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Molecular orientation and conductivity in highly drawn poly(p-phenylene vinylene)
,Synthetic Metals, 20 (1987) 85 - 95
85
MOLECULAR ORIENTATION AND CONDUCTIVITY IN HIGHLY DRAWN POLY(p-PHENYLENE VINYLENE) D A V I D R. G A G N O N * , F R A N K R O B E R T W. L E N Z
E. K A R A S Z , E D W I N L. T H O M A S
and
Department of Polymer Science and Engineering, University of Massachusetts, Amherst, MA 01003 (U.S.A.) (Received October 17, 1986 ; in revised form December 4, 1986 ; accepted December 5, 1986)
Abstract Films of poly(p-phenylene vinylene), PPV, have been prepared with various degrees of uniaxial orientation by thermal conversion and drawing of cast films of a precursor polymer. Molecular orientation was determined by X-ray diffraction and infrared dichroism as a function of the draw ratio. The crystallite size transverse to the stretch direction was also estimated by X-ray line broadening. Electrical conductivities of the PPV films, doped with AsFs, show a substantial enhancement in conductivity in the orientation direction, which correlates well with the degree of molecular orientation. Conductivity in the direction perpendicular to the stretch direction decreases with increased orientation. These data may be understood in terms of the relative contributions of intrachain and interchain transport of the electronic charge carriers.
Introduction Recent advances in conjugated conducting polymer syntheses through soluble precursor polymers permit the preparation of materials having wellordered and uniform morphologies previously inaccessible in direct syntheses [1 -3]. Highly conjugated polymers are generally insoluble, infusible and intractable because of the rigidity of their backbones, and the nascent polymers resist major morphological changes. For example, polyacetylene, PA, probably the most widely studied conducting polymer, when synthesized by the usual Shirakawa route [4] is an inhomogeneous random mat of semicrystalline fibrils, which renders interpretation of macroscopic conductivity on a molecular charge transport scale difficult [5, 6]. Other potentially conducting polymers in forms such as pressed powders of poly(p-phenylene vinylene) obtained by the Wittig m e t h o d [7], or electrochemically prepared *Present address: 3M Center, Building 219-1-01, CRTDL, St. Paul, MN 55144, U.S.A. 0379-6779/87/$3.50
© Elsevier Sequoia/Printed in The Netherlands
86 films of poly(thiophene) [8] or poly(pyrrole) [9], have similar morphological inhomogeneities that compromise the analysis of their charge transport properties. Because of these morphological limitations, fundamental questions regarding the mobility of the charge carriers along the conjugated backbone (intrachain transport), and the contribution of interchain mobility to conduction have yet to be answered satisfactorily. In the case of PA, attempts to orient the fibrillar films by uniaxial extension have resulted in the orientation of the fibrils alone without appreciably changing the molecular structure [10]. The conductivity of oriented fibrillar films increased in the draw direction by up to a factor of about eight, which can probably be explained in terms of the alignment and consolidation of the already anisotropic fibrils. Other methods to obtain oriented PA during the Shirakawa synthesis, such as shear flow polymerization of the film onto a rotating stator [11] or epitaxial polymerization on to crystalline substrates [12, 13], have been reported, but in each case the orientation was of the fibrillar morphology. The significant advantage of the precursor method is that an amorphous film of the appropriate precursor polymer may be cast from solution prior to the thermal or chemical conversion of the latter into homogeneous films of the conjugated polymer [1, 3]. Concomitantly, during the thermal conversion process the films may be highly stretched, thus enabling continuous films of the conjugated polymer to be prepared with varying degrees of uniform molecular orientation [14]. This makes it possible to study conductivity and charge transport structure/property relationships without the ambiguity associated with the inhomogeneous fibrillar of powder morphologies. The most widely used precursor routes to conjugated polymers have been in the preparation of PA following the method of Edwards and Feast [1] and in the preparation of PPV by the method of Wessling and Zimmerman [3]. The respective high molecular weight [15, 16] precursor polymers (> l 0 s Daltons) are relatively easily synthesized in both cases, and may be cast into films with good mechanical integrity. The films can then be stretched to over ten times their initial length during the elimination step, thus imparting extremely high molecular orientation and very large electrical anisotropy after reaction with an appropriate dopant [17, 18]. In order to account for the anisotropic electrical properties in the oriented polymers, it is first necessary to understand the crystallographic and morphological structure. In this report, we discuss the optimization of the processing and the resulting degree of molecular orientation of PPV films, as a function of the draw ratio. The materials were synthesized via the water-soluble precursor polymer, poly(p-xylylene-~-dimethylsulfonium chloride) following an improved procedure of Wessling and Zimmerman. Orientation was characterized using both X-ray diffraction and infrared dichroism, and is compared with the conductivity of AsFs-doped samples measured parallel and perpendicular to the stretch direction at a series of draw ratios. The crystal struc-
87 ture of highly oriented PPV obtained from electron diffraction data has been published elsewhere [19].
