Polylactide/montmorillonite nanocomposites: Structure ...ve4ep.com/a/19june/статьи 18_06_10/333.pdfHowever, dielectric properties of poly-lactide-based nanocomposites containing

EUROPEAN

European Polymer Journal 43 (2007) 2819–2835

www.elsevier.com/locate/europolj

POLYMERJOURNAL

LO

GY

Macromolecular Nanotechnology

Polylactide/montmorillonite nanocomposites:Structure, dielectric, viscoelastic and thermal properties

M. Pluta a,*, J.K. Jeszka a, G. Boiteux b

a Department of Polymer Physics, Centre of Molecular and Macromolecular Studies, Polish Academy of Sciences,

Sienkiewicza 112, 90-363 Lodz, Polandb Universite Lyon 1, Laboratoire des Materiaux Polymeres et des Biomateriaux, CNRS URA 507, F-69622 Villeurbanne, France

Received 21 December 2006; received in revised form 21 March 2007; accepted 9 April 2007Available online 27 April 2007

MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

Abstract

Polylactide-based systems composed of an organoclay (Cloisite� 30B) and/or a compatibilizer (Exxelor VA1803) pre-pared by melt blending were investigated. Two types of not compatibilized nanocomposites containing 3 wt% or 10 wt% ofthe organoclay were studied to reveal the effect of the filler concentration on the nanostructure and physical properties ofsuch systems. The 3 wt%-nanocomposite was also additionally compatibilized in order to improve the nanoclay dispersion.Neat polylactide and polylactide with the compatibilizer processed in similar conditions were used as reference samples.The X-ray investigations showed the presence of exfoliated nanostructure in 3 wt%-nanocomposite. Compatibilizationof such system noticeably enhanced the degree of exfoliation of the organoclay. Viscoelastic spectra (DMTA) showedan increase of the storage and loss moduli with the increase of the organoclay content and dispersion. Dielectric propertiesof the nanocomposites show a weak influence of the nanoclay on segmental (aS) and local (b)-relaxations in PLA, exceptfor the highest nanoclay content. Above Tg a strong increase of dc conductivity related to ionic species in the clay isobserved. It gives rise also to the Maxwell–Wagner–Sillars interfacial polarization and both real and imaginary parts ofe strongly increase. In the temperature dependence of low frequency dielectric constant and mechanical moduli(at 1 Hz) an additional maximum around 80–90 �C is observed due to cold crystallization of PLA.� 2007 Elsevier Ltd. All rights reserved.

Keywords: Polylactide; Nanocomposites; Compatibilization; Thermal properties; Viscoelastic properties; Dielectric properties

1. Introduction

Polylactide (PLA) is one of the most importantbiodegradable and biocompatible polymers in a

0014-3057/$ - see front matter � 2007 Elsevier Ltd. All rights reserveddoi:10.1016/j.eurpolymj.2007.04.009

* Corresponding author. Tel.: +48 (42) 6803237; fax: +48 (42)6847126.

E-mail address: [email protected] (M. Pluta).URL: http://www.cbmm.lodz.pl (M. Pluta).

group of degradable plastics. It can be derived fromrenewable resources, it is friendly for the environ-ment and exhibits interesting physical properties,which can be further modulated by filling withselected additives by simple blending in the moltenstate. All these features make PLA attractive alter-native for synthetic plastic materials of petrochemi-cal origin which degrade slowly (even a few hundredyears) [1]. Therefore, PLA is the subject of growing

.

mailto:[email protected]://www.cbmm.lodz.pl

2820 M. Pluta et al. / European Polymer Journal 43 (2007) 2819–2835

MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

scientific and practical interest in the last years. Inthe literature one can find numerous papers con-cerning polylactide filled with layered silicates of dif-ferent nature and properties [2–11]. If thesecomponents are compatible, a true nanocompositesystem can be formed. Such nanocomposites exhibitimproved physical properties, comparing to those ofthe unfilled polymer matrix: mechanical strength,barrier properties, thermal resistance and dimen-sional stability, even at low filler concentration(1–5 wt%). However, dielectric properties of poly-lactide-based nanocomposites containing layeredsilicate have not been explored yet.

Polylactide (Scheme 1) is a type A polyester(there is a component of the dipole moment parallelto the chain). In amorphous, racemic samples threerelaxation maxima are observed in the dielectricspectra: at ca. �80 �C (b-relaxation process, localmotions), at ca. 60 �C (aS-relaxation, dynamic glasstransition) and at ca. 85 �C (aN-relaxation, normalmode) (all at 1 Hz) [12–17]. In the samples whichcrystallise the normal mode (aN) is suppressed[15,17]. In the investigations of dynamic mechanicalproperties the maximum related to crystalline frac-tion is usually referred to as the a-relaxation processand consequently the b-process corresponds to aS(and Tg) and the c-process corresponds to the b-relaxation (local movements).

In our previous studies on PLA nanocomposites[5,11,18] we concentrated on optimization of thepreparation of the nanocomposites (nanoclaydelamination), their mechanical properties and mor-phology. This paper is aimed at the study of thephysical properties of PLA-layered silicate nano-composites with a special attention being paid tothe dielectric properties and their relationship withmechanical and thermal properties. Dielectric spec-troscopy covers broader range of frequency thandynamic mechanical measurements and is sensitiveto movements of the elements of the polymer chainwhich possess dipole moments. It is thereforeappropriate to study polylactide. Dielectric mea-surements were used to study molecular dynamicsin polymer nanocomposites based on polyiso-prene (PI) [19], polyamide, polystyrene (PS) [20],

Scheme 1. Molecular structure of PLA.

poly(methyl methacrylate) (PMMA) [21], and poly-propylene (PP) [22,23] but only in few cases fullpotential of this technique was used. In the case ofPS it was claimed that organoclay decreases dielec-tric permittivity [20,21]. In PP grafted with maleicanhydrite an increase of dielectric permittivity,attributed to Maxwell–Wagner–Sillars interfacialpolarization, was observed [23].

The nanocomposites investigated in this studycontained organomodified montmorillonite(organoclay), Cloisite� 30B, selected basing onthe previous investigations’ [11,18]. The effect ofthe organoclay concentration was studied for thenanocomposites containing 3 wt% or 10 wt% oforganoclay. To improve the filler dispersion a com-patibilizer – elastomeric ethylene copolymer func-tionalized with maleic anhydride was also used.Reference samples, PLA and PLA with the com-patibilizer, were investigated as well. The systemswere characterized using various techniques: broad-band dielectric spectroscopy (BDS), melt rheology,X-ray diffraction (XRD), dynamic periodical defor-mation (DMTA), differential scanning calorimetry(DSC).

2. Experimental

2.1. Materials

Polylactide was kindly provided by Cargill-Dow.It contains 95.9% of L-lactide and 4.1% D-lactide(residual lactide content is 0.1%). Its melt flow indexis 6.7 g/10 min (210 �C @ 2.16 kg). Organically trea-ted montmorillonite Cloisite� 30B from SouthernClay Products (Gonzales, TX) was used as a filler.This organoclay contains methyl-bis(2-hydroxy-ethyl) tallowalkyl ammonium cations (Scheme 2,where T denotes tallow consisting of �65% C18,�30% C16 and �5% C14.) Its content is 29.2 wt%(as determined with TGA, in N2 at 20 �C/min).The organic modifier increases the interlayer dis-tance (i.e. the gallery thickness) to 1.84 nm. Exxe-lorTM VA1803 (ExxonMobil Chemical) was usedfor compatibilization. Exxelor modifiers (amor-phous maleic anhydrite functionalised elastomeric

Scheme 2. Chemical formula of the organomodifier.

