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Mechanical Deformation Behavior in Highly Anisotropic Elastomers Made from Ferroelectric Liquid Crystalline Polymers Hanna M. Brodowsky,* Undine-C. Boehnke, and Friedrich Kremer Physik anisotroper Fluide, Universita¨ t Leipzig, Linnestraβe 5, D-04103 Leipzig, Germany
Elisabeth Gebhard and Rudolf Zentel Institut fu¨ r Materialwissenschaften und Fachbereich Chemie, BUGH Wuppertal, Gauss Straβe 20, D-42097 Wuppertal, Germany Received January 5, 1998. In Final Form: July 3, 1998 Cross-linked ferroelectric liquid crystalline polymers are studied by atomic force microscopy. Polysiloxane copolymers have been synthesized with mesogenic and photo-cross-linkable side groups, the latter connected either directly to the backbone via a short spacer or as terminal groups on a part of the mesogens. Although the polymers are otherwise identical, the restrictions imposed on the network formation process by the anisotropy of the smectic mesophase are different for the two positions of the cross-linkable group: In the first case (“intralayer cross-linking”), a predominantly two-dimensional network is formed in the backbone layers separating the smectic layers; in the second case (“interlayer cross-linking”), a primarily threedimensional network is established which is dependent on the mesophase of cross-linking. These elastomers are prepared as thin freely suspended films in homeotropic orientation. The topography consists of plateaus separated by steps of characteristic height, corresponding to the surfaces and edges of smectic layers. If a film of “intralayer cross-linked” elastomer (network formation in the microphase-separated backbone layers) is elongated, the layers can slide on one another, showing occasional tears but no surface roughening (roughness 0.5 nm at 30% elongation). In an “interlayer cross-linked” film (network formation via the mesogens), the three-dimensional network introduces forces perpendicular to the direction of the mechanical deformation, leading to a characteristic depression pattern on the surface which indicates a distortion of the smectic order.
Introduction Elastomers prepared from ferroelectric liquid crystalline (FLC) polymers combine the elasticity of rubbers with the ordered structure and mobility of liquid crystalline phases. They offer a scope of interesting ferroelectric, piezoelectric, and pyroelectric properties.1-5 Recently, LC elastomers were found which show a piezoeffect stronger than that of piezoceramics combined with the material properties of a soft rubber.6 One advantage of elastomeric piezosubstances is the fact that they can be formed into any desired shape. In particular, there are many applications requiring thin films. Besides, their inherent anisotropy leads to an unusual mechanical behavior. To test the mechanical properties as well as to deduce the crystallization of a polymer or the network structure of an elastomer, atomic force microscopy (AFM) imaging of mechnically deformed7-9 polymer samples is proving to * To whom correspondence should be addressed. (1) (a) Meier, W., Finkelmann, H. Macromol. Chem., Rapid. Commun. 1990, 11, 599. (b) Vallerien, U.; Kremer, F.; Fischer, E. W.; Kapitza, H.; Zentel, R.; Poths, H. Macromol. Chem., Rapid. Commun. 1990, 11, 593. (2) Meier, W.; Finkelmann, H. Macromolecules 1993, 26, 1811. (3) Zentel, R. Angew. Chem. Adv. Mater. 1989, 101, 1437. (4) Hikmet, R. A. M. Macromolecules 1992, 25, 5759. (5) Kelly, S. M. J. Mater. Chem. 1995, 5, 2047. (6) Eckert, T.; Finkelmann, H.; Keck, M.; Lehmann, W.; Kremer, F. Macromol. Chem. 1996, 17, 767. (7) (a) Hild, S.; Gutmannsbauer, W.; Lu¨thi, R.; Fuhrmann, J.; Gu¨ntherodt, H.-J. J. Polym. Sci. B 1996, 34, 1953. (b) Hild, S.; Rosa, A.; Marti, O. Submitted to Scanning Probe Microsc. Polym. (8) Drechsler, D.; Karbach, A.; Fuchs, H. Proc. SXM-2 (JVST B), Vienna, Sept. 1996. Surf. Interface Anal. 1997, 25, 537. (9) Suzuki, A.; Yamazaki, M.; Kobiki, Y.; Suzuki, H. Macromolecules 1997, 30, 2350.
