Polymer Nanogels Grafted from Nanopatterned Surfaces Studied by

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Polymer Nanogels Grafted from Nanopatterned Surfaces Studied by AFM Force Spectroscopy Soyeun Park† and Wolfgang Frey*,‡ † ‡

Department of Physics, Texas Tech University, Lubbock, Texas 79409, United States Department of Biomedical Engineering, The University of Texas at Austin, Austin, Texas 78712, United States

bS Supporting Information ABSTRACT: Nanopatterned cross-linked polymers are important for applications with controlled mechanical properties. Grafted linear and cross-linked polydimethylacrylamide gels on micro- and nanopatterns were created using iniferter-driven quasi-living radical polymerization combined with conventional photolithography and nanosphere lithography. Micropatterned linear polymers reproduce the expected scaling behavior at moderate grafting density. The addition of cross-linker to the polymerization solution leads to an increased tendency of early termination as determined by AFM force spectroscopy. Similarly, nanopatterned linear polymers show reduced thickness in agreement with the expected scaling relationship for nanoisland grafts that have reduced lateral confinement. The addition of cross-linker reintroduces some of the lateral confinement for the length of polymers reported here. The mechanical properties of both the micro- and nanopatterned linear as well as cross-linked polymers were analyzed using an algorithm to objectively determine the contact point in AFM force spectroscopy and two independent Hertz-based analysis approaches. The obtained Young’s moduli are close to those expected for homogeneous thick polymer films and are independent of pattern size. Our results demonstrate that polymeric nanopillars with controlled elastic modulus can be fabricated using irreversible cross-linkers. They also highlight some of the factors that must be considered for successful fabrication of grafted nanopillars of defined mechanical and structural properties.

’ INTRODUCTION Surface polymer films attract considerable attention due to their use in a variety of applications, from friction control to biomaterial and sensor surface modification. They are also interesting for their unique interfacial properties.1
3 Robust surface anchoring and a high packing density are achieved by the grafting-from approach using surface-initiated polymerization, in which polymer chains are grown from the surface-anchored initiators.4
6 A controlled living radical polymerization such as the atom transfer radical polymerization or the quasi-living iniferter-driven photoactivated polymerization achieves a well-controlled polymer growth with defined chain lengths and low polydispersity; such polymerization techniques also open the possibility for the synthesis of block copolymers.7
12 Because of the ability to observe and exploit surface-dominated phenomena at the microand nanoscales, polymer grafts at those scales are of growing interest.13,14 Nanopatterned polymer grafts are widely applicable in areas such as nanoactuators, biosensors, proteomic chips, nanofluidic devices, and engineered tissues.1,13,15
17 Patterned polymer grafts have been fabricated using nanolithographic and template approaches, such as dip-pen lithography, microcontact printing, photolithography, and electron-beam lithography.12,18
26 Distinct morphological and thermodynamic properties of nanopatterned polymer grafts have been investigated both theoretically and experimentally by varying the grafting density, the r 2011 American Chemical Society

polymer length, and the lateral dimensions of the initiator patterns. These properties have been shown to contribute to a broadening of the collapse transition of stimuli-responsive polymer chains.18,20
22,27
29 Despite recent advances, it remains a challenge to fabricate nanopatterned polymer grafts with controlled mechanical properties. Various approaches such as the block copolymers and polymer
metal hybrid structures were attempted to obtain controlled mechanical properties of micro- and nanopatterned polymer grafts.30
32 While the mechanical properties of bulk polymer gels are normally controlled by the cross-linking density, it is expected that the addition of cross-linkers to surface-grafted polymers modulates not only their mechanical properties but also permeability and swelling behaviors.32,33 One of the problems in achieving well-defined mechanical behavior of surfacegrafted polymers is the difficulty in characterizing their mechanical properties. Atomic force microscopy (AFM), surface force apparatus,34 scanning electron microscopy, near-field scanning optical microscopy,35 secondary ion mass spectrometry imaging,36
38 and localized surface plasmon resonance measurements39 have been used to obtain structural, conformational, and chemical Received: March 31, 2011 Revised: May 23, 2011 Published: June 16, 2011 8956

