Easy synthesis of organic–inorganic hybrid nanomaterials: study of DC conduction mechanism for light dependent resistors

Abstract

A surfactant assisted chemical oxidation method has been employed for facial synthesis of polypyrrole/tin oxide (PPy/SnO₂) hybrid nanoneedles. The charge transport properties of the prepared hybrid nanomaterials have been analyzed under different conduction mechanisms for their possible application for light dependent resistors (LDRs). The scanning electron microscopy studies revealed that the increased concentration of additive SnO₂ quantum dots alters the surface morphology from nanowall-like to nanoneedle-like and the formation of PPy/SnO₂ nanocomposites is confirmed X-ray diffraction studies. The strong coupling between the PPy and SnO₂ results in the transformation of PPy from highly oxidized states to oxidized states and is confirmed through Raman analysis. The simple band conduction model and Kivelson's power law based conduction mechanism could not be applied to explain the conduction mechanism due to different reasons. The log σ_dc curves for all samples were well fitted for γ = 1/4, suggesting the applicability of Mott's three dimensional variable range hopping VRH model for charge transport. The room temperature dark conductivity of the samples was found to decrease with an increase in SnO₂ concentration. The formation of p–n junctions between p-type PPy and n-type SnO₂ changes the bandgap of PPy and the work function. This modifies the electronic structure of PPy which brings a synergistic effect in the photosensitivity of the PPy/SnO₂ nanocomposites. The room temperature photoresponse of the PPy/SnO₂ samples was found to increase from 2.85 to 6.25% at 100 mW cm⁻² illumination intensity with an increase in the SnO₂ doping concentration from 0 to 20%.

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I. Introduction

Developing countries are wasting large amounts of electrical energy in the unnecessarily overriding of lightening equipment at homes and streets/roads due to their manual operation. At present the cost of the light dependent resistor (LDR) used in the automatic switches is high. Moreover, most of the studies on LDR devices and photodiodes are devoted to the cadmium based chalcogenides (CdTe, CdS, CdSe) or silicon based materials.^1–5 The precursor materials used in cadmium or silicon based devices are highly toxic which results in long term hazardous environmental effects and requires highly sophisticated and high temperature synthesis processes. For example, silicon based technology is very costly and highly sophisticated in terms of growth process, equipments and precursor materials. The dopant gases (diborane or phosphene) used in silicon are highly toxic and special cares are required to avoid the leakage in ambient environment.⁵ Some metal oxides (TiO₂ and ZnO, etc.) and metal nitrides (InN) are also used for photovoltaic applications and UV detections.^5–10 In recent years, conducting polymers are emerging out as potential candidate for solar cells,¹¹ light emitting diodes,^12,13 gas sensors,^14,15 electromagnetic interference shielding¹⁶ and other applications. Poly-3-hexylthiophene (P3HT) is being extensively studied for photovoltaic applications,¹¹ whereas, polypyrrole (PPy) and polyaniline (PANI) are analyzed for gas sensor applications.^14,15 However, these materials still have some stability related issues under ambient conditions, which can be resolved by modifying the synthesis process and making them in the nanocomposite form with other inorganic materials. This new hybrid form may not only provide the environmental stability, but also improve their properties due to enhanced surface area or increased number of active sites for interaction at nano-level. Out of several studied conducting polymers, PPy is considered as one of the most stable conjugated polymeric system with low cost and easy synthesis process with large possibility of foreign element (SnO₂, ZnO, Fe₂O₃, CNTs and Ni, etc.) incorporation for various applications.^17–23 The long term stability and non-toxic nature of PPy can make it a potential candidate for these applications. The nanocomposites of PPy/SnO₂ are well known for their sensing applications for both reducing¹⁹ and oxidizing²⁰ gases. PPy/SnO₂ nanocomposites also exhibit many applications in batteries²¹ and supercapacitors²² as an electrode material and their electrical properties are well reported.²³ Although there is a dearth of data on the study of optical properties of PPy and its nanocomposites for light dependent resistor (LDR) applications. Investigation of photoconductive properties of a material is an important tool in understanding the nature of photo-excitations and quality of material for its possible application in light dependent resistors (LDRs). The photoconductivity is generally governed by generation of electron–hole pairs through the absorption of light photon which results in an increase in charge carrier concentration and thereby electrical conductivity of the material. The rise and fall in the photocurrent curves depend upon the nature and distribution of the traps and recombination centers in the forbidden gap of the material. Thus, the study of photoconductive property of a material plays a vital role in understanding the nature of electronic excitations, nature and distribution of traps and recombination centers so that the materials can use for different applications in electronic and optoelectronic devices.²⁴ A good photoconductor not only have efficient charge separation ability but also have efficient transport of charge carriers to electrodes. Thus, to improve the charge separation ability and charge transport property of PPy, the preparation of hybrid nanostructures of PPy with SnO₂ are focused in the present study. The encroachment in photoconduction is considered to occur from the interfacial interaction between the PPy and SnO₂, however, the mechanism of interaction is quite complex and requires special attention. Therefore, before gaining the insight on the photoconduction behavior and charge separation ability of PPy, the understanding of the dc conduction behavior or transport of charge carries to the electrodes in nano/micro scaled PPy is necessary. Conjugated polymers generally obey the Mott's variable range hopping (VRH) conduction mechanism of disordered materials in their bulk forms. However, it may be interesting to study the conduction behavior of PPy and its nanocomposites with SnO₂ quantum dots (SnO₂ QDs) at nanoscale or submicron-scale. Therefore, in the present study, we synthesized the nanocomposites of PPy with SnO₂ QDs by surfactant directed solution based chemistry method. The temperature dependent DC conduction behavior of the nanocomposites has been analyzed under Kivelson's power law and Mott's variable range hopping conduction models. The photoconduction behavior of the prepared samples has also been analyzed as a function of time and illumination intensity for their possible application in light dependent resistors.

