Influence of Hydrothermal Pressure during Crystallization on the

Aug 30, 2010 - Center for Micro/Nano Science and Technology, National Cheng Kung UniVersity, Tainan 70101, Taiwan. ReceiVed: July 2, 2010; ReVised ...
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J. Phys. Chem. C 2010, 114, 15625–15632

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Influence of Hydrothermal Pressure during Crystallization on the Structure and Electron-Conveying Ability of TiO2 Colloids for Dye-Sensitized Solar Cells Po-Tsung Hsiao,† Ming-De Lu,‡ Yung-Liang Tung,‡ and Hsisheng Teng*,†,§ Department of Chemical Engineering, National Cheng Kung UniVersity, Tainan 70101, Taiwan, Green Energy & EnVironment Research Laboratories, Industrial Technology Research Institute, Hsinchu 31040, Taiwan, and Center for Micro/Nano Science and Technology, National Cheng Kung UniVersity, Tainan 70101, Taiwan ReceiVed: July 2, 2010; ReVised Manuscript ReceiVed: August 13, 2010

This study synthesizes TiO2 anatase colloids at 250 °C under varying pressures of 57-120 bar by adjusting the residual volume in the autoclaving chamber. Transmission electron microscopy and X-ray diffraction analyses showed that the amorphous phase content of TiO2 powders and films obtained from calcining the colloids increased with the pressure during crystallization, whereas the Ti vacancy in the crystalline phase decreased. This illustrated a trade-off between lattice distortion and vacancy reduction as a result of an increase in pressure during crystallization. X-ray absorption fine structure spectroscopic analysis showed that the coordination number of the Ti4+ sites in the TiO2 increased with the pressure during crystallization to reach a maximum value at 100 bar and then decreased with further increases in pressure. A dye-sensitized solar cell assembled with a TiO2 film from a 100-bar synthesis exhibited the highest solar energy conversion efficiency. Electrochemical impedance spectroscopy analysis showed that the 100-bar film had the highest charge collection efficiency for photogenerated electrons. From these results, we concluded that the TiO2 crystallization pressure affects the density of defects in the produced TiO2 films and, therefore, the electronconveying performance in DSSCs. Introduction Titania (TiO2) colloids, or nanoparticles, can be used in a variety of applications for photovoltaic, photocatalytic, and gassensing systems.1-10 Preparing nanosized TiO2 colloids very commonly involves sol-gel chemistry, which includes hydrolysis, precipitate formation, sol peptization, and, finally, hydrothermal crystallization.11 The final hydrothermal crystallization step is generally conducted in an autoclave and involves a mechanism of crystal growth at the cost of smaller crystals, that is, the Ostward ripening. The pressure during autoclaving influences the crystal growth mechanism and, therefore, the crystal structure of the resulting colloids.12-16 Using TiO2 as an electron conductor in a photovoltaic device, such as a dyesensitized solar cell (DSSC), the crystal structure of TiO2 governs the electron-transport rates and thus the photovoltaic performance of the cells.17-19 How the pressure during hydrothermal crystallization affects the crystal structure of TiO2 colloids and their electron-conveying ability requires in-depth investigation. Titania exists in three polymorphs: anatase, rutile, and brookite. Rutile is the thermodynamically stable phase, whereas anatase and brookite are metastable and their growth may be governed by crystallization kinetics.13 The crystal structure in anatase has more inter-TiO6 octahedron connections and is the preferred phase for applications in DSSCs.20-23 To obtain TiO2 anatase colloids from sol-gel synthesis, titanium alkoxide is very commonly used as a precursor. During hydrothermal crystallization in an autoclave, CO2 gas evolves as a biproduct * To whom correspondence should be addressed. E-mail: hteng@ mail.ncku.edu.tw. Fax: 886-6-2344496. † Department of Chemical Engineering, National Cheng Kung University. ‡ Industrial Technology Research Institute. § Center for Micro/Nano Science and Technology, National Cheng Kung University.

Figure 1. Schematic configuration of an autoclave for hydrothermal crystallization of TiO2 colloids.

from decomposition of alkoxide, and therefore, the internal pressure of the autoclave increases. Under these circumstances, the residual volume in the autoclave, that is, the difference between the volume of the autoclave and the volume of reacting liquid (Figure 1), determines the pressure during TiO2 colloid crystallization. Our preliminary experiments have observed that the residual volume had substantially influenced the pressure during hydrothermal crystallization, although this aspect was rarely discussed in previous studies. One of the merits of the autoclaving treatment is the pressurized environment that helps to improve colloidal crystallization. The saturated vapor pressure in an autoclave is dependent on the temperature and irrelevant to the residual volume. CO2 evolution provides the additional pressure, which increases with a decrease in residual volume. In a pressurized system, the rutile phase is thermodynamically preferred because the structure of rutile is more compact (with a density of 4.275 g cm-3) than that of anatase (with a density of 3.893 g cm-3). Crystallization of TiO2 anatase colloids in a pressurized autoclave may very likely induce transformation of the metastable

