Conformation of Single Polymer Chain in Rubbed Thin Film Observed

Article pubs.acs.org/Langmuir

Conformation of Single Polymer Chain in Rubbed Thin Film Observed by Fluorescence Imaging Toru Ube,†,§ Akihiko Shin,† Hiroyuki Aoki,*,†,‡ and Shinzaburo Ito†,‡ †

Department of Polymer Chemistry, and ‡Advanced Biomedical Engineering Research Unit, Kyoto University, Nishikyo, Kyoto 615-8510, Japan ABSTRACT: The conformation of poly(methyl methacrylate) (PMMA) chains in a thin film after the rubbing process was investigated through the direct observation of the single chains by scanning near-field optical microscopy (SNOM) and excitation polarization modulation microscopy (EPMM). The rubbing at room temperature hardly changed the dimension on the whole chain scale in spite of the increase in orientational order on the segmental scale. The increase in the chain dimension along the rubbing direction was observed in the film rubbed at the higher temperature, which showed a surface morphology with fine groove. The extension ratio of the whole chain in the rubbed film was much smaller than that in the uniaxially stretched film. This indicates that the rubbing process mainly induces the conformational change on the length scale of the monomer unit rather than for the whole chain.

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INTRODUCTION The surface properties of polymer materials such as friction, wettability, and adhesion play important roles in their practical use. These macroscopic properties are dominated by the structure of the polymer chain near the surface. An understanding of the polymer chain near the surface is essential to controlling and improving the performance of polymer materials. The rubbing process is an important surfacemodification method. It is well known that liquid-crystal molecules are strongly aligned on a rubbed polymer surface, and this has been applied to flat panel displays.1 The orientational order of polymer chains in a film increases by the rubbing process, which has been evaluated by birefringence2−4 and infrared absorption3−5 measurements. In recent years, surface-sensitive experimental techniques such as grazing incidence X-ray scattering,6 near-edge X-ray absorption fine structure,7−9 and sum-frequency generation vibrational spectroscopy have been developed,10 and they revealed that rubbing induces a higher orientation near the surface than in the bulk. The surface morphology of the rubbed film has been investigated by AFM.11−14 It was shown that the groovelike structure was formed during the rubbing process. With these methods, the effect of rubbing has been investigated in terms of the surface morphology and orientation on the monomeric scale. However, less information has been obtained for the behavior of polymer chains under rubbing compared to that under tensile and shear deformations. For example, the conformation of a single polymer chain on the whole chain scale, which is characterized by parameters such as the radius of gyration, has not been investigated so far for the chain in the rubbed film. Approaches from the various length scales are essential to describing the exact chain behavior.15 Furthermore, the orientation obtained from previous measurements suffers from blurring by averaging over a large number of molecules. The direct observation of the single chain would be a more © 2012 American Chemical Society

effective approach to revealing the response of the polymer chain to rubbing, which could be inhomogeneous. To detect in situ features of polymer chains located inside a bulk medium, a single chain must be distinguished from the surrounding ones. Fluorescence labeling is an established method, and it has been applied to the observation of single DNA molecules.16,17 However, conventional fluorescence microscopy suffers from the low spatial resolution of a half wavelength of light because of the diffraction limit. Therefore, the application of optical microscopy to single macromolecular imaging has been limited to the observation of huge biomacromolecules such as DNA. Scanning near-field optical microscopy (SNOM) is an emerging scanning probe technique that allows optical measurements at high resolution beyond the diffraction limit.18,19 The light incident on a sub-wavelengthsized aperture generates an optical near field restricted in the space of the aperture size. This allows one to illuminate a specimen and to obtain the optical response from the nanometric area. The fluorescence labeling technique can be combined with SNOM. This enables us to observe the conformation of the flexible single chain at high resolution directly, which is fluorescently labeled and contrasted with the surrounding unlabeled polymers.20−26 Furthermore, the topographic image is simultaneously obtained during the SNOM measurement. The chain conformation can be discussed in contrast to the surface morphology. In addition to the observation of the whole chain by SNOM, the orientation of the dyes in the fluorescently labeled polymers can be evaluated by the use of polarized light because the absorption and fluorescence anisotropies reflect the orientation of the transition dipole. The polarization measurement under a Received: June 21, 2012 Revised: September 3, 2012 Published: September 5, 2012 13871

