Liquid crystalline rod-coil block copolymers by stable free radical polymerization: Synthesis, morphology, and rheology

Article abstract of DOI:10.1021/ma021573u

We have synthesized and characterized rod-coil diblock, triblock, and starblock copolymers of controlled molecular weight by stable free radical polymerization, where the rod part consists of "mesogen jacketed liquid crystalline segments" (MJLCP). The MJLCP segmented examined in our studies is poly{2,5-(4-butylbenzol)oxystyrene} (PBBOS) while the coil part is poly(styrene). SAXS studies of the block copolymers at room temperature predominantly show only a first-order maximum in their SAXS spectra. Above the T_g of the two blocks (～150°C) no phase-separated morphology could be detected in temperature-dependent SAXS measurements of the block copolymers. Slow cooling regenerates the morphology, indicating very slow ordering kinetics. The effects of variation in molecular architecture and increasing topological constraint in going from homopolymer to star block copolymer on melt rheology (i.e., viscosity, relaxation times, and activation energies) are reported. In the master curves of all the block copolymers and homopolymers studied, nonterminal behavior was observed at low shear frequencies along with the absence of an entanglement plateau. In the PBBOS homopolymer an increase in molecular weight resulted in a decrease in the modulus of the nematic phase that was associated with a greater extent of shear orientation. On varying the molecular architecture from homopolymer to diblock to starblock copolymer, a decrease in the complex viscosity was observed that is attributable to the presence of polystyrene as the second block and the effect of topological constraints on molecular relaxation. The activation energies also follow the same trend as complex viscosity and relaxation times.

Full text of DOI:10.1021/ma021573u

Macromolecules 2003, 36, 3357-3364
3357
Liquid Crystalline Rod-Coil Block Copolymers by Stable Free Radical
Polymerization: Synthesis, Morphology, and Rheology
P a d m a Gop a la n , Yu a n m in g Zh a n g, Xu efa Li, Ulr ich Wiesn er , a n d
Ch r istop h er K. Ober *
Department of Materials Science and Engineering, Cornell University, Ithaca, New York 14853
Received October 10, 2002; Revised Manuscript Received February 20, 2003
ABSTRACT: We have synthesized and characterized rod-coil diblock, triblock, and starblock copolymers
of controlled molecular weight by stable free radical polymerization, where the rod part consists of
“mesogen jacketed liquid crystalline segments” (MJ LCP). The MJ LCP segment examined in our studies
is poly{2,5-(4-butylbenzoyl)oxystyrene} (PBBOS) while the coil part is poly(styrene). SAXS studies of
the block copolymers at room temperature predominantly show only a first-order maximum in their SAXS
spectra. Above the Tg of the two blocks (∼150 °C) no phase-separated morphology could be detected in
temperature-dependent SAXS measurements of the block copolymers. Slow cooling regenerates the
morphology, indicating very slow ordering kinetics. The effects of variation in molecular architecture
and increasing topological constraint in going from homopolymer to star block copolymer on melt rheology
(i.e., viscosity, relaxation times, and activation energies) are reported. In the master curves of all the
block copolymers and homopolymers studied, nonterminal behavior was observed at low shear frequencies
along with the absence of an entanglement plateau. In the PBBOS homopolymer an increase in molecular
weight resulted in a decrease in the modulus of the nematic phase that was associated with a greater
extent of shear orientation. On varying the molecular architecture from homopolymer to diblock to
starblock copolymer, a decrease in the complex viscosity was observed that is attributable to the presence
of polystyrene as the second block and the effect of topological constraints on molecular relaxation. The
activation energies also follow the same trend as complex viscosity and relaxation times.
In tr od u ction
crystalline. Their liquid crystallinity may be attributed
to the strong steric interactions among the side groups
and between the side groups and the main chain
segments.5b In recent work the authors have synthe-
sized poly[di(4-heptyl)vinyl terephthalate] and reported
the formation of a hexagonal columnar LC phase. This
supramolecular hexagonal columnar phase, according
to the authors, is formed by pendent nonmesogenic
templates and has the highest liquid crystalline order
observed in MJ LCPs so far. The unique molecular
design of MJ LCP’s induces liquid crystallinity even in
polymers which lack the conventional characteristics of
rigid segments and the presence of flexible spacers.5c
Our interest in MJ LCPs is to design rod-coil block
copolymers with the rod part consisting of the MJ LCP.
The unique molecular design of MJ LCPs could poten-
tially lead to different morphologies than those observed
to date in rod-coil systems. To synthesize these rod-
coil polymers, we must consider that most of the
mesogenic groups contain functional groups (such as
ester or amide links) which limits synthetic strategy.
Typically, anionic polymerization has been used to
synthesize rod-coil block copolymers.5d This choice
limits the selection of monomers to anionically poly-
merizable monomers and excludes many highly func-
tional monomers owing to significant side reactions.
