Polymer/Clay Nanocomposites 1588830640

Citation preview

Encyclopedia of Nanoscience and Nanotechnology

www.aspbs.com/enn

Polymer/Clay Nanocomposites Masami Okamoto Toyota Technological Institute, Tempaku, Nagoya, Japan

CONTENTS 1. Introduction 2. Structure of Clay and Its Modification with Surfactants 3. Preparative Methods and Structure of Polymer/Clay Nanocomposites 4. Characterization of Polymer/Clay Nanocomposites 5. Types of Polymers Used for Polymer/Clay Nanocomposite Preparation 6. Materials Properties of Polymer/Clay Nanocomposites 7. Crystallization of Polymer/Clay Nanocomposites 8. Melt Rheology of Polymer/Clay Nanocomposites 9. Processing Operations of Polymer/Clay Nanocomposites 10. Conclusions Glossary References

1. INTRODUCTION Polymer/clay nanocomposites (PCNs) are a new class of materials which have attracted much attention from both scientists and engineers in recent years due to their excellent properties such as high dimensional stability, heat deflection temperature, gas barrier performance, reduced gas permeability, optical clarity, flame retardancy, and enhanced mechanical properties when compared with the pure polymer or conventional composites (micro- and macrocomposites) [1–15]. The first successful PCN appeared about 10 years ago through the pioneering efforts of a research team from Toyota Central Research & Development Co. Inc. (TCRD) in the form of a Nylon 6/clay hybrid [1
2]. ISBN: 1-58883-064-0/$35.00 Copyright © 2004 by American Scientific Publishers All rights of reproduction in any form reserved.

Earlier attempts to prepare polymer/clay composites are found in almost half-a-century old patent literatures [16
17]. In such cases, incorporation of 40 to 50 wt% clay mineral (bentonite, hectorite, etc.) into a polymer was attempted but ended with unsatisfactory results: The maximal modulus enhancement was only around 200%, although the clay loading was as much as 50 wt%. The failure was obvious because they failed to achieve good dispersion of clay particles in the matrix, in which clay minerals existed as agglomerated tactoids. Such a poor dispersion of the clay particles could improve the material rigidity but certainly sacrifice the strength, the elongation at break, and the toughness of the material [16
17]. A prime reason for this impossibility of improving the tactoid dispersion into well-dispersed exfoliated monolayers of the clay is obviously due to the intrinsic incompatibility of hydrophilic layered silicates with hydrophobic engineering plastics. One attempt at circumventing this difficulty was made by Unitika Ltd. [18] about 30 years ago in preparing Nylon 6/clay composites (not nanocomposites) via in-situ polymerization of -caprolactam with montmorillonite, but the results were not very good. The first major breakthrough of the problem was brought about in 1987, when Fukushima and Inagaki of TCRD, via their detailed study of polymer/layered silicate composites, persuasively demonstrated that lipophilization, by replacing inorganic cations in galleries of the native clay with alkylammonium surfactant, successfully made them compatible with hydrophobic polymer matrices [19]. The modified clay was thus called lipophilized clay, organo-phillic clay, or simply organo-clay (organoclay). Furthermore, they found that the lipophilization enabled one to expand clay galleries and exfoliate the silicate layers into single layers of a nanometer thickness. Six years later, in 1993, Usuki, Fukushima, and their colleagues at TCRD successfully prepared, for the first time, exfoliated Nylon 6/clay hybrid (NCH) via in-situ polymerization of -caprolactam, in which alkylammonium-modified organoclay was thoroughly dispersed in advance [1
2]. The resulting composite of the loading of only 4.2 wt% clay possessed a doubled modulus, a 50% enhanced strength, and an increase in heat distortion by 80
C compared to the neat Nylon 6, as shown in Table 1. This invention opened up a

Encyclopedia of Nanoscience and Nanotechnology Edited by H. S. Nalwa Volume 8: Pages (791–843)

Polymer/Clay Nanocomposites

792

Properties Clay content (wt%) Specific gravity Tensile strength (MPa) Tensile modulus (GPa) Impact (kJ/m2 ) HDT (
C at 1.8 MPa)

Nylon 6 nanocomposite

Neat Nylon 6

42 115 107 21 28 147

0 114 69 11 23 65

new era of engineering materials, which we may call “polymer/clay nanocomposites.” Thus, along the stream of development in PCN technologies, many studies have been devoted to PCNs since their intrinsically excellent properties of polymer should have attractive potential for continuous expansion of application versatility. Apart from this, rheological behavior, especially elongational and shear flow behavior in the molten state and crystallization behavior under supercooled state of PCNs, has not well been reported yet, although such knowledge should be indispensable in relation with their performance in processing operations. One objective of this chapter is to focus on a profound understanding of PCNs for their innovations in practical material production. For this purpose, it is indispensable to illuminate the nanostructure as well as rheological properties of PCNs to assess appropriate processing conditions for designing and controlling their hierarchical nanostructure, which must be closely related to their material performance.

2. STRUCTURE OF CLAY AND ITS MODIFICATION WITH SURFACTANTS The commonly used clays for the preparation of PCNs belong to the same general family of 2:1 layered or phyllosilicates. (see Table 2.) Their crystal structure consists of layers made up of two silica tetrahedral fused to an edgeshared octahedral sheet of either aluminum or magnesium hydroxide. The layer thickness is around 1 nm and the lateral dimensions of these layers may vary from 30 nm to several micrometers and even larger depending on the particular layered silicate. Stacking of the layers leads to a regular van der Waals gap between the layers called the interlayer or Table 2. Chemical formula and characteristic parameter of commonly used 2:1 phyllosilicates.

2:1 Phyllosilicates

Chemical fourmulaa

Montmorillonite

Mx (Al4−x Mgx Si8 O20 (OH)4 Mx (Mg6−x Lix Si8 O20 (OH)4 Mx Mg6 (Si8−x Alx  Si8 O20 (OH)4

Hectorite Saponite a

CEC (mequiv /100 gm)

Particle length (nm)

110

100–150

120

200–300

gallery. Isomorphic substitution within the layers (for example, Al+3 replaced by Mg+2 or by Fe+2 , or Mg+2 replaced by Li+1 ) generates negative charges that are counterbalanced by alkali and alkaline earth cations situated inside the galleries, as shown in Figure 1. The most commonly used layered silicates are montmorillonite (MMT) hectorite and saponite having different chemical formulas, respectively, Mx (Al4−x Mgx Si8 O20 (OH)4 , Mx (Mg6−x Lix Si8 O20 (OH)4 , and Mx (Si8−x Alx Si8 O20 (OH)4 (x = 03–1.3). The type of clay is characterized by a moderate surface charge (cation exchange capacity) (CEC of 80–120 mequiv/100 gm) and layer morphology. These clays are only miscible with hydrophilic polymers, such as poly(ethylene oxide) (PEO) [20] and poly(vinyl alcohol) (PVA) [21]. To improve miscibility with other polymer matrices, one must convert the normally hydrophilic silicate surface to organophilic, which makes possible intercalation of many engineering polymers. Generally, this can be done by ion-exchange reactions with cationic surfactants including primary, secondary, tertiary, and quaternary alkyl ammonium or alkylphosphonium cations. The role of alkylammonium or alkylphosphonium cations in the organosilicates is to lower the surface energy of the inorganic host and improve the wetting characteristics with the polymer matrix and results in a larger interlayer spacing. One can evaluate that about 100 alkylammonium salt molecules are localized near the individual silicate layers (∼8 × 10−15 m2 ) and active surface area (∼800 m2 /g). Additionally, the alkylammonium or alkylphosphonium cations could provide functional groups that can react with the polymer matrix or in some cases initiate the polymerization of monomers to improve the strength of the interface between the inorganic and the polymer matrix [22
23]. Vaia et al. [24] have shown that alkyl chains can vary from liquidlike to solidlike, with the liquidlike 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 due to the relatively small energy differences between the trans and

Tetrahedral

Basal spacing

Table 1. Mechanical and thermal properties of Nylon 6/clay hybrid.

~1nm

Octahedral

Tetrahedral

Exchangeable cations

Al, Fe, Mg, Li OH O

866

Li, Na, Rb, Cs

50–60

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

Figure 1. Structure of 2:1 phyllosilicates. Adapted with permission from [32], E. P. Giannelis et al., Adv. Polym. Sci. 138, 107 (1999). © 1999, Springer-Verlag.

Polymer/Clay Nanocomposites

793 Polymerization can be initiated either by heat or radiation, by the diffusion of a suitable initiator, or by an organic initiator or catalyst fixed through cation exchange inside the interlayer before the swelling step by the monomer.

(a)

(b)

(c) Figure 2. Alkyl chain aggregation models. (a) Short chain lengths: the molecules are effectively isolated from each other. (b) Medium lengths: quasi-discrete layers form with various degrees of in-plane disorder and interdigitation between the layers. (c) Long lengths: interlayer order increases leading to a liquid–crystalline polymer environment. Open circles represent the CH2 segments while cationic head groups are represented by filled circles. Reprinted with permission from [24], R. A. Vaia et al., Chem Mater. 6, 1017 (1994). © 1994, American Chemical Society.

gauche conformers; the idealized models described earlier assume all trans conformations. In addition, for the longer chain length surfactants, the surfactants in the layered silicate can show thermal transition akin to melting or liquid– crystalline to liquidlike transitions upon heating.

3. PREPARATIVE METHODS AND STRUCTURE OF POLYMER/CLAY NANOCOMPOSITES So far there have been much literature available devoted to developing PCNs with different combinations of organoclays and matrix polymers such as epoxy polymer resin [5], polyurethanes (PU) [6], PEI [7], polybenzoxazine [8], polypropylene (PP) [9
14
15
25], polystyrene (PS) [10
11], poly(methyl methacrylate) (PMMA) [11
12], and poly(-caprolactone) (PCL) [13] liquid crystalline polymers (LCP) [26] by employing somewhat different technologies appropriate to each. The technologies are broadly classified into three main categories.

3.1. Intercalation of Polymer or Prepolymer from Solution This is based on a solvent system in which polymer or prepolymer is soluble and the silicate layers are swellable. The layered silicate is first swollen in a solvent, such as water, chloroform, 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 PCNs.

3.2. In-situ Intercalative Polymerization Method In this method, the organoclay is swollen within the liquid monomer or a monomer solution so that the polymer formation can occur in between the intercalated sheets.

3.3. Melt Intercalation Method This method involves annealing, statically or under shear, a mixture of the polymer and organoclay above the softening point of the polymer. This method has great advantages over either in-situ intercalative polymerization or polymer solution intercalation. First, this method is environmentally benign due to the absence of organic solvents. Second, it is compatible with current industrial process, such as extrusion and injection molding. The melt intercalation method allows the use of polymers which were previously not suitable for in-situ polymerization or the solution intercalation method. Other possibilities are exfoliation–adsorpsion [21
27] and template synthesis [28
29]. Nowadays this solvent-free method is much preferred for practical industrial material production for its high efficiency and possibility of avoiding environmental hazards.

3.4. Structure of PCNs The in-situ polymerization is employed for the first time in NCH production [1
2], and the melt intercalation is the direct blending of organoclay into modified polymer matrix such as used in PP/clay nanocomposite [9
14
15
25
30]. Since Vaia et al. [31] found that the melt compounding of polymers with clay is possible without using organic solvent, nanocomposite preparations via this method have been widely used in practice, especially for polyolefin-based nanocomposites. This process involves annealing, statically or under shear, a mixture of the polymer and organoclay above the softening point of the polymer. During the anneal, the polymer chains diffuse from the bulk polymer melt into the galleries between the silicate layers. Depending on the degree of penetration of the matrix into the organically modified layered silicate galleries, nanocomposites are obtained with structures ranging from intercalated to exfoliated. Polymer penetration resulting in finite expansion of the silicate layers produces intercalated nanocomposites consisting of well-ordered multilayers with alternating polymer/silicate layers and a repeat distance of few nanometers (intercalated; see Fig. 3) [30]. On the other hand, extensive polymer penetration resulting in disordered and eventual delamination of the silicate layers produces near to exfoliated nanocomposites consisting of individual silicate layers dispersed in polymer matrix (exfoliated) [32]. Under some conditions, the intercalated nanocomposites exhibit flocculation because of the hydroxylated edge–edge interaction of silicate layers (intercalated and flocculated). The length of the oriented collections in the range of 300–800 nm is far larger than original clay (mean diameter  150 nm) [33
34]. Such flocculation presumably is governed by an interfacial energy between polymer matrix and organoclays and is controlled by ammonium cation–matrix polymer interaction. The polarity of the matrix polymer is of fundamental importance in controlling the nanoscale structure.

Polymer/Clay Nanocomposites

794 2000

Al, Fe, Mg,Li

L 1 nm

Original OMLS

OH O Li, Na, Ra, Cs

Tetrahedral

1500

Octahedral

1000

Tetrahedral

500

One Clay Platelet L: 100 – 200 nm in case of MMT

dd clay clay

0

Exchangeable cations

Intercalated

clay clay

1500

Lclay L clay

The structure of 2:1 layered silicates

Intensity /A.U.

Form factors of dispersed clay

Intercalated

200 nm

1000 500 0

Intercalated-and-flocculated 1500 1000

Intercalated-and-flocculated Intercalated

Intercalated-and-flocculated

Exfoliated

Figure 3. Schematic illustration of three different types of thermodynamically achievable polymer/clay nanocomposites. Reprinted with permission from [327], S. Sinha Ray et al., Macromolecules 36, 2355 (2003). © 2003, American Chemical Society.

200 nm

500 0

Exfoliated 1500 1000 500

4. CHARACTERIZATION OF POLYMER/CLAY NANOCOMPOSITES Generally the structure of the PCNs has typically been established using a wide-angle X-ray diffraction (WAXD) analysis and transmission electron microscope (TEM) observation. Due to its easiness and availability WAXD is most commonly used to probe the PCN structure and sometimes to study the kinetics of the polymer melt intercalation. By monitoring the position, shape, and intensity of the basal reflections from the distributed silicate layers, the PCN structure, either intercalated or exfoliated, may be identified. For example, in case of exfoliated nanocomposites, the extensive layer separation associated with the delamination of the original silicate layers in the polymer matrix results in the eventual disappearance of any coherent X-ray diffraction from the distributed silicate layers. On the other hand, for intercalated nanocomposites, the finite layer expansion associated with the polymer intercalation results in the appearance of a new basal reflection corresponding to the larger gallery height. Although WAXD offers a convenient method to determine the interlayer spacing of the silicate layers in the original layered silicates and in the intercalated nanocomposites (within 1–4 nm), little can be said about the spatial distribution of the silicate layers or any structural inhomogeneities in the PCNs. Additionally, some layered silicates initially do not exhibit well-defined basal reflection. Thus, peak broadening and intensity decreases are very difficult to study systematically. Therefore, conclusions concerning the mechanism of nanocomposite formation and their structure based solely on WAXD patterns are only tentative. On the other hand, TEM allows a qualitative understanding of the internal structure, spatial distribution of the various phases, and defect structure through direct visualization. However, special care must be exercised to guarantee a representative cross section of the sample. The WAXD patterns and corresponding TEM images of three different types of nanocomposites are presented in Figure 4.

Exfoliated

0 2

4

6

8

200 nm

10

2Θ/degrees

Figure 4. (Left) WAXD patterns and (Right) TEM images of three different types of nanocomposites.

5. TYPES OF POLYMERS USED FOR POLYMER/CLAY NANOCOMPOSITE PREPARATION 5.1. Vinyl Polymers These include the vinyl addition polymers derived from common monomers like methyl methacrylate (MMA) [11
33
35– 44], methyl methacrylate copolymers [12
33
45
46], other acrylates [47–49], acrylic acid [50
51], acrylonitrile [52–55], styrene (S) [10
11
31
56–80], 4-vinylpyridine [81], acrylamide [82
83], and tetrafluoro ethylene [84]. In addition, selective polymers like PVA [21
85–87], poly(N -vinyl pyrrolidone) [88
89], poly(vinyl pyrrolidinone) [90–93], poly(vinyl pyridine) [94], poly(ethylene glycol) (PEG) [95], poly(ethylene vinyl alcohol) [96], poly(vinylidene fluoride) [97], poly(p-phenylenevinylene) [98], polybutadiene [99], poly(styrene-co-acrylonitrile) (SAN) [100], ethyl vinyl alcohol copolymer [101], polystyrene–polyisoprene diblock copolymer [102
103], and others [104] have also been used.

5.1.1. PMMA/Clay Nanocomposites Okamoto et al. [11
33] used organically modified smectite clays for the preparation of PMMA and PS nanocomposites. Organically modified smectite clays (SPN and STN) were prepared by replacing Na+ -smectite with QA (quaternary), oligo (oxypropylene) diethylmethylammonium cation (SPN) or methyltrioctilammonium cation (STN) by exchange reaction. In a typical synthesis, both lipophilized smectite clays were (SPN and STN) dispersed in MMA and S via ultrasonication at 25
C for 7 h to obtain suspensions. After that t-butyl peroxy-2-ethylhexanate and/or 1,1-bis(t-butyl peroxy)

Polymer/Clay Nanocomposites

795

cyclohexane as an initiator was added to the suspensions and then free-radical polymerization was carried out in the dark at 80
C for 5 h (for MMA) and at 100
C for 16 h (for S) in a silicon oil bath. For comparison authors also prepared PMMA and/or PS including QA as the reference under the same conditions and procedure. WAXD analyses were performed directly from the suspensions of MMA/SPN, MMA/STN, and S/SPN, and corresponding nanocomposites. From WAXD patterns of MMA/STN suspension (see Fig. 5a), the higher order peaks of interlayer spacing corresponding to d 002 and d 003 are clearly observed along with basal spacing d 001 peak, suggesting MMA intercalated into the STN gallery without the loss of layer structure, while the corresponding nanocomposite, PMMA/STN, exhibits rather broad Braggs peaks, indicating the formation of disordered intercalated structure. In contrast, for MMA/SPN suspension (see Fig. 5b), the absence of any Bragg diffraction peaks indicates that the clay has been completely exfoliated or delaminated in the suspension. A similar pattern was observed in case of corresponding PMMA/SPN nanocomposite but with a small remnant shoulder as shown in Figure 5b. Further studies by the same group [33] have demonstrated the effect of the nature of co-monomers on the structure 6

(a)

+

(001) 2.96nm

MMA/STN PMMA/STN

1.81nm

4

(001) (002) 2.66nm

2

(002)

(003) (003) +

Intensity / 103cps

0

(b)

4.20nm

2.02nm

3

+ MMA/SPN PMMA/SPN

4.55nm

2

of PMMA nanocomposites prepared via in-situ free-radical co-polymerization of MMA in presence of lyophilized smectite clays (each containing 10 wt%). They used three different types of co-monomers (each 1 mol%) such as N , N -dimethylaminopropyl acrylamide (PAA), N , N dimethylaminoethyl acrylate (AEA), and acrylamide (AA) for the free-radical polymerization of MMA. Figure 6 represents the results of WAXD patterns of various nanocomposites (each containing 10 wt% of SPN clay). The same behavior of the stacked is observed but the thickness of the aggregation slightly decreased compared to that of PMMA/SPN10. In case of PMMA-PAA(1)/SPN10 (see Fig. 7c), individual silicate layers connected through the edge are clearly observed in the PMMA-PAA(1) matrix and large anisotropy of the dispersed clay is observed. In contrast, the PMMA-AA(1)/SPN10 nanocomposite (see Fig. 7d) exhibited less stacking of 4–5 silicate layers with a distance of about 5 nm as a fine dispersion in the PMMAAA(1) matrix. The coherent orders of the silicate layers in this system are higher than that in other systems and are consistent with the WAXD patterns (see Fig. 6). They also prepared PMMA/SPN10 nanocomposite with a high content (3 mol%) of AA and AEA co-monomers and then tried to find out the effect of nanocomposite morphology on co-monomer amount. The length versus thickness schemes of randomly dispersed silicate layers in the nanocomposites nicely demonstrate the characteristic effects of the polar group in each co-monomer on the morphology (see Fig. 8). Incorporation of 1 mol% of AEA co-monomer possessing a dimethyl amine group appears to lead to a slight edge– edge interaction. On the other hand, the introduction of the AA co-monomer having an amide group appears to play an important role in delaminating the silicate layers. However, incorporation of 3 mol% of AEA and AA lead to the stacking of the layers compared to the corresponding 1 mol% copolymer matrix systems. We believe this behavior may be due to the formation of strong hydrogen bonding between the polar groups. For the PAA co-monomer having both polar groups, a much stronger flocculation takes place.

1 0

(c)

(c)

(b)

(d)

+ St/SPN PS/SPN

(001)

+3.65nm

20

(a)

(002)

10

0

+

3.95nm

0

2

4

6

8

10

12

2Θ / degrees Figure 5. WAXD patterns of various monomer/organoclay suspensions and corresponding polymer/clay nanocomposites. The dashed lines indicate the location of the silicate 00l reflection of organoclay from suspensions and nanocomposites. The asterisk indicates the position of 00l reflections from suspensions and nanocomposites. The arrows indicate a small shoulder or a weak peak. Reprinted with permission from [11], M. Okamoto et al., Polymer 41, 3887 (2000). © 2000, Elsevier Science.

200nm

Figure 6. Bright Field TEM images of: (a) PMMA/SPN, (b) PMMAAEA (1 mol%)/SPN, (c) PMMA-PAA (1 mol%)/SPN, and (d) PMMAAA (1 mol%)/SPN. Each contains 10 wt% SPN. Reprinted with permission from [33], M. Okamoto et al., Polymer 42, 1201 (2001). © 2001, Elsevier Science.

Polymer/Clay Nanocomposites

796

5.1.2. PS/Clay Nanocomposites

10

(a)

8

(001)

(002)

PMMA/SPN10

4.38nm

6 4 2 0 8

(b)

PMMA-AEA(1)/SPN10

(c)

PMMA-PAA(1)/SPN10

(d)

PMMA-AA(1)/SPN10

Intensity / 103cps

6 4 2 0 8 6 4 2 0 8 6

(002) +

4

(003) +

2 0

0

2

8

6

4

10

12

2Θ / degrees Figure 7. WAXD patterns of various nanocomposites. Reprinted with permission from [33], M. Okamoto et al., Polymer 42, 1201 (2001). © 2001, Elsevier Science.

500 PMMA-PAA(1)/SPN10 PMMA-AEA(3)/SPN10

Length / nm

400

PMMA-AEA(1)/SPN10

300 PMMA/SPN10

200

PMMA-AA(1)/SPN10

100 PMMA-AA(3)/SPN10

0 0

20

40

60

80

100

120 140

Thickness / nm Figure 8. Plots of length vs thickness of the dispersed clay particles in various copolymer matrices estimated from TEM images. The estimated values are located within the shaded area. Reprinted with permission from [33], M. Okamoto et al., Polymer 42, 1201 (2001). © 2001, Elsevier Science.

The same method recently was used by Zeng et al. [77] for the preparation of PS-based nanocomposites. S-monomer was first intercalated into the interlayer space of organoclay. Upon the intercalation, the complex was subsequently polymerized in the confinement environment of the interlayer space with a free-radical initiator, 2,2-azobis isobutyronitrile. Akelah and Moet [57
60] have used this in-situ intercalative polymerization technique for the preparation of PS-based nanocomposites. They modified Na+ MMT and Ca+2 -MMT with vinylbenzyltrimethyl ammonium cation by the ion exchange reaction and these modified MMTs were used for the preparation of nanocomposites. They first disperse and swell modified clays in various solvent and co-solvent mixtures such as acetonitrile, acetonitrile/toluene, and acetonitrile/THF by stirring for 1 h under N2 atmosphere. To the stirred solution S and N -N ′ -azobis (isobutyronitrile) (AIBN) were added, and polymerization of S was carried at 80
C for 5 h. The resulting composites were isolated by precipitation of the colloidal suspension in methanol, filtered off, and dried. In this way, intercalated PS/MMT nanocomposites were produced and the extent of intercalation completely depends on the nature of solvent used. Although the PS is well intercalated, a drawback of this procedure remains that the macromolecule produced is not a pure PS but rather a copolymer between S and vinylbenzyltrimethylammonium cations. For the preparation of PS based nanocomposites, Doh and Cho [63] have more commonly used MMT. They compared the ability of several tetra-alkylammonium cations incorporated in Na+ MMT through the exchange reaction to promote the intercalation of PS through the free-radical polymerization of S initiated by AIBN at 50
C. They found that the structural affinity between S monomer and the surfactant of modified MMT plays an important role in the final structure and the properties of nanocomposites. This concept has, however, been nicely achieved by Weimer et al. [67] for the preparation of PS/MMT nanocomposites. They modified Na+ -MMT by anchoring an ammonium cation bearing a nitroxide moiety known for its ability to mediate the controlled/“living” free-radical polymerization of S in bulk. The absence of WAXD peaks in the low angle area together with the TEM observations of silicate layers randomly dispersed within the PS matrix attest to the complete exfoliation of the layered silicate. PS was also the first polymer used for the preparation of nanocomposite using the melt intercalation technique with alkylammonium cation modified MMT [31]. In a typical preparative method, PS first was mixed with host organoclay powder, the mixture was pressed into a pellet, and then the pellet was heated in a vacuum at 165
C. This temperature is well above the bulk glass transition temperature of PS ensuring the presence of a polymer melt. The WAXD patterns of the hybrid before heating show peaks characteristic of the pure organoclay and during heating the organoclay peaks were progressively reduced while a set of new peaks corresponding to the PS/clay appeared. After 25 h, the hybrid shows the WAXD patterns corresponding predominantly to that of intercalated structure. The same authors also carried out the same experiment under the same experimental conditions using Na+ -MMT, but WAXD patterns did

Polymer/Clay Nanocomposites

797

not show any intercalation of PS into the silicate galleries emphasizing the importance of polymer/clay interactions. They also attempt to intercalate PS in the toluene with the same organoclay as used for melt intercalation resulting in intercalation of the solvent instead of PS. Therefore, direct melt intercalation enhances the specificity for the polymer by eliminating the competing host–solvent and polymer– solvent interaction. The schematic illustration of polymer chains intercalated in organoclay is presented in Figure 9. Syndiotactic polystyrene (s-PS)/organoclay nancomposites have also been prepared by the solution intercalation technique by mixing pure s-PS and organoclay with adsorbed cetyl pyridium chloride [75]. The WAXD analyses and TEM observations clearly established the near exfoliate structure of the prepared nanocomposites.

5.1.3. SAN/Clay Nanocomposites Kim et al. [100] used this method for the preparation of SAN/clay nanocomposites using PCL as a compatibilizer. They used a two-step mixing sequence for the preparation of SAN nanocomposites. PCL/clay master batches with different degrees of intercalation were first prepared and then were melt-mixed with SAN, where PCL is miscible with SAN. The intercalation behavior of PCL in the master batches was investigated in terms of the type of organoclay and mixing conditions. Longer mixing time and lower mixing temperature were required for the preparation of PCL master batches with exfoliated structure. As the degree of exfoliation of organoclay becomes better, the stiffness reinforcement effect of the organoclay increases in both PCL/clay master batches and their blends with SAN.

5.1.4. PVA/Clay Nanocomposites More recently, Strawhecker and Manias [87] have used this method in attempts to produce PVA/MMT nanocomposite films. PVA/MMT nanocomposite films were cast from MMT/water suspension where PVA was dissolved. Room temperature distilled water was used to form a suspension of Na+ -MMT. The suspension was first stirred for 1 h and then sonicated for 30 min. Low viscosity, fully polymer - silicon hybrid

hydrolyzed atactic PVA was then 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 and again sonicated for 30 min, and finally films were cast in a closed oven at 40
C for 2 days. The recovered cast films were then characterized by both WAXD and TEM. Both the d-spacing and their distribution decrease systematically with increasing MMT wt% in the nanocomposites. The TEM photograph of 20 wt% clay containing nanocomposite reveals the coexistence of silicate layers in the intercalated and the exfoliated states.

5.1.5. Block Copolymer/Clay Nanocomposites Krishnamoorti et al. [102
103] prepared the block copolymer-based layered silicate nanocomposites. Disordered polystyrene–polyisoprene block copolymer/layered silicate nanocomposites were prepared by solution mixing of appropriate quantities of finely ground dimethyldioctadecylammonium cation modified MMT (2C18-MMT) and an anionically synthesized monodisperse polystyrene-1,4polyisoprene (7 mol% 3
4 and 93 mol% 1,4) diblock copolymer (PSPI18) in toluene at room temperature. The homogeneous solution was dried extensively at room temperature and subsequently annealed at 100
C in a vacuum oven for ∼12 h to remove any remaining solvent and to facilitate complete polymer intercalation between the silicate layers. The WAXD patterns of PSPI18/2C18MMT or PS/2C18-MMT clearly indicate the formation of intercalated structure, whereas composite prepared with 1,4-polyisoprene showed no change in gallery height. The intercalation of PS into the silicate layers may be due to the slight Lewis base character imparted by the phenyl ring in PS, leading to the favorable interactions with the 2C18MMT layers. Further, the interlayer gallery spacing for the PSPI18/2C18-MMT composites is independent of the silicate loading. All the hybrids exhibit clear regular layered structure, demonstrated by the presence of the d001 and higher order diffraction peaks. This independence in gallery height on the silicate loading is consistent with the results obtained by Vaia and co-workers on model PS-based nanocomposite systems.

