J. Phys. Chem. C 2008, 112, 5213-5220
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Evaluation of Nitrogen Doping of Tungsten Oxide for Photoelectrochemical Water Splitting Brian Cole,*,† Bjorn Marsen,† Eric Miller,† Yanfa Yan,‡ Bobby To,‡ Kim Jones,‡ and Mowafak Al-Jassim‡ Hawaii Natural Institute, School of Ocean and Earth Science and Technology, UniVersity of Hawaii at Manoa, Honolulu, Hawaii, and National Renewable Energy Laboratory, Golden, Colorado ReceiVed: September 21, 2007; In Final Form: NoVember 30, 2007
Thin films of tungsten oxide were fabricated by reactive RF magnetron sputtering for applications related to the direct solar splitting of water. To investigate band gap reduction by nitrogen doping, films were deposited with nitrogen introduced into the sputtering ambient at partial pressures in the range of 0-6 mTorr N2. For dilute doping (pN2 < 2 mTorr), the films exhibit a small increase in band gap, and show a degradation of the carrier collection efficiency. Structurally, X-ray diffraction indicates a decrease in crystallinity with nitrogen incorporation. For higher levels of nitrogen doping (pN2 > 3 mTorr), the diffraction pattern shows the evolution of a new phase that shows an increase in scattering power relative to the pure WO3. For these samples, a reduction of the optical band gap to 3 mTorr). The interpretation of this shift in OCP is complex. For a high quality material, this shift would indicate a movement of the conduction band minimum to more negative values, where additional potential shifts relative to the pure WO3 would be required to generate photocurrents. Complicating the analysis, however, a high level of bulk or surface recombination can result in a reduction in the photogenerated voltage, unfavorably altering the OCP. For the 6 mTorr sample, not only is the OCP shifted to higher bias, the slope for the photocurrent onset is shallow. This behavior is consistent with an increase in resistance for the film, attributed to either a decrease in carrier density or mobility. As nitrogen is a group V atom, a substitutional doping on a tungsten site can be expected to impact the carrier density as nitrogen will compensate for n-type nature of WO3. Any defects
Figure 2. Absorption overlap with the AM1.5 global solar spectrum. Shaded region is absorption for the pure WO3. The hatched region is additional absorption as a result of nitrogen doping.
associated with nitrogen doping will impact the mobility. It was found (Table 2) that as nitrogen was incorporated, the resistance of the film increased. The resistance was shown to increase by a factor of 36 for the 6mTorr sample relative to the pure WO3. Comparing the Jsat to the Jmax as a function of nitrogen incorporation, we find a continual decrease in the collection efficiency. Although a significant band gap reduction was realized, the additional carriers generated upon illumination cannot be extracted. From these initial experiments, it is hypothesized that the electronic transport properties are degraded as a result of a general decrease in the crystallinity, i.e., smaller grains, more grain boundaries, and an associated higher concentration of defects. This will be examined in detail in the following sections that will attempt to identify structural aspects of the films. C. Morphology and Structure of WO3 Films. Figure 4 shows the morphology of the films in plan view for the case of
Figure 3. Current vs potential scans measured under AM1.5G simulated solar spectrum in 0.33 M H3PO4.
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Cole et al.
Figure 4. SEM images showing evolution of morphology for dilute nitrogen doping.
dilute doping. For each sample, the crystallites are well defined with a platy morphology. We estimate the grain size to be in the regime of 100 nm, with no appreciable variation as nitrogen is incorporated. To investigate the structure, the samples were evaluated by X-ray diffraction (Figure 5). The pure WO3 has the characteristic X-ray pattern of monoclinic tungsten oxide (PDF card# 43-1035). The monoclinic WO3 falls into the general classification of a perovskite structure with octahedral coordination of tungsten and oxygen atoms. The lattice can be described as a distorted cubic ReO3 structure, assembled via corner sharing of the octahedra. This distortion results in a doubling of the lattice parameter along all three axes to generate the unit cell.23 The three peaks at ∼23° 2θ are then attributed to the (002), (020), and (200) lattice planes of monoclinic WO3. Ideally, the scattering intensity from each set of lattice planes should be equivalent. However, we observe a significant increase in intensity for the (020) planes. It is expected that some part of this increase is due to preferred orientation that can be expected for a thin film deposition. To test for preferred orientation, the film was removed from the substrate and remeasured as a powder sample. As suspected, the powder measurement verified the presence of preferred orientation along [020] but not enough to account for entire increase in intensity for the (020) planes. It is suspected that the observed increase in intensity for the (020) planes relative to the (002) and (200) can be attributed to a small amount of {102} crystallographic shear, as is commonly found in crystalline WO3.24,25 Table 3 gives the integrated intensity and FWHM for the (002), (020), and (200) reflections as the nitrogen is varied from 0 to 6 mTorr. Each sample has comparable thickness, and therefore we choose to compare raw counts, as this may highlight relative changes between samples. As nitrogen is incorporated into the material, broadening is observed indicating that the material is becoming less crystalline. At partial pressures
Figure 5. XRD for dilute nitrogen doping. (a) survey over 15-60° 2θ and (b) detailed examination of the (002), (020), and (200) reflections indicating a decrease in crystallinity as nitrogen in incorporated into the film.
