Phase morphology and molecular dynamics of a polyurethane ionomer reinforced with a liquid crystalline filler

European Polymer Journal 39 (2003) 2167–2174 www.elsevier.com/locate/europolj

Phase morphology and molecular dynamics of a polyurethane ionomer reinforced with a liquid crystalline filler A.G. Charnetskaya a, G. Polizos b,*, V.I. Shtompel a, E.G. Privalko a, Yu.Yu. Kercha a, P. Pissis b,* a

b

Institute of Macromolecular Chemistry, National Academy of Sciences of Ukraine, 02160 Kyiv, Ukraine Department of Physics, National Technical University of Athens, Zografou Campus, GR 15780 Athens, Greece Received 20 March 2003; received in revised form 20 June 2003; accepted 22 June 2003

Abstract Solution-blended binary composites of ionic segmented polyurethane (SPU-I) and liquid crystalline oligomer (LCO) were characterized by wide-angle (WAXS) and small-angle (SAXS) X-ray scattering, differential scanning calorimetry (DSC), thermally stimulated depolarization currents (TSDC) and dielectric relaxation spectroscopy (DRS). Both components mutually influenced their states of aggregation in blends (most significantly, promoting smearingout of interfaces between stiff and soft chain fragments of SPU-I into broad interfacial regions of intermediate composition). Apparently, the blend with w ¼ 0:10 happened to be most favorable for crystallization of the LCO, while the degree of microphase separation for SPU-I became lower and the distribution of stiff domains by sizes became broader, the higher the LCO content. The overall molecular mobility of SPU-I in blends was significantly reduced. This reduction included the intensity of the secondary and the primary relaxations, and of the interfacial Maxwell–Wagner– Sillars (MWS) relaxation, whereas the transition temperatures remained essentially composition-invariant. The Arrhenius-like behavior for the dc conductivity concomitant to the non-Arrhenius (i.e., Vogel–Tammann–Fulcher) frequency dependence for the a relaxation in blends suggested a decoupling of conductivity from the motion of the SPU-I soft chain segments.
2003 Elsevier Ltd. All rights reserved. Keywords: Liquid crystalline/polyurethane ionomer composties; Microphase separation; Molecular mobility; Dielectric relaxation

1. Introduction The degree of microphase separation (DMS) is one of the key parameters controlling physical properties of segmented polyurethanes (SPU) [1]. The polyaddition reaction by which stiff segments (mostly aromatic diisocyanates), are bonded to soft segments (oligomeric ethers, esters, dienes, etc.) via chain extenders (short-

*

Corresponding authors. Tel.: +30-210-7722983, fax: +30210-7722932 (G. Polizos); Tel.: +30-210-7722986, fax: +30-2107722932 (P. Pissis). E-mail addresses: [email protected] (G. Polizos), [email protected] (P. Pissis).

chain diols and/or diamines) is a stochastic process. This aspect, as well as rather wide molar weight distributions of component segments, is the natural explanation for both a broad dispersion of the size of stiff microdomains separated from a continuous soft phase, and a fairly low DMS. This parameter can be varied either through chemical modification (e.g., by changing the intrinsic flexibility of stiff and soft chain fragments of SPU, or by ionization), or through physical modification (by incorporation of appropriately chosen additives). In this respect, liquid crystalline fillers proved particularly effective as concerns regulation of the DMS of SPU [2,3]. Thus, the aim of the present paper is the further elucidation of this latter aspect through comparative studies of the effect of

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2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0014-3057(03)00136-8

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a liquid crystalline oligomer (LCO) on the DMS and on molecular dynamics, as revealed by dielectric techniques, of the SPU ionomer.

