Article pubs.acs.org/JPCC
Control of Recombination Pathways in TiO2 Nanowire Hybrid Solar Cells Using Sn4+ Dopants James A. Dorman,† Jonas Weickert,† Julian B. Reindl,† Martin Putnik,† Andreas Wisnet,‡ Matthias Noebels,† Christina Scheu,‡ and Lukas Schmidt-Mende*,† †
Department of Physics, University of Konstanz, Universitätsstr. 10, Konstanz D-78457, Germany Department of Chemistry and Center for NanoScience (CeNS) Munich, Ludwig Maximilian University, Butenandtstr. 11, 81377 Munich, Germany
‡
ABSTRACT: Hybrid nanostructures have shown increasing potential as a replacement for Si solar cells due to the availability of low-cost material combinations. However, up to now, hybrid solar cells, where photon absorption occurs in a semiconducting polymer and charge separation occurs at a metal oxide-polymer interface, show limited efficiencies. One limitation is caused by a relative low charge carrier mobility in the metal oxide. Here we addressed this issue and describe the use of a Sn:TiO2|TiO2 core−shell nanowire array to increase the charge-carrier mobility in the core of the nanowires while decreasing the charge-carrier recombination at the metal oxide−polymer interface due to fast electron extraction from this interface, driven by a cascaded conduction band energy from shell to core of the nanowires. These doped cores with an undoped shell structure resulted in impressive efficiency improvement in hybrid solar cells of 33% over the reference TiO2-based device. Additionally, this device structure resulted in a 17% increase in recombination lifetimes based on both photovoltage decay measurements and impedance spectroscopy. Recombination mechanisms are proposed for the core and core−shell systems to highlight the various effects of the Sn4+-doped TiO2 nanowire arrays. Doped core−shell structures have the potential for application in the hybrid-type devices without the limitations that are seen with the current dual metal oxide structures due to the seamless interface of the metal oxide host for direct transport of the electrons into the highly mobile core material.
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both hole and electron mobilities within the system,10 meaning the mobilities must be balanced for optimum device performance,11 they must not be viewed as a simultaneous optimization task. Therefore, it is important to understand the electron dynamics within the TiO2 system to reduce the effect of charge recombination while increasing the charge extraction efficiency. An alternative to the mesoporous TiO2 films that has been studied recently are rutile TiO2 nanowire arrays12 synthesized on fluorine-doped tin oxide (FTO) using a hydrothermal growth method, as first reported by Liu et al.13 The growth of these 1D arrays are specific to FTO due to the crystal facets of the substrate limiting the growth of the films along the [001] axis.13 Because of the single-crystalline nature of the nanowires, the mobility of the electrons is on the order of 1 cm2 V−1 s−1.14 The increase in mobility is due to a lack of grain boundaries and wire orientation, increasing the electron mobility while shuttling the electron to the electrode, respectively.15 These
INTRODUCTION While hybrid inorganic/organic based solar cells have been able to achieve impressive power conversion efficiencies in the past few years, between 13 and 15% for dye-sensitized1 and mesostructured absorbers,2,3 they are still limited due to many recombination processes that occur at the materials interface. The organic portion of these hybrid devices can vary from a dye electrolyte combination to a pure polymer or metal organic perovskites. However, the standard for the inorganic framework is a mesoporous titanium oxide (TiO2) film,1,2,4−6 providing a high-surface-area material for both light absorption and charge separation. Unfortunately, after the charges are separated, they must be rapidly transported away from the interface between the hole transporter material and TiO2 surface before recombination occurs.7 Furthermore, the electron mobility of TiO2 is relatively low, 10−1 cm2 V−1 s−1, for single-crystal anatase,8 which is further decreased due to the high number of grain boundaries that the charges must cross. These low electron mobilities, and subsequent higher recombination probabilities, have been cited as one of the major limitations of using TiO2 for all types of metal-oxidebased solar cells,9 that is, hybrid, dye-sensitized, and perovskitetype solar cells. While the electron dynamics are a function of © 2014 American Chemical Society