greater than 0.5 mTorr, the (002) reflection becomes extinct and the (200) is severely broadened. The broadening can be
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Figure 6. SEM images showing the microstructure for higher dopant concentrations. Close inspection of the 6 mTorr sample shows that the agglomerates contain acicular crystals having preferred orientation within the agglomerate.
TABLE 3: Analysis of Diffraction Data for Nitrogen Doped WO3 p (N2) (mTorr) 0 0.5 1 2 3 4 6
θ002
I002
θ020
I020
θ200
I200
22.85 22.86
537 104
23.49 23.51 23.50 23.48 23.35 23.78 24.02
1271 2351 3204 2550 1996 1610 874
24.20 24.10 23.98 23.89 23.93
223 553 825 1271 2033
evaluated via the Scherrer formula to provide insight on the crystal size along all three axes
Dhkl )
0.9λ B cos θB
where D is the size of the grain along a particular axis, λ is the wavelength of the X-ray radiation (1.542 Å), B is the full width half-maximum of the reflection, and θB is the Bragg diffraction angle. The results for this calculation are included in Table 3. As indicated by the Scherrer analysis, the crystal size along [200] has a relatively limited number of planes, which is consistent with the plate-like grains observed in the electron microscopy of Figure 4. Figure 6 shows the morphology for the samples over the larger range of doping concentration. To illustrate the evolution, the 0 mTorr sample is also included here. It is quite clear, that at a critical dopant concentration of 3 mTorr N2, a change in the physical chemistry occurs. A mixture of phases is observed, with agglomerates of nanocrystalline material interdispersed with larger grains of the variety seen for dilute doping. Comparing the end members in the series, the large, well-defined grains found for pure tungsten oxide have devolved into agglomerates
θ100
23.01 22.99 23.02
I100
2937 3989 4459
D002 (Å)
D020 (Å)
D200 (Å)
514 417
600 427 388 287 211 56 43
287 207 140 116 76
D100 (Å)
507 605 680
of nanoparticles. Close inspection of the 6 mTorr sample shows that these nanoparticles have an interesting morphology. The agglomerates are formed from acicular crystals that are roughly orientated along a common direction. Figure 7 shows the evolution of the X-ray pattern over this range of higher nitrogen pressure. Here the relative changes in X-ray data correlate well with observed changes in the morphology. At 3 mTorr, the contributions from the monoclinic (002), (020), and (200) planes are diffuse and severely broadened. However, a new peak at 23.0° 2θ appears, and grows in intensity with higher nitrogen concentration. At 6 mTorr, this new phase shows a single class of reflections that is distinct from the monoclinic WO3. The reflection at 23.0° 2θ has significant intensity and a narrow width measuring 0.12°. An index of (100) is assigned to this plane. The second peak at 47.1° 2θ corresponds to the higher order (200) reflection. The absence of (110) and (111) reflections would indicate a high degree of preferred orientation. Comparing integrated intensities, we observe that the scattering power for the (100) planes is greater than the (002), (020), and (200) reflections of the pure WO3. This is counterintuitive to the apparent decrease in crystallinity observed via SEM. This interesting diffraction pattern will be analyzed in greater detail in the discussions section.
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Figure 7. XRD for the higher concentration nitrogen samples. (a) survey over 15-60° 2θ and (b) close examination showing the evolution of a prominent reflection at 23.0° 2θ.
