Biodegradable Polymer/LayeredSilicate Nanocomposites ......Biodegradable Polymer/Layered Silicate Nanocomposites A Review 3 ﬂammability [6], and increased biodegradability of biodegradable

CHAPTER 8

BiodegradablePolymer/Layered SilicateNanocomposites: A Review

Masami OkamotoAdvanced Polymeric Nanostructured Materials Lab,Graduate School of Engineering, Toyota Technological Institute,Hisakata 2-12-1, Tempaku, Nagoya 468 8511, Japan

CONTENTS

1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 22. History of PLS Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . 33. Structure of LS and its Modification with Surfactants . . . . . . . . . 44. Preparation Methods and Structure of

PLS Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 54.1. Intercalation of Polymers or

Prepolymers from Solution . . . . . . . . . . . . . . . . . . . . . . . . 54.2. The In Situ Intercalative Polymerization Method . . . . . . . . 54.3. The Melt Intercalation Method . . . . . . . . . . . . . . . . . . . . . 64.4. Structure of PLS Nanocomposites . . . . . . . . . . . . . . . . . . . 6

5. Characterization of PLS Nanocomposites . . . . . . . . . . . . . . . . . 76. Biodegradable Polymer/LS Nanocomposites . . . . . . . . . . . . . . . 8

6.1. PCL/LS Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . 96.2. PVA/LS Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . 106.3. PLA/LS Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . 106.4. PBS/LS Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . 136.5. PHB/LS Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . 146.6. Starch/LS Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . 156.7. Plant Oil/LS Nanocomposites . . . . . . . . . . . . . . . . . . . . . 156.8. Chitosan/LS Nanocomposites . . . . . . . . . . . . . . . . . . . . . 17

7. Materials Properties of Biodegradable Polymer/LSNanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 187.1. Mechanical Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . 19

ISBN: 1-58883-053-5/$00.00Copyright © 2005 by American Scientific PublishersAll rights of reproduction in any form reserved.

1

Handbook of Biodegradable PolymericMaterials and Their Applications

Edited by Surya Mallapragada and Balaji NarasimhanVolume 1: Pages (1–45)

2 Biodegradable Polymer/Layered Silicate Nanocomposites A Review

8. Crystallization of Biodegradable Polymer/LSNanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31

9. Melt Rheology of Biodegradable Polymer/LSNanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 349.1. Dynamic Oscillatory Shear Measurement . . . . . . . . . . . . . 349.2. Steady Shear Flow . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 379.3. Elongational Flow and Strain-Induced Hardening . . . . . . . 38

10. Processing Operations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3911. Creating Porous Ceramic Materials Via PLA/LS

Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4112. Current Status and Future Prospects of

Biodegradable Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . 43References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43

1. INTRODUCTION

Advanced technology in petrochemical-based polymers has brought many benefits tomankind. However, it has become clear that nondegradable plastic materials used for dispos-able applications are significantly disturbing and damaging the Earth’s ecosystem. The envi-ronmental impact of persistent plastic wastes is of increasing global concern, and alternativedisposal methods are limited. Incineration of these plastic wastes always produces a largeamount of carbon dioxide that contributes to global warming; in some cases, toxic gases arealso produced, which contributes to global pollution. On the other hand, satisfactory landfillsites are also limited. Another disadvantage of nondegradable plastic materials is that theEarth has finite resources in terms of fossil origin fuel.

For these reasons, there is an urgent need to develop renewable, source-based, environ-mentally benign polymeric materials (biopolymers), especially for use in short-term packag-ing and disposable applications. Such materials would not involve the use of toxic or noxiouscomponents in their manufacture, and could allow for composting into naturally occurringdegradation products.

The ideal biopolymer is of renewable biological origin and biodegradable at the end ofits life. Biopolymers include polysaccharides such as cellulose and starch; carbohydoratepolymers produced by bacteria and fungi [1]; and animal protein-based biopolymers suchas wool, silk, gelatin, and collagen. On the other hand, poly(vinyl alcohol) (PVA), poly(�-caprolactone) (PCL), and poly(butylene succinate) (PBS) are examples of polymers thathave synthetic origin but are biodegradable.

In today’s commercial venues, biopolymers have proven to be relatively expensive andavailable only in small quantities. This has led to limited applications to date. However,there are signs that this is changing, with increasing environmental awareness and morestringent legislation regarding recyclability and restrictions on waste disposal. Cargill Dowhas a polylactide (PLA) in production (Natureworks™), and Metabolix has been developingpolyhydroxyalkanoate (PHA) (Biopol™).

Thus, the increasing appreciation of the various intrinsic properties of biopolymers, cou-pled with the knowledge of how such properties can be improved to achieve compatibilitywith thermoplastics processing, manufacturing, and end-use requirements, has fueled tech-nological and commercial interest in biopolymers.

Of particular interest is a recently developed nanocomposite technology consisting of apolymer and layered silicate. This combination often exhibits remarkably improved mechan-ical and various other properties [2] when compared to pure polymers or conventionalcomposites (both micro- and macrocomposites). These improvements can include highmoduli [3], increased strength and heat resistance [4], decreased gas permeability [5] and

Biodegradable Polymer/Layered Silicate Nanocomposites A Review 3

flammability [6], and increased biodegradability of biodegradable polymers [7]. On the otherhand, these materials have also proven to be unique model systems for the study of thestructure and dynamics of polymers in confined environments [8].

The main reason for these improved properties is interfacial interaction between matricesand organically modified layered silicates (OMLSs), as opposed to conventional composites.Layered silicates (LSs) have layer thickness in the order of 1 nm and very high aspect ratios(e.g., 10–1,000). A small weight percentage of OMLSs that are properly dispersed throughoutthe matrix thus creates a much larger surface area for polymer-filler interfacial interactionsthan do conventional composites [9]. Although the intercalation chemistry of polymers (whenmixed with appropriately modified layered silicate and synthetic layered silicates [10, 11])has been understood for a long time, the study of polymer/LS (PLS) nanocomposites hasrecently gained greater momentum.

This review is intended to highlight the major developments in this area during the lastdecade. The different techniques used to prepare biodegradable PLS nanocomposites, theirphysicochemical characterization, the improved materials properties that those materials candisplay, and the processing and probable applications of biodegradable PLS nanocompositeswill be reported in detail.

2. HISTORY OF PLS NANOCOMPOSITESThe first successful PLS nanocomposite appeared about ten years ago, through a pioneeringeffort of a research team at Toyota Central Research & Development Co., Inc. (TCRD).This material was a nylon 6/LS hybrid [12, 13].

Earlier attempts at preparing polymer/clay composites are described in patent literatureof the mid-1990s [14, 15]. In such cases, incorporation of 40–50 wt% LS (clay mineral)(bentonite, hectorite, etc.) into a polymer was attempted but produced unsatisfactory results:The maximal modulus enhancement was only around 200%, although the clay loading wasas much as 50 wt%. The poor results were due to failure to achieve good dispersion of clayparticles in the matrix, in which clay minerals existed as agglomerated tactoids. Such poordispersion of the clay particles could improve the material rigidity, but certainly sacrificedthe strength, the elongation at break, and the toughness of the material [14, 15].

A primary reason why it is impossible to improve tactoid dispersion into well-dispersedexfoliated monolayers of clay is due to the intrinsic incompatibility of hydrophilic layeredsilicates with hydrophobic engineering plastics. One attempt at circumventing this difficultywas made by Unitika Ltd. [16] about thirty years ago; those researchers prepared nylon 6/claycomposites (not nanocomposites) by in situ polymerization of �-caprolactam with montmo-rillonite, but the results were disappointing.

The first major breakthrough in this area was in 1987, when Fukushima and Inagaki ofTCRD, in their detailed study of polymer/layered silicate composites, persuasively demon-strated that lipophilization caused by replacing inorganic cations in galleries of native claywith alkylammonium surfactants successfully made them compatible with hydrophobic poly-mer matrices [17]. The modified clay was thus called lipophilized clay, organophilic clay, orsimply organoclay (OMLS). Furthermore, they found that the lipophilization enabled expan-sion of the clay galleries and exfoliation of the silicate layers into single layers 1 nm thick.

In 1993, Usuki, Fukushima, and their colleagues at TCRD successfully prepared, for thefirst time, an exfoliated nylon 6/LS hybrid (NCH) by in situ polymerization of �-caprolactam,in which alkylammonium-modified organoclay was thoroughly dispersed in advance [12, 13].The resulting composite (when only 4.2 wt% clay was loaded) possessed a doubled modulus;strength enhanced by 50%; and an 80�C increase in heat distortion compared to neat nylon 6,as shown in Table 1. This produced a new generation of engineering materials, which wecall “polymer/LS nanocomposites.”

Thus, along the stream of development in PLS nanocomposite technologies, many studieshave been devoted to PLS nanocomposites with intrinsically excellent polymer propertiesthat should have an attractive potential for continuous expansion of application versatility.


Table 1. Mechanical and thermal properties of a nylon 6/LS hybrid.

Nylon 6 NeatProperties Nanocomposite Nylon 6

Clay content (wt%) 4.2 0Specific gravity 1.15 1.14Tensile strength (Mpa) 107 69Tensile modulus (Gpa) 2.1 1.1Impact (kJ/m2) 2.8 2.3HDT (�C at 1.8 Mpa) 147 65

3. STRUCTURE OF LS AND ITS MODIFICATIONWITH SURFACTANTS

The clays commonly used for the preparation of PLS nanocomposites are in the same generalfamily of 2:1 layered or phyllosilicates (see Table 2). Their crystal structure consists of lay-ers composed of two silica tetrahedrals fused to an edge-shared octahedral sheet of eitheraluminium or magnesium hydroxide. The layer thickness is ∼1 nm and the lateral dimen-sions of these layers may vary from 30 nm to several microns and even larger, dependingon the particular layered silicate. Stacking of the layers leads to a regular van der Waal’sgap between the layers, called the interlayer or gallery. Isomorphic substitution within thelayers (for example, Al+3 replaced by Mg+2 or by Fe+2, or Mg+2 replaced by Li+1) generatesnegative charges that are counterbalanced by alkali and alkaline earth cations situated insidethe galleries, as shown in Figure 1.

The most commonly used layered silicates are montmorillonite (MMT, Mx(Al4−xMgx)·Si8O20(OH)4) hectorite (Mx(Mg6−xLix)Si8O20(OH)4), and saponite (Mx(Si8−xAlx)Si8O20·(OH)4(x = 0.3–1.3). This type of clay is characterized by a moderate surface charge (cationexchange capacity: CEC of 80–120 mequiv/100 gm) and layer morphology. These claysare only miscible with hydrophilic polymers, such as poly(ethylene oxide) (PEO) [18] andpoly(vinyl alcohol) (PVA) [17]. To improve miscibility with other polymer matrices, one mustconvert the normally hydrophilic silicate surface to organophilic, which enables intercalationof many engineering polymers. Generally, this can be done by ion-exchange reactions withcationic surfactants, including primary, secondary, tertiary, and quaternary alkylammoniumor alkylphosphonium cations. The role of alkylammonium or alkylphosphonium cations inthe organosilicates is to lower the surface energy of the inorganic host and improve thewetting characteristics with the polymer matrix; this results in larger interlayer spacing. Onecan demonstrate that about 100 alkylammonium salt molecules are localized near the indi-vidual silicate layers (∼8 × 10−15 m2) and active surface area (∼800 m2/g). Additionally,the alkylammonium or alkylphosphonium cations could provide functional groups that canreact with the polymer matrix or, in some cases, initiate the polymerization of monomers toimprove the strength of the interface between the inorganic and the polymer matrix [10, 19].

Vaia and Giannelis [20] have shown that alkyl chains can vary from liquid-like to solid-like,with the liquid-like structure dominating as the interlayer density or chain length decreases(see Fig. 2), or as the temperature increases by using Fourier transform infrared spectroscopy(FTIR). This is understandable because of the relatively small energy differences betweenthe trans and gauche conformers; the idealized models described earlier assume all transconformations. In addition, for the longer chain length surfactants, the surfactants in thelayered silicate can show thermal transition akin to melting or liquid-crystalline to liquid-liketransitions upon heating.

Table 2. Chemical formulas and characteristic parameters of commonly used 2:1 phyllosilicates.

2:1 Phyllosilicates Chemical Formulaa CEC (mequiv/100 gm) Particle Length (nm)

Montmorillonite Mx(Al4−xMgx)Si8O20(OH)4 110 100–150Hectorite Mx(Mg6−xLix)Si8O20(OH)4 120 200–300Saponite MxMg6(Si8−xAlx)Si8O20(OH)4 86.6 50–60

aM = monovalent cation; x = degree of isomorphous substitution (between 0.3 and 1.3).


Bas

al s

paci

ng

Tetrahedral

Tetrahedral

Exchangeable cations

Octahedral

Al, Fe, Mg, Li

Li, Na, Rb, Cs

O

OH

~1 nm

Figure 1. Structure of 2:1 phyllosilicates.

4. PREPARATION METHODS AND STRUCTURE OFPLS NANOCOMPOSITES

There is considerable literature available devoted to developing PLS nanocomposites with dif-ferent combinations of OMLS and matrix polymers, employing somewhat different technolo-gies appropriate to each. The technologies are broadly classified into three main categories.

4.1. Intercalation of Polymers orPrepolymers from Solution

This technology is based on a solvent system in which polymers or prepolymers are solubleand the silicate layers are swellable. The layered silicate is first swollen in a solvent such aswater, chloroform, or toluene, etc. When the polymer and layered silicate solutions are mixed,the polymer chains intercalate and displace the solvent within the interlayer of the silicate.Upon solvent removal, the intercalated structure remains, resulting in PLS nanocomposites.