be a valuable tool. In this paper, AFM is used to study the effect of mechanical stress on thin films of oriented FLC elastomers; two molecularly differing network structures are compared. To the best of our knowledge, this is the first AFM study of the deformation behavior in unsupported polymer films of submicrometer thickness. Two different FLC polymers are discussed (Figure 1). The substances, statistical copolymers with one side chain in 3.7 polysiloxane units have a degree of polymerization of about 30 and a polydispersity of 1.66. They have either a cross-linkable group attached to the polymer backbone via a short spacer (polymer A) or a terminal cross-linkable group replacing the chiral end group in 10% of the mesogens (polymer B). Their syntheses are described in refs 10 and 11; the behavior of these and similar polymers has been characterized in a number of studies (electooptic switching experiments,10-14 time-resolved Fourier transform infrared spectroscopy,15,16 dielectric spectroscopy,13 small angle X-ray scattering17). In the smectic mesophases the polymers perform a microphase separa(10) Brehmer, M.; Zentel, R. Macromol. Rapid. Commun. 1995, 16, 659. (11) Brehmer, M.; Zentel, R.; Wagenblast, G.; Siemensmeyer, K. Macromol. Chem. Phys. 1994, 195, 1891. (12) Brehmer, M.; Zentel, R.; Giesselmann, F.; Germer, R.; Zugenmaier, P. Liq. Cryst. 1996, 21, 589. (13) Kocot, A.; Wrzalik, R.; Vij, J. K.; Brehmer, M.; Zentel, R. Phys. Rev. B 1994, 50, 16346. (14) Poths, H.; Zentel, R. Liq. Cryst. 1994, 16, 749. (15) Shilov, S. V.; Skupin, H.; Kremer, F.; Gebhard, E.; Zentel, R. Liq. Cryst. 1997, 22, 203. (16) Shilov, S. V.; Skupin, H.; Kremer, F.; Wittig, T.; Zentel, R. Phys. Rev. Lett. 1997, 79, 1686. (17) Diele, S.; Oelsner, S.; Kuschel, F.; Higsen, B.; Ringsdorf, H.; Zentel, R. Macromol. Chem. 1987, 188, 1993.
10.1021/la980021v CCC: $18.00 © 1999 American Chemical Society Published on Web 12/16/1998
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Figure 1. Structures of the two FLC polymers and corresponding phase sequences (Determined by differential scanning calorimetry and optical polarization microscopy with transition temperatures in K. SX*, unidentified higher ordered chiral smectic phase; SC*, chiral smectic C* phase; SA, smectic A phase; i, isotropic phase) along with cartoons of the polymers after cross-linking (acrylate chain marked as dashed line).
tion and the backbones arrange themselves in layers between the smectic layers of the mesogenic side groups. 18
The polymers may be photo-cross-linked to form soft elastomers. They were synthesized so that, for polymer A, the acrylate, limited by the length of the spacers, should predominantly interconnect the polysiloxane backbones within one microphase-separated layer. A preferably twodimensional network is expected (cartoon in Figure 1). This type of cross-linking ought to stabilize an externally set order, while leaving the mesogenic mobility undisturbed. In polymer B, the mesogens with a cross-linkable terminal group should form a network connecting separated layers; as they are fixed on both ends, they hinder the switching chiral mesogens in their motion. A ferroelectric switching should lead to an elastic stress in this soft elastomer. It is a candidate for a rubberlike piezomaterial.11 Recent AFM studies19 of the temperature behavior of these two FLC elastomers have shown that indeed the network formation is of such a two-dimensional, or intralayer, type for polymer A and of a threedimensional, or interlayer, type for polymer B. In this paper, the mechanical deformation of thin oriented films of these elastomers is studied by atomic force microscopy. Experimental Section To enable photo-cross-linking of the polymers, 1 wt % photoinitiator (Lucirin TPO, BASF AG) and 0.5 wt % inhibitor (3-tert-butyl-4-hydroxy-5-methylphenyl sulfide) are added20-22 (the latter to prevent thermal cross-linking). (18) Kunchenko, A. B.; Svetogorsky, D. A. J. Phys. 1986, 47, 2015. (19) Brodowsky, H. M.; Boehnke, U.-C.; Kremer, F.; Gebhard, E.; Zentel, R. Langmuir 1997, 13, 5378. (20) Reibel, J. Ph.D. Thesis, Mainz, 1994.
