Synthesis Techniques for Polymer...

Edited by

Vikas Mittal

Synthesis Techniques for Polymer

Nanocomposites

Polymer Nano-, Micro- &Macrocomposite Series

Mittal, V. (ed.)

Surface Modification of

Nanotube FillersSeries: Polymer Nano-, Micro- &

Macrocomposite (Volume 1)

2011

Print ISBN: 978-3-527-32878-9

Mittal, V. (ed.)

In-situ Synthesis of Polymer

NanocompositesSeries: Polymer Nano-, Micro- &


2012

Print ISBN: 978-3-527-32879-6

Mittal, V. (ed.)

Characterization Techniques

for Polymer NanocompositesSeries: Polymer Nano-, Micro- &


2012

Print ISBN: 978-3-527-33148-2

Mittal, V. (ed.)

Modeling and Prediction of

Polymer Nanocomposite

PropertiesSeries: Polymer Nano-, Micro- &


2013

Print ISBN: 978-3-527-33150-5

Mittal, V. (ed.)

Thermoset NanocompositesSeries: Polymer Nano-, Micro- &


2013

Print ISBN: 978-3-527-33301-1

Edited by Vikas Mittal

Synthesis Techniques for Polymer

Nanocomposites

The Editor

Dr. Vikas Mittal

The Petroleum Institute

Chemical Engineering Department

Room 2204, Bu Hasa Building

Abu Dhabi

United Arab Emirates

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V

Contents

Preface XI

List of Contributors XIII

1 Synthesis of Polymer Nanocomposites: Review of Various

Techniques 1

Joel Fawaz and Vikas Mittal

1.1 Introduction 1

1.2 Synthesis Methods 4

1.2.1 Melt Intercalation 4

1.2.2 Exfoliation Adsorption 9

1.2.2.1 Solution Intercalation 9

1.2.2.2 Emulsion Polymerization 11

1.2.3 In Situ Polymerization 16

1.2.4 Nontraditional Methods 23

References 26

2 Masterbatch Approach to Generate HDPE/CPE/Graphene

Nanocomposites 31

Ali U. Chaudhry and Vikas Mittal

2.1 Introduction 31

2.2 Experimental 33

2.2.1 Materials 33

2.2.2 Preparation of Graphite Oxide and Graphene Oxide 34

2.2.3 Nanocomposite Generation 35

2.2.4 Material Characterization 36

2.3 Results and Discussion 37

2.4 Conclusions 47

Acknowledgments 48

References 48

VI Contents

3 Preparation and Applications of Hydroxyapatite Nanocomposites

Based on Biodegradable and Natural Polymers 51

Pau Turon, Luis J. del Valle, Carlos Alemán, and Jordi Puiggalí

3.1 Introduction 51

3.2 Preparation of HAp Nanocrystals 52

3.3 Preparation of HAp Nanocomposites 58

3.4 Applications of HAp/DNA Nanocomplexes as Gene Carriers 61

3.5 Tissue Engineering Applications of HAp Nanocomposites Based on

Biodegradable Polymers 65

3.6 Applications of HAp Nanocomposites Based on Biodegradable

Polymers as Drug Delivery Systems 72

3.7 Miscellaneous Applications of HAp Nanocomposites Based on

Biodegradable Polymers 76

3.8 Concluding Remarks 79

Acknowledgments 80

References 80

4 Synthetic Methods for Nanocomposites Based on Polyester Resins 87

Michał Kedzierski

4.1 Introduction 87

4.2 Nanocomposites with Zero-Dimensional Nanofillers 89

4.2.1 Silicon-Containing Nanospheres 89

4.2.2 Metal Oxides 91

4.2.3 Other 0-D Nanoparticles 93

4.3 Nanocomposites with One-Dimensional Nanofillers 93

4.3.1 Carbon Nanotubes and Nanofibers 93

4.3.2 Cellulose Nanofibers 96

4.3.3 Other 1-D Nanofillers 97

4.4 Nanocomposites with Two-Dimensional Nanofillers 97

4.4.1 Layered Aluminosilicate Clays 97

4.4.1.1 Mixing Methods 98

4.4.1.2 Effects of the Clay Modification 99

4.4.1.3 Nanocomposites with MMT Introduced during the Synthesis of

Pre-polymer 102

4.4.1.4 Various Properties and Multiphase Nanocomposites 103

4.4.1.5 Vinyl Ester–Clay Nanocomposites 106

4.4.2 Layered Double Hydroxides 106

4.4.3 Graphene-Based Nanofillers 107

4.5 Conclusions 109

Abbreviations 110

References 110

Contents VII

5 Synthesis Fabrication and Characterization of Ag/CNT-Polymer

Nanocomposites 115

Vijaya K. Rangari and Sanchit Dey

5.1 Introduction 115

5.2 Experimental Procedure 118


5.3.1 XRD analysis 119

5.3.2 Transmission Electron Microscopy 119

5.3.3 TGA Analysis of Nanoparticles 121

5.3.4 Thermal Response of the Polymer Composites 121

5.3.5 Compression Test Results of Polymer Composites 124

5.3.6 Flexure Test Results of Polymer Composites 125

5.4 Conclusion 127

Acknowledgments 127

References 127

6 Preparation and Characterization of PVDF-Based

Nanocomposites 131

Derman Vatansever Bayramol, Tahir Shah, Navneet Soin, and Elias Siores

6.1 Synthesis of Poly(vinylidene fluoride) (PVDF) 131

6.2 Structure and Piezoelectric Properties of PVDF 131

6.2.1 Relationships and Equations 135

6.2.1.1 The Piezoelectric Charge Constant and Piezoelectric Voltage

Constant 136

6.3 Processing of PVDF for Energy Harvesting Applications 137

6.4 Processing of PVDF Based Materials: Polymer/Polymer,

Polymer/Nanofiller, Polymer/Ionomer Blends 138

6.5 PVDF Based Nanocomposites for Energy Harvesting

Applications 139

6.6 Conclusion 140

References 141

7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer

Nanocomposites 145

Luana Persano, Andrea Camposeo, AnnaMaria Laera, Francesca Di

Benedetto, Vincenzo Resta, Leander Tapfer, and Dario Pisignano


7.2 Thermal-Assisted In Situ Synthesis: Material Choice and

Nanocomposite Characterization 146

7.2.1 Precursor Molecules 146

7.2.1.1 Metal Salts 147

7.2.1.2 Organometallic Compounds 147

7.2.2 Thermal Synthesis and Composites Characterization 151

7.2.2.1 Microstructural Characterization 152

7.2.2.2 Optical Spectroscopy Experiments 154

VIII Contents

7.3 Fabrication of Nanocomposites and Patterning 155

7.3.1 Nanocomposites by Photoirradiation 157

7.3.1.1 UV and Visible Irradiation 157

7.3.1.2 Multiphoton Irradiation 160

7.3.2 Nanocomposites by Electron-BeamWriting 160

7.3.3 Nanocomposite Polymer Fibers 165

7.3.3.1 Photo-Assisted Synthesis 167

7.3.3.2 Thermal-Assisted Synthesis 169

7.4 Conclusions 171

Acknowledgments 172

References 172

8 Synthesis of Polymer Nanocomposites by Water-Assisted

Extrusion 179

Naïma Sallem-Idrissi, Michel Sclavons, and Jacques Devaux


8.2 Nanocomposites Structure and Characterization 180

8.2.1 Clays 180

8.2.2 Organomodification of Layered Silicates 181

8.2.3 Nanocomposites Structure and Characterization 182

8.3 Nanocomposites Preparation 183

8.3.1 Intercalation from Solution 183

8.3.2 In Situ Polymerization 183

8.3.3 Melt Compounding 184

8.3.3.1 Melt Blending of Polymer/Organoclay Nanocomposites 184

8.3.3.2 Melt Blending of Polymer/Pristine Clay Nanocomposites 186

8.4 Nanocomposite Properties 195

8.4.1 Thermal Stability 195

8.4.2 Flame Retardancy 197

8.5 Toward Fully Green Composites? 198

References 201

9 In Situ Preparation of Conducting Polymer Nanocomposites 211

Liping Yang, Cher Ling Toh, and Xuehong Lu


9.1.1 Electrically Conductive Polymer Nanocomposites andTheir

Applications 212

9.1.2 PercolationTheory 213

9.1.3 Factors Affecting the Electrical Conductivity of

Nanocomposites 214

9.1.3.1 Physical Properties of the Fillers 214

9.1.3.2 Filler Distribution and Dispersion 216

9.1.3.3 Physical Properties of Polymer Matrices 216

9.1.3.4 Filler Orientation and Alignment 217

9.1.3.5 Nanocomposite Fabrication Methods and Conditions 218

Contents IX

9.2 In Situ Preparation of Conductive Nanocomposites 219

9.2.1 In Situ Polymerization Strategy 219

9.2.1.1 Step Growth 220

9.2.1.2 Chain Growth 224

9.2.1.3 Aligning Conductive Fillers in in situ Polymerization Processes 227

9.2.2 In Situ Formation of Conducting Polymer Nanocomposites 228

9.2.2.1 In Situ Formation of rGO-Based Polymer Nanocomposites 228

9.2.2.2 In Situ Formation of Metallic Conductive Pathways 232

9.3 Challenges and Outlook 233

References 235

10 Near IR Spectroscopy for the Characterization of Dispersion in

Polymer–Clay Nanocomposites 241

Ana VeraMachado, JoanaMargarida Barbas, and Jose Antonio Covas


10.2 Morphology and Properties 241

10.3 Preparation Methods 243

10.4 Characterization Techniques 243

10.5 Dispersion by Melt Mixing 247

10.6 Online and Inline Monitoring of Dispersion 249

10.7 Conclusions 259

References 259

11 Synthesis of Polymer Nanocomposites in Supercritical CO2 267

Yuvaraj Haldorai and Jae-Jin Shim


11.2 Background on Supercritical CO2 268

11.3 Physical and Chemical Properties of scCO2 270

11.4 Preparation of Polymer/Inorganic Filler Nanocomposites in

Supercritical CO2 272

11.4.1 Ex SituMethod 272

11.4.1.1 Solution Blending 272

11.4.1.2 Melt Blending 272

11.4.2 In SituMethod 276

11.4.2.1 Synthesis of Nanocomposites by Dispersion Polymerization 277

11.4.2.2 Synthesis of Nanocomposites by Other Techniques 281

11.5 Conclusions 286

References 286

Index 291

XI

Preface

Nanocomposites are high-value nanomaterials with applications in diverse

fields. Owing to the requirement of the dispersion of filler at nanoscale (less

than 100 nm), a large number of synthesis routes have been developed. The

choice of the synthesis route depends on the nature of the polymer and filler

and correspondingly results in the required composite properties. Occasion-

ally, combinations of methods are also employed to enhance the composite

microstructure. Thus it is of importance to combine these synthetic methods

into a meaningful text that would provide guidelines for the readers to make the

choice of correct synthesis route.

Chapter 1 reviews the various synthesis routes to generate the polymer

nanocomposites, for example, melt intercalation, solution mixing, in-situ poly-

merization, and so on. Chapter 2 provides details on the masterbatch approach

for the synthesis of polyolefin nanocomposites with compatibilizer. Chapter 3

focuses on the different synthetic approaches that can be applied to prepare

nanohydroxyapatite crystals with controlled morphology and the procedures to

generate composites based on nanohydroxyapatite and biodegradable polymers

of natural or synthetic origin. Chapter 4 also describes various synthetic methods

for generating nanocomposites based on polyester resins. Chapter 5 elaborates

on the use of microwave radiation to produce the metal nanoparticles on the

outer surface of CNTs, which are subsequently used as filler in the fabrication

of multifunctional polymer nanocomposites for various cutting-edge applica-

tions. Chapter 6 reviews the preparation and characterization of PVDF-based

nanocomposites (polymer/polymer blends, polymer/nanoparticle blends, and

ternary blends) and focuses on the preparation and characterization of PVDF-

based nanocomposites for energy harvesting applications. Chapter 7 explains

in-situ synthesis and patterning methods, also in combined modes, based on

photon, and electron beam assisted procedures to generate nanocomposites.

Chapter 8 describes water assisted extrusion process for the generation of

nanocomposites, which is not only an affordable method (no fillers’ organophilic

modification is needed), but also less hazardous to health.

Chapter 9 concentrates on the conducting nanocomposites (with insulating

polymer matrices) prepared via in-situ polymerization or in-situ processing

XII Preface

methods. Chapter 10 discusses the use of NIR spectroscopy for the characteriza-

tion of dispersion in polymer nanocomposites, with a focus on the application of

inline techniques to monitor the preparation of polymer-clay nanocomposites by

melt compounding. Chapter 11 analyzes the synthesis of polymer nanocompos-

ites by ex-situ and in-situ methods in scCO2 by providing a general overview of

the techniques and strategies used for the preparation of nanocomposites.

Abu Dhabi Vikas Mittal

November 2014

XIII

List of Contributors

Carlos Aleman

Universitat Politècnica de

Catalunya

Departament d’Enginyeria

Quımica

Avinguda Diagonal 647

08028 Barcelona

Spain

JoanaMargarida Barbas

University of Minho

Institute of Polymers and

Composites (IPC/I3N)

Campus de Azurém

4800-058 Guimarães

Portugal

Derman Vatansever Bayramol

Namık Kemal University

Department of Textile

Engineering

Silahtaraga Mah. Universite

1. Sok No:13

59850 Corlu-Tekirdag

Turkey

Andrea Camposeo

National Nanotechnology

Laboratory of Istituto

Nanoscienze-CNR

via Arnesano

73100 Lecce

Italy

Ali U. Chaudhry


Department of Chemical

Engineering

Bu Hasa Building

Room 2204

2533 Abu Dhabi


Jose Antonio Covas

University of Minho



Campus de Azurém

4800-058 Guimarães

Portugal

Luis J. del Valle


Catalunya


Quımica


08028 Barcelona

Spain

Jacques Devaux

UCL-IMCN/BSMA

Croix du Sud 1

L7.04.02

1348 Louvain-la-Neuve

Belgium

XIV List of Contributors

Sanchit Dey

Tuskegee University

Department of Materials Science

and Engineering

100 James Center

Tuskegee, AL 36088

USA

Francesca Di Benedetto



Nanoscienze-CNR

via Arnesano

73100 Lecce

Italy

and

ENEA

Technical Unit of Material

Technologies Brindisi

Strada Statale 7 Appia km. 706

72100 Brindisi

Italy

Joel Fawaz



Engineering

Bu Hasa Building

Room 2204

2533 Abu Dhabi


Yuvaraj Haldorai

Yeungnam University

Supercritical Fluids and Nano

Processes Laboratory

School of Chemical Engineering

214-1 Dae–dong, Gyeongsan

712-749 Gyeongbuk

Republic of Korea

and

Department of Energy and

Materials Engineering

Dongguk University-Seoul

30, Pildong-ro 1gil, Jung-gu

Seoul, 100-715

Republic of Korea

Michał Kedzierski

Industrial Chemical Research

Institute

Department of Polyesters

Epoxide Resins and

Polyurethanes

Rydygiera Street 8

01 793 Warsaw

Poland

AnnaMaria Laera

ENEA




72100 Brindisi

Italy

Xuehong Lu

Nanyang Technological

University

School of Materials Science and

Engineering

639798

Singapore

List of Contributors XV

Ana VeraMachado

University of Minho



Campus de Azurém

4800-058 Guimarães

Portugal

VikasMittal



Engineering

Bu Hasa Building, Room 2204

2533 Abu Dhabi


Luana Persano



Nanoscienze-CNR

via Arnesano

73100 Lecce

Italy

Dario Pisignano



Nanoscienze-CNR

via Arnesano

73100 Lecce

Italy

and

Università del Salento

Dipartimento di Matematica e

Fisica “Ennio De Giorgi”

via Arnesano

73100 Lecce

Italy

Jordi Puiggalı


Catalunya


Quımica


08028 Barcelona

Spain

Vijaya K. Rangari

Tuskegee University

Department of Materials Science

and Engineering

100 James Center

Tuskegee, AL 36088

USA

Vincenzo Resta

ENEA




72100 Brindisi

Italy

and

University of Salento

Department of Engineering for

Innovation

CEDAD-Center for

Dating and Diagnostics

via Monteroni

73100 Lecce

Italy

Naıma Sallem-Idrissi

UCL-IMCN/BSMA

Croix du Sud 1

L7.04.02


Belgium

XVI List of Contributors

Michel Sclavons

UCL-IMCN/BSMA

Croix du Sud 1

L7.04.02


Belgium

Tahir Shah

University of Bolton

Institute for Materials Research

and Innovation

Bolton, BL3 5AB

UK

Jae-Jin Shim

Yeungnam University

Supercritical Fluids and Nano

Processes Laboratory

School of Chemical Engineering

214-1 Dae-dong, Gyeongsan

712-749 Gyeongbuk

Republic of Korea

Elias Siores



and Innovation

Bolton, BL3 5AB

UK

Navneet Soin



and Innovation

Bolton, BL3 5AB

UK

Leander Tapfer

ENEA




72100 Brindisi

Italy

Cher Ling Toh

Nanyang Technological

University

School of Materials Science and

Engineering

639798

Singapore

Pau Turon

B. Braun Surgical S.A.

Carretera de Terrasa 121

08191 Rubı (Barcelona)

Spain

Liping Yang

A*STAR (Agency for Science,

Technology and Research)

