Nanowear on Polymer Films of Different Architecture - Langmuir (ACS

Feb 6, 2007 - In this paper, we describe atomic force microscope (AFM) friction experiments on different polymers. The aim was to analyze the influenc...
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Langmuir 2007, 23, 3150-3156

Nanowear on Polymer Films of Different Architecture R. Berger,† Y. Cheng,† R. Fo¨rch,† B. Gotsmann,§ J. S. Gutmann,†,‡ T. Pakula,† U. Rietzler,† W. Scha¨rtl,‡ M. Schmidt,‡ A. Strack,‡ J. Windeln,§ and H.-J. Butt*,† Max Planck Institute for Polymer Research, Ackermannweg 10, D-55128 Mainz, Germany, Institute for Physical Chemistry, UniVersity of Mainz, 55099 Mainz, Germany, and IBM Research GmbH, Zurich Research Laboratory, 8803 Ru¨schlikon, Switzerland ReceiVed July 14, 2006. In Final Form: NoVember 20, 2006 In this paper, we describe atomic force microscope (AFM) friction experiments on different polymers. The aim was to analyze the influence of the physical architecture of the polymer on the degree and mode of wear and on the wear mode. Experiments were carried out with (1) linear polystyrene (PS) and cycloolefinic copolymers of ethylene and norbornene, which are stabilized by entanglements, (2) mechanically stretched PS, (3) polyisoprene-b-polystyrene diblock copolymers, with varying composition, (4) brush polymers consisting of a poly(methyl methacrylate) (PMMA) backbone and PS side chains, (5) PMMA and PS brushes grafted from a silicon wafer, (6) plasma-polymerized PS, and (7) chemically cross-linked polycarbonate. For linear polymers, wear depends critically on the orientation of the chains with respect to the scan direction. With increasing cross-link density, wear was reduced and ripple formation was suppressed. The cross-linking density was the dominating material parameter characterizing wear.

Introduction Wear of a nanoscopic contact with a polymer surface is not only important for a better understanding of friction between inorganic and polymeric surfaces, it is also relevant for designing nano- and micro-electromechanical systems. Friction and wear are among the most critical problems to be solved to fabricate commercially viable microdevices with moving parts. The specific motivation for the project described here was to develop a suitable coating for scanned-probe data storage1-5 and to better understand scanned probe lithography and patterning.4,6-8 When a single nanocontact slides over a polymer surface, different wear modes are observed. On thermoplasts such as polystyrene (PS),9-16 polyacetylene,17 polyester,18 polycarbonate * To whom correspondence should be addressed. E-mail: butt@ mpip-mainz.mpg.de. † Max-Planck-Institute for Polymer Research. ‡ University of Mainz. § Zurich Research Laboratory. (1) Mamin, H. J.; Rugar, D. Appl. Phys. Lett. 1992, 61, 1003-1005. (2) Terris, B. D.; Rishton, S. A.; Mamin, H. J.; Ried, R. P.; Rugar, D. Appl. Phys. A 1998, 66, S809-S813. (3) Vettiger, P.; Cross, G.; Despont, M.; Drechsler, U.; Du¨rig, U.; Gotsmann, B.; Ha¨berle, W.; Lantz, M. A.; Rothuizen, H. E.; Stutz, R.; Binnig, G. K. IEEE Trans. Nanotechnol. 2002, 1, 39-55. (4) Vasilev, C.; Heinzelmann, H.; Reiter, G. J. Polym. Sci., Part B: Polym. Phys. 2004, 42, 1312-1320. (5) Drechsler, U.; Bu¨rer, N.; Despont, M.; Du¨rig, U.; Gotsmann, B.; Robin, F.; Vettiger, P. Microelectron. Eng. 2003, 67-68, 397-404. (6) Majumdar, A.; Oden, P. I.; Carrejo, J. P.; Nagahara, L. A.; Graham, J. J.; Alexander, J. Appl. Phys. Lett. 1992, 61, 2293-2295. (7) Jin, X.; Unertl, W. N. Appl. Phys. Lett. 1992, 61, 657-659. (8) Kunze, U.; Klehn, B. AdV. Mater. 1999, 11, 1473-1475. (9) Leung, O. M.; Goh, M. C. Science 1992, 255, 64-66. (10) Meyers, G. F.; DeKoven, B. M.; Seitz, J. T. Langmuir 1992, 8, 23302335. (11) Woodland, D. D.; Unertl, W. N. Wear 1997, 203-204, 685-691. (12) Schmidt, R. H.; Haugstad, G.; Gladfelter, W. L. Langmuir 1999, 15, 317-321. (13) Sills, S.; Overney, R. M. Phys. ReV. Lett. 2003, 91, 095501. (14) Schmidt, R. H.; Haugstad, G.; Gladfelter, W. L. Langmuir 2003, 19, 10390-10398. (15) Vettiger, P.; Binnig, G. Sci. Am. 2003, Jan., 35-41. (16) Sills, S.; Gray, T.; Overney, R. M. J. Chem. Phys. 2005, 123, 134902. (17) Elkaakour, Z.; Aime´, J. P.; Bouhacina, T.; Odin, C.; Masuda, T. Phys. ReV. Lett. 1994, 73, 3231-3234. (18) Jing, J.; Henriksen, P. N.; Wang, H.; Marteny, P. J. Mater. Sci. 1995, 30, 5700-5704.

