Local Tuning of Conductivity in Amorphous ... - ACS Publications

13972

J. Phys. Chem. C 2007, 111, 13972-13980

Local Tuning of Conductivity in Amorphous Titanium Oxide Films by Selective Electron Beam Irradiation P. Kern,*,† R. Widmer,‡ P. Gasser,† and J. Michler† Mechanics of Materials and Nanostructures, and Nanotech@Surfaces, Empa, Feuerwerkerstrasse 39, CH-3602 Thun, Switzerland ReceiVed: February 12, 2007; In Final Form: July 20, 2007

Amorphous TiO2 (a-TiO2) films electrodeposited on steel substrates have recently been found highly e-beam sensitive, allowing for reduction and crystallization with high lateral resolution. We here report on the strong correlation between the e-beam-induced changes in oxidation state and film structure, as observed by transmission electron microscopy and Micro-Raman spectroscopy, and the local film conductivity (σ), measured by conductive atomic force microscopy (c-AFM). c-AFM at constant applied potentials reveals a low σ of 1.9 × 10-4 Ω-1 cm-1 for a-TiO2 and a further decrease of σ by up to a factor 5.5 × 10-3 at locally reduced spots, rather than an enhancement in conductivity consistently reported in literature upon reduction of crystalline TiO2. However, σ significantly rises with the first appearance of anatase phase in micro-Raman spectra, and makes a second large jump coinciding with the formation of TiO, with σ being 2 × 10+4 times higher than in the surrounding a-TiO2. We further observe a rectifying effect due to the presence of micrometer-sized anatase/steel Schottky barrier diodes at locally crystallized spots. The absolute probe current rather than the current density is found responsible for e-beam-induced crystallization, underlining the thermal nature of crystallization at electron energies disqualifying atomic displacement. Structural changes are strictly limited to the irradiated area with anatase nanocrystals being extended nearly throughout the 60 nm thick film, not showing preferential orientation. The results demonstrate the possibility of e-beam-induced confined tuning of σ in the amorphous oxide film in both directions. This observation is novel and will allow for highresolution patterning of electrical properties in amorphous semiconducting oxides.

1. Introduction Because of interesting physical, chemical, optical, and electrical properties, and for this reason manifold important applications in photocatalysis,1,2 solar cells,3,4 gas and humidity sensors,5,6 antireflexion coatings,7 optical filters8 or waveguides,9 as well as for biomaterials,10,11 titanium dioxide (TiO2) thin films have been among the most widely studied semiconductor materials within the past decade. In crystalline form, TiO2 exists in the tetragonal phases anatase and rutile, as well as in the less frequently encountered orthorhombic phase brookite. Crystalline TiO2 is an n-type semiconductor with band gap of 3.1-3.4 eV.12,13 Very little has been reported on properties of amorphous TiO2 (a-TiO2) films, mostly due to difficulties in characterizing the poorly defined amorphous state and the fact that its electrical properties are generally considered less attractive than those of crystalline TiO2. A band gap of 4.4 eV13 and a dielectric constant for TiO2/ Si of 3.014 (anatase, 3.9; rutile, 5.1) have been reported for a-TiO2. By controlling the amount of hydrogen in the plasma gas, Ohmori and co-workers15 plasma sprayed titanium oxide coatings with different oxygen content on steel and found a continuous increase of electrical conductivity with ongoing state of reduction. The monoxide TiO, immediately oxidizing to TiO2 in atmosphere,16 is reported to have a resistivity in the range of 170-900 µΩ cm,17,18 similar to the well-studied conductive * Corresponding author. Tel.: + 41 33 228 36 26. Fax: +41 33 228 44 90. E-mail: [email protected]. † Mechanics of Materials and Nanostructures. ‡ Nanotech@Surfaces.

RuO2 or highly doped semiconductors. This low resistivity and the excellent barrier properties against interdiffusion of Si and Al make this material a perspective thin film material for low resistance contact metallization in microelectronics.17,19 Moreover, TiO has raised particular interest for improving fatigue lifetime22 when associated with lead zirconate titanate (PZT) films, and it is used as seed substrate for the epitaxy of PZT.21 Changing the electrical properties of a semiconducting film on a local scale has potential interesting micro-technical and electronic applications. Such local modifications of the electrical properties of crystalline titanium oxide have been achieved by ion implantation22-25 and tripled YAG laser irradiation.26-28 Implantation of ions generally resulted in a gradual transformation into the amorphous phase and preferential sputtering of O leading to formation of Ti2O3, while the resistivity rapidly decreased to a saturation value, in the case of Nb+ implantation up to 107 times less than in the initial oxide.25 Le Mercier et al. reported the formation of Ti2O3 and TiO2-x and a decrease of the electrical resistance by a factor 103 26 and even 108 26 after laser irradiation of single-crystal rutile. Electron beam (e-beam) exposure effects on titanium oxide have previously been investigated exclusively on crystalline TiO2 and mostly in the transmission electron microscope (TEM) at accelerating voltages above 100 keV, conditions that allow for atomic displacement and e-beam sputtering. McCartney and co-workers29,30 observed the formation of TiO at current densities between 5 and 50 A/cm2 and reduction beyond the monoxide after irradiation of rutile single-crystal samples at extremely high current densities of 103-104 A/cm2 at 100-

10.1021/jp0711803 CCC: $37.00 © 2007 American Chemical Society Published on Web 08/29/2007

