Relaxation Behavior of Poly(ester carbonate) Block Copolymer Across the Melting Region

556

Relaxation Behavior of Poly(ester carbonate) BlockCopolymer Across the Melting Region

Cristina Alvarez,1 Maria J. Capitµn,1 Azar Alizadeh,2 Zbigniew Roslaniec,3 Tiberio A. Ezquerra* 1

1 Instituto de Estructura de la Materia, CSIC, Serrano 119, Madrid 28006, Spain2 General Electric, CR&D, Niskayuna, NY 12309, USA3 Institute of Materials Engineering, Technical University of Szczecin, Szczecin, Poland

Keywords: copolymers; dielectric spectroscopy; thermoplastic elastomers;

IntroductionThermoplastic elastomers are a class of block copolymersin which one segment, usually referred to as “soft”, is inthe rubbery state, imparting flexibility to the material,while the other, called “hard”, is either in the glassy orsemicrystalline state providing high elastic modulus,strength and mechanical stability to the material.[1] Thesematerials are of interest for a great variety of industrialprocesses including extrusion, injection and compositepreparation. On this basis, special attention has beendevoted to poly(ether-ester) copolymers (PEE) in whichthe hard segment corresponds to poly(butylene terephtha-late) (PBT) and the soft segment to poly(tetramethyleneoxide) (PO4).[1, 2] The ability to crystallize of the PBTsegments controls the appearance of microphase sepa-rated crystalline hard segment domains. The resultingsemicrystalline copolymers may exhibit combined prop-

erties of vulcanized rubbers and thermoplastics.[1, 3] Gen-erally, the chemical and physical characteristics of theindividual segments in a block copolymer have profoundeffects on both the phase behavior and the structure of theresulting microphases. For this reason, it is important awell-understanding of the relationship between physicalproperties, chemical composition and microstructure. Ofspecial interest is to characterize the changes of the vis-coelastic properties when those elements providingmechanical stability to the system are modified or even-tually destroyed. This happens, for example, when themelting temperature of the crystalline microdomains isapproached. Different mechanisms taking place in theamorphous phase can lead to energy dissipation in poly-meric materials.[4] At temperatures below the glass transi-tion temperature, Tg, mechanical energy can be dissipatedto some extent by the so called secondary relaxation,

Full Paper: The dielectric behavior of a family of blockcopolymers based on poly(butylene terephthalate) (PBT)and aliphatic polycarbonate (PC), with the composition ofPBT varying from 100 wt.-% to 40 wt.-%, has been inves-tigated by broad-band dielectric spectroscopy in the fre-quency range from 10–1 to 109 Hz. At temperatures abovethe glass transition, the merging of local c process and thecooperative b process was studied for PBT and 40/60 and60/40 PBT-PC copolymers. The experimental data can besatisfactorily analyzed considering a sum of two Havri-liak-Negami functions indicating that both b and c relaxa-tions can be treated as independent processes. Dielectricmeasurements for the 40/60 copolymer were extendedover and above the melting region to characterize theinfluence of crystal melting on the dynamic behavior. It isshown that the progressive loss of the crystalline lamellarstack microstructure, as characterized by differential scan-ning calorimetry and small- and wide-angle X-ray scatter-ing, provokes a dramatic change in the dynamic behaviorof both the b and c process.

Macromol. Chem. Phys. 2002, 203, No. 3 i WILEY-VCH Verlag GmbH, 69469 Weinheim 2002 1022-1352/2002/0302–0556$17.50+.50/0

DSC scans for PBT-PC copolymers. PBT/PC at differentratios. Heating rate: 20 K N min–1.

Macromol. Chem. Phys. 2002, 203, 556–564

Relaxation Behavior of Poly(ester carbonate) Block Copolymer Across the Melting Region 557

which involves local motions of the macromolecularchain.[5, 6] Additionally, polymers exhibit a primaryrelaxation above Tg involving segmental motions.[5, 6]

Typically for thermoplastic elastomers in the temperatureregion of technological interest both relaxation processescoexist or eventually merge. This fact makes the study ofthe relaxation behavior relatively complex.