5.2. Condensation Polymers and Rubbers

polymer

silicon layer alkyl-ammonium cations

Figure 9. Schematic illustration of polymer chains intercalatedin organoclay. Reprinted with permission from [31], R. A. Vaia et al., Chem. Mater. 5, 1694 (1993). © 1993, American Chemical Society.

Several technologically important polycondensates have also been used in the nanocomposite preparation with layered silicate. These include Nylon 6 [1
2
105–125], several others polyamides [126–132], PCL [13
133–140], poly(ethylene terephtalate) (PET) [141–147], poly(butylene terephthalate) (PBT) [148], polycarbonate (PC) [149
150], PEO [20
151– 168], ethylene oxide copolymers [169], poly(ethylene imine) [170], poly(dimethyl siloxane) (PDMS) [171–176], LCP [26
177] polybenzoxazole (PBO) [178], butadiene copolymers [179–181], epoxidized natural rubber [182
183], epoxy polymer resins (EPR) [184–204], phenolic resins [205], PU [206–209], polyurethane uera (PUU) [210], polyimides [211–226], poly(amic acid) [227–229], polysulphone [230], and polyetherimide [231
232].

Polymer/Clay Nanocomposites

798 5.2.1. Nylon/Clay Nanocomposites

Clay

TCRD first reported [1] the ability of
-amino acid [COOH–(CH2 n−1 –NH+ 2 , with n = 2
3
4
5
6
8
11
12
18] modified Na+ -MMT to be swollen by the -caprolactam monomer at 100
C and subsequently to initiate its ringopening polymerization to obtain Nylon 6/MMT nanocomposites. For the intercalation of -caprolactam, they chose the ammonium cation of -amino acids because these acids catalyze ring-opening polymerization of -caprolactam. The number of carbon atoms in
-amino acids has a strong effect on swelling behavior as reported in Figure 10, indicating the extent of intercalation of -caprolactam monomer is high when the number of carbon atoms in

-amino acid is high. Figure 11 represents the conceptual view of swelling behavior of , -amino acid modified Na+ MMT by -caprolactam. In a typical synthesis, 12-aminolauric acid modified MMT (12-MMT) and -caprolactam were mixed in a motor. The content of 12-MMT ranged from 2 to 70 wt%. These were then heated at 250–270
C for 48 h to polymerize -caprolactam, using 12-MMT as catalyst. A small amount of 6-aminocaproicacid was added to the mixture to confirm the ring-opening polymerization of -caprolactam, and 12-MMT content in the mixture become less than 8 wt%. A typical procedure used was as follows: A 3 L three-necked separable flask, coupled with a mechanical stirrer, was used as the reaction vessel. In the vessel, 509 g of -caprolactam, 29.7 g of 12-MMT, and 66 g of 6-aminocaproic acids were placed under nitrogen. The mixture was heated at 100
C in an oil bath while being stirred for 30 min, followed by heating at 250
C for 6 h. The products were mechanically

18

Intensity

12

11 8

HO O

+

NH3

NH+3

O OH

Caprolactam

Clay

Figure 11. Swelling behavior of -amino acid modified MMT by -caprolactam. Reprinted with permission from [1], A. Usuki et al., J. Mater. Res. 8, 1174 (1993). © 1993, Materials Research Society.

crushed, washed with 2 L of water at 80
C for 1 h, and dried at 80
C overnight. A conceptual scheme for the synthetic method is presented in Figure 12 [105]. Further works have demonstrated that intercalative polymerization of -caprolactam could be realized without the necessity to render the MMT surface organophilic. In attempts to carry out the whole synthesis in one pot [107], the system has proved to be sensitive to the nature of the acid used to promote the intercalation of -caprolactam. Reichert et al. have also used same method for the preparation of Nylon 12/MMT nanocomposites [127]. For the preparation of Nylon 6/clay nanocomposites people generally used an in-situ intercalative polymerization technique. Liu et al. [114] first used this technique for the preparation of commercially available Nylon 6 with C18 MMT nanocomposites by using a twin-screw extruder. They prepared nanocomposites with MMT content from 1 to 18 wt%. WAXD patterns and TEM observations respectively indicated that nanocomposites prepared with MMT less than 10 wt% lead to the exfoliated structure but more than 10 wt% MMT leads to the formation of intercalated structure. WAXD and differential scanning calorimetry (DSC) analyses also showed that exfoliated structures strongly influenced the nature of the Nylon 6 crystallization, favoring the formation of -crystals in addition to the crystals of the -form observed in the native Nylon 6 matrix, and they also have a strong heterophase nucleation effect. After that VanderHart et al. [121] prepared Nylon 6/clay nanocomposites using the melt intercalation method. Recently, Fornes et al. [119] have reported the preparation of Nylon 6/clay nanocomposites using a twin-screw extruder under molten state. They used three different molecular grades of Nylon 6 for the preparation of nanocomposites with bis(hydroxyethyl)(methyl)-rapeseed quaternary ammonium [(HE)2 M1 R1 ] modified MMT and tried to

6 5 4

polymerization

3 n=2 1.0

5.0

10.0

2Θ (Co-Kα) Figure 10. XRD patterns of -amino acid [NH2 (CH2 )n−1 COOH] modified Na+ -MMT. Reprinted with permission from [1], A. Usuki et al., J. Mater. Res. 8, 1174 (1993). © 1993, Materials Research Society.

caprolactam

clay mineral a layer of clay

nylon6

Figure 12. Schematic illustration for synthesis of Nylon 6/clay nanocomposite. Reprinted with permission from [1], A. Usuki et al., J. Mater. Res. 8, 1179 (1993). © 1993, Materials Research Society.

Polymer/Clay Nanocomposites

find effects of matrix molecular weights on structure, properties, rheology, etc. Nanocomposites were prepared using a Haake, co-rotating, intermeshing twin-screw extruder, which was operated at 240
C with a screw speed of 280 rpm and a feed rate of 980 g/h. They reported the WAXD patterns of (HE)2 M1 R1 based MMT and the nanocomposites based on the three Nylon 6 matrixes of low molecular weight (LMW), medium molecular weight (MMW), and high molecular weight (HMW), having an approximate MMT concentration of 1.5 wt%. TEM images of various nanocomposites corresponding to the WAXD patterns are presented in Figure 13. WAXD patterns and TEM observations collectively revealed a mixed structure for the LMW based nanocomposites, while the MMW and HMW based nanocomposites revealed wellexfoliated structures. The average number of platelets per stack is shown to decrease with increasing matrix molecular weight. The mechanical properties of the nanocomposites were consistent with the morphological structure found via WAXD and TEM analyses. In further study [123], they examined the effect of organoclay structure on Nylon 6 nanocomposite morphology and properties. In order to understand this, a series of organic

(a) 100 nm

(b) 100 nm

(c) 100 nm Figure 13. Bright field TEM images of melt compounded nanocomposites containing ∼3 wt% MMT based on (a) HMW, (b) MMW, and (c) LMW Nylon 6. Reprinted with permission from [119], T. D. Fornes et al., Polymer 42, 9929 (2001). © 2001, Elsevier Science.

799 amine salts were ion exchanged with Na+ -MMT to form organoclays varying in amine structure or exchange level relative to the MMT. Each organoclay was melt-mixed with a HMW Nylon 6 using a twin-screw extruder operated under the same conditions as described previously; some organoclays were also mixed with LMW Nylon 6. The structure and corresponding nomenclature of various amine compounds that were used for the modification of Na+ -MMT using an ion exchanged method are presented in Figure 14. They conclude that three distinct surfactant effects were identified that lead to greater extents of exfoliation, higher stiffness, and increased yield strengths for nanocomposites based on the HMW Nylon 6: (a) one long alkyl tail on the ammonium ion rather than two, (b) methyl groups on the amine rather than 2-hydroxy-ethyl groups, and (c) an equivalent amount of amine surfactant on the clay opposed to an excess amount. Very recently, Gilman et al. [79] reported the preparation of Nylon 6- and PS-based nanocomposites of MMT modified with trialkylimidazolium cation in order to expect high stability of OMLS at high processing temperature. Figure 15 represents various kinds of imidazolium salts used for the modification of MMT. For the preparation of nanocomposites they used a miniextruder, which was operated at 10
C above the melting point of the polymer with a residence time of 3–5 min and screw speed of 200–300 rpm. WAXD analyses and TEM observations respectively established the formation of exfoliated structure in case of Nylon 6-based nanocomposite whereas there was intercalated structure with the PS/MMT system.

5.2.2. PCL/Clay Nanocomposites For the preparation of PCL-based nanocomposites, Messersmith and Giannelis [133] modified MMT with protonated aminolauric acid and dispersed the modified MMT in liquid -caprolactone before polymerizing at high temperature. The nanocomposites were prepared by mixing up to 30 wt% of the modified MMT with dried and freshly distilled -caprolactone for a couple of hours followed by ring-opening polymerization under stirring at 170
C for 48 h. The same authors [13] have also reported on the -caprolactone polymerization inside a Cr+3 -exchanged fluorohectorite at 100
C for 48 h. Pantoustier et al. [137
138] used this in-situ intercalative polymerization method. They used both pristine MMT and -amino dodecanoic acid modified MMT for the comparison of prepared nanocomposite properties. For nanocomposite synthesis, in a polymerization tube, the desired amount of pristine MMT was first dried under vacuum at 70
C for 3 h. A given amount of -caprolactone was then added under nitrogen and the reaction medium was stirred at room temperature for 1 h. A solution of initiator [Sn(Oct)2 or Bu2 Sn(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 h at room temperature. After polymerization, a reverse ion-exchange reaction was used to isolate the PCL chains from the inorganic fraction of the nanocomposite. A colloidal suspension was obtained by stirring 2 g of the nanocomposite in 30 mL of THF for 2 h at room

Polymer/Clay Nanocomposites

800

Me

(a)

Cl–

–

–

CH3 N+

+

–

– CH3

HT (C18) – N

–

–

T (C18)

Cl–

CH3

CH3

+

N

N

– CH3

Me

R

CH3

X- – CL-, BF4-

Dimethyl alkyl imidazolium salts

M3(HT)1

M3T1

propyl – CH3CH2OH Cl

CH3

CH3

(HE)2M1R1 Cl–

CH3CH2OH

N+

C+ (C12)

– CH3

(HE)2M1T1

+

(HE)2M1C1

(HSO4)+

– – N+ – CH3

–

N+

–

CH3

–

HT(C15)

–

HT (C18)

H

–

(HSO4)–

CH3

HT(C18)

M2H1(HT)1

M1H1(HT)2

(b) Effect

Organoclays M3(HT)1

vs.

M2(HT)2-95

M2H1(HT)1

vs.

M1H1(HT)2

Hydroxy-ethyl vs. Methyl

M3T1

vs.

(HE)2M1T1

MER Loading

M2(HT)2-95

vs.

M2(HT)2-125

Quaternary vs. Tertiary

M3(HT)1

vs.

M2H1(HT)1

M2(HT)2-95

vs.

M1H1(HT)2

M3T1

vs.

M3(HT)1

Number of Alkyl Tails

Saturation Length of Alkyl Tail

Figure 15. Structures of various imidazolium salts used to treat Na+ -MMT. Reprinted with permission from [79], J. W. Gilman et al., Chem. Mater. 14, 3776 (2002). © 2002, American Chemical Society.

– N+ – CH3 CH3CH2OH

CH3CH2OH

H

Cl–

–

– –

hexadecyl

M2(HT)2

–

T (C11)

decyl

HT(C18)

CH3CH2OH

CH3CH2OH

R=

– CH3

N+

–

–

–

HT (C18)

–

N

butyl

–

– –

–

R (C22)

+

Cl–

(HE)2M1C+1 (HE)2M1T1 (HE)2M1R1

Figure 14. (a) Molecular structure and nomenclature of amine salts used to organically modify Na+ -MMT by ion exchange. The symbols M (methyl), T (tallow), HT (hydrogenated tallow), HE:2 (hydroxy-ethyl), R (rapeseed), C (coco), and H (hydrogen) designate the substitutents on the nitrogen. (b) Organoclays used to evaluate the effect of structural variations of the amine cations on nanocomposite morphology and properties. Reprinted with permission from [123], T. D. Fornes et al., Polymer 43, 5915 (2002). © 2002, Elsevier Science.

temperature. Separately, a solution of 1 wt% of LiCl in THF was prepared. The nanocomposite suspension was added to 50 mL of the LiCl solution and left to stir at room temperature for 48 h. The resulting solution was centrifuged at 3000 rpm for 30 min. The supernatant was then decanted and the remaining solid washed by dispersing in 30 mL of THF followed by centrifugation. The combined supernatant was concentrated and precipitated from petroleum ether. The polymerization of CL with pristine MMT gives PCL with a molar mass of 4800 g/mol and a narrow distribution. For comparison, authors also conducted the same experiment without MMT, but there is no polymerization of CL. These results demonstrate the ability of MMT to catalyze and to control CL polymerization at least in terms of molecular weight distribution which remains remarkably narrow. For the mechanism of polymerization, authors assume that the CL is activated through the interaction with acidic site on the clay surface and the polymerization is likely to proceed via the activated monomer mechanism by the cooperative function of Lewis acidic aluminum and Bronstrated acidic silanol functionalities on the initiator walls. On the other hand, in the polymerization of CL with the protonated

-amino dodecanoic acid modified MMT, the molar mass is 7800 g/mol with a monomer conversion of 92% and again a narrow molecular weight distribution. The WAXD patterns of both nanocomposites indicate the formation of intercalated structure. In another recent publication [139
140], the same group prepared PCL/MMT nanocomposites by using in-situ ring-opening polymerization of CL using dibutyl tin dimethoxide as an initiator/catalyst.

5.2.3. PET/Clay Nanocomposites There are many reports on the preparation and characterization of PET/clay nanocomposites [141–147] using this method. Unfortunately no reports give a detailed description of the preparative method. There is one report concerning the preparation of a PET nanocomposite by in-situ polymerization of a dispersion of organoclay in water; however, characterization of the resulting composite was not reported [141]. This report claims that water serves as a dispersing aid for the intercalation of monomers into the galleries of the organoclay and discloses that a wide variety of small molecules can serve as dispersing aids in place

Polymer/Clay Nanocomposites

801

of, or in combination with, water. There is another route developed by Nanocor, Inc. [167] that was based on novel exfoliation of clays into ethylene glycol, one of the basic monomers for PET. The key to this approach includes finding a suitable surfactant that gives exfoliation of clay into ethylene glycol and also finding polymerization conditions that permit polymerization while retaining the exfoliation of the clay. Recently Imai et al. [147] reported the preparation of higher modulus PET/expandable fluorine mica nanocomposites with a novel reactive compatibilizer. Details regarding synthetic route are presented in [147]. Davis et al. [146] first reported the preparation of PETbased nanocomposites using the melt intercalation method. They used 1,2-dimethyl-3-N -alkyl imidazolium salt modified MMT (hexadecyl–MMT) for nanocomposite preparation with PET. PET/MMT nanocomposites were compounded via melt blending in a corotating mini twin-screw extruder operated at 285
C. WAXD analyses and TEM (see Fig. 16) observations respectively established that the formation of mixed delaminated/intercalated structure is achieved in the nanocomposites.

5.2.4. PBT/Clay Nanocomposites Like PET nanocomposites, this method was successfully applied by Chisholm et al. [148] for the preparation PBT/clay nanocomposites. They used sulfonated PBT for the preparation of nanocomposites. Because of the ionic nature of the −SO3 Na groups and the expected insolubility of the −SO3 Na groups in the polyester matrix, it was thought that the presence of the −SO3 Na groups may provide a thermodynamic driving force for the production of nanocomposites derived from MMT. But after preparation and characterizations of nanocomposites it was found that degree of intercalation was not strongly dependent on the amont of −SO3 Na groups; however, the mechanical properties increased significantly with increasing −SO3 Na content. (a)

500 nm

50 nm

(b)

500 nm

50 nm

Figure 16. TEM images of: (a) CD12 showing high levels of dispersion and exfoliation, average tactoids of four sheets per stack and (b) CD13 showing similar levels of dispersion and delamination as compared to CD12. Reprinted with permission from [146], C. H. Davis et al., J. Polym. Sci. B 40, 2661 (2002). © 2002, Wiley-VCH.

This behavior indicates that with high −SO3 Na content the number of interactions increases between the clay particles and the matrix via strong specific interactions involving the −SO3 Na groups.

5.2.5. PC/Clay Nanocomposites Recently, Huang et al. [149] reported the synthesis of a partially exfoliated bisphenol A PC nanocomposite using carbonate cyclic oligomers and ditallowdimethyl-exchanged MMT. WAXD patterns indicate that exfoliation of this organoclay occurs after mixing with the cyclic oligomers in a brabender mixer for 1 h at 180
C. Subsequent ring-opening polymerization of the cyclic oligomers converts the matrix into the liner polymer without the disruption of the nanocomposite structure. The TEM image revealed that a little exfoliation is obtained, although no indication of layer correlation is observed in WAXD. Mitsunaga et al. [150] also reported intercalated PC/clay nanocomposites prepared through the melt intercalation method in the presence of compatibilizer. The morphology of these nanocomposites and degradation of the PC matrix after nanocomposites preparation could be controlled by varying surfactants used for the modification of clay and compatibilizers. The intercalated PC/clay nanocomposites exhibited remarkable improvements of mechanical properties when compared with PC without clay.

5.2.6. PEO/Clay Nanocomposites In 1992, Aranda and Ruiz-Hitzky [20] first reported the preparation of PEO/MMT nanocomposites by using this method. They have carried out a series of experiments to intercalate PEO (Mw = 105 g/mol) into Na+ -MMT using different polar solvents [e.g., water, methanol, acetonitrile, mixtures (1:1) of water/methanol and methanol/acetonitrile]. In this method the nature of solvents is very crucial to facilitate the insertion of polymers between the silicate layers, polarity of the medium being a determining factor for intercalations. The high polarity of water causes swelling of Na+ -MMT provoking the cracking of the films. Methanol is not suitable as a solvent for high molecular weight PEO, whereas water/methanol mixtures appear to be useful for intercalations, although cracking of the resulting materials is frequently observed. PEO intercalated compounds derived from the homoionic M+n -MMT and M+n -hectorite can satisfactorily be obtained using anhydrous acetonitrile or methanol/acetonitrile mixture as solvents. The resulting PEO/clay nanocomposites show good stability toward treatment with different solvents (e.g., acetonitrile, methanol, ethanol, water, etc.) in experiments carried out at room temperature for long time periods (>24 h). In addition, the lack of PEO replacement by organic compounds having high affinity toward the parent silicate, such as dimethyl sulfoxide and crown ethers, indicates again the high stability of PEO-intercalated nanocomposites. On the other hand, treatment with salt solutions provokes the replacement of the interlayer cations without disruption of PEO. For example, Na+ ions in PEO/Na+ -MMT are easily + replaced by NH+ 4 or CH3 (CH2 2 NH3 ions, after treatment (2 h) at room temperature with aqueous solution of their

802 chloride, perchlorate, and thiocyanate salts (1 N solutions), in a reversible process. After that, Wu and Lerner [151] reported the intercalation of PEO in Na+ -MMT and Na+ -hectorite by using this method in acetonitrile, allowing stoichiometric incorporation of one or two polymer chains in between the silicate layers and increasing the intersheet spacing from 0.98 to 1.36 and 1.71 nm respectively. Study of the chain conformation using two-dimensional double-quantum nuclear magnetic resonance on 13 C enriched PEO intercalated in Na+ -hectorite reveals that the conformation of the “-OC– CO-” bonds of PEO is 90 ± 5% gauche, inducing constraints on the chain conformation in the interlayer [157]. Recently, Choi et al. [162] prepared PEO/MMT nanocomposites by a solvent casting method using chloroform as a co-solvent. WAXD analyses and TEM observations respectively established the intercalated structure of these nanocomposites. Various other authors [163
168] have also used same method and solvent for the preparation PEO/clay nanocomposites. Vaia et al. [152] also applied the same method to intercalate PEO in Na+ -MMT layers. Intercalation of PEO in layered silicate was accomplished by heating the PEO with the Na+ -MMT at 80
C. The WAXD patterns before any heating contain peaks characteristic of both Na+ -MMT and crystalline PEO. After heating to 80
C, the intensity of the peaks corresponding to the unintercalated silicate and crystalline PEO is progressively reduced while a set of new peaks corresponding to the PEO-intercalated MMT are observed signifying the completion of intercalation. Very recently Shen et al. [167] reported the preparation of PEO/OMLS nanocomposites using melt intercalation technique. In order to discover the effect of thermal treatment on the amount of PEO and PE–PEG diblock copolymer intercalated into the layers of Na+ -MMT and on ionic conductivity of PEO/Na+ -MMT, Liao et al. [164] prepared PEO/Na+ MMT and PE–PEG diblock copolymer/Na+ -MMT nanocomposites using a melt intercalation technique. It was found that PEO can be intercalated into the layers of Na+ -MMT by simple mechanical blending and part of PE in PE–PEG diblock copolymers was also intercalated into the layers of Na+ -MMT. The interclated amount increases with thermal treatment time, which ultimately improves the ionic conductivity of the PEO/Na+ -MMT nanocomposites. Recently, this method was successfully applied by Artzi et al. [101] for the preparation of ethylene-vinyl alcohol copolymer/clay nanocomposites.

5.2.7. PDMS/Clay Nanocomposites PDMS/clay nanocomposites were synthesized by sonicating a mixture of silanol-terminated PDMS and a commercial organoclay at room temperature for 2 min [171]. Delamination of the silicate particles in the PDMS matrix was accomplished by suspending the organosilicate in PDMS at room temperature and sonicating. WAXD analyses of various nanocomposites revealed no distinct features in the low 2 ranges indicating the formation of exfoliated nanocomposites.

Polymer/Clay Nanocomposites

5.2.8. LCP/Clay Nanocomposites Vaia and Giannelis [26] reported the reversible intercalation between organoclay and LCP in the nematic state. Melt intercalation of a model main chain liquid crystalline co-polymer based on 4,4′ -dihydroxy-a-methylstilbene and a 50:50 mole ratio mixture of heptyl/nonyl alkyl dibromide was accomplished by annealing a powder mixture of the polymer and organoclay within the nematic region of the polymer. In another report, Chang et al. [177] have prepared nanocomposites of thermotropic liquid crystalline polyester (TLCP) and Cloisite 25A using a melt intercalation method above the melt transition temperature of the TLCP. Liquid crystallinity of the nanocomposites was observed when organoclay content was up to 6 wt%.

5.2.9. PBO/Clay Nanocomposites Zhu et al. [76] used phosphonium salt for the first time for the modification of clay and then tried to find the differences between organo ammonium and phosphonium salt treatments of clay fillers in nanocomposites toward thermal stability. This technique was successfully applied by Hsu and Chang [178] in order to prepare PBO/clay nanocomposite from a PBO precursor, polyhydroxyamide (PHA), and an organoclay. The PBO precursor was made by the low temperature polycondensation reaction between isophthaloyl chloride and 2,2-bis(3-amino-4-hydroxyphenyl) hexafluoropropane with an inherent viscosity of 0.5 dl/g. For the preparation of PBO/clay nanocomposite, the organoclay was first dispersed in dimethyacetamide in which PHA was dissolved. The PHA/clay film was obtained from solution casting and dried at 80
C under vacuum. Finally PBO/clay nanocomposite was obtained by curing the film at 350
C to form a benzoxazole ring. The WAXD pattern and TEM observation respectively established that silicate layers were delaminated in PBO/clay nanocomposite film.

5.2.10. Thermoset Epoxy/Clay Nanocomposites The studies of thermoset epoxy (EPR) systems were concerned with the ring-opening polymerization of epoxides to form polyether nanocomposites. This chemistry was followed by studies of both rubbery and glassy thermoset EPR/clay nanocomposites using different types of amine curing agents. The mechanisms leading to the monolayer exfoliation of clay layers in thermoset epoxy systems have been greatly elucidated. In addition, the polymer/clay interfacial properties have been shown to play a dominant role in determining the performance benefits derived from nanolayer exfoliation. Messersmith and Giannelis [184] first reported the preparation of epoxy resin based nanocomposites of organoclay. They have analyzed the effect of different curing agents and curing conditions in the formation of nanocomposites based on the diglycidyl ether of bisphenol A (DGEBA) and a MMT modified by bis(2-hydroxyethyl) methyl hydrogenated toallow alkylammonium cation. They found that this modified clay dispersed readily in DGEBA when sonicated for a short time period as determined by the increase in viscosity at relatively low shear rates and the transition of the

Polymer/Clay Nanocomposites

803

suspension from opaque to semitransparent. The increase in viscosity was attributed to the formation of a so-called “house-of-cards” structure in which edge-to-edge and edgeto-face interactions between dispersed layers form a percolation structure. WAXD patterns of uncured clay DGEBA also indicate that intercalation occurred during and that this intercalation improves going from room temperature to 90
C. After that Wang and Pinnavaia [186] reported the preparation of PCNs of epoxy resin in the DGEBA, and the concomitant delamination of acidic forms of MMT at elevated temperature using the self-polymerization technique. They used a series of acidic cation such as H+ , NH+ 4, and acidic onium ions of the type [H3 N(CH2 n−1 COOH]+ , [H3 N(CH2 n NH2 ]+ , [H3 N(CH2 n NH3 ]+2 (n = 6 and 12) for the modification of MMT, carried out the polymerization– delamination process over the temperature range of 198–287
C, and found that the EPR-clay delamination temperature (PDT) was dependent on the heating rate and nature of the cation used for the modification of clay. In general, the PDT increased with decreasing cation acidity and basal spacing of the clay. The delamination of MMT in the polymerized epoxy resin was confirmed by X-ray powder diffraction (XRD) as shown by the powder patterns in Figure 17a and b where [H3 N(CH2 11 COOH]+ -MMT remains crystalline over the temperature range 25–229
C. However, no diffraction peaks were observed for a 5:95 (w/w) clay/polyether nanocomposite formed from [H3 N(CH2 11 COOH]+ -MMT at 229
C (see Fig. 17c). Only a very diffuse scattering characteristic of the amorphous polyether appears in the XRD pattern of the composite. The absence of a 17-Å peak for [H3 N(CH2 11 COOH]+ -MMT suggests that the clay particles have been exfoliated and the 9.6-Å-thick clay layers dispersed at the molecular level. TEM provides unambiguous evidence for the delamination of the MMT in the polyether

(a) [H3N(CH2)11 COOH]+-Mont. (25 °C)

17.0 Å

(b) [H3N(CH2)11 COOH]+-Mont. (299 °C)

Relative Intensity

(c) 5%[H3N(CH2)11 COOH]+-Mont./ Epoxy Nanocomposite

4.51 Å

3.36 Å (a) (b) (c)

0

10

20

30

40

50

Diffraction Angle (2θ) Figure 17. XRD powder patterns for a freeze-dried [H3 N(CH2 )11 COOH]+ -MMT, (b) [H3 N(CH2 )11 COOH]+ -MMT freeze-dried and then heated at 229
C, and (c) clay-polyether nanocomposite containing 5 wt% [H3 N(CH2 )11 COOH]+ -MMT. Reprinted with permission from [186], M. S. Wang and T. J. Pinnavaia, Chem. Mater. 7, 468 (1995). © 1995, American Chemical Society.

matrix. PDT values and thermodynamic data for MMTpolyether nanocomposites formed from bifunctional onium + ion MMT, onium ion NH+ 4 , and H MMT are respectively presented in Tables 3 and 4. In another study, Lan and Pinnavaia [185] have reported the preparation of PCNs with a rubber-epoxy matrix obtained from DGEBA derivatives cured with a diamine so as to reach subambient glass transition temperatures. It has been shown that depending on the alkyl chain length of modified MMT, an intercalated and partially exfoliated or a totally exfoliated nanocomposite can be obtained. Lan et al. [187] have also studied other parameters such as the nature of alkyl ammonium cations present in the gallery and the effect of the cation exchange capacity of the MMT when DGEBA was cured with m-phenylene diamine. The same kind of study was also conducted by Zilg et al. [194] who cured DGEBA with hexahydrophthalic acid anhydride in the presence of different types of clays and also modified them with a wide variety of surfactants. Recently, Kornmann et al. [201] reported the synthesis of epoxy-based nanocomposites using two different types of MMT clays with different CECs in order to investigate the influence of the CEC of the MMT clay on the synthesis and structure of nanocomposites. The CEC of any clay is a very important factor during nanocomposite synthesis because it determines the amount of surfactants which can be intercalated between the silicate layers. In this context, the swelling behavior is of critical importance to the final nanocomposite structure. For the preparation of nanocomposites both MMT clays were modified with octadecylammonium cation, and a nanocomposite synthesis procedure was the same as that used by previous authors. From WAXD patterns, it has been found that an MMT with a low CEC is exfoliated already during swelling in the epoxy resin prior to curing. A possible mechanism explaining this phenomenon is homopolymerization of the EPR during the swelling phase, causing diffusion of new epoxy molecules into the clay galleries. The large amount of space available between the layers favors the diffusion. The swelling duration of the clay with high CEC is shown to be critical for the synthesis of an exfoliated nanocomposite. Regarding the structure, TEM observations reveal interlamellar spacing of 90 Å (low CEC) and 110 Å (high CEC). However, multiplatelets of nonexfoliated layers were also observed. In summary, these nanocomposites contain clay/polymer composite particles consisting of inhomogeneously distributed silicate layer aggregates with polymer between these layers. A group of researchers [203] from Australia has recently reported the morphology, thermal relaxation, and mechanical properties of PCNs of high-functionality epoxy resins. Three different types of resins used were bifunctional DGEBA, trifunctional triglycidyl p-amino phenol (TGAP), and tetrafunctional tetraglycidyldiamino diphenylmethane (TGDDM) and all were cured with diethyltoluene diamine (DETDA). The structures of all resins and curing agents are presented in Table 5. MMT modified with octadecylammonium cation was used for the preparation of nanocomposites. The PCNs were prepared using the previously reported method. In a typical synthesis, organoclay was first dispersed in the resin at 80
C using a stirrer at 500 rpm.

