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Dielectric relaxation spectra of two side chain liquid crystalline homopolymers and their copolymers
Eur. Pdym. J. Vol. 32, No. 12, pp. 1361-1369, 1996 Copyright 0 1996 ElsevierScienceLtd
Pergamon PII: s0014-3057(%)00108-5
Printed in Great Britain. All rights reserved 0014-3057/96 $15.00 + 0.00
DIELECTRIC RELAXATION SPECTRA OF TWO SIDE CHAIN LIQUID CRYSTALLINE HOMOPOLYMERS AND THEIR COPOLYMERS G. M. DAY,’ W. R. JACKSON’ and G. P. SIMON** ‘Department of Chemistry and ‘Department of Materials Engineering, Monash University, Wellington Road, Clayton, Victoria, 3168, Australia (Received 13 September 199.5; accepted in final form 13 November 1995) Abstract-Dielectric relaxation spectra have been measured for a complete series of siloxane-based side chain liquid crystalline copolymers (with two mesogenic side groups of similar structure but having flexible spacer groups of different length) and the effect of the copolymer chemical structure on thermal properties and dielectric motion is presented. The qualitative behaviour observed in the spectra of two of the polymers with regard to their change in peak height of the loss relaxation with temperature is discussed in terms of influences on relaxation strength and alignment with temperature. The speed of motion of the mesogens is compared between different members of the copolymer series by normalising the experimental temperature by either the glass transition temperature of the material or its clearing point (liquid crystalline to isotropic temperature). The effect of an increased content of the side group with a very long flexible spacer is to increase the frequency of motion due to greater mesogen mobility if the data is compared at temperatures normalised by the clearing point. However, if the temperature is normalised by the glass transition, the peak frequencies of motion are much closer (although the material with a longer spacer group still moves slightly faster), indicating the important role the motion of the polymer backbone plays in determining the mesogenic mobility. The activation energy of motion is found to be less for the copolymer with the longer spacer length and consequently in the copolymers with greater content of the long spacer unit side chain. Copyright 0 1996 Elsevier Science Ltd
INTRODUCTION
Dielectric Relaxation Spectroscopy (DRS) is a useful technique for the study of the rotational motions of macromolecules having a permanent dipole moment [ 11. Due to the polar nature of many of the mesogenic pendant groups used, side chain liquid crystalline polymers (SCLCPs) are excellent materials for study using DRS [2-51. Structure property relationships between chemical structure and dielectric relaxation have been determined for a range of side chain liquid crystalline homopolymers. Many have concentrated on liquid crystalline homopolymers with acrylate or methacrylate backbones with different combinations of flexible spacers or mesogenic side groups [2, S-81. Several studies on siloxane-based liquid crystalline homopolymers have also been undertaken and these materials have been well characterised in terms of their fundamental molecular motions [3, $9-101. Research on the dielectric properties of liquid crystalline copolymers has been limited [ 1l-151 with the effect of copolymer composition on thermal properties mainly reported. Some work concerning the dielectric properties of siloxane-based side chain liquid crystalline copolymers having two very different mesogenic pendant groups has concentrated on the effect of changing the ratios of the two pendant groups [ 12, 131 on the speed and activation energy of *To whom all correspondence
should be sent.
motion. In these systems, one of the homopolymers has also tended not to show a liquid crystalline phase 1131. The work reported here uses DRS to compare the mobility of a complete series of siloxane-based side chain liquid crystalline copolymers, with the two mesogenic side groups having a similar structure and flexible spacers based on the vinylacetic acid spacer (structures are shown in Fig. 1) developed in our laboratory [16, 171. The structure of the mesogenic units attached to the siloxane backbone (RI and R2) differs only in the length of the flexible spacer and both the homopolymers demonstrate liquid crystalline phases. Comparisons of dielectric mobility between the unaligned samples of the two respective homopolymers la and le, as well as that of the copolymers lb, lc and Id, will be presented. EXPERIMENTAL
Marerials
The polymers synthesised for this study used methods described elsewhere [ 16, 171.The structures of the polymers are shown in Fig. 1 and their thermal characteristics from differential scanning calorimetry (DSC) presented in Table 1 (T# is the glass transition temperature, while T, represents the isotropic temperature at which the material passes from the liquid crystalline state to the isotropic liquid state, also known as the clearing point). All samples were run on a Perkin Elmer DSC7 at a scan rate of lOC/min and the temperatures calibrated with an indium standard. The glass transition temperatures could not be observed for samples
1361
G. M. Day et ul
I362
Copolymer
R, : -
series
1
W ,l,CO,GH
R2 : -W
2,30 -&LO
z-@N
,),C02(CH &, 0
Fig. 1. Structures
of the polymers
lb and Ic, although it is reasonable to assume they would be close to that of the homopolymer and sample le. All materials showed conical fan textures, typical of a smectic A structure in hot stage transmission optical microscope experiments with crossed polars above r,.
I
of motions of the mesogenic pendant group since the siloxane backbone is itself relatively non-polar [3]. Previous studies on amorphous [2, 191 and semicrystalline [9] SCLCP materials have shown that it is necessary to exceed the glass transition temperature or melting point (the latter only relevant to semicrystalline materials) in order to liberate the strong dielectric relaxations of the mesogenic groups in the l&lo5 Hz frequency range. The dielectric spectra of the polymers in this study all show loss peak maxima temperatures well above the TB values observed by thermal analysis (or expected T, values in the case of copolymers 1b and lc).
Dielectric measurement
Variable frequency dielectric spectroscopy was performed using a Gen Rad 1689 Digibridge and techniques similar to those previously described [4, 181. The data presented involves G/w @F), where G is the sample conductance and o is the angular frequency of measurement (w = 2nf, where ,fis between 10 and IO5Hz). G/o is a dielectric parameter related to the imaginary component of the dielectric loss permittivity, t” = (G/w)/C,; C,, is the empty capacitance of the sample and hence determination of t” requires knowledge of the sample thickness. Samples of - 1 cm diameter were prepared by squeezing the material between the two plates at a temperature above the clearing point of the sample. This technique precludes a direct measurement of the thickness of the sample, preventing the calculation of the dielectric constant (6’) and the dielectric loss permittivity (6”). Since only peak position and shape is of interest in this work, G/o vs log f is presented, as has been done often elsewhere [4,9, 18-201in DRS of SCLCP materials.