Block copolymers based on poly(butylene terephtha-late) (PBT) and aliphatic polycarbonate (PC) exhibit ahigher melting point depression and lower crystallinitythan equivalent poly(ester-ether) block copolymers basedon PBT and poly(tetramethylene oxide) (PO4).[3, 7–9] Thisis mainly due to the Flory-Huggins interaction parameterwhich is expected to be lower for PBT-PC than for PBT-PO4 on the basis of the solubility parameters.[10] Theinteraction between the hard and soft segments slowsdown the kinetics of crystallization and decreases themelting temperature when hard segment concentrationdecreases. This makes these systems suitable candidatesto investigate the effect of thermal stability on the visco-elastic properties of the materials.

In the present paper, we report results related to thevariation of the dynamic properties with temperature, asrevealed by dielectric spectroscopy, in model poly(estercarbonate) block copolymers. The main aim of the workis to follow the primary and secondary relaxation, whichin this kind of block copolymers are referred to as b and c

relaxation, respectively,[11, 12] across the melting region inorder to characterize the impact of the mechanical stabi-lity loss on the dynamic properties.

Experimental PartA family of block copolymers based on poly(butylene ter-ephthalate) (PBT) and an aliphatic polycarbonate (PC, Dura-carb 1.2 from PPG Inc. (USA), M

—n = 1000, q = 1.076 g/cm3),

where the composition of the hard segment (PBT) was variedfrom 100 wt.-% to 40 wt.-%, was investigated. The chemicalstructure and composition of the block copolymers are illus-trated in Figure 1 and Table 1, respectively. For the synthesisof the copolymers a three-stage polycondensation methodwas used as described elsewhere.[7, 13] The samples werecompression moulded between thin Teflon films at 20 Kabove their corresponding calorimetric melting temperature(see Table 1). The thickness of the resulting films wasapproximately equal to 200 lm and the samples were kept atroom temperature subsequent to quenching.

Calorimetric measurements were carried out using a Per-kin-Elmer DSC-7 differential scanning calorimeter. The tem-perature calibration was performed by recording the onsetmelting temperatures of indium and zinc standards. The melt-ing behavior of the samples was monitored at a constant heat-ing rate of 20 K N min–1 and under a nitrogen atmosphere. Toreduce differences among samples due to thermal lag, discoidsamples of approximately 11 mg weight were employed.

Wide-angle X-ray scattering (WAXS) experiments wereconducted at room temperature by means of a Seifert dif-

fractometer (XRD 3000), which operates in a reflection h/2hgeometry with Ni-filtered Cu Ka radiation. Small-angle X-ray diffraction (SAXS) measurements were also performedat room temperature, using a Rigaku (Rotaflex, Ru-200)rotating anode with Ni-filtered Cu Ka wavelength combinedwith a 2D Photek 5126512 pixels CCD detector. The CCDhas an active surface of 45 mm diameter and is providedwith a beam stop to avoid the high intensity of the directbeam. The beam focus size on the anode (2006200 lm2)was reduced by using double collimating slits down to 150lm diameter at the sample position. SAXS experiments wereconducted in the transmission geometry. Finally, the 2D datawere integrated along a radial cut in order to obtain the one-dimensional intensity function weighted by the pixel contri-bution. This function was, in turn, corrected by the corre-sponding air scattering intensity and the Lorentz factor.[14]

Some additional SAXS measurements were also performedwith synchrotron radiation in the polymer beam line A2 atHASYLAB (Hamburg, Germany) as described elsewhere.[13]

Complex dielectric permittivity measurements (e* = e9 –ie99) were performed over a broad frequency window, 10–1

Hz a F a 109 Hz, in the temperature range of 123 K a

T a 373 K. To cover the above frequency interval, two dif-ferent experimental set-ups were used: in the 10–1–106 Hzrange, a Novocontrol (Huntsangen, Germany) system inte-grating a SR 830 lock-in amplifier with a dielectric interfacewas employed, and in the 106–109 Hz range, dielectric meas-urements were obtained by means of a HP4291 (Hewlett

Figure 1. Scheme of the investigated copolymers: m valuesgiven in Table 1, l L 7 and x L 1.12, R = C6H12 isomers. Thegiven data are average values derived from stoichiometric ratioand NMR analysis.

