marras nanostruct vs microst

EUROPEAN

European Polymer Journal 43 (2007) 2191–2206

www.elsevier.com/locate/europolj

POLYMERJOURNAL

OL

OG

Y

Macromolecular Nanotechnology

Nanostructure vs. microstructure: Morphologicaland thermomechanical characterization ofpoly(L-lactic acid)/layered silicate hybrids

Sotirios I. Marras a,b, Ioannis Zuburtikudis b,*, Costas Panayiotou a,*

a Department of Chemical Engineering, Aristotle University of Thessaloniki, 54124 Thessaloniki, Greeceb Department of Industrial Design Engineering, TEI of Western Macedonia, 50100 Kozani, Greece

Received 9 January 2007; received in revised form 23 February 2007; accepted 5 March 2007Available online 13 March 2007

MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

N

Abstract

Nano- and micro-composites of poly(L-lactic acid) (PLLA) with various loadings of natural and hexadecylamine-modified montmorillonite were prepared by the solvent casting method to study the effect of nanostructure on the thermo-mechanical properties of the hybrid materials. The changes on structure and surface of montmorillonite, induced by theion-exchange modification process, were characterized by X-ray diffraction (XRD) analysis and zeta-potential determina-tion, while the morphology of the hybrids and the dispersion of the clay into the polymer matrix were examined by XRD,transmission electron microscopy and atomic force microscopy. The results showed that, although at low clay contentexfoliation dominates, for filler loadings greater than 5 wt% both exfoliation and intercalation of the clay filler areobserved. Thermal degradation studies of the materials produced using thermogravimetry revealed the introduction ofa small amount of organo-modified silicate significantly improves their thermal stability. Differential scanning calorimetryshowed the thermal behavior of the polymer matrix strongly depends on the nature and content of the silicate filler.Scanning electron microscopy of the deformed surfaces affirmed a different deformation process mechanism betweenthe two types of composites.� 2007 Elsevier Ltd. All rights reserved.

Keywords: Biodegradable polymer; Poly(L-lactic acid); Nanocomposites; Thermal properties; Mechanical properties

1. Introduction

Polymer nanocomposites are a class of materialsthat has already gained a big popularity among

0014-3057/$ - see front matter � 2007 Elsevier Ltd. All rights reserved

doi:10.1016/j.eurpolymj.2007.03.013

* Corresponding authors. Tel.: +30 24610 40161; fax: +3024610 39682 (I. Zuburtikudis), tel.: +30 2310 996223; fax: +302310 996232 (C. Panayiotou).

E-mail addresses: [email protected] (I. Zuburtikudis), [email protected] (C. Panayiotou).

researchers both in academia and industry due totheir exceptional properties, which are superior tothose of the virgin polymers and of the conventionalcomposites [1]. These hybrid materials consist of apolymer matrix in which inorganic particles withat least one dimension in the nanometer scale arehomogeneously dispersed. The nano-sized particlesthat are typically less than 5 wt% in proportioninteract with the polymer matrix and improve sub-stantially many physical properties of the pristine

.

mailto:[email protected]



2192 S.I. Marras et al. / European Polymer Journal 43 (2007) 2191–2206

MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

polymer including increased mechanical characteris-tics [2], enhanced barrier resistance [3], decreasedflammability [4,5], and increased biodegradabilityin the case of biodegradable polymers [6].

Biodegradable and biocompatible polymers havealso attracted significant attention in the last decadein general and, more specifically, they have becomethe polymer matrix of choice for many researchersin the field of nanocomposites; they are environmen-tally friendly, as they degrade under the action ofphysicochemical and biological factors in a humidenvironment, while they can be employed in manybiomedical and pharmaceutical applications suchas sutures and biologically active controlled-releasedevices [7]. In this direction, poly(lactic acid)(PLA) is rapidly gaining recognition as one of themost popular biopolymers. It is a linear aliphaticpolyester derived from 100% renewable resourcessuch as corn and sugar beets.

PLA can be prepared either by direct condensa-tion of lactic acid or by the ring-opening polymeri-zation of the cyclic lactide dimer [8]. The directcondensation path is an equilibrium reaction caus-ing difficulties in removing trace amounts of waterin the late stages of polymerization. For this reason,only a low molecular mass polymer can be obtainedby this approach. Three stereoisomers of lactideexist: L-lactide, D-lactide and mesolactide. The ste-reochemical structure of the lactide monomer deter-mines the stereochemical composition of theresulting polymer. Sequentially, it is the architectureof the resulting polymer (determined by the propor-tions of L, D, and the molecular mass) that dictatesits properties such as crystallinity, melting point,and mechanical strength. Thus, PLA can be totallyamorphous or up to 40 % crystalline depending onthe proportions of D and L-lactides present [8]. Asa consequence, PLA polymers can display a broadrange of mechanical properties, their glass transi-tion temperature may range from 50 �C to 80 �C,while their melting point can vary from 130 �C to180 �C.

As regards the nano-sized inorganic particles thatare added in the polymer matrix, phyllosilicatessuch as montmorillonite clay are of particular inter-est because of their abundance, low cost and theirgeometrical features. They consist of numerous lay-ers stacked in parallel, otherwise called ‘‘platelets’’,that are nm-thick and hundreds-nm- or lm-long;that is, each platelet has a high aspect ratio, whichenhances the possibility for energy transfer fromthe inorganic phase to the organic one and vice

versa [9]. Natural clays mixed with polymers leadto the formation of nanocomposites very rarely.Their hydrophilic nature obstructs their homoge-neous dispersion in the organic polymer phase.However, it is well known [10] for immiscibleorganic-inorganic systems, that the weak interactionbetween the organic and the inorganic componentsresults in poor mechanical and thermal properties.In order to render the hydrophilic phyllosilicatesmore organophilic, hydrated cations (Ca++, Na+)of the interlayer space are exchanged with cationicsurfactants such as alkylammonium or alkylphos-phonium. This ion-exchange reaction lowers thesurface energy of the inorganic material and makesit compatible with organic liquids [11] or polymers[12].