Institute of Chemical and

Engineering Sciences

1 Pesek Road

627833 Jurong Island

Singapore

1

1

Synthesis of Polymer Nanocomposites:

Review of Various Techniques

Joel Fawaz and Vikas Mittal

1.1

Introduction

Polymer nanocomposites are hybrid organic–inorganicmaterials with at least one

dimension of the filler phase less than 100 nm [1]. Polymer nanocomposites are

synthesized via various methods that can be categorized into four major routes:

melt intercalation, template synthesis, exfoliation adsorption, and in situ poly-

merization intercalation [1–6]. On the basis of the method and materials used,

three types of microstructure can be obtained: unintercalated (or microcompos-

ite), intercalated (and/or flocculated), or exfoliated (or delaminated), as shown in

Figure 1.1.

Melt intercalation is the typical standard approach for synthesizing thermo-

plastic polymer nanocomposites. It involves annealing the polymer matrix at high

temperatures, adding the filler, and finally kneading the composite to achieve uni-

formdistribution, as illustrated in Figure 1.2. It has the advantage of being environ-

mental friendly because of the lack of solvent usage. In addition, it is considered

compatible with industrial processes such as injection molding and extrusion,

which makes it more convenient to utilize and, thus, more economical. However,

the high temperatures used in the process can damage the surface modification of

the filler. For example, organoclaysmodifiedwith alkyl ammoniumusually decom-

pose at temperatures higher than 140 ∘C; however, the processing temperature of

melt intercalation is in the range of 190–220 ∘C [4]. Therefore, optimization of

the processing conditions is a very important factor that plays a big role in achiev-

ing good dispersion and exfoliation. For instance, operating at lower temperatures

or using more thermally stable modifications can avoid degradation [1]. Weak

electrostatic forces among the filler interlayers and compatibility with the poly-

mer matrix allow the polymer to crawl into the interlayers forming intercalated or

exfoliated nanocomposites [6].

Exfoliation adsorption, also called polymer or prepolymer intercalation from

solution, is based on a solvent in which the polymer or prepolymer is soluble.The

layered silicate, for instance, is first swollen and dispersed in solvent beforemixing

it with the polymer solution.The polymer chains then intercalate and displace the

Synthesis Techniques for Polymer Nanocomposites, First Edition. Edited by Vikas Mittal.© 2015 Wiley-VCH Verlag GmbH & Co. KGaA. Published 2015 by Wiley-VCH Verlag GmbH & Co. KGaA.

2 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques

Layered silicate Polymer

(a) (b) (c)

Figure 1.1 Types of composite microstructures: (a) Unintercalated (Phase separated (micro-

composite)), (b) intercalated (Intercalated (nanocomposite)), and (c) exfoliated (nanocompos-

ite). Reproduced from [6] with permission from Elsevier.

NH3+

NH3+

NH3+ NH3

+NH3

+ NH3+

Organophilicclay

Thermoplasticpolymer

Intercalation

Blending+

annealing

Figure 1.2 The melt intercalation process. Reproduced from [3] with permission from

Elsevier.

solvent within the silicate interlayers. Eventually, on removal of the solvent, a mul-

tilayer structure is formed as the sheets reassemble trapping the polymer chains,

as shown in Figure 1.3 [2, 5, 6]. This approach is widely used for water-soluble

polymers to produce intercalated nanocomposites based on polymers with low or

no polarity such as poly (vinyl alcohol), poly (ethylene oxide), poly (vinylpyrroli-

done), or poly (acrylic acid) [3, 6]. However, unlikemelt intercalation, this method

is environmentally unfriendly because of the usage of large amounts of solvents.

Emulsion polymerization is considered to be under this method as monomers,

usually methyl methacrylate and styrene, are dispersed in water along with an

emulsifier and different silicate concentrations [5]. The monomer is polymerized

with a part of silicate embedded inside the polymer particle and a part adsorbed

on the particle surface, forming a nanocomposite.

1.1 Introduction 3

Clay dispersion

Polymer solution

Polymer intercalation in thegalleries of dispersed clay

Solvent evaporation andnanocomposite recovery

Figure 1.3 The exfoliation adsorption process. Reproduced from [3] with permission from

Elsevier.

In situ polymerization involves the swelling of the filler in liquid monomer or

monomer solution as the low-molecular-weight monomer seeps in between the

interlayers causing the swelling [5]. Polymerization starts either using heat, radi-

ation, initiator diffusion, or by organic initiator or catalyst fixed through cationic

exchange [6]. The monomers then polymerize in between the interlayers forming

intercalated or exfoliated nanocomposites. The advantage of this approach lies

in the better exfoliation achieved compared to melt and exfoliation adsorption

methods [4]. Figure 1.4 illustrates the synthesis of nylon-6/clay nanocomposite

via in situ polymerization in which clay is dispersed in caprolactammonomer and

under polymerization conditions, an exfoliated nanocomposite is formed.

Template synthesis, also known as sol-gel technology, is based on an opposite

principle than the previous methods. This approach involves the formation of the

inorganic filler in an aqueous solution or gel containing the polymer and the filler

building blocks [1, 3–6]. The polymer serves as a nucleating agent and promotes

the growth of the inorganic filler crystals. As those crystals grow, the polymer is

trapped within the layers and thus forms the nanocomposite. It is mainly used

for the synthesis of double-layer hydroxide-based nanocomposite and is much

Caprolactam Clay mineral

A layer of clay Nylon 6

Polymerization

Figure 1.4 Schematic example of in situ polymerization process involving the synthesis of

nylon-6/clay nanocomposite. Reproduced from [2] with permission from Elsevier.


less developed for the synthesis of layered silicates. This is because of the high

temperature used during synthesis that degrades the polymer and the resulting

aggregation tendency of the growing inorganic crystals [3, 5]. Therefore, this pro-

cess is not commonly used.

Till now, only a brief introduction to each method has been provided. However,

this chapter focuses on discussing the recent studies conducted in each of the three

main respective synthesis methods. Different types of fillers, such as carbon nan-

otubes (CNTs), silicates, and clay and graphene/graphite oxide, are inclusive in

this review.This chapter also analyzes nontraditional methods such as microwave

assisted and redox reactions. The readers are referred to these review papers for

further reading [1–9].

1.2

Synthesis Methods

1.2.1

Melt Intercalation

As discussed earlier, melt intercalation is considered environmental friendly and a

much better substitution for solution mixing, if permittable. However, processing

conditions, surface modification of fillers, and compatibility of filler and polymer

matrix all play important roles in determining how well the dispersion can be

achieved. Alig et al. [10] discussed the relation between processing conditions

and morphologies obtained for CNT nanocomposites. Moreover, the authors

explained the dispersion process by breaking it into four steps: (i) Wetting of

initial agglomerates by the polymer, (ii) infiltration of polymer chains into the

initial agglomerates to weaken them, (iii) dispersion of agglomerates by rupture

and erosion, and (iv) distribution of individualized nanotubes into the matrix.

Similarly, Pavlidou and Papaspyrides [3] explained the thermodynamics behind,

and the effects of multiple conditions on, melt intercalation for polymer/layered

silicates. The entropy loss, associated with the confinement of a polymer melt, is

balanced with an entropy gain that is associated with layer separation and greater

conformational energy of aliphatic chains of alkylammonium cations. Therefore,

it is generally agreed that melt intercalation depends on the surface energies of

polymer and modified layered silicates [3].

Junior et al. [11] reported the synthesis of recycled high-impact polystyrene

(PS)/organoclay nanocomposites by melt intercalation. The processing was done

in an interpenetrating corotating twin screw extruder with screw diameter of

20mm and L/D ratio of 36. Two different speeds and two types of clay fillers

(Viscogel S4 and S7 montmorillonite clays), each with different surfactant, were

used. Temperature varied between 150 and 190 ∘C in the processing zones.

The high-impact PS was milled before mixing in order to increase the surface

area and facilitate dispersion. It was reported that the higher mixing speed of


600 rpm yielded nanocomposites with better dispersion than the ones processed

at 450 rpm.

Poly(ε-caprolactone) (PCL)/organo-modifiedmontmorillonites (MMTs) nano-

composites are synthesized in a corotating twin screw extruder whose length

is 1200mm and L/D ratio of 48 [12]. The extrusion was conducted at 140 ∘C at

250 rpm and 3 kg h−1 polymer flow. However, masterbatches of different types of

organoclay were prepared to be fed into the extruder rather than following direct

addition. Mixed intercalated or exfoliated structures were obtained with different

clay material as the nanocomposite prepared with C30B® clay mineral yields

an intercalated/exfoliated structure whereas Nanofils5® and Nanofils2® give

rise to intercalated nanocomposite Figure 1.5 shows the transmission electron

microscope (TEM) images used to characterize the nanocomposites at 3wt%

loading. However, rheological tests showed that better dispersion was obtained

for the nonpolar Nanofils2® and this was reflected in the enhancement of the

respective thermal and mechanical properties.

Maiti et al. [13] reported the preparation of PCL–multiwalled carbon nan-

otubes (MWCNTs) mixture via melt blending followed by the synthesis of

polycarbonate/ε-PCL–MWCNT nanocomposite. A masterbatch of PCL–

MWCNT with 3.5wt% MWCNT loading was first prepared via melt blending

using internal mixer at 65 ∘C and 60 rpm for 10min. Then, the masterbatch

was melt mixed with pure PC at 280 ∘C and 60 rpm for 10min. This procedure

100 . 0 KU X25 K 200 nm

100 . 0 KU X100 K 50 nm 100 . 0 KU X100 K 50 nm

(a)

(b) (c)

Figure 1.5 TEM images of PCL nancomposites at 3wt% of: (a) Nanofil5®, (b) C30B®, and(c) Nanofil2®. Reproduced from [12] with permission from Elsevier.


yielded a homogeneous dispersion of CNTs at low loadings as analyzed in

scanning electron microscope (SEM). Moreover, through this method, chemical

modification of CNTs was not needed as the percolation threshold obtained

was at 0.14wt%. This suggested that an interconnected network was successfully

achieved at a low CNT loading.

Other studies conducted by Annala et al. [14] and Wang et al. [15] utilized

the masterbatch process to improve the properties of the final nanocompos-

ites. Annala et al. [14] reported the synthesis of poly(methyl methacrylate)

(PMMA)/MWCNT and PS/MWCNT using in situ polymerized masterbatches

that were to be used in corotating twin screw mini-extruder with the capacity

of 16 cm3 and screw length of 150mm. Different mixing speed and time were

investigated to determine the optimum conditions for better properties. Similarly,

Wang et al. [15] synthesized phthalocyanine (Pc)/MWCNT nanocomposites by

placing the prepared masterbatch in a preheated mold at 250 ∘C and cured at

controlled elevated temperatures for 4 h. In both situations, good dispersion of

the CNTs was achieved. However, it was noted that depending on the properties

of the system, the feeding method of CNTs can affect the properties of the final

composite [14].

Tan et al. [16] reported a novel approach of synthesizing rubber/clay nanocom-

posites via latex compounding and melt mixing. In this approach, well-exfoliated

masterbatches and intercalated/exfoliated nanocomposites were achieved by

using Ca-MMT modified with bis[3-triethoxysilylpropyl-]tetrasulfide (TESPT).

This modification enhanced the interface by reacting with the surface groups of

Ca-MMT. The masterbatch was first prepared by latex compounding in which

the cooled organic clay aqueous suspension was mixed with natural rubber (NR)

latex. The mixture was vigorously stirred, co-coagulated in 10% calcium chloride

and eventually washed and dried. The masterbatchs were added to a 6-inch

two-roll mill along with styrene butadiene rubber (SBR) and epoxidized natural

rubber (ENR) to be melt mixed to achieve the nanocomposite. Figure 1.6 shows

the X-ray diffraction (XRD) patterns for the pristine Ca-MMT, the masterbatch,

and the nanocomposite. It can be noted that an exfoliated structure was obtained

in the masterbatch following the absence of peaks. Moreover, this led to an

exfoliated/intercalated structure as some of the initial clay in the masterbatch

was intercalated by the rubber chains.

A novel approach of melt spinning layered double hydroxide (LDH)/high-

density polyethylene (HDPE) nanocomposites prepared by melt extrusion was

reported by Kutlu et al. [17]. LDHs were hydrophobically modified by carboxylic

acid salts of different alkyl chain lengths to improve the lack of compatibility

between LDH and polymer matrix. Those modified LDHs were first mixed with

PE-g-maleic anhydride (MA) to improve the miscibility of LDH and PE followed

by the dilution of masterbatches with HDPE. Then, they were processed in a

microcompounder at 190 ∘C, 100 rpm and 5–10min mixing time. Different

modifiers yielded different interlayer arrangements. Polymer chains were stated

to diffuse into LDH galleries because of the high-shearing force, and partial

exfoliation was achieved, as supported by XRD and TEM analysis. Myristic acid


2 4 6 8 10

2θ (°)

Inte

nsity (

a.u

.)

(a)

(c)

(b)

5.8° (d001=1.5 nm)1.8° (d001=4.8 nm)

6.8° (d001=1.3 nm)

5.5° (d001=1.6 nm)

5.8° (d001=1.5 nm)

Figure 1.6 XRD patterns for: (a) pristine Ca-MMT, (b) NR/modified Ca-MMT masterbatch,

and (c) rubber/clay nanocomposite. Reproduced from [16] with permission from Elsevier.

modified LDH/HDPE nanocomposite showed the highest exfoliation degree at

1wt% filler level as well as the best processing conditions and mechanical prop-

erties of the fiber elements. On the other hand, Mezghani et al. [18] reported the

synthesis of linear low density polyethylene (LLDPE)/MWCNT nanocomposite

fibers prepared via melt extrusion and spun through a spinneret die. The effects

of CNT loadings on the properties of LLDPE/MWCNT nanocomposite were

investigated and it was noted that on slight addition of CNT, the properties are

generally enhanced.