Stable free radical polymerization (SFRP) would be an
ideal technique to apply to the synthesis of a wider
range of functionalized rod-coil block copolymers8 as
it offers the ease of free radical polymerization with
other features of living polymerization such as the
synthesis of a variety of star and block architectures.9
Recently, Zhou et al. have used SFRP to synthesize rod-
coil diblock copolymers of poly{styrene-b-2,5-bis[4-
methoxyphenyl]oxycarbonyl}styrene}.10 They also stud-
ied the self-assembly behavior in dilute solution by both
Side chain LC polymers (SCLCP’s) with mesogenic
groups laterally attached to the polymer backbone
constitute a distinct class of liquid crystals. Unlike the
conventional SCLCPs with end-on attachment of meso-
gens, lateral attachment of mesogens hinders the rota-
tional motion along the long axis of the rods. Weissflog
et al.1 synthesized the first such SCLCP. Since then, a
number of laterally attached side chain LCPs have been
reported including poly(siloxanes),2 poly(acrylates),3 and
norbornene derivatives.4 All these polymers have a
flexible linker connecting the mesogen to the backbone.
The family of SCLCPs with laterally attached mesogenic
groups was further extended by Zhou et al. to “mesogen-
jacketed liquid crystal polymers” (MJ LCPs).5a MJ LCPs
have mesogenic groups attached laterally to the polymer
backbone through their center of gravity without any
spacers or with very short spacers. The lack of spacers
leads to an extended LC region in the polymer and
properties similar to semirigid main chain LC polymers.
In most of the SCLCPs, if there are no flexible spacers,
the motion of the mesogenic side groups is directly
coupled to the motion of main chain segments, prevent-
ing any LC orientation.6 However, MJ LCPs form a
nematic liquid crystalline phase even in the absence of
a critical length of flexible spacer. Many MJ LCPs
reported to date are based on 2,5-disubstituted vinyl-
hydroquinone. For example, 2,5-bis[(4-butylbenzoyl)oxy]-
styrene (BBOS) has rigid and bulky benzoate groups at
the 2,5-positions of the phenyl ring.5a Despite these
bulky substituents, BBOS polymerizes efficiently under
both conventional5a and stable free radical7 polymeri-
zation conditions. MJ LCPs of the type poly[di(cyclo-
hexyl)vinyl terephthalate] have been reported, which do
not have rigid-rod mesogenic groups but are liquid
10.1021/ma021573u CCC: $25.00 © 2003 American Chemical Society
Published on Web 04/08/2003

3358 Gopalan et al.
Macromolecules, Vol. 36, No. 9, 2003
static and dynamic light scattering techniques and
reported a core and shell structure in xylene, a nonsol-
vent for the rod.
It is interesting to examine the melt flow behavior of
the MJ LCPs, given their unique molecular architecture.
To date, neither mechanical property studies nor rheo-
logical studies have been performed on MJLCPs. SCLCPs
with laterally fixed mesogenic groups are relatively new
and have not been studied widely. To the best of our
knowledge, the only work on the rheology of liquid
crystals with side-on fixed mesogenic groups has been
reported by Berghausen and co-workers.13 They exam-
ined a nematic side-group liquid crystalline polymer
with polymethacrylate backbone and a hydroquinone
bis(benzoate) mesogen coupled laterally through a
spacer of 11 methylene units. Compared to the end-on
mode SCLCPs, these researchers found that the side-
on type could not be aligned by large-amplitude oscil-
latory shear and that the master curves of G′ and G′′
did not show any plateau modulus, indicating lack of
entanglements for the molecular weights examined.
Our earlier studies on the polymerization behavior
of laterally attached mesogens such as 2,5-bis(4-butyl-
benzoyl)oxystyrene (BBOS) under SFRP conditions
showed unusually high reactivity compared to other
styrenics and by varying the reaction time good control
of molecular weight could be achieved.7,11 On the basis
of these observations, here we report the synthesis of
selected molecular architectures including diblock, star-
block, and ABA triblock polymers of MJ LCP segments
with polystyrene. The bulk morphology of these systems
is studied using cross-polarized optical microscopy, wide-
and small-angle X-ray scattering (WAXS and SAXS),
and transmission electron microscopy (TEM). Further-
more, their basic rheological properties were examined
in order to study the effects of molecular architecture
and increasing topological constraint on their dynamic
behavior.
F igu r e 1. Structure of [4-butylbenzoyloxy]styrene (BBOS).
the reaction of 2-vinyl-1,4-dihydroxybenzene and 4-butylben-
zoyl chloride according to the method described by Zhou et
al.5a 1H NMR (CDCl3): δ ) 0.95, 6H for CH3a, 1.35, 4H for
CH2b, 1.65, 4H for CH2c, 2.70, 4H for CH2d, 5.10-5.90, 2H for
dCH2c, 6.60-6.95, 1H for -CHfd and 7.00-8.30, 11H for the
phenylene rings. As reported earlier,5a BBOS exhibits a smectic
phase between its melting temperature 68 °C and clearing
temperature 95 °C.
