Scan-Induced Patterning in Glassy Polymer Films: …Scan-Induced Patterning in Glassy Polymer Films:...

Scan-Induced Patterning in Glassy Polymer Films: UsingScanning Force Microscopy To Study Plastic Deformation

at the Nanometer Length Scale

Ronald H. Schmidt,† Greg Haugstad,*,‡ and Wayne L. Gladfelter*,†

Department of Chemistry, University of Minnesota, 207 Pleasant Street SE, andIT Characterization Facility, University of Minnesota, 100 Union Street SE,

Minneapolis, Minnesota 55455

Received December 27, 2001. In Final Form: October 17, 2002

The response of thin glassy films of polystyrene and poly(vinyl acetate) to a raster-scanned, sliding SFMtip was investigated. Several of the previously proposed mechanisms of the familiar scan-induced patternsare discussed. Increases in film volume and frictional response are quantified, and suggest that the observedtip-induced plastic deformation may relate to a second-order phase transition (glass-to-rubber) beneaththe sliding tip. Analysis of the scan-induced patterns suggests a crazing mechanism for the observedplastic deformation. The susceptibility of the film to plastic deformation was examined as a function ofscan geometry, applied load, and the gain of the feedback loop that maintains a constant applied load. Anempirical quantity called the roughening susceptibility is defined and shown to be linear with respect tovariations in the scan conditions. The roughening susceptibility is highly robust in quantifying thedependences on scan history and load. This finding will be further exploited in the second paper of theseries, analyzing rate and temperature dependences and their relationship to the glass transition.

IntroductionThis paper is concerned with the damage done to thin,

supported, glassy polymer films when they are scannedby a nanometer-scale tip sliding across the sample surface.Qualitatively similar images have been reported in theliterature for a wide variety of polymer/substrate sys-tems: polystyrene on mica,1 Si,2 glass,3 and metal;4 self-supported injection molded polycarbonate5-8 and poly-(methyl methacrylate)6 surfaces; a polystyrene/poly-(ethylene oxide) blend on mica;9 a protein (monoclonalIgM) on mica;10 poly(ethylene terephthalate) on mica;11

polyacetylene on mica and WSe2;12 polyimide on ITO-coated glass;13 PEO on mica;14 nylon on ITO-coated glass.15

While scan-induced patterning appears to be ubiquitouswith respect to the variety of polymer/substrate systemsfor which the phenomenon has been observed, a criticalreview of the literature reveals that relatively few, and

in some cases contradictory, quantitative studies havebeen published in the field. In the current paper weattempt to address this problem by quantifying severalaspects of the patterning process under a wide variety ofexperimental conditions. The methodology and phenom-enological approach presented herein will be applied tothe analysis of temperature- and rate-dependent pat-terning in a subsequent paper, where we will also discussthe issue of temperature-assisted plastic deformationusing more fundamental concepts. The scan-inducedpatterns considered in this work are to be differentiatedfrom the more intricate patterns reported for a tip scanningin contact with a polymer melt at temperatures higherthan the glass transition.16,17

Results are presented that characterize the scan-induced morphology as a function of sample loading andrasterscangeometry.Within theregionofparameterspaceinvestigated, the response of the film is highly linear withrespect to the loading conditions and the number ofsuccessive visits of the sliding tip to the same region ofthe sample surface. Although the exact mechanism ofdeformation within the polymer film remains unknown,evidence is presented that refutes several proposedmechanisms. We present frictional data that show thatthe response of the patterned region is consistent with a“generalized” stick-slip mechanism related to a second-order phase transition in the polymer beneath the slidingtip. It is also shown that the scanning process inducesdamage in polystyrene films analogous to the familiarphenomenon of crazing in brittle bulk polymer systems.

Experimental SectionSolutions of monodisperse atactic polystyrene (PS; Aldrich)

and high-dispersity atactic poly(vinyl acetate) (PVAc; Aldrich)were prepared by dissolving the polymer in toluene. The solutions

* To whom correspondence should be addressed.† Department of Chemistry.‡ IT Characterization Facility.(1) Leung, O. M.; Goh, C. Science 1992, 255, 64-66.(2) Meyers, G. F.; DeKoven, B. M.; Seitz, J. T. Langmuir 1992, 8,

2330-2335.(3) Woodland, D. D.; Unertl, W. N. Wear 1997, 203-204, 685-691.(4) Pickering, J. P.; Vancso, G. J. Appl. Surf. Sci. 1999, 148, 147-

154.(5) Hamada, E.; Kaneko, R. J. Phys. D 1992, 25, A53-A56.(6) Kaneko, R.; Hamada, E. Wear 1993, 162-164, 370-377.(7) Iwata, F.; Matsumoto, T.; Sasaki, A. Nanotechnology 2000, 11,

10-15.(8) Khurshudov, A.; Kata, K. Wear 1997, 205, 1-10.(9) Nie, H. Y.; Motomatsu, M.; Mizutani, W.; Tokumoto, H. J. Vac.

Sci. Technol., B 1995, 13, 1163-1166.(10) Lea, A. S.; Pungor, A.; Hlady, V.; Andrade, J. D.; Herron, J. N.;

Voss, E. W., Jr. Langmuir 1992, 8, 68-73.(11) Jing, J.; Henriksen, P. N.; Wang, H.; Marteny, P. J. Mater. Sci.

1995, 30, 5700-5704.(12) Elkaakour, Z.; Aime, J. P.; Bouhacina, T.; Odin, C.; Masuda, T.

Phys. Rev. Lett. 1994, 73, 3231-3234.(13) Pidduck, A. J.; Haslam, S. D.; Bryan-Brown, G. P.; Bannister,

R.; Kitely, I. D. Appl. Phys. Lett. 1997, 71, 2907-2909.(14) Nie, H. Y.; Motomatsu, M.; Mizutani, W.; Tokumoto, H. J. Vac.

Sci. Technol., B 1997, 15, 1388-1393.(15) Ruetschi, M.; Grutter, P.; Funfschilling, J.; Guntherodt, H. J.

Science 1994, 265, 512-514.

(16) Schmidt, R. H.; Haugstad, G.; Gladfelter, W. L. Langmuir 1999,15, 317-321.

(17) Schmidt, R. H.; Haugstad, G.; Gladfelter, W. L. In Microstructureand Microtribology of Polymer Surfaces; Tsukruk, V. V., Wahl, K. J.,Eds.; ACS Symposium Series, Vol. 741; American Chemical Society:Washington, DC, 2000; pp 227-238.