Local Tuning of Conductivity in a-TiO2 Films 400 keV. Su and co-workers31 did not observe a noticeable structural or electronic change upon TEM irradiation of TiO2 anatase powder at 200 keV and 3 A/cm2. Irradiating at low energy and density (1 keV, µA/cm2) in an Auger microscope, Wang et al.32 reported an energy- and dose-dependent controlled number of Ti3+ defects in TiO2 (110). Potential changes in local electrical properties at irradiated zones were not investigated. Recently, working under scanning electron microscope (SEM) conditions at 20 keV, we have reported for the first time a high sensitivity of amorphous TiO2 films toward controlled e-beaminduced local reduction and crystallization.33 In contrast to previous studies on crystalline TiO2, electrolytic a-TiO2 films react significantly more sensitively toward electron stimulated oxygen desorption (ESD), allowing one to create well-defined topographical features within the film.34 Gradual reduction to Ti2O3, TiO, and even metallic Ti on sub-micrometer scale is observed as shown by Auger, WDX, and Raman analysis. Parallel to reduction phenomena, local crystallization occurs after already 1 s of irradiation using a probe current of 5 µA and a current density of 4.8 A/cm2. In-situ measurements of the global sample temperature33 and estimations of the additional temperature rise in the beam center indicate that crystallization is possible at temperatures as low as 150 °C, well below the atmospheric crystallization temperature (440 °C).35 In this study, we report on a strong impact of the previously observed e-beam reduction and crystallization phenomena on the local resistivity and semiconducting properties of irradiated a-TiO2 films, as studied by current imaging with conducting atomic force microscopy (c-AFM) at constant applied potentials. Structural changes in the irradiated zones are analyzed by microRaman spectroscopy and TEM imaging and are correlated with changes in local electrical properties. Well-controlled uniform irradiation conditions with beam currents between 500 nA and 5 µA are achieved using an electron probe microanalyzer (EPMA). 2. Experimental Section Electrochemical Deposition of TiO2. Galvanic electrolytic deposition on mirror polished stainless steel substrates (AISI 316L, φ ) 15 mm, ∆d ) 1 mm) was performed in 0.5 M TiCl4 + 0.5 M H2O2 in 3:1 vol % MeOH/H2O at -75 mA cm-2 up to 5 C deposition charge (50 s). Experimental details as well as the deposition mechanism are described elsewhere.36 Asdeposited uniform peroxotitanium hydrate films (TiO3(H2O)x) were annealed at 200 °C with a heating/cooling rate of 6 °C/ min and a holding time of 30 min. During heating, the films lose excess humidity and transform into amorphous TiOx (x being close to 2), with residual C and N contaminations of approximately 2 at. %, as measured by X-ray photoelectron spectroscopy.34 The final oxide has a thickness of approximately 60 nm. Thermal crystallization of the films does not occur below 440 °C.36 Electron Beam Irradiation. The TiO2/steel samples were fixed to the sample holder with conductive silver glue. Surface contaminations were removed by low-intensity plasma cleaning using air as feeding gas (Gala Instrumente, Plasma Prep 2). No effect on this cleaning step on film properties was observed. Controlled e-beam exposure was performed with an electron probe microanalyzer (EPMA) Jeol JXA-8800RL at 20 keV. The EPMA allows for independent adjustment of probe current and size through defocusing of the beam probe by electron-optical lenses, resulting in nearly uniform current density distribution in the beam. Irradiations were performed at probe currents IP of 500 nA (D ) 3.8 µm), 1 µA (D ) 5.0 µm), and 5 µA (D )

J. Phys. Chem. C, Vol. 111, No. 37, 2007 13973 11.5 µm), where D is the approximate probe diameter. After exposure, the TiO2 surfaces were again plasma cleaned to remove potential e-beam deposited hydrocarbon contaminations. Micro-Raman. Laser Raman spectra were collected in backscattering geometry with a Renishaw Ramascope 2000 using a 633 nm HeNe laser and a spot size of approximately 3 m. The spectra were collected in the center of irradiated spots, well-visible under the optical microscope due to the local blue appearance from Ti2O3 formation. Transmission Electron Microscopy. A TEM lamella of about 20 µm width and 100 nm thickness through the center of an irradiated spot (IP ) 5 µA, D ) 11.5 µm, t ) 5 s) was prepared by focused ion beam (FIB) milling. Prior to milling, the oxide was locally protected by e-beam deposited Pt in the lamella region, followed by ion-beam deposited Pt to form a thicker protective layer. TEM analysis was performed at 300 kV with a Philips CM30 microscope. For indexing of diffraction patterns of crystallized TiO2, the average camera length was calculated from the nanocrystalline pattern of Pt (111) with d ) 2.265 Å, Pt (200) with d ) 1.962 Å, and Pt (220) with d ) 1.387 Å acquired at the same magnification. Conductive Atomic Force Microscopy. c-AFM experiments were performed with a commercial AFM/STM Nanoscope III Multimode (Veeco) in contact mode with an applied external force of 30 nN and a scan speed of 100 µm/s. TiN (30 nm) coated contact ultrashsharp silicon cantilevers (NT-MDT) with a typical curvature radius of 35 nm were used. The potential between the steel substrate and the AFM probe was applied by an external source (ES 030-5, Delta Elektronika) through the signal access module of the Nanoscope AFM. All applied potentials are given with respect to the steel substrate. The current was acquired using the Nanoscope electronics (1 V ) 100 nA), the maximum readable current being 1000 nA. A small offset of +9.156 mV in the current sensing was measured before bringing sample and cantilever into contact and has been systematically corrected for. The WSxM 4.0 develop software (Nanotech Electronica S.L., Spain) was used for image evaluation. General c-AFM Considerations. Various scanning probe techniques such as scanning tunneling microscopy (STM), conductive atomic force microscopy (c-AFM),43-45 scanning electrochemical microscopy (SECM) or combined SECMAFM,46 as well as current sensing scanning near-field optical microscopy (SNOM)47 have recently been applied for microscopic analysis of electrical materials properties. c-AFM has shown to be particularly useful for probing resistances on sub100 nm dimensions also on highly resistive materials such as p-n junctions,43 Ti/TiO2/Pt anodes,45 or metal oxide surfaces.48 Recently, Szot and co-workers have reported the possibility of switching the electrical resistance of dislocations in singlecrystalline SrTiO3, measured by c-AFM under vacuum conditions.49 Contrary to stationary point-contact measurements of local I/V curves, scanning resistance microscopy or current imaging experiments performed with c-AFM not only require well conducting but also hard and abrasion-resistant probes. A compromise between contact resistance and wear resistance of the coated cantilever is therefore necessary. Thomson and coworkers have studied contact resistances of Au, Pt, and Ag coated as well as doped Si tips.50 They reported contact resistances in the range of 3-5 MΩ for pure or doped Si cantilevers and 30-450 Ω for Au, Pt, and Cr coated cantilevers under resulting tip forces between 7 and 90 µN. Unfortunately, Au48 or Pt45 coatings on Si or Si3N4 probes have insufficient mechanical stability for providing reproducible measurements