From the structural point of view, two types ofpolymer/phyllosilicate nanocomposites are possible,the so-called ‘‘intercalated’’ and the ‘‘exfoliated’’one, although in most of the real cases and unlessspecial care has taken, the usual situation is a mix-ture of ‘‘intercalated’’ and ‘‘exfoliated’’ nanocom-posites [1]. In an intercalated system, the polymeris inserted into the silicate galleries, i.e. the inter-layer space, but the stacking of the silicate layersis preserved. In an exfoliated system, the silicate lay-ers are fully separated and randomly dispersed inthe polymer matrix due to the extensive polymerinsertion [13]. This good dispersion of the clay plate-lets gives rise to a bigger property enhancement,since the nanostructure effect can be more pro-nounced in the exfoliated nanocomposite than inthe intercalated one.

Polymer/layered-silicates nanocomposites can beprepared by the following methods: (i) Solution-casting, (ii) in situ polymerization, (iii) melt interca-lation. In the first case, the inorganic material isthoroughly mixed and dispersed in a polymer solu-tion resulting in the insertion of polymer chains intothe galleries of the clay, which swells. In situ poly-merization involves monomer intrusion into the claygalleries followed by polymerization. In the meltintercalation process, the molten polymer is blendedwith the clay, so that the polymer/layered silicateinteractions be optimized and the insertion of thepolymer chains is facilitated [1].

PLA/layered silicate nanocomposites have beenprepared by several methods such as melt extrusion[10,14–16] and solution casting [17–19]. The aim hasbeen to produce a hybrid material with improvedproperties over pristine PLA. However, manyaspects of the composite’s structure-property rela-

S.I. Marras et al. / European Polymer Journal 43 (2007) 2191–2206 2193

MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

tionship have not been elucidated yet. Understand-ing this structure-property relationship thoroughlyis the key in designing materials with tailor-madeproperties. The effect of clay content and the com-posite’s morphology on the physical properties ofthe hybrid material are two key issues towards thisdirection.

In the present study, micro- and nanocompositesof poly(L-lactic acid) (PLLA) with various composi-tions of natural and organophilic montmorilloniteclay were prepared by the solution casting methodusing chloroform as the solvent. The aim was toinvestigate the role of nanostructure on the thermaland mechanical behavior of the hybrid materials bycomparing the thermo-mechanical properties of thenanocomposites with the corresponding ones of themicrocomposites.

2. Experimental

2.1. Materials

PLLA (Mn = 81,000, Mw/Mn = 1.9 and 0% con-tent in D-lactide stereoisomer was obtained fromGalactic S.A. (Belgium), while sodium montmoril-lonite (NaMMT) (CEC = 92.6 meq/100 g) was sup-plied by Southern Clay Products (Texas, USA).Hexadecylamine (C16A), chloroform (CHCl3) andconcentrated HCl were purchased from Sigma–Aldrich. All materials were used as received andwithout any further purification.

2.2. Preparation of the organophilic layered silicate

Organically modified montmorillonite(C16MMT) was prepared through an ion-exchangereaction between sodium montmorillonite and hexa-decylammonium salt (C16A). Quantity of C16A,equal to 1.4 times the cation exchange capacity(CEC) of the inorganic material and an equivalentamount of HCl were dissolved in a 1:1 solution ofethanol and water at 75 �C. The inorganic materialwas dispersed in water at 75 �C and then, this aque-ous suspension (3 wt%) was added to the alkylam-mine solution. The mixture was stirred vigorouslyfor 24 h at the aforementioned temperature. Theprecipitate was collected by filtration and washedwith a hot 1:1 solution of ethanol and water untilthe AgNO3 test verified the absence of chloride.Eventually, the cation-exchanged montmorillonitewas dried and ground to a powder, which was keptin a desiccator for further use.

2.3. Preparation of hybrid materials

Nanocomposites were fabricated by using thesolution casting technique. Initially, the organo-modified montmorillonite was dispersed in chloro-form. The clay suspension was sonicated for 1 hwith a Heat Systems-Ultrasonics, W-375 sonicator.Then, the polymer was dissolved in it and the wholedispersed-clay polymer solution (5 wt%) was furthersonicated for 1 h. Extended increase of the solutiontemperature was avoided by using a pulse mode.The mixture was then cast onto Petri dishes andwas kept for 48 h in a chloroform atmosphere inorder to accomplish a slow solvent evaporation,which eliminates polymer skin formation. Chloro-form was selected as the PLLA solvent instead ofdichloromethane, because of its lower vapor pres-sure when compared with dichloromethane and thisled to the formation of a more homogeneous film[20]. That way, nanocomposite membranes withorganoclay content ranging from 1 to 30 wt% andwith thickness between 100 and 150 lm wereobtained; these films were eventually dried in vac-uum at 50 �C for 24 h.

As written in the introduction, in order toachieve a better understanding of the nanostructureeffect on the thermo-mechanical properties of thehybrid materials, microcomposites were also pro-duced by the above described procedure using, how-ever, unmodified sodium montmorillonite. In thiscase, a smaller number of final hybrid compositionswere prepared (1–8 wt% clay), just for the requiredcomparisons.

2.4. Characterization

The electrokinetic potential (f) of both the natu-ral, i.e. unmodified clay, and the organophilic onewere measured at different pH values using Smolu-chowski’s equation and with the aid of an auto-mated microelectrophoresis instrument (ZetaPlus,Brookhaven Instruments Corporation). At least 20runs were conducted for each sample and the aver-age values were recorded. For the preparation of thesamples, sufficient amount (2 mg/ml) of dry claywas dispersed in de-ionized water and the suspen-sion was shaken intensively for 24 h. pH adjustmenttook place by dropwise addition of HCl or NaOHsolutions. All measurements were conducted at23 �C ± 0.2 �C.

The structure of the hybrid materials being eithernano- or micro-scaled was investigated by XRD

0 3 6 9 12 15

*

C16MMT

Rel

ativ

e In

tens

ity

2θ (degrees)

NaMMT 1.2 nm

1.82 nm

*

Fig. 1. X-ray diffraction patterns of natural and organo-modifiedmontmorillonite.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

using a Siemens D 500 diffractometer and Ni-fil-tered Cu Ka radiation (k = 0.154 nm). The scanningrange was varied from 2h = 2� to 30�.