898 Langmuir 2003, 19, 898-909

10.1021/la015769j CCC: $25.00 © 2003 American Chemical SocietyPublished on Web 01/08/2003

were filtered twice through 0.2 µm porosity syringe filters intoclean scintillation vials. Silicon(100) wafers (4 in. diameter) wereused as substrates. Other than using a nitrogen jet to removedust from the wafers, no special procedure was used to treat thesubstrates prior to spin coating the films. As a result, the siliconsubstrate is covered by a thin (0.6-2 nm), amorphous layer ofSiO2.18 The films were prepared by spin coating the solutionsonto the wafers. By varying the solution concentration in therange of 0.25-2% w/v and the spin rate from 1000 to 5000 rpm,films could be prepared whose thickness ranged from ∼3 to 150nm. Films thinner than 2-3 nm were discontinuous, while filmsmuch thicker than 150 nm generally contained fairly long scale(∼1 mm) lateral variations in thickness that were visible to thenaked eye as optical fringes. Many of the films investigated wereunannealed, so that although they were several days to severalweeks old at the time that the measurements were made, theymay possess anisotropy related to spin coating and/or containtrace amounts of plasticizing solvent.

Data were collected at ambient conditions using a DigitalInstruments Multimode SPM with a 150 µm scanner and aMolecular Imaging PicoSPM with scanners ranging from 30 to50 µm. A Nanoscope III controller and Nanoscope system softwarewere used to control both SPMs and to acquire all data. Exceptwhere otherwise stated, all imaging was performed using afeedback loop to maintain a constant vertical deflection of thecantilever (and hence a constant applied load) as the tip wasscanned across the sample surface. Gold-coated Si3N4 cantilevers,supplied by Digital Instruments, (knorm ) 0.58 N/m) were usedfor imaging. The trajectory of the sliding tip is described in detailelsewhere.16,19 Briefly, the tip is scanned in a sawtooth patternat an angle slightly offset from the horizontal (“fast-scan”) axisof the images. A total of 128, 256, or 512 fast-scan cycles (i.e., aleft-to-right scan followed by a right-to-left scan) can be acquiredalong the vertical (“slow-scan”) axis of the imaging window. Thefast-scan frequency refers to the number of fast-scan cyclesperformed per second.

Measuring Film Thickness with AFM. Various techniques(X-ray reflectivity, ellipsometry, RBS/FRES) were employed todetermine film thickness. All of these methods yield thicknessmeasurements averaged over macroscopic lateral dimensions. Amore convenient method of determining film thickness, and onethat measures the thickness at relatively small lateral lengthscales, is described below.

A small region of the film (typically 0.1 × 0.1 to 1 × 1 µm) wasscanned at a high load (∼25-100 nN) to sweep the polymer outof the scanned region. Multiple raster scans were usuallyperformed, and the scan angle was switched periodically from0° to 90°, until the scanned region was observed to have a flattopography. The applied load was then reduced to a much smallervalue (typically 2.5 nN), and a larger region of the film wasimaged, as shown in Figure 1. (In some cases friction imageswere collected at high temperature and/or low scan velocity,producing material contrast between the undisplaced film andthe bared substrate.) The thickness of the film was determinedfrom a horizontal cross section of the image. In the cross sectionshown, the film thickness was taken to be the depth approachedasymptotically as the tip scanned from left to right. Thediscrepancy between this steady-state value of the depth andthe maximum apparent depth (at the left edge of the hole) isattributed to the piezo creep effect.20 Holes with smaller lateraldimensions therefore require slower scan velocities to ensurethat the steady-state depth is reached before the tip encountersthe far edge of the hole, thereby minimizing the errors associatedwith piezo creep.

For low molar mass (13 kg/mol) ultrathin PS films (i.e., filmsthinner than ∼15 nm), holes that reach the substrate can beproduced at ambient temperature. The film shown in Figure 1had a molar mass of 212 kg/mol and was presumably tougherdue to the presence of entanglements; consequently, a slightlyhigher temperature (50 °C) was required to excavate the polymer.Thicker films (>50 nm) generally require temperatures near the

bulk glass transition to produce a hole that reaches the substrate.Sample heating was accomplished using a copper resistiveheating stage attached to the Molecular Imaging PicoSPM.

Results

Qualitative Description of Scan-Induced Defor-mations. Figure 2 shows several series of 1 × 1 µm imagesof an 8.5 nm, 212 kg/mol PS film acquired at roomtemperature. Each column represents a single series ofimages, and the applied load is higher moving from leftto right. During the acquisition of a given series of images,the tip was allowed to raster repeatedly over the sameregion, with 256 fast-scan cycles performed per raster scan.To eliminate sample tilt and the curvature associated withthe pendulum arc described by the sweeping piezo tube,a third-order polynomial was fit to each scan line of data,and the fit was then subtracted from the raw data to givethe flattened images shown in the figure. (For the sakeof consistency and simplicity, we used the same procedurefor flattening all images, regardless of image size or themicroscope used for acquisition.) The gray scale representstopography, with the darkest colors associated with thelowest features and the brightest colors corresponding tothe highest features. The difference in height betweenthe brightest and darkest features is 5 nm.

Initially relatively small, isotropic height fluctuationsare present, as in all amorphous polymer surfaces. As

(18) Okorn-Schmidt, H. F. IBM J. Res. Dev. 1999, 43, 351-365.(19) Command Reference ManualsVersion 4.10; Digital Instruments,

Inc.: Santa Barbara, CA 93103.(20) The Piezo Book; Burleigh Instruments, Inc.: Fishers, NY 14453-

0755.

Figure 1. Top: topography image of PS on silicon. The centralregion was scanned with a high load, resulting in completeremoval of the film. Bottom: The topographical cross sectioncorresponds to the dashed line in the image above. Due to piezocreep, the depth at the leading edge of the hole is overestimated;the true depth of the hole (which corresponds to the filmthickness) is approached asymptotically.

Scan-Induced Patterning in Glassy Polymer Films Langmuir, Vol. 19, No. 3, 2003 899

more scans are performed over the same region, thesurface morphology changes. Bundles and trenches thatare oriented perpendicular to the tip trajectory begin toform. The height and depth of the bundles and trenches,respectively, grow with the number of raster scansperformed. At higher applied loads, fewer scans aregenerally required to produce a given amplitude of thesesurface waves. Close inspection (compare scans 4 and 16taken under 30 nN applied load) also reveals that, withrepeated scanning, the height fluctuations are more inphase, or better correlated, along the vertical axis.