4.2. The In Situ Intercalative Polymerization Method

In this method, the OMLS is swollen within the liquid monomer or a monomer solutionso the polymer formation can occur between the intercalated sheets. Polymerization can beinitiated either by heat or radiation, by the diffusion of a suitable initiator, or by an organic

(a) (b) (c)

Figure 2. Alkyl chain aggregation models: (a) short chain lengths, where the molecules are effectively isolated fromeach other; (b) medium chain lengths, where quasi-discrete layers form with various degrees of in-plane disor-der and interdigitation between the layers; and (c) long chain lengths, where interlayer order increases, leadingto a liquid-crystalline polymer environment. Open circles represent the CH2 segments while cationic head groupsare represented by filled circles. Adapted from Ref. 2, M. Okamoto, “Encyclopedia of Nanoscience and Nano-technology,” with permission from American Scientific Publishers.


initiator or catalyst fixed through cation exchange inside the interlayer before the monomerswelling step.

4.3. The Melt Intercalation Method

This method involves annealing, statically or under shear, a mixture of the polymer andOMLS above the softening point of the polymer. This method has great advantages overeither in situ intercalative polymerization or polymer solution intercalation. First, this methodis environmentally benign due to the absence of organic solvents. Second, it is compatiblewith current industrial processes such as extrusion and injection molding. The melt inter-calation method allows the use of polymers which were previously not suitable for in situpolymerization or the solution intercalation method.

Other possibilities are exfoliation-adsorption [21, 22] and template synthesis [23]. Nowa-days, this solvent-free method is much preferred for practical industrial material productiondue to its high efficiency and its possible avoidance of environmental hazards.

4.4. Structure of PLS Nanocomposites

As described by Kojima et al. [12] and in a 1950 U.S. patent [13], in situ polymerization wasemployed for the first time in NCH production, and the melt intercalation was the directblending of OMLS into a modified polymer matrix such as used in polypropylene (PP)/LSnanocomposites [24]. Since Giannelis and his colleagues [25] found that melt-compoundingof polymers with clay is possible without using organic solvents, nanocomposite preparationsusing this method have been widely used, especially for polyolefin-based nanocomposites.This process involves annealing, statically or under shear, a mixture of the polymer andOMLS above the softening point of the polymer. During the annealing, the polymer chainsdiffuse from the bulk polymer melt into the galleries between the silicate layers. Dependingon the degree of penetration of the matrix into the organically modified layered silicate gal-leries, nanocomposites are obtained with structures ranging from intercalated to exfoliate.Polymer penetration resulting in finite expansion of the silicate layers produces intercalatednanocomposites consisting of well-ordered multilayers with alternating polymer/silicate lay-ers and a repeat distance of few nanometers (intercalated, see Fig. 3) [24]. On the other

1 nmL

dclay

Lclay

ξclay

One Clay PlateletL: 100 – 200 nm in case of MMT

Form factors of dispersed clay

Intercalated Intercalated-and-flocculated Exfoliated

The structure of 2:1 layered silicates

Al, Fe, Mg, Li

Li, Na, Ra, Cs

Exchangeable cations

Tetrahedral

Octahedral

Tetrahedral

OH

O

Figure 3. Schematic illustration of three different types of thermodynamically achievable polymer/clay nano-composites.


hand, extensive polymer penetration resulting in disordered and eventual delamination ofthe silicate layers produces nearly exfoliated nanocomposites consisting of individual silicatelayers dispersed in the polymer matrix (exfoliated) [26]. Under some conditions, the inter-calated nanocomposites exhibit flocculation because of the hydroxylated edge–edge interac-tion of silicate layers (intercalated-and-flocculated). The length of the oriented collections(in the range of 300–800 nm) is far larger than that of the original clay (mean diameter∼150 nm) [27, 28]. Such flocculation presumably is governed by an interfacial energy betweenthe polymer matrix and OMLS and is controlled by ammonium cation-matrix polymer inter-action. The polarity of the matrix polymer is of fundamental importance in controlling thenanoscale structure.

5. CHARACTERIZATION OF PLSNANOCOMPOSITES

The structure of PLS nanocomposites has typically been established using a wide-anglex-ray diffraction (WAXD) analysis and transmission electron microscope (TEM) observa-tions. Due to its availability and ease of use, WAXD is most commonly used to probe PLSnanocomposite structures and, sometimes, to study the kinetics of the polymer melt interca-lation. By monitoring the position, shape, and intensity of the basal reflections from the dis-tributed silicate layers, the PLS nanocomposite structure (either intercalated or exfoliated)may be identified. For example, in exfoliated nanocomposites, the extensive layer separationassociated with the delamination of the original silicate layers in the polymer matrix resultsin the eventual disappearance of any coherent x-ray diffraction from the distributed silicatelayers. On the other hand, for intercalated nanocomposites, the finite layer expansion asso-ciated with the polymer intercalation results in the appearance of a new basal reflectioncorresponding to the larger gallery height.

WAXD offers a convenient method to determine the interlayer spacing of the silicate lay-ers in the original layered silicates and in the intercalated nanocomposites (within 1–4 nm),but little can be said about the spatial distribution of the silicate layers or any structuralinhomogeneities in the PLS nanocomposites. Additionally, some layered silicates initiallydo not exhibit well defined basal reflection. Thus, peak broadening and intensity decreasesare very difficult to study systematically. Therefore, only tentative conclusions can be drawnconcerning the mechanism of nanocomposite formation and their structure based solely onWAXD patterns.

On the other hand, TEM allows a qualitative understanding of the internal structure,spatial distribution of the various phases, and defect structure of nanocomposites throughdirect visualization. However, special care must be exercised to ensure that a representativecross section of the sample is evaluated. The WAXD patterns and corresponding TEMimages of three different types of nanocomposites are presented in Figure 4.

The solid-state nuclear magnetic resonance (NMR) method of quantifying the level ofclay exfoliation is also a very important facet of nanocomposite characterization. The mainobjective in solid-state NMR measurement is to connect the measured longitudinal relax-ation (TH

1 s) of protons and 13C nuclei with the quality of clay dispersion. The extent ofand the homogeneity of the dispersion of the silicate layers within the polymer matrix arevery important for determining physical properties. The surfaces of naturally occurring lay-ered silicates such as MMT are mainly made of silica tetrahedrals, while the central planeof the layers contains octahedrally coordinated Al3+ (see Fig. 1 and Table 2) with frequentnonstoichiometric substitutions, where an Al3+ is replaced by Mg2+ and, somewhat less fre-quently, by Fe3+. Typical concentrations of Fe3+ (spin = 5/2) in naturally occurring claysproduce nearest neighbor Fe–Fe distances of about 1.0–1.4 nm [29]. At such distances, thespin exchange interaction between the unpaired electrons on different Fe atoms is expectedto produce magnetic fluctuations in the vicinity of the Larmor frequencies for protons or 13Cnuclei [29]. The spectral density of these fluctuations is important because the TH

1 of protonsand 13C nuclei within ∼1 nm of the clay surface can be directly shortened. For protons,if that mechanism is efficient, relaxation will also propagate into the bulk of the polymerby spin diffusion. Thus, this paramagnetically induced relaxation will influence the overall


200 nm

200 nm

200 nm

0

500

1000

1500

2 6 10

Exfoliated

0

500

1000

1500Intercalated-and-flocculated

1000

1500Intercalated

0

500

1000

1500

2000

Original OMLS

2Θ/degrees

Inte

nsity

/A.U

.Intercalated

Intercalated-and-flocculated

Exfoliated

500

0

4 8

Figure 4. (a) WAXD patterns, and (b) TEM images of three different types of nanocomposites.

measured TH1 to an extent that will depend both on the Fe concentration in the clay layer

and, more importantly, on the average distances between clay layers. The latter dependencesuggests a potential relationship between measured TH

1 values and the quality of the claydispersion. If the clay particles are stacked and poorly dispersed in the polymer matrix, theaverage distances between polymer-clay interfaces are greater, and the average paramagneticcontribution to TH

1 is weaker. VanderHart et al. [30] also employed the same arguments tounderstand the stability of a particular OMLS under different processing conditions.

6. BIODEGRADABLE POLYMER/LS NANOCOMPOSITES

Recently, some groups have undertaken the preparation and characterization of the materi-als properties of various kinds of biodegradable polymer/LS nanocomposites having proper-ties suitable for a wide range of applications. To date, reported biodegradable polymers forthe preparation of nanocomposites are:

• poly(�-caprolactone) (PCL) [31–35]• poly(vinyl alcohol) (PVA)[36, 37]• poly(lactide) (PLA) [38–48],• poly(butylene succinate) (PBS) [49–51],• unsaturated polyester [52],• poly(hydroxy butyrate) (PHB) [53, 54],• aliphatic polyester [55–58],• thermoplastic starch [59, 60]• other renewable resources [61, 62]


6.1. PCL/LS Nanocomposites

For the preparation of poly(�-caprolactone) (PCL)-based nanocomposites, Messersmithand Giannelis [31] modified MMT with protonated aminolauric acid and dispersed themodified MMT in liquid �-caprolactone (CL) before polymerizing at high temperature.The nanocomposites were prepared by mixing up to 30 wt% of the modified MMT with driedand freshly distilled CL for a couple of hours, followed by ring-opening polymerization understirring at 170�C for 48 hours. The same researchers [32] have also reported on �-caprolactonepolymerization inside a Cr+3-exchanged fluorohectorite at 100�C for 48 hours.

Pantoustier and his colleagues [33] used this in situ intercalative polymerization method.They used both pristine MMT and �-amino dodecanoic acid-modified MMT for the com-parison of prepared nanocomposite properties. For nanocomposite synthesis in a polymer-ization tube, the desired amount of pristine MMT was first dried under vacuum at 70�Cfor three hours. A given amount of CL was then added under nitrogen and the reactionmedium was stirred at room temperature for one hour. A solution of initiator (Sn(Oct)2) orBu2Sn(Ome)2) in dry toluene was added to the mixture in order to reach a [monomer]/[Sn]molar ratio equal to 300. The polymerization was then allowed to proceed for 24 hoursat room temperature. After polymerization, a reverse ion-exchange reaction was used toisolate the PCL chains from the inorganic fraction of the nanocomposite. A colloidal sus-pension was obtained by stirring 2 g of the nanocomposite in 30 mL of THF for two hoursat room temperature. Separately, a solution of 1 wt% of LiCl in THF was prepared. Thenanocomposite suspension was added to 50 mL of the LiCl solution and left to stir at roomtemperature for 48 hours. The resulting solution was centrifuged at 3,000 rpm for 30 minutes.The supernatant was then decanted and the remaining solid was washed by dispersing in30 mL of THF, followed by centrifugation. The combined supernatant was concentrated andprecipitated from petroleum ether.

The polymerization of CL with pristine MMT gives PCL with a molar mass of 4,800 g/moland a narrow distribution. For comparison, the researchers also conducted the same exper-iment without MMT, but there was no polymerization of CL. These results demonstratethe ability of MMT to catalyze and control CL polymerization, at least in terms of molecu-lar weight distribution that remains remarkably narrow. For the mechanism of polymeriza-tion, the researchers assumed that the CL is activated through the interaction with acidicsite on the clay surface and the polymerization is likely to be proceeding via the activatedmonomer mechanism by the cooperative function of Lewis acidic aluminum and Bronstedacidic silanol functionalities on the initiator walls (see Fig. 5). On the other hand, in the poly-merization of CL with the protonated �-amino dodecanoic acid-modified MMT, the molarmass was 7,800 g/mol with a monomer conversion of 92% and, again, a narrow molecularweight distribution. The WAXD patterns of both nanocomposites indicate the formationof intercalated structure. In another very recent publication [34, 35], same group prepared

100 nm

Figure 5. TEM image of a PCL nanocomposite containing 3 wt% MMT-Alk. Reprinted with permission from [66],B. Lepoittevin et al., Polymer 43, 4017 (2002). © 2002, Elsevier Science Ltd.


PCL/MMT nanocomposites by using in situ ring-opening polymerization of CL using dibutyltin dimethoxide as an initiator/catalyst.

6.2. PVA/LS Nanocomposites

More recently, Strawhecker and Manias [36] used this method in attempts to producepoly(vinyl alcohol) (PVA)/MMT nanocomposite films. PVA/MMT nanocomposite films werecast from MMT/water suspensions where PVA was dissolved. Room temperature distilledwater was used to form a suspension of Na+-MMT. The suspension was first stirred forone hour and then sonicated for 30 minutes. Low viscosity, fully hydrolyzed atactic PVA wasthen added to the stirring suspensions such that the total solid (silicate plus polymer) was≤5 wt%. The mixtures were then heated to 90�C to dissolve the PVA, again sonicated for30 minutes, and, finally, films were cast in a closed oven at 40�C for two days. The recov-ered cast films were then characterized by both WAXD and TEM. Both the d-spacing andtheir distribution decreased systematically with increasing MMT wt% in the nanocomposites.The TEM photograph of 20 wt% clay containing nanocomposites reveals the coexistence ofsilicate layers in the intercalated and the exfoliated states.

6.3. PLA/LS Nanocomposites

Okamoto and his colleagues [38, 40] first reported the preparation of intercalated PLA/LSnanocomposites. For nanocomposite (PLACN) preparation, C18-MMT and PLA were firstdry-mixed by shaking them in a bag. The mixture was then melt-extruded by using a twin-screw extruder operated at 190�C to yield very light gray colored strands of PLACNs.Nanocomposites loaded with very small amounts of oligo-PCL (Mw = 2,000 g/mol) as acompatibilizer were also prepared in order to understand the effect of oligo-PCL on themorphology and properties of PLACNs [38]. The compositions of various nanocompositesof PLA with C18-MMT are summarized in Table 3. WAXD patterns of a series of nanocom-posites are shown in Figure 6. Figure 7 shows TEM photographs of nanocomposites corre-sponding to the WAXD patterns. On the basis of WAXD analyses and TEM observations,they calculated form factors (see Table 4) [i.e., average length (Lclay), thickness (dclay)] ofthe stacked intercalated silicate layers, and the correlation length (�clay) between them (seeFig. 3). These data clearly established that silicate layers of the clay were intercalated andrandomly distributed in the PLA matrix. Incorporation of very small amounts of oligo-PCLas a compatibilizer in the nanocomposites led to a better parallel stacking of the silicatelayers and also much stronger flocculation due to the hydroxylated edge–edge interaction ofthe silicate layers. Due to the interaction between clay platelets and the PL-matrix in pres-ence of very small amounts of oligo-PCL, the disk–disk interaction plays an important role

Table 3. Composition and characteristic parameters of various PLACNs based on PLA, oligo-PCL, andC18-MMT.