Shanks and Cerezo [19] reported the synthesis of poly(propylene-g-maleic

anhydride) (PPMA)/expanded graphite oxide (EGO) nanocomposites. This was

done in HAAKE heated kneading mixer for 30min at 200 ∘C and 60 rpm. Because

of the unpolar nature of PP (polypropylene), a compatibilizer containing polar

groups such as MA was required to improve compatibility between the two

systems. There was no change in the d-spacing of graphite layers in PPMA/EGO

nanocomposites at different EGO loadings, as reported by XRD results. The

graphite layers were said to be ordered and multilayered in the final composite.

Unnikrishnan et al. [20] reported the synthesis of PMMA/organoclay nanocom-

posites using a 69-cm3 batch mixer with roller rotors. Before blending, the dif-

ferent organoclays (C30B®, C10A®, and C93A®) and PMMA pellets were dried

for 12 h for better processing. Temperature was set to 180 ∘C at a rotor speed

of 50 rpm for 30min. It was noted that with the addition of maleic anhydride,

as a grafting agent, better intercalation was achieved as investigated in the TEM

images. The grafting agent improved the interfacial region between the PMMA

and the clay minerals, which led to the intercalation of the polymer chains in

between the clay layers. PMMA/C30B® nanocomposite was reported to have an

optimum, as well as the highest, d-spacing of 4.16 nm.

Thermoplastic Polyurethane (TPU)/C15A® clay nanocomposites were reported

to be synthesized by Barick and Tripathy [21] in HAAKE extruder at 185 ∘C and


100 rpm rotor speed for 6min. It was detected by XRD that exfoliated structures

were obtained at low loadings of clay minerals because of either high disorder

state or the exfoliation of the silicate layers. However, the peak position at

d001 = 16.5Å and d002 = 36.64Å of the clay is shifted to 19.5 and 40.5Å in 9wt%

loaded nanocomposite, respectively.This indicated the intercalation was achieved

above 5wt% loading. Because of the absence of functional groups on C15A®and high shear stresses from melt processing, mixed exfoliation/intercalation

nanocomposites were obtained. Moreover, it was visible and supported in TEM

that with increasing clay loading, small clusters of clay particles were observed

giving rise to intercalated structures.

Poly(ethylene oxide) (PEO)/clay nanocomposites were reported using Li-MMT

[22] and Na-MMT [23]. Erceg et al. [22] reported the synthesis of different con-

centration of PEO/Li-MMT viamelt intercalation at 90 ∘C for 8 h in vacuum oven.

The maximum value of interlayer distance of Li-MMT was reported, according

to SAXS, to be 1.88 nm (18.8Å) for 70/30 PEO/Li-MMT nanocomposite. This

increase amounts to 56.7% of Li-MMT original value, indicating an intercalated

structure. On the other hand, Na+-modified MMT was used in the synthesis

of PEO/clay nanocomposites, as reported by [23]. XRD results showed that the

gallery size remained the same (8.3Å) at different PEO loadings when prepared

via melt intercalation unlike when prepared via solution intercalation, as shown

in Figure 1.7. This was explained to be because of the stretching of PEO chains

as they enter the silicate gallery at low PEO loading in solution intercalation.

However, at higher loading, PEO chains reduce their length to accommodate

more PEO chains, thus expanding the gallery to 8.3Å for concentrations higher

than 15%. In melt intercalation, the PEO chains diffuse into the silicate gallery

03

4

5

6

7

8

9

5 10 15 20

Melt intercalationSolution intercalation

Shouder value

25 30 35 40

PEO content (wt%)

Galle

ry s

ize (

Å)

Figure 1.7 Gallery size of PEO/MMT nanocomposites prepared from melt and solution

intercalation at different PEO loadings. Reproduced from [23] with permission from Elsevier.


while maintaining their helical structure, achieving the final gallery spacing from

the start.

1.2.2

Exfoliation Adsorption

Solution intercalation method can be generally divided into several substeps [24]:

(i) dispersion of nanotubes in a solvent by agitation, (ii) mixing of nanotubes and

polymer solutions by agitation, and (iii) controlled evaporation of solvent and/or

precipitation of nanocomposite. Unlike in melt intercalation, the driving force

behind exfoliation adsorption is the entropy gained by the desorption of solvent

[2, 3].This compensates the decreased entropy of the confined intercalated chains.

This method is considered good for the intercalation of polymers with little or no

polarity [2].

1.2.2.1 Solution Intercalation

Elastomer/graphene nanocomposites were prepared by solution intercalation,

as demonstrated in Figure 1.8 [25]. Graphene platelets (∼3 nm in thickness)

700 °C for 1 min

Raw GICs SBR (gum)

Dissolving in THFThermal shock

Add

THF

Ultrasonication

Mechanical mixing

Ultrasonication

THF evaporation

Precipitation and drying

85 °C using around-bottom flaskwith condenser

Figure 1.8 Synthesis flowchart for SBR/graphene nanocomposite by solution mixing.

Reproduced from [25] with permission from Elsevier.


were obtained from graphite-intercalated compound (GIC) by exposing them

to thermal shock and treating them in tetrahydrofuran (THF) solvent while

being ultrasonicated. The suspension was then added to the SBR mixture and

mechanically mixed at 200 rpm followed by sonication for 1 h below 30 ∘C.Evaporation of the solvent was done till 60 ∘C by mechanical stirring in which

60% was evaporated and at 60 ∘C, ethanol was used to precipitate, collect, wash,

and dry the nanocomposite powder. According to XRD and TEM, intercalated

structures were obtained. Moreover, the authors compared those results with

those obtained from melt mixing, and better exfoliation and dispersion was

achieved in the former. This is because more interlayer spacing is available for

polymer to intercalate. This was validated with the lower percolation threshold

and higher mechanical properties obtained.

Bian et al. [26] reported the synthesis of poly(propylene carbonate) (PPC)/

modified graphite oxide (MGO) nanocomposites via solution intercalation.

MGO was first dispersed in 25ml dimethylformamide (DMF) for 30min and

then mechanically stirred for 10min. PPC was then added to the dispersion

and stirred for 24 h at 40 ∘C. Evaporation of the solvent was done in a Petri

dish under vacuum at room temperature. The modification of GO (graphite

oxide) was necessary considering the incompatibility of hydrophobic PPC

with the hydrophilic GO. Therefore, hydroxyl groups were grafted on the GO

surface in order to enhance the interfacial adhesion and promote nanocomposite

formation. According to XRD results, a d-spacing of 1.7 nm was achieved in

PPC/MGO nanocomposites, which is 1.4 nm greater than that in natural graphite

powder (= 0.335 nm). This indicated that intercalated/exfoliated structures were

obtained. Moreover, enhanced thermal and mechanical properties were obtained

as a result of good dispersion of MGO in PPC matrix.

PS/modified laponite clay nanocomposites were synthesized as reported by

[27]. Modification of laponite was performed by an ion-exchange reaction with

the cationic surfactant cetyltrimethyl ammonium bromide (CTAB). This was

done to enhance the compatibility between the clay mineral and the hydrophobic

polymer matrix. Good compatibility was achieved as PS chains intercalate into

the interlayer spacings of laponite as observed by SEM. However, with increasing

laponite, clay loading, aggregation, and agglomeration were observed in the

nanocomposite.

Gu et al. [28] reported the synthesis of elastomer/organo-MMT nanocompos-

ite via solution intercalation. First, the organo-modified MMT was dispersed in

a solvent oil before adding it to the cis-1,4-polybutadiene rubber (BR) solution.

The mixture was stirred for 30min at 60 ∘C and then the solvent was evaporated.

The nanocomposite powder was then compounded and cured for specimen

preparation. Intercalated structures were obtained as determined by XRD and

TEM results in which d-spacing increased from 1.55 nm, for the original MMT,

to 3.63 nm in the BR/organo-MMT nanocomposite.

Polyamide (PA)/MWCNTsnanocomposites synthesized via solutionmixing are

reported in the literature [24, 29]. Functionalized CNTs better disperse the filler

in the polymer matrix, as compared to pristine CNTs [29]. Moreover, the use of


initiators to create polymer grafted nanotubes would also help in dispersion [24].

This is because of the enhanced interfacial interaction between the polymermatrix

and CNTs. In both cases, good dispersion of CNTs was achieved throughout the

polymer matrix.

Another use of MWCNTs as filler materials was reported by Marroquin et al.

[30]. The authors reported the synthesis of a novel material based on chitosan.

Fe3O4/MWCNT/chitosan nanocomposites were prepared by solution mixing

according to the schematic in Figure 1.9. Fe3O4 andMWCNTwere ultrasonicated

for 1 h in distilled water before adding chitosan and acetic acid. The mixture was

magnetically stirred for 2 h followed by ultrasonication for 30min. The mixture

was degassed and vacuum dried to obtain the nanocomposite films. Intercalation

with good dispersion was achieved as noted from XRD results following the

disappearance of the peak in theMWCNT signal at 2𝜃 = 26∘ from nanocomposite

signals. Fe3O4 acted as an antiplasticizer agent that led to higher crystallinity and

thus better electrical and mechanical properties.

Zeng et al. [31] and Chen et al. [32] reported the synthesis of PMMA/MWCNT

nanocomposite foams via solution mixing. Solvent casting and antisolvent pre-

cipitation methods were used by Zeng et al. [31] to prepare the foams in order to

investigate themethodology impact on foammorphology and properties.The for-

mer involves evaporating the solvent whereas the latter utilizes another solvent to

precipitate the nanocomposite from the main solvent. In both cases, uniform dis-

persion ofMWCNTs increased the bubble density and reduced cell size. However,

much notable results were reported for the modified antisolvent precipitation

method that involves suspending CNTs in a solvent before adding to the polymer

solution [31, 32].

In addition to foams, Shirazi et al. [33] used solution casting and solvent

evaporation methods to synthesize polyvinyl alcohol (PVA)/MWCNT nanocom-

posite membranes. On the other hand, Chen et al. [34] used the coprecipitation

process to graft poly(3,4-ethylenedioxythiophene) hollow spheres (b-PEDOT)

on MWCNTs and to wrap MnO2 nanograins on the b-PEDOT. MnO2/b-

PEDOT/MWCNTs hybrid nanocomposite was synthesized as a result and was

used to prepare a microsupercapacitor device.

1.2.2.2 Emulsion Polymerization

PS/carbon black (CB) nanocomposites were prepared by emulsion polymeriza-

tion [35]. Synthesis was carried out by first manually mixing CB with styrene

monomer at room temperature. A viscous paste was formed as carbon absorbed

the monomer. A surfactant was added to reduce the viscosity of the system. This

was followed by the addition of Azobisisobutyronitrile (AIBN) initiator to pre-

pare emulsified monomer droplets. In order to disperse the system, a surfactant

solution was added in the presence of ultrasound. Eventually, the dispersion was

sent to the reactor for polymerization to take place. The conditions were set to

be 60 ∘C, 350 rpm mixing speed, and 120min reaction time. According to TEM

results, as shown in Figure 1.10, twomain results were obtained: particle diameter

close to 50 nm and high polydispersity and a layer of CB surrounding the polymer


MWNTs Fe3O4

Fe3O4

Chitosan

-Ultrasonication

-Stirring

-Heating/Vacuum

++

Figure 1.9 Schematic of Fe3O4/MWCNT/chitosan nanocomposite synthesis by solution

mixing. Reproduced from [30] with permission from Elsevier.


100 nm

100 nm50 nm

100 nm

(a) (b)

(c) (d)

Figure 1.10 TEM images of PS/CB nanocomposite at: (a) 15 k×, (b) 27.5 k×, (c) 38 k×, and(d) 50 k×. Reproduced from [35] with permission from Elsevier.

particles, which is because of carbon primary aggregates being modified during

the dispersion stage.

Hassan et al. [36] and Hu et al. [37] reported the synthesis of PS/graphene

nanocomposites. Using sodium dodecyl sulfate (SDS) as a surfactant and sta-

bilizing agent, and ultrasonication, graphene sheets can be obtained from the

expanded graphite (EG) that are in turn prepared from the thermal shock of

GIC [36]. Graphene nanosheets also can be obtained using hydrazine hydrate

in the reaction mixture to reduce GO sheets into graphene [37]. Graphene

dispersion was then mixed with styrene monomer, potassium persulfate (KPS)

initiator, sodium bicarbonate (NaHCO3) buffer, water, and SDS in a reactor [36].

Conditions were set to 70 ∘C, 350 rpm, and 3 h reaction time [36]. Figure 1.11

illustrates the synthesis procedure in [37]. Good dispersion and exfoliation was

achieved in the final nanocomposite.

Another graphene nanocomposite was prepared by Kuila et al. [38] using

PMMA as the polymer matrix. The polymerization procedure is similar to that

reported by Hu et al. GO solution was ultrasonicated before adding SDS aqueous

solution. AIBN and styrene monomer were added to the stirred dispersion.

Hydrazine monohydrate was added to the mixture that underwent reflux for


OHO

O

O O

COOH HOOC

HO

OH O

O

OH

COOH

OO

COOH

OH

COOH

COOHHO

HOOC

HO

OH O

O

OH

COOHHydrazine hydrate

ReductionOO

COOH OH

COOH

COOHHO

COOH

COOH

Styrene, SDS

Ultrasonication for 15 min In situ polymerization

K2S2O8 (KPS)

OHCOOH

Graphene oxide nanosheets

Graphene oxide nanosheets-polystyrene microspheres

Graphene nanosheets-polystyrenemicrospheres

Styrene-linked graphene oxide nanosheets

HOOC

OH

HOHO

Figure 1.11 Schematic of PS/graphene nanocomposite synthesis. Reproduced from [37]

with permission from Elsevier.

additional 16 h to reduce GO to graphene sheets. Eventually, the mixture was

precipitated with dilute hydrochloric acid (HCL) and vacuum dried to obtain the

nanocomposite. When characterized by XRD, the nanocomposite signals did not

show the GO peak. This indicated that GO was successfully reduced to graphene

sheets and that their periodic structure was destroyed. According to TEM, the

graphene layers were distributed uniformly forming a continuous network.

Polyaniline (PANI)/activated carbon (AC) nanocomposites were synthesized by

Oh and Kim [39] using dodecyl benzenesulfonic acid (DBSA). DBSA was used as

surfactant and dopant that participated positively in the synthesis of PANI/AC

nanocomposites. AC and DBSA aqueous solution were sonicated before adding

the aniline monomers followed by intiator. Once the polymerization completed,

ethanol was added to precipitate the nanocomposite. The nanocomposite struc-

ture can be represented by the schematic in Figure 1.12. It was noted from SEM

that with increasingDBSA concentration, the roughness of DBSA-PANI films that

cover the surface of AC increases.

Similar to CNTs, inorganic halloysite nanotubes (HNTs) were used as fillers to

HIPS nanocomposites [40]. HNTswere uniformly dispersed in thematrix because

of PS nanospheres formation on the surface ofHNTs, as shown in Figure 1.13.This

was prepared by first dispersing the dried HNTs in aqueous SDS. Ammonium

persulfate and styrene monomers were added to the stirred solution. Polymer-

ization was done under argon blanket at 70–75 ∘C and 400 rpm for 18 h. HNTs

were also used as filler in epoxy matrix reported by Ye et al. [41]. However, in this

case, HNTs were not uniformly dispersed in the hybrid material that contained


Activated carbon Aniline monomer PANI

DBSA DBSA–anilinium cation complex

Figure 1.12 Schematic of PANI/AC nanocomposite synthesis. Reproduced from [39] with

permission from Elsevier.

