Cluster-Assembled Nanostructured Titanium Oxide Films with Tailored

Sep 23, 2009 - Francesca BorghiBianca ScaparraCostanza PaternosterPaolo MilaniAlessandro Podestà. Langmuir 2018 34 (35), 10230-10242. Abstract | Full...
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Cluster-Assembled Nanostructured Titanium Oxide Films with Tailored Wettability A. Podesta`,*,† G. Bongiorno,†,‡ P. E. Scopelliti,† S. Bovio,† and P. Milani*,†,‡ CIMAINA and Dipartimento di Fisica, UniVersita` degli Studi di Milano, Via Celoria 16, 20133 Milano, Italy, and Fondazione Filarete, V.le Ortles 22/4, 20139 Milano, Italy

C. Semprebon and G. Mistura Dipartimento di Fisica G.Galilei and CNISM, UniVersita` di PadoVa, Via Marzolo 8, 35131 PadoVa, Italy ReceiVed: June 24, 2009; ReVised Manuscript ReceiVed: August 28, 2009

We report on the production of cluster-assembled nanostructured titanium oxide (ns-TiOx) surfaces with tailored wettability properties. The contact angle of the ns-TiOx surfaces changes in a controllable and reproducible way by changing the morphology through the roughness and/or with annealing the films in air at different temperatures. With increasing the annealing temperature, reversible switching from hydrophobic to hydrophilic behavior can be obtained, due to the removal of organic contaminants and to the increase of surface hydroxyl groups, while control of surface morphology allows fine-tuning of the contact angle within a given wetting regime. Thanks to the possibility of regulating and controlling the surface wettability, ns-TiOx thin films are very promising substrates for protein microarray applications where a reproducible control on protein spot dimension is required. Introduction The understanding and the control of the chemical and morphological factors influencing the wettability of solid surfaces is necessary in view of the fabrication of biocompatible systems for protein and cell microarrays.1,2 At present chemical functionalization is the most widely used approach for the modification of the wettability of surfaces: typically, hydrophilic surfaces are prepared by chemically increasing the number of hydroxyl OH groups on the surface, while hydrophobic surfaces are obtained by binding hydrophobic moieties (silanes, as well as fluorinated groups) or by coating with hydrophobic polymers. Chemical modification allows switching from superhydrophilic (θ ≈ 0°) to hydrophobic (θ ≈ 150°) behavior. Although quite flexible and versatile, chemical functionalization may not be compatible with applications requiring biocompatibility so that different strategies taking into account surface morphology should be developed. Since the pioneering works of Wenzel and Cassie and Baxter,3,4 many authors have pointed out the critical influence of morphological parameters such as surface roughness and pattern at the micro- and nanoscale on the wetting properties.5-10 In the simple Wenzel picture, the increase of surface specific area and roughness induced by surface patterning or roughening enhances the intrinsic wetting character of the surface. Special surface geometries, such as regular micrometer-scale holes, strips, and spikes on silicon, were successfully employed for the realization of almost fully water-repellent surfaces, exploiting the effect of air pockets formation under the liquid drop, typical of materials showing θ > 90° for flat surfaces.5,8 Following the same scheme, intrinsically hydrophilic surfaces can be made more wettable by increasing their specific area, although the mechanism of air pockets formation and the breaking-up of the * To whom correspondence should be addressed. [email protected] and [email protected]. † Universita` degli Studi di Milano. ‡ Fondazione Filarete.

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simplified Wenzel model for complex surfaces make it possible under particular conditions to obtain a morphology-driven hydrophilic-to-hydrophobic transition.9,10 Complex and expensive fabrication technologies can be used to design with great precision the wettability of surfaces. This is obtained by a top-down approach based on material removal to create a pattern through lithographic parallel processes, or by serial etching processes using laser or electron beams. These technologies have been developed for the fabrication of MEMS and can be applied to silicon or polymeric materials.11,12 An alternative route is the exploitation of the morphological complexity of random rough surfaces at a micrometer and nanoscale obtained by chemical etching,13 mechanical milling (see ref 14), or nanoparticle assembling.15,16 A potential advantage of materials resulting from the bottom-up assembling of nanoparticles is that it is possible to modify the surface roughness without simultaneously affecting the chemical nature of the surface, as typically happens with crystalline materials, where surface roughening will result in the exposure of new lattice planes with different intrinsic surface energies. In the case of materials obtained by assembling of nanoparticles, the overall surface chemistry is basically the same as that of the single nanoparticles, which does not change when the film gets rougher, because the intrinsic nanostructure is conserved. Despite the high efficiency and low cost of the synthesis techniques that are used to produce random rough surfaces, it is difficult in general to control the morphology of random rough surfaces and this is the reason why very few studies have been published on the modification of random rough surface to control the wettability as an alternative to top-down patterning and chemical functionalization. A powerful technique allowing the synthesis of nanostructured thin films with fine control over the surface morphology is supersonic cluster beam deposition (SCBD).15 The surface morphology of films deposited by SCBD has been shown to evolve with the deposition time according to simple scaling laws, which make it easy to tailor parameters such as root-mean-