2. Experimental 2.1. Materials The initial SPU (hereafter referred to as SPU-O) was prepared by a pre-polymer method in the mixed solution of dimethyl formamide and methyl ethyl ketone. In the first stage, the pre-polymer was prepared at 80 C from oxytetramethylene glycol oligomer (molar mass 1000) and a mixture of 2,4- and 2,6-toluene diisocyanates (molar ratio 1:4). In the second stage, the pre-polymer chain was extended at 25 C with a stoichiometric quantity (3 mol) of N-methyl diethanol amine (MDEA) to obtain SPU-O (mass concentration of stiff chain fragments: 51%). The ionic SPU-I with the concentration of ionic centers of 8% was obtained by quaternization of tertiary nitrogens in the MDEA fragments of SPU-O with hydrochloric acid (1.1 mol of HCL per 1 mol of MDEA) [4]. The final chemical structure of the SPU-I can be represented as

The LCO (cholesteric ester of caprylic acid) was blended with the dimethyl formamide solution of SPU-I; thin films for physical characterization (LCO mass fractions w ¼ 0, 0.02, 0.10 and 0.20) were cast from the solution (initial concentration
19 wt.%) on glass slides and evacuated overnight at room temperature to constant weight. 2.2. Methods Wide-angle X-ray scattering (WAXS) patterns in the range of scattering angles (2h) 5–40 were measured with a DRON-2,0 diffractometer (copper radiation, k ¼ 0:154

nm; nickel filtering of the reflected beam; step-by-step scanning regime; recording of scattered radiation with a scintillation counter and digital conversion). As usual, scattering curves were normalised by thickness and Xray absorption [5]. Small-angle X-ray scattering (SAXS) data were obtained with the Kratky camera (KRM-type diffractometer). The primary beam intensity was controlled with a monitoring channel in the scattering angles range from 3 to 5. Copper Ka radiation and nickel filtering of the primary beam were used. Recording of the scattering radiation with a scintillation counter and digital conversion were performed using the step-by-step scanning regime. The geometrical parameters of the X-ray beam in the specimen plane and the detector position (30 mm for the length of the homogeneous portion of X-ray beam, and 290 mm for the specimen-detector distance) were chosen so as to satisfy the conditions of an ‘‘infinite’’ slit collimation [5]. Thermal transitions in the temperature interval 170– 470 K were recorded with the Mettler-Toledo DSC 9804 (heating rate: 10 K/min). Thermally stimulated depolarization currents (TSDC) were measured in the temperature interval 100– 300 K as described in detail elsewhere [6,7]. A disk-like

specimen (diameter: 13 mm) placed between brass plates of a parallel capacitor was polarized by a dc electric field (of 4.0 kV/cm) and cooled down to the liquid nitrogen temperature; then the field was switched off, the sample short-circuited and the depolarization current was recorded by an electrometer during heating to room temperature at a constant rate (3 K/min). In dielectric relaxation spectroscopy (DRS), complex dielectric permittivity, e ¼ e0  ie00 , of disc-like specimens (diameter: 20 mm) sandwiched between goldcoated brass electrodes was measured over the frequency window 102 –106 Hz in the temperature interval 150– 350 K using the Schlumberger frequency response ana-

A.G. Charnetskaya et al. / European Polymer Journal 39 (2003) 2167–2174

3. Results and discussion 3.1. Phase morphology In the first heating run of SPU-I (Fig. 1) one observes the low-temperature endothermic inflection centered around 205 K, the second inflection around 280 K, and the broad, bimodal endothermic effect of 47.8 J/g, with the main peak at 430 K and the shoulder at 403 K. On the second heating, the first transition remains nearly unchanged, whereas the next two transitions are replaced by a very broad process setting on at 285 K and extending over nearly 70 K. These data are the evidence for a complex phase morphology of the initial SPU-I, in which the soft segments form a continuous matrix with the glass transition temperature Tg1 ¼ 205 K, and the phase-segregated stiff domains consist of poorly ordered microregions (presumably, containing ionic groups) with Tg2 ¼ 280 K, and of well-ordered (but still noncrystalline [8], see also the WAXS data below) microregions with softening temperatures of Ts1 ¼ 403 K and Ts2 ¼ 430 K. In this context, the appearance of a single broad transition spanning the temperature interval from Tg2 to Ts1 in the second heating run is the result of a more homogeneous distribution of ionic groups throughout the disordered microphase of stiff chain fragments.