Table 1. PBT weight percentage, degree of copolymerization,melting enthalpy (DSC), long space (SAXS), glass transitiontemperature, and melting temperature in poly(ester carbonate)copolymers.

Sample PBTwt:-%

m DHm

J N gÿ1

Lnm

Tg

KTm

K

100/0 100 77.1 310.7 498.380/20 80 20 51.0 8.4 298.7 450.960/40 60 8 50.5 10 279.8 391.350/50 50 5 27.2 13.3 273.4 354.440/60 40 3 47.5 16.0 261.5 334.30/100 0 46.3 259.1 342.9

558 C. Alvarez, M. J. Capitµn, A. Alizadeh, Z. Roslaniec, T. A. Ezquerra

Packard, Hyogo, Japan) coaxial line reflectometer. In thiscase, the complex permittivity was calculated by measuringthe reflection coefficient at a particular reference plane.[15]

These two instruments were integrated in a Novocontrolbroadband dielectric spectrometer. The temperature in thisspectrometer is controlled by nitrogen jet with a temperatureerror of l0.1 K during every single sweep in frequency. Cir-cular gold electrodes were sputtered onto the free surfaces ofthe films with a diameter of 3 cm for the measurements up to106 Hz and 0.5 cm for measurements at F A 106 Hz. In orderto study the relaxation behavior of 40/60 PBT-PC copolymerabove the melting temperature, Kapton spacers were used toavoid the collapse of the sample during heating.

Results

Thermal Behavior

The DSC heating traces of the moulded and stored PBTand PBT/PC copolymers are shown in Figure 2. It can beobserved that all copolymer samples present an uniqueglass transition temperature, Tg, as revealed by a step-liketransition at low temperatures. The value of the glasstransition temperature, as taken from the middle point ofthe transition, decreases by decreasing the PBT content(see Table 1). In Figure 2, it can also be observed that thecopolymers exhibit a multiple melting behavior over arelatively broad temperature range. This melting rangenarrows continuously with decreasing the PBT content ofthe copolymer. The final melting temperatures of thesamples, which correspond to the melting of the moststable crystals, are tabulated in Table 1. The continuousdecrease of the final melting temperature with decreasingthe PBT content of the copolymers is a consequence ofthe limited size of the crystallites, which, in turn, is dueto the increment of imperfections as the composition ofPBT segments decreases.[1, 13]

In accordance with previous studies, this transition inthe PBT sample is accompanied by a small endothermic

peak which has been reported to be due to a physical age-ing effect.[2, 16] This effect corresponds to the relaxation ofthe glass toward the equilibrium which occurs during thestorage of the sample at room temperature. The endother-mic peak disappears in the immediate heating DSC scanafter cooling down from melt as illustrated in Figure 3. InFigure 3, the heating traces of the copolymers immedi-ately after cooling from the melt to 173 K at a rate of20 K N min–1 are collected (the samples in this case werenot stored at room temperature). These second heatingscans reveal that the copolymers with a hard segmentcomposition equal or lower than 50 wt.-% can beobtained in a completely amorphous state upon quench-ing from the melt. From this observation one can con-clude that for these samples the endothermic peaksdetected in the 1st DSC thermograms (Figure 2) corre-spond to the melting of crystals formed upon storage atroom temperature.

Wide-Angle X-Ray Scattering

The analysis of the WAXS diagrams of the samples (Fig-ure 4) shows that copolymers with a high PBT contentcrystallize with a structure equivalent to the a-poly-morphic form of the PBT homopolymer.[17] However, asthe composition of PBT in the copolymer is diminished,the corresponding crystalline phase is more severelyaffected by the presence of soft segment of PC. Theseeffects are clearly revealed by the change in both thenumber and position of the Bragg peaks for the highestPC content copolymers (see Figure 4). In fact, for the 40/60 copolymer, the positions and/or relative intensities ofthe peaks appearing at 11.9, 17.7, 18.7, 22.5, and 24.98can not be purely assigned neither to the a nor the b struc-tures of PBT.[17] Instead, the mentioned peaks could be

Figure 2. DSC scans for PBT-PC copolymers. PBT/PC at dif-ferent ratios. Heating rate: 20 K N min–1.