Special Issue: Michael Grätzel Festschrift Received: December 26, 2013 Revised: March 14, 2014 Published: March 14, 2014 16672
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substrates in a Teflon-lined autoclave with 15 mL of H2O, 15 mL of concentrated HCl (Sigma-Aldrich), and 525 μL of titanium tert-butoxide (Sigma-Aldrich, >99%) after stirring for 15 min. The autoclaves were placed in an oven for 3 h at 180 °C for nanowire growth. After the reaction was complete, the autoclaves were rapidly cooled in a water bath. The nanowire arrays were then removed from the solution and washed in DI H2O three times to remove any excess salt ions. Sn4+ doping was achieved through the addition of tin tert-butoxide (SigmaAldrich, >99%) dissolved in concentrated HCl in a 2:1 ratio, with volumes between 20 and 160 μL for various dopant concentrations in solution. The Sn:HCl precursor was added directly to the growth solution, after the Ti precursor, using the same reaction conditions as previously described. The Sn:TiO2 growth performed here is slightly modified from that reported by Xu et al.26 Dopant concentrations referred to in this study are based on the mol percentage in the initial reaction solution because the concentration in the nanowires is too low for accurate measurements. TiCl4 Surface Treatment. To remove surface defects and form the core−shell architecture, a TiCl4 (Sigma-Aldrich) surface treatment was performed after the hydrothermal growth. The thickness of the TiO2 shell can be controlled by varying the TiCl4 concentration in the growth solution. For this study, the TiCl4 concentration was set at 100 mM to form a thick shell layer, which completely covered the wire surface. All samples were submerged in the TiCl4 solution and placed in a water bath for a uniform heating profile. The solution was heated to 70 °C and allowed to react for 3 h. Finally, the shell was converted to the oxide form by annealing the wires at 450 °C for 30 min. The wire arrays were allowed to cool naturally over a period of ∼2 h to reduce any stress that may form after rapid cooling. Solar-Cell Fabrication. Solar-cell fabrication proceeded by immersing the surface-treated nanowire arrays in squaraine dye (SQ2) (Mitsubishi Chemicals) dissolved in ethanol (0.2 M) for 3 h. The SQ2 dye was chosen to extract charge from both the interfacial and hole transport materials.27,28 Excess dye was removed by rinsing the nanowire arrays with isopropanol and dried under a N2 gas. The samples were then pretreated in chlorobenzene (Sigma-Aldrich) before a poly(3-hexylthiophene-2,5-diyl) (P3HT, RMI-001EE Rieke materials), 40 mg/ mL in chlorobenzene, was spin coated. The P3HT-coated samples were heated to 120 °C for 5 min in ambient air to remove any excess solvent from the film. Finally, a 130 nm Ag film was thermally evaporated through a shadow mask to complete the device fabrication. Device Characterization. The structure and composition of the nanowire arrays were characterized using a scanning electron microscopy (SEM) (Zeiss Neon 40 EsB) operated at 5 keV accelerating voltage. High-resolution transmission electron microscopy measurements were collected using a Jeol JEM 2011 operated at 200 kV and a FEI Titan (S)TEM 80−300 operated at 300 kV. Samples were suspended on a 400-mesh carbon-coated TEM grid. Additionally, the crystal structure was characterized using the Al Kα emission line of an X-ray diffraction (XRD). Nanowire films were scanned at a rate of 1°/s for 2θ between 20 and 70°. UV−vis measurements were performed in a Cary 5000 series UV−vis-NIR spectrometer (Agilent Technologies). Samples were placed in the center of an attached integrating sphere and masked using a Teflon sample holder with a 0.78 cm2 hole for uniform sample illumination. The absorption spectra were scanned from 800 to
