Structure, morphology and electronic properties of Trans-(CH)x

Synthetic Metals, 6 (1983) 243 - 263

243

STRUCTURE, M O R P H O L O G Y AND ELECTRONIC PROPERTIES OF

TRANS-( CH)~ C. R. FINCHER, Jr.*, D. MOSES*, A. J. HEEGER* and A. G. MACDIARMID**

Laboratory for Research on the Structure of Matter, University of Pennsylvania, Philadelphia, PA 19104 (U.S.A.) (Received January 18, 1983)

Summary The structure and morphology of trans-(CH)x films prepared by the Shirakawa technique are established. It is shown that the trans-(CH)x chains lie parallel to the fibril axes. The results rule o u t the proposed lamellar morphology and chain folding as principal structural features. Electrical and optical properties indicate that the principal 7r-electron transport is along the trans-(CH)x chains. The effects of dopant uniformity on the s e m i c o n d u c t o r metal transition are reviewed. It is shown that the data for d o p e d trans-(CH)~ are not consistent with the percolation model.

1. Introduction Organic polymers which have the electronic properties of semiconductors or metals when doped with donor or acceptor species have recently emerged as an important new class of materials [1]. The p r o t o t y p e example of these conducting polymers is polyacetylene, (CH)x. The simplest conjugated polymer, polyacetylene, consists of chains of CH units as shown in Fig. 1. For each CH unit, three of the four carbon valence electrons are in sp2-hybridized orbitals; two of the o-type bonds construct the polymer backbone while the third forms a bond with the hydrogen side group. The fourth electron goes into the 7r-electron system (Pz orbital) which is the origin of the electronic activity and the semiconducting or metallic properties. The 120 ° bond angle between the three o-bonds can be satisfied b y two possible arrangements of the carbon atoms, trans-(CH)x and cis-(CH)x as shown in Fig. 1. Both cis and trans forms can be prepared as silvery flexible films, which can be made free-standing or can be polymerized directly as *Permanent address: Institute for Polymers and Organic Solids, University of California, Santa Barbara, CA 93106, U.S.A. **Permanent address: Department of Chemistry, University of Pennsylvania, Philadelphia, PA 19104, U.S.A. 0379-6779]83]$3.00

© Elsevier Sequoia/Printed in The Netherlands

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(CH)x. thin films on a variety of substrates. The trans isomer is the thermodynamically stable form. Although it is possible to obtain pure cis-(CH)x through careful synthesis at low temperatures (--78 °C), isomerization from c/s to trans occurs in a few minutes upon raising the temperature above 150 °C. Interest in this semiconducting polymer was stimulated by the discovery and successful demonstration of doping, with associated control of the electrical properties, over a remarkably wide range [2 - 4]. Overall, the conductivity of films of (CH)x can be varied over more than twelve orders of magnitude, from that of an insulator (<10 -9 ~ - I cm I) through semiconductor to a metal (>103 ~2-i cm-1). The doping is carried o u t after synthesis; (CH)~ films can be chemically or electrochemically [5] d o p e d at room temperature with a variety of donors or acceptors to form n-type or p-type semiconductors [6]. Doping to higher levels (above ~1%) results in a semic o n d u c t o r - m e t a l transition giving a whole new class of "metallic polymers" with a wide range of electronegativity. The existing experimental data already show that these materials have potential for use in a number of areas of future technology. Experimental studies at the University of Pennsylvania have demonstrated that d o p e d polyacetylene might be useful in such diverse applications as: (1) the replacement of increasingly scarce, conventional conductors b y synthetic metals; (2) the development of light-weight, high-energy-density batteries; (3) low-cost solar photovoltaic materials. Values for the d.c. electrical conductivity greater than 3 X 103 ~ - i cm-1 have already been achieved with only partially crystalline samples, and analysis of the transport [7] and optical data [8] implies that a further increase of at least one order of magnitude should be possible. Controlled electrochemical doping and "undoping" have been demonstrated, and p r o t o t y p e high-energydensity rechargeable batteries have been constructed using (CH)~ as both the active cathode and anode [9, 10]. Photovoltaic phenomena have been observed in heterojunctions [ 11 ], Schottky-barrier junctions [ 12] and photoelectrochemical junctions [13]. Thus, although these and other potential applications will require considerable future work before ultimate technological value can be determined, the properties appear promising. Electron microscopy studies [14] show that the as-formed (CH)~ films consist of randomly oriented fibrils (typical fibril diameter ~ 2 0 0 •).

245 The bulk density is 0.4 g/cm 3 compared with ~ 1 . 2 g/cm 3 as obtained by flotation techniques. Thus, the polymer fibrils fill only a b o u t one third of the total volume, and the effective surface area is quite high ( ~ 6 0 m2/g). Both higher [15] and lower [16, 17] density forms have been synthesized and characterized; however, the broad base of studies carried o u t in our laboratories has focussed primarily on the standard films synthesized in the manner initially developed by Shirakawa and colleagues [ 18]. The (CH)x films can be stretch-oriented in excess of three times their original length with concomitant partial alignment of the fibrils [ 1 9 - 2 1 ] . The oriented films exhibit highly anisotropic electrical [19] and optical properties [20] suggestive of an anisotropic electronic structure. Both the electrical and optical studies indicate that the preferred axis for electrical transport is along the orientation direction, e.g., in the metallic regime, the measured conductivity is approximately twenty times greater along the direction of stretch-orientation than the value measured perpendicular to this direction. The description of the complex morphology of polyacetylene and the relation of this morphology to the observed electronic properties have been a subject of interest and controversy [ 2 2 - 26]. Although the fibrillar morphology has been widely reported by many research groups, Wegner and colleagues [22 - 24] have asserted that the fundamental morphology consists of lamellae stacked (in a manner analogous to poker chips) to form the apparent fibrils. They argue, in addition, that within the proposed lamellae, the polymer structure is chain-folded with the principal chain axis perpendicular to the fibril axis. In their view, the principal transport is perpendicular to the (CH)x chains. The analysis of the morphology and its relation to electronic properties is additionally complicated by potential variations due to the use of different catalyst systems, different aging times [27] for the catalyst prior to use, different experimental conditions during synthesis, etc. Thus, in order to reach firm conclusions, morphology, structure, and electronic studies must be carried o u t on samples prepared under identical conditions. Once these questions are settled for a given set of conditions, the conclusions can be tested for other materials prepared with other catalysts and/or techniques, etc. In this paper we establish the structure and morphology of stretchoriented films prepared in our laboratory using the basic Shirakawa synthetic m e t h o d (see below for details). X-ray scattering measurements [28] on trans(CH)x prove that the polymer chains are parallel to the oriented fibrils seen by electron microscopy. These results rule o u t the proposed lamellae and chain folding as principal structural features. We then show that in such oriented films both electrical and optical properties indicate that the principal 7r-electron transport is along (rather than perpendicular to) the trans(CH)x chains. All measurements on as-grown films (i.e., n o t oriented) are consistent with the same conclusions. We then consider the effects of dopant uniformity on the metal-insulator transition and show that the data for doped trans-(CH), are not consistent with the percolation model.