(PC),19-21 and poly(tert-butyl acrylate),22 ripple (also called bundle) formation usually perpendicular to the scan direction is observed. Keeping scan velocity, load, and molecular weight constant, the spacing between ripples and the overall roughness increases with temperature.13,14 When the temperature exceeds the glass-transition temperature, the whole scanned area is homogeneously worn out and material accumulates at the rim.22,23 The rippling wear mode changes to the piling wear mode, also called tearing wear mode. While the influence of parameters such as temperature, sliding velocity,13,14,16 and molecular weight on wear have been studied extensively, significantly less is known about the influence of the physical structure of the polymer. In this paper, we describe atomic force microscope (AFM) friction experiments on different polymers (Figure 1). The aim was to analyze the influence of the physical architecture of the polymer on wear. We started with simple linear polymer chains. Such thermoplasts are stabilized by entanglements between the chains. The first linear polymer used was PS as the standard reference. To see the influence of chain orientation, we also stretched PS samples mechanically. In addition, films made from a cycloolefinic copolymer (COC) of ethylene and norbornene with the trade name Topas were analyzed. By a selected use of a catalyst, the statistics and the stereoselectivity of the norbornene insertion can be controlled and the properties of the amorphous (within a certain composition range) copolymer can be customized. These properties include high transparency, low birefringence, low water adsorption, excellent water vapor barrier properties, and high heat decomposition temperatures.24-26 Its glass-transition temperature Tg can be adjusted by choosing the (19) Iwata, F.; Matsumoto, T.; Sasaki, A. Nanotechnology 2000, 11, 10-15. (20) Kaneko, R.; Hamada, E. Wear 1993, 162-164, 370-377. (21) Khurshudov, A.; Kato, K. J. Vac. Sci. Technol., B 1995, 13, 1938-1944. (22) Wang, X. P.; Loy, M. M. T.; Xiao, X. Nanotechnology 2002, 13, 478483. (23) Gotsmann, B.; Du¨rig, U. Langmuir 2004, 20, 1495-1500. (24) Haselwander, F. A.; Heitz, W.; Kru¨gel, S. A.; Wendorff, J. H. Macromol. Chem. Phys. 1996, 197, 3435-3453. (25) Rische, T.; Waddon, A. J.; Dickinson, L. C.; MacKnight, W. J. Macromolecules 1998, 31, 1871-1874. (26) Kolarik, J.; Krulis, Z.; Slouf, M.; Fambri, L. Polym. Eng. Sci. 2005, 817-826.

10.1021/la0620399 CCC: $37.00 © 2007 American Chemical Society Published on Web 02/06/2007

Nanowear on Polymer Films

Figure 1. Different kinds of polymer architectures used.