Quantifying the Frictional Response in the Pat-terned Region. In Figure 3, we consider the simulta-neously acquired lateral force and topography imagescorresponding to scan 16 at 12 nN applied load from Figure2. Topography and lateral force data (which have beenoffset in the figure by an amount equal to the mean heightand mean lateral force, respectively) are shown for thesame scan line in the image. Waves of similar repeatdistance are apparent in both the topography and lateralforce traces. It is known that the lateral force between ascanning tip and a sample surface consists of two terms:a friction (dissipative) force as well as a component thatis proportional to the topographical gradient along thefast-scan axis (dZ/dX, provided that dZ/dX is not sub-

stantially larger than 0.3).21,22 To determine whether theobserved fluctuations in the lateral force image are relatedto the topographical gradient, the lateral force data areplotted as a function of both Z and dZ/dX in Figure 4.While there appears to be little correlation between Z andlateral force, a fairly strong correlation is evident betweendZ/dX and lateral force. It has been shown that thecontribution of dZ/dX to the lateral force can be removedapproximately by subtracting the lateral force dataacquired during the right-to-left (retrace) scan from thedata acquired during the left-to-right (trace) scan.21 Onealso must be careful to account for the hysteresis in thepiezo scanner.20 For the features acquired during the traceand retrace scans to properly line up, it was necessary toshift the retrace friction data 6 pixels to the right (adistance equal to 58 nm for the 512 pixel, 5 µm scan lines).When the shifted retrace image is subtracted from thetrace image, we are left with an image in which the

(21) Haugstad, G.; Gladfelter, W. L.; Weberg, E. B.; Weberg, R. T.;Weatherill, T. D. Langmuir 1994, 10, 4295-4306.

(22) An additional term that is proportional to (dZ/dX)2 becomessignificant for dZ/dX > 0.3. If the pattern shown in Figure 3 isapproximated as a sinusoidal wave of amplitude 2.5 nm and repeatdistance 70 nm, |dZ/dX| has a maximum value of 0.2. Thus, for this casethe quadratic term can be neglected.

Figure 2. Topography images of 1 × 1 µm acquired with an 8.5 nm film for several applied loads. Continuing to scan over thesame region roughens the surface, and the effect is stronger for higher loads.

900 Langmuir, Vol. 19, No. 3, 2003 Schmidt et al.

contribution of dZ/dX to the lateral force has beeneffectively removed, so that remaining variations in theimage correspond instead to intrinsic variations in thetip-sample friction force. Evidence of a residual alterationin the intrinsic frictional response of the polymer surfaceis presented in Figure 5. The data were taken with a Si3N4tip sliding on a 10 nm thick PVAc film cast on a siliconwafer at ambient temperature. The central region un-derwent 10 512-line raster scans with a 23 nN appliedload and a fast-scan frequency of 11 Hz. The scan size wasthen increased to 5 µm and the scan frequency wasdecreased to 3 Hz to image the friction response of thepatterned region. Eight images were acquired withdifferent applied loads, ranging from 23 nN during thefirst scan to 3 nN, and the scan size was increased againto 10 × 10 µm. In the 10 × 10 µm scan shown in the figure,very little modification is visible in the topographical data,which may be attributed to the 5 × 5 µm scans, and no“friction box” 23 is visible within the central 5×5 µm region.The central, patterned region (region B) clearly exhibitsfluctuations in tip-sample frictional force; these fluctua-tions are presumably not random, but they could not beclearly correlated with either Z or dZ/dX. We turn ourattention now to the distribution of friction response inthe image. The histograms shown in Figure 6 reveal asingle Gaussian peak in region A and two partiallyoverlapping peaks in the friction response of region B.The histogram corresponding to region B was fit to adouble-Gaussian distribution function using a nonlinearleast-squares curve-fitting routine, and the resulting fitis shown in the figure (the solid line corresponds to thedouble-Gaussian fit; the two dashed lines correspond tothe individual Gaussian curves which comprise that fit).The mean values derived from the Gaussian fits are plottedin Figure 6 as a function of the applied load.

The friction response within region B is interpreted asbeing partitioned between two states, which tentativelywill be termed type I friction and type II friction. (Laterwe will argue that these can be assigned to less/morerubbery domains, respectively.) Inspection of the frictionvs load plot (Figure 6) reveals that the coefficient of frictionfor the less dissipative component (type I) closely matchesthe coefficient of friction observed in the unperturbedregion A, with the remaining pixels exhibiting a largercoefficient of friction (type II). Two islands of high frictionare apparent in the 5 × 5 µm scan in Figure 5. Analysisof their friction distributions was difficult due to theirsmall size and the presence of noisy scan lines, but it wasclear that, with repeated scanning, these regions graduallyreturned to exhibiting solely type I friction. The high-friction islands are no longer visible in the 10 × 10 µmfriction image.

The first study of alterations in the frictional responseof polymer films caused by a sliding tip was reported byHaugstad et al. for gelatin films cast on mica.23 When asmall region of the film was scanned and a larger imagewas subsequently taken, the initially scanned regionexhibited higher friction than the surrounding, unper-turbed region. Unlike the phenomenon presented in thisstudy, no change was apparent in the topographical image.Eventually, with repeated scanning, the frictional re-sponse of the inner region returned to its initial state (i.e.,to the state of the surrounding region). The tendency to

(23) Haugstad, G.; Gladfelter, W. L.; Weberg, E. B.; Weberg, R. T.;Jones, R. R. Langmuir 1995, 11, 3473-3482.

Figure 3. Top: topography and lateral force images corre-sponding to scan 16 at 12 nN applied load from Figure 2. Bottom:topography and friction sections along a single horizontal scanline. Both sections were acquired in the trace (left-to-right)scan direction.

Figure 4. Top: A poor correlation is observed between thelateral force data in Figure 3 and the height variations.Bottom: The lateral force is strongly correlated with thegradient of the topography.


create these friction boxes, and the tendency for them torelax back to the original state during subsequent imaging,exhibited a strong dependence on the applied load andthe velocity of the scanning tip. Haugstad et al. postulatedthat the effect of the tip-sample interaction was to placethe region initially scanned into a more dissipative state,effectively moving that subsystem toward the glasstransition and hence toward the maximum in the vis-coelastic loss tangent. The gradual return of the gelatinto its initial state corresponds to a relaxation of the morehighly dissipative conformers present within the sub-system to the equilibrium conditions set by kT.

To examine whether a relaxation of the type II frictioncomponent occurred within region B, we consider thefraction of the pixels that are present in the two states asdefined by the following relationships:

where AI and AII are the areas under the single Gaussiancurves as exemplified in Figure 6. The friction force wasdistributed approximately evenly between type I and typeII friction (i.e., xI ) xII ) 0.5) for the 5 µm image shownin Figure 5, and this distribution remained unchangedwith repeated scanning of the polymer within region B.This suggests that the high-friction moieties in region Bare fundamentally different from the friction boxes seenby Haugstad et al. or from the high-friction islands visiblein region A of the 5 × 5 µm scan, which relaxed to typeI friction with repeated scanning.