Polystyrene Macroinitiators. Three different TEMPO-based
initiators were used as shown in Figure 2. These are (a) unimer
with one initiating site which generates diblock copolymers,
(b) diarm initiator with two initiation sites which generates
triblock copolymers, and (c) star initiator with three initiating
sites which generates starblock copolymers.
In all the block copolymers synthesized in this series,
styrene was used as the first block; therefore, the first step is
the synthesis of macroinitiator from the three different initia-
tors. In a typical procedure styrene was mixed with the
required amount of one of the above initiators in a Schlenk
flask containing a magnetic stir bar. The reaction mixture was
purged with nitrogen and subjected to four freeze-thaw cycles
to remove any dissolved oxygen and sealed under vacuum. The
flask was placed directly into an oil bath preset at 130 °C.
Typically, polymerizations were carried out for 50-75 h
depending on the desired molecular weight. The macroinitiator
was isolated by repeated precipitation in methanol and dried
under vacuum overnight.
Homopolymer of BBOS (PBBOS). Homopolymer of BBOS
was synthesized by typically mixing 300 mg (0.65 mmol) of
BBOS with 6.86 mg of uniarm initiator (0.026 mmol) in a 2
mL ampule containing a magnetic stir bar. The polymerization
mixture was purged with N2, subjected to four cycles of freeze-
thaw, and sealed under vacuum. The sealed ampule was
placed in a preheated oil bath at 125 °C and polymerized for
8 h. The reaction mixture was precipitated in methanol and
purified by column using dichloromethane as solvent to obtain
a BBOS homopolymer of Mn 5800 and PDI 1.30.
Exp er im en ta l Section
P olym er Ch a r a cter iza tion . A Varian XL 400 NMR using
chloroform-d as solvent with tetramethylsilane as internal
standard was used to obtain the NMR spectra of all the
compounds. Thermal analysis was carried out under nitrogen
atmosphere by means of Perkin-Elmer DSC 7 and TGA 7
instruments at a heating rate of 10 °C/min. Measurements of
both molecular weight and polydispersity of all polymer
samples were carried out using Waters Ultrastyragel HT
columns operating at 40 ° C. THF was used as the solvent,
and the GPC was operated at 1 mL/min. Solution concentra-
tions of 1 mg/mL for PBBOS, 2 mg/mL for PDAS, and solution
volumes of 50 µL were employed. Molecular weights were
calculated from GPC elution volume data using monodisperse
polystyrene standards.
Mor p h ology Exa m in a tion s. Bulk samples were cast from
5 wt % toluene solution by slow evaporation over 1 week and
then annealed at 150 °C for 4 days under high vacuum followed
by slow cooling over 6 h. Cross-polarized optical microscopy
(POM) and a hot stage were used to examine the temperature
dependence of liquid crystallinity in the samples. Morphologies
were examined by a combination of wide- and small-angle
X-ray scattering (WAXS and SAXS), at room temperature and
at elevated temperatures, as well as by transmission electron
microscopy (TEM). For TEM studies, a piece of bulk sample
was embedded into epoxy, which was cured at 70 °C for 7 h.
Microtomy was then carried out with a Leica ultramicrotome
to produce the cross-sectioned thin film specimen (∼70 nm
thick). A J EOL 1200 TEM with an accelerating voltage of 120
keV was used for imaging.
Block Copolymers. The block copolymers were synthesized
using uniarm, diarm, and star polystyrene macroinitiators and
polymerizing them in a predetermined ratio with BBOS. Block
copolymerization was carried out in a solvent and not in bulk.
Solvents such as xylene, benzene, and o-dichlorobenzene were
examined, and it was found that the polymerization was best
controlled in o-dichlorobenzene, as it maintains good solubility
for the blocks throughout the polymerization and has a low
transfer constant. In a typical procedure copolymerization was
carried out in two-necked Pyrex ampules sealed under reduced
pressure. Different ratios of the polystyrene macroinitiator
were mixed with the liquid crystalline monomer in o-dichlo-
robenzene and polymerized at 125 °C for 2-3 days. For
example, 41.8 mg of star PS macromonomer (Mn 19 000, PDI
1.18) was mixed with 200 mg of BBOS (0.4 mmol) and 0.5 mL
of o-dichlorobenzene. This mixture was subjected to four cycles
of freeze-thaw and reacted for 48 h. The polymerization
mixture was precipitated in methanol and subsequently puri-
fied by column chromatography using dichloromethane as
eluant to remove residual monomer. The block copolymer was
further extracted with acetone to remove dead polystyrene
macroinitiator. This gave a starblock copolymer of Mn 40 980.