To investigate sub-structure for nitrogen doping, two samples were examined by transmission electron microscopy (pure WO3, and 6 mTorr nitrogen doping). Figure 8 shows the results for the HR-TEM. For the pure WO3 sample, the lattice is observed to have minimal defects, with the electron diffraction pattern indicating a high level of crystallinity. For the 6 mTorr nitrogen sample, the high-resolution imaging shows a large amorphous contribution. In distinct regions, lattice planes are observed only along a single direction. The electron diffraction pattern within this area shows smearing in the direction perpendicular to these planes, indicating that the material has a highly defective lattice. Discussion Tungsten oxide derives its n-type character from the presence of oxygen vacancies within the material. These vacancies provide electrons, which are considered free and move through the material via polaronic transport.26 The favorable transport properties result in appreciable photocurrents for the highly crystalline pure WO3 thin films. A well-known observation for WO3 is that as the substoichiometry increases, the additional vacancies are accommodated within the lattice by a transformation where the corner shared octahedra collapse to edge sharing, thereby eliminating the vacancy. In the bulk crystal, these edge sharing point defects have high mobility and can coalesce to form planes of shear, termed crystallographic shear (CS) planes. The formation of CS planes in WO3 has been well documented.27-29 Dependent on the degree of substoichiometry, these dislocations are energetically favored to form over preferred directions, and at regular intervals. As an example, for a small degree of
Cole et al. substoichiometry (WO3-x where x < 0.03), the preferred CS plane is over the [102] direction.30 The presence of these CS planes will impact the X-ray data. The lattice shift associated with the CS plane will provide extinction for the (002) and (200) reflections. Layering of the tungsten planes via corner shared octahedral along [020] is unaffected. This type of defect is consistent with the X-ray data for the pure WO3 samples, indicating that some CS is present. The formation of CS planes has important implications on the transport properties in tungsten oxide. Lattice defects in the form of crystallographic shear planes have been shown to degrade the electron transport properties. The degradation is attributed to trapping of carriers at the CS plane, and also a change in mobility across the defect plane from polaronic transport to hopping.24 For our polycrystalline thin films, we expect these same mechanisms for defect accommodation to exist, gathering to form grain boundaries. An increase in defect concentration associated with nitrogen incorporation then leads to smaller grains with an associated degradation of transport properties across the grain boundaries. This decrease in mobility correlates well with the experimentally observed increase in resistivity for the films as nitrogen in incorporated. This is also evident in the PEC curves for the nitrogen samples as the shallow slope for the photocurrent onset is indicative of poor carrier transport. For higher levels of nitrogen doping (6 mTorr), it is difficult to reconcile the evolution of an apparent highly crystalline X-ray pattern with the nanocrystalline WO3 agglomerates imaged in SEM. Evaluating the XRD pattern, only two peaks of appreciable intensity are found. Indices of (100) and (200) have been assigned to these peaks. The following interpretation is made: (1) The increase in integrated intensity for the (100) reflection as nitrogen is incorporated suggests a significant decrease in distortion for the unit cell. The tungsten atoms in the pure monoclinic phase form puckered planes, which provides for some extinction of the diffracted intensity. As the distortion (puckering) is reduced, an increase in integrated intensity can be expected as tungsten atoms are doing a better job of lying on a plane. (2) The absence of (110) and (111) planes indicates a high degree of preferred orientation along [100]. However, making a powder sample from the film has no effect on the diffraction pattern, indicating that preferred orientation is not the cause for the absence of (110) and (111) planes. (3) The single lattice parameter of 3.87 Å for the 6 mTorr sample would suggest that the film might be a cubic sodium tungsten bronze (NaxWO3),31 having a high degree of preferred orientation along [100]. We eliminate this possibility for the following reasons: (a) The presence of Na is not observed via XPS, (b) the dramatic color changes associated with the incorporation of Na (from deep blue, to black, to red, and finally bronze as x increases from 0 to 1 is not observed, and (c) preferred orientation is not suspected for this film. (4) The narrow width of the peak would indicate a large correlation length (680 Å as evaluated via the Scherrer formula). As previously mentioned, close inspection of the agglomerate in the 6mTorr sample shows the presence of acicular crystals orientated along a common direction. These crystals have a significant number of lattice planes along this single axis of preferred growth, which is consistent with the observed X-ray pattern. Is the new phase cubic? It has been theorized that, in materials such as tungsten oxide and zirconia, the high-symmetry cubic
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Figure 8. TEM images of pure WO3 (a,b) and the film deposited under a partial pressure for nitrogen of 6 mTorr (c,d). High-resolution lattice planes are imaged for (b) pure WO3 and (d) 6 mTorr nitrogen. The insets in (b) and (d) show SAED patterns corresponding to the high-resolution images shown here.