M. Pluta et al. / European Polymer Journal 43 (2007) 2819–2835 2821

MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

ethylene copolymers) are typically designed for useas impact modifiers, compatibilizers, couplingagents and adhesion promoters [24].

2.2. Sample preparation

After drying at 105 �C for 4 h under reducedpressure, PLA was melt-blended with other compo-nents in a counterrotating internal mixer (Brab-ender OHG, Duisburg, Germany). Unfilled PLAwas also melt-processed to have a reference mate-rial. Melt processing was carried out at the rotationspeed of 50 rpm for 20 min, in a dry nitrogen atmo-sphere to prevent thermo-oxidative degradation ofPLA. The temperature was set at 180 �C, however,it increased by about 10 �C during blending due toshearing processes. Concentration of the Cloisite�

30B was 3 wt% or 10 wt% (inorganic fraction).Two polymer systems compatibilized with 3 wt%Exxelor VA1803, 3 wt% Cloisite� 30B containingnanocomposite, and polylactide were also studied.Sample abbreviations, composition and molecularweight of PLA matrices are given in Table 1. Spec-imens for the investigations were prepared by com-pression molding (185 �C) to the thickness of100 lm for the dielectric studies and to 0.5 mm forstructural characterization using other techniques.The specimens were melt-quenched to obtain amor-phous PLA.

2.3. Characterization

Molecular weight was determined for the neatPLA and all melt-processed samples, after removingadditives, by size exclusion chromatography (SEC)method in methylene chloride as described in [18].The X-ray diffraction (XRD) technique was usedto characterize nanocomposites on the nanostruc-ture level. The measurements were performed in

Table 1Description of the investigated samples

Systems Compositions Mw Mw/Mn

PLA Processed PLA 94400 1.48N3M 97 wt%PLA + 3 wt%Cloisite� 30B 73700 1.25N10M 90 wt%PLA + 10 wt%Cloisite� 30B 89000 2.00PLA-C 97 wt%PLA + 3 wt%Exxelor

VA1803118000 1.68

N3M-C 94 wt%PLA + 3 wt%Cloisite�

30B + 3 wt%Exxelor VA180372000 1.25

Molecular masses of PLA determined using SEC in methylenechloride (neat polylactide has Mw = 126000, Mw/Mn = 1.48).

the transmission mode (coupled h/2h) in the 1.2�to 8� range of 2h. A wide-angle goniometer was cou-pled to a sealed-tube source of filtered CuKa radia-tion, operating at 50 kV and 30 mA (PhilipsPW3830). The slit system enabled collection of dif-fracted beam with the divergence angle of less than0.05�.

Rheological properties were studied with anadvanced research-grade rheometer (ARES; Rheo-metric Scientific) using a parallel plate geometry(diameter 25 mm). This technique is sensitive to fol-low the dispersion of the filler particles in polymermatrix [18,25]. The compression molded sampleswere placed between hot plates and stabilized at170 �C for about 5 min before the measurement.Dynamic frequency sweep was performed in theregion of linear viscoelastic response (LVR), witha strain 0.8% for the nanocomposites and 2% forsamples not containing montmorillonite, startingfrom high frequency, 512 rad/s down to 0.02 rad/s.The LVR region was experimentally establishedbecause it is dependent on the sample compositionand degree of dispersion of the nanoclay [18].Experimental data were related to the actual gapvalue (�0.9 mm).

Thermal properties of the samples were investi-gated with a DSC 2920, TA Instruments, undernitrogen atmosphere. The crystallization behaviorof the PLA matrix from the initially glassy, amor-phous state was investigated at a heating rate of3 �C/min, following the melt-quenching scan.

Dynamic mechanical properties of 0.5 mm thicksamples were measured with an MkIII DMTAapparatus (Rheometric Scientific, Inc.) in a dual-cantiliver bending mode. The dynamic storage andloss moduli (E 0 and E00) were determined at a con-stant frequency of 1 Hz as a function of temperaturefrom �90 �C to 140 �C at a heating rate of 2 �C/min.

The dielectric properties were studied in parallelplate geometry. The samples were provided withevaporated circular Al electrodes (30 mm in diame-ter). Before Al deposition the specimens were driedovernight in vacuum at 40 �C (i.e. below Tg to avoidthe cold crystallization of the initially amorphousPLA matrix). The complex dielectric function wasmeasured at constant temperatures in the frequencyrange from 10�1 Hz to 106 Hz by a NovocontrolConcept 40 a-analyzer interfaced to the sample bya broadband dielectric converter (BDC, Novocon-trol) or a lock-in amplifier (Stanford Research810). The temperature of the sample was varied

Fig. 1. X-ray diffractograms recorded for the nanocompositesN3M, N10M and N3M-C. Reference diffractograms of theircomponents: unfilled PLA, Cloisite� 30B and 3 wt% ExxelorVA1803 nanocomposite (NEx3M) are shown for comparison.The diffractograms are vertically shifted for clarity.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

from �100 �C to +120 �C in steps of 3 �C and wascontrolled with a stability of DT = 0.1 �C (Novo-control Quatro system controller BDS 1330). Aver-age heating rate was ca. 0.5 �C/min.

3. Results and discussion

3.1. Size exclusion chromatography (SEC)

The melt processing leads to some decrease ofmolar weight (Mw) due to degradation of the poly-lactide (Table 1) although protective nitrogen atmo-sphere has been used. This is due to susceptibility ofPLA to degradation in the molten state observedalso by other groups [26,27]. It can be seen thatthe Mw of the unfilled PLA is decreased by �25%,while that of the compatibilized PLA (PLA-C) onlyby �6% as compared to the neat PLA. It suggestssome stabilizing effect of the compatibilizer towardsPLA degradation or some chemical interactionsbetween these components. The melt processing inthe presence of the nanoclay contributes to furtherdecrease of the Mw, even if the compatibilizer wasused.

3.2. XRD analysis

Fig. 1 shows X-ray diffractograms of the nano-composites N3M, N10M and N3M-C. Diffracto-grams of unfilled PLA, of Cloisite� 30B powderand of Exxelor VA1803 filled with 3 wt% Cloisite�

30B (sample NEx3M) are also shown for compari-son. The unfilled PLA sample shows typical back-ground scattering with an intensity increasingbelow 1.2�. The diffractogram for Cloisite� 30Bhas a distinct maximum around 2h = 4.7�(d001 = 1.8 nm). This maximum is not observed forthe nanocomposites. The nanocomposites revealonly very small bulge around 2h � 2.5� (d001 =3.5 nm). This feature is characteristic of a good dis-persion of the organoclay, achieved by an intercala-tion followed by tactoids formation and thenexfoliation of the nanoplatelets in the PLA matrix.These features were observed by TEM for systemssimilar to N3M, discussed in [18]. Compatibilizationof 3 wt% nanocomposite (sample N3M-C) results ina further increase of the organoclay dispersion asthe bulge in the diffractogram is decreased and thescattering at the lowest 2h is stronger than that forthe not compatibilized counterpart (N3M). In orderto investigate the interaction between the organo-clay and compatibilizer, an additional nanocompos-

ite (Exxelor VA1803 + 3 wt% Cloisite� 30B) wasalso prepared. Diffractogram for this system, in con-trast to the other ones, reveals a strong intensityincrease at low 2h (typical for the exfoliated struc-tures) and a small diffraction maximum at2h � 6.0�, i.e. larger than for the Cloisite� 30B(2h = 4.7�). It corresponds to the decreased inter-layer thickness (to about 1.5 nm). This indicates acollapse of some fraction of the nanoplatelets inthe assembled regions. Diffractogram of the unfilledcompatibilizer (not shown in Fig. 1) is similar tothat of the unfilled PLA sample.