appropriate amount of ethylene and norbonene.25,27 The main motivation to use a COC is its special mechanical properties. Up to Tg, the elastic storage modulus and its microhardness are constant and at Tg they decrease sharply.28,29 To study the potential influence of loops on friction and wear, different samples without loops at the polymer surface were prepared. To reduce the number of loops and to orient the polymer chains normal to the surface, an ultrathin diblock copolymer film was formed. Because of the low film thickness, the lower energy block of the copolymer will preferentially segregate to the air interface. In this segregated layer, the individual chains have a preferred orientation normal to the surface and adopt a stretched conformation. Another way of avoiding loops at the surface is to use “brush polymers”. We used a brush polymer consisting of a poly(methyl methacrylate) (PMMA) backbone and PS side chains. Brush polymers are relatively stiff rods since the side chains sterically hinder the backbone to bend. In this way, only the side chains are at the surface and do not expose loops. Polymers grafted from a surface also expose few loops on the surface. In contrast to the diblock copolymer and brush polymer films, they are additionally stabilized by the covalent bond to the solid substrate. We produced and analyzed PS and PMMA films. Cross-linking has an influence on wear. Cross-linked epoxy20 and PS,30 for example, showed significantly less wear than noncross-linked thermoplasts. To study the influence of cross-linking, films were made by plasma-polymerization with styrene as the monomer. In particular, the influence of the plasma power was studied. With increasing power density, the cross-linking density in the films increased. In addition to plasma polymerization of PS, we made chemically cross-linked films of polycarbonate. For un-cross-linked films, the mechanical properties of bisphenol A based polycarbonate are better than those of un-cross-linked polystyrene (e.g., lower hardness and lower Youngs modulus, ref 31). For this reason, cross-linked polycarbonate is expected to show less wear than cross-linked polystyrene. Materials and Methods Atomic Force Microscopy. Samples were analyzed with an atomic force microscope (Dimension 3100, NanoScope IIIa controller, Veeco, USA) with microfabricated tips in contact mode using (27) Forsyth, J. F.; Scirivani, T.; Benavente, R.; Martesin, C.; Perena, J. M. J. Appl. Polym. Sci. 2001, 82, 2159-2165. (28) Ekizoglou, N.; Thorshaug, K.; Cerrada, M. L.; Benavente, R.; Pe´rez, E.; Perena, J. M. J. Appl. Polym. Sci. 2003, 89, 3358-3363. (29) Blochowiak, M.; Pakula, T.; Butt, H. J.; Bruch, M.; Floudas, G. Macromolecules 2006, in press. (30) Gotsmann, B.; Duerig, U. T.; Sills, S.; Frommer, J.; Hawker, C. J. Nanoletters 2006, 6, 296-300.

Langmuir, Vol. 23, No. 6, 2007 3151 rectangular silicon cantilevers (Cont-W, Nanoworld, Switzerland). The spring constant for each individual spring was experimentally determined by the thermal noise method. Before each measurement, the cantilevers with tips were cleaned in a plasma cleaner (Ar plasma, PDC-001, Harrick Scientific Corporation) for 30 s. To measure wear, the analysis involved several steps. First, polymer films were imaged with an AFM in contact mode with the lowest force possible. The standard scan range was 5 × 5 µm2. In addition, films were inspected by an optical microscope to ensure that the samples were homogeneous on the larger length scale. Second, a wear test was performed by scanning 100 times an area of 2.5 × 2.5 µm2 at a scan velocity of 25 µm/s and at a defined force (load of 10 nN if not mentioned otherwise). Finally, an AFM image of 5 × 5 µm2, which contained the previously scanned area, was taken at low force. Most of the tips used were later imaged in a lowvoltage scanning electron microscope (SEM, LEO 1530 Gemini, Oberkochen, Germany) to see if tip wear had occurred and to obtain the radius of curvature of the tip. By using a low-voltage SEM, the tips did not have to be coated with a metal. Polystyrene. As a reference, we studied linear polystyrene films on silicon wafers (Si(100), with a natural oxide layer, Wacker, Burghausen, Germany). The polymer was synthesized by anionic polymerization (molecular weight Mw ) 111 kDa and 70 kDa, Mw/ Mn ) 1.08). The polymer was spin-coated from toluene solution with a concentration of 20 mg/mL at a rotation speed of 2500 rpm for 60 s on an open spin-coater. They were not cross-linked. Cycloolefinic Copolymer (COC). We chose a COC of Mw ) 106 kDa (Mw/Mn ≈ 1.8) with a norbornene content of 61.5 mol %. Its Tg as determined by differential scanning calorimetry (DSC) was 181 °C.29 Films were formed by spin-coating from toluene solution (rotation speed 2500 rpm) on silicon wafers and not cross-linked. Mechanically Stretched PS. Stretched samples were made by first pressing granular PS (Mw ) 212 kDa, Mw/Mn ) 1.07) in a rectangular pressing mold (PW 40EH, Paul-Otto Weber Maschinenund Apparatebau GmbH, Germany) of 10 × 60 mm2 size at an applied force of 20 kN at 160 °C. Each sample contained 0.8 g of material leading to a thickness of ≈1.1 mm. To stretch PS blocks, they were clamped at both ends in an extensometer (Instron 6022, Instron Deutschland GmbH, Germany) with an initial gap of 30 mm. The sample was heated to 100 °C and was then pulled apart at constant velocity of 0.5 mm/min to a total length of 75 mm leading to a draw ratio of 250%. Diblock Copolymers. Three diblock copolymers consisting of polyisoprene-b-polystyrene (PI-b-PS, Mw ) 22.7, 23.9, 27.6 kDa, Mw of PI block 10, 12, 8 kDa, respectively, Mw/Mn ) 1.05-1.08) were synthesized. Films were prepared on silicon wafers by spincoating a 9.5-9.7 mg/mL toluene solution at 2000 rpm. The dry film thicknesses were 42, 42, and 39 nm for the Mw ) 22.7, 23.9, and 27.6 kDa diblock copolymer films as measured by X-ray reflectivity (Figure 2). The film thickness was obtained by modeling of the block copolymer film using a Parratt formalism32 and a minimal number of layers. The modeled reflectivity is shown with the experimental data in Figure 2. It is well-known that diblock copolymers in an ultrathin supported film geometry preferentially segregate during film formation to expose the low-energy block at the air interface.33,34 This preferential segregation is not affected by the microstructure of the diblock copolymer. For the used PI-b-PS diblock copolymers, this leads to an exposure of the PI chains at the free film surface. Thus, the PI-chains are preferentially oriented perpendicular to the polymer surface, irrespective of the used diblock copolymer, resulting in a very smooth surface as evidenced by the large amount of oscillations in the X-ray reflectivity curve (Figure 2). (31) Briscoe, B. J.; Fiori, L.; Pelillo, E. J. Phys. D 1998, 31, 2395-2405. (32) Parratt, L. G. Phys. ReV. 1954, 95, 359-369. (33) Knoll, A.; Horvat, A.; Lyakhova, K. S.; Krausch, G.; Sevink, G. J. A.; Zvelindovsky, A. V.; Magerle, R. Phys. ReV. Lett. 2002, 89, 035501. (34) Karim, A.; Singh, N.; Sikka, M.; Bates, F. S.; Dozier, W. D.; Felcher, G. P. J. Chem. Phys. 1994, 100, 1620-1629.