II. Experimental details

A. Synthesis of SnO₂ QDs

A simple hydrothermal process has been used for the synthesis of SnO₂ QDs. In a typical synthesis process, 50 mM salt solution of sodium stannate trihydrate was prepared in 200 ml double distilled water and refluxed for 2 h at 120 °C. Thereafter, 50 mM solution of capping agent {cetyltrimethylammonium bromide (CTAB)} was added to it followed by addition of 1 M NH₃. A separately prepared 1 M KOH solution was added drop-wise to it and the reaction was carried out for 2 h at the same temperature and then cooled down to room temperature. Finally, the prepared material was filtered out, dried and annealed at 300 °C. The synthesis of SnO₂ nanocrystals in the aqueous solution follows the chemical reactions

NH₃ + H₂O → NH₄⁺ + OH⁻ (1)

Sn⁴⁺ + 4OH⁻ → Sn(OH)₄ (2)

Sn(OH)₄ → SnO₂ + 2H₂O (3)

B. Synthesis of polypyrrole/SnO₂ nanocomposites

Polypyrrole/tin oxide (PPy/SnO₂) nanocomposites were prepared by chemical oxidation polymerization method at room temperature (18 °C). Three samples of PPy/SnO₂ nanocomposites were prepared with three different concentrations of SnO₂ QDs viz. 5, 10 and 20% by weight and designated as sample S1, S2 and S3, respectively. A pure PPy sample (S0) was also prepared. For the synthesis of sample S1, 60 mM of CTAB was dissolved in double distilled water, followed by addition of 2 M HCl solution. Thereafter, 5 wt% SnO₂ QDs was added to the surfactant solution and stirred for 45 min to ensure the homogeneous mixing of SnO₂ QDs. A separately prepared 60 mM oxidant solution of ammonium persulphate (APS) was mixed in the above solution. After that distilled pyrrole monomer (30 mM) was added drop-wise to it and the reaction was carried out for 12 h. The black precipitates of PPy/SnO₂ were separated out using a Whatman filter paper and washed several times with distilled water and methanol alternately. The black precipitates were dried at 60 °C and labeled as sample S1. Similar synthetic route was used for the synthesis of other nanocomposites (S2 and S3) with 10 and 20 wt% of SnO₂ and pure PPy (S0) without SnO₂ doping.

C. Sample characterization

X-ray diffraction studies of SnO₂ QDs as well as on the PPy/SnO₂ nanocomposites have been carried using D8 discover (ASX-Bruker), X-ray diffractometer using Cu-Kα radiation. High resolution transmission electron microscopy (HRTEM) (FEI, Tecnai G2 F30-STWIN) study has been carried out of the SnO₂ sample for particle size determination. Scanning electron microscope (SEM) studies of PPy/SnO₂ nanocomposites has been carried out for observing the surface morphology using Zeiss EVO 50. The molecular structure of samples was analyzed through Raman spectroscopy using Renishaw InVia Reflex micro Raman spectrometer. The temperature dependent DC conductivity has been measured in the temperature range of ∼77 to 400 K using Keithley's 2410 source measure unit (SMU). Before performing the electrical measurements, gold electrodes (3 mm dia.) were deposited on both sides of the pellets (7 mm in dia. and 1 mm thick) by thermal evaporation technique under a high vacuum (∼10⁻⁶ mbar) condition. The electrical measurements were carried out in the sandwich structure (Au/polymer/Au) by the two probe method. A constant voltage was applied across the samples and the change in electric current was measured by Keithley 2410. The temperature is controlled by the Eurotherm 3216 temperature controller. The room temperature photoresponse of samples has also been recorded in the wavelength range of ∼400–700 nm using Keithley 2410 SMU in the resistance mode at different illumination intensities ranging from 30 to 100 mW cm⁻². The UV-Vis spectra of the samples dispersed in isopropyl alcohol were recorded at Thermoscientific Evolution-300.