Polymer/Clay Nanocomposites

804

Table 3. PDT values and thermodynamic data for polyether/clay nanocomposites formed from bifunctional onium ion-MMT.

Interlayer cation [H3 N(CH2 11 COOH]+ [H3 N(CH2 5 COOH]+ [H3 N(CH2 12 NH3 ]2+ [H3 N(CH2 6 NH3 ]2+ [H3 N(CH2 12 NH2 ]+ [H3 N(CH2 6 NH2 ]+

Initial basal spacing (Å)

PDTa (
C)

Heat of reaction (J/g)

Heat of polymb (kJ/mol)

170 ± 01 133 ± 00 134 ± 01 131 ± 01 135 ± 00 132 ± 01

229 ± 1 248 ± 1 271 ± 1 273 ± 2 281 ± 2 287 ± 2

572 ± 16 565 ± 06 566 ± 08 568 ± 07 563 ± 07 557 ± 03

228 ± 6 225 ± 2 225 ± 3 226 ± 3 224 ± 3 222 ± 2

Note: The clay:polymer composition was 5:95 (w/w). a PDT is the onset temperature for epoxide polymerization–clay delamination at a heating of 20
C/min. Heat of reaction for two epoxide equivalents. Source: Reprinted with permission from [186], M. S. Wang and T. J. Pinnavaia, Chem. Mater. 7, 468 (1995). © 1995, American Chemical Society. b

After mixing the resin–organoclay blend for 30 min the curing agent was added and the system kept under vacuum for another 60–90 min at around 70
C. The blends were then cured for 2 h at 100
C, 1 h at 130
C, 12 h at 160
C followed by a postcure for 2 h at 200
C. The morphology of the cured samples was investigated using WAXD and different microscopy techniques. Figure 18a represents the WAXD patterns of the MMT concentration series showing that the organoclay with an initial d-spacing of 2.3 nm is mainly exfoliated in the DGEBAbased system. On the other hand, high content (10 wt%) organoclay shows intercalated structure, while for DGEBAbased systems, resins of higher functionality show distinctive peaks even at low organoclay loading, indicating that these nanocomposites have a lower degree of exfoliated structure. WAXD patterns are shown in Figure 18b for TGAP and Figure 18c for TGADDM based nanocomposites of MMT. In case of any nanocomposite system, the peak observed around 2.5 nm correlates to the (002) plane and therefore represents only half the distance of the d-spacing. Figure 19a and b represents atomic force microscopy (AFM) phase contrast images of the DGEBA nanocomposite containing 5 wt% layered silicate. Individual layers cannot be seen by AFM as they usually are by the TEM. A striated structure, however, can be seen with increasing phase intervals at the top of the picture. So from the AFM images it is established that silicate layers are not homogeneously distributed in the matrix, and some stacked layers are present. Very recently, Chen et al. [204] synthesized epoxy-MMT nanocomposite using a surface initiated method in order to understand the interlayer expansion mechanism and

thermal–mechanical properties of these nanocomposites. MMT modified with bis-2-hydroxyethyl methyl tallow ammonium cation (C30B) was used as OMLS for nanocomposite synthesis. 3
4-Epoxycyclohexylmethyl-3,4epoxycyclohexane carboxylate was used as epoxy monomer, and hexahydro-4-methylphthalic anhydride (HHMPA), ethylene glycol (EG), and benzyldimethylamine (BDMA) were respectively used as curing agent, initiator, and catalyst during synthesis. In a typical preparative method, the epoxy monomer HHMPA was mixed in a molar ratio of epoxide groups to HHMPA of 1:0.87. The resulting mixture was denoted as the resin by the authors. To the resin was added a specific weight percent of the resin weight either as EG, BDMA, or C30B. All materials were then blended using an orbital mixture, until the blend become bubble free and homogeneously mixed. The blended samples were then immediately cured first isothermally for up to 8 h at temperatures ranging from 70 to 140
C, followed by 8 h at 180
C, and finally 12 h at 220
C under vacuum. The curing mechanism for an epoxy-anhydride system with an alcohol initiator is shown in Figure 20. Amine catalysts like BDMA were added to the mixture to accelerate the reaction by facilitating the ring opening of epoxy groups. Several published papers indicate that intragallery onium ions can catalyze the epoxy curing reaction and thus lead to favorable conditions for obtaining exfoliated PCNs. To understand fully in case of this system, Chen et al. [204] verify that the cross-linking reactions in the presence of C30B were due to hydroxy initiation and not due to catalytic reactions. For this reason, the extent of reaction of a resin containing C30B was compared to the extent of reaction for a neat resin and resins containing either EG

Table 4. PDT values and thermodynamic data for polyether/clay nanocomposites formed from onium ion, NH+4 , and H+ MMT. Interlayer cation [H3 N(CH2 11 CH3 ]+ [H3 N(CH2 5 CH3 ]+ NH2+ 4 H+

Initial basal spacing (Å)

PDTa (
C)

Heat of reaction (J/g)

Heat of polymb (kJ/mol)

159 ± 02 149 ± 01 125 ± 01 139 ± 01

198 ± 1 287 ± 1 247 ± 1 231 ± 1

550 ± 3 554 ± 6 554 ± 5 555 ± 12

219 ± 2 220 ± 3 220 ± 2 221 ± 5

a The PDT is the onset temperature for epoxide polymerization-clay delamination at a heating rate of 20
C/min and a composite a clay:polymer compositiom of 5:95 (w/w). b Heat of polymerization for two epoxide equivalents. Source: Reprinted with permission from [186], M. S. Wang and T. J. Pinnavaia, Chem. Mater. 7, 468 (1995). © 1995, American Chemical Society.

Polymer/Clay Nanocomposites

805

Table 5. Epoxy resins and hardener as used for nanocomposite synthesis.

10% clay 7.5% clay 5% clay 2.5% clay

d (002) =24 A

Formula CH3

O

DGEBA

CH2

CHCH2O

Intensity

Substance

(a)

O OCH2CH

C

CH2

CH3 O O

TGAP

CH2CH

CH2

CHCH2O

CH2

N CH2CH

CH2

0

2

4

6

O

8

10

12

14

2θ [°]

O O CH

CH2

CH2 N

CH2

CH

CH2

CH

CH2

(b)

CH2

CH2

CH2

CH

O

O CH3

CH3 NH2

DETDA CH2CH3

CH3CH2 NH2

10% clay 7.5% clay 5% clay 2.5% clay

d(001)=40.2 A

N

H2N

CH3CH2

NH2

d(002)=20.5 A

Intensity

TGDDM

CH2

d(002)=21.8 A

CH2CH3

Source: Reprinted with permission from [203], O. Becker et al., Polymer 43, 4365 (2002). © 2002, Elsevier Science.

0

2

4

6

8

10

12

14

2θ [°] (c)

10% clay 7.5% clay 5% clay 2.5% clay

d(002)=24.8 A

Intensity

or BDMA. As seen in Figure 21, the curing kinetics of the resin-C30B system more closely resembles the curing kinetics of the resin-EG system than the resin-BDMA system. This is a direct indication that the nanocomposites predominantly cured via initiation by the surfactant hydroxy groups and not by catalytic means. Figure 21 also shows that the curing rate of the pristine resin is significantly lower than the other mixtures. This highlights an important prerequisite for interlayer expansion, which is that extragallery polymerization rates should be slower than intragallery polymerization rates. Time-resolved high-temperature XRD was used to probe the expansion behavior of the silicate layers during curing of the PCNs, and Figure 22 represents a plot of the results of a PCN containing 15 wt% C30B held isothermally at 70
C. In Figure 23, the changes in d-spacing are plotted against the isothermal cure time for various clay loadings and cure temperature. On the basis of various characterization methods, the authors proposed an exfoliation mechanism of surface-initiated epoxy nanocomposites consisting of three stages. In the first stage, the interlayer expansion induced by intragallery polymerization must overcome any polymer chains that bridge the silicate layers. The interlayer expansion cannot proceed beyond the first stage if the number of bridging units becomes too great. The second stage was characterized by a steady and linear increase in interlayer spacing and accounts for the majority of the total expansion realized. During this stage, the silicate layers could be monitored via isothermal DSC experiments. Also, for samples that exhibited a large increase in interlayer expansion, it was found that the activation energy associated with the interlayer expansion was less than the activation energy associated with the curing. The reverse was true for samples that showed no increase in interlayer spacing. In the

d(002)=25.6 A d(002)=22.4

0

2

4

6

8

10

12

14

2θ [°] Figure 18. WAXD patterns of (a) DETDA cured DGEBA nanocomposites, (b) DETDA cured TGAP nanocomposites, and (c) DETADA cured TGDDM nanocomposites containing 0–10 wt% organoclay. Reprinted with permission from [203], O. Becker et al., Polymer 43, 4365 (2002). © 2002, Elsevier Science.

third stage, the interlayer expansion slowed then stopped, and in some cases decreased slightly. This was ascribed to the evolving modulus of the extragallery polymer such that the interlayer expansion stopped when the modulus of the extragallery polymer became equal to or exceeded the modulus of the intragallery polymer.

5.2.11. PU/Clay Nanocomposites Chen et al. [207] have used a PCL-based nanocomposite synthesis technique for the preparation of novel segmented PU/clay nanocomposites articulated on diphenylmethane diisocyanate, butanediol, and preformed polycaprolactone diol. Even if the mechanism proposed for the chemical link

Polymer/Clay Nanocomposites

806 (b)

(a)

25 wt % BDMA

600

400

200

Extent of reaction, p

7 wt % EG 5 wt % C30B

0

200

400

600

Prisline rosin

0 0

3.31 µm

0

200

400

600 nm

Figure 19. Phase contrast AFM images of DETDA cured DGEBA containing 5 wt% organoclay. Reprinted with permission from [203], O. Becker et al., Polymer 43, 4365 (2002). © 2002, Elsevier Science.

between the nanofiller surface and the polymer does not appear appropriate, they succeeded in producing a material where the nanofiller acts as a multifunctional chain extender inducing the formation of star-shaped segmented polyurethane. Recently, Wang and Pinnavaia [206] reported the preparation of polyurethane-MMT nanocomposites using a direct in-situ intercalative polymerization technique. More recently, Yao et al. [209] reported the preparation of a novel kind of PU/MMT nanocomposite using a mixture of modified 4
4′ -di-phenymethylate diisocyanate (MMDI), modified polyether polyol (MPP), and Na+ -MMT. In a typical synthetic route a known amount of Na+ -MMT was first mixed with 100 mL of MPP and then stirred at 50
C for 72 h. Then the mixture of MPP and Na+ -MMT was blended with a known amount of M-MDI and stirred for 30 s at 20
C, and finally curing was conducted at 78
C for 168 h. Biomedical PUU/MMT (modified with dimethyl ditallow ammonium cation) nanocomposites were prepared by adding organoclay suspension in toluene dropwise to the solution of PUU in N , N -dimethylacetamide (DMAC) [210]. The mixture was then stirred overnight at room temperature. The solution was degassed, and then films were cast on to round glass Petri dishes. The films were airdried for 24 h and subsequently dried under vacuum at 50
C for 24 h. WAXD analyses indicated the formation

0

50

100

150

200

Isothermal time, min Figure 21. Extent of reaction of a neat resin and resins containing either C30B, EG, or BDMA. The inset shows the same data but for longer reaction. Reprinted with permission from [204], J. S. Chen et al., Polymer 43, 4895 (2002). © 2002, Elsevier Science.

of intercalated nanocomposites but they did not report any TEM photographs.

5.2.12. Polyimide/Clay Nanocomposites Yano et al. [209] have conducted the preparation of polyimide/MMT nanocomposites from a DMAC solution of poly(amic acid) and a DMAC dispersion of MMT modified with dodecylammonium cation. Table 6 shows the dispersibility of various kinds organically modified MMTs in DMAC and the average diameter of organophilic MMTs obtained from the dynamic light scattering experiment. In case of 12CH3 -MMT, MMT appeared to disperse in DMAC homogeneously and the average diameter of the dispersed MMT particles was the smallest of all. Another interesting aspect is that as the carbon number of the surfactant increases, the hydrophilicity of the ogranophilic MMT decreases. Table 6 also indicates that 10–12 carbon atoms are appropriate for organophilic MMT to be dispersed in DMAC. In another report [226] polyimide/MMT nanocomposites were prepared using a solvent cast method from solution of poly(amic acid) precursors and the dodecyl-MMT using

50

e

tim al rm he min

ot

Is

200 150 100 0

1

2

3

4

5

6

2-theta Figure 20. Schematic illustration of generalized curing reaction involving the epoxy monomer, HHMPA, EG, and BDMA. Reprinted with permission from [204], J. S. Chen et al., Polymer 43, 4895 (2002). © 2002, Elsevier Science.

Figure 22. TT-XRD of a resin containing 15 wt% C30B held isothermally at 70
C. Reprinted with permission from [204], J. S. Chen et al., Polymer 43, 4895 (2002). © 2002, Elsevier Science.

Polymer/Clay Nanocomposites

d-spacing, angstroms

(a)

807

75

Table 6. Dispersibility and average diameter of organically modified MMT in DMAC.

65

Dispersibility Average of organophilic diametera MMT in DMAC (m)

Intercalated salts

55 120 C 100 C 80 C 70 C

45

35 0

100

200

300

400

500

Isothermal time, min

d-spacing, angstroms

(b)

75

65

55 100 C 80 C 70 C

45

35 0

100

200

300

400

500

Isothermal time, min

d-spacing, angstroms

(c)

75

65

n-Octyltrimethylammonium chloride Ammonium salt of dodecylammine (12CH3 -MMT) Ammonium salt of 12-aminododecanoic acid (12COOH-MMT) n-Decyltrimethylammonium chloride (C10A-MMT) n-Dodecyltrimethylammonium chloride n-Hexadecyltrimethylammonium chloride n-Dioctadecyldimethylammonium chloride n-Trioctylmethylammonium chloride n-Benzyltrimethylammonium chloride

not dispersible dispersible

— 0.44

partly dispersible

3.75

,,

0.61

not dispersible ,,

— —

,,

—

,, ,,

— —

a Values of average diameter are much bigger than 200 nm, because an average diameter from light scattering measurement includes solvent around a substance.

perature Tg , thermal decomposition temperature, a drastic decrease in solvent uptake, etc., compared to the virgin PEI but they did not check the stability of intercalated salt in organoclay at this high temperature of mixing.

5.3. Polyolefins

55 120 C 100 C 80 C 70 C

45

35 0

100

200

300

400

500

Isothermal time, min Figure 23. Changes in d001 as a function of the curing time and temperature: (a) 5, (b) 10, and (c) 15 wt% silicate loading. The dashed lines denote the quatitative detection limit of the TT-XRD setup. Reprinted with permission from [204], J. S. Chen et al., Polymer 43, 4895 (2002). © 2002, Elsevier Science.

N -methyl-2-pyrrolidone as a solvent. The cured films of the rigid-rod polyimide/MMT nanocomposites as characterized by FTIR, TEM, and WAXD were exfoliated nanocomposites at low MMT content and partially exfoliated structure at high MMT content.

5.2.13. Polyetherimide/Clay Nanocomposites PEI/MMT nanocomposites were prepared by melt blending of hexadecylamine modified MMT and PEI at 350
C [231
232]. WAXD patterns of various nanocomposites show no peaks but TEM observations show relatively stacked silicate layers not homogeneously dispersed in the polymer matrix. According to the authors, due to the strong interaction between PEI and organoclay, the nanocomposites exhibited a substantial increase at glass transition tem-

PP [9
14
15
25
30
233–266], PE [267–273], polyethylene oligomers [274], copolymers like poly(ethylene-co-vinyl acetate) (EVA) [275], and ethylene propylene diene methylene linkage rubber (EPDM) [276, 277] have also been used.

5.3.1. PP/Clay Nanocomposites Tudor et al. [235] first used this in-situ intercalative polymerization method for the preparation of PP/clay nanocomposites. They have demonstrated the ability of soluble metallocene catalyst to intercalate inside silicate layers and to promote the coordination polymerization of propylene. Accordingly, a synthetic hectorite (Laponite RD) was first treated with methylaluminoxane (MAO) in order to remove all the acidic protons and to prepare the interlayer spacing to receive the transition metal catalyst. It has to be noted that MAO is commonly used in association with metallocenes to produce coordination catalysts active in olefin polymerizations. During this first treatment step, WAXD analysis showed no noticeable increase of the layer spacing, although the diffraction peak broadened slightly, but the increase in Al content and complete disappearance of Si–OH signals from infrared spectra agree with the MAO reaction/adsorption inside the layered silicate galleries. Upon the addition of the metallocene catalyst ([Zr(-C5 H5 Me(THF)]+ ), a cation exchange reaction occurs between Na+ in MAO treated hectorite and the metallocene catalyst as demonstrated by an increase in the interlayer spacing of 0.47 nm, consistent with the size of the species. Details can be seen in Figure 24. Using a synthetic fluorinated mica-type layered silicate that is deprived from any protons in the galleries, the catalyst was even

Polymer/Clay Nanocomposites

808

9.6Å OH Groups

i

9.6Å

= Na+

Methylaluminoxane [MeAl–(µ–O)]n

ii iii

Zr

14.4

Cl Me

= zr(η-C5H5)2Me+ iv

=

Stearyl ammonium n

Figure 24. Schematic illustration of the modification and ion exchange of Laponite with [Zr(-C5 H5 )2 Me(thf)]BPh4 and propene polymerization. Details regarding the reagent and conditions are shown in relevant reference. Reprinted with permission from [235], J. Tudor et al., Chem. Commun. 2031 (1996). © 1996, Royal Chemical Society.

intercalated Directly within the silicate layers without the need for MAO treatment. These two modified layered silicates catalyzed with reasonably high activity with the polymerization of propylene when contacted with an excess of MAO, producing PP oligomers. Unfortunately, the authors did not report any characterization of these composites, so one cannot make claims about the morphology. In another recent publication, Sun and Garces [259] have reported the preparation of PP/clay nanocomposites by in-situ intercalative polymerization with metallocene/clay catalysts. Usuki et al. [239] first reported a novel approach to prepare PP/clay nanocomposites using functional oligomer with polar telechelic OH groups (PP-OH) as compatibilizer. In this approach, first PP-OH intercalates between the layers of 2C18 -MMT, the second step was the melt mixing of PP-OH/2C18 -MMT with PP, and the nanocomposite with intercalated structure was obtained. Further study by the same group [14] reported the preparation of PP/MMT nanocomposites by melt blending of PP, maleic anhydride grafted PP oligomer (PP-MA), and clays modified with stearylammonium using a twin-screw extruder. In their study, they used two different types of maleic anhydride modified PP oligomer with different amounts of maleic anhydride groups and two types of organically modified clays to understand the miscibility effect of the oligomers on the dispersibility of the organoclay in PP matrix, and to study the effect of hybridization on their mechanical properties when compared with neat PP and PP/clay nanocomposites without oligomers. WAXD analyses and TEM observations respectively established the intercalated structure for all nanocomposites. On the basis of WAXD patterns and TEM images, they proposed a possible mechanism for dispersion of intercalated clay layers in the PP matrix. Figure 25 shows a schematic presentation of the mixing process of the three components (i.e., PP, PP-MA, and OMLS) into the nanocomposites. Authors believe that the driving force of the intercalation originates from the maleic anhydride group and the oxygen groups of the silicate through hydrogen bond. In a recent report, Hasegawa et al. [240] found that PP-MA was able to intercalate into the intergalleries of OMLS, as like the functional oligomer, and they described a facile approach for the preparation of PP/clay nanocomposite using a PP-MA and organically modified clay by a melt intercalation technique. In a typical preparative

Silicate Layer of Clay

PP-MA Oligomer Maleic Anhydride Group

PP

Figure 25. Schematic representation of the dispersion process of the organoclay in the PP matrix with the aid of PP-Mas. Reprinted with permission from [239], A. Usuki et al., J. Appl. Polym. Sci. 63, 137 (1997). © 1997, Wiley-VCH.

method, PP/clay nanocomposite pellets were prepared by a melt blending of pellets of PP-MA and the powder of C18 -MMT at 200
C using a twin-screw extruder. The same authors [245] also prepared polyethylenepropylene rubber (EPR)/C18 -MMT nanocomposites by melt blending of EPR-g-MA and C18 -MMT powder using a twinscrew extruder, and then tried to compare the morphology of EPR based nanocomposites with PP/clay nanocomposites. Recently, Nam et al. [30] prepared PP/MMT nanocomposites using the same method as used by previous authors. For example a mixture of PP-MA (0.2 wt% MA) and C18 -MMT was melt extruded at 200
C in a twinscrew extruder. They prepared PCNs with three different amounts of clay content (inorganic part) of 2
4, and 7.5 wt%, which were correspondingly abbreviated as PPCN2, PPCN4, and PPCN7.5 respectively. The WAXD patterns of C18 -MMT, PP-MA, and various PPCNs are presented in Figure 26. WAXD patterns clearly established near to exfoliate structure is formed with PPCN2, disordered intercalated nanocomposite is formed in case of PPCN4, while PPCN7.5 represents an ordered intercalated structure. These features are clearly observed with bright field

Polymer/Clay Nanocomposites

809

3000 Organophilic clay

(001)

2000

1000 0 PPCN 2 2000

*

1000

Intensity / A.U.

0 PPCN 4 3000 2000

*

1000 0

PPCN 7.5 3000

* 2000 1000 0 PP-MA 2000 1000 0 0

2

4

6

8

10

2Θ / degrees

Figure 26. WAXD patterns for organoclay, PP-MA, and various PPCNs. The dashed lines indicate the location of the silicate (001) reflection of organoclay. The asterisks indicate a remanant shoulder for PPCN2 or a small peak for PPCN4. Reprinted with permission from [30], P. H. Nam et al., Polymer 42, 9633 (2001). © 2001, Elsevier Science.

TEM images of various PPCNs as shown in Figure 27. In order to understand the hierarchical structure of PPCNs they also used polarizing optical microscopy, light scattering, and small angle X-ray diffraction for the characterization of PPCNs along with WAXD and TEM. On the basis of these analyses they demonstrated the dispersed clay and interfibril structure of PPCNs. The schematic illustration of this structure is presented in Figure 28. In another publication [253], Liu and Wu have reported the preparation of PPCN via grafting-melt compounding using a new type of co-intercalated organophilic clay which has a larger interlayer spacing than the ordinarily orgonophilic clay that is only modified with alkyl ammonium cation. One of the co-intercalation monomers is unsaturated so it could tether on the PP backbone by virtue of grafting reaction. The co-intercalated organophilic clay (EM-MMT) was prepared as follows: 130 g hexadecylammonium modified MMT (C16 -MMT) and 20 g epoxypropyl methacrylate were mixed in a Hake Reocorder 40 mixer for 1 h. Before mixing with clay, the initiator of grafting reaction, dibenzoyl peroxide, and donor agent were disolved in epoxypropyl methacrylate. The nanocomposites were prepared using a twin-screw extruder with a screw speed of 180 rpm and operated at temperature around 200
C. WAXD patterns

Figure 27. Bright field TEM images of PPCNs: (a) 2, (b) 4, and (c) 7.5 wt% MMT. The dark lines are the cross sections of silicate layers and the bright areas are the PP-MA matrix. Reprinted with permission from [30], P. H. Nam et al., Polymer 42, 9633 (2001). © 2001, Elsevier Science.

and TEM observations established that the larger interlayer spacing and strong interaction caused by grafting can improve the dispersion effect of silicate layers in the PP matrix. Recently, Manias et al. [254] reported the preparation of PP/organically modified MMT, having a coexisting intercalated and exfoliated structure by the melt intercalation technique. They prepared nanocomposites two different ways: (a) by introducing functional groups in PP and using common alkylammounium MMT and (b) by using neat/unmodified PP and a semifluorinated surfactant modification for the MMT. Kaempfer et al. [264] reported the preparation of new PCNs via melt compounding of syndiotactic polypropylene (sPP) containing organoclay and in-situ formed core/shell nanoparticles. Melt compounding of sPP with organohectorite, obtained via cation exchange of fluohectorite with octadecylammonium cation, in a co-rotating twin-screw extruder represents an attractive new route to reinforced sPP with considerably higher stiffness. The matrix reinforcement is achieved by in-situ formation of silicate nanoparticles via exfoliation combined with simultaneous in-situ encapsulation of the resulting nanosilicates in a thin shell of iPP-graft-MA. The resulting anisotropic core/shell type nanoparticles, containing stacks of organohectorite layer as

Polymer/Clay Nanocomposites

810 (a) PPCN2 clay particle crosshatch lamellae Lclay

dlamellae

ξ clay dclay

Llamellae

Lclay ~ 193 dclay ~ 5.2 ξ clay ~ 62 d(001) ~ 3.24 Llamellae ~ 14.8 dlamellae ~ 7.2

d(001) individual silicate layer

fibril (assembly of lamellae)

The extent of exfoliation and intercalation completely depends on the hydrophilicity of polyethylene grafted with MA and the chain length of the organic modifier in the clay. When the number of methylene groups in alkylamine (organic modifier) was larger than 16, the exfoliated nanocomposite was obtained, and the maleic anhydride grafting level was higher than about 0.1 wt% for the exfoliated nanocomposite with the clay modified with dimethyl dehydrogenated tallow ammonium cation or octadecylammonium cation. Very recently, EPDM/clay nanocomposites have been prepared by mixing EPDM with OMLS via a vulcanization process [276]. They used thiuram and dithiocarbamate for the vulcanization accelerator.

5.4. Specialty Polymers (b) PPCN7.5 clay particle crosshatch lamellae dlamellae

Lclay

ξ

Llamellae clay

dclay

d(001)

fibril (assembly of lamellae) individual silicate layer

Lclay ~ 127 dclay ~ 10.2 ξ clay ~ 35 d(001) ~ 2.89 Llamellae ~ 15.0 dlamellae ~ 7.4

Figure 28. The schematic illustration for dispersed clay structure and the interfibriller structure for PPCNs with: (a) 2 and (b) 7.5 wt% MMT. Reprinted with permission from [30], P. H. Nam et al., Polymer 42, 9633 (2001). © 2001, Elsevier Science.

core and iPP-graft-MA as shell, represent a very effective new class of nucleating agents for sPP crystallization.

5.3.2. PE/Clay Nanocomposites PE/clay nanocomposites have also been prepared by the in-situ intercalative polymerization of ethylene using so-called polymerization filling technique [272]. Pristine MMT and hectorite were first treated with trimethylaluminum-depleted methylaluminoxane before being contacted by a Ti-based constrained geometry catalyst. The nanocomposite was formed by addition and polymerization of ethylene. In the absence of a chain transfer agent, ultrahigh molecular weight polyethylene was produced. The tensile properties of these nanocomposites were poor and essentially independent of the nature and content of the silicate. Upon hydrogen addition, the molecular weight of the polyethylene was decreased with parallel improvement of mechanical properties. The formation of exfoliated PCNs was established using WAXD and TEM analyses. In another report, Heinemann et al. [268] prepared (co)polyolefin/clay nanocomposites using this method. MA grafted polyethylene (PE-MA)/clay nanocomposites were also prepared by a melt intercalation technique [271].

In addition to the aforementioned conventional polymers, several interesting developments have also taken place in the preparation of nanocomposites of layered silicates with some speciality polymers including the N hetrocyclic polymers like polypyrrole (PPY) [278–283], poly(N -vinylcarbazole) (PNVC) [284
285], and polyaromatics such as polyaniline (PANI) [286–300], poly(p-phenylene vinylene) [301], and related polymers [302]. PPY and PANI are known to display electric conductivity [303], and PNVC is well known for its high thermal stability and characteristic optoelectronic properties [304–307]. Some research has also been initiated with liquid crystalline polymer based nanocomposites [308–312] and hyperbranch polymers (HBP) [313].