Shape and height
The dielectric
qf the dielectric
loss spectrum
loss cunles
for the homopolymer
la (100 mol % RI) is shown in Fig. 2. It can be seen that the peak does not become prominent until the temperature reaches 4O”C, which is close to the T, for la and the loss peak decreases in height with increasing temperature (up to T = 65). The lack of a relaxation peak in the liquid crystalline state seems to be due to the low frequency of motion at these temperatures and its disappearance in the low frequency conductivity tail. Other workers have shown that peak height tends to remain relatively constant for loss peaks measured above r [21]. The copolymer 1b (Fig. 3) containing 75 mol % RI and 25 mol % R? behaved very differently to the homopolymer la in that the loss peak height showed a significant increase with increasing temperature, in particular above the clearing point. Copolymers lc and Id (Figs 4 and 5, respectively) and homopolymer le (Fig. 6) show relaxation spectra only within the frequency window of measurement in the liquid crystalline state and
RESULTS AND DISCUSSION
The unaligned dielectric loss spectra (G/w vs logf) for all homo- and copolymers la-le are presented in Figs 2-6. All the spectra showed an increasing loss tail at lower frequencies. This so-called “conductivity tail” may be due to some residual impurities in the samples which lead to conductivity. The dielectric loss maxima for most spectra first become visible at temperatures of between 40 and 50°C (35°C for Id). The loss spectra shown is attributed to a combination
Table 1. Thermal data for the ohase transitions of the nolvmers 1 Ratio Compd la lb 1C Id le
7-m
APwt
( C)
Pa
AH, (J/g)
20 22 26 34 18
3.4 4.6 6.3 5.1 6.5
RI
RZ
100 75 50 25 0
0 2s 50
-7 *
15
12
78
100
-4
106
_
*Glass transition not observed. tBaseline peak width of clearing transition from DSC
40 52 71
&IQ 10.9 13.4 19.4 16.2 17.2
Dielectric relaxation of liquid crystalline polymers
1363
___e___ Tt35 ___Q___ T=40 +
T=45
+
T=50
+
T=55
-o-T=60 +
T=65
e
T=70
___ +--- Tt76
1
2
3
4
log f (Hz)
5
Fig. 2. G/w (PF) vs loglof for polymer la (7’ in “C).
thus the effect of heating the sample through the clearing point cannot be determined and these relaxation peaks are constant in height. The decrease in loss peak height of la (Fig. 2) could be explained if the original sample of la was either partially or fully homeotropically aligned and the decrease in height could be ascribed to decreasing homeotropic orientation. Homeotropic alignment occurs when the mesogens are aligned perpendicular to the electrodes and thus the measuring field of the dielectric bridge largely interacts with the longitudi-
nal dipoles of these mesogenic units. Such homeotropic alignment could be induced electrically by application of electric fields of low frequency for SCLCPs with mesogens of positive dielectric anisotropy, as are those in this study [19] or by induced surface alignment [22]. The fact that such alignment, should it exist, persists so deeply into the isotropic region (the peak is still decreasing in height at 76”C, some 36°C above the clearing point) could be explained in part by the broadness of the biphasic region (coexistence of liquid crystalline and
w6-l(PF)
2
3
4
Fig. 3. G/w @F) vs loglOf for polymer 1b (T in “C).
5
+
T=40
+
T=45
--8-
T=50
+
T=55
*
T=60
+
T=65
1364
G. M. Day PI al
+T=40 +T=45 +
T=50
---I---
1dS.i
*
T=60
+
T=65
___Q___ T=70 ___+YJ___ T=76 ___+___ T=61
1
I
I
I
2
3
4
5
logf (Hz)
Fig. 4. GW @F) vs loglii / for polymer Ic (T in C).
amorphous phase as the material is heated into the isotropic melt). However. the entire biphasic transition region is some 20°C for sample la. as shown in Table 1. An alternative explanation for the continuing decrease in the height of the loss peaks (which is related to the relaxation strength and thus the effective number of dipoles per unit volume) is that it is due to increased thermal randomisation of the dipoles, a similar explanation as is proposed for the decrease in the c( relaxation in isotropic
thermoplastic relaxation spectra with temperature [231. The increase in the height of the loss peaks with increasing temperature for copolymer lb (Fig. 3)
could be explained if the original unaligned sample was partially or wholly planarly aligned (mesogenic direction parallel to the electrode surface) and that the increase in height of the loss peak with increasing temperature could be due to decreasing planar orientation. Planar orientation could be expected in
+
T=30
+
T=35
-x--T=40 +
T=45
+
T=50
+
T=55
._.Q_._T=60 .__a___T=65 .__+___T=70
0
i
I
I
I
1
2
3
4
Fig. 5. G:o @F) vs log10f for polymer Id (T in C).