In the 13C spectra additional peaks for the ester carbonyl at
166 ppm between 10 and 50 ppm for the butyl substituent
along with the 1H NMR and GPC data confirm the block
nature (Figure 3).
Rh eologica l Mea su r em en ts. Samples for rheological mea-
surements were prepared into 13 mm diameter (1 mm thick-
ness) disks at 100 °C and 5 psi pressure in a hot press. These
disks were annealed at 110 °C for a day under vacuum. All
the samples were held at 180 °C for 10 min prior to the
Syn th esis. BBOS Monomer. The monomer 2,5-bis[(4-bu-
tylbenzoyl)oxy]styrene, BBOS (Figure 1), was synthesized by

Macromolecules, Vol. 36, No. 9, 2003
Rod-Coil Block Copolymers 3359
F igu r e 2. Synthetic scheme for polystyrene macroinitiators from (a) unimer or uniarmed, (b) diarmed, and (c) triarmed
alkoxyamine initiators.
F igu r e 4. Representative GPC plots for diblocks, triblocks,
and starblocks of PS with PBBOS along with the correspond-
ing macroinitiators.
F igu r e 3. (a) 13C NMR and (b) 1H NMR of PS-b-PBBOS
diblock copolymer.
important factor for narrow polydispersity. Hence,
BBOS monomer was purified by column chromatogra-
phy and recrystallized twice from petroleum ether. All
the polymerizations were carried out in o-dichloroben-
zene as solvent. Better control over the polymerization
was obtained by using 50 wt % solutions of the monomer
and macroinitiator for polymerization. Higher wt %
concentration of monomer led to faster polymerization
but larger polydispersity in the sample. BBOS block
copolymers were found to be soluble both at room
temperature and at 125 °C; it is also possible to directly
dissolve the PS macroinitiators in a BBOS melt at 95
°C (above the clearing temperature of BBOS) and
perform the polymerization without solvent at 125 °C;
however, the polydispersities obtained are higher (1.15-
1.45) compared to that of polymer formed by solution
polymerization.
experiments. A Rheometrics ARES rheometer was used for the
measurements. The experimental temperature was controlled
by using forced N2 gas convection with integral and propor-
tional plus derivative controller. The temperature control was
within (0.5 °C. A parallel-plate geometry (diameter ) 13 mm,
gap ) 1 mm) was used in studying the temperature depen-
dence of the apparent shear viscosity at a fixed rim shear rate.
Measurements were performed at temperatures spaced every
5 °C in the temperature range 100-200 °C. Dynamic temper-
ature sweep experiments were also performed using the same
setup.
Resu lts a n d Discu ssion
Syn th esis. As in all radical polymerizations, oxygen
can adversely affect polymer formation of MJ LC mono-
mers. The critical step in successful polymerization is
removal of oxygen via freeze-thaw cycle followed by N2
purge to remove any dissolved oxygen. Purity of the
liquid crystalline monomer was also found to be an
Figure 4 shows representative GPC plots for diblocks,
triblocks, and starblocks of BBOS with styrene. Com-
pared to the starting PS macroinitiator, a clean uni-

3360 Gopalan et al.
Macromolecules, Vol. 36, No. 9, 2003
Ta ble 1. Sp ecifica tion s of Va r iou s Block Cop olym er s of
P S a n d P BBOS Syn th esized in Ou r Stu d y
PS
Mn of block
copolymer
(g/mol)b
sample macroinitiator
mol %
IDa
Mn (g/mol)
PDI rod unitc LCd
BD1
BD2
BD3
BD4
BT1
BT2
BS1
BS2
6140
6140
10350
12370
3570
8870
19070
8620
19 890
13 840
18 500
26 155
30 400
18 160
40 980
16 530
1.35
1.30
1.22
1.15
1.34
1.26
1.14
1.21
34
22
15
20
63
19
21
17
x
x
Xe
x
x
x
x
x
a
BD stands for PS-b-PBBOS diblock, BT for PBBOS-b-PS-b-
PBBOS triblock, and BS for PS-b-PBBOS starblock copolymer.
b
Measured by GPC using polystyrene as standard and THF as
eluant. c Calculated based on GPC molecular weights. Phase
d
behavior as observed by polarized optical microscopy. e X: observed
only by WAXS.
F igu r e 5. WAXS plots of intensity vs 2θ for some of the
diblocks, triblocks, and starblocks of poly(styrene) with poly-
[(4-butylbenzyloxy)styrene] (PBBOS).
modal shift to higher molecular weight is seen in all
three block copolymers shown in the figure. The living
nature of these chains was confirmed by sequential
addition of a second monomer like styrene. The unimo-
dal distribution in GPC along with the presence of
aromatic peaks from both BBOS and PS block in the
1H and 13C NMR spectra (Figure 3) confirms the block
nature.
