Chapter 13 Viscoelastic Measurements in Thin Polystyrene Melts as Derived from Scanning Force Microscopy-Induced Nanoflow Patterns 1
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Ronald H. Schmidt , Wayne L. Gladfelter , and Greg Haugstad
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Department of Chemistry and Center for Interfacial Engineering, University of Minnesota, Minneapolis, MN 55455
The response of a thin-film polystyrene melt to a raster-scanned SFM tip was investigated. At high temperatures the scanning process induced intricate pattern formation whose quantitative characteristics were compared as a function of tip trajectory, temperature, and scan rate. Analysis of the images revealed that the features are strongly correlated with the geometry of the raster scan pattern. The dependence of the patterns on temperature and scan rate was consistent with time-temperature superposition as described by the Williams-Landel-Ferry (WLF) equation. WLF analysis implies an increased glass transition temperature derived from elevated pressure beneath the tip. The latter provides an estimate of the radius of the affected film region near the tip.
Tremendous interest has arisen in recent years regarding the issues of chain mobility and glass transition at polymer surfaces (1-4). Keddie et al used ellipsometric measurements of thermal expansivity to study T of polystyrene on silicon as a function of film thickness (1). The observed T asymptotically approached the bulk value of 100°C as thickness was increased, but for thinner films (s 10's of nm), T was significantly depressed. Clearly, such effects are of critical importance in the application of ultrathin polymer films as lubricants and protective coatings in magnetic data storage devices and micromechanical systems (5). Until recently, however, direct mechanical measurements of polymer behavior at the nanoscopic scale (which is necessary for mechanical analysis of ultrathin films) and at elevated temperatures has not been possible. SFM has proven to be a powerful tool for measuring time-dependent mechanical properties of polymer films at ambient temperatures. Because piezoelectric scanners are incompatible with substantial temperature elevation, however, the SFM community has only begun to examine the role of temperature in g
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© 2000 American Chemical Society Tsukruk and Wahl; Microstructure and Microtribology of Polymer Surfaces ACS Symposium Series; American Chemical Society: Washington, DC, 1999.
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228 material response. Fundamentally, the response of "soft" materials to external forces is often dominated by entropie (i.e., temperature-dependent) effects (6,7). The approach to thermodynamic equilibrium for polymeric systems is typically slow, so that the response to mechanical perturbation is a function of both temperature and rate (i.e., is viscoelastic). Rigorous studies of thin-film polymer nanotribology therefore must include methodologies to quantify the interrelated roles of these two parameters.
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Methodology In order to access sample temperatures significantly different from ambient conditions without affecting the performance of the scanner, the SFM's piezo transducer and accompanying electronics must remain completely isolated from the sample. The SFM used for this study was designed and constructed by Molecular Imaging to meet this stringent requirement, thereby allowing us to resistively heat the sample. An aluminum-coated silicon cantilever, manufactured by Nanotechnology-MDT, with a spring constant of 3.8 N/m and a tip radius of fg results in rubbery or fluid-like behavior. As shown in Table \,fg = 0.032 for PS. Using this value offg, we fixed the value of Cj in order to estimate an effective glass transition temperature of 113°C from our measurements. In interpreting this result, we must consider two complications in SFM of ultrathin films: interfacial effects and elevated pressure. Several studies have demonstrated that as the thickness of a polystyrene film approaches the polymer coil dimensions, Tg decreases significantly below the bulk value of 100°C (1-3). In a system similar to what we investigated (i.e., 20 nm PS film on a Si substrate), Keddie et al measured a Tg of 87°C. On the other hand, substantial increases in Tg have been reported in bulk PS when the sample was subjected to elevated hydrostatic pressure (15). Assuming that the 26°C increase in Tg is exclusively due to pressure effects, we estimate— based on the effect of hydrostatic pressure on the Tg of bulk PS— that the effective tip pressure on the surface is 84 MPa (15). Given the effective pressure and the applied load (110 nN), we can calculate an effective perturbation radius of 20 nm, which is remarkably close to the tip's radius of curvature (approximately 10 nm). 0
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142 100 (bulk T (14))* 100 113 G
Q(°C) (2.303 C;)' 43 0.10 0.032 =fg* 51 * 7.6 0.014 14 0.032 1
Table 1. Comparison of WLF analysis to literature values for bulk polystyrene. * indicates literature values (14). Summary. Both temperature and scan frequency affect the patterns created by movement of a nanoscopic tip in contact with a polymer surface. At experimental time scales which are faster than the relaxation time associated with the elongation of a polymer coil, the polymer's response to the tip-induced shearing forces is elastic. Experimental time scales slower than the characteristic relaxation time results in alignment of the polymer with the tip trajectory and net translational movement of polymer towards the center of the scan area. The time-temperature dependence of the patterns are well-described by the WLF equation which is typically used to describe viscoelastic behavior. Analysis of our data further suggests that the Tg of the polymer is elevated in the region confined between the tip and the substrate. Acknowledgments. This work was supported by grants from the Center for Interfacial Engineering and the donors of the Petroleum Research Fund, administered
Tsukruk and Wahl; Microstructure and Microtribology of Polymer Surfaces ACS Symposium Series; American Chemical Society: Washington, DC, 1999.
238 by the American Chemical Society. We also wish to thank Molecular Imaging for their technical support and instrument development.
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Literature Cited (1) Keddie, J. L.; Jones, R. A. L.; Cory, R. A. Europhys. Lett. 1994, 27, 59-64. (2) Forrest, J. Α.; Dalnoki-Veress, K.; Stevens, J. R.; Dutcher, J. R. Phys. Rev. Lett. 1996, 77, 2002-2005. (3) Forrest, J. Α.; Dalnoki-Veress, K.; Dutcher, J. R. Phys. Rev. E. 1997, 56, 57055716. (4) Mayes, A. M. Macromolecules 1994, 27, 3114-3115. (5) Bhushan, B. Handbook ofMicro/Nanotribology; CRC Press: Boca Raton, 1995. (6) Tant, M . R.; Wilkes, G. L. Polymer Engineering and Science 1981, 21, 874-895. (7) Matsuoka, S. Polymer Engineering and Science 1981, 21, 907-921. (8) Leung, Ο. M.; Goh, M . C. Science 1992, 255, 64-67. (9) Lea, A. S.; Pungor, Α.; Hlady, V.; Andrade, J. D.; Herron, J. N.; Voss Jr., E. W. Langmuir 1992, 8, 68-73. (10) Meyers, G. F.; DeKoven, B. M.; Seitz, J. T. Langmuir 1992, 8, 2330-2335. (11) Elkaakour, Z.; Aime, J. P.; Bouhacina, T.; Odin, C.; Masuda, T. Phys. Rev. Lett. 1994, 73, 3231-3234. (12) Jing, J.; Henriksen, P. N.; Wang, H.; Marteny, P. J. Mat. Sci. 1995, 30, 57005704. (13) Woodland, D. D.; Unertl, W. N. Wear 1997, 203-204, 685-691. (14) Ferry, J. D. Viscoelastic Properties ofPolymers; Wiley & Sons: New York, 1980. (15) Stevens, J. R.; Coakley, R. W.; Chau, K. W.; Hunt, J. L. J. Chem. Phys. 1985, 84, 1006-1014.
Tsukruk and Wahl; Microstructure and Microtribology of Polymer Surfaces ACS Symposium Series; American Chemical Society: Washington, DC, 1999.