The shape of the diffractogram obtained for the10%-nanocomposite (sample N10M) indicates alsoa high degree of the organoclay dispersion. How-ever, this diffractogram has a slightly larger bulgearound 2.5� than the less filled N3M sample. Thiscan mean that the content of the organoclay in theN10M nanocomposite is too high to obtain homo-geneous dispersion of the organoclay and full sepa-ration of the silicate nanoplateles in the PLAmatrix. Probably in this nanocomposite system con-centration of the organoclay is above a percolationlevel and some montmotrillonite network is formed.

3.3. Rheological properties

Polylactide filled with the organoclay and/orcompatibilizer exhibits different rheological proper-ties as compared with neat PLA. This is illustrated


LO

GY

in Fig. 2 showing dependencies of log G 0, log G00 andlog g* vs. log x for the considered samples. Theunfilled PLA is featured by a typical increase ofthe G 0 and G00 with deformation frequency x(Fig. 2a). At lower x the G 0 is smaller than G00, indi-cating a classical liquid-like behavior of the moltenpolymer. The difference between G 0 and G00

decreases gradually with raising deformation fre-quency. Finally, the curves cross at x � 115 rad/s,and above this frequency G 0 is higher than G00

revealing domination of the solid response. It isworth to notice, that at the highest frequencies allthe samples exhibit comparable values of G 0 andcomparable values of G00, with G 0 > G00. This indi-cates that at these frequencies the response of thesamples is weakly affected by the nanostructureand determined mainly by the properties of thepolymer matrix. Similar relations at high frequen-

Fig. 2. Comparison of frequency dependencies of the storage modulusPLA, compatibilized PLA (PLA-C), nanocomposites (N3M and N10M)the same for all the samples to illustrate/highlight changes in the rheol

cies have been found for other nanocomposites(e.g. [28]).

Polylactide with the compatibilizer (sample PLA-C) exhibits rheological behavior similar to that ofunfilled PLA. For the former sample the differencebetween G 0 and G00 is somewhat smaller, and thecrossover point appears at slightly lower x �100 rad/s (Fig. 2b). The complex viscosity of theunfilled PLA and the compatibilized sample exhibita Newtonian behavior (g* is independent of fre-quency) with values about 4000 Pa s and 5000Pa s, respectively, then, above 10 rad/s a shear-thin-ning takes place. This behavior is typical forhomopolymers.

Samples containing the organoclay – nanocom-posites N3M, N3M-C and N10M, exhibit morepronounced changes of the rheological parameters.One can see that the higher organoclay dispersion

(G 0), loss modulus (G00) and complex viscosity (g*) of the unfilledand compatibilized nanocomposite (N3M-C). The vertical scale isogical behavior.

MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

Fig. 3. DSC heating thermograms recorded for the samples at aheating rate of 3 �C/min. The scans were performed directly aftermelt-quenching at a cooling rate of 3 �C/min. The curves arenormalized to the mass of PLA.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

(samples N3M and N3M-C) and the higher organo-clay concentration (samples N3M and N10M) thelarger increase of G 0 and G00 at low frequencies isobserved (compare Fig. 2c with Fig. 2e and d withFig. 2e, respectively). Moreover, for the 3 wt%-nanocomposites (N3M and N3M-C) two crossoverpoints are observed: x1 � 0.02 rad/s and x2 �180 rad/s for N3M; x1 � 50 rad/s and x2 � 150rad/s for N3M-C, respectively (arrows in Fig. 2.Between x1 and x2 the behavior of the samples isliquid-like (G 0 < G00). Below x1, G 0 exceeds G00 andshows a plateau which indicates the so called‘‘pseudo-solid-like behavior’’ of the material atlow deformation frequencies [28].

The plateau region is larger for the nanocompos-ite N3M-C than for the nanocomposite N3M due tobetter dispersion of the organoclay. The dispersiondependent plateau/pseudo-solid-like response wasalso observed for similar PLA/organoclay systemsin which exfoliation degree was increased by prolon-gation of the compounding time [18]. For the nano-composite N10M with higher organoclay contentthe plateau region is enlarged which leads to therelation G 0 > G00 (‘‘pseudo-solid-like’’ response) inthe whole frequency range.

Viscosity g* of all the nanocomposites (N3M,N3M-C, N10M) does not show a Newtonian behav-ior in the deformation frequency range studied. Theviscosity decreases with increasing frequency, show-ing a shear-thinning response, the more pronouncedthe higher is dispersion and/or concentration of theorganoclay (the initial g* slope is: �1.02 for N10M,�0.98 for N3M-C and �0.95 for N3M, while it isclose to zero for PLA and PLA-C samples). Theincrease of the moduli and of the viscosity, seen dis-tinctly at lower frequency range, reflects a reinforce-ment of the molten polymer matrix by the filler. Thereinforcing effect results from the interactionsbetween the components due to hydrogen bondingof hydroxyl groups in the organic ‘‘surfactant’’ inthe organoclay and carbonyl groups of PLA chainsegments. The interactions are the stronger the lar-ger is the interface area, thus being related to higherdispersion and/or concentration of the organoclayin the PLA matrix. Similar concentration/dispersiondependent rheological spectra were observed forpolystyrene-clay nanocomposites [29].

3.4. Calorimetric characterization

Fig. 3 shows DSC heating thermograms of sam-ples of initially amorphous PLA and of the nano-

composites. The measurements were performedimmediately after melt-quenching scans, so the sam-ples have the same thermal history without agingcycle. The samples reveal the following thermalevents with increasing temperature: glass–rubbertransition (at Tg), cold crystallization process (char-acterized by Tcc and DHcc) and the melting processwith two components (characterized by Tm1, DHm1and Tm2, DHm2). Tcc and DHcc denote crystalliza-tion temperature and enthalpy, respectively. Tm1,and Tm2, DHm1 and DHm2 denote the temperaturesof the melting peaks and the corresponding meltingenthalpies. Comparing thermograms and calorimet-ric parameters collected in Table 2 one can see thatfilling of PLA modifies, more or less, the individualthermal processes. Tg is somewhat decreased, from57.7 �C for PLA to 56.9 �C and to 53.0 �C for thenanocomposites N3M and N10M, respectively. ThisTg decrease can be explained by a plasticizing effectof the clay surfactant (�30% in the organoclay).Compatibilized PLA (sample PLA-C) exhibitTg = 58.6 �C, higher than unfilled PLA (57.7 �C),suggesting some reinforcement of the polylactidematrix by the compatibilizer. The compatibilizednanocomposite N3M-C shows slightly lowerTg = 57.1 �C than PLA (DTg = �0.6 �C) indicatingsome balancing effect of the compatibilizer and thenanoclay surfactant.

