Substrate Remote Control of Polymer Film Surface Mobility

Article pubs.acs.org/Macromolecules

Substrate Remote Control of Polymer Film Surface Mobility Igor Siretanu,†,‡ Jean Paul Chapel,† and Carlos Drummond*,† †

Centre de Recherche Paul Pascal, UPR8641, CNRS, Université de Bordeaux, Avenue Schweitzer, 33600 Pessac Cedex, France Laboratoire de Chimie des Polymères Organiques, CNRS, Université Bordeaux 1, 33607 Pessac Cedex, France

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ABSTRACT: Polymer segments at the surface of glassy polymer films remain mobile at temperatures below the glass transition temperature, Tg. This mobility, which is usually attributed to the access to a larger free volume by segments at the surface, opens pathways for polymer surface structuration by the effect of a destabilizing force. By studying the destabilization of polystyrene films under the influence of ions dissolved in degassed water at temperatures well below Tg, we have observed that this mobility can be strongly affected by a substrate buried down distances of the order of the chain size below the film surface. This effect is particularly important if there is a strong interaction between the polymer and the substrate or in the presence of pinning points for the polymer chains. These results can be qualitatively interpreted in terms of the sliding model for Tg reduction in thin polymer films. This effect allows remotely controlling the structuration of the polymer surface.

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on the solid surface.8 This only happens in the presence of certain ions if dissolved gases are removed from the aqueous phase; otherwise, no modification of the polymer film is observed. Atomic force microscopy micrographs illustrating this novel phenomenon are presented in Figure 1. The typical size of the self-assembled patterns depends on pH, temperature, and the amount of dissolved gas in the aqueous phase. This patterning process is a consequence of an electrohydrodynamic instability at the nanoscale9 due to the adsorption of the water− ions at the polymer−water interface. Regardless of its unexpected origin, this process clearly reveals the mobility of the upmost surface layer of otherwise glassy polymer films. In this work we report on how this phenomenon is modified by the molecular weight of the polymer and the polymer− substrate interaction.

INTRODUCTION A large number of studies have addressed the properties of polymer thin films and their relationship to their bulk counterparts. In particular, the glass transition of polymer thin films has been extensively investigated. Since the seminal paper by Keddie et al.1 showing the influence of film thickness on the glass transition temperature Tg was published, many studies dealing with chain mobility under confinement have been reported.2 An aspect investigated in great detail has been the influence of the boundary conditions on the glass transition. A majority of the experimental results show that attractive walls tend to reduce the mobility of the polymer and increase the glass transition temperature, Tg. On the contrary, less attractive or free boundaries enhance the mobility and hence reduce Tg. In fact, as suggested by de Gennes3 and later experimentally demonstrated by Ellison and Torkelson4 by studying the fluorescence of selectively marked PS layers at different depths, thin films’ behavior can be better understood in terms of a depth-dependent distribution of the actual polymer mobility (or Tg). The “globally” measured Tg value will then depend on the way the average is performed through the entire specimen. Other studies support the idea of a heterogeneous dynamic of thin glassy polymer films.5,6 In particular, the existence of a mobile layer several nanometers thick on top of glass-forming polymer films is gaining general acceptance in the polymer physics community:7 polymer segments near the surface are obviously less constraint than segments inside the bulk of the film. It is reasonable to expect that this excess mobility will persist if the film is in contact with a low-viscosity nonswelling media. As an example of this mobile surface layer, we have recently shown that when polystyrene films are put in contact with degassed water solutions of acidic or basic pH for a few minutes, a long-lasting nanostructuration spontaneously forms © 2011 American Chemical Society

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MATERIALS AND METHODS

Polystyrene (PS) of four different molecular weights (7, 160, 250, and 500 kDa) were investigated. PS of 500 kg/mol (Tg 103 °C) and 59 kg/ mol (Tg 99 °C) were obtained from Sigma-Aldrich. PS of 250 kg/mol (Tg 103 °C) was obtained from ACROS Organics. PS of 7 kg/mol (Tg 90 °C), synthesized by ATRP, was a gracious gift from Dr. Antoine Bousquet. Film Preparation. We prepared films of PS on four different substrates by spin-coating. First, hydrophilic surfaces of freshly cleaved molecularly smooth mica and thoroughly cleaned oxide-coated silicon wafers were used. A second set of films were prepared on surfaces of the same materials which were modified with an adhesion primer layer in order to modify the interaction of the polymer layer with the substrate. Mica surfaces were coated with a monolayer of 10 nm Received: September 28, 2011 Revised: December 5, 2011 Published: December 22, 2011 1001

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flow of nitrogen gas. In some rare cases water penetrated between the polymer film and the substrate. These tests were not further considered.