The differences of the clays surface morphologywas visualized by using a Jeol JSM-5900 LV scan-ning electron microscope (SEM). The samples werecoated with a carbon thin film under vacuum in anAgar auto carbon coater.

The dispersion of the silicate layers was also eval-uated by transmission electron microscopy (TEM)bright field imaging using a Jeol 120CX microscopeoperating at 120 kV and a JEOL 2011 operating at200 kV (for higher resolution). Specimens suitablefor planar view observations were mechanicallythinned down to 10 lm followed by further thinningto electron transparency by Ar ion beam milling at alow incident angle of 5�.

Atomic force microscopy (AFM) measurementswere performed on a Veeco, CP-II microscope.The instrument was operated in intermittent-con-tact mode to eliminate lateral forces, which candamage the sample or the tip and reduce image res-olution. Height (topography) and phase imageswere recorded at room temperature and ambientatmosphere. A commercial silicon tip-cantileverwas used with a spring constant 20–80 N/m, a reso-nance frequency in the 200–400 Hz range and a tipradius of less than 10 nm. For this analysis, smoothcross-sections were made by room temperaturemicrotoming. Usually, the ‘‘simply cut’’ cross sec-tions are too rough to obtain the fine details. Thesamples were imaged 10 days after micro-cleavage,showing that the sample cross-sections were stablefor long time.

The thermal properties of the produced hybridswere investigated with the aid of a ScimadzuDSC-50 differential scanning calorimeter (DSC).Temperature scans were carried out at a heatingrate of 10 �C/min under constant nitrogen flow of20 cm3/min. The samples were heated to 200 �Cand remained at that temperature for 5 min inorder to release their thermal history and, after-wards, they were quenched to �5 �C. The determi-nation of the glass transition temperature (Tg),crystallization temperature (Tc) and melting point(Tm) took place during the second run. Thermo-gravimetric analyses (TGA) of the hybrids wereperformed on a Shimadzu TGA-50 with a heatingramp of 10 �C/min up to 750 �C under both flow-ing nitrogen (20 cm3/min) and static air. TGA runsin isothermal conditions were also conducted at280 �C in air.

The mechanical properties of the produced com-posites were studied in tensile loading conditionsusing a Thumler GmbH Z3-500 tensile testeraccording to ASTM D1708 method. The crossheadspeed was 5 mm/min. At least six measurementswere taken for each sample and the results wereaveraged to obtain a mean value. The surface mor-phology of the tensile deformed composites andtheir fracture surface were examined with a scan-ning electron microscope (SEM). The sample prep-aration for the SEM examination was similar tothat described above.

3. Results and discussion

3.1. Organoclay characteristics

The unit layers of montmorillonite are stackedtogether face-to-face forming the crystal lattice.The basal spacing (i.e. the distance between two suc-cessive plane layers) of both the natural (sodium)montmorillonite and the organically (hexadecylam-monium cation) modified montmorillonite wasinvestigated by XRD and the results are presentedin Fig. 1. It appears that the basal spacing of the lay-ered mineral is 1.2 nm. Insertion of hexadecylam-monium ions into the interlayer spaces of thenatural clay led to an increase of this spacing from1.2 nm to 1.82 nm indicating a bilayer arrangementof the alkylammonium molecules [21,22].

In Fig. 2, the electron photomicrographs takenby the SEM describe the surface characteristics ofNaMMT and C16MMT. As observed, the granuleof modified montmorillonite exhibits a more open

Fig. 2. SEM micrographs of (a) un-modified and (b) modified montmorillonite.


OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

structure, while individual blocks of layers can bedistinguished in it.

Fig. 3 illustrates the zeta potential of the clay as afunction of the solution’s pH. As it is clearly shown,the zeta potential of the natural clay is negative overthe whole range of pH values. Moreover, it does notchange significantly with pH. These observationsare in accordance with Swartzen et al. [23], whohave claimed that the surface charge of NaMMTin aqueous solutions is always negative and pHindependent due to the isomorphic substitutions ofAl3+ by Mg2+ in the octahedral sheet and those ofSi4+ by Al3+ in the tetrahedral sheet. However, atthe edges of the particles where the aluminumhydrosilicate layers are disrupted, polar sites aresituated due to broken bonds. These sites are condi-tionally charged (positively or negatively) by direct

2 4 6 8 10 12-60

-45

-30

-60

-30

0

30

60

Zet

a P

oten

tial (

mV

)

pH

NaMMT

C16MMT

Fig. 3. Zeta potential of natural and organophilic montmoril-lonite as a function of pH.

MA

CR

H+ or OH� transfer from the aqueous phasedepending on the pH. Addition of C16A leads toan increase of the surface charge of MMT, whichbecomes positive in neutral and acidic solutionsand decreases as the concentration of the hydroxylions is increased. Xu and Boyd [24] proposed thatthe alkylammonium cation is initially adsorbed inthe interlayer space by cation exchange and thiscauses extensive clay aggregation. Increasing thecation loading, C16A is adsorbed on the externalsurfaces of the aggregates via both cation exchangeand hydrophobic bonding; the hydrophobic bond-ing causing positive charge development on the sur-faces and ultimately clay dispersion. Two types ofamine adsorption reactions on montmorillonitewere also pointed out by Zhang et al. [25]: cation-exchange reaction and adsorption at non-exchange-able sites with the former to be the dominatingreaction, when the amount of amine is less than100% of the clay CEC. These views are in accor-dance with all of the experimental findings describedin Fig. 3 and explain them.