More quantitative information on the structure and DMS in SPU-I can be derived from the analysis of X-ray data. The diffuse maximum at 2h ¼ 20:8 on the WAXS curve of SPU-I (Fig. 2a) is the evidence for its noncrystalline state with a characteristic small-scale order (interchain spacing) of 0.420 nm, while the well-defined interference maximum at 2h ¼ 500 on the SAXS curve (Fig. 2b) can be attributed to a rather high DMS, manifesting itself as formation of a well-organized spatial macrolattice of stiff domains with a characteristic large-scale order (long period) of 10.5 nm. A sharp endothermic effect of 70.7 J/g with maximum at Tm ¼ 384 K and a small inflection at Ti ¼ 493 K on the DSC trace of the LCO in the first heating run (Fig. 1) can be attributed to the crystal-mesophase, and to the mesophase-isotropic melt transitions, respectively. High crystallinity of the initial LCO (of the order of 85%) is also confirmed by the presence of several sharp diffraction maxima (at 2h ¼ 9160 , 1370 , 16580 and 19040 ) on

1800

Scattering intensity / arb. units

lyzer (FRA 1260) supplemented by a buffer amplifier of variable gain (Chelsea dielectric interface) and the Hewlett-Packard 4284A Precision LCR meter [6,7].

2169

1600 1400 1200 1000 800

20

600

10

400

0

200

100 0

5

10

15

100

20 10 2 0 150

200

20

25

30

35

40

Scattering angle / deg 12000

Scattering intensity / arb. units

Heat flow / arb. units

endo

(a)

250

300

350

400

450

500

550

600

650

(b) Fig. 1. DSC traces for the first (––) and second (– – –) heating runs. Numbers at the curves: LCO contents in the blends.

2 8000

10 20

6000

4000

2000

0

T/ K

0

10000

0

20

40

60

80 100 120 140 160 180

Scattering angle / min Fig. 2. WAXS (a) and SAXS (b) patterns.

A.G. Charnetskaya et al. / European Polymer Journal 39 (2003) 2167–2174

the WAXS curve (Fig. 2a). In the second heating run, the melting peak becomes significantly broader, its maximum is shifted to about 365 K, and the total area (that is, crystallinity) is decreased by nearly 30%. In the first heating DSC runs for all binary blends, one can easily identify thermal transitions characteristic to each individual component (Fig. 1), although the experimental DSC traces for blends cannot be quantitatively reproduced by the simple additivity scheme. In fact, a common feature for all blends is the slight depression of the LCO melting point (by 2–3 K), suggesting limited interactions at the molecular scale; however, the intensity of the corresponding peak for sample with w ¼ 0:10 is much higher than for other blends. Moreover, with increasing LCO content the softening endotherm for SPU-I significantly broadened, its low-temperature shoulder and the main peak gradually shifted to lower temperatures (until the former eventually disappeared at w ¼ 0:20), while the corresponding heat effects tended to increase (except the sample with w ¼ 0:10 for which the total heat effect proved to be close to the additive sum of contributions from the heat of LCO melting and from the heat of softening of the SPU-I stiff domains). The weakening of approximately composition-invariant transition temperatures for each component in the blends concomitant to composition-dependent changes of the shape and intensities of the relevant endothermic effects suggest mutual influence of components on their states of aggregation (most probably, promoting smearing-out of interfaces between stiff and soft chain fragments of SPU-I into broad interfacial regions of intermediate composition). Apparently, the blend with w ¼ 0:10 happens to be most favorable for crystallization of the LCO, while the DMS for SPU-I becomes lower and the distribution of stiff domains by sizes becomes broader, the higher the LCO content. These conclusions are supported by the available Xray data. In fact, all crystalline reflections of the individual LCO can be identified on the WAXS patterns of the blends, especially at the highest LCO loading (Fig. 2a), although the reflections at the largest angles (i.e., at 16580 and 19040 ) are slightly shifted to 17390 and 19330 , respectively. The angular position of the SAXS maximum for SPU-I in the blends remains nearly unchanged, while its intensity gradually decreases, the