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Figure 3. Storage (G′), loss modulus (G′′), and the loss angle (tan δ) versus temperature of the cross-linked polycarbonate. Figure 2. X-ray reflectivity versus wavevector transfer for the three PI-b-PS diblock copolymer films. The continuous lines represent the modeling by a single layer model on the basis of the Parratt formalism.32 Brush Polymers. Films of two types of brushes were made (for details of synthesis see ref 35): PS-6030 (Mw ) 6030 kDa, Mw/Mn ) 1.37) and PS-2930 (Mw ) 2930 kDa, Mw/Mn ) 1.28). Both had a PMMA backbone and PS side chains (grafting density 100%). Side chains had a molecular weight of Mw ) 4.5 kDa (Mw/Mn ) 1.06). Films were prepared by spin-coating in air from toluene solutions (20 mg/mL, 2500 rpm). Grafted Polymers. We produced and analyzed PS and PMMA films. The individual chains are polymerized from an initiator which is covalently bound to the silicon wafer. PMMA and PS were synthesized via surface-initiated ATRP and RAFT, respectively (for details, see ref 36). Dry films were 25-100 nm thick. Plasma Polymerized PS. Polystyrene films were also made by plasma polymerization on silicon wafer using a home-built, cylindrical, capacitively coupled 13.56 MHz reactor, which has been described before.37 The monomer chosen was styrene providing PS-like films of 25-100 nm thickness. All monomers were deposited in a continuous wave plasma process at a pressure of 5 Pa and at different input powers. Cross-linked polycarbonate films were prepared by thermal polymerization of diethylleneglycol bis(allylcarbonate) (ADC, supplied by Great Lakes Chemical Corporation) using di-benzoly peroxide as initiator. In a first test, the degassed solution of the neat monomer and starter was spin-coated onto silicon wafers and was thermally polymerized at 90 °C immediately after spin-coating. However, films could not easily be prepared as the monomer dewetted the silicon wafer. Even after silanating the wafer with octadecyltrimethylsilane, spontaneous dewetting occurred. For this reason, we prepared thick solution cast films for further investigation. Crosslinked polycarbonate films were prepared as follows: (1) Diethyleneglycol bis(allylcarbonate) monomer (ADC) and the di-benzoyl peroxide (BPO) starter were mixed to obtain a solution containing 1-2% BPO. (2) The solution was stirred until the BPO was complete dissolved. (3) Residual oxygen was removed from the reaction mixture via freeze-pump-thaw cycles. (4) The reaction mixture was cast onto a wafer and was polymerized at 93 °C under vacuum for 12 h followed by a second post curing step at 110 °C for 1 h. We found that the quality of the BPO starter greatly affects the surface properties of the cast and cured films. Therefore, fresh BPO was used. (35) Hugenberg, N.; Loske, S.; Mu¨ller, A. H. E.; Scha¨rtl, W.; Schmidt, M.; Simon, P. F. W.; Strack, A.; Wolf, B. A. J. Non-Cryst. Solids 2002, 307-310, 765-771. (36) Bumbu, G. G.; Kircher, G.; Wolkenhauer, M.; Berger, R.; Gutmann, J. S. Macromol. Chem. Phys. 2004, 205, 1713-1720. (37) van Os, M. T.; Menges, B.; Fo¨rch, R.; Knoll, W.; Vancso, G. J. Chem. Mater. 1999, 11, 3252-3257.