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RESULTS AND DISCUSSION As mentioned before, the exposure of the PS films to degassed acid water during few minutes provokes the formation of bumps several nanometers tall and several tens of nanometers wide. As can be observed in Figures 1 and 2, the self-assembled

Figure 1. Height AFM micrographs of films of 250 kDa polystyrene films measured in tapping mode in air. The films of indicated thickness were spun-coated on bare and silanized silicon wafers (fractional OTS surface coverage 0.7) and then exposed to a degassed solution of nitric acid in double distilled water at pH 1.5 at room temperature. A typical height profile for each image is presented. The presence of asperities of regular nanometric size is clearly observed on some of the films exposed to the degassed solution for 5 min. On the contrary, no modification was detected when identical films were exposed to the same solution under identical conditions (or much longer times) before removing the dissolved gases. All the films were featureless before treatment.

Figure 2. Average height (circles) and diameter (squares) of selfassembled bumps observed on the surface of thin polymer films of PS 250 kDa of different thickness treated with degassed water at pH 1.5. The films were deposited on (a) bare and ceria-coated mica and (b) silicon wafers and silanized silicon wafers (fractional OTS surface coverage 0.7). Closed symbols: coated substrates; open symbols: bare substrates.

surface pattern depends on the type of substrate used. The typical radii and heights of the observed bumps are similar in bare mica or oxidized Si wafers and are independent of polymer film thickness down to 10 nm. On the contrary, the characteristic size of the bumps is substantially reducedor they are not presentfor films spun on ceria-coated mica or hydrophobized wafers for polymer films thinner than a certain thickness. The same results were observed after longer times (few hours) of exposure of the polymer films to degassed water (kept out of contact with air). Although we cannot discard the possibility of a very slow evolution of the surface which would lead to surface structuration at longer times, this would not challenge the fact that there is a dramatic deceleration of the structuration process due to the presence of the primer layers. As mentioned above, the idea that polymer segments close to the surface are more mobile that bulk segments is now broadly accepted, although the properties of the mobile layer (thickness, variation with temperature, and molecular weight) are still debated. On the contrary, it is not obvious why this excess mobility should be impaired by the influence of the underlying substrate buried down several tens of nanometers below the polymer film surface. It is noteworthy (cf. Figure 2) that above certain film thickness the size of the formed bumps seems to achieve a terminal value which is independent of PS molecular weight. On the contrary, the threshold thickness for surface structuration is determined by polymer size, as can be observed in Figure 3a,b. The larger the molecular weight, the longer the range of the effect of the substrate on the structuration of the polymer layer. In all cases the observed effect has a range too largeup to 100 nm for the largest polymer investigatedto be interpreted in terms of the dispersion interaction between the solid substrate and water through the polymer film. The reason for the long-range effect of the substrate on the polymer structuration must be related to the connectivity of the polymer chains. It is remarkable that by normalizing the film thickness by the mean-square end-to-end

diameter ceria nanoparticles by dip-coating. A close-packed monolayer of ceria nanoparticles get adsorbed on mica by this procedure, which strongly increases the adhesion between the subsequently deposited polymer layer and the substrate.10 SiO2-coated silicon wafers were hydrophobized at different degrees by the Langmuir−Blodgett coating technique11 with octadecyltrichlorosilane (OTS) by varying the imposed surface pressure while film transfer. It has been reported that slipping of PS is modified in silane-modified silicon wafers and depends on silane grafting density.12 After spin-coating, the films were annealed at 95 °C for 12 h, unless indicated otherwise. They were subsequently exposed to degassed double distilled water at pH 1.5 as described below. Film Characterization. The thickness of the films was determined by ellipsometry (Nanofilm); roughness and morphology were assessed by atomic force microscopy in tapping mode in air (multimode and Icon, Veeco) before and after treatment with degassed water. Extremely smooth films, with rms roughness smaller than 0.5 nm were typically obtained after spin-coating. After degassed water treatment the generated structures were characterized by the average size and height of the observed self-assembled polymer bumps. The films were also characterized by water contact angle (IDC Concept). The contact angle of water on the films was around 90° before and after structuration, with very small hysteresis. Water Degassing and Film Treatment. Millipore water with a conductivity of 18 MΩ cm−1 was used for preparing the acid solutions. The pH of the aqueous phase was adjusted to 1.5 by adding small amounts of nitric acid (Aldrich) as necessary. Carefully cleaned Teflon bar stirrers were introduced in the solutions to be degassed to induce the nucleation of gas bubbles. The solutions were subjected to agitation under pressure of 0.2 mbar for 2 h. The appearance of macroscopic bubbles in the aqueous phase was observed only during the first 30 min of degassing. After degassing was finished, the air pressure on the flask was gently increased back to atmospheric pressure. The degassed solutions were put right away in contact with the polymer surfaces for few minutes after stopping the pumping, unless otherwise indicated. The PS films were then dried with a gentle 1002