higher the LCO content (Fig. 2b). The latter effect can be attributed to both the lower electron density of the LCO, as well as to the lower DMS and concomitant broadening of the distribution of stiff domains by sizes. 3.2. Molecular dynamics Fig. 3 shows the representative TSDC thermogram of the sample with w ¼ 0:10. Similar thermograms were obtained also with the other two samples studied by dielectric techniques (Table 1). Four relaxations are observed in each thermogram, designated as c, b, a and Maxwell–Wagner–Sillars (MWS) in the order of increasing temperature [6,7]. c and b, at about 120 and 170 K, respectively, are secondary relaxations, assigned to crankshaft motions of the (CH2 )n sequences with participation of the attached polar carbonyl groups and to motions of the carbonyl groups with attached water molecules, respectively [7,9]. The a peak, close to Tg1 in DSC, is due to the primary a relaxation associated with the glass transition of the soft-segments matrix, whereas the MWS peak is due to the interfacial MWS relaxation, associated with the heterogeneous, microphase-separated morphology of SPUs [6–9]. The a and MWS relaxations, which are expected to be more intimately related to morphology than the local c and b relaxations [7], were studied in some detail, the results being listed in Table 1. The peak temperatures of

0.9 0.8

x5

0.7 0.6

I (pA)

2170

0.5 0.4 0.3 0.2 0.1 0.0 100

120

140

160

180

200

220

240

260

280

T (K)

Fig. 3. TSDC thermogram obtained with the sample with w ¼ 0:10.

Table 1 TSDC results for the a and the MWS relaxations w

Ta (K)

TMWS (K)

Ja (A/m2 )

JMWS (A/m2 )

FHW (K)

W (eV)

0.00 0.02 0.10

208 207 208

256.0 259.5 255.5

2.9 · 108 2.3 · 108 1.2 · 108

3.3 · 107 3.7 · 107 0.7 · 107

14.1 11.0 12.3

1.0 1.2 1.1

Ta and TMWS are the peak temperatures, Ja and JMWS the normalised peak current densities, FHW the full half-width of the MWS peak and W the activation energy of the MWS peak.

A.G. Charnetskaya et al. / European Polymer Journal 39 (2003) 2167–2174

The overall decrease of molecular mobility in the composites is not confined to room temperature (Fig. 4), but it shows up in the whole temperature range of DRS measurements, 150–350 K, in particular also at T < Tg . Fig. 5 shows e0 ðT Þ and e00 ðT Þ at a fixed frequency of 100 Hz. The data have been recorded isothermally in the frequency range 102 –106 Hz with temperature steps of 5 or 10 K and replotted in Fig. 5. In addition to the overall, non-additive decrease of molecular mobility in the composites, two relaxations (loss peaks) are observed in Fig. 5, in agreement with the results of TSDC measurements: the c relaxation at about 150 K and the a relaxation at about 220 K. The b relaxation (Fig. 3) is suppressed in the DRS measurement, as these are performed in dry nitrogen atmosphere and the sample is relatively dry. The properties of the c and a relaxations were studied in some detail by DRS measurements in frequency scans

8.5 8.0 7.5 7.0

w=0 w=0.02 w=0.10

6.5 6.0

ε'

the two relaxations, Ta and TMWS respectively, do not practically change with composition, whereas the corresponding normalised current density maxima, Ja and JMWS , are significantly reduced for the sample with w ¼ 0:10. Ja and JMWS have been normalised to the same polarizing field (4.35 · 105 V/m) and they are representative for the relaxation strength e of each peak, i.e., for the number of relaxing units contributing to the peak [10]. The full half-width, FHW, and the activation energy, W , in Table 1 refer to the MWS peak. The peak becomes more narrow in the composites. W has been calculated by the initial rise method [10]. We comment on the W values later in relation to the DRS results. The transition temperatures in Table 1 are independent of composition, whereas the intensity of a and of MWS does not behave additively (i.e., it does not change in proportion to the amount of SPU in the sample), these results being in agreement with those obtained by DSC. The significant decrease of the intensity of the MWS relaxation in the composite with w ¼ 0:10 suggests a significant decrease of DMS, in agreement with the results of DSC and SAXS. We will comment further on these results later in relation to corresponding results obtained by DRS. A comparative plot of the frequency dependence of the dielectric permittivity e0 ðf Þ at 298 K is shown in Fig. 4 for the three samples studied by DRS. An overall decrease of molecular mobility is observed in the composites, confirmed also by plots of e00 ðf Þ and rac ðf Þ (where e00 the dielectric loss and rac the frequencydependent conductivity) at the same temperature, not shown here. The decrease of e0 in the composites is by far more than it would correspond to additivity. The high values of e0 at low frequencies are related with conductivity effects, to be studied in more detail later.