10.1021/jp905930r CCC: $40.75  2009 American Chemical Society Published on Web 09/23/2009

Cluster-Assembled Nanostructured TiOx Films square roughness, specific area, as well as in-plane correlation length.17 We have recently demonstrated that controlling the self-affine surface morphology of nanostructured cluster-assembled carbon films allows modulating their wettability across the hydrophilic/hydrophobic regime.18,19 Cluster-assembled nanostructured titania (ns-TiOx) films are of great interest for biomedical applications. Recently we have demonstrated that ns-TiOx films deposited by SCBD have excellent biocompatibility by performing long-term experiments with a range of cancer and primary cells.20 ns-TiOx films deposited by SCBD resulting from a random stacking of nanoparticles are characterized, on the nanoscale, by granularity and porosity mimicking those of extracellular matrix structures.20 Their large nanoscale porosity, along with the abundance of adsorption sites and defects (see ref 21), makes ns-TiOx a promising candidate as a substrate for the adsorption and stable docking of proteins, and for the realization of high-throughput protein microarrays. Protein microarray is a new powerful technology for highthroughputstudiesofprotein-protein,protein-lipid,protein-DNA, protein-drugs, and protein-small molecules interaction,1 which has great potential in basic research,22 diagnostics (see refs 23-25), and drug discovery applications.26 Protein microarray production is based on spotting of protein solution droplets on a substrate in order to obtain a high-density array with thousands of different immobilized protein probes. Controlled and reproducible substrate wettability is crucial to reduce sample-tosample variations and to have reliable, comparable, and reproducible results.27,28 We show in this paper that by combining the control of the surface morphology by varying the film thickness with that of the surface chemical state by postdeposition mild thermal annealing, we can tune the wettability of ns-TiOx films across a wide range of contact angles, from 0 to ∼140°. While thermal treatments allow switching from one wetting regime to the other (from hydrophilic to hydrophobic), fine control of surface roughness allows the fine modulation of the contact angle within a given wetting regime. The influence of controlled ns-TiOx wettability on protein spot dimension has been characterized. Materials and Methods Film Synthesis and Postdeposition Annealing. Nanostructured TiOx films were deposited with use of an SCBD apparatus equipped with a pulsed microplasma cluster source (PMCS) as described in detail in refs 15, 29, and 30. The PMCS operation principle is based on the ablation of a titanium rod by a helium or argon plasma jet, ignited by a pulsed electric discharge.29,30 After the ablation, Ti and TiOx ions thermalize with helium or argon and condense to form clusters. The mixture of clusters and inert gas is then extracted in vacuum through a nozzle to form a seeded supersonic beam, which is collected on a set of round glass coverslips (diameter 15 mm, thickness 0.13-0.16 mm, Electron Microscopy Sciences) intercepting the beam. The clusters kinetic energy is low enough to avoid fragmentation and hence a nanostructured film is grown. Cluster mass distribution is affected by the carrier gas: typical mass distribution with He as carrier gas is log-normal and it peaks at several hundreds atom per cluster, corresponding to a diameter of roughly 1 nm. The use of Ar shifts the distribution toward larger masses.15 The as-deposited material is poor of oxygen; however, it is oxidized very quickly as soon as it is exposed to air. It follows that the value of x in ns-TiOx is very close to (but less than) 2. The slight substoichiometry is due to defects located on the surface of nanoparticles (undercoordinated Ti atoms,