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Figure 3. Effect of polymer molecular weight on the self-assembled structure. Average height (a) and diameter (b) of self-assembled bumps of thin polymer films deposited on silanized silicon wafers (fractional surface coverage 0.7) treated with degassed water at pH 1.5. The results for the different polymer sizes collapse on a single curve if the film thickness is normalized by the polymer end-to-end distance, REE (c, d).

distance of the polymer, REE, the results for the different molecular weights tested collapse on a single master curve, as can be observed in Figure 3c,d. This collapse indicates that a mechanism related to the polymer size is responsible for the observed drop of polymer surface mobility. By exploring different annealing conditions, we observed that the effect of the two primer layers investigated is not completely equivalent, but rather process dependent. While ceria-coated mica suppressed the pattern formation in thin films regardless of the annealing process, silanized silicon wafers were effective structuration inhibitors only after annealing of the films. Otherwise, the modification of the supported film resembled the one observed for the case of bare silicon wafers. In addition, we observed that varying the grafting density of OTS on the Si wafer has a strong influence on the inhibitor effect of the primer silane layer, as can be observed in Figure 4.

The more puzzling question emerging from the results reported in this work is how the effect of the solid substrate can extend over tens of nanometers up to the polymer surface. We can attempt to rationalize our results at a qualitative level with the help of the sliding mechanism sketched by de Gennes,3 formulated to explain the reduction of Tg in free-standing or supported polymer thin films. This model proposes that the reduced density at the polymer surface confers a greater mobility to the polymer segments: monomers at the polymer surface are less sterically constrained to move. This is probably also valid for the case of contact of the polymer film with a lowviscosity poor solvent, the scenario studied in this work. de Gennes proposed that this enhanced free volume then propagates at a certain speed through the polymer chain to depths of the order of the size of the polymer molecules, effectively fluidizing the polymer layer several nanometers inside the film. The free surface will then act as source and sink of “free volume”. The idea behind the model is that the energy barrier for kink (free volume) transport is lower than for segmental motion; polymer loops that start and end at the free surface can then be effectively fluidized, with the chains moving along their own paths. This process is not to be mistaken with the reptation of entangled polymer chains in a melt, which is obviously not possible at temperatures below Tg.13 This motion is not available for chains with only one “contact” with the free surface: sliding motion would imply entering into volume occupied by other chains. Milner and Lipson ML14 have built on this idea to develop their “delayed glassification model”, to quantitatively explain how the effect of a free surface can then fluidize a polymer layer, and have looked into the issue of how deep the influence of the free surface propagates into the sample. They labeled a certain loop (starting and ending at the surface) as “fast” when the time necessary to cover the distance between the source and the sink of free volume is shorter than a reasonable experimental time. ML studied in detail the probability for a chain segment at a given depth in the film to be part of a fast loop in order to quantify the plasticizing effect of the free surface.14 Here we are interested in the analogous problem of an otherwise fast loop being in contact or not with the supporting substrate. If the polymer−substrate interaction is sufficiently strong, the presence of the wall can slow down or stop the free volume propagation through the

Figure 4. Height tapping mode AFM micrographs measured in air of 50 nm thick 250 kDa polystyrene films after exposure to a degassed solution of nitric acid in double distilled water for 5 min. pH 1.5; T 25 °C. The films were deposited on silanized silicon wafers with OTS surface coverage indicated on each micrograph. Scale bars correspond to 200 nm. The propagation of a kink (indicated by the black arrows) is slowed by the pinning effect of the OTS layer at intermediated grafting densities, stopping the structuration of the polymer surface.