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Langmuir information on surface-grafted polymers. AFM has been the most widely used method to characterize mechanical properties of micro- and nanopatterned polymer grafts. Two limitations, however, hinder the mechanical analysis of thin surface polymer films by AFM. First, when using the conventional contact and tapping modes, the AFM tip significantly influences the thickness and lateral dimensions of nanopatterned polymer grafts, as has been demonstrated both experimentally21,22,28 and numerically.28,29 Even minimizing the tip
sample interaction18,22 in an effort to obtain the unperturbed topography of polymer grafts resulted in distorted topographic images along the fast scan direction and reflected insufficient tracking of the background.40 In contrast, a molecular dynamics study of the interaction between patterned polymer grafts and an AFM tip demonstrated that the unperturbed configuration of the nanopatterned polymer chains can be reconstructed from the force
distance (f
d) curves at nearzero penetration.28 Therefore, by extrapolating zero-force contact, two-dimensional (2D) AFM force spectroscopy has been used successfully to reconstruct the unperturbed true topography of soft biological samples, regardless of the tip
sample interactions.40,41 Additionally, 2D AFM force spectroscopy has been used as an analytical tool to characterize adhesiveness and wettability42 as well as polymer chain lengths and grafting densities.43
45 Second, there is some controversy over whether AFM force spectroscopy can be used to obtain mechanical properties, such as the Young’s modulus, from ultrathin films. Although the Hertz model is widely used to determine the Young’s modulus of thicker samples (>500 nm),46
51 the strong influence of the substrate for thin soft samples violates core assumptions of the model. However, several semiempirical models based on the Hertz model have been developed to convert 2D AFM force spectroscopic information into quantitative mechanical information for very thin films on solid surfaces.52
55 One recent model is based on the asymptotic analysis of surface polymer films and has correctly predicted the mechanical properties of dry surface polymer films.55 Here we report on the influence of an irreversible cross-linker on the topographic and mechanical properties of nanopatterned polymers grown by surface-initiated iniferter polymerization. Micropatterned polymer grafts were used to extract structural information such as approximations for the grafting density and contour length. The 2D AFM force spectroscopy was used to accurately determine the topographic properties of the nanopatterned polymer grafts. Two independent approaches, both based on the Hertz model, were adapted to extract the Young’s modulus of the micro- and nanopatterned polymer.

’ MATERIALS AND METHODS Materials. Deionized water with a resistivity of at least 18.0 MΩ (EPure, Barnstead Thermolyne, Dubuque, IA) was used in all experiments. Plain microscope glass slides (Erie Scientific Co., Portsmouth, NH) and polystyrene nanospheres of 400 nm diameter (Interfacial Dynamics Corp., Portland, OR) were used. Tungsten boats for thermal evaporation, carbon crucibles and silicon dioxide (SiO2) pieces for electron-beam evaporation were purchased from Kurt J. Lesker Co. (Clairton, PA). Gold shots (Au, Alfa Aesar, Ward Hill, MA), chromium pieces (Cr, R. D. Mathis, Long Beach, CA), and p-chloromethylphenyltrimethoxysilane (CMPTMS, Gelest, Morrisville, PA) were also used. Sodium N,N-diethyldithiocarbamatetrihydrate, 1-hexadecanethiol (HDT), N,N-dimethylacrylamide (DMA), N,N0 -methylenebis(acrylamide) (BA),

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Figure 1. (a) Schematic illustration of the fabrication steps for surfaceinitiated polymerization from nanoislands. (i) Creation of a mono- or bilayer of nanospheres on an Au thin film. (ii) SiO2 deposition and nanosphere removal to obtain a 2D array of SiO2 nanoislands. (iii) Protection of the exposed Au-areas by a 1-HDT monolayer and selective anchoring of the DC-iniferter-silane (dithiocarbamate-CMPTMS) on the SiO2 nanoislands. (iv) Surface-initiated quasi-living radical polymerization from the SiO2 nanoislands by illumination with collimated UV light in the presence of a highly concentrated monomer solution. (b) Schematic depiction of the polymerization setup. Quasi-living polymerization was carried out by UV illumination under a nitrogen atmosphere. A thin layer of the concentrated monomer solution with various amounts of cross-linker was placed above the DC-glass. Micron-sized stripe patterns were transferred using a photomask aligner. 1,4-dioxane, and toluene were purchased from Sigma-Aldrich (Milwaukee, WI); methanol, ethanol, sulfuric acid, hydrogen peroxide, and ammonium hydroxide were purchased from Fisher Scientific (Pittsburgh, PA). All chemicals were used as received. Preparation of Nanopatterned Templates. Nanopatterned substrates for surface-initiated polymerization were created using a modified nanosphere lithography (NSL) technique. Thin gold films were prepared by thermally evaporating 2 nm of Cr followed by 50 nm of Au on a piranha (H2SO4:H2O2 (3:1))-cleaned microscope glass slide (Denton Vacuum, Moorestown, NJ). The Au-covered surface was cleaned in TL1 solution (NH4OH:H2O2:H2O (1:1:5)) for 10 min at 75
C. NSL was used to create the desired nanopatterns on Au surfaces (Figure 1a-i).56 Briefly, polystyrene nanospheres were used to form either a monolayer or bilayer of densely packed spheres on the Au surface using a custom-built capillary deposition machine. Films of 2 nm of Cr followed by 25 nm of SiO2 were deposited on the surface using an electron-beam evaporator (Edwards Auto500, Crawley, UK). The nanospheres were then removed by sonication in methanol (Aquasonic Model-75D, VWR, West Chester, PA). The resulting nanopatterns of 2 nm Cr and 25 nm of SiO2 served as templates for the selective polymerization from the nanoislands (Figure 1a-ii). Preparation of DC-Glass. Following the protocol originally developed by Matsuda et al.,10 dithiocarbamate-derivatized glass surfaces (DC-glass) were used to initiate polymerization. For the selective 8957