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toluene solution of unlabeled PMMA and PMMA-Pe (0.02 wt % with respect to the unlabeled polymer) was spin-coated onto a glass substrate to form a film with a thickness, d, of 7−170 nm. Rubbing was performed at 30, 60, and 80 °C by unidirectionally pulling a velvet cloth for 30 cm (80 °C) or 250 cm (30 and 60 °C) at a speed of 1 cm/ s under a load of 4 g/cm2. Excitation Polarization Modulation Microscopy. The orientation anisotropy of fluorescent dyes was measured by EPMM, which consists of an inverted fluorescence microscope (TE-2000, Nikon) equipped with an EMCCD camera (Cascade II, Roper Scientific). A linearly polarized 442 nm laser beam was passed through an electrooptic modulator (EOM) (EO-AM-NR-C4, Thorlabs) and focused on the back focal plane of an objective lens (100×, 1.4 NA, Nikon) for Köhler illumination. The fluorescence signal was collected through a filter cube (BV-2A, Nikon) that contains a dichroic mirror (455 nm) and a long-pass filter (470 nm). The polarization direction of the excitation beam was alternately modulated in the parallel and perpendicular directions with respect to the rubbing axis at a frequency of 0.25 Hz by the EOM. The fluorescent images under parallel and perpendicular polarization illumination were acquired synchronously with the polarization modulation. SNOM Measurement. The SNOM measurement was performed with a commercially available instrument (α-SNOM, WITec) using a hollow cantilever probe with a subwavelength aperture of 60 nm. The laser beam at a wavelength of 441 nm (BCL-015−440, CrystaLaser) was focused onto the back side of the aperture to generate the optical near-field. While the sample surface was scanned in contact mode with the cantilever, the perylene fluorescence was collected with a microscope objective (0.80NA, 60×, Nikon) from the back side of the substrate, passed through a long-pass filter (LP02-442RS-25), and detected with a photomultiplier tube (H8631, Hamamatsu Photonics). The SNOM measurement was carried out under ambient conditions. All SNOM images were taken with the same probe.

microscope reveals the spatial distribution of the orientational anisotropy.26,27 In this study, the conformation of the single polymer chain in the rubbed poly(methyl methacrylate) (PMMA) film is investigated through direct observation. We rubbed films at various temperatures to investigate the effect of the rubbing on the chain deformation in relation to the chain mobility. The conformations for the whole chain and monomeric scales are evaluated with SNOM and excitation polarization modulation microscopy (EPMM), respectively. The chain orientation induced by the rubbing is discussed in contrast to the surface morphology.

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EXPERIMENTS

Sample Preparation. Perylene-labeled PMMA (PMMA-Pe, Figure 1) was synthesized by the random copolymerization of methyl methacrylate and 3-perylenyl methyl methacrylate as described elsewhere.23,28 The fraction of the labeled unit was evaluated to be 0.77% by UV−vis absorption (U3500, Hitachi). The unlabeled PMMA was synthesized by atom-transfer radical polymerization.29 The weightand number-averaged molecular weights, Mw and Mn, respectively, were determined by GPC measurement as shown in Table 1. A mixed

Table 1. Characterization of PMMA PMMA-Pe PMMA

Mw/106

Mn/106

Mw/Mn

1.99 2.12

1.58 1.33

1.26 1.60

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RESULTS AND DISCUSSION

Figure 2a shows an EPMM fluorescence image of the thinnest PMMA film (d = 7 nm) before rubbing. Each fluorescent spot was confirmed to be an individual PMMA-Pe chain from the statistical analysis.30 The orientation of the dye molecules in each labeled chain was examined in terms of (Ix − Iy)/(Ix + Iy), where Ix and Iy are the intensities of the fluorescence from each labeled chain excited by linearly polarized light in the x and y directions, respectively. The x axis is defined as the rubbing

Figure 1. Chemical structure of random copolymer PMMA-Pe.

Figure 2. (a) Fluorescence and (b−e) anisotropy images of single polymer chains obtained by EPMM in the PMMA films (d = 7 nm) (a, b) before rubbing and after rubbing at (c) 30, (d) 60, and (e) 80 °C. 13872