13974 J. Phys. Chem. C, Vol. 111, No. 37, 2007

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TABLE 1: Chosen EPMA e-beam Irradiation Conditions: IP ) Probe Current, D ) Probe Diameter, t ) Exposure Time, ∆T ) Calculated Temperature Rise in the Beam Center (see Eq 1)a no.

Ip (µA)

D (µm)

t (s)

J (A cm-2)

dose (C/cm-2)

∆T (°C)

mechanism

1 2 3 4 5

0.5 1 1 5 5

3.8 5 5 11.5 11.5

5 1-20 80 1-20 60

4.4 5.1 5.1 4.8 4.8

22 5-102 306 5-96 289

35 55 55 127 127

reduction (a-TiO2-x, a-Ti2O3) reduction (a-TiO2-x, a-Ti2O3) reduction (a-TiO2-x, a-Ti2O3) + crystallization (anatase) reduction (a-TiO2-x, a-Ti2O3) + crystallization (anatase) reduction (a,c-TiO) + crystallization (c-TiO)

a

The “dominating mechanism” summarizes the observed changes in oxidation state and structure within the a-TiO2 film (see text).

during scanning in contact mode;50,51 even so, successful extensive scanning with Au coated tips has been reported by Luo and co-workers.48 As alternatives, NbN coated Si3N4 tips have been proposed.51 Working under static conditions with an externally applied load of 6 nN, Wiederhold and co-workers reported a contact resistance of 272 Ω for Ti/TiN coated Si cantilevers.52 For the present c-AFM current imaging experiments, we have chosen TiN coated Si probes due to the intrinsic hardness and good electrical conductivity of TiN.52 The influence of applied external load and scan speed on contact resistance is still badly understood, and contact resistances are usually measured in static mode. Typical applied external forces for static c-AFM experiments are greater than 1 µN44,50 with the contact area estimated to be 30 nm at 100 µN force.44 Thomson et al. pointed out that tip-sample forces of 1 µN are required for making contacts with resistances lower than 10 MΩ, and that typical AFM contact forces in the nN range were insufficient for establishing a stable electrical contact.50 For scanning in contact mode, however, large contact forces are in detriment of the tip lifetime. Despite the claimed need by some authors for larger applied forces, the chosen force of approximately 30 nN in the present study yielded a stable electrical contact and allowed extended scanning without noticing a significant increase in contact resistance with time.

The irradiation conditions chosen in the present study, being a combination of the independent parameters beam current, probe diameter (spot-size), and exposure time, are listed in Table 1 together with the calculated T and the dominating chemical and structural changes within the oxide. For calculation of ∆T, κ of steel rather than TiO2 is taken, due to the small film thickness with respect to the micrometer-sized interaction volume. Conditions no. 2,3 and no. 4,5 are different in terms of IP, while the current density J is kept almost identical in all conditions, being equal to J in previous irradiations experiments with IP ) 5 µA.34 Figure 1a presents the micro-Raman spectra for the 1 µA series (no. 2,3) as a function of exposure time. After 20 s of irradiation, the same amorphous background seen before irradiation is observed, whereas a peak close to the characteristic Eg mode (≈146-148 cm-1) of TiO2 anatase structure appears after 60 s of exposure, indicating a small amount of anatase present within the oxide film. The small deviation in wavenumber (155 cm-1) is attributed to the limited resolution due to the small volume of anatase present. The 5 µA series in Figure 1b reveals strong anatase signature already after 1 s of exposure, with the anatase modes39 at 147 (Eg), 395 (B1g), 516 (A1g+B1g),

