Structure and electroconductivity of a sheared liquid-crystalline polyacetylene derivative: Poly(5-p-(trans-4-pentacyclohexyl)phenoxy-1-pentyne)

Article abstract of DOI:10.1021/ma010684a

The crystal structure of the compound 5-p-(trans-4-pentacyclohexyl) phenoxy-1-pentyne (PCH503A) was determined by X-ray diffraction. The structure and orientational behavior of the polymer (PPCH503A), prepared from the monomer compound, was examined. As a result, it was found that the polymers form an alternating comblike conformation of side chains and are stacked into a smectic hexagonal structure. When shear stress was applied to the polymers, polymeric chains were oriented with their axes of mesogenic groups parallel to the shearing direction; that is, the polymer backbones were aligned perpendicular to the shearing direction. The oriented polymer exhibited a high anisotropy of electroconductivity, i.e., σ_∥ (the electroconductivity parallel to the backbone) became 1.7 × 10^-5 S/cm upon doping iodine, and the ratio of σ_∥ to σ_⊥ (the electroconductivity perpendicular to the backbone) w about 10.

Full text of DOI:10.1021/ma010684a

Macromolecules 2002, 35, 1307-1313
1307
Structure and Electroconductivity of a Sheared Liquid-Crystalline
Polyacetylene Derivative:
Poly(5-p-(trans-4-pentacyclohexyl)phenoxy-1-pentyne)
Hid ek a tsu Ku r od a ,† Hir om a sa Goto,‡ Ka zu o Ak a gi,‡ a n d Ak iyosh i Ka w a gu ch i*,†
Department of Applied Chemistry, Faculty of Science and Engineering, Ritsumeikan University,
Nojihigashi, Kusatsu, Shiga 525-8577 J apan; and Institute of Materials Science,
University of Tsukuba, Ibaraki 305-8573 J apan
Received April 20, 2001; Revised Manuscript Received August 6, 2001
ABSTRACT: The crystal structure of the compound 5-p-(trans-4-pentacyclohexyl) phenoxy-1-pentyne
(PCH503A) was determined by X-ray diffraction. The structure and orientational behavior of the polymer
(PPCH503A), prepared from the monomer compound, was examined. As a result, it was found that the
polymers form an alternating comblike conformation of side chains and are stacked into a smectic
hexagonal structure. When shear stress was applied to the polymers, polymeric chains were oriented
with their axes of mesogenic groups parallel to the shearing direction; that is, the polymer backbones
were aligned perpendicular to the shearing direction. The oriented polymer exhibited a high anisotropy
of electroconductivity, i.e., σ| (the electroconductivity parallel to the backbone) became 1.7 × 10-5 S/cm
upon doping iodine, and the ratio of σ| to σ (the electroconductivity perpendicular to the backbone) was
about 10.
In tr od u ction
n-pentyl group as a terminal moiety, we bagan analysis
of the crystal structure of the monomer, and on the basis
of the result, we investigated the liquid crystalline
texture of the polymer by X-ray and electron diffraction.
Furthermore, since it was not examined very much in
ref 2, dependence of the molecular orientation on shear
stress and on temperature was studied, together with
the electroconductivity of the sheared sample.