10.1021/jp1061013  2010 American Chemical Society Published on Web 08/30/2010

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anatase phase to the stable rutile. Thus, phase transition must be taken into account when varying the pressure during crystallization. Previous studies have reported that the anatase-to-rutile transformation passed through a persistent amorphous phase and the activity of titania was dependent on the amorphous content due to the formation of localized electronic states.24,25 For the applications of TiO2 in photovoltaics, the amorphicity of the TiO2 may govern the characteristics of electron transport.26,27 By employing the X-ray diffraction technique in combination with the Rietveld refinement simulation, the present study intends to quantitatively characterize the crystallinity of TiO2 obtained from varying pressures during crystallization. We also employed X-ray absorption spectroscopy to probe the average coordination environment of the Ti4+ sites in the TiO2 specimens. In fact, X-ray absorption spectroscopy that is atomspecific and capable of probing the short-to-medium range structure around an absorbing atom is a complement of X-ray diffraction that requires a long ordered domain for coherent scattering with no detection of an amorphous phase.28 Within this scope, we presented the results of an extensive investigation on the structure and electron-conveying ability of TiO2 anatase nanocrystalline films obtained from TiO2 colloids with varying pressures during crystallization. The results indicated that exerting pressure increased the integrity of the TiO2 framework stoichiometry while it also increased the tendency of phase-transition to cause amorphous phase formation. The corresponding influence of pressure during crystallization on the behavior of electron transport in DSSCs is interpreted using electrochemical impedance spectroscopy. This study demonstrated that optimizing the pressure during crystallization for TiO2 colloid formation is essential to assuring efficient electron conduction in DSSCs. Experimental Section This study synthesized TiO2 colloids via sol-gel hydrolysis precipitation of titanium isopropoxide (Ti(OC3H7)4), which was followed by hydrothermal crystallization in an autoclave under varying pressures. In hydrolysis, titanium isopropoxide and acetic acid were mixed at a CH3COOH/Ti molar ratio of 1/1 at room temperature. This solution was then mixed with water and stirred for 1 h to form a white precipitate. After adding an appropriate amount of nitric acid, the mixture was heated to 80 °C and stirred vigorously for 2 h, to achieve peptization. After peptization, the resultant slurry was hydrothermally treated in a titanium-lined autoclave at 250 °C for 12 h to give a TiO2 colloid solution.29 The vapor pressure of water at 250 °C was 40 bar. By varying the residual volume in the autoclave, the pressure was sustained at three different pressures of 57, 100, and 120 bar during the hydrothermal treatment. To prepare TiO2 suspensions, the TiO2 solution was alternatively washed with ethanol and centrifuged for three cycles to obtain a solution containing 40 wt % TiO2 in ethanol. The TiO2/ ethanol solution was then blended with ethyl cellulose powders, followed by evaporation of the ethanol with a rotary evaporator to obtain 18 wt % TiO2 paste. Finally, the paste was treated with a three-roll mill to improve the dispersion of the colloids. Each sample of viscous TiO2 paste was calcined at 450 °C in air for 30 min, transforming into a nanoparticle powder for structural analysis. To prepare TiO2 films, the paste was coated onto a fluorinedoped SnO2 (FTO) conducting glass substrate by using a screenprinting technique. The TiO2-coated substrate was subsequently calcined at 450 °C for 30 min to form a mesoporous electrode.