III. Results and discussion

A. Morphological and structural investigations

Fig. 1(a) shows the HRTEM image of SnO₂ sample. A uniform distribution of SnO₂ quantum dots (QDs) having particle size in the range of 2–5 nm is observed throughout the sample. The diffused ring in the selected area electron diffraction (SAED) indicates the amorphous nature of the sample or the formation of smaller sized particles/crystals (inset Fig. 1(a)). The observed rings correspond to (hkl) planes (110), (101) and (201) of the SnO₂ rutile phase. Fig. 1(b)–(d) elucidates the SEM micrographs of PPy/SnO₂ nanocomposites. It is observed that the sample prepared with 5 wt% doping of SnO₂ (Fig. 1(b)) has flake-like nanowalls of size greater than 10 μm with the wall thickness of ∼300–400 nm. However, with increase in SnO₂ doping concentration from 5 to 10 wt%, the morphology of the samples get significantly changed from nanowalls to nanoneedles emanating from these flakes (Fig. 1(c)). The length of these nanoneedles is ∼50 μm. It can be observed that several nanoneedles originating from different portions of the flake are converged into a single rod with sharp tip. The concentration of nanoneedles was found to increase with increase in doping concentration of SnO₂ QDs (Fig. 1(d)). Thus, the presence of SnO₂ QDs greatly influences the surface morphology of the nanocomposites. This can be attributed to the change in the shape of micelles in the presence of SnO₂ QDS because of different ionic characters of surfactant (cationic nature of CTAB) and SnO₂ QDs (n-type character).

Fig. 1 (a) HRTEM micrograph of SnO₂ QDs and SEM micrographs of PPy/SnO₂ nanocomposite samples with (b) 5 wt% doping of SnO₂ (S1), (c) 10 wt% doping of SnO₂ (S2), and (d) 20 wt% doping of SnO₂ in PPy (S3).

XRD measurements have been carried out for the structural analysis of the nanocomposites. Fig. 2 shows the XRD patterns of SnO₂ QDs and PPy/SnO₂ nanocomposites S1, S2 and S3, which consists of three broad peaks located at 2θ of ∼27, 34 and 52° associated with the (hkl) planes (110), (101) and (211), respectively. This indicates the formation of SnO₂ rutile phase (JCPDS 41-1145).^19,25,26 The XRD pattern of PPy/SnO₂ nanocomposite samples also have similar peaks as observed in the pure SnO₂ QDs along with a broad hump near (110) plane of SnO₂. This could be due to the overlapping of broad hump of amorphous structure of PPy located at ∼27° with (110) plane of SnO₂ that results into a broad hump shifted towards left as compared to (110) plane of SnO₂. However, with increase in concentration of SnO₂, this peak shifted towards (110) peak of SnO₂ accompanied with decrease in full width at half maximum value of this band. Moreover, the relative intensities of the peaks corresponding to (101) and (211) planes as compared to mix peaks of PPy and SnO₂ (110) located at ∼27° are found to increase with the SnO₂ concentration. The presence of SnO₂ peaks in the XRD patterns indicates the formation of PPy/SnO₂ nanocomposites.

Fig. 2 XRD patterns of SnO₂ QDs, PPy (S0) and PPy/SnO₂ nanocomposites (S1–S3).

Fig. 3 elucidates the UV-Vis spectra of pure SnO₂ QDs, pure PPy (S0) and PPy/SnO₂ nanocomposites with 5 and 10 wt% of SnO₂ (S1 and S2). The optical band gaps of samples are calculated from the Tauc's relation as,²⁷ α = A(hν − E_g)ⁿ, here, hν is the photon energy and E_g is the band gap in eV. A is the proportionality constant and power index n = 1/2 for direct band gap, α is the absorbance coefficient. It has been observed that the band gaps of pure SnO₂ QDs, PPy (S0) and nanocomposites samples S1 and S2 are 3.58, 2.25, 2.41 and 2.83 eV, respectively. The band gap of PPy is found to increase with increase of SnO₂ doping concentration. It is observed that the band gaps of the pure PPy and PPy/SnO₂ QDs lies in the visible region of the radiation spectrum suggesting their possible application in white light detectors, whereas, band gap of SnO₂ QDs falls in UV region which can make it useful for UV sensors.

Fig. 3 UV-Vis spectra of SnO₂ QDs, PPy (S0) and PPy/SnO₂ nanocomposites (S1, S2).

Fig. 4 elucidates the Raman spectra (600–1800 cm⁻¹) of all the PPy samples with and without doping of SnO₂, whereas, the inset of Fig. 4 shows Raman spectrum of SnO₂ nanoparticles. Raman spectrum of SnO₂ nanoparticles consists of all the major bands of SnO₂ rutile phase at ∼475, 634 and 778 cm⁻¹ and can be attributed to the first-order Raman active modes, E_g, A_1g and B_2g, respectively. In these Raman active modes, Sn atoms remain in the stationary state and the oxygen atoms vibrate. The band located at ∼585 cm⁻¹ can be ascribed to the amorphous SnO₂ hydrous oxide.^27,28

Fig. 4 Raman spectra of samples S0–S3 and inset shows Raman spectrum of SnO₂ QDs.