Polyacetylene is an archetypical conducting polymer,
but it further warrants research for practical use
because it is difficult to process the polymer and because
the polymer is easily oxidized. Chemically modifying
polyacetylene by introducing an alkyl or aromatic sub-
stituent into its backbone, many researchers have tried
to make the polymers soluble and fusible. In particular,
if the substituent is a liquid crystalline mesogen, the
polyacetylene derivatives thus modified are soluble and
fusible, and they form a liquid crystalline texture by self-
assembling of the mesogen groups.1-4 Because of their
liquid crystalline nature, the derivatives are advanta-
geous because they can be oriented easily by applying
external force, such as shear stress1-3 or an electric or
magnetic field.1,2 Previous papers have reported and
compared the structure and electroconductivity of liquid
crystalline polyacetylene homologues to the one in the
present paper: poly(5-p-(trans-4-pentacyclohexyl)phe-
noxy-1-pentyne) (PPCH503A). When external force is
applied to the liquid crystalline state, the liquid crystal-
line groups are aligned with their molecular axis in the
direction parallel to the external force, and they are
organized into a layered structure of smectic form. In
the ordered texture, the electroconductivity is higher in
the direction perpendicular to the applied force, i.e., in
the direction parallel to the polymer backbone compris-
ing conjugated double bonds. Thus, it was found that
the electrical properties are largely improved by aligning
polymer chains. However, the layered structure is not
always explained in the molecular dimension, although
a textural model is demonstrated from results of X-ray
diffraction and modeling from a molecular mechanics
simulation.5 In the present work, focusing on the
present homologue with a spacer -(CH2)3O- and an
Exp er im en ta l Section
A monomer compound, 5-p-(trans-4-pentacyclohexyl)phe-
noxy-1-pentyne (PCH503A), was synthesized and polymerized
according to the scheme in ref. 1. Seven grams of p-(trans-4-
pentylcyclohexyl)phenol and 2.9 g of 5-chloro-1-pentyne were
dissolved in sodium ethoxide. After adding 0.46 g of potassium
iodide to the reaction mixture, as a catalyst, the mixture was
refluxed for 1 week. A 1.5 g amount of the thus-prepared
monomer was dissolved in 9.6 mL of THF, with 0.022 g of [Rh-
(NBD)Cl]2 catalyst, and this was polymerized by stirring the
reaction for 24 h at room temperature. The polymerization
reaction was terminated by pouring the reaction mixture into
a large amount of methyl alcohol. The polymer product was
filtered off, washed with methyl alcohol, and then dried under
a vacuum. The molecular weight and polydispersity (Mw/Mn)
of the polymer were 1.9 × 104 and 1.44, respectively.
A single crystal of PCH503A was prepared, by slow evapo-
ration of solvent for 3 weeks, from a solution in ethyl alcohol.
X-ray diffraction data of the single crystal were accumulated
by a four-circle auto-measurement instrument, a Rigaku AFC-
5R diffractometer. The X-ray beam was Cu KR monochroma-
tized with a graphite monochromator. All of thus-accumulated
data on PCH503A were processed using the teXsan graphic
crystal software package.
We examined the relation between the molecular orientation
and the strain rate and/or shearing temperature, when the
polymer was solidified from the melt or liquid crystalline state
after shearing. Using a system as shown in Figure 1, the
structural change was monitored when shear stress was
applied to the polymer melt. The hot-stage consists of two glass
plates: one is settled in a temperature-controllable hotplate,
and the other is able to slide on it. The rate of shear strain
was regulated by changing the drawing rate of the upper glass
plate, by use of a motor. The polymer was dissolved at a
* To whom correspondence should be addressed. Fax: +81-(0)-
77-561-2659. E-mail: akiyoshi@se.ritsumei.ac.jp.
† Ritsumeikan University.
‡ University of Tsukuba.
10.1021/ma010684a CCC: $22.00 © 2002 American Chemical Society
Published on Web 01/12/2002

1308 Kuroda et al.
Macromolecules, Vol. 35, No. 4, 2002
Ta ble 1. Cr ysta llogr a p h ic Da ta of P CH503A
chem formula
fw
C
22H32
O
312.49
cryst dimens, mm
cryst color, habit
no. of reflcns
cryst syst
space group
lattice params
a, nm,
b, nm
c, nm
R, deg
â, deg
0.30 × 0.20 × 0.40
colorless, plate
2818
triclinic
P1 (No. 2)
1.1251(2)
1.7182(2)
0.5447(2)
93.319(2)
103.731(2)
72.580
γ, deg
Z value
d
µ(Cu KR), cm-1
R (Rw) factor, %
2
calcd, g/cm3
1.060
4.72
7.1 (12.3)
unit cell, respectively. Two molecules with identical
conformation are packed in the cell in an antiparallel
way. The end pentyl group of the compound has a trans
conformation. The cyclohexyl group is in a stable chair
conformation. The plane consisting of C13, C14, C16, and
C17 of the cyclohexyl group, is almost perpendicular to
the plane of the phenyl group. The chain conformation
from C3 to C6 is gauche, and as a whole, the molecule is
extended in an almost straight fashion. The linear
distance from C2 to C22 is 1.895 nm, corresponding to
the length of the side chain of the polymer.