246 2. Synthesis and orientation of t r a n s - ( C H ) x Cis-rich (CH)x (~85% cis isomer; ~0.1 mm thick) was synthesized by a m e t h o d based on that described by Shirakawa [18]. All solvents were thoroughly dried by standard procedures and rigid attention was given to excluding even traces of air and/or water during all phases of catalyst preparation, polymerization, washing, and drying of the film. All operations were performed either on the high vacuum line or in an argon atmosphere in an inert atmosphere box or by using Schlenk tube techniques. Traces of oxygen were removed from the argon by the use of B.A.S.F. catalyst (R3-11); water was removed by passage of the gas through KOH pellets, 4 A molecular sieve, silica gel, and finally through P4010 dispersed on glass wool. Acetylene gas was purified by bubbling through wash bottles of concentrated sulfuric acid and then through a column of KOH pellets and finally through a column of P4010 dispersed on glass wool. Infrared spectra of the C2H2 showed that all traces of acetone and other impurities had been removed. The (n-C4HgO)4Ti (Alfa-Ventron) was distilled under reduced pressure in an argon atmosphere and was stored under vacuum. The 95% (C2Hs)3A1 (Alfa-Ventron) was used as received. The preparation of the catalyst involved both Schlenk tube and high vacuum system techniques. Approximately 20 ml of dry toluene was distilled into the b o t t o m of the reactor in which the polymerization was later to be carried out. Then 1.70 ml of (n-C4HgO)4Ti was added by means of a hypodermic syringe. After shaking, 2.70 ml of 95% (C2Hs)3A1 (Alfa-Ventron) was added using a hypodermic syringe and the mixture was again shaken. It was then allowed to age at room temperature for about 60 min. The morphology of the (CH)x obtained is critically dependent on the aging time [ 27]. The reactor was next cooled to --78 °C, evacuated, and shaken to remove dissolved ethylene and possibly other volatile substances formed during aging. When shaking did not produce further bubbling of the catalyst, the reactor vessel was again shaken, this time with a back and-forth-motion. This was made possible by a flexible stainless steel bellows connection joining it to the vacuum line. This motion resulted in an even coating of the relatively viscous catalyst solution on the walls of the reactor at --78 °C. Gaseous C2H2 (200- 700 Torr) was then i m m e d i a t e l y (while there was still a relatively thick film of the catalyst on the walls) admitted to the reactor. After 5 - 15 min, the rate of uptake of C2H 2 had decreased greatly and the C2H2 pressure had decreased by ca. 150 Torr. The pressures were measured in a total volume of ca. 2.75 1 (0.5 1 for the reactor, ~ 2 liters for the C2H 2 storage bulb and ~0.25 1 for the manifold volume). After the catalyst solution had been removed by a hypodermic needle under argon, 20 ml of pentane was distilled into the reactor which was then shaken gently to wash catalyst from the film. This solution was then poured into an attached side-arm flask, and then the pentane was distilled back into the reactor at --78 °C leaving catalyst residue behind in the side arm. This process was repeated during several hours until the pentane washings were completely colorless.

247 On removal from the reactor, the side of the dry film which had been touching the reactor glass walls was a bright, silvery-coppery color; the side facing the inside of the reactor was dull grey in color. Gentle rubbing of the dull side made it silvery. The film which formed on the surface of the catalyst solution was always discarded since its elemental analysis showed it contained significant amounts of ash. A representative elemental analysis (Galbraith Laboratories, Inc.) of cis-(CH)x film from the walls of the reactor is: C: found, 92.13%; calcd., 92.26%; H: found, 7.75%; calcd., 7.74%; total (C plus H) 99.88%. Isomerization to trans-(CH)x was accomplished by heating under vacuum to a temperature above 150 °C. Cis-(CH)~ was stretched either in an argon atmosphere or in the presence of laboratory air (exposure times < 5 min) [21] by mounting the film (ca. 3 mm wide) in a jig consisting of two pairs of clamps, each m o u n t e d in a sliding support. Each end of the specimen was placed between a pair of clamps (ca. 10 mm apart) and the film was then stretched by slowly moving the clamps apart. The degree of stretching is given by the ratio l/lo where l0 is the original length and l is the length of the film after stretching. This cis film could be further elongated during isomerization b y mounting it in a tube and attaching a weight, which was only 80% of that needed to cause fracture as ascertained in a preliminary experiment, to its lower end. The tube was evacuated and heated at a rate of ca. 5 °C/min until a temperature of 200 °C was reached. An additional 20% elongation (approx.) could be obtained in this manner during isomerization. The stretched trans-(CH)x used in the X-ray scattering studies (Section 3) were prepared by isomerizing previously stretched cis-(CH)x films at approximately 150 °C.