Composition (wt%)

Sample PLA oligo-PCL C18-MMTb Mw × 10−3(g/mol) Mw/Mn Tg(�C) Tm(�C) cc (%)

PLACN1 97 3 [2.0] 178 1.81 60.0 169 50.65PLACN2 95 5 [3.0] 185 1.86 60.0 170 39.01PLACN3 93 7 [4.8] 177 1.69 59.8 170 43.66PLACN4 94.8 0.2 5 [3.3] 181 1.76 58.6 170 41.47PLACN5 94.5 0.5 5 [3.3] 181 1.76 57.6 169 32.91PLACN6 93 2.0 5 [2.8] 180 1.76 54.0 168 —PLACN7 92 3.0 5 [2.4] 181 1.77 51.0 168 —PLAa 100 187 1.76 60.0 168 36.24PLA1 99.8 0.2 180 1.76 58.0 168.5 46.21PLA2 99.5 0.5 180 1.76 57.0 168.8 52.51PLA3 98 2.0 180 1.76 54.7 169 —

aMw and PDI of extruded PLA (at 190�C) are 180 × 103 (g/mol) and 1.6, respectively. bValue in the parenthesesindicates the amount of clay (inorganic part) content after burning. cThe degree of crystallinity.


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Figure 6. WAXD patterns for C18-MMT and various PLACNs: (a) without oligo-PCL, and (b) with oligo-PCL.The dashed line in each figure indicates the location of the silicate (001) reflection of C18-MMT. The asterisksindicate the (001) peak for C18-MMT dispersed in PLA matrices. Reprinted with permission from [28], S. Sinha Rayet al., Macromolecules 35, 3104 (2002). © 2002, American Chemical Society.

in determining the stability of the clay particles and, hence, the enhancement of mechanicalproperties of such nanocomposites.

In their further research [41, 42, 48], this group prepared a series of PLACNs with varioustypes of organoclay in order to investigate the effect of organoclay on the morphology, prop-erties, and biodegradability of PLACNs. Four different types of pristine layered silicates wereused and each of them was modified with a different type of surfactant. Detailed specifications

(a) (b)

(c) (d)

2 µm 2 µm

500 nm 500 nm

Figure 7. TEM bright field images: (a) PLACN2 (×10,000); (b) PLACN4 (×10,000); (c) PLACN2 (×40,0000); and(d) PLACN4 (×40,000). The dark entities are cross sections of intercalated organoclay, and the bright areas arethe matrices. Reprinted with permission from [28], S. Sinha Ray et al., Macromolecules 35, 3104 (2002). © 2002,American Chemical Society.


Table 4. Comparison of form factors between PLACN2 and PLACN4obtained from WAXD patterns and TEM observations.

Form factors PLACN2 PLACN4

WAXDd001 (nm) 3.03 2.98dclay (nm) 13 10

TEMdclay (nm) 38 ± 17.25 30 ± 12.5Lclay (nm) 448 ± 200 659 ± 145Lclay/dclay 12 22�clay (nm) 255 ± 137 206 ± 92

of the various types of organoclay they used are presented in Table 5. On the basis ofWAXD analyses and TEM observations, the researchers successfully formed four differenttypes of PLACNs. Ordered intercalated-and-flocculated nanocomposites were obtained whenODA was used as the organoclay; disordered intercalated structures resulted in the case ofPLA/SBE4 nanocomposites; PLA/SAP4 nanocomposites showed near to exfoliate nanocom-posites; and the coexistence of stacked intercalated and exfoliated nanocomposite structureswas evident with PLA/MEE4 nanocomposites. Thus, the nature of OMLS has a strong effecton the final morphology of PLA-based nanocomposites.

In a very recent work, Okamoto and Maiti [43] prepared a series of PLACNs with threedifferent types of pristine layered silicate such as saponite, MMT, and mica, and each ofthem was modified with alkylphosphonium salts having various chain lengths. Their firstgoal was to determine the effect of alkylphosphonium modifiers of different chain lengthson the properties of organoclay and how the different clays behave differently having sameorganic modifier. Second, they wanted to determine the effects of dispersion, intercalation,and aspect ratio of clay on materials properties. From the resulting WAXD patterns, it wasclearly observed that the d-spacing (001) increases with increasing modifier chain lengthand, for a fixed modifier, it increases with increasing lateral dimensions of the clay particles.These researchers concluded that there are two reasons for this type of behavior: the CECvalue, and the lateral size of various pristine layered silicates. In both cases, layered silicatesfollowed the order mica > MMT > saponite. The CEC factor is more important than thelateral size of the silicates to control the d-spacing/stacking of silicate layers. Since mica hasa large lateral size and also a high amount of surfactant molecules due to its high CEC value,surfactant chains inside the integrally have restricted conformation due to physical jamming.These researchers believe there is less physical jamming in saponite due to its lower CECand smaller lateral size. The results for OMLS, based on TEM and WAXD analyses, areschematically illustrated in Figure 8.

Figure 9 compares the WAXD patterns of nanocomposites with different clay dimensionshaving the same clay (n-hexadecyl tri-n-butyl phosphonium bromide (C16)-modified) content(3 wt%). For MMT-based nanocomposites, the peaks are sharp and the crystallite sizes areslightly smaller than those of the corresponding organoclay, indicating an almost ordered

Table 5. Specifications and designations of OMLS used for the preparation of PLACNs.

Particle length CEC Organic salt used forClay Codes Pristine Clay (nm) mequiv/100 gm the modification of clay Suppliers

ODA MMT 150–200 110 Octadecyl Nanocor Inc., USAammonium cation

SBE MMT 100–130 90 Trimethyloctadecyl Hojun yoko Co.,ammonium cation Japan

MEE Synthetic 200–300 120 Dipolyoxyethylene CO-OP Chemicals,F-mica alkyl(coco) methyl Japan

ammonium cationSAP Saponite 50–60 86.6 Tributylhexadecyl CO-OP Chemicals,

phosphonium cation Japan


Smectite Montmorillonite Mica

200 nm

2.44

nm

50 nm

1.87

nm

70 nm

2.13

nm

Figure 8. Schematic representation of organoclays (OMLSs) with C16 ions. Reprinted with permission from [43],P. Maiti et al., Chem. Mater. 14, 4654 (2002). © 2002, American Chemical Society.

structure of MMT in the nanocomposites. The peaks of the nanocomposites prepared withmica clay are very sharp, similar to those of corresponding organoclay, and the slightly largercrystallite sizes indicate that the number of stacked silicate layers is the same as that of theoriginal organoclay. However, some amount of PLA is intercalated between the galleries,giving rise to a larger crystallite size. On the basis of WAXD patterns and crystallite size,stacking of silicate layers in the organoclays and in various nanocomposites prepared withthree different organoclays is presented schematically in Figure 10.

More recently, Dubois et al. [44, 46] reported the preparation of plasticized PLA/MMTnanocomposites. The OMLS they used was MMT modified with bis-(2-hydroxyethyl)methyl(hydrogenated tallowalkyl) ammonium cations. WAXD analyses have confirmed the forma-tion of intercalated nanocomposites (see Fig. 11).

6.4. PBS/LS Nanocomposites

Poly(butylene succinate) (PBS) is also an aliphatic thermoplastic polyester with many inter-esting properties, including biodegradability, melt processability, and thermal and chemicalresistance. Although the above properties suggest potential applications of PBS, some ofits other properties such as softness, gas barrier properties, flexural properties, etc. are fre-quently not adequate for a wide-range of applications.

K. Okamoto and M. Okamoto [49, 50] first reported the preparation of PBS /MMTnanocomposites (PBSCNs) by simple melt extrusion of PBS and OMLS, having proper-ties suitable for many applications. MMT modified with octadecylammonium chloride wasused as organoclay for nanocomposite preparation. In recent publications [50, 51], the sameresearchers also reported the details of their studies on structure-property relationshipsinvolving PBSCNs. Figure 12 represents the WAXD patterns of various PBSCNs. A TEMimage of the representative PBSCN is shown in Figure 13.

20

500

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1500

2000

4 62Θ /deg

Inte

nsity

/a.u

8 10

Smectite

MMT

Mica

C16

Figure 9. WAXD patterns of smectite (SAP), MMT, and mica nanocomposites with C16 organoclay (OMLS) andsome clay content (3 wt%). Reprinted with permission from [43], P. Maiti et al., Chem. Mater. 14, 4654 (2002).© 2002, American Chemical Society.


organoclays nanocomposites

Smectite

MMT

Mica

Figure 10. Schematic presentation of silicate layers in organoclay (OMLS) and in various nanocomposites.Reprinted with permission from [43], P. Maiti et al., Chem. Mater. 14, 4654 (2002). © 2002 American ChemicalSociety.

In 2002, Lee and his colleagues [55] reported the preparation of biodegradable aliphaticpolyester (APES)/organoclay nanocomposites using a melt intercalation method. Two kindsof organoclays, Cloisite 30B and Cloisite 10A with different ammonium cations located inthe silicate galleries, were chosen for the nanocomposite preparation. The WAXD analysesand TEM observations showed a higher degree of intercalation in the case of APES/Cloisite30B nanocomposites as compared to that of APES/Cloisite 10A nanocomposites. Accordingto the researchers, this behavior may be due more hydrogen bonding interaction betweenAPES and the hydroxyl group in the galleries of Cloisite 30B nanocomposites than in theAPES/Cloisite 10A nanocomposites.

6.5. PHB/LS Nanocomposites

Poly(hydroxy butyrate) (PHB)/LS nanocomposites were successfully prepared by Maiti andhis colleagues [54]. PHB forms well-ordered intercalated nanocomposites with OMLSs.Although the dispersion of clay particles was not so good (see Fig. 14), the nanocompositesexhibited a higher storage modulus (an increase up to 40% when compared to pure PHB).Better biodegradation behavior was also observed for these nanocomposites as compared tothe neat PHB.

(a) (b)

100 nm 50 nm

Figure 11. TEM images of a fully exfoliated Cloisite 30B-based nanocomposite, showing (a) fine distribution of theclay in the matrix, and (b) delamination of silicate layers. Reprinted with permission from [46], M.-A. Paul et al.,Macromol. Rapid Commun. 24, 561 (2003). © 2003, Wiley-VCH Verlag GmbH & Co.


1500

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(001) qC16-sap

PBSCN1

PBSCN2

PBSCN5

PBSCN6

PBS

PBSCN3

PBSCN4

(a)

(b)

Figure 12. (a) WAXD patterns for pure C18-MMT powder and corresponding PBSCNs. The dashed line indicatesthe location of the silicate (001) reflection of C18-MMT. The asterisks indicate the (001) peak for C18-MMT dis-persed in the PBS matrix. (b) WAXD patterns for pure qC16-MMT powder and corresponding PBSCNs. The dashedline indicates the location of the silicate (001) reflection of qC16-SAP. The asterisks indicate the (001) peak forqC16-MMT dispersed in the PBS matrix. Reprinted with permission from [50], S. Sinha Ray et al., Macromolecules36, 2355 (2003). © 2003, American Chemical Society.

6.6. Starch/LS Nanocomposites

Park and his colleagues [59, 60] reported on their efforts to develop environmentally friendlypolymer hybrids. Biodegradable thermoplastic starch/LS nanocomposites were prepared bythe melt intercalation method. Natural montmorillonite (Na+ MMT; Cloisite Na+) and oneorganically modified MMT with methyl tallow bis-2-hydroxyethyl ammonium cations locatedin the silicategallery (Cloisite 30B) were chosen for the nanocomposite preparation. Starchwas prepared from natural potato starch by gelatinizing and plasticizing it with water andglycerol. The dispersion of the silicate layers in the sarch hybrids was characterized byusing WAXD and TEM. They observed that the starch/Cloisite Na+ nanocomposites showedhigher tensile strength and thermal stability, and better barrier properties to water vapor ascompared to starch /Cloisite 30B nanocomposites and the pristine starch, due to the for-mation of the intercalated nanostructure (see Fig. 15). The effect of clay contents on thetensile, dynamic mechanical, and thermal properties as well as on the barrier properties ofthe nanocomposites were investigated.

6.7. Plant Oil/LS Nanocomposites

Uyama et al. [61] investigated “green” nanocomposites prepared by an acid-catalyzed curingof epoxidized plant oils in the presence of OMLS. Nanocomposites with a homogeneousstructure of organic and inorganic components were obtained, in which the clay was


(a) (b) (e) (f)

(c) (d) (g) (h)

PBSCN1 PBSCN4

PBSCN3 PBSCN6

200 nm

200 nm 200 nm 200 nm 100 nm

100 nm 200 nm 100 nm

Figure 13. TEM bright field images of PBSCNs: (a) PBSCN1 (×100,000), (b) PBSCN1 (×200,000), (c) PBSCN3(×40,000), (d) PBSCN3 (×100,000), (e) PBSCN4 (×100,000), (f) PBSCN4 (×200,000), (g) PBSCN6 (×100,000),and (h) PBSCN6 (×200,000) in which the dark entities are the cross sections of the intercalated or exfoliatedsilicate layers. Reprinted with permission from [50], S. Sinha Ray et al., Macromolecules 36, 2355 (2003). © 2003,American Chemical Society.

intercalated and randomly distributed in the polymer matrix (see Fig. 16). The reinforce-ment effect of the addition of the clay was confirmed by dynamic viscoelasticity analysis.Furthermore, the nanocomposites exhibited flexible properties. These researchers also foundgood biodegradability of the cured polymer from epoxidized soybean oil. These nanocom-posites are anticipated to become a new class of coating materials derived from inexpensiverenewable resources, which will contribute to global sustainability.