Mor p h ology. Liquid Crystalline Phase Behavior. As
reported earlier, BBOS monomer has a melting tem-
perature of 68 °C and a clearing temperature of 98 °C.
It exhibits a smectic LC phase between 68 and 98 °C.
The polymer as such is found to be nematic above its
Tg, i.e., 130 °C, with no detectable clearing temperature.5a
Most of the diblock, starblock, and ABA triblock
copolymers of BBOS and styrene listed in Table 1
showed birefringence when observed by POM. On heat-
ing the block copolymers above 130 °C, birefringence
develops with no observable clearing temperature.
Diblock copolymers for which no liquid crystallinity was
observable by POM had less than 15 mol % of the LC
block. This observation is in agreement with similar
reports by Zhou et al.7 where they found Schlieren
textures for block copolymers with 26 and 36 mol % of
LC units. The presence or absence of birefringence as
observed by POM is recorded in Table 1. Wide-angle
X-ray studies on solvent-cast block copolymer films give
more definitive information about the nematic liquid
crystalline phase. Figure 5 shows intensity, I, vs 2θ plots
for some of the block copolymers. In all cases, i.e.,
irrespective of chain architecture, we see a peak between
19 and 20 Å (2θ ) 4-5°) that corresponds to an
interchain packing distance that is close to that ob-
served for fibers drawn from BBOS homopolymers
(17.2-18 Å). A second broader peak between 4 and 5 Å
(2θ ) 19-20°) represents the intermolecular distance
between the mesogenic units.7
Microphase Separation. Small-angle X-ray scattering
(SAXS) was used to examine the microphase-separated
morphology of the diblock and starblock copolymers at
various temperatures. Results for two representative
samples are shown in Figure 6. Their molecular weights
and compositions are listed in Table 2. These polymers
were prepared using the same methods as those poly-
mers reported in Table 1. At room temperature (see the
top curves in Figure 6a,b), both the block and starblock
copolymers show only one peak in their SAXS spectra,
at 0.39 nm-1 for the diblock and at 0.29 nm-1 for the
starblock copolymer. This suggests that the block co-
F igu r e 6. Temperature-dependent SAXS data of (a) a diblock
and (b) a starblock PS-PBBOS copolymer. Temperatures of
SAXS measurements are labeled above the corresponding
curves. The top curves labeled with 23 °C (initial) are the room
temperature results of initial samples prepared through very
slow solvent casting procedure described in the Experimental
Section. The bottom curves labeled with 23 °C are the results
after the samples were heated to 150 °C and cooled to room
temperature subsequently.
Ta ble 2. Molecu la r Weigh t a n d P DI of Hom op olym er s,
Diblock , a n d Sta r block Cop olym er s Sa m p les Used in
Rh eologica l Stu d ies
sample ID
Mn (g/mol)
composition
homopolymer
homopolymer
PS-b-PBBOS 10300-8150
PS-b-PBBOS 9050-9910
PDI
PBBOS1
PBBOS2
diblock
9 000
44 700
18 500
18 960
1.40
1.35
1.20
1.22
starblock
polymers are phase-separated with a structural d spac-
ing of 16 and 21.5 nm for the diblock and starblock
copolymer, respectively. The absence of higher order
peaks indicates that, even after slow and careful solvent
casting, the composition profile across the interface is
rather smooth, and/or no significant long-range order
of the microphase separated structures could be estab-

Macromolecules, Vol. 36, No. 9, 2003
Rod-Coil Block Copolymers 3361
F igu r e 7. TEM of an unstained starblock copolymer BS1 (see
Table 1). Because of packing density contrast between PS and
PBBOS, PBBOS domain shows as relatively dark area in the
micrograph. The scale is indicated at the right corner of
micrograph.
lished in the block copolymers. On heating to 120 °C,
the peak intensity gradually decreases in both block
copolymers. At 150 °C, which is above the Tg of both
blocks, the SAXS peak disappears in both samples. This
suggests that, between 120 and 150 °C, the order-
disorder transition occurs in both block copolymer
samples. Fast cooling back to room temperature did not
regenerate the peak, whereas slow cooling over a time
period of 5-6 h did. This indicates that the microphase-
separated morphology is indeed thermodynamically
stable for the copolymers even though the ordering
kinetics is very slow. The marginal temperature window
between the TODT of the copolymers (below 150 °C) and
the Tg of the PBBOS block (130 °C) can be used to
rationalize the extremely slow ordering kinetics. While
the diblock copolymer shows a constant SAXS peak
position on heating the sample from 23 to 120 °C (Figure
6a), a marginal change in the d spacing from 215 to 235
Å was observed for the starblock copolymer (Figure 6b).