VolumeChangeinPatternedRegions. Hamada andKaneko have suggested that scan-induced deformationsare accompanied by a volume increase which they at-tributed to the formation of voids within the film.5 Theevidence supporting this conclusion is based on an analysisof a single scan line in a scan-scratched injection-moldedpolycarbonate surface. A more extensive and quantitative

analysis of the volume increase that accompanies bundleformation is described below. After a series of scans wereperformed over the same region with a relatively highapplied load, the contact force was reduced and a largerregion was scanned, as shown in Figure 7. Region Acorresponds to the pattern shown in Figure 2 acquiredwith a 30 nN applied load. The scan size was increasedto a 3 × 3 µm region after the original 1 × 1 µm scans were

Figure 5. Top: 5 µm topography and friction images of a PVAcfilm on silicon. The friction image corresponds to the difference(including sign) of the trace and retrace lateral force images.Bottom: After the final 5 × 5 µm images were acquired, thescan size was increased to 10 µm. No alteration in the surfacewhich can be attributed to the 5 µm scans is visible in theseimages.

xI )AI

AI + AIIand xII ) 1 - xI (1)

Figure 6. Histograms of the friction data reveal a singleGaussian peak in region A (top) and two poorly resolved peaksin region B (middle). Bottom: friction vs load data. The valuesplotted correspond to the mean friction derived from theGaussian fit as shown by the dashed lines in Figure 5.


performed. Several scans were performed at variousapplied loads (ranging from 1.5 to 23 nN), and the scansize was increased again to 5 × 5 µm with an applied loadof 9 nN used to acquire the image shown. This processgenerated topographical “boxes”, in which the regionsaltered by the scanning tip are clearly visible as well-defined regions of brighter color compared to the lessperturbed surroundings. It is expected that the scanconditions used to acquire the final image that is shownwill also eventually alter the sample surface. Because theapplied load was greatly diminished and the distancebetween successive visits by the tip along the slow-scanaxis is much larger than what was used to generate theoriginal surface modification, this is expected to be agradual process requiring many raster scans. (Evidencewill be presented later that justifies this statement.) Thus,we assume that the outer region of the image is unalteredby the scanning tip. Topographical cross sections of theimage are shown in the figure. The top section correspondsto a single horizontal scan line drawn through the middleof the image. As is evident in the image, the heightfluctuations are largest in the center, smallest in the outerregion, and intermediate in the intervening region. Theseheight fluctuations, which were also seen in the topo-graphical cross sections reported by Hamada and Kaneko,5make it difficult to compare the mean height of thepatterned region to that of the surrounding unperturbedregion. We found that these height fluctuations are easilyremoved by examining the arithmetic mean of many crosssections. In the middle graph, the mean of approximately50 horizontal cross sections that intersected the center1 × 1 µm region is plotted. While the height fluctuationsin the outer region and the central 3 × 3 µm region arelargely removed by averaging, the central 1 × 1 µm regionstill shows large-scale fluctuations in the height. Thesefluctuations are largest at the left and right edges of theregion and tend to decay in amplitude toward the centerof the region. The ability of the scan-induced heightfluctuations to cancel out depends on a relatively weakcorrelation (or poor alignment) of the bundles and trenches

between adjacent scan lines. Thus, the middle graphdemonstrates that the surface waves are relatively poorlycorrelated in region B and better correlated at the edgesof region A than they are toward the center of that region.

A more useful way to examine the mean height is toaverage the cross sections along the vertical axis, as shownin the bottom graph in Figure 7. The height fluctuationscorresponding to region A are nearly eliminated, and wenow can clearly recognize that the mean topography isgreater in region B than it is in region C, and that thehighest mean topography is present in region A. Thus, wehave demonstrated that the tip-induced modifications sofar considered have a tendency to raise the average heightof the sample surface.

Repeat Distance between Bundles. Figure 8 (top)shows the mean one-dimensional (i.e., measured alongthe X-axis, or perpendicular to the bundle orientation)power spectra of the images in the rightmost column ofFigure 2. The spectra are offset along the vertical axis forclarity. The peaks in the spectra correspond to funda-mental and higher harmonics of the “wavelength” or repeatdistance of the scan-induced pattern. A plot of λ vs thenumber of raster scans24 reveals that, under the scanconditions considered, the repeat distance reaches a nearlysteady state value after only four raster scans (bottom ofFigure 8).

Discussion

Roughening Susceptibility. The RMS roughness ofthe images shown in Figure 2 are plotted as a function ofthe number of raster scans in Figure 9. The roughness,or equivalently the wave amplitude, grows linearly for allof the scan conditions shown. The slope of the plots ofroughness vs the number of scan lines will hereafter becalled the roughening susceptibility. The rougheningsusceptibility was observed to increase monotonically with

(24) Because the peaks observed in the power spectra are broad andtherefore yield a determination of λ lacking precision, the λ values plottedin the bottom of Figure 8 were instead determined in real space.

Figure 7. Left: 5 × 5 µm topography and friction images. After sixteen 1 µm scans were performed in the center of the image,the scan size was increased to 3 µm for several scans prior to the acquisition of the images shown. The previously scanned regionsexhibit higher topography and friction. Right: Cross sections of the topography image are used to quantify the mean heightdifference between the patterned region and the surroundings. The best results are obtained from the average of many verticallines.


the applied load. The data in Figures 2 and 9 correspondto images acquired with 256 fast-scan cycles per image;similar measurements were also performed on imagescollected with 128 and 512 fast-scan cycles per raster scan.The results from all of these runs are compiled in Figure10. To compare roughening susceptibilities collected withdifferent numbers of fast-scan cycles, we divide theroughening susceptibilities plotted in Figure 10 (top) by

the number of scan cycles per image and replot all of theresults in Figure 10 (bottom). We see reasonable agree-ment between these normalized roughening susceptibili-ties acquired with different scan line densities. Thus, weconclude that the roughening susceptibility dependslinearly on the scan line density, as well as the appliedload. This is to our knowledge the first quantitative reportof these fundamental phenomenological relationships.