Figure 3. DSC (2nd run) scans for PBT-PC copolymers. PBT/PC at different ratios. Heating rate: 20 K N min–1.


associated with two possible structures: i) a distorted butpure PBT crystalline lattice, or ii) a mixed crystallinePBT/PC phase with the appropriate segment weight ratio.Our ground for the second hypothesis is based on aremarkable feature in the X-ray diagrams of the copoly-mers with middle-to-high PC contents, where the relativeintensity in the vicinity of 21.08 is increased. Thisincrease, on the other hand, is accompanied by the fadingof the valley intensity between the (�1111) and (100)reflections for copolymers with a PC content equal orhigher than 50 wt.-%. The presence of the excess inten-sity in the vicinity of 21.08 is consistent with the inclu-sion of the PC segments in the crystalline phase. Finally,the inclusion of the PC segments in the crystalline latticeof the copolymer would cause a severe depression of itsmelting temperature, which is in a good agreement withthe experimental observations (Table 1).

Small-Angle X-ray Scattering

Figure 5 illustrates the SAXS patterns of the investigatedcopolymers. In all cases, a maximum in the scatteredintensity, as a function of the reciprocal lattice vectorq = 4psinh/k being 2h the scattering angle, is detectedindicating the existence of a crystalline lamellar stackmicrostructure. As the PC content increases, two effectsare evidenced: i) a peak sharpening, and ii) a shift of themaximum towards lower angles. The values of the longspacing calculated as L = 2p/qmax after the Lorentz cor-rection are collected in Table 1. The observed increase in

the L-values with increasing PC content is due to thethickening of the interlamellar amorphous phase separat-ing the crystalline lamellae. As mentioned earlier, thecalorimetric results (Figure 2 and 3) clearly show a con-tinuous depression of the melting temperature, and thus areduction of the thickness of the lamellar crystals, withincreasing PC content. Additionally, the heat of fusion ofthese crystals, i.e. the overall crystallinity of the samples,decreases with increasing the PC content of the copoly-mer (Table 1). Therefore, the sharpening effect could beunderstood as due to the formation of a more homoge-neous crystal lamellae dispersion in the material as PCcontent increases.

Broad-Band Dielectric Spectroscopy

The dielectric loss (e99) values measured for the PBT, 40/60 and 60/40 PBT/PC copolymers as a function of fre-quency at three temperatures (above the correspondingcalorimetric glass transition temperature (Tg)) are illu-strated in Figure 6. The dielectric spectrum of all com-pounds reveals the existence of two dielectric processes:the b relaxation, associated with the glass transition, anda faster secondary process that is referred as c relaxa-tion.[9–12] The frequency of the maximum loss, Fmax, forboth processes shifts towards higher frequencies as thetemperature is increased. In all cases, both processesmerge at high temperatures. In addition, at lower frequen-cies, a strong increment of e99, associated with an electri-cal conductivity, is observed. It is to be noted that for

Figure 4. WAXS diffractograms (vertically shifted) as a func-tion of the scattering angle 2h of the investigated copolymers.Vertical dotted lines indicate characteristic reflections of the a-form of PBT.

Figure 5. SAXS patterns of the 40/60, 50/50, 60/40, and 80/20PBT-PC copolymers as a function of the scattering vectorq = 4p/k sinh, being 2h the scattering angle and k = 0.15 nm theX-ray wave length.


both relaxations, the intensity of e99 in the pure PBT islower than that of copolymers. This is, in turn, due to thecorresponding higher degree of crystallinity of the homo-polymer.

In order to determine the specific effect of crystal melt-ing on the relaxation behavior of the 40/60 copolymer,the dielectric measurements for this sample wereextended over and above the melting region (Figure 7).The choice of this particular sample was based on its rela-tively low temperature melting range, which allows boththe b and c relaxations to remain within the experimentalfrequency window even in the molten state.