3. Results and Discussion 3.1. Electron Beam-Induced Chemical and Structural Changes. Under the present experimental conditions with electrons having a kinetic energy of 20 keV, atomic displacement and e-beam sputtering effects are improbable.37 As discussed previously,33 e-beam-induced localized changes in amorphous titanium oxide films under vacuum conditions occur through inelastic electron-electron scattering effects leading to oxide reduction via electron stimulated oxygen desorption at relatively low current densities, and to crystallization due to e-beam heating in the radiation damaged area in the case of sufficiently intense exposure. Both effects were discussed in detail elsewhere.34 Varying current density and dose at a constant absolute probe current IP of 5 µA, no threshold value of current density was found for reduction of a-TiO2, while crystallization occurred at 4.8 A/cm2 but not at 0.8 A/cm2. The global substrate heating at beam currents e5 A is negligible.34 The additional temperature rise ∆T in the beam center of relatively large e-beams can thus be estimated by33

∆T )

3UIP(1 - η) 2πκ(RE/2 + D)

(1)

where U is the accelerating voltage, η is the percentage of backscattered electrons (Fe: η ) 0.29), RE is the electron range (Fe: RE ) 1.2 µm at 20 kV), D is the probe diameter, and κ is the thermal conductivity (steel, κ ) 22 W m-1 K-1; TiO2, κ ) 11.7 W m-1 K-1). For D , RE/2, eq 1 reduces to the formula given by a derivation in ref 38.

Figure 1. Laser micro-Raman spectra of electrolytic TiO2 films on AISI 316L steel. (a) Non-irradiated film and after irradiation during 20 and 60 s at a probe current IP ) 1 µA corresponding to conditions no. 2 and 3 in Table 1. (b) After irradiation during 0-60 s at IP ) 5 µA corresponding to conditions no. 4 and 5 in Table 1.

Local Tuning of Conductivity in a-TiO2 Films

Figure 2. (a) FIB ion generated secondary electron image of an EPMA patterned region. e-beam exposure was performed at 5 µA during 5 s, corresponding to condition no. 4 in Table 1. The channelling effect inside the bright spots indicates the presence of crystalline structure. The film was protected with e-beam (eb) and ion-beam (ib) assisted Pt deposition prior to focused ion beam milling of the TEM lamella. (b) Composed TEM bright-field image of the transitional region between e-beam irradiated and partially crystalline (left) and non-irradiated amorphous (right) TiO2 at the edge of the Pt covered spot in (a).

and 636 cm-1 (Eg) becoming more intense up to 20 s of exposure. After 60 s, the weakening of the anatase modes is attributed to reduction to the TiO monoxide, known to be Raman inactive.40 We expect this phase to be crystalline, because it is formed from a crystalline phase giving previously rise to the characteristic anatase signal. This spectrum also indicates the existence of amorphous Ti2O3 in the range of 180-360 cm-1.41 As proposed in ref 34, crystallization cannot occur without partial oxide reduction, and it seems reasonable to assume that reduced crystalline TiO2-x phases may not be distinguished from stoichiometric anatase in Raman spectra, both giving rise to anatase signal until reduction of the crystalline phase to the monoxide is completed. Figure 1 underlines the importance of the absolute probe current and hence, ∆T, as an additional parameter to current density and dose, governing oxide crystallization. At IP ) 1 µA and J ) 5.1 µA/cm2 (∆T ) 55 °C), hardly any crystallization occurs, while at 5 µA and J ) 4.8 µA/cm2 (∆T ) 127 °C), crystallization is observed already after 1 s of exposure. This is clear evidence for the importance of thermal effects in the crystallization mechanism of radiation damaged zones, even under vacuum conditions. Previously reported crystallization of amorphous ZrO2 films kept at liquid helium temperature under TEM conditions42 was most possibly due to the generation of Frenkel defects at electron energies allowing for atomic displacement to happen, acting as nucleation centers for crystallization. This mechanism is not plausible at the present relatively low accelerating voltage. These new findings allow for a generalization of previously discussed necessary conditions34 for localized crystallization to occur: IP g 1 µA, J g 1 A/cm2, and t g 1 s. 3.2. TEM Analysis. The FIB micrograph in Figure 2a, being an ion-beam (Ga+) generated secondary electron image, represents an extraction of an e-beam patterned area and a platinum protected spot before ion-beam milling of the TEM lamella through the center of this spot. While hardly any contrast is obtained in normal secondary electron microscopy imaging mode, the large contrast between the irradiated spots (condition no. 2 in Table 1 with t ) 5 s) and the surrounding a-TiO2 film is indicative of the local e-beam-induced chemical and structural changes, strongly affecting the generation of secondary electrons. Clearly, fewer electrons are generated in the amorphous