3. X-ray scattering studies of trans-(CH}x : structure and morphology [28] The X-ray scattering experiments were performed upon a triple axis X-ray spectrometer using, primarily, Cu 1.54 A radiation (or Mo 0.707 A where needed for larger values of Q), and a pyrolytic graphite m o n o c h r o m a t o r and analyzer. The sample was prepared by the method of Shirakawa [18], stretch-oriented [21], isomerized, and cut into about ten strips of approximately 12 × 4 mm 2 (total thickness ~ 1 mm). The ten strips of (CH)x were then carefully aligned, clamped together, and placed in an evacuated Be can. In the scattering geometry, the direction of orientation ( a n d , h e n c e , the average fibril axis) lies in the scattering plane; we denote this the L direction. Because the fibrils are not oriented w i t h r e s p e c t to rotation about ~', each reflection results in a ring around the L-axis with a radius equal to the transverse Q and intercepts the scattering planes at K values equal to the radius of the ring (L and /~ form a rectangular coordinate system which spans the scattering plane). As a result of the partial orientation, the X-ray pattern consists of arcs (rather than the uniform rings of a powder pattern). The spread of the fibril orientation can be inferred from the intensity distribution along the arcs; we find an approximately Gaussian shape with full width at half maximum of a b o u t 35 °. .

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248 Scanning electron microscope pictures of the films, as grown and after stretch alignment, are shown in Fig. 2. These films were prepared and stretch oriented using procedures identical with those used for the X-ray samples. The characteristic branched and twisted fibrils of the unstretched polymer, discussed earlier by Shirakawa et al. [14], are clearly visible in Fig. 2(a). Elongation results in partial alignment, as shown in Fig. 2(b). The observed degree of fibril orientation is consistent with that inferred from the X-ray data.

(a) (b) Fig. 2. Electron mierograph showing fibrillar morphology of polyaeetylene. (a) as-grown; (b) stretch-oriented. We assume that the backbone structure remains essentially unchanged when the chains are arranged in a three-dimensional (3d) lattice; i.e., the trans-(CH)x chain structure is that shown in Fig. 1 with a repeat dimension of approximately 2.46 A. As a result, one expects the principal Bragg reflections to appear at Q = n (2~r/2.46), with odd-n forbidden (or at least very weak) for the trans-(CH)x chain with (nearly) uniform bond lengths. Figure 3 shows the peak at 5.1 A -1 for ~ along the orientation direction, corresponding to the carbon-carbon spacing along the backbone {Q = 2~/1.23 = 5.1 A-l). This (002) reflection is unobservable with the sample rotated by 90 ° with ~ transverse to ~. Thus, the polymer chains are parallel to the orientation direction and parallel to the fibril axis. For the intense interchain reflection at Q = 1.72 )~-1, a combination of (110) and {020), we find (in agreement with earlier results) [29] the opposite behavior; Fig. 4 shows strong scattering for ~ perpendicular to ~, while for ~ parallel to ~ the intensity is more than ten times weaker. The demonstration by X-ray scattering {Figs. 3 and 4) that the trans(CH)x chains are parallel to the orientation direction, and the demonstration by electron microscopy (Fig. 2) that the polymer fibrils are parallel to the orientation direction, provide irrefutable evidence that the polymer chains lie parallel to the fibril axis. We conclude that the complex lamellar morphology {involving chain folding in such a way that the backbone of the polymer is primarily perpendicular to the axis of the fibril) proposed by Wegner [22 -

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24] is definitely impossible for the oriented samples prepared in our laboratory. Note that these X-ray data do not come from special selected areas. In the scattering configuration, the beam passes through a macroscopic scattering volume distributed over approximately ten oriented films of trans(CH)x. The small number of reflections precluded a full refinement of the structure of the polymer by standard methods. However, because the basic structure of a single polymer chain of trans-(CH)x is independently known, the limited scattering results are sufficient to allow determination of both the interchain packing and the intrachain distortions [28]. The arrangement of the chains when projected onto a_>plane normal to the chain direction can be obtained from a scan normal to L (Fig. 5). After indexing these as (h k 0) type reflections, it was concluded that the projection has pgg symmetry. Very good agreement between calculated and measured intensities was found for a setting angle Q = 55° with respect to (010), close to the value found by Baughman from packing calculations [30]. Detailed analysis of the X-ray data (positions and intensities) has been carried out in order to determine the crystal structure of trans-(CH)~. We conclude that the structure (see Fig. 6) is P21/n which is consistent with not only all of the X-ray data, but also agrees with packing calculations [ 30] and the data of other experiments. The observed positions of the Bragg reflections limit fi (the monoclinic angle) to the range 91 ° < fi < 93 °. Within this range, best fits to the scattering profiles yield values of ~ -~91.5 and u0 -~0.03 A (the amplitude of the longitudinal distortion which leads to bond alternation). Given the space group, the lattice constants (a = 4.24 /~, b = 7.32 /~ and c = 2.46 A), the setting angle (55°), ~, and the amplitude of the dimerization distortion, the essential

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aspects of the polymer structure are determined• The lattice constants are very close to those reported by Shimamura et al. [31]. Although they initially suggested an orthorhombic (P,~,,) space group (which is incompatible with our data and which would not allow for a dimerized chain structure), the X-ray and electron diffraction data are consistent. The data of Figs. 3, 4 and 5 indicate a high degree of crystallinity when compared with other published data. Although an absolute determination of the a m o u n t of crystalline vs. amorphous material is difficult, it is straightforward to make a relative comparison with previous results. By comparing these data with the results of Akaishi et al. [29] or Baughman et al. [30], one sees that the background scattering outside the crystalline reflections is two to three times lower, implying that the sample used in this work contained a larger fraction of crystalline material. Since Akaishi e t al. [29] infer ~80% crystallinity from analysis of their data, the films of trans-(CH)x used in our experiments would be in excess of 90% crystalline. In spite of the higher crystallinity, there is still a significant a m o u n t of disorder, as evidenced by the small number of observable reflections and their relatively large widths• The width of the interchain reflections, e.g., at Q = 1.72 A -x, is (FWHM)F = 0.10 A - I , and is independent of the magnitude of the m o m e n t u m transfer. This is consistent with a finite crystallite size in the perpendicular direction, perhaps resulting from the ~ 2 0 0 A fibril dimension (half of the chains are within ~ 3 0 A of the fibril boundary, and all are within 100 A). On the other hand, the width along the chain axis is not ~ independent; the (004) at Q = 10.2 A -1 is about four times wider than the relatively narrow {002), suggesting a IQ]2 dependence. Therefore, while one might estimate a chain length of about 100 A (~ -~ 2~r/Ffor a finite array of point scatterers), this is only a lower limit. The disorder implied by the large width of the equatorial reflections is consistent with a large effective " D e b y e - W a l l e r " factor found (by us and by

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others [ 30 ]) in the analysis of intensities. Although the intensities decreased as exp(--BQ2}, B = 0.07 A 2, B did not change significantly between 300 K and 20 K. We conclude that considerable static disorder exists in the 3d lattice of the chain projection, as might be introduced by occasional crosslinks of the poly.mer and/or by the finite fibril dimensions. The behavior of the widths for ~ along the chain must be considered in a completely different manner. A familiar example which gives a IQI2 width dependence is the

252 l d lattice [31, 32]. This is not to imply that the (CH)~ polymer is a truly l d lattice, but rather that a combination of effects such as the bends and curves of a fibril, etc., could give rise to the observed width, even with a very large actual chain length.