Cross-linked polycarbonate was characterized by dynamic mechanical spectroscopy (Rheometrics Scientific ARES Spectrometer). Stripes of 31-mm length, 10-mm width, and 0.42-mm thickness were placed into a torsion rectangular sample holder. Isochronal temperature scans were recorded at constant oscillatory excitation frequency of 10 rad/s in a temperature range of -20 °C to 200 °C at a heating rate of 2 °C/min. The strain needed for an optimal torque was determined at -20 °C prior to the temperature sweep. During the temperature sweep, losses in torque were automatically compensated by adjusting the applied strain. The dynamic mechanical characterization of bulk samples showed a broad transition in the temperature range between 60 °C and 110 °C (Figure 3). The broad peak in the loss tangent (tan δ) indicates that a whole series of relaxation processes with very different activation temperatures are simultaneously occurring. Such a behavior is indicative for a range of different structural elements such as cross-linking points with three or four linkages and a large variation of distances between cross-linking points. Consequently, the sample does not show a single well-defined glass-transition temperature but softens gradually over a large temperature range.

Results and Discussion Films of linear PS were smooth and displayed a roughness of only rms ) 0.4 nm (2.5 × 2.5 µm2). When scanning PS films, ripples appeared with a spacing of 70-80 nm. With an increasing number of scans, the height of these ripples increased and also a small increase in ripple spacing to 80-90 nm was observed. When scanning long enough, the ripples were pushed to the side and material piled up at the rim of the scanned area. The number of scans required to initiate a certain effect decreased with decreasing scan speed (Figure 4) and load. The inner 1 µm2 square was scanned at a scan speed of 5 µm/s. Strong rippling with a spacing of 80-90 nm was observed even after two scans. After three scans, we zoomed out and scanned an area of 2.5 × 2.5 µm2 with a scan speed of 25 µm/s. Only after 5-10 scans ripples became visible and after 24 scans the spacing (62-68 nm) was still lower than on the slowly scanned region. As a result of scanning, the roughness increased significantly to rms ) 2.2 nm on 2.5 × 2.5 µm2 (Table 1). This agrees with what has been observed before with PS: ripple formation, the growth of spacing and height with the scan number,9,10,22,38 the final removal of the top layer in the piling wear mode,23,39 and the decrease with scan velocity.12,17,22 COC films were relatively smooth (rms ) 0.5 nm). After scanning a few times, ripples appeared and as with PS grew in (38) Schmidt, R. H.; Haugstad, G.; Gladfelter, W. L. Langmuir 2003, 19, 898-909. (39) Aoike, T.; Uehara, H.; Yamanobe, T.; Komoto, T. Langmuir 2001, 17, 2153-2159.

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Figure 4. AFM image of a polystyrene surface (Mw ) 70 kDa) after scanning the inner square of 1 µm2 three times at a load of FL ) 10 nN with a scan speed of V ) 5 µm/s. Then, we zoomed out and scanned a square of 2.5 × 2.5 µm2 (FL ) 10 nN, V ) 25 µm/s) 24 times. Finally, the area shown was scanned at low load. Height scale 5 nm. Table 1. Mean rms Roughnesses of Polymer Films before and after Scanning 100 Times at a Load of 10 nN with a Scan Speed of 25 µm/sa sample

roughness before rms (nm)

polystyrene COC stretched PS

0.4 0.5 3

PI-b-PS PS-PS-6030 PS-PS-2930 grafted PS grafted PMMA plasma polym. PS cross-linked PC

0.6 0.9 0.5 1 0.4 0.4-0.7 0.6

a

roughness after rms (nm)

Figure 5. AFM image of a COC film after scanning the central region of 2.5 × 2.5 µm2 100 times with a load of 10 nN at V ) 25 µm/s. Height scale 20 nm, scan area 4.4 × 4.4 µm2. The graph shows the mean spacing between ripples versus the number of the scan for the same experiment.