J. Phys. Chem. C, Vol. 111, No. 37, 2007 13975 oxide. The dark spots within the bright irradiated spots are arising from ion channelling effects pointing indicative of the presence of crystalline structure. The four irradiated spots in Figure 2a correspond to the four spots at the lower left part analyzed by AFM in Figure 6b. A sharp transition between the irradiated zone and the surrounding amorphous oxide is evidenced in the TEM brightfield image in Figure 2b representing the edge of the spot in Figure 2a. The oxide thickness is approximately 60 nm at the non-irradiated part, and a very slight decrease in thickness from oxygen desorption can be observed in the exposed part. To the left in the EPMA irradiated zone, some contrast due to crystals penetrating the whole oxide thickness can be observed, responsible for the anatase Raman signal in Figure 1b (5 s). No signs of crystalline phase are detected to the right side. The lateral localization of crystallization is remarkable considering that this process is thermally activated and that the underlying steel substrate is a good thermal conductor, expected to heat also the area adjacent to the directly exposed spot. The observation underlines the role of the radiation damage limited to the irradiated zone, which, together with the vacuum conditions, is considered responsible for the important decrease of crystallization temperature as compared to atmospheric crystallization of the a-TiO2 film. Figure 3a represents a TEM bright-field image around the center of the irradiated spot with crystalline phase clearly visible. The steel substrate is to the left and the e-beam as well as the ion-beam deposited nanocrystalline Pt film are to the right side of the oxide film. The dark-field images in Figure 3b-d of the spots 1, 2, and 3 in the diffraction pattern (Figure 3e) show three different crystal orientations 200 (b), 101 (c), and 103 (d), respectively, in a still partially amorphous matrix. Figure 3f represents the diffraction pattern characteristic of nanocrystalline Pt, used for determination of the camera constant. Clearly, e-beam-induced crystals in a-TiO2 do not show any preferential orientation. A higher magnification of an e-beam-induced nanocrystal in Figure 3g and h indicates the presence of a remaining thin amorphous layer between this crystal and the steel substrate, which is also observed in Figure 3b-d. This indicates that the orientation of the e-beam-induced anatase nanocrystals is not influenced by the underlying steel crystals. 3.3. Local Modification of Electrical Properties. 3.3.1. Current Imaging at e-Beam Irradiated Spots. In Figure 4, current and topography images simultaneously acquired by c-AFM are presented for the 1 µA and the 5 µA probe current series (Table 1). The local volume loss visible in the topography is direct evidence for the oxygen desorption from oxide reduction during e-beam exposure. The probe diameter D in Table 1 is derived from the average width and height of the indentations in Figure 4. As can be seen, the beam shape in Figure 4b is slightly asymmetric. In Table 2, the quantitative values obtained from c-AFM experiments are summarized. The applied potential was +3.64 V for experiments in Figure 4a (5-8) and 4b (5-7) and +2.704 V for the scan in Figure 4a (8), the higher conductivity in the latter case leading to a current overflow at +3.64 V. The average total measured resistance R in irradiated spots in Table 2 is the sum of the contact resistance Rc and the film (or TiO2/ steel junction) resistance Rf. For calculating the conductivity (σ) of the TiO2/steel layer in Table 2, Rc was subtracted from the measured resistance. With the exception of the weak irradiation condition no. 1 in Table 1, a remaining local oxide thickness of 50 nm was considered in Table 2, assuming a loss of about 10 nm in the originally 60 nm thick film as suggested from topography images in Figure 4. Rc is found to be


Kern et al.

Figure 3. TEM bright-field (a) and dark-field (b-d) images of the central zone of the spot in Figure 2a, e-beam irradiated at 5 µA during 5 s, corresponding to condition no. 4 in Table 1. (e) Diffraction pattern of the partially crystalline TiO2 film with three indicated reflexes corresponding to anatase (200) in Figure 3b (1), anatase (101) in Figure 3c (2), and anatase (103) in Figure 3d (3). (f) Diffraction pattern of nanocrystalline Pt for determination of the camera length used for indexing of anatase structure in (e). Parts (f) and (g) show the bright-field and dark-field image of a nanocrystal at higher magnification.

approximately 1 MΩ from measurements on gold and steel reference substrates under static conditions with an applied contact force of 30 nN (Table 2). Because of the above-discussed compromise in tip coating material as well as the small externally applied force, the present Rc is significantly larger than reported values for metal coated cantilevers. However, because the total measured resistance of a-TiO2 is still a factor 4000 larger, the contact resistance is not problematic in the present study. Scanning the tip at 100 µm/s for current imaging in Figure 4 did only marginally increase Rc. In Table 1, the film resistivity F and the inverse conductivity σ are estimated by F ) (R - Rc)*(A/d), with A being the tip-surface contact area and d being the film thickness. Considering the simple Hertz solution53 for the contact area, we estimate a contact radius of about 5 nm to be reasonable under the present tip geometry, load, and film elasticity conditions. With Rc ) 1 MΩ, the resistivity of the a-TiO2 film at an applied potential of +3.64 V (R ) 4 × 10+10 Ω) is approximately 5.2 × 10+3 Ω cm. This is comparable to resistivities reported for sputter deposited anatase films54 and for sintered TiO2-anatase powder pellets.12 No values for amorphous TiO2 were found for comparison.

Plasma sprayed TiO2 coatings on steel have conductivities in the range of 1-100 cm-1,15 several orders of magnitude larger than the present electrolytic TiO2 films, most probably as a result of large amounts of intrinsic impurities. The rather high resistivity of the present amorphous TiO2 is explained by the previously reported low amount of contaminants in these electrolytically deposited films.34 In Figure 4a, a clear decrease in current corresponding to higher resistance at irradiated spots with respect to the surrounding a-TiO2 is observed after irradiation during 1, 5, and 20 s at IP ) 1 µA. Already after 1 s, R has increased by a factor 3 and continues to increase after 20 s of exposure to reach a factor 150 times the resistance measured on a-TiO2. Under these irradiation conditions, oxide reduction occurs as previously shown by AFM, wavelength dispersive X-ray analysis, and Auger spectroscopy,34 while crystallization is not possible as clearly visible in Figure 1a. After 60 s at IP ) 1 µA (Figure 4a (8)), which coincides with the onset of crystallization in the Raman spectrum in Figure 1a, the resistance clearly drops below the non-irradiated state (R ) 2.5 × 10+9 Ω), and a well-localized