J. Phys. Chem. C, Vol. 113, No. 42, 2009 18265 oxygen vacancies) and at the boundaries between nanoparticles. The stoichiometry of ns-TiOx is not influenced by the carrier gas. The postdeposition thermal treatments have been carried out in a tube furnace in ambient air atmosphere. The heating cycle comprised three phases: an initial climb ramp from room to the final temperature at a rate of 240 deg/h, followed by a plateau at the maximum temperature (100, 200, and 400 °C) over a time period of 2 h, and terminated by the cooling ramp, at a rate of about 100 d/h. When the temperature fell to about 100 °C, the samples were extracted from the furnace and let cool to room temperature in air, at this point being ready for the contact angle characterization. For the production of protein arrays, standard optical microscopy glass slides have been coated by ns-TiOx films with thickness of 50 and 200 nm, using Ar as carrier gas, and annealed for 3 h in a muffle furnace at temperatures of 200 and 250 °C. Controlling Film Morphology. Film roughness is a monotonically increasing function of film thickness, provided the deposition rate is constant. The deposition rate is about 10-15 nm/min when He is used as the carrier gas and about 15-20 nm/min when Ar is used. To produce films with uniform thickness across the whole substrate area, a much larger area is raster-scanned by the source during deposition, resulting in smaller deposition rates. Typically, in a few hours, ns-TiOx films with a thickness up to 50 nm can be grown with He, while a thickness up to 400 nm can be obtained with Ar. Maximum values for film roughness are 10 and 30 nm, accordingly. AFM Characterization of Morphology of ns-TiOx Films. We have used a Nanoscope IV multimode atomic force microscope (Veeco Instruments) for the quantitative characterization of surface morphology of ns-TiOx films. The AFM was equipped with rigid cantilevers (resonance frequency 250-300 kHz) with single-crystal silicon tips (nominal radius 5-10 nm), and operated in Tapping Mode in a dry nitrogen atmosphere. Typically, several (4-6) 2 µm × 1 µm images (2048 × 1024 points) were acquired on each sample, and flattened by lineby-line subtraction of first and second-order polynomials in order to get rid of the tilt of the sample and of the scanner bow. From flattened AFM images, the average microscale root-mean-square roughness and specific area parameters were calculated. The specific area is the ratio of the surface area to the projected area. The specific area calculated from AFM images is always underestimated because of the inability of the AFM tip to detect overhangs and because of its finite size. Root-mean-square roughness is usually a more robust parameter. Film thickness was calculated from AFM images acquired across a sharp step produced by masking the film during the deposition. Contact Angle Measurements. Contact angles of water have been measured with a homemade apparatus consisting of a syringe pump, a video camera, and motorized sample and camera stages, all of them controlled via a PC. Small drops (volume ∼0.5 mL) of Milli-Q water were produced with the syringe pump and gently deposited on the surface. For each image, the overall drop profile was fitted with an elliptic curve and the error related to the fitting procedure was typically less than (1°.31 To obtain statistically sound results, at least five drops for each sample were typically analyzed. The representative contact angle θ was then taken as the mean of these different determinations and the corresponding standard deviation was around (2°, unless otherwise stated. Contact angles have been measured on the same samples that underwent the systematic AFM characterization of surface

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Figure 1. Three-dimensional and top-views of AFM topographies of (a, b) a thin ∼15 nm thick ns-TiOx film deposited by using He as carrier gas and (c, d) a thick 150 nm thick ns-TiOx film deposited by using Ar as the carrier gas. The vertical color scales in parts b and d are 20 and 120 nm, accordingly.