The phenomenological Havriliak-Negami (HN) formu-lation was used to describe the frequency dependence ofthe complex dielectric constant (e* = e9 – ie99) in the vici-nity of dipolar absorption:[18]

e� ¼ ev þe0 ÿ ev

½1þ ðixsHNÞb�cð1Þ

where e0 and ev are the relaxed (x = 0) and unrelaxed(x = v) values of the dielectric constant, sHN is the cen-tral relaxation time of the relaxation time distributionfunction, and b and c being [0 a (b, c) a 1] are shapeparameters which describe the symmetric and asymmetric

broadening of the relaxation time distribution function.An addtional contribution of the conductivity process istaken into account by adding a term: –ir/(evacx

s) to Equa-tion (1). Here, r is related to the direct current electricalconductivity, evac is the dielectric constant of vacuum, andthe value of the coefficient 0 a s a 1 depends on the con-duction mechanism. At low temperatures (T a Tg), thedielectric data (e9and e99) from the isothermal scans werefitted by a single HN function since only the c relaxationwas observed. Recent studies show that primary and sec-ondary relaxations can be treated as being independentprocesses.[19–21] Therefore, above Tg a sum of two HNfunctions was used to describe both b and c relaxations.For both relaxations, a symmetric shape was obtainedfrom the fitting procedure (c = 1). The dash lines in Fig-ure 6 exemplify the separate contributions of both pro-

Figure 6. Isothermal plot of the dielectric loss component, e99,versus frequency in the merging region for investigated copoly-mers. The continuous solid lines represent the results of fittingexperimental data to the sum of two HN equations with a con-ductivity term. The dash lines at the highest temperature showthe separate contributions of b and c processes.

Figure 7. Isothermal plot of the dielectric loss component, e99,versus frequency in the melting region for (f) 40/60 PBT-PCcopolymer (Tm = 334.3 K). The continuous solid lines representthe results of fitting experimental data to the sum of two HNequations with a conductivity term. The dash lines show theseparate contributions of b and c processes.


cesses at the higher temperatures. In the melting region(T A Tm), a sum of two HN functions was also used to fiteach relaxation of the isothermal data set as illustrated inFigure 7. It is worthy to remark that when the merging ofthe both relaxations is strong, the relaxation time value ofthe c process was extrapolated from its Arrhenius beha-vior at low temperatures. Additionally, c = 1 value forboth c and b relaxations was assumed because it wasshown to provide an accurate description of the experi-mental data. In Figure 8, the relaxation times obtainedfrom the Havriliak-Negami fitting, sNH, for both the b andc processes are represented as a function of the reciprocaltemperature for the different samples. The temperaturedependence of the relaxation times for the c process fol-lows an Arrhenius-type behavior, which is typical forrelaxations corresponding to local motions of small mole-cular fragment.[22] No change of the apparent activationenergy value (Ea L 45 kJ N mol–1) with PBT content in thecopolymer is observed. The b process exhibits a charac-teristic non-Arrhenius temperature dependence similar tothat observed for the relaxation associated with the coop-erative segmental motions appearing above the glass tran-sition temperature[22] In this case, it is well known that thetemperature dependence of relaxation times, sHN, followsa Vogel-Fulcher-Tamman (VFT) behavior:

sHN ¼ s0 expDT0

ðT ÿ T0Þð2Þ

where s0 is the relaxation time in the high temperaturelimit, T0 is referred as the Vogel temperature, and D is thefragility strength parameter.[23] In order to obtain accuratefits, and in accordance with a recent discussion,[24] a valueof s0 of 10–13 s was assumed. The continuous lines in Fig-

ure 8 represent the best fits of the experimental sHN valuesto the VFT equation. The corresponding parameters arecollected in Table 2. The dielectric strength, De (beingDe = e0 – ev), and the shape parameter, b, for both relaxa-tions corresponding to the different copolymers aredepicted in Figure 9 as a function of a normalized tem-perature, Tg /T.

Discussion

The c-Process

As mentioned above, the low temperature process can beascribed to a local motion in the amorphous phase, inagreement with previous observations in similar sys-

Figure 8. sHN as a function of reciprocal temperature for b andc processes. Symbols are experimental data: (h, H) 60/40 and(f, F) 40/60 PBT-PC copolymers. For the copolymers, solid andopen symbols represent the b and c processes, respective-ly.6and + represent b and c processes for PBT. Solid lines arefitting to the VFT equation.

Table 2. VFT parameters obtained from Equation (2) for the brelaxation.