4. Anisotropic electrical conductivity [19] of partially oriented trans-(CH}~ The room temperature electrical conductivity of as-grown polycrystalline films heavily doped to [CH(AsFs)0.14]x was reported [33] as 560 (~t cm) -1 , comparable with the best single crystals of the metallic organic charge transfer salts such as T T F - T C N Q [34]. This high conductivity is particularly interesting in view of the fibrillar morphology discussed above. As a result of the large interchain spacings (see Fig. 6), we expect the interchain electronic transfer integrals to be small, ~0.1 eV, i.e., less than, or comparable with, the intermolecular transfer integrals along the b-axis in T T F - T C N Q where the intermolecular spacing is 3.6 A. On the other hand, molecular spectroscopic studies of short chain polymers lead to the conclusion that the intrachain transfer integrals for carbon atoms separated by ~ 1 . 4 A are of the order 2.5 - 3 eV. Thus we anticipate a highly anisotropic band structure with r-electrons delocalized along the (CH)x chains and relatively weak interchain coupling. As a result the electrical transport in (CH)~ would be expected to be correspondingly anisotropic with highest conductivity along the chain direction. The temperature dependence of the parallel and perpendicular conductivities and the anisotropy for an oriented film (l/lo = 2.91) d o p e d into the metallic regime with AsF s are shown in Fig. 7 [19]. Again, the films used in the transport studies were prepared and stretch-oriented using procedures identical with those used for the X-ray and electron microscopy studies. The solid points result from 4-probe measurements on two separate (I] and J_) films; the " x " points result from use of the Montgomery method. The three samples were taken from the same initial film and doped simultaneously to a final composition [CH(AsFs)0.10]x. The results from the two independent sets of measurements are consistent; the conductivity parallel to the orientation direction is greater than the perpendicular value by more than an order of magnitude. Moreover, the room temperature parallel conductivity is in excess of 2000 (~2 cm) -1 ; the average of the two measurements yields 2150 ( ~ cm) -1. On cooling, oil decreased slowly; however, the conductivity remains high even at the lowest temperatures consistent with metallic behavior. A more detailed examination of the data shows that oal remains approximately constant, increasing slightly (~0.5%) down to 260 K, whereas al decreases monotonically. Recent experiments have extended the data into the milli-Kelvin range [35]. The doped (CH)~ samples remain conducting even at these extremely low temperatures. At temperatures below 0.03 K, a log T increase in resistance is observed, suggestive of the onset of quantum localization effects.

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The transport results are therefore fully consistent with the conclusions reached in Section 3 regarding structure and morphology. The principal direction for electronic transport is parallel to the (CH)x chains. This conclusion is valid for both undoped and doped polyacetylene, since even the undoped oriented polymer exhibits an anisotropy of more than 10:1 with higher values along the orientation direction. This does not imply that interchain transport is unimportant. As the structural studies clearly demonstrate [28], the trans-(CH)x chains are imperfect. Thus, although the individual chain structural coherence length may be significantly greater than a typical mean free path (at least at room temperature), eventually interchain transfer will limit the conductivity as shown by Park et al. [7]. The general conclusion is that the transfer is limited by a combination of intra-chain scattering, interchain transfer (e.g., at a defect or cross-link), and interfibril contact. Nevertheless, the observed electrical anisotropy for oriented films demonstrates that the principal direction for electron transport is along the (CH)x chains, consistent with the delocalized ~-electron functions.

254

5. Anisotropic optical properties [8, 20] of partially oriented trans-(CH)x When examined visually, the bright surface of an oriented film has a reflection similar to aluminum foil, b u t somewhat darker. Through a polarizer, the reflection polarized parallel to the fibril, and polymer chain orientation is bright and uncolored, whereas the reflection polarized perpendicular is weak and pastel orange. The polarized reflectance data for undoped trans-(CH)x, obtained from oriented films, are shown in Fig. 8. R Hshows a broad maximum throughout the visible and decreases in the infrared, consistent with expectations for a semiconductor. The perpendicular reflectance is small over the entire range; 4% at low frequencies with a very weak maximum centered at 1.65 eV. Qualitatively, the data imply the optical properties of a semiconductor for electronic m o t i o n parallel to the trans-(CH)x chains, and those of an insulator for electronic motion perpendicular to the chain. Quantitative Kramers-Kronig analysis [8] of the Rii data demonstrates that all the electron oscillator strength is involved in the interband transition polarized parallel to the trans-(CH)x chains. The optical anisotropy is large and appears to be intrinsic; Rll/R l goes through a minimum of 4.7 at 1.65 eV, increases to about 10 in the region of the interband transition, and then decreases somewhat at higher energies. Similar anisotropy is observed after doping [8, 36]. The polarized reflectance of metallic [CH(AsFs)0.12]~ is shown in Fig. 9. Rll is characteristic of a short mean-free-path metal; the reflectance is high and increases toward unity in the infrared. For perpendicular polarization, however, the reflectance remains low with no signature of metallic behavior. Only at the longest wavelengths is an increase of R± observed, consistent with the smaller d.c. conductivity in this direction. Again, Kramers-Kronig analysis [8] of The Rll data demonstrates that all the ~-electron oscillator strength is I01~

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Fig. 8. Polarized r e f l e c t a n c e data for u n d o p e d t r a n s - ( C H ) x , o b t a i n e d f r o m o r i e n t e d films. Fig. 9. Polarized r e f l e c t a n c e data for heavily d o p e d metallic [CH(AsFs)0.12]x , o b t a i n e d f r o m o r i e n t e d films.

