TiO2 Nanotubes: Interdependence of Substrate ... - ACS Publications

Dec 6, 2011 - Highly ordered TiO2 nanotube arrays have received great interest in the field of nanoscience and nanotechnology. In view of applications...
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TiO2 Nanotubes: Interdependence of Substrate Grain Orientation and Growth Characteristics Silvia Leonardi,† Andrea Li Bassi,§,|| Valeria Russo,§ Fabio Di Fonzo,|| Odysseas Paschos,† Thomas M. Murray,^ Harry Efstathiadis,^ and Julia Kunze*,†,‡ †

Department of Physics E19, Technische Universit€at M€unchen, James-Franck-Strasse 1, 85748, Garching, Germany Institute for Advanced Study, Lichtenbergstrasse 2 a, 85748 Garching, Germany § Department of Energy and NEMAS - Center for NanoEngineered MAterials and Surfaces, Politecnico di Milano, via Ponzio 34/3, 20133 Milan, Italy Center for Nano Science and Technology@PoliMi, Italian Institute of Technology, via Pascoli 70/3, 20133 Milan, Italy ^ College of Nanoscale Science and Engineering, University at Albany SUNY, 251 Fuller Road, Albany, New York 12203, United States

)

‡

ABSTRACT: Highly ordered TiO 2 nanotube arrays have received great interest in the field of nanoscience and nanotechnology. In view of applications in devices, it is of paramount importance to control the growth over large areas. The results described in this paper show how the growth characteristics of anodic TiO2 oxide films can be tailored by controlling the orientation of the underlying substrate. The correlation between the growth characteristics of anodic, self-organized TiO2 nanotubes and the crystallographic orientation of the underlying titanium substrate grains has been studied using electron backscatter diffraction (EBSD), scanning electron microscopy (SEM), and micro-Raman spectrometry. The preferred formation site for self-organized, amorphous TiO2 nanotubes is Ti(111), whereas on Ti(001) grains no nanotube growth is observed. Instead, a compact oxide film is formed, exhibiting a mixed anatase and rutile nanocrystalline character with a large degree of structural disorder. As an intermediate case, on Ti(010) and on Ti(110) nanotubes form but are capped by a strongly chemically etched overlayer. It has been observed that TiO2 nanotubes exclusively form on grains with an orientation that allows for the formation of a thick valve metal oxide. This insight will enable controlled and effective growth over large areas of nanotubular films on titanium and presumably on other transition metals that did not show growth of nanotubular oxides before.

1. INTRODUCTION Compact anodic oxide films on titanium have been subject of intense investigation. Anodization of titanium at moderate voltages in aqueous electrolytes leads to the formation of amorphous films, while at higher voltages and specific environments the formation of all three crystalline structures of TiO2, rutile, anatase, and brookite has been reported.1 6 Several authors describe a solution composition effect and an influence of the electrolyte pH on the resulting crystallographic structure of anodic TiO2 layers.7 The oxide structure has also been described to be determined by the substrate structure and preparation procedure.8,9 Differences in the crystallographic structure are often caused by relatively small differences in the parameters used during anodization.4,5 In 1999, Zwilling et al.10 reported for the first time the electrochemical formation of self-organized nanotubular structures on titanium in fluoride-containing electrolytes. Subsequent tailoring of the growth conditions led to an increased aspect ratio of the nanotubes and an optimization of the efficiency of nanotube formation.11 13 Those self-organized structures stimulated intense research activities due to their high potential in applications r 2011 American Chemical Society

relating to catalyst systems, sensors, wettability tuning of surfaces, biomedicine, electrochromic devices, and solar energy conversion. An overview is given in ref 14. A field-assisted dissolution type process has been considered for formation of pores in anodic titania,15 similar to that of anodic alumina.16,17 However, a combination of flow and dissolution processes may occur in nanotubular anodic titania, which has been considered in tracer studies.18 The investigation of the growth mechanism of nanotubular TiO2 layers on sputter-deposited Ti thin films could support the oxide flow model.19 A photoresist masking method of thin Ti films was used to prepare SEM cross sections to directly obtain information on oxide morphology, layer thickness, and metal substrate loss. Therefore, not only features of the initial growth stages but also oxide expansion factors could be accurately determined which confirmed substantial contribution to steady-state tube growth by a plastic oxide flow mechanism.20 It was also reported that either a disordered Received: September 29, 2011 Revised: November 21, 2011 Published: December 06, 2011 384