Figure 3 shows an electron diffraction pattern of a
sheared thin film. Sharp spots up to the third order are
observed in the direction parallel to the shearing direc-
tion. The diffraction feature explains that the samples
are composed of regularly stacked layers, the normal
of which is parallel to the shearing direction. The long
spacing is evaluated at 3.726 nm from the spots and is
comparable to the distance between the centers of
gravity of adjacent neighboring layers. From X-ray
analyses of the monomer crystal, it is found that the
length of the sequence of C2 to C22, depicted in Figure
2a, is 1.895 nm. The long spacing is nearly equal to twice
the C2 to C22 length. This equality explains that the side
chains are fully extended in a conformation similar to
that of monomer, and that their axes are oriented
almost perpendicular to the backbone. Strictly speaking,
however, the long spacing is a little shorter than twice
the C2 to C22 length, when the thickness of the backbone
chain and the van der Waals radii of atoms at the end
groups are taken into account. The side chains may be
a little tilted from normal to the layer plane. Akagi et
al. proposed that the polymerization proceeds through
a head-to-head or a tail-to-tail reaction.2 On the basis
of their idea of the polymerization mechanism, they put
forward such a polymer configuration and stacking plan
as shown in Figure 4: The fully extended, linear side
groups are arrayed such that they protrude in the
direction perpendicular to the linear backbone, alter-
natively on both sides along it, and the alternating
comblike molecules are stacked into a layer structure.
This model explains the long period obtained from the
X-ray diffraction profile. The present ED is basically
interpreted by this model. It is found, from DSC and
IR measurements, that when the polymer is heated in
the temperature range 167-175 °C, the cis form is
easily isomerized into the trans form.6,7 A specimen
annealed above the temperature exhibits an ED pattern
similar to that in Figure 3, giving an identical long
period of 3.726 nm. As the long periods of both forms
F igu r e 1. Apparatus for measuring the birefringence of
polymers by changing temperature and rate of shear: (a)
overview, and (b) sketch of the hot stage.
concentration of 0.9 wt % in decahydronaphthalene (decalin).
Some drops of the solution were put between two slide glasses
kept at a given temperature. After the solvent was evaporated,
the upper glass plate was slid at that temperature and cooled.
The degree of molecular orientation of the resultant sample
was graded at high, medium, or low, by visual evaluation of
birefringence, by monitoring with a polarizing microscope
equipped with a CCD camera.
Very thin samples for electron microscopy were prepared
using another small-scale, temperature-controllable copper
device, the setup and usage of which was the same as for the
hot-stage of Figure 1. By confirmation of high birefringence
with a polarizing microscope, highly oriented samples were
selected out of the sheared samples. The upper glass plate was
easily detached from the shearing device. Subsequently, it was
dipped in a dilute hydrofluoric acid solution to float off a thin
film from it. The thin film was picked up on a Cu grid for
electron microscopy.
Samples for measurement of electroconductivity were also
prepared at 145 °C, in the manner described above. Electro-
conductivity of the samples was measured by a two-probe
method. The sample was doped with iodine vapor at room
temperature.
Resu lts a n d Discu ssion
Crystallographic data of PCH503A that were obtained
by X-ray diffraction analyses are summarized in Table
1. Parts a and b of Figure 2 show the molecular
conformation of PCH503A and packing models in the

Macromolecules, Vol. 35, No. 4, 2002
Liquid-Crystalline Polyacetylene Derivative 1309
F igu r e 2. (a) Molecular conformation of PCH503A in a crystal, and (b) the packing model of molecules in the unit cell: the
projection on the a-c basal plane (left) and the three-dimensional view (right).
are identical, the stacking models for cis and trans forms
of PPCH503A are considered to be as shown in Figure
4. We shall inquire into the conformations of the
backbone and side chain for both cis and trans stacking
patterns, taking account of the following requirements
experimentally obtained here: (1) the side chains have
an almost fully extended conformation, (2) their axes
are oriented nearly perpendicular to the backbone, and
(3) the long periods for both cis and trans backbone
conformations are identical. These requirements are
fulfilled when the cis and trans forms of molecules take
the conformations shown in the magnified models in
Figure 4, parts a and b, respectively. The distance from
the central backbone to the end atom of the side chain
may differ between cis and trans forms, depending on
the rotation angle around the C2-C3 or C′2-C′3 bond.