255 involved in the metallic transport along the trans-(CH)x chains. The resulting data for 0(09) leads to an estimate of the intrinsic d.c. metallic conductivity characteristic of an individual fibril: 2 × 104 ~2 cm -1 at room temperature. Note that this estimate is obtained from analysis of the frequency dependent (or wavelength dependent) conductivity at wavelengths in the infrared. Therefore, the results represent an average over all microscopic structural features (<103 A); the data are consistent with inter-fibril contacts as the major source of series limited transport. The optical anisotropy throughout the visible and infrared observed through reflectance measurements from oriented films again demonstrates that the principal direction for transport in both semiconducting and metallic (doped) trans-(CH)x is along the (CH)x chains.

6. Morphology, structure, and transport: non-oriented trans-(CH)x The morphology, structure, and transport in oriented films was established in Sections 3, 4 and 5. Are the same conclusions valid for the as-grown films? The answer is yes, unless the stretch-orienting procedure introduces structural and/or morphological changes. The electron micrograph (Fig. 2) shows no obvious change after orientation. The appearance is simply that of random vs. partially aligned fibrillar structures. The typical fibril dimensions do not change on orientation. The electron diffraction studies of Shimamura et al. [31] provide confirming evidence that the structure is the same on oriented and nonoriented polymer. They observe no additional diffraction peaks in oriented regions as compared with non-oriented regions. Moreover, their diffraction data from individual fibrils indicate independently that the trans-(CH)x chains are parallel to the fibril axes. X-ray studies of oriented and non-oriented films show identical Bragg reflections. The only difference is that the powder pattern rings of the asgrown film become arcs in the oriented polymer as described in Section 3. Our conclusion is that for films prepared using the Shirakawa m e t h o d (see Section 2 for details) there appears to be no major change in structure or morphology resulting from the orientation process. Thus, the trans-(CH)x chains lie parallel to the fibrils in the as-grown films as well. This conclusion is consistent with optical and transport studies of non-oriented trans-(CH)x. Although the parallel conductivity increases with alignment, the dependence is smooth and scales smoothly with the degree of alignment. Absorption measurements characteristic of an unoriented film are shown in Fig. 10. Also shown are values for ~(c0) calculated from the Kramers-Kronig analysis of the polarized reflectance data obtained from an oriented film [8, 37]. As presented, the calculated curve for ~(w) has been divided by 2 for direct comparison of the two curves. The shape and position are essentially identical. The factor of two discrepancy arises in part from uncertainty in the absolute value of a(c~) for the as-grown films, and in part

256

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Fig. 10. A b s o r p t i o n c o e f f i c i e n t of trans-(CH)x. T h e d a s h e d curve results f r o m direct m e a s u r e m e n t ; t h e solid curve was calculated f r o m t h e K r a m e r s - K r o n i g analysis of t h e r e f l e c t a n c e data.

from uncertainty in correcting for the orientation (i.e., if the fibrils in the asgrown film are truly random, one would expect a factor of three difference). The excellent agreement between the frequency dependences of the two curves implies that the optical properties of as-grown and oriented films are identical; the only difference being in the degree of orientation.

7. The effect of non-uniform doping on electrical transport in trans-(CH)x: studies of the s e m i c o n d u c t o r - m e t a l transition [38] Due to the strong dependence of the resistivity on the dopant concentration, non-uniformity across the thickness of a sample leads to an apparent anisotropy in the resistivity {measured parallel to, and perpendicular to, the plane of the film). Thus the magnitude of the apparent anisotropy (for nonoriented samples) offers a convenient means for quantifying the degree of dopant non-uniformity in trans-(CH)x. This technique can be used to evaluate the uniformity achieved in [CH(AsFs] x by conventional vapor phase doping as compared with the slow doping technique developed earlier for susceptibility studies. From the transport data, we find that the better the uniformity in the distribution of dopant, the more abrupt is the semic o n d u c t o r - m e t a l transition at a critical concentration, Yc, of about 0.001. Doped samples, prepared with the slow doping technique developed earlier [39], yielded o11/~1 -~ 2 at all doping levels. Thus, although the absolute resistivity increased by m a n y orders of magnitude on doping, the anisotropy remained essentially constant and equal to that found in undoped trans-(CH)~. The constant anisotropy implies uniform doping throughout the film thickness. If the center were undoped (or if there were a significant concentration gradient), the resistivity of the center portion would be high, while that of the outer surfaces could be low. Since aOy) is a strong function, such a dopant non-uniformity would lead to a large apparent anisotropy. Such an anisotropy is indeed obtained on samples doped in the more conventional manner [38]. For such samples we typically find oil/ o , ~ 25 - 35.

257

The concentration dependence of the conductivity, a(y), for uniformly doped [CH(AsFs)y]x is presented in Fig. 11. Compared with earlier data [7] the results indicate a somewhat more sharply defined SM transition as a function of y and show higher electrical conductivities in the transitional region. For example, at y = 0.002, o = 0.5 ~ - i cm 1, more than three orders of magnitude larger than obtained earlier using more rapid doping. Similarly, at y ~ 0.02, a = 30 ~ - 1 cm-1, again more than an order of magnitude greater than samples prepared with known non-uniformity (i. e., apparent anisotropy of ~ 2 5 ) . At higher (and lower) concentration, the conductivities were typical of doped polyacetylene. The concentration dependence of the thermopower, S(y), for uniformly doped [CH(AsFs)y]x is shown in Fig. 12. Again, the transition to "metallic" behavior is more abrupt as a function of y and occurs at a lower concentration than reported earlier for samples prepared with standard doping procedures. For example, at y = 0.0004, S = +840 p V / K and at y = 0.0025, S = 20 pV/K, whereas in earlier studies such small values were not achieved until y -~ 0.01. For the y = 0.0025 sample (i.e., just above the transition), the thermopower is small and quasi-linear with temperature, consistent with the sharp transition. This transition is also observed in the temperature dependence of conductivity. In the dilute regime, o ( T ) ~ T n with n ~- 13 [40]. This strong dependence persists until y -~ 0.001. At higher concentrations the depen-