morphology, at room temperature as well as after each annealing step at 100, 200, and 400 °C. Protein Arrays. Bovine Serum Albumin (BSA) conjugated with Alexa 647 dye solution (provided by Invitrogen) is prepared at 2 µM concentration in PBS buffer. Protein solution droplets (40 nL) are spotted on ns-TiOx coated slides, in an array 10 × 24 format, using a BioJet 3000 Plus spotter from BioDot. After 1 h of incubation slides are washed in 2% BSA and rinsed two times with PBS and two times with Milli-Q water. The fluorescence signal of the slides is read with a GenePix 4000B microarray scanner (635 nm laser wavelengths, 100% laser power, 580 PMT gain). Images were analyzed with GenePix Pro 4.0 software. Results and Discussion Surface Morphology of ns-TiOx Films. Figure 1 shows the comparison between AFM topographies acquired on two nsTiOx films, with thickness 15 and 150 nm, respectively. The vertical scale in the three-dimensional maps (panels a and c) is the same for both images in order to facilitate the comparison. In panels b and d the top-view of the same AFM maps are shown, in order to highlight the granularity and the lateral correlations developed in the two films. Figure 1 clearly indicates that the morphology of ns-TiOx films deposited by SCBD consists of a fine raster of nanometer-sized grains, with high specific-area, and porosity at the subnanometer scale.32 This peculiar morphology is the consequence of the low-energy deposition regime typical of supersonic cluster beam deposition,15 and of the ballistic deposition regime, where sizedispersed clusters impinge on the surface and stick without being significantly diffuse.17 The size of the smallest units observed in panels b and d of Figure 1 is at the nanometre level, larger units resulting from the coalescence of single TiOx clusters, occurring already in the cluster beam. It can be noticed that in the thick film (Figure 1c,d) the smallest topographic features appear clustered in larger features, determining both larger height fluctuations of the interface and spatial correlations in the in-plane directions extending well beyond the size of a single cluster. Such multiscale, self-affine morphology, which extends in the case of these films only up to a couple of decades because of the reduced overall thickness of the films, is recognized to play a role in determining a strong coupling between wettability and morphological properties.6,7 Figure 2 shows the correlation between the morphological parameters of ns-TiOx films characterized by AFM and the film

Figure 2. Correlation between the morphological parameters (roughness and specific area) and the film thickness of ns-TiOx films, deposited with both He and Ar as carrier gas. Error bars, when not visible, are smaller than the data markers.

thickness. Remarkably, morphological parameters of ns-TiOx films are similarly correlated, irrespective of the carrier gas used in the cluster source. The latter influences the deposition rate, i.e., the maximum thickness obtainable in a typical deposition session, and therefore the range of roughness and specific area values obtainable. The size distribution of Ti clusters is also affected by the choice of the carrier gas. A systematic analysis of the evolution of the surface morphology of ns-TiOx films shows that the growth mechanisms of ns-TiOx films and therefore the scaling of morphology are not influenced by the carrier gas, nor do they depend on the fine details of the nanoparticles size distribution function.33 Once the operational parameters of the cluster source are fixed (among these, the carrier gas, Ar or He), the film thickness depends only on the deposition time. The deposition time can be used therefore as the driving parameter to control the rootmean-square roughness and the specific area of ns-TiOx films, according to the calibration curve shown in Figure 2. Effect of Annealing on Film Morphology. We have investigated whether thermal annealing of ns-TiOx films at temperatures up to 400 °C has some influence on the morphology of films. This is an important point in order to assess the relative importance of surface morphology and surface chemistry in determining the wetting properties of ns-TiOx films. We have compared the root-mean-square roughness and specific area

Cluster-Assembled Nanostructured TiOx Films

Figure 3. Contact angles and corresponding cos(θ) values measured on ns-TiOx films produced in different deposition and postdeposition conditions as a function of root-mean-square roughness. The contact angle measured on a single-crystal rutile TiO2 sample at room temperature is also shown as reference. The dotted lines are a linear fit of data. Error bars, when not visible, are smaller than the data markers.

values measured on films with different thickness before and after annealing at temperature up to 400 °C (data not shown) and found no evidence of relevant morphological changes induced by thermal annealing. The maximum relative change of roughness observed on the roughest film was below 3%. Surface morphology and surface chemistry are therefore two independently controllable properties of ns-TiOx films deposited by SCBD. Considered the strong influence that these parameters have on the contact angle of water on ns-TiOx films, their mutual independence is a key factor for the production of films with tailored wettability. Contact Angles vs. Morphology. Figure 3 shows the contact angle θ and cos(θ) values measured on films with different rootmean-square roughness, which underwent different thermal treatments. The polished 〈100〉 surface of a single-crystal rutile TiO2 sample has also been characterized and the corresponding morphological and wetting parameters have been included in Figure 3 as a reference. The surface morphology of the rutile sample is fairly ideal, with a roughness of only a few angstroms and a specific area of ∼1, values typical of nearly atomically smooth surfaces. Our data show that postdeposition thermal annealing changes the overall wetting character of ns-TiOx films: while as-deposited films are hydrophobic, annealing at 200 °C makes them mildly hydrophilic, and annealing at 400 °C turns them into superhydrophilic films. The improved wettability of ns-TiOx films upon annealing at moderate temperatures can be explained in terms of removal of physisorbed hydrophobic organic contaminants and of the recovering of OH groups bonded to undercoordinated Ti atoms.34 Morphology has an important impact on the wetting behavior of ns-TiOx. Controlling surface roughness in the range 3-30 nm allows tuning the contact angle from 140° to 90° in the hydrophobic regime, from 70° to 25° in the hydrophilic regime. Remarkably, in the superhydrophilic regime it is still possible to avoid complete wetting by keeping the roughness parameter below 5 nm, i.e., by depositing films with thickness below 20 nm. Combining morphological and chemical surface properties of ns-TiOx it is therefore possible to tune the wetting properties of these films with high elasticity, spanning almost the whole range of contact angles, from 0° (complete wetting) to 140° (almost superhydrophobic).