500 nm

Figure 14. TEM bright field images of PHB nanocomposites. Reprinted with permission from [54], P. Maiti et al.,Polm. Mater. Sci. Eng. 88, 58 (2003). © 2003, Polymeric Materials Science & Engineering Division of the AmericanChemical Society.


(a) (b)

(c) (d)

Figure 15. TEM images of starch/clay nanocomposites of different types of OMLS. (a) starch 95/Cloisite Na+,(b) starch 95/Cloisit e6A, (c) starch 95/Cloisite10A, (d) starch 95/Cloisite 30B. Reprinted with permission from [59],H. M. Park et al., Macromol. Mater. Eng. 287, 8, 553 (2002). © 2002, Wiley-VCH Verlag GmbH & Co.

6.8. Chitosan/LS Nanocomposites

Ruiz-Hitzky and his colleagues [62] reported on the intercalation of the cationic biopolymerchitosan in Na+-MMT, providing compact and robust three-dimensional nanocompositeswith interesting functional properties (see Fig. 17).

These researchers used CHN chemical analysis, x-ray diffraction, Fourier transforminfrared spectroscopy, scanning transmission electron microscopy, energy-dispersion x-rayanalysis, and thermal analysis to characterize the nanocomposites, confirming the adsorptionin mono- or bilayers of chitosan chains (depending on the relative amount of chitosan withrespect to the cationic exchange capacity of the clay). The first chitosan layer is adsorbedthrough a cationic exchange procedure, while the second layer is adsorbed in the acetate saltform. Because the deintercalation of the biopolymer is very difficult, the -NH+

3 Ac− species


(A)

(B)

2 µm 200 nm

2 µm 200 nm

Figure 16. TEM images of ESO (epoxidized plant oils)-clay nanocomposites with clay content of (a) 5%, and(b) 15%. Reprinted with permission from [61], H. Uyama et al., Chem. Mater. 15, 2492 (2003). © 2003, AmericanChemical Society.

belonging to the chitosan second layer act as anionic exchange sites and, in this way, suchnanocomposites become suitable systems for the detection of anions. These materials havebeen successfully used in the development of bulk-modified electrodes exhibiting numerousadvantages such as easy surface renewal, ruggedness, and long-term stability. The resultingsensors were applied in the potentiometric determination of several anions, showing a higherselectivity toward monovalent anions. This selectivity behavior could be explained by thespecial arrangement of the polymer as a nanostructured bidimensional system.

The interlayer space in the nanocomposites prepared from chitosan-clay ratios of 0.25:1and 0.5:1 can be related to the thickness of one chitosan sheet and, thus, to its intercalationas a monolayer covering the interlayer surface of the clay, as shown in Figure 18. Abovesuch chitosan-clay ratios, the increase of the basal spacing can be explained as the uptakeof two chitosan layers by the clay.

7. MATERIALS PROPERTIES OF BIODEGRADABLEPOLYMER/LS NANOCOMPOSITES

PLS nanocomposites consisting of a polymer and clay (modified or not) frequently exhibitremarkably improved mechanical and materials properties as compared to those of pris-tine polymers containing small amounts (≤5 wt%) of layered silicate. Improvements can

OH

OH

OH

C–O OH

HOHO

HOHO

NH3+NH3

+

NH3+

NH

CH3

O O

O O

O O

O O

O

Figure 17. Chitosan chemical structure.


Figure 18. Intercalation of chitosan into Na+-MMT. Reprinted with permission from [61], M. Darder et al., Chem.Mater. 15, 3774 (2003). © 2003, American Chemical Society.

include higher moduli, increased strength and heat resistance, decreased gas permeabilityand flammability, and increased biodegradability of biodegradable polymers.

7.1. Mechanical Properties

7.1.1. Dynamic Mechanical Analysis (DMA)DMA results are expressed by three main parameters: (1) the storage modulus (G′) corre-sponding to the elastic response to the deformation; (2) the loss modulus (G′′), correspond-ing to the plastic response to the deformation, and (2) tan (i.e., the (G′′/G′) ratio), usefulfor determining the occurrence of molecular mobility transitions such as the glass transi-tion temperature (Tg). DMA analysis has been studied to track the temperature dependenceof G′, G′′, and tan of pure PLA upon nanocomposite formation with five different typesof OMLS. Figure 19 shows the temperature dependence of G′, G′′, and tan of pure PLA


109

108 ω = 6.28 rad.s–1

Strain = 0.05%

PLAPLA/C

18-MMT

PLAPLA/qC2

18-MMT

(a) (b)

(c) (d)

ω = 6.28 rad.s–1

Strain = 0.05%

ω = 6.28 rad.s–1

Strain = 0.05%ω = 6.28 rad.s–1

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0 40 80Temperature/°C Temperature/°C

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13(OH)-Mica

Heating rate = 2 °C/min Heating rate = 2 °C/min

Heating rate = 2 °C/min

Heating rate = 2 °C/min Heating rate = 2 °C/min

Figure 19. Temperature dependence of storage modulus (G′), loss modulus (G′′), and tan of pure PLA andcorresponding nanocomposites: (a) PLA/C18-MMT4; (b) PLA/qC2-MMT4; (c) PLA/qC18-MMT4; (d) PLA/qC16-SAP4; (e) PLA/qC13(OH)-mica 4. Reprinted with permission from [63], S. Sinha Ray et al., Macromol. RapidCommun. 24, 815 (2003). © 2003, Wiley-VCH Verlag GmbH & Co.

and various PLACNs. For all nanocomposites, significant enhancement of G′ can be seenin the investigated temperature range, indicating that all OMLS have strong influence onthe elastic properties of pure PLA. Below Tg, the enhancement of G′ is also clear for allPLACNs [63]. In the temperature range of −20—0�C, the increments in G′ are 37% forPLA/C18-MMT, 52% for PLA/qC2

18-MMT, 45% for PLA/qC18-MMT, 31% for PLA/qC16-SAP(saponite), and 23% for PLA/qC13(OH)-mica nanocomposites compared to that of purePLA [63]. In the temperature range of 80–90�C, all nanocomposites exhibit much greaterenhancement in G′ (103% for PLA/C18-MMT, 105% for PLA/qC2

18-MMT, 96% for PLA/qC18-MMT, and 111% for PLA/qC13(OH)-mica) than that of pure PLA, with the exception ofPLA/qC16-SAP nanocomposite with a 45% increment. This is due to the mechanistic rein-forcement by clay particles at high temperature. Above Tg, when materials become soft,the reinforcement effect of clay particles becomes prominent due to the restricted move-ment of the polymer chains and, hence, strong enhancement of modulus appeared in thisstudy [32]. The restriction of polymer chain movement by the clay particles is high in the


case of PLA/qC13(OH)-mica because of the low value of �clay (see Table 4 and Figure 3).For this reason, PLA/qC13(OH)-mica nanocomposites show high increments in G′ at hightemperature ranges compared to that of other PLACNs. Furthermore, at room tempera-ture (25�C), PLA/C18-MMT, PLA/qC18-MMT, PLA/qC2

18-MMT, and PLA/qC16-SAP showedhigher increments in G′ of 38%, 44%, 51%, and 30%, respectively, than that of pure PLA,while that of PLA/qC13(OH)-mica showed only 26% higher. These increments come fromthe extended intercalation, the higher degree of crystallization, and also the high aspect ratioof dispersed clay particles in MMT-based nanocomposite systems.

On the other hand, above Tg the enhancement of G′′ is significant in all nanocompositesin comparison with that of below Tg, indicating that plastic response to the deformation ofpure PLA is prominent in the presence of OMLS when materials become soft. However, thepresence of OMLS does not lead to a significant shift and broadening of the tan curvesfor all PLACNs compared to that of pure PLA. This behavior has been ascribed to therestricted segmental motions in the organic-inorganic interface neighborhood of intercalatedPLACNs.

The increment in G′ directly depends upon the aspect ratio of dispersed clay particles,which is also clearly observed in PBSCNs. The temperature dependence of G′ of PBS and var-ious PBSCNs are shown in Figure 20). The nature of enhancement of G′ in PBSCNs with tem-perature is somewhat different from well established intercalated polypropylene/LS nanocom-posites [24] and well known exfoliated nylon 6/LS nanocomposite systems (N6CNs) [12, 64].In the latter system, there is a maximum of 40–50% increment of G′ as compared to thatof matrices at well below Tg; above Tg, there is a strong enhancement (>200%) in G′. Thisbehavior is common for nanocomposites evaluated to date, and the reason is the strong rein-forcement effect of the clay particles above Tg when materials become soft. But with PBSCNs,

–40 –20 0 20 40 60 80 100

tan

δ

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10–1

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108

109

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G′′ /

Pa

Heating rate = 2°/minStrain = 0.05%ω = 6.28 rad/s

PBSPBSCN1PBSCN2PBSCN3PBSCN4

G′ /P

a

Figure 20. Temperature dependence of G′, G′′ and tan for neat PBS and various PBSCNs prepared withC18-MMT. Reprinted with permission from [50], S. Sinha Ray et al., Macromolecules 36, 2355 (2003). © 2003,American Chemical Society.


the order of enhancement of G′ is almost the same below and above Tg; this behavior may bedue to the extremely low Tg (−29�C) of the PBS matrix. At the temperature range of −50�Cto −10�C, the increments in G′ are 18% for PBSCN1, 31% for PBSCN2, 67% for PBSCN3,and 167% for PBSCN4 compared to that of neat PBS. Furthermore, at room temperature,PBSCN3 and PBSCN4 show higher increments in G′ (82% and 248%, respectively) than thatof neat PBS, while those of PBSCN1 and PBSCN2 are 18.5% and 44% higher, respectively.At 90�C, only PBSCN4 exhibits very strong enhancement of G′ compared to that of the otherthree PBSCNs.

In Figure 21, Okamoto summarizes the clay content dependence of G′ of various typesof nanocomposites obtained well below Tg. The Einstein coefficient (kE) derived by usingHalpin and Tai’s theoretical expression (modified by Nielsen) is shown in Figure 21, andrepresents the aspect ratio (Lclay/dclay) of dispersed clay particles without intercalation.

From Figure 21, it can be clearly observed that PBSCNs show very high increment in G′

compared to other nanocomposites having the same clay content in the matrix. PPCNs arewell known for intercalated systems; N6CNs are already well established exfoliated nanocom-posites; PLACNs will soon be established intercalated-and-flocculated nanocomposites; andPBSCNs are intercalated-and-extended flocculated nanocomposites systems [50, 51]. Due tothe strong interaction between hydroxylated edge–edge groups, the clay particles are some-times flocculated in the polymer matrix. As a result of this flocculation, the length of the clayparticles increases enormously, as does the overall aspect ratio. For the preparation of high-molecular-weight PBS, di-isocyanate end-groups are generally used as chain extenders [65].These isocyanate end-group chain extenders make urethane bonds with hydroxy-terminated,low-molecular-weight PBS, and each high-molecular-weight PBS chain contains two suchbonds (see the schematic illustration in Figure 22). These urethane-type bonds lead to stronginteraction with the silicate surface by forming hydrogen bonds and, hence, strong floccula-tion (see Fig. 23). For this reason, the aspect ratio of dispersed clay particles is much higherin PBSCNs compared to all other nanocomposites, which results in high enhancement ofthe modulus.

7.1.2. Tensile PropertiesThe effect of clay content on the tensile properties of PLC/LS nanocomposites has beenstudied (see Tables 6 and 7). The Young’s modulus of nanocomposites is higher comparedto neat PCL, as result of their intercalated/exfoliated structure. For instance, the Young’smodulus is significantly increased from 216 MPa for pure PCL to more than 390 MPa forthe composite that contains 10 wt% of MMT-(OH)2 [66]. PCL is a ductile polymer able tosustain large deformations (700% at break). Adding nanoclays only slightly decreases theelongation at break, as shown in Tables 6 and 7. PCL remains ductile with an elongation at

1

10

1001010.1

G′ n

anoc

ompo

site

/G′ m

atrix

Vol % of clay

PLACN1PLACN2PLACN3PLACN4PLACN5PLACN6

70 15160

N6CN1.6N6CN3.7PBSCN1PBSCN2PBSCN3PBSCN4PBSCN5PBSCN6

T=20 °C

T=0 °C

T=–50 °C

Figure 21. Plots of G′nanocomposite/G

′matrix versus volume percent of clay for various nanocomposites. The Einstein

coefficient (kE) is shown with the number in the box. The lines show the calculated results from Halpin and Tai’stheory with various kEs.


O

OC

N

H

OCN

O

H

PBS, Mw is 50300 gm/mol.

Figure 22. Formation of urethane bondings in high-molecular-weight PBS.

break higher than 550%. However, higher clay content (10 wt%) has a detrimental effect,as confirmed by an ultimate elongation lower than 10% for the two types of OMLS.

7.1.3. Flexural PropertiesFlexural modulus and strength of pure PLA and various PLACNs measured at 25�C aresummarized in Table 8. Compared to that of pure PLA, there is a significant increase inflexural modulus for all PLACNs except PLA/qC16-SAP nanocomposites. The high modulusvalue in PLA/qC13(OH)-mica and the low modulus value in PLA/qC16-SAP may be dueto the different aspect ratios of dispersed clay particles in the PLA matrix. On the otherhand, flexural strength is also remarkably increased in PLA/C18-MMT, PLA/qC18-MMT, andPLA/qC2

18-MMT, but there is not such a large increase in the cases of PLA/qC13(OH)-micaor PLA/qC16-SAP. This behavior may be due to some brittleness appearing in materials inthe presence of qC13(OH)-mica or qC16-SAP, as revealed by distortion values (see Table 8).