Experimental
The synthesis of the monomer, p-xylylene-bis(dimethylsulfonium chloride), and the precursor polyelectrolyte, poly(p-xylylene~x-dimethylsulfonium chloride) has been described in detail elsewhere [20, 16]. Solutions of the polyelectrolyte were cast into free-standing films in vacuo at or below room temperature. Oriented PPV films were obtained by the uniaxial stretching of approximately 1.5 cm X 5 cm X 20 gm (thickness) precursor films at a temperature between 115 °C and 180 °C in a simple zone heating apparatus. The latter consisted of two glass slides separated by 2 mm thick Teflon spacers and wrapped with a nichrome wire heater. The sample temperature was monitored by a thermocouple placed at the center of the heating zone. The portion of the film that was within the heating zone ( approximately 1.5 cm) deformed with necking at temperatures greater then 115 °C. One portion of the 'neck' could be maintained within the heating zone as the film was manually stretched until the neck propagated to the end of the sample. Then the film was heated until the draw ratio was uniform throughout the film length. The draw ratio was determined by measuring the distance between bench marks. Final thermal elimination giving quantitative conversion to PPV was performed by heating under vacuum at approximately 370 °C for at least four hours. The orientation functions for samples of varying draw ratios were calculated using a Siemens D-500 X-ray Diffractometer fitted with a Huber four-circle goniometer employing Cu K s radiation with a receiving slit width of 0.15 mm. Because PPV exhibits considerable chain axis paracrystallinity [19], meridional reflections are streaked normal to the chain axis direction, prohibiting use of the normal (00/) t y p e reflections for orientation analysis. For this reason the strong (110) reflection was utilized. The crystailite size transverse to the orientation direction was estimated from Scherrer analysis of the linewidth of the (210) reflection on samples of various draw ratios. Corrections for instrumental broadening were carried out by measuring the width of the (200} reflection of hexamethylene tetramine. Polarized infrared dichroism studies were performed with an IBM IR-38 Fourier transform instrument fitted with a Perkin-Elmer gold wire grid polarizer on a rotatable mount. To minimize the effect of instrumental vertical polarization of the i.r. beam, the PPV sample was m o u n t e d with its stretch direction at 45 °C to the vertical, and the polarizer directions used were - 45 ° to vertical (90 ° to sample) and 45 ° to vertical (0 ° to sample). The vinyl C--H o u t of plane bending absorption at 963 cm -1 was used in these dichroism studies.
88 AsF s doping and conductivity measurements were made as described previously [18]. The concentration of AsF s in the doped samples was nearly constant at 50 + 5 wt.%, represented on a molar basis by [(CsH6)~79(AsF6)o.21]..
Results and discussion Upon removal of water during the casting of the polyelectrolyte precursor, a small amount of sulfonium elimination occurs to give a film represented by eqn. (1), in which the degree of elimination is defined by the m :n ratio of the repeat units:
__
CH3---~CH3
Since water stabilizes the sulfonium group, some further elimination in the cast films occurs with time even at room temperature. For example, a freshly cast film had an m :n ratio of above 4:1, which stabilized to a ratio of approximately 1:2 after three weeks at room temperature. The m :n composition is important when considering the details of the conversion and orientation. The precursor film could only be stretched to high draw at a temperature above a b o u t 115 °C, which coincides with that of the first major elimination process [16]. Upon heating, the film apparently becomes 'plasticized' by the volatile gases evolved. The a m o u n t of these gases is determined by the m : n ratio; whereas the temperature protocol determines the elimination rate, which is directly related to the instantaneous concentration of these 'plasticizing' gases. Above a temperature of a b o u t 180 °C, the elimination rate is so rapid that the products are generated faster than they can diffuse o u t of the film and voids are formed during the draw. Macroscopic voids are also noticeable when precursor films have initial thicknesses greater than 20 pm. Thus the optimal conditions for obtaining high draw and uniform orientation are to start with a precursor film less than 20 #m thick with an m : n ratio of approximately 2:1 (nearly freshly cast), and to heat the film under tension rapidly to a temperature between 115 °C and 180 °C. Final elimination of the drawn precursor films to PPV is performed in vacuum at a temperature between 360 - 370 °C for four hours. This results in a quantitative conversion to PPV. The crystal structure for highly Oriented PPV films prepared as described above has been derived from electron diffraction data [19]. PPV has a monoclinic cell and exhibits significant paracrystalline disorder along the chain axis direction. For highly oriented samples, meridional layer lines consist of diffuse streaks of intensity resulting from partial axial translational disorder of the PPV chains.