Hsiao et al. The apparent area of the films was 0.28 cm2, and the thickness of the TiO2 layers was 17 µm. To assess the change in the crystal structure of TiO2 nanoparticles after thermal necking into films, we also subjected the fragments from scraping the TiO2 films to structural analysis. Powder X-ray diffraction (XRD) using a Rigaku RINT2000 diffractometer with Cu KR radiation at 40 kV and 40 mA was the technique used to identify the phase of TiO2 specimens. Data were collected with a step interval of 0.02° and a measuring time of 10 s per point in the 2θ range of 20-70°. Crystalline structures were refined with the Rietveld technique using the Generalized Structure and Analysis Software (GSAS) package.30-32 The microstructure of the TiO2 nanocrystalline film was explored with a Hitachi FE-2000 high-resolution transmission electron microscope (HRTEM) operated at 200 kV. The structural features and chemical environment of the Ti4+ sites on the TiO2 specimens were characterized with X-ray absorption spectra, which were recorded at room temperature in a transmission mode on the Wiggler beamline of the Taiwan Synchrotron Radiation Research Center. A double-crystal Si(111) monochromator was used for Ti K-edge (4.966 keV) experiments, and X-rays were generated from the 1.5 GeV electron storage ring with a current range of 120-200 mA. The spectrum scans, including X-ray absorption near-edge structure (XANES) and extended X-ray absorption fine structure (EXAFS), were collected from -200 to 1000 eV relative to the energy at the transition edges. Analysis of the EXAFS data was performed with UWXAFS 3.0 and FEFF 7.0 programs.33 Fourier transform was used on k3-weighted EXAFS oscillation in the range of 2.6-12 Å-1. Each shell fitting was conducted in the R-space. To prepare a dye-covered TiO2 film for a DSSC, a calcined TiO2 film (17 µm) was immersed in a 0.5 mM N719 dye solution in a mixture of acetonitrile and tert-butyl alcohol with a 1/1 volume ratio for 24 h. To determine the amount of dye covered on each TiO2 film, we placed the sensitized film into a 10 mM KOH solution to desorb the dye and then measured the amount of the desorbed dye with absorption spectroscopy (U-4100, Hitachi) at 502.2 nm. This analysis showed that all the films had a similar amount of dye loading. The dye-covered TiO2 film was then assembled with a Pt-coated conducting glass using a 60 µm thick thermoplastic frame in a sandwich type. The electrolyte composition for the DSSC was as follows: 0.1 M LiI, 0.05 M I2, 0.6 M 1,2-dimethyl-3-n-propylimidazolium iodide, and 0.5 M 4-tert-butylpyridine in acetonitrile. Photovoltaic measurements of the DSSCs employed an AM 1.5 solar simulator (sp91160A-4739, Newport), and the intensity of the simulated light was calibrated as 100 mW cm-2 by using a reference Si solar cell. The electrochemical impedance spectra of the cells were measured with a potentiostat equipped with a frequency response analyzer (IM6, Zahner), with the frequency range of 0.05-105 Hz. The bias potential was set at 0.7 V with an ac potential amplitude of 10 mV under an AM 1.5 solar illumination of 100 mW cm-2. Results and Discussion Figure 2 shows the HRTEM images of the TiO2 nanocrystalline film specimens obtained from TiO2 colloids crystallized at varying pressures with subsequent calcination at 450 °C. The codes TC1, TC2, and TC3 denote the specimens obtained from crystallization at 57, 100, and 120 bar, respectively. All three specimens were composed primarily of crystalline TiO2 nanoparticles of 15-20 nm in size. The crystalline nanoparticles showed visible lattice fringes. The inset of each panel in Figure

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Figure 3. Rietveld refinement plots for the powder (a-c) and film (d-f) specimens. Experimental data and calculated curves are indicated by crosses and continuous lines, respectively. All of the TiO2 specimens contain the pure phase of anatase, and the tick marks correspond to anatase. The difference curve is displayed near the bottom of the graph.

Figure 2. HRTEM images of TiO2 films: (a) TC1, crystallized at 57 bar; (b) TC2, crystallized at 100 bar; (c) TC3, crystallized at 120 bar. The insets show the fast Fourier transform patterns projected along the anatase [111] zone axis.

2 shows the selected-area fast Fourier transform (FFT) pattern of the crystalline TiO2 that was identified as being in the anatase phase. The FFT patterns along the anatase [111] zone axis show clear diffraction spots from the (01j1) and (1j01) planes, both of which were calculated to have a spacing distance of 3.517 Å. This analysis indicates that varying the hydrothermal pressure did not change the crystalline phase. In addition to the crystalline TiO2 nanoparticles, some amorphous particles without visible lattice fringes were present in the specimens. As stated above, the presence of the amorphous TiO2 particles may affect the electron-transport behavior in the TiO2 films and, therefore, the performance of DSSC. Figure 3 compares the crystal structures of the TiO2 powders and films. Figure 3a-c shows the XRD patterns and the corresponding Rietveld simulation results of the TiO2 powders. All three TiO2 specimens were of phase-pure anatase and had nearly identical XRD patterns. Rietveld refinement analysis on the patterns gave the unit cell parameters of the anatase constituting the crystalline nanoparticles. The refined data listed in Table 1 are essentially correct because of the satisfactory reliability factors (Rp and Rwp). The factors Rp and Rwp stand for the regression sum of relative errors and the regression sum

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TABLE 1: Anatase Unit Cell Parameters for the TiO2 Powders and Films of Varying Pressures during Crystallization (57, 100, and 120 bar for TC1, TC2, and TC3, Respectively) Obtained from Rietveld Refinements Using the GSAS Package TiO2 specimen

´ a (Å)

´ c (Å)

TC1 TC2 TC3

3.784 3.784 3.785

9.497 9.499 9.500

TC1 TC2 TC3

3.784 3.784 3.785

9.497 9.499 9.500

standard anatase

3.785

9.513

Ti vacancy

O vacancy

Ti-Ti (Å)