100 μm 1 μm

(a) (b)

Figure 1.13 (a,b) SEM images of HIPS/HNT nanocomposites. Reproduced from [40] with

permission from Elsevier.

carbon fibers. Instead, HNT-rich regions were obtained and were considered as

rigid composite particles with highHNT content.This was determined from SEM

images, as shown in Figure 1.14. The hybrid material was prepared by dispersing

HNTs in acetone while mechanically stirred. Epoxy resin, followed by a curing

agent, was added to the degassed mixture. The laminates were then placed in alu-

miniummold to be cured in a hot pressing agent.They were precured at 80 ∘C for

2 h and postcured at 160 ∘C for another 4 h.

Ultrasound can be used to synthesize nanocomposites in emulsion polymer-

ization. Examples are reported by Cetintas and Uyanık [42] and Bhanvase et al.

[43]. For instance, to synthesize PS/clay nanocomposites, potassium hydroxide

and SDS were dissolved in water in three neck round-bottom flask [42]. Mean-

while, styrene monomer and clay minerals were stirred in an ultrasound bath at

0 ∘C.The two solutions were then mixed together and potassium peroxodisulfate

initiator was added. Eventually, the temperature was raised to 50 ∘C to start the

polymerization reaction that lasted 24 h. Finally, the nanocomposite was obtained

by precipitation, washing, and vacuum drying. Exfoliated nanocomposites were

prepared as determined by XRD results. This was supported by Bhanvase et al.


15 kV X100 HKUST HKUST SEI 5.0 kV X2.000 WD 8.3 mm10 μm

HKUST SEI 5.0 kV X7.000 WD 7.8 mm1 μmHKUST SEI 5.0 kV X10000 WD 8.3 mm1 μm

100 μm

(a) (b)

(c) (d)

Figure 1.14 (a–d) SEM images of epoxy/HNT/carbon fiber hybrid nanocomposites.


[43] as their poly(methyl methacrylate-co styrene)/montmorillonite [P(MMA-co-

St)/O-MMT] nanocomposite was found to be exfoliated with the use of ultra-

sound.This was determined by XRD as no peaks appeared in the nanocomposite,

as shown in Figure 1.15.

1.2.3

In Situ Polymerization

Several advantages are attributed to in situ polymerization. First of all,

thermoplastic- and thermoset-based nanocomposites can be synthesized

via this route [3]. In addition, it permits the grafting of polymers on filler surface,

which can generally improve properties of the final composite. Partially exfoliated

structures can be attainable with this method because of the good dispersion and

intercalation of fillers in the polymer matrix. Abedi and Abdouss [4] state that

in situ polymerization is the most suitable preparation method for polyolefin/clay

nanocomposites because of its lack of rigorous thermodynamic requirement

compared to the other methods.

Guo et al. [44] reported the synthesis of graphene, GO, and functionalized

GO – Epoxy nanocomposites via in situ polymerization. The synthesis was


4

1000

2000

3000

4000

5000

6000

7000

8000

6 8 10 12 14

2θ (°)

A

1.73 nm

1.67 nm

A - Bare MMT

B - Modified MMT

C - Poly(MMA-co-Styrene)

D - Poly(MMA-co-Styrene)/MMT

B

C

D

Inte

nsity (

CP

S)

Figure 1.15 XRD signals for: (A) pristine clay, (B) O-MMT, (C) poly(MMA-co-St) polymer, and

(D) poly(MMA-co-St)/O-MMT nanocomposite with 4% O-MMT loading. Reproduced from [43]


carried out by first dispersing the filler in acetone by ultrasonication. The dis-

persion was then added to the epoxy matrix before placing it in a vacuum oven

at 50 ∘C. m-Phenylenediamine was added when 80% of the solvent evaporated,

accompanied by vigorous stirring. Eventually, the mixture was poured into a

stainless steel mold, dried at 60 ∘C for 5 h to remove the residual solvent, precured

in an oven at 80 ∘C for 2 h, and postcured at 120 ∘C for two additional hours to

obtain the composites. TEM images, in Figure 1.16, show that better dispersion

was achieved in epoxy/graphene and epoxy/functionalized GO nanocomposites

compared to epoxy/GO composites. Bundles of GO were visible following Van

der Waals and hydrogen bond interactions between GO sheets. On the other

hand, absence of polar groups and better interfacial interactions were the reasons

behind better dispersion and hair-like structure for other composites.

However, Huang et al. [45] reported good dispersion of GO in PP matrix

as evaluated in TEM and SEM. In order to do so, Zeigler-Natta (ZN) cata-

lyst was incorporated into GO sheets in the process shown in Figure 1.17.

Grignard reagent (RMgCl) was used prior to adding titanium tetrachloride to

synthesize GO-supported ZN catalyst. This catalyst was then added at 60 ∘C to

hexane–propylene liquid mixture that is subjected to vigorous stirring. Triethyl

aluminium (AlEt3) and dimethoxydiphenylsilane (DDS) initiators were added

to the mixture to initiate the polymerization reaction. The final composite was

obtained by filtering, washing, and drying.

Other reports of GO composites include PMMA/GO [46] and polypyrrole

(PPy)/GO [47]. Exfoliated structures were obtained for both nanocomposites,

as suggested by XRD studies. However, according to TEM, agglomeration of

GO sheets in PMMA/GO nanocomposite was visible at higher loadings above


EP/1%GO

EP/1%FGO

EP/1%Graphene

200 nm(a)

(b)

(c)

200 nm

200 nm

Figure 1.16 (a–c) TEM images of epoxy/graphite nanocomposites. Reproduced from [44]

with permission from American Chemical Society.


HOHO

OH OH

CI

O

O

CI

CICI

CICICICICI

CI CI

CICI

CI

RR

O

CICI

MgOMgCI

OMgCIMg

CICI CI

CI

CICICICICI

Ti Ti

OMgCI

CIMgO

CICI

TiTi

OH

CI CI CI CI

OR

RC3H6

AIEt3

CI

CICICICI Ti Ti

Ti Ti

Mg

Mg

OMgCI OMgCI

TiCI4

OMgCI OMgCI

OMgCIOMgCI

CIMgOCIMgO

R

RR

R

GORMgCI/GO

TiCI4/(RMgCI/GO)PP/GO nanocomposites

OH

RMgCIO

O O

O

Figure 1.17 Schematic of PP/GO nanocomposite synthesis. Reproduced from [45] with

permission from American Chemical Society.

1wt% [46]. PPy/GO composites were synthesized via liquid–liquid interfacial

polymerization, as shown in Figure 1.18. The reason behind the authors using

this method instead of the conventional in situ polymerization method was its

slower and controllable attributes. Moreover, bulk quantities can be prepared by

this method.

Intercalated and exfoliated PE/graphite nanocomposites were reported by

Fim et al. [48]. GIC was first exposed to thermal shock to obtain the EG. In

turn, the suspension of EG/ethanol was treated in an ultrasound bath to attain

graphite nanosheets (GNSs). Methylaluminoxane (MAO) was used to treat GNS

surfaces and as a cocatalyst along with bis(cyclopentadienyl)zirconium dichloride

(Cp2ZrCl2). The polymerization conditions were as follows: 70 ∘C, toluene as

solvent, 2.8 bar ethylene pressure, and 30min. Table 1.1 summarizes the XRD

data for the nanocomposites. It is noted that with thermal and ultrasound treat-

ment, graphite sheets exfoliated, increasing their interlayer spacing. Moreover,

crystal size decreased following agitation and dispersion of graphite, eventually

reducing the number of stacked graphene sheets. The 5.6wt% graphite loading

nanocomposite yielded good dispersion with higher interlayer spacing and

smaller crystal size. This is because of the polymer chains growing in between

the GNSs.

Graphene was used in preparing many nanocomposites via in situ polymer-

ization such as nylon-6 (PA-6) [49] and poly(butylene terephthalate) (PBT)

[50] – graphene composites. Moreover, ring opening polymerization was used

to prepare those nanocomposites. In both cases, good dispersion of graphene

was achieved because of the enhanced interfacial interactions [49, 50]. Table 1.2

summarizes XRD results for PBT/graphene nanocomposites. It is noted that at


Graphite

Graphite oxide

Water

Water

Chloroform ChloroformPyrrole

Before polymerization

InterfaceAfter

polymerization

After 24 h

Polymerizationat interface

Product

Graphene oxidesheets

Ultrasonication

30 min

H2SO4

GO, FeCI3

KMnO4

Figure 1.18 Schematic of liquid-liquid interfacial polymerization of PPy/GO nanocompos-

ites. Reproduced from Ref. [47] with permission from Elsevier.

Table 1.1 XRD results of graphite, GNS, and PE/graphite nanocomposites.

Sample 2𝜽 (∘) d002 (nm) Crystal size, C (nm)

Graphite flake 26.67 0.333 58.38

GNS 26.52 0.336 28.15

PE/graphite 1.2% 26.53 0.336 24.77

PE/graphite 5.6% 26.42 0.338 14.58

Reproduced from [48] with permission from Wiley Interscience.

1wt%, d-spacing decreased and this was attributed to the strong π–π interactionsbetween graphene sheets that did not permit polymer intercalation.

Clay nanocomposites prepared by in situ polymerization are reported using

many polymers such as PAs [51], PP [52], polybenzoxazine (PBz) [53], and

polysulfone (PSU) [54]. Puffr et al. [51] reported the synthesis of PA-6, PA-8,

PA-12, and MPA12 (N-methyl-polyamide 12)/organo-MMT nanocomposites.

The MMT was modified by cationic exchange in which 12-aminododecanoic

acid (ADA) was used to intercalate the clay mineral. The intercalated MMT

with lactam monomers and ADA were blended together as a solid mixture,

melted, and then sent to the glass ampoules for polymerization to take place

at 260 ∘C. XRD results showed that the nanocomposites produced were exfo-

liated or with d-spacing higher than 6 nm. Regarding PP/clay nanocomposites,


Table 1.2 XRD results of graphene and PBT/graphene nanocomposites.

Sample 2𝜽 (∘) d002 (Å)

Graphene 26.403 3.373

PBT/graphene 0.5% 26.348 3.380

PBT/graphene 0.75% 26.326 3.383

PBT/graphene 1% 26.408 3.372


different clay-supported magnesium/titanium ZN catalysts were used and were

investigated by Dias et al. [52]. Slurry polymerizations at 70 ∘C and 2 bars

were conducted to synthesize the nanocomposites. It was determined that the

performance of the catalyst to yield exfoliated/intercalated structures depends

on the clay mineral and the synthesis conditions. PBz/organo-modified MMT

nanocomposites were synthesized by thermal ring-opening polymerization [53].

The intercalated benzoxazine (Bz)-MMT clay was first prepared by ion-exchange

reaction and was then dispersed in fluid Bz monomers by mechanical stirring,

as shown in Figure 1.19. The cast films were cured at 240 ∘C for 3 h in air oven

for polymerization to take place. XRD and TEM results revealed that partially

exfoliated/intercalated structures were obtained. Similarly, Dizman et al. [54]

reported the synthesis of exfoliated/intercalated PSU/organo-modified MMT

nanocomposites. They were achieved via in situ photo-induced cross-linking

polymerization. Sixteen Philips 8W/06 lamps emitting light at 𝜆> 350 nm were

used as a source of irradiation. Figure 1.20 shows the TEM images of PSU/MMT

nanocomposites in which “e” refers to exfoliation and “i” to intercalation.

Another composite synthesized via in situ polymerization is poly(ethylene

terephthalate) (PET)/LDH by Cui et al. [55]. Terephthalate-intercalated LDH

were first dispersed in ethylene glycol and then mixed with dimethyl tereph-

thalate (DMT) and manganese acetate and magnesium acetate as catalysts. The

synthesis was carried out in two steps: ester interchange reaction at 190–230 ∘Cand polycondensation reaction at 280 ∘C. Partially exfoliated structures were

achieved as revealed by morphological studies.

+Na+

Na+

NaNaNa+

Na Na Na Na

+BPy

+BPy

+BPyBPy

(BPy+)

N

O

N OHBr−

+ 11

N OH

O

5H2O, 3 days

(Na-MMT) (qBPy-MMT)

Polybenzoxazine/MMT

nanocomposite

Fluid benzoxazine

BPy+

BPy+

BPy+

BPy

Na+

Na

+Na

+Na

+Na Na

+

+

+

Figure 1.19 Schematic of PBz/MMT nanocomposite synthesis. Reproduced from [53] with

permission from Wiley Periodicals.


i

e

e

e e

e

e

e

e

e

e

e

e

ii

i

e

e

i

i

e

i

i

i

i

i

ii

i

ii

i

e

e

e

e

ee

e

e

e

e

e

e

e

e

e

e

20 nm

(a) (b) (c)

(a) (b) (c)

50 nm 50 nm 50 nm

20 nm 20 nm

Figure 1.20 TEM images of PSU/MMT nanocomposites at: (a) 1wt% (b) 3wt% (c) 5wt%

in high magnification at top and low magnification at below images. Reproduced from [54]

with permission from WILEY-VCH Verlag GmbH & Co. KGaA.

Dash et al. [56] reported the synthesis of poly(anthranilic acid) (PAnA)/

MWCNT composites via in situ chemical oxidative polymerization. The CNTs

were first functionalized using H2SO4 and HNO3 to provide carboxylic acid

groups at the surface. Then, the functionalized MWCNTs were sonicated in

a 1.2-M HCl solution for 2 h before adding aniline and anthranilic acid to

the suspension. Ammonium persulfate reagent in HCl solution was added to

the mixture and mechanically stirred. The copolymer products obtained were

filtered, washed, and vacuum dried. SEM analysis showed that the diameter of

the nanocomposite increased with increasing MWNT loading as PAA coated

itself on the outer surface of the nanotubes.This coating happened because of the

strong interactions between the comonomer (i.e., aniline) and the functionalized

MWNTs, as suggested by the authors. Using a similar procedure, Li and Kim [57]

reported the synthesis of PANI/MWCNT composites for sensor applications.

Core and shell structures were visible in SEM images, which signal the typical

structure of polymer-grafted nanocomposites.

Wu and Liu [58] prepared PS/MWCNTs via solution-free radical in situ poly-

merization. Without any pretreatment of MWCNTs, they were combined with

styrenemonomers, toluene, and AIBN initiators.Themixture was heated at 90 ∘Cfor 11 h and the product was precipitated and vacuum dried. Fourier transform

infrared (FTIR) spectroscopy analysis concluded the successful grafting of PS onto

the walls of CNTs. Qualitative relationships between initiator and temperature


Table 1.3 Effect of polymerizing conditions on monomer conversion and polymer grafting

percentages for PS nanocomposites.

MWCNTs-PS Polymerizing temperature (∘C) AIBN added (g) C%of St PG%

1 90 0.01 9.9 2.9

2 90 0.02 30.5 4.9

3 90 0.05 39.0 15.6

4 90 0.10 55.2 4.2

5 90 0.15 58.3 0.8

6 90 0.20 59.7 0.8

7 80 0.5 34.1 2.2

8 70 0.5 19.0 1.5

9 60 0.5 13.1 0.9

10 50 0.5 9.0 0.6

Reproduced from [58] with permission from Taylor & Francis.

with monomer conversion and polymer grafting were established by the authors,

as shown in Table 1.3.

1.2.4

Nontraditional Methods

In order to facilitate better dispersion of the filler in the polymer matrix for

improved properties of final composites, researchers investigated different routes

based on the traditional methods mentioned earlier. For instance, in situ poly-

merization can be customized to be redox [59, 60] or catalytic chain transfer [61]

or even photo-induced polymerizations [54]. Others include microwave-induced

synthesis [62, 63], one-pot synthesis [64–66], template-directed synthesis [67],

electrochemical synthesis [68], self-assembly synthesis [69, 70], and intermatrix

synthesis (IMS) [71–74].