Of course, this results in a change of the long period.
The difference in the side chain length would be offset
by loose stacking between neighboring layers, and
consequently, the long spacing would remain un-
changed. This is acceptable in consideration of the
rather disordered stacking of layers as discussed later
on.
We should notice here the following about the ED
pattern: Meridian reflections only up to the 003 spot
are observed even if the ED patterns are deliberately
taken, and no discrete reflections are observed in the
equatorial direction. In the X-ray diffraction profile, a
reflection of the third order in the small angle region is
also barely observable.1 First, we shall consider the
diffraction feature that only reflections up to the third
order are observed in electron or X-ray diffraction.
Diffraction theory tells us that how high an order of
diffraction may be observable will depend on the order
of stacking, in the present case, of the alternating
comblike polymers.8,9 According to Hosemann’s para-
crystalline theory,8 we evaluated the g value, which is
a measure of the stacking disorder. It means the
fluctuation of the distance between the centers of gravity
of layers, and is defined as ∆l/l, where l is the average

1310 Kuroda et al.
Macromolecules, Vol. 35, No. 4, 2002
bone carbons, to which the mesogens extending on the
same side are anchored, is estimated at 0.5 nm at a
maximum in the trans conformation. The four-carbon
“topological” repeat is too short for the side chains to
accommodate in the coplanar way. To avoid their steric
hindrance, possibly it might be accompanied by a slight
twisting of the backbone. It could result in the tilting
of moiety chains to the layer surface, leading to a shorter
long period and to loose side-by-side packing of the side
chains. Although the explanation is speculative, the
packing model of the smectic hexagonal mode shown in
Figure 6 is probably set up to explain the diffraction
feature. Since the present solidus state is not “crystal-
line” in the strict sense, the change from liquid to solidus
state is described not as “crystallization”.
The polymer sample was sheared by changing the
rate of shear at various temperatures, followed by
cooling. Through visual evaluation, we estimated the
birefringence of samples thus obtained. Figure 7 shows
an example of polarizing optical micrographs of the
sheared sample. The photographs indicate large bire-
fringence, i.e., a high degree of molecular orientation.
When such high birefringence as shown in Figure 7 was
observed, we defined this as a state of “high molecular
orientation”. In the same way as shown for a main-chain
thermotropic liquid crystalline polymer,10 the results are
mapped as a diagram of molecular orientation in the
temperature-shear rate space in Figure 8. Two things
are to be noted from the present experimental view-
point. First, limit lines are not always as definite as
shown in Figure 8 but have some range, because the
birefringence was not so definitely determined, since it
is measured by visual observation. Second, the birefrin-
gence evaluation was done using samples in which
solidification occurred by cooling, not in situ during
shearing. The relaxation behavior of strained polymers
after cessation of shear is not taken into account.11 The
upper line (1) shows a limiting line showing whether
birefringence is observed; i.e., solidification is acceler-
ated by shear stress, accompanying molecular orienta-
tion in the sample. In the region above the upper limit
line, the sample remains molten, even under shear, and
exhibits no birefringence during shearing. The lower
line (3) shows the limit that the following situation is
experimentally accessible: The sample still remains
liquid crystalline when cooled at a given temperatures
the molten phase is not retained but is transformed into
the liquid crystalline phase at the temperaturesand
solidification does not occur until shear stress is applied.
In other words, even if no shear stress is applied to the
sample, solidification takes place as soon as the sample
is brought to temperatures below line 3, and conse-
quently, no molecular orientation takes place. Line 2
denotes the borderline indicating the phase before being
sheared; above the line, the material is in the molten
phase, and below it, it is in the liquid crystalline phase.
The polymer exists in the liquid crystalline state over
the temperature range indicated by a vertical, arrowed
bar. Line 2 just corresponds to the upper limit of the
temperature range. In the limited region between lines
1 and 2, the molten polymer chains are aligned when
the shear stress is applied. Below line 2, the liquid
crystalline state is realized. The molecularly oriented
texture in the area between lines 2 and 3 is produced
from the liquid crystalline phase. From the phase
diagram, we see that the liquid crystalline phase is not
always transformed into a uniaxially oriented state of
F igu r e 3. (a) Electron diffraction pattern of a sheared thin
film of PPCH503A. The sheared direction is vertical. The
sample was prepared at 145 °C and at a rate of shear over 40
s-1
.
distance between the centers, i.e., the long spacing.