Rp/Rwp (%/%)

powders 0.057 0.042 0.037

0.027 0.016 0.030

3.035 3.036 3.037

8.14/10.74 8.39/11.17 8.69/11.38

films 0.058 0.043 0.037

0.039 0.029 0.046

3.035 3.036 3.037

8.09/10.72 8.71/11.43 7.50/10.21

0

3.040

model compound 0

of weighted squared errors, respectively.30 Table 1 shows little difference in the unit cell parameters of all the three specimens from those of the standard anatase given at the bottom of Table 1. However, the Ti and O vacancies were not so small as to be negligible. The vacancy in Ti decreased with an increase in pressure during crystallization, reflecting the increase of the chemical intactness of the nanoparticles. The O vacancy decreases with the pressure during crystallization to reach a minimum at 100 bar, indicating that pressurization improved the intactness of the crystal. However, a further pressure increase led to an O vacancy increase that may be associated with anatase lattice distortion or amorphous phase formation, which will be discussed later. Anatase, being constructed of distorted, edge sharing TiO6 octhedra, has the highest degree of polyhedral edge-sharing of all the titania polymorphs. Figure 4 illustrates that the short Ti-Ti contact propagates in the c direction with the strong repulsion between Ti4+ cations, compensated for by a shortening of shared edges and an expansion of unshared edges.34 Introducing Ti vacancy relieved Ti-Ti repulsion across the shared edges, resulting in more regular octahedra and contraction along the c direction.34-36 In fact, the TiO2 powder data in Table 1 shows a concomitant dilation of the length of the c axis (from 9.497 to 9.500 Å) and Ti-Ti bond distance (from 3.035 to 3.037 Å) when the pressure during crystallization was increased to reduce the Ti vacancy (from 0.057 to 0.037). As to amorphous phase formation, previous studies had reported that the amorphous content, as well as the Ti occupancy, increased as the calcination temperature increased due to the increased tendency of anataseto-rutile transformation.24,25 However, the amorphous phase was invisible in XRD patterns. This study analyzed the content of the amorphous phase for the TiO2 specimens using the XRD refinement technique with the introduction of a known proportion of an internal standard crystalline powder (see the Supporting Information).28 Quantitative analysis revealed that the contents of amorphous TiO2 are 22, 26,

Figure 4. [010] projection of the anatase structure featuring the arrangement of TiO6 octahedra.

and 29 wt % for TC1, TC2, and TC3, respectively, for both the powder and the film specimens. Similar to the reported temperature effect,24,25 increasing pressure during crystallization led to amorphous phase formation because of the persistent presence of amorphous TiO2 during anatase-to-rutile transformation. Figure 3d-f shows the XRD data of the TiO2 films that were obtained by thermally necking the TiO2 nanoparticles. These films were also of the phase-pure anatase. Table 1 shows that the anatase unit cell parameters of the films were almost identical to those of the powders, except that the films had a slight higher O vacancy. The higher vacancy can most likely be attributed to disorder created in the crystal lattice during the thermal necking of nanoparticles into films.37 Following the trend for the powder specimens, the TC2 film, which was composed of the 100-bar-crystallized colloids, had the lowest O vacancy. This study analyzed the entire coordination environment of the Ti4+ sites on the TiO2 specimens with X-ray absorption fine structure spectroscopy. Figure 5a shows the Ti K-edge XANES patterns of the TiO2 films. The patterns of the powder specimens were similar to those of the film specimens. In the pre-edge region, each specimen showed three obvious peaks labeled A1, A3, and B (Figure 5b-d). An additional peak labeled A2 is present as a weak shoulder on the low-energy side of the A3 peak. The A2 feature of TiO2 is only present in the XANES of the anatase phase.38 These pre-edge peaks correspond to the transition of the core electrons to higher energy levels that are made up of crystal-field split Ti 3d orbitals hybridized with Ti 4p orbitals.39,40 Feature A1 and features A2 and A3 are attributed to quadrupolar and dipolar transitions, respectively, to the t2g orbital of TiO6 octahedron. Feature B represents a dipolar transition to the eg orbital. The intensity of the B peak and the intensity ratio of A2/A3 serve as a sensitive probe for the degree of distortion in the TiO6 octahedron.41,42 The present study decomposed the pre-edge patterns using Gaussian fittings, and Figure 5b-d shows the individual peaks and simulated patterns in dashed lines. The B peak intensities for the TC1, TC2, and TC3 films were 0.37, 0.33, and 0.43, respectively, and the intensity ratios of A2/A3 were 0.31, 0.25, and 0.4. The TC2 film had the lowest values for both the B peak intensity and the A2/ A3 ratio, indicating the lowest degree of structural distortion in the TC2 film. Figure 6 shows the Fourier transforms (FTs) of k3χ(k) EXAFS for the TiO2 powders and films with the radial distance ranging within 3.6 Å. The spectra in the figure have been corrected for phase shifts. Table 2 shows the bond distance (R) and coordination number (CN) of two shells (Ti-O and Ti-Ti) obtained by FEFF simulation in a range of 1-3.6 Å. The CNs of the specimens were smaller than those of standard anatase, which are also provided in Table 2. The lowering of the CNs can be attributed to the presence of the amorphous phase that was

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Figure 5. (a) Ti K-edge XANES spectra of the TiO2 film specimens. The pre-edge features of the XANES spectra: (b) TC1, (c) TC2, (d) TC3. The dashed lines in panels b-d indicate the Gaussian-fitting curve that is composed of individual peaks A1, A2, A3, and B.