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Figure 1. SEM micrographs of TiO2 nanotubes grown in 1 M (NH4)H2PO3 + 0.5 wt % NH4F at 20 V for 3 h on nonelectropolished Ti substrates (a c) and on electropolished Ti substrates (d f).

nanoporous layer or a self-organized nanotubular layer can be formed on titanium depending on the experimental parameters.21,22 The nanoporous layer was shown to have characteristics of a rutile TiO2 structure, whereas the nanotubular films were mainly amorphous. Recently, Shin et al.23 demonstrated extremely uniform and self-organized regular arrays of anodic TiO2 nanotubes by a combination of electropolishing and a two-step anodization. Even though a large number of studies have been devoted to TiO2 nanotubes, so far the interdependence of nanotube growth and substrate orientation has not yet been considered. To the best of our knowledge, only one paper describes retarded nanotube growth on the Ti(001) surface.24 The results, even though interesting, do not provide a detailed understanding of the interdependence of nanotube growth and grain structure, resulting in deriving only qualitative conclusions. It is well-known that the electrochemical reactivity depends on the thermodynamic properties of the substrate and the oxide film as well as on the ionic conductivity and electronic properties of the latter. In the case of TiO2 on Ti, it has been shown that the ionic and electronic properties of the oxide depend on the crystallographic orientation of the substrate grains and that the density of surface atoms determines the properties of the oxide films.25 28 A combination of microelectrochemistry and electron backscatter diffraction (EBSD) has been employed by K€onig et al.28 to study the correlation

between crystallographic orientation and electrochemical behavior of anodic oxide on polycrystalline titanium. In their work the oxide layer thickness and oxide crystallinity were analyzed, and they found that on Ti(001) thin crystalline oxide films were formed, whereas on Ti(hk0) thicker less crystalline films grew. A microelectrochemistry and EBSD study revealed that on Ti(001) anodic oxygen evolution takes place at potentials >4 V, whereas on Ti(hk0) and misoriented grains oxide forms exclusively.29 This showed the influence of the substrate orientation on the rate of electron-transfer reactions and confirmed the results reported earlier.25 28

2. EXPERIMENTAL METHODS 2.1. Reagents, Solutions, and Electrode Materials. Polycrystalline titanium metal sheets (99.6% purity, Advent Ltd., England) with a thickness of 1 mm and an edge length of 1 cm were used as substrates for anodic nanotubes formation. A solution of 1 M (NH4)H2PO3 + 0.5 wt % NH4F (Merck, 9.998% purity) was used as electrolyte during the anodization treatment. All chemicals were purchased from chemical suppliers and used without further purification. Prior to the anodic treatment, the polycrystalline titanium substrates were mechanically polished with wet SiC grinding paper, grade 4000. They were degreased by sonicating in acetone, isopropanol, and methanol, thoroughly rinsed with Milli-Q 385

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Figure 2. (a) SEM micrograph of the area chosen for EBSD investigations. (b) Inverse pole figure (IPF) of the area between the markers shown in (a). (c) Color legend for IPF with drawings explaining the respective crystallographic orientations.

water, and finally dried in a nitrogen stream. In order to remove the amorphous Beilby layer which forms upon mechanical polishing, an electropolishing procedure has been carried out in a mixture of methanol (purity g99.9%, Merck), perchloric acid (HClO4, 40%) (suprapure, Merck), and butoxyethanol (purity g99%, Alfa Aesar).30 The electropolishing solution was cooled with liquid nitrogen and held at a temperature between 20 and 40 C. The titanium samples were contacted from the back and wrapped with a Teflon mask in order to expose a well-defined circle with 9 mm diameter to the solution only. A voltage of 60 V was applied for 5 min between the sample and the counter electrode (CE), which was a gold plate with an edge length of 1 cm. Afterward, the samples were thoroughly rinsed with Milli-Q UV (DI) water and blown dry with Ar. After two electropolishing cycles, the Teflon mask was removed and the samples were cleaned in ethanol in an ultrasonic bath for 10 min. 2.2. Electrochemical Measurements and Instrumentation. All electrochemical anodization treatments were carried out at room temperature in a Teflon electrochemical cell with a conventional two-electrode configuration with a platinum gauze as counter electrode. The Teflon cell was cleaned in Caro’s acid (3:1 mixture of H2SO4 and H2O2) and thoroughly rinsed with DI water prior to every experiment. A dc power supply (PS2403D, Fluke) was employed to apply the respective anodization voltages. A potential ramp from 0 to 20 V with a sweep rate of ∼1 V/s was applied followed by retaining the potential constant at 20 V for 3 h. After anodization the samples were thoroughly rinsed with DI water and dried in a nitrogen stream. 2.3. Surface Analysis. 2.3.1. Electron Backscattered Diffraction (EBSD). A field-emission scanning electron microscope (FEI Nova NanoSEM600 Dualbeam, HKL Channel 5) was used to

