Letter pubs.acs.org/Langmuir
Tapping-Mode AFM Study of Tip-Induced Polymer Deformation under Geometrical Confinement Hong Zhang,†,‡ Yukio Honda,‡ and Shinji Takeoka*,† †
Department of Life Science and Medical Bioscience, Graduate School of Advanced Science and Engineering, Waseda University (TWIns), 2-2 Wakamatsu-cho, Shinjuku-ku, Tokyo 162-8480, Japan ‡ Consolidated Research Institute for Advanced Science and Medical Care, Waseda University (ASMeW), 513 Waseda Tsurumaki-cho, Shinjuku-ku, Tokyo 162-0041, Japan S Supporting Information *
ABSTRACT: The morphological stability of polymer films is critically important to their application as functional materials. The deformation of polymer surfaces on the nanoscale may be significantly influenced by geometrical confinement. Herein, we constructed a mechanically heterogeneous polymer surface by phase separation in a thin polymer film and investigated the deformation behavior of its nanostructure (∼30 nm thickness and ∼100 nm average diameter) with tapping-mode atomic force microscopy. By changing different scan parameters, we could induce deformation localized to the nanostructure in a controllable manner. A quantity called the deformation index is defined and shown to be correlated to energy dissipation by tip−sample interaction. We clarified that the plastic deformation of a polymer on the nanoscale is energy-dependent and is related to the glass-to-rubber transition. The mobility of polymer chains beneath the tapping tip is enhanced, and in the corresponding region a rubberlike deformation with the lateral motion of the tip is performed. The method we developed can provide insight into the geometrical confinement effects on polymer behavior.
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INTRODUCTION Compared to metals and ceramics, polymers are rather compliant and soft. Moreover, the mechanical properties of polymers can be significantly affected by temperature and interactions with the surroundings. It has been widely recognized that the glass-transition temperature (Tg) of a thin polymer film differs from that of the bulk because of the geometrical confinement effects (i.e., the enhanced mobility of polymer chains at free surfaces and the depressed mobility at substrates1−3). If polymer chains are restrained in two or three dimension, then it is reasonable to expect nanostructured polymer to exhibit more complex behaviors that are derived from the multidimensional confinement. Although the nature of polymer confinement is still unclear, many polymer-based nanodevices have been developed, which usually involve a wide variety of nanostructures. Therefore, understanding the polymer deformation in nanoconfined geometries is critically important to the utilization of actual nanodevices. Atomic force microscopy (AFM) is a powerful tool for the characterization of structures and studying the properties of solid surfaces. In a seminal report, Goh et al. first recognized that the formation of a periodic pattern on the nanoscale was due to the interaction between the AFM tip and the polymer surface.4 The resulting ripplelike structures were oriented perpendicularly to the scanning direction. Further studies © 2013 American Chemical Society
reported that the magnitude of these ripplelike structures was significantly influenced by several parameters such as the number of scans,5−9 molecular weight of polymer,5 applied load,6−8 scan speed,7,8 and temperature.9 To some extent, these results are due to the applied load from the operating AFM in contact mode, which produces a large lateral force and an unfavorable pressure on the order of the polymer yield stress. Thus, to avoid such deformation, it is more circumspect to use tapping-mode AFM when scanning is performed on polymer surfaces. In tapping-mode AFM, a tip is excited by an external signal at a frequency close to its resonance value with free amplitude. The frequency and amplitude of the tip change because of the tip−sample interaction. The cantilever is then brought close to the sample surface until the amplitude reaches a given set point value. In general, topographical imaging is performed while the amplitude is kept constant by a feedback loop. The phase shift between excitation and the response of the tip can also be recorded to map compositional variations as complementary features to the topography, which may reflect the mechanical properties of the sample and the local energy dissipation. With Received: October 29, 2012 Revised: January 16, 2013 Published: January 16, 2013 1333
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Table 1. Summary of the Tapping-Mode AFM Scan Parameters Used in This Study entry
set point (V)
free amplitude (V)
attenuation ratio, γsp
drive amplitude (mV)
drive frequency (kHz)
0 1 2 3 4