7.1.4. Heat Distortion TemperatureHeat distortion temperature (HDT) of a polymeric material is an index of heat resistancetoward applied load. Most PLS nanocomposite studies report HDT as a function of claycontent, characterized by ASTM D-648. Kojima and his colleagues [13] first showed thatthe HDT of pure nylon 6 increases up to 80�C after nanocomposite preparation with MMT.In another investigation [67], they reported clay content dependence of HDT in nylon 6/LSnanocomposites. There is a marked increase in HDT, from 65�C for the neat nylon 6 to150�C for 4.7 wt% nanocomposites; beyond the weight percent of MMT, the HDT of thenanocomposites levels off. They also conducted HDT tests of various nylon 6/LS nanocom-posites prepared with clay having different lengths, and found that HDT also depends uponthe aspect ratio of dispersed clay particles [13].

Since the degree of crystallinity of nylon 6/LS nanocomposites is independent of the amountand nature of clay, the HDT of nylon 6/LS nanocomposites is related to the presence of stronginteractions between matrix and silicate surfaces by forming hydrogen bonds (see Fig. 24).Although nylon 6 in nanocomposites stabilizes in a different crystal phase (�-phase) than thatfound in pure nylon 6, this different crystal phase is not responsible for the higher mechanicalproperties of nylon 6/clay nanocomposites because the �-phase is a very soft crystal phase.

OSi

OSi

OSi

OSi

OSi

OSi

OSi

OSi

OSi

OSi

O Si OSi

OSi

OSi

OSi

OSi

OSi

OSi

OSi

OSi

NH

NH

NH

HO

HO

OH

OH

NH

NH

Figure 23. Formation of hydrogen bonds between PBS and clay, which leads to flocculation of the dispersedsilicate layers.


Table 6. Tensile properties of PCL/LS nanocomposites containing MMT-Alk.

ElongationSample OMLS (wt%) Young’s Modulus (Mpa) at Break (%) Tensile Strength (Mpa)

PCLC9 1 262 ± 13 659 ± 27 33 ± 1PCLC10 3 282 ± 9 528 ± 58 26 ± 3PCLC11 5 307 ± 18 598 ± 43 28 ± 1PCLC12 10 371 ± 15 9 ± 1 18 ± 1

Source: Reprinted with permission from [66], B. Lepoittevin et al., Polymer 42, 4017 (2001). © 2001, ElsevierScience Ltd.

On the other hand, the increased mechanical properties of nylon 6/LS nanocomposites thataccompany increasing clay content is due to the mechanical reinforcement effect.

Okamoto et al. examined the HDT of neat PLA and various PLA/LS nanocompos-ites (PLACNs) with different load conditions. As seen in Figure 25(a) [48], in PLACN7(inorganic clay content = 5 wt%) there is a marked increase of HDT with an intermedi-ate load of 0.98 MPa, from 76�C for the neat PLA to 98�C for PLACN4 (inorganic claycontent = 3 wt%). The value of HDT gradually increases with increasing organoclay content;in PLACN7, the value increases up to 111�C.

On the other hand, imposed load dependence on HDT is clearly observed in PLACNs.Figure 25(b) shows the typical load dependence in PLACN7. The increase of HDT of neatPLA due to nanocomposite preparation is a very important property improvement, not onlyfrom the industrial point of view but also pertaining to molecular control on the silicatelayers (i.e., crystallization through interaction between PLA molecules and SiO4 tetrahedrallayers in the MMT). When there is a high load (1.81 MPa), it is very difficult to achievehigh HDT enhancement without strong interaction between the polymer matrix and organo-MMT [13]. With all of the PLACNs studied here, the values of the melting temperature(Tm) do not change significantly as compared to that of neat PLA. Thus, the improvement ofHDT with an intermediate load (0.98 MPa) originates in the better mechanical stabilityof PLACNs due to mechanical reinforcement by the dispersed clay particles, higher levels ofcrystallinity (c), and intercalation. This is qualitatively different from the behavior of a nylon6/LS nanocomposite system, where the MMT layers stabilize in a different crystalline phase(�-phase) [67] than that found in neat nylon 6, with the strong hydrogen bonding betweenthe silicate layers and nylon 6 as a result the discrete lamellar structure on both sides of theclay (see Fig. 24).

7.1.5. Izod-Impact PropertiesTable 9 shows the Izod-impact strength values for PCL/LS nanocomposites prepared withMMT-Na+, MMT-Alk of MMT-(OH)2 as a function of clay content [66]. The Izod-impactstrength systematically decreases with increasing clay content. It drops from 48 J/m forunfilled PCL to 33 and 13 J/m when 1 wt% and 10 wt%, respectively, of the MMT-(OH)2nanoclay are incorporated.

7.1.6. Thermal StabilityThe thermal stability of polymeric materials is usually studied by thermogravimetric analysis(TGA). The weight loss due to the formation of volatile degradation products is monitoredas a function of temperature ramp. When the heating is conducted under an inert gas

Table 7. Tensile properties of PCL/LS nanocomposites containing MMT-(OH)2.

ElongationSample OMLS (wt%) Young’s Modulus (Mpa) at Break (%) Tensile Strength (Mpa)

PCLC1 1 259 ± 11 705 ± 47 36 ± 2PCLC2 3 272 ± 16 563 ± 62 25 ± 4PCLC3 5 313 ± 23 560 ± 46 24 ± 3PCLC4 10 399 ± 23 7 ± 1 17 ± 0.5

Source: Reprinted with permission from [66], B. Lepoittevin et al., Polymer 42, 4017 (2001). © 2001, ElsevierScience Ltd.


Table 8. Comparison of practical materials properties of pure PLA and various nanocomposites.

Practical PLA/ PLA/ PLA/ PLA/ PLA/Materials Properties PLA C18-MMT qC2

18-MMT qC18-MMT qC16-SAP qC13(OH)-mica

Flexural modulus 4.84 5.66 5.43 5.57 4.5 6.11(GPa at 25�C)

Flexural strength 86 132 102 134 93 94(/MPa at 25�C)

Distortion at break (%) 1.9 3.2 3.9 3.1 1.5 1.5HDT (�C with 0.98 MPa 76 94 91.3 93 98 93

load)O2 gas permeability 200 172 171 177 120 71(mL/mm) (m2/day.MPa)O2 gas permeability 200 168 167 178 169 68(mL/mm) (m2/day.MPa)a

aCalculated on the basis of the Nielsen theoretical equation (Eq. (2)) in this chapter.Source: Reprinted with permission from [63], S. Sinha Ray et al., Macromol. Rapid Commun. 24, 815 (2003). © 2003, Wiley-VCH

Verlag GmbH & Co.

N

H

O

N

H

O

N

H

O

n

Si

OO

Si

O

Si

O

Si

O

Si

O

Si

O

Si

17.2 Å

HH

OO

Figure 24. Schematic illustration of the formation of hydrogen bonds in nylon 6/MMT nanocomposites.

0

Load = 0.98 MPa

HD

T/°

C

HD

T/°

C

70

80

90

100

110

120(a)

Organoclay/wt.%0.4 0.8 1.2 1.6 2

PLAPLACN7

Load /MPa

60

80

100

120

140

160(b)

2 4 6 8

Figure 25. (a) OMLS (wt %) dependence of HDT of neat PLA and various PLACNs. (b) Load dependence ofHDT of neat PLA and PLACN7. Reprinted with permission from [48], S. Sinha Ray et al., Polymer 44, 857 (2003).© 2003, Elsevier Science Ltd.

Table 9. Izod-impact properties of PCL/LS nanocomposites contain-ing MMT-Na, MMT-Alk, and MMT-(OH)2.

Izod-impact (J/m)OMLS (wt%) MMT-Na MMT-Alk MMT-(OH)2

1 33 ± 5 28 ± 6 33 ± 33 22 ± 2 22 ± 2 18 ± 35 19 ± 1 15 ± 1 13 ± 110 15 ± 1 16 ± 3 13 ± 1

Source: Reprinted with permission from [66], B. Lepoittevin et al., Polymer42, 4017 (2001). © 2001, Elsevier Science Ltd.


flow, a nonoxidative degradation occurs, while the use of air or oxygen allows following theoxidative degradation of the samples. Generally, the incorporation of clay in the polymermatrix enhances thermal stability by acting as a superior insulator and mass transport barrierto the volatile products generated during decomposition.

The thermal stability of the PCL/LS nanocomposites has also been studied by TGA.Generally, the degradation of PCL fits a two-step mechanism [35]: First, a statistical rup-ture of the polyester chains by pyrolysis of ester groups with the release of CO2, H2O,and hexanoic acid; and second, �-caprolactone (cyclic monomer) is formed as a result ofan unzipping depolymerization process. The thermograms of nanocomposites prepared withorganoclay and pure PCL recovered after clay extraction are presented Figure 26. Bothintercalated and exfoliated nanocomposites show higher thermal stability when comparedto pure PCL or the corresponding microcomposites. The nanocomposites reached a highof 25�C in decomposition temperature at 50% weight loss. The shift of the degradationtemperature may be ascribed to: (1) a decrease in oxygen; (2) a decrease in the permeabil-ity/diffusivity of volatile degradation products due to the homogeneous incorporation of claysheets; (3) a barrier of high aspect ratio fillers; and (4) char formation. The thermal stabilityof nanocomposites systematically increases with increasing clay; however, above a loading of5 wt%, the thermal stability is no longer improved.

A completely different behavior is observed in synthetic biodegradable aliphatic polyester(BAP)/clay nanocomposite systems. Here, the thermal degradation temperature and thermaldegradation rate are systematically increased with an increasing amount of organoclay, upto 15 wt% [68]. Like PS/LS nanocomposites, a small amount of clay also increased theresidual weight of BAP/OMMT because of the restricted thermal motion of the polymer inthe silicate layers. The residual weight of various materials at 450�C increased in the orderBAP < BAP03 < BAP06 < BAP09 < BAP15 (here, the number indicates the weight percentof clay). This type of improved thermal properties is also observed in other systems likeSAN [69], the intercalated nanocomposites prepared by emulsion polymerization.

Many researchers believe the role of clay in nanocomposite structure might be the mainreason for the difference in TGA results of these systems compared to the so far reportedsystems. The clay acts as a heat barrier that could enhance the overall thermal stability ofthe system, as well as assisting in the formation of char after thermal decomposition. Thus,in the beginning stage of thermal decomposition, the clay could shift the decompositiontemperature higher. However, after that, this heat barrier effect would result in a reversethermal stability. In other words, the stacked silicate layers could hold accumulated heat

250 300 350

25

75

100

50

400 450 500

PCL

1 wt%

3 wt%

5 wt%

10 wt%

temperature °C

wei

ght (

%)

0

Figure 26. Temperature dependence of weight loss under an air flow for neat PCL and PCL nanocompositescontaining 1, 3, 5, and 10 wt% (relative to inorganics) of MMT-Alk (heating rate: 20 K/min). Reprinted withpermission from [35], B. Lepoittevin et al., Polymer 43, 1111 (2002). © 2002, Elsevier Science Ltd.


that could be used as a heat source to accelerate the decomposition process, in conjunctionwith the heat flow supplied by an outside heat source.

7.1.7. Gas Barrier PropertiesNanoclays are believed to increase the barrier properties by creating a maze or “tortuouspath” (see Fig. 27) that retards the progress of gas molecules through the matrix resin.The direct benefit of the formation of this type of path is clearly observed in polyimide/claynanocomposites, which show dramatically improved barrier properties with simultaneousdecrease in the thermal expansion coefficient [70, 71]. The polyimide/LS nanocompositesrevealed a several-fold reduction in the permeability of small gases (e.g., O2, H2O, He,CO2, and the organic vapor ethylacetate) with the presence of a small fraction of OMLS.For example, at 2 wt% clay loading, the permeability coefficient of water vapor was decreasedten-fold for synthetic mica relative to pristine polyimide. By comparing nanocompositesmade with layered silicates of various aspect ratios not only was the permeability noted todecrease with increasing aspect ratio.

Okamoto and Yamada [72] measured the O2 gas permeability for exfoliated PLA/syntheticmica nanocomposites. Figure 28 shows a plot of the relative permeability coefficient valueas a function of the weight percent of clay, PPLACN/PPLA, where PPLACN and PPLA are thenanocomposite and pure PLA permeability coefficients, respectively. The curve fitting wasachieved by using the Nielsen theoretical expression [77], allowing the prediction of gaspermeability as a function of the length and width of the filler particles as well as theirvolume fraction within the PLA matrix.

In the Nielsen model [73], where platelets of length (� Lclay) and width (� Dclay) of theclay (which are dispersed parallel in the polymer matrix), the tortuosity factor (�) can beexpressed as:

� = 1 +(

Lclay

2Dclay

)�clay (1)

where �clay is the volume fraction of dispersed clay particles. Therefore, the relative perme-ability coefficient (PPLS nano/PNeat) is given by:

PPCNnano

PNeat= �−1 = 1

1 + !Lclay/2Dclay"�clay#(2)

where PPCNnano and PNeat are the permeability coefficients of PCN and neat polymer,respectively.

7.1.8. BiodegradabilityAnother most interesting and exciting aspect of nanocomposite technology is the signif-icant improvement in the biodegradability of biodegradable polymers after nanocompos-ite preparation with OMLS. Aliphatic polyesters are among the most promising materialsfor the production of environmentally friendly biodegradable plastics. Biodegradation ofaliphatic polyester is well known, in that some bacteria degrade them by producing enzymeswhich attack the polymer. Tetto et al. [74] first reported on the biodegradability of PCL-based nanocomposites, where PCL/LS nanocomposites showed improved biodegradabilitycompared to pure PCL. According to these researchers, the improved biodegradability of

Conventional composites“Tortuous path” in layered silicate

nanocomposites

Figure 27. Formation of “tortuous paths” in polymer/clay nanocomposites.