89
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0.8
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2
I
4
,
~
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8
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Fig. 1. Herman's molecular orientation function, f, of PPV films determined by X-ray diffraction (open circles) and by infrared dichroism (filled circles)as a function of the
draw ratio, ~. The degree of molecular orientation in samples with varying draw ratios is shown in Fig. 1. The data represented by open circles were calculated from the measurement of the (110) X-ray diffraction intensity measured as a function of the azimuthal angle, ¢ [21]. Since stacking of multiple films could have added artificial disorder, measurements were made on single films less than 10/~m thick. The (110) reflection was chosen because of its large intensity. The orientation function, f, was estimated from
f=
3 cos2<~b}-- 1 2
(2)
90 where the average aximuthal angle, (¢>, was taken to be the full angular width at half the maximum (FWHM) of the intensity distribution of the reflection. The more exact integral m e t h o d of calculating the cosine average from lineshape analysis was n o t possible because of uncertainties in the lineshape of the intensity distribution at high ¢ angles, especially with films of lower draw ratio. The FWHM m e t h o d was less sensitive to lineshape ambiguity and provided sufficient precision for the analysis to show the relation of f to the draw ratio, ),. The values reported refer to the average orientation of the normal to the (110) plane relative to the perpendicular to the macroscopic orientation direction. Because of the rigidity of the chains, it has been assumed that the (110) orientation yields an acceptable relative measure of the c-axis orientation with respect to the draw direction. Infrared dichroism was also used to characterize the molecular orientation. The dichroic ratio, D, is defined by the ratio of the absorptions parallel, (All) and perpendicular (A±) to the reference direction. The orientation function in this case is calculated from
f=
[I -llt 0+211 \ D + 2 / \ D 0 - - 1/]
(3)
where Do = 2 cot 2 ~, and ~ is the angle of the absorption transition m o m e n t relative to the reference direction [21]. Absorbances associated with stretching vibrational modes are typically chosen for i.r. dichroism studies because ~ can then be unequivocally established. However, in PPV, the C--H stretch region centered at 3021 cm -1 was found to contain overlapping bands, which makes absorption measurements uncertain. In contrast, the trans-vinylene C--H o u t of plane bending absorption at 963 cm -1 is strong and is not overlapped. Although the angle a is n o t known absolutely, the trans-vinyl C--H bend transition m o m e n t is considered to be nearly perpendicular to the plane of the double bond, and would thereby be nearly perpendicular with respect to the PPV chain axis. Assuming the 963 cm -I transition m o m e n t is perpendicular to the chain direction, we will again define the axis perpendicular to the orientation direction as the frame of reference; thus c~ -~ 0 ° and f becomes simply D--1 f -~ ~ D+2
(4)
The infrared orientation function (fiR) is shown in Fig. 1 by filled circles and the X-ray orientation function (fx-ray) by the open circles. There is reasonably good agreement between the t w o measurements of orientation: both orientation functions increase rapidly with increasing draw ratio up to ~ 6, above which additional drawing has a less significant effect. The values of the orientation functions are relatively high even at low draw ratios. The efficiency of the draw process can be understood when one considers that as the flexible polyelectrolyte chain is being stretched it is
91 irreversibly converted to a rigid conjugated chain. It should also be noted that the total sample volume decreases by a b o u t 50% as the volatile elimination products diffuse o u t of the film, so that the effective draw ratio is actually higher than that measured. The observed difference between fir and fx-ray can be explained in two ways. The structure of the chain is n o t known in sufficient detail to exclude the possibility that the effective transition m o m e n t is n o t precisely perpendicular to the chain axis. The measured ratios of fir and fx-ray can conceivably be used to calculate the angle a. For X = 8.3, this calculation yields = 8 °, for X = 2.7, a = 19 °. The former is in fact in excellent agreement with recent interpretations of the transition m o m e n t from measurements by Bradley et al. [22, 23]. An alternative, non-exclusive possibility is that the X-ray intensities are sensitive more to the chains that are well ordered or in larger crystal domains, while the i.r.-derived orientation function is insensitive to local disorder. In this interpretation fx-r,y should be considered as the upper b o u n d of the true orientation. An average lateral crystallite size was obtained from a Scherrer analysis of the 20 breadth of the equatorial reflections [24]. The (210) reflection was chosen because it is the lowest order well-separated reflection. Crystallite size, t, as a function of the draw ratio, was calculated from t -
KA
{4) ~3~o=cos 0 assuming flco= = [t32 -t3std 2] in where K is the Scherrer constant, A is the Cu Ka X-ray wavelength, 0 is the Bragg angle, ~ is the full width at half maxim u m (FWHM) of the 20 scan; ~std is the FWHM for the standard (hexamethylene tetramine) used to approximate the instrumental broadening assuming a Gaussian lineshape. The precise value of K (~ 1) has been in dispute [25]. Since the object of this experiment was to examine the trend of crystallite size as a function of draw ratio, K was arbitrarily taken to be 1.0. Figure 2 shows the calculated lateral crystallite size as a function of draw ratio. It must be emphasized that the value of t is subject to the error associated with the Scherrer constant and with the fact that the treatment also does n o t account for lattice reflection broadening due to lattice distortions of the second kind (i.e., paracrystallinity). These uncertainties are, however, largely independent of X, thus it may be concluded that the crystallite size does n o t increase with increasing draw ratio to the same extent as does molecular orientation. The increase in the parallel conductivity with draw ratio discussed below must therefore be primarily the effect of the increase in molecular orientation. The axial conductivity (oll) of AsFs-doped oriented PPV as a function of draw ratio is shown in the upper curve of Fig. 3. The trend is similar to that seen with the orientation function. Thus log a increases rapidly with draw until X ~ 6 - 8 and then becomes rather insensitive to further stretching
92 130
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~
120
110
!