5
logf (Hz)
Dielectric relaxation of liquid crystalline polymers
1365
(PF)
WI
5 \
\ +T=!Xl
4
--t-T=66 *T=65 -o-T=71
3
___+___ T = 76
2
1
I
i 1
I 2
I
I 4
3 Fig. 6. G/o (PF) vs log,,f
these systems due to template alignment in the microscratches in the electrode surfaces and the shearing, squeezing flow that the mesogens experience when the material is pressed between electrodes during sample preparation, aligning the liquid crystalline units parallel to the electrodes, in the direction of flow. Support for the idea that this material is at least planarly aligned also comes from the very broad loss peaks in Fig. 3 with widths at half height of greater than logf = 3, a common feature of planarly aligned materials [24] where the dipole being relaxed is mainly that which is perpendicular to the mesogenic group. It is difficult to see why this copolymer would differ in terms of alignment due to surface forces or squeezing flow compared to, say, the homopolymer la, although it is noted that the rheological behaviour of these materials is somewhat complicated [25]. Some contribution to increases in peak height and relaxation strength of SCLCPs, such as that seen for copolymer 1b, could be due to disruption of antiparallel packing that often occurs in liquid crystalline systems with strong terminal dipoles (such as the cyano groups used in this work). This packing causes cancellation of the dipoles (and thus reduced dipole strength). On heating, these anti-parallel associations are disturbed and may cause an increase in dielectric strength and peak height due to an increase in the number of available dipoles [26]. As explained above, the persistence in such an increase beyond the nominal clearing point may be due to the width of the biphasic region. The shapes of dielectric curves for SCLCP materials have been well-described in the literature [2,3] and have been attributed to different motions of the polar mesogenic pendant group. The two basic relaxations which have been observed are labelled 6 and tl. The 6 relaxation is related to motion of the longitudinal dipole due to the terminal cyano group
for
5
logf (Hz)
polymer le (T in “C)
and the mesogens short axis, while a relaxation is largely due to motions of the transverse dipole (resulting mainly from the motion of the ester group) around the long axis of the mesogen. The a relaxation in this context should not be thought to necessarily be the same as the a relaxation observed in isotropic thermoplastics above the DSC glass transition temperature. It should be noted that there is some controversy in the assignation of the a relaxation in SCLCPs dielectric spectra. In methacrylate systems it is believed to be largely related to main-chain motion, being ascribed to the relaxation of the carbonyl group in the side chain closest to the polymer backbone [2]. However, in polysiloxane materials it has been more often associated with the motion of the dipole transverse to the mesogen [2,3]. The a relaxation occurs at a higher frequency (faster motion) than the 6 relaxation because it does not require a large scale molecular rearrangement, being due mainly to spinning of the mesogen around its long axis. The correctness of the a and 6 assignment can be demonstrated if the samples are aligned homeotropically or planarly using an electric field [2,22], which was not successfully achieved in this study. For many side chain liquid crystalline polysiloxanes with a conventional alkyl chain based flexible spacer, it has been shown that dielectric loss curves of unaligned samples consist of a superposition of 6 and a peaks, although the peak maximum was effectively that of the much stronger, low frequency S relaxation (i.e. the motion of the longitudinal dipole moment). The a peak was usually determined by deconvolution, for example, by assuming a Fuoss-Kirkwood empirical relaxation function of the broad, overlapping 6-u peaks [3,24,27]. Recent use of a vinylacetic acid based flexible spacer in SCLCP materials has seen the dramatic reduction of clearing temperatures (cf. to SCLCP materials with an alkyl chain based spacer)
G. M. Day et al
In f,,, (Hz) 15 ela -e--lb --x.-lc .__+___ ld *le
2.9
3
3.1
Fig. 7. In 1,. vs I T ior the polymers
3.2
3.3
l/T
1.
[28]and the unaligned dielectric loss spectra of these and copolymers of Series 1 at the same absolute materials has shown distinct frequency separation temperature, as this does not take into account between u and 6 relaxation peaks [lo]. The homo- and the relationship of the measuring temperature to copolymers in this work (la-le) (which also involve key events in the polymer spectra (glass transition a vinylacetic acid based spacer in both mesogenic temperature, clearing point) [2,3,29, 301. Rather, pendant groups) do not exhibit distinct 6 and ry such comparisons are best made by comparing relaxation peaks (see Figs 2-6). A possible reason for dielectric loss spectra positions at the same reduced this could be that the high frequency c( peaks may temperature (T,,J. Potentially, Tredcan thus be scaled exist outside the frequency window at the experimental by the clearing point (= TJZ , where Tcxp is the experimental temperature and all temperatures are temperatures used. The loss spectra of the homopolyin K) or as a function of glass transition (= TCxp/Ta). mer le recorded at 40°C and 45°C show what appears to be the beginning of a peak due to c( relaxation at It has been stated that for several liquid crystalline high frequencies (Fig. 6). It should be noted that since polysiloxanes at temperatures more than 20 K below T,, it is better to scale data by T, since at these only one loss curve is observed for all the copolymers lbld, it appears that the copolymerised species temperatures of measurement the glass forming dynamics of the polymer backbone dominate the are statistically distributed along the polymer chain speed of motion [30]. Such a temperature scaling leading to something of an averaging of the dielectric relaxation of the homopolymer processes. This is as approach is similar to cooperativity or fragility plots [31] that have been employed in understanding seen in dielectric spectra of other SCLCP copolymers intra- and inter-chain coupling between conventional [13], although in another reported system, two distinct peaks were observed [12] and attributed to thermoplastic chains above the glass transition temperature [32]. “blocky” SCLCP copolymers. Figure 8 contains the same data as Fig. 7 but Position of the dielectric loss curves plotted in terms of peak position (In fmax) as a function The positions of dielectric peak position as a of l/T,,, , where T,& = T,,,/T,. It can be seen that this function of temperature (In fm vs l/r) are shown in separates the peaks to a greater degree. That is, rather Fig. 7. It can be seen that the spectra are relatively than lying close to each other as they are in Fig. 8, Arrhenius-like (linear). they are quite strongly separated, with the homopolySuch behaviour is usually observed for the 6 mer le with 100 mol % Rz (the longest spacer group) relaxation [2, 3,241, whilst the tl relaxation tends to lying at the highest frequency (fastest motion) for a show spectra that are more curved (WLF effect). given reduced temperature. Conversely, polymer la Copolymers lb and lc showed slight curvature, as with the shortest spacer group is moving the slowest would be expected, particularly in the case of 1b at a given reduced temperature, with the copolymers which, as discussed earlier, gave indications of planar lying in between in monotonic compositional order. alignment (and thus may be largely an a relaxation Inspection of Figs 7 and 8 at constant temperature process) by its increase in relaxation strength with and reduced temperature, respectively, allows comtemperature. parisons to be made as to the efficacy of scaling the It is not totally valid to compare peak positions data (and hence yields information about the nature (speed of motion) between the members of the homoof the motion). Even though glass transitions could
r
Dielectric relaxation of liquid crystalline polymers
In f,(Hz) 15
0
10
0 0
X
0
ox
cl
X4
0
X
0
0 0
0
0
la
o
lb
x
lc
l
Id
0
le
0 4
x
0
0
1 0
4
x
0
5
0
4
1367
0
4
x
0
0
4
0
4
X
4
0
0 0.8
1
0.9
1.1
1.2
1.3 lflrad
Fig. 8. In& vs l/T& for the polymers 1 (7’,, = T,,&T).