The influence of Cloisite� 30B content on thecold crystallization of the PLA matrix is discernible.The Tcc decreases from 104.6 �C for unfilled PLA to98.3 �C for N3M and to 92.0 �C for N10M, while

Table 2Calorimetric parameters of the investigated materials (heating rate 3 �C/min)

Sample Tg (�C) Cold crystallization Melting

Tcc (�C) DHcc (J/g) Tm1 (�C) DHm1 (J/g) Tm2 (�C) DHm2 (J/g) DHmTotal (J/g)

PLA 57.7 104.6 29.1 148.8 12.8 156.8 16.9 29.0N3M 56.9 98.3 29.2 147.3 8.1 155.4 21.5 29.6N10M 53.0 92.0 31.4 – – 154.0 33.4 33.4PLA-C 58.6 105.2 29.0 149.4 14.7 156.8 14.7 29.4N3M-C 57.1 107.8 30.3 149.6 15.6 156.3 15.9 31.5


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

the crystallization enthalpies are comparable(�30 J/g) (Table 2). The decrease of the cold crystal-lization temperature can be ascribed to a nucleatingeffect of the filler. The concentration dependentnucleating effect has also been observed in similarcompositions of PLA and Cloisite� 30B [10]. More-over, the organoclay surfactant can also contributeto lowering of the cold crystallization temperature,because of its plasticizing effect. The effect of plast-icizers on the crystallization of PLA is discussed indetail in other papers [30–32]. Compatibilized sam-ples PLA-C and N3M-C reveal the cold crystalliza-tion at higher temperatures than the PLA sample,with an exothermic peaks at 105.2 �C and107.8 �C, respectively. The crystallization enthalpyis equal to �30 J/g – like those of the other samplesdiscussed (Table 2). These results indicate that thecompatibilizer reduces mobility of PLA macromol-ecules. This effect is stronger in the compatibilizednanocomposite (sample N3M-C) due to improveddispersion of the organoclay (accordingly to X-rayand rheological data). The influence of the organo-clay dispersion on crystallization of PLA was dis-cussed in detail in [18].

Melting of the samples is reflected by an endo-therm with two components, at Tm1 (around148.5 ± 1.2 �C) and at Tm2 (around 155.5 ±1.3 �C). Two melting peaks reflect the process ofmelting of crystallites having different sizes and/orperfection of ordering. In the case of non-compati-bilized samples (PLA, N3M and N10M) the sizeand temperature position of the low temperaturemelting component decreases with the organoclaycontent, which corresponds to a decrease the coldcrystallization temperature (Tcc) of these samples.For the 10 wt% nanocomposite (N10M) the meltingcomponent at Tm1 is seen as a low temperatureshoulder of the main melting peak at Tm2. Thismeans, the lower is Tcc the lower is the contributionof the melting process at Tm1, while that at Tm2relatively increases. For the N10M sample the

total melting enthalpy DHmTotal � 33.0 J/g, i.e. it ishigher than the crystallization enthalpy DHcc �31.4 J/g (Table 2). This is connected with the factthat smaller, less perfect crystallites, formed at lowerTcc undergo during the measurement a furtherrecrystallization at higher temperature, above Tcc,contributing to the melting enthalphy at Tm2. Therecrystallization process was also revealed by anon-reversing signal of the TMDSC (not shownhere). Compatibilized samples, PLA and the nano-composite N3M-C, are featured by a slightly largerlow temperature melting component (DHcc �14.6 J/g) which appears at a higher Tm1 � 149.5 �C,than for the unfilled PLA (DHcc = 12.8 J/g, Tm1 �148.8 �C). This is due to the cold crystallization athigher temperature in the case of the compatibilizedsamples (Fig. 3). Therefore more stable crystallitesare formed, which, give raise only to one meltingprocess on further heating, without recrystallizationcontribution (as indicated by the non-reversing sig-nal of the TMDSC – not shown here).

It is interesting to note that the crystallizationenthalpy (DHcc) of PLA is similar for all the samples(within an experimental error) and it is close to thetotal melting enthalpy (DHmTotal), (Table 2). Thisindicates that the additives used (organoclay, com-patibilizer) does not affect significantly the finalcrystallinity of the PLA matrix developed duringheating, but have some influence on the crystal sizeand perfection (as Tcc changes).

3.5. Viscoelastic properties

Fig. 4a shows the storage modulus (E 0) of the dis-cussed samples determined by dynamic mechanicalthermal analysis (DMTA) at 1 Hz, plotted vs. tem-perature. For all the samples the following charac-teristic E 0 changes with raising temperature areobserved: a gradual decrease in the region �90 �Cto 50 �C, a rapid drop below 55–60 �C due to theglass–rubber transition (Tg), an increase in the cold

Fig. 4. Temperature dependencies of E 0 (a) and E00 (b) recordedfor the investigated systems of different composition.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

crystallization range (around 100 �C) owing to rein-forcing by the crystallites being formed, and then adecrease as a result of the pre-melting process. It isworth to mention that below Tg E

0 slightlyincreases, before dropping down. This effectdecreases with filling and it was observed also forother amorphous PLA samples [33]. An increaseof E 0 prior to Tg can be ascribed to a relaxation ofinternal stresses, frozen in during melt-quenchingof the material.

Filling PLA with the organoclay and/or the com-patibilizer, increases E 0 of the system in the wholetemperature range. Below the glass transition tem-perature E 0 increases in the order PLA < PLA-C � N3M-C < N3M < N10M. It means, that E 0increases with the organoclay content, but it is alsoenhanced by the compatibilizer. This indicates thatthe compatibilizer (which is an elastomer with Tg

around �57 �C) being dispersed in the PLA matrixreinforces it slightly ðE0PLA-C > E0PLAÞ. This suggestssome interaction of both components. The presenceof the organoclay in the compatibilized system(N3M-C) results also in the increase of E 0. The max-imum around 100 �C in E 0 (and E00) correspond tothe cold crystallization, revealed also by the DSCscans (Fig. 3).

The dependencies of E00 on temperature areshown in Fig. 4b. The filled systems show theincreased mechanical loss in comparison with theunfilled PLA. A small maximum is observed around�60 �C, probably related to local movements isreferred to as a c-relaxation in mechanical measure-ments and a b-relaxation in the dielectric conven-tion. The maximum in E00 about 60 �C correspondsto segmental relaxation (mechanical b-relaxationor aS in the dielectric convention), related to theglass–rubber transition in PLA. The maximum ofthe E00 around 100 �C (a-process according to themechanical convention) reflects an increase of themechanical loss due to the cold crystallization. Itsintensity increases with the increasing organoclaycontent.

3.6. Broadband dielectric spectroscopy

Fig. 5 shows frequency dependencies of the imag-inary component of the complex dielectric functione00 for initially amorphous samples of unfilled PLAand the compatibilized nanocomposite (N3M-C)at selected temperatures. The maximum, clearlyobserved above 60 �C for PLA and also for thenanocomposite, is ascribed to the aS-relaxation[12–17]. Above 80 �C the crystallization of initiallyamorphous PLA begins, which leads to a ca. two-fold decrease of the aS maximum at 90 �C and101 �C (observed also by other groups [14,17]).The temperature range in which PLA crystallizationtakes place in dielectric measurements is differentfrom that observed in DSC (Fig. 3) because theheating rate in BBDS experiment is much lower(compare also Figs. 9a and 10a below).