3.2. Morphology of PLLA Hybrids

Fig. 4 presents XRD results for the compositesformed. As mentioned earlier, the mean basal spac-ing of C16MMT is 1.82 nm. The diffraction patternsof the composites reveal that the characteristic peak,referred to the d(0 0 1) of C16MMT, is shifted towardslower 2h values and is broadened indicating thatpolymer chains have diffused into the silicate galler-ies expanding the clay structure (‘intercalation’).The corresponding calculated d(0 0 1) values changein a consistent way with the clay content of thehybrids produced; that is, decrease of the clay

2 4 6 8 10

*

*

*

2.6 nm

3.27 nm

2.45 nmC16MMT

15% C16MMT

30% C16MMT

9% C16MMT

5% C16MMT

Rel

ativ

e in

tens

ity

2θ (degrees)

2θ (degrees)

Pure PLLA

3% C16MMT

1.82 nm *

0 5 10 15 20 25 30

NaMMT

Rel

ativ

e in

tens

ity

PLLA+3% NaMMT

a

b

Fig. 4. X-ray diffraction patterns of (a) PLLA nanocompositeswith various loadings of C16MMT and (b) PLLA microcompositewith 3 wt% of NaMMT.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

content increases the d(0 0 1) value. This means thatdecrease of the clay loading facilitates polymerinsertion, which in turn increases the interlayerspacing. However, as more polymer enters the sili-cate galleries, the clay layers lose their parallelstacking and become partially disordered. Thebroadening of the peaks may be the result of super-position of reflections corresponding to variousinterlayer distances. At low clay contents, largeramounts of polymer that enter the galleries maypush the platelets so far apart that the platelet sep-aration exceeds the XRD sensitivity. The lack ofpeak detection for clay loadings less than 5 wt%was probably due to this high separation (‘exfolia-tion’) of the clay platelets. At low filler loading,exfoliation is the predominant mechanism of claydispersion. Increase of the clay loading impedescomplete exfoliation of such a high aspect ratio sil-

icate sheets due to the restricted area remainingavailable in the polymer matrix and hence, thedegree of intercalation enhances [26].

The intensity of the characteristic peak of theclay (Id(0 0 1)) increases as the percentage of theC16A-modified clay in the hybrid material becomeshigher. This is because the number density of theintercalated clay particles in the polymer matrixincreases, as the clay content arises, and many moreX-rays are diffracted.

Mixing of the polymer with the natural, unmod-ified clay did not result in the formation of a nano-composite material, as no shift of the clay’scharacteristic peak was observed (Fig. 4b). Usually,blending polymer matrices with NaMMT leads tothe synthesis of conventional composites, since mostof the polymers do not migrate into the hydratedNa+ interlayer space due to its hydrophobicity [12].

Information arising from the XRD analysis is ingood agreement with TEM observations. Fig. 5aillustrates the structure of nanocomposites with15 wt% organophilic clay. Dark lines represent thecross sections of the clay layers and the gray areacorresponds to the polymer matrix. Inspection ofthis picture reveals a homogeneous dispersion ofthe clay platelets throughout the matrix and in thenanometer-scale level. Extensive TEM observationsrevealed a coexistence of silicate particles in theintercalated (Fig. 5c) and the exfoliated (Fig. 5d)state for the full range of clay loadings examined.

As it is shown in Fig. 5b, TEM investigations ofthe composite prepared by the addition of NaMMTevidenced the existence of immiscible clay tactoidswith sizes in the micrometer range. The lack of affin-ity between the hydrophilic silicate and the hydro-phobic polymer causes agglomeration of themineral in the polymer matrix.

The high dispersion of the organoclay in PLLAwas also verified by AFM observations. The topo-graphy and phase images presented in Fig. 6a andb, respectively describe the morphology of a hybridprepared by the addition of 9 wt% of organoclay.The height differences in the topography imageand the light phase variations in the phase imageindicate the presence of small and large intercalatedclay tactoids dispersed into the matrix.

3.3. Thermal properties of PLLA hybrids

Fig. 7a shows TGA data for pure PLLA and itshybrids, which were taken in a N2 atmosphere andin the temperature range of 165 �C, (where the sam-

Fig. 5. TEM micrographs of PLLA with (a) 15 wt% C16MMT and (b) 3 wt% NaMMT. High resolution TEM micrographs of PLLA with15 wt% C16MMT showing areas with intercalated (c) and exfoliated (d) clay platelets.

Fig. 6. AFM (a) topography and (b) phase images of PLLA nanocomposite containing 9 wt% C16MMT.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

ples are considered to be free of humidity), up to750 �C. With a heating rate of 10 �C/min, PLLAvolatilizes completely in a single step beginning atabout 270 �C. The introduction of organophilicinorganic material improves the thermal stabilityof the polymer increasing the temperature ofdecomposition initiation, especially when the inor-

ganic content is small. This may be attributed tothe labyrinth (‘barrier’) effect originated by the highaspect ratio clay platelets, which are dispersed in thepolymer matrix delaying the escape of volatile deg-radation products within the nanocomposite [27].Increase of the clay loading destabilizes the polymermatrix. This can be mainly assigned to the increased

200 300 400 500 600 700 800

0

20

40

60

80

100

160 200 240 280 32096

97

98

99

100

15%

30%

9%

5%

1%

0%

3%

1.5%

Wei

ght (

%)

Temperature (oC)

Pure PLLA

1% C16MMT

1,5% C16MMT

3% C16MMT

5% C16MMT

9% C16MMT

15% C16MMT

30% C16MMT

200 300 400 500 600 700 800

0

20

40

60

80

100

200 240 28096

97

98

99

100

0%

5%

3%

1%

9%

Wei

ght (

%)

Temperature (oC)

Pure PLLA 1% NaMMT 3% NaMMT 5% NaMMT 9% NaMMT

Fig. 7. TGA data for PLLA hybrids prepared with various loadings of (a) C16MMT and (b) NaMMT, in inert (N2) atmosphere.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

presence of the ammonium cation, which thermallydecomposes to produce ammonia and the corre-sponding olefin following the Hofmann mechanism[27]. As it is revealed in Fig. 8, C16MMT displays alower decomposition initiation temperature fromboth the natural MMT and the pristine polymerdue to the ammonium cation incorporation. There-fore, when the organophilic clay is present in higherthan about 9 wt% in the polymer matrix, the corre-sponding hybrid materials exhibit a decreased onsetdecomposition temperature compared to the pris-tine polymer due to the degradation of the organicmodifier.

To support this conclusion in a more sound wayand to have a concrete view about the influence ofnanostructure on the polymer thermal stability,

microcomposites were tested in the same TGAequipment and at the same conditions, as the onesused for the hybrid nanocomposite materials. Asillustrated in Fig. 7b, the incorporation of NaMMTinto the matrix does not affect the thermal degrada-tion of the polymer significantly.