Figure 11 demonstrates the reproducibility of thisrescaling result. In two trials, thirty raster scans wereperformed over the same region, but the scan line densitywas changed after scan 10 and returned to the initial valueafter scan 20. The roughness plotted against the numberof raster scans is continuous, but the slope of the dataclearly changes abruptly when the scan line density ischanged (top of the figure). By plotting the roughnessagainst the total number of scan lines (i.e., the numberof raster scans times the scan line density), all of the datacollapse onto a single straight line (bottom of the figure).This is further confirmation that, for a given set of loadingconditions, the roughening susceptibility appears todepend only on the number of visits by the scanning tip.It further suggests the robustness of this parameter forcross-study comparison.

Figure 8. Top: power spectral densities for the imagescorresponding to the right-most column of Figure 2. The spectraare offset along the vertical axis for brevity. Bottom: repeatdistance vs the number of scans for the same set of images.

Figure 9. RMS roughness vs the number of scans for thevarious loads shown in Figure 2. The solids lines representlinear fits to the data. The slopes of such fits define a parametercalled the roughening susceptibility. Note the monotonicincrease in roughening susceptibility with applied load.

Figure 10. Top: roughening susceptibilities plotted as afunction of applied load for three different scan line densities.Note that the roughening susceptibility tends to increase withboth the applied load and the scan line density. The solid linesrepresent linear fits forced through the origin. Bottom: Thesame data are rescaled by the scan line density. All of the datacollapse onto a single line.


Another parameter that was investigated was the gainsetting of the feedback loop, which controls how effectivelythe feedback circuit responds to changes in the sampletopography and sample tilt. At high gain settings, thefeedback circuit is fully engaged, and the applied load isapproximately constant across the entire scanned area.At lower gain settings, the tip remains in contact andoverall sample tilt is accommodated, but regions of steeplyincreasing height (which occur at the leading edge of thebundles) experience larger cantileverdeflectionsandhencehigher applied loads, while regions of steeply decreasingheight (at the trailing edge of the bundles) experiencesmaller applied loads. The average applied load acrossthe imaged region, however, remains nearly constant. Theeffect of varying the feedback gain setting during theacquisition of a series of images is shown in Figure 12.Images 1-10 were acquired with the feedback loop fullyengaged (integral gain19 2), scans 11-20 were performedwith the feedback loop nearly disabled (integral gain 0.25),and scans 21-30 were collected with the feedback fullyreengaged. The apparent drop in the RMS roughnessduring scans 11-20 is an artifact since the signal used togenerate the topography image is based on the voltage tothe Z-piezo, which in turn is dependent on the output ofthe feedback circuit. When the feedback circuit is reen-gaged (scans 21-30), we see that the RMS roughness is

precisely as large as expected from a linear extrapolationof the first 10 scan cycles. This result is remarkable inlight of the fact that, during the intermediate scans forwhich the feedback loop was nearly disabled, the appliedload was not uniform across a given peak or trough. Thus,we have further confirmation of the linear relationshipbetween surface roughening and the average applied loadacross the scanned area, as well as the cross-studyrobustness of the bundle formation phenomenon.

Mechanisms of Tip-Induced Deformations. Schal-lamach Waves. Meyers2 noted a similarity in theappearance of the scan-induced morphology in thick (∼2µm) PS films to the buckling phenomenon that occurs atthe surface of an elastomer when it is slid across a second,hard surface under a compressive load (i.e., “Schallamachwaves”). At the macroscopic scale, Schallamach waveshave been observed for both a hard sphere sliding on arubber track and a rubber sphere sliding on a hard track.25

The buckling is attributed to “waves of detachment” drivenby a tangential stress gradient along the contact zone ofthe sliding interface. Energy loss occurs during thepropagation of the waves of detachment, which Schal-lamach attributed to the breaking of adhesive bondsbetween the two surfaces. The buckling of the rubbersurface, however, is reversible; when the shear forces areremoved from the rubber slider, the surface springs backto its original, smooth state.

Meyers concluded that the surface of a PS film was notglassy, but rather that it behaved as an elastomer.However, contrary to what would be expected for anelastically deformed surface, the bundles were not ob-served to relax to the initial unperturbed state when thetip was withdrawn and the sample was left unperturbedfor up to several hours. Furthermore, it should berecognized that the spacing of Schallamach’s waves ofdetachment is an order of magnitude smaller than thesize of the diameter of contact between the sphere andtrack, whereas the repeat distance of scan-induced pat-terns is typically much larger than the radius of curvatureof an AFM tip. Meyers attempted to compare the loaddependence of wave spacing for his data to those ofSchallamach waves. However, Schallamach’s data were

(25) Schallamach, A. Wear 1971, 17, 301-312.

Figure 11. Top: RMS roughness plotted against the numberof scans for two trials in which the scan line density was changedafter scan 10 and returned to its initial value after scan 20.Note the abrupt change in slope. Bottom: the same data rescaledalong the horizontal axis by the total number of fast-scan cyclesperformed. All of the data collapse onto a single line.

Figure 12. RMS roughness plotted against the number of scansfor an experiment in which the feedback was nearly disabledafter scan 10 and reenabled after scan 20. Note the abruptchanges in the apparent surface roughness. The line representsa linear fit to the solid circles.


taken at loads more than 7 orders of magnitude largerthan those used in the AFM experiment, and Meyers failedto take into account the relative sizes of the contact areasfor the two experiments. It is likely that the stresses (i.e.,load divided by area) in the AFM experiment are muchlarger than those experienced by the sphere and track inSchallamach’s work. In summary, despite the superficialresemblance of the scan-induced patterns to Schallamachwaves, we do not consider this model a valid explanation.

Mechanisms of Tip-Induced Deformations. Stick-Slip Behavior. Several authors have postulated someform of stick-slip behavior at the tip-polymer interfaceto explain the repeating nature of the patterns.10,12 It isassumed that a mound of polymer accumulates in advanceof the scanning tip as it slides across the surface. Thepresence of this mound results in an increased lateralforce on the tip whose vector points in opposition to thesliding direction, and as the mound grows larger, so doesthe lateral force. Eventually, this lateral force overcomesthe tip-sample adhesive interaction, and the tip slipsover the mound and begins to form a new one. The extentof the mound along the slow-scan axis is finite, andtherefore, the positions of the mounds along adjacent scanlines are correlated; the net result is that the observedmounds appear to be fairly continuous and in-phase alongthe slow-scan axis.