As the name implies, one-pot synthesis refers to a sequence of reactions being

carried out in the same reactor. As it refers to a location, this mode can be

inclusive of other synthesis methods. For instance, Hwang et al. [66] reported

the synthesis of tin (Sn)-embedded carbon-silica polymer nanocomposites. Even

though it is a one-pot synthesis, the preparation was conducted via self-assembly

method. Through the selective interaction of resol (carbon precursor), tetraethy-

lorthosilicate (TEOS), and tributylphenyltin (Sn precursor) with an amphiphilic

diblock copolymer, poly(ethylene oxide-b-styrene) (PEO-bPS), unique struc-

tures of nanowires, or nanoparticles, were achieved, as shown in Figure 1.21.

It was reported that Sn was uniformly embedded in the rigid carbon-silica

matrix.

Self-assembly, as the name implies, dictates the spontaneous arrangement

of the existing components following local interactions among the compo-

nents. As a result, ordered structures can be obtained as illustrated by Liu


Silicate =

oligomer

Si

SiSi

Si

O

O

O OH

OH

CH2OH

HOH2C

H3C

H3C

H3CBr

m

O

O

On

CH2OH

H2

OHOH

C

CH2

Sn

Resol = TBPT = -b-PS

hydrophilic

Selectiveincorporation

hydrophobic

Block copolymer/precursors in THF

Sn nanoparticles/CS (Sn-49-CS)

Sn nanowires/OMCS (Sn-8.5-CS)

Self assembly

Sn

Sn

700 °C, N2

PFO

Figure 1.21 Schematic of Sn/carbon-silica composite synthesis. Reproduced from [66] with

permission from American Chemical Society.

et al. [69]. Graphene-polymer composites were prepared by π–π stacking as

shown in Figure 1.22. Pyrene-terminated Poly(2-N,N′-(dimethyl amino ethyl

acrylate)) (PDMAEA) and pyrene-terminated poly(acrylic acid) (PAA) were

first dissolved in aqueous graphene solution and sonicated. Excess polymer

was removed by centrifugation at 14 000 rpm for 30min. The precipitate was

collected and redispersed in water to obtain the final composites. Layer-by-layer

graphene-polymer composites were prepared by electrostatic interactions via

self-assembly.

IMS is used to synthesize polymer stabilized metal nanoparticles (PSMNPs)

[74]. In order to use this method, the polymer matrix must possess some

functional groups capable of binding the nanoparticles. There are two versions

in which IMS can be performed to prepare PSMNP-based sensors: in situ and ex

situ [72]. The former deals with depositing the neat polymer onto the electrode

surface followed by metal loading and metal reduction either by chemical or

electrochemical means. The latter deals with dissolving the PSMNP-polymer

nanocomposite in a solvent to form an ink. This ink can then be easily deposited

on the electrode surface. Ruiz et al. [72] reported the synthesis of monometallic

Pd-PSMNPs in sulfonated poly(ether ether ketone) (SPEEK) using intermatrix

approach. It was noted that the properties of membranes prepared depended on

the preparation route and reduction method. Conversely, Domènech et al. [74]

stated that SPEEK possesses high hydrophilicity which limits its applications.

Therefore, sulfonated polyethersulfone with Cardo group (SPES-C) was used as a

polymer matrix to prepare catalytic membrane reactors by phase-inversion.


O

O

SS S S

DMAEA or AA

Self-assembly

or

R = N R = HOO O O

AIBN/65–70 °CS

O

O

O

OS

S S S

S

Graphene sheet

R

nO

O

++

+

+

++

+

+

+

+

+

+

++

−

−−

−−

−

−

−

−

−

−

−+

Figure 1.22 Schematic of graphene-polymer composite synthesis. Reproduced from [69]

with permission from American Chemical Society.

In situ electrochemical synthesis is considered useful for the quick detection

of current–voltage characteristics. Ameen et al. [68] reported the synthesis of

PANI/graphene nanocomposites via this method. This synthesis was performed

in a three-electrode system: fluorinated tin oxide glass (FTO) as working elec-

trode, platinum wire as counter electrode, and reference electrode. Graphene

oxide and aniline monomers were dispersed in HCl. This permitted the aniline

to form its salt with a positive charge and to be adsorbed on to the surface of

graphene oxide. Following the electrostatic interactions between the components,

a homogeneous mixture of graphene oxide/aniline was obtained. The suspension

was spread on FTO substrates by spin coat and then dried in a vacuum oven. The

applied potential of −1.0 to +1.0V with scan rate of 0.02V s−1 was used for the

oxidation and polymerization of aniline on the surface of graphene oxide and

the simultaneous reduction to graphene. Figure 1.23 shows the Field-Emission

scanning electron microscope (FESEM) and TEM images of graphene and the

composite.

Microwave-assisted method has considerable advantages such as rapid vol-

umetric heating, high reaction time, enhanced reaction selectivity, and energy

saving behavior [62]. Cellulose–silver nanocomposites were prepared using

microcrystalline cellulose and silver nitrate in ethylene glycol as a solvent [62]. In

addition, ethylene glycol is useful as a reducing agent and a microwave absorber.

Through this route, silver nanoparticles were formed in situ on the cellulose

surface. According to SEM, silver particles were homogeneously dispersed in the

cellulose substrate.


18 10.0 kV 5.3 mm×30.0 k SE(U) 1.00 um 1610.0 kV 5.4 mm×10.0 k SE(U) 5.00 um

0.1 μm 0.2 μm

PANI/Gr

(a) (b)

(c) (d)

Figure 1.23 FESEM images of: (a) graphene, (b) PANI/graphene composite and TEM images

of, (c) graphene, and (d) PANI/graphene composite. Reproduced from [68] with permission

from Elsevier.

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31

2

Masterbatch Approach to Generate HDPE/CPE/Graphene

Nanocomposites

Ali U. Chaudhry and Vikas Mittal

2.1

Introduction

Polyolefins have superior rank among commodity plastics owing to their use in a

variety of applications. Especially, high-density polyethylene (HDPE) has a wide

range of properties including low cost, ease of recycling, good processability,

nontoxicity, biocompatibility, and good chemical resistance. The demands of

modern-day applications, however, are not met solely by neat polymers. There-

fore, in addition to the voluminous applications of neat HDPE, it is required to

improve the performance of the polymer in terms of properties such as stiffness

and rigidity by forming composites [1, 2]. In general, composites with polyolefin

matrices are formed to fulfill various requirements for different applications

where cost and weight reduction, dimensional stability, opacity, heat stability,

and processability are required. The advances in nanofillers and nanofibers have

made it possible to produce lightweight nanocomposites with better physical

and mechanical properties at a low filler concentration than conventional

composites. This is owing to the much higher number of interfacial contacts

of inorganic particles with the polymer chains in such nanocomposites, which

generate a completely different interfacial morphology as compared to the bulk

polymer [3, 4].

Graphene, which consists of one atomic thick sheets of covalently sp2-bonded

carbon atoms in a hexagonal arrangement, has already received the attention

of researchers for the generation of polymer nanocomposites [5]. Its choice

as a filler is its excellent electrical and mechanical properties, which are sig-

nificantly better than other inorganic filler materials. A single defect-free

graphene layer has Young’s modulus of≈ 1.0 TPa, intrinsic strength≈ 42Nm−1,

thermal conductivity≈ 4840–5300W (m⋅K)−1, electron mobility exceeding

25 000 cm2 V−1⋅s−1, excellent gas impermeability, and specific surface area

of≈ 2630m2/g [5]. All these properties make this material even superior to

carbon nanotubes (CNTs) for use in polymer nanocomposites. A number of

studies on polymer nanocomposites based on graphene have been published

in a short span of time since its development [5–13]. The parent material for


32 2 Masterbatch Approach to Generate HDPE/CPE/Graphene Nanocomposites

graphene is graphite, which is present prolifically in nature. Graphene (and

graphene oxide) can be produced from graphite by different methods, such

as thermal expansion of chemically intercalated graphite, micromechanical

exfoliation of graphite, chemical vapor decomposition, and chemical reduction

method of graphene oxide [7].Themost commonly usedmethod is the exfoliation

of intercalated graphite oxide by introducing graphite oxide suddenly to a higher

temperature [14].

The macroscopic properties of polymer nanocomposites are dependent mainly

on the thermodynamic factors such as interfacial compatibility of polymer and

filler phases, polarity match between the filler surface and the polymer chains,

and so on. In addition, nanoscale dispersion and the distribution of the filler also

depends on the size, shape, dispersion techniques and equipment, time of mixing

and applied shear, and so on. The full advantage of nanofillers could be achieved

only by considering these factors, which could lead to uniform transfer of superior

properties of nanofiller to host polymer matrix [4, 15–19].

Significant research effort has focused to attain the full potency of nanofillers

using different mixing techniques, modification of polymer backbone, or filler

surface, use of compatibilizer (functional polymers) and coupling agents, and

so on. In the case of graphene, it has very low surface energy as compared to

graphite oxide, which is a precursor to graphene.The presence of the less number

of functional groups (such as carboxyl, epoxide, and hydroxyl) on the surface

of pristine graphene leads to lower compatibility with polar polymer matrices,

resulting in poor dispersion and lower enhancement in polymer properties [5].

Similarly, the dispersion of polar graphene oxide in nonpolar polymers is not

optimal owing to the absence of positive interactions between them. One of the

routes to overcome this limitation is the functionalization of filler surface, which

results in significant enhancement of the mechanical and electrical properties

of polymer nanocomposites. Bing et al. grafted amine-functionalized multi-

walled carbon nanotubes (MWCNTs) with polyethylene by reactive blending

using maleic anhydride [20]. The improved stiffness, strength, ductility, and

toughness of the polymer was attributed to the uniform dispersion of nanofiller

and improved interfacial adhesion owing to grafted polyethylene on CNT.

Similar results were achieved by adding functionalized CNTs to polypropylene

matrix [18]. Ramanathan et al. prepared nanocomposites of functionalized

sheets of graphene and poly(methyl methacrylate) (PMMA) by sonication and

high-speed shearing of expanded graphene [21]. Partially oxygenated wrinkled

sheets showed a shift of 30 ∘C in Tg of PMMA, which was superior to that

obtained using single-walled carbon nanotubes (SWCNTs) and expanded

graphite platelets.

The other method described in the literature is the use of compatibilizer. In

the case of polyethylene, the lack of polar groups in its backbone is a consider-

able hurdle in homogenous dispersion and exfoliation of nanofillers. Introduc-

tion of amphiphilic compatibilizer that has polar and nonpolar groups, which

act as bridges between filler and host polymer, has resulted in improved filler

dispersion. Valdes et al. reported that introduction of ethylene acid copolymer

2.2 Experimental 33

compatibilizer in linear low density polyethylene (LLDPE) and clay nanocompos-

ite improved exfoliation of clay particles, which resulted in better thermal prop-

erties [22]. Masterbatch technique was used for the preparation of composites.

Similarly, Kim et al. used maleic anhydride as a bridge for the nanocomposites

of low density polyethylene (LDPE) and exfoliated graphite nanoplatelets [14].

Different dispersion techniques, that is, solution andmelt blending, alongwith dif-

ferent arrangements of screws, were used in the study. The better results in terms

of filler dispersion were shown by solution mixing followed by counterrotating

screw arrangements.

Chlorinated polyethylene (CPE) has also been reported to be efficient as adhe-

sion promoter and compatibilizer between polymer blends and fillers in compos-

ites [23]. In poly(vinyl chloride) (PVC) and wood flour composites using CPE

(chlorine content ranged from 25 to 42%) as a compatibilizer, improvements in

processing, melt strength, and elongation at break were observed. In a similar

work, Simon et al. showed the effect of acid-base interaction between chlorine

and hydroxyl group on the adhesion of chlorinated polypropylenewith polypropy-

lene [24]. Significant changes in the mechanical properties of blends of varying

amounts of CPE with HDPE have also been reported by Maksimov et al. [25].

In the current study, two types of CPE (25 and 35% chlorine content) have been

used as a compatibilizer in order to study the dispersion of graphene oxide in

HDPE and its effect on the resulting nanocomposite properties. Solution blend-

ing technique was used for the blending of compatibilizer and graphene oxide.

The obtainedmasterbatches were thenmelt mixed with HDPE.The effect of chlo-

rine content in the compatibilizer as well as the amount of compatibilizer on the

morphology, mechanical, and rheological properties of the polyethylene graphene

oxide nanocomposites were studied.

2.2

Experimental

2.2.1

Materials

CPE grades Weipren® 6025 (25% chlorine content, named as CPE25) and CPE

135A (35% chorine content, named as CPE35) were obtained from Lianda

Corporation, USA, and Weifang Xuran Chemicals, China, respectively. Matrix

polymer, that is, HDPE BB2581 was received fromAbuDhabi Polymers Company

Limited (Borouge), UAE. The polymer materials were used as obtained. The

specifications of the polymers as received from the suppliers are also reported

in Table 2.1. Graphite powder (325mesh) was procured from Alfa Aesur GmbH

and Co., Germany. Concentrated sulfuric acid (H2SO4, 95–98%), sodium

nitrate (NaNO3), and potassium permanganate (KMnO4) were supplied by S. D.

Fine Chemicals Ltd., India, Eurostar Scientific Ltd., UK, and Fisher Scientific,

UAE, respectively.


Table 2.1 Specifications of the polymers as received from the suppliers.

Property CPE25 CPE35 HDPE

Appearance White granules White powder Transparent pellets

Specific gravity, ASTM D792 1.1–1.3 1.1–1.16 0.958

Mooney viscosity ML (1+ 4) at

135 ∘C, ASTM D4603

82 — —

Melting point (∘C), ASTM D7138 — 180–190 147

Heat of fusion (J g−1), ASTM

D3418

45 2 —

MFR 190 ∘C/2.16 kg, g/10min,

ASTM D1238

— — 0.35

Heat deflection temperature

(0.45Nmm−2) (∘C), ASTM D648

— — 80

2.2.2

Preparation of Graphite Oxide and Graphene Oxide

Graphene oxide was prepared through thermal exfoliation of precursor graphite

oxide [26] using modified Hummer’s method [27]. A short description of the

preparation of graphite oxide and graphene oxide is as follows: 5 g of graphite

powder was added with 125ml concentrated H2SO4. Subsequently, 2.5 g of

NaNO3 was added to this mixture. The mixture was kept in an ice-bath (5 ∘C)under stirring. After 30min, 15 g of KMnO4 was added to the mixture, and the

temperature was allowed to rise gradually to 35 ∘C. The mixture was stirred for

2 h under these conditions. This was followed by the addition of deionized water

till the temperature increased to 100 ∘C. After further stirring for 15min while

maintaining the same temperature, the mixture was quenched and diluted by

pouring it into 1.5 l deionized water. 30% H2O2 was slowly added to the dilute

solution until the evolution of bubbles of hydrogen stopped. The solution was

then filtered using Buchner funnel to remove the nongraphite oxide (GO) waste.

The residues were dispersed in 2 l deionized water andwere addedwith dilute HCl

(6%) (2 l) to remove the SO42− ions. The dispersion was filtered and the filtrate

was analyzed for SO42− and Cl− ions using BaSO4 and AgNO3, respectively

(generation of white precipitates). The cleaning and filtration was continued until

no SO42− and Cl− ions can be observed in the filtrate. The washed GO was dried

under vacuum at 60 ∘C for 24 h [27]. Graphene oxide was generated via thermal

exfoliation of dried GO. The process was carried out by placing 1 g GO in a long

quartz tube with 25mm internal diameter and sealed at one end. The other end

of the quartz tube was closed using a rubber stopper. The sample was flushed

with nitrogen, followed by the insertion of the tube in a tube furnace preheated to

1050 ∘C.The tube was held in the furnace for 30 s [26].The density of the obtained

graphene oxide was measured by tapped density tester to be 0.0161 gml−1.