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direction of the film, and the y axis is perpendicular to it. (Ix − Iy)/(Ix + Iy) for each labeled chain indicates the average orientation of the dye moieties. In our previous study, we examined the relationship between the orientational order parameters of the dye and the backbone segment for PMMAPe by comparing the results of EPMM and birefringence measurements.26 It was shown that the orientation of the dye correlates with that of the backbone segment. The positive and negative values of (Ix − Iy)/(Ix + Iy) indicate the orientation of the backbone in the x and y directions, respectively. For the complete orientation in the x and y axes, the values of (Ix − Iy)/ (Ix + Iy) are +1 and −1, respectively. Thus, the segmental orientation of individual chains can be obtained from the EPMM image. The values of (Ix − Iy)/(Ix + Iy) for the individual chains are represented by pseudocolors in Figure 2b−e. Whereas the conformation of each chain was not clearly observed in the fluorescence image because of the diffractionlimited resolution, the segmental orientation for a single PMMA chain can be examined through the excitation anisotropy. For the sample film before rubbing, Figure 2b shows PMMA-Pe chains with various colors, indicating the random orientation of the chain segment. In Figure 2c−e, however, many of the observed chains are depicted in red and orange, indicating the preferential orientation in the rubbing direction. Figure 3 shows the histograms of (Ix − Iy)/(Ix + Iy). In the film before rubbing, (Ix − Iy)/(Ix + Iy) showed a symmetric distribution at a peak position of zero, indicating the isotropic orientation. We compared the experimental value of (Ix − Iy)/(Ix + Iy) in the film before rubbing with that obtained from the random walk simulation.31 The result of the simulation is shown by the solid curve in Figure 3a. The experimental data was in good agreement with the simulation,

indicating the random distribution in the initial state. In the rubbed films, the distribution was shifted to a higher (Ix − Iy)/ (Ix + Iy) value. This indicates that the orientational order of fluorescent dyes in each polymer chain increased by rubbing. The orientational order parameter of fluorescent dyes in each chain is defined as p=

3⟨cos2 θd⟩chain − 1 2

(1)

where θd is the angle between the x axis and the molecular axis of the dye and the bracket ⟨ ⟩chain represents the statistical average over the dye moieties for each chain. The value of p for each chain can be evaluated from the excitation anisotropy as32

p=

Ix − Iy Ix + 2Iy

(2)

We define the ensemble-averaged order parameter, p̅, as the average of p over all observed chains weighted by the emission intensity. The rubbing temperature dependence of p̅ is shown in Figure 4. The value of p̅ was zero before rubbing. After rubbing, p̅ took positive values, indicating that the chain segment was oriented along the rubbing direction. The film rubbed at 80 °C showed a higher anisotropy than those at 30 and 60 °C, reflecting that the molecular mobility increases as the temperature approaches Tg. This is consistent with the literature, which investigated the optical phase retardation of the polyimide films rubbed at various temperatures.13 We next investigated the excitation anisotropy of the chains in rubbed PMMA films with various thicknesses (7−170 nm) using EPMM. Figure 5 shows the thickness dependence of p̅ in films rubbed at 30 °C. The significant change in p̅ was not observed in films thicker than 30 nm. This implies that the rubbing mainly induces the orientation of chain segments within several nanometers from the surface and hardly affects the segments located in deeper regions of the film. In addition to the segmental orientation, which was evaluated by EPMM, we next investigated the conformation on the whole chain scale and the surface morphology by SNOM. Figure 6a− d shows the fluorescent SNOM images of the PMMA films (d = 7 nm). The perylene-labeled PMMA chains embedded in the unlabeled matrix were observed as the bright spots in the fluorescent image. The SNOM images showed much higher resolution than the fluorescent images obtained with diffraction-limited optics (Figure 2). The topographic images simultaneously obtained during the SNOM measurements are

Figure 3. Histograms of excitation anisotropy of chains obtained by EPMM for the films (d = 7 nm) (a) before rubbing and after rubbing at (b) 30, (c) 60, and (d) 80 °C. The solid curve in panel a shows the random distribution obtained from the computer simulation.

Figure 4. Rubbing temperature dependence of the excitation anisotropy of PMMA films (d = 7 nm). 13873

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shown in Figure 6e−h. The groove along the rubbing direction with a height of 2−5 nm was observed in the rubbed films. The films were partially scratched and peeled off by rubbing at 60 and 80 °C. The film rubbed at 80 °C showed markedly different morphology from the others. There were fine grooves with a height of 3−15 nm. This suggests that the surface of the film was largely deformed by the shear force. In this film, some of the chains had a highly stretched conformation along the rubbing direction (Figure 6d).

Figure 5. Excitation anisotropy of PMMA films after rubbing at 30 °C plotted against the thickness. The solid curve was drawn as a guide for the eye.