2171

5.5 5.0 4.5 4.0 3.5 3.0 2.5 140 160 180 200 220 240 260 280 300 320 340 360

(a)

T (K)

20 w=0 w=0.02 w=0.10

10

0

12

ε ''

ε'

16

w=0 w=0.02 w=0.10

10

-1

8

10

4 -1

10

0

10

1

10

2

10

3

10

4

10

5

6

10

10

f (Hz) 0

Fig. 4. Frequency dependence of dielectric permittivity, e ðf Þ, at 298 K for the samples indicated on the plot.

-2

140 160 180 200 220 240 260 280 300 320 340 360

(b)

T (K)

Fig. 5. Temperature dependence of dielectric permittivity, e0 ðT Þ (a), and of dielectric loss, e00 ðT Þ (b), at 100 Hz for the samples indicated on the plot.

A.G. Charnetskaya et al. / European Polymer Journal 39 (2003) 2167–2174

at several temperatures. The main results, however, can be already seen in the temperature plots of Fig. 5. The temperature/frequency position of the relaxations does not change with composition, in agreement with the TSDC results. For the a relaxation this result is in agreement also with the DSC data for the glass transition temperature Tg1 being independent of composition. The shape of the response does not change with the composition for neither of the two relaxations, as indicated by normalised plots, e00 =e00max vs. f =fmax , not shown here. The Arrhenius plot for the c relaxation gives for the activation energy W ¼ 0:40 eV and for the preexponential frequency factor f0 ¼ 6  1015 Hz, independent from composition and in the range of values obtained also for other SPUs [6]. The main result in Fig. 5 is the significant reduction of the magnitude of the two relaxations in the composites, in particular for the sample with w ¼ 0:10. For that composite e00max decreases by a factor of about 2 for the c relaxation and by a factor of about 4 (as compared to a factor of about 2.5 in TSDC measurements, see Table 1) for the a relaxation, with respect to the initial polymer. This large and by far overproportional (with respect to additivity) reduction of the intensity of the c and a relaxations suggests some interactions between LCO and SPU and modification of the morphology in the composites. Probably, smearing-out of interfaces between stiff and soft fragments of SPU-I (see above) may be responsible for these effects. We will come back to this point later. It is interesting to note, however, at this stage that the reduction of the intensity of the cooperative a relaxation is larger than that of the local c relaxation. The large increase of e0 and e00 at high temperatures in Fig. 5 is related to conductivity effects. These effects were further studied as the motion of ions, giving rise to these effects, may be used as a probe of local morphology.

10

10

-7

w=0 w=0.02 w=0.10

-8

10

ð1Þ

where e0 the permittivity of free space. M 00 ðf Þ has been obtained by 1 ¼ M 0 ðf Þ þ iM 00 ðf Þ e ðf Þ e0 ðf Þ e00 ðf Þ þ i ¼ 02 2 2 e ðf Þ þ e00 ðf Þ e0 ðf Þ þ e002 ðf Þ

M  ðf Þ ¼

ð2Þ

At low frequencies rac ðf Þ in Fig. 6 becomes independent of frequency and the plateau value gives the dc conductivity rdc . This was determined at several temperatures from plots similar to those shown in Fig. 6. Fig. 7 shows the Arrhenius plot of rdc . We observe that at each temperature rdc is slightly larger in the composite with w ¼ 0:02 and significantly reduced in the composite with w ¼ 0:10, as compared to pure SPU. The data in Fig. 7 can be described by straight lines, indicating that the temperature dependence of rdc is described by the Arrhenius equation   W ð3Þ rdc ðT Þ ¼ r0 exp  kT

w=0 w=0.02 w=0.10

-10.0

-9

M''

σac(S/cm)

10

rac ðf Þ ¼ e00 ðf Þ2pf e0

0.09

w=0 w=0.02 w=0.10

0.06 10

Conductivity effects are typically studied within the conductivity formalism and the modulus formalism [7,11]. The modulus formalism was also employed to study dipolar and conductivity effects in LC/polymer composites based on (hydroxypropyl)cellulose and photo-polymerisable acrylic acids [12]. Fig. 6 shows a comparative plot of the frequency dependence of ac conductivity rac ðf Þ and of the imaginary part of electric modulus M 00 ðf Þ at 343 K for the three samples studied. rac (actually the real part of the complex conductivity) has been calculated by

0.03

-10

log( σdc(S/cm))

2172

-10.5

-11.0

-11.5

-11

-12.0

0.00 10

-2

10

-1

10

0

10

1

10

2

10

3

10

4

10

5

10

6

f (Hz)

-12.5 2.90

2.95

3.00

3.05

3.10

3.15

-1

1000/T (K ) Fig. 6. Ac conductivity rac (open symbols) and imaginary part of electric modulus M 00 (filled symbols) against frequency f at 343 K for the samples indicated on the plot.