1.5 1.5 0.5, 1.5, 2.5, 3.5, 4.5 2.5 2.0
6.6 6.6 6.6 7.4, 7.2, 6.6, 5.2, 3.5 6.6, 6.6, 6.6, 6.1, 2.9
0.23 0.23 0.08, 0.23, 0.38, 0.53, 0.68 0.34, 0.35, 0.38, 0.48, 0.71 0.30, 0.30, 0.30, 0.33, 0.69
400 400 400 600, 500, 400, 300, 200 400
136.5 136.5 137.0 137.0 136.0, 136.5, 137.0, 137.5, 138.0
Information Figure S1). The resonance frequency of the cantilever (136.156 kHz) was confirmed by the thermal power spectral density (PSD), and the spring constant was calibrated to be 7.5 N m−1, in general agreement with the nominal values (75−175 kHz and 4.0− 22.3 N m−1). The average sensitivity (i.e., AmpInvOLS) of an optical detection system for a cantilever was calculated to be 124.2 nm V−1. All of the AFM scan parameters are listed in Table 1. The scan speed was 25.04 μm s−1, and the scan size was 5 × 5 μm2 unless otherwise mentioned. According to the Whangbo et al. description of tapping conditions, the attenuation ratio (γsp) of the set point amplitude to the free oscillation was calculated as well, indicating that imaging was mainly performed at moderate tapping (γsp = 0.4−0.7) and hard tapping (γsp < 0.4) in this study.17
tapping-mode AFM, the tip makes intermittent mechanical contact with the sample surface. Although it can minimize the applied load and essentially eliminate the lateral force, it can also bring about polymer surface deformation. As a preliminary experiment, polystyrene films with a thickness of ∼30 nm were prepared on a silicon oxide substrate. By repetitive scanning, ripplelike structures gradually developed on the surface (Supporting Information Movie S1). In the past few years, efforts have been made toward the investigation of imaging the internal structure of polymer films by AFM10 and interpreting the property changes of polymer materials that are related to the tip−sample interaction.11−13 Our previous studies have focused on the fabrication, properties, and applications of flat, smooth thin polymer films.14,15 In a recent study, we constructed a nanostructured thin film derived from a polymer blend (i.e., polystyrene (PS)/ poly(methyl methacrylate) (PMMA)) and investigated its phase separation behavior with the help of a freestanding technique.16 Herein, the deformation behavior of such a mechanically heterogeneous surface was investigated with tapping-mode AFM. Compared to the homopolymer thin film, the compositional continuity of the polymer blend thin film is broken up by the geometrical confinement. The advantage of our method is that the magnitude of deformation can be quantitatively characterized. In combination with the theoretical analysis of tapping-mode AFM, imaging the force and energy dissipation by the tip−sample interaction can be determined as well, which allowed us to provide insight into the deformation of polymer on the nanoscale.
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RESULTS AND DISCUSSION We prepared thin polymer films with phase-separated nanostructures by using a PS/PMMA blend, which has been well described elsewhere.16 In a typical sample, the morphological features of an as-cast film show numerous PS craters isolated by a matrix of PMMA (Figure 1a). The variation in
EXPERIMENTAL SECTION
Film Preparation. The polymers used in this study were PS (Mw = 170 000, Mw/Mn = 1.06, d = 1.05 g/cm3, and Tg(bulk) = 100 °C) and PMMA (Mw = 120 000, Mw/Mn = 1.8−2.0, d = 1.19 g/cm3, and Tg(bulk) = 105 °C), which were purchased from Chemco Scientific Co., Ltd. (Osaka, Japan) and Sigma-Aldrich (St Louis, MO), respectively. The substrates were silicon (100) wafers covered with 200 nm of thermally grown silicon oxide purchased from KST World Co. (Fukui, Japan), which were cut to a size of 20 × 20 mm2. The substrates were cleaned at 120 °C in a piranha solution of sulfuric acid and 30% hydrogen peroxide (3:1 v/v) for 15 min and then thoroughly rinsed with deionized water (resistivity 18 MΩ cm) and dried with compressed nitrogen gas. Polymer blend solutions were prepared by dissolving PS and PMMA in a 1:9 weight ratio in analytical-grade ethyl acetate as a common solvent. The mixture was stirred overnight prior to film preparation in order to ensure complete dissolution. The total polymer concentration in the solutions was 1.0 wt %. The thin polymer films were prepared by spin-casting at 5000 rpm for 60 s using an MS-A100 spin coater (Mikasa Co., Ltd., Tokyo, Japan). All procedures for film preparation were conducted at room temperature (25 °C) and normal relative humidity. AFM and Cantilever. The AFM used in this study was an MFP3D instrument (Asylum Research, Santa Barbara, CA) controlled by 0909091 + 1214 and Igor Pro 6.12A software (WaveMetrics, Lake Oswego, OR). All AFM procedures were conducted at room temperature (25 °C) with an AC200TS silicon cantilever (Olympus, Tokyo, Japan) with a pyramid tip of radius ∼9 nm (Supporting