6 8 10

Theoretical curve based on

Lclay/Dclay = 275

Experimental value

O2

TR

/ml.m

m.m

–2.d

ay–1

.MP

a–1

OMSFM/wt %

0

50

100

150

200

0 2 4

Figure 28. Oxygen gas permeability of neat PLA and various PLACNs as a function of organoclay content measuredat 20�C and 90% relative humidity. The filled circles represent the experimental data. Theoretical fits are based onthe Nielsen tortuosity model.

PCL after nanocomposite formation may be due to the catalytic role of the OMLS in thebiodegradation mechanism. But it is still unclear how clay increases the biodegradation rateof PCL.

In 2002, Lee et al. [55] reported on the biodegradation of aliphatic polyester-basednanocomposites under compost. Figure 29(a) and (b), respectively, show the clay contentdependence of biodegradation of APES-based nanocomposites prepared with two differ-ent types of clays. These researchers assumed that the biodegradation was retarded due toimprovement of the barrier properties of the aliphatic APSE after nanocomposite prepara-tion with clay. However, they provided no data about permeability.

Very recently, Yamada and Okamoto et al. [42, 48] first reported on the biodegradability ofneat PLA and corresponding nanocomposites prepared with trimethyl octadecylammonium-modified MMT (C3C18-MMT) with details mechanism. The compost used was prepared fromfood waste and tests were carried out at 58�C ± 2�C. Figure 30(a) shows an actual pictureof the samples of neat PLA and PLACN4 (C3C18-MMT = 4 wt%) recovered from compostwith time. The decreased molecular weight (Mw) and residual weight percentage (Rw) of theinitial test samples with time are shown in Figure 30(b). The biodegradability of neat PLAis significantly enhanced after PCN preparation. Within one month, the decrease in Mw andthe extent of weight loss are almost the same for both PLA and PLACN4. However, afterone month a sharp change occured in weight loss of PLACN4, and within two months it wascompletely degraded in compost. The degradation of PLA in compost is a complex processinvolving four main phenomena: (1) water absorption; (2) ester cleavage and formation ofoligomer fragments; (3) solubilization of oligomer fragments; and (4) diffusion of solubleoligomers by bacteria [75]. Therefore, the factor that increases the hydrolysis tendency ofPLA ultimately controls the degradation of PLA.

(b)(a)

Time (Day) Time (Day)0

05 10 15 20

20

40

60

80

100

25 3530 0 5 10 15 20 25 3530

Bio

degr

adab

ility

(w

t%)

0

20

40

60

80

100

Bio

degr

adab

ility

(w

t%)

APESAPES30B (97/3 wt %)APES30B (95/5 wt %)APES30B (90/10 wt %)APES30B (80/20 wt %)APES30B (70/30 wt %)

APES/10A (97/3 wt %)

APES/10A (90/10 wt %)APES/10A (90/20 wt %)APES/10A (70/30 wt %)APES

APES/10A (95/5 wt %)

Figure 29. Biodegradability of APES nanocomposites with: (a) Closite 30B and (b) Closite 10A. Reprinted withpermission from [55], S. R. Lee et al., Polymer 43, 2495 (2002). © 2002, Elsevier Science Ltd.


PLA

PLACN4

after 32 days after 60 daysafter 50 days

PLA

PLA

(b)

(a)

0

20

40

60

80

100

0

50

100

150

0 10 20 30 40 50 60 70

PLACN4

PLACN4

Rw /%

MW

x 1

0–3 /(

gm/m

ol)

Time /days

Figure 30. (a) Actual picture of biodegradability of neat PLA and PLACN4 recovered from compost with time.The initial shape of the crystallized samples was 3 × 10 × 0%1 cm3. (b) Time dependence of residual weight, Rw andof matrix, Mw of PLA and PLACN4 under compost at 58 ± 2�C. Reprinted with permission from [48], S. SinhaRay et al., Polymer 44, 857 (2003). © 2003, Elsevier Science Ltd.

These researchers concluded that the presence of terminal hydroxylated edge groups onthe silicate layers may be one of the factors responsible for this behavior. In the case ofPLACN4, the stacked (approximately four layers) and intercalated silicate layers are homo-geneously dispersed in the PLA matrix (from TEM imaging) and these hydroxy groupsstart heterogeneous hydrolysis of the PLA matrix after absorbing water from the compost.Because this process takes some time to start, the weight loss and degree of hydrolysis ofPLA and PLACN4 are almost the same within up to one month [see Fig. 30(b)]. However,after one month there is a sharp weight loss in the PLACN4 compared to the PLA. Thatmeans that one month is a critical time at which to start heterogeneous hydrolysis; due tothis type of hydrolysis, the matrix breaks into very small fragments and disappears with com-posting. This assumption was confirmed by conducting the same type of experiment withPLACN prepared with dimethyl dioctdecyl ammonium salt-modified synthetic mica, whichhas no terminal hydroxylated edge group. The degradation tendency was almost the sameas that of neat PLA [76].

Yamada and Okamoto et al. [41, 76] also conducted respirometric tests to study degrada-tion of the PLA matrix in a compost environment at 58�C ± 2�C. For this test, the compostwas made from bean curd refuse, food waste, and cattle feces. Rather than weight loss, whichreflects the structural changes in the test sample, CO2 evolution provides an indicator ofthe ultimate biodegradability of PLA in PLACN4. The PLA in PLACN4 was prepared with(N (cocoalkyl)N , N -[bis(2-hydroxyethyl)]-N -methylammonium-modified synthetic mica); inother words, the samples were mineralized. Figure 31 shows the time dependence of the


0 10 20 30 40 50

Neat PLA PLACN4

0

20

40

60

80

100(a)

0 4 6 8 10 12 14

Neat PLA PLACN4

Mw

×10

–3–1

0

50

100

150

200(b)

2

Time/Days Time/Days

Deg

ree

of B

iode

grad

atio

n/%

/g.m

ol

Figure 31. Degree of biodegradation (i.e., CO2 evolution), and (b) time-dependent change of matrix Mw of neatPLA and PLACN4 (MEE clay = 4 wt%) under compost at 58 ± 2�C. Reprinted with permission from [41],S. Sinha Ray et al., Macromol. Rapid Commun. 23, 943 (2002). © 2002, Wiley-VCH Verlag GmbH & Co.

degree of biodegradation of neat PLA and PLACN4, indicating that the biodegradabilityof PLA in PLACN4 is significantly enhanced. The presence of OMLS may thus cause adifferent mode of attack on the PLA component, which might be due to the presence ofhydroxy groups. Details of the mechanism of biodegradability are presented in the relevantliterature [41, 76].

K. Okamoto and M. Okamoto also investigated the biodegradability of neat PBS beforeand after nanocomposite preparation with three different types of OMLS. They used alky-lammonium or alkylphosphonium salts for the modification of pristine layered silicates, andthese surfactants are toxic for microorganisms.

Figure 32(a) shows actual pictures of samples of neat PBS and various nanocompos-ites recovered from compost after 35 days. This figure clearly shows that many cracksappeared in the nanocomposite samples compared to that of neat PBS. This confirms theimproved degradability of nanocomposites in compost. This kind of fracture is advantageousfor biodegradation because it creates much more surface area for further attack by microor-ganisms. It should be noted that the extent of fragmentation is directly related to the natureof the OMLS used for nanocomposite preparation. These researchers also conducted gelpermeation chromatography (GPC) measurement of the samples recovered from compost,and they found that the extent of molecular weight loss was almost the same for all samples(see Table 10). This result indicates that the extent of hydrolysis of PBS in a pure state orin OMLS-filled systems is the same as in compost.

PBS PBS/C18-MMT

PBS/qC18-MMT PBS/qC16-SAP

PBS PBS/C18-MMT

PBS/qC18-MMT PBS/qC16-SAP

(a) (b)

Figure 32. Biodegradability of neat PBS and various nanocomposite sheets (a) under compost, and (b) under soilfields. Reprinted with permission from [51], K. Okamoto et al., J. Polym. Sci. Part B: Polym. Phys. 41, 3160 (2003).© 2003, John Wiley & Sons, Inc.


Table 10. GPC results of various samples recovered from compost after 35 days.

Samples Mw × 10−3(g/mol) Mn × 10−3(g/mol) Mow × 10−3(g/mol) Mw/Mo

w

PBS 16 3.8 101 0%16PBS/C18-MMT 17 6.6 104 0%16PBS/qC18-MMT 17 4.4 112 0%15PBS/qC16-SAP 8.7 1.2 91 0%096

Mow is the molecular weight before composting.

Except for the PBS/qC16-SAP system, the degree of degradation was not different forother samples. This observation indicates that MMT or alkylammonium cations and otherproperties have no effect on the biodegradability of PBS. The accelerated degradation of thePBS matrix in the presence of qC16-SAP may be due to the presence of the alkylphosphoniumsurfactant. This type of behavior was also observed in the case of PLA/OMLS nanocompositesystems.

Yamada and Okamoto et al. also observed nature of degradation of PBS and variousnanocomposites under soil fields. These experiments were conducted for one, two, and sixmonths. After one and two months, there was no change in the nature of the sample surfaces,but after six months black or red spots appeared on the surface of nanocomposite samples.Figure 32(b) shows the results of degradation of neat PBS and various nanocomposite sheetsrecovered from soil fields after six months. These researchers concluded that these spots onthe sample surfaces were due to fungus attack, because when they put these samples intoa slurry they observed clear fungus growth. These results also indicate that nanocompositesexhibited the same or higher levels of biodegradability compared to PBS matrices.

8. CRYSTALLIZATION OF BIODEGRADABLEPOLYMER/LS NANOCOMPOSITES

Crystallization of PLS nanocomposites might be a good method for controlling the structureand various properties of nanocomposites.

Okamoto and Nam reported in detail on the crystallization behavior of PLA/C18-MMTwith 4 wt% of C18-MMT as a representative system [77]. To understand the crystallizationkinetics of pure PLA before and after nanocomposite preparation at low Tc (≤120�C), weused time-resolved LS photometry, which is a powerful tool for estimating the overall crystal-lization rate and its kinetics in supercooled crystalline polymer liquid [44]. Details regardingLS experiments can be found elsewhere [78]. For the kinetics of crystallization, integratedscattering intensity can be employed; the invariant Q is defined as:

Q =∫ �

0I!q"q2dq (3)

where q [scattering vector = !4*/+LS" sin!,LS/2"] and I!q" is the intensity of the scatteredlight at q [78]. In the Hv mode, the invariant Q can be described by the mean-square opticalanisotropy 2�:

Q ∝ 2� ∝ �s!-r − -t"2 (4)

where �s is the volume fraction of the spherulites, and -r and -t are the radial and tangentialpolarizabilities of the spherulites, respectively. We constructed a plot of reduced invariantQ /Q

� versus time (t) with Q�

being Q at an infinitely long time of crystallization (up tofull solidification of the melt).

Figure 33 shows the time variation of the invariant Q /Q� taken for pure PLA and

PLA/C18-MMT at 110�C. The overall crystallization rate was determined from the slope ofQ /Q

� (d(Q /Q

� )/dt) in the crystallization region, as indicated by the solid line in Figure 33,

and we plotted in Figure 34. It is clear that the overall crystallization rate increases inPLA/C18-MMT, in comparison to pure PLA, as well as the rate increases in PLA/C18-MMTfor a particular Tc. The same trend is also observed over the wide range of Tcs studied here.


400300200Time/sec.

Qδ/

Q∞ δ

10000.0

0.2

0.4

0.6

0.8

1.0

PLACN4

neat PLA

Tc=110 °C

Figure 33. Time variation of reduced invariant Q /Q� during isothermal crystallization at quiescent state at

Tc = 110�C. The solid line represents the slope (overall crystallization rate). Reprinted with permission from [63],S. Sinha Ray et al., Macromol. Rapid Commun. 24, 815 (2003). © 2003, Wiley-VCH Verlag GmbH & Co.

It should be noted that the equilibrium melting temperatures (T 0m) of the PLA/C18-MMT and

pure PLA are the same. The equilibrium melting temperatures were measured by isothermalcrystallization at various temperatures by constructing a Hoffman-Weeks [48] plot, as shownin Figure 35. Both PLA/C18-MMT and pure PLA show the same value of T 0

m of 179.5�C;this would nullify the effect of supercooling .T ! ≡ T 0

m − Tc" on the overall crystallizationrate, linear growth rate, G, etc. The overall crystallization rate with Tc is observed typicalrate curve as usual for semicrystalline polymer. However, the crystallization rate of PLACN4is enhanced for every measured temperature, especially at higher Tcs. From the onset time(t0), we can estimate the induction time of the crystallization until start of crystallization.The observed value of t0 at 110�C was 74 s for pure PLA and 56 s for PLA/C18-MMT. At allTcs measured here, the t0 value for PLA/C18-MMT was lower than that of pure PLA. Thisreduction of t0 in PLA/C18-MMT is attributed to the presence of clay as the nucleating agent.

A typical example of the time variation of the diameter of a spherulite (D) for purePLA and PLA/C18-MMT at higher Tcs is shown in Figure 36(a), and the linear growth rate[G = 1/2!dD/dt"] of the spherulites is summarized in Figure 36(b). For both pure PLAand PLA/C18-MMT, G decreases with increasing Tc in the temperature range of 120–140�C.However, for PLA/C18-MMT, G shows a slightly higher value compared to that of purePLA. This observation indicates that the dispersed clay particles do not have much effect onthe crystallization and no big acceleration of G in the crystallization of the PLA/C18-MMT.This behavior suggests that the diffusion rate of bulk PLA molecules is not enhanced with

PLACN4neat PLA

10–1

10–2

10–3

10–4

70 80 90 100 110 120 130 140

Crystallization Temperature/°C

d(Q

δ/Q

∞ δ)

dt

Figure 34. Tc dependence of the overall crystallization rate of pure PLA and PLACN4. Reprinted with permissionfrom [63], S. Sinha Ray et al., Macromol. Rapid Commun. 24, 815 (2003). © 2003, Wiley-VCH Verlag GmbH & Co.