Fig. 7. Arrhenius plot of dc conductivity rdc for the samples indicated on the plot.

A.G. Charnetskaya et al. / European Polymer Journal 39 (2003) 2167–2174

where r0 a constant, k BoltzmannÕs constant and W the activation energy of conductivity. It is interesting to note that very often in SPUs rdc ðT Þ above Tg is described by a Vogel–Tammann–Fulcher equation similar to the a relaxation [6,7]. The validity of the Arrhenius equation for the systems under investigation suggests a decoupling of conductivity from the motion of the polymeric chains. The values of W obtained from Fig. 7 are 1.3 eV for the initial SPU and 1.2 eV for both composites (±0.1 eV). Returning to Fig. 6, rac ðf Þ deviates from the plateau value with increasing frequency. In addition, a structure is observed in rac ðf Þ, more pronounced for the composite with w ¼ 0:10. This is better followed within the modulus formalism. For the sample with w ¼ 0:10 a peak and a shoulder are observed in M 00 ðf Þ in Fig. 6, at about 1 and 50 Hz, respectively. They correspond to the deviation of rac ðf Þ from the plateau value (the ‘‘knee’’ in rac ðf Þ) and the structure in rac ðf Þ, respectively. The latter is the pattern of the MWS interfacial polarization, as it is suggested by comparison with e00 ðf Þ (not shown here) and it will be confirmed later. Often the terms conductivity relaxation and conductivity current relaxation are used for the peak and the shoulder, respectively, of M 00 ðf Þ in Fig. 6, also in ionically conducting SPUs [13]. With decreasing temperature both features in M 00 ðf Þ for the composite with w ¼ 0:10 shift to lower temperatures (Fig. 8). The intensity of the shoulder does not change with temperature, which provides strong support for associating the shoulder with the MWS relaxation [13,14]. The temperature shift of the peak and the shoulder in Fig. 8 is described by an Arrhenius-type equation similar to Eq. (3) and the values of activation energies W obtained are W ¼ 1:1 and 0.9 (±0.1) eV, respectively. Returning again to Fig. 6, single peaks are observed in M 00 ðf Þ for the samples with w ¼ 0 and 0.02. Measurements at lower temperatures, similar to those shown

T=298 K T=303 K T=308 K T=313 K T=318 K T=323 K T=328 K T=333 K T=338 K T=343 K

0.08

M''