I
I
I
I 8
I 10
I 12
I 14
f
10C
90
8(
I 2
I 4
I 6
Fig. 2. Transverse crystallite correlation length, t, of PPV films as a function of draw ratio k.
above a draw ratio of about 10. This correlation provides evidence that the axial conductivity is quite sensitive to the degree of orientation of the conjugated polymer chains. On the other hand, the conductivity transverse to the draw direction ( o i l shown in the lower curve of Fig. 3, decreases with increasing draw ratio, although the dependence on X is far less. A possible explanation for the behavior of o± is that the film becomes fibrillated during the drawing process. However, scanning electron microscopy did not reveal any significant fibrillization until the draw ratio was above 10, and with
93
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Fig. 3. Conductivity of AsFs-doped PPV parallel (Oil) and perpendicular (o±) to the stretching direction as a function of draw ratio k. Fig. 4. The conductivity shown in Fig. 3 as a function of the orientation function, f.
careful stretching, unfibrillated films could be prepared even up to k = 12. It may be noted that with the present stretching apparatus, the PPV films drawn to k > 6 are t o o narrow to investigate conductivity in the transverse configuration with high precision. Figure 4 shows the axial conductivity as a function of the X-ray derived orientation function. Again oll increases strongly with f and continues to increase even after the maximum orientation value of 0.96 has been reached. This also indicates that the second-order orientation function cannot completely determine the effect of orientation upon conductivity. This behavior is also c o m m o n l y observed in studies of the relation of modulus and thermal conductivity to stretching in many polymers. Higher-order determinations of orientation are required to characterize this effect more precisely. A further consideration is that the orientation was measured on u n d o p e d samples and it is n o t inconceivable that the dopant molecules have a significant effect on the local structure of the macromolecule. Future studies will characterize the microstructures of PPV samples after doping. Conclusion X-ray and infrared dichroism investigationshave demonstrated that the molecular orientation of stretched P P V films increases strongly as a function
94
of the macroscopic draw ratio. This effect is enhanced because there occurs an irreversible conversion to extremely rigid conjugated PPV sequences, which retard chain relaxation during the stretching process in addition to the normal deformation orientation process. There is a modest increase in the lateral crystallite size with increased stretching. The increase in conductivity parallel to the orientation direction provides evidence pointing to the dominant role of intrachain diffusion of the charge carriers along the PPV backbone. Certainly if interchain conduction (via hopping or some other mechanism) were the sole path of charge transport, there should be an increase in conductivity perpendicular to the orientation direction due to an increase in the lateral order, and a very small or negative effect parallel to the chain direction. It is obvious that a single chain does n o t extend throughout the entire length of the sample, so interchain transport is an important part of the mechanism; this study seems to indicate that interchain transport is the rate-limiting step. Independent measurements of the anisotropy in charge carrier mobility (e.g., by Hall effect studies) are needed to confirm this interpretation.
Acknowledgements This work was supported by A F O S R Grant 85-0044. One of us (D.R.G.) thanks the P.I.A. for additional support. We would also like to thank Professor S. L. Hsu for helpful discussions concerning the infrared dichroism studies, and Dr. Thierry Granier for discussions concerning the crystal structure of PPV.
References 1 J . H . Edwards and W. J. Feast, Polyrn. Commun., 21 (1980) 595. 2 H. Kahlert and G. Leising, MoL Cryst. Liq. Cryst., 117 (1985} 1. 3 R. A. Wessling and R. G. Zimmerman, (a) U.S. Pat. 3 401 152 (1968); (b} U.S. Pat. 3 706 677 {1972}. 4 J. C. W. Chien,Polyacetylene, Academic Press, New York, 1984. 5 H. Shirakawa and S. Ikeda,Polym. J., 2 (1971} 23. 6 C. R. Fincher, D. Moses, A. J. Heeger and A. G. MacDiarmid, Synth. Met., 6 {1983) 243. 7 K. D. Gourley, C. P. Lillya, J. R. Reynolds and J. C. W. Chien, Macromolecules, 17 (1984) 1025. 8 G. Tourillon and F. Garnier, J. Polym. Sci. (Phys.), 22 (1984) 33. 9 M. Salmon, A. F. Diaz, A. J. Logan, M. Krounbi and J. Bargon, Mol. Cryst. Liq. Cryst., 83 (1982) 265. 10 Y. W. Park, M. A. Druy, C. K. Chiang, A. G. MacDiarmid, A. J. Heeger, H. Shirakawa and S. Ikeda, J. Polym. Sci. (Lett.), 17 (1979} 195. 11 W. H. Meyer, MoL Cryst. Liq. Cryst., 77 (1981) 147. 12 T. Woerner, A. G. MacDiarmid, A. Feldblum and A. J. Heeger, J. Polym. Sci. (Lett.), 22 (1984} 119. 13 M. Ozaki, Y. Ikeda and T. Arakawa, Mol. Cryst. Liq. Cryst., 117 (1985} 39.
95 14 D. R. Gagnon, J. D. Capistran, F. E. Karasz and R. W. Lenz, Am. Chem. Soc. Polym. Prepr., 25 (2) (1984) 284. 15 K. Harper and P. G. James, Mol. Cryst. Liq. Cryst., 11 7 (1985) 55. 16 R. W. Lenz, F. E. Karasz, J. D. Capistran, D. R. Gagnon and S. Antoun, submitted to
Macromolecules. 17 G. Leising, H. Kahlert and O. Leitner, in B. Kramer et al. (eds.), Springer Series in Solid S tate Physics, Springer -Verlag, Berlin, 1985, Vol. 61. 18 D. R. Gagnon, J. D. Capistran, F. E. Karasz and R. W. Lenz,Polym. Bull., 12 (1984) 93. 19 T. Granier, E. L. Thomas, D. R. Gagnon, F. E. Karasz and R. W. Lenz, J. PolymerSci. (Phys.), in press. 20 J. D. Capistran, D. R. Gagnon, S. A. George, R. W. Lenz and F. E. Karasz, Am. Chem. Soc. Polym. Prepr., 25 (2) (1984) 282. 21 B. E. Read, in I. M. Ward (ed.), Structure and Properties of Oriented Polymers, Wiley, New York, 1975, Ch. 4. 22 D. D. C. Bradley, R. H. Friend, T. Hartman, E. A. Marseglia, M. M. Sokolowski and P. D. Townsend, ICMS '86, Kyoto, June, 1986, Synth. Met., 17 (1987) 473. 23 D. D. Bradley, R. H. Friend, H. Lindenberger and S. Roth, Polymer, in press. 24 M. Kakudo and N. Kasai, X-ray Diffractions by Polymers, Elsevier, New York, 1972. 25 B. E. Warren,Phys. Rev., 59 (1941)693.