The phase transition’s model is one of two knownmechanisms for intrinsic stick-slip behavior.26-29 It wasdeveloped to account for the stick-slip friction observedwhen simple isotropic fluids are confined between twocrystalline surfaces. Simulations have suggested thatwhen the sliding plates are very closely spaced so that thethickness of the interfacial film approaches the dimensionsof the molecular length scale, the film undergoes a first-order phase transition between crystalline and liquidlikestates as the plates shear.26,27 While the interfacial filmis solidlike, the two crystalline plates cannot slide pastone another. Eventually, as shear energy is accumulatedwithin the interfacial region, the film melts, allowing theplates to slip. The film then freezes again, and the platesremain stuck until the next melting event. As the slidingvelocity of the plates is increased, the time betweensuccessive stick-slip events decreases, and the magnitudeof the stick-slip spikes decreases as well. Thompson andRobbins explained this common observation of decreasingcoefficient of friction with velocity by postulating thatvarious moieties within the interfacial region can possessone of two values: a “static value” whose coefficient offriction is given by Fs and a “kinetic value” with a coefficientof friction given by Fk, where Fk < Fs.26 As the slidingvelocity is increased, the number of kinetic eventsincreases at the expense of the number of static events.Because the overall observed friction is the sum of allfrictional events (some of which stimulate the static statesand some of which stimulate the kinetic ones), the overallobserved friction decreases with velocity. Eventually acritical velocity is reached, above which no stick-slipoccurs. A semiempirical relationship that describesthis critical velocity has been given by Yoshizawa andIsraelachivili:30,31

where Fs and Fk are the static and kinetic coefficients offriction, respectively,

k is the cantilever spring constant, andτ is the characteristic time for molecular rearrange-

ments.We turn our attention now to the measured friction

between the sliding AFM tip and the polymer bundles todetermine whether evidence of stick-slip behavior canbe found.

The strong correlation between lateral force and dZ/dX(Figure 4) implies that tip-sample slippage, at least atlength scales corresponding to the pattern’s repeatdistance, does not appear to play a role in the observedmorphology of the scan-induced pattern. Before stick-slip behavior at the tip-polymer interface is dismissed,however, it must be recognized that the mechanism ofstick-slip proposed by Thompson and Robbins dependson the ability of the interfacial region to undergo a first-order phase transition (i.e., melting-freezing). Amorphouspolymers do not undergo first-order phase transitions,but instead, experience more gradual changes in theirthermodynamic state functions (i.e., second-order phasetransitions32). For example, rather than experiencing alatent heat or an abrupt volume change (both of whichrepresent discontinuities in the first derivative of thechemical potential) as a polymer is heated or cooledthrough the glass transition, changes in the heat capacityand changes in the thermal expansivity (i.e., disconti-nuities in the second derivative of the chemical potential)are observed. Continuing with this theme, we postulatean analogous gradual change in the population of con-formational states with varying degrees of energy dis-sipation within the interfacial region. Rather than ex-periencing an instantaneous jump from stick to slipbehavior, one might instead expect a continuous change(temporal or spatial) of the frictional response as thepolymer beneath the tip moves between domains of glassy,rubbery, or transitional dynamics. Thus, one might assignthe dissimilar bundle and trough friction states in Figure6 to more rubbery or glassy dynamics. A detailed exami-nation of the temperature and/or rate dependence offriction would be required to locate precisely the differentfriction states along the transition from glassy to rub-bery.33,34

Mechanisms of Tip-Induced Deformations. A Frac-ture Mechanics Description. Perhaps the most well-developed proposed mechanism to date of scan-inducedpatterns in glassy polymer films comes from Elkaakouret al.,12 who studied the formation of bundles whilecontinually scanning the tip back and forth over the sameline (i.e., in slow-scan-disabled mode) on a thick (∼14 µm)film of polyacetylene on mica. The process was describedin terms of a crack propagating within the polymer filmin the region directly in front of the tip. Growth of thecrack is controlled by an energy balance similar to Griffith’scriterion, in which the release rate of the strain energyaccumulated in advance of the scanning tip is balancedby the increased surface energy associated with the walls

(26) Thompson, P.; Robbins, M. Science 1990, 250, 792-794.(27) Robbins, M. O.; Thompson, P. A. Science 1991, 253, 916.(28) Israelachvili, J. N.; Gee, M. L.; McGuiggan, P.; Thompson, P.;

Robbins, M. In Dynamics in Small Confining Systems; Drake, J. M.,Klfafter, J., Kopelman, R., Eds.; MRS Publications: 1990; pp 3-6.

(29) Israelachvili, J. N.; McGuiggan, P.; Gee, M.; Homola, A.; Robbins,M.; Thompson, P. J. Phys.: Condens. Matter 1990, 2, SA89-SA98.

(30) Yoshizawa, H.; Israelachvili, J. N. J. Phys. Chem. 1993, 97,11300-11313.

(31) Yoshizawa, H.; Chen, Y. L.; Israelachvili, J. J. Phys. Chem. 1993,97, 4128-4140.

(32) Hiemenz, P. C. Polymer Chemistry; Marcel Dekker: New York,1984.

(33) Hammerschmidt, J. A.; Moasser, B.; Gladfelter, W. L.; Haugstad,G.; Jones, R. R. Macromolecules 1996, 29, 8996-8998.

(34) Hammerschmidt, J. A.; Haugstad, G.; Gladfelter, W. L. Mac-romolecules 1999, 32, 3360.

Vc )Fs - Fk

5kτ(2)


of the crack. Elkaakour invoked the notion of slippage toexplain the periodic nature of the patterning process. Toour knowledge, no one has published direct evidence ofslippage between the scanning tip and the polymer surface.However, as discussed earlier, analysis of the height ofthe patterned regions suggests that the scanning processinduces voids in the polymer film. Using a technique calledultrasonic force microscopy (the tip was driven at afrequency of 2 MHz with an amplitude less than 1 nm),Iwata et al. compared the elastic contact stiffness of scan-induced bundles in an injection-molded polycarbonatesurface to that of the unperturbed surroundings.7 Iwataand co-workers found that the bundles were less stiff thanthe undamaged surface, which they also interpreted asevidence of the presence of microvoids or cracks in thedamaged region. The higher degree of energy dissipationin the patterned regions of the present study (i.e., thetype II friction seen in Figures 5 and 6) may be attributedto the greater mobility of the polymer molecules in thevicinity of such voids.

Our findings are consistent with the mechanism ofElkaakour et al.,12 in that the formation of cracks withina film is expected to increase the measured volume of theperturbed region, which with the assumption of incom-pressibility of the film is a consequence of mass conserva-tion. (Another possibility is delamination between the filmand the substrate.)