0.06

0.04

0.02

for the composite with w ¼ 0:10 in Fig. 8, show shoulders on the low-frequency side, corresponding to the conductivity relaxation. With increasing temperature they overlap more strongly with the MWS relaxation and finally cannot be resolved. Thus, the M 00 ðf Þ peaks for the samples with w ¼ 0 and 0.02 in Fig. 6 correspond to the MWS relaxation and the reason that conductivity relaxation and MWS relaxation are resolved for the sample with w ¼ 0:10 but not for the other two samples at 343 K is the lower rdc value for the former. As a consequence of this interpretation, the MWS relaxation in the M 00 ðf Þ representation appears at approximately the same frequency for the three samples studied, e.g., at about 100 Hz at 343 K. This result is in agreement with the TSDC result for the MWS peak in Table 1. From the temperature shift of the MWS M 00 ðf Þ peaks the activation energy of the MWS relaxation for the samples with w ¼ 0 and 0.02 has been determined to 1.0 and 0.9 (±0.1) eV, respectively. It is interesting to note that extension of the Arrhenius straight lines to frequencies of about 103 Hz, the characteristic frequency of TSDC measurements [10], gives temperatures in the range of 250–260 K (also for the sample with w ¼ 0:10), i.e., similar to those measured by TSDC for the MWS peak (Table 1). This provides additional support for the common origin of the M 00 ðf Þ peaks in Fig. 6 and the MWS TSDC peaks in Fig. 3 and Table 1. Thus, the temperature of both dc conductivity and interfacial MWS relaxation is described by the Arrhenius equation. The corresponding values of activation energies, obtained for dc conductivity within the conductivity formalism and the modulus formalism, and for MWS within the modulus formalism and by TSDC, are for the three samples studied in the range 0.9–1.3 eV. In general, similar values have been obtained for dc conductivity and for MWS relaxation, which is expected within the various models for MWS relaxation [13,14], providing further support for our interpretation. Finally, we turn our attention to the shape of the M 00 ðf Þ peak for the MWS relaxation. From the plots in Fig. 6 the full half-width (FHW) of the peaks is 1.6 decades for the sample with w ¼ 0:02, more than 2 decades for the initial SPU, and about 1.6 decades (but less accurate) for the sample with w ¼ 0:10. These values do not practically change with temperature, providing further support that the underlying mechanism is MWS relaxation [13,14] and allow to calculate approximate values of the Kohlrausch–Williams–Watts stretchedexponential shape parameter bKWW from the relationship [15]

0.00 -1

10

0

10

1

10

2

10

3

10

4

10

5

10

6

10

f (Hz)

Fig. 8. Imaginary part of electric modulus M 00 against frequency f for the composite with w ¼ 0:10 at several temperatures indicated on the plot.

2173

bKWW ¼

1:14 FHW

ð4Þ

to 0.57 for the pure SPU sample and to 0.71 for both composites. Two comments are here in order. First, these results are in very good agreement with the

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decrease of the half-width of the MWS TSDC peak in the composites. Second, the shape of the M 00 ðf Þ peak in ionically conducting systems has often been discussed in terms of the diversity of conducting paths [16], so that the results obtained here would suggest a homogenization of conducting paths for the MWS relaxation in the composites. Summarizing, the results for molecular dynamics and for charge carrier motion, obtained by the dielectric measurements, suggest that the overall molecular mobility is significantly reduced in the composites. This reduction includes the intensity of the secondary and the primary relaxations and of the interfacial MWS relaxation, whereas the transition temperatures remain invariant. These results suggest some mixing of LCO with SPU at the molecular level, in particular with the softsegment phase responsible for the a relaxation (which is consistent with the slight depression of the LCO melting point detected by the DSC) and reduction of DMS in SPU. This latter result is in agreement also with the significant reduction of dc conductivity in the composite with w ¼ 0:10. Interactions between the two components have been observed also in LC/polymer blends based on cholesteryl palmitate and poly(ethylene adipate) or polytetrahydrofuran by thermal analysis techniques [17]. It is interesting to note also that DSC and high resolution mechanical spectroscopy in microcomposites of a UV curable pre-polymer and LC molecules (cyanobiphenyl or benzoates mixtures) show a depression of the matrix Tg caused by LC molecules that remain dissolved in the matrix and act as a plasticizer [18]. Non-additivity in the composite properties and reduction of the initial DMS of the polymer matrix has been observed also in several other micro- and nanocomposites [19].

4. Conclusions 1. Solution blending of SPU-I and LCO results in the mutual influence of components on their states of aggregation (most significantly, promoting smearing-out of interfaces between stiff and soft chain fragments of SPU-I into broad interfacial regions of intermediate composition). Apparently, the blend with w ¼ 0:10 happens to be most favorable for crystallization of the LCO, while the DMS for SPU-I becomes lower and the distribution of stiff domains by sizes becomes broader, the higher the LCO content. 2. The overall molecular mobility of SPU-I in blends is significantly reduced. This reduction includes the intensity of the secondary and the primary relaxations, and of the interfacial MWS relaxation, whereas the transition temperatures remain essentially compositioninvariant.

3. The Arrhenius-like behavior for the dc conductivity concomitant to the non-Arrhenius (i.e., Vogel– Tammann–Fulcher) frequency dependence for the a relaxation in blends suggests a decoupling of conductivity from the motion of the SPU-I soft chain segments.

Acknowledgement Thanks are due to Prof. V.P. Privalko for his interest in this work.

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