The conformation of the single PMMA chain was quantitatively evaluated from the fluorescence intensity distribution.22 The fluorescence intensity is proportional to the number of fluorescent dye molecules randomly introduced into the PMMA-Pe chain; therefore, the intensity at each pixel is proportional to the number of chain segments therein. The first moment of the fluorescence intensity distribution denotes the position of the center of mass, 1 r0 = ∑ riIi I i (3) where Ii is the fluorescence intensity at the ith pixel, ri is the position vector, and I is the total fluorescence intensity from the single chain. The second moment of the fluorescence intensity distribution along the x axis is calculated as 1 R xx 2 = ∑ (xi − x0)2 Ii I i (4)

Figure 6. (a−d) Fluorescent and (e−h) topographic images of single polymer chains in the PMMA films (d = 7 nm) obtained by SNOM: (a, e) before rubbing and after rubbing at (b, f) 30, (c, g) 60, and (d, h) 80 °C.

where a2 is the variance of the point-spread function. We defined the extension ratio at the single-chain level, λc, as

where xi and x0 are the x coordinates of the ith pixel and the center of mass, respectively. Rxx indicates the dimension of the fluorescent spot along the x axis. SNOM shows a much higher resolution than does conventional microscopy; however, it still suffers from the limit of finite resolution of ∼100 nm, which is caused by the finite dimension of the aperture. Therefore, the observed value of Rxx is somewhat larger than the true dimension of the chain, Rxx * .33 The point-spread function in the fluorescent SNOM measurement is well approximated as a Gaussian function. The fluorescent image is expressed by the convolution of the spacial distribution of the chain segment and the point-spread function. It is mathematically derived that the second moment of the convoluted function is a sum of the second moments of the original functions.34 Therefore, R xx 2 = R *xx 2 + a 2

λc 2 =

⟨R *xx 2 ⟩ ⟨R *xx 2 ⟩0

(6)

where ⟨R*xx ⟩0 and ⟨R*xx ⟩ denote the averages of the true dimensions of the chains along the x axis before and after rubbing. A number of chains (100−150) were analyzed to evaluate λc for each sample. The details of the analysis method were described in our previous paper.26 Figure 7 shows histograms of R*xx normalized by the average value before rubbing, ⟨R*xx⟩0. The distribution function of R*xx/⟨R*xx⟩0 was not significantly changed by the rubbing at 30 °C. This indicates that the global conformation of the whole chain was not changed by the rubbing near room temperature. The film rubbed at 60 °C also showed a similar distribution function to that of the unrubbed film. In the film rubbed at 80 2

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Figure 8. Chain-extension ratio plotted against the average excitation anisotropy for the thin films rubbed at (●) 30, (▲) 60, and (■) 80 °C. Open circles represent the bulk film uniaxially stretched at 160 °C.

whole chain scale below bulk Tg. Even in this case, the value of λc was still smaller than that for the uniaxial extension above Tg. This indicates that the rubbing process mainly induces the conformational change on the length scale of the monomer unit rather than the whole chain.

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CONCLUSIONS The conformation of the single polymer chain in the rubbed thin film was investigated through direct observation by EPMM and SNOM. The orientational anisotropy of the dyes in the labeled chains increased by rubbing, indicating that the conformation on the length scale of the monomer unit was rearranged. However, the conformation on the whole chain scale was not significantly changed by the rubbing near room temperature. The increase in the chain dimension along the rubbing direction was observed in the film rubbed at 80 °C, which was associated with the surface morphology with the fine groove. This conformational change on the whole chain scale below bulk Tg is expected to be caused by the reduced number of entanglements and the stress-enhanced chain mobility. The extension ratio of the whole chain in the rubbed film was much smaller than that in the uniaxially stretched film. This indicates that the rubbing process mainly changes the local conformation on the length scale of the monomer unit rather than the global conformation of the whole chain.

Figure 7. Histograms of the chain dimension parallel to the rubbing direction, which was normalized by the initial average value for the films (d = 7 nm) (a) before rubbing and after rubbing at (b) 30, (c) 60, and (d) 80 °C.