2.2 Experimental 35

2.2.3

Nanocomposite Generation

The nanocomposites were prepared by either solution mixing followed by melt

mixing or direct melt mixing, as described earlier [14]. The synthesis scheme is

described in Figure 2.1. Table 2.2 also describes the compositions of different

composites. For the solution mixing method, CPE was stirred in p-xylene (3%

solid content) at 100 ∘C under reflux until the solution became limpid. Graphene

oxide (weighed according to the requirement for masterbatches) was suspended

CPE + p-xylene

at 100 °C, 2 h

Graphene oxide + p-xylene

at room temperature for 1 h,

sonication 10 min

CPE and graphene oxide + p-xylene

at 100 °C, 30 min, sonication

Stirring and gradual cooling to

room temperature

Drying at room temperature

overnight and at 40 °C for 24 h

Masterbatch + HDPE

melt mixing

Figure 2.1 Schematic representation of the synthesis of CPE–graphene oxide master-

batches by solution mixing and subsequently HDPE nanocomposites by melt mixing.


Table 2.2 Compositions of the composites (in wt%).

Code Polymer/nanocomposite HDPE (%) CPE25 (%) CPE35 (%) Graphene oxide (%)

1 HDPE 100 — — —

2 HDPE/G 99.5 — — 0.5

3 HDPE/1%CPE25/G 98.5 1 — 0.5

4 HDPE/2%CPE25/G 97.5 2 — 0.5

5 HDPE/5%CPE25/G 94.5 5 — 0.5

6 HDPE/10%CPE25/G 89.5 10 — 0.5

7 HDPE/5%CPE35/G 94.5 — 5 0.5

8 HDPE/10%CPE35/G 89.5 — 10 0.5

in a little milliliters of p-xylene for 1 h at room temperature. In between, the

suspension was sonicated for 10min. The graphene oxide suspension was then

added to polymer solution at 100 ∘C and further sonicated for 15min. The

mixture was stirred and gradually brought to room temperature. The solution

was kept overnight at room temperature followed by 40 ∘C for 24 h in order

to remove any solvent residues, which resulted in dried CPE/graphene oxide

masterbatches.

To form nanocomposites, melt mixing of CPE/graphene oxide masterbatches

with HDPE was carried out at 190 ∘C using a mini twin screw extruder (Mini-

Lab HAAKE Rheomex CTW5, Germany). The screw length and screw diameter

were 109.5 and 5/14mm conical, respectively. Batch size of 5 g was used and the

shearmixing was performed for 5min at 60 rpm. Direct melt mixing of HDPE and

graphene oxide was also similarly performed. Pure HDPE was also processed by

subjecting it to similar shear and thermal conditions.

Disc- and dumbbell-shaped test specimens were injection molded using a mini

injection molding machine (HAAKEMiniJet, Germany) at a processing tempera-

ture of 190 ∘C.The injection pressurewas 700 bar for 6 s, whereas holding pressure

was 400 bar for 3 s. The temperature of the mold was kept at 50 ∘C.

2.2.4

Material Characterization

Calorimetric properties of nanocomposites were recorded on a Perkin-Elmer

Pyris-1 differential scanning calorimeter under nitrogen atmosphere. The scans

were obtained from 50 to 190 ∘C at a heating rate of 20 ∘Cmin−1. The heat

enthalpies (used to calculate the extent of crystallinity) were measured with an

error of ±0.1% and were confirmed by repeating the runs.

Rheological properties such as storagemodulus (G′), loss modulus (G′′), viscos-

ity, and elasticity of the nanocompositesweremeasured usingAR2000Rheometer

from TA Instruments. The measuring temperature and gap opening were 185 ∘Cand 1.6mm, respectively. Disc-shaped samples of diameter 25mm and thickness

2mm were used. Strain sweeps were recorded at 𝜔= 1 rad s−1 from 0.1 to 100%


strain.The shear stability of the samples was observed up to 10% strain. Hence, as

a safe approach, frequency sweeps (dynamic testing) were recorded at 4% strain

from 𝜔= 0.1 to 100 rad s−1 [28].

Tensile testing of composites was performed on universal testing machine

(Testometric, UK). The sample dimensions for tensile test were: sample length

73mm, gage length 30mm, width 4mm, and thickness 2mm. A loading rate

of 4mmmin−1 was used and the tests were carried out at room temperature.

Tensile modulus and yield stress were calculated using built-in softwareWin Test

Analysis. An average of three values was reported.

Transmission electron microscopy (TEM) of graphene oxide, masterbatch,

and the nanocomposite samples was performed using EM 912 Omega (Zeiss,

Oberkochen BRD) and Philips CM 20 (Philips/FEI, Eindhoven) electron

microscopes at 120 and 200 kV accelerating voltage, respectively. Sections of

70–90 nm thickness were microtomed from the block of the specimen and were

subsequently supported on 100mesh grids sputter coated with a 3-nm thick

carbon layer.

2.3

Results and Discussion

In the current study, graphene oxide–polyethylene nanocomposites were gener-

ated using solution mixing and melt mixing processes. Two CPE compatibilizers

differing in chlorination extent were used in order to study their effect on filler

dispersion as well as resulting polymer properties. Solution mixing was used

to generate masterbatches of CPE with graphene oxide, which can help to

better disperse the filler in the matrix polymer when the masterbatch is melt

compounded with it. The EDX analysis of the graphene oxide surface revealed

C/O ratio of 20, indicating the presence of polar surface groups (hydroxyl,

epoxide, carboxyl, etc.) which can interact with the polar compatibilizers used in

the study.

Table 2.3 and Figures 2.2–2.4 describe the calorimetric analysis of the pure poly-

mer, compatibilizers as well as polymer nanocomposites. The melt enthalpy of

pure crystallineHDPEwas taken as 293 J g−1 andwas used to determine the extent

of crystallinity in the polymer [28].

CPE25 compatibilizer was semicrystalline in nature as indicated by the crys-

talline melting peak in the DSC thermogram in Figure 2.2. Peak melting tempera-

ture of 130 ∘Cwas observed. On the other hand, CPE35 was amorphous in nature

as no melting transition was observed. Thus, the compatibilizers were different

not only in the extent of chlorination, but also in morphology.

The peak melting temperatures in the nanocomposites were always higher

than they were in the pure polymer, indicating the impact of graphene oxide on

polymer morphology. The impact was observed even on adding the amorphous

compatibilizer to the system, though it was less in magnitude as compared to

the system compatibilized with semicrystalline compatibilizer. In case of CPE35


Table 2.3 Calorimetric analysis of the pure polymers and polymer nanocomposites.

Code Polymer/nanocomposite 𝚫H (J g−1) Peak melting temperature (∘C) Crystallinity (%)

1 HDPE 147 142 50

2 HDPE/G 142 144 49

3 HDPE/1%CPE25/G 153 143 52

4 HDPE/2%CPE25/G 149 143 51

5 HDPE/5%CPE25/G 151 147 52

6 HDPE/10%CPE25/G 150 145 51

7 HDPE/5%CPE35/G 147 143 48

8 HDPE/10%CPE35/G 150 144 46

9 CPE25 47 130 —

10 CPE35 — — —

60

0

1

2

3

4

90 120 150

Temperature (°C)

HDPE

HDPE/G

CPE25

CPE35

Heat flow

(m

W m

g−1

)

180

Figure 2.2 DSC thermograms of HDPE, CPEs, and HDPE/G composite.

composites, an increase of 1–2 ∘C in peak melting temperature as compared

to pure HDPE was observed, whereas, this increase for the same amount of

CPE25 compatibilizer was 3–5 ∘C. The impact of graphene oxide (especially at

higher compatibilizer content) was also confirmed by DSC analysis of blends of

HDPE with CPE. In HDPE/CPE35 blends, decrease in peak melting temperatures

was observed, whereas only a marginal increase was observed for the CPE25

compatibilized HDPE system. Also, the peak melting temperature increased on

increasing the compatibilizer content (Figure 2.3) indicating that the enhanced


60 80 100

HDPE/5% CPE25/G

HDPE/10% CPE25/G

HDPE/5% CPE35/G

HDPE/10% CPE35/G

120 140 160 180

Temperature (°C)

He

at

flo

w,

(mW

mg

−1)

Figure 2.3 DSC thermograms of HDPE/CPE/G composites with 5 and 10wt% compatibilizer

content.

filler dispersion would have taken place, subsequently enhancing the thermal

resistance of the crystals [29]. The degree of crystallinity of the polymer was

also observed to be affected by graphene oxide as well as compatibilizer. The

composite without any compatibilizer had an extent of crystallinity of 49%, which

was marginally lower than the pure polymer crystallinity of 50%. It indicated

that the graphene oxide platelets slightly hindered the chain mobility and hence

their packing into the crystal structure [14]. Addition of amorphous CPE35 also

resulted in a further decrease in the degree of crystallinity of polymer, the mag-

nitude of which increased on increasing the content of compatibilizer. Thus, the

increased number of amorphous chains in the matrix resulted in the hindering

of the crystalline packing of HDPE chains. In the case of CPE25 compatibilizer,

the extent of crystallinity was always higher than the pure polymer irrespective of

the compatibilizer content. However, changes in the melt transition curves were

observed on increasing the compatibilizer content beyond 5wt% (Figure 2.4),

indicating changes in the crystallization behavior. There is a possibility that CPE

crystallized separately from HDPE owing to either its incompatibility with HDPE

or its interaction with graphene oxide surface, which led to its separation from

the matrix polymer.

Network structure of the polymer nanocomposites was evaluated with shear

rheology and the storage, loss, and complex moduli of the samples as a function

of angular frequency are demonstrated in Figures 2.5–2.7. Strain sweep was

conducted and samples were found to be safe up to 10% strain. Frequency sweep

of the samples was performed with controlled shear strain at 4% using frequency


60 90 120 150 180

Temperature (°C)

HDPE/1% CPE25/G

HDPE/2% CPE25/G

HDPE/5% CPE25/G

HDPE/10% CPE25/G

Heat flow

, (m

W m

g−1

)

Figure 2.4 DSC thermograms of HDPE/CPE25/G composites with 1, 2, 5, and 10wt% com-

patibilizer content.

0.11000

10 000

100 000

G′,

Pa

1000

10 000

100 000

G′,

Pa

1 10 100

HDPEHDPE/G

HDPE/5% CPE25/G

HDPE/1% CPE25/GHDPE/2% CPE25/GHDPE/5% CPE25/GHDPE/10% CPE25/G

HDPE/10% CPE25/G

HDPE/5% CPE35/GHDPE/10% CPE35/G

Angular frequency, rad s−1


0.1 1 10 100

(a) (b)

Figure 2.5 (a,b) Storage modulus of HDPE and HDPE nanocomposites as a function of

angular frequency.

range of 0.1–100 rad s−1. As is evident in Figure 2.5a, the storage modulus of pure

HDPE was the lowest among all the samples at all frequencies. On addition of

0.5wt% graphene oxide without compatibilizer, an order of magnitude increase

in the storage modulus occurred. For example, at a frequency of 10 rad s−1, the

storage modulus for HDPE was 15 730 Pa, which was enhanced to 112 000 Pa

with the addition of only 0.5wt% graphene oxide.The rate of increase in modulus

decreased on increasing the angular frequency; however, both the samples


0.11000

10 000

100 000G

″, P

a

10 000

100 000

G″,

Pa

1 10 100

HDPEHDPE/GHDPE/5% CPE25/G


HDPE/10% CPE25/GHDPE/5% CPE35/G

HDPE/10% CPE35/G


0.1 1 10 100

Angular frequency, rad s−1(a) (b)

Figure 2.6 (a,b) Loss modulus of HDPE and HDPE nanocomposites as a function of angular

frequency.

1000

10 000

100 000

G*,

Pa

10 000

100 000

G*,

Pa

0.1 1 10 100


0.1 1 10 100


HDPEHDPE/GHDPE/5% CPE25/GHDPE/10% CPE25/GHDPE/5% CPE35/GHDPE/10% CPE35/G


(a) (b)

Figure 2.7 (a,b) Complex modulus of HDPE and HDPE nanocomposites.

became independent of frequency and showed sudden shear thinning at the fre-

quency of∼20 rad s−1. In the case of graphene oxide nanocomposite, such a behav-

ior could be because of the rupture of interface between the polymer and graphene

oxide surface at higher frequencies. Literature studies have suggested this phe-

nomenon to be a result of the alignment of filler platelets in the direction of flow

at high shear or the slipping between the polymer and filler during high shear flow

[30]. The compatibilized nanocomposites exhibited good low-frequency depen-

dence followed by gradual decline in the modulus enhancement because of shear

thinning effect. CPE25 composites had higher storage moduli than the CPE35

containing nanocomposites. On increasing the amount of compatibilizer from

5 to 10wt%, the storage modulus was observed to decrease owing to the exten-

sive plasticization of the matrix (Figure 2.5a). Figure 2.5b also shows the effect of

1–10%CPE25 on the storagemodulus of the nanocomposites. Composites with 1

and 2wt% compatibilizer content were comparable in behavior and had modulus


value of 104 000 Pa at 10 rad s−1 frequency, which was similar to 112 000 Pa for the

nanocomposite without compatibilizer. Also, the storage modulus curves over-

lapped with each other for composites with 5 and 10wt% compatibilizer at higher

frequency value of 100 rad s−1.

The loss modulus results of the pure polymer and nanocomposites are pre-

sented in Figure 2.6. On comparison with the storage moduli of the samples,

it was observed that in pure HDPE, G′′ was always higher than G′ at any

frequency. It indicated that the polymer chains had dominant viscous behavior

with long relaxation times. In the case of nanocomposites, G′′ >G′ at lower

frequency indicated the dominance of the viscous part. Later, a transition was

observed, after which G′ >G′′ for the whole range of frequency indicated a

strong elastic character of the material with shorter relaxation times. In the

case of polymer nanocomposite without compatibilizer, such transition was

observed below 3 rad s−1 frequency. In the case of compatibilized systems, the

transition frequency increased on increasing the compatibilizer content. In

CPE25 nanocomposites, the frequency increased from 2.5 to 6 rad s−1, when

the compatibilizer content was raised from 1 to 10wt%. Similarly, in the case of

CPE35 containing nanocomposites, the transition frequency increased from 3 to

8 rad s−1 on increasing the compatibilizer content from 5 to 10wt%. It indicated

that the material behavior became more strongly viscous on increasing the com-

patibilizer content with CPE35 having stronger effect than CPE25 compatibilizer.

The loss modulus in the case of CPE25-containing composites was higher than

the CPE35-containing composites. Also, the magnitude of the loss modulus

decreased on increasing the content of the compatibilizer in the composite

(Figure 2.6a). Figure 2.6b also shows the impact of CPE25 compatibilizer on

the loss modulus when its content was enhanced from 1 to 10wt%. The loss

modulus at 10 rad s−1 was observed to be 88 590 Pa for 1% compatibilizer content,

which was reduced to 59 010 Pa at a compatibilizer amount of 10wt% in the

composite. The resulting overall complex moduli of the samples as a function of

angular frequency are also shown in Figure 2.7. It is also worth noting that the

improvements in the rheological properties with the addition of a small amount

of graphene are very significant when compared to other filler systems. For

example, only a slight increase in the storage and loss modulus of polypropylene

nanotube nanocomposites was observed at 1wt% nanotube content as compared

to the order of magnitude increase at 0.5wt% graphene oxide reported in the

current study [31].