collect an EBSD grain orientation map (inverse pole figure, IPF). The standard measurement configuration (45 pretilted specimen mount, Ted Pella Item 16355, 70 angle between electron beam and surface) with tilt correction was used to record a microstructural map 50 μm  50 μm in area with a step size of 0.5 μm. The operating conditions for EBSD collection were an acceleration voltage of 30 kV, a working distance of 12 15 mm, and a beam current of 0.64 nA. Samples were fixed to the specimen mount with conductive carbon sticky tabs to allow easy removal for further analysis. The crystallographic orientation of the grains is wholly determined by a set of three Euler angles j1, ϕ, and j2. In electrochemistry, the orientation of the surface plane exposed to the electrolytic phase is the most important information and usually expressed by means of the Miller indices {hkl}. To transform the texture data (Euler space) into the (hkl) planes, a simple matrix multiplication can be used,28,31 which is included in the EBSD software on the SEM instrument (EBSD software with a Nordliss camera). 2.3.2. Scanning Electron Microscopy (SEM). SEM micrographs were acquired using a field emission SEM (Zeiss Supra 40) using the detector for secondary electrons. 2.3.3. Raman Spectroscopy. All spectra have been recorded with a Renishaw InVia micro-Raman spectrometer equipped with an Ar+ laser (excitation wavelength 514.5 nm).

3. RESULTS The present work explores the role of the orientation of the underlying Ti metal in determining the resultant film morphology (nanotubular, compact, or compact with unordered or ordered porosity) upon anodic polarization in fluoride-containing electrolytes. Material science, surface science, and electrochemistry 386

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Figure 3. SEM micrographs of specific grains after anodization (in 1 M (NH4)H2PO3 + 0.5 wt % NH4F at 20 V for 3 h) of the area marked in Figure 2a: (a) oxide on grain number 4 (red) corresponding to Ti (001), (b) oxide on grain number 16 (green) corresponding to Ti(010), and (c) oxide on grain number 13 (blue) corresponding to Ti(110).

are combined to gain deeper insight into nanotubular growth and its dependence on the titanium grain orientation. The results presented here aim at providing insights for fabricating self-organized nanotubular structures and to open pathways for further investigation on other transition metals that so far have not shown selfnanostructuring behavior. Figure 1 shows SEM micrographs of TiO2 nanotube layers grown on nonelectropolished Ti substrates (panels a c) in comparison to TiO2 nanotube layers grown on electropolished Ti substrates (panels d f). Remarkable differences can be noted concerning the tube formation characteristics when comparing the differently pretreated substrates. At relatively low fluoride concentrations, on the nonelectropolished Ti after a 3 h anodization only a partially ordered nanoporous structure can be achieved. TiO2 nanotubes start forming randomly and without order in some spots on the surface. A thick compact oxide capping layer (oxide covering the mouths of the nanotubes) is predominant all over the surface. Such morphology is found on the whole surface. No dependency on the crystallographic properties of the substrate grains is visible. The oxide layer grown under the same conditions on the electropolished Ti substrate shows an interdependence of oxide film morphology and grain structure of the polycrystalline Ti substrate (Figure 1d f). In order to understand this phenomenon and to elucidate the reason for the observed interdependence, EBSD experiments were performed prior to anodic nanotube growth.