A series of measurements similar to those shown inFigure 7 were performed for two different PS films ofknown thickness (8.5 and 17 nm) and with a molar massof 220 kg/mol. Regions of 3×3 µm were scanned repeatedlywith applied loads of 5-30 nN, and the scan size wasincreased to 9 × 9 µm after the load was decreased to 5nN. Only one scan was performed at the latter conditions,and afterward, the size of the imaged region was increasedyet again to confirm that the previous scan had not alteredthe sample surface (i.e., that the area corresponding toregion B in Figure 7 was unaltered and therefore did notexhibit a step in height relative to region C). The heightof the steps was observed to correlate with the RMSroughness of the central perturbed region, as shown inFigure 13. A discussion of this correlation follows.

Inspection of the topographical cross section shown inFigure 3 suggests that the scan-induced bundles can be

described approximately by a sinusoidal wave of amplitudeA and repeat distance λ. Given that the mean height ofthe images is zero, the RMS roughness is given by thestandard deviation in height19 determined along m periodsof a sinusoidal wave:

Thus, assuming that the scan-induced height variationsare much larger than intrinsic height fluctuations (whichare presumably still present in the scan-induced patterns),the amplitude of the scan-induced waves is equal to x2times the RMS roughness of the flattened image.

We will consider now several hypothetical models forthe cross-sectional morphology of the bundles. Model A inFigure 14 describes the perturbed region as a sinusoidalfluctuation in the film thickness which, in the absence ofvoids, results in a mean thickness equal to that of theoriginal unperturbed film. In this model, the scanningprocess simply moves material around the surface, withno rupture within the film or at the film-substrateinterface; no increase in average surface height occurs.On the basis of the preceding findings, this model can bedismissed. Model B in Figure 14 explains the scanningprocess as the delamination of the film from the substrate.Assuming that the thickness of the buckled film in theperturbed region is not altered and that the buckles canbe described as a sinusoidal wave, the height discrepancybetween the mean topography of the perturbed andunperturbed regions should correspond to the amplitudeof the waves, or equivalently to x2 times the RMSroughness of the perturbed region. In other words, if thismodel were correct, the linear fit shown in Figure 13 wouldhave a slope of unity. Instead, the line has a slope of 0.51.This leaves us with two possibilities: (1) the bundlesdeviate from sinusoidal waves and/or the film’s thicknessis smaller at the peaks and troughs (i.e., regions of highcurvature), or (2) the volume increases within the per-turbed regions by some mechanism other than delami-nation.

It must be noted that other authors have seen similarscan-induced patterns in much thicker samples than whatwe consider here. Elkaakour et al. performed measure-ments on 14 µm films of polyacetylene cast onto mica andWSe2 substrates.12 Other authors saw qualitatively similarpatterns on bulk samples (injection-molded surfaces) ofpolycarbonateandpoly(methylmethacrylate).5-8 For these

Figure 13. Height discrepancy between the patterned regionand surroundings plotted as a function of the amplitude of thepattern (i.e., calculated from eq 3). Data are shown for twodifferent PS film thicknesses. The solid line represents a linearfit forced through the origin.

Figure 14. Models of the scan-induced deformation: (A)sinusoidal thickness fluctuations in the absence of voids, (B)delamination, (C) crazing within the film. The horizontal linesrepresent the mean height of the patterned region relative tothat of the unperturbed surroundings.

RRMS ) [ 1mλ∫0

mλA2 sin2(2π

λX) dX]1/2

) x22

A (3)


cases, delamination is not a realistic explanation of theobserved patterns. Therefore, we return to the idea thatthe scanning tip produces defects within the polymer inthe vicinity of the sample surface. We suggest that thedefects are not completely vacant, however, but containfibrils of polymer that occupy approximately half of thevolume within a void matrix, which explains the slope ofthe graph in Figure 13. This volume discrepancy isconsistent with the formation of crazes.35,36 Long thoughtto occur only in amorphous polymer glasses such aspolystyrene, poly(methyl methacrylate), and polycarbon-ate, crazing also has been found in semicrystallinepolymers during the past 20 years (though not reportedfor polyacetylene).37-39 Investigations of crazes in bulksamples reveal that the fibrils have diameters on the orderof 10 nm, and that these fibrils stretch as the craze growswider. This results in the storage of an elastic stress withinthe craze that tends to prevent further opening of thedefect. Model C in Figure 14 shows a craze within eachof the mounds generated by the scanning tip, and thefibrils are shown perpendicular to the tear within the film,as would be expected on the basis of scanning electronmicrographs of crazes in bulk samples.36,40 Measurementssimilar to those described in Figures 7 and 13 were alsoperformed on a 10 nm thick unannealed PVAc film. Giventhe low bulk glass transition temperature of PVAc (32 °C)and the fact that the sample was unannealed, this filmwas perhaps somewhat rubbery in its unperturbed state.Nevertheless, the mean height of the patterned regionwas again raised compared to that of the surroundingregion.

It is generally accepted that the presence of a free surface(i.e., a polymer-air or polymer-vacuum interface) tendsto increase polymer mobility and decrease the glasstransition temperature.41-46 A detailed discussion of thevarious models that have been developed to describe theeffect of film thickness on the glass transition temperatureand polymer mobility can be found in a review by Forrestand Dalnoki-Veress.45 At a simple level, one may intui-tively ascribe a depressed Tg, or enhanced polymermobility, at the polymer-air interface to a larger freevolume than what is encountered in bulk samples at thesame temperature. More precisely, it has been suggestedthat polymer molecules near an interface adopt confor-mations different from those predicted by self-avoidingrandom walk statistics that govern the behavior in bulksystems.47 The effect of an altered conformational statenear the polymer-air interface may be to reduce thenumber of entanglements,48 or to reduce the local density

(although differences from the bulk value have not beenobserved49,50), thereby resulting in an enhanced polymermobility. Other models describe the Tg behavior in thinfilms in terms of a characteristic length scale for dynamicsin glassy systems, suggesting that a layer near the surfaceexhibits faster dynamics, and hence a Tg lower than thebulk value.46 While several studies have shown anincreased glass transition temperature for different thinfilm/substrate systems than what have been consideredin this paper, this has been attributed to attractiveinteractions between the film and substrate51-53 and toan increased polymer density52 in the vicinity of thesubstrate. In studies of the glass transition in thin, self-supporting polymer films (i.e., films without a substrateand with two polymer-air interfaces), the Tg of the filmhas always been observed to be less than the bulk value.43,44

Because the generation of voids in the bundles isexpected to increase the effective polymer-air interfacialarea, it is reasonable to postulate that the bundles willexhibit a lower effective glass transition temperature thanthe trenches or the surrounding, unperturbed regions ofthe film.54 One may therefore expect the more rubberymoieties to exhibit a greater tendency to dissipate me-chanical energy compared to the more glassy domains ofthe film. In fact, this is exactly what was observed in thefriction data presented in Figures 5-7, where the meanfriction in the patterned region is elevated relative to thatof the surrounding film and two discreet domains of frictionforce were encountered within the pattern itself. Wetherefore may assign the type I friction to more glassy/less rubbery domains, and type II friction to more rubbery/less glassy domains. The broad distribution of the frictionforce in the patterned region (Figure 6) is consistent withthe continuous transition from glassy to rubbery behaviorproposed earlier in this paper. Scan velocity is of courseimportant in locating the system precisely along thistransition. In this study no attempt was made to quantifyhow far along the viscoelastic spectrum the more rubberydomains were translated relative to less rubbery domains.