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Figure 2. Scheme of the sample preparation procedure. (1) Drawing of films: The drawing/stretching apparatus consists of two movable steel blades guided by a 5 × 7 mm2 temperaturecontrolled brass frame (in the cartoon, only one edge is shown in the back). The FLC polymer is placed on the blades and heated to reduce the viscosity, usually to a few degrees below the SA-i transition temperature. If the blades are drawn apart, the smectic layers orient parallel to the film-air interface, forming a homeotropically ordered film of submicrometer thickness suspended from the two blades and the frame edges. (2) The film may be UV cross-linked and then (3) be cut loose from the frame edges, until it is only suspended from the two blades, to which it remains attached due to the cross-linking reaction. (4) If the blades are drawn further apart, the film is stretched. (5) For AFM imaging, it is transferred to a substrate. Freely Suspended Films. The drawing apparatus (Figure 2) consists of two movable 5.5 mm steel blades that are arranged with parallel edges. Their corners slide on a temperaturestabilized 5 × 7 mm2 brass frame. The FLC polymer is placed onto the blades as they touch each other and heated to reduce the viscosity, usually to a few degrees Celsius C below the sA-i transition temperature. As the blades are drawn apart, the LC forms a film, suspended from the two blades and the frame edges. The film is stabilized by the smectic ordering yet flexible due to the fluidity of the substances: At the surfaces, the smectic layers orient parallel to the film-air interface, forming a homeotropically oriented film.23,24 The maximum size of the film is given by the frame; the thickness of the film can be controlled by the interference colors in optical reflection microscopy with an accuracy of 10 nm. Immediately after drawing, the film is very inhomogeneous. However, it smooths out on a time scale of 10 h due to flow of material within the film. Resulting domains are (21) Reibel, J.; Brehmer, M.; Zentel, R.; Decher, G. Adv. Mater. 1995, 7, 849. (22) Gebhard, E.; Brehmer, M.; Zentel, R.; Reibel, J.; Decher, G.; Brodowsky, H. M.; Kremer, F. In The Wiley Polymer Networks Group Review; Nijenhuis, K., Mijs, W., Eds.; John Wiley and Sons Ltd.: Chichester, U.K., 1997; Vol. I. (23) Pieranski, P.; Beliard, L.; Tuornellec, J.-P.; Leoncini, X.; Furtlehner, C.; Dumoulin, H.; Riou, E.; Jouvin, B.; Fe´nerol, J.-P.; Palaric, P.; Heuving, J.; Cartier, B.; Kraus, I. Physica A 1993, 194, 364. (24) Bahr, C.; Fliegner, D. Phys. Rev. Lett. 1993, 70, 1842.
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Figure 3. Reflection micrographs of a cross-linked film of polymer A: (a) undeformed film; (b) the same film as in part a mechanically stretched to 14% relative elongation. The vivid colors are due to interference effects and indicate the local film thickness. To demonstrate that the mechanical stretching is uniform, independent of the local film thickness, the photograph of part a was stretched by an affine image-processing routine to 14% relative elongation and then flipped vertically: the shapes of the thickness domains in the resulting image (c) represent a good mirror image to those of part b; the slightly changed color is due to a thickness change related to the stretching. typically 60 µm wide; the thickness ranges from 8 to beyond 1000 nm (or 2-200 smectic layers) with an average of around 150 nm (Figure 3a). The freely suspended films may then be cross-linked by exposure to UV light (366 nm, 12 W, 3 cm distance of the film to the light source Konrad Benda NU6K1) for 30 min.25 The mechanical stabilization of the cross-linking reaction is evident both on the macroscopic (flow processes, visible by reflection microscopy, are stopped; the film may be cut without rupture) and on the microscopic scales (the elastomers may be imaged by contact mode AFM even at temperatures where the non-crosslinked substances are liquids). Deformation. It is possible to cut the cross-linked film until it is only suspended from two opposing blades. Due to the crosslinking reaction, the film adheres to these blades; if they are drawn apart, it reacts to the mechanical stress with a deformation (Figure 3b). The stretching direction lies within the film plane. At 10% elongation, inelastic deformation has already set in: if the blades are approached again, the film hangs loosely between the blades. Elongations of more than 30% without rupture were rare; at 50% at the latest, film failure was observed. Despite the large thickness variation of the films, the lateral stretching is uniform within the film, independent of the local film thickness (cf. Figure 3c). The films are transferred onto partially gold coated glass substrates as described in ref 26. A solid substrate is (25) Decher, G.; Maclennan, J.; Reibel, J.; Sohling, U. Adv. Mater. 1995, 3, 617. (26) Maclennan, J.; Decher, G.; Sohling, U. Appl. Phys. Lett. 1991, 59, 917. (27) Brodowsky, H. M.; Boehnke, U.-C.; Kremer, F. Rev. Sci. Instrum. 1996, 76, 4198.