Figures 2.8–2.10 demonstrate viscosity, elasticity, and complex viscosity of

the samples as a function of angular velocity. Contrary to the modulus, all of

these quantities were observed to decrease on increasing angular frequency.

Following increase in the frequency, polymer structure shows temporary

network of entanglements, which leads to more flexibility (lower viscosity).

In composites, this allowed more deformation energy to be stored resulting

in elastic dominance, which leads to increased modulus. Thus, it simultane-

ously reduced the contribution of lost deformation energy and hence viscous

behavior [17]. Lowest values were observed in the case of HDPE, whereas


0.1 1 10 100


0.1 1 10 100


1000

10 000

100 000η′

, P

a.s

η′,

Pa.s

1000

10 000

100 000

HDPE

HDPE/G

HDPE/5% CPE25/G

HDPE/10% CPE25/G

HDPE/5% CPE35/G

HDPE/10% CPE35/G

HDPE/1% CPE25/G

HDPE/2% CPE25/G

HDPE/5% CPE25/G

HDPE/10% CPE25/G

(a) (b)

Figure 2.8 (a,b) Viscosity curves of HDPE and HDPE nanocomposites as a function of angu-

lar frequency.

0.1 1 10 100


0.1 1 10 100


1000

10 000

η″,

Pa

.s

1000

10 000

η″,

Pa

.s


HDPE/1% CPE25/G

HDPE/2% CPE25/G

HDPE/5% CPE25/G

HDPE/10% CPE25/G

(a) (b)

Figure 2.9 (a,b) Elasticity of HDPE and HDPE nanocomposites as a function of angular fre-

quency.

0.1 1 10 100


0.1 1 10 100


1000

10 000

100 000

η*,

Pa

.s

1000

10 000

100 000

η*,

Pa

.s



(a) (b)

Figure 2.10 (a,b) Complex viscosity of HDPE and HDPE nanocomposites.


composite with graphene oxide without compatibilizer exhibited highest values.

For example, at a frequency of 10 rad s−1, the viscosity of the pure HDPE was

2204 Pa⋅s, which was enhanced to 9137 Pa⋅s in the case of HDPE-graphene oxide

nanocomposite (Figure 2.8a). Similarly, the elasticity increased from 1573 Pa⋅sfor pure polymer to 11 230 Pa⋅s for HDPE–graphene oxide nanocomposite

(Figure 2.9a) indicating the strong impact on both viscosity and elasticity by

adding only a small amount of graphene oxide. The slope of decrease in these

quantities as a function of angular frequency also decreased after a frequency

of ∼10 rad s−1. CPE25 composites had higher viscosity and elasticity than the

CPE35 containing nanocomposites. On increasing the amount of compatibilizer

from 5 to 10wt%, the viscosity and elasticity were observed to decrease further

(Figures 2.8a and 2.9a). Figures 2.8b and 2.9b also show the effect of increasing

the amount of CPE25 content from 1 to 10% on viscosity and elasticity of the

nanocomposites. Composites with 1 and 2wt% compatibilizer content were

comparable in behavior and had viscosity and elasticity values of 8859 and

10 360 Pa⋅s at 10 rad s−1 frequency, which was similar to the nanocomposite

without compatibilizer. On the other hand, these values decreased to 5901 and

6402 Pa⋅s in the case of 10wt% compatibilizer system.

Similar to the shear moduli, 𝜂′ >𝜂′′ was true for HDPE at all angular frequency

values indicating that the viscous contribution dominated the effect of elastic-

ity in the pure polymer (Figure 2.9a). In the case of HDPE/G composite, 𝜂′ >𝜂′′

was true only at lower angular frequency values. At ∼3 rad s−1, the transition in

behaviorwas observed indicating the dominance of elasticity component at higher

shear frequencies. Similarly, for nanocomposites with compatibilizer, the transi-

tion between 𝜂′ and 𝜂′′ occurred between 3 and 6 rad s−1 frequency. The com-

bined effect of viscosity and elasticity of the pure polymer and nanocomposites is

demonstrated in the form of complex viscosity in Figure 2.10.

Morphology of the masterbatches as well as nanocomposites was also analyzed

through microscopy as shown in Figure 2.11 (for 5% compatibilizer content).

The CPE35 masterbatch (Figure 2.11a) was observed to have better graphene

oxide dispersion as compared to CPE25 containing masterbatch (Figure 2.11b).

Although complete nanoscale delamination of the graphene oxide platelets was

not observed in the composites, the composites with CPE35 compatibilizer

had much better filler dispersion as compared to the corresponding CPE25

nanocomposites. Graphene oxide stacks of varying thicknesses (single layers to

multiple layers) can be observed for CPE35 containing composites in Figure 2.11c

and d, whereas, the stack thickness wasmuch higher for CPE25 nanocomposite as

shown in Figure 2.11e. As the compatibilizers differ in the extent of chlorination,

the resulting morphology can be related to the interaction of polar chlorine atoms

with the graphene oxide surface. Higher extent of chlorination in the matrix

resulted in higher magnitude of interfacial interactions between the polymer

and the filler surface, resulting in increased extent of filler delamination. Thus,

although increased chlorination content decreased the polymer crystallinity,

it increased the susceptibility of filler platelets to delaminate in the polymer

matrix.


250 nm

250 nm

100 nm100 nm

100 nm

(a) (b)

(c) (d)

(e)

Figure 2.11 TEM micrographs of (a) HDPE/CPE35 masterbatch, (b) HDPE/CPE25 master-

batch, (c, d) HDPE/5% CPE35/G and (e) HDPE/5%CPE25/G nanocomposites. The black lines

are the intersection of graphene oxide platelets.


Table 2.4 Tensile properties of the pure polymer and polymer-graphene oxide

nanocomposites.

Sr. No. Polymer/nanocomposite Young’s modulusa)

(MPa)

Peak stressb)

(MPa)

Peak strainc)

(%)

1 HDPE 1063 62 8.4

2 HDPE/G 1150 58 7.8

3 HDPE/1%CPE25/G 974 56 8.5

4 HDPE/2%CPE25/G 1181 69 7.8

5 HDPE/5%CPE25/G 1228 75 7.6

6 HDPE/10%CPE25/G 1160 65 7.8

7 HDPE/5%CPE35/G 905 48 8.7

8 HDPE/10%CPE35/G 1105 68 8.0

a) Relative probable error 2%.

b) Relative probable error 2%.

c) Relative probable error 5%.

Table 2.4 reports the tensile properties of HDPE and its graphene oxide

composites. The tensile modulus for pure polymer was observed to be 1063MPa,

which enhanced to 1150MPa in graphene oxide – HDPE nanocomposite without

compatibilizer. The addition of semicrystalline compatibilizer CPE25 increased

the modulus gradually till 5 wt%, after which a reduction in the modulus was

recorded. Similar observations have also been made earlier [32, 33], where the

modulus was described to be affected by a balance between an increase in modu-

lus owing to enhanced filler dispersion by the compatibilizer and a simultaneous

decrease in modulus following the plasticization of the matrix caused by it.

Till 5 wt% content of compatibilizer in the composite, the delamination effect

dominated, resulting in the increment of 16% in the modulus as compared to

pure HDPE. On further increasing the compatibilizer content, the plasticization

effect dominated the performance, resulting in the decrease in tensile modulus.

The increased extent of the interfacial interactions (hence filler delamination)

also resulted in the different behavior of CPE25 and CPE35 compatibilizers, as

CPE35-containing composites exhibited an increase in modulus even at 10%

compatibilizer content. However, the modulus of CPE35-containing composites

was lower than the corresponding CPE25 composites owing to amorphous nature

of CPE35. It should also be noted that though the increments in the modulus

are not tremendous, these enhancements are significant owing to a very low

amount of graphene oxide used to achieve them. The importance of masterbatch

approach was also confirmed by comparing the tensile modulus value of the

HDPE/G/5% CPE25 with the similar composite generated only by melt mixing.

A value of 969MPa was obtained for such melt mixed composite, which was

much lower than 1228MPa for the composite generated with masterbatch

approach. The peak stress also showed similar behavior as tensile modulus.

On the one hand, addition of graphene oxide to HDPE without compatibilizer

2.4 Conclusions 47

resulted in a slight decrease of peak stress because of restriction in segmental

mobility via mechanical interlocking with graphene oxide tactoids. On the other

hand, addition of CPE25 gradually enhanced the strength till 5 wt% content

resulting in an increment of 21% as compared to pure polymer. The CPE35-

containing composites showed an increase of 10% in strength at a compatibilizer

content of 10wt%. The composites still remained rigid as the peak strain was not

significantly affected with the addition on compatibilizers. A comparison of the

mechanical properties of the generated HDPE–graphene oxide nanocomposites

with other systems like HDPE–clay nanocomposites also revealed their high

potential. For example, HDPE–clay nanocomposites without compatibilizer

showed an increase of 4% in Young’s modulus at 2wt% filler content [34], which

was much lower than the graphene oxide nanocomposites even with 0.5wt%

filler content.

2.4

Conclusions

Nanocomposites of HDPE, graphene oxide, and two different CPE compatibiliz-

ers were generated using masterbatch (by solution mixing of CPE and graphene

oxide) and melt mixing methods. Addition of even 0.5wt% of graphene oxide

and different amounts of compatibilizers significantly impacted the morphol-

ogy and properties of polymer. The addition of graphene oxide caused slight

reduction in the polymer crystallinity because of reduction in chain mobility

and packing. The compatibilizer with a lower extent of chlorination exhibited

semicrystalline nature and did not decrease the overall crystallinity of poly-

mer in the composites, whereas the compatibilizer with higher chlorination

content was amorphous and led to a decrease in the polymer crystallinity. The

rheological characterization concluded that the addition of CPE improved the

processing of HDPE nanocomposites, rather than pure HDPE and HDPE/G

nanocomposites, which showed sudden shear thinning at low frequency. The

CPE25-containing composites were superior in performance than the corre-

sponding CPE35 composites. The performance also reduced on increasing

the amount of compatibilizer. The compatibilizer with higher chlorination

content also resulted in better interfacial interactions with graphene oxide

leading to higher extent of filler delamination. Interplay of increased mechanical

performance owing to filler delamination and decreased properties following

matrix plasticization affected the tensile response of the nanocomposites.

CPE25-containing composites had the maximum improvement of 16 and 21% in

modulus and strength at 5wt% compatibilizer content. The CPE35-containing

composites exhibited an increase in the properties even at 10wt% compatibi-

lizer content because of the higher magnitude of interfacial interactions. The

mechanical properties in CPE35 containing composites, however, were lower

than the corresponding CPE25 composites because of the amorphous nature of

CPE35 compatibilizer.


Acknowledgments

Thedefinite version of this chapter has been published in Polymer Engineering and

Science, 2013, 53(1), 78–88, Copyright SPE Wiley.

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51

3

Preparation and Applications of Hydroxyapatite

Nanocomposites Based on Biodegradable and

Natural Polymers

Pau Turon, Luis J. del Valle, Carlos Alemán, and Jordi Puiggalí

3.1

Introduction

Hydroxyapatite (HAp) is a bioceramic material that makes up the majority of the

inorganic components of human bones, tendons, and teeth. Biologically formed

HApgenerally appears as nanocrystals that are logically produced undermild con-

ditions of pressure and temperature. By contrast, HAp can also be found in dead

nature as mineral deposits that have usually been grown over many years under

extreme conditions of pressure and temperature.

The almost ubiquitous presence of HAp in the body in either dissolved or

solid form supports biocompatibility, bioactivity, and osteoconductivity prop-

erties [1] of synthetic HAp. Biodegradable and biobased polymers filled with

synthetic HAp have great interest as drug delivery systems and gene therapy.

Furthermore, synthetic HAp nanocomposites are nowadays among the most

important implantable materials used in biomedical applications and particularly

for hard tissue regeneration. Although these materials can closely mimetize the

structure of human hard tissues, improvement of their properties is still desirable

[2]. Hence, efforts are now focused on varying composition and processing

conditions since chemical, structural, and morphological properties become

highly influenced [3, 4]. Specifically, several processes such as precipitation

[5, 6], hydrothermal [7, 8], sonochemical [9], sol–gel [10, 11], and emulsion and

microemulsion [12, 13] have been investigated and reported for the synthesis of

HAp with controlled morphology. For example, one-dimensional (1-D) nano-

materials (nanowires, nanorods, nanobelts, and nanotubes) with well-controlled

dimensions, composition, and crystallinity have attracted special interest due to

their excellent properties and superior applications respect to bulk counterparts

[14]. HAp nanorods raised great interests for biomaterials scientists because of

their potential applications in bionanocomposites [15, 16].

Excellent reviews can be found in the literature about synthesis and applica-

tions of HAp-containing materials. The interest of these composites is growing

exponentially and, for example, during the last 2 years we can found rele-

vant works concerning generic applications of biomimetic synthetic calcium


52 3 Preparation and Applications of Hydroxyapatite Nanocomposites

phosphates (CaP) [17], preparation of hybrid nanocomposite scaffolds that mimic

the complex nanostructured architecture of bones [18], specific characteristics

of bioactive materials when used in bulk or as coatings [19], strategies focused

into the design and choice of the appropriate scaffold for tissue engineering

applications [20, 21], preparation of multifunctional nanoscaffolds for tissue

engineering applications using electrospinning [22], design of novel bioactive

materials based on chemical reactivity in body fluid [23], and development of

homogeneously mineralized self-assembled peptide-amphiphile nanofiber–HAp

composites [24].

3.2

Preparation of HAp Nanocrystals

CaPs exist in different forms exhibiting distinct crystal structures and Ca/P

ratios. The most known are amorphous calcium phosphate (ACP), tricalcium

phosphates (TCPs), tetracalcium phosphates (TTCPs), brushite (B), monetite

(M), octa-calcium phosphate (OCP), and HAp [25, 26].

HAp is chemically represented as Ca10(PO4)6(OH)2 and corresponds to the

most stable and least soluble of all calcium orthophosphates. Pure HAp crystal-

lizes in the monoclinic P21/b space group (a= 0.984 nm, b= 2a, c= 0.688 nm,

and 𝛾 = 120∘) that at temperatures above 250 ∘C experiments a transition toward

a hexagonal phase (a= b= 0.943 nm, c= 0.689 nm, and 𝛾 = 120∘) defined by

a P63/m space group [26, 27]. Some impurities, like partial substitution of

hydroxide (e.g., by fluoride or chloride ions) stabilize the hexagonal structure

of HAp at ambient temperature. For this reason, the very rare single crystals of

natural HAp always exhibit an hexagonal space group.

HAp can be prepared by mixing aqueous solutions containing the corre-

sponding stoichiometric amounts of calcium and phosphate ions. However, it

is difficult to get the exact stoichiometry (i.e., a Ca/P ratio equal to 1.67) since

other Ca/P ratios can be stabilized in the formed crystals depending on the

synthetic method and conditions employed. Various HAp phases can thus be

formed, which can be categorized into calcium-deficient HAp, oxy-HAp, and

carbonate-substituted HAp.

The following two chemical reactions describe the most widely used aqueous

chemical precipitation routes:

10 Ca(OH)2 + 6 H3PO4 → Ca10(PO4)6(OH)2 + 18 H2O (3.1)

10 Ca(NO3)2 + 6 (NH4)2HPO4 + 2 H2O → Ca10(PO4)6(OH)2+ 12 NH4NO3 + 8 HNO3 (3.2)

The crystal growth process becomes strongly influenced by the way of mixing

(e.g., quick or slow addition of one reactant over the other), pH conditions, pres-

ence of surfactants and chelating agents. Synthesis can also be performed using

ethanol instead of water [28].