5.4.1. PNVC/Clay Nanocomposites Biswas and Sinha Ray [284] first reported the preparation of PNVC/MMT nanocomposites by direct polymerization of N -vinylcarbazole (NVC) (solid or in solution) in the presence of MMT without the use of any free-radical initiator. Melt polymerization of NVC in MMT (at 70
C) as well as solution (in benzene) polymerization of NVC in the presence of MMT at 50
C resulted in the formation of PNVC/MMT nanocomposite with intercalated structure. After repeated benzene extraction of prepared nanocomposites, intercalated PNVC could not be removed, while all the surface-adsorbed PNVC was extracted with benzene. WAXD analyses confirmed intercalation of PNVC in MMT interlayer galleries. According to the authors, the initiation in the NVC/MMT system was suggested to be cationic involving Brønstrated acid sites in MMT arising from the dissociation of interlayer water molecules coordinated to the exchangeable cations [284]. Yet another possibility, especially with NVC, was that the transition metal oxides (Fe2 O3 /TiO2 ) present in MMT could also lead to cationic initiation of NVC. The same authors subsequently reported that direct interaction of MMT with pyrrole led only to ca. 5% yield of PPY in 3 h while ANI could not be polymerized by MMT [281]. According to them, such a trend is possibly not surprising since NVC is relatively more susceptible to cationic polymerization compared to the latter monomers. Further work by Sinha Ray and Biswas, in the study of a NVC/MMT polymerization/nanocomposite formation

Polymer/Clay Nanocomposites

system, the addition of FeCl3 was considered to be interesting. Results of a recent study [285] indicated that in a NVC-MMT polymerization/nanocomposite formation system addition of FeCl3 increased the percentage loading of PNVC in the composite.

5.4.2. PANI/Clay and PPY/Clay Nanocomposites Polymerization vis-a-vis nanocomposite formation in PY/MMT-water and ANI/MMT-water systems was possible after using FeCl3 and (NH4 2 S2 O8 [281
287] as oxidant respectively in the two systems. Kim et al. [286] first used this method for the preparation of PANI/MMT nanocomposite using docecylbenzenesulfhonic acid (DBSA) and camphorsulphonic acid (CSA) as dopant. In a typical synthetic method, Na+ -MMT was first dispersed in an aqueous medium and then sonicated by using an ultrasonic generator. The DBSA or CSA dopants were dissolved in distilled water and mixed with ANI monomer solution at a 1:1 molar ratio; then the emulsion solutions were mixed in a four-neck reactor by stirring while the temperature was kept at 25
C. The oxidant initiator, (NH4 2 S2 O8 solution, was dropped into the reactor. The WAXD analysis clearly indicates the formation of intercalated nanocomposites. In another recent publication, Kim et al. [283] reported the preparation of intercalated PPY/Na+ -MMT nanocomposite via an inverted emulsion pathway method.

5.4.3. HBP/Clay Nanocomposites Very recently, Plummer et al. [313] have used this method for the preparation of HBP/MMT nanocomposites. The chemical structure of the HBP is presented in Figure 29. The

Figure 29. Chemical structure of the dendrimer analog of the second pseudo-generation HBP. Reprinted with permission from [313], C. J. G. Plummer et al., Chem. Mater. 14, 486 (2002). © 2002, American Chemical Society.

811 PCNs were prepared by introducing the required amount of Na+ -MMT to 10 g of HBP dispersed in 75 mL of boiling deionized water. The mixture was stirred in air at 50
C with a magnetic stir bar. After evaporation of half the water, the resulting gel was transferred to an open silicone rubber mold and dried in air for 2 days at 50
C. The remaining solid was then dried for another 2 days at 120
C under vacuum, ground, and pressed into 25-mm diameter, 1-mm thick disks at 60
C for WAXD analyses. At high Na+ MMT contents, WAXD analyses indicated 2.5–2.8, 2.8–3, and 3.6–3.9 nm silicate layer basal spacing for the second, third, and fourth pseudo-generation HBPs, respectively, as opposed to 1.06 nm for the as-received Na+ -MMT. The corresponding WAXD peaks disappeared as the clay content was reduced to below 20 wt% for all the HBPs, consistent with the previous study by Strawhecker and Manias [87]. TEM images of nanocomposite containing 20 wt% Na+ MMT revealed stacks of 5–10 silicate layers with a relatively well-defined spacing, interspersed with exfoliated silicate layers. At 10 wt% MMT, however, exfoliation was confirmed to dominate.

5.5. Biodegradable Polymers Recently, some groups have started the preparation, characterization, and materials properties of various kinds of biodegradable polymers/nanocomposites having properties suitable for a widerange of applications. So far reported biodegradable polymers for the preparation of nanocomposites are polylactide (PLA) [34
314–324], poly(butylene succinate) (PBS) [325–328], PCL [13
133– 140], unsaturated polyester [329], polyhydroxy butyrate [330], and aliphatic polyester [331–334].

5.5.1. PLA/Clay Nanocomposites Sinha Ray et al. [34
315] first reported the preparation of intercalated PLA/layered silicate nanocomposites. For nanocomposite (PLACNs) preparation C18 -MMT and PLA were first dry 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 color strands of PLACNs. Nanocomposites loaded with a very small amount of oligo-PCL as a compatibilizer were also prepared in order to understand the effect of oligo-PCL on the morphology and properties of PLACNs [34]. The compositions of various nanocomposites of PLA with C18 -MMT are summarized in Table 7. WAXD patterns of a series of nanocomposites are shown in Figure 30. Figure 31 represents the TEM photographs of nanocomposites corresponding to the WAXD patterns. On the basis of WAXD analyses and TEM observation, they calculated form factors (see Table 8), that is, average length (Lclay ), thickness (dclay ) of the stacked intercalated silicate layers, and the correlation length (lay  between them (see Fig. 3). These data clearly established that silicate layers of the clay were intercalated and randomly distributed in PLA matrix. Incorporation of very small amount of oligo-PCL as a compatibilizer in the nanocomposites lead to a better parallel stacking of the silicate layers and also much stronger flocculation due to the hydroxylated edge–edge interaction of the silicate layers. Owing to the interaction between clay platelets and

Polymer/Clay Nanocomposites

812

Table 7. Composition and characteristic parameters of various PLACNs based on PLA, oligo-PCL, and C18 -MMT. Composition (wt%) Sample

PLA

PLACN1 PLACN2 PLACN3 PLACN4 PLACN5 PLACN6 PLACN7 PLAa PLA1 PLA2 PLA3

97 95 93 948 945 93 92 100 998 995 98

b

Oligo-PCL

C18 -MMT

02 05 20 30

3 20 5 30 7 48 5 33 5 33 5 28 5 24

02 05 20

Mw × 10−3 (g/mol)

Mw /Mn

Tg (
C)

Tm (
C)

cc (%)

178 185 177 181 181 180 181 187 180 180 180

1.81 1.86 1.69 1.76 1.76 1.76 1.77 1.76 1.76 1.76 1.76

60.0 60.0 59.8 58.6 57.6 54.0 51.0 60.0 58.0 57.0 54.7

169 170 170 170 169 168 168 168 1685 1688 169

5065 3901 4366 4147 3291 — — 3624 4621 5251 —

a

Mw and PDI of extruded PLA (at 190
C) are 180 × 103 (g/mol) and 1.6 respectively. Values in brackets indicate the amount of clay (inorganic part) content after burning. c The degree of crystallinity. b

the PLA matrix in the presence of a very small amount of oligo-PCL, the disk–disk interaction plays an important role in determining the stability of the clay particles and hence the enhancement of mechanical properties of such nanocomposites. In their further research [316–318
324], they prepared a series of PLACNs with various types of organoclay in order to investigate the effect of organoclay on the morphology, properties, and biodegradability of PLACNs. Four different types of pristine layered silicates were used and each of them was modified with a different type of surfactant. Detailed specifications of various types of organoclay was used by them are presented in Table 9. On the basis of WAXD analyses and TEM observations, the authors concluded the formation of four different types of PLACNs. Ordered intercalated-and-flocculated nanocomposites were obtained when ODA was used as organoclay, disordered intercalated structure was found in case of PLA/SBE4 nanocomposite, PLA/SAP4 nanocomposite shows near to exfoliate nanocomposite, while the coexistence of stacked intercalated and exfoliated nanocomposite structure is found with PLA/MEE4 nanocomposite. So the nature of OMLS has a strong effect on the final morphology of PLA-based nanocomposites. In a very recent work, Maiti et al. [319] have prepared a series of PLACNs with three different types of pristine layered silicate such as saponite, MMT, and mica, and each of them was modified with alkylphosphonium salts having various chain lengths. In their work they first try to find the effect of alkylphosphonium modifier of different chain lengths on the properties of organoclay and how the different clays behave differently having same organic modifier. Second, they study the effects of dispersion, intercalation, and aspect ratio of clay on materials properties. From the WAXD patterns it is clearly observed that the d-spacing (001) increases with increasing modifier chain length and for a fixed modifier it increases with increasing lateral dimension of the clay particle. We believe there are two reasons to observe this type of behavior: one is the CEC value and the other is the lateral size of various pristine layered silicates, and in both cases layered silicates follow the order

mica > MMT > saponite. According to the authors, out of two factors, the former factor is more important in controlling the d-spacing/stacking of silicate layers than that of the latter. Since mica has a high lateral size and also a high amount of surfactant molecules due to its high CEC value, surfactant chains inside the intergallery have restricted conformation due to physical jamming. According to them this physical jamming is smaller in case of saponite due to its lower CEC and smaller in lateral size. Therefore, the situation for OMLS, based on TEM and WAXD analyses, is schematically illustrated in Figure 32. Figure 33 compares the WAXD patterns of nanocomposites with different clay dimensions having the same clay [n-hexadecyl tri-n-butyl phosphonium bromide (C16 modified] content (3 wt%). For MMT-based nanocomposite, the peaks are sharp and crystallite sizes are slightly less than those of the corresponding organoclay, indicating an almost ordered structure of MMT in nanocomposite. The peaks of the nanocomposites prepared with mica clay are very sharp, similar to those of corresponding organoclay, and slightly larger crystallite sizes indicate that the number of stacked silicate layers is the same to that of original organoclay but 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 with three different organoclays is presented schematically in Figure 34. More recently Dubois et al. [320
322] reported the preparation of plasticized PLA/MMT nanocomposites. The OMLS used was MMT modified with bis-(2-hydroxyethyl)methyl (hydrogenated tallowalkyl) ammonium cations. WAXD analyses have confirmed the formation of intercalated nanocomposites.

5.5.2. PBS/Clay Nanocomposites Like PLA, PBS is also an aliphatic thermoplastic polyester with many interesting properties, including biodegradability, melt processability, and thermal and chemical resistance. Although these properties show the potential applications of PBS, some of the other properties such as softness, gas

Polymer/Clay Nanocomposites

813 (b)

(a) (a) C18-mmt

(001)

2000 1500 1000

2µm

2µm

(c)

500

(d)

0

PLACN1

2000 1500

Intensity /A.U.

1000

*

500

500nm

Figure 31. TEM bright field images: (a) PLACN2 (×10000), (b) PLACN4 (×10000), (c) PLACN2 (×400000), and (d) PLACN4 (×40000). The dark entities are the cross section of intercalated organoclay, and the bright areas are the matrices. Reprinted with permission from [34], S. Sinha Ray et al., Macromolecules 35, 3104 (2002). © 2002, American Chemical Society.

0

PLACN2

2000 1500 *

1000

500nm

500 0

barrier properties, flexural properties, etc. are frequently not enough for a wide range of applications. Sinha Ray et al. [325
326] first reported the preparation of PBS/MMT nanocomposites (PBSCNs) by simple melt extrusion of PBS and OMLS, having properties suitable for a wide range of applications. MMT modified with octadecylammonium chloride was used as organoclay for nanocomposite preparation. In recent publications [327
328], the same authors also reported the details of structure–property relationships in case of PBSCNs. Recently, Lee et al. [332] have reported the preparation of biodegradable aliphatic polyester (APES)/organoclay nanocomposites using a melt intercalation method. Two kinds of organoclays, Cloisite 30B and Cloisite 10A with different ammonium cations located in the silicate galleries, were chosen for the nanocompoites preparation. The WAXD analyses and TEM observations respectively showed a higher degree of intercalation in case of APES/Cloisite 30B nanocomposites as compared to that of APES/Cloisite 10A nanocomposites. According to the authors, this behavior may be due the hydrogen-bonded interaction between APES and hydroxyl group in the galleries of Cloisite 30B nanocomposites than for the APES/Cloisite nanocomposites.

PLACN3

2000 1500

*

1000 500 0 2

0

4

6

8

10

2Θ/degrees (b) PLACN2

2000 1500 *

1000 500

Intensity /A.U.

0

PLACN4

2000 1500

*

1000 500 0

PLACN5

2000 1500

*

1000

Table 8. Comparison of form factors between PLACN2 and PLACN4 obtained from WAXD patterns and TEM observations.

500 0

0

2

4

6

8

10

2Θ/degrees Figure 30. 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 asterisks indicate the (001) peak for C18 -MMT dispersed in PLA matrices. Reprinted with permission from [34], S. Sinha Ray et al., Macromolecules 35, 3104 (2002). © 2002, American Chemical Society.

Form factors WAXD d001 (nm) dclay (nm) TEM dclay (nm) Lclay (nm) Lclay /dclay clay (nm)

PLACN2

PLACN4

3.03 13

2.98 10

38 ± 1725 448 ± 200 12 255 ± 137

30 ± 125 659 ± 145 22 206 ± 92

Polymer/Clay Nanocomposites

814 Table 9. Specifications and designation of organoclay used for the preparation of PLACNs. Clay codes

Pristine clay

Particle length (nm)

ODA SBE MEE

MMT MMT synthetic F-mica

150–200 100–130 200–300

SAP

Saponite

CEC (mequiv/100 gm) 110 90 120

50–60

Organic salt used for the modification of clay octadecylammonium cation trimethyloctadecylammonium cation dipolyoxyethylene alkyl(coco) methylammonium cation tributylhexadecylphosphonium cation

866

6. MATERIALS PROPERTIES OF POLYMER/CLAY NANOCOMPOSITES PCNs consisting of a polymer and clay (modified or not) frequently exhibit remarkably improved mechanical and materials properties as compared to those of pristine polymers containing a small amount of (≤5 wt%) layered silicate. Improvements can include high moduli, increased strength and heat resistance, decreased gas permeability and flammability, and increased biodegradability of biodegradable polymers. The main reason for these improved properties in PCNs is interfacial interaction between matrix and layered silicate as opposed to conventional filler reinforced systems.

6.1. Mechanical Properties 6.1.1. Dynamic Mechanical Analysis Dynamic mechanical analysis (DMA) measures the response of a given material to a cyclic deformation (here in tension– torsion mode) as a function of temperature. DMA results are expressed by three main parameters: (a) the storage modulus (G′ ) corresponding to the elastic response to the deformation; (b) the loss modulus (G′′ ) corresponding to the plastic response to the deformation, and (c) tan , that is, the (G′′ /G′ ) ratio, useful for determining the occurrence of molecular mobility transitions such as the glass transition temperature (Tg ). In the temperature dependence of storage modulus G′ and loss factor tan of the relevant PCNs, below the glass transition region, both samples exhibit high G′ and restriction of a substantial drop in G′ with increasing temperature. G′ in the glassy region below Tg is approximately 50–100% higher in the nanocomposite compared with the blend systems without clays. The shift of tan reaches to higher temperatures and a decrease of the value indicates an increase in nanocomposite Tg . In some cases, the magnitude of the shift estimated by tan is about 10
C or more [12]. These indicate that the interaction between polymer and silicate layers at the interface of layers and polymer

Suppliers Nanocor Inc., USA Hojun Yoko Co., Japan CO-OP Chemicals, Japan CO-OP Chemicals, Japan

matrix could suppress the mobility in the polymer segments near the interface. That is, in both systems the intercalated nanocomposites are formed. The large aspect ratio of the structural hierarchy in the nanoscale level might lead to such a large enhancement in G′ throughout the glassy and rubbery region. The temperature dependence of G′ of the PPCNs is reported (Fig. 35) [30]. For all PPCNs, there is a strong enhancement of moduli in the investigated temperature range, which indicates that the plastic and as well as elastic response of PP toward deformation is strongly influenced in the presence of organoclay. Below, Tg , the enhancement of G′ is clear in the intercalated PPCNs. The clay content dependence of G′ is also reported (Fig. 36) [30]. G′PPCN and G′PP-MA are the moduli of the intercalated PPCNs and the PP-MA, respectively. The large reinforcement in G′ is observed in the figure. The essential factor governing the enhancement of mechanical properties is the aspect ratio of the dispersed clay particles. According to the Halpin–Tai theoretical expression on the enhancement of G′ [335], the Einstein coefficient kE was estimated by selecting an appropriate value for the best fit to the experimentally obtained G′PPCN /G′PP-MA versus volume fraction of the clay plots. The estimated values of kE are about 60 for PPCN4 and about 31 for PPCN2 and PPCN7.5. In the intercalated PPCNs, the explanation for the enhancement of G′ by only the kE factor as discussed in the previous case of PMMA/clay nanocomposites [33] is hampered because each PPCN exhibits a different value of kE , despite the different clay contents. In the case of intercalated PPCNs, the enhancement of G′ is

200 nm 70 nm

2.13 nm

1.87 nm

Smectite

2.44 nm

50 nm

Montmorillonite

Mica

Figure 32. Schematic representation of organoclays with C16 ion. Reprinted with permission from [319], P. Maiti et al., Chem. Mater. 14, 4654 (2002). © 2002, American Chemical Society.

Figure 33. WAXD patterns of smectite, MMT, and mica nanocomposites with C16 organoclay and same clay content (3 wt%). Reprinted with permission from [319], P. Maiti et al., Chem. Mater. 14, 4654 (2002). © 2002, American Chemical Society.

Polymer/Clay Nanocomposites

815 10 PPCN 2 PPCN 4 PPCN 7.5

80 60 31

MMT

T = 25°C ω = 6.283 rad/s Strain = 0.05 %

1 0.1

organoclays

nanocomposites

Figure 34. Schematic presentation of silicate layers in organoclay and in various nanocomposites. Reprinted with permission from [319], P. Maiti et al., Chem. Mater. 14, 4654 (2002). © 2002, American Chemical Society.

due to both the degree of the intercalation and the aspect ratio of the dispersed clay particles. The two-dimensional aspect ratios of the dispersed clay particles Lclay /dclay estimated from TEM observation are reported to be 37 for PPCN2, 20 for PPCN4, and 12 for PPCN7.5, respectively. In the PPCN4, despite the lower value of Lclay /dclay compared to PPCN2 and low addition of the clay compared to PPCN7.5, this nanocomposite shows the highest value of kE , suggesting that much higher 10

G' / Pa

10

10

9

8

10

7

ω = 6.283 rad/s Strain = 0.05 % o

G ''/ Pa

10

6

10

8

10

7

10

6

10

5

-1

10

-2

Heating rate = 2 C / min

PP matrix PPCN2 PPCN4 PPCN7.5

tan δ

10

10

-3

-50

10

1

100

Vol % of clay

Mica

10

10

G'PPCN / G'PP-MA

Smectite

0

50

100

150

Temperature / oC

Figure 35. Temperature dependence of G′ , G′′ , and tan for PP-MA matrix and various PPCNs. Reprinted with permission from [30], P. H. Nam et al., Polymer 42, 9633 (2001). © 2001, Elsevier Science.

Figure 36. Plots of G′PPCN /G′PP-MA vs vol% of clay for PPCNs. The value of Einstein coefficient kE was shown in the box. The theoretical lines show the results calculated by the Halpin–Tai expression with various kE . Reprinted with permission from [30], P. H. Nam et al., Polymer 42, 9633 (2001). © 2001, Elsevier Science.

efficiency of the intercalation for the reinforcement is attained. The details were also reported in PBS/clay nanocomposites [325]. The remarkable improvement in G′ related to the strong interaction between matrix and organoclay is clearly observed in case of Nylon 6/clay nanocomposites [336]. Table 10 represents the temperature dependence of G′ and the thermal expansion coefficient (  of the Nylon 6 matrix and various nanocomposites (N6CNs). They summarized the details in a dynamic temperature RAM test for neat Nylon 6 and various N6CNs. All N6CNs show a very high increment of moduli at all temperature ranges. Increased G′ related to the dimension of the dispersed clay particles is further demonstrated in case of PLACNs [34]. In order to understand the effect of compatibilizer on morphology and mechanical properties, the authors also prepared PLACNs with a very small amount of oligoPCL. The details of composition and designation of various types of nanocomposites are presented in Table 7. The enhancement in G′ clearly appears in different magnitudes at investigated temperature ranges for all PLACNs. This behavior indicates that the elastic properties of PLA are significantly affected in the presence of C18 -MMT. Below Tg , the enhancement of G′ is also clear in case of various intercalated PLACNs. On the other hand, all PLACNs show much higher increments of G′ at high temperature ranges compared to that of PLA matrices. This is due to both mechanical reinforcement by the clay particles and extended intercalation at high temperature [260]. Above Tg , when materials become soft, the reinforcement effect of the clay particles becomes prominent, due to restricted movement of the polymer chains, and hence strong enhancement of G′ . At the other extreme, PLACN4 and PLACN5 exhibit strong enhancement of G′ as compared to that of the PLACN2 with comparable clay loading and PLA/o-PCL matrices containing 0.2 and 0.5 wt% oligo-PCL (see Table 8 and Fig. 31b). The presence of small amounts of oligoPCL does not lead to a big shift and broadening of the tan curves. However, a large increment in G′ above Tg became clear, indicating that the large anisotropy of the dispersed particles due to the flocculation enhanced the loss component.

Polymer/Clay Nanocomposites

816

Table 10. Summary of DMA test for Nylon 6 and various N6CNs under different temperature ranges. Sample

Term

0
C

25
C

50
C

100
C

150
C

200
C

Nylon 6

× 105 /cm/cm.
C G′ (GPa)

95 104

12 094

14 052

22 016

31 011

48 0065

N6CN1.6

× 105 /cm/cm.
C G′ (GPa)

89 12

97 11

11 08

16 041

22 027

60 001

N6CN3.7

× 105 /cm/cm.
C G′ (GPa)

64 19

66 18

87 12

85 074

14 052

67 018

N6CN4.4

× 105 /cm/cm.
C G′ (GPa)

71 14

77 13

98 095

11 056

15 038

42 014

G′/Pa

particles above Tg when materials become soft. But in case of PBSCNs, the order of enhancement of G′ is almost same below and above Tg , and this behavior may be due to the extremely low Tg (−29
C) of PBS matrix. At the temperature range of −50 to −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 respectively show higher increments in G′ of 82% and 248% than that of neat PBS, while those of PBSCN1 and PBSCN2 are 18.5% and 44% higher. At 90
C, only PBSCN4 exhibits very strong enhancement of G′ compared to that of the other three PBSCNs. In Figure 38, Okamoto summarized the clay content dependence of G′ of various types of nanocomposites obtained well below Tg . The Einstein coefficient kE derived

10

PBS PBSCN1 PBSCN2 PBSCN3 PBSCN4

9

10 8

10

G′′/Pa

The G′ values of various PLACNs and corresponding matrices without clay at different temperature ranges are summarized in Table 11. PLACNs with a very small amount of oligo-PCL (PLACN4 and PLACN5) exhibit very high enhancement of mechanical properties as compared to that of PLACNs with comparable clay loading (PLACN2). The essential factor governing the enhancement of mechanical properties of the nanocomposites is the aspect ratio of the dispersed clay particles [335]. From the TEM figures (see Fig. 31) it is clearly indicated that, in the presence of a very small amount of oligo-PCL, flocculation of the dispersed clay particles took place again due to the edge–edge interaction of the clay particles. The two-dimensional aspect ratios of the dispersed clay particles Lclay /dclay estimated from TEM observation are 22 for PLACN4 and 12 for PLACN2 (see Table 8) and there is hence strong enhancement of mechanical properties. The increment in G′ directly depends upon the aspect ratio of dispersed clay particles and is also clearly observed in case of PBSCNs. The temperature dependence of G′ of PBS and various PBSCNs are reported (Fig. 37). The nature of enhancement of G′ in PBSCNs with temperature is somewhat different from well-established intercalated PPCNs [30] and well-known exfoliated N6CN systems [105]. In the latter system, there is a maximum of 40–50% increment of G′ compared to that of matrix at well below Tg , and above Tg there is a strong enhancement (>200%) in G′ . This behavior is common for so far reported nanocomposites and the reason is the strong reinforcement effect of the clay

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

8

10 7

Table 11. G′ value of various PLACNs and corresponding matrices without clay at different temperature ranges. 10 6

Storage modulus G′ (GPa) −20
C

40
C

100
C

145
C

PLACN1 PLACN2 PLACN3 PLACN4 PLACN5 PLACN6 PLACN7 PLA PLA1 PLA2 PLA3

232 290 414 371 304 208 186 174 173 168 167

207 265 382 321 260 197 176 160 160 155 162

016 025 027 043 032 023 016 013 013 012 012

009 010 019 016 016 008 007 006 006 006 006

10 -1

tan δ

Samples

10 -2 -40

-20

0

20

40

60

80

100

Temperature / °C

Figure 37. Temperature dependence of G′
G′′ , and their ratio tan for PBSCNs (prepared with C18 -MMT) and neat PBS. Reprinted with permission from [328], K. Okamoto et al., J. Polym. Sci. B, in press.

Polymer/Clay Nanocomposites

817

10

G' nanocomposite/G'matrix

PLACN1 PLACN2 o PLACN3 T=20 C PLACN4 PLACN5 PLACN6

160

70

15 O

O

Si

N H O

N H Si

O

Si

O

Si

O

Si

O

Si

O

N H

N H Si

O

Si

O

OH Si OH

N6CN1.6 N6CN3.7 PBSCN1 PBSCN2 PBSCN3 PBSCN4 PBSCN5 PBSCN6

o

T=0 C

o

T=-50 C

1

1

0.1

Si

100

10

Vol % of clay

Figure 38. Plots of G′nanocomposite /G′matrix vs vol% of clay for various nanocomposites. The Einstein coefficient kE is shown with the number in the box. The lines show the calculated results from the Halpin–Tai theory with various kE .

using Halpin and Tai’s theoretical expression modified by Nielsen is shown in the figure and represents the aspect ratio (Lclay /dclay  of dispersed clay particles without intercalation. From this figure, it is clearly observed that PBSCNs show very high increment in G′ compared to other nanocomposites having the same content of clay in the matrix. PPCNs are well known for intercalated systems, N6CNs are well-established exfoliated nanocomposites, PLACNs are going to establish intercalated-and-flocculated nanocomposites, while PBSCNs are intercalated-andextended flocculated nanocomposites systems [327
328]. Due to the strong interaction between hydroxylated edge– edge groups, the clay particles are sometimes flocculated in the polymer matrix. As a result of this flocculation the length of the clay particles increases enormously and hence the overall aspect ratio. For the preparation of high molecular weight PBS, di-isocyanate end groups are generally used as a chain extender [336]. 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 such kinds of bonds (see schematic illustration in Fig. 39). These urethane type bonds lead to strong interaction with the silicate surface by forming hydrogen bonds and hence strong flocculation (see Fig. 40). For this reason, the aspect ratio of dispersed clay particles is much higher in case of PBSCNs compared to all nanocomposites, and hence there is a high enhancement of modulus.

O

HO

Si

O

Si

O

O Si

Si

O

Si

O

Si

O

Si

N H O

Si

O

Si

O

Si

HO

Figure 40. Formation of hydrogen bonds between PBS and clay, which leads to the flocculation of the dispersed silicate layers.

6.1.2. Tensile Properties The tensile modulus of a polymeric material expressing the stiffness has shown to be remarkably improved when nanocomposites are formed with layered silicates. N6CNs prepared through the intercalative ring opening polymerization of -caprolactam, leading to the formation of exfoliated nanocomposites, exhibit a drastic increase in the tensile properties at rather low filler content. The main reason for the drastic improvement of modulus in case of N6CNs is to the strong interaction between matrix and silicate layers via formation of hydrogen bonds as shown in Figure 41. In case of nanocomposites, the extent of improvement of modulus directly depends upon the average length of the dispersed clay particles, hence the aspect ratio. Figure 42 represents the dependence of tensile modulus (E) measured at 120
C for exfoliated N6CNs with various clay contents obtained by in-situ intercalative polymerization of -caprolactam in the presence of protonated aminododecanoic acid-modified MMT and saponite [2
114]. Moreover, the difference in the extent of exfoliation, as observed for Nylon 6-based nanocomposite synthesized by in-situ intercalative polymerization of -caprolactam using Na+ -MMT and various acids, strongly influenced the final modulus of nanocomposites. In Table 12 we summarized the tensile modulus of 1potNCH together with neat Nylon 6 and N6CN (=NCH) [107]. The excellent modulus in case of 1pot-NCH can be considered to have its origin in the uniformly dispersed silicate layers. Furthermore, 1pot-NCH has much improved mechanical properties compared with NCH. The polymer matrix in nanocomposites prepared by 1pot synthesis is only the homopolymer of Nylon 6, whereas NCH prepared via intercalative ring-opening polymerization is a copolymer of Nylon 6 and a small amount of Nylon 12. The presence of Nylon 12 may give the low modulus. Various kinds of acids have been used to catalyze the polymerization. One 17.2 Å

O C O

H

O

O N N C

O

H

O

O Si

N n

O

H

H

O

N

N

Si

O Si

H O

O Si

Si

O Si

O Si

PBS, Mw is 50300 gm/mol. Figure 39. Formation of urethane bondings in high molecular weight PBS.