TABLE 2: Coordination Number (CN) and Bond Distance (R) of the Ti4+ Sites in the TiO2 Powders and Films of Varying Pressures during Crystallization (57, 100, and 120 bar for TC1, TC2, and TC3, Respectively) Obtained from EXAFS Analysis σ2 (Å2)

Ti-O Ti-Ti Ti-O Ti-Ti Ti-O Ti-Ti

powders 5.38 3.24 5.54 3.15 5.33 2.95

1.962 3.028 1.952 3.027 1.944 3.026

0.005 0.003 0.005 0.003 0.005 0.003

Ti-O Ti-Ti Ti-O Ti-Ti Ti-O Ti-Ti

films 5.18 3.02 5.33 2.95 5.07 2.66

1.962 3.028 1.953 3.027 1.944 3.026

0.005 0.003 0.006 0.003 0.007 0.003

TC1

TC3

TC1 TC2 TC3

detected by XRD. The CNs were smaller for the films relative to those of the powders, indicating that the thermal necking between nanoparticles had enhanced the structural distortion.37,43 For both the powder and the film specimens, the 100-barcrystallized TC2 gave the highest CN for Ti-O, in agreement with the results obtained from the XRD and XANES analyses. This indicated that a pressure increase for crystallization reduced the ionic vacancy to promote the CN of Ti-O, whereas an overpressurization leads to significant structural distortion because of extensive amorphous phase formation. The CN of the Ti-Ti shell decreased with the pressure during crystallization because of the increasing amorphous phase content. The decreasing Ti-O bond length with the pressure could be associated with the tendency of phase transition toward rutile because the Ti-O bond had contracted for anatase transformation to rutile.

R (Å)

shell

TC2

Figure 6. Fourier transformed k3χ(k) EXAFS spectra of the TiO2 powder and film specimens. The dashed line curves denote the best fitting of the spectra in the range of 1-3.6 Å.

CN

TiO2 specimen

It must be pointed out that the structural parameters obtained from the X-ray absorption analysis comprise those of the crystalline and amorphous phases, differing from the refined XRD data where the information of non-Bragg diffracting amorphous TiO2 cannot be obtained. By assuming that the crystalline phase had a CN of 6 for Ti-O, the EXAFS data in combination with the amorphous phase contents could be used to calculate the CN of the amorphous phase for each TiO2 specimen. Figure 7 shows the variation in CN with the pressure during crystallization for the powder and film specimens. Both types of specimens showed a similar trend in variation of the CN. At a low pressure of 57 bar, the CN of the amorphous phase was low because of the high ionic vacancies in TiO2. When the pressure was increased to 100 bar, a significant increase in the CN occurred because of the reduced ionic

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Figure 7. Variation of coordination number (CN) of Ti-O with the pressure during crystallization for the TiO2 amorphous phase in the powder and film specimens.

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Figure 9. Photocurrent-voltage characteristics of DSSCs assembled with the TiO2 films (TC1, TC2, and TC3) under AM 1.5 solar illumination at 100 mW cm-2. All the TiO2 films had a thickness of 17 µm. η and ff denote the solar energy conversion efficiency and fill factor, respectively.

Figure 8. Schematic showing how a maximum CN for Ti4+ ions in TiO2 could be reached by optimizing the pressure during crystallization, based on the trade-off between vacancy reduction and distortion promotion when the pressure during crystallization is increased.

vacancies. With a further increase in the pressure, the CN of the amorphous phase decreased because the ionic vacancies reached an asymptotic value while the tendency of anatase lattice distortion still increased with the pressure. Figure 8 shows a schematic to summarize the trade-off between vacancy reduction and distortion promotion because of the pressure increase during crystallization. Optimization of pressure during crystallization was necessary to give a maximum CN for Ti ions in the TiO2 powders or films. How the structure varied with the pressure would eventually influence the electron-conveying ability has yet to be explored. To evaluate the ability to convey electrons, the present study sensitized the TiO2 films of 17 µm in thickness with a ruthenium complex dye (N719) and assembled them into DSSCs. Figure 9 shows the photocurrent-voltage characteristics of the DSSCs under AM 1.5-type solar illumination at 100 mW cm-2. The TC2 cell exhibited the highest short-circuit current density (Jsc), outperforming the TC3 cell by ca. 12%. The light-to-electric energy conversion efficiency (η) and the fill factor (ff) for these cells were also given in the figure. The TC2 cell showed the highest conversion efficiency of 7.7% because of its large Jsc. This efficiency was ca. 14% higher than that of the TC3 cell. The value of Jsc is associated with electron-transport dynamics in the TiO2 film. The high CN of the Ti4+ sites in the TC2 film must have played a role in facilitating electron transport in the TC2 film. We subjected the DSSCs to electrochemical impedance spectroscopy analysis. Figure 10a shows the Nyquist impedance