not be observed for lb and lc and thus normalised data with respect to T, could not be determined, if it is assumed that the glass transitions of 1b and lc were approximately in between those of the other polymers as shown in Table 1 (and thus are all of similar magnitude), comparison of data at constant Tti = TLxp/Tgis effectively the same as comparing them at constant experimental temperature. Thus the data of Fig. 7 can be thought to be qualitatively similar to being plotted against constant T, = TJT,. It can be seen that data of Fig. 7 (nominally TRd= Tcip/Tg) overlies each other to a greater degree than data of Fig. 8 (at constant Tti = T,JT). That is, it appears as though normalising the peak frequency positions with respect to the glass transition temperature is more effective than scaling using the clearing point. Such comparisons of scaling can best be seen by tabulation of the data at constant reduced temperatures (vertical cross-sections in Figs 7 and 8). Comparison of In fm for the polymers 1 at a Tmd= T47 of _ 1 is shown in the top part of Table 2 and for T,, = Tcxp/Tgof 1.27 is shown in the lower part of Table 2. As mentioned above, not all
Table 2. Comparison of In& values for polymers Polymer la lb IC
Id le
(T>,)
1.00 1.01 l.Oll 0.98 0.92
1at a constant
Td
IiT.4
Te, (“C)
In fm (Hz)
1.00 0.99 1.I0 1.02 I .09
40 55 70 70 76
4.15 1.83 9.21 9.21 10.13
Td
Polymer la
Id le
( T.rp/Ts) 1.27 I .27 1.27
l/T&
0.787 0.787 0.787
T., (‘C)
In .h (Hz)
65 70 70
8.75 9.21 9.21
of the polymers 1 had an observable T, and thus Table 2 is incomplete. Despite this, at a constant Td = T,,JT, = 1, the data in Table 2 shows that an increase in the amount of the mesogenic pendant group with the very long spacer (RJ leads to an increase in the value of lnf, , which indicates a faster S motion. Comparison of the two homopolymers la and le clearly shows that there is more rapid 6 motion by the polymer le with the longer spacer (Rz), as may be expected with increased mobility due to greater decoupling of the mesogen from the restricting, main-chain motion. The mesogen is able to move more freely and thus faster. The copolymers show intermediate behaviour with the different compositions of Id and le having little effect on speed of motion. This behaviour can also be seen by inspection of Fig. 8. However, when the data of Tmd= T,.,/T, is compared at a constant value for samples la, Id and le, it can be seen in Table 2 that the peak position at T, = 1.27 is approximately independent of composition, corresponding to the approximate overlaying of the curves in Fig. 7. It should be noted, however, that even when the data is scaled by the glass transition temperature, there is a slight increase in speed for copolymers and homopolymers Id and le which contain the longer, flexible spacer unit.
Table 3. Activation energies (G,) for polymers 1 Polymer
.% (Id mol-‘)
la (>E) lb (tTJ lb (z-T,)
166 168
Ic (< Ti)
175 141 137
Id (CT,) le (tK)
146
I368
G. M. Day YI al.
The fact that the data effecttvely scaled by the glass transition temperature are similar between the homopolymers and at least one copolymer is indicative of the role that the polymer backbone plays in the motion of the mesogens. Despite the side chain being sufficiently decoupled from the main-chain motion to allow the liquid crystalline phase to be attained, the motion of the mesogens remain tied to that of the main chain. This has been illustrated in other ways, for example by comparison of dynamic mechanical data (which measures backbone mobility) and dielectric relaxation (which examines the motion of the polar mesogen), demonstrates their interconnectedness [9]. Coupled backbone and mesogen motion from alignment measurements and variable frequency dielectric studies have also been demonstrated by Birenheide et ul. [33] who showed that the activation energy of the temperature dependence of polymer viscosity was similar to that of the S relaxation (motion around the mesogenic groups short axis) and hence were related. Given, however, that much of the data presented here were measured at temperatures close to or above the isotropic temperature, it is reasonable to assume that normalising data with respect to the isotropic temperature is also justifiable and such a scaling exaggerates the effect of increasing the spacer length in speeding up the mesogenic motion. Thus, provided that the data are at temperatures close to or above the clearing point, Tred= T,,,!T, may also be a valid way of comparing relaxational dielectric data between copolymers and not solely for comparison of static dielectric behaviour, as indicated by mean field theory [30]. If an Arrhenius relationship is assumed for the In _&, vs l/T data, then activation energies of motion (&J in the liquid crystalline and isotropic states can be calculated from Fig. 7. Calculated activation energies for the polymers 1 are shown in Table 3. where it is also indicated whether the data were obtained above or below the clearing point. It is difficult to make comparisons from the data of Fig. 7 across the whole composition range. as temperature-frequency data could not be obtained for all materials totally in the liquid crystalline or in the isotropic state. However, a few comments from the activation energies determined can be made. It can be seen from sample 1b that the activation energy of the mesogen in the liquid crystalline state is greater than that of the isotropic state, as is usually observed, and is ascribed to hindrance of motion due to the effect of the liquid crystalline potential field [34]. Comparison of activation energies of samples 1b. Ic. Id and le (all in the liquid crystalline states) indicates a lower activation energy for the copolymers with greater content of the longer, flexible spacer unit. Clearly the energy barrier to overcome motion is less in these materials. This is similarly true of comparison of la and sample lb in the isotropic state when both are above their clearing points. Thus, even in the absence of a liquid crystalline potential, copolymers with greater content of long-spacer mesogens are able to move with less energetic hindrance. This is in agreement with Simon and Coles [29] who showed that an increased length of spacer unit in a homologous series of polysiloxane SCLCP
systems
resulted
in an overall decrease
in activation
energy.