85
MOLECULAR ORIENTATION AND CONDUCTIVITY IN HIGHLY DRAWN POLY(p-PHENYLENE VINYLENE) D A V I D R. G A G N O N * , F R A N K R O B E R T W. L E N Z
E. K A R A S Z , E D W I N L. T H O M A S
and
Department of Polymer Science and Engineering, University of Massachusetts, Amherst, MA 01003 (U.S.A.) (Received October 17, 1986 ; in revised form December 4, 1986 ; accepted December 5, 1986)
Abstract Films of poly(p-phenylene vinylene), PPV, have been prepared with various degrees of uniaxial orientation by thermal conversion and drawing of cast films of a precursor polymer. Molecular orientation was determined by X-ray diffraction and infrared dichroism as a function of the draw ratio. The crystallite size transverse to the stretch direction was also estimated by X-ray line broadening. Electrical conductivities of the PPV films, doped with AsFs, show a substantial enhancement in conductivity in the orientation direction, which correlates well with the degree of molecular orientation. Conductivity in the direction perpendicular to the stretch direction decreases with increased orientation. These data may be understood in terms of the relative contributions of intrachain and interchain transport of the electronic charge carriers.
Introduction Recent advances in conjugated conducting polymer syntheses through soluble precursor polymers permit the preparation of materials having wellordered and uniform morphologies previously inaccessible in direct syntheses [1 -3]. Highly conjugated polymers are generally insoluble, infusible and intractable because of the rigidity of their backbones, and the nascent polymers resist major morphological changes. For example, polyacetylene, PA, probably the most widely studied conducting polymer, when synthesized by the usual Shirakawa route [4] is an inhomogeneous random mat of semicrystalline fibrils, which renders interpretation of macroscopic conductivity on a molecular charge transport scale difficult [5, 6]. Other potentially conducting polymers in forms such as pressed powders of poly(p-phenylene vinylene) obtained by the Wittig m e t h o d [7], or electrochemically prepared *Present address: 3M Center, Building 219-1-01, CRTDL, St. Paul, MN 55144, U.S.A. 0379-6779/87/$3.50
© Elsevier Sequoia/Printed in The Netherlands
86 films of poly(thiophene) [8] or poly(pyrrole) [9], have similar morphological inhomogeneities that compromise the analysis of their charge transport properties. Because of these morphological limitations, fundamental questions regarding the mobility of the charge carriers along the conjugated backbone (intrachain transport), and the contribution of interchain mobility to conduction have yet to be answered satisfactorily. In the case of PA, attempts to orient the fibrillar films by uniaxial extension have resulted in the orientation of the fibrils alone without appreciably changing the molecular structure [10]. The conductivity of oriented fibrillar films increased in the draw direction by up to a factor of about eight, which can probably be explained in terms of the alignment and consolidation of the already anisotropic fibrils. Other methods to obtain oriented PA during the Shirakawa synthesis, such as shear flow polymerization of the film onto a rotating stator [11] or epitaxial polymerization on to crystalline substrates [12, 13], have been reported, but in each case the orientation was of the fibrillar morphology. The significant advantage of the precursor method is that an amorphous film of the appropriate precursor polymer may be cast from solution prior to the thermal or chemical conversion of the latter into homogeneous films of the conjugated polymer [1, 3]. Concomitantly, during the thermal conversion process the films may be highly stretched, thus enabling continuous films of the conjugated polymer to be prepared with varying degrees of uniform molecular orientation [14]. This makes it possible to study conductivity and charge transport structure/property relationships without the ambiguity associated with the inhomogeneous fibrillar of powder morphologies. The most widely used precursor routes to conjugated polymers have been in the preparation of PA following the method of Edwards and Feast [1] and in the preparation of PPV by the method of Wessling and Zimmerman [3]. The respective high molecular weight [15, 16] precursor polymers (> l 0 s Daltons) are relatively easily synthesized in both cases, and may be cast into films with good mechanical integrity. The films can then be stretched to over ten times their initial length during the elimination step, thus imparting extremely high molecular orientation and very large electrical anisotropy after reaction with an appropriate dopant [17, 18]. In order to account for the anisotropic electrical properties in the oriented polymers, it is first necessary to understand the crystallographic and morphological structure. In this report, we discuss the optimization of the processing and the resulting degree of molecular orientation of PPV films, as a function of the draw ratio. The materials were synthesized via the water-soluble precursor polymer, poly(p-xylylene-~-dimethylsulfonium chloride) following an improved procedure of Wessling and Zimmerman. Orientation was characterized using both X-ray diffraction and infrared dichroism, and is compared with the conductivity of AsFs-doped samples measured parallel and perpendicular to the stretch direction at a series of draw ratios. The crystal struc-
87 ture of highly oriented PPV obtained from electron diffraction data has been published elsewhere [19].