The observed Raman spectra of PPy/SnO₂ nanocomposites are quite similar to that of PPy sample without doping. However, a small shift in the peak position from ∼1592 to 1583 cm⁻¹ of the band associated with the C=C stretching vibration was observed. The shift of this band towards the lower wavenumber is due to lower oxidation level of the sample with higher concentration of SnO₂. Generally, in a highly oxidized sample with high conductivity, this band is observed towards higher wavenumber, whereas, in the case of lower oxidized or reduced samples this band is observed towards the lower wavenumber. When this band is observed at higher wavenumber (∼1592 cm⁻¹), it can be attributed to the overlapping of two bands associated with oxidized (polaronic at ∼1590 cm⁻¹) and highly oxidized (bipolaronic at ∼1610 cm⁻¹) states of PPy. However, when this band is observed at lower wavenumber side (1580 cm⁻¹), it can be attributed to the overlapping of bands associated with oxidized and neutral (∼1560 cm⁻¹) states of PPy. Thus, the shift in the peak position of this band towards lower wavenumber side indicates the reduction oxidation of PPy with SnO₂ and hence, there is a decrease in its electrical conductivity.²⁹ The broad hump below the C=C band have two diminished but well visualized peaks at 1380 and 1330 cm⁻¹ associated with oxidized ring deformation (C–N stretching) and C–H in plane bending of neutral states, respectively.^{15,16,30–32} The band observed at ∼1265 cm⁻¹ can be attributed to C–H in plane deformation. Polarons and bipolarons are considered to be major charge carriers in the polymeric materials and the bands associated with these charge carriers are observed in 900–1100 cm⁻¹ range in Raman spectrum of PPy. The presence of all the fundamental bands of PPy in the PPy/SnO₂ nanocomposites indicates polymerization of pyrrole.^{15,16,30–32} There are two major bands (located at ∼1050 and 980 cm⁻¹) in 900–1100 cm⁻¹ range of Raman spectra of all the samples. The band located at 1050 cm⁻¹ is asymmetric towards higher wavenumber side which indicates the presence of another band at ∼1080 cm⁻¹, whereas, the bands located at 980 cm⁻¹ is asymmetric towards lower wavenumber with a shoulder at ∼940 cm⁻¹. The bands located at 1050 and 1080 cm⁻¹ can be attributed to polaronic and bipolaronic states of C–H out of plane deformations, whereas, the bands observed at ∼980 and 940 cm⁻¹ are ascribed to the polaronic and bipolaronic states of the ring deformation of PPy. The intensity of the bands is found to be different for different doping concentration of SnO₂ which indicates the change in charge carrier concentration of the sample with doping concentration. Thus, to know the effect of SnO₂ concentration on charge carrier concentration, the Raman spectra of all the samples are separately scanned in the range of 890–1120 cm⁻¹ and deconvoluted into four components by using the Gaussian (80%) + Lorentzian (20%) fitting (Fig. 5(a–d)). It has been observed that the intensity (area under the curve) of the bipolaronic bands (∼940 and 1080 cm⁻¹) decreases, whereas, the intensity (area under the curve) of polaronic bands (∼980 and 1050 cm⁻¹) increases with the SnO₂ doping. To estimate the change in the charge carrier concentration with SnO₂ doping, the ratio of the area of the bipolaronic bands (A_B = A₉₄₀ + A₁₀₈₀) to polaronic (A_P = A₉₈₀ + A₁₀₅₀) bands has been calculated from the area under the bipolaronic and polaronic bands (as shown in Fig. 5(e)). It is observed that the A_B/A_P ratio decreases with increase in SnO₂ doping concentration, which indicates a transition from the bipolaronic to polaronic state. This suggests the decrease in charge carrier concentration with increase in SnO₂ doping concentration. The polarons and bipolarons are the charge carriers in conducting polymers and are responsible for the electrical conductivity. Thus, we have calculated the room temperature electrical conductivity of the samples and observed that the electrical conductivity of the samples decreases with increase in SnO₂ doping concentration as well as with decrease in bipolaron to polaron ratio (as shown in Fig. 4(e)). Apart from the SnO₂ concentration, the electrical conductivity can also be affected by the surface morphology of the samples as one dimensional nanostructures can provide easy path for the charge carriers. The improved mobility may result in the higher electrical conductivity. Thus, there is a competition between the effect of surface morphology and SnO₂ doping concentration. The nanoneedles type morphology of samples try to improve the mobility or electrical conductivity, whereas, the wide band gap SnO₂ nanoparticles tries to retard the motion from the valence band to the conduction band and hence the electrical conductivity. However, in the present study, the effect of SnO₂ doping dominates over the morphological effects and hence a decrease in the electrical conductivity is observed with increase in SnO₂ doping concentration even though the morphology change from nanowalls like flakes to nanoneedles.

Fig. 5 (a–d) Deconvoluted Raman spectra of samples S0–S3 in the range 890–1120 cm⁻¹, and (e) variation of the area ratio of the bipolaron to polaron bands (A_B/A_P) and room temperature dc conductivity versus SnO₂ doping concentration.