hydrothermally grown nanowires have diameters between 80 and 100 nm, based on growth conditions, and scatter light very well, creating an opaque film on the FTO.13 Despite their appealing properties, these nanowire arrays have been used in all different types of hybrid solar cells with limited success thus far. Specifically, these arrays have been combined with poly(3hexylthiophene-2,5-diyl) (P3HT) or CH3NH3PbI3 perovskites and produced conversion efficiencies of 1.5 and 9.4%, respectively.16,17 Although the latter value is impressive, it is still significantly lower than power conversion efficiencies achieved on TiO2 nanoparticle films, which exceed 14%.2 Alternatively, a single-crystal, anatase, mesoporous thin film was combined with the perovskite absorber and showed conversion efficiencies of 7.3%.18 Comparison of the various crystal structures shows that while the mobility of the structure is important for high-efficiency solar cells, the recombination must be addressed for further improvements. It is believed that the recombination processes can be reduced by controlling the position of the electrons within the nanowire arrays. First proposed by Law et. al, it is possible to direct the electrons toward the core of the nanowire by creating a cascading conduction band, where the core of the wire has a slightly lower conduction band than the shell.19 To date, devices with these cascading energy levels have been produced using a highly conducting core, such as ZnO or SnO2, with a TiO2 shell layer.19−22 However, device performance has been limited for core−shell dye-sensitized solar cells with efficiencies around 2 to 3%.19,20,23 An alternative approach to the formation of the cascading band structure is to dope the TiO2 core with Sn4+, which is an attractive dopant due to the resulting high mobility, lower conduction band, increased electron density, and similar crystal structure.24 This combination has allowed the fabrication of high-sensitivity gas sensors using a 5 mol % Sn concentration in TiO2.25 More recently, Xu et. al have shown that the synthesis of Sn-doped TiO2 nanowire arrays is possible using a modified hydrothermal growth process, resulting in photocurrents greater than 2 mA cm−2.26 However, no work has been performed on the Sn-doped TiO2 core in solar-cell devices, specifically with a TiO2 shell layer, forming this cascading conduction band and directing the electrons away from the interface. This paper focuses on the impact of these cascading energy levels in the Sn-doped TiO2 core−shell nanowires arrays on the charge recombination and overall photon conversion efficiency of the devices. The TiO2 and Sn:TiO2|TiO2 core−shell nanowire arrays are synthesized using a two-step process consisting of a rutile core, via the hydrothermal growth, and a thin shell, deposited from a TiCl4 bath. Electron−hole recombination is probed using multiple spectroscopic techniques to accurately elucidate the charge kinetics and dynamics for the nanowire arrays. A competing recombination pathway is proposed for the doped core−shell structure due to the cascading electronic structure to visualize how the recombination processes are controlled. Finally, the effects of the Sn4+ dopant on the device performance are reported, resulting in a significant efficiency improvement.
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EXPERIMENTAL METHODS Hydrothermal Growth. TiO2 nanowires were grown on FTO substrates (Solaronix, 15 Ω/□) cut into 1.4 × 1.4 cm2 squares. The FTO substrates were first ultrasonicated in water with dishwashing detergent, acetone, and isopropanol for 30 min. Nanowire growth proceeded by placing the FTO 16673
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300 nm at a rate of 1 nm s−1 and slit size of 5 nm and referenced to 100 and 0% transmission. Additionally, the spectra were normalized so that the absorption at 800 nm was 0. Current density−voltage (J−V) and external quantum efficiency (EQE) measurements were performed using a Keithley 2400 SourceMeter controlled through a self-written LabView program. Cells were illuminated via a LOT-Oriel LS0106 solar simulator (AM 1.5G, 100 mW cm−2) through a shadow mask with a resulting active area of 0.125 cm2. Light intensities were calibrated with a certified Si reference solar cell (Fraunhofer Institute) with a KG5 filter. To control the illumination wavelength for the EQE measurements, light from a 150 W Xe lamp was passed through a LOT-Oriel Omni 150 monochromator. All EQE spectra were normalized to the measured J−V current for accurate comparison. For the lightintensity measurements, the illumination light was passed through a series of neutral density filters before illumination to obtain light intensities between 10 and 95 mW cm−2. For photovoltage decay (PVD) and photocurrent decay (PCD) measurements, a pulsed laser (18 Hz, 532 nm) was focused through the shadow mask onto the sample. The sample was background-illuminated with a LOT-Oriel LS0106 solar simulator with variable light intensities (10−95 mW cm−2). Signals were recorded with a Tekscope DPO 7254 digital oscilloscope. Termination resistances 1 MΩ and 50 Ω were used for PVD and PCD measurements, respectively. Data were fitted with monoexponential decays to estimate the characteristic decay lifetimes. Impedance spectroscopy (IS) measurements were conducted with an Ecochemie Autolab potentiostat/galvanostat. The samples were placed in a dark, grounded metal box, and the spectra were collected with and without a white light diode at ∼100 mW cm−2. Additionally, the measurements were performed at biases between −0.2 to 0.6 V with a small AC perturbation using frequencies between 0.1 Hz and 1 MHz. The resulting spectra were modeled using ZView software based on circuit model from Boix et. al.29