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258 dence of o upon T becomes extremely weak. For example, if one evaluates the slope, kBd(lrlp)/d ( T -1) at room temperature, the value drops from 0.31 eV below y = 0.001 to 0.02 eV at y -~ 0.002. Thus, major changes in the magnitude and temperature dependence of both the t h e r m o p o w e r and conductivity occur at y -~ 0.001. The anisotropy measurements on non-oriented films lead to the conclusion that the slow doping procedure [38, 39] does result in uniform doping across the thickness of the trans-(CH)x film. Such measurements are relatively insensitive to intra-fibril concentration gradients. However, the slow doping procedure with cyclic cryopumping allows considerable time for the establishment of a uniform equilibrium concentration throughout the 200 A dia. fibrils. The intra-fibril uniformity is confirmed by the essentially complete disappearance of the (neutral soliton) Curie-law susceptibility in such samples at intermediate doping levels [39]. We point out here that by uniform doping, we mean that the density averaged over any small macroscopic volume does not depend on the region of the sample over which the average is taken. Thus, even for uniform doping, statistical fluctuations will always give rise to local variations on a small enough scale. The observation that the SM transition is much sharper in samples which are macroscopically uniform (all/o ± -- 2) as compared with samples with known non-uniformity (as evidenced b y an apparent anisotropy, o11/ a, ~ 25) is particularly important. Smearing of the transition by concentration gradients is to be expected if the behavior (for example, S vs. y, or × vs. y, etc.) is abrupt. This was demonstrated explicitly for the Pauli susceptibility by Epstein et al. [41]. The transport data obtained from uniformly doped samples (Fig. 11) together with the magnetic susceptibility results, define the SMtransition. For y ~ 0.002, the t h e r m o p o w e r and conductivity data (vs. temperature and pressure) indicate hopping transport. Using the soliton as a basis, a quantitative understanding of the hopping transport in trans-(CH)x at dilute doping levels has been demonstrated. Kivelson's theory [42] of transport by inter-soliton hopping accounts for the magnitude, temperature dependence [40], pressure dependence [40], and frequency dependence [43] of the electrical conductivity, and for the magnitude and sign of the temperature independent thermopower [40]. Alternatively, attempts to explain the results in terms of variable range hopping predicted conductivity values off by fifteeri orders of magnitude and predicted pressure and frequency dependences which are not consistent with the experimental results [40, 43]. In the heavily doped regime, y > 0.07, the Pauli susceptibility [39], linear term in the heat capacity [44], linear temperature dependence of the thermopower [7], and infrared absorption data [3], all imply metallic behavior with a density of states at the Fermi energy (N(EF) - 0 . 1 states/ e V - C atom) consistent with a broad energy band. Optical absorption [8] and electron energy loss [45, 46] experiments suggest that the energy gap of pure (CH)x has closed (or at least reduced to a value less than a few tenths of

259 an eV). Thus, in the high concentration limit, relatively simple and traditional metallic behavior is observed. The question of whether the bond alternation goes to zero in this regime or whether the reduced gap is filled in by a combination of disorder and interchain coupling remains to be settled [47]. The experimental results are inconsistent with the percolation model proposed by Tomkiewicz et al. [48, 49] and by Wegner [22]. The essential feature of such a percolation transition is the presence of disconnected metallic regions at concentrations well below the critical value (Yc). One would therefore expect ×p to increase linearly at concentrations well below Yc at which point the regions would connect leading to metal-like transport. Tomkiewicz et al. [48, 49] observed a Pauli susceptibility which increased linearly with dopant concentration (measurements carried o u t primarily on cis-(CH)x samples). However, subsequent studies of the magnetic properties of carefully doped trans-(CH)x (by Penn group [39], in a joint Penn-IBM study [50], and b y the Stuttgart group [51]) demonstrated the nonmagnetic character of the doped polymer, [CH(AsFs)]~, for concentrations below about y = 0.07. In contrast to the predictions of the percolation model, the susceptibility was found to increase abruptly at y -~ 0.07 to values characteristic of the Pauli spin susceptibility of a broad band metal, × p ~ 3 × 10 -3 emu/mole. Frequency dependent conductivity studies have been useful in checking for the proposed "metallic islands". Grant and Krounbi [52] found the room temperature conductivity to be frequency independent (10 H z 10 MHz) for all levels of doping with AsFs. Similarly, Epstein et al. [53] investigated the frequency dependence of o from d.c. through 500 MHz at intermediate doping levels. In all cases, they found that for doped (CH)~ the conductivity was frequency independent up to their maximum frequency. On the other hand, any model of serial segregation into metallic and insulating regions would predict an increasing conductivity at higher frequencies. A frequency dependence has been observed for undoped trans-(CH)x with magnitude and function dependence in agreement with Kivelson's theory of inter-soliton hopping [43]. Mihaly et al. [54] extended the data through microwave measurements (9 GHz) and obtained the ratio oa.~./a~¢, as a function of iodine concentration. In contrast to the behavior expected for metallic islands, Mihaly et al. [54] found oa.c./e~¢, to decrease from a b o u t 103 at y = 0 to around unity at y = 0.03. These results rule o u t the possibility of non-uniformity in shape of metallic islands. Moreover, these studies were carried o u t on samples prepared with conventional doping techniques where dopant non-uniformity is known to exist. We conclude that the proposed metallic islands do not exist in trans-(CH)~, whatever non-uniformity exists is either macroscopic across the film thickness, or primarily radial through the cross-section of the fibrils. The intermediate regime, 0.002 < y < 0.07, is particularly interesting The t h e r m o p o w e r and spin susceptibility results are compared directly in

260 this regime in Fig. 13. We emphasize that identical, uniformly
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The transition from the dilute regime, characterized by Kivelson's inter-soliton hopping [ 42 ], to the non-magnetic intermediate regime is most clearly demonstrated in the thermopower data. The small values and quasilinear (with temperature) behavior are suggestive of delocalized carrier transport. However, these carriers cannot arise from simple band states. In this concentration range, there is no evidence of a finite density of states at the Fermi energy. Moreover, optical studies [55, 56] show the simultaneous existence of the mid-gap transition (originating from the soliton bound state) and the interband transition. Thus a generalized soliton picture, involving delocalized carriers, appears to be implied by the combination of optical, magnetic, and transport data.