°C, however, the peak shifted to the higher value and the distribution became broader. This suggests that the chains were stretched along the rubbing direction. The deformation on the whole chain scale accompanied the large change in the surface morphology, which was observed in the topographic image. The average chain extension ratio, λc, is plotted against p̅ in Figure 8. The data for the uniaxial extension of the bulk film at 160 °C in our previous study25 are also shown by the open symbols therein. λc indicates the deformation on the whole chain scale, whereas p̅ indicates the orientation on the segmental scale. In the uniaxial extension process above Tg, both the orientation on the segmental scale and the dimension of the whole chain increased with strain because the chain conformation is deformed according to the affine transformation. However, the rubbing process near room temperature hardly changes the dimension of the whole chain in spite of the increase in the orientational anisotropy of the fluorescent dyes. At room temperature, the motion of the main chain is frozen. It is also plausible that the chains interact strongly with the substrate, which restricts the molecular motion. Therefore, it is unlikely that the rearrangement of the chain contour on the whole chain scale is induced by rubbing. Only the orientation on the segmental scale was induced near the surface at low temperature. In the rubbing process at the higher temperature, the motion of the main chain is thermally activated; therefore, the individual PMMA chains were slightly elongated in the rubbing direction. The deformation on the whole chain scale was observed even below the bulk Tg of PMMA. In a thin film, there are fewer entanglements compared to the number in the bulk state, resulting in a higher chain mobility.35 Furthermore, recent studies on the plastic flow of bulk PMMA have shown that the chain mobility is also enhanced by applying the stress.36,37 These decreased numbers of entanglements and the applied stress are expected to allow the deformation on the

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Present Address §

Research and Development Initiative, Chuo University, 1-1327 Kasuga, Bunkyo-ku, Tokyo 112-8551, Japan. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by grants-in-aid from the Ministry of Education, Culture, Sports, Science and Technology, Japan (MEXT). We also acknowledge the Innovative Techno-Hub for the Integrated Medical Bioimaging Project of the Special Coordination Funds for Promoting Science and Technology from MEXT. T.U. is grateful for a research fellowship from the Japan Society for the Promotion of Science for Young Scientists. 13875

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(24) Ube, T.; Aoki, H.; Ito, S.; Horinaka, J.; Takigawa, T.; Masuda, T. Affine deformation of single polymer chain in poly(methyl methacrylate) films under uniaxial extension observed by scanning near-field optical microscopy. Polymer 2009, 50, 3016−3021. (25) Ube, T.; Aoki, H.; Ito, S.; Horinaka, J.; Takigawa, T.; Masuda, T. Relaxation of single polymer chain in poly(methyl methacrylate) films under uniaxial extension observed by scanning near-field optical microscopy. Macromolecules 2011, 44, 4445−4451. (26) Ube, T.; Aoki, H.; Ito, S.; Horinaka, J.; Takigawa, T. Relaxation of single polymer chain in binary molecular weight blends observed by scanning near-field optical microscopy. Soft Matter 2012, 8, 5603− 5611. (27) Gupta, V. K.; Kornfield, J. A.; Ferencz, A.; Wegner, G. Controlling molecular order in “hairy-rod” Langmuir-Blodgett films: a polarization-modulation microscopy study. Science 1994, 265, 940− 942. (28) Aoki, H.; Tanaka, S.; Ito, S.; Yamamoto, M. Nanometric inhomogeneity of polymer network investigated by scanning near-field optical microscopy. Macromolecules 2000, 33, 9650−9656. (29) Grimaud, T.; Matyjaszewski, K. Controlled/“living” radical polymerization of methyl methacrylate by atom transfer radical polymerization. Macromolecules 1997, 30, 2216−2218. (30) Aoki, H.; Kunai, Y.; Ito, S.; Yamada, H.; Matsushige, K. Twodimensional phase separation of block copolymer and homopolymer blend studied by scanning near-field optical microscopy. Appl. Surf. Sci. 2002, 188, 534−538. (31) Each PMMA-Pe chain was labeled by 150 perylene molecules. The value of (Ix − Iy)/(Ix + Iy) for a single chain was calculated as the averaged orientation for 150 bond vectors, which was randomly generated. The calculation was performed for 106 chains to construct the probability distribution function of (Ix − Iy)/(Ix + Iy). (32) Ohmori, S.; Ito, S.; Onogi, Y.; Nishijima, Y. Fluorescence method for studying surface orientation of polymer film using vacuumdeposition technique. Polym. J. 1987, 19, 1269−1278. (33) Rudnick, J.; Gaspari, G. The shapes of random-walks. Science 1987, 237, 384−389. (34) Bracewell, R. N. The Fourier Transform and Its Applications, 2nd ed.; McGraw-Hill: New York, 1986. (35) Brown, H. R.; Russell, T. P. Entanglements at polymer surfaces and interfaces. Macromolecules 1996, 29, 798−800. (36) Lee, H.; Paeng, K.; Swallen, S. F.; Ediger, M. D. Dye reorientation as a probe of stress-induced mobility in polymer glasses. J. Chem. Phys. 2008, 128, 134902. (37) Lee, H.; Paeng, K.; Swallen, S. F.; Ediger, M. D. Direct measurement of molecular mobility in actively deformed polymer glasses. Science 2009, 323, 231−234.

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