J. Phys. Chem. C, Vol. 113, No. 42, 2009 18267 The measured cos(θ) values for the samples as-deposited and annealed at 200 °C scale linearly to a good approximation with the surface roughness. The dotted lines in Figure 3 represent linear fits to the data. The trend observed for the samples annealed at 200 °C is compatible with the Wenzel equation, predicting a positive slope for the cos(θ) vs. specific area curve, i.e., the enhancement of the intrinsic hydrophilic character of the surface.35 The same trend in our case is expected for the cos(θ) vs. roughness curve, because the two quantities are linearly correlatedssee Figure 2. In the case of as-deposited films, we observe that the poor hydrophilicity of the flat sample (θRutile ) 82°) is reverted, and the surface becomes more and more hydrophobic as the roughness increases. In other words, the surface wettability obeys an effective Wenzel rule, but with a critical angle that is smaller than 90°; in the case of as-deposited ns-TiOx samples the critical angle is actually smaller than θRutile ) 82°. Such effective Wenzel behavior has been recently observed for cluster-assembled nanostructured carbon films deposited by SCBD and wetted by water,19 and for functionalized porous silicon surfaces.36,37 This effective Wenzel regime, according to some authors, is a consequence of the effect of a complex random surface morphology, which can be regarded as a random composition of grooves, with radial and circular symmetry.9,10,38 Another factor that can provide an enhancement of hydrophobicity of rough surface is the formation of air pockets,4,7 which is possible also in intrinsically hydrophilic surfaces provided some overhang is present, at least at the smallest scales.6 This condition is certainly satisfied in nanoporous ns-C and ns-TiOx films. Moreover, theoretical works have recently predicted that the multiscale (self-affine) character of a surface, like that of our nanostructured surfaces, promotes transition toward superhydrophobicity, irrespective to the intrinsic wettability of the surface.6,7 Remarkably, both ns-C and ns-TiOx films possess a self-affine morphology.18 Extrapolation toward zero roughness (and toward unity specific area) of the contact angle data provides two different values of the intrinsic contact angle for the films annealed at 200° deposited with different carrier gases, while the contact angle measured on rutile at room temperature represents an excellent extrapolation for both types of as-deposited ns-TiOx films (the weighted linear fit of data provides a contact angle at zero roughness of 80((2)° and 79((10)° for the films grown with He and Ar, accordingly, in good agreement with the contact angle of rutile). This could be due to the fact that the surface chemistry, and therefore the contribution to the contact angle that is not morphology-dependent, is determined by the degree of surface contamination by organic moieties and by the effective hydroxylation of the surface; these properties in turn are well-stabilized in the case of as-deposited films, which did not undergo any postdeposition treatments, while in the case of annealed samples, the effectiveness of thermal treatment in removing organics from the surface can strongly depend upon the surface morphology, and especially on granularity and porosity. Moreover, we observe in Figure 3 that for all films the slope of the linear trends depends on the carrier gas used during deposition, despite the fact that the statistical morphological properties of films deposited with different carrier gases evolves similarly with thickness (see Figure 2), and that the surface chemistry is expected to be rather independent of the carrier gas. One possible explanation for the observed difference in contact angles is the different porosity (different granularity) of the films deposited with different carrier gases, which can

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Figure 5. Influence of (a, c) the postdeposition annealing temperature (nominal film thickness 50 nm, AD for as-deposited) and (b, d) the surface roughness on the protein spot diameter on ns-TiOx films. The ns-TiOx films in panels b and d have been annealed at 250 °C, then exposed to air for 8 h before spotting, in order to partially recover their original wettability. The size of the smallest spots in panels a and b is 600 and 800 nm accordingly. Figure 4. Recovery curves for the contact angle of ns-TiOx films. The contact angles of the (a) thinnest and (b) roughest films have been measured immediately before (day -1) and after (day 0) the thermal annealing at 200 and 400 °C, and then regularly up to 20-40 days after the annealing date.