After electropolishing, the surface area that was to be analyzed with EBSD was marked using focused ion beam (FIB) (see crosses in Figure 2). The texture of a selected area of the Ti substrate (Figure 2a, area inside of the FIB marks) was characterized by EBSD immediately after the electropolishing procedure and prior to the anodic treatment. Figure 2b shows the inverse pole figure (IPF) recorded with EBSD which can be read like a microstructural map of the Ti substrate. The IPF shows the orientation of the titanium metal if we consider a thickness of the native oxide of ∼1.3 5.4 nm.32 The color coding provides a measure of the grain orientation based on the normal to the surface, with each solid color corresponding to an individual grain. A color legend (orientation distribution map) for the inverse pole plot is shown in Figure 2c. The morphology of all grains that were investigated using SEM was marked with a number in the corresponding inverse pole plot map (see Figure 2b). Figure 3 depicts SEM micrographs of the grains corresponding to the three main orientations at the corners of the orientation distribution map after anodization, with red corresponding to a (001) orientation, green corresponding to a (010) orientation, and blue corresponding to a (110) orientation. The oxide films formed on the (001) oriented grain are compact and show almost no traces of chemical etching. The oxide films formed on the (010) and the (110) oriented grains are more open and are characterized by deep cavities that are 387

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Figure 4. SEM micrographs of specific grains after anodization (in 1 M (NH4)H2PO3 + 0.5 wt % NH4F at 20 V for 3 h) of the area marked in Figure 2a: (a) oxide on grain number 17 (yellow) corresponding to Ti (102), (b) oxide on grain number 12 (magenta) corresponding to Ti(112) tilted toward Ti(001), and (c) oxide on grain number 10 (light purple) corresponding to Ti(112) tilted toward Ti(110).

Figure 5. SEM micrographs of specific grains after anodization (in 1 M (NH4)H2PO3 + 0.5 wt % NH4F at 20 V for 3 h) of the area marked in Figure 2a: (a) oxide on grain number 2 (purple) corresponding to Ti (111) tilted toward Ti(110) and (b) oxide on grain number 15 (dark purple) corresponding to ∼Ti(111). 388

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observed which, after anodization, is characterized by TiO2 nanotubes that are partially covered by compact oxide features (Figure 5a). The grain corresponding to the dark purple area of the map shows uncovered TiO2 nanotubes (Figure 5b). The tubes have a diameter of 100 nm and are ∼1.75 μm long (see Figure 1f). The grains with the orientation Ti(111) (dark purple area of the orientation distribution map) are the only grains on which open self-organized titania nanotubes could be observed after anodic oxidation in fluoride containing electrolyte with low fluoride concentration. It is noteworthy that on grains which show more than one color in the IPF due to the presence of very small crystallographic domains the oxide layer shows a mixed morphology corresponding to the characteristics observed on the different components. A clear example is reported in Figure 6 where the morphology of the oxide layer grown on grain 1 results from a mixture of the oxide layer morphologies of grains corresponding to green (etched compact overlayer) and dark purple areas (open self-organized TiO2 nanotubes) of the map. Micro-Raman spectroscopy was used to determine differences in the structure and phase composition of the different TiO2 layers obtained on different Ti grains after anodization in fluoride containing electrolytes. Figure 7 shows micro-Raman spectra of the oxide films formed on the differently oriented grains. The grains shown in Figure 7a are characterized by completely closed films (red area, grain number 4) or closed films with cavities from chemical etching (light magenta and magenta areas, grains 11 and 10). The micro-Raman spectra corresponding to these grains are characterized by the presence of the anatase Eg mode at ∼145 150 cm 1 (for all grains) and by broad bands at about 430 440 and 610 cm 1, roughly corresponding to the positions of the rutile Eg mode (447 cm 1) and A1g mode (612 cm 1), i.e., in grains 4, 10, and 12, or by the copresence of features at about 400, 520, and 620 630 cm 1, roughly corresponding to the positions of the anatase B1g modes (398 and 519 cm 1) and Eg mode (639 cm 1), i.e., in grains 5 and 11, or as weak shoulders on other grains (see e.g. refs 33 and 34). These spectra indicate the presence of a disordered titanium oxide structure, probably composed by nanocrystalline domains with a large fraction of disordered grain boundaries or amorphous oxide matrix (especially for grain 12). The presence of small nanocrystalline anatase or rutile domains is suggested by the shift of the anatase Eg peak with respect to its position in bulk crystals, which is 144 cm 1, and by its broadening, up to 25 30 cm 1,34 and also by the downshift and strong broadening (up to 60 cm 1) of the rutile Eg mode, see e.g. grains 4 and 10.35 The grains depicted in Figure 7b have a more open morphology which can be TiO2 nanotubes (grains 15 and 2) or an open capping layer formed by chemical etching (grain 16). The Raman spectra of the open nanotubes and of the capping layer exhibit three very broad features around 150 200, 450, and 610 cm 1, suggesting a very disordered, amorphous structure (a very weak, broad peak at ∼155 cm 1 is observed only for grains 2 and 15). Such broad bands have been frequently observed in amorphous titania, assigned to the Ti O bending (200 and 450 cm 1) and Ti O stretching (610 cm 1) vibrations.36 38 The same bands have been found in a recent Raman study on TiO2 nanotubes.39