Although the exact molecular origin for this change in
d spacing is not known, its occurrence in only the
starblock copolymers is clearly related to differences in
the molecular architecture.
F igu r e 8. Temperature dependence of complex viscosity η*
(Eta*) and loss-tangent tanδ in (a) homopolymer PBBOS2 and
(b) PS-PBBOS2 diblock copolymer. The data were measured
at a frequency of 1rad/s. Depending on temperature, the rim
strain amplitude was varied from 1% to 10% in order to obtain
reasonable signal/noise ratios.
mers examined have about 50% mass fraction of MJ LCP
blocks and similar net LC block molecular weights as
PBBOS1. MJ LCPs are especially difficult systems to
work with as they have a clearing temperature above
their thermal degradation temperature, and hence
completely erasing their thermal history is not feasible.
Therefore, we have kept the thermal treatment identical
for all the samples so that comparison from one sample
to the other can give us reliable qualitative and some
quantitative information. Each sample was measured
twice, and the results were found to be reproducible.
After the rheological measurements were completed,
molecular weights of the samples were checked by GPC
and found to be the same as before the experiments,
which indicates that no significant degradation occurred
during the measurements. Small shear amplitudes were
used in the experiments. As a result, WAXS measure-
ments on the samples showed no net orientation after
the rheological measurements.
Temperature-Dependent Experiments. First, to assess
effects of the thermal transitions on their rheological
behavior, temperature-sweep experiments were per-
formed on all PBBOS homopolymers and PS-b-PBBOS
block copolymers. Over the temperature range from 95
to 180 °C, both homopolymers present similar features
in the temperature dependence of their rheological
behavior. The diblock and starblock copolymers, on the
other hand, show almost identical temperature-sweep
results. The temperature dependence of the complex
shear viscosity, η*, and loss tangent, tan δ, of a PBBOS
homopolymer sample and a PS-b-PBBOS block copoly-
mer are shown in parts a and b of Figure 8, respectively.
No abrupt change in the complex viscosity was observed
By using transmission electron microscopy (TEM), the
local structure of the block copolymers was examined
in real space. Most of the unstained samples showed
oval aggregates with no long-range order corroborating
the SAXS results. A typical TEM micrograph for an
unstained block copolymer (BS1 in Table 1) is shown
in Figure 7. The aggregates vary in size (5-60 nm)
depending on the block composition. However, a precise
size of these aggregates could not be calculated from
TEM due to poor contrast and the lack of a preferential
staining agent. In many regions the oval aggregates
seem to be breaking up from lamellar regions and are
most likely closer to the hockey puck two-dimensional
micelle model as proposed by William and Fredrick-
sons.12
Melt Rh eology. As highlighted in Table 2, the two
MJ LCP homopolymers examined were PBBOS1 and
PBBOS2 with molecular weights of 9000 and 44 700
g/mol, respectively. The diblock and starblock copoly-

3362 Gopalan et al.
Macromolecules, Vol. 36, No. 9, 2003
in any of the samples, indicating the absence of a
nematic-isotropic transition. This agrees with our POM
observations that none of the samples go through a
clearing temperature in the measured temperature
range. Around 130 °C, a slight change of slope for the
complex viscosity was observed in the hompolymers
(Figure 8a). Correspondingly, the tan δ data present a
broad peak at that temperature. For the temperature
dependence of the loss tangent (tan δ) of the block
copolymers (Figure 8b), there is a peak emerging at the
low-temperature end, i.e., 90 °C besides the similar
feature around 130 °C. These peaks at 90 and 130 °C
correspond to the Tg’s of PS and PBBOS, respectively,
as also determined by DSC measurements. The pres-
ence of two distinct Tg’s in the block copolymers is in
agreement with our SAXS results that show microphase
separation in these block copolymers at temperatures
below 150 °C. As discussed in the previous section, the
temperature-dependent SAXS measurements indicate
that, in both the diblock and star block copolymers, the
order-disorder transition of the microphase-separated
structure should occur between 120 and 150 °C. Com-
mon rheological features related to such a transition,
namely a significant decrease of moduli and increase
of tan δ, do not unambiguously show up in our mea-
surements. The reason might be that the glass transi-
tion of PBBOS blocks is in the vicinity of the order-
disorder transition. A glass transition usually bears
similar and even stronger rheological features than an
order-disorder phase transition. Therefore, the two
thermal transitions might not be distinguished in the
rheological measurements of our block copolymers.
F igu r e 9. Master curve of storage (G′) and loss modulus (G′′)
for (a) PBBOS1 and (b) PBBOS2 with reference temperature
(Tref) at 130 °C.