Sample s0

sD T0

K

100/0 1 N 10–13 4.1 279.960/40 1 N 10–13 5.6 236.040/60 1 N 10–13 7.53 209.0

Figure 9. Characteristic parameters of the dielectric losscurves obtained from fitting to the HN equation for the b and cprocesses as a function of normalized temperature Tg(K)/T(K):(a) symmetric broadening parameter b, and (b) dielectricstrength De. Error bars a provided on the right side. Symbols areexperimental data: (h, H) 60/40 and (f, F) 40/60 PBT-PC copo-lymers. For the copolymers, solid and open symbols representthe b and c processes, respectively.6and + represent b and cprocesses for PBT. Solid lines are a guide for the eye. The glasstransition temperatures are 279.8 K, 261.5 K, and 310.7 K for60/40, 40/60, and PBT, respectively.


tems.[9, 12] As shown in Figure 8, the relaxation time forthis process follows an Arrhenius-type behavior with anactivation energy of L45 kJ N mol–1 for all the investi-gated samples. Due to the fact that the reported activationenergies for both carbonate and ester groups are verysimilar, it is difficult to distinguish between these pro-cesses in terms of their activation energies.[5, 12] The abovevalue is consistent with the assignment of the low tem-perature process to a local motion associated either withthe ester group in PBT or with a combination of ester andcarbonate groups in the PBT-PC copolymer.[9] The c-pro-cess can be well described by a symmetric (c = 1) HN-function with a similar broadening b-value in all theinvestigated systems (Figure 9 (a)). At fixed temperature,the dielectric strength values (Figure 9 (b)) slightlyincrease with increasing PC block concentration asexpected by the progressive decreasing of crystallinitywhich tends to increase the amount of dipoles in theamorphous phase involved in the relaxation. Below Tg,Dec values slightly and linearly decrease with decreasingreciprocal temperature. A deviation from this linear beha-vior is observed when the glass transition temperature, Tg,is approached. This latter effect is most remarkable forthe copolymers. Dec values exhibit a clear increase abovecertain temperatures Tx L 0.97 Tg, denoting a change inthe temperature dependence of this magnitude (the cross-over temperatures are indicated by arrows in Figure 9(b)). This effect has been observed in other polymericsystems and has been explained in terms of an increase inthe amplitude of the motion as more free volumebecomes available around Tg.[19, 21]

The c-b Merging

The b-process in this kind of copolymers is associated withthe segmental motions in the amorphous phase appearingabove Tg.[9] Accordingly, the corresponding relaxationtimes follow a VFT temperature dependence (Figure 8).The strong-fragile scheme proposed by Angell[23] allowsone to classify glass-forming systems on the basis of thetemperature dependence exhibited by the viscosity or therelaxation time of the equilibrium liquid. The fragilitystrength parameter, D, controls the degree of departurefrom the Arrhenius law exhibited by the relaxation time(Equation (2)). This approach can be used to distinguishbetween strong (high D values) and fragile glass formers(low D values). In polymeric systems, Ngai et al.[25] asso-ciated fragility with the degree of intermolecular coopera-tivity. In our case, the values obtained for the PBT and 40/60 and 60/40 PBT-PC copolymers (Table 2), suggest thatthe system tends to be more fragile when the PBT segmentcontent increases. A higher PBT content may induce anincrease of the chain rigidity, causing an increasing hin-drance for the segmental motions in the amorphous phaseas fewer conformations become available.

At temperatures above Tg, both the b and c processescoexist (see Figure 6) and can be treated as independentprocesses.[19–21] In Figure 9(b), it can be observed that Deb