account for different evolution of air pockets, and therefore for the observed different contact angles. Changes in wettability induced by thermal treatments are not permanent. ns-TiOx films recover their initial wetting character on a time scale of several days upon exposure to ambient air. Contact angle recovery has been characterized by measuring the contact angle immediately before and after annealing, and then on a regular base for several days up to a few months after annealing. In Figure 4 we report recovery curves for the smoothest and the roughest films produced in this experiment. The thermal treatment is done at “day 0” and the monitoring of contact angle proceeds until the original contact angle value (the one recorded before the treatment, at day -1) is recovered. A typical complete recovery curve spans 30-40 days, although the typical time scale of wettability recovery is below 24 h. We have observed that the recovery of the contact angles after thermal annealing is a reproducible and stable process. If kept in protected environment (dry nitrogen, or vacuum), recovery can be slowed down, and the thermally induced hydrophilicity can be maintained for several days. In any case, the recovery kinetics is slow enough to ensure that during an experiment involving ns-TiOx films started immediately after the annealing process the surface wettability does not change, an important point for applications As already discussed, the main effect of annealing is to remove weakly bound organic moieties that make the surface hydrophobic, and to recover the original surface density of hydroxyl groups, which make the surface of metal oxides hydrophilic. After annealing, surface contamination initiates again, accelerated by the high reactivity of the TiO2 surface (in turn enhanced by the high effective area and high density of defects and undercoordinated Ti sites of the nanostructured material). The process tends toward saturation, when the initial wettability is recovered. For annealing temperatures not exceeding 400 °C, no significant morphological and structural modifications of ns-TiOx films take place; therefore the original surface chemistry can be reproducibly recovered. For temper-

atures higher than 400 °C, some irreversible modifications of the film can be expected. It turns out that moderate thermal treatments are a simple and effective way of reversibly switching ns-TiOx films from the mildly hydrophobic to the hydrophilic regime, up to complete wetting. Annealing allows the surface chemistry to be suitably modified without the need of any complex and irreversible chemical functionalization processes. When the surface morphology is concurrently controlled, the wettability of ns-TiOx films can be fine-tuned across the whole range of contact angles that are interesting for applications. Protein Microarrays. To check the possibility of regulating and controlling protein spot dimensions in a protein microarray experiment with ns-TiOx as substrate, fluorescent protein droplets are spotted in a 24 × 10 matrix on a ns-TiOx-coated glass slide, as described in the Materials and Methods section. Surface-drop contact area, and consequently final spot dimension, depends on surface wettability, which can be tailored, as demonstrated in the previous sections, either by performing thermal treatments or by controlling surface morphology, or by a combination of both. In Figure 5a,b ns-TiOx substrates with nominal thickness 50 nm that have undergone different thermal treatments are compared with respect to the protein spot diameter. The similarity of the thickness assures that the surface roughness of the films is also similar. Spot diameter can be decreased by about 20% by reducing the annealing temperature from 250 to 200 °C, and by about 45% by avoiding thermal treatment. The more hydrophilic the surface, the larger the spot diameter. Thermal treatment also influences the amount of immobilized proteins.39 This effect can be mainly attributed to the release of atmospheric organic contaminants. For higher annealing temperature there are more released contaminants, and therefore more free sites for protein immobilization. In Figure 5c,d the effect of surface morphology on protein spot diameter is shown for films that, after a thermal treatment at 250 °C, were allowed to partially recover their original wettability (hydrophobic) by exposing them to air for 8 h. The spot diameter can be reduced by 20% by increasing the nominal film thickness from 50 to 200 nm, which corresponds to a measured increase of surface roughness from 11 to 19 nm. Here the property of surface roughness to inhibit wettability of intrinsically hydrophobic surface has been exploited.