Frequency-Dependent Measurements. To investigate
the dynamic fingerprint of the different molecular
architectures, the frequency dependence of the rheo-
logical properties was measured at temperatures from
100 to 180 °C. For all the samples, master curves could
be composed using time-temperature superposition
(TTS) to retrieve the rheological responses over a wide
frequency range: from 10-3 to 105 rad/s. The tempera-
ture dependence of the shift factors (aT) could be fitted
with the well-known WLF relation14 (data not shown).
creases from 9000 g/mol in PBBOS1 to 44 700 g/mol in
PBBOS2, the viscosity drops significantly. Berghausen
and co-workers have reported a similar phenomenon.13
Considering the rigid backbone conformations in MJ L-
CP, an increase of molecular weight results in a higher
aspect ratio of the rodlike macromolecule, which is
expected to enhance the nematic liquid crystallinity of
the materials. Therefore, a greater extent of the MJ LCP
backbone alignment in the higher molecular weight
sample will tend to lower the viscosity of the material
under shear. Second, both block copolymers present
lower complex viscosities than the homopolymer PB-
BOS1 even though the net molecular weight of the LC
block is comparable in all three samples (∼9000 g/mol).
The reduction of viscosity is most likely due to the
presence of the polystyrene block. PS is in the rubbery
state above its Tg (90 °C) and hence leads to a net
plasticization of the polymer. It is interesting to note
that the diblock copolymer has a marginally higher
complex shear viscosity than its starblock counterpart.
Since both block copolymers have similar net molecular
weights of the MJ LCP block (∼9000 g/mol), the differ-
ence is only that the LC part is split into three arms in
the PS-PBBOS star block copolymer. Each arm has
thus only one-third MJ LCP molecular weight of that in
the diblock copolymer. Following the arguments used
in comparison of the complex viscosity data between
PBBOS1 and PBBOS2, namely greater extent of MJ L-
CP backbone alignment for higher molecular weight
MJ LCPs, contrary to what is observed experimentally,
the diblock copolymer is expected to show lower complex
shear viscosity than its starblock counterpart.
Master curves of storage (G′) and loss modulus (G′′)
for PBBOS1 and PBBOS2 are shown in parts a and b
of Figure 9, respectively. No G′/G′′ crossover and/or
plateau region were observed at intermediate frequen-
cies. Thus, the entanglement molecular weight was not
attained in any of the samples even though the molec-
ular weight of PBBOS2 is 44 700 g/mol. This reflects
the rigid chain conformations of the present MJ LCP’s.
In both the homopolymer and the block copolymer
samples, nonterminal behavior was observed at lower
shear frequencies, consistent with the liquid crystalline
nature of the present materials. In contrast to the
previous study reported by Berghausen and co-work-
ers,13 no discontinuity was observed in the master
curves. This difference is due to the fact that their
system goes through a sharp nematic-to-isotropic phase
transition which is absent in our system.
The effect of molecular architecture on chain dynam-
ics becomes evident on superimposing the complex shear
viscosity, η*, and loss tangent, tan δ, data for homopoly-
mers (PBBOS1 and PBBOS2), diblock and starblock
copolymers as shown in parts a and b of Figure 10,
respectively. There are two obvious trends for the
complex viscosity of the samples (Figure 10a). First,
when the molecular weight of the homopolymer in-
To thus examine the dynamics in the materials in
more detail, the master curves of the loss tangent (tan
δ) data are shown for all samples in Figure 10b. All the

Macromolecules, Vol. 36, No. 9, 2003
Rod-Coil Block Copolymers 3363
Ta ble 3. Activa tion En er gies of P BBOS Hom op olym er s
a n d Th eir Diblock a n d Sta r block Cop olym er s w ith
P olystyr en ea
sample
activation energy (kJ /mol)
PBBOS1
PBBOS2
diblock
200.86 ( 16.0
173.75 ( 8.5
152.08 ( 9.8
137.80 ( 6.67
starblock
a
The values are calculated from temperature dependence of the
time-temperature-superposition (TTS) shift factors (aT).
F igu r e 11. Cartoon depicting the molecular architecture of
PBBOS1, diblock, and starblock copolymer with trends in their
complex shear viscosity and relaxation times.
In general, the terminal relaxation times of a branched
macromolecule are much shorter than a linear one of
same molecular weight. The higher the number of
branches, the larger the difference will be.14,15a This
effect has been well studied in poly(styrene) star poly-
mers where, for low molecular weight melts, the zero
shear viscosity is found to be lower than in linear
polymers of the same molecular weights.15b,c Similar
studies are reported for star-shaped poly(isoprene).15d
In our system the presence of three LC chain ends per
molecule compared to one LC and one PS chain end in
the diblock copolymer brings about a reduction in the
relaxation time in the starblock.