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RESULTS AND DISCUSSION The nanowire arrays were synthesized using the hydrothermal method and measured to have diameters of roughly 80−100 nm and lengths of ∼1 μm after a 3 h growth period (Figure 1a). Additionally, the nanowire arrays were indexed to the rutile crystal phase (JCPDS no. 21-1276) using XRD (Figure 1b). The wires are single crystalline and form a compact TiO2 layer at the FTO surface before the wires are elongated in the [001] growth direction. Furthermore, no impurities can be seen in the EDX spectrum, indicating that high-quality nanowires were achieved. Finally, no changes in the nanowire structures were observed after the incorporation of low concentrations of the Sn precursor into the growth solution. A TEM image is shown in Figure 1c to highlight the nonconformal but completely encapsulated with up to a 2 nm TiCl4 shell layer with the presence of nanoparticles around 5 nm in diameter. The thickness of this layer can be varied by adjusting the concentration and growth time. It is believed that this combined shell layer plays a key role in the device performance, discussed later, because the thin shell layer separates the core from the P3HT, forming an intermediate energy level in the conduction band, while increasing the surface area of the structure with the formation of the nanoparticles. However, the importance of the shell conformity and surface area is not fully
Figure 1. (a) Side view of the TiO2 nanowire arrays with a height of ∼1 μm and diameters around 80−100 nm. (b) Nanowires can be index to the rutile crystal phase (JCPDS no. 21-1276). The sample was scanned while still on the FTO substrate, which was also indexed for reference. (c) HRTEM image shows the TiO2 coating that is achieved after the 100 mM TiCl4 treatment. The film is nonconformal but covers the wires completely with thicknesses between 1 and 2 nm. The HRTEM image highlights the lattice planes of both the rutile core (3.24 Å, JCPDS no. 21-1276) and the anatase shell layer (3.51 Å, JCPDS no. 21-1272).
understood and will be studied at a later date in combination with atomic layer deposition. The core and shell layer are indexed to rutile (3.24 Å) and anatase (3.51 Å, JCPDS no. 211272) based on the lattice spacing, respectively. No difference was observed in the morphology of the Sn:TiO2 and pure TiO2 nanowire structures after the hydrothermal growth. Because of the low concentration of incorporated Sn, as reported by Xu et. al for a pure SnCl4 additive, it was not possible to measure the concentration of the Sn4+ in the nanowire array. However, based on the XPS results reported by Xu, it is believed that the 16674
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indicate that the Sn4+ drastically increases the device performance due to an increase in photocurrent, while the open-circuit voltage (VOC) is relatively unchanged for the uncoated samples. The pure TiO2 device resulted in the worst device performance due to low VOC and short-circuit current (JSC), while the incorporation of the Sn4+ dopant increased the generated current while maintaining the measured VOC. The improvement in the JSC is attributed to an increase in charge density within the wires due to the Sn4+ dopant, which is shown in the IS results discussed later. However, the device performance can be further improved after the formation of a core−shell Sn:TiO2| TiO2 device, which produces currents around 7 mA cm−2 and VOC values around 0.65 V, which is in the range of mesoporous TiO2−P3HT devices.28 The increase in photocurrent is attributed purely to the Sn4+ because no structural change was observed, as mentioned previously. Furthermore, the increase was shown to be independent of the morphology based on the absorption spectra (Figure 2b). It can be seen in the Figure that there is little difference between the SQ2 absorption (550−720 nm) and TiO2 absorption (