According to the criterion that loss of reflection may be
characterized by the range in which the intensity
maxima do not exceed the background by more than
10∼20%, the g value is estimated at about 0.08 in the
present case, where reflections up to the third order are
observable. This value corresponds to ∆l ∼ 0.3 nm. The
difference in side chain length between cis and trans
forms would be trivial in comparison with the long
period fluctuation.
Figure 5 shows the wide-angle X-ray diffraction
profile of PPCH503A polymer. Sharp, small-angle X-ray
diffractions up to the third order and a rather sharp,
wide-angle single peak are observed.1 A similar diffrac-
tion feature is commonly observed in other liquid
crystalline polyacetylenes with different long chain
moieties from the present polymer.3,4 Since the wide
angle reflection (2θ ) 18.68°) is quite different in
appearance from the so-called amorphous halo, it is
reasonable to consider that the reflection should be
attributed to a diffraction of a crystalline nature. From
the peak angle, the lattice spacing is evaluated at 0.47
nm. In consideration of the layered texture, it is natural
to assume that this reflection should be caused by the
lateral packing of mesogen chains. The diffraction
feature, with a single peak, implies that the packing
mode is nearly hexagonal, although less-ordered. It is
expected that when the chain packing is hexagonal,
higher order reflections at x3 and, possibly, twice the
fundamental wave vector should be observed. However,
no such reflections are actually observed. Paying atten-
tion to the broad profile again, it is explained as
follows: The fluctuation of side-by-side packing among
side chains is so large, that is, roughly, g value ≈ 0.2,
that only the first order of diffraction could be observ-
able. It explains that they are very loosely packed
laterally. On the assumption of hexagonal packing, the
average distance between the neighboring mesogen
stems is calculated at 0.54 nm, with a fluctuation of
∼0.1 nm. Furthermore, assuming the side chains as rods
(cf. Figure 2a), their diameter is calculated at about 0.6
nm, including van der Waals radii of hydrogen atoms.
On the other hand, the separation of every four back-

Macromolecules, Vol. 35, No. 4, 2002
Liquid-Crystalline Polyacetylene Derivative 1311
F igu r e 4. Stacking models of (a) cis and (b) trans forms of comblike polymers and magnified sketches of their core parts. The
side chain molecules are twisted about the C2-C3 bond, so that the H1 and H2 hydrogen atoms are directed upward from the
figure plane (the plane containing the backbone), so as to release repulsion with the neighboring H3 hydrogen. About the C′2-C′3
bond, the side chains would be rotated reversewise to the above rotation, so that the two mesogens would be disposed in the
antipode position with respect to the intervening double bond.
the strain rate and/or temperature in the processing of
polymers and plastics. In liquid crystalline substances,
it is expected, as well, that the larger the rate of shear,
the higher the solidification temperature. In reality, the
orientation temperature of molecules increases with an
increased rate of shear, while the rate of increase
becomes dull at a rate of shear above 20 s-1. Here,
however, the following experimental uncertainty is to
be noticed: When shear stress over 50 s-1 was applied,
polymer molecules solidified so quickly that they ad-
hered to the slide glass, and hence some parts of the
sample were broken due to sliding of the glass during
observation.
Figure 9 shows a typical example of the changes in
electroconductivity (σ) of the oriented and unoriented
cis samples (which were prepared at 145 °C) with doping
of iodine. σ| and σ denote the electroconductivities in
the direction parallel to and perpendicular to the
extended backbone, respectively, that is, those in the
directions perpendicular to and parallel to the shearing
direction, respectively. For the first hour in Figure 9,
the sample was not exposed to iodine vapor, and after
that, iodine vapor was introduced. The change in
electroconductivity of unoriented sample with doping
iodine is shown as curve 3 in Figure 9. Before doping,
F igu r e 5. Wide-angle X-ray diffraction profile of PPCH503A.
texture. Whether or not the oriented texture is obtained
is largely dependent on the relaxation behavior of
strained chains.11
The lowest rate of shear to induce a molecular
orientation is 3 s-1 at about 115 °C, and with an increase
in the rate of shear, the solidification temperature
becomes higher. Acceleration of crystallization by shear
or elongational strain, which leads to an increase of
crystallization temperature, is common in the crystal-
lization of polymers, and hence it is important to control

1312 Kuroda et al.
Macromolecules, Vol. 35, No. 4, 2002
F igu r e 8. Map of molecular orientation in the temperature-
shear rate space. In the shaded region, molecular orientation
takes place. The vertical arrow bar indicates the temperature
range over which the liquid crystalline state occurs under
quiescent conditions.