Figure 10. (a) Nyquist impedance plots of DSSCs assembled with the TiO2 films (17 µm) of varying pressures during crystallization under AM 1.5 solar illumination at 100 mW cm-2. Bias voltages were set at 0.7 V, with the frequency range being 0.05-105 Hz. The solid lines represent curves calculated using the parameters in Table 3. (b) General equivalent circuit of DSSCs. rt, rct, and cµ correspond to the electrontransport resistance, interfacial charge-transfer resistance, and chemical capacitance in the TiO2 film. Rs is the sheet resistance of the conducting glass, RPt and CPt are the charge-transfer resistance and capacitance at the Pt/electrolyte interface, respectively, and ZN is the Nernst impedance.

spectra for the cells at a fixed bias of 0.7 V under 100 mW cm-2 solar illumination. Each spectrum comprises three arcs, which are associated with the charge transfer at the Pt/electrolyte interface, the electron transport through the mesoporous TiO2 film, and the Nernstian diffusion of triiodide in the electrolyte, respectively, in the high-, middle-, and low-frequency regimes of the spectrum.44-46 An equivalent circuit (Figure 10b) corresponding to the transmission line model was employed to interpret the impedance spectra.47,48 The elements of the circuit relating to TiO2 film are the electron-transport resistance, Rt () rtL), the interfacial charge recombination resistance, Rct () rct/ L), and the chemical capacitance produced by the accumulation

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TABLE 3: Equivalent-Circuit Parameters for the DSSCs Assembled with TiO2 Films at Varying Pressures during Crystallization (57, 100, and 120 bar for TC1, TC2, and TC3, Respectively) Obtained at 0.7 V under 100 mW cm-2 Solar Illumination TiO2 specimen

Cµ (µF)

rt (Ω µm-1)

rct (Ω µm)

ηcc (%)

TC1 TC2 TC3

844 789 947

0.28 0.25 0.30

264 255 269

70 72 68

of electrons in the TiO2 film, Cµ () cµL), with L being the TiO2 film thickness. Other elements in the equivalent circuit include the sheet resistance of the conducting glass, Rs, the chargetransfer resistance and interfacial capacitance at the Pt/electrolyte interface, RPt and CPt, respectively, and, finally, the Nernst impedance, ZN, which describes the diffusion behavior of the redox species in the electrolyte. The electron transport in the TiO2 film is the major concern. Table 3 shows the values of Cµ, rt, and rct obtained from the equivalent-circuit simulation on the middle-frequency arc. The TC2 TiO2 film had the lowest rt value. This can be attributed to the fact that the TC2 film had a high CN for Ti that led to a low density of defect states. A low CN would lead to formation of localized Ti3+ defect states that act as trap states for conduction electrons.49,50 Therefore, the TC2 film is the most effective in electron transport. The TC2 film also had the lowest value in chemical capacitance, which corresponds to the total density of the electrons in the conduction band and in the trap states.44,51 In fact, the chemical capacitance increased with a decrease in CN of Ti-O. This supported the above argument that the density of trap states was closely related to the CN of Ti-O determined by EXAFS analysis. Meanwhile, the TC2 cell also had the smallest rct value, probably due to effective diffusion toward the TiO2/electrolyte interface for recombination.52 Both the rt and the rct values were to be incorporated to evaluate the efficiency of charge collection. Because the photogenerated electrons diffuse forward and recombine with I3-, the electron collection rate at the FTO substrate is

1 1 1 1 1 ) ) Rcc Rt Rct rtL rct /L

(1)

where Rcc is the apparent resistance for electron collection. Accordingly, the electron collection efficiency (ηcc) can be written as53,54

ηcc )

1 Rcc 1 1 + Rcc Rct

)1-

rt 2 L rct

(2)

Table 3 shows the calculated ηcc values, which have an order of TC2 > TC1 > TC3, identical to the order for the CN of Ti-O. The ηcc results indicated that the defect states not only impede the electron transport in TiO2 film but also increase the probability of losing electrons to recombination. Therefore, a fine-tuning of pressure during crystallization for TiO2 colloids is critical for minimizing the defect density of the electronconveying TiO2 film in DSSCs.