Figure 4. Tapping mode AFM images of cross-linked films of polymer A: scale bars, 1 µm; height scale, 25 nm valid for all images; top, Undeformed film; center, stretched to (12% ( 2) relative elongation (The arrow indicates a tear; see text.); bottom, film stretched to (30% ( 4) relative elongation. The surface of the film remains even despite the mechanical deformation. approached from the bottom side and the pressure reduced in the chamber below the film. The film bends downward; once it touches the substrate at one point, it attaches completely to reduce the surface energy. It is then imaged with a commercial atomic force microscope (Nanoscope IIIa, Dimension 3000, Digital Instruments, Santa Barbara, CA) calibrated prior to the measurements,27 unless otherwise stated, in the so-called tapping mode (cantilever resonance frequency ca. 180 kHz, Nanosensors). The delicate nature of the films discussed in this paper has so far impeded either imaging a freely supported film in the drawing/ stretching apparatus or removing a transferred film from a substrate undamaged, so the direct comparison of the same part of the film before stretching and at several elongations could not be made up to now.
Results and Discussion Polymer A. Figure 4a shows a typical image of the top surface of a transferred freely suspended film at room
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Figure 5. Tapping mode AFM images of cross-linked films of polymer B: scale bars, 1 µm; height scale, 25 nm valid for all images; top, undeformed film; center, film stretched to (12% ( 2) relative elongation with different sampling areas of the same film (About half the surface of this film was covered by an obvious depression pattern; the surface topology of the rest was unchanged by the stretching.); bottom, film stretched to (30% ( 4) relative elongation. The depression pattern caused by the mechanical deformation is evident.
temperature. Relatively flat plateaus with a root-meansquare roughness Rrms of 0.5 nm are separated by steps
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of multiples of 4.2 nm height. These are interpreted as the surfaces and edges of smectic layers For relative elongations of (12 ( 2)% and (30 ( 4)%, the major part of the surfaces remains even (Rrms 0.5 nm), the aspect of the film is not much changed except for a number of small holes (100 nm in diameter) and for occasional tears in the surface layers (Figure 4). The holes, which were never seen in an unstretched film, are usually situated at the edge of a plateau; they are too small for the tip to probe their shape or depth, but they are more than one layer deep. The tears have more diffuse edges, both in the form of the edge line and in profile, than the edges of smectic layers in an unelongated film. Very often they are situated close to a significantly thicker, yet homeotropically ordered, part of film, as is seen in the lower left-hand corner in Figure 4 (the order is evident from steps of the thickness of a smectic layer visible in the scan profile). In a 12% elongated film, these tears are usually only 1 layer deep, but they deepen as the film is elongated further. Apart from these defects, the deformation causes no change in the surface topography, the surface remains even for deformations up to 30%, and the characteristic surface patterns remain unchanged without developing a preferential direction in or perpendicular to the stretching direction. Although relative elongations up to 50% were occasionally observed, the films always broke before they could be transferred to a substrate; this can be explained by instabilities due to the defects discussed above. The fact that the film stays even may be explained by the nature of the intralayer cross-linking: As the crosslinking takes place predominantly within the backbone layers to form a two-dimensional network with relatively few knots connecting the layers, vertical forces, which would lead to a roughening of the surface, are negligible. Either existing connections between layers break or the end of the polymer chain may move within the backbone layer (the polysiloxane chains are 30 units long, with on average three cross-linkable groups). This is supported by the width of some of the tears (sometimes more than 1 µm wide): obviously the cross-linked backbone layers may slide over each other with the “liquid” mesogen layers between them. Polymer B. For Polymer B, where the network formation occurs via 10% of the mesogens that contain a terminal cross-linkable acrylamide group instead of the chiral group, the resulting three-dimensional network reacts differently to the mechanical stress: Elongation of a crosslinked film leads to a characteristic roughening of the surface. In an undeformed, homeotropically ordered film of polymer B, the surface of a plateau is even, as for polymer A (Figure 5). For small deformations, this remains unchanged, exept for small holes that appear close to the steps, as in polymer A. However, beyond 12% elongation, an isotropic, irregular depression pattern is formed. The pattern is not influenced by the edges of the smectic layers but passes over them, indicating that it is not an effect of a single layer but of the “bulk” of the film. At 12% deformation, about half the surface is covered by this depression pattern; the rest is as even as an undeformed film (Figure 5 center images). In the places where depressions are observed, they are around 4 µm wide with a height variation of about 5 nm. The surfaces of the depressions are smooth; on a 1 µm scale, no difference in the surface topography, as compared to that of a nonelongated film, was detected. At 30% elongation, the surface is evenly covered by depression patterns visible
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Figure 6. Root-mean-square roughness Rrms determined for images of different sizes to characterize the surface patterns resulting from the deformation. Filled symbols and solid lines, polymer A: 9, undeformed film; 1, 10% elongation; [, 30% elongation. Hollow symbols and dashed lines, polymer B: 0, undeformed film; 3, 12% elongation, even area; 4, 12% elongation, area of the same film covered by a depression pattern; ], 30% elongation. For undeformed films of either polymer, this Rrms is 0.5 nm for images larger than 1 µm2 and diminishes as the size of the image approaches that of the structural details. In polymer A, the surface patterns and therefore the Rrms are not changed by the deformation process. In contrast, in polymer B, the stretching leads to a change in topology that sets in at about 10% deformation. For 12% deformation, there are two distinct surface morphologies: one unaltered by the stretching and another displaying an additional surface pattern on a 2 µm scale leading to a higher Rrms on this length scale. For 30% deformation, the surface is uneven on all length scales.