B. DC conduction behaviour of PPy/SnO₂ nanocomposites

The temperature dependent DC conduction behavior of PPy (S0) and PPy/SnO₂ nanocomposites (S1–S3) has been investigated in the temperature range of 77–400 K. It is observed from Fig. 6(a) that the electrical conductivity of all the samples increases with increase in temperature. This suggests the semiconducting behavior of the samples which follows the relation^18,33where, σ₀, E_A, k_B and T are the pre-exponential factor, activation energy, Boltzmann's constant, and temperature, respectively. It is observed that the temperature dependence of the conductivity becomes strong with increase in the doping concentration of SnO₂ QDs. This may be due to increase in width of barriers between the conducting regions and decrease in higher conducting phase (polymeric region) with doping concentration. Thus, the overall conduction behavior of the samples may be governed by the dynamic interaction between the conducting (polymeric) and non-conducting (metal oxide) phases. The temperature dependent conduction behavior of the present samples has been analyzed with the possibility to follow different conduction mechanisms. The Kivelson's inter-chain hopping conduction between the neutral and charge solitons states model^34,35 which has been modified to inter-polarons hopping model for the conjugated polymers can be one of the possible models for the charge transport in these conjugated polymeric samples. In such conducting polymeric samples, the inter-polaron conduction will take place by the hopping between the polaron/bipolaron and their counterions. In this model, the temperature dependent dc conductivity obeys the power law of temperature given by the expression^34,35

σ_dc = AT^β (5)

log σ_dc = log A + β log T (6)

where, A is a constant and the index of power ‘β’ can has a value of ten or above. The conductivity data has been plotted as log σ_dc versus log T for all the samples with and without SnO₂ doping and the results are shown in Fig. 6(b). To evaluate the value of parameter ‘β’ all the curves have been fitted linearly (eqn (6)). The values of the slopes ‘β’ are ∼2.99, 3.72, 4.78 and 5.51 for samples S0, S1, S2 and S3, respectively. According to the Kivelson's model that parameter ‘β’ can has value about ten or above. However, in some conditions, it can be of the order of five. Thus, Kivelson's model can be applied for sample S3, but the temperature dependence of parameter ‘β’ rules out this possibility in sample S3 as well as in other samples. Thus, the lower values of parameter ‘β’ and its temperature dependence conflict with possible applicability of Kivelson's model in the present investigation. The band conduction may be other possible conduction mechanism whose applicability can be tested by plotting the activation energy as a function of temperature. The logarithmic DC conductivity (log σ_dc) plotted as a function of reciprocal of temperature show a nonlinear behavior for all samples (S0–S3). This discards the possibility of single activation energy at all temperatures. The temperature dependent activation energy has been evaluated from the slope of the curves of log σ_dc versus 1000/T (Fig. 6(a)) at various temperatures by using the expression¹⁸

Fig. 6 Variation of log σ_dc as a function of (a) 1000/T and (b) log T of the samples S0–S3.

The log E_A versus log T curves are shown in Fig. 7. The activation energies for all samples have shown temperature dependent behavior, which rules out the simple band conduction and nearest neighbor hopping conduction for these samples according to the Holstein considerations. According to Holstein³⁶ hopping or conduction model, the jumping of charge carriers or small polarons between the neighboring sites can principally occur when the unperturbed electronic energy of the adjacent sites coincide. Under such condition, the activation energy becomes equal to the minimum potential energy of the system and is substantially lower than the polaron binding energy. This indicates that the energy required to establish a neighbor coincidence must always be lower than the energy required for the thermal dissociation of the polaron. Thus, in the classical limits (hν ≪ k_BT), we may define a transition temperature (T_T), at which the polaron bandwidth becomes equal to ħ times the mean life time of the localized state. Thus, the temperature dependent conduction process can be divided into two parts. The higher temperature region can be described by the random site jump, whereas, the lower temperature region can be expressed by the polaron motion. The value of T_T lies in the range of 0.4 to 0.55 of the Debye temperature (θ). Here θ/2 can be defined as the temperature at which the log σ_dc versus 1/T curve deviates from the linearity in high temperature region or activation energy shows a temperature dependent behavior. Thus, in the high temperature region (>θ/2), the charge transport properties are dominated by the thermally activated hopping of small polarons where the band conduction is discarded by temperature dependent activation energy. Moreover, such polaronic hopping conduction can involve a single phonon or multiphonon interaction process as suggested by Holstein.³⁶ However, at lower temperatures, the multiphonon processes are gradually replaced by the single optical phonon absorption and emission processes. In this temperature region, only polarons contribute to the jump frequency³⁶ and the activation energy can be given by the Holstein relation as