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functionalization of the SiO2 nanoislands, the Au surface of the prepared nanopatterned template was masked with an inert thiol layer prior to the CMPTMS immobilization; the nanopatterned templates were incubated in a 5 mM ethanolic solution of 1-HDT for 24 h, rinsed in ethanol, and dried with dry nitrogen. CMPTMS was vapor-deposited onto the 1-HDT-treated nanopatterned templates for 8 h, then rinsed with toluene and ethanol, and annealed at 100
C for 1 h (Figure 1a-iii). Silanized glass slides were incubated in a 5% (w:v) ethanolic solution of sodium N,N-diethyldithiocarbamate trihydrate for 24 h protected from light. Samples were thoroughly rinsed with ethanol and DI water, dried with dry nitrogen, and stored in a dark desiccator until used. Micropatterned polymers were synthesized from a homogeneous DC-glass without nanopatterns using near-UV lithography. Surface-Initiated Polymerization. Polymerization was performed from the surface-immobilized initiator (dithiocarbamate) immersed in the 5 M dimethylacrylamide monomer in dioxane solution with varying amounts of cross-linker present during the polymerization under nitrogen protection in a custom-made chamber (Figure 1b). A 5 M DMA solution in dioxane with various concentrations of BA crosslinker was freeze
thaw cycled three times to remove oxygen. 1 mL of the monomer solution was dropped onto the DC-glass placed in the custom-built chamber filled with dry nitrogen. A photolithography mask was placed on top of the monomer solution to transfer the micropattern, while a quartz slide was used to seal the chamber and collimated near-UV light from a mask aligner (USH-205DP, 200 W, Ushio, Tokyo, Japan) irradiated the DC-glass at room temperature. After polymerization, the surface was sonicated in dioxane, rinsed in ethanol and DI water, and then dried with nitrogen. Atomic Force Microscopy. All AFM experiments were performed with an MFP 3D (Asylum Research, Santa Barbara, CA). Topographic images were obtained in contact mode for all samples under either ambient or aqueous conditions using silicon nitride probes (Microlevers, Veeco Metrology, Santa Barbara, CA). Quasi-contact and zero-force thickness, fixed force distances, maximum rupture forces, and maximum rupture distances and all data for the elasticity analysis for all micro- and nanopatterned polymer grafts were extracted from approach and retraction force
distance (f
d) curves (see Supporting Information for details). All f
d curves were obtained using a nonfunctionalized silicon nitride tip in aqueous environment. Prior to obtaining f
d curves, the cantilever force constant was calibrated using the thermal noise calibration tool in the MFP 3D software. Nanopatterned polymers were imaged with a 2D array (40  40 or 30  30 curves) of f
d curves over an area of 1  1 μm2, using a custom script and automatic tip retraction at a given deflection. Polymer thickness was determined with contact mode in the dry hd and hydrated state hw and compared to the zero-force polymer thickness hf according to hf ¼ jzcmax
z0 j

ð1Þ

Here, the sample-to-tip contact point z0 was defined as the zero-force point determined using a fitting routine as described in the Supporting Information. The maximum compression point zcmax is defined as the extrapolated contact point between the tip and the underlying substrate, coinciding with the maximum measured indentation. Assuming that the polymer allows for the tip to reach close to the anchoring surface, which is a good assumption for the thin films studied here, the zero-force polymer thickness hf therefore automatically does not contain any topographic information on the underlying substrate. Rupture analysis was performed on events defined as a snap-back in the retraction f
d curve when the cantilever deflection abruptly changes from a negative value to zero. We consider only rupture events where the rupture force is greater than 5 times the rms noise of the cantilever without contact to the sample (see Supporting Information for details).

Figure 2. Topography of PDMA synthesized by 15 min UV illumination through a stripe mask. (a) Phase contrast image. (b) Contact AFM topographic images in the dry (top) and hydrated (bottom) states. Line profile taken from the same scans (dry: solid line; hydrated: dotted line). (c) Thickness of linear PDMA grafted from micropatterns as a function of polymerization time in the dry state hd (light gray) and in the hydrated state hw (dark gray) measured with contact AFM and with zero-force hf (black).

Elastic Moduli. The elastic modulus, i.e. the Young’s modulus E, of polymer grafts was derived from force
indentation (f
δ) curves using the objectively determined contact point z0 and two different fitting approaches that avoid the shortcomings of the simple Hertz model for thin layers on solid supports (see Supporting Information for details). Briefly, for the first approach, a sequence of fits to the f
δ curve with an increasing indentation range Δδ was performed. The result is an elastic modulus that remains relatively stable for the initial fits, until E dramatically increases at larger indentations, indicating the violation of the Hertz model assumptions and the dominating influence of the substrate. The Hertz model for a parabolic tip indenting the thin film gives the indentation force as57 f ¼

4 E pffiffiffiffiffiffiffi3ffi Rδ 3 1
ν2

ð2Þ

Here, f is the indentation force, ν is the Poisson ratio (assumed to be 0.5), R is the radius of curvature of the tip, and δ is the indentation. The radius of the cantilever tip (44 ( 5 nm) was determined by scanning electron microscopy (Hitachi S-4300 SE/N at 2 kV, Hitachi High Technologies America, Inc.) from a sample of tips taken from the same batch of tips that was used for the experiments and agrees with the manufacturer’s specifications (40 nm) to within 10%. The cantilever spring constant (∼0.03 N/m) was determined prior to the experiment as described above. The second approach is based on the semiempirical model by Cappella,55,58 where we made the assumption that the modulus of tip and substrate is so large compared to the polymer that it can be neglected (see Supporting Information for details).