Cluster-Assembled Nanostructured TiOx Films These results show that ns-TiOx allows controlling the protein spot dimension with high flexibility, giving the possibility to manage the trade-off between array density, amount of immobilized proteins, and correct proteins folding in protein microarray applications. Indeed, wettability not only defines the ultimate array density, but also influences the conformational state of immobilized proteins. In fact hydrophobic interactions can destabilize protein structures, because of the possible exposition to the surface of the inner hydrophobic residues, which may cause protein unfolding and loss of protein functionality. With ns-TiOx coatings it is possible to optimize surface wettability in order to obtain the best array performance. Since ns-TiOx coated slides also have high protein immobilization capability,39 they are promising substrates for several microarray approaches, for example, for reverse phase protein microarray (RPM),24 in which a large amount of protein has to be functionally immobilized onto the surface from cell or tissue lysates. ns-TiOx properties can be exploited also for the production of chemically coated microarray slides. For most microarray applications chemical coating is used to promote protein immobilization,40 and ns-TiOx surfaces can be easily functionalized with several chemical groups such as silane, epoxy, or amine groups. In principle, a suitable choice of the surface morphology could allow adjusting the surface wettability that can be influenced by the chemical functionalization. This opens the possibility to optimize the performance and to better control the density of these arrays. Conclusions We have shown that it is possible to grow cluster-assembled nanostructured TiOx films by supersonic cluster beam deposition with tailored wettability. In particular, the contact angle of water on ns-TiOx can be tuned reliably across the 0-140° range by exploiting a combination of inexpensive, quick, postdeposition thermal treatments in air at moderate temperatures (e400 °C), and control of surface morphology, which can be reproducibly adjusted by controlling the film thickness during deposition. These two channels for controlling film wettability turned out to be independent from each other. Changes in wettability induced by thermal annealing are reversible, the complete wettability recovery occurring in ambient air within the first 24 h. The surface wetting properties are therefore stable in air on the typical time scales of experiments. Overall, this approach represents a powerful and relatively inexpensive strategy for the production of nanostructured materials with tailored wetting properties, which can provide more elasticity than standard approaches based heavily on surface chemical functionalization and/or surface regular micromachining and patterning (the latter hardly implementable on titania surfaces). We have also shown that the control of surface wettability allows controlling the protein spot diameter, i.e., the protein spot density, of ns-TiOx-based protein microarrays. This gives the possibility to better optimize array performance, regulating surface wettability depending on the needed array density. Moreover, the control over surface wettability allows more reliable and reproducible results to be obtained in microarray experiments. Acknowledgment. We thank Luca Giorgetti for support in the array experiments. This work was supported by Fondazione Cariplo under project “Sviluppo di Film Fotocatalitici Nanostrutturati per la Conversione di Energia su Micropiattaforme” (ref no. 2006-0660) and by the University of Padova under project CPDA077281/07.