Figure 41. Schematic illustration of formation of hydrogen bonds in Nylon 6/MMT nanocomposite.

Polymer/Clay Nanocomposites

818

6 HMW

6

Modulus (GPa)

Tensile modulus (GPa)

8

4

MMW

5

LMW

4

3

2

0

10

20 2

Clay content (wt%)

0

1

2

3

4

5

6

7

MMT (%)

Figure 42. Dependence of tensile modulus (E) on clay content measured at 120
C. Reprinted with permission from [114], L. M. Liu et al., J. Appl. Polym. Sci. 71, 1133 (1999). © 1999, Wiley.

can observe variation of the modulus of the nanocomposites [107]. On the other hand, the WAXD peak intensity (Im  that is inversely related to the exfoliation of clay particles also depends on the nature of the acid used to catalyze the polymerization process. For an increase in the Im values, a parallel decrease in the modulus is observed, indicating that exfoliated layers are the main factor responsible for the stiffness improvement, while intercalated particles, having a less important aspect ratio, rather play a minor role. The effect of MMT content and N6 molecular weight on the tensile modulus of nanocomposites prepared using MMT modified with (HE)2 M1 R1 is shown in Figure 43 [119
123]. The addition of organoclay leads to substantial improvement in stiffness for the composites based on each three Nylon 6 (e.g., LMW, MMW, and HMW). Interestingly, the stiffness increases with increasing matrix molecular weight at any given concentration even though the moduli of the neat Nylon 6 are all quite similar. In Table 13 we summarized the moduli and other mechanical properties of the virgin materials and selected (HE)2 M1 R1 Nylon 6/clay nanocomposites. The slightly larger modulus of 2.82 GPa for LMW may the result of a higher degree of crystallinity resulting in faster crystallization kinetics during the cooling of the specimen during injection molding. Similar trends with respect to the level of organoclay content and molecular weight are evident in the yield strength result. The dependence of yield strength on MMT content and molecular weight is shown in Figure 44. Yield strength increases with the content of MMT. However, the levels of strength improvement for the pure Nylon 6 are nearly

Figure 43. Effect of MMT content on tensile modulus for LMW-, MMW-, and HMW-based nanocomposites. Reprinted with permission from [119], T. D. Fornes et al., Polymer 42, 9929 (2001). © 2001, Elsevier Science.

identical. The HMW- and MMW-based nanocomposites show a steady increase in strength with content of clay, while the LMW-based nanocomposites show a less pronounced effect. The differences in strength improvement with respect to molecular weight are very prominent at the highest clay content. The increase in strength relative to the virgin matrix for the HMW composite is nearly double that of the LMW composite. The relationship between MMT content and elongation at break for the different matrices is shown in Figure 45 for two different rates of extension. Figure 45a shows that the virgin polyamides are very ductile at a test rate of 0.51 cm/min. With increasing clay content the ductility gradually decreases; however, the HMW and MMW based composites show reasonable levels of ductility at MMT concentration as high as 3.5 wt%. Elongation at breaks for the LMW based nanocomposites decreases rapidly at low MMT content at around 1 wt%. The larger reduction in the LMW-based systems may be due to the presence of stacked silicate layers. On the other hand, at the higher testing rate of 5.1 cm/min, shown in Figure 45b, similar trends are observed, but the absolute levels of elongation at break values are significantly lower. Interestingly, the strain at break for LMW composites is relatively independent of rate of extension, similar to what has been observed in case of glass fiber reinforced composites. Even at the highest clay content, the HMW composite exhibits ductile fracture, whereas the LMW- and MMW-based nanocomposites fracture in a brittle manner at the highest clay content.

Table 12. Mechanical properties of 1pot-NCH synthesized in presence of phosphoric acid. Mechanical properties

Method used

Nylon 6

Tensile modulus (GPa) 23
C 120
C

ASTM-D 638M JIS-K 7113

111 019

Tensile strength (MPa) 23
C 120
C

ASTM-D 638M JIS-K 7113

686 266

ASTM-D 648

65

HDT (
C at 1.82 MPa) a

Prepared by in-situ intercalative method.

NCHa (MMT = 47 wt%) 187 061 972 323 152

1pot-NCH (MMT = 41 wt%) 225 067 102 347 160

Polymer/Clay Nanocomposites

819

Table 13. Mechanical properties of some Nylon 6/(HE)2 M1 R1 nanocomposites. Elongation at break (%)

N6/(HE)2 M1 R1 Nanocomposites

Modulus (GPa)

Yield strength (MPa)

Izod impact strength (J/m)

Straina (%)

LMW 0.0 wt% MMT 3.2 wt% MMT 6.4 wt% MMT

282 365 492

692 789 836

40 35 22

232 12 24

MMW 0.0 wt% MMT 3.1 wt% MMT 7.1 wt% MMT

271 366 561

702 866 952

40 35 24

269 81 25

101 18 5

393 383 393

HMW 0.0 wt% MMT 3.2 wt% MMT 7.2 wt% MMT

275 392 570

697 849 976

40 33 26

34 119 41

129 27 61

439 447 462

Crosshead speed 0.51 cm/min 5.1 cm/min 28 11 48

360 323 320

a

Strain at yield point measured during modulus and yield strength testing using a crosshead speed of 0.51 cm/min. Source: Reprinted with permission from [119], T. D. Fornes et al., Polymer 42, 9929 (2001). © 2001, Elsevier Science.

100

Yield Strength (MPa)

HMW MMW

90

LMW

80

increasing clay content. The tensile strength also increases with increasing clay content up to 4 wt% and then levels off. In case of nanocomposites if the systems are not thermodynamically favorable, these properties will be changed during processing because the nanocomposite structure will be changed during processing. A recent work by Reichert et al. [337] showing a systematic study of the dependencies

Elongation at Break (%)

(a)

300

Test Rate = 0.51 cm/min

250 200 HMW 150 MMW

100 50 LMW 0

0

1

2

3

4

5

6

7

MMT (%) (b) Elongation at Break (%)

In case of PPCNs, most of the studies report tensile properties as a function of clay content [254]. Figure 46 shows an Instron study of a neat-PP/f-MMT (modified by alkyltrichlorosilane) composite compared to a PP/2C18 MMT “conventional” composite. In case of PPVNs, there is a sharp increase of tensile modulus for very small clay loading (≤3 wt%) followed by a much slower increase beyond clay loading of 4 wt%, and this is the characteristic behavior of PCNs. With increase clay content, the strength does not change markedly compared to the neatPP value and there is only a small decrease in the maximum strain at break. PP systems—conventionally filled, with no nanometer-level dispersion by similar fillers—do not exhibit as strong improvement in their tensile modulus. On the other hand, as the PP/layered silicate interaction is improved, for example when MA functional groups are incorporated in the polymer, the stresses are much more effectively transferred from the polymer matrix to the inorganic filler, and thus a higher increase in tensile properties is expected. Figure 47a and b respectively represents MMT content dependence of tensile modulus and strength of various PPCNs prepared by melt extrusion of PP-MA and C18 MMT. The modulus of PPCNs systematically increases with

140 Test Rate = 6.1 cm/min

120 100 HMW

80 60

MMW 40 LMW 20

70

0 0

1

2

3

4

5

6

7

MMT (%)

Figure 44. Effect of MMT content on yield strength for LMW-, MMW-, and HMW-Nylon 6 based nanocomposites. Reprinted with permission from [119], T. D. Fornes et al., Polymer 42, 9929 (2001). © 2001, Elsevier Science.

0

1

2

3

4

5

6

7

MMT (%) Figure 45. Effect of MMT content on elongation at break for LMW-, MMW-, and HMW-Nylon 6 based nanocomposites at a crosshead speed of (a) 0.51 and (b) 5.1 cm/min. Reprinted with permission from [119], T. D. Fornes et al., Polymer 42, 9929 (2001). © 2001, Elsevier Science.

Polymer/Clay Nanocomposites 1100 1000 900 800 700 600

Yield Stress (MPa)

35

30

25

Strain at Break (%)

20 800 700 600 500 400 0

1

2

3

4

5

Inorganic concentration φmml (vol %)

Figure 46. Tensile characetization of the PP/f-MMT nanocomposites (
) by Instron. For comparison, conventionally filled PP/2C18 -MMT “macro” composites are also shown (). Reprinted with permission from [254], E. Manias et al., Chem. Mater. 13, 3516 (2001). © 2001, American Chemical Society.

1.6

PP/o-mml

1.5

on compatibilizer functionality and organic modification revealed that considerable enhancement of tensile properties could be achieved only when appropriate PP-MA compatibilizers are used to pretreat the OMLS in conjugation with specific organic modification of the MMT. Similar materials under different processing conditions showed much smaller improvements in the practical materials properties [337]. The evolution of the tensile modulus for the epoxy matrix with three different types of layered silicates loading is presented in Figure 48 [191]. A C18 -MMT, a C18 A-magadiite, and magadiite modified with methyl-octadecylammonium cation (C18 A1M-magadiite) were used for nanocomposite preparation. This figure shows much significant increase of the modulus for the MMT-based nanocomposites with the filler content of 4 wt%. According to the present authors, this behavior is due the difference in layer charge in magadiite and MMT. Organomagadiites have a higher layer charge density and subsequently higher alkylammonium content than organo-MMT. As the alkylammonium ions interact with the epoxy resin while polymerizing, dangling chains are formed, and more of this chain is thus formed in the presence of organomagadiites. These dangling chains are known to weaken the polymer matrix by reducing the degree of network cross-linking, then compromising the reinforcement effect of the silicate layer exfoliation. 4

Tensile Strength (MPa)

Young's Modulus (MPa)

820

1.4

3

2

1

0

1.3

C18A-montmorillonite C18A1M-magadiite C18A-magadiite

2

0

1.2

6

8

10

12

16

1.1

(a)

1.0 2.0

4

Magadiite Loading (wt % SiO2)

PP-MA/o-mml

1.8 1.6 1.4 1.2

(b)

1.0 0

2

4

6

8

10

organo-mmt concentration φ0-mml (wt %)

Tensile Modulus (MPa)

Relative Modulus

A

B

14 12 10 8 6

C18A-montmorillonite C18A1M-magadiite C18A-magadiite

4 2 0

2

4

6

8

10

12

Magadiite Loading (wt % SiO2) Figure 47. Relative moduli of various PP-based nanocomposites, each normalized by modulus of the respective neat PP. (a) PP-based nanocomposites with: f-MMT (
), C18 -MMT (% ), and 2C18 -MMT (). (b) PP-MA/2C18 -MMT nanocomposite (
) and PP hybrids with various PP-g-MA pretreated organically modified MMT: C18 -MMT (right triangle open), C18 -MMT (
△), and C8 -MMT (%
). Reprinted with permission from [254], E. Manias et al., Chem. Mater. 13, 3516 (2001). © 2001, American Chemical Society.

Figure 48. A comparison of (A) the tensile strengths and (B) tensile moduli for epoxy nanocomposites prepared from C18 A-MMT, C18 Amagnitide, and C18 A1M-magnitide. The silicate loading was determined by calcinning the composites in air at 650
C for 4 h at a heating rate of 2
C/min. Reprinted with permission from [191], Z. Wang and T. J. Pinnavaia, Chem. Mater. 10, 1820 (1998). © 1998, American Chemical Society.

Polymer/Clay Nanocomposites

821

For thermoset matrices, a significant enhancement in the tensile modulus is observed for an exfoliated structure when alkylammonium cations with different chain length modified MMTs were used for nanocomposites preparations, with the exception that MMT modified with butylammonium only gives an intercalated structure with a low tensile modulus. Zilg et al. [194] have reported the correlations between polymer morphology, silicate structures, stiffness, and toughness of thermoset nanocomposites as a function of layered silicate type and content. According to them, the main factor for the matrix stiffness improvement resides in the formation of supramolecular assemblies obtained by the presence of dispersed anisotropic laminated nanoparticles. They also described a stiffening effect when the MMT is modified by a functionalized organic cation (carboxylic acid or hydroxyl groups) that can strongly interact with the matrix during curing. The tensile properties of APES/Cloisite 30B and APES/ Cloisite 10A nanocomposites at various clay contents are presented in Table 14 [338]. In comparison to the APES, the tensile strength and modulus have been improved with little decrease of elongation at break. APES/Cloisite 30B nanocomposites exhibit much higher tensile strength and modulus compared to the APES/Cloisite 10A nanocomposites. This is also attributed to the strong interaction between APES matrix and Cloisite 30B. These results further confirm the importance of strong interaction between matrix and clay, which ultimately leads to the better overall dispersion as already observed by TEM observations.

6.1.3. Flexural Properties Nanocomposite researchers are generally interested on tensile properties of final materials and there are very few reports concerning the flexural properties of neat polymer and its nanocomposite with OMLS. Very recently, Sinha Ray et al. [324] reported detailed measurements of flexural properties of neat PLA and various PLACNs. They conducted flexural properties measurements with injectionmolded samples and according to the ASTM D-790 method. Table 15 shows the flexural modulus, flexural strength, and distortion at break of neat PLA and various PLACNs measured at 25
C. There is a significant increase in flexural modulus for PLACN4 compared to that of neat PLA followed by a much slower increase with increasing OMLS content, and a maximum of 21% in case of PLACN7. On the other hand, flexural strength and distortion at break remarkably increase with PLACN4 then gradually decrease with organoclay loading. According to the authors this behavior Table 14. Tensile properties of APES/Closite 30B nanocomposites. Closite 30B content (wt%) 0 1 3 5 10 20 30

Modulus (kgf/cm2 )

Strength (kgf/cm2 )

Elongation at break (%)

1067 1123 1144 1182 1295 1444 1738

1317 1390 1441 1498 1577 1908 2135

1245 1225 1195 1140 1090 1130 1225

Table 15. Comparison of materials properties between neat PLA and various PLACNs prepared with trimethyl octadecylammonium modified MMT. Materials properties

PLA

PLACN4

PLACN5

PLACN7

Bending modulus (GPa) Bending strength (MPa) Distortion at break (%)

48 86 19

55 134 31

56 122 26

58 105 2

may be due to high organoclay content leading to a brittleness of materials.

6.1.4. Heat Distortion Temperature Heat distortion temperature (HDT) of a polymeric material is an index of heat resistance toward applied load. Most of the PCNs studies report HDT as a function of clay content, characterized by ASTM D-648. Kojima et al. [2] first show that HDT of pure N6 increases up to 80
C after PCN preparation with MMT. In their further work [106] they report clay content dependence of HDT of NCH. In case of NCH there is a marked increase in HDT, from 65
C for the neat N6 to 150
C for 4.7 wt% nanocomposite; beyond that wt% of MMT, the HDT of nanocomposite levels off. They also conducted HDT tests of various NCHs prepared with clays having different lengths and found that HDT also depends upon the aspect ratio of dispersed clay particles [2]. Like all other mechanical properties, 1potNCH also shows higher HDT than that of NCH prepared by in-situ intercalative polymerization. Since the degree of crystallinity of Nylon 6/clay nanocomposite is independent of the amount and nature of clay, the HDT of Nylon 6/clay nanocomposite is due to the presence of a strong interaction between matrix and silicate surface by forming hydrogen bonds (see Fig. 41). Although Nylon 6 in nanocomposite stabilizes in a different crystal phase (-phase) than that found in pure Nylon 6, this different crystal phase is not responsible for higher mechanical properties of Nylon 6/clay nanocomposites because -phase is a very soft crystal phase. On the other hand, increased mechanical properties of NCH with increasing clay content is due to the mechanical reinforcement effect. The nanodispersion of MMT in the PP matrix also promotes a higher HDT [254]. In Table 16 we summarized the HDT of PP and its PCNs based on f-MMT and alkylammonium modified MMT. Like previous systems, there is a significant increase in HDT, from 109
C for the neat PP to 152
C for a 6 wt% of clay; after that the HDT of nanocomposite levels off. This improvement in HDT of neat PP after nanocomposite preparation originates from the better mechanical stability of the nanocomposite compared to the neat PP since there is no increment of melting point of neat PP after PCN preparation. Sinha Ray et al. examined the HDT of neat PLA and various PLACNs with different load conditions. As seen in Figure 49a [324], in case of PLACN7 (inorganic clay content = 5 wt%), there is marked increase of HDT with intermediate load of 0.98 MPa, from 76
C for the neat PLA to 98
C for PLACN4 (inorganic clay content = 3 wt%). The value of HDT gradually increases with increasing organoclay

Polymer/Clay Nanocomposites

822 Table 16. HDT of PP/clay nanocomposites and the respective unfilled PP. HDT (
C)

Organically modified MMT (wt%)

PP/f-MMT

0 3 6 9

PP/alkyl-MMT

109 ± 3 144 ± 5 152 ± 5 153 ± 5

109 ± 3 130 ± 7a 141 ± 7b

mechanical stability of the PLACNs due to mechanical reinforcement by the dispersed clay particles, higher value of the degree of crystallinity c , and intercalation. This is qualitatively different from the behavior of the NCH system, where the MMT layers stabilize in a different crystalline phase (phase) [106] than that found in the neat Nylon 6, with the strong hydrogen bondings between the silicate layers and Nylon 6 as a result of the discrete lamellar structure on both sides of the clay (see Fig. 41).

a

C18 -MMT filler, extruder processed. 2C18 -MMT filler, twin-head mixer. Source: Reprinted with permission from [254], E. Manias et al., Chem. Mater. 13, 3516 (2001). © 2001, American Chemical Society. b

content, and in case of PLACN7, the value increases up to 111
C. On the other hand, imposed load dependence on HDT is clearly observed in case of PLACNs. Figure 49b shows the typical load dependence in case of PLACN7. The increase of HDT of neat PLA due to nanocomposite preparation is a very important property improvement, not only from the industrial point of view but also molecular control on the silicate layers, that is, crystallization through interaction between PLA molecules and SiO4 tetrahedral layers in the MMT. In case of high load (1.81 MPa), it is very difficult to achieve high HDT enhancement without strong interaction between polymer matrix and organo-MMT [2]. In case of all PLACNs studied here, the values of the melting temperature Tm do not change significantly as compared to that of neat PLA. So the improvement of HDT with intermediate load (0.98 MPa) originates from the better (a)

120

HDT /°C

110 100 90 80 Load =0.98 MPa 70

0

2

4

6

8

Organoclay /wt.% (b) 160 PLA PLACN7

HDT /°C

140 120 100 80 60 0.4

0.8

1.2

1.6

2

Load /MPa Figure 49. (a) Organoclay (wt%) dependence of HDT of neat PLA and various PLACNs. (b) Load dependence of HDT of neat PLA and PLACN7. Reprinted with permission from [324], S. Sinha Ray et al., Polymer 44, 857 (2003). © 2003, Elsevier Science.

6.2. Thermal Stability The thermal stability of polymeric materials is usually studied by thermogravimetric analysis (TGA). The weight loss due to the formation of volatile degradation products is monitored as a function of temperature ramp. When the heating is operated under an inert gas flow, a nonoxidative degradation occurs while the use of air or oxygen is allowed following the oxidative degradation of the samples. Generally the incorporation of clay in the polymer matrix enhanced the thermal stability by acting as a superior insulator and mass transport barrier to the volatile products generated during decomposition. Blumstein [35] first reported the improved thermal stability of a PCN that combined PMMA and MMT. PMMA intercalated (d-spacing increase of 0.76 nm) between the galleries of MMT clay resisted the thermal degradation under conditions that would otherwise completely degrade pure PMMA. These PMMA nanocomposites were prepared by free-radical polymerization of MMA intercalated in the clay. TGA data reveal that both linear PMMA and crosslinked PMMA intercalated into MMT layers have a 40– 50
C higher decomposition temperature. Blumstein argues that the stability of the PMMA nanocomposite is due not only to its different structure but also restricted thermal motion of the PMMA in the gallery. Recently, there have been lots of reports concerned with the improved thermal stability of nanocomposites prepared with various types of organoclay and polymer matrices [339– 341]. Recently Zanetti et al. [275] published the details TG analyses of nanocomposites based on EVA. The inorganic phase was florohectorite or MMT, both exchanged with octadecylammonium cation. According to them the deacylation of EVA in nanocomposites is accelerated and may occur at temperatures lower than those for the pure polymer or corresponding microcomposite due to catalysis by the strongly acid sites created by thermal decomposition of the silicate modifier. These sites are active when there is an intimate contact between the polymer and the silicate. Slowing down of the volatilization of the deacylated polymer in nitrogen may be due to the labyrinth effect of the silicate layers in the polymer matrix [342]. In air, the nanocomposite presents a significant delay of weight loss that may derive from the barrier effect due to diffusion of both the volatile thermo-oxidation products to the gas and oxygen from the gas phase to the polymer. According to Gilman et al. this barrier effect increases during volatilization owing to ablative reassembly of the reticular of the silicate on the surface [343]. The thermal stability of the PCL-based nanocomposites has also been studied by TGA. Generally, the degradation

Polymer/Clay Nanocomposites

823

of PCL fits a two-step mechanism [140]: first there is a statistical rapture of the polyester chains by pyrolysis of ester groups with the release of CO2 , H2 O, and hexanoic acid and in the second step, -caprolactone (cyclic monomer) is formed as a result of an unzipping depolymerization process. The thermograms of nanocomposites prepared with organoclay and pure PCL recovered after clay extraction are presented Figure 50. Both intercalated and exfoliated nanocomposites show higher thermal stability when compared to the pure PCL or the corresponding microcomposites. The nanocomposites exhibit a 25
C rise in decomposition temperature at 50% weight loss. The shift of the degradation temperature may be ascribed to a decrease in oxygen and volatile degradation product permeability/diffusivity due to the homogeneous incorporation of clay sheets, to a barrier of these high aspect ratio fillers, and to char formation. The thermal stability of the nanocomposites systematically increases with increasing clay; however, above a loading of 5 wt%, the thermal stability is not improved anymore. But completely different behavior is observed in case of synthetic biodegradable aliphatic polyester (BAP)/clay nanocomposite systems, in which the thermal degradation temperature and thermal degradation rate are systematically increased with an increasing amount of organoclay up to 15 wt% [344]. Like PS-based nanocomposites, a small amount of clay also increased the residual weight of BAP/OMMT because of the restricted thermal motion of the polymer in the silicate layers. The residual weight of various materials at 450
C increased in the order BAP < BAP03 < BAP06 < BAP09 < BAP15 (here number indicates wt% of clay). This kind of improved thermal property is also observed in other systems like SAN [345], the intercalated nanocomposites prepared by emulsion polymerization. Many researchers believe the role of clay in the nanocomposite structure might be the main reason for the difference in TGA results of these systems compared to the systems reported thus far. The clay acts as a heat barrier, which could enhance the overall thermal stability of the system, as well as assisting in the formation of char after thermal decomposition. Thereby, in the beginning stage of 100

weight (%)

75

PCL

50

1 wt% 3 wt%

25

5 wt% 10 wt% 0 250

300

350

400

450

500

temperature °C

Figure 50. Temperature dependence of the weight loss under an air flow for neat PCL and PCL nanocomposites containing 1, 3, 5, and 10 wt% (relative to inorganics) of MMT-Alk (heating rate: 20 K/min). Reprinted with permission from [140], B. Lepoittevin et al., Polymer 43, 1111 (2002). © 2002, Elsevier Science.

thermal decomposition, the clay could shift the decomposition temperature higher. However, after that, this heat barrier effect would result in a reverse thermal stability. In other words, the stacked silicate layers could hold accumulated heat that could be used as a heat source to accelerate the decomposition process, in conjunction with the heat flow supplied by the outside heat source.

6.3. Fire Retardant Properties The cone calorimeter is the most effective bench-scale method for studying the fire retardant properties of polymeric materials. Fire-relevant properties, measured by the cone calorimeter, such as heat release rate (HRR), peak HRR, and smoke and CO yield, are vital to the evaluation of the fire safety of materials. Gilman et al. reviewed the flame retardant properties of nanocomposites in detail [346
347]. Table 17 represents the cone calorimeter data of three different kinds of polymers and their nanocomposites with MMT. From Table 17 we can see that all MMT-based nanocomposites reported here show reduced flammability. Peak HRR is reduced by 50–75% for Nylon 6, PS, and PP-MA nanocomposites [347]. According to the authors the MMT must be nanodispersed for it to affect the flammability of the nanocomposites. However, the clay need not be completely delaminated for it to affect the flammability of the nanocomposite. In general, the nanocomposite flame retardant mechanism is that a high-performance carbonaceous-silicate char builds up on the surface during burning; this insulates the underlying material and slows the mass loss rate of decomposition products. For a PPCN with 4 wt% organoclay [347], there is a 75% reduction in flammability compared to the neat matrix (see Fig. 51).

6.4. Gas Barrier Properties Nanoclays are believed to increase the barrier properties by creating a maze or “tortuous path” (see Fig. 52) that retards the progress of the gas molecules through the matrix resin. The direct benefit of the formation of this type of path is clearly observed in polyimide/clay nanocomposite by showing dramatically improved barrier properties with simultaneous decrease in thermal expansion coefficient [211
213]. The polyimide/layered silicate nanocomposites revealed a several-fold reduction in the permeability of small gases (e.g. O2 , H2 O, He, CO2 , and the organic vapor ethylacetate) with the presence of a small fraction of organoclay. For example, at 2 wt% clay loading, the permeability coefficient of water vapor was decreased tenfold for the synthetic mica relative to pristine polyimide. By comparing nanocomposites made with layered silicates of various aspect ratios the permeability was noted to decrease with increasing aspect ratio. The O2 gas permeability was measured for the exfoliated PLA/synthetic mica nanocomposites prepared by Sinha Ray et al. [348]. The relative permeability coefficient value (i.e., PPLACN /PPLA where PPLACN and PPLA stand for the nanocomposite and pure PLA permeability coefficient, respectively) has been plotted as a function of the wt% of clay. The curve fitting has been achieved by using Nielsen theoretical

Polymer/Clay Nanocomposites

824 Table 17. Cone calorimeter data of various polymers and their nanocomposites with organoclay.

Sample (structure)

% Residue yield (±05)

Peak HRR (kW/m2 ) ()%)

Mean HRR (kW/m2 ) ()%)

Mean Hc (MJ/kg)

Mean SEA (m2 /kg)

Mean CO yield (kg/kg)

1 3

1010 686 (32)

603 390 (35)

27 27

197 271

001 001

6

378 (63)

304 (50)

27

296

002

0 3

1120 1080

703 715

29 29

1460 1840

009 009

4

567 (48)

444 (38)

27

1730

008

3 5 6

491 (56) 1525 450 (70)

318 (54) 536 322 (40)

11 39 44

2580 704 1028

014 002 002

12

381 (75)

275 (49)

44

968

002

Nylon 6 N6 nanocomposite 2% (delaminated) N6 nanocomposite 5% (delaminated) PS PS-silicate mix 3% (immiscible) PS-nanocomposite 3% (intercalated/delaminated) PSw/DBDPO/Sb2 O3  30% PP-MA PP-MA nanocomposite 2% (intercalated/delaminated) PP-MA nanocomposite 4% (intercalated/delaminated)

a Heat flux, 35 kW/m2 . Hc , specific heat of combustion; SEA, specific extinction area. Peak heat release rate, mass loss rate, and SEA data, measured at 35 kW/m2 , are reproducible to within ±105. The carbon monoxide and heat of combustion data are reproducible to within ±15%. Source: Reprinted with permission from [347], J. W. Gilman et al., Chem. Mater. 12, 1866 (2000). © 2000, American Chemical Society.

expression [349] allowing the prediction of gas permeability in function of the length and width of the filler particles as well as their volume fraction with in the PLA matrix (see Fig. 53). In the Nielsen model [349], platelets of length (Lclay ) and width (Dclay ) of the clay are dispersed parallel in polymer matrix; then the tortuosity factor (+) can be expressed as (1)

+ = 1 + Lclay /2Dclay ,clay

where ,clay is the volume fraction of dispersed clay particles. Therefore, the relative permeability coefficient (PPCN /PNeat  is given by PPCN /PNeat = + −1 = 1/1 + Lclay /2Dclay ,clay 

(2)

where PPCN and PNeat are the permeability coefficient of PCN and neat polymer, respectively. PP intercalated (2% silicate) PP pure (PP-g-MA, 0.4% MA) PP intercalated (4% silicate)

Heat Release Rate (kW/m2)

1600 1400

Flux = 35 kW/m2

1200 1000 800 600

The H2 O-vapor permeability for the PUU/clay nanocomposites is presented by Xu et al. in terms of Pc /Po , that is, the permeability coefficient of the nanocomposite (Pc ) relative to that of the neat PUU (Po  [350]. The nanocomposite formation results in a dramatic decrease in H2 Ovapor transmission through the PUU sheet. The solid lines are based on the argument of the tortuosity model for the aspect ratio of 300 and 1000. A comparison between the experimental values and the theoretical model prediction suggests a gradual change in the effective aspect ratio of the filler (Lclay /Dclay .