Nonperturbative scanning-probe friction measurementsof semicrystalline poly(vinyl alcohol) and gelatin filmshave revealed that highly crystalline domains do notexhibit the glass-to-rubber transitional behavior (elevatedfriction) seen in more amorphous domains.55 Therefore,it would be illuminating to compare the patterningbehavior of predominantly crystalline and amorphousdomains of the same polymer film, i.e., examine bothvolume and dissipative changes.56 By applying the meth-odologies presented in this paper to a wider variety of

(35) Kramer, E. J.; Berger, L. L. Adv. Polym. Sci. 1990, 91/92, 1-68.(36) Strobl, G. The Physics of Polymers; Springer-Verlag: Berlin,

1995.(37) Friedrich, K.; Kausch, H. Adv. Polym. Sci. 1983, 52/53, 225-

274.(38) Jang, B. Z.; Uhlmann, D. R.; Vander Sande, J. B. Polym. Eng.

Sci. 1985, 25, 98-112.(39) Narisawa, I.; Ishikawa, M. Adv. Polym. Sci. 1990, 91/92, 353-

391.(40) Kambour, R. P. Macromol. Rev. 1973, 7, 1.(41) Keddie, J. L.; Jones, R. A. L.; Cory, R. A. Europhys. Lett. 1994,

27, 59-64.(42) Keddie, J. L.; Jones, R. L.; Cory, R. A. Faraday Discuss. 1994,

98, 219-230.(43) Forrest, J. A.; Dalnoki-Veress, K.; Stevens, J. R.; Dutcher, J. R.

Phys. Rev. Lett. 1996, 77, 2002-2005.(44) Forrest, J. A.; Dalnoki-Veress, K.; Dutcher, J. R. Phys. Rev. E

1997, 56, 5705-5716.(45) Forrest, J. A.; Dalnoki-Veress, K. Adv. Colloid Interface Sci.

2001, 94, 167-196.(46) Forrest, J. A.; Mattson, J. Phys. Rev. E 2000, 61, 53-56.(47) Jones, R. L.; Kumar, S. K.; Ho, D. L.; Briber, R. M.; Russell, T.

P. Nature 1999, 400, 146-149.(48) Brown, H. R.; Russell, T. P. Macromolecules 1996, 29, 798-800.

(49) Forrest, J. A.; Dalnoki-Veress, K.; Dutcher, J. R. Phys. Rev. E1998, 58, 6109-6114.

(50) Wallace, W. E.; Beck Tan, N. C.; Wu, W. L. J. Chem. Phys. 1998,108, 3798-3804.

(51) van Zanten, J. H.; Wallace, W. E.; Wu, W. L. Phys. Rev. E 1996,53, R2053-R2056.

(52) Grohens, Y.; Brogly, M.; Labbe, C.; David, M.; Schultz, J.Langmuir 1998, 14, 2929-2932.

(53) Tsui, O. K. C.; Russell, T. P. Macromolecules 2001, 34, 5535-5539.

(54) One might contend that chain extension and alignment imposedon the polymer during the formation of a craze would more likely resultin an increased local Tg. It has been shown, however, that macroscopicrubbing of a polymer imposes alignment of molecules near the surfaceand that the relaxation of this anisotropy (determined from opticalbirefringence measurements) is not inhibited: the Tg of the rubbedfilms was depressed by an amount very similar to that reported in studiesof unstressed polystyrene film.57 Thus, in this case the presence of afree interface apparently had a greater impact on Tg than did increasedorientation or chain extension.

(55) Haugstad, G.; Hammerschmidt, J. A.; Gladfelter, W. L. InInterfacialPropertieson theSubmicrometerScale;Frommer,J.,Overney,R. M., Eds.; American Chemical Society: Washington, DC, 2001; Vol.781.


polymers, the potential universality of our proposedmechanism of scan-induced pattern formation may beexplored.

SummaryWe have characterized the susceptibility of PS and PVAc

films to damage when a sharp tip is scanned in contactwith the surface in a raster pattern. The damage causedby the tip can be approximated as a sinusoidal waveoriented perpendicular to the tip trajectory. An analysisof the mean height within patterned regions showed anincreased volume that is consistent with the formation ofcrazes within the polymer film. The volume that is gainedwhile forcing plastic deformation is interpreted to con-tribute toward a second-order phase transition in the film,including more/less rubbery domains within the patternedregions as contrasted by friction force imaging. Aiding

these findings was a newly defined parameter, rougheningsusceptibility, the slope of the observed linear increase ofroughness under repeated raster scanning. This param-eter, as well as the wavelength of the scan-inducedpatterns, was characterized as a function of the numberof raster scans performed over the same region, the appliedload, the scan line density, and the gain of the feedbackcircuit used to maintain a constant applied load, generallydemonstrating the robustness of our methodology andremoving a curtain of mystery regarding the impact ofthese parameter settings. Moreover, the similarity be-tween the patterns we characterized and those reportedby other investigators for a wide variety of polymer/substrate systems suggests that the methodologies wedeveloped can be applied to the analysis of many glassyand semicrystalline polymers. In a subsequent paper, wewill address the temperature and rate dependence of thesescan-induced patterns and further exploit the methodologyreported herein.

Acknowledgment. We thank the Center for Inter-facial Engineering and the donors of the PetroleumResearch Fund, administered by the American ChemicalSociety, for supporting this work.

LA015769J

(56) Several authors have examined patterning in nominally sem-icrystalline polymers. The crystalline nature of the PET films inves-tigated by Jing et al.11 and the PEO films studied by Nie and co-workers14

was confirmed by the initial AFM scan of the surfaces (i.e., before theonset of significant damage by the scanning tip). Nie et al. found thatcrystalline films of PET were much more difficult to pattern thanamorphous PET films. For the case of the polyacetylene film investigatedby Elkaakour et al.,12 however, the degree of crystallinity is unknown.

(57) Schwab, A. D.; Agra, D. M. G.; Kim, J.; Kumar, S.; Dhinojwala,A. Macromolecules 2000, 33, 4903-4909.


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