The degradation properties of the polymer/claynanocomposites were also investigated under oxida-tive conditions, as shown in Fig. 9a. Oxygen inducesthermal volatilization of PLLA at a lower tempera-ture than nitrogen does. However, the organoclayaddition delays the polymer mass loss significantlyand, thus, increases the decomposition temperatureof the hybrid material. In fact, the presence of theorganophilic inorganic material enhances char for-mation and hinders diffusion of the volatile decom-

0 300 600 900

0

40

80

Pure PLLA

C16

MMT

Wei

ght (

%)

Temperature (oC)

NaMMT

Fig. 8. TGA data for natural (NaMMT) and organophilic(C16MMT) montmorillonite and pristine PLLA.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

position products within the nanocomposite [4].Moreover, the silicate platelets act as an additionalbarrier, which most likely works against oxygen dif-fusion towards the bulk of the polymer [27].

TGA data taken in oxidant atmosphere forPLLA microcomposites prepared with various load-ings of NaMMT (Fig. 9b) reveal that the introduc-tion of NaMMT to the polymer matrix resulted in asmaller enhancement of the thermal stability of themicrocomposite due to the poorer dispersion ofthe unmodified inorganic material into the polymermatrix. The effect of the clay type and clay contenton the onset decomposition temperature (that is, thetemperature required for 1 wt% material mass loss)are described in Fig. 10a and b. Fig. 10a shows datataken in the presence of nitrogen (inert atmosphere),while Fig. 10b presents data taken in air (oxidantatmosphere). It appears that increased thermalstability in comparison with that of the purepolymer is observed only in the case of nano-composites. Enhanced thermal stability in air ofPLLA/MMT nanocomposites was also reported in[10].

In isothermal conditions (280 �C) under air, thepolymer/clay nanocomposites lost weight with areduced rate compared to the pure polymer,(Fig. 11a). As illustrated in Fig. 11b, the nanohy-brids, prepared by the incorporation of 3–5 wt%of organically modified clay, require almost threetimes the time that the pure polymer needs for a50% mass loss. This greatly retarded material vola-tilization may be attributed to the formation of acarbonaceous-silicate char structure, which is pro-

duced on the surface of the nanocomposite duringburning [4]. In the presence of oxygen, the intimatecontact between the clay platelets and the polymerfavors this char formation [28], which acts as anexcellent insulator and mass transport barrier thatdelays the diffusion of oxygen from the air to thepolymer mass and slows down the escape of the vol-atile products generated during decomposition.Reduction in composite volume during volatiliza-tion improves this barrier property due to reassem-bly of the dispersed silicate layers [27].

The synergy between the polymer and the well-dispersed organo-modified silicate layers on thethermal stability of the composites can be confirmedby investigating the thermal behavior of the corre-sponding microcomposites at the same conditionsas before (oxidant atmosphere, 280 �C). Here, vola-tilization was not delayed by the introduction of theNaMMT powder in the polymer matrix. This has todo with the very low clay dispersion and indicatesthe weak interfacial interaction between the ultra-fine inorganic NaMMT particles and the matrix.A comparison of the nanocomposite (with 3 wt%C16MMT) and the microcomposite (with 3 wt%NaMMT) is presented in Fig. 12, where it is clearlyshown that only the former, the nanocomposite,exhibits a retarded weight loss.

Fig. 13a presents the DSC heating curves of thequenched nanocomposites. In all samples, an exo-thermic peak can be observed at the temperaturerange from 95 to 120 �C. This peak represents thecrystallization of PLLA. The pure polymer crystal-lizes at 107 �C. Introduction of the clay nanoparti-cles facilitates the PLLA crystallization processduring heating, leading to a decrease of the crystal-lization temperature. The silicate layers providelarge surface area due to their nanosize and theirhigh aspect ratio and, hence, serve as additionalnucleation sites enhancing the crystallization rateof the matrix molecules [14,16,19]. This effectbecomes stronger with the increase of the organom-odied clay content, because of the increased pres-ence of nucleation sites. The hybrid with 5 wt%organoclay loading displays the lowest crystalliza-tion temperature (Tc). Further increase of this con-tent retards the crystal growth process increasingthe crystallization temperature [29].

At low clay contents, the distance between thedispersed platelets is long enough, so that the addi-tional nucleation sites that are created with theincrease of the loading can easily incorporate sur-rounding polymer. However, when the clay loading

200 300 400 500 600 700 800

0

20

40

60

80

100

160 200 240 28096

97

98

99

100

30%

0%

1%

1.5%3%

15%5%

9%

Wei

ght (

%)

Temperature (ºC)

Pure PLLA

1% C16MMT

1.5% C16MMT

3% C16MMT

5% C16MMT

9% C16MMT

15% C16MMT

30% C16MMT

200 300 400 500 600 700 800

0

20

40

60

80

100

180 210 240 270 30096

97

98

99

100

9%

5%

0%

3%1%

Wei

ght (

%)

Temperature (ºC)

Pure PLLA

1% NaMMT

3% NaMMT

5% NaMMT

9% NaMMT

Fig. 9. TGA data for PLLA hybrids prepared with various loadings of (a) C16MMT and (b) NaMMT, in oxidant (air) atmosphere.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

becomes high enough, the platelets can retard crys-tallization acting as a physical hindrance to the dif-fusing polymer molecules that travel to the growingcrystallite. Moreover, the high miscibility betweenthe polymer matrix and the organic modifier and,hence, the interaction between them, partiallyimmobilizes polymer chains attached to the clay lay-ers and obstructs them from taking part in the flowprocess and, thus, the crystallization process. There-fore, the higher the concentration of the organoclay,the larger the amount of polymer which interactswith the organic modifier [14].

The incorporation of the inorganic materialaffects also the melting point (Tm) of the polymer.Tm is slightly decreased with increasing clay content.Besides, it is concluded that the presence of the

nanoclay did not influence the glass transition tem-perature (Tg) of the polymer. This is in accordancewith previous observations [10], where it wasreported that Tg was independent of the fillerconcentration.

In contrast to the organophilic montmorillonite,the addition of the unmodified one to the polymerdid not significantly affect the aforementioned ther-mal parameters, as illustrated in Fig. 13b. Thisresult indicates the poor interaction of NaMMTwith the matrix.