6.5. Ionic Conductivity Solvent-free electrolytes are of much interest because of their charge-transport mechanism and their possible applications in electrochemical devices. With this background, Vaia et al. [152] have considered the preparation of PEO/clay nanocomposites to fine tune ionic conductivity of PEO. An intercalated nanocomposite prepared by melt intercalation of PEO (40 wt%) into Li+ -MMT (60 wt%) has shown to enhance the stability of the ionic conductivity at lower temperature when compared to a more conventional PEO/LiBF4 mixture. This improvement in conductivity is explained by the fact that PEO is not able to crystallize when intercalated, hence eliminating the presence of crystallites, nonconductive in nature. The higher conductivity at room temperature compared to conventional

400 200 0 0

Conventional composites 120

240

360

480

“Tortuous path” in layered silicate nanocomposites

600

Time (seconds) Figure 51. Heat release rate during cone-calorimetry combustion of heat PP-MA and PPCNs. Reprinted with permission from [347], J. W. Gilman et al., Chem. Mater. 12, 1866 (2000). © 2000, American Chemical Society.

Figure 52. Formation of tortuous path in polymer/clay nanocomposites.

825

Experimental value 150

100

50

0 0

2

4

6

8

10

OMSFM /wt % Figure 53. Oxygen gas permeability of neat PLA and various PLACNs as a function of organoclay content measured at 20
C and 90% relative humidity. The filled circles represent the experimental data. Theoretical fits based on Nielsen tortuousity model.

PEO/LiBF4 electrolytes with a single ionic conductor character makes those nanocomposites new promising electrolyte materials. The same type of behavior in ionic conductivity is also observed in case of poly[bis(methoxy-ethoxy) ethoxy phosphazene]/Na+ -MMT nanocomposite as prepared by Hutchison et al. [351]. In a recent report, Okamoto et al. [12] have reported the correlation between internal structure and ionic conductivity behavior of PMMA/clay and PS/clay nanocomposites having various dispersed morphologies of the clay layers by using an impedance analyzer in the temperature range of 90–150
C. The nanocomposites having finer dispersion of the clay layers exhibit higher ionic conductivity rather than the other systems such as PMMA/clay nanocomposite with stacking layer structure. The activation energy of the conductivity in finer dispersed morphology systems becomes larger than the other systems and the corresponding organoclay solids.

6.6. Optical Transparency Although layered silicates are micrometer in lateral size, they are just 1 nm thick. Thus, when single layers are dispersed in a polymer matrix, the resulting nanocomposite is optically clear in the visible region. Strawhecker and Manias have reported the ultraviolet (UV)/visible transmission spectra of pure PVA and PVA/Na+ -MMT nanocomposites with 4 and 10 wt% MMT [87]. The spectra show that the visible region is not affected at all by the presence of the silicate layers and retains the high transparency of the PVA. For the UV wavelengths, there is strong scattering and/or absorption, resulting in very low transmission of the UV light. This behavior is not surprising as the typical MMT lateral sizes are 50–1000 nm. Like PVA, various other polymers also show optical transparency after nanocomposite preparation with organoclay [254].

6.7. Biodegradability of Green Polymeric Nanocomposites Another most interesting and exciting aspect of nanocomposite technology is the significant improvement of biodegradability of biodegradable polymers after nanocomposite preparation with organoclay. There is a need for

the development of green polymeric materials that would not involve the use of toxic or noxious components in their manufacture and could be degraded in the natural environment or easily recycled. Aliphatic polyesters are among the most promising materials for the production of environmentally friendly biodegradable plastics. Biodegradation of aliphatic polyester is well known, in that some bacteria degrade them by producing enzymes, which attack the polymer. Tetto et al. [352] first reported some results about the biodegradability of nanocomposites based on PCL, where the authors found that the PCL/clay nanocomposites showed improved biodegradability compared to pure PCL. According to them, the improved biodegradability of PCL after nanocomposite formation may be due to the catalytic role of the organoclay in the biodegradation mechanism. But still it is unclear how the clay increases the biodegradation rate of PCL. Recently, Lee et al. [332] reported the biodegradation of aliphatic polyester based nanocomposites under compost. Figure 54a and b respectively represents the clay content dependence of biodegradation of APES based nanocomposites prepared with two different types of clays. They assumed that the retardation of biodegradation is due to the improvement of the barrier properties of the aliphatic APES after nanocomposites preparation with clay. However, there are no data about permeability.

100

Biodegradability (wt%)

Theoretical curve based on Lclay/Dclay = 275

200

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

80 60

(a)

40

20

0 0

5

10

15

20

25

30

35

25

30

35

Time (Day) 100

Biodegradability (wt%)

O2 TR /ml.mm.m-2.day-1.MPa-1

Polymer/Clay Nanocomposites

APES/10A (97/3 wt %) APES/10A (95/5 wt %) APES/10A (90/10 wt %) APES/10A (90/20 wt %) APES/10A (70/30 wt %) APES

80

(b)

60

40

20

0 0

5

10

15

20

Time (Day) Figure 54. Biodegradability of APES nanocomposites with: (a) Closite 30B and (b) Closite 10A. Reprinted with permission from [332], S. R. Lee et al., Polymer 43, 2495 (2002). © 2002, Elsevier Science.

Polymer/Clay Nanocomposites

826 Very recently, Sinha Ray et al. [318
324] first reported the biodegradability of neat PLA and corresponding nanocomposites prepared with trimethyl octadecylammonium modified MMT (C3 C18 -MMT) with detailed mechanisms. The compost used was prepared from food waste and tests were carried out at temperature of 58 ± 2
C. Figure 55a shows the real picture of the recovered samples of neat PLA and PLACN4 (C3 C18 -MMT = 4 wt%) from compost with time. The decreased Mw and residual weight percentage Rw of the initial test samples with time are also reported in Figure 55b. The biodegradability of neat PLA is significantly enhanced after PCN preparation. Within one month, both extent of Mw and extent of weight loss are at almost the same level for both PLA and PLACN4. However, after one month, a sharp change occurs in weight loss of PLACN4, and within two months, it is completely degraded in compost. The degradation of PLA in compost is a complex process involving four main phenomena, namely: water absorption, ester cleavage and formation of oligomer fragments, solubilization of oligomer fragments, and finally diffusion of soluble oligomers by bacteria [353]. Therefore, the factor that increases the hydrolysis tendency of PLA ultimately controls the degradation of PLA. They expect the presence of terminal hydroxylated edge groups of the silicate layers may be one of the factors responsible for (a)

after 50 days

after 32 days

after 60 days

PLA

Degree of Biodegradation / %

100

PLACN4 (b) PLACN4

100

80 60

50

40 20

PLACN4

0

0

PLA

10

20

30

40

50

60

60 40

70

Time /days

Neat PLA PLACN4

20

0

10

20

30

40

50

Time / Days 200

(b) Mw×10-3 /g.mol-1

Rw /%

100

(a) 80

0

150

MW x 10–3 /(gm/mol)

PLA

0

this behavior. In case of PLACN4, the stacked (∼4 layers) and intercalated silicate layers are homogeneously dispersed in the PLA matrix (from TEM image) and these hydroxy groups start heterogeneous hydrolysis of the PLA matrix after absorbing water from compost. This process takes some time to start. For this reason, the weight loss and degree of hydrolysis of PLA and PLACN4 are almost same up to one month (see Fig. 55b). However, after one month there is a sharp weight loss in case of PLACN4 compared to that of PLA. That means one month is a critical value to start heterogeneous hydrolysis, and due to this type of hydrolysis the matrix becomes very small fragments and disappears with the compost. This assumption was confirmed by conducting the same type of experiment with PLACN prepared by using dimethyl dioctdecyl ammonium salt modified synthetic mica which has no terminal hydroxylated edge group, and the degradation tendency was almost the same with neat PLA [354]. They also conducted respirometric testing to study degradation of the PLA matrix in compost environment at 58 ± 2
C. For this test the compost used was made from bean-curd refuse, food waste, and cattle feces. Unlike weight loss, which reflects the structural changes in the test sample, CO2 evolution provides an indicator of the ultimate biodegradability of PLA in PLACN4 (prepared with N (cocoalkyl)N
N -[bis(2-hydroxyethyl)]-N methylammonium modified synthetic mica), that is, mineralization, of the samples. Figure 56 shows the time dependence of the degree of biodegradation of neat PLA and PLACN4, indicating that the biodegradability of PLA in PLACN4 is enhanced significantly. The presence of organoclay may thus cause a different mode of attack on the PLA

Neat PLA PLACN4

150

100

50

0

0

2

4

6

8

10

12

14

Time / Days Figure 55. (a) Real picture of biodegradability of neat PLA and PLACN4 recovered from compost with time. Initial shape of the crystallized samples was 3 × 10 × 01 cm3 . (b) Time dependence of residual weight Rw and of matrix Mw of PLA and PLACN4 under compost at 58 ± 2
C. Reprinted with permission from [324], S. Sinha Ray et al., Polymer 44, 857 (2003). © 2003, Elsevier Science.

Figure 56. Degree of biodegradation (i.e., CO2 evolution), and (b) time-dependent change of matrix Mw of neat PLA and PLACN4 (MEE clay = 4 wt%) under compost at 58 ± 2
C. Reprinted with permission from [317], S. Sinha Ray et al., Macromol. Rapid Commun. 23, 943 (2002). © 2002, Wiley–VCH.

Polymer/Clay Nanocomposites

827

component, which might be due to the presence of hydroxy groups. Details degrading the mechanism of biodegradability are presented in relevant literature [317
354].

spherulites (i.e., the number of heterogeneous nuclei N ) was given by

6.8. Other Properties

where Dm is the maximum diameter of the spherulite (i.e., the attainable diameter before impingement). The calculated values of N at 130
C were 4 × 10−8 for PP-MA, 50 × 10−8 for PPCN2, and 200 × 10−8 m−3 for PPCN7.5, respectively. The time variation of the volume fraction of the spherulites increases in proportion to NG3 (overall crystallization rate). This fact suggests that the overall crystallization rate of the PPCNs is about one or two orders of magnitude higher than that of matrix PP-MA without clay.

Another property of polymer that is strongly affected by the incorporation of layered silicates is a sharp increase of the scratch resistance [254], dimensional stability [355], and solvent resistance [355] via nanocomposite technology.

7. CRYSTALLIZATION OF POLYMER/CLAY NANOCOMPOSITES

N = 3/4. Dm /2−3

(3)

7.1. Spherulitic Texture and Growth

7.2. Formation of
-Form in PPCNs

Crystallization of PCNs might be a good tool for controlling the structure od PCNs and thereby the various properties. Maiti et al. [260] have reported an example of the time variation of the diameter of the spherulite D for PP-MA and PP/clay nanocomposites at 135
C (see Fig. 57). A linear growth of D is seen in a range of t scale for PP-MA, PPCN2, and PPCN7.5. The linear growth rate G =1/2 dD/dt, defined as the initial slope of the plots, slightly increases with increasing clay content. From the extrapolation of D vs t plots, we estimated the onset time t0 , which corresponds to the induction time of the crystallization. The t0 of both PPCNs decreases with clay content compared to the PP-MA matrix without clay. The reduction of t0 in the PPCNs is attributed to the nature of the clay as the nucleating agent. For PPCNs, G shows almost the same value compared to PP-MA without clay. In spite of the increase in clay content, the dispersed clay particles have not much effect on the crystallization and no big acceleration of G in the crystallization of the PPCNs is seen. In the changes in t0 with crystallization temperture Tc , the PPCNs show remarkably short time especially at high Tc , suggesting that the dispersed clay particles have some contribution to enhance the nucleation as mentioned. The primary nucleation density of the

Maiti et al. [261] have reported the temperature dependence of the fraction of -form f . The value of f was calculated from the area of the specific peak, compared to the total area of and -forms. At low Tc ( LMW, and hence the resulting shear stresses exerted by the pure polymers also follows same order. Therefore, during melt mixing, the level of stress exerted on the organoclay by the LMW polyamide is significantly lower than those developed in presence of HMW or MMW polyamides. As a result the breakup of clay particles is much easier in case of HMW polyamide and ultimately improved the clay particle dispersion. Figure 64 shows schematic suggestions of various roles that shear stress may play during the melt compounding of nanocomposites. Therefore, the role of polymer molecular weight is believed to stem from the fact that a melt viscosity exerts the taller stacks into shorter ones. The final step in exfoliation involves peeling the platelets of the stacks one by one, and this takes time and requires a strong matrix–organoclay interaction to cause spontaneous wetting.

8.2. Steady Shear Flow The steady shear rheological behaviors of neat PBS and various PBSCNs are shown in Figure 65. The steady viscosity of PBSCNs is enhanced considerably at all shear rates with time and at a fixed shear rate increases monotonically with increasing silicate loading [327]. On the other hand, all intercalated PBSCNs exhibit strong rheopexy behavior, 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 attains a plateau after a certain time, and the time required to attain this plateau decreases with increasing shear rates. The possible

Shear

Stacks of silicate platelets of tactoids Organoclay particle (~8 µm)

(a) Shear

Stress = η γ• Shearing of platelet stacks leads to smaller tactoids (b) Shear

Diffusion

Platelets peel apart by combined diffusion/shear process (c) Figure 64. Stemwise mechanism of clay platelets exfoliation during melt compounding: (a) organoclay breakup, (b) intercalated organoclay tactoid breakup, and (c) platelet exfoliation. Reprinted with permission from [119], T. D. Fornes et al., Polymer 42, 9929 (2001). © 2001, Elsevier Science.

reasons for this type of behavior may be the planer alignment of the clay particles toward the flow direction under shear. When the shear rate is very slow (0.001 s−1 ), clay particles take longer to attain complete planer alignment along the flow direction, and this measurement time (1000 s) is too short to attain such alignment and hence shows strong rheopexy behavior. On the other hand, under high shear rates (0.005 or 0.01 s−1 ) this measurement time is considerable enough to attain such alignment, and hence nanocomposites show time independent shear viscosity after a certain time. Figure 66 shows shear rate dependence of viscosity for neat PBS and corresponding nanocomposites measured at 120
C. The neat PBS exhibits almost Newtonian behavior at all shear rates, whereas nanocomposites exhibited nonNewtonian behavior. At very low shear rates, shear viscosity of nanocomposites initially exhibits some shear-tickening behavior and this corresponds to the rheopexy behavior as we observed at very low shear rates (see Fig. 65). After that all nanocomposites show very strong shear thinning behavior at all shear rates and this behavior is analogous to the results obtained in case of dynamic oscillatory shear measurements [324]. Additionally, at very high shear rates, the viscosities of nanocomposites are comparable to that of neat PBS. These observations suggest that the silicate layers are strongly oriented toward the flow direction at high shear rates, and

Polymer/Clay Nanocomposites

833 106

104

PBS

PBS

Temp. = 120 °C

Temperature =120°C

PBSCN2 5

3

10

η/Pa.s

10 Shear rate =0.001s-1 Shear rate =0.005s-1 Shear rate = 0.01s-1

PBSCN3 PBSCN4 PBSCN6

4

10

103

102 PBSCN2

105

102 10-4

104

10-3

10-2

10-1

100

. γ/s-1

101

Figure 66. Shear viscosity as a function of shear rates for the shear rate sweep test. Reprinted with permission from [327], S. Sinha Ray et al., Macromolecules 36, 2355 (2003). © 2003, American Chemical Society.

103 PBSCN3

105

/Pa.s

8.3. Elongational Flow and Strain-Induced Hardening

104

Okamoto et al. [256] first conducted elongation tests of PP/clay nanocomposites (PPCN4) under molten state at constant Hencky strain rate ˙0 using an elongation flow optorheometry [367], and also they attempted to control the alignment of the dispersed silicate layers with nanometer dimensions of an intercalated PPCNs under uniaxial elongational flow. Figure 67 shows double logarithmic plots of transient elongational viscosity E ˙0 2 t against time t observed for

103 PBSCN4

105

104

103

108

PBSCN6

N6CN3.7 106

104

225°C

Eo/s–1 1.0 0.5 0.1 0.05 0.03 0.01 0.005 0.001 3+ηo (cone-plate: 0.001s–1)

104 103 101

102

103

104

102

Figure 65. Time variation of shear viscosity for PBSCN. Reprinted with permission from [327], S. Sinha Ray et al., Macromolecules 36, 2355 (2003). © 2003, American Chemical Society.

shear thinning behavior at high shear rates is dominated by that of neat polymer. The PCNs always exhibit significant deviation from Cox– Merz relation [366], while all neat polymers nicely obey the empirical Cox–Merz relation, which requires that for ˙ = , the viscoelastic data should obey the relationship ˙ = ∗ 

. We believe there are two possible rea  sons for the deviation from the Cox–Merz relation in case of nanocomposites: first of all this rule is only applicable for homogenous systems like homopolymer melts but nanocomposites are heterogeneous systems. For this reason this relation is nicely obeyed in case of neat polymer [327]. Second, the structure formation is different when nanocomposites are subjected to dynamic oscillatory shear and steady shear measurements.

h/ Pa.s

Time /s

(a) 100 PPCN4 106

150°C

104

102

(b) 100 10–1

100

101

102

103

Time / s Figure 67. Time variation of elongational viscosity E ˙0 2 t for (a) N6CN3.7 melt at 225
C and for (b) PPCN4 at 150
C. The solid ˙ t, taken at a low line shows three times the shear viscosity, 3E 2 shear rate ˙ = 0001 s−1 on a cone-plate rheometer.

Polymer/Clay Nanocomposites

834 a Nylon 6/clay system (N6CN3.7) and PPCN4 (MMT = 4 wt%) with different Hencky strain rates ˙0 ranging from 0.001 to 1.0 s−1 . The solid curve represents time develop˙ t, at 225
C with ment of threefold shear viscosity, 30 2 −1 a constant shear rate ˙ = 0001 s . In E ˙0 2 t at any ˙0 , N6CN3.7 melt shows a weak tendency of strain-induced hardening as compared to that of PPCN4 melt. A strong behavior of strain-induced hardening for the PPCN4 melt originated from the perpendicular alignment of the silicate layers to the stretching direction as reported by Okamoto et al. [256]. From TEM observations (see Fig. 60), the N6CN3.7 forms a fine dispersion of the silicate layers of about 100 nm in Lclay , 3 nm thickness in dclay , and clay of about 20–30 nm between them. The clay value is one order of magnitude lower than the value of Lclay , suggesting the formation of spatially linked like structures of the dispersed clay particles in Nylon 6 matrix. For N6CN3.7 melt, the silicate layers are densely dispersed into the matrix and hence difficult to align under elongational flow. Under flow fields, the silicate layers might translationally move, but not rotationally in such a way that the loss energy becomes minimum. This tendency was also observed in PPCN7.5 melt having a higher content of MMT (=75 wt%) [365]. On the other hand, one can observe two features for the shear viscosity curve. First, the extended Trouton rule, ˙ t  E ˙0 2 t, does not hold for both N6CN3.7 and 30 2 PPCN4 melts, as opposed to the melt of ordinary homopolymers. The latter, E ˙0 2 t, is more than 10 times larger ˙ t. Second, again unlike ordinary than the former, 30 2 ˙ t of N6CN3.7 melt increases continpolymer melts, 30 2 uously with t, never showing a tendency of reaching a steady state within the time span (600 s or longer) examined here. This time-dependent thickening behavior may be called antithixotropy or rheopexy. Via slow shear flow (˙ = 0001 s−1 , ˙ t of N6CN3.7 exhibits a much stronger rheopexy 30 2 behavior almost two orders of magnitude higher than that of PPCN4. This reflects a fact that the shear-induced structural change involved a process with an extremely long relaxation time as well as for other PCNs having rheopexy behavior [325
327], especially under weak shear field.

8.4. Alignment of Silicate Layers The orientation of silicate layers and Nylon 6 crystallites in injection molded N6CN3.7 using WAXD and TEM is examined [111]. Kojima et al. have found three regions of different orientations in the sample as a function of depth. Near the middle of the sample, where the shear forces are minimal, the silicate layers are oriented randomly and the Nylon 6 crystallites are perpendicular to the silicate layers. In the surface region, shear stresses are very high, so both the clay layers and the Nylon 6 crystallites are parallel to the surface. In the intermediate region, the clay layers, presumably due to their higher aspect ratio, still orient parallel to the surface and the Nylon 6 crystallites assume an orientation perpendicular to the silicate. Very recently, Medellin-Rodriguez et al. [117] reported that the molten N6CN samples showed planar orientation of silicate layers along the flow direction, which is strongly dependent

on shear time as well as clay loading, reaching a maximally orienting level after being sheared for 15 min with ˙ = 60 s−1 . Okamoto et al. conducted TEM observations for the sheared N6CN3.7 with ˙ = 00006 s−1 for 1000 s [368]. The edges of the silicate layers laying along the z-axis [marked with the arrows (a)] or parallel alignment of the silicate edges to the shear direction (x-axis) [marked with the arrows (b)] rather than assuming random orientation in the Nylon 6 matrix is observed, but in fact, one cannot see these faces in this plane (Fig. 68). Here, it should be emphasized that the planar orientation of the silicate faces along the x–z plane does not take place prominently. For the case of rapid shear flow, the commonly applicable conjecture of the planar orientation of the silicate faces along the shear direction was first demonstrated to be true by Kojima et al. [111]. In uniaxial elongational flow (converging low) for a PPCN4, the formation of a house-of-cards structure is found by TEM analysis [256]. The perpendicular (but not parallel) alignment of disklike clay particles with large anisotropy toward the flow direction might sound unlikely but this could be the case especially under an elongational flow field, in which the extensional flow rate is the square of the converging flow rate along the thickness direction, if the assumption of affine deformation without volume change is valid. Obviously under such conditions, the energy dissipation rate due

y (shear gradient)

x (flow)

x

y

O z (neutral)

z

(a)

(b)

100nm Figure 68. TEM micrograph in the x–z plane showing N6CN3.7 sheared at 225
C with ˙ = 00006 s−1 for 1000 s. The x-, y-, and z-axes correspond respectively to flow, shear gradient, and neutral direction.

Polymer/Clay Nanocomposites

to viscous resistance between the disk surface and the matrix polymer is minimal, when the disks are aligned perpendicular to the flow direction. Moreover, Lele et al. [265] recently reported the in-situ rheo-X-ray investigation of flow-induced orientation in syndiotactic PP/layered silicate nanocomposite melt. Some 20 years ago van Olphen [369] pointed out that the electrostatic attraction between the layers of natural clay in aqueous suspension arises from higher polar force in the medium. The intriguing features such as yield stress thixotropy and/or rheopexy exhibited in aqueous suspensions of natural clay minerals may be taken as a reference to the present PCNs.

8.5. Electrorheology Electrorheological fluids (ERFs), sometimes referred to as “smart fluids,” are suspensions consisting of polarizable particles dispersed in insulating media [370]. A mismatch in conductivity or dielectric constant between the dispersed particle and the continuous medium phase induces polarization upon application of an electric field. The induced particle dipoles under the action of an electric field tend to attract neighboring particles and cause the particles to form fibril-like structures, which are aligned to the electric field direction [371]. Among various materials [372–374], semiconducting polymer is one of the novel intrinsic ER systems since it has the advantage of a wide range of working temperature, reduced abrasion of device, low cost, and relatively low current density. As a result, development of a high-performance ER fluid followed by conducting polymer optimization and tuning has been the subject of considerable interest for practical applications as a new electomechanic interface. Nevertheless, the yield stress and modulus of ER fluids are lower than those of magnetorheological fluids. Thus the performance of conducting polymer-based ER fluids is still insufficient for the successful development of specific application devices. On this basis of this information, Kim et al. [286] first introduced nanocomposite as ERFs using PANI/clay nanocomposites with intercalated structure. Though PANI/clay intercalated nanocomposites are a new material for application of ER materials, yield stresses of the system showed less than 100 Pa at 1.2 kV/mm (20 wt% suspensions). This value is a little lower than the yield stress of a pure PANI particle system [374]. In other words, no synergistic effect of clay on yield stress was shown. Recently, Park et al. [375] observed remarkable enhancement of yield stress for electrorheological fluids in PANIbased nanocomposites of clay. In further study [376], they fabricated three kinds of ERFs containing different contents of PANI/clay nanocomposite and pure PANI particles in order to investigate the effect of nanocomposite particles on the enhancement of yield stress more systematically. They observed that there is an optimum content ratio between nanocomposite and pure PANI particles to produce minimum yield stress. Details regarding data collection and explanations are presented in [376].

835

9. PROCESSING OPERATIONS OF POLYMER/CLAY NANOCOMPOSITES The flow-induced internal structural change occurs in both shear and elongational flow, but they almost differ from each other, as judged from the previous results on E ˙0 2 t and ˙ t (see Fig. 67). Thus, with these rheological features 30 2 of the PCNs and the characteristics of each processing operation, which process type should be selected for a particular nanocomposite for the enhancement of its mechanical properties? For example, the strong strain-induced hardening in E ˙0 2 t is requisite for withstanding the stretching force ˙ t sugduring the processing, while the rheopexy in 30 2 gests that for such PCN a promising technology is processing in confined space such injection molding where shear force is crucial.

9.1. Foam Processing Using Supercritical CO2 Via batch processing in an autoclave, the foam processing on PPCNs having different contents of clay by using supercritical CO2 as foaming agent under 10 MPa at various temperatures is reported [257]. Figure 69 shows the typical results of scanning electron microscope (SEM) images of the fracture surfaces of the PPCN4 and PP-MA without clay. Both foams exhibit the polygon closed-cell structures having pentagonal

(a)

(b)

Figure 69. SEM micrographs for (a) PPCN4 and (b) PP-MA foamed at 134.7
C. Reprinted with permission from [257], M. Okamoto et al., Nano Lett. 1, 503 (2001). © 2001, American Chemical Society.

Polymer/Clay Nanocomposites

836 and hexagonal faces, which express the most energetically stable state of polygon cells. The morphological parameters of the cells are listed in Table 20. The function for determining cell density Nc is defined as [377] 1− Nc  The mean wall thickness

5f 5p

(4)

10−4 d 3

is given by


=d 

1 1−

5f 5p



−1

(5)

where 5p
5f , and d are the density of the polymer (prefoamed materials), the density of the foam (postfoamed samples) in g/cm3 , and the mean cell size in mm, respectively. The PPCN4 foam shows smaller d and larger Nc compared to PP-MA foam, suggesting that the dispersed clay particles act as nucleating sites for cell formation and lowering of d with increasing clay content [262]. The final 5f is controlled by the competitive process in the cell nucleation, its growth, and coalescence. Cell nucleation, in a heterogeneous nucleation system such as PPCN4 (MMT = 4 wt%) foam, took place in the boundary between the matrix PP-MA and the dispersed clay particles. The cell growth and coalescence are strongly affected by the modulus and the loss modulus (viscosity component) of the materials during processing. As mention in Section 8.3, the strain-induced hardening behavior is probably strong enough to increase the extensional viscosity under biaxial flow and to protect the cell from its breakage at high temperature. Therefore, the straininduced hardening leads to the high cell wall thickness in the PPCN4 foam. The alignment of the dispersed clay particles under biaxial flow in the foam processing is also observed (Fig. 70). Due to the biaxial flow of material during the foam process, the clay particles either turned their face [marked with the arrows (A) in Fig. 70] or had fixed face orientation [marked with the arrows (B) in Fig. 70] and aligned along the flow direction of materials (i.e., along the cell boundary). The interesting point here is that such aligning behavior of the clay particles may help cells to withstand the stretching force from breaking the thin cell wall, in other words, to improve the strength of foam in mechanical properties. The clay particles seem to act as a secondary cloth layer to protect the cells from being destroyed by external forces. The compression modulus K ′ of the PPCN foams appears higher than Table 20. Morphological parameters and compression modulus of PP-MA and PPCN foams. Foam samples

5f (g/cm−3 )

d (m)

NC × 10−6 (cell/cm−3 )

(m)

K′a (MPa)

PP-MA PPCN2 PPCN4 PPCN7.5

006 006 012 013

1553 1330 934 339

249 394 964 220

56 46 119 27

044 172 195 280

a

cell boundary

At 25
C. Source: Reprinted with permission from [257], M. Okamoto et al., Nano Lett. 1, 503 (2001). © 2001, American Chemical Society.

200nm cell boundary

Figure 70. TEM micrographs for PPCN4 foamed at 134.7
C (monocell wall). Reprinted with permission from [257], M. Okamoto et al., Nano Lett. 1, 503 (2001). © 2001, American Chemical Society.

that of PP-MA foam even though its at same 5f level (see Table 20). This may create the improvement of mechanical properties for polymeric foams through PCNs. Recently, some literature has become available [378] related to the reactive extrusion foaming of various nanocomposites.