Finally, flow activation energies were calculated from
a plot of ln aT vs 1/T for all the samples in their high-
temperature nematic range. It allows us to compare the
MJ LCPs with other similar LCPs in a more quantitative
manner. The results for all the samples are shown in
Table 3. The values are higher than what has been
reported by Berghausen and co-workers for side group
polymers with laterally attached mesogens14 and by
Zentel and co-workers for end-on attached side group
LCP’s.17 The differences are probably due to the fact
that there are no flexible linkers in MJ LCPs to decouple
the motion of the side chains from that of the polymer
backbone. This leads to the bulky mesogenic rods
wrapping themselves around the polymer backbone, the
so-called “mesogen jacketed effect”. The values of the
activation energies follow the same trend as exhibited
by the complex shear viscosity η* and the relaxation
time derived from the tan δ data, namely PBBOS1 >
diblock > starblock > PBBOS2; see also Figure 11.
These results suggest that there are two ways to make
MJ LCPs more attractive from a processing point of
view: either by enhancing nematic liquid crystallinity
through higher molecular weights or by incorporating
branching points, e.g., through a starlike architecture,
in the molecular design.
F igu r e 10. Master curves of (a) complex viscosity η* and (b)
loss-tangent tan δ for PBBOS1 (4), PBBOS2 (3), diblock (0),
and starblock (]) copolymer with reference temperature at 130
°C.
samples present a broad tan δ maximum at the lower
frequency end. A maximum of loss tangent is usually
related to dynamic transitions from one microscopic
relaxation mode to another. In the case of our MJ LCP
homopolymer and block copolymers, on the basis of our
previous discussion, this tan δ maximum should cor-
respond to the dynamic transition from relaxations of
individual macromolecules to that of nematic LC su-
pramolecular structures. Therefore, it may serve as a
time scale mark for us to compare the effect of molecular
architecture on microscopic relaxation processes in the
MJ LCP materials. Compared to the lower molecular
weight homopolymer PBBOS1, the maximum of tan δ
in all other three samples significantly shifts to higher
frequencies. This indicates that the corresponding mo-
lecular relaxation times in these samples are signifi-
cantly shorter than those in the homopolymer PBBOS1.
Furthermore, the molecular relaxation times follow the
same order as the magnitudes of the complex shear
viscosities of these materials, namely PBBOS1 > diblock
> starblock > PBBOS2. This observation suggests that
there may be a “correlation” between the reduction of
the complex shear viscosity and the decrease of the
molecular relaxation times.
In our previous discussion, we have suggested that,
compared with PBBOS1, the enhanced liquid crystal-
linity in PBBOS2 and PS plasticization in the block
copolymers give rise to the reduction of the complex
shear viscosity in these materials. Comparing diblock
and starblock copolymers, differences of the molecular
architecture can explain relatively shorter molecular
relaxation times in the latter. A low molecular weight
star polymer can be regarded as a branched molecule.

3364 Gopalan et al.
Macromolecules, Vol. 36, No. 9, 2003
Con clu sion s
support by the Cornell Center for Materials Research
(CCMR), a Materials Research Science and Engineering
Center of the National Science Foundation (DMR-
0079992), is appreciated. Financial support of this work
by the Xerox Research Center of Canada is appreciated.
Our studies show that well-defined macromolecular
architectures such as diblock, triblock, and starblock
copolymers can been synthesized by SFRP using a
monofunctional, difunctional, or trifunctional polysty-
rene macroinitiator, respectively, end-capped with a
TEMPO-based alkoxyamine initiator. Morphological
studies on these blocks by POM and high-temperature
WAXS show a nematic LC phase for most of the block
copolymers which are stable up to the decomposition
temperature. SAXS and TEM show the formation of
microphase-separated nanostructures for the block co-
polymers. No significant long-range order of these
structures could be detected.
The effect of varying the molecular architecture in
going from homopolymers to diblock and starblock
copolymers is reflected in physical properties such as
complex viscosity η* and relaxation times. No discon-
tinuity was observed in the master curves due to the
lack of a nematic to isotropic phase transition in these
systems. No entanglement plateau was observed for the
compositions of the block copolymers examined. The
flow activation energies were found to be higher than
in systems with flexible linkers between the backbone
and the mesogen. There is a reduction in complex shear
viscosity in going from the homopolymer BBOS to
diblock to starblock copolymers attributed to the pres-
ence of the PS block leading to a plasticizing effect and
to topological factors arising from the branched molec-
ular architecture. Molecular relaxation times derived
from loss tangent δ data follow the same trend as the
shear viscosity. The results provide a useful synthetic
strategy to tune the physical properties and hence make
MJ LCPs and their block copolymers more attractive
from a processing point of view.
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Ack n ow led gm en t. The authors thank Dr. Craig
Hawker of the IBM Almaden Research Center for
providing us with the macroinitiators used in this study
and Anurag J ain for his help and suggestions in
carrying out the rheological measurements. Funding by
the Air Force Office of Scientific Research through its
MURI program is gratefully acknowledged. Partial
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