Figure 1. Number of consecutive scans induced a surface deformation. (a) AFM image after one scan. (b) In situ AFM image after nine consecutive scans. (c) Lateral profile corresponding to the line shown in panels a and b. The green line shows an ∼1.5 nm height difference between the original PS and PMMA regions. (d) Height distribution histogram of each AFM image after nine scans. (e) Cartoon of a tapping tip moving across a nanostructured polymer surface resulting in the formation of a shoulder structure. (f) Schematic representation of the definition of the deformation index. After nine scans, the portion of the height distribution histogram corresponding to the shoulder structures is exhibited in pink. 1334
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set point (Figure 2b and Supporting Information Movie S3), drive amplitude (Figure 2e and Supporting Information Movie
height between PS and PMMA (∼1.5 nm) was due to differences in their solubility during film preparation. The thickness of the films was ∼30 nm, and the average diameter of the PS regions was measured to be ∼100 nm (Supporting Information Figure S2). We intentionally operated AFM in hard tapping (γsp = 0.23). By repetitive scanning, all of the PS regions were gradually deformed and ploughed out of the PMMA matrix, whereas the PMMA matrix is deformationresistant (Figure 1b and Supporting Information Movie S2). A previous report also showed that the PMMA surface was completely stable to contact-mode AFM scanning, and no ripplelike structures were exhibited.18 It was indicated that the obtained morphology was not reliable when any factors were changed for tapping-mode AFM.11 In this study, scanning electron microscopy (SEM) was also utilized to confirm the tipinduced deformation, and the obtained morphologies from AFM images were clarified to be reliable (Supporting Information Figure S3). Although the mechanical properties of PS and PMMA are very similar in the bulk,19 the observed deformation is limited to PS nanostructures. It has been well demonstrated that with decreasing film thickness below ∼40 nm on the silicon oxide substrate the Tg of PS decreases, whereas the opposite is true for PMMA.3 The difference can be interpreted from the stronger interaction between PMMA chains and the substrate, which restricts the polymer chains’ mobility. By using a PS/ PMMA blend thin film, our AFM results supported these previous studies in a compelling way and showed that the surface Tg (i.e., the molecular mobility on the surface) plays an important role in determining the deformation of the thin polymer film. As the number of scans increased, the amount of piled up PS increased (Figure 1c), indicating that the deformation was a cumulative process akin to the formation of ripplelike structures. The height scale of all AFM images was set to be 7 nm (from −3.5 to 3.5 nm) in this study. Height distribution histograms of AFM images for repetitive scanning are shown in Figure 1d, whose variation properly describes the deformation. The surface roughness (root-mean-square, rms) and the surface skewness were introduced to describe the deviation and the asymmetry of the height distribution histogram, respectively, as shown in Supporting Information 1. Herein, the PMMA matrix restricted the formation of ripplelike structures, forcing the PS to be ploughed out of the inner regions and to form a shoulder structure (Figure 1e). Therefore, the magnitude of deformation can be quantitatively characterized by counting the height of the displaced PS based on the surface of the PMMA matrix as a standard level. As such, the average height of the PS shoulders was calculated by a weighted average from the height distribution histogram, referred as the deformation index (eq 1, Figure 1f).
Figure 2. Series of scan parameters inducing surface deformation. (a) Number of scans influences the surface deformation. AFM was equally divided into two parts (top and bottom regions of the image). The top region shows the morphology after 1 scan, and the bottom region shows the morphology after 10 scans. (b) Influence of the set point. The AFM image was equally divided into five parts from the top to the bottom where the set point values were 0.5, 1.5, 2.5, 3.5, and 4.5 V, respectively. (e) Influence of drive amplitude. The AFM image was equally divided into five parts from the top to the bottom where the drive amplitude was set to 600, 500, 400, 300, and 200 mV, respectively. (f) Influence of drive frequency. The AFM image was equally divided into five parts from the top to the bottom where the drive frequencies were 136.0, 136.5, 137.0, 137.5, and 138.0 kHz, respectively. (c, d, g, and h) Cross-sectional 3D images generated from a, b, e, and f, respectively.