for PBT increases with temperature as frequentlyobserved for semicrystalline polymers.[26] This fact hasbeen explained by a temperature assisted mobilization ofthe amorphous phase which in semicrystalline polymersis dynamically restricted by the presence of crystals. Anopposite behavior is observed for the 60/40 and 40/60copolymers. Deb for these samples, while semicrystalline,seem to follow qualitatively the Fuoss-Kirkwood equa-tion proposed for amorphous systems.[22] This can beunderstood by recalling the fact that in this kind of ther-moplastic elastomers, the amorphous phase is not aseffectively restricted by the crystalline phase as it is forsemicrystalline homopolymers.[27, 28] This is triggered bythe loose microstructure of these systems consisting of ahomogeneous dispersion of lamellar crystals embeddedon a highly flexible amorphous phase.[27] In our case,such a picture is supported by the SAXS experiments(Figure 5), which show that the crystallites become moreand more separated as the amount of PC in the copolymerincreases. As the temperature increases above Tg, Deb ofthese copolymers tends to decrease concurrently with anincrease of Dec. This behavior has been observed in sev-eral amorphous polymers[20] but not in block copolymers.Moreover, extrapolation of the Deb results shown in Fig-ure 9 (b) would indicate a Tonset temperature (Deb = 0) ofabout 1.40 Tg for the block copolymers. This temperaturecorresponds to the point where both b and c relaxationsshould merge completely and only a common relaxationwould be observed.[20, 21] However, in these block copoly-mers, it can be observed that Deb values show a cleardeparture from the decreasing behavior at temperaturesabove 1.15 Tg. Attempts to describe the experimentalresults by forcing Deb values extrapolated from the pre-vious linear behavior did not provide reasonable fits. Forthe 40/60 copolymer the dielectric measurements crossthe melting range allowing one to characterize the influ-ence of crystal melting on the dynamic behavior. Thedielectric loss (e99) values measured for this copolymer asfunction of the frequency for four temperatures in themelting range are illustrated in Figure 7. The dielectricspectrum for each temperature reveals the presence ofboth b and c relaxations above Tm in the measured fre-quency range. The impact of the crystal melting on therelaxation behavior is reflected by the dramatic change inthe trend of Deb with temperature. For T A Tm (1.27 Tg)Deb undergoes a dramatic increment as a function of tem-perature. Concurrently, the increase of Dec with tempera-ture becomes stronger. This could be due to an increaseof the number of relaxing dipoles. SAXS experimentsperformed in the melting temperature range describe thestructural changes across the melting region of the sam-ple. The SAXS patterns (Figure 10) show a continuous


shift peak maximum towards lower q values by increas-ing the heating temperature. This behavior corresponds toan increase of the long period which goes from L175 �at 1.16 Tg to L248 � at 1.29 Tg. Above this temperature,the accuracy in the peak position decreases due to thepeak broadening and to the proximity of the beam stop.The integrated area under the SAXS peak decreases dra-matically at temperatures above 1.29 Tg, indicating themelting of the copolymer (see inset Figure 10). Theseresults in conjunction with the DSC data suggest a pro-gressive reduction of the crystalline fraction. Therefore,the corresponding increment of the amorphous phase isresponsible for the observed increment of the dielectricstrength for both b and c relaxations in the meltingregion.

ConclusionsThe dielectric broad-band experiments in poly(ester car-bonate) block copolymers reveal the existence of tworelaxation processes. The c relaxation is an Arrheniusactivated process associated with a local motion of esterand carbonate groups and exhibits similar apparent acti-vation energy values for all the investigated copolymers.The b relaxation appears above Tg and is associated tosegmental motions with relaxation times governed by aVogel-Fulcher type temperature dependence. Both pro-cesses can be considered as being independent in the mer-ging region and therefore described by a superposition oftwo Havriliak-Negami functions. For the 40/60 copoly-mer the relaxation studies were extended over and abovethe melting region. A full analysis of the temperaturebehavior of relaxation times, relaxation strengths, andshape parameters of both processes was performed evenacross and above the melting region. A strong difference

in the temperature dependence of the dielectric strength,De, for both c and b relaxation was observed upon com-paring the homopolymer, PBT, and the copolymers.While for PBT Deb increases linearly with reciprocal tem-perature, the opposite behavior was found for the copoly-mers. The impact of the crystal melting on the relaxationbehavior is reflected by the dramatic change in the trendof both Deb and Dec with temperature undergoing a dra-matic increment as a function of temperature.

Acknowledgement: C. A. thanks the “Comunidad de Madrid”for the tenure of a fellowship. The authors are indebted to MCYT(grant FPA2000-0950), Spain, for generous support of thisinvestigation.