The most important difference in the dielectricspectrum of the nanocomposite, as compared withthe unfilled PLA, is a strong increase of e00 at lowfrequencies, the more important the higher is thetemperature. Its frequency dependence and compar-ison with the results obtained for pure Closite� 30Bsample, show that is should be ascribed to directcurrent (dc) conductivity and interfacial polariza-tion (Maxwell–Wagner–Sillars effect – MWS), most

Fig. 5. Dielectric spectra of PLA (a) and of the nanocompositeN3M-C (b) at several temperatures around and above Tg.

Fig. 6. Comparison of e00 spectra of the investigated systems at80 �C. Solid lines represent fits using HN functions for aS-relaxation and MWS process and rDC � x�0.85. Thin solid andbroken lines show fit components for PLA and N3M,respectively.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

probably related to the presence of ionic speciesused to intercalate the montmorillonite during itsorganomodification (vide infra). Similar behaviorwas observed for the non-compatibilized nanocom-posite (N3M).

The frequency dependence of the imaginary partof dielectric permittivity (e*) in the investigated sys-tems can be fitted using a superposition of Haviri-lak–Negami (HN) empirical functions [34]

e� � e1 ¼De

ð1þ ðixsHNÞaÞcð1Þ

where e1 is the real part of e* for frequencies muchhigher than the maximum frequency of a given pro-cess, De is relaxation strength, x denote frequency,sHN is so called Havirilak–Negami relaxation timeand exponents a and c represent broadening of thedistribution of the relaxation times.

A contribution of the dc conductivity and elec-trode polarization can be taken into account assum-ing its frequency dependence in a usual form [34]

e00 ¼ rdce0xm

where e0 denotes the dielectric permittivity of freespace and rdc – the dc conductivity (at frequencyapproaching zero). The exponent m should be equalto 1 in the case of pure electronic conductivity butusually is smaller than one because electrode andinterfacial polarization effects also come into play.In our systems at 80 �C m was equal to 0.8 (alsofor pressed nanoclay pellets) and increased withtemperature up to 0.95 at 120 �C.

In Fig. 6 e00 spectra vs. frequency for all the sam-ples at 80 �C (>Tg) are compared. This temperatureallows to observe all relevant processes in practi-cally still amorphous samples. The lines goingthrough the experimental points represent fittingresults and thin lines show the components contrib-uting to the fits for PLA and N3M. The componentsof other fittings were omitted for clarity. It can beseen that the e00 spectra of PLA and PLA-C can bewell fitted using a superposition of two main pro-cesses: aS-relaxation with the maximum around104 Hz and rdc contribution at low frequencies.An additional weak process around 3 Hz may berelated to the normal mode (aN) (which shouldappear around this temperature [12,16]). Its smallintensity may be caused by crystallization of PLAwhich begins below 80 �C (lower heating rate as

Fig. 7. Comparison of the tand spectra of PLA, PLA-C and thenanocomposites (N3M, N3M-C and N10M) recorded at �70 �C(a) and tand spectra of the nanocomposites (N3M, N10M) atvarious temperatures (b) Corresponding spectra of montmoril-lonite at �10 �C, �40 �C and �70 �C (divided by 10) are shownfor comparison (lines with small symbols).


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

compared with DSC) and appearing crystallites hin-der chain relaxation.

The intensity of the aS process is weakly depen-dent on sample composition, the contribution ofrdc increases in the nanocomposites. In the nano-composites, in spite of the high conductivity, wecan clearly distinguish an additional strong maxi-mum in e00. This maximum is observed also in otherrepresentations which reduce the effect of dc con-ductivity – modulus representation and de 0/d(log f)(not shown). It can be assigned to the Maxwell–Wagner–Sillars effect as e 0 also strongly increasesat low frequencies in this temperature range (videinfra, Figs. 9 and 10). MWS process was alsoreported in polymer-silicate nanocomposites basedon polyisoprene [19] and PP [23]. The fitting param-eters are collected in Table 3.

The distribution of the aS-process is similar asfound by other authors (a-coefficient equal to 0.55[12] or 0.66–0.76 [12,16] and c = 0.65–0.8 [16]). Inthe case of the nanocomposites, especially N10M,the fitting of the aS-process is not so reliable becauseslightly different sets of parameters for the overlap-ping processes give similar errors.

The relaxation time of the maximum assigned tothe Maxwell–Wagner–Sillars polarization decreaseswith increasing conductivity of the nanocompositesas expected. Surprisingly its intensity slightlydecreases for higher organoclay loading, whichcan indicate somewhat worse dispersion and separa-tion of the filler nanoplatelets in the polymer matrix.DC conductivity of the nanocomposites is roughlyproportional to the organoclay content and hasthe same frequency dependence as that of neat Clo-site� 30B sample (pressed discs) at low frequencies(rdc = 1.5 · 10�8 Ohm/m, detailed data not shown).

The influence of montmorillonite on tand at alow temperature (�70 �C) is shown in Fig. 7a. Atthis temperature the b-relaxation maximum is wellseen. This process is attributed to local movementsof small chain elements (carboxyl groups). The spec-

Table 3H–N parameters and rdc used for fitting of e00 spectra presented in Fig

aS

a c s De

PLA 0.62 0.65 2.4 · 10�4 2.55PLA-C 0.62 0.65 2.8 · 10�4 2.2N3M 0.62 0.56 2.8 · 10�4 2.55N3M-C 0.62 0.56 2.8 · 10�4 2.5N10M 0.57 0.45 2.8 · 10�4 3.1

tra of the nanocomposites with 3% of organoclay(N3M and N3M-C) are similar to those of theunfilled PLA and PLA-C, respectively. Theyappear, however, at somewhat higher frequencieswhich indicates that the presence of the additives

. 6 using Eq. (1)

MWS rdc

a c s De

0.56 1 0.23 0.25 3.35 · 10�12

0.62 1 0.9 0.23 1.2 · 10�11

0.66 1 0.15 3.25 3.25 · 10�10

0.66 1 0.15 3.3 3.2 · 10�10

0.7 1 0.02 2.5 1.3 · 10�9


only slightly increases local mobility of PLA chains.The b-relaxation is highly distributed (a coefficientin Eq. (1) equal to ca. 0.3) and symmetric (c = 1).

By contrast the 10 wt%-nanocomposite (N10M)has a different tand spectrum. It shows considerablyenhanced dielectric loss with a pronounced tandmaxima at frequency regions below 2 Hz and above106 Hz which mask much weaker b-relaxation peakwith maximum around 102–103 Hz. The position ofthe low frequency tand maximum is practically tem-perature-independent, while the high frequency one

Fig. 8. The temperature dependence of the r

moves, for both N10M and N3M, to higher fre-quencies as the temperature is increased (Fig. 7a).

These strong maxima should be related to thepresence of low molecular weight species (organicnanoclay modifier and possibly strongly bound,not removed H2O). The remarkable differencebetween N3M and N10M and Closite� 30B suggestthat a particular percolative filler nanostructure isformed in the N10M nanocomposite. The fact thatneither of these relaxation processes is observed inneat Closite� 30B (Fig. 7b) proves good dispersion

elaxation times of a- and b-relaxation.

MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

Fig. 9. Temperature dependence of dielectric permittivity e 0 (a)and tand (b) at 1 kHz for PLA and for the nanocomposites(initially amorphous – 1st run and crystalline samples 2nd run).


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

of the nanofiller in spite of its high concentration, inagreement with XRD data (cf. Fig. 1). The percola-tion paths and/or nanostructure are considerablyaffected by heating and crystallization as thespectrum of the crystalline N10M sample (secondrun – not shown) is qualitatively different from thatof the amorphous sample (but also different fromthat of unfilled, semicrystalline PLA).

Fig. 8 shows the temperature dependence of therelaxation times for the aS and b processes deter-mined from frequencies of maximum loss

s ¼ 12pfmax

For the b-maximum the temperature dependenceobeys the Arrhenius law with an activation energy9.9 kcal/mol (41.5 · 103 J/K/mol) in reasonableagreement with the activation energy 10.5 kcal/molgiven in [12]. It can also be seen that neither theorganoclay nor the compatibilizer have significanteffect on the activation energy, only on the pre-exponential factor is decreased.

The temperature dependence of the aS-relaxationfor PLA and the nanocomposites obeys Vogel–Ful-cher–Tammann law

s ¼ s0 expB

T � T 0

� �

with s0 = 13, B = 800 K and T0 = 302 K in reason-able agreement with previous results s0 = 12.5–14.3[12] or 11.6–12.6 [16], B = 520–760 K [12] or 259–289 K [12] and 452–564 [16] and T0 = 277–285 [16](for materials of different molecular weight).

In the case of the composites (especially N10M)the relaxation times at higher temperatures cannotbe unambiguously determined, because of overlap-ping of Maxwell–Wagner–Sillars maximum and dcconductivity.

Fig. 9 shows a comparison of e 0 and tand vs. tem-perature above 0 �C for unfilled PLA and the nano-composites at frequency 1 kHz. The spectra forinitially amorphous and for semicrystalline samples(the second run on the same sample after heating to120 �C and annealing for 1 h) are presented to showthe effect of crystallisation. The maximum of e 0

around 80 �C, observed in the initially amorphoussamples, appears as a result of the cold crystalliza-tion and it is discussed in more detail below. It isnot observed in the semicrystalline samples in whichcold crystallization does not occur any more. Amonotonic increase of e 0 above 50 �C in the nano-composites is due to the Maxwell–Wagner effect

related with increasing dc conductivity or possiblyalso with electrode polarization.

Tand shows maxima for both crystalline andamorphous samples. In crystalline samples theirintensity is decreased and it is shifted by ca. 5 �Cto higher temperatures. This is due to limited move-ments of the chains partially incorporated in thecrystallites, in agreement with [15,17]. The maxi-mum for PLA with the compatibilizer (PLA-C) isalso slightly shifted to higher temperatures. Crystal-lization does not have a significant effect on anincrease of tand at higher temperatures due to dcconductivity.

Fig. 10 shows a comparison of the temperaturedependence of e 0 and tand at fixed frequency 1 Hz(the same as used in the mechanical measurements).At lower frequency the aS-relaxation is observed at

Fig. 10. Temperature dependence of dielectric permittivity(a) and tand (b) at 1 Hz for PLA and the nanocomposites.

Fig. 11. Temperature dependence of tand for PLA-C (a) andN3M-C (b) at different frequencies.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

lower temperatures (c.f. Fig. 9) and it is clearly seenthat the dielectric spectra of the amorphous samplesin the range 50–90 �C are influenced by the coldcrystallization. The increase at high temperaturesis more important at 1 Hz (please mind the logarith-mic scale) which confirms that it is related mostly tothe Maxwell–Wagner–Sillars effect and dc conduc-tivity. The compatibilizer does not show any signif-icant influence on the dielectric properties (noadditional relaxation process is observed) but itspresence slows down the crystallization in agree-ment with the results obtained by DSC (Fig. 3)and DMTA (Fig. 4).

The plot of tand vs. temperature (Fig. 9b) showsclearly the maximum corresponding to the aS-relax-ation. Its position �60 �C is closer to Tg determined

from the DSC data (Table 2), because of lower fre-quency. Its intensity also decreases after crystalliza-tion. The second maximum, which is not observedin the crystalline samples (data not shown for clar-ity) is a result of fast crystallization (the relaxationstrength decreases). Its maximum temperature isslightly lower as compared with the mechanical data(cf. Fig. 4b).

It can be seen that the presence of the organoclayresults in a strong increase of the dielectric loss at1 Hz already above the room temperature. The aS-relaxation is in the same position (around 60 �C)but it is much broader. This effect is observed alsoin the mechanical relaxation. It should be relatedin part to the organoclay itself, which exhibits a

Table 4Comparison of the influence of the nanofiller and compatibilizeron the maximum temperature of segmental relaxation (dynamicglass transition) as determined using dynamic mechanical anal-ysis and dielectric spectroscopy and Tg determined by DSC (all in�C)

Sample DMA(tand at1 Hz)

BBDS(tand at1 Hz)

BBDS(tand at1 kHz)

DSC(inflectionpoint)

PLA 60.3 61.2 74.3 57.7N3M 58.9 60.4 73.2 56.9N10M 57.3 58.5 72.0 53.0PLA-C 59.0 61.2 74.3 58.6N3M-C 59.3 59.7 73.5 57.1


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

strong loss in this temperature range. However, theobserved spectrum is not a simple superposition.Dipolar species present in the organoclay (amine)are now in a different medium (as compared withneat Closite� 30B) and their relaxation is controlledto some extend by the mobility of adjacent polymerchains. The maximum attributed to the crystalliza-tion of PLA around 83–90 �C is also observed inthe nanocomposites and it is much stronger (mindthe logarithmic scale), as the crystallizationdecreases also ion mobility and thus rdc). It shouldalso be noted that the influence of the compatibilizeron the position of this maximum is similar as in thecase of PLA. These observations confirm an interac-tion between the compatibilizer and the polymermatrix.

In Fig. 11 one can see an evolution of the temper-ature dependence of tan d for PLA-C and the nano-composite at different frequencies. The maximumrelated to the aS-process shifts towards higher tem-peratures as the frequency increases while the max-imum related to crystallization is in the sameposition. Thus, both maxima merge together for fre-quencies higher than 100 Hz.

4. Discussion

Various experimental techniques used enabled usto address the problem of the influence of the nano-filler on physical properties of the polymer matrix,in particular on chain relaxation. Mechanical prop-erties and viscosity are sensitive primarily to largescale chain movements while dielectric spectroscopyprovide information mainly on local dipolar groupsfluctuations. The presence of rigid silicate planes inthe nanocomposite must restrict movements of thesurrounding polymer chains. The better is the dis-persion of the clay (delamination) and the higheris the clay content the stronger effect should beobserved.

The most significant influence is expected for thelarge scale movements – on the rheological proper-ties in the terminal region and on the normal mode(aN relaxation). We indeed observed significant stiff-ening of the composite as compared with neatpolymer matrix in the dynamic mechanical measure-ments (Fig. 4) but the most evident effect is thesolid-like behavior of the complex viscosity at lowfrequencies (Fig. 2) (in spite of some lowering ofthe molecular mass of PLA – Table 1). It was notpossible to observe a shift of the aN relaxation inthe dielectric spectra in our systems because the

crystallization of PLA takes place in the same tem-perature range (above 80 �C), in which this modewas observed [12] and it is suppressed. In the nano-composites there is also a strong MWS maximumobserved in this temperature range.