Supersaturated solution Prenucleation clusters Postcritical ACP clusters CAp nucleation

Aggregation

Continue growth

(001)

(1–

10

)

(10

0)

(01

0)

Final CAp CrystalsCrystalline core

amorphous shellCa2+

PO43−

CO32−

Na+

OH−

Cry

stal

gro

wth

Figure 3.1 Schematic illustration of the nucleation and growth mechanisms of carbonated

apatite (CAp) nanocrystals. Reproduced with permission from Ref. [32] © 2013, Royal Society

of Chemistry.

In general, the size ofHAp particles becomes larger and the dispensability worse

when solvent power decreases (e.g., using ethanol instead of water) [29].The effect

of temperature on the shape of crystals is rather complicated and may depend on

the solution chemistry. Thus, the morphology changes from needles at 40 ∘C to

spheroidal at 100 ∘C when the process is performed according to Reaction 3.1,

while the trend is reversed when HAp is prepared by Reaction 3.2. Changes in the

morphology with temperature have been analyzed taking into account the driving

force for the HAp precipitation and the supersaturation level of Ca2+ and PO43−

ions with respect to HAp [30].

ACP is often the first precipitated phase in the synthesis of HAp by rapidmixing

of aqueous solutions [31]. This transient amorphous phase slowly converts to the

more stable HAp crystalline phase by thermal treatment [32] (Figure 3.1).

Therefore, chemical precipitation followed by hydrothermal processing is con-

ventionally employed in the laboratory-scale synthesis of HAp.This hydrothermal

process stands out among the different techniques developed up to now since it

has several advantages that include environmental concerns, easily controllable

reaction conditions, relatively large scale and high yield, and use of water as

the reaction medium. The process allows achieving several types of dominant

morphologies (e.g., nanorods, nanowires, microsheets, bur-like microspheres,

and microflowers) [33], depending on the pH value of the reaction solution

(Figure 3.2).

Adsorption of OH− ions onto the crystal surface is necessary for the crystal-

lization and growth of HAp. A high adsorption is expected under basic condi-

tions which should lead to isotropic or weak-anisotropic growth (i.e., formation of

short nanorods or nanoparticles), whereas the limited adsorption attained at low


500 nm

(a) (b)

(c) (d)

(e) (f)

(g) (h)

200 nm

100 μm 5 μm

20 μm

10 μm 1 μm

5 μm

2 μm

Figure 3.2 SEM images showing different

morphologies of HAp samples obtained at

different pH values: (a,b) nanorods, pH 7.0;

(c,d) bur-like microspheres, pH 5.0; (e,f )

microflowers, pH 4.5; and (g,h) microsheets,

pH 4.0. Reproduced with permission from

Ref. [33] © 2009, American Chemical Society.


Figure 3.3 Representation of the formation mechanism of hierarchically nanostructured

HAp consisting of nanosheets assembled from nanorods. Reproduced with permission from

Ref. [36] © 2012, PMC.

pH resulted in anisotropic growth (e.g., two-dimensional microsheets and three-

dimensional microflowers).

The addition of compounds able to complex calcium ions may play also a cru-

cial role in the crystal morphology. In this way, trisodium citrate was revealed

highly effective since its chelating effect was weakened under the high tempera-

ture and pressure conditions of the hydrothermal treatment. Calcium ions could

therefore be gradually released to the reaction medium slowing down nucleation

and subsequent crystal growth. Furthermore, selective adsorption of citrate on

the surface of HAp-growing particles could change the relative surface energy of

different crystal facets and therefore could influence on the growth rates along cer-

tain orientations [33–35]. Thus, a preferential adsorption of citrate on the crystal

facets perpendicular to the anisotropic growth direction could lead to formation

of nanorods instead of nanowires.

The use of potassium sodium tartrate as chelating ligand and templatemolecule

has also been explored [36]. This compound may play a key role in the forma-

tion of hierarchically nanostructured HAp. A plausible mechanism based on soft-

template and self-assembly was proposed for the formation and growth of such

nanostructures (Figure 3.3).

Morphology of HAp grown in a solution system based on simulated body fluid

(SBF) has been studied in detail [37]. Nanometric low-dimensional forms, such

as sheets and needles elongated in the c crystallographic axis, were produced

with phosphate-surplus (or calcium deficient) HAp in the solution at human

body temperature. The change of pH from 6.5 to 7.0 leads to an increase on the

growth rate and a change from nanoneedle to nanosheet morphology (Figure 3.4).

On the other hand, micrometric bulky hexagonal shapes and faceted plates of

semi-stoichiometric HAp were grown under hydrothermal conditions at pH

7.0 and 7.4, respectively. The variation of the morphology was explained on

the basis of the change of the growth mode of HAp crystals depending on the


(100)

Low High

At high temperature

Stoichiometric

Micrometric Nanometric

Phosphate-rich

At low temperature

Simulatedbody fluid

Hexagonal rod plate

(100)

HighLowGrowth rate

Growth rate

SheetNeedle

(110)

Figure 3.4 Schematic illustration showing the morphological variation of HAp grown in

SBF-based solution. A slight increase on the pH of the medium strongly influences the

growth rate that becomes enhanced. Based on Ref. [37].

supersaturated conditions. It was also assumed that the adsorption of phosphate

species to the specific faces miniaturizes the HAp crystals and changes their

morphology to low-dimensional forms.

A better control on the crystal morphology attained through the hydrothermal

process can be achieved using surfactant molecules. In fact, biological processes

concerning HAp provide some clues about how uniform and well-defined

crystalline structures can be obtained through macromolecular control and

cell organization. Surface layers of the growing crystals can incorporate soluble

additives provided that there is a degree of complementarity in charge and size

between the guest ions and the interstices in the structure of the crystal boundary

layers [38]. Several macromolecules, such as stearic acid, monosaccharides,

and related molecules have been explored to control crystal morphology and

specifically cetyltrimethylammonium bromide (CTAB) has been revealed highly

effective [39]. Interactions between phosphate anionic groups and ionized CTAB

molecules (Figure 3.5) gave rise to rod-like morphologies since crystal growth

along the c crystallographic axis becomes favored. However, abundant hydroxyl

groups exist in solution when the pH of the medium increases (e.g., higher than

9). These hydroxyls are able to compete for establishing interactions with CTAB

cations and rounded morphologies become in this case favored.

The use of mixtures between cationic (e.g., CTAB) and anionic (e.g., sodium

dodecyl sulfate, SDS) surfactants appears as an additional tool to control the mor-

phology of nano-HAp [40] (e.g., from rod-like to sheet-like crystals for anionic-

rich and cationic-rich surfactant media, respectively). It is clear that molecular

assembly of themixed surfactants can form a variety of structures (e.g., cylindrical

micelles, vesicles, and planar bilayers) depending on their mixing ratio. Columbic

interactions between the hydrophilic parts of anionic and cationic surfactants


HO−

HO− HO− HO−

3−

HO−HO− HO−

O−

O−

O−

O−

P+

HO−

CH3

CH3

CH3

N+

Figure 3.5 Scheme showing interactions between surfactant cation, phosphate anion, and

hydroxyl groups.

can lead to new composites, whose properties are greatly different from those of

single-component systems [41].

Other organic compounds such as D-sorbitol, polyethylene glycol, gelatin and

urea have also been studied to control the size and morphology of synthetic

HAp [42]. Thus, thinner and longer HAp nanorods can be produced by adding

D-sorbitol as a modifier when synthesis temperature was low due to the tem-

plating effect of the linear structured D-sorbitol. This effect was lost at higher

temperatures since in this case the interaction between the linear D-sorbitol

and the resultant HAp crystallites was weakened. By contrast, addition of PEG

(poly(ethylene glycol)) molecules has only a remarkable effect enhancing nanorod

morphology when synthesis temperature is high. In this case, PEG molecules

have a high flexibility in aqueous solution and should favorably induce the axis

orientation growth of HAp via an interaction between the ether bonds of PEG

and HAp nanocrystallites, resulting in the formation of nanorods with long

lengths. Zhan et al. obtained HAp rods with length of hundreds of micrometers

by solution precipitation in the presence of gelatin and urea [43].

Nevertheless, alternative methods that allow a facile and effective anisotropic

growth of 1-D HAp nanorods via a hydrothermal route without using any

template/surfactant reagent have also been proposed. For example, single crys-

talline HAp nanorods with several hundred nanometers in length and tens of

nanometers in width can be prepared in weak acid environment in the presence

of sodium bicarbonate [44].

The sol–gel method has recently been developed to get HAp with an improved

chemical homogeneity with respect to that obtained using other conventional

methods (e.g., hydrothermal synthesis) [45, 46]. The sol–gel product is charac-

terized by a nanosize of primary particles that allows to improve contact region

and stability of interfaces. The method is a mild process that does not require

high temperatures and pHs. Moreover, the process allows to control crystallite

size and also to get thin film coatings in a rather simple way [47, 48]. The method

offers a molecular-level mixing of calcium and phosphorus precursors, being

until now proposed different combinations (e.g., calcium diethoxide (Ca(OEt)2)

and triethyl phosphate [49]; (Ca(NO3)2⋅4H2O) and phenyldichlorophosphite

(C6H5PCl2) [50]; calcium nitrate and phosphonoacetic acid [48]; N-butyl acid


phosphate mixed with calcium nitrate tetrahydrate [51]; and triethyl phosphite

or other phosphorous alcoxides and calcium nitrate [52]).

Surfactant-based emulsion systems are also highly promising for the synthe-

sis of nanoparticle with controlled morphology. Addition of surfactant stabilizes

the emulsion by reducing the surface tension of the immiscible liquid and by for-

mation of nanosized liquid droplets where synthesis of nanopowder takes place.

Morphology and particle size of the final powder depends therefore on the droplet

size and shape. Type (e.g., ionic and nonionic) and concentration of the surfac-

tant present in the liquid medium strongly influence the emulsion process. Sev-

eral attempts have been made to control morphology of the HAp nanopowders

using ionic (e.g., sodium bis(2-ethylhexyl)sulfosuccinate and dioctyl sulfosucci-

nate sodium salt) [53] and nonionic (e.g., poly(oxyethylene)5 nonylphenol ether

and poly(oxyethylene)12 nonylphenol ether) [54] surfactant-based systems. Inter-

estingly, the ratio between organic and aqueous phases plays also a crucial role

in the final morphology (Figure 3.6) [55]. At high water content, needle-shaped

crystals are favored since HAp nucleation and growth take place within the con-

straints of cylinder-like surfactant micelles. Less hydrogen bonding interactions

can be established between surfactant molecules and the aqueous core when sys-

tems have low water content. Dynamic exchange with other micelles favors the

formation of micelles with larger size and roughly spherical shape that finally lead

to nanoparticles with low aspect ratio.

Balance between provision of thematerials necessary for crystal growth and the

frequency of crystals nucleation is also a key factor to control the agglomeration

of nanocrystals produced by water in oil emulsion systems [56]. Aggregates com-

posed of plate-like nanocrystals andmonodisperse tiny ones can be obtained with

high-frequency nucleation and slow reactant provision, whereas monodisperse

tiny nanocrystals are obtained with the opposite conditions.

HAp nanoparticles have also been grown in solution by a rapid, economical, and

environment-friendlymethod using ultrasonic irradiation.The use of ultrasounds

has proved to improve considerably the crystal growth rate and furthermore son-

ication may also act as an additional energy source for the nucleation process. In

this way, HAp crystals can be formed in much shorter time than using conven-

tional processes and smaller and more uniform nanoparticles with higher surface

area can be obtained [57].

3.3

Preparation of HAp Nanocomposites

The high surface area of nano-hydroxyapatite (nano-HAp) leads to a greater

proportion of its ions to be located on its surface rather than in the interior. As

a consequence biological performance, such as cell adhesion, osteointegration,

cell proliferation, and differentiation can be enhanced. All these benefits are, for

example, ideal to the growth of new tissues within a short period of time [58].

3.3 Preparation of HAp Nanocomposites 59

Organic phase

Aqueous phase SurfactantInorganic nuclei

surfactant complex

Crystal nucleationand growth

Micellefusion

Aqueous phase

Organicphase

SurfactantCrystal nucleationand growth

Inorganic nucleisurfactant complex

Ca2+ PO43−

Ca2+ PO43−

Ca2+

PO43−Ca

2+

PO43− +

(a)

(b)

Figure 3.6 Formation of HAp nanopar-

ticle in the core of the cylindrical (a) and

spherical (b) reverse micelles in system with

aqueous/organic phase rates of 1 : 5 and

1 : 15, respectively. In the last case micelles of

nonuniform size were produced by the ran-

dom fusion of smaller micelles. Reproduced

with permission from Ref. [55] © 2009,

Elsevier.

A nanocomposite material consists at least of two chemically identified phases,

which are separated by interface(s). Different factors have a strong influence on

final properties and therefore their control becomes essential. Some of these fac-

tors concern only to the filler (e.g., shape, size, size distribution, and inherent

properties), other to the polymer matrix (e.g., molecular weight and other physic-

ochemical properties) and other to the polymer/filler mixture (e.g., concentra-

tion and dispersion of filler particles in the polymer matrix and the state of the

filler/matrix interface). For biocomposites, characteristics like biocompatibility

and nontoxicity of both filler and matrix, and degradation rate of matrix should

also to be taken into account.


Figure 3.7 Typical nanocomposite arrangement where the inorganic HAp nanoparticles are

embedded into the organic polymer matrix.

HAp nanocomposites generally consists on the arrangement obtained by

embedding the particles into the polymer matrix (Figure 3.7).

Twomain techniques can bementioned about the preparation of nanocompos-

ites based on nano-HAp and organic polymeric matrices:

1) Thermo-mechanical methods that incorporate the bioceramic particles into

the polymer matrix using conventional plastics processing technologies (e.g.,

compounding to get a homogeneous distribution,milling, and injectionmold-

ing) [59–61]. Specifically, the melt extrusion method has been shown to be a

good way to prepare homogeneous ceramic/polymer blends.

2) Physicochemical methods that incorporate the particles by in situ precipita-

tion of mineral crystals in the polymer matrix (coprecipitation or biomimetic

process) or alternatively by dispersion of bioceramic particles in the polymer

solution with subsequent consolidation (solvent casting).Themain drawback

of the solvent casting method is the risk of potentially toxic organic solvent

residues. Solvent and concentration of the polymer solution play a funda-

mental role to get a uniform distribution of nanoparticles. In the same way,

processing variables such as gelation rate, stirring time, andmixingmode have

a great influence on the process.

Modification of HAp surface appears highly interesting to get better distribu-

tion of particles and also to improve the adhesion with the polymermatrix in such

a way that debonding could be delayed. Wetting is important in the bonding or

adherence of the filler surface and the polymer and depends on the hydrophilicity

or polarity of the filler and the available polar groups of the polymer. Great efforts

are focused to select appropriate surface modifiers with nontoxic and biocompat-

ible characteristics that not change the properties of the nanoparticles [62].

Hexanoic and dodecanoic acids have been used to get a hydrophobic sur-

face while strong hydrogen bonding interactions are established between the

carboxylic groups and the P-OH groups of HAp [63]. Surfactant molecules such


as oleic acid, stearic acid, and sodiumdodecyl sulfate have also been employed

[64–66]. HAp surface can also be modified by esterification reactions between

acidic phosphate sites and added alcohols like dodecyl alcohol [67]. A grafting

ring-opening polymerization of L-lactide onto the surface of HAp nanoparticles

has been revealed highly effective to provide nanofillers able to be well dispersed

in a polylactide (PLA