261

8. Conclusion In this paper we have established the structure and morphology of stretch-oriented trans-(CH) x films prepared in our laboratory using the basic Shirakawa synthetic method. X-ray scattering measurements on trans-(CH) x prove that the polymer chains are parallel to the oriented fibrils seen by electron microscopy. These results rule out the perpendicular lamellar morphology with chain folding proposed by Wegner as principal structural features. Both the electrical and optical properties of such oriented films indicate that the principal It-electron transport is along (rather than perpendicular t o ) t h e trans-(CH)x chains. In the case of as-grown (non-oriented) films, all structural, transport, and optical data are consistent with the same conclusions. The effects of dopant uniformity on the s e m i c o n d u c t o r - m e t a l transition were reviewed. It was shown that more uniform doping leads to a sharper transition. Finally, it was shown that the transport and magnetic data are not consistent with the percolation model. All the data presented were obtained from samples prepared in our laboratory under identical conditions. Thus, the conclusions regarding the structure and morphology and their relation to the electronic properties are firm. These conclusions can now be tested for other materials prepared with other catalysts and techniques, etc. We emphasize, however, that the ability to generate partially oriented films of trans-(CH)x was critically important to progress in this area. Through broad-based studies of such oriented films, it has been possible to establish the basic s t r u c t u r e - p r o p e r t y relationships of doped trans-( CH)x.

Acknowledgements The research summarized in this review and the preparation of the review were supported by DARPA/ONR on a grant monitored by ONR. The X-ray studies were carried o u t under funding from the NSF-MRL program (DMR 7923647).

References 1 For references see the following review papers: A. J. Heeger and A. G. MacDiarmid, in L. Alcacer (ed.), The Physics and Chemistry o f L o w Dimensional Solids, Reidel, Dordrecht, 1980, p. 353, (the following paper in this volume (p. 393) focuses on the chemical aspects); A. J. Heeger and A. G. MacDiarmid, Chem. Scr., 1 7 (1981) 115. 2 H. Shirakawa, E. J. Louis, A. G. MacDiarmid, C. K. Chiang and A. J. Heeger, J. Chem. Soc., Chem. Commun., (1977) 578. 3 C. K. Chiang, C. R. Fincher, Jr., Y. W. Park, A. J. Heeger, H. Shirakawa, E. J. Louis, S. C. Gau and A. G. MacDiarmid, Phys. Rev. Lett., 39 (1977) 1098. 4 C. K. Chiang, M. A. Druy, S. C. Gau, A. J. Heeger, E. J. Louis, A. G. MacDiarmid, Y. W. Park and H. Shirakawa, J. Am. Chem. Soc., 100 (1978) 1013.

262 5 P. J. Nigrey, A. G. MacDiarmid and A. J. Heeger, J. Chem. Soc., Chem. Commun., (1979) 494. 6 C. K. Chiang, S. C. Gau, C. R. Fincher, Jr., Y. W. Park, A. G. MacDiarmid and A. J. Heeger, Appl. Phys. Left., 33 (1978) 18. 7 Y. W. Park, A. J. Heeger, M. A. Druy and A. G. MacDiarmid, J. Chem. Phys., 73 (1978) 946; Y. W. Park, A. Denenstein, C. K. Chiang, A. J. Heeger and A. G. MacDiarmid, Solid State Commun., 29 (1979) 747. 8 C. R. Fincher, Jr., M. Ozaki, M. Tanaka, D. L. Peebles, L. Lauchlan, A. J. Heeger and A. G. MacDiarmid, Phys. Rev. B, 20 (1979) 1589. 9 P. J. Nigrey, D. MacInnes, D. P. Nairns, A. G. MacDiarmid and A. J. Heeger, J. Electrochem. Soc., 128 {1981) 1651. 10 D. MacInnes, M. A. Druy, P. J. Nigrey, D. P. Nairns, A. G. MacDiarmid and A. J. Heeger, J. Chem. Soc., Chem. Commun., (1981) 317. 11 M. Ozaki, D. L. Peebles, B. R. Weinberger, A. J. Heeger and A. G. MacDiarmid, J. Appl. Phys., 51 (1980) 4252. 12 B. R. Weinberger, S. C. Gau and Z. Kiss, Appl. Phys. Lett., 38 (1981) 555; M. Ozaki, D. L. Peebles, B. R. Weinberger, C. K. Chiang, S. C. Gau, A. J. Heeger and A. G. MacDiarmid, Appl. Phys. Lett., 35 (1979) 83. 13 S. N. Chen, A. J. Heeger, Z. Kiss, A. G. MacDiarmid, S. C. Gau and D. L. Peebles, Appl. Phys. Lett., 36 (1980) 96. 14 T. Ito, H. Shirakawa and S. Ikeda, J. Polym. Sci., PartA-1 Polym. Chem., 12 (1974) 11; 13 (1975) 1943. 15 Y. Kobayashi, personal communication. 16 G. E. Wnek, J. C. W. Chien, F. E. Karasz, M. A. Druy, Y. W. Park, A. G. MacDiarmid and A. J. Heeger, J. Polymer Sci., Polymer Lett. Ed., 17 (1979) 779. 17 H. Shirakawa and S. Ikeda, Synth. Met., 1 (1979 - 80) 175. 18 H. Shirakawa and S. Ikeda, Polym. J., 2 (1971) 231; H. Shirakawa, T. Ito and S. Ikeda, Polym. J., 4 (1973) 460; T. Iko, H. Shirawaka and S. Ikeda, J. Polym. Sci., Polym. Chem. Ed., 13 {1975) 1943. 19 Y. W. Park, M. A. Druy, C. K. Chiang, A. G. MacDiarmid, A. J. Heeger, H. Shirakawa and S. Ikeda, J. Polym. Sci., Polym. Lett. Ed., 1 7 (1979) 195. 20 C. R. Fincher, Jr., D. L. Peebles, A. J. Heeger, M. A. Druy, Y. Matsumura and A. G. MacDiarmid, Solid State Commun., 27 (1978) 489. 21 M. A. Druy, Ph.D. Thesis, Univ. Pennsylvania, 1981 (unpublished); M. A. Druy, C.-H. Tsang, N. Brown, A. J. Heeger and A. G. MacDiarmid, J. Polym. Sci., Polym. Phys. Ed., 17 (1979) 779. 22 G. Wegner, Angew. Chem., Int. Ed. Engl., 20 (1981) 361. 23 G. Lieser, G. Wegner, W. Muller and V. Enklemann, Makromol. Chem., Rapid Commun., 1 {1980)621. 24 G. Lieser, G. Wegner, W. Muller and V. Enklemann, Makromol. Chem., Rapid Commun., 1 (1980)627. 25 F. E. Karasz, J. C. W. Chien, R. Galkiewicz, G. E. Wnek, A. J. Heeger and A. G. MacDiarmid, Nature (London), 282 (1979) 286. 26 A. J. Epstein, H. Rommelman, R. Fernquist, H. W. Gibson, M. A. Druy and T. Woerner, Polymer, i~ press. 27 M. Aldissi, Ph.D. Thesis, Montpellier, 1981 (unpublished). 28 C. R. Fincher, Jr., C.-E. Chen, A. J. Heeger, A. G. MacDiarmid and J. B. Hastings, Phys. Rev. Lett., 48 (1982) 100. 29 T. Akaishi, K. Miyasaka, K. Ishikawa, H. Shirakawa and S. Ikeda, J. Polym. Sci., Polym. Phys. Ed., 18 (1980) 745. 30 R. H. Baughman, S. L. Hsu, L. R. Anderson, G. P. Pez and A. J. Signorelli, in W. Hatfield (ed.), Molecular Metals, Plenum Press, N.Y., 1979, p. 189; R. H. Baughman, personal communication. 31 K. Shimamura, F. E. Karasz, J. A. Hirsch and J. C. W. Chien, Makromol. Chem., Rapid Commun., 2 (1981) 473.