’ RESULTS Micropatterned Linear PDMA Grafts. Linear PDMA chains were grafted from dithiocarbamate (DC)-modified glass by iniferter driven quasi-living polymerization.10 Figure 2 shows a stripe pattern of linear PDMA with 33 μm wide lines separated by 8958

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Figure 3. Influence of the cross-linker concentration (0
1.0% (w:w)) on micropatterned PDMA-BA grafts synthesized for 10 min. Typical approach (a) and retraction (b) force f vs scanner displacements z curves. Also shown in (a) is the scanner displacement corresponding to a fixed deflection force of 0.1 nN for each curve. (c) Thickness of polymer-graft films of varying BA concentration as measured with contact scanning in the dry hd and hydrated state hw and with f
d analysis hf in the hydrated state. Also shown are the percentage of the retraction curves that show rupture events. More than 300 f
d curves were taken for each sample.

28 μm gaps created from a UV-photo mask with 40 μm lines and 20 μm gaps. A clear difference can be seen in phase contrast between the growth in UV-illuminated areas and the bare surface of nonilluminated areas (Figure 2a). The contact AFM topographic image confirms the selective growth of polymer chains to the illuminated areas (Figure 2b). PDMA swells in water, and consequently, in the UV-illuminated areas the thickness of the grafted polymer increases strongly relative to dry conditions, while no swelling is observed in the non-UV-illuminated areas (Figure 2b,c). The uniform thickness across a stripe of polymer chains indicates that the lack of lateral confinement at the edges of the stripe did not have a significant effect on the polymer configuration inside the stripe (see line profile in Figure 2b). As expected, the hydrated thickness of linear PDMA chains gained from contact AFM topographic images hw was slightly lower than the zero-force thickness hf derived from the f
d analysis. Both the dry and hydrated thicknesses of micropatterned PDMA increase with polymerization time (Figure 2c). This increase is linear for the dry thickness, indicating a relatively high grafting density. However, for the hydrated thickness a linear increase of the thickness with polymerization time is found with a nonzero y-intercept. This suggests that the chains have some lateral freedom. Although visibly thick (∼micrometer) polymer patterns could be synthesized with longer exposure times, only relatively short exposure times (5
15 min) were investigated here. Micropatterned Cross-Linked PDMA-BA Grafts. Bis(acrylamide)-cross-linked PDMA (PDMA-BA) was grafted from micropatterns with iniferter polymerization in the presence of varying concentrations of the cross-linker (bis(acrylamide) (BA)) in the monomer solution. Cross-linking the polymer leads to decreasing film thickness with increasing BA concentration as seen both in contact imaging hd, hw, and from the f
d analysis hf (Figure 3c). The absolute decrease in the dry films is less pronounced than in the hydrated films, but the relative change is similar to the change in wet conditions. Typical approach f
d curves for these micropatterned gels with varying degrees of cross-linking show qualitatively an increase in the elastic response as indicated by the fixed force distances in Figure 3a. Upon retraction (Figure 3b), the number of rupture events observed increases and reaches 100% of the retractions for 0.75% (w:w) (PDMA-BA075) concentration (Figure 3c), probably caused by a higher degree of entanglement of the tip with the polymer chains for higher cross-linker concentrations. In the masked, nonilluminated areas, only a linear force increase in the

approach and no rupture event in the retraction curves are seen (not shown), indicating a polymer-free surface. F
d Rupture Event Analysis. The number of measurable rupture events shows a maximum for PDMA-BA075 and slightly decreases in the highest cross-linker concentration (1.0% (w:w) PDMA-BA100) (Figure 3c). The same biphasic dependence on the cross-linker concentration was found for the mean of the maximum rupture force (Figure 4). The micropatterned PDMABA075 shows significantly higher rupture forces (230 ( 150 pN) than any other sample with 44 ( 16 pN for pure PDMA, 59 ( 73 pN for 0.5% (w:w) (PDMA-BA050), and 146 ( 191 pN for PDMA-BA100. The maximum rupture distance shows a similar biphasic behavior as the maximum rupture force and the number of rupture events, but the dependence on the cross-linker concentration is more complex (Figure 5). The most likely maximum rupture distance is with 428 nm again the greatest for PDMA-BA075, while PDMA-BA050 and PDMA-BA100 show about half that value (179 and 189 nm); the linear polymer grafts show too few rupture events to allow for an analysis. A close look, however, reveals a bimodal distribution of the maximum rupture distance for each of the cross-linked polymer grafts. To deconvolute the two modes, a two-Gaussian fit of the histograms was performed. As the cross-linker concentration increases, the higher mode is maintained between 650 and 800 nm, while the lower mode depends more strongly on the cross-linker concentration and mirrors the biphasic dependence of other features in the rupture events. Interestingly, the lower mode of PDMA-BA100 is smaller than for PDMA-BA050 and the higher mode is slightly bigger, thereby separating the two modes completely. Furthermore, the area under the peak in the lower mode of rupture distances increases with the cross-linker concentration; these areas are, relative to the PDMA-BA100 lower mode area, 37% for PDMABA050 and 86% for PDMA-BA075. These two modes suggest a structural difference in the gel composition. The higher mode most likely stems from weakly cross-linked polymer chains that are able to extend easily, while the lower mode suggests a more strongly cross-linked core that is confined closer to the substrate. Although the rupture events of the linear polymer chains are often too weak to be distinguished from the noise of f
d curves, the longest maximum rupture distances measured were quite consistent—about 1000 nm for 10 min polymerization time— regardless of the cross-linker concentration; they were 909, 1008, 919, and 961 nm for PDMA, PDMA-BA050, PDMA-BA075, and PDMA-BA100, respectively. This result reaffirms that the overall 8959