Very high or very low silane grafting densities produced no effect on the structuration: in these cases the self-assembled structures were observed on polymer films down to 10 nm thick. On the contrary, intermediate grafting densities effectively diminished the surface structuration of thin polymer films, provided that the film has been annealed before treatment with degassed acid water. 1003

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Figure 5. Characteristic relaxation times. PS relaxation at the surface (this work, blue circles) increases more slowly with decreasing temperatures than reported bulk α-relaxation (red squares, ref 16) or extrapolated values from polymer behavior at temperatures above Tg (black line, ref 17). Inset: time dependence of the of bump height at selected temperatures and typical 1 × 1 μm2 AFM height micrographs measured in air before and after relaxation of the structure.

longest fast loop allowed, which they modeled by using a temperature dependence of the WLF form15 by extrapolating from the behavior of bulk PS above Tg. However, it has been experimentally shown that the dynamic of near-surface polymer chains in glassy PS is substantially faster than the one predicted by extrapolation from the behavior of liquid PS or the one measured in bulk glassy PS.7 To further investigate this point, we have conducted an extensive study of PS surface relaxation using the structuration technique discussed in this work. As we have reported in a previous publication,8 the bump size decreases with time (and eventually disappears) following a thermally activated relaxation process well described by singleexponential time decay, as illustrated in the inset of Figure 5. The measured characteristic relaxation times of the surface modification indicate that the dynamic of the polymer surface slows down with decreasing temperature at a lesser pace than the polymer bulk in a glassy state. Similarly, it is clear that extrapolating the molten state behavior to temperatures below Tgthe method chosen by MLgrossly underestimates the mobility of polymer chains close to the surface of the film. The characteristic relaxation times of the polymer near to the surface would be between the ones measured at the surface and at the bulk of the film, both much shorter than the values extrapolated from the behavior at T > Tg. Significantly longer fast loops close to the surface than the ones considered in ref 14 should then be anticipated. A second reason can be evoked to explain the long spatial range of fast loops mentioned above: the cooling of the polymer films after annealing might have been carried out too quickly (in terms of the relaxation time of the system). A very long time may be necessary to reach thermodynamic equilibrium of the films at room temperature. Upon cooling, the polymer chains might effectively be frozen in conformations that would correspond to a higher temperature (an effective or “structural” temperature). Indeed, it has been recently reported that the relaxation times of spin-coated polymer films,

chain, effectively freezing a given loop. If this probability is negligible, the polymer chains at the free surface should not be affected by the presence of the wall. It is obvious that for polymer films substantially thicker than REE the effect of the free surface will not be modified by the presence of the solid boundary. On the contrary, for films thinner than REE the presence of the boundary might affect the propagation of a kink along a polymer chain, as sketched in Figure 4. The penetrating effect of the substrate will be less pronounced in samples of lower molecular weight. For the two substrate modification methods discussed in this work (ceria-coated mica and silanized Si wafer) the friction coefficient for the monomers in contact with the modified surface is higher than with the bare substrate, slowing down the polymers movement near the substrate and reducing the structuration. The origin of the slow dynamic seems to be different for each case: for the ceria-coated mica it is due to the enhanced dispersion interaction between polymer chains and the ceria nanoparticles. On the contrary, for the silanized wafers the enhanced friction is probably due to the interpenetration between the polymer molecules and the silane layer, which is enhanced upon annealing, but it is more difficult at very high OTS coverage (or absent at low coverage). In that sense, the defects on the OTS layer acts as “pinning points” that avoid the sliding of the polymer chains. Even though both effects are short-ranged (a fraction of a nanometer) their influence propagates to the polymer surface due to the longrange effects of chain connectivity. Although conceptually appealing, the sliding model of de Gennes in the ML formulation14 predicts a probability of polymer segments belonging to a fast loop that decays very quickly with depth at room temperature. If that is true, polymer films thicker than a few nanometers should not be sensitive to the presence of the solid boundary, which is clearly inconsistent with our results. At least two reasons can be evoked to rationalize this discrepancy. First, it is possible that the ML model overestimates the temperature decay of the length of the 1004