263 32 A. Guinier, X-Ray Diffraction, Freeman, San Francisco, 1963, Ch. 9; see also R. Spal, C.-E. Chen, T. Egami, A. J. Heeger and A. G. MacDiarmid, Phys. Rev. B, 21 (1980) 3110 and references therein. 33 The value 560 ~ - 1 cm-1 was given in ref. 3. For as-grown films, doping to the metallic range leads to conductivities in the range 102 - 103 ~ - 1 cm-1 depending on the dopant and the detailed method of doping. 34 A. J. Heeger, in J. T. Devreese and V. E. van Doren (eds.), Highly Conducting OneDimensional Solids, Plenum Press, New York, 1979, p. 69. 35 C. M. Gould, D. M. Bates, H. M. Bozler, A. J. Heeger, M. A. Druy and A. G. MacDiarmid, Phys. Rev. B, 23 (1981)682. 36 M. Tanaka, personal communication. 37 D. L. Peebles, Thesis, Univ. Pennsylvania, 1980 (unpublished). 38 D. Moses, A. Denenstein, J. Chen, P. McAndrew, T. Woerner, A. J. Heeger and A. G. MacDiarmid, Phys. Rev. B, 25 (1982) 7652. 39 S. Ikehata, J. Kaufer, T. Woerner, A. Pron, M. A. Druy, S. Sivak, A. J. Heeger and A. G. MacDiarmid, Phys. Rev. Lett., 45 (1980) 1123. 40 D. Moses, J. Chen, A. Denenstein, M. Kaveh, T.-C. Chung, A. J. Heeger and A. G. MacDiarmid, Solid State Commun., 40 (1981) 1007. 41 A. J. Epstein, H. Rommelmann, M. A. Druy, A. J. Heeger and A. G. MacDiarmid, Solid State Commun., 38 (1981) 683. 42 S. Kivelson, Phys. Rev. Lett., 46 (1981) 1344. 43 A. J. Epstein, H. Rommelmann, M. Abkowitz and H. W. Gibson, Phys. Rev. Lett., 47 (1981) 1549. 44 D. Moses, A. Denenstein, A. Pron, A. J. Heeger and A. G. MacDiarmid, Solid State Commun., 36 (1980) 219. 45 J. J. Ritsko, E. J. Mele, A. J. Heeger, A. G. MacDiarmid and M. Ozaki, Phys. Rev. Lett., 44 (1980) 1351. 46 J. J. Ritsko, Phys. Rev. Lett., 46 (1981) 849. 47 E. J. Mele and M. J. Rice, Phys. Rev. B, 23 (1981) 5397. 48 Y. Tomkiewicz, T. D. Schultz, H. B. Brom, T. C. Clarke and G. B. Street, Phys. Rev. Lett., 43 (1979) 1532. 49 Y. Tomkiewicz, T. D. Schultz, H. B. Brom, A. R. Taranko, T. C. Clarke and G. B. Street, Phys. Rev. B, 24 (1981) 4348. 50 T. C. Clarke, A. R. Taranko, Y. Tomkiewicz, J. Flood, M. A. Druy, S. Ikehata, T. Woerner, A. J. Heeger and A. G. MacDiarmid, to be published. 51 M. Peo, H. Forster, K. Marke, J. Hocker, J. A. Gardner, S. Roth and K. Dansfeld, Proc. Int. Conf. on Low-Dimensional Conductors, Boulder, Colo., Aug., 1981, to be published in Mol. Cryst. Liq. Cryst. 52 P. M. Grant and M. Krounbi, Solid State Commun., 36 (1980) 291. 53 A. J. Epstein, H. W. Gibson, P. M. Chaikin, W. G. Clark and G. Gruner, Phys. Rev. Lett., 45 (1980) 1730. 54 G. Mihaly, G. Vancso, S. Pekker and A. Janossy, Synth. Met., 1 (1980) 357. 55 N. Suzuki, M. Ozaki, S. Etemad, A. J. Heeger and A. G, MacDiarmid, Phys. Rev. Lett., 45 (1980) 1209; Erratum Phys. Left., 45 (1980) 1983. 56 A. Feldblum, J. Kaufman, T.-C. Chung, A. J. Heeger and A. G. MacDiarmid, Phys. Rev. B, 26 (1982) 815.