2.2 nm 5 4-5 (0°, 45°) 10-15 nm (90°) destroyed destroyed destroyed 2.5 0.5 0.7-2 0.8

The roughness was determined over areas of 2.5 × 2.5 µm2.

height and also spacing with the number of the scan (Figure 5). Typical distances between ripples were 70-120 nm. In parallel, the roughness increased and reached rms ≈ 5 nm after 100 scans. Thus, COC did not show superior wear behavior over PS and like other thermoplasts showed rippling wear. Stretched PS. Stretched PS sample showed a higher roughness (rms ≈ 3 nm) than spin-coated films. When scanning the surface of stretched PS (100 scans, FL ) 10 nN, V ) 25 µm/s) with the fast scan direction being parallel to the stretch direction, the surface showed significant wear (Figure 6). Ripples with a spacing of typically 80-90 nm appeared on the surface, which is similar to results obtained on spin-coated PS films. When rotating the sample to 45° with respect to the fast scan direction, the result was similar to the result observed in parallel orientation. When scanning perpendicular to the stretch direction (90°), wear increased drastically. Holes appeared on the surface and eventually the scanned area was removed. A possible explanation of the increased wear when scanning perpendicular to the stretch direction is that more polymer chains might be cut by the tip. While scanning, the tip crosses more chains per unit length and the ends of the chains are more tightly attached because of the orientation. That individual chains can be broken by an AFM tip has been demonstrated by Cappella et al.40 In the case of a stretched polymer, a breaking of chains (40) Cappella, B.; Sturm, H.; Weidner, S. M. Polymer 2002, 43, 4461-4466.

Figure 6. AFM images of 250% stretched polystyrene after scanning the inner square of 2.5 × 2.5 µm2 100 times (FL ) 10 nN, V ) 25 µm/s). The angle of the fast scan direction with respect to the stretch direction was changed from being parallel (0°) over 45° to being perpendicular (90°). Height scales are 25 nm (0°, 45°) and 40 nm (90°).

can also lead to a relaxation of the two ends which can further widen the mechanically created hole. Diblock Copolymer. The roughness of the fresh PI-b-PS films was about rms ≈ 0.6 nm. When scanning, whole layers were removed and the surface was completely destroyed (Figure 7). After 100 scans, the polymer layer was completely removed and the substrate was reached. In the case shown in Figure 7, the layer was 37-nm thick. Qualitatively, the same was observed with brush polymers. With a roughness rms ≈ 0.8 nm, these films were not as smooth as many of the other films. After

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Figure 7. The 5 × 5 µm2 AFM images of a film of polyisopreneb-polystyrene (top, Mw ) 24 kg/mol) and a polystyrene brush polymer PS-6030 (bottom) after scanning the central 2.5 × 2.5 µm2 100 times (FL ) 10 nN, V ) 25 µm/s). Height scales: 120 (top) and 140 nm (bottom).

Figure 9. AFM images of PS films formed by plasma polymerization at different power densities of 20, 80, and 200 W. The AFM images show the films after scanning the central 2.5 × 2.5 µm2 100 times (FL ) 10 nN, V ) 25 µm/s). Film thicknesses, image size, and height scales were 26 nm, 4.4 × 4.4 µm2, 20 nm for 20 W; not measured, 4.2 × 4.2 µm2, 5 nm for 80 W; 55 nm, 4.3 × 4.3 µm2, 8 nm for 200 W.

Figure 8. AFM image of grafted PS and PMMA polymer layers after scanning the central 2.5 × 2.5 µm2 100 times (FL ) 10 nN, V ) 25 µm/s). The image sizes are 3.9 × 3.9 µm2 and 4.1 × 4.1 µm2 and the height scales are 20 and 4 nm, respectively.

scanning, whole layers of the film were removed leaving a square hole which went down to the substrate; in the experiment shown in Figure 7, this was 100 nm. We did not observe significant differences between PS-2930 and PS-6030. Both films were completely removed probably because the grafting leads to an increased stiffness which prevents entanglement. The individual PS side chains are too short to form loops or to entangle; the entanglement molecular weight of PS is Me ) 18.1-19.1 kDa.41,42 For the diblock copolymer, entanglement was reduced by the (41) Fetters, L. J.; Lohse, D. J.; Milner, S. T.; Graessley, W. W. Macromolecules 1999, 32, 6847-6851. (42) Brandrup, J.; Immergut, E. H.; Grulke, E. A. Polymer Handbook; 4th ed.; John Wiley & Sons: USA, 1999.