Experimental
The synthesis of the monomer, p-xylylene-bis(dimethylsulfonium chloride), and the precursor polyelectrolyte, poly(p-xylylene~x-dimethylsulfonium chloride) has been described in detail elsewhere [20, 16]. Solutions of the polyelectrolyte were cast into free-standing films in vacuo at or below room temperature. Oriented PPV films were obtained by the uniaxial stretching of approximately 1.5 cm X 5 cm X 20 gm (thickness) precursor films at a temperature between 115 °C and 180 °C in a simple zone heating apparatus. The latter consisted of two glass slides separated by 2 mm thick Teflon spacers and wrapped with a nichrome wire heater. The sample temperature was monitored by a thermocouple placed at the center of the heating zone. The portion of the film that was within the heating zone ( approximately 1.5 cm) deformed with necking at temperatures greater then 115 °C. One portion of the 'neck' could be maintained within the heating zone as the film was manually stretched until the neck propagated to the end of the sample. Then the film was heated until the draw ratio was uniform throughout the film length. The draw ratio was determined by measuring the distance between bench marks. Final thermal elimination giving quantitative conversion to PPV was performed by heating under vacuum at approximately 370 °C for at least four hours. The orientation functions for samples of varying draw ratios were calculated using a Siemens D-500 X-ray Diffractometer fitted with a Huber four-circle goniometer employing Cu K s radiation with a receiving slit width of 0.15 mm. Because PPV exhibits considerable chain axis paracrystallinity [19], meridional reflections are streaked normal to the chain axis direction, prohibiting use of the normal (00/) t y p e reflections for orientation analysis. For this reason the strong (110) reflection was utilized. The crystailite size transverse to the orientation direction was estimated from Scherrer analysis of the linewidth of the (210) reflection on samples of various draw ratios. Corrections for instrumental broadening were carried out by measuring the width of the (200} reflection of hexamethylene tetramine. Polarized infrared dichroism studies were performed with an IBM IR-38 Fourier transform instrument fitted with a Perkin-Elmer gold wire grid polarizer on a rotatable mount. To minimize the effect of instrumental vertical polarization of the i.r. beam, the PPV sample was m o u n t e d with its stretch direction at 45 °C to the vertical, and the polarizer directions used were - 45 ° to vertical (90 ° to sample) and 45 ° to vertical (0 ° to sample). The vinyl C--H o u t of plane bending absorption at 963 cm -1 was used in these dichroism studies.
88 AsF s doping and conductivity measurements were made as described previously [18]. The concentration of AsF s in the doped samples was nearly constant at 50 + 5 wt.%, represented on a molar basis by [(CsH6)~79(AsF6)o.21]..
Results and discussion Upon removal of water during the casting of the polyelectrolyte precursor, a small amount of sulfonium elimination occurs to give a film represented by eqn. (1), in which the degree of elimination is defined by the m :n ratio of the repeat units:
__
CH3---~CH3
Since water stabilizes the sulfonium group, some further elimination in the cast films occurs with time even at room temperature. For example, a freshly cast film had an m :n ratio of above 4:1, which stabilized to a ratio of approximately 1:2 after three weeks at room temperature. The m :n composition is important when considering the details of the conversion and orientation. The precursor film could only be stretched to high draw at a temperature above a b o u t 115 °C, which coincides with that of the first major elimination process [16]. Upon heating, the film apparently becomes 'plasticized' by the volatile gases evolved. The a m o u n t of these gases is determined by the m : n ratio; whereas the temperature protocol determines the elimination rate, which is directly related to the instantaneous concentration of these 'plasticizing' gases. Above a temperature of a b o u t 180 °C, the elimination rate is so rapid that the products are generated faster than they can diffuse o u t of the film and voids are formed during the draw. Macroscopic voids are also noticeable when precursor films have initial thicknesses greater than 20 pm. Thus the optimal conditions for obtaining high draw and uniform orientation are to start with a precursor film less than 20 #m thick with an m : n ratio of approximately 2:1 (nearly freshly cast), and to heat the film under tension rapidly to a temperature between 115 °C and 180 °C. Final elimination of the drawn precursor films to PPV is performed in vacuum at a temperature between 360 - 370 °C for four hours. This results in a quantitative conversion to PPV. The crystal structure for highly Oriented PPV films prepared as described above has been derived from electron diffraction data [19]. PPV has a monoclinic cell and exhibits significant paracrystalline disorder along the chain axis direction. For highly oriented samples, meridional layer lines consist of diffuse streaks of intensity resulting from partial axial translational disorder of the PPV chains.