Figure 6. SEM micrographs of specific grains after anodization (in 1 M (NH4)H2PO3 + 0.5 wt % NH4F at 20 V for 3 h) of the area marked in Figure 2a: (a) oxide on mixed grain number 1 (∼Ti(010) with Ti(111)); (b) zoom into frontier between grain 2 and grain 1, the latter partially showing open TiO2 nanotubes (NT) that are covered by a chemically etched capping layer.

due to chemical etching of the compact TiO2 film. It is visible that the openings are not due to a self-organized growth of TiO2 nanotubes. The etched oxide layer on the (110) oriented grain (blue) appears more closed compared to the one formed on the (010) oriented grain (green) due to a more displaced arrangement of the etching holes. When moving along the upper edge of the orientation distribution map toward Ti(010), the yellow color which represents an orientation of (102) is reached. The anodic oxide formed on top of the corresponding grain is compact and shows minor traces of chemical attack (hole in the center of Figure 4a2). Moving along the lower edge of the orientation distribution map toward Ti(110) and stopping at the magenta area corresponds to grains with an orientation of (112) tilted toward (001). The oxide film observed on top of these grains is compact. Chemical etching of the film is more pronounced than on (001) and (102) oriented grains, as can be seen in Figure 4b2 by the presence of numerous etching holes. Moving further toward (110) (see grain corresponding to the light purple area of the color map in Figure 4c) reveals grains on which a compact oxide film is found that is severely chemically attacked by the fluorides from the electrolyte solution. The etching holes are deep and start merging to larger craters. Moving further toward (110) finally reveals a crystallographic orientation of the substrate at which an oxide consisting of selforganized TiO2 nanotubes grows upon anodization (Figure 5). In this purple region of the orientation distribution map, a grain is

4. DISCUSSION As discussed in the Introduction, it has been demonstrated that TiO2 nanotube growth starts with the formation of a 389

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Figure 7. Raman spectra and SEM micrographs of specific grains after anodization (in 1 M (NH4)H2PO3 + 0.5 wt % NH4F at 20 V for 3 h) of the area marked in Figure 2a: (a) spectra and micrographs of grains 4, 5, 10, 11, and 12; (b) spectra and micrographs of grains 2, 15, and 16.

compact oxide film. This film is first chemically attacked by the fluorides in solution, and then small tubes start forming. First, they are rather unordered until a stable growth sets in, leading to a self-organized process producing vertically ordered high-aspect ratio TiO2 nanotubes.20 Thus, the nanotube growth must depend on the structure of the compact oxide that first forms during the initial nanotube growth stages. As a consequence, nanotube growth must also depend on the orientation of the metal support, since the structure and quality of anodic compact TiO2 strongly depend on the orientation of the underlying Ti metal as it has been reported by K€onig et al.28 They found that the best crystallized oxide films are formed on Ti(001), whereas on Ti(010) and Ti(110) thicker and less ordered oxide films were observed— both findings can be confirmed by the above presented results, as explained in the following paragraph. On (001) oriented Ti substrate grains (red), compact anodic TiO2 films without any holes resulting from chemical etching of the fluoride-containing solution are observed, whereas the oxide films on Ti(010) (green) and Ti(110) (blue) are strongly attacked chemically and show tremendous etching cavities. This finding is well in agreement with observations reported in the literature.8,9,25 29 The more crystalline oxide film on Ti(001) withstands the chemical etching of the fluorides in solution much better than the less crystalline thicker oxide films on Ti(010) and Ti(110) (Figure 3). The Ti(001) plane differs from the Ti(hk0) planes significantly by a doubled atomic density, resulting in the growth of an oxide showing a higher donor concentration26,28 which influences the nature of