100

190

185

175

165

160

170

180

120 140

Tc/°CT

m/°

C160 180

neat PLAPLACN4

Figure 35. Tm versus Tc (Hoffman-Weeks) plots of pure PLA and PLACN4. Reprinted with permission from [63],S. Sinha Ray et al., Macromol. Rapid Commun. 24, 815 (2003). © 2003, Wiley-VCH Verlag GmbH & Co.

the addition of clay at every measured temperature; thus, the overall crystallization rate isaffected only by nucleation of the clay particles [79, 80].

Figure 37 demonstrates that the number of heterogeneous nuclei (N ) can be estimatedfrom a rough approximation [using Eq. (4)]. The calculated value of N at 130�C was9%3 × 10−7 /m−3 for pure PLA and 55%7× 10−7 /m−3 for PLA/C18-MMT. The time variationof the volume fraction of the spherulites increases in proportion to NG3 ( � the overallcrystallization rate). This fact suggests that the overall crystallization rate of PLA/C18-MMTat high temperature (Tc = 130�C) is about a one-half order of magnitude higher than thatof matrix PLA without clay. The difference in N between pure PLA and PLA/C18-MMTat Tc = 130�C is higher than at low Tc. This suggests that PLA/C18-MMT exhibits hetero-geneous nucleation kinetics, which depend on more originating from the well dispersedclay particles in the matrix at high temperature. It should be noted that the spherulites ofPLA/C18-MMT have a lower ordering than those of pure PLA, due to the presence of dis-persed clay particles in the spherulites [77]. Hence, if an aggregation of clay particles (whichare not nucleated during crystallization) exists inside the spherulite, then the occurrence ofthe regular orientation of lamella stacks inside the spherulite may be disrupted (see Fig. 38).

Figure 38 shows WAXD profiles of neat PLA and PLACN4 after crystallization at 110�Cfor 1.5 hours. The neat PLA exhibits a very strong reflection at 20 = 17%1 degrees due todiffraction from the (200) and/or (110) planes, and another reflection peak at 20 = 19%5degrees occurring from the (203) plane. On the other hand, in PLACN4 these peaks areshifted toward the lower diffraction angle accompanied by another small peak at 20 = 15%3degrees. After calculation, it was confirmed that this reflection is due to the (010) diffrac-tion plane. These profiles indicate neat PLA crystals are typical orthorhombic crystals [81];

00 0.09

0.10

0.11

0.12

20

40

60

80

100

120

140

500 1000

Crystallization Time/sec

Sph

erul

ite D

iam

eter

/µm

G/µ

ms–1


2000 110 120 130 140 1501500

neat PLAPLACN4

neat PLAPLACN4

c

(a) (b)

T =130 °C

Figure 36. (a) Spherulite diameter as a function of crystallization time at Tc = 130�C; and (b) linear growth rate ofpure PLA and PLACN4 as a function of Tc. Reprinted with permission from [63], S. Sinha Ray et al., Macromol.Rapid Commun. 24, 815 (2003). © 2003, Wiley-VCH Verlag GmbH & Co.


1e–980 90 100


N/µ

m–3

110 120 130 140 150

1e–8

1e–7

1e–6

1e–5

1e–4

1e–3

1e–2

1e–1

neat PLA

PLACN4

Figure 37. Nucleation density (N ) of pure PLA and PLACN4 as a function of Tc. Reprinted with permissionfrom [63], S. Sinha Ray et al., Macromol. Rapid Commun. 24, 815 (2003). © 2003, Wiley-VCH Verlag GmbH & Co.

however, the PLACN4 sample crystallized in a defect-ridden crystalline form. This unstablegrowth of crystallites of PLA in the presence of MMT particles may be due to the interca-lation of PLA chains into the silicate galleries.

9. MELT RHEOLOGY OF BIODEGRADABLEPOLYMER/LS NANOCOMPOSITES

The measurement of rheological properties of any polymeric material in a molten state is cru-cial to gain fundamental understanding of the processability of that material. In polymer/LSnanocomposites, measurements of melt rheological properties are not only important tounderstand the nature of processability of these materials, but also to determine thestrength of polymer-layered silicate interactions and the structure-property relationship inthe nanocomposites, because melt rheological behaviors are strongly influenced by theirnanoscale structure and interfacial characteristics.

9.1. Dynamic Oscillatory Shear Measurement

Generally, the rheology of polymer melts strongly depends on the temperature at which mea-surement is carried out. It is well known that for the thermorheological simplicity, isothermsof G′!�"1G′′!�" and complex viscosity !�2∗�!�"" can be superimposed by horizontal shifts

100

2000

4000

6000

8000

10000

15

(010) *

(203)

PLACN4

neat PLA

C18MMT

*

(110)or

(200)

*

(105)*

202Θ /degrees

Inte

nsity

/A.U

.

3025

Figure 38. Typical WAXD patterns of pure PLA and PLACN4 samples crystallized at 110�C for 1.5 hours.Reprinted with permission from [260], J. Y. Nam et al., Macromolecules 36, 7126 (2003). © 2003, American ChemicalSociety.


along the frequency axis:

bTG′!aT�1 Tref" = bTG′!�1 T "

bTG′′!aT�1 Tref" = bTG′′!�1 T "

�2∗�!aT�1 Tref" = �2∗�!�1 T "

where aT and bT are the frequency and vertical shift factors, respectively, and Tref is thereference temperature. All isotherms measured for pure PLA and for various PLA/C18-MMTcan be superimposed along the frequency axis.

In polymer samples, it is expected that at the temperatures and frequencies at whichthe rheological measurements were carried out, the polymer chains should be fully relaxedand exhibit characteristic homopolymer-like terminal flow behavior (i.e., the curves can beexpressed by a power law of G′!�" ∝ �2 and G′′!�" ∝ �".

The linear dynamic viscoelastic master curves for the neat PLA and various PLACNs areshown in Figure 39 [48]. These curves were generated by applying the time-temperaturesuperposition principle and shifting to a common temperature (Tref), using both the fre-quency shift factor (aT) and the modulus shift factor (bT ). The moduli of the nanocompos-ites increased with increasing clay loading at all frequencies (�). At high frequencies, thequalitative behavior of G′!�" and G′′!�" was essentially same and unaffected by frequen-cies. However, at low frequencies G′!�" and G′′!�" increased monotonically with increasingclay content. In the low frequency region, the curves can be expressed by the power lawof G′!�" ∝ �2 and G′′!�" ∝ � for neat PLA, suggesting that this is similar to those ofthe narrow Mw distribution homopolymer melts. On the other hand, for aT < 5 rad.s−1,the viscoelastic response [particularly G′!�"] for all of the nanocomposites displayed sig-nificantly diminished frequency dependence as compared to the matrices. In fact, for allPLACNs, G′!�" became nearly independent at low aT� and exceeded G′′!�", characteristicof materials exhibiting a pseudosolid-like behavior [82]. The terminal zone slope values ofboth neat PLA and PLACNs were estimated at the lower aT� region ( < 10 rad.s−1), andare presented in Table 11. The lower slope values and the higher absolute values of thedynamic moduli indicate the formation of “spatially linked” structures in the PLACNs in themolten state [83]. Because of this structure or highly geometric constraints, the individualstacked silicate layers are incapable of freely rotating. Hence, by imposing small aT�, therelaxations of the structure are almost completely prevented. This type of prevented relax-ation due to the highly geometric constraints of the stacked and intercalated silicate layersleads to the pseudosolid-like behavior observed in PLACNs. This behavior probably corre-sponds to the shear-thinning tendency, which dramatically appears in the viscosity curves(aT� < 5 rad.s−1) (�2∗� versus aT �" [84]. Such features are highly dependent on the shearrate in the dynamic measurement due to the formation of shear-induced alignment of thedispersed clay particles [85].

The temperature dependence frequency shift factors (aT, Williams-Landel-Ferry type [86])used to generate the master curves shown in Figure 39 are shown in Figure 40. The depen-dence of the frequency shift factors on the silicate loading suggests that the temperature-dependent relaxation process observed in the viscoelastic measurements are somehowaffected by the presence of silicate layers [82]. In case of nylon 6/LS nanocomposites, wherethe hydrogen bonding of the already formed hydrogen-bonded molecule to the silicate sur-face, the system exhibits a high level of flow activation energy [estimated from slope inFigure 40(a)], near one order higher in magnitude compared to that of neat nylon 6 [87].

The shift factor (bT) shows significant deviation from a simple density effect, but it wouldbe expected that the values would not vary far from unity [86]. One possible explanation isinternal structure development occurring in PLACNs during measurement (shear process).The alignment of the silicate layers probably supports the PCN melts to withstand the shearforce, thus leading to the increase in the absolute values of G′!�" and G′′!�".

Figure 41 represents the clay content-dependent (weight percent) flow activation energy(Ea) of pure PLA and various PLA/C18-MMTs obtained from Arrhenius fits of mastercurves. It is clearly observed that Ea values significantly increased in PLA/C18-MMTs con-taining 3 wt% of C18-MMT and then correlates fairly well with increasing C18-MMT content.


103

104

aTω /rad.s–1

|η* |

/Pa.

s

10–2 10–1 100 101 102

101

102

103

104

10–1

100

101

102

103

104

105

b T G

′′ /P

a 1

2

Tref = 175 °C

PLAPLACN4PLACN5PLACN7

b TG

′ /Pa 1

0.5

2

Figure 39. Reduced frequency dependence of storage modulus, loss modulus, and complex viscosity of neat PLAand various PLACNs. Reprinted with permission from [63], S. Sinha Ray et al., Macromol. Rapid Commun. 24, 815(2003). © 2003, Wiley-VCH Verlag GmbH & Co.

Table 11. Terminal slopes of G′ and G′′ versus aT� for PLAand various PLACNs.

System G′ G′′

PLA 1%3 0%9PLACN4 0%2 0%5PLACN5 0%18 0%4PLACN7 0%17 0%32

10–1

100

101

2.08 2.12 2.16 2.2 2.24


b T

1/T × 103/K

Tref =175 °C

(b)

10–2

10–1

100

2.08 2.12 2.16 2.2 2.24

Tref = 175 °C


a T

1/T ×103/K

(a)

Figure 40. (a) Frequency shift factors (aT) and (b) modulus shift factors (bT) as a function of temperature.


–1

160

180

200

220

240

0 1 2

MMT/wt.%E

a/K

J.m

ol–1

3 4 5

Tref = 175 °C

6

Figure 41. Flow activation energy of pure PLA and various PLA/C18-MMT nanocomposites as a function of MMTcontent. Reprinted with permission from [63], S. Sinha Ray et al., Macromol. Rapid Commun. 24, 815 (2003).© 2003, Wiley-VCH Verlag GmbH & Co.

This result indicates that in the presence of MMT, it is very difficult for the materials to flow.This behavior is also ascribed to the formation of spatially linked structures in PLA/C18-MMTs in molten states.

9.2. Steady Shear Flow

The steady shear rheological behaviors of neat PBS and various PBSCNs are shown inFigure 42. The steady viscosity of PBSCNs is enhanced considerably with time at all shearrates, and at a fixed shear rate the steady viscosity increases monotonically with increasingsilicate loading [50]. On the other hand, all intercalated PBSCNs exhibit strong rheopexybehavior, and this becomes prominent at low shear rates, while neat PBS exhibits a time-independent viscosity at all shear rates. With increasing shear rates, the shear viscosity attainsa plateau after a certain time, and the time required to attain this plateau decreases withincreasing shear rates. This type of behavior may be due to the planer alignment of the clayparticles toward the flow direction under shear. When the shear rate is very slow (0.001 s−1),clay particles take a longer time to attain complete planer alignment along the flow direction,and this measurement time (1,000 s) is too short to attain such alignment. Therefore, strongrheopexy behavior results. On the other hand, under high shear rates (0.005 s−1 or 0.01 s−1),this measurement time is long enough to attain such alignment, and hence, nanocompositesshow time-independent shear viscosity after certain periods of time.

Figure 43 shows shear rate dependence of viscosity for neat PBS and correspondingnanocomposites measured at 120�C. Neat PBS exhibits almost Newtonian behavior at allshear rates, whereas nanocomposites exhibit non-Newtonian behavior. At very low shearrates, shear viscosity of nanocomposites initially exhibits some shear-thickening behavior thatcorresponds to the rheopexy behavior we observed at very low shear rates (see Fig. 42).After that initial period, all nanocomposites show very strong shear-thinning behavior at allshear rates; this behavior is analogous to the results obtained with dynamic oscillatory shearmeasurements [48]. Additionally, at very high shear rates the viscosities of nanocompositesare comparable to that of neat PBS. These observations suggest that the silicate layers arestrongly oriented toward the flow direction at high shear rates, and shear-thinning behaviorat high shear rates is dominated by that of neat polymers.

PLS nanocomposites always exhibit significant deviation from the empirical Cox-Merzrelation [88], while all neat polymers obey that relation. (The Cox-Merz relationrequires that for �̇ = �, the viscoelastic data should obey the relationship 2!�̇" = �2∗�!�").They speculated there are two possible reasons for nanocomposite deviation from the Cox-Merz relation: First, this rule is only applicable for homogenous systems like homopolymermelts, but nanocomposites are heterogeneous systems. For this reason, this relation workswell in the case of neat polymers [50]. Second, the structure formation is different whennanocomposites are subjected to dynamic oscillatory shear and steady shear measurements.


PBS

Shear rate =0.001s–1

Shear rate =0.005s–1

Shear rate = 0.01s–1

Temperature =120 °C

103

102

104

104

105PBSCN2

103

104

105PBSCN3

103

η/P

a.s

104

105PBSCN4

103

104

PBSCN6

103

10 102 103

Time/s0 101

Figure 42. Time variation of shear viscosity for PBSCN. Reprinted with permission from [50], S. Sinha Ray et al.,Macromolecules 36, 2355 (2003). © 2003, American Chemical Society.