J. Phys. Chem. C, Vol. 113, No. 42, 2009 18269 References and Notes (1) Macbeath, G.; Schreiber, S. L. Science 2000, 289, 1760–1763. (2) Ziauddin, J.; Sabatini, D. M. Nature 2001, 411, 107–110. (3) Wenzel, R. N. J. Phys. Colloid Chem. 1949, 53, 1466. (4) Cassie, A. B. D.; Baxter, S. Trans. Faraday Soc. 1944, 40, 546. (5) Bico, J.; Marzolin, C.; Quere, D Europhys. Lett. 1999, 47, 220– 226. (6) Herminghaus, S. Europhys. Lett. 2000, 52, 165–170. (7) Yang, C.; Tartaglino, U.; Persson, B. N. J. Eur. Phys. J. E 2008, 25, 139–152. (8) Patankar, N. A. Langmuir 2003, 19, 1249–1253. (9) Zhou, X. B.; De Hosson, J. Th. M. J. Mater. Res. 1995, 10, 1984– 1992. (10) Palasantzas, G.; De Hosson, J. Th. M. Acta Mater. 2001, 49, 3533– 3538. (11) Judy, J. W. Smart Mater. Struct. 2001, 10, 1115–1134. (12) Liu, C AdV. Mater. 2007, 19, 3783–3790. (13) Wang, M.-F.; Raghunathan, N.; Ziaie, B Langmuir 2007, 23, 2300– 2303. (14) Calka, A.; Wexler, D. Nature 2002, 419, 147–151. (15) Wegner, K.; Piseri, P.; Tafreshi, V. H.; Milani, P. J. Phys. D: Appl. Phys. 2006, 39, R439–459. (16) Velev, O. D.; Gupta, S AdV. Mater. 2009, 21, 1897–1905. (17) Barabasi, A.-L.; Stanley, H. E. Fractal Concepts in Surface Growth; Cambridge University Press: New York, 1995. (18) Milani, P.; Podesta`, A.; Piseri, P.; Barborini, E.; Lenardi, C.; Castelnovo, C. Diamond Relat. Mater. 2001, 10, 240–247. (19) Ostrovskaya, L.; Podesta`, A.; Milani, P.; Ralchenko, V. Europhys. Lett. 2003, 63, 401–404. (20) Carbone, R.; Marangi, I.; Zanardi, A.; Giorgetti, L.; Chierici, E.; Berlanda, G.; Podesta, A.; Fiorentini, F.; Bongiorno, G.; Piseri, P.; Pelicci, P. G.; Milani, P. Biomaterials 2006, 27, 3221–3229. (21) Caruso, T.; Lenardi, C.; Agostino, R. G.; Amati, M.; Bongiorno, G.; Mazza, T.; Policicchio, A.; Formoso, V.; Maccallini, E.; Colavita, E.; Chiarello, G.; Finetti, P.; Sutara, F.; Skala, T.; Piseri, P.; Prince, K. C.; Milani, P. J. Chem. Phys. 2008, 128, 094704. (22) Zhu, H.; Bilgin, M.; Bangham, R.; Hall, D.; Casamayor, A.; Bertone, P.; Lan, N.; Jansen, R.; Bidlingmaier, S.; Houfek, T.; Mitchell, T.; Miller, P.; Dean, R. A.; Gerstein, M.; Snyder, M. Science 2001, 293, 2101–2105. (23) Speer, R.; Wulfkuhle, J.; Espina, V.; Aurajo, R.; Edmiston, K. H.; Liotta, L. A.; Petricoin, E. F. Cancer Genomics Proteomics 2007, 4, 157– 164. (24) Petricoin, E. F.; Zoon, K. C.; Kohn, E. C.; Barrett, J. C.; Liotta, L. A. Nat. ReV. Drug DiscoVery 2002, 1, 683–695. (25) Zhu, H.; et al. Proc. Natl. Acad. Sci. 2006, 103, 4011–4016. (26) Krishnanand, D. K. Expert Opin. Drug DiscoVery 2007, 2 (11), 1467–1476. (27) Smith, J. T.; Benjamin, L. V.; Monty, W. R. Langmuir 2002, 18 (16), 6289–6293. (28) Ho, C. E.; Tseng, F. G.; Lin, S. C.; Su, C. J.; Liu, Z. Y.; Yu, R.-J.; Chen, Y. F.; Huang, H.; Chieng, C.-C. J. Micromech. Microeng. 2005, 15, 2317–2325. (29) Barborini, E.; Piseri, P.; Milani, P J. Phys. D: Appl. Phys. 1999, 32, L105-L109. (30) Tafreshi, H. V.; Piseri, P.; Benedek, G.; Milani, P. J. Nanosci. Nanotechnol. 2006, 6, 1140–1149. (31) Semprebon, C; Mistura, G.; Orlandini, E.; Bissacco, G.; Segato, A.; Yeomans, J. M. Langmuir 2009, 25, 5619. (32) Kholmanov, I. N.; Barborini, E.; Vinati, S.; Piseri, P.; Podesta`, A.; Ducati, C.; Lenardi, C.; Milani, P. Nanotechnology 2003, 14, 1168–1173. (33) Podesta`, A. To be submitted for publication. (34) Takeda, S.; Fukawa, M.; Hayashi, Y.; Matsumoto, K. Thin Solid Films 1999, 339, 220–224. (35) We could not measure the contact angle of the rutile sample after annealing at 200 °C, but both the extrapolated value toward zero-roughness and the contact angle measured on rutile at room temperature (below 90 °C) suggest that the intrinsic wetting character of samples treated at 200 °C is hydrophilic. (36) Dattilo, D.; Armelao, L.; Maggini, M.; Fois, G.; Mistura, G. Langmuir 2006, 22, 8764–8769. (37) Dattilo, D.; Armelao, L.; Fois, G.; Mistura, G.; Maggini, M. Langmuir 2007, 23, 12945–12950. (38) Oliver, J. P.; Huh, C.; Mason, S. G. Colloids Surf. 1980, 1, 79– 104. (39) Giorgetti, L.; Bongiorno, G.; Podesta`, A.; Berlanda, G.; Scopelliti, P. E.; Carbone, R.; Milani, P. Langmuir 2008, 24, 11637–11644. (40) Angenendt, P.; Glo¨kler, J.; Sobek, J.; Lehrach, H.; Cahill, D. J. J. Chromatogr., A 2003, 1009, 97–104.

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