Fig. 14a and b present the thermal properties (Tg,Tc, Tm) of PLLA as a function of both the clay typeand the clay content. It is evident that the organo-modifier plays a very important role in the intimatecontact between the inorganic material and the

0 3 6 9 12 15260

265

270

275

280

285

290

295

Ons

et d

ecom

posi

tion

tem

pera

ture

(º C

)

Clay content (wt, %)

Pure PLLA PLLA+C16MMT

PLLA+NaMMT

0 5 10 15 20 25 30210

220

230

240

250

260

270

Ons

et d

ecom

posi

tion

tem

pera

ture

(º C

)

Clay content (wt, %)


PLLA+NaMMT

a b

Fig. 10. Effect of clay nature and clay content on the onset decomposition temperature (1 wt% of mass loss) of the hybrid materials, ininert (a) and oxidant (b) atmosphere.

0 200 400 600

0

20

40

60

80

100

0%

9%

5%

1%

15%

30%

3%

1.5%

Wei

ght (

%)

Time (min)

0 5 10 15 20 25 3060

90

120

150

180

210

Tim

e re

quire

d fo

r 50

% o

f wei

ght l

oss

(min

)

C16

MMT content (wt,%)

a

b

Fig. 11. (a) Effect of C16MMT content on the weight loss ofPLLA nanocomposites in oxidant atmosphere and isothermalconditions (280 �C). (b) Time required for 50% degradation of thenanocomposites.

0 200 400 600 800

0

20

40

60

80

100

3% NaMMT

3% C16MMT

Wei

ght (

%)

Time (min)

Pure PLLA

Fig. 12. TGA data in isothermal conditions (280 �C) and oxidantatmosphere for pure PLLA, nanocomposite (with 3 wt%C16MMT) and microcomposite (with 3 wt% NaMMT) PLLA.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

polymer and lowers the crystallization temperatureand the melting point of the polymer matrix.

3.4. Mechanical behavior of PLLA hybrids

Table 1 illustrates the effect of the filler type andcontent on the composite’s mechanical behavior.The introduction of the organoclay significantlystrengthens the hybrid material. A maximum valuefor the strength is observed at a clay loading of3 wt%. Further increase of the filler content hindersits dispersibility in the polymer matrix encouragingpremature failure initiation. It is worth noting thatthe hybrid prepared by the addition of untreatedmontmorillonite showed higher strength comparedto the pristine polymer. Similar results were alsopresented by Cho and Paul [2] who studied thebehavior of Nylon 6 composites.

The elongation at break of the hybrids does notchange significantly at very low organo-modified

30 60 90 120 150 180 210

Pure PLLA

5% C MMT

30% C MMT

15% C MMT

9% C MMT

3% C MMT

1.5% C MMT

1% C MMT

End

o

Temperature (oC)

30 60 90 120 150 180 210

Pure PLLA

1% NaMMT

9% NaMMT

5% NaMMT

3% NaMMT

End

o

Temperature (oC)

Fig. 13. DSC heating thermographs of quenched specimens of PLLA (a) nanocomposites and (b) microcomposites with various loadingsof C16MMT and NaMMT, respectively.

0 5 10 15 20 25 30

90

95

100

105

110

Tc

(o C)

Clay content (wt,%)


PLLA+NaMMT

0 5 10 15 20 25 30

52

56

60

172

176

180

Tg


PLLA+NaMMT

Tem

pera

ture

(o C

)

Clay content (wt,%)

Tm

a b

Fig. 14. Effect of clay type and clay content on (a) the crystallization temperature (Tc), (b) the glass transition temperature (Tg) and themelting point (Tm) of PLLA.

Table 1Mechanical characteristics of the prepared hybrids

Sample Yield strength (MPa) Strength at break (MPa) Elongation at break (%) Young’s modulus (GPa)

Pure PLLA 46.2 ± 0.75 35.2 ± 1.1 62 ± 4 2.05 ± 0.05

Nanocomposite PLLA

1% C16MMT 49.7 ± 0.9 41.6 ± 0.6 66 ± 9 2.29 ± 0.073% C16MMT 52,5 ± 1 44 ± 0.53 52 ± 5 2.55 ± 0.065% C16MMT 50 ± 0.6 43.1 ± 0.7 34 ± 2 2.81 ± 0.099% C16MMT 48.5 ± 1.1 42.2 ± 0.8 6.5 ± 0.3 3.18 ± 0.11

Microcomposite PLLA

1% NaMMT 48.2 ± 0.8 39 ± 0.7 50 ± 5 2.23 ± 0.13% NaMMT 49 ± 1.2 42 ± 0.9 30 ± 3 2.38 ± 0.15% NaMMT 48.4 ± 0.6 41.2 ± 0.3 10 ± 1 2.5 ± 0.069% NaMMT 45.4 ± 0.5 39.5 ± 0.5 4.5 ± 0.02 2.51 ± 0.08


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

clay content (around 1 wt%). Increase of the fillercontent restricts the polymer ductility leading to arapid decrease of the elongation of the hybrid mate-rial. In fact, at low organoclay loadings (<5 wt%)exfoliation of the silicate layers prevails. Theappearance of inorganic nanoparticles may restrictthe molecular mobility of the surrounding matrixmaterial leading to hybrid embrittlement [30].Increase in clay surface area due to their nanosizeand high aspect ratio can contribute to toughnessenhancement provided the clay/matrix interface issufficiently strong. Therefore, the rise in plasticdeformation is expected to be larger with decreasingscale of reinforcement [31]. As the filler quantityincreases, the amount of the delaminated sheets(i.e. platelets) decreases. However, the presence ofnon-exfoliated aggregates triggers a substantialdecrease in hybrid’s elongation at break. Gersappe[32] suggested that failure in polymer nanosompos-ites is strongly dependent on the mobility of thenanofiller particle rather than its surface area. Thismobility is a complex function of the filler size, theattraction between the polymer and the filler, andthe thermodynamic state of the matrix.