J. Phys. Chem. C, Vol. 111, No. 37, 2007 13977

Figure 4. (a) c-AFM topography (1a (1-4)) and current (1a (5-8)) images after irradiation of a-TiO2 at a probe current of 1 µA during 1 s (1,5), 5 s (2,6), 20 s (3,7), and 60 s (4,8), corresponding to conditions no. 2 and 3 in Table 1. (b) c-AFM topography (1b (1-4)) and current (1b (5-8)) images after irradiation of a-TiO2 at a probe current of 5 µA during 1 s (1,5), 5 s (2,6), 20 s (3,7), and 60 s (4,8), corresponding to conditions no. 4 and 5 in Table 1. The applied external potential was 3.64 V except in Figure 4b (4,8) where 2.70 V was applied. The quantitative results from Figure 4 are given in Table 2.

path of increased conductivity is seen in the center of the irradiated spot. In agreement with a growing amount of crystalline phase evidenced from Micro-Raman data in Figure 1b, irradiation with IP ) 5 µA for 1 s further lowers the minimum local resistance (Figure 4b (5)), with R ) 4.4 × 10+8 Ω after 20 s of exposure, being approximately 110 times lower than in the surrounding a-TiO2. A further very significant drop in resistance in Figure 4b (8) coincides with the appearance of the monoxide TiO in Figure 1b, known to have a low resistivity close to metallic behavior.17 The measured resistance at this spot (2.70 × 10+6 Ω) is approximately 15 000 times lower than in the nonirradiated oxide (Table 2) and only a factor 2.7 larger than Rc measured on the reference steel substrate, the latter measured at a lower potential of 0.5 V, however. A conductivity saturation value of 0.3 Ω-1 cm-1 after high dose Ar ion implantation into rutile single crystals, and a maximum conductivity of 30 Ω-1 cm-1 after implantation of W and Sn ions were reported,22 close

Figure 5. Measured current (b) and total resistance R (X) as a function of applied external potential from c-AFM measurements of the nonirradiated a-TiO2 film on AISI 316L. The dotted line represents a linear fit through current data expected for ohmic behavior. Local film breakdown in the highly resistive a-TiO2 film was not observed below +14 V.


Kern et al.

Figure 6. c-AFM topography and current images on patterned a-TiO2 film electrodeposited on polished AISI 316L substrate. Topography (a) and current (b,c) images after e-beam patterning with a probe current of 5 µA for 5 s/spot (condition no. 4 in Table 1). The applied potentials were +2.21 V (b) and -2.21 V (c) with respect to the steel substrate. Topography (e) and current (f,g) images after e-beam patterning with a probe current of 500 nA for 5 s/spot (condition no. 1 in Table 1). The applied potentials were +3.64 V (f) and -3.64 V (g) with respect to the steel substrate. Parts (d) and (h) are three-dimensional representations of current images (b) and (f), respectively.

to the present conductivity measured after local e-beam reduction to the monoxide phase. Yet a further increase of conductivity would be expected under conditions where local e-beam reduction to the metallic phase occurs.34 The average current and the total resistance R in the imaged area as a function of applied potential are shown in Figure 5 for the non-irradiated a-TiO2 film. The current increases with potential, as expected for ohmic behavior, while the calculated resistance shows some fluctuations probably due to inhomogeneities within the film. The current images in Figure 4 were taken at potentials where the amorphous matrix exhibits highly resistive properties, the average current flowing through the cantilever/a-TiO2 contact being below 0.12 nA at potentials up to +4 V. Dielectric breakdown and irreversible damage to the film were not observed below +14 V. At +17 V, local breakdown at defects (pores) occurred, leading to irreversible

changes in local conductivity (not shown). It was beyond the scope of the present work to perform static and temperaturedependent I/V measurements on the a-TiO2 film or on locally modified spots for possible in-depth analysis of local semiconductor properties. In summary, in comparison to the highly resistive nonirradiated a-TiO2 film, c-AFM measurements indicate a further increase of the total measured resistance under conditions where reduction but no crystallization occurs, whereas a drop of the resistance is observed as soon as irradiation conditions are sufficiently intense to trigger partial local crystallization parallel to the ongoing oxide reduction. Crystallization results in paths of higher conductivity within an amorphous matrix, as the single nanocrystals extend almost throughout the total film, as seen in Figure 3. A further large drop in resistance corresponds to the formation of TiO phase. The present experiments show an


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TABLE 2: Applied Potential, Total Measured Resistance R, Estimated Film Thickness at Irradiated Spot, as Well as Calculated Resistivity and Conductivity Derived from c-AFM Experimentsa sample TiO2-x TiO2-x TiO2-x TiO2-x TiO2-x TiO2-x TiO2-x TiO2-x TiO2-x TiO2-x AISI 316L-ref Au-ref

IP (µA) 0 0.5 1 1 1 1 5 5 5 5 0 0

t (s)

potential (V)

resistance (Ω)