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DIELECTRIC RELAXATION SPECTRA OF TWO SIDE CHAIN LIQUID CRYSTALLINE HOMOPOLYMERS AND THEIR COPOLYMERS G. M. DAY,’ W. R. JACKSON’ and G. P. SIMON** ‘Department of Chemistry and ‘Department of Materials Engineering, Monash University, Wellington Road, Clayton, Victoria, 3168, Australia (Received 13 September 199.5; accepted in final form 13 November 1995) Abstract-Dielectric relaxation spectra have been measured for a complete series of siloxane-based side chain liquid crystalline copolymers (with two mesogenic side groups of similar structure but having flexible spacer groups of different length) and the effect of the copolymer chemical structure on thermal properties and dielectric motion is presented. The qualitative behaviour observed in the spectra of two of the polymers with regard to their change in peak height of the loss relaxation with temperature is discussed in terms of influences on relaxation strength and alignment with temperature. The speed of motion of the mesogens is compared between different members of the copolymer series by normalising the experimental temperature by either the glass transition temperature of the material or its clearing point (liquid crystalline to isotropic temperature). The effect of an increased content of the side group with a very long flexible spacer is to increase the frequency of motion due to greater mesogen mobility if the data is compared at temperatures normalised by the clearing point. However, if the temperature is normalised by the glass transition, the peak frequencies of motion are much closer (although the material with a longer spacer group still moves slightly faster), indicating the important role the motion of the polymer backbone plays in determining the mesogenic mobility. The activation energy of motion is found to be less for the copolymer with the longer spacer length and consequently in the copolymers with greater content of the long spacer unit side chain. Copyright 0 1996 Elsevier Science Ltd
INTRODUCTION
Dielectric Relaxation Spectroscopy (DRS) is a useful technique for the study of the rotational motions of macromolecules having a permanent dipole moment [ 11. Due to the polar nature of many of the mesogenic pendant groups used, side chain liquid crystalline polymers (SCLCPs) are excellent materials for study using DRS [2-51. Structure property relationships between chemical structure and dielectric relaxation have been determined for a range of side chain liquid crystalline homopolymers. Many have concentrated on liquid crystalline homopolymers with acrylate or methacrylate backbones with different combinations of flexible spacers or mesogenic side groups [2, S-81. Several studies on siloxane-based liquid crystalline homopolymers have also been undertaken and these materials have been well characterised in terms of their fundamental molecular motions [3, $9-101. Research on the dielectric properties of liquid crystalline copolymers has been limited [ 1l-151 with the effect of copolymer composition on thermal properties mainly reported. Some work concerning the dielectric properties of siloxane-based side chain liquid crystalline copolymers having two very different mesogenic pendant groups has concentrated on the effect of changing the ratios of the two pendant groups [ 12, 131 on the speed and activation energy of *To whom all correspondence
should be sent.
motion. In these systems, one of the homopolymers has also tended not to show a liquid crystalline phase 1131. The work reported here uses DRS to compare the mobility of a complete series of siloxane-based side chain liquid crystalline copolymers, with the two mesogenic side groups having a similar structure and flexible spacers based on the vinylacetic acid spacer (structures are shown in Fig. 1) developed in our laboratory [16, 171. The structure of the mesogenic units attached to the siloxane backbone (RI and R2) differs only in the length of the flexible spacer and both the homopolymers demonstrate liquid crystalline phases. Comparisons of dielectric mobility between the unaligned samples of the two respective homopolymers la and le, as well as that of the copolymers lb, lc and Id, will be presented. EXPERIMENTAL
Marerials
The polymers synthesised for this study used methods described elsewhere [ 16, 171.The structures of the polymers are shown in Fig. 1 and their thermal characteristics from differential scanning calorimetry (DSC) presented in Table 1 (T# is the glass transition temperature, while T, represents the isotropic temperature at which the material passes from the liquid crystalline state to the isotropic liquid state, also known as the clearing point). All samples were run on a Perkin Elmer DSC7 at a scan rate of lOC/min and the temperatures calibrated with an indium standard. The glass transition temperatures could not be observed for samples
1361
G. M. Day et ul
I362
Copolymer
R, : -
series
1
W ,l,CO,GH
R2 : -W
2,30 -&LO
z-@N
,),C02(CH &, 0
Fig. 1. Structures
of the polymers
lb and Ic, although it is reasonable to assume they would be close to that of the homopolymer and sample le. All materials showed conical fan textures, typical of a smectic A structure in hot stage transmission optical microscope experiments with crossed polars above r,.
I
of motions of the mesogenic pendant group since the siloxane backbone is itself relatively non-polar [3]. Previous studies on amorphous [2, 191 and semicrystalline [9] SCLCP materials have shown that it is necessary to exceed the glass transition temperature or melting point (the latter only relevant to semicrystalline materials) in order to liberate the strong dielectric relaxations of the mesogenic groups in the l&lo5 Hz frequency range. The dielectric spectra of the polymers in this study all show loss peak maxima temperatures well above the TB values observed by thermal analysis (or expected T, values in the case of copolymers 1b and lc).
Dielectric measurement
Variable frequency dielectric spectroscopy was performed using a Gen Rad 1689 Digibridge and techniques similar to those previously described [4, 181. The data presented involves G/w @F), where G is the sample conductance and o is the angular frequency of measurement (w = 2nf, where ,fis between 10 and IO5Hz). G/o is a dielectric parameter related to the imaginary component of the dielectric loss permittivity, t” = (G/w)/C,; C,, is the empty capacitance of the sample and hence determination of t” requires knowledge of the sample thickness. Samples of - 1 cm diameter were prepared by squeezing the material between the two plates at a temperature above the clearing point of the sample. This technique precludes a direct measurement of the thickness of the sample, preventing the calculation of the dielectric constant (6’) and the dielectric loss permittivity (6”). Since only peak position and shape is of interest in this work, G/o vs log f is presented, as has been done often elsewhere [4,9, 18-201in DRS of SCLCP materials.