89
1.01
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J
i
'
I
'
I
,
I
'
I
,
0.9
0.8
f
o.6
O. 5 L
2
I
4
,
~
6
8
10
Fig. 1. Herman's molecular orientation function, f, of PPV films determined by X-ray diffraction (open circles) and by infrared dichroism (filled circles)as a function of the
draw ratio, ~. The degree of molecular orientation in samples with varying draw ratios is shown in Fig. 1. The data represented by open circles were calculated from the measurement of the (110) X-ray diffraction intensity measured as a function of the azimuthal angle, ¢ [21]. Since stacking of multiple films could have added artificial disorder, measurements were made on single films less than 10/~m thick. The (110) reflection was chosen because of its large intensity. The orientation function, f, was estimated from
f=
3 cos2<~b}-- 1 2
(2)
90 where the average aximuthal angle, (¢>, was taken to be the full angular width at half the maximum (FWHM) of the intensity distribution of the reflection. The more exact integral m e t h o d of calculating the cosine average from lineshape analysis was n o t possible because of uncertainties in the lineshape of the intensity distribution at high ¢ angles, especially with films of lower draw ratio. The FWHM m e t h o d was less sensitive to lineshape ambiguity and provided sufficient precision for the analysis to show the relation of f to the draw ratio, ),. The values reported refer to the average orientation of the normal to the (110) plane relative to the perpendicular to the macroscopic orientation direction. Because of the rigidity of the chains, it has been assumed that the (110) orientation yields an acceptable relative measure of the c-axis orientation with respect to the draw direction. Infrared dichroism was also used to characterize the molecular orientation. The dichroic ratio, D, is defined by the ratio of the absorptions parallel, (All) and perpendicular (A±) to the reference direction. The orientation function in this case is calculated from
f=
[I -llt 0+211 \ D + 2 / \ D 0 - - 1/]
(3)
where Do = 2 cot 2 ~, and ~ is the angle of the absorption transition m o m e n t relative to the reference direction [21]. Absorbances associated with stretching vibrational modes are typically chosen for i.r. dichroism studies because ~ can then be unequivocally established. However, in PPV, the C--H stretch region centered at 3021 cm -1 was found to contain overlapping bands, which makes absorption measurements uncertain. In contrast, the trans-vinylene C--H o u t of plane bending absorption at 963 cm -1 is strong and is not overlapped. Although the angle a is n o t known absolutely, the trans-vinyl C--H bend transition m o m e n t is considered to be nearly perpendicular to the plane of the double bond, and would thereby be nearly perpendicular with respect to the PPV chain axis. Assuming the 963 cm -I transition m o m e n t is perpendicular to the chain direction, we will again define the axis perpendicular to the orientation direction as the frame of reference; thus c~ -~ 0 ° and f becomes simply D--1 f -~ ~ D+2
(4)
The infrared orientation function (fiR) is shown in Fig. 1 by filled circles and the X-ray orientation function (fx-ray) by the open circles. There is reasonably good agreement between the t w o measurements of orientation: both orientation functions increase rapidly with increasing draw ratio up to ~ 6, above which additional drawing has a less significant effect. The values of the orientation functions are relatively high even at low draw ratios. The efficiency of the draw process can be understood when one considers that as the flexible polyelectrolyte chain is being stretched it is
91 irreversibly converted to a rigid conjugated chain. It should also be noted that the total sample volume decreases by a b o u t 50% as the volatile elimination products diffuse o u t of the film, so that the effective draw ratio is actually higher than that measured. The observed difference between fir and fx-ray can be explained in two ways. The structure of the chain is n o t known in sufficient detail to exclude the possibility that the effective transition m o m e n t is n o t precisely perpendicular to the chain axis. The measured ratios of fir and fx-ray can conceivably be used to calculate the angle a. For X = 8.3, this calculation yields = 8 °, for X = 2.7, a = 19 °. The former is in fact in excellent agreement with recent interpretations of the transition m o m e n t from measurements by Bradley et al. [22, 23]. An alternative, non-exclusive possibility is that the X-ray intensities are sensitive more to the chains that are well ordered or in larger crystal domains, while the i.r.-derived orientation function is insensitive to local disorder. In this interpretation fx-r,y should be considered as the upper b o u n d of the true orientation. An average lateral crystallite size was obtained from a Scherrer analysis of the 20 breadth of the equatorial reflections [24]. The (210) reflection was chosen because it is the lowest order well-separated reflection. Crystallite size, t, as a function of the draw ratio, was calculated from t -
KA
{4) ~3~o=cos 0 assuming flco= = [t32 -t3std 2] in where K is the Scherrer constant, A is the Cu Ka X-ray wavelength, 0 is the Bragg angle, ~ is the full width at half maxim u m (FWHM) of the 20 scan; ~std is the FWHM for the standard (hexamethylene tetramine) used to approximate the instrumental broadening assuming a Gaussian lineshape. The precise value of K (~ 1) has been in dispute [25]. Since the object of this experiment was to examine the trend of crystallite size as a function of draw ratio, K was arbitrarily taken to be 1.0. Figure 2 shows the calculated lateral crystallite size as a function of draw ratio. It must be emphasized that the value of t is subject to the error associated with the Scherrer constant and with the fact that the treatment also does n o t account for lattice reflection broadening due to lattice distortions of the second kind (i.e., paracrystallinity). These uncertainties are, however, largely independent of X, thus it may be concluded that the crystallite size does n o t increase with increasing draw ratio to the same extent as does molecular orientation. The increase in the parallel conductivity with draw ratio discussed below must therefore be primarily the effect of the increase in molecular orientation. The axial conductivity (oll) of AsFs-doped oriented PPV as a function of draw ratio is shown in the upper curve of Fig. 3. The trend is similar to that seen with the orientation function. Thus log a increases rapidly with draw until X ~ 6 - 8 and then becomes rather insensitive to further stretching
92 130
II
I
I
~
120
110
!