electron and ion exchange between the Ti surface and the electrolyte solution and which affects the dissolution rate at the interface, the oxide layer formation, and the properties of the layer.28 The oxide formed on Ti(102) (yellow) is still compact and shows very few holes that result from chemical etching only. The number of etching cavities increases when the sample is tilted from (001) toward (110); on Ti(112) (magenta) local etching holes are already clearly visible all over the compact oxide film. Tilting more toward Ti(110) (light purple) causes more etching, the local cavities develop into larger craters, and the morphology of the compact oxide is visibly altered (Figure 4). This observation shows that the quality of the oxide decreases when the orientation of the underlying Ti substrate is changed from (001) over (112) to (110). When tilting reaches the purple part in the orientation distribution map, open self-organized TiO2 nanotubes can be observed (Figure 5). Self-organized TiO2 nanotubes can only be observed on Ti(112) grains, tilted toward Ti(110). The dark purple spot on the orientation distribution map which corresponds to an orientation of the substrate of Ti(111) is the only orientation on which nanotubes without any compact oxide cover can be observed. If the substrate orientation is too far tilted toward Ti(110) (Figure 5a), the tubes are not appearing completely open anymore. The Raman results reported in Figure 7 help to support findings on the growth mechanism already reported in the literature.22 Therein, polished and preanodized titanium sheets were anodized in fluoride-containing electrolyte. Regions with compact 390

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5. CONCLUSIONS It has been demonstrated that TiO2 nanotube formation strongly depends on the orientation of the underlying Ti substrate. No nanotubes are found on Ti(001); the oxide films formed here are compact and show nanocrystalline/disordered character with mixed anatase and rutile phases. The preferred formation site for self-organized, amorphous TiO2 nanotubes is Ti(111). Nanotubes most likely also form on Ti(010) and on Ti(110), but in these cases they are capped by a strongly chemically etched overlayer. It can be concluded from the findings reported in this paper that TiO2 nanotubes only form on grains with an orientation that allows for the formation of a thick valve metal oxide. Such oxides enable the penetration of fluorides through the film to the metal/ oxide interface where nanotube growth is initiated. The results reported in this paper are well in agreement with earlier studies and clearly support some of the aspects from the literature that so far could only be assumed. Two main conclusions can be drawn from the present work. From a fundamental point of view, the above-described novel understanding can enable controlled and effective growth of nanotubular films on titanium and on other transition metals that so far did not show evidence of nanotubular oxide growth. If grain-dependent growth of anodic oxide is observed, it can be assumed that on high index planes of a given transition metal the growth of nanotubular oxides is possible. From an engineering point of view, single crystalline substrates could be employed for large area manufacturing of devices based on nanotube arrays. Hence, these findings have the potential to open up a complete new path in electrochemical nanotechnology.

Figure 8. Schematic drawing of different crystallographic orientations of Ti grains42 on top of which oxide films with varying morphologies have been examined in this paper.