Shape and height
The dielectric
qf the dielectric
loss spectrum
loss cunles
for the homopolymer
la (100 mol % RI) is shown in Fig. 2. It can be seen that the peak does not become prominent until the temperature reaches 4O”C, which is close to the T, for la and the loss peak decreases in height with increasing temperature (up to T = 65). The lack of a relaxation peak in the liquid crystalline state seems to be due to the low frequency of motion at these temperatures and its disappearance in the low frequency conductivity tail. Other workers have shown that peak height tends to remain relatively constant for loss peaks measured above r [21]. The copolymer 1b (Fig. 3) containing 75 mol % RI and 25 mol % R? behaved very differently to the homopolymer la in that the loss peak height showed a significant increase with increasing temperature, in particular above the clearing point. Copolymers lc and Id (Figs 4 and 5, respectively) and homopolymer le (Fig. 6) show relaxation spectra only within the frequency window of measurement in the liquid crystalline state and
RESULTS AND DISCUSSION
The unaligned dielectric loss spectra (G/w vs logf) for all homo- and copolymers la-le are presented in Figs 2-6. All the spectra showed an increasing loss tail at lower frequencies. This so-called “conductivity tail” may be due to some residual impurities in the samples which lead to conductivity. The dielectric loss maxima for most spectra first become visible at temperatures of between 40 and 50°C (35°C for Id). The loss spectra shown is attributed to a combination
Table 1. Thermal data for the ohase transitions of the nolvmers 1 Ratio Compd la lb 1C Id le
7-m
APwt
( C)
Pa
AH, (J/g)
20 22 26 34 18
3.4 4.6 6.3 5.1 6.5
RI
RZ
100 75 50 25 0
0 2s 50
-7 *
15
12
78
100
-4
106
_
*Glass transition not observed. tBaseline peak width of clearing transition from DSC
40 52 71
&IQ 10.9 13.4 19.4 16.2 17.2
Dielectric relaxation of liquid crystalline polymers
1363
___e___ Tt35 ___Q___ T=40 +
T=45
+
T=50
+
T=55
-o-T=60 +
T=65
e
T=70
___ +--- Tt76
1
2
3
4
log f (Hz)
5
Fig. 2. G/w (PF) vs loglof for polymer la (7’ in “C).
thus the effect of heating the sample through the clearing point cannot be determined and these relaxation peaks are constant in height. The decrease in loss peak height of la (Fig. 2) could be explained if the original sample of la was either partially or fully homeotropically aligned and the decrease in height could be ascribed to decreasing homeotropic orientation. Homeotropic alignment occurs when the mesogens are aligned perpendicular to the electrodes and thus the measuring field of the dielectric bridge largely interacts with the longitudi-
nal dipoles of these mesogenic units. Such homeotropic alignment could be induced electrically by application of electric fields of low frequency for SCLCPs with mesogens of positive dielectric anisotropy, as are those in this study [19] or by induced surface alignment [22]. The fact that such alignment, should it exist, persists so deeply into the isotropic region (the peak is still decreasing in height at 76”C, some 36°C above the clearing point) could be explained in part by the broadness of the biphasic region (coexistence of liquid crystalline and
w6-l(PF)
2
3
4
Fig. 3. G/w @F) vs loglOf for polymer 1b (T in “C).
5
+
T=40
+
T=45
--8-
T=50
+
T=55
*
T=60
+
T=65
1364
G. M. Day PI al
+T=40 +T=45 +
T=50
---I---
1dS.i
*
T=60
+
T=65
___Q___ T=70 ___+YJ___ T=76 ___+___ T=61
1
I
I
I
2
3
4
5
logf (Hz)
Fig. 4. GW @F) vs loglii / for polymer Ic (T in C).
amorphous phase as the material is heated into the isotropic melt). However. the entire biphasic transition region is some 20°C for sample la. as shown in Table 1. An alternative explanation for the continuing decrease in the height of the loss peaks (which is related to the relaxation strength and thus the effective number of dipoles per unit volume) is that it is due to increased thermal randomisation of the dipoles, a similar explanation as is proposed for the decrease in the c( relaxation in isotropic
thermoplastic relaxation spectra with temperature [231. The increase in the height of the loss peaks with increasing temperature for copolymer lb (Fig. 3)
could be explained if the original unaligned sample was partially or wholly planarly aligned (mesogenic direction parallel to the electrode surface) and that the increase in height of the loss peak with increasing temperature could be due to decreasing planar orientation. Planar orientation could be expected in
+
T=30
+
T=35
-x--T=40 +
T=45
+
T=50
+
T=55
._.Q_._T=60 .__a___T=65 .__+___T=70
0
i
I
I
I
1
2
3
4
Fig. 5. G:o @F) vs log10f for polymer Id (T in C).
