn2+ surface atoms [54]. This additional in plane oxygen removal is likely responsible for the added oxygen adsorption and enhanced gas detection observed. But in samples deposited at lower temperature (2.4% L), substoichiometric phase is higher and response to CO2 is somewhat less compared to 2.4% H films having higher OP to OB ratio. Except for samples doped in 3.6 wt% Zn concentration deposited at higher temperature of 310 °C, the OP vacancies are found to dominate over OB vacancies. This observation points out that, defect concentration related to in-plane and bridging ‘O’ vacancies is influential in the CO2 gas sensing by Zn-SnO2. Decrease in the disorder as the Zn content increases, must be related to the quenching of oxygen vacancy population at the particle surface due to the substitution of Sn4+ by Zn2+ ions. Accordingly, the photoluminescence analysis exposes the reason for the high CO2 sensing response of Zn doped SnO2. It is also found that the dislocation density along with oxygen vacancies is a factor that determines the sensor response in Zn-doped SnO2 films, and the response peaks at when they are of higher values. High surface to volume ratio influences the sensing only when it is supplemented by adequate amount of surface defects.

24

3.6 CO2 sensing mechanism On deposition in air, oxygen is absorbed at the surface oxygen vacancies of SnO2 and the ionized VO2+ will switched to VO0 by the use of conduction electrons and the material resistance is high. VO2+ + 2e′ ½O2 → VO0, where VO2+, is the doubly ionized oxygen vacancy and VO0, a neutral oxygen vacancy. With sufficient energy, helpful for dissociation at temperatures above 180°C, the VO0 will liberate electrons which are subsequently ionized to VO2+, causing the electrons to release back to the semiconductor to produce a decrease in resistance. On exposure to CO2, the ionised oxygen vacancy VO2+ at the surface capture electrons, which subsequently are absorbed from the material to evidence for an increase in resistance, to effect for gas sensing. CO2(ads) + VO2+ → CO22+ + VO0 CO22+ + 2e′ (CB) → CO2(ads) As a consequence, oxygen vacancies and their position are key factors that determine the sensing characteristics. Zn-doped SnO2 NPs exhibit higher gas-sensing response to CO2 than pure SnO2 NPs and 2.4wt.% Zn at 310°C shows the highest response. It is attributed to the large amount of oxygen vacancies arising from the substitution of Sn4+ by Zn2+ in SnO2 lattice. Gas response is found to be related to the PL intensity of oxygen-vacancy-related defects in as-deposited gas sensors [55]. The enhanced oxygen vacancies facilitate the oxygen adsorption on the surface of SnO 2. This approach in essence, help to modify the gas-sensing properties by tuning the amount of oxygen vacancies which is an overriding factor in gas-sensing process and influences the gas molecule adsorption and catalytic reaction at surface [56, 57]. Therefore, the increased gas-sensing responses of Zn-doped SnO2 NPs, with Zn concentration from 0 to 2.4 wt% should be ascribed to the enhanced oxygen vacancies at the surface of SnO2 crystallites. Appearance of the characteristic Raman lines together with their shift confirms the oxygen vacancies and the coexistence of sub-stoichiometric phases. On doping SnO2 lightly with zinc, the oxygen vacancies seem to be partially quenched; thereby reducing anion oxygen vacancies, but ZnSn center in SnO2 promotes the increase in concentration of chemisorbed oxygen. In PL, the decreased intensity cannot be solely ascribed to the modification in the crystal field, but to the quenching of oxygen defect states through interaction of the dopant atoms with the support oxide. Increased gas-sensing responses of Zn-doped SnO2 nanoparticulate films with 25

lower Zn concentration is ascribed to the enhanced oxygen vacancies at the surface of SnO 2 crystallites. Detailed luminescence analysis of films containing NPs revealed that the defects related to emissions at 495 nm and 536 nm, corresponding to the in-plane ‘O’ vacancies and that at 556-559 nm and 609 nm related to bridging ‘O’ vacancies play a crucial role in the CO2 sensing process. The extract of findings is tabulated with sensor response in Table 3. Table 3 CO2 Sensor responses and the ratio of in-plane to surface bridging oxygen vacancies of films deposited at different temperatures. Sample Description CO2 Sensor Ratio OP/OB response % 43 undoped at 270°C 1.22 65 2.4 wt.% Zn at 270°C 1.49 3.6 wt.% Zn at 270°C Undoped at 310°C

76 39

1.66 0.96

2.4 wt.% Zn at 310°C 3.6 wt.% Zn at 310°C

94.4 65

2.01 1.45

A careful luminescence and Raman analysis in the divalent cation doped nanoparticulate SnO2 films illustrate the dominant role of in-plane oxygen vacancies in gas sensing action. With 2.4 wt% doping at temperature of 310 °C, the in-plane oxygen vacancies are found to lead over bridging oxygen vacancies by a factor two to obtain a peak sensor response of 94.4 %. 4. Conclusion In the present work, it is shown that by the doping of Zn in SnO2, the crystal structure is not disturbed. No secondary phases are observed in the patterns of XRD, suggesting that zinc occupies the tin sites in SnO2. The lattice constants are calculated and interestingly there is a slight change in the lattice parameters with increasing dopant concentration. It is important to note that with increasing dopant concentration, c/a values are found to increase gradually. From XRD and HRTEM it is inferred that, as the preferred orientation of (110) get intensified, the CO2 sensing get enhanced. The grain size reduction causes lattice distortion which is likely to produce more surface defects and oxygen vacancies. The ZnSn, oxygen interstitials (Oi) or charged oxygen vacancies VO+ and VO2+ are believed to be the reason for an increase in the size of the lattice. It is logical that Zn2+ doping in SnO2 happens to introduce new Sn2+ states at Sn4+ sites and in fact more oxygen vacancies. Thus it is concluded that the combination of the doping effects and lattice expansion led to the increase in CO2 response of Zn2+-doped SnO2 nanocrystals. 26

With 2.4 wt% doping at temperature of 310 °C, the in-plane oxygen vacancies are found to exceed bridging oxygen vacancies by a factor two to obtain a peak sensor response of 94.4 %. Hence on light doping, the surface defects especially the ratio in-plane oxygen vacancies to bridging oxygen vacancies, together with (110) crystallite orientation, play an important role in enhancing the CO2 sensing performance. The results evidently depict that Zn2+ doping in SnO2 introduces new surface and structural defects in the deposited films. Defects, in particular, oxygen vacancies, appear to be critical in understanding the gas sensing action. However, regarding the sensing temperature the position of oxygen vacancy appears to be vital along with impurity states. A more detailed study on the role of oxygen vacancy-type within the band gap energy levels and the gas sensing behavior is underway and will be reported in future communications.

Acknowledgments The author (DS) gratefully acknowledges the Faculty Improvement Fellowship from the UGC (SWRO/FIP 12th Plan/ KLMG 038 TF-06), and the author (PK) is grateful to KSCSTE (822/DIR/2014-15/KSCSTE dated 09.02.2015) for the financial assistance. The authors thank Mr. S. Sarath at Amrita Centre for Nanosciences for his assistance in XPS analyses.

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