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(10) Chapel, J.-P.; Morvan, M. PCT Int. Appl. WO 2007126925 A2 20071108, 2007. (11) Peterson, I. R. J. Phys. D 1990, 23, 379−395. (12) Xu, L.; Sharma, A.; Joo, S. W. Macromolecules 2010, 43, 7759− 7762. (13) de Gennes, P. G. Scaling Concepts in Polymer Physics; Cornell University Press: Ithaca, NY, 1979. (14) Milner, S. T.; Lipson, J. E. G. Macromolecules 2010, 43, 9865− 9873. Lipson, J. E. G.; Milner, S. T. Macromolecules 2010, 43, 9874− 9880. (15) Williams, M. L.; Landel, R. F.; Ferry, J. D. J. Am. Chem. Soc. 1955, 77, 3701−3707. (16) Dhinojwala, A.; Wong, G. K.; Torkelson, J. M. J. Chem. Phys. 1994, 100, 6046−6054. (17) Ferry, J. D. Viscoelastic Properties of Polymers, 3rd ed.; John Wiley and Sons: New York, 1980. (18) Raegen, A.; Chowdhury, M.; Calers, C.; Schmatulla, A.; Steiner, U.; Reiter, G. Phys. Rev. Lett. 2010, 105, 227801. (19) Reiter, G.; Hamieh, M.; Damman, P.; Sclavons, S.; Gabriele, S.; Vilmin, T.; Raphaël, E. Nature Mater. 2005, 4, 754−758. (20) Barbero, D. R.; Steiner, U. Phys. Rev. Lett. 2009, 102, 248303. (21) Thomas, K. R.; Chenneviere, A.; Reiter, G.; Steiner, U. Phys. Rev. E 2011, 83, 021804.

measured from dewetting experiments, strongly depends on the processing history of the sample.18,19 It has also been reported that an extremely long time (of the order of weeks) may be necessary to reach final equilibrium in thin viscous polymer films, even at temperatures above Tg.20,21 This could also explain the weak temperature dependence of the relaxation at low temperatures (Figure 5). A more detailed report of thermal aspects related to the process described here will be published elsewhere. The effect described in this work can be used to direct the structuration of thin polymer films by controlled patterning of the supporting substrate. As an example, Figure 6 shows how a

Figure 6. Height tapping mode AFM micrograph measured in air of 50 nm thick 250 kDa polystyrene films after exposure to a degassed solution of nitric acid in double distilled water for 5 min. pH 1.5; T 25 °C. The film was spin-coated on a half-silanized silicon wafer. Scale bar corresponds to 2 μm. The presence of asperities of regular nanometric size bumps is clearly observed on the bare silicon wafer. On the contrary, no modification was detected in the region which was coated with OTS.

polymer film coating a half-silanized silicon wafer gets partially structured after treated with degassed water at pH 1.5. Only the region of the film above the bare silicon wafer gets structured; more complicated surface morphologies can be conceived by following this strategy. In conclusion, we have shown that the structuration of the surface of glassy PS films by contact with degassed water can be affected by the presence of a solid boundary (up to film thickness of the order of REE) which may slow down or even prevent the polymer film restructuration if there is a strong interaction between the polymer chains and the substrate.

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REFERENCES

(1) Keddie, J.; Jones, R.; Cory, R. Europhys. Lett. 1994, 27, 59−64. (2) Alcoutlabi, M.; McKenna, G. B. J. Phys.: Condens. Matter 2005, 17, R461−524. (3) de Gennes, P. G. C. R. Acad. Sci. Paris, Ser. IV 2000, 1, 1179− 1186; Eur. Phys. J. E 2000, 2, 201−203. (4) Ellison, C. J.; Torkelson, J. M. Nature Mater. 2003, 2, 695−700. (5) Rotella, C.; Napolitano, S.; De Cremer, L.; Koeckelberghs, G.; Wübbenhorst, M. Macromolecules 2010, 43, 8686−8691. (6) Inoue, R.; Kawashima, K.; Matsui, K.; Kanaya, T.; Nishida, K.; Matsuba, G.; Hino, M. Phys. Rev. E 2011, 83, 02180. (7) Fakhraai, Z.; Forrest, J. A. Science 2008, 319, 600−604. (8) Siretanu, I.; Chapel, J.-P.; Drummond, C. ACS Nano 2011, 5, 2939−2947. (9) Schaffer, E.; Thurn-Albrecht, T.; Russell, T. P.; Steiner, U. Nature 2000, 403, 874−877. 1005

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