Conclusions This study demonstrated that the pressure during crystallization of TiO2 colloids from sol-gel synthesis influenced the coordination of Ti ions and the electron-conveying performance of the consequent TiO2 films in DSSCs. The TiO2 anatase powder and film obtained from calcination of sol-gel synthesized TiO2 colloids contained a substantial proportion of amorphous phase TiO2, which is regarded as a persistent intermediate in the transformation of anatase to rutile. Quantitative XRD analysis showed that the fraction of the amorphous phase increased with the pressure during crystallization, whereas the Ti vacancy in the anatase crystallites decreased. This indicated that high pressure during crystallization had enhanced the structural intactness of the anatase as well as induced phase transformation. As a result of the competition between the enhanced structural intactness and phase transformation, an optimal pressure during crystallization of 100 bar gave the lowest O vacancy for the anatase phase in the consequent calcined TiO2 powders or films. X-ray absorption fine structure spectroscopic analysis also showed that a high CN for Ti-O, corresponding to a low defect density in the powders or films, occurred for the TiO2 specimens synthesized at the optimal pressure during crystallization. The DSSCs assembled with the low-defect TiO2 films exhibited a high solar energy conversion efficiency, principally resulting from a high photocurrent. Electrochemical impedance spectroscopic analysis showed that the efficiency of charge collection in TiO2 films had increased with the increasing CN of Ti-O. This study showed that DSSCs using TiO2 synthesized at an optimized pressure could outperform those using TiO2 of other pressures by ca. 14% in terms of solar energy conversion efficiency. Acknowledgment. This research is supported by the National Science Council of Taiwan (98-2221-E-006-110-MY3, 98-2622E-006-012-CC2, 98-3114-E-007-011, 98-3114-E-007-005, and 98-2221-E-006-112-MY2); the Bureau of Energy, Ministry of Economic Affairs, Taiwan (98-D0204-2); and the Green Energy & Environment Research Laboratories of the Industrial Technology Research Institute, Taiwan. We thank Dr. Jyh-Fu Lee of the Taiwan Synchrotron Radiation Research Center for his help with the analysis of the X-ray absorption data. Supporting Information Available: Quantitative determination of the amorphous phase content in TiO2 specimens. This material is available free of charge via the Internet at http:// pubs.acs.org. References and Notes (1) O’Regan, B.; Gra¨tzel, M. Nature 1991, 353, 737. (2) Adachi, M.; Murata, Y.; Takao, J.; Jiu, J.; Sakamoto, M.; Wang, F. J. Am. Chem. Soc. 2004, 126, 14943. (3) Wang, K. P.; Teng, H. S. Phys. Chem. Chem. Phys. 2009, 11, 9489. (4) Li, T. L.; Teng, H. S. J. Mater. Chem. 2010, 20, 3656. (5) Banerjee, S.; Mohapatra, S. K.; Das, P. P.; Misra, M. Chem. Mater. 2008, 20, 6784. (6) Sawatsuk, T.; Chindaduang, A.; Sae-Kung, C.; Pratontep, S.; Tumcharern, G. Diamond Relat. Mater. 2009, 18, 524. (7) Fang, X. M.; Zhang, Z. G.; Chen, Q. L.; Ji, H. B.; Gao, X. N. J. Solid State Chem. 2007, 180, 1325. (8) Chaudhary, Y. S.; Chinthalapelly, D.; Bhat, U. M.; Nayak, P. K.; Khushalani, D. J. Mater. Chem. 2008, 18, 3636. (9) Bidaye, P. P.; Khushalani, D.; Fernandes, J. B. Catal. Lett. 2010, 134, 169. (10) Chen, Q.; Xu, D.; Wu, Z.; Liu, Z. Nanotechnology 2008, 19, 365708. (11) Wang, C. C.; Ying, J. Y. Chem. Mater. 1999, 11, 3113. (12) Barbe´, C. J.; Arendse, F.; Comte, P.; Jirousek, M.; Lenzmann, F.; Shklover, V.; Gra¨tzel, M. J. Am. Ceram. Soc. 1997, 80, 3157.

15632

J. Phys. Chem. C, Vol. 114, No. 37, 2010

(13) Testino, A.; Bellobono, I. R.; Buscaglia, V.; Canevali, C.; D’Arienzo, M.; Polizzi, S.; Scotti, R.; Morazzoni, F. J. Am. Chem. Soc. 2007, 129, 3564. (14) Tsai, C. C.; Teng, H. S. Chem. Mater. 2004, 16, 4352. (15) Yin, H.; Wada, Y.; Kitamura, T.; Kambe, S.; Murasawa, S.; Mori, H.; Sakata, T.; Yanagida, S. J. Mater. Chem. 2001, 11, 1694. (16) Das, J.; Freitas, F. S.; Evans, I. R.; Nogueira, A. F.; Khushalani, D. J. Mater. Chem. 2010, 20, 4425. (17) Cass, M. J.; Walker, A. B.; Martinez, D.; Peter, L. M. J. Phys. Chem. B 2005, 109, 5100. (18) Nakada, S.; Saito, Y.; Kubo, W.; Kitamura, T.; Wada, Y.; Yanagida, S. J. Phys. Chem. B 2003, 107, 8607. (19) Wang, K. P.; Teng, H. S. Appl. Phys. Lett. 2007, 91, 173102. (20) Hsiao, P. T.; Wang, K. P.; Cheng, C. W.; Teng, H. S. J. Photochem. Photobiol., A 2007, 188, 19. (21) Xu, Y. M.; Fang, X. M.; Zhang, Z. G. Appl. Surf. Sci. 2009, 255, 8743. (22) Kurata, T.; Mori, Y.; Isoda, S.; Jiu, J. T.; Tsuchiya, K.; Uchida, F.; Adachi, M. Curr. Nanosci. 2010, 6, 269. (23) Lee, C. K.; Lyu, M. D.; Liu, S. S.; Chen, H. C. J. Taiwan Inst. Chem. Eng. 2009, 40, 463. (24) Lim, S. H.; Phonthammachai, N.; Liu, T.; White, T. J. J. Appl. Crystallogr. 2008, 41, 1009. (25) Lim, S. H.; Ritter, C.; Ping, Y.; Schreyer, M.; White, T. J. J. Appl. Crystallogr. 2009, 42, 917. (26) Paronyan, T. M.; Kechiantz, A. M.; Lin, M. C. Nanotechnology 2008, 19, 115201. (27) Zhang, D. S.; Yoshida, T.; Minoura, H. AdV. Mater. 2003, 15, 814. (28) Orlhac, X.; Fillet, C.; Deniard, P.; Dulac, A. M.; Brec, R. J. Appl. Crystallogr. 2001, 34, 114. (29) Ito, S.; Murakami, T. N.; Comte, P.; Liska, P.; Gra¨tzel, C.; Nazeeruddin, M. K.; Gra¨tzel, M. Thin Solid Films 2008, 516, 4613. (30) Rietveld, H. M. J. Appl. Crystallogr. 1969, 2, 65. (31) McCusker, L. B.; Von Dreele, R. B.; Cox, D. E.; Loue¨r, D.; Scardi, P. J. Appl. Crystallogr. 1999, 32, 36. (32) Toby, B. H. J. Appl. Crystallogr. 2001, 34, 210. (33) Stern, E. A.; Newville, M.; Ravel, B.; Yacoby, Y.; Haskel, D. Physica B 1995, 208, 117. (34) Lim, S. H.; Ferraris, C.; Schreyer, M.; Shih, K.; Leckie, J. O.; White, T. J. J. Solid State Chem. 2007, 180, 2905.