where, E′_A is the calculated activation energy at different temperatures by using the room temperature activation energy (E_A) and phonon frequency ν₀ ≈ 10¹³ Hz. The calculated activation energy at various temperatures has also been plotted in Fig. 7 for sample S3, which shows a temperature dependent behavior. Such temperature dependent behavior of the activation energy (E′_A) discards the possibility of the polaronic conduction dominated by multiphonon processes. This supports the possibility of single phonon assisted Mott's variable range hopping (VRH) conduction mechanism in the present samples. To further investigate the behavior of the dc conductivity of these samples, the activation energies for these samples have been analyzed by using the following expression^{18,33,37–40}where, T_Mott is the Mott's characteristic temperature, which can possibly be related to the disorderness and found to depend on the electronic structure, hopping barrier and energy distribution of the electronic states. The parameter ‘γ’ is related with the dimensionality (d) of hopping (γ = 1/1 + d), where d can be 1, 2 or 3 for one, two or three dimensional hopping, respectively. The value of γ can be obtained from the slope (1 − γ) of log E_A versus log T curve using the relation (7), whereas, intersection of log E_A axis gives T_Mott. However, it can be seen from Fig. 7 that the slopes of the activation energy curves have temperature dependent nature. This changes with change in temperature which rules out the possibility of single dimensional hopping conduction mechanism. In higher temperature region, the samples may follow three dimensional (3D) hopping with γ ≈ 1/4, whereas, in mid temperature range, some deviation from the 3D hopping is observed and thus, samples may obey two dimensional hopping conduction. However, a large deviation from 3D hopping was observed in samples S1 and S3 in the lower temperature region. In such disordered materials, the conduction behavior can be given by random thermal movement of the charge carriers, where the density of states no longer remains uniform. The charge transport properties of the disordered materials like conjugated polymer or their hybrid nanostructures are strongly influenced by the distribution of energy states and doping concentration. The hopping conduction may be assisted by the Gaussian or exponential types of distribution of the energy states near the Fermi level. The electrostatic and steric engrossment of the dopant sites with the surrounding organic polymeric matrix may cause a random distribution of the density of electronic sites represented by Gaussian or exponential distribution. The random placement of the charged molecules in the polymer matrix may cause an increment in the Gaussian disorderness due to the dipolar interaction. The random distribution of these electronic states near the Fermi level may be the reason for temperature dependent deviation in the dimensionality of conduction mechanism. Thus, to make a further approximation on the temperature dependent hopping conduction in a wide range of temperatures, the measured dc conductivity has been plotted as a function of T^−1/2, T^−1/3 and T^−1/4 (Fig. 8(a–c)) and analyzed in the light of Mott's VRH model,⁴¹ where the dc conductivity (σ_dc) can be given by

Fig. 7 Variation of log E_A versus log T of the samples S0–S3.

Fig. 8 Variation of log σ_dc versus (a) T^−1/2 (b) T^−1/3 and (c) T^−1/4 for samples S0–S3.

It can be expected from eqn (11) that for some suitable value of γ, log σ_dc versus T^−γ curve should be a straight line and the linear fitting factor gives the dimensionality of the hopping. Therefore, all curves of log σ_dc versus T^−1/2, T^−1/3, T^−1/4 in Fig. 8(a–c) are fitted with the linear line. It is observed that the log σ_dc curves for all the samples are well fitted for γ = 1/4, which suggest that in a wide range of temperature the present samples might be analyzed for three dimensional hopping conduction. The conduction behavior of the samples can be further analyzed in terms of Mott's parameters estimated for three dimensional hopping (γ = 1/4).

The Mott's characteristic temperature (T_Mott) and pre-exponential factor (σ₀) can be evaluated from the slope and intersection of log σ_dc versus T^−1/4 curves, respectively, and these can be given by the expressions^{18,33,37–40}

σ₀ = R²e²ν₀N(E_F) (12)

where, N(E_F) is the density of states at the Fermi level, λ (=1/r_p) is the coefficient of exponential decay in the hopping process, r_p is the polaron radius and can have a value of ∼3 Å, as electrons are considered to be delocalized to the extent of π-orbitals on the monomer unit. R is the average hopping distance between two sites, which is given by

The average polaron hopping energy (W) for the present samples can be evaluated by knowing the values of N(E_F) and R by using the expression

The Mott's parameters T_Mott, N(E_F), R and W evaluated by using the expressions (13)–(15) for the three dimensional hopping in the PPy samples doped with different concentrations of SnO₂ QDs are summarized in Table 1.

Table 1 Mott's VRH parameters for samples S0–S3

Samples	S0	S1	S2	S3
TMott (K)	3.94 × 10^6	9.16 × 10^6	2.49 × 10^7	4.36 × 10^7
N(EF) (cm^−3 eV^−1)	1.98 × 10^21	8.49 × 10^20	3.13 × 10^20	1.79 × 10^20
R (cm)	12.04 × 10^−8	14.88 × 10^−8	19.09 × 10^−8	21.96 × 10^−8
W (meV)	69.15	85.45	109.68	126.19
αR	4.01	4.96	6.36	7.32
EA (meV)	74.94	106.63	125.92	161.43
σdc (S cm^−1) at 300 K	6.27 × 10^−2	1.52 × 10^−2	1.61 × 10^−3	5.29 × 10^−4
AB/AP	1.1499	0.8802	0.7877	0.6937