9.2. Slow Shear Flow Processing Very slow shear processing for the Nylon 6/clay system (N6CN3.7) (MMT = 37 wt%) is examined with the expectation that it would result in an excellent material having enhanced mechanical properties of the PCN [368]. As anticipated, N6CN3.7 subjected to shear processing shows strong enhancement in relative modulus as compared to the corresponding Nylon 6 matrix (Table 21). For example, after having been sheared with ˙ = 00006 s−1 for 1000 s, the modulus of N6CN3.7 at 30
C was 2.6 times higher, while for the presheared N6CN3.7 the modulus was only 1.8 times higher than that of neat Nylon 6. To improve the modulus, the silicate layers seem to act as an internal bone layer to protect the materials from being bent by external forces.

9.3. Electrospinning Processing Fibers and nanofibers of N6CN (diameter of 100– 500 nm) were electrospun from HFIP solution and collected as nonwoven fabrics or as aligned yarns [379]. The Table 21. Bending modulus for N6CN3.7 and neat Nylon 6. Sample T a
C 30 150 a

Modulus of N6CN3.7 (GPa)

Modulus of Nylon 6 (GPa)

Preshear

Postshear

Preshear

Postshear

269 176

400 222

151 046

154 052

Measuring temperature.

Polymer/Clay Nanocomposites

837

electrospinning process resulted in highly aligned MMT particles and Nylon 6 crystallites. The cylindrical shaped fibers, nanofibers, and ribbon shaped fibers were also found in the products (Fig. 71). The electrospinning can be expected to align other nanofillers such as carbon nanotubes.

9.4. Porous Ceramic Materials via PCNs Very recently, a new route for the preparation of porous ceramic material from thermosetting epoxy/clay nanocomposite was first demonstrated by Brown et al. [380]. This route offers an attractive potential for diversification and application of the PCNs. Sinha Ray and co-workers have reported the results on the novel porous ceramic material via burning of the PLA/clay system (PLACN) [315]. The SEM image of the fractured surface of porous ceramic material prepared from simple burning of the PLACN in a furnace up to 950
C is shown in Figure 72. After complete burning, as seen in the figure, the PLACN becomes a white mass with a porous structure. The bright lines in the SEM image correspond to the edge of the stacked silicate layers. In the porous ceramic material, the silicate layers form a house-of-cards structure, which consists of large plates having lengths of ∼1000 nm and thicknesses of ∼30–60 nm. This implies that the further stacked platelet structure is formed during burning. The material exhibits the open-cell-type structure having 100–1000 nm diameter voids, BET surface area of 31 m2 g−1 , and low density of porous material of 0.187 g ml−1 estimated by the buoyancy method. The BET surface area value of MMT is 780 m2 /g and that of the porous ceramic material is 31 m2 /g, suggesting about 25 MMT plates stacked together. When MMT is heated above 700
C (but below 960
C) first all OH groups are eliminated from the structure and thus MMT is decomposed into that of a nonhydrated aluminosilicate. This transformation radically disturbs the crystalline network of the MMT, and the resulting diffraction pattern is indeed often typical of an amorphous (or noncrystalline) phase.

Figure 71. TEM micrograph of a ribbon shaped nanofiber. Reprinted with permission from [379], P. H. Fong et al., Polymer 43, 775 (2002). © 2002, Elsevier Science.

Figure 72. SEM image of porous ceramic material after coated with platinum layer (∼10 nm thickness). Reprinted with permission from [315], S. Sinha Ray et al., Nano Lett. 2, 423 (2002). © 2002, American Chemical Society.

The estimated rough value of compression modulus K is in the order of ∼12 MPa, which is five orders of magnitude lower than the bulk modulus of MMT (∼102 GPa) [33]. In the stress–strain curve, the linear deformation behavior is nicely described in the early stage of the deformation (i.e., the deformation of the material closely resembles that of ordinary polymeric foams) [381]. This open-cell-type porous ceramic material consisting of the house-of-cards structure is expected to provide strain recovery and excellent energy dissipation mechanism after unloading in the elastic region up to 8% strain; probably each plate bends like a leaf spring. This porous ceramic material is an elastic and very lightweight new material. This new route for the preparation of porous ceramic material via burning of nanocomposites can be expected to pave the way for a much broader range of applications of the PCNs.

10. CONCLUSIONS Development of the PCNs is one of the latest evolutionary steps of polymer technology. PCNs offer attractive potential for diversification and application of conventional polymeric materials. Since the possibility of direct melt intercalation first demonstrated by Vaia et al. [31], the melt intercalation method has become the mainstream preparation for intercalated polymer nanocomposites without in-situ intercalative polymerization. It is a quite effective technology for the case of the PCN industry. Some of PCNs are already commercially available and applied in industrial products. Biodegradable polymer based PCNs seem to have a very bright future for a wide range of applications such as high performance biodegradable materials. Undoubtedly, the unique properties originating from the controlled nanostructure pave the way for a much broader range of applications. Although a significant amount of work has already been done on various aspects of PCNs, a lot of research still remains to be carried out in order to understand the structure–property relationship in various PCNs. On the other hand, we have to conduct rheological measurements of various PCNs under molten states in detail, in order to

838 discover the processing conditions of these materials, and this is the final goal of any polymeric material. Finally, PCNs show concurrent improvement in various materials properties at very low clay content together with the ease of preparation through simple processes such as melt intercalation, directly by melt extrusion, or injection molding, opening a new dimension for plastics and composites.

GLOSSARY Clay General family of 2:1 layered- or phyllosilicates. Their crystal structure consists of layers made up of two silica tetrahedral fused to an edge-shared octahedral sheet of either aluminium or magnesium hydroxide. The layer thickness is around 1 nm and the lateral dimensions of these layers may vary from 30 nm to several microns and even larger depending on the particular layered silicate. Exfoliation Extensive polymer penetration resulting in disordered and eventual delamination of the silicate layers produces near to exfoliated nanocomposites consisting of individual silicate layers dispersed in polymer matrix. Flocculation Conceptually this is same with intercalated nanocomposites, however, silicate layers are sometimes flocculated due to hydroxylated edge–edge interaction of the silicate layers (intercalated and flocculated). Green polymeric materials These would not involve the use of toxic or noxious components in their manufacture, and could be degraded in natural environment or easily recycled. Aliphatic polyesters are among the most promising materials for the production of environmentally friendly biodegradable plastics. Intercalation Polymer penetration resulting in finite expansion of the silicate layers produces intercalated nanocomposites consisting of well-ordered multilayers with alternating polymer/silicate layers and a repeat distance of few nanometers. Montmorillonite One of the most commonly used layered silicates (M1/3 (Al5/3 Mg1/3 )Si4 O10 (OH)2 ). This clay is only miscible with hydrophilic polymers, such as poly (ethylene oxide) and poly (vinyl alcohol). Nanocomposite Phase mixing of at least two dissimilar materials (e.g., polymer and inorganic filler) occurs on a nanometer scale. The nanocomposite exhibits remarkable improvement of materials properties. Organo-clay To improve miscibility with other polymer matrices, one must convert the normally hydrophilic silicate surface to organophilic, which makes possible intercalation of many engineering polymers. Generally, this can be done by ion-exchange reactions with cationic surfactants including primary, secondary, tertiary, and quaternary alkyl ammonium or alkylphosphonium cations.

REFERENCES 1. A. Usuki, Y. Kojima, A. Okada, Y. Fukushima, T. Kurauchi, and O. Kamigaito, J. Mater. Res. 8, 1174 (1993). 2. Y. Kojima, A. Usuki, M. Kawasumi, A. Okada, Y. Fukushima, T. Karauchi, and O. Kamigaito, J. Mater. Res. 8, 1185 (1993).

Polymer/Clay Nanocomposites 3. P. C. LeBaron, Z. Wang, and T. Pinnavaia, J. Appl. Clay Sci. 15, 11 (1999). 4. E. Manias, A. Touny, L. Wu, K. Strawhecker, B. Lu, and T. C. Chung, Chem. Mater. 13, 3561 (2001). 5. T. Lan, D. K. Padmananda, and T. J. Pinnavaia, Chem. Mater. 7, 2144 (1995). 6. Y. I. Tien and K. H. Wei, Polymer 42, 3213 (2001). 7. J. C. Huang, Z. Zhu, J. Yin, X. Qian, and Y. Sun, Polymer 42, 873 (2001). 8. T. Agag and T. Takeichi, Polymer 41, 7083 (2000). 9. G. Galgali, C. Ramesh, and A. Lele, Macromolecules 34, 852 (2001). 10. X. Fu and S. Qutubuddin Polymer 42, 807 (2001). 11. M. Okamoto, S. Morita, H. Taguchi, Y. H. Kim, T. Kotaka, and H. Tateyama, Polymer 41, 3887 (2000). 12. M. Okamoto, S. Morita, and T. Kotaka, Polymer 42, 2685 (2001). 13. P. Messersmith and E. Giannelis, J. Polym. Sci., Polym. Chem. 33, 1047 (1995). 14. M. Kawasumi, N. Hasegawa, M. Kato, A. Usuki, and A. Okada, Macromolecules 30, 6333 (1997). 15. N. Hasegawa, M. Kawasumi, M. Kato, A. Usuki, and A. Okada, J. Appl. Polym. Sci. 67, 87 (1998). 16. National Lead Co., U.S. Patent 2
531
396, 1950. 17. Union Oil Co., U.S. Patent 3
084
117, 1963. 18. Unitika Ltd., Japanese Kokai Patent Application 109998, 1976. 19. Y. Fukushima and S. Inagaki, J. Incl. Phen. 5, 473 (1987). 20. P. Aranda and E. Ruiz-Hitzky, Chem. Mater. 4, 1395 (1992). 21. D. J. Greenland, J. Colloid Sci. 18, 647 (1963). 22. A. Blumstein, J. Polym. Sci. A 3, 2653 (1965). 23. Krishnamoorti, R. A. Vaia, and E. P. Giannelis, Chem. Mater. 8, 1728 (1996). 24. R. A. Vaia, R. K. Teukolsky, and E. P. Giannelis, Chem. Mater. 6, 1017 (1994). 25. M. J. Solomon, A. S. Almusallam, K. F. Seefeld, S. Somwangthanaroj, and P. Varadan, Macromolecules 34, 1864 (2001). 26. R. A. Vaia and E. P. Giannelis, Polymer 42, 1281 (2001). 27. J. Billingham, C. Breen, and Yarwood, J. Vibr. Spectrosc. 14, 19 (1997). 28. C. O. Oriakhi, I. V. Farr, and M. M. Lerner, Clay Clay Minerals 45, 194 (1997). 29. M. Alexandre and P. Dubois, Mater. Sci. Eng. R 28, 1 (2000). 30. P. H. Nam, P. Maiti, M. Okamoto, T. Kotaka, N. Hasegawa, and A. Usuki, Polymer 42, 9633 (2001). 31. R. A. Vaia, H. Ishii, and E. P. Giannelis, Chem. Mater. 5, 1694 (1993). 32. E. P. Giannelis, R. Krishnamoorti, and E. Manias, Adv. Polym. Sci. 138, 107 (1999). 33. M. Okamoto, S. Morita, Y. H. Kim, T. Kotaka, and H. Tateyama, Polymer 42, 1201 (2001). 34. S. Sinha Ray, P. Maiti, M. Okamoto, K. Yamada, and K. Ueda, Macromolecules 35, 3104 (2002). 35. A. Blumstein, S. L. Malhotra, and A. C. Watterson, J. Polym. Sci. A 8, 1599 (1970). 36. D. C. Lee and L. W. Jang, J. Appl. Polym. Sci. 61, 1117 (1996). 37. G. Chen, X. Chen, Z. Lin, W. Ye, and K. Yao, J. Mater. Sci. Lett. 18, 1761 (1999). 38. G. Chen, K. Yao, and J. Zhao, J. Appl. Polym. Sci. 73, 425 (1999). 39. A. Tabtiang, S. Lumlong, and R. A. Venables, Eur. Polym. J. 36, 2559 (2000). 40. A. Tabtiang, S. Lumlong, and R. A. Venables, Polym. Plast. Technol. Eng. 39, 293 (2000). 41. S. Bandyopadhyay and E. P. Giallelis, Polym. Mater. Sci. Eng. 82, 208 (2000). 42. X. Huang and W. J. Brittain, Macromolecules 34, 3255 (2001). 43. C. Zeng and L. J. Lee, Macromolecules 34, 4098 (2001). 44. N. Salahuddin and M. Shehata, Polymer 42, 8379 (2001).

Polymer/Clay Nanocomposites 45. C. Forte, M. Geppi, S. Giamberini, G. Ruggeri, C. A. Veracini, and B. Mendez, Polymer 39, 2651 (1998). 46. F. Dietsche, Y. Thomann, R. Thomann, and R. Mulhaupt, J. Appl. Polym. Sci. 75, 369 (2000). 47. T. Seckin, Y. Onal, I. Aksoy, and M. E. Yakinci, J. Mater. Sci. 31, 3123 (1996). 48. F. Dietsche and R. Mulhaupt, Polym. Bull. 43, 395 (1999). 49. Z. Chen, C. Huang, S. Liu, Y. Zhang, and K. Gong, J. Appl. Polym. Sci. 75, 796 (2000). 50. J. Billingham, C. Breen, and Yarwood, J. Vib. Spectrosc. 14, 19 (1997). 51. J. Lin, J. Wu, Z. Yang, and M. Pu, Macromol. Rapid Commun. 22, 422 (2001). 52. R. Blumstein, A. Blumstein, and K. K. Parikh, Appl. Polym. Symp. 25, 81 (1994). 53. Y. Sugahara, S. Satakawa, K. Kuroda, and C. Kato, Clays Clays Miner. 36, 343 (1988). 54. F. Bergaya and F. Kooli, Clay Miner. 26, 33 (1991). 55. Y. S. Choi, K. H. Wang, M. Xu, and I. J. Chung, Chem. Mater. 14, 2936 (2002). 56. C. Kato, K. Kuroda, and H. Takahara, Clays Clay Miner. 29, 294 (1981). 57. A. Akelah, in “Polymers and Other Advanced Materials: Emerging Technologies and Business Opportunities” (P. N. Prasad, J. E. Mark, and F. J. Ting, Eds.), p. 625. Plenum, New York, 1995. 58. R. A. Vaia, K. D. Jant, E. J. Kramer, and E. P. Giannelis, Chem. Mater. 8, 2628 (1996). 59. R. Krishnamoorti, R. A. Vaia, and E. P. Giannelis, Chem. Mater. 8, 1728 (1996). 60. A. Akelah and M. Moet, J. Mater. Sci. 31, 3589 (1996). 61. M. Sikka, L. N. Cerini, S. S. Ghosh, and K. I. Winey, J. Polym. Sci. B 34, 1443 (1996). 62. M. Laus, M. Camerani, M. Lelli, K. Sparnacci, F. Sandrolini, and O. F. Francescangeli, J. Mater. Sci. 33, 2883 (1998). 63. J. G. Doh and I. Cho, Polym. Bull. 41, 511 (1998). 64. T. L. Porter, M. E. Hagerman, B. P. Reynolds, and M. E. Eastman, J. Polym. Sci. B 36, 673 (1998). 65. N. Hasegawa, H. Okamoto, M. Kawasumi, and A. Usuki, J. Appl. Polym. Sci. 74, 3359 (1999). 66. M. W. Noh and D. C. Lee, Polym. Bull. 42, 619 (1999). 67. M. W. Weimer, H. Chen, E. P. Giannelis, and D. Y. Sogah, J. Am. Chem. Soc. 121, 1615 (1999). 68. X. Fu and S. Qutubuddin, Polymer 42, 807 (2000). 69. G. Chen, S. Liu, S. Zhang, and Z. Qi, Macromol. Rapid Commun. 21, 746 (2000). 70. Y. T. Lim and O. O. Park, Macromol. Rapid Commun. 21, 231 (2000). 71. B. Hoffman, C. Dietrich, R. Thomann, C. Friedrich, and R. Mulhaupt, Macromol. Rapid Commun. 21, 57 (2000). 72. C. Zilg, R. Thomann, M. Baumert, J. Finter, and R. Mulhaupt, Macromol. Rapid Commun. 21, 1214 (2000). 73. H. D. Wu, C. R. Tseng, and F. C. Chang, Macromolecules 34, 2992 (2001). 74. P. Xiao, M. Xiao, and K. Gong, Polymer 42, 4813 (2001). 75. C.-R. Tseng, J.-Y. Wu, H.-Y. Lee, and F.-C. Chang, Polymer 42, 10063 (2001). 76. J. Zhu, A. B. Morgan, F. J. Lamelas, and C. A. Wilkie, Chem. Mater. 13, 3774 (2001). 77. Q. H. Zeng, D. Z. Wang, A. B. Yu, and G. Q. Lu, Nanotechnology 13, 549 (2002). 78. F. L. Beyer, N. C. B. Tan, A. Dasgupta, and M. E. Galvin, Chem. Mater. 14, 2983 (2002). 79. J. W. Gilman, W. H. Awad, R. D. Davis, J. Shields, R. H. Harris, Jr., C. Davis, A. B. Morgan, T. E. Sutto, J. Callahan, P. C. Trulove, and H. C. DeLong, Chem. Mater. 14, 3776 (2002). 80. M. Y. Gelfer, H. S. Hyun, L. Liu, B. S. Haiao, B. Chu, M. Rafailovich, M. Si, and V. Zaitsev, J. Polym. Sci. B 41, 44 (2003).

839 81. H. Z. Friedlander and C. R. Frink, J. Polym. Sci. B 2, 457 (1964). 82. N. A. Churochkina, S. G. Starodoubtsev, and A. R. Khokhlov, Polym. Gels Networks 6, 205 (1998). 83. D. Gao, R. B. Heimann, M. C. Williams, L. T. Wardhaugh, and M. Muhammad, J. Mater. Sci. 34, 1543 (1999). 84. A. Wheeler, U.S. Patent 2
847
391, 1958. 85. N. Ogata, S. Kawakage, and T. Ogihara, J. Appl. Polym. Sci. 66, 573 (1997). 86. H. Matsuyama and J. F. Young, Chem. Mater. 11, 16 (1999). 87. K. E. Strawhecker and E. Manias, Chem. Mater. 12, 2943 (2000). 88. C. W. Francis, Soil Sci. 115, 40 (1973). 89. K. A. Carrado and L. Xu, Chem. Mater. 10, 1440 (1998). 90. Y. Komori, Y. Sugahara, and K. Kuroda, Chem. Mater. 11, 3 (1999). 91. R. Levy and C. W. Francis, J. Colloid Interface Sci. 50, 442 (1975). 92. A. Nisha, M. K. Rajeswari, and R. Dhamodharan, J. Appl. Polym. Sci. 76, 1825 (2000). 93. C. M. Koo, H. T. Ham, M. H. Choi, S. O. Kim, and I. J. Chung, Polymer 44, 681 (2003). 94. K. G. Fournaris, M. A. Karakassides, and D. Petridis, Chem. Mater. 11, 2372 (1999). 95. R. L. Parfitt and D. J. Greenland, Clay Miner. 8, 305 (1970). 96. X. Zhao, K. Urano, and S. Ogasawara, Colloid Polym. Sci. 267, 899 (1989). 97. L. Priya and J. P. Jog, J. Polym. Sci. B 41, 31 (2003). 98. C. O. Oriakhi, X. Zhang, and M. M. Lerner, Appl. Clay Sci. 15, 109 (1999). 99. N. Nugay, Kusefoglus, and B. Erman, J. Appl. Polym. Sci. 66, 1943 (1997). 100. S. W. Kim, W. H. Jo, M. S. Lee, M. B. Ko, and J. Y. Jho, Polymer 42, 9837 (2001). 101. N. Artzi, Y. Nir, M. Nakris, and A. Seigmann, J. Polym. Sci. B 40, 1741 (2002). 102. J. Ren, A. S. Silva, and R. Krishnamoorti, Macromolecules 33, 3739 (2000). 103. C. A. Mitchell and R. Krishnamoorti, J. Polym. Sci. B 40, 1434 (2002). 104. N. Schamp and Huylebroeck, J. Polym. Sci. Polym. Symp. 42, 553 (1973). 105. A. Usuki, Y. Kojima, M. Kawasumi, A. Okada, Y. Fukushima, T. Kurauchi, and O. Kamigaito, J. Mater. Res. 8, 1179 (1993). 106. Y. Kojima, A. Usuki, M. Kawasumi, A. Okada, T. Kurauchi, and O. Kamigaito, J. Polym. Sci. A 31, 983 (1993). 107. Y. Kojima, A. Usuki, M. Kawasumi, A. Okada, T. Kurauchi, and O. Kamigaito, J. Polym. Sci. A 31, 1755 (1993). 108. Y. Kojima, A. Usuki, M. Kawasumi, A. Okada, T. Kurauchi, and O. Kamigaito, J. Appl. Polym. Sci. 49, 1259 (1993). 109. A. Okada and A. Usuki, Mater. Sci. Eng. C 3, 109 (1995). 110. Y. Kojima, A. Usuki, M. Kawasumi, A. Okada, T. Kurauchi, O. Kamigaito, and K. Kaji, J. Polym. Sci. B 32, 625 (1994). 111. Y. Kojima, A. Usuki, M. Kawasumi, A. Okada, T. Kurauchi, O. Kamigaito, and K. Kaji, J. Polym. Sci. B 33, 1039 (1995). 112. A. Usuki, A. Koiwai, Y. Kojima, M. Kawasumi, A. Okada, T. Kurauchi, and O. Kamigaito, J. Appl. Polym. Sci. 55, 119 (1995). 113. “Polymer–Clay Nanocomposites” (T. J. Pinnavaia and G. W. Beall, Eds.). Wiley, New York, 2000. 114. L. M. Liu, Z. N. Qi, and X. G. Zhu, J. Appl. Polym. Sci. 71, 1133 (1999). 115. S. H. Wu, F. Y. Wang, C.-C. M. Ma, W. C. Chang, C. T. Kuo, H. C. Kuan, and W. J. Chen, Mater. Lett. 49, 327 (2001). 116. D. M. Lincoln, R. A. Vaia, Z. G. Wang, and B. S. Hsiao, Polymer 42, 1621 (2001). 117. F. J. Medellin-Rodriguez, C. Burger, B. S. Hsiao, B. Chu, R. A. Vaia, and S. Phillips, Polymer 42, 9015 (2001). 118. H. R. Dennis, D. L. Hunter, D. Chang, S. Kim, J. L. White, J. W. Cho, and D. R. Paul, Polymer 42, 9513 (2001).

840 119. T. D. Fornes, P. J. Yoon, H. Keskkula, and D. R. Paul, Polymer 42, 9929 (2001). 120. J. S. Shelley, P. T. Mather, and DeVries, Polymer 42, 5849 (2001). 121. D. L. VanderHart, A. Asano, and J. W. Gilman, Chem. Mater. 13, 3781 (2001). 122. Q. Wu and X. Liu, Polymer 43, 1933 (2002). 123. T. D. Fornes, P. J. Yoon, D. L. Hunter, H. Keskkula, and D. R. Paul, Polymer 43, 5915 (2002). 124. M. R. Kamal, N. K. Borse, and A. Garcia-Rejon, Polym. Eng. Sci. 42, 1883 (2002). 125. M. N. Bureau, J. Denault, K. C. Cole, and G. D. Enright, Polym. Eng. Sci. 42, 1897 (2002). 126. Y. Ding, D. J. Jones, P. Maireles-Torres, and J. Roziere, Chem. Mater. 7, 562 (1995). 127. P. Reichert, J. Kressler, R. Thomann, R. Mulhaupt, and G. Stoppelmann, Acta Polym. 49, 116 (1998). 128. B. Hoffman, J. Kressler, G. Stoppelmann, C. Friedrich, and G. M. Kim, Colloid Polym. Sci. 278, 629 (2000). 129. E. Giza, H. Ito, T. Kikutani, and N. Okui, J. Polym. Eng. 20, 403 (2000). 130. G. M. Kim, D. H. Lee, B. Hoffmann, J. Kresler, and G. Stoppelmann, Polymer 42, 1095 (2001). 131. S. V. Nair, L. A. Goettler, and B. A. Lysek, Polym. Eng. Sci. 42, 1872 (2002). 132. X. Liu, Q. Wu, Q. Zhang, and Z. Mo, J. Polym. Sci. B 44, 63 (2003). 133. P. B. Messersmith and E. P. Giannelis, Chem. Mater. 5, 1064 (1993). 134. G. Jimenez, N. Ogata, H. Kawai, and T. Ogihara, J. Appl. Polym. Sci. 64, 2211 (1997). 135. R. Krishnamoorti and E. P. Giannelis, Macromolecules 30, 4097 (1997). 136. G. Jimenez, N. Ogata, H. Kawai, and T. Ogihara, J. Appl. Polym. Sci. 64, 2211 (1997). 137. N. Pantoustier, M. Alexandre, P. Degee, C. Calberg, R. Jerome, C. Henrist, R. Cloots, A. Rulmont, and P. Dubois, Polymer 9, 1 (2001). 138. N. Pantoustier, B. Lepoittevin, M. Alexandre, D. Kubies, C. Calberg, R. Jerome, and P. Dubois, Polym. Eng. Sci. 42, 1928 (2002). 139. B. Lepoittevin, N. Pantoustier, M. Devalckenaere, M. Allexandre, D. Kubies, C. Calderg, R. Jerome, and P. Dubois, Macromolecules 35, 8385 (2002). 140. B. Lepoittevin, M. Devalckenaere, N. Pantoustier, M. Alexandre, D. Kubies, C. Calberg, R. Jerome, and P. Dubois, Polymer 43, 1111 (2002). 141. Y. Ke, C. Long, and Z. Qi, J. Appl. Polym. Sci. 71, 1139 (1999). 142. D. J. Sekelik, S. Stepanov Enazarenko, D. Schiraldi, A. Hiltner, and E. Baer, J. Polym. Sci. B 37, 847 (1999). 143. J. C. Matayabas, Jr., S. R. Turner, B. J. Sublett, G. W. Connell, and R. B. Barbee, (Eastman Chemical Co.), PCT Int. Appl. Wo 98/29499, 1998. 144. T. Takekoshi, F. F. Khouri, J. R. Campbell, T. C. Jordan, and K. H. Dai, (General Electric Co.), U.S. Patent 5
530
052, 1996. 145. M. Okamoto, Y. Shinoda, T. Okuyama, A. Yamaguchi, and T. Sekura, J. Mater. Sci. Lett. 15, 1178 (1996). 146. C. H. Davis, L. J. Mathias, J. W. Gilman, D. A. Schiraldi, J. R. Shields, P. Trulove, T. E. Sutto, and H. C. Delong, J. Polym. Sci. B 40, 2661 (2002). 147. Y. Imai, S. Nishimura, E. Abe, H. Tateyama, A. Abiko, A. Yamaguchi, T. Aoyama, and H. Taguchi, Chem. Mater. 14, 477 (2002). 148. B. J. Chisholm, R. B. Moore, G. Barber, F. Khouri, A. Hempstead, M. Larsen, E. Olson, J. Kelley, G. Balch, and J. Caraher, Macromolecules 35, 5508 (2002). 149. X. Huang, S. Lewis, W. J. Brittain, and R. A. Vaia, Macromolecules 33, 2000 (2000). 150. M. Mitsunaga, K. Hironaka, and M. Okamoto, Polym. Preprints Japan 51, 2645 (2002); Macromol. Mater. Eng. 288, 543 (2003). 151. J. Wu and M. M. Lerner, Chem. Mater. 5, 835 (1993).