phase can be stabilized at room temperature via the introduction of static disorder to the lattice.32,33 This “static disorder” results from defects that are assumed to randomly distribute through the lattice and are localized to a particular site. These pinned defects are shown to reduce the distortion of the unit cell. In the case of zirconia, a higher symmetry cubic phase stabilized at room temperature was claimed, albeit with a diffraction pattern that more closely matched a simulated cubic pattern with a high temperature correction factor.32 For the nitrogen doped WO3, the increase in diffraction intensity would suggest that the distortion is significantly reduced, and by this theory of static disorder, a move toward cubic symmetry. The absence of 110 and 111 planes is explained by the morphology of the crystallites, where the high aspect ratio results in an insufficient number of planes along directions other than [100] to provide for diffraction. We therefore conclude that preferred growth of acicular crystals along [100] with a highly defective lattice within the plane perpendicular to [100] is consistent with the summation of all results. It follows that the formation of the highly defective crystallites negatively impacts the PEC. Although we realize a decrease in band gap for nitrogen doping, the associated degradation of transport properties prevents the extraction of the carriers generated upon illumination. The increase of lattice defects associated with nitrogen doping results in the degradation of the transport properties, resulting in poor saturated photocurrents and a shallow slope for the photocurrent onset. The observed increase in OCP with nitrogen doping is therefore attributed to the increase in defects and the associated increase in carrier recombination. Conclusions Nitrogen doping of tungsten oxide was evaluated for applications as a photoanode for the solar powered splitting of water.
For very small amounts of nitrogen incorporation, a minimal impact on the band gap was observed, however the saturated photocurrents were severely degraded. The decreased photocurrents were attributed to a degradation of transport properties that prevented the extraction photogenerated carriers from the photoanode. For higher concentrations of nitrogen doping, a change in the physical chemistry of the tungsten oxide occurred. The growth of agglomerates having acicular nanocrystals was observed, with the X-ray pattern indicating a decrease in the distortion for WO3 octahedra. Evaluation of the nitrogen doped WO3 by TEM indicates a highly defective lattice. These defects represent a significant barrier to the realization of an efficient photoanode based on the nitrogen doped WO3 material. Acknowledgment. This work was supported by the U.S. Department of Energy via the University of Nevada Las Vegas Research Foundation, Grant RF-05-SGHR-005. The authors thank Zey Tong for experimental assistance. References and Notes (1) Granqvist, C. G. Sol. Energy Mater. Sol. Cells 2000, 60, 201. (2) Smith, D. J.; Vatelino, J. F.; Falconer, R. S.; Wittman, E. L. Sens. Actuators B 1993, 13, 264. (3) Butler, M. A. J. Appl. Phys. 1977, 48 1914. (4) Hodes, G.; Cahen, D.; Manassen, J. Nature 1976, 260, 312. (5) Desilvestro, J.; Gratzel, M. J. Electroanal. Chem. 1987, 238, 129. (6) Miller, E. L.; Marsen, B.; Cole, B.; Lum, M. Electrochem. Solid State Lett. 2006, 9, G248. (7) Miller, E. L.; Rocheleau, R. E.; Deng, X. M. Int. J. Hydrogen Energy 2003, 28, 615. (8) Santato, C.; Ulmann, M.; Augustynski, J. J. Phys. Chem. B 2001, 105, 936. (9) Berek, J. M.; Sienko, M. J. J. Solid State Chem. 1970, 2, 109. (10) Butler, M. A.; Nasby, R. D.; Quinn, R. K. Solid State Comm. 1976, 19, 1011. (11) Scaife, D. E. Solar Energy 1980, 25, 41. (12) Gratzel, M. Nature 2001, 414, 338.
5220 J. Phys. Chem. C, Vol. 112, No. 13, 2008 (13) Miller, E. L.; Marsen, B.; Paluselli, D.; Rocheleau, R. Electrochem. Solid State Lett. 2005, 8, A247. (14) Green, M.; Hussain, Z. J. Appl. Phys. 1991, 69, 7788. (15) Sato, S. Chem. Phys. Lett. 1986, 123, 126. (16) Asahi, R.; Morikawa, T.; Ohwaki, T.; Aoki, K.; Taga, Y. Science 2001, 293, 269. (17) Livraghi, S.; Paganini, M. C.; Giamello, E.; Selloni, A.; Valentin, C. D.; Pacchioni, G. J. Am. Chem. Soc. 2006, 128, 15666. (18) Cole, B.; Marsen, B.; Miller, E. Mater. Res. Soc. Symp. Proc. 2007, 974E, CC09-04. (19) Paluselli, D.; Marsen, B.; Miller, E. L.; Rocheleau, R. E. Electrochem. Solid State Lett. 2005, 8, G301. (20) Marsen, B.; Cole, B.; Miller, E. L. Sol. Energy Mater. Sol. Cells 2007, doi:10.1016/j.solmat.2007.08.008 (21) Ahn, K.; Yan, Y.; Al-Jassim, M. J. Vac. Sci. Technol. B 2007, 25, L23.
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