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Figure 4. Histograms of maximum rupture forces determined for the micropatterned linear PDMA (a) and cross-linked PDMA-BA050 (b), PDMABA075 (c), and PDMA-BA100 (d). The maximum rupture forces increase as the cross-linker concentration increases with a biphasic signature having a maximum for PDMA-BA075. Notice that the maximum rupture force of the linear PDMA is no bigger than 120 pN while the cross-linked PDMA-BA displays significantly bigger rupture forces up to 1800 pN.

kinetics of polymer-chain growth was not affected by the addition of cross-linker but rather was regulated by the UV-light exposure time; however, the number of early termination increases with cross-linker concentration, leading to the decreased overall polymer film thickness seen in Figure 3c. Linear and Cross-Linked PDMA on Nanopatterns. Following the bottom-up fabrication strategy illustrated in Figure 1a, polymerization was selectively initiated on the 2D arrays of nanoislands created by NSL in a “grafting from” process. Figure 6 shows representative examples of nanopatterns before (Figure 6a,b) and after polymerization (Figure 6c
f). In dry conditions, nanopatterns show added thickness after polymerization but this increase of less than 10 nm is much smaller than on the micropattern, and the thickness is clearly larger on the pattern “defect lines” due to their larger area. The dry PDMA thickness on the nanoislands was 8 ( 1 and 19 ( 4 nm for 10 min (Figure 6c) and 15 min (Figure 6d) polymerization, respectively, and 5 ( 4 nm on a bilayer after 15 min polymerization (Figure 6e). Similarly, only a thickness of 3 ( 2 nm was found for PDMA-BA075 (Figure 6f) after 15 min polymerization, while on pattern “defect lines” the thickness was clearly larger due to their larger area: 22 ( 4 nm (Figure 6c), 41 ( 12 nm (Figure 6d), 37 ( 8 nm (Figure 6e), and 23 ( 5 nm (Figure 6f). A strong decrease in PDMA thickness grafted from nanopatterns is expected as the nanopattern does not provide the lateral confinement compared to a micropattern due to the dominance of the edges.29 Accordingly, the edges of the patterns with grafted polymer become more rounded and the patterns appear larger in

area by about 40% than the original pattern. The difference in the appearance of the nanoisland-grafted PDMA between a mostly confined pattern at 10 min (Figure 6c) and a strong overgrowth at 15 min polymerization (Figure 6d) is most likely a result of chain entanglement leading to a higher resistance to the scanning AFM tip, considering a polymer contour length of ∼1 μm at 10 min polymerization, as shown above for the micropattern. Somewhat surprisingly, a bilayer pattern shows good apparent confinement to the nanoislands even at 15 min polymerization, possibly due to the larger pattern separation (434 ( 8 nm in a bilayer compared to 227 ( 13 nm in a monolayer). Also unexpectedly, despite tries with much longer polymerization times, PDMA-BA100 did not yield a detectable thickness in contact AFM in the dry state. In order to verify whether the addition of BA cross-linker to PDMA has a more pronounced influence on nanopatterned than on micropatterned grafts, two-dimensional (2D) f
d mapping in wet conditions was used to obtain images of the quasi-contact topography hw (Figure 7a
d), the zero-force topography hf (Figure 7e
h), and the fixed force distance (Figure 7i
l). The quasi-contact AFM topographic images were obtained using the return deflection set point as a measure. Since the measured topography represents the combined thickness of the mostly compressed polymer chains and the nanoislands, the thicknesses of the polymers can be determined, and these thicknesses agree with the numbers found from contact AFM within the experimental error. In contrast to the contact (Figure 6) and quasi-contact (Figure 7a
d) topography, the zero-force thickness hf maps 8960

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Figure 5. Histograms of the maximum rupture distance for the micropatterned linear PDMA (a) and cross-linked PDMA-BA050 (b), PDMA-BA075 (c), and PDMA-BA100 (d). The solid line in (b)
(d) is a two-peak Gaussian fit to each histogram with high and low modes marked with arrows.

(Figure 7e
h) revealed the topography of soft gel pillars. A considerably larger thickness was measured for the hydrated polymer compared to the dry or quasi-contact AFM experiments in all cases (Figure 7e
h). The thickness of hydrated PDMA and PDMA-BA nanoisland grafts was around 70% of that on the micropattern for all cross-linker concentrations, with a slight decrease with increasing BA concentration (see Figure S4 in the Supporting Information). For comparison, we used nonpatterned polymer films, which, due to the lack of a mask, have slightly higher thickness than shown in Figure 3c (see also Figure S4 in the Supporting Information); this more closely resembles the situation of the nanopatterns that also do not have an optical mask. However, the polymer thickness in-between the patterns in all cases is not zero as expected from the contact AFM images (Figure 6). The thickness is about 50% of the pattern height for PDMA but decreases with cross-linker concentration, reaching only ∼20% for PDMA-BA075, thereby indicating some increased confinement achieved by the cross-linking. PDMABA100, however, is very different because the height differences between the pattern and the background nearly disappear. While this observation is consistent with the contact AFM measurements (Figure 6) and also the AFM rupture experiments on the micropattern (Figures 4 and 5), it is not quite clear why the distribution is so homogeneous for PDMA-BA100. Upon closer analysis, excluding PDMA-BA100, the influence from the edges of the nanopatterns is visible as a ring of greater thickness for both PDMA and PDMA-BA and clearly outlining the NSL nanopattern of somewhat depressed thickness. Typically, the thickness of the inner nanoislands is about 15% less