5
logf (Hz)
Dielectric relaxation of liquid crystalline polymers
1365
(PF)
WI
5 \
\ +T=!Xl
4
--t-T=66 *T=65 -o-T=71
3
___+___ T = 76
2
1
I
i 1
I 2
I
I 4
3 Fig. 6. G/o (PF) vs log,,f
these systems due to template alignment in the microscratches in the electrode surfaces and the shearing, squeezing flow that the mesogens experience when the material is pressed between electrodes during sample preparation, aligning the liquid crystalline units parallel to the electrodes, in the direction of flow. Support for the idea that this material is at least planarly aligned also comes from the very broad loss peaks in Fig. 3 with widths at half height of greater than logf = 3, a common feature of planarly aligned materials [24] where the dipole being relaxed is mainly that which is perpendicular to the mesogenic group. It is difficult to see why this copolymer would differ in terms of alignment due to surface forces or squeezing flow compared to, say, the homopolymer la, although it is noted that the rheological behaviour of these materials is somewhat complicated [25]. Some contribution to increases in peak height and relaxation strength of SCLCPs, such as that seen for copolymer 1b, could be due to disruption of antiparallel packing that often occurs in liquid crystalline systems with strong terminal dipoles (such as the cyano groups used in this work). This packing causes cancellation of the dipoles (and thus reduced dipole strength). On heating, these anti-parallel associations are disturbed and may cause an increase in dielectric strength and peak height due to an increase in the number of available dipoles [26]. As explained above, the persistence in such an increase beyond the nominal clearing point may be due to the width of the biphasic region. The shapes of dielectric curves for SCLCP materials have been well-described in the literature [2,3] and have been attributed to different motions of the polar mesogenic pendant group. The two basic relaxations which have been observed are labelled 6 and tl. The 6 relaxation is related to motion of the longitudinal dipole due to the terminal cyano group
for
5
logf (Hz)
polymer le (T in “C)
and the mesogens short axis, while a relaxation is largely due to motions of the transverse dipole (resulting mainly from the motion of the ester group) around the long axis of the mesogen. The a relaxation in this context should not be thought to necessarily be the same as the a relaxation observed in isotropic thermoplastics above the DSC glass transition temperature. It should be noted that there is some controversy in the assignation of the a relaxation in SCLCPs dielectric spectra. In methacrylate systems it is believed to be largely related to main-chain motion, being ascribed to the relaxation of the carbonyl group in the side chain closest to the polymer backbone [2]. However, in polysiloxane materials it has been more often associated with the motion of the dipole transverse to the mesogen [2,3]. The a relaxation occurs at a higher frequency (faster motion) than the 6 relaxation because it does not require a large scale molecular rearrangement, being due mainly to spinning of the mesogen around its long axis. The correctness of the a and 6 assignment can be demonstrated if the samples are aligned homeotropically or planarly using an electric field [2,22], which was not successfully achieved in this study. For many side chain liquid crystalline polysiloxanes with a conventional alkyl chain based flexible spacer, it has been shown that dielectric loss curves of unaligned samples consist of a superposition of 6 and a peaks, although the peak maximum was effectively that of the much stronger, low frequency S relaxation (i.e. the motion of the longitudinal dipole moment). The a peak was usually determined by deconvolution, for example, by assuming a Fuoss-Kirkwood empirical relaxation function of the broad, overlapping 6-u peaks [3,24,27]. Recent use of a vinylacetic acid based flexible spacer in SCLCP materials has seen the dramatic reduction of clearing temperatures (cf. to SCLCP materials with an alkyl chain based spacer)
G. M. Day et al
In f,,, (Hz) 15 ela -e--lb --x.-lc .__+___ ld *le
2.9
3
3.1
Fig. 7. In 1,. vs I T ior the polymers
3.2
3.3
l/T
1.
[28]and the unaligned dielectric loss spectra of these and copolymers of Series 1 at the same absolute materials has shown distinct frequency separation temperature, as this does not take into account between u and 6 relaxation peaks [lo]. The homo- and the relationship of the measuring temperature to copolymers in this work (la-le) (which also involve key events in the polymer spectra (glass transition a vinylacetic acid based spacer in both mesogenic temperature, clearing point) [2,3,29, 301. Rather, pendant groups) do not exhibit distinct 6 and ry such comparisons are best made by comparing relaxation peaks (see Figs 2-6). A possible reason for dielectric loss spectra positions at the same reduced this could be that the high frequency c( peaks may temperature (T,,J. Potentially, Tredcan thus be scaled exist outside the frequency window at the experimental by the clearing point (= TJZ , where Tcxp is the experimental temperature and all temperatures are temperatures used. The loss spectra of the homopolyin K) or as a function of glass transition (= TCxp/Ta). mer le recorded at 40°C and 45°C show what appears to be the beginning of a peak due to c( relaxation at It has been stated that for several liquid crystalline high frequencies (Fig. 6). It should be noted that since polysiloxanes at temperatures more than 20 K below T,, it is better to scale data by T, since at these only one loss curve is observed for all the copolymers lbld, it appears that the copolymerised species temperatures of measurement the glass forming dynamics of the polymer backbone dominate the are statistically distributed along the polymer chain speed of motion [30]. Such a temperature scaling leading to something of an averaging of the dielectric relaxation of the homopolymer processes. This is as approach is similar to cooperativity or fragility plots [31] that have been employed in understanding seen in dielectric spectra of other SCLCP copolymers intra- and inter-chain coupling between conventional [13], although in another reported system, two distinct peaks were observed [12] and attributed to thermoplastic chains above the glass transition temperature [32]. “blocky” SCLCP copolymers. Figure 8 contains the same data as Fig. 7 but Position of the dielectric loss curves plotted in terms of peak position (In fmax) as a function The positions of dielectric peak position as a of l/T,,, , where T,& = T,,,/T,. It can be seen that this function of temperature (In fm vs l/r) are shown in separates the peaks to a greater degree. That is, rather Fig. 7. It can be seen that the spectra are relatively than lying close to each other as they are in Fig. 8, Arrhenius-like (linear). they are quite strongly separated, with the homopolySuch behaviour is usually observed for the 6 mer le with 100 mol % Rz (the longest spacer group) relaxation [2, 3,241, whilst the tl relaxation tends to lying at the highest frequency (fastest motion) for a show spectra that are more curved (WLF effect). given reduced temperature. Conversely, polymer la Copolymers lb and lc showed slight curvature, as with the shortest spacer group is moving the slowest would be expected, particularly in the case of 1b at a given reduced temperature, with the copolymers which, as discussed earlier, gave indications of planar lying in between in monotonic compositional order. alignment (and thus may be largely an a relaxation Inspection of Figs 7 and 8 at constant temperature process) by its increase in relaxation strength with and reduced temperature, respectively, allows comtemperature. parisons to be made as to the efficacy of scaling the It is not totally valid to compare peak positions data (and hence yields information about the nature (speed of motion) between the members of the homoof the motion). Even though glass transitions could
r
Dielectric relaxation of liquid crystalline polymers
In f,(Hz) 15
0
10
0 0
X
0
ox
cl
X4
0
X
0
0 0
0
0
la
o
lb
x
lc
l
Id
0
le
0 4
x
0
0
1 0
4
x
0
5
0
4
1367
0
4
x
0
0
4
0
4
X
4
0
0 0.8
1
0.9
1.1
1.2
1.3 lflrad
Fig. 8. In& vs l/T& for the polymers 1 (7’,, = T,,&T).