0 5 1 5 20 60 1 5 20 60 0 0

+3.64 +3.64 +3.64 +3.64 +3.64 +2.70 +3.64 +3.64 +3.64 +2.70 (0. 5 (0. 5

4.0 × 10 5.6 × 10+10 1.2 × 10+11 1.2 × 10+11 6.1 × 10+12 2.5 × 10+09 1.2 × 10+09 4.7 × 10+08 4.4 × 10+08 2.7 × 10+06 1.0 × 10+06 7.7 × 10+05 +10

thickness (nm)

resistivity (Ω cm)

conductivity (Ω-1 cm-1)

σ/σ0

60 55 50 50 50 50 50 50 50 50

5.2 × 10 8.0 × 10+03 1.9 × 10+04 1.9 × 10+04 9.5 × 10+05 4.0 × 10+02 1.8 × 10+02 7.4 × 10+01 7.0 × 10+01 2.7 × 10-01

1.9 × 10 1.2 × 10-04 5.4 × 10-05 5.2 × 10-05 1.1 × 10-06 2.5 × 10-03 5.5 × 10-03 1.4 × 10-02 1.4 × 10-02 3.7 × 10+00

1.0 × 10+00 6.5 × 10-01 2.8 × 10-01 2.7 × 10-01 5.5 × 10-03 1.3 × 10+01 2.9 × 10+01 7.1 × 10+01 7.5 × 10+01 2.0 × 10+04

+03

-04

a The conductivity change with respect to the non-irradiated a-TiO2 film is also given. Resistivity and conductivity are calculated after subtraction of the contact resistance Rc (1 ΜΩ), measured on AISI 316L and Au references. e-Beam irradiation conditions are given in Table 1 (IP ) probe current, t ) exposure time).

excellent correlation between structural and electrical properties. The decrease in resistivity as a result of local crystallization is qualitatively explained by the higher band gap of amorphous TiO2 (4.4 eV13) as compared to the crystalline phase (3.1-3.4 eV13). However, the increase in film resistance under pure reductive conditions is unexpected. Reduction of crystalline TiO2 has consistently been reported to result in decreased resistivity. For example, controlling the oxygen content in plasma-sprayed TiO2 rutile, Ohmori et al.15 measured a continuous drop of resistivity with oxygen loss in the coating. Also, Le Mercier et al.26 reported a decrease in electrical resistance by a factor 103 upon formation of mixed valence compounds in laser irradiated rutile single crystals. Today, nothing has been reported on local changes of conductivity in amorphous titanium oxide. In the present conditions, an increase of the resistance from e-beam deposited contaminations can be excluded, because the irradiated films are effectively plasma cleaned prior to local electrical measurements. Moreover, hydrocarbon contaminations are generally well-conductive. We thus attribute the increase in resistivity to an e-beam-induced increase in ordering and a diminution of defect density in the amorphous film, overbalancing the decrease in resistivity from the ongoing parallel oxide reduction. With respect to the irradiation-induced reduction mechanism, the measured higher resistivity of amorphous TiO2 as compared to the crystalline phase might explain the previously observed significantly increased sensitivity of amorphous TiO2 toward e-beam-induced oxygen desorption according to the known Knotek-Feibelman mechanism. In our previous experiments, e-beam exposure of partially crystalline TiO2 formed by anodic oxidation did not show any oxygen desorption, supposedly because its resistivity is not sufficient to allow an interatomic Auger decay necessary for oxygen desorption to happen.34 3.3.2. Local Diode BehaVior. In Figure 6, c-AFM images of EPMA patterned regions after irradiation at conditions no. 4 (t ) 5 s, Figure 6a) and no. 1 (Figure 6e) in Table 1, resulting in local crystallization and pure reduction, respectively, are presented. The topographies in Figure 6a and e reveal scratches arising from substrate polishing. Figure 6b and f shows current images acquired at +2.21 and +3.64 V, respectively, indicating spots of higher conductivity (R ) 4.71 × 10+8 Ω) in partially crystallized regions (6b) and slightly higher resistivity (R ) 4.1 × 10+9 Ω) after irradiation with IP ) 500 nA in partially reduced regions (6f) when compared to the surrounding a-TiO2 (R ) 3.95 × 10+9 Ω). Scratches in the substrate are completely covered by the electrodeposited a-TiO2 film and do not give

rise to local conductivity changes. When reversing the potential to -3.64 V in Figure 6c and g, no current is passing the imaged areas, exhibiting rectifying behavior and demonstrating the possibility of e-beam-induced formation of confined anatase/ steel Schottky barrier diodes at the partially crystallized spots (6b,c). To the best of our knowledge, this is the first time local diode behavior is observed by confined crystallization of an amorphous, semiconducting matrix. The shown behavior in Figure 6 implies little effect of the crystallographic orientation of the nanocrystals and the thin amorphous layers between nanocrystals and substrate on the local diode behavior (refer to section 3.2). Figure 6d and h shows three-dimensional representations of Figure 6b and f. The dark edges around the conductivity spots in Figure 6b, indicating higher resistance, are believed to be an artifact arising from a smaller tip-sample contact area at the slope of the e-beam-induced indentations. It is important to note that the electrical properties of the chosen TiN coated cantilever cannot be responsible for the observed rectifying effect, because a linear I-V characteristic in both biasing directions was obtained on the steel and the gold reference. Nevertheless, for detailed static studies of the observed diode behavior, a metal coated cantilever giving rise to a more ideal ohmic contact would be advantageous.55 4. Conclusion Local changes in oxidation state and structure in amorphous TiO2 thin films on steel substrates have been achieved by controlled e-beam irradiation at 20 keV using an electron probe microanalyzer. We identified the absolute probe current IP, additionally to the previously reported importance of current density and irradiation dose, as a key parameter for triggering confined crystallization. While no threshold values are identified for oxide reduction, at least partial crystallization to anatase phase is found possible under the following conditions: IP g 1 µA, J g 1 A/cm2, and t g 1 s. At the present electron energies disqualifying atomic displacement to happen, the dominating role of IP at otherwise identically chosen current densities underlines the importance of the estimated temperature rise in the irradiated zone for triggering local crystallization. A first sign of anatase formation under conditions provoking an estimated temperature rise ∆T of about 55 °C (T ≈ 80 C) is observed after 60 s of irradiation. TEM imaging through an irradiated spot leading to anatase signal in micro-Raman spectra shows the presence of anatase crystals penetrating the total film thickness. The transitional region between crystalline and amorphous TiO2 is well-defined