I
I
I
I 8
I 10
I 12
I 14
f
10C
90
8(
I 2
I 4
I 6
Fig. 2. Transverse crystallite correlation length, t, of PPV films as a function of draw ratio k.
above a draw ratio of about 10. This correlation provides evidence that the axial conductivity is quite sensitive to the degree of orientation of the conjugated polymer chains. On the other hand, the conductivity transverse to the draw direction ( o i l shown in the lower curve of Fig. 3, decreases with increasing draw ratio, although the dependence on X is far less. A possible explanation for the behavior of o± is that the film becomes fibrillated during the drawing process. However, scanning electron microscopy did not reveal any significant fibrillization until the draw ratio was above 10, and with
93
l l J d l l t l l l l l
I
1o
id
~10 2
u U) u
b
b
J
IC
10
~
2
f
4
6
I
8
I
I
10
I
I
12
I
14
10
j-JJ"
02
0.4
0.6
0.8
1.0
Fig. 3. Conductivity of AsFs-doped PPV parallel (Oil) and perpendicular (o±) to the stretching direction as a function of draw ratio k. Fig. 4. The conductivity shown in Fig. 3 as a function of the orientation function, f.
careful stretching, unfibrillated films could be prepared even up to k = 12. It may be noted that with the present stretching apparatus, the PPV films drawn to k > 6 are t o o narrow to investigate conductivity in the transverse configuration with high precision. Figure 4 shows the axial conductivity as a function of the X-ray derived orientation function. Again oll increases strongly with f and continues to increase even after the maximum orientation value of 0.96 has been reached. This also indicates that the second-order orientation function cannot completely determine the effect of orientation upon conductivity. This behavior is also c o m m o n l y observed in studies of the relation of modulus and thermal conductivity to stretching in many polymers. Higher-order determinations of orientation are required to characterize this effect more precisely. A further consideration is that the orientation was measured on u n d o p e d samples and it is n o t inconceivable that the dopant molecules have a significant effect on the local structure of the macromolecule. Future studies will characterize the microstructures of PPV samples after doping. Conclusion X-ray and infrared dichroism investigationshave demonstrated that the molecular orientation of stretched P P V films increases strongly as a function
94
of the macroscopic draw ratio. This effect is enhanced because there occurs an irreversible conversion to extremely rigid conjugated PPV sequences, which retard chain relaxation during the stretching process in addition to the normal deformation orientation process. There is a modest increase in the lateral crystallite size with increased stretching. The increase in conductivity parallel to the orientation direction provides evidence pointing to the dominant role of intrachain diffusion of the charge carriers along the PPV backbone. Certainly if interchain conduction (via hopping or some other mechanism) were the sole path of charge transport, there should be an increase in conductivity perpendicular to the orientation direction due to an increase in the lateral order, and a very small or negative effect parallel to the chain direction. It is obvious that a single chain does n o t extend throughout the entire length of the sample, so interchain transport is an important part of the mechanism; this study seems to indicate that interchain transport is the rate-limiting step. Independent measurements of the anisotropy in charge carrier mobility (e.g., by Hall effect studies) are needed to confirm this interpretation.
Acknowledgements This work was supported by A F O S R Grant 85-0044. One of us (D.R.G.) thanks the P.I.A. for additional support. We would also like to thank Professor S. L. Hsu for helpful discussions concerning the infrared dichroism studies, and Dr. Thierry Granier for discussions concerning the crystal structure of PPV.
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95 14 D. R. Gagnon, J. D. Capistran, F. E. Karasz and R. W. Lenz, Am. Chem. Soc. Polym. Prepr., 25 (2) (1984) 284. 15 K. Harper and P. G. James, Mol. Cryst. Liq. Cryst., 11 7 (1985) 55. 16 R. W. Lenz, F. E. Karasz, J. D. Capistran, D. R. Gagnon and S. Antoun, submitted to
Macromolecules. 17 G. Leising, H. Kahlert and O. Leitner, in B. Kramer et al. (eds.), Springer Series in Solid S tate Physics, Springer -Verlag, Berlin, 1985, Vol. 61. 18 D. R. Gagnon, J. D. Capistran, F. E. Karasz and R. W. Lenz,Polym. Bull., 12 (1984) 93. 19 T. Granier, E. L. Thomas, D. R. Gagnon, F. E. Karasz and R. W. Lenz, J. PolymerSci. (Phys.), in press. 20 J. D. Capistran, D. R. Gagnon, S. A. George, R. W. Lenz and F. E. Karasz, Am. Chem. Soc. Polym. Prepr., 25 (2) (1984) 282. 21 B. E. Read, in I. M. Ward (ed.), Structure and Properties of Oriented Polymers, Wiley, New York, 1975, Ch. 4. 22 D. D. C. Bradley, R. H. Friend, T. Hartman, E. A. Marseglia, M. M. Sokolowski and P. D. Townsend, ICMS '86, Kyoto, June, 1986, Synth. Met., 17 (1987) 473. 23 D. D. Bradley, R. H. Friend, H. Lindenberger and S. Roth, Polymer, in press. 24 M. Kakudo and N. Kasai, X-ray Diffractions by Polymers, Elsevier, New York, 1972. 25 B. E. Warren,Phys. Rev., 59 (1941)693.