Hsiao et al. (35) Bastow, T. J.; Moodie, A. F.; Smith, M. E.; Whitfield, H. J. J. Mater. Chem. 1993, 3, 697. (36) Grey, I. E. N.; Wilson, C. J. Solid State Chem. 2007, 180, 707. (37) Hsiao, P. T.; Teng, H. S. J. Am. Ceram. Soc. 2009, 92, 888. (38) Luca, V.; Djajanti, S.; Howe, R. F. J. Phys. Chem. B 1998, 102, 10650. (39) Wu, Z. Y.; Ouvrard, G.; Gressier, P.; Natoli, C. R. Phys. ReV. 1997, B55, 10382. (40) Parlebas, J. C.; Khan, M. A.; Uozumi, T.; Okada, K.; Kotani, A. J. Electron Spectrosc. Relat. Phenom. 1995, 71, 117. (41) Pillep, B.; Fro¨ba, M.; Phillips, M. L. F.; Wong, J.; Stucky, G. D.; Behrens, P. Solid State Commum. 1997, 103, 203. (42) Pickup, D. M.; Neel, E. A. A.; Moss, R. M.; Wetherall, K. M.; Guerry, P.; Smith, M. E.; Knowles, J. C.; Newport, R. J. J. Mater. Sci.: Mater. Med. 2008, 19, 1681. (43) Hsiao, P. T.; Tung, Y. L.; Teng, H. S. J. Phys. Chem. C 2010, 114, 6762. (44) Kern, R.; Sastrawan, R.; Ferber, J.; Stangl, R.; Luther, J. Electrochim. Acta 2002, 47, 4213. (45) Febregat-Santiago, F.; Bisquert, J.; Palomares, E.; Otero, L.; Kuang, D.; Zakeeruddin, S. M.; Gra¨tzel, M. J. Phys. Chem. C 2007, 111, 6550. (46) Adachi, M.; Sakamoto, M.; Jiu, J.; Ogata, Y.; Isoda, S. J. Phys. Chem. B 2006, 110, 13872. (47) Bisquert, J. J. Phys. Chem. B 2002, 106, 325. (48) Frbregat-Santiago, F.; Bisquert, J.; Garcia-Belmonte, G.; Boschloo, G.; Hagfeldt, A. Sol. Energy Mater. Sol. Cells 2005, 87, 117. (49) Mackrodt, W. C.; Simson, E. A.; Harrison, N. M. Surf. Sci. 1997, 384, 192. (50) Ko, K. H.; Lee, Y. C.; Jung, Y. J. J. Colloid Interface Sci. 2005, 283, 482. (51) Bisquert, J. Phys. Chem. Chem. Phys. 2003, 5, 5360. (52) Kopidakis, N.; Benkstein, K. D.; van de Lagemaat, J.; Frank, A. J. J. Phys. Chem. B 2003, 107, 11307. (53) Wang, Q.; Ito, S.; Gra¨tzel, M.; Febregat-Santiago, F.; Mora-Sero´, I.; Bisquert, J.; Bessho, T.; Imai, H. J. Phys. Chem. B 2006, 110, 25210. (54) van de Lagemaat, J.; Park, N.-G.; Frank, A. J. J. Phys. Chem. B 2002, 104, 2044.

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