13980 J. Phys. Chem. C, Vol. 111, No. 37, 2007 despite a thermal energy dissipation occurring around the irradiated spot. This is evidence for the well-confined electron radiation damage limiting the lateral expansion of the structural change. No preferred crystal orientation was found in partially crystallized regions. c-AFM results demonstrate the possibility of e-beam-induced localized tuning of the conductivity in amorphous TiO2 films in both directions, dropping by up to a factor 5.5 × 10-3 in the case of reduction while increasing by up to a factor 2 × 10+4 upon crystallization and formation of TiO. This dual effect has, to the authors’ best knowledge, never been observed in the thoroughly investigated crystalline TiO2 films, powders, or single crystals. The feasibility of formation of micrometer-sized anatase/steel Schottky barrier diodes is demonstrated by c-AFM when inversing the anodic to a cathodic potential on a pattern of e-beam crystallized spots. While local conductivity spots in a highly resistive a-TiO2 matrix are observed under application of a positive potential on the steel substrate, hardly any current is passing when reversing the potential. The same phenomenon is found for a pattern of amorphous, locally reduced resistivity spots in the a-TiO2 matrix. The reported effects will allow for high-resolution patterning of electrical properties in amorphous semiconducting oxides in combination with electron lithography tools. A local increase in resistivity occurs already at small beam currents, and, therefore, high-resolution patterning can possibly be achieved using a field-emission electron gun. For local crystallization and paths of higher conductivity to be obtained, considerably larger currents are needed in the present system, in detriment of lateral resolution. However, crystallization with better-focused, weaker e-beams is considered possible when substrates with low thermal conductivity (e.g., glass or indium tin oxide) are chosen. Acknowledgment. We thank Ch. Ja¨ggi (EMPA Thun, CH) for performing Raman measurements. Y. Mu¨ller (EMPA Du¨bendorf, CH) is acknowledged for help with EPMA irradiation experiments. V. Friedli and S. Hoffmann (EMPA Thun, CH) are thanked for helpful discussions regarding diode behavior. References and Notes (1) Ashai, R.; Morikawa, T.; Ohwaki, T.; Aoki, K.; Taga, Y. Science 2001, 293, 269. (2) Linsebigler, A. L.; Lu, G. Q.; Yates, J. T. Chem. ReV. 1995, 95, 735. (3) Maruska, H. P.; Gosh, A. K. Sol. Energy Mater. 1979, 1, 237. (4) O’Regan, B.; Lenzmann, F.; Muis, R.; Wienke, J. Chem. Mater. 2002, 14, 5023. (5) Moon, W. T.; Lee, K. S.; Jun, Y. K.; Kim, H. S.; Hong, S. H. Sens. Actuators 2006, B115, 123 (6) Katayama, K.; Hasegawa, K.; Takahashi, Y.; Akiba, T. Sens. Actuators 1990, A24, 55. (7) Bange, K.; Ottermann, C. R.; Anderson, O.; Jeschkowski, U.; Laube, M.; Feile, R. Thin Solid Films 1991, 197, 279. (8) Sawada, Y.; Taga, Y. Thin Solid Films 1984, 116, L55. (9) Siefering, K. L.; Griffin, G. L. J. Electrochem. Soc. 1990, 137, 1206. (10) Textor, M.; Sittig, C.; Frauchiger, V.; Tosatti, S.; Brunette, D. M. Properties and Biological Significance of Natural Oxide Films on Titanium and its Alloys. In Titanium in Medicine; Brunette, D. M., Tengvall, P., Textor, M., Thomsen, P., Eds.; Springer: Berlin, Heidelberg, 2001; pp 176203. (11) Ja¨ggi, C.; Kern, P.; Michler, J.; Patscheider, J.; Tharian, J.; Munnink, F. Surf. Interface Anal. 2006, 38, 182. (12) Earle, M. D. Phys. ReV. 1942, 61, 56. (13) Afanas’ev, V. V.; Stesmans, A.; Chen, F.; Li, M.; Campbell, S. A. J. Appl. Phys. 2004, 95, 7936. (14) Kim, J. W.; Kim, D. O.; Hahn, Y. B. Korean J. Chem. Eng. 1998, 15, 217. (15) Ohmori, A.; Park, K. C.; Inuzuka, M.; Arata, Y.; Inoue, K.; Iwamoto, N. Thin Solid Films 1991, 201, 1.

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