S4), and drive frequency (Figure 2f and Supporting Information Movie S5)) during scanning. The change in morphology was continuously observed after varying each scan parameter. Combined with the results of repetitive scanning (Figure 2a and Supporting Information Movie S2), we believe that the surface deformation can be performed in a controllable manner by using our strategy. It was found that the magnitude of surface deformation increased with increasing the number of scans, set point, drive amplitude, and drive frequency. The results also showed that the influence of each mentioned scan parameter of deformation was somewhat different. For example, as the set point increases, the heights of both elevated PS and the inner PS regions increased (Figure 2d). However, as the drive amplitude increases, the amount of piled up PS increased with the depression of the inner PS regions (Figure 2g).
3.5
deformation index ≡
∑H = H HI 0
3.5
∑H = H I 0
(1)
where H and I refer to data elements of height and the corresponding intensity, H0, is the surface height of the PMMA matrix. For example, the deformation index was found to increase from ∼0.7 to ∼1.3 nm after nine consecutive scans. Given that the tip−sample interaction forms the basis of the surface deformation, we decided to investigate whether such a deformation could be induced more generally. Thus, we changed the scan parameters of tapping-mode AFM (i.e., the 1335
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Figure 3. APD curve measurement and determination of tip−sample interaction. (a) APD curve measurement, (e) calculation of imaging force, and (i) energy dissipation upon changing the number of scans. (b, f, and j) Changing the set point. Set point values are color coded as green, 0.5; orange, 1.5; cyan, 2.5; pink, 3.5; and blue, 4.5 V. (c, g, and k) Changing the drive amplitude. Drive amplitude values are color coded as green, 600; orange, 500; cyan, 400; pink, 300; and blue, 200 mV. (d, h, and l) Changing the drive frequency. The drive frequency values are color coded as green, 136.0; orange, 136.5; cyan, 137.0; pink, 137.5; and blue, 138.0 kHz. Vertical lines in e−l show the set point amplitude used in each case.
Table 2. Results for the Determination of the Deformation Index and Tip−Sample Interaction entry 1 2 3 4
AFM results Figure Figure Figure Figure
1a,b and MovieS2 2b and MovieS3 2e and MovieS4 2f and MovieS5
deformation index (nm) 0.71, 0.58, 1.24, 0.76,
0.96, 0.81, 1.07, 0.79,
1.16, 0.94, 0.95, 0.85,
1.26, 1.01, 0.82, 0.88,
1.28 1.06 0.69 1.03
imaging force (μN)
energy dissipataion (pW)
4.76 5.70, 4.77, 3.83, 2.90, 1.97 4.57, 4.33, 3.83, 2.50, 0.95 4.30, 4.30, 4.30, 3.86, 0.82
7.03, 21.09, 35.15, 49.21, 63.27 1.49, 9.38 18.30, 25.17, 26.79 47.54, 30.96, 18.30, 9.45, 4.65 6.24, 9.92, 13.83, 16.73, 19.41
this study could be adjusted in the range from ∼1 to ∼50 pW (Figure 3i−l). The details on calculating the tip−sample interaction are shown in Supporting Information 2 and 3. All of the results are listed in Table 2. We plotted the deformation index as a function of the tip− sample interaction. Interestingly, it was found that the deformation index is strongly correlated with energy dissipation (Figure 4b). The results showed that the greater the energy that was introduced into the PS regions, the greater the amount of PS that was piled up. Combined with AFM images, in this study we found that ∼20 pW energy dissipation was sufficient to pile up an observable quantity of PS out of the PMMA matrix. Because of the ultrasmall volume of each inner PS region, the
It should be noted that in tapping-mode AFM the tip is moving in the trace (left to right) and subsequently in the retrace (right to left) directions for each line. All of the presented AFM images were collected by scanning lines from left to right. The asymmetry of the resulting morphology possibly reflected the foremost direction of tip motion. AFM images collected by the retrace channel were simultaneously taken and found to be similar to those collected by the trace channel (i.e., the shoulder structure would not be erased or change the shape by AFM scanning with the same scan parameters except from right to left (Supporting Information Movie S6)). We reasoned that this observation provided an important clue that the deformation of polymer can be significantly affected by geometrical confinement. We have verified that the change in morphology persists for at least 1 month at room temperature, which suggest that the deformation is essentially a plastic deformation. To obtain a deeper understanding concerning the influence of each scan parameter on the deformation, an amplitude-phase distance (APD) curve measurement was carried out in PS regions, where the amplitude and phase were recorded as the tip−sample distance was decreased. In each case, at least 20 APD curves were recorded and the results were averaged (Figure 3a−d). The negative phase shifts (