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It can be observed from Table 1 that the value of Mott's characteristic temperature increases with increase in SnO₂ doping concentration and decrease in the conductivity values of the samples. This indicates the strong localization of charge carriers in the sample having higher doping concentration (S3) accompanied with lower density of states {N(E_F)} at the Fermi level. The values of density of states (N(E_F)) in these samples are found to be in the range of 10²⁰ to 10²¹ cm⁻³ eV⁻¹. Thus, decrease in electrical conductivity of the samples can be attributed to the decrease in density of states with increased doping concentration of SnO₂ QDs. The average hopping distance (R) and the average hopping energy (W) are also dependent on the doping concentration of the SnO₂ QDs and are found to increase with increase in doping concentration. This indicates the requirement of higher hopping energy by the charge carrier to hop at a distant site. The Mott's parameters also satisfies the requirement of αR ≫ 1 and W ≫ k_BT, for three dimensional VRH to distant sites for charge transport. Thus, the estimated Mott's parameters and the temperature dependent activation energy suggest that the Mott's VRH model is well suited for the investigation of conduction behavior of PPy/SnO₂ nanocomposites.

C. Photoconduction behavior of PPy/SnO₂ nanocomposites

The photoconduction behavior of the samples has been investigated keeping in view of their applications in LDR. Fig. 9 shows the change in the electrical resistance of samples as a function of time in dark as well as under illumination of ∼100 mW cm⁻². The repeated cycles of illumination (30 s) and recovery (30 s) in dark have been recorded to verify the reproducibility of the results. The photoresponse in the present study is defined by the relative change in electrical resistance aswhere, R_{[small script l]} and R_d represent the resistances with and without illumination, respectively. The electrical resistance of all the samples was found to decrease when illuminated with ∼100 mW cm⁻² light intensity. The change in the electrical resistance of samples can be attributed to the photoexcited charge carriers assisted conduction and can be given by the expression⁴²where, q and μ are the electronic charge and mobility, [small script l], d and t are the length, width and thickness of the sample and Δn is the change in charge carrier concentration under illumination.

Fig. 9 Variation of photoconductive response for samples S0–S3 versus time at the illumination intensity of 100 mW cm⁻².

At a constant temperature, the electronic states of the prepared sample are in equilibrium with Fermi level in dark condition. However, the rise in temperature may disturb this thermal equilibrium of the electronic state and gives rise to a new equilibrium. Moreover, the applied electric field further perturbs the electronic equilibrium to set a net flow of charge carriers through the samples and it gives rise to dark current in the absence of external illumination. The interaction of light photon with material creates bound electron–hole pairs [hν → e⁻ + h⁺] and some of the electrons with sufficient energy jump over the potential barrier to make a transition from the valence band to conduction band.^43,44 The electrical resistance of the samples decrease with increase in illumination time due to increase in electron–hole pairs with exposure time (Fig. 9). However, as soon as light is turned off, all samples almost regain their original states. The nanoneedles like structures in samples having higher SnO₂ concentration may provide a larger surface area or active sites for the electron–hole pair generation via photon–phonon interaction, which may lead to the improvement in the photoresponse. The dynamic resistance of samples under light and dark conditions can be given by the relations

here, R₀ is the steady state resistance in dark, τ_res and τ_rec are the time constants of response and recovery, respectively. t_{[small script l]} (∼30 s for the present case) is the time required to achieve the lowest value of the resistance (R_{[small script l]} = R_min) and t_r (∼30 s for the present case) is the recovery time or the time required to achieve the original value of resistance (R_d = R₀). The response time is defined by the time taken by the detector to acquire 90% of the maximum resistance. The response and recovery times for the samples S0 and S3 have also been measured at 90% of the maximum response and 10% of the maximum response (fall to 90% response) after switching off light, respectively. The response and recovery times for S0 are 22.5 and 26 s, whereas, for S3 these times are 21 and 24.5 s, respectively. The photoresponse is governed by several types of transitions such as valence band to conduction band generally called band to band transition, defect state to conduction band or valence band to unoccupied states, etc. On the other hand, the recovery depends on the recombination process via de-excitation of the electrons to the ground state. The recombination of electrons can be protracted by different processes like the recombination of holes with negatively charged oxygen molecules available at the surface that creates the deficiency of holes for recombination of unpaired electrons. The recombination can only be possible after adsorption of new oxygen molecules which slow down the recovery of samples.