F igu r e 6. Three-dimensional model of the smectic texture of
PPCH503A. Neighboring comblike chains in a layer are shifted
one another in the direction parallel to the backbone, by half
the distance between neighboring side chains along the chain,
and thus are packed in a nearly hexagonal mode.
Figu r e 9. Change in electroconductivity of PPCH503A against
doping time: σ|(1) and σ (2) of an oriented sample and σunorient
(electroconductivity of unoriented sample) (3).
Ta ble 2. Electr ocon d u ctivity of P P CH503A befor e a n d
a fter Dop in g
σ (S/cm)
σ| (S/cm)
σ
unorient (S/cm)
intact
doped
8.5 × 10-9
4.7 × 10-8
2.5 × 10-8
1.5 × 10-6
1.7 × 10-5
1.2 × 10-8
of molecular orientation, and it suggests that electro-
conductivity is very sensitive to molecular orientation.
In summary, it was found that a side-chain-type
polymer, poly(5-p-(trans-4-pentacyclohexyl)phenoxy-1-
pentyne), forms an alternating comblike conformation
and that when shear strain was applied to the polymer,
it forms a smectic texture, where the normal to the
smectic layers is parallel to the shearing direction, that
is, the long axes of the side-chains are aligned along the
direction. Using X-ray structural data of the monomer,
it was found, experimentally, that the side chains are
extended in an almost straight fashion in the texture.
The textural structure of the sheared polymers was
greatly affected by the rate of shear. Here, the region
of the temperature and rate of shear where the molec-
ular orientation was performed was depicted as a
temperature-shear rate space diagram. The molecular
orientation does not always orient over the entire range
of liquid crystalline state when shear stress is applied;
it depends on the temperature and the rate of shear.
The highly oriented sample exhibited high electrocon-
ductivity. The diagram is expected to be useful to
determine the processing condition needed to obtain a
molecularly ordered structure exhibiting high perfor-
mance of such physical properties as electroconductivity.
F igu r e 7. (a) Polarizing optical microscopic images of a
sample prepared as follows: sheared at 145 °C, at 25 s-1, and
cooled. The arrows of a and p denote the directions of the
analyzer and polarizer, respectively. (b) Polarizer and analyzer
rotated by 45° counterclockwise from the position in Figure
8a. Between the analyzer and polarizer, a sensitive color plate
with R ) 530 µm is inserted. The shearing direction is vertical.
σ and σunorient (the electroconductivity of the unoriented
sample) are almost the same, and both are slightly lower
than σ|. Immediately after the sample is doped, σ|
rapidly increases by 3 orders of magnitude while σ
gradually increases by 2 orders of magnitude over
almost 3 h. σunorient decreases slightly with doping.
Electroconductivities before doping and after prolonged
doping are summarized in Table 2. For the oriented
sample that was prepared by applying the magnetic
field of 0.7∼1.0 T, σ| and σ have been measured.1 σ| in
the present sample is significantly larger, roughly by 1
order of magnitude, than that of the magnetically
oriented sample, while σ s of both samples are compa-
rable. This may be caused by the difference in the degree
Ack n ow led gm en t. This work was partly supported
by a Grant-in-Aid for Scientific Research on Priority
Areas, “Mechanism of Polymer Crystallization” (No.
12127207), from the Ministry of Education, Culture,
Sports, Science, and Engineering, and partly by the
Ritsumeikan Fund for Research and Faculty, 2000. The

Macromolecules, Vol. 35, No. 4, 2002
Liquid-Crystalline Polyacetylene Derivative 1313
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authors are grateful to Professor M. Cakmak of the
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criticism of the manuscript.
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MA010684A