Polymer/Clay Nanocomposites 152. R. A. Vaia, S. Vasudevan, W. Krawiec, L. G. Scanlon, and E. P. Giannelis, Adv. Mater. 7, 154 (1995). 153. S. Wong, S. Vasudevan, R. A. Vaia, E. P. Giannelis, and D. Zax, J. Am. Chem. Soc. 117, 7568 (1995). 154. R. A. Vaia, B. B. Sauer, O. K. Tse, and E. P. Giannelis, J. Polym. Sci. B 35, 59 (1997). 155. H. L. M. Hatharasinghe, M. V. Smalley, J. Swenson, C. D. Willians, R. K. Heenan, and S. M. King, J. Phys. Chem. B 102, 6804 (1998). 156. L. Hernan, J. Morales, and J. Santos, J. Solid State Chem. 141, 327 (1998). 157. D. J. Harris, T. J. Bonagamba, and K. Schmidt-Rhor, Macromolecules 32, 6718 (1999). 158. W. Chen, Q. Xu, and R. Z. Yuan, J. Mater. Sci. Lett. 18, 711 (1999). 159. J. Bujdak, E. Hackett, and E. P. Giannelis, Chem. Mater. 12, 2168 (2000). 160. G. Schmidt, A. I. Nakatani, P. D. Butler, A. Karim, and C. C. Han, Macromolecules 33, 7219 (2000). 161. Y. Xiao, K. A. Hu, Q. C. Yu, and R. J. Wu, J. Appl. Polym. Sci. 80, 2162 (2001). 162. H. Y. Choi, S. G. Kim, Y. H. Hyun, and M. S. Jhon, Macromol. Rapid Commun. 22, 320 (2001). 163. Y. H. Hyun, S. T. Lim, H. Y. Choi, and M. S. Jhon, Macromolecules 34, 8084 (2001). 164. B. Liao, M. Song, H. Liang, and Y. Pang, Polymer 42, 10007 (2001). 165. H. W. Chen and F. C. Chang, Polymer 42, 9763 (2001). 166. H. W. Chen, C. Y. Chiu, H. D. Wu, I. W. Shen, and F. C. Chang, Polymer 43, 5011 (2002). 167. Z. Shen, G. P. Simon, and Y. B. Cheng, Polymer 43, 4251 (2002). 168. S. K. Lim, J. W. Kim, I. Chin, Y. K. Kwon, and H. J. Choi, Chem. Mater. 14, 1989 (2002). 169. H. R. Fischer, L. H. Gielgens, and T. P. M. Koster, Acta Polym. 50, 122 (1999). 170. L. Wei, M. Rocci-Lane, P. Brazis, C. R. Kanneworf, Y. I. Kim, W. Lee, J. H. Choy, and M. G. Kanatzidis, J. Am. Chem. Soc. 122, 6629 (2000). 171. S. D. Burnside and E. P. Giannelis, Chem. Mater. 7, 1597 (1995). 172. S. Wang, C. Long, X. Wang, Q. Li, and Z. Qi, J. Appl. Polym. Sci. 69, 1557 (1998). 173. H. Takeuchi and C. Cohen, Macromolecules 32, 6792 (1999). 174. S. D. Burbside and E. P. Giannelis, J. Polym. Sci. B 38, 1595 (2000). 175. L. Bokobza and N. Nugay, J. Appl. Polym. Sci. 81, 215 (2001). 176. M. A. Osman, A. Atallah, M. Muller, and U. W. Suter, Polymer 42, 6545 (2001). 177. J. H. Chang, B. S. Seo, and D. H. Hwang, Polymer 43, 2969 (2002). 178. S. L. C. Hsu and K. C. Chang, Polymer 43, 4097 (2002). 179. A. Akelah, M. A. El-Borai, M. F. A. El-Aal, A. Rehab, and M. S. Abou-Zeid, Macromol. Chem. Phys. 200, 955 (1999). 180. L. Zhang, Y. Wang, Y. Wang, Y. Sui, and D. Yu, J. Appl. Polym. Sci. 78, 1873 (2000). 181. Y. Wang, L. Zhang, C. Tang, and D. Yu, J. Appl. Polym. Sci. 78, 1879 (2000). 182. A. K. Manna, D. K. Tripathy, P. P. De, S. K. De, M. K. Chatterjee, and D. G. Pfeiffer, J. Appl. Polym. Sci. 72, 1895 (1999). 183. Y. T. Vu, J. E. Mark, L. H. Pham, and M. Engelhardt, J. Appl. Polym. Sci. 82, 1391 (2001). 184. P. B. Messersmith and E. P. Giannelis, Chem. Mater. 6, 1719 (1994). 185. T. Lan and T. J. Pinnavaia, Chem. Mater. 6, 2216 (1994). 186. M. S. Wang and T. J. Pinnavaia, Chem. Mater. 7, 468 (1995). 187. T. Lan, P. D. Kaviratna, and T. J. Pinnavaia, Chem. Mater. 7, 2144 (1995). 188. Z. Wang, T. Lan, and T. J. Pinnavaia, Chem. Mater. 8, 2000 (1996). 189. H. Shi, T. Lan, and T. J. Pinnavaia, Chem. Mater. 8, 1584 (1996). 190. T. J. Pinnavaia, T. Lan, Z. Wang, H. Shi, and P. D. Kaviratna, in “Nanotechnology: Molecularly Designated Materials” (G.-M. Chow and K. E. Gonsalves, Eds.), Vol. 622, p. 250. American Chemical Society, Washington, DC, 1996. 191. Z. Wang and T. J. Pinnavaia, Chem. Mater. 10, 1820 (1998).

Polymer/Clay Nanocomposites 192. J. Massam and T. J. Pinnavaia, in “Chemical and Pyrolytic Routes to Nanostructured Powders and Their Industrial Application” (G. Beaucage, G. Burns, D.-W. Hua, and J. E. Mark, Eds.), Vol. 520, p. 223. Materials Research Society, Warrendale, PA, 1998. 193. D. C. Lee and J. W. Jang, J. Appl. Polym. Sci. 68, 1997 (1998). 194. C. Zilg, R. Mulhaupt, and J. Finter, Macromol. Chem. Phys. 12, 8000 (2000). 195. X. Kornmann, H. Lindberg, and L. A. Berglund, Polymer 42, 1303 (2001). 196. L. Jiankun, K. Yucai, Q. Zongneng, and Y. Xiao-Su, J. Polym. Sci. B 39, 115 (2001). 197. I. J. Chin, T. T- Albrecht, H. C. Kim, and J. Wang, Polymer 42, 5947 (2001). 198. X. Kornmann, H. Lindberg, and L. A. Berglund, Polymer 42, 4493 (2001). 199. A. S. Zerda and A. J. Lesser, J. Polym. Sci. B 39, 1137 (2001). 200. I. J. Chin, T. T. Albrecht, H. C. Kim, T. P. Russell, and J. Wang, Polymer 42, 5947 (2001). 201. X. Kornmann, R. Thomann, R. Mulhaupt, J. Finter, and Berglund, Polym. Eng. Sci. 42, 1815 (2002). 202. W. Feng, A. Ait-Kadi, and B. Rield, Polym. Eng. Sci. 42, 1827 (2002). 203. O. Becker, R. Varley, and G. Simon, Polymer 43, 4365 (2002). 204. J. S. Chen, M. D. Poliks, C. K. Ober, Y. Zhang, U. Wiesner, and E. P. Giannelis, Polymer 43, 4895 (2002). 205. M. H. Choi, I. J. Chung, and J. D. Lee, Chem. Mater. 12, 2977 (2000). 206. Z. Wang and T. J. Pinnavaia, Chem. Mater. 10, 3769 (1998). 207. T. K. Chen, Y. I. Tien, and K. H. Wei, J. Polym. Sci. A 37, 2225 (2000). 208. Y. I. Tien and K. H. Wei, Polymer 42, 3213 (2001). 209. K. J. Yao, M. Song, D. J. Hourston, and D. Z. Luo, Polymer 43, 1017 (2002). 210. R. Xu, E. Manias, A. J. Snyder, and J. Runt, Macromolecules 34, 337 (2001). 211. K. Yano, A. Usuki, A. Okada, T. Kurauchi, and O. Kamigaito, J. Polym. Sci. A 31, 2493 (1993). 212. T. Lan, P. D. Kaviratna, and T. J. Pinnavaia, Chem. Mater. 6, 573 (1994). 213. K. Yano, A. Usuki, and A. Okada, J. Polym. Sci. A 35, 2289 (1997). 214. Z.-K. Zhu, Y. Yang, J. Yin, X. Y. Wang, Y. C. Ke, and Z. N. Qi, J. Appl. Polym. Sci. 73, 2063 (1999). 215. Y. Yang, Z. K. Zhu, J. Yin, X. Y. Wang, and Z. E. Qi, Polymer 40, 4407 (1999). 216. H. L. Tyan, K. H. Wei, and T. E. Hsieh, J. Polym. Sci. B 38, 2873 (2000). 217. J. C. Huang, Z. K. Zhu, J. Yin, X. F. Qian, and Y. Y. Sun, Polymer 42, 873 (2001). 218. A. Gu and F. C. Chang, J. Appl. Polym. Sci. 79, 289 (2001). 219. A. Gu, S. W. Kuo, and F. C. Chang, J. Appl. Polym. Sci. 79, 1902 (2001). 220. S. H. Hsiao, G. S. Liou, and L. M. Chang, J. Appl. Polym. Sci. 80, 2067 (2001). 221. H. L. Tyan, C. M. Leu, and K. H. Wei, Chem. Mater. 13, 222 (2001). 222. J. C. Huang, Z. K. Zhu, X. D. Ma, X. F. Qian, and J. Yin, J. Mater. Sci. 36, 871 (2001). 223. T. Agag, T. Koga, and T. Takeichi, Polymer 42, 3399 (2001). 224. A. B. Morgan, J. W. Gilman, and C. L. Jackson, Macromolecules 34, 2735 (2001). 225. C. M. Leu, Z. W. Wu, and K. H. Wei, Chem. Mater. 14, 3016 (2002). 226. R. Magaraphan, W. Lilayuthalert, Sirivat, and J. W. Schwank, Compos. Sci. Technol. 61, 1253 (2001). 227. H. L. Tyan, Y. C. Liu, and K. H. Wei, Polymer 40, 4877 (1994). 228. J. Kim, R. Ahmed, and S. J. Lee, J. Appl. Polym. Sci. 80, 592 (2001).

841 229. J. H. Chang, D. K. Park, and K. J. Ihn, J. Polym. Sci. B 39, 471 (2001). 230. G. S. Sur, H. L. Sun, S. G. Lyu, and J. E. Mark, Polymer 42, 9783 (2001). 231. J. Lee, T. Takekkoshi, and E. P. Giannelis, Mater. Res. Soc. Symp. Proc. 457, 513 (1997). 232. J. C. Huang, Z. K. Zhu, X. F. Q, and Y. Y. Sun, Polymer 42, 873 (2001). 233. Y. Kurokawa, H. Yasuda, and A. Oya, J. Mater. Sci. Lett. 15, 1481 (1996). 234. N. Furuichi, Y. Kurokawa, K. Fujita, A. Oya, H. Yasuda, and M. Kiso, J. Mater. Sci. 31, 4307 (1996). 235. J. Tudor, L. Willington, D. O’Hare, and B. Royan, Chem. Commun. 2031 (1996). 236. Y. Kurokawa, H. Yasuda, M. Kashiwagi, and A. Oya, J. Mater. Sci. Lett. 16, 1670 (1997). 237. M. R. Nyden and J. W. Gilman, Com. Theor. Polym. Sci. 7, 191 (1997). 238. M. Kato, A. Usuki, and A. Okada, J. Appl. Polym. Sci. 66, 1781 (1997). 239. A. Usuki, M. Kato, A. Okada, and T. Kurauchi, J. Appl. Polym. Sci. 63, 137 (1997). 240. N. Hasewaga, H. Okamoto, M. Kato, and A. Usuki, J. Appl. Polym. Sci. 78, 1918 (2000). 241. A. Oya, Y. Kurokawa, and H. Yasuda, J. Mater. Sci. 35, 1045 (2000). 242. J. W. Lee, Y. T. Lim, and O. O. Park, Polym. Bull. 45, 191 (2000). 243. Q. Zhang, Q. Fu, L. Jiang, and Y. Lei, Polym. Int. 49, 1561 (2000). 244. J. M. Garces, D. J. Moll, J. Bicerano, R. Fibiger, and D. G. McLeod, Adv. Mater. 12, 1835 (2000). 245. N. Hasegawa, H. Okamoto, M. Kawasumi, M. Kato, A. Tsukigase, and A. Usuki, Macromol. Mater. Eng. 280/281, 76 (2000). 246. S. Hambir, N. Bulakh, P. Kodgire, R. Kalgaonkar, and J. P. Jog, J. Polym. Sci. B 39, 446 (2001). 247. M. Zanetti, G. Camino, P. Reichert, and R. Mulhaupt, Macromol. Rapid Commun. 22, 176 (2001). 248. G. Galgali, C. Ramesh, and A. Lele, Macromolecules 34, 852 (2001). 249. J. M. Gloaguen and J. M. Lefebvre, Polymer 42, 5841 (2001). 250. J. M. Garcia-Martinez, O. Laguna, S. Areso, and E. P. Collar, J. Appl. Polym. Sci. 81, 625 (2001). 251. C. I. Park, O. O. Park, J. G. Lim, and H. J. Kim, Polymer 42, 7465 (2001). 252. P. Reichert, B. Hoffman, T. Bock, R. Thomann, R. Mulhaupt, and C. Friedrich, Macromol. Rapid Commun. 22, 519 (2001). 253. X. Liu and Q. Wu, Polymer 42, 10013 (2001). 254. E. Manias, A. Touny, K. E. Strawhecker, B. Lu, and T. C. Chung, Chem. Mater. 13, 3516 (2001). 255. E. Manias, Mater. Res. Soc. Bull. 26, 862 (2001). 256. M. Okamoto, P. H. Nam, P. Maiti, T. Kotaka, N. Hasegawa, and A. Usuki, Nano Lett. 1, 295 (2001). 257. M. Okamoto, P. H. Nam, M. Maiti, T. Kotaka, T. Nakayama, M. Takada, M. Ohshima, A. Usuki, N. Hasegawa, and H. Okamoto, Nano Lett. 1, 503 (2001). 258. S. Hambir, N. Bulakh, P. Kodgire, R. Kalgaonkar, and J. P. Jog, J. Polym. Sci. B 39, 446 (2001). 259. T. Sun and J. M. Garces, Adv. Mater. 14, 128 (2002). 260. P. Maiti, P. H. Nam, M. Okamoto, T. Kotaka, N. Hasegawa, and A. Usuki, Macromolecules 35, 2042 (2002). 261. P. Maiti, P. H. Nam, M. Okamoto, T. Kotaka, N. Hasegawa, and A. Usuki, Polym. Eng. Sci. 42, 1864 (2002). 262. P. H. Nam, P. Maiti, M. Okamoto, T. Kotaka, T. Nakayama, M. Takada, M. Ohshima, A. Usuki, N. Hasegawa, and H. Okamoto, Polym. Eng. Sci. 42, 1907 (2002). 263. S. Hambir, N. Bulakh, and J. P. Jog, Polym. Eng. Sci. 42, 1800 (2002).

842 264. D. Kaempfer, R. Thomann, and R. Mulhaupt, Polymer 43, 2909 (2002). 265. A. Lele, M. Mackley, G. Galgali, and C. Ramesh, J. Rheol. 46, 1091 (2002). 266. Q. Zhang, Y. Wang, and Q. Fu, J. Polym. Sci. B 41, 1 (2003). 267. H. G. Jeon, H. T. jung, S. W. Lee, and S. D. Hudson, Polym. Bull. 41, 107 (1998). 268. J. Heinemann, P. Reichert, R. Thomsaa, and R. Mulhaupt, Macromol. Rapid Commun. 20, 423 (1999). 269. V. P. Privalko, F. J. B. Calleja, D. I. Sukhorukov, E. G. Privalko, R. Walter, and K. Friedrich, J. Mater. Sci. 34, 497 (1999). 270. J. Rong, J. Jing, H. Li, and M. Sheng, Macromol. Rapid Commun. 22, 329 (2001). 271. K. H. Wang, M. H. Choi, C. M. Koo, Y. S. Choi, and I. J. Chung, Polymer 42, 9819 (2001). 272. M. Alexandre, P. Dubois, T. Sun, J. M. Graces, and R. Jerome, Polymer 43, 2123 (2002). 273. T. G. Gopakumar, J. A. Lee, M. Kontopoulou, and J. S. Parent, Polymer 43, 5483 (2002). 274. M. A. Osman, G. Seyfang, and U. W. Suter, J. Phys. Chem. 104, 4433 (2000). 275. M. Zanetti, G. Camino, R. Thomann, and R. Mulhaupt, Polymer 42, 4501 (2001). 276. A. Usuki, A. Tukigase, and M. Kato, Polymer 43, 2185 (2002). 277. S. D. Wangle and J. P. Jog, J. Polym. Sci. B 41, 1014 (2003). 278. A. I. Nazzal and G. B. Street, J. Chem. Soc. Chem. Commun. 375 (1985). 279. Y. Sun and E. Ruckenstein, Synth. Met. 72, 261 (1995). 280. L. Wang, P. Brazis, M. Rocci, C. R. Kannewurf, and M. G. Kanatzidis, in “Organic/Inorganic Hybrid Materials” (R. M. Laine, C. Sanchez, J. F. Brinker, and E. Giannelis, Eds.), Vol. 519, p. 257. Materials Research Society, Warrendale, PA, 1998. 281. S. Sinha Ray and M. Biswas, Mater. Res. Bull. 35, 1187 (1999). 282. B.-H. Kim, J.-H. Jung, J. Joo, J.-W. Kim, and H.-J. Choi, in “Nanocomposites 2001” Proceedings. 283. J. W. Kim, F. Liu, H. J. Choi, S. H. Hong, and J. Joo, Polymer 44, 289 (2003). 284. M. Biswas and S. Sinha Ray, Polymer 39, 6423 (1998). 285. S. Sinha Ray and M. Biswas, J. Appl. Polym. Sci. 73, 2971 (1999). 286. J. M. Kim, S. G. Kim, H. J. Choi, and M. S. Jhon, Macromol. Rapid Commun. 20, 450 (1999). 287. M. Biswas and S. Sinha Ray, J. Appl. Polym. Sci. 77, 2948 (2000). 288. L. Dai, Q. Wang, and M. Wan, J. Mater. Sci. Lett. 19, 1645 (2000). 289. D. Lee, S. H. Lee, K. Char, and J. Kim, Macromol. Rapid Commun. 21, 1136 (2000). 290. Q. Wu, Z. Xue, Z. Qi, and F. Wang, Polymer 41, 2029 (2000). 291. S. Uemura, M. Yoshie, N. Kobyashi, and T. Nakahira, Polym. J. 32, 987 (2000). 292. B. H. Kim, J. H. Jung, J. Joo, J. W. Kim, and H. J. Choi, J. Korean Phys. Soc. 36, 366 (2000). 293. J. M. Yeh, S. J. Liou, C. Y. Lai, P. C. Wu, and T. Y. Tsai, Chem. Mater. 13, 1131 (2001). 294. B. Feng, Y. Su, J. Song, and K. Kong, J. Mater. Sci. Lett. 20, 293 (2001). 295. J. W. Kim, S. G. Kim, H. J. Choi, M. S. Suh, M. J. Shin, and M. S. Jhon, Int. J. Mod. Phys. 15, 657 (2001). 296. B. H. Kim, J. H. Jung, J. W. KIM, H. J. Choi, and J. Joo, Synth. Met. 17, 115 (2001). 297. H. J. Choi, J. W. Kim, J. Joo, and B. H. Kim, Synth. Met. 21, 1325 (2001). 298. M. S. Cho, H. J. Choi, K. Y. Kim, and W. S. Ahn, Macromol. Rapid Commun. 23, 713 (2002). 299. B. H. Kim, J. H. Jung, S. H. Hong, J. Joo, A. J. Epstein, K. Mizoguchi, J. W. Kim, and H. J. Choi, Macromolecules 35, 1419 (2002). 300. G. Nascimento, V. R. L. Constantino, and M. L. A. Temperini, Macromolecules 35, 7535 (2002).

Polymer/Clay Nanocomposites 301. B. Winkler, L. Dai, and A. W.-H. Mau, J. Mater. Sci. Lett. 18, 1539 (1999). 302. S. Ke, M. Ying, C. Ya-An, C. Zhao-Hui, J. Xue-hai, and Y. Jianniah, Chem. Mater. 13, 250 (2000). 303. “Handbook of Conducting Polymers,” 2nd ed. (T. A. Skotheim, R. L. Elsenbaumer, and J. R. Reynolds, Eds.), Ch. 25. Dekker, New York. 304. R. C. Pennwell, B. N. Ganguly, and T. W. Smith, J. Polym. Sci. Macromol. Rev. 13, 63 (1973). 305. M. Biswas and S. K. Das, Polymer 23, 1706 (1982). 306. M. Biswas and A. Majumder, J. Appl. Polym. Sci. 41, 1489 (1991). 307. M. Biswas and P. Mitra, J. Appl. Polym. Sci. 41, 1989 (1991). 308. G. Lagaly, Clay Miner. 16, 1 (1981). 309. M. Kawasumi, N. Hasegawa, A. Usuki, and A. Okada, Mater. Sci. Eng. C 6, 135 (1998). 310. R. A. Vaia and E. P. Giannelis, Polymer 42, 1281 (2001). 311. W. Zhou, J. E. Mark, M. R. Unroe, and F. E. Arnold, J. Macromol. Sci. Pure. Appl. Chem. A 38, 1 (2001). 312. J. H. Chang, B. S. Seo, and D. H. Hwang, Polymer 43, 2969 (2002). 313. C. J. G. Plummer, L. Garamszegi, Y. Leterrier, M. Rodlert, and J.-A. E. Manson, Chem. Mater. 14, 486 (2002). 314. N. Ogata, G. Jimenez, H. Kawai, and T. Ogihara, J. Polym. Sci. B 35, 389 (1997). 315. S. Sinha Ray, K. Okamoto, K. Yamada, and M. Okamoto, Nano Lett. 2, 423 (2002). 316. S. Sinha Ray, M. Okamoto, K. Yamada, and K. Ueda, Polym. Preprints Japan 155 (2002). 317. S. Sinha Ray, K. Yamda, A. Ogami, M. Okamoto, and K. Ueda, Macromol. Rapid Commun. 23, 493 (2002). 318. S. Sinha Ray, M. Okamoto, K. Yamada, and K. Ueda, Nano Lett. 2, 1093 (2002). 319. P. Maiti, K. Yamada, M. Okamoto, K. Ueda, and K. Okamoto, Chem. Mater. 14, 4654 (2002). 320. M. Pluta, A. Caleski, M. Alexandre, M.-A. Paul, and P. Dubois, J. Appl. Polym. Sci. 52, 1497 (2002). 321. S. Sinha Ray, K. Yamada, M. Okamoto, A. Ogami, and K. Ueda, Chem. Mater. 15, 1456 (2003). 322. M.-A. Paul, M. Alexandre, P. Degee, C. Henrist, A. Rulmont, and P. Dubois, Polymer 44, 443 (2003). 323. J.-H. Chang, Y. Uk-An, and G. S. Sur, J. Polym. Sci. B 41, 94 (2003). 324. S. Sinha Ray, K. Yamada, M. Okamoto, and K. Ueda, Polymer 44, 857 (2003). 325. S. Sinha Ray, K. Okamoto, and M. Okamoto, J. Nanosci. Nanotech. 2, 171 (2002). 326. S. Sinha Ray, K. Okamoto, and M. Okamoto, in “Nanocomposites 2002” Proceedings, ECM. 327. S. Sinha Ray, K. Okamoto, and M. Okamoto, Macromolecules 36, 2355 (2003). 328. K. Okamoto, S. Sinha Ray, and M. Okamoto, J. Polym. Sci. B 41, 3160 (2003). 329. X. Kornmann, L. A. Berglund, J. Sterete, and E. P. Giannelis, Polym. Eng. Sci. 38, 1351 (1998). 330. H. J. Choi, J. H. Kim, and J. Kim, Macromol. Symp. 119, 149 (1997). 331. S. H. Park, H. J. Choi, S. T. Lim, T. K. Shin, and M. S. Jhon, Polymer 42, 5737 (2001). 332. S. R. Lee, H. M. Park, H. L. Lim, T. Kang, X. Li, W. J. Cho, and C. S. Ha, Polymer 43, 2495 (2002). 333. R. K. Bharadwaj, A. R. Mehrabi, C. Hamilton, C. Trujillo, M. Fan Murga, A. Chavira, and A. K. Thompson, Polymer 43, 3699 (2002). 334. S. T. Lim, Y. H. Hyun, and H. J. Choi, Chem. Mater. 14, 1839 (2002). 335. L. E. Nielsen, “Mechanical Properties of Polymers and Composites.” Dekker, New York, 1975. 336. T. Yasuda and E. Takiyama, U.S. Patent 5
391
644, 1995.

Polymer/Clay Nanocomposites 337. P. Reichert, H. Nitz, S. Klinke, R. Brandsch, R. Thomann, and R. Mulhaupt, Macromol. Mater. Eng. 275, 8 (2000). 338. T. Mitsunaga, K. Okada, and Y. Nagase, in “PPS 2002 Proceedings,” Taipei, p. 139. 339. M. Biswas and S. Sinha Ray, Adv. Polym. Sci. 155, 167 (2001). 340. E. P. Giannelis, Appl. Organomet. Chem. 12, 675 (1998). 341. J. M. Gilman, Appl. Clay Sci. 15, 31 (1999). 342. G. Camino, R. Sgobbi, S. Colombier, and C. Scelza, Fire Mater. 24, 85 (2000). 343. J. W. Gilman, T. Ksahiwagi, E. P. Giannelis, E. Manias, S. Lomakin, J. D. Lichtenhan, and P. Jones, in “Chemistry and Technology of Polymer Additives” (S. Al-Malaika, A. Golovoy, and C. A. Wilkie, Eds.), Ch. 14. Blackwell Science, Oxford, 1999. 344. S. T. Lim, Y. H. Hyun, and H. J. Choi, Chem. Mater. 14, 1839 (2002). 345. J. T. Yoon, W. H. Jo, M. S. Lee, and M. B. Ko, Polymer 42, 329 (2001). 346. J. W. Gilman, Appl. Clay Sci. 15, 31 (1999). 347. J. W. Gilman, C. L. Jackson, A. B. Morgan, R. Harris, Jr., E. Manias, E. P. Giannelis, M. Wuthenow, D. Hilton, and S. H. Phillips, Chem. Mater. 12, 1866 (2000). 348. S. Sinha Ray, K. Yamada, M. Okamoto, A. Ogami, and K. Ueda, Polymer 44, 6631 (2003). 349. L. Nielsen, J. Macromol. Sci. Chem. A 1, 929 (1967). 350. R. Xu, E. Manias, A. J. Snyder, and J. Runt, Macromolecules 34, 337 (2001). 351. J. C. Hutchison, R. Bissessur, and D. F. Shiver, Chem. Mater. 8, 1597 (1996). 352. J. A. Tetto, D. M. Steeves, E. A. Welsh, and B. E. Powell, in “ANTEC’99,” p. 1628. 353. J. W. Liu, Q. Zhao, and C. X. Wan, Space Med. Med. Eng. 14, 308 (2001). 354. S. Sinha Ray, K. Yamada, M. Okamoto, and K. Ueda, Macromol. Mater. Eng. 288, 936 (2003). 355. J. Massam, Z. Wang, T. J. Pinniavaia, T. Lan, and G. Beall, Polym. Mater. Sci. Eng. 78, 274 (1998). 356. B. B. Lotz, S. Graff, and J. C. Wittmann, J. Polym. Sci. B 24, 2017 (1986). 357. R. Khare, J. J. de. Pablo, and A. Yethiraj, Macromolecules 29, 7910 (1996).

843 358. P. Maiti and M. Okamoto, Macromol. Mater. Eng. 288, 440 (2003). 359. R. Krishnamoorti and E. P. Giannelis, Langmuir 17, 1448 (2001). 360. R. Krishnamoorti, J. Ren, and A. S. Silva, J. Chem. Phys. 15, 4968 (2001). 361. R. Krishnamoorti and K. Yurekli, Current Opinion Colloid Interface Sci. 6, 464 (2001). 362. J. Ren, A. S. Silva, and R. Krishnamoorti, Macromolecules 33, 3739 (2000). 363. M. Okamoto, H. Taguchi, H. Sato, T. Kotaka, and H. Tatayama, Langmuir 16, 4055 (2000). 364. M. L. Williams, R. F. Landel, and J. D. Ferry, J. Amer. Chem. Soc. 77, 3701 (1955). 365. P. H. Nam, Master Thesis, Toyota Technological Institute, 2001. 366. W. P. Cox and E. H. Merz, J. Polym. Sci. 28, 619 (1958). 367. T. Kotaka, A. Kojima, and M. Okamoto, Rheol. Acta 36, 646 (1997). 368. M. Okamoto, P. H. Nam, and P. Maiti, unpublished results, 2003. 369. H. van Olphen, “An Introduction to Clay Colloid Chemistry.” Wiley, New York, 1977. 370. H. Block and J. P. Kelly, J. Phys. D 21, 1661 (1998). 371. W. M. Winslow, J. Appl. Phys. 20, 1137 (1992). 372. T. C. Halsey, Science 258, 761 (1992). 373. M. Trau, S. Sankaran, D. A. Saville, and I. A. Aksay, Nature 374, 437 (1995). 374. J. W. Kim, M. H. Noh, H. J. Choi, D. C. Lee, and M. S. Jhon, Polymer 41, 1229 (2000). 375. J. H. Park, Y. T. Lim, and O. O. Park, Macromol. Rapid Commun. 22, 616 (2001). 376. Y. T. Lim, J. H. Park, and O. O. Park, J. Colloid Interface Sci. 245, 198 (2002). 377. D. Klempner and K. C. Frisch, “Handbook of Polymeric Foams and Foam Technology.” Hanser, Munich/Vienna, 1991. 378. T. Abiko, H. Kawato, T. Kanai, T. Oda, and H. Saito, in “Nanocomposites 2002 Proceedings,” 379. H. Fong, W. Liu, C. S. Wang, and R. A. Vaia, Polymer 43, 775 (2002). 380. J. M. Brown, D. B. Curliss, and R. A. Vaia, in “Proceedings of the PMSE Spring Meeting,” San Francisco, 2000, p 278. 381. “Cellular Solids” (L. J. Gibson and M. F. Ashby, Eds.), p. 8. Pergamon, New York, 1988.