microphase separation and the layered structure. The individual blocks were deliberately too short to lead to significant entanglement. Longer chains entangle but they do not prevent the formation of loops at the surface. In the case of brush polymers, the increased stiffness prevents an entanglement of the backbones. Again, the side chains are too short to entangle. Both types of polymers are not suitable to reduce wear. These experiments show that loops do not play an important role in friction between a nanoscale tip and a polymer. Grafted Polymers. Grafted PMMA films were relatively smooth with a roughness of rms ≈ 0.4 nm over 2.5 × 2.5 µm2. PS films were slightly rougher (rms ≈ 1 nm) and showed granular features on the surfaces. After scanning, PMMA films were only slightly modified and the central square is even difficult to identify on AFM images (Figure 8). Rippling could hardly be observed at loads of up to 40 nN. Even when wear was observed, rippling was often not dominant and removal of material was the typical wear mode. If ripples could be seen, they were typically narrowly spaced (20-30 nm). On PS, scanning had a stronger effect (Figure 8). The roughness increased to typically 2.5 nm and ripples perpendicular to the scanning direction started to form. With a spacing of 45-65 nm, the periodicity was shorter than on PS films made of linear chains. However, grafting PS to the surface did not much reduce wear on PS. We have no knowledge of the

Nanowear on Polymer Films

reason for the decreased effect on PMMA, however, it agrees with earlier results obtained with nongrafted polymer layers.20 Accordingly, PMMA has been used for thermomechanical writing.1,6,43 Plasma Polymerized PS. Polymerization at different input powers enabled us to synthesize films of different cross-link densities and subsequently different chemical and physical properties of films from the same precursor.44 At high input power, the precursor molecules are completely dissociated and there is little control over chemical structures. Thus, the resulting films are often little more than highly cross-linked hydrocarbon deposits. The degree of dissociation of the molecules and thus the degree of cross-linking in the deposit decreases with decreasing input power. PS films made by plasma polymerization were as smooth (rms ≈ 0.4 nm) as films made from linear PS by spin-coating, at least when using a low power. With increasing power, a patchlike structure formed (Figure 9 bottom) on the surface leading to an increased roughness of rms ≈ 0.7 nm at 200 W (Figure 10). The scratch test showed that the wear mode of plasmapolymerized PS is qualitatively different from wear of linear polymer chains. At the relatively low power of 20-30 W, significant abrasion occurred. Material was removed and wear particles piled up at the rim of the scanned square (Figure 9 top). Rippling and the formation of bundles was not observed. In the case shown (Figure 9 top), the hole was ≈10 nm deep after 100 scans. Since the film is much thicker, this is caused by a gradual removal of material. With increasing power, the wear mode changed from peeling to rippling. At intermediate plasma powers (80-100 W), ripples formed. At high powers (above 150 W), no effect of the tip on the structure of the polymer film was observed. Cross-Linked Polycarbonate. The roughness of freshly prepared thick cross-linked polycarbonate films was rms ≈ 0.6 nm. The surface was not affected by scans at a load of 10 nN. Only when increasing the force to 30-50 nN was wear observed (Figure 11). If wear occurred, nanoparticles piled up at the rim of the scanned area. Ripples were not observed. This is in contrast to AFM wear experiments on non-cross-linked PC where rippling and an increase in volume were observed.20,21 Summarizing the experimental results, it was observed that cross-linking stabilizes the polymer film, prevents rippling, and reduces wear. It is the dominating parameter which characterizes the material with respect to wear. Why is the cross-linking density so crucial to reduce wear? Several models which have been proposed to interpret rippling and wear are consistent with this observation. Qualitatively, ripple formation can be explained by assuming that polymer accumulates in advance of the scanning tip as it slides across the surface.38 The presence of this mound results in an increased lateral force on the tip. As the mound grows larger, so does the lateral force. Eventually, this lateral force overcomes the adhesion plus load of the tip and the tip slides over the mound. This process is repeated over and over again. For this process to occur, chains are moved and partially pulled forward. Chain pullout becomes impossible with cross-linking and as a result the wear mechanism (43) Hinz, M.; Kleiner, A.; Hild, S.; Marti, O.; Du¨rig, U.; Gotsmann, B.; Drechsler, U.; Albrecht, T. R.; Vettiger, P. Eur. Polym. J. 2004, 40, 957-964. (44) Fo¨rch, R.; Zhang, Z.; Knoll, W. Plasma Processes Polym. 2005, 2, 351372. (45) Washiyama, J.; Kramer, E. J.; Hui, C. Y. Macromolecules 1993, 26, 2928-2934. (46) Schallamach, A. Wear 1971, 17, 301-312. (47) Yang, A. C. M.; Wu, T. W. J. Mater. Sci. 1993, 28, 955-962. (48) Hamilton, G. M.; Goodman, L. E. J. Appl. Mech. 1966, 371-376.