9.3. Elongational Flow and Strain-Induced Hardening

Okamoto and his colleagues [89] first conducted elongation tests of PP/LS nanocomposites(PPCN4) in the molten state at constant Hencky strain rates (�̇0), using elongation flowoptorheometry [90]. They also attempted to control the alignment of the dispersed silicatelayers with nanometer dimensions of intercalated PPCNs under uniaxial elongational flow.

Figure 44(a) shows double-logarithmic plots of transient elongational viscosity (2E) againsttime (t) observed for PLA/C18-MMT containing 4 wt% of OMLS at 170�C with differ-ent Hencky strain rates (�̇0) ranging from 0.01–1 s−1. We observed a very strong tendency

102

103

104

105

106

PBSPBSCN2PBSCN3PBSCN4PBSCN6

η/P

a.s

10–4 10–3 10–2 10–1 100 101

.γ/s–1

Temp. = 120 °C

Figure 43. Shear viscosity as a function of shear rates for the shear rate sweep test. Reprinted with permissionfrom [50], S. Sinha Ray et al., Macromolecules 36, 2355 (2003). © 2003, American Chemical Society.


0.1104 10–1

100

101

105

106

107

108

1 10 100 1000 10–3 10–2 10–1

PLACN3

Hen

cky

stra

in

Elo

ngat

iona

l vis

cosi

ty/P

a.s

Hencky strain rate/s–1

Hencky strain rate/s–1

Time/s

Temperature = 170 °C

PLACN3

Temperature = 170 °C

100 101

1

0.5

0.05

0.1

0.01

(a) (b)

Figure 44. (a) Time variation of elongational viscosity for PLA/C18-MMT(4) melt at 170�C; (b) Strain rate depen-dence of uprising Hencky strain. Reprinted with permission from [63], S. Sinha Ray et al., Macromol. RapidCommun. 24, 815 (2003). © 2003, Wiley-VCH Verlag GmbH & Co.

toward strain-induced hardening in the PLA/C18-MMT melt. In the early stage, 2E graduallyincreases with t but almost independent of �̇0, which we generally call the linear region ofthe viscosity curve. After a certain time (t2E

), which we call the uprising time [marked withupward arrows in Figure 44(a)], strongly dependent on �̇0, we saw rapid upward deviationof 2E from the curves of the linear region. On the other hand, we tried but were unable toaccurately measure the elongational viscosity of pure PLA. We concluded that very low shearviscosity of pure PLA is the main reason for this, because the minimum viscosity range ofour instrument was greater than 104 Pa.s. However, we confirmed that neither strain-inducedhardening in elongation nor rheopexy in shear flow took place in pure PLA having the samemolecular weight and polydispersity as PLA/C18-MMT.

As in PP/LS systems, the extended Trouton rule [320!�̇; t) � 2E!�̇0; t)] also does not holdfor PLA/C18-MMT melts, as opposed to the melt of pure polymers. These results indicatethat in PLA/C18-MMT, the flow-induced internal structural changes also occur in elongationflow [89], but the changes are quite different from shear flow. The strong rheopexy observedin shear measurements for PLA/C18-MMT at very slow shear rates reflects the fact that theshear-induced structural change involved a process with an extremely long relaxation time.

Regarding elongation-induced structure development, Figure 44(b) shows Hencky strainrate dependence of the uprising Hencky strain (�2E

" = �̇0 × t2Etaken for PLA/C18-MMT at

170�C. The �2Eincreases systematically with the �̇0. The lower the value of �̇0, the smaller

the value of �2E. This tendency probably corresponds to the rheopexy of PLA/C18-MMT

under slow shear flow.

10. PROCESSING OPERATIONSIn the preceding sections, dynamic measurements indicated the formation of spatially linkedstructures in PLA/C18-MMT melts. Shear measurements revealed very strong rheopexybehavior in PLA/C18-MMT under very slow shear fields. Under uniaxial elongation flow,PLA/C18-MMT exhibited very high viscosity and a tendency toward strong strain-inducedhardening, which probably originated in the perpendicular alignment of the silicate lay-ers to the stretching direction. This strain-induced hardening behavior is an indispensablecharacteristic for foam processing, due to its capacity to withstand the stretching force expe-rienced during the latter stages of bubble growth. To evaluate the performance potential ofbiodegradable PLA/C18-MMT in foam applications, Okamoto et al. conducted foam process-ing of one representative nanocomposite, PLA/C18-MMT, using a newly developed pressurecell technique. The goal was to devlop an advanced biodegradable foam with excellent mate-rials properties [91].

Figure 45 shows typical scanning electron microscope (SEM) images of the freeze frac-ture surfaces of foamed neat PLA and two different foamed nanocomposites. PLA andPLA/MMT(ODA)5 were foamed at 140�C and PLA/MMT(SBE)5 was foamed at 165�C.


Neat PLA 100 µm PLA/ODA5 10 µm PLA/SBE5 5 µm000000 000000000000

Figure 45. SEM images of freeze fracture surfaces of neat PLA and two different nanocomposite foams. Reprintedwith permission from [91], Y. Fujimoto et al., Macromol. Rapid Commun. 24, 457 (2003). © 2003, Wiley-VCHVerlag GmbH & Co.

All of the foams exhibited desirable closed-cell structures. Homogeneous cells were formedin both nanocomposites, while the neat PLA foam showed nonuniform cell structures havinglarge cell size (∼230 /m). Also, the nanocomposite foams showed smaller cell size (d) andgreater cell density (Nc) compared to neat PLA foam, suggesting that the dispersed silicateparticles act as nucleating sites for cell formation [92, 93].

For the nanocomposite foams, we calculated the distribution function of cell size fromSEM images, and these are presented in Figure 46. The nanocomposite foams conformedwell to Gaussian distributions. In the case of PLA/MMT(SBE)5 [see Fig. 46(b)], we cansee that the width of the distribution peaks, which indicates the cell size dispersity, becamenarrow, accompanied by finer dispersion of silicate particles. From the SEM images, wequantitatively calculated various morphological parameters of two different nanocompositefoams, which are summarized in Table 12.

From Table 12, we can see that the PLA/MMT(SBE)5 (nanocellular) foam has a smallerd value ( � 360 nm) and a huge Nc! � 1%2 × 1014 cell · cm−3" compared to that ofPLA/MMT(ODA)5 (microcellular) foam (d � 2%59 /m and Nc � 3%56 × 1011 cell · cm−3).These results indicate that the nature of the dispersion has vital role in controlling the size ofcells during foaming. On the other hand, the very high value of Nc in the PLA/MMT(SBE)5foam indicates that the final 5f is controlled by the competitive process in the cell nucle-ation (its growth and coalescence). Cell nucleation in nanocomposite systems took placein the boundary between the matrix polymer and the dispersed silicate particles. For thisreason, cell growth and coalescence are strongly affected by the characteristic parame-ter (see Table 13) and the storage and loss modulus ( � viscosity component) of thematerials during processing. This may create nanocellular foams without the loss of mechan-ical properties in polymeric nanocomposites. Okamoto et al. described a novel foamingapproach for biodegradable polylactide/layered silicate nanocomposites that results in con-trolled structure of nanocomposite foams (from microcellular to nanocellular).

00

5

10

20

30

35

25

15

0

5

10

20

30

PLA/ODA5 PLA/SBE525

15

1 2 3 4 5

Cell size/µm

Fra

ctio

n/%

Fra

ctio

n/%

6 7 8 9 10 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0

Cell size/µm

Figure 46. Cell size distribution of two different nanocomposite foams. Average values of d in /m and variances6 2

d in /m2 in the Gaussian fit through the data are 2.59 and 0.551, respectively, for PLA/ODA5 foam, and 0.36 and0.011, respectively, for PLA/SBE5 foam. Reprinted from with permission [91], Y. Fujimoto et al., Macromol. RapidCommun. 24, 457 (2003). © 2003, Wiley-VCH Verlag GmbH & Co.


Table 12. Morphological parameters of two different nano-composite foams.

Parameters PLA/ODA5 PLA/SBE5

5f (g.cm−3) 0%46 0%57d(/m) 2%59 0%36Nc × 10−11 (cell.cm−3) 3%56 1172 (/m) 0%66 0%26d(�clay) 10%1 4%5d(Lclay) 5%78 1%8 (Lclay) 1%47 1%3

11. CREATING POROUS CERAMIC MATERIALS VIAPLA/LS NANOCOMPOSITES

Very recently, a new route for the preparation of porous ceramic materials from thermo-setting epoxy/LS nanocomposites was first demonstrated by Brown et al. [94]. This routeoffers attractive potential for diversification and applications of PLS nanocomposites (CNs).Okamoto and coworkers reported on their development of a novel porous ceramic materialvia burning of the PLA/LS system (PLACN) [40]. In the PLACN containing 3.0 wt.% inor-ganic clay. Figure 47 shows a SEM image of the fractured surface of the porous ceramicmaterial prepared from simple burning of PLACN in a furnace up to 950�C. After completeburning, as seen in the figure, the PLACN became a white mass with a porous structure.The bright lines in the SEM image correspond to the edges of the stacked silicate layers.In the porous ceramic material, the silicate layers form a “house of cards” structure con-sisting of large plates having lengths of ∼1,000 nm and thicknesses of ∼30–60 nm. Thisimplies that the further stacked platelet structure is formed during burning. The materialexhibits an open cell-type structure having 100–1,000-nm diameter voids, a BET surfacearea of 31 m2/g−1, and low density of porous material (0.187 g/mL−1 estimated by the buoy-ancy method). The BET surface area value of MMT is 780 m2/g and that of the porousceramic material is 31m2/g, suggesting approximately 25 MMT plates stacked together. WhenMMT is heated above 700�C (but below 960�C), first all of the OH groups are eliminatedfrom the structure and, thus, MMT is decomposed into a nonhydrated aluminosilicate. Thistransformation radically disturbs the crystalline network of the MMT, and the resultingdiffraction pattern is indeed often typical of an amorphous (noncrystalline) phase. A roughestimate of the compression modulus (K) is on the order of ∼1.2 MPa, which is five ordersof magnitude lower than the bulk modulus of MMT (∼102 GPa) [27]. In the stress-straincurve, the linear deformation behavior is well described in the early stage of the defor-mation (i.e., the deformation of the material closely resembles that of ordinary polymeric

Table 13. Comparison of some characteristic parameters of PLA/ODA5and PLA/SBE5.

Parameters PLA/ODA5 PLA/SBE5

Mw × 10−3 (g.mol−1) 161 163Mw/Mn 1.58 1.61Tg(�C) 59.2 59.7Tm(�C) 169.8 169.3d001(nm) 3.03 2.85Lclay(nm) 448 ± 200 200 ± 25�clay(nm) 255 ± 137 80 ± 20D(nma) 38 12.37Lclay(D) 12 18D(d001) 12.5 4.3G′

PLACN/G′bPLA 1.65 1.43

aCalculated on the basis of Scherrer equation. bG′PLACN and G′

PLA are the stor-age modulus of PLACN and PLA, respectively, at 25�C. Mw and Mw/Mn for neatPLA are 177,000 g/mol and 1.58, respectively.


Figure 47. SEM image of porous ceramic material after being coated with a platinum layer (∼10 nm thickness).Reprinted with permission from [40], S. Sinha Ray et al., Nano Letts. 2, 423 (2002). © 2002, American Chemi-cal Society.

foams) [95]. This open cell-type porous ceramic material with its “house-of-cards” structureis expected to provide strain recovery (up to 8% strain) and an excellent energy dissipationmechanism after unloading in the elastic region, probably having each plate bend like a leafspring (see Fig. 48). This new porous ceramic material possesses elastic features and is verylightweight. This new route for the preparation of porous ceramic materials via burning ofnanocomposites can be expected to pave the way for much a broader range of applicationsfor PLS nanocomposites.

200

Deformation rate = 5%/min

Deformation rate = 5%/min

150

100

50

00 5 10

Strain/%

Stre

ss/K

Pa.

30

60

45

15

0

Stre

ss/K

Pa.

2015

0 2 4

Strain/%

A

B

6 8

1st run2nd run3rd run

Temp = 25 °C

Temp = 25 °C

Figure 48. Stress-strain curve (a) and the strain recovery behavior (b) of porous ceramic material under compres-sion test. The author conducted the compression test using porous ceramic material of 2 × 2 × 1%5 mm3 size.Reprinted with permission from [40], S. Sinha Ray et al., Nano Letts. 2, 423 (2002). © 2002, AmericanChemical Society.


Figure 49. Schematic illustration of the main concept of PLA/OMLS nancomposites.

12. CURRENT STATUS AND FUTURE PROSPECTS OFBIODEGRADABLE NANOCOMPOSITES

Development of PLS nanocomposites is one of the latest evolutionary steps in polymer tech-nology. Nanocomposites offer attractive potential for diversification and new applications ofconventional polymeric materials.

Since the possibility of direct melt intercalation was first demonstrated by Giannelis andhis colleagues [25], melt intercalation has become a mainstream method of preparing inter-calated polymer nanocomposites without in situ intercalative polymerization. It is a quiteeffective technology within the PLS nanocomposite industry.

Biodegradable polymer-based nanocomposites have a great deal of future promise forpotential applications as high-performance biodegradable materials. These are entirely newtypes of materials based on plant and other natural materials (OMLS). When disposed of incompost, these are safely decomposed into CO2, water, and humus through the activity ofmicroorganisms. The CO2 and water could become corn or sugar cane again through plantphotosynthesis (see Fig. 49). Undoubtedly, these unique properties originating in controllednanostructures pave the way to a much broader range of applications (already commer-cially available through Unitika Ltd., Japan), and open a new dimension for plastics andcomposites.

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