Received: June 19, 2001Revised: September 10, 2001

Accepted: September 25, 2001

[1] N. R. Legge, G. Holden, H. E. Schroeder, “ThermoplasticElastomers, A Comprehensive Review”, Hanser Publishers,Munich 1987.

[2] J. L. Castles, M. A. Vallance, J. M. McKenna, S. L.Cooper, J. Polym. Sci., Polym. Phys. 1985, 23, 2119.

[3] F. J. Baltµ-Calleja, Z. Roslaniec, “Block Copolymers”,Marcel Dekker, Inc., New York 2000.

[4] I. M. Ward, D. W. Hadley, “An introduction to the mechan-ical properties of solid polymers”, Wiley, New York 1993.

[5] P. Hedvig, “Dielectric Spectroscopy of Polymers”, AdamHilger Ltd., Bristol 1977.

[6] A. R. Blythe, “Electrical Properties of Polymers”, Cam-bridge University Press, Cambridge 1979.

[7] Z. Roslaniec, H. Wojcikiewicz, Polymery (Warsaw) 1988,33, 360.

[8] Z. Roslaniec, Polymery (Warsaw) 1992, 37, 328.[9] Z. Roslaniec, T. A. Ezquerra, F. J Baltµ-Calleja, Colloid

Polym. Sci. 1995, 273, 58.[10] Z. Roslaniec, Polymer 1993, 34, 1249.[11] A. M. North, R. A. Pethrick, A. D. Wilson, Polymer 1978,

19, 923.[12] J. Runt, L. Du, L. M. Martynowicz, D. M. Brezny, M.

Mayo, Macromolecules 1989, 22, 3908.[13] F. J. Baltµ-Calleja, S. Fakirov, Z. Roslaniec, M. Krumova,

T. A. Ezquerra, D. R. Rueda, J. Macromol. Sci., Phys.1998, B37(2), 219.

[14] F. J. Baltµ-Calleja, G. G. Vonk, “X-ray scattering of syn-thetic polymers”, Elsevier Science, New York 1989.

[15] T. A. Ezquerra, G. Kremer, G. Wegner, “Progress in Elec-tromagnetic Research”, in: Dielectric Properties of Hetero-geneus Materials, Vol. 6, A. Priou, Ed., Elsevier, Amster-dam 1992, Chapter 7.

[16] Stephen Z. D. Cheng, R. Pan, B. Wunderlich, Makromol.Chem. 1988, 189, 2443.

[17] S. A. Nitzsche, Y. K. Wang, S. L. Hsu, Macromolecules1992, 25, 2397.

[18] S. Havriliak, S. Negami, Polymer 1967, 8, 161.[19] R. Bergman, F. Alvarez, A. Alegría, J. Colmenero, J.

Chem. Phys. 1998, 109, 7546.

Figure 10. SAXS diffractograms as a function of the scatteringvector q for 40/60 PBT-PC copolymer. Inset shows integratedintensity as a function of reciprocal temperature.


[20] F. Garwe, A. Schönhals, H. Lockwenz, M. Beiner, K.Schröter, E. Donth, Macromolecules 1996, 29, 247.

[21] R. Casalini, D. Fioretto, A. Livi, M. Lucchesi, P. A. Rolla,Phys. Rev. B 1997, 56, 3016.

[22] N. G. McCrum, B. E. Read, G. Williams, “Anelastic andDielectric Effects in Polymeric Solids” Dover, New York1991.

[23] R. Richert, C. A. Angell, J. Chem. Phys. 1998, 21, 9016.[24] C. A. Angell, Polymer 1997, 38, 6261.

[25] K. L. Ngai, C. M. Roland, Macromolecules 1993, 26,6824.

[26] A. Nogales, T. A. Ezquerra, F. J. Batallµn, B. Frick, E.López-Cabarcos, F. J. Baltµ-Calleja, Macromolecules1999, 32, 2301.

[27] T. A. Ezquerra, Z. Roslaniec, E. López-Cabarcos, F. J.Baltµ-Calleja, Macromolecules 1995, 28, 4516.

[28] R. W. Seymour, J. R. Overton, L. S. Corley, Macromole-cules 1975, 8, 331.

Relaxation Behavior of Poly(ester carbonate) Block Copolymer Across the Melting Region

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