than the ring surrounding the nanoislands. Most likely, this ring is due to polymers grafted to the sides of the ∼27 nm high SiO2 patterns. The ring coincides with a similar ring in the “fixed force” maps (Figure 7i
l). The image contrast in the “fixed force” maps reflects a qualitative measure of the elastic response, which is confirmed by a quantitative elasticity analysis below.

’ DISCUSSION Using optical lithography and nanolithography, we created surface-grafted polymer chains and gels by iniferter-driven grafting-from polymerization. The choice of the nanopattern technique was for ease of fabrication—pattern defects had little effect to the purpose of the investigation here—but other nanopattern techniques such as step-and-flash nanolithography could be used in a completely analogous way.59 Contact-AFM topographic images and force mapping confirmed the growth of linear and cross-linked polymer chains on micro- and nanopatterned surfaces. Several theoretical and experimental studies have shown that surface-grafted linear polymer chains can organize in a mushroom-like or a brush-like conformation depending on the grafting density.2,60,61 Surface-initiated living-radical polymerization helps to achieve a high grafting density, in which the polymer chains are confined to a mostly extended brush-like conformation. At high grafting density, the polymer thickness increases with the power of 1/3 with the grafting density. Assuming a simple approximation for the polymer behavior,3,62,63 the film thickness in a good solvent increases linearly with the polymerization time N in the dense brush region but only with N1/2 in the mushroom region where little or no confinement exists. 8961

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Figure 6. AFM topographic images obtained in dry conditions from polymers on nanopatterns before (a, b) and after (c
f) polymerization. Nanopatterns were created from monolayer (a, c, d, f) and bilayer (b, e) NSL patterns. Linear PDMAs were synthesized by 10 min (c) and 15 min (d, e) polymerization. PDMA-BA075 was synthesized for 15 min (f). The scale bar represents 2 μm. The vertical scale bar applies to all images.

We found a linear dependence of the dry and wet film thickness on the polymerization time for the micropatterns (Figure 2c), suggesting that the polymer chains were in the brush configuration, although the fact that the y-intercept of the thickness vs polymerization time for the hydrated thickness is not zero suggests that this most likely is a diluted brush configuration. This conclusion is supported by the fact that the hydrated thickness is much larger than the radius of gyration of an ideal polymer, which we have deduced from the maximum contour length observed in AFM rupture experiments. Assuming a monomer backbone length of 0.25 nm, a contour length of 1000 nm obtained from the 10 min polymerization leads to a radius of gyration RG ≈ 6.5 nm. Given the molecular weight MW1 of 99 g/mol and the density F of 962 g/L for DMA, the dry thickness hd suggests a grafting density σ = hdFNA/MWn of about 0.04 chains/nm2, independent of cross-linker concentration. Here NA is Avogadro’s constant. Such a grafting density would fall between 1/(πRG2) ≈ 0.008 chains/nm2 for the mushroom regime and ∼1 chains/nm2 for a fully stretched brush regime. The measured hydrated thickness relative to the estimated contour length of about 0.2 at a dimensionless grafting density of σ* ≈ 0.027 agrees well with the scaling relationship of the relative thickness versus dimensionless grafting density σ* developed by Yamamoto et al.,4 given an approximate monomer volume of DMA of 0.25 nm  0.68 nm2. For the nanopatterns, such an estimate is more difficult to make. In contrast to the micropatterned polymer grafts, the dry nanopatterned polymers, under otherwise identical polymerization conditions, were significantly thinner. That result is somewhat surprising even considering the lack of lateral confinement.18,21,29 For the analysis of the nanopatterned polymer grafts, the AFM force spectroscopic imaging and analysis was therefore

essential and provided a very different picture. The hydrated films of PDMA and PDMA-BA are thinner than the micropatterned polymers, but only to a degree consistent with the size of the nanopattern. For the nanopatterns, the polymer film thickness is decreased by an additional ∼30% relative to the micropattern thickness. We therefore believe that the grafting density on the nanopattern and micropattern is similar. This is supported by a comparison to the scaling relationship demonstrated numerically and experimentally.21,29 Using a Kuhn length of ∼0.6 nm for acrylamide,64 dimensionless parameters introduced by Patra and Linse29 can be calculated. These are σ = 0.014, N = 1667, Δ/N = 0.11, hNP/(Nσ1/3) = 0.64, and h¥/(Nσ1/3) = 0.94. The results for the linear polymer grafts are in good agreement with the scaled relationship demonstrated by Patra and Linse based on molecular simulations; the slightly higher value of 0.64 as compared to 0.55 in their simulations is most likely due to the presence of raised structures (∼27 nm) in our nanopattern that lead to some confinement due to grafting to the sides. The agreement also shows that AFM measurements of the polymer zero-force thickness corrects for the errors in previous measurements as analyzed by Linse’s group.28,29 Using the experimental scaled relationship of Lee et al.21 and hydrated values for the thickness on the nanopattern and micropattern, a grafting dilution of 0.2 would be needed to bring the results into agreement. This is not unreasonable considering the larger silane surface attachment group compared to that of thiols. The addition of cross-linker to micro- and nanopatterned polymer grafts leads to an increase of the number of early terminations of chains as seen from the decrease of the polymer thickness with increasing cross-linker concentration. Most likely, this is caused by an increased likelihood of chain
chain loops facilitated by the very short cross-linker. Such an increased 8962