In the case of the conventional composites, theelongation is reduced substantially. The poor inter-action between the inorganic material and thematrix leads to a low clay dispersion and inadequatepolymer chain linkage to the mineral surface.Consequently, large clay aggregates formed canserve as load concentrating defects resulting in animportant reduction of the material ductility. Fur-thermore, these agglomerates can promote the trap-ping of air during sample preparation leading to theformation of voids that act as damage initiationsites [33].

Even though the ASTM 1708 method is not rec-ommended for the determination of the Young’smodulus, an estimation of the composite’s morphol-ogy on it can be obtained. As shown in Table 1, thepresence of organoclay significantly improves themodulus of elasticity of the composite materialand this modulus increases with filler content. Thus,a 50% rise of the nanocomposite’s modulus can beobtained at a clay loading of 9 wt%. This may beattributed to the structure of the silicate layers,which have a high aspect ratio and a large surfacearea, as well as to the interaction of these ultrafineinorganic particles with the matrix. Strawheckerand Manias [34] suggested that polymer stiffensthrough its affinity and adhesion to the filler sur-faces. Furthermore, the increase of the effective par-

ticle volume fraction caused by the increase of theinterlamellar spacing within each particle may playan important role in the increase of polymer stiffness[35,36].

The introduction of the natural clay enhances theelastic modulus of the hybrid material to a lesserextent than the organoclay does. Conventional com-posites displayed limited modulus increment withthe increase of the inorganic clay content up to5 wt%. Beyond this loading, the additional clay par-ticles were introduced in polymer regions that hadbeen already affected by other montmorilloniteaggregates and, thus, the Young’s modulus increasewas even smaller.

SEM observations of the tensile deformedcomposites’ microstructure in the vicinity of thepropagated crack affirmed a different deforma-tion process mechanism between the two types ofthe composites. In the case of nanocomposites(Fig. 15a and c) ‘micro’ cracks were formed on thesurface due to high miscibility of the organo-modi-fied inorganic particles with the matrix and excellentadhesion with the polymer. In contrast, the conven-tional composites (Fig. 15b and d) displayed largecavities on the deformed surface, suggesting a defi-cient bonding between the non-dispersed clay parti-cles and the polymer.

In a nanocomposite, the particle interlayer spac-ing is similar to or even smaller than the polymercrystalline lamellae thickness and, therefore, thepolymer confined between adjacent intercalated sil-icate layers is usually non-crystallizable [37]. Forthis reason, the matrix polymer molecules insertedinto the silicate galleries exhibit different propertiescompared to those of the bulk pristine polymer[38]. Moreover, the interaction between the silicatelayers is relatively weaker than that of the PLLA/organoclay interfacial bonding strength and thecohesive strength of PLLA. Therefore, during thedeformation process, microcrack initiation takesplace inside the weakly bonded clay [30,37,39,40].Accordingly, crack propagation occurs by splittingclay layers, breaking matrix ligaments and mergingmicrocracks. This argument is also supported by thefact that micro-voids were formed between the sili-cate layers of some samples during the preparationfor TEM observations, as shown in Fig. 16. Similarobservations were also mentioned by Nair et al. [31].Moreover, Chan et al. [41] reported that in polypro-pylene/CaCO3 nanocomposites failure initiationoccurred at the amorphous region of the polymerconfined between the CaCO3 particles.

Fig. 15. SEM micrographs of the surface morphology in the vicinity of a crack of tensile deformed PLLA hybrids containing 9 wt%C16MMT ((a) and (c)) and 9 wt% NaMMT ((b) and (d)).

Fig. 16. TEM micrograph of the nanocomposite material containing 5 wt% C16MMT. The white areas represent voids formed betweensilicate layers during the sample preparation process.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

When a conventional composite is subject to aload, stress concentrates in the vicinity of the largenon-dispersed particles due to the elastic modulusdifference between the clay and the polymer. When

the stress concentration reaches a critical value, voidformation takes place most likely by debonding atthe interface separating the unmodified clay andthe matrix. This debonding is possible because the

Fig. 17. SEM micrographs of the fracture surface of PLLA nanohybrids containing (a) 1 wt% and (b) 9 wt% of organo-modified MMT.The samples were broken after immersing them in liquid nitrogen.


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

adhesion in this interface is not due to strong inter-actions or chemical bonds. Larger inhomogeneitiesconcentrate much higher stresses locally and, there-fore, larger voids are formed around them. Decohe-sion occurring at the NaMMT/matrix interface wasalso reported recently [42,43] from studies of theproperties of Nylon 6 nanocomposites.

The fracture surface of the nanocomposites, bro-ken by using liquid nitrogen, has been examined bySEM. As it can be inferred from Fig. 17, increase ofthe organoclay content causes an increase in thefracture surface roughness of the material. When apropagating crack meets clay layers oriented per-pendicular to the crack propagation direction, itmay bend along the region adjoining the clay-matrixinterface. As the clay concentration increases, thedistance between intercalated clay particlesdecreases leading to a more tortuous path for thecrack [40,44,45]. Consequently, the surface rough-ness arising from the increased path length will alsoincrease, in accordance with the SEM observations.

4. Conclusions

The influence of the micro- and nano-structureon the thermomechanical characteristics ofpoly(L-lactic acid)/montmorillonite (PLLA/MMT)hybrid materials, which were prepared by the solu-tion casting method, has been investigated. XRD,TEM and AFM studies of the nanohybrids’ mor-phology showed that, at low clay content, exfolia-tion dominates. However, when the filler contentexceeds 5 wt%, complete exfoliation of the highaspect ratio organo-modified MMT particles intothe limited polymer mass is obstructed and a coex-

istence of exfoliated and intercalated clay particlesis observed.

The nanocomposites display enhanced thermalstability compared to microcomposites and to purepolymer, especially in oxidant atmosphere. The dis-persed inorganic platelets assist the formation of aprotective layer on the surface of the polymer,which acts as an insulator and a mass transport bar-rier retarding the decomposition rate of the polymermatrix. Moreover, the addition of organo-modifiedclay leads to a decrease in the crystallization temper-ature (Tc) and melting point (Tm) of the matrix. Thedispersed inorganic platelets serve as nucleatingagents facilitating polymer crystallization. However,high clay loadings retard the overall bulk crystalli-zation kinetics by hindering the mobility of polymerchains. The addition of the natural, unmodified claydoes not have any significant effect on both thermalparameters, Tc and Tm, underlying the poor interac-tion between the untreated clay and the polymermatrix. In addition, there is not any consistent trendregarding the effect of the filler on the glass transi-tion temperature (Tg) of the polymer in both thenanocomposite and the microcomposite.