The smaller is the scale of molecular motions theless evident is the effect of space confinement. Wemust also remember that low molecular weightorganic molecules with long alkyl chains, presenton the clay sheets, increase local mobility at theclay/polymer interface. Thus, in the case of segmen-tal relaxation related to the glass transition (termedas in discussions of dielectric properties and b inmechanical properties) the situation is not so obvi-ous. We observed a small effect of the filler on thetemperature of this relaxation as measured usingdifferent methods. The data are collected in Table4. It can be seen that the values obtained for differ-ent materials do not differ much. DSC gives thevalues lower by ca. 2 �C because they correspondto the frequency of ca. 10�2 Hz. Mechanical anddielectric measurements at 1 Hz give similar results.In the case of BBDS an error of the results obtainedfor the nanocomposites (especially for N10M) at1 Hz is of the order of 1–2 �C because of the strongeffects related to conductivity at low frequencieswhich make fitting ambiguous. However, the sametrend is observed at higher frequencies (where inturn some error might be due to an overlap withPLA crystallization). In summary we can concludethat the presence of the nanofiller shifts the segmen-tal relaxation to lower temperatures by ca. 1 �C at3% filler content and by ca. 3 �C at 10 wt% content.The effect of the compatibilizer is small and notobvious as different methods give contradictoryresults.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

Space confinement can increase or decrease chainmobility and Tg, depending on interactions with thesolid surface as discussed extensively e.g. in [34] andreferences therein. However, neither the experi-ments on confinement in nanopores nor those onspin-coated thin films, strictly correspond to the sit-uation in the investigated layered nanocomposites.It should be noted that taking into account specificsurface of the clay (750 m2/g) and its content, evenat perfect delamination an average distance betweenthe nanoplates can be estimated to be at most ca70 nm or 20 nm for N3M and N10M respectively.It is much bigger than estimated radius of gyration,so a small effect is not surprising, especially takinginto account that in the real system the dispersionand delamination are not perfect. Increased hin-drance of the chain segment motions should alsoresult in a broadening of the relaxation maximumwhich indeed was also observed in the temperaturedependencies (Fig. 11) and as an increase of Havir-ilak–Negami exponents (a coefficient equal to ca.0.6 for PLA and PLA-C, 0.55 for N3M and N3M-C and 0.53 for N10M. The effect for the nanocom-posites is, however, again obscured by the conduc-tivity, so it is difficult to separate the symmetricand asymmetric broadening effects.

The short range local relaxations giving rise tothe b-process in the dielectric spectra are notexpected to be influenced by the nanofiller andindeed only a small shift was observed. The mostsignificant is the effect of the compatibilizer inPLA (but not in N3M-C where it is probablylocated at the nanofiller/PLA interface). The b-pro-cess is practically not visible in mechanical relaxa-tion except for the systems with the compatibilizer.It is not clear if a small maximum observed at�61 �C and �66 in PLA-C and N3M-C respectivelyis due to Tg of the compatibilizer or to b-process inPLA. Evidently in the dielectric spectra it is due tob-relaxation in all cases, as evidenced by the temper-ature dependencies of the relaxation times (Fig. 8b).At low temperatures the most striking is the strongincrease of tand in N10M while in N3M no signifi-cant difference as compared with PLA is found. Weattribute it to a percolation of the nanoclay network(responsible also for the stiffening observed in rheo-logical measurements – Fig. 2), which facilitates dif-fusion of low molecular weigh ions on longerdistances along the filler/PLA interface.

As expected for the two-phase system a maxi-mum due to the Maxwell–Wagner effect wasobserved in the dielectric spectra of the nanocom-

posites. Similar effect was observed also in nano-composites based on other polymers: polyisoprene[19] polyamide-6 [35,36] poly(propylene-graft-maleic anhydride) [23] relation of its parameters todelamination/exfoliation is however not obvious.

5. Conclusions

The X-ray investigations exhibited exfoliatednanostructure in 3 wt%-nanocomposite. Compati-bilization noticeably enhanced the degree of exfolia-tion of the organoclay due to combined interactionsof the organoclay surfactant with polylactide chainsand maleic anhydrite groups of the compatibilizer.In the 10 wt%-nanocomposite mixed – intercalatedand exfoliated nanostructures were detected due tohigh concentration of the filler. Rheological andmechanical properties suggest that a sort of silicatenetwork was formed.

Using different techniques we were able to studythe effect of the nanoclay on relaxation processes ina broad range of temperature and frequency. Rheo-logical measurements, which are sensitive to themovements of whole chains (terminal processes)show a significant increase of viscosity and rein-forcement of the molten PLA matrix due to thepresence of the dispersed organoclay. This effectwas enhanced by the compatibilization andincreased with the organoclay content. Rheologicalproperties of the unfilled PLA and compatibilizedPLA were comparable indicating negligible influ-ence of the compatibilizer on PLA in the moltenstate.

Viscoelastic spectra (DMTA data) showed agradual increase of the storage and loss moduli withthe increase of the organoclay content andimproved dispersion. The mechanical relaxationprocesses (a, b and c) were affected to some extendby the material composition and nanostructure. Thenanostructure affected the cold crystallization pro-cess of PLA matrix as detected by DMTA, DSCand BBDS. In the systems with the compatibilizercrystallization of PLA occurred at highertemperature.

The dielectric properties of all the nanocompos-ites in the high temperature range are dominatedby the processes related to ionic species present inthe nanoclay – dc conductivity, Maxwell–Wagner–Sillars effect and electrode polarization. At hightemperatures both real and imaginary parts of e*

strongly increase, especially at low frequencies. Anadditional maximum of e00 at around 80–90 �C


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

resulting from the cold crystallization is observed.Dielectric studies show a relatively weak effect ofthe nanofiller on aS and b relaxation processes inPLA matrix, except for the highest organoclay con-tent (N10M system), most probably because theaverage distance between PLA segments and claysheets is too big to observe space confinementeffects.

Acknowledgments

This work was partially supported by the Minis-try of Science and Information Society Technolo-gies (Poland) through the Center of Molecular andMacromolecular Studies, PAS, under Grant No.PBZ–KBN-070/T09/2001, 2003–2006.

The author are indebted to Dr. I. Stevenson(LMBP) helping with dielectric measurements andgratefully acknowledges the Cargill-Dow Polymers,LLC, for supplying PLA.

References

[1] Sinclair RG. The case for polylactic acid as a commoditypackaging plastic. J Macromol Sci Pure Appl Chem1996;A33(5):585–97.

[2] Ogata N, Jimenez G, Kawai H, Ogihara T. Structure andthermal/mechanical properties of poly(L-lactide) blend. JPolym Sci, Part B: Polym Phys 1997;35(2):389–96.

[3] LeBaron PC, Wang Z, Pinnavaia TJ. Polymer-layeredsilicate nanocomposites: an overview. Appl Clay Sci1999;15(1–2):11–29.

[4] Alexandre M, Dubois P. Polymer-layered silicate nanocom-posites: preparation, properties and uses of a new class ofmaterials. Mater Sci Eng R-Rep 2000;28(1–2):1–63.

[5] Pluta M, Galeski A, Alexandre M, Paul MA, Dubois P.Polylactide/montmorillonite nanocomposites and micro-composites prepared by melt blending: Structure andsome physical properties. J Appl Polym Sci 2002;86(6):1497–506.