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Figure 7. Polymer property maps derived from 2D f
d mapping: quasi-contact thickness hw (a
d), zero-force thickness hf (e
h), fixed force distance (f = 0.6 nN) (i
l), and elastic modulus (m
p). Images are from PDMA (a, e, i, m), PDMA-BA050 (b, f, j, n), PDMA-BA075 (c, g, k, o), and PDMA-BA100 (d, h, l, p) synthesized for 10 min. Each image is 1 μm  1 μm, corresponding to a resolution of 25 nm in each direction, and the vertical scales of the line scans are in nanometers. The modulus scales are 0
1000 Pa and 0
5000 Pa in (m) and (n
p), respectively. Because of the long acquisition time, drift may occur and can result in some lateral distortion of the images.

bimolecular termination has been seen before for high-density grafted-from polymers.65,66 A higher number of early chain terminations is consistent with the fact that the maximum contour

length and the longest maximum rupture distance, seen in AFM rupture experiments, did not depend on the cross-linker concentration (Figure 5). Also in support of an increased bimolecular 8963

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Figure 8. Elastic moduli obtained from micropatterned PDMA-BA using the Hertz and the Cappella models and from the nanopatterned PDMA-BA using the Cappella model. For comparison, the elastic moduli of polyacrylamide gels with 5.5% and 7.5% acrylamide concentration were calculated from the shear modulus presented in the work by Yeung et al.71 The elastic moduli are presented as a function of the crosslinker concentration.

termination is the bimodal distribution of rupture distances, which we interpret as a highly cross-linked inner core with many loops and a less cross-linked outer corona. Finally, further support comes from the result that the lower mode of the maximum rupture distances becomes shorter with increased cross-linker concentration (Figure 5). Additionally, there appears to be a balance between polymer extension and termination events leading to an optimum homogeneity at 0.75% crosslinker concentration, as seen in the frequency, magnitude, and distance of the rupture events on the micropatterns as well as the nanopatterns. The introduction of cross-linker leads to only a small increase in confinement and reduction in the lateral extension of the polymer from the nanoislands, as seen by the slight reduction of polymer thickness between the nanopatterns (Figure 7e
h). Each nanoisland showed a thickness profile of a lower core—approximately covering the area of original nanoislands where chains are anchored—and a thicker ring extending from the nanopattern. Such a ring is most likely due to chains anchoring to the sides of the underlying nanopattern that has a topography with a height of 27 nm. Polymer gels are typically characterized by the elastic modulus, i.e., Young’s modulus. The fixed force measurements for the micro- and nanopatterned polymers indicate qualitatively that the modulus should increase with cross-linking density and show a lateral inhomogeneity of the elasticity for the nanopatterns (see Figure 7i
1). The Hertz model is often used to determine the elastic modulus of a material. However, the extreme proximity of the substrate in polymer grafts of less than 200 nm makes this approach highly questionable. The Hertz model assumes an infinite thickness of an elastic material relative to the indenting depth and the radius of the indenter. Another practical problem arises from the exact determination of the contact point in AFM measurements of soft materials, which is needed to calculate indentation. Nevertheless, Hertz-based models have been used successfully on polymer films of less than 200 nm thickness using semiempirical descriptions of the effective film modulus.30,47,48,55,58,67
70 We have described a procedure to objectively determine the contact point (see Figure S1 in the Supporting Information). The polymer thickness derived from this contact point agrees very well with those from contact-AFM measurements for micropatterned polymer gels, although it is always slightly larger as would be expected for the “zero force” condition. Most importantly