The presence of organoclay results in a substan-tial enhancement of the nanocomposite’s mechanicalproperties. On the contrary, the microcompositesformed by the addition of the natural clay do notshow as much improvement in their mechanicalproperties as the nanocomposites do. The observa-tion of the deformed surfaces in the vicinity ofthe crack by the use of SEM reveals a different defor-mation process mechanism between the nano-composites and the microcomposites. In the firstcase, micro-cracks are formed possibly due to the


MA

CR

OM

OL

EC

UL

AR

NA

NO

TE

CH

NO

LO

GY

delamination between intercalated clay layers, whilethe microcomposites display a big number of largecavities suggesting that a decohesion occurs at theinterface between PLLA and MMT agglomerates.

Acknowledgments

The financial support of the Greek Ministry ofEducation and the European Union through theresearch grant Archimedes (MIS 86383) is gratefullyacknowledged.

The help of Veeco Co. is also acknowledged andespecially that of Drs. Hartmut Stadler, NataliaErina and Rian Li, who took the AFM images.

References

[1] Ray SS, Okamoto M. Prog Polym Sci 2003;28:1539.[2] Cho JW, Paul DR. Polymer 2001;42:1083.[3] Lu CS, Mai YW. Phys Rev Lett 2005;95:088303.[4] Gilman JW, Jackson CL, Morgan AB, Harris R, Manias E,

Giannelis EP, et al. Chem Mater 2000;12:1866.[5] Jang BN, Costache M, Wilkie CA. Polymer 2005;46:10678.[6] Ray SS, Bousmina M, Okamoto K. Macromol Mater Eng

2005;290:759.[7] Ikada Y, Tsuji H. Macromol Rapid Comm 2000;21:117.[8] Drumright RE, Gruber PR, Henton DE. Adv Mater

2000;12:1841.[9] Shi H, Lan T, Pinnavaia TJ. Chem Mater 1996;8:1584.

[10] Pluta M, Galeski A, Alexandre M, Paul MA, Dubois P. JAppl Polym Sci 2002;86:1497.

[11] Moraru VN. Appl Clay Sci 2001;19:11.[12] LeBaron PC, Wang Z, Pinnavaia TJ. Appl Clay Sci

1999;15:11.[13] Giannelis EP. Adv Mater 1996;8:29.[14] Di YW, Iannace S, Di Maio E, Nicolais L. J Polym Sci Part

B-Polym Phys 2005;43:689.[15] Nam PH, Ninomiya N, Fujimori A, Masuko T. Polym Eng

Sci 2006;46:39.[16] Nam JY, Ray SS, Okamoto M. Macromolecules 2003;

36:7126.[17] Chang JH, An YU, Cho DH, Giannelis EP. Polymer 2003;

44:3715.[18] Krikorian V, Pochan DJ. Chem Mater 2003;15:4317.

[19] Lee JH, Park TG, Park HS, Lee DS, Lee YK, Yoon SC,Nam JD. Biomaterials 2003;24:2773.

[20] Strawhecker KE, Kumar SK, Douglas JF, Karim A.Macromolecules 2001;34:4669.

[21] Lagaly G. Solid State Ionics 1986;22:43.[22] Hackett E, Manias E, Giannelis EP. J Chem Phys 1998;

108:7410.[23] Swartzen Sl, Matijevi E. Chem Rev 1974;74:385.[24] Xu SH, Boyd SA. Environ Sci Technol 1995;29:312.[25] Zhang ZZ, Sparks DL, Scrivner NC. Environ Sci Technol

1993;27:1625.[26] Paul MA, Alexandre M, Degee P, Henrist C, Rulmont A,

Dubois P. Polymer 2003;44:443.[27] Zanetti M, Camino G, Reichert P, Mulhaupt R. Macromol

Rapid Comm 2001;22:176.[28] Berta M, Lindsay C, Pans G, Camino G. Polym Degrad

Stabil 2006;91:1179.[29] Fornes TD, Paul DR. Polymer 2003;44:3945.[30] Gam KT, Miyamoto M, Nishimura R, Sue HJ. Polym Eng

Sci 2003;43:1635.[31] Nair SV, Goettler LA, Lysek BA. Polym Eng Sci 2002;

42:1872.[32] Gersappe D. Phys Rev Lett 2002;89:058301.[33] Singh RP, Zhang M, Chan D. J Mater Sci 2002;37:781.[34] Strawhecker KE, Manias E. Chem Mater 2000;12:2943.[35] Kornmann X, Thomann R, Mulhaupt R, Finter J, Berglund

L. J Appl Polym Sci 2002;86:2643.[36] Sheng N, Boyce MC, Parks DM, Rutledge GC, Abes JI,

Cohen RE. Polymer 2004;45:487.[37] Dennis HR, Hunter DL, Chang D, Kim S, White JL, Cho

JW, Paul DR. Polymer 2001;42:9513.[38] Vaia RA, Sauer BB, Tse OK, Giannelis EP. J Polym Sci Part

B-Polym Phys 1997;35:59.[39] Dasari A, Yu ZZ, Mai YW, Hu GH, Varlet JL. Compos Sci

Technol 2005;65:2314.[40] Wang K, Chen L, Wu JS, Toh ML, He CB, Yee AF.

Macromolecules 2005;38:788.[41] Chan CM, Wu JS, Li JX, Cheung YK. Polymer 2002;

43:2981.[42] Masenelli-Varlot K, Reynaud E, Vigier G, Varlet J. J Polym

Sci Part B-Polym Phys 2002;40:272.[43] Xie SB, Zhang SM, Zhao B, Qin HL, Wang FS, Yang MS.

Polym Int 2005;54:1673.[44] Zerda AS, Lesser AJ. J Polym Sci Part B-Polym Phys 2001;

39:1137.[45] Lee J, Yee AF. Polymer 2000;41:8363.

marras nanostruct vs microst

Science