porous morphology and regions with nanotubular morphology (self-organized TiO2 nanotubes) were observed. The development of nanotubes was considered to be related to the penetration of fluoride ions to the titanium/film interface, which hinders the formation of crystalline film and hence stabilizes a mainly amorphous film structure that promotes growth of ordered nanotubes.22 Recently, the existence of a fluoride-rich layer forming at the metal/oxide interface upon anodization in fluoride containing electrolyte has been confirmed by high-resolution Auger electron spectroscopy.40 The fact that the nanotube layers show no Raman signals that correspond to any crystalline phase of TiO2 and that they thus are mainly amorphous is in line with previous findings.39,41 It is noteworthy that both the open nanotubular films and the grains covered by an open capping layer (e.g., grain 16) show very similar micro-Raman signals. This finding indicates that self-organized TiO2 nanotubes are present under the porous capping layers on Ti(010) and on Ti(110). It can also be assumed that no nanotubes have formed under the compact oxide on Ti(001) and under the partially etched compact oxides on Ti(102) and on Ti(112). In the near future, these indications will be verified by investigating the cross-section morphologies through respective grains using focused ion beam. The above-discussed results can be further interpreted with the help of the structural details of all the involved crystallographic planes (Figure 8), since completely open nanotubes exclusively form on one specific orientation of the underlying titanium substrate which is Ti(111). This orientation is located between Ti(112) and Ti(110) on the orientation distribution map (deep purple). Ti(112) is among the most open planes of the hexagonal close-packed Ti. Ti(111) is not close-packed and characterized by voids in the uppermost surface plane.42 The findings suggest that the compact oxide formed on Ti(111) allows for an easy and homogeneous penetration of fluoride ions through the film to the metal/film interface. This oxide must thus be rich in defects that are evenly distributed all over its surface. Another important point is the fact that the grains with oxides that show a more crystalline character, like those grown on Ti(001), almost show no valve metal like behavior, as has been described earlier.28 Therefore, self-organized valve metal oxide nanotubes exclusively form on grains that allow for the formation of a thick compact, less crystalline valve metal oxide prior to nanotube initiation and growth.

’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected]. Phone: +49-89-289-12526. Fax: +49-89-289-12536.

’ ACKNOWLEDGMENT The authors thank the DFG and the Technische Universit€at M€unchen Institute for Advanced Study, funded by the German Excellence Initiative, for financial support. A.L.B., V.R., and F.D.F. acknowledge funding from the Fondazione Cariplo (project no. 2009/2527 “Nanostructured MATerials for innovative Hybrid Solar cells MATHYS”). J.K. thanks Ulrich Stimming for acting as her host and for helpful scientific discussions. ’ REFERENCES (1) Kelly, J. J. Electrochim. Acta 1979, 24, 1273–1282. (2) Aladjem, A. J. Mater. Sci. 1973, 8, 688–704. (3) Aladjem, A.; Brandon, D. G.; Yahalom, J.; Zahavi, J. Electrochim. Acta 1970, 15, 663–671. (4) Ohtsuka, T.; Guo, J.; Sato, N. J. Electrochem. Soc. 1986, 133, 2473–2476. (5) Ohtsuka, T.; Otsuki, T. Corros. Sci. 2003, 45, 1793–1801. (6) Yamaguchi, S. J. Electrochem. Soc. 1961, 108, 302–302. (7) Yamabi, S.; Imai, H. Chem. Mater. 2002, 14, 609–614. (8) Kozlowski, M. R.; Tyler, P. S.; Smyrl, W. H.; Atanasoki, R. T. Surf. Sci. 1988, 194, 505–530. (9) Froelicher, M.; Hugot-LeGoff, A.; Jovancicevic, V. Thin Solid Films 1981, 82, 81–88. (10) Zwilling, V.; Darque-Ceretti, E.; Boutry-Forveille, A.; David, D.; Perrin, M. Y.; Aucouturier, M. Surf. Interface Anal. 1999, 27, 629–637. 391