not be observed for lb and lc and thus normalised data with respect to T, could not be determined, if it is assumed that the glass transitions of 1b and lc were approximately in between those of the other polymers as shown in Table 1 (and thus are all of similar magnitude), comparison of data at constant Tti = TLxp/Tgis effectively the same as comparing them at constant experimental temperature. Thus the data of Fig. 7 can be thought to be qualitatively similar to being plotted against constant T, = TJT,. It can be seen that data of Fig. 7 (nominally TRd= Tcip/Tg) overlies each other to a greater degree than data of Fig. 8 (at constant Tti = T,JT). That is, it appears as though normalising the peak frequency positions with respect to the glass transition temperature is more effective than scaling using the clearing point. Such comparisons of scaling can best be seen by tabulation of the data at constant reduced temperatures (vertical cross-sections in Figs 7 and 8). Comparison of In fm for the polymers 1 at a Tmd= T47 of _ 1 is shown in the top part of Table 2 and for T,, = Tcxp/Tgof 1.27 is shown in the lower part of Table 2. As mentioned above, not all
Table 2. Comparison of In& values for polymers Polymer la lb IC
Id le
(T>,)
1.00 1.01 l.Oll 0.98 0.92
1at a constant
Td
IiT.4
Te, (“C)
In fm (Hz)
1.00 0.99 1.I0 1.02 I .09
40 55 70 70 76
4.15 1.83 9.21 9.21 10.13
Td
Polymer la
Id le
( T.rp/Ts) 1.27 I .27 1.27
l/T&
0.787 0.787 0.787
T., (‘C)
In .h (Hz)
65 70 70
8.75 9.21 9.21
of the polymers 1 had an observable T, and thus Table 2 is incomplete. Despite this, at a constant Td = T,,JT, = 1, the data in Table 2 shows that an increase in the amount of the mesogenic pendant group with the very long spacer (RJ leads to an increase in the value of lnf, , which indicates a faster S motion. Comparison of the two homopolymers la and le clearly shows that there is more rapid 6 motion by the polymer le with the longer spacer (Rz), as may be expected with increased mobility due to greater decoupling of the mesogen from the restricting, main-chain motion. The mesogen is able to move more freely and thus faster. The copolymers show intermediate behaviour with the different compositions of Id and le having little effect on speed of motion. This behaviour can also be seen by inspection of Fig. 8. However, when the data of Tmd= T,.,/T, is compared at a constant value for samples la, Id and le, it can be seen in Table 2 that the peak position at T, = 1.27 is approximately independent of composition, corresponding to the approximate overlaying of the curves in Fig. 7. It should be noted, however, that even when the data is scaled by the glass transition temperature, there is a slight increase in speed for copolymers and homopolymers Id and le which contain the longer, flexible spacer unit.
Table 3. Activation energies (G,) for polymers 1 Polymer
.% (Id mol-‘)
la (>E) lb (tTJ lb (z-T,)
166 168
Ic (< Ti)
175 141 137
Id (CT,) le (tK)
146
I368
G. M. Day YI al.
The fact that the data effecttvely scaled by the glass transition temperature are similar between the homopolymers and at least one copolymer is indicative of the role that the polymer backbone plays in the motion of the mesogens. Despite the side chain being sufficiently decoupled from the main-chain motion to allow the liquid crystalline phase to be attained, the motion of the mesogens remain tied to that of the main chain. This has been illustrated in other ways, for example by comparison of dynamic mechanical data (which measures backbone mobility) and dielectric relaxation (which examines the motion of the polar mesogen), demonstrates their interconnectedness [9]. Coupled backbone and mesogen motion from alignment measurements and variable frequency dielectric studies have also been demonstrated by Birenheide et ul. [33] who showed that the activation energy of the temperature dependence of polymer viscosity was similar to that of the S relaxation (motion around the mesogenic groups short axis) and hence were related. Given, however, that much of the data presented here were measured at temperatures close to or above the isotropic temperature, it is reasonable to assume that normalising data with respect to the isotropic temperature is also justifiable and such a scaling exaggerates the effect of increasing the spacer length in speeding up the mesogenic motion. Thus, provided that the data are at temperatures close to or above the clearing point, Tred= T,,,!T, may also be a valid way of comparing relaxational dielectric data between copolymers and not solely for comparison of static dielectric behaviour, as indicated by mean field theory [30]. If an Arrhenius relationship is assumed for the In _&, vs l/T data, then activation energies of motion (&J in the liquid crystalline and isotropic states can be calculated from Fig. 7. Calculated activation energies for the polymers 1 are shown in Table 3. where it is also indicated whether the data were obtained above or below the clearing point. It is difficult to make comparisons from the data of Fig. 7 across the whole composition range. as temperature-frequency data could not be obtained for all materials totally in the liquid crystalline or in the isotropic state. However, a few comments from the activation energies determined can be made. It can be seen from sample 1b that the activation energy of the mesogen in the liquid crystalline state is greater than that of the isotropic state, as is usually observed, and is ascribed to hindrance of motion due to the effect of the liquid crystalline potential field [34]. Comparison of activation energies of samples 1b. Ic. Id and le (all in the liquid crystalline states) indicates a lower activation energy for the copolymers with greater content of the longer, flexible spacer unit. Clearly the energy barrier to overcome motion is less in these materials. This is similarly true of comparison of la and sample lb in the isotropic state when both are above their clearing points. Thus, even in the absence of a liquid crystalline potential, copolymers with greater content of long-spacer mesogens are able to move with less energetic hindrance. This is in agreement with Simon and Coles [29] who showed that an increased length of spacer unit in a homologous series of polysiloxane SCLCP
systems
resulted
in an overall decrease
in activation
energy.
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