[6] Ray SS, Yamada K, Okamoto M, Fujimoto Y, Ogami A,Ueda K. New polylactide/layered silicate nanocomposites. 5.Designing of materials with desired properties. Polymer2003;44(21):6633–46.

[7] Ray SS, Okamoto M. New polylactide/layered silicatenanocomposites. 6. Melt rheology and foam processing.Macromol Mater Eng 2003;288(12):936–44.

[8] Ray SS, Okamoto M. Polymer/layered silicate nanocom-posites: a review from preparation to processing. ProgPolym Sci 2003;28(11):1539–41.

[9] Ray SS, Bousmina M. Biodegradable polymers and theirlayered silicate nanocomposites: In greening the 21st centurymaterials world. Prog Mater Sci 2005;50(8):962–1079.

[10] Di YW, Iannace S, Di Maio E, Nicolais L. Poly(lactic acid)/organoclay nanocomposites: Thermal, rheological propertiesand foam processing. J Polym Sci, Part B: Polym Phys2005;43(6):689–98.

[11] Pluta M, Paul MA, Alexandre M, Dubois P. Plasticizedpolylactide/clay nanocomposites. I. The role of filler contentand its surface organo-modification on the physico-chemicalproperties. J Polym Sci, Part B: Polym Phys2006;44(2):299–311.

[12] Mierzwa M, Floudas G, Dorgan J, Knauss D, Wegner J.Local and global dynamics of polylactides. A dielectricspectroscopy study. J Non-Cryst Solids 2002;307:296–303.

[13] Henry F, Costa LC, Devassine M. The evolution ofpoly(lactic acid) degradability by dielectric spectroscopymeasurements. Eur Polym J 2005;41(9):2122–6.

[14] Mijovic J, Sy JW. Molecular dynamics during crystallizationof poly(L-lactic acid) as studied by broad-band dielectricrelaxation spectroscopy. Macromolecules 2002;35(16):6370–6.

[15] Ren JD, Adachi K. Dielectric relaxation in blends ofamorphous poly(DL-lactic acid) and semicrystalline poly(L-lactic acid). Macromolecules 2003;36(14):5180–6.

[16] Ren JD, Urakawa O, Adachi K. Dielectric and viscoelasticstudies of segmental and normal mode relaxations inundiluted poly(D,L-lactic acid). Macromolecules 2003;36(1):210–9.

[17] Fitz BD, Andjelic S. Real-time monitoring of segmentaldynamics during crystallization of poly(L(-)-lactide) bysimultaneous DRS/SALS technique. Polymer 2003;44(10):3031–6.

[18] Pluta M. Melt compounding of polylactide/organoclay:Structure and properties of nanocomposites. J Polym Sci,Part B: Polym Phys 2006;44(23):3392–405.

[19] Mijovic J, Lee HK, Kenny J, Mays J. Dynamics in polymer-silicate nanocomposites as studied by dielectric relaxationspectroscopy and dynamic mechanical spectroscopy. Mac-romolecules 2006;39(6):2172–82.

[20] Wang HW, Chang KC, Chu HC, Liou SJ, Yeh JM.Significant decreased dielectric constant and loss of polysty-rene-clay nanocomposite materials by using long-chainintercalation agent. J Appl Polym Sci 2004;92(4):2402–10.

[21] Wang HW, Shieh CF, Chang KC, Chu HC. Synthesis anddielectric properties of poly(methyl methacrylate)-clay nano-composite materials. J Appl Polym Sci 2005;97(6):2175–81.

[22] Kim DH, Cho KS, Mitsumata T, Ahn KH, Lee SJ.Microstructural evolution of electrically activated polypro-pylene/layered silicate nanocomposites investigated byin situ synchrotron wide-angle X-ray scattering and dielec-tric relaxation analysis. Polymer 2006;47(16):5938–45.

[23] Bohning M, Goering H, Fritz A, Brzezinka KW, Turky G,Schonhals A, et al. Dielectric study of molecular mobility inpoly(propylene-graft- maleic anhydride)/clay nanocompos-ites. Macromolecules 2005;38(7):2764–74.

[24] Exxelor-datasheet. , 2006.[25] Wagener R, Reisinger TJG. A rheological method to

compare the degree of exfoliation of nanocomposites.Polymer 2003;44(23):7513–8.

[26] Degee P, Dubois P, Jerome R. Bulk polymerization oflactides initiated by aluminium isopropoxide. 3. Thermalstability and viscoelastic properties. Macromol Chem Phys1997;198(6):1985–95.

[27] Wachsen O, Platkowski K, Reichert KH. Thermal degra-dation of poly-L-lactide – Studies on kinetics, modelling andmelt stabilisation. Polym Degrad Stabil 1997;57(1):87–94.

[28] Ren JX, Silva AS, Krishnamoorti R. Linear viscoelasticityof disordered polystyrene–polyisoprene block copolymer

http://mymatweb.com


based layered-silicate nanocomposites. Macromolecules2000;33(10):3739–46.

[29] Zhao J, Morgan AB, Harris JD. Rheological characteriza-tion of poly styrene-clay nanocomposites to compare thedegree of exfoliation and dispersion. Polymer 2005;46(20):8641–8660.

[30] Kulinski Z, Piorkowska E. Crystallization, structure andproperties of plasticized poly(L-lactide). Polymer 2005;46(23):10290–300.

[31] Pluta M. Morphology and properties of polylactide modifiedby thermal treatment, filling with layered silicates andplasticization. Polymer 2004;45(24):8239–51.

[32] Jacobsen S, Fritz HG. Plasticizing polylactide – The effect ofdifferent plasticizers on the mechanical properties. PolymEng Sci 1999;39(7):1303–10.

[33] Pluta M, Galeski A. Plastic deformation of amor-phous poly(L/DL-lactide). Biomacromolecules, doi:10.1021/bm061229v, in press.

[34] Kremer F, Schoenhals A. Broadband Dielectric Spectros-copy. Heidelberg: Springer Verlag; 2002.

[35] Noda N, Lee YH, Bur AJ, Prabhu VM, Snyder CR, RothSC, et al. Dielectric properties of nylon 6/clay nanocom-posites from online process monitoring and off-linemeasurements. Polymer 2005;46(18):7201–17.

[36] Davis RD, Bur AJ, McBrearty M, Lee YH, Gilman JW,Start PR. Dielectric spectroscopy during extrusion process-ing of polymer nanocomposites: a high throughput process-ing/characterization method to measure layered silicatecontent and exfoliation. Polymer 2004;45(19):6487–93.

MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

http://dx.doi.org/10.1021/bm061229vhttp://dx.doi.org/10.1021/bm061229v

Polylactide/montmorillonite nanocomposites: Structure, dielectric, viscoelastic and thermal propertiesIntroductionExperimentalMaterialsSample preparationCharacterization

Results and discussionSize exclusion chromatography (SEC)XRD analysisRheological propertiesCalorimetric characterizationViscoelastic propertiesBroadband dielectric spectroscopy

DiscussionConclusionsAcknowledgmentsReferences

Documents

Polylactide/montmorillonite nanocomposites: Structure ...ve4ep.com/a/19june/статьи 18_06_10/333.pdfHowever, dielectric properties of poly-lactide-based nanocomposites containing