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(11) Gong, D.; Grimes, C. A.; Varghese, O. K.; Hu, W. C.; Singh, R. S.; Chen, Z.; Dickey, E. C. J. Mater. Res. 2001, 16, 3331. (12) Macak, J. M.; Tsuchiya, H.; Schmuki, P. Angew. Chem. 2005, 44, 7463–7465. (13) Albu, S.; Ghicov, A.; Macak, J. M.; Schmuki, P. Phys. Status Solidi RRL 2007, 1, 181–183. (14) Macak, J. M.; Tsuchiya, H.; Ghicov, A.; Yasuda, K.; Hahn, R.; Bauer, S.; Schmuki, P. Curr. Opin. Solid State Mater. Sci. 2007, 11, 3–18. (15) Macak, J. M.; Tsuchiya, H.; Schmuki, P. Angew. Chem. 2005, 44, 2100. (16) Raja, K. S.; Misra, M.; Paramguru, K. Electrochim. Acta 2005, 51, 154–165. (17) Hoar, T. P.; Mott, N. F. J. Phys. Chem. Solids 1959, 9, 97–99. (18) LeClere, D. J.; Valota, A.; Skeldon, P.; Thompson, G. E.; Berger, S.; Kunze, J.; Schmuki, P.; Habazaki, H.; Nagata, S. J. Electrochem. Soc. 2008, 155, C487–C494. (19) Berger, S.; Macak, J. M.; Kunze, J.; Schmuki, P. Electrochem. Solid-State Lett. 2008, 11, C37–C40. (20) Berger, S.; Kunze, J.; Schmuki, P.; LeClere, D.; Valota, A.; Skeldon, P.; Thompson, G. Electrochim. Acta 2009, 54, 5942–5948. (21) Kunze, J.; Seyeux, A.; Schmuki, P. Electrochem. Solid-State Lett. 2008, 11, K11–K13. (22) Seyeux, A.; Berger, S.; LeClere, D.; Valota, A.; Skeldon, P.; Thompson, G. E.; Kunze, J.; Schmuki, P. J. Electrochem. Soc. 2009, 156, K17–K22. (23) Shin, Y.; Lee, S. Nano Lett. 2008, 8, 3171–3173. (24) Crawford, G. A.; Chawla, N. Scr. Mater. 2009, 60, 874–877. (25) Schultze, J. W.; Kudelka, S. Interface 1997, 6, 28. (26) Kudelka, S.; Schultze, J. W. Electrochim. Acta 1997, 42, 2817– 2825. (27) Kudelka, S.; Michaelis, A.; Schultze, J. W. Electrochim. Acta 1996, 41, 863–870. (28) K€onig, U.; Davepon, B. Electrochim. Acta 2001, 47, 149–160. (29) Davepon, B.; Schultze, J. W.; K€onig, U.; Rosenkranz, C. Surf. Coat. Technol. 2003, 169, 85–90. (30) Arsov, L. D. Electrochim. Acta 1985, 30, 1645–1657. (31) Kallend, J. S.; Kocks, U. F.; Rollett, A. D.; Wenk, H.-R. Mater. Sci. Eng., A 1991, 132, 1–11. (32) Schultze, J. W.; Lohrengel, M. M. Electrochim. Acta 2000, 45, 2499–2513. (33) Ohsaka, T.; Izumi, F.; Fujiki, Y. Raman Spectrosc. 1978, 7, 321–324. (34) Li Bassi, A.; Cattaneo, D.; Russo, V.; Bottani, C. E.; Barborini, E.; Mazza, T.; Piseri, P.; Milani, P.; Ernst, F. O.; Wegner, K.; Pratsinis, S. E. J. Appl. Phys. 2005, 98, 074305–074308. (35) Mazza, T.; Barborini, E.; Piseri, P.; Milani, P.; Cattaneo, D.; Li Bassi, A.; Bottani, C. E.; Ducati, C. Phys. Rev. B 2007, 75, 045416. (36) Fernandez-Garcia, M.; Wang, X.; Belver, C.; Hanson, J. C.; Rodriguez, J. A. J. Phys. Chem. C 2007, 111, 674–682. (37) Zhang, J.; Li, M.; Feng, Z.; Chen, J.; Li, C. J. Phys. Chem. B 2006, 110, 927–935. (38) Kern, P.; J€aggi, C.; Utke, I.; Friedli, V.; Michler, J. Appl. Phys. Lett. 2006, 89, 02102. (39) Likodimos, V.; Stergiopoulos, T.; Falaras, P.; Kunze, J.; Schmuki, P. J. Phys. Chem. C 2008, 112, 20574. (40) Berger, S.; Albu, S. P.; Schmidt-Stein, F.; Hildebrand, H.; Schmuki, P.; Hammond, J. S.; Paul, D. F.; Raichlmaier, S. Surf. Sci. 2011, 605, L57–L60. (41) Albu, S. P.; Ghicov, A.; Aldabergenova, S.; Drechsel, P.; LeClere, D. J.; Thompson, G. E.; Macak, J. M.; Schmuki, P. Adv. Mater. 2008, 20, 4135–4139. (42) Figure drawn with: http://surfexp.fhi-berlin.mpg.de.

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dx.doi.org/10.1021/jp209418n |J. Phys. Chem. C 2012, 116, 384–392