It is observed from Fig. 9 that the photoresponse of samples increase with increase in doping concentration of SnO₂ QDs in PPy matrix. The conversion of light energy into electrical energy by the photocells (or detected by the photo detectors) follows two processes viz. (i) absorption of radiation via an electrical excitation and (ii) separation of charge carriers. The separation of photo-generated charge carriers in a conventional p–n junction device is governed by the depletion region (junction built in potential) at the p–n junction.⁴⁵ The hybrid nanostructures of PPy/SnO₂ have shown photoresponse in the white light region and improved with SnO₂ doping. The enhanced photoresponse of PPy/SnO₂ nanocomposites can be attributed to the formation of a large number of p–n junctions. The pure PPy is a p-type electron donor organic material, whereas, pure SnO₂ is an n-type material with the electron accepting nature. These SnO₂ QDs are embedded throughout the amorphous matrix of PPy like seeds in watermelon. During the synthesis of nanocomposites, PPy chains wrap over SnO₂ QDs to form a large number of p–n junctions throughout the materials. The electron of SnO₂ QDs annihilates the holes of PPy molecules at the grain boundaries to form depletion layers, which make PPy electrically insulating, indicated by decreased dc conductivity. These p–n junctions formed in the donor–acceptor structure may augment depletion barrier height, resulting in local electric field at the junction (called junction potential) leading to enhanced photoresponse. The mechanism of enhanced photoresponse in PPy/SnO₂ nanocomposites can be tentatively understood. SnO₂ has a very poor conductivity at room temperature due to very small electron density and the positive charge carriers are distributed more sparsely towards SnO₂ in the depletion region of PPy/SnO₂ nanocomposites as compared to the PPy region. This may cause a wider depletion region towards the SnO₂ as compared to the PPy sheath. On exposure to white light, large numbers of electron–hole pairs are generated (corresponding to large number of p–n junctions) and are separated by the junction potentials. Free holes move to the PPy side, whereas, free electrons to the SnO₂. This results in an increase in hopping conductivity of nanocomposites. Once the generation of electron–hole pairs is counterbalanced by the recombination, no further change in the electrical conductivity occurs for particular illumination intensity. However, the change in the intensity may further change the electron–hole pair concentration and hence the electrical conductivity. The difference in the photoresponse for different doping concentration of SnO₂ can be attributed to an increase in p–n junction concentration and junction potentials. The photoresponse of sample has also been measured for different illumination intensities ranging from ∼30 to 100 mW cm⁻² (Fig. 10). The intensity of light source is increased by increasing the applied voltage at the source. It has been observed that the photoresponse of the samples increases almost linearly with increase in illumination intensity (Fig. 11(a)) which may be due to increase in concentration of p–n junctions. The response time shortens with illumination intensity. The response of sample S3 is found to increase from 2.64 to 6.13 with increase in light intensity from 30 to 100 mW cm⁻². The photoresponse of samples is found to increase almost linearly with SnO₂ doping concentration (Fig. 11(b)) and becomes slightly saturated at higher concentration. This may be due to dominant effect of SnO₂ over PPy. Thus, the non-toxic nature, reproducibility and reasonable photo-sensing response of PPy/SnO₂ nanocomposites can make them a potential candidate for white light detectors or LDR applications.

Fig. 10 Variation of photoconductive response for samples S0–S3 versus time at different illumination intensities varying from 30 to 100 mW cm⁻².

Fig. 11 Variation of photoconductive response for samples S0–S3 with (a) illumination intensity, and (b) SnO₂ doping concentration.

IV. Conclusions

Tin oxide QDs have been synthesized by surfactant assisted hydrothermal process, whereas, polypyrrole and its nanocomposite structures have been prepared by chemical oxidation polymerization method. The HRTEM studies suggest the formation of SnO₂ QDs, whereas, the SEM studies revealed the formation of PPy/SnO₂ nanowalls for lower doping concentration of SnO₂, which changed to long length needle like structures emanating from the nanowalls at higher doping concentrations. XRD studies confirmed the formation of nanostructured SnO₂ and PPy/SnO₂ nanocomposites. The temperature dependent conductivity measurements suggested that the conduction following the single phonon assisted three dimensional variable range hopping mechanism. The electrical conductivity of samples is found to decrease with increase in SnO₂ concentration in dark, which suggested the interfacial interaction between the PPy and SnO₂. The coupling between PPy and SnO₂ modify the band structure of PPy. The change in optical band gap of PPy with doping concentration of SnO₂ has been investigated by UV-Vis spectroscopic studies. The optical band gaps of pure SnO₂ QDs, PPy (S0) and nanocomposites samples S1 and S2 are found to ∼3.58, 2.25, 2.41 and 2.83 eV, respectively. The obtained band gaps of pure PPy and its nanocomposites are lying in visible region, which make them potential candidates for white light detectors. Due to formation of p–n junctions between PPy and SnO₂, the electronic structures of composite get changed, which causes to change in the photosensitivity of PPy/SnO₂ at different illumination intensities ranging from 30–100 mW cm⁻². The photoconductive response of the samples improved from 2.85 (S0) to 6.25% (S3) with increase in SnO₂ doping concentration from 0 to 20% at 100 mW cm⁻² light intensity. The doping concentration dependent white light photoresponse has been observed in the samples, which may lead to make them potential candidates for futuristic non-toxic LDRs.

Acknowledgements

One of the authors (IR) is grateful to the Principal, Kirori Mal College, University of Delhi, Delhi-110007 (India) for his support during the work. Mr R. K. Tripathi is grateful to the Ministry of New and Renewable Energy, Government of India for providing the senior research fellowship during the course of this work.

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