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these objectively determined contact points are much further from the substrate than if intuitively picked. On the basis of a clearly defined contact point, we obtained the Young’s modulus using two methods: (1) from an indentation range very near the contact point (only for the micropatterns) and (2) using the semiempirical model of Silbernagl and Cappella58 (see the Supporting Information for details on both methods). We chose the Silbernagl and Cappella model because of its reliance on only limiting behavior of the elastic response as the tip indents. We also chose a direct determination of the elastic modulus rather than an indirect method, such as lateral friction,32 in order to avoid the interference by other parameters and because the nanopattern topography would make interpretation of lateral force data difficult. For the micropatterned PDMA-BA (0%, 0.5%, 0.75%, and 1.0% (w:w)), the elastic modulus was 764 ( 98 Pa, 1282 ( 90 Pa, 1345 ( 106 Pa, and 1446 ( 112 Pa, respectively. For the nanopatterns the Young’s moduli differ only slightly from those on the micropatterns, although with more noise; modulus maps are shown in Figure 7m
p. The elastic modulus of PDMA-BA increases linearly with cross-linker concentration. These results— the magnitude as well as the slope of the cross-linker dependence— are surprisingly consistent with previously reported elastic moduli obtained from polyacrylamide gels, which are shown together with our data in Figure 8.71 Note that for the concentrations of cross-linker (0
1%) and grafting density investigated here the concentrations of PDMA-BA lie in the same range of polyacylamide gels with acrylamide concentrations between 5.5% and 7.5%. The results are also comparable to 4.5 kPa measured for thick films by Harmon et al.49 at their lowest crosslinking concentration, which is still slightly higher than the one used here. The elasticity maps of the nanopatterned grafts show a high modulus for the areas between the nanopatterns where polymers are not grafted but extended from the patterned areas. These areas are very thin, and any elastic behavior is strongly dominated by the substrate; that influence increases as the cross-linker density increases and the residual thickness in these areas decreases. The elasticity maps also indicate that around the nanopatterns there is a soft ring, coinciding with the increased height around the nanopatterns in the topographic images. Such a softer area supports our earlier conclusion that there is a structural difference between a more-cross-linked inner core and a less-cross-linked corona.

’ CONCLUSIONS We have successfully fabricated grafted PDMA and PDMABA polymers in the form of micro- and nanopatterns using iniferter-driven quasi-living radical polymerization by combining conventional photolithography and nanosphere lithography. The introduction of a cross-linker to the polymerization solution leads to an increased tendency of early termination most likely by looping and a core
corona distribution of the polymer with a stiffer, more highly cross-linked core, enriched in early terminations, and a highly swellable softer corona. Nanopatterned linear polymers show reduced thickness and cover a larger area than the underlying pattern in agreement with the expected scaling relationship for nanopatterned grafts that lack the lateral confinement of homogeneous films. The addition of cross-linker reintroduces some of the lateral confinement, but the confinement is not complete for the length of polymers reported here. 8964

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Langmuir There is an optimum cross-linker concentration that results in part from the opposing trends of increased confinement and increased early termination. We also developed an algorithm to objectively determine the contact point in AFM force spectroscopy, and using this contact point successfully determined the elastic modulus for micropatterned polymers of less than 200 nm thickness. Our results demonstrate the fabrication of nanopillars of controlled elastic modulus and highlight some of the factors to be considered to successfully fabricate grafted nanopillars of defined mechanical and structural properties.

’ ASSOCIATED CONTENT

bS

Supporting Information. Detailed information is provided on (1) the different methods to determine the thickness of swollen polymer grafts (Figure S1), (2) the rupture force analysis, (3) the two methods to determine the elastic modulus with sample results (Figures S2 and S3). Additionally, the thickness as a function of cross-linker concentration for nanopatterned polymers and non-patterned polymer films is compared (Figure S4). This material is available free of charge via the Internet at http://pubs.acs.org.

’ AUTHOR INFORMATION Corresponding Author

*Tel: 512-471-7434; Fax: 512-471-0616; e-mail: wfrey@ mail.utexas.edu.

’ ACKNOWLEDGMENT We thank David Kahn and Dipika Patel for their help with the grafting procedure. This work was supported in part by a grant to W.F. from NIH (5R21EB003038) and to S.P. from Higher Education Assistance Funds (HEAF) of the State of Texas. The scanning electron microscopy images were obtained at the Texas Tech University Imaging Center supported by NSF MRI (0421032). ’ REFERENCES (1) Stuart, M. A. C.; Huck, W. T. S.; Genzer, J.; Muller, M.; Ober, C.; Stamm, M.; Sukhorukov, G. B.; Szleifer, I.; Tsukruk, V. V.; Urban, M.; Winnik, F.; Zauscher, S.; Luzinov, I.; Minko, S. Nature Mater. 2010, 9, 101. (2) Wu, T.; Efimenko, K.; Genzer, J. J. Am. Chem. Soc. 2002, 124, 9394. (3) Milner, S. T. Science 1991, 251, 905. (4) Yamamoto, S.; Ejaz, M.; Tsujii, Y.; Fukuda, T. Macromolecules 2000, 33, 5608. (5) Prucker, O.; Ruhe, J. Macromolecules 1998, 31, 592. (6) Prucker, O.; Ruhe, J. Macromolecules 1998, 31, 602. (7) Edmondson, S.; Osborne, V. L.; Huck, W. T. S. Chem. Soc. Rev. 2004, 33, 14. (8) Ejaz, M.; Yamamoto, S.; Ohno, K.; Tsujii, Y.; Fukuda, T. Macromolecules 1998, 31, 5934. (9) Pyun, J.; Kowalewski, T.; Matyjaszewski, K. Macromol. Rapid Commun. 2003, 24, 1043. (10) Nakayama, Y.; Matsuda, T. Langmuir 1999, 15, 5560. (11) Benetti, E. M.; Zapotoczny, S.; Vancso, J. Adv. Mater. 2007, 19, 268. (12) Zapotoczny, S.; Benetti, E. M.; Vancso, G. J. J. Mater. Chem. 2007, 17, 3293. (13) Mrksich, M.; Whitesides, G. M. Trends Biotechnol. 1995, 13, 228.

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