on all length scales (4 nm height at 3 µm width, 1.3 nm height at 300 nm width, 0.5 nm height at 30 nm width). This behavior may be explained by the concept of an interlayer network. The mechanical stress that the sample is submitted to is purely within the film plane, but due to the three-dimensional character of the network structure, transversal effects arise. For small deformations, they may be suppressed by a rearranging of the mesogens and backbones, as the system tries to keep a smectic order, but beyond a threshold of ca. 12%, deformations result in a changed surface topography. In a recent publication, Nishikawa et al.28 studied the viscoelastic behavior of a bulk sample of a mechanically oriented SmA elastomer. The modulus along the plane normal was found to be comparable to that for the lowmolecular-weight smectic liquid crystal, while that within the layer plane was 2 orders of magnitude lower and induced by the network. The network structure may be related to that of polymer B, but as the system is mechanically oriented during the cross-linking reaction and not, as are the samples of the present paper, cross(28) Nishikawa, E.; Finkelmann, H.; Brand, H. R. Macromol. Rapid Commun. 1997, 18, 65.
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linked in a macroscopically oriented near-equilibrium state, the network architecture is probably more complex. Besides, the stretching behavior of the film bulk was studied with stress versus strain measurements, which provide different information on anisotropic elastomers than microscopy of the surface topography. However, a common result of both studies is that the smectic layering is not destroyed by a deformation within the film plane, indicating that the in plane mobility of the mesogens is high (“liquid-like” according to ref 28) despite the crosslinking. Comparison. To characterize the surface patterns resulting from the deformation, in Figure 6 the surface Rrms for both polymers and a variety of sampling sizes is plotted logarithmically. All values were obtained under equivalent scanning conditions, most of them with the same tip. For undeformed films of either polymer, this Rrms is 0.5 nm for images larger than 1 µm2. As each image is flattened to remove a slope, the Rrms diminishes as the size of the image is decreased to that of the structural details. In polymer A, the surface patterns, and therefore the Rrms, are not changed by the deformation process on any length scale. In contrast, in polymer B, the stretching leads to a change in topology: For 12% deformation, there is an additional surface pattern on a 5 µm scale, while on a smaller scale the surface topology remains unchanged: the Rrms calculated on small areas is unchanged but jumps to a much higher value for a 5 × 5 µm2 surface. For 30% deformation, the surface is uneven compared to that of an untreated film on all length scales. Conclusion To analyze the impact of the molecular structure on the network formation in ferroelectric liquid crystalline polymers, two copolymers are studied, which are identical but for the molecular position of the cross-linkable group: (i) polymer A with cross-linkable groups attached to the backbone via a short spacer and (ii) Polymer B with the cross-linkable group in the terminal position of a mesogenic side group. These molecular particularities result in two distinct network structures: For polymer A, intralayer cross-linking results in two-dimensional networks in each backbone layer, separated by liquid-like FLC side group layers. In polymer B a three-dimensional network is formed. When mechanical stress is imposed on thin films in homeotropic orientation by stretching, the two elastomers react differently to the deformation: In polymer A each layer is stretched, but as there are no vertical connections in this intralayer network, no vertical distortions occur, and the layers may slide on each other. In polymer B the system reacts with mesogen and backbone reorientation, as well as, beyond 10% deformation, a distortion of the smectic layering. Acknowledgment. The authors wish to thank the Innovationskolleg “Pha¨nomene an den Miniaturisierungsgrenzen” (Projekt F) for financial support. The European Community is thanked for an HCM grant which facilitated the cooperation between both groups. LA980021V