Langmuir, Vol. 23, No. 6, 2007 3155

Figure 10. Roughness (rms) of plasma-polymerized PS films depending on the power.

Figure 11. AFM image of a cross-linked polycarbonate film after scanning the central 2.5 × 2.5 µm2 100 times with a load of 50 nN at V ) 25 µm/s.

has to switch from chain pullout to crazing, just as it is described for polymer fracture (e.g., ref 45). Rippling has been interpreted as Schallamach waves. Schallamach suggested his theory to explain the formation of periodic structures observed when rubber slides on a hard smooth substrate.46 In his model, he replaced the rubber by equidistant vertical ridges connected by springs. The ridges were assumed to deform under shear and obey, like the springs, Hooks’ law. This model can also be used to explain the formation of ripples on a polymer surface when being scanned by a nanoscale tip.11,39 Cross-linking and the resulting increased shear stiffness would correspond to stiffer springs in the model, which leads to a completely different periodicity or practically to an end of ripple formation. In another model, it is assumed that cracks are formed once a certain shear stress has been exceeded.17,19,21,39,47 Usually, Hertz theory is used to quantify this approach. A spherical, infinitely hard tip of radius of curvature R is in contact with a flat, elastic polymer under a given load FL. Attractive forces between the tip and the polymer in particular adhesion are neglected. The tip forms a circular contact region with the surface. When the tip slides, shear is applied. The maximal shear stress is present at the back of the contact circle. It is given by48

σmax )

[

FL1/3 E 4 2 R 3(1 - ν2)

][ 2/3

]

3µ(4 + ν) 1 - 2ν ≈ + 8 π

()

0.5‚FL1/3

E R

2/3

Here, E is Young’s modulus of the polymer, ν is Poisson’s ratio,

3156 Langmuir, Vol. 23, No. 6, 2007

Berger et al.

and µ is the friction coefficient. The approximation was obtained with typical values µ ) 0.4 and ν ) 0.3. Inserting E ) 3.5 GPa (for high molecular weight PS11), R ) 20 nm (determined from scanning electron micrographs), and FL ) 10 nN, we estimate a maximal shear stress of σmax ≈ 0.7 GPa. This exceeds the tensile strength of most polymers49 and so crack formation is likely. Cross-linking increases the tensile strength because the van der Waals force between neighboring polymer chains is replaced by covalent bonds. In this case, the dependency on scan speed could come into play by the time required to break bonds. To break chemical bonds, an activation energy barrier has to be overcome.50-52 By pulling on a bond, the initial state increases its energy until thermal fluctuations are sufficient to overcome the activation barrier. This leads to a logarithmic dependency on the pulling rate. Several authors have already assumed the presence of an activated process in friction force measured on polymers.14,23 Ripple formation is caused by the release of strain after breaking bonds and the subsequent relaxation of the tip. Another approach focuses on the energy dissipation during scanning.13,14,16,30 Energy can be dissipated by relaxation processes in the polymer. Such relaxations occur at different length scales

from the monomer up to a length of the whole polymer. In crosslinked systems, the maximal backbone relaxation is restricted between two cross-links. All models qualitatively agree with the experimental observation that cross-linking density reduces wear. Future quantitative experiments are required to distinguish between the different models and to find the range of applicability of each model.

(49) Hamada, E.; Kaneko, R. Ultramicroscopy 1992, 42-44, 184-190. (50) Bell, G. I. Science 1978, 200, 618-627. (51) Heymann, B.; Grubmu¨ller, H. Phys. ReV. Lett. 2000, 84 (26), 61266129. (52) Tees, D. F. J.; Waugh, R. E.; Hammer, D. A. Biophys. J. 2001, 80, 668682.

Acknowledgment. We thank J. Thiel, G. Kircher, and T. Wagner for the synthesis of many of the polymers and U. Duerig, T. Altebaeumer, and A. Knoll for fruitful discussions.

Conclusions All experimental observations show that the dominant material parameter with respect to wear of polymers by a nanosized rigid body is the cross-link density. The most wear-resistant films were high-power plasma-polymerized PS and chemically crosslinked polycarbonate. Ordering of linear polymer chains with respect to the scan direction has significant influence on wear. Wear is more pronounced when the chains are oriented perpendicular rather than parallel to the scan direction. When the chains are oriented normal to the surface (diblock copolymer films, brush polymers), they are usually destroyed by the tip because of a lack of lateral stability. Only grafted polymer layers showed low wear caused by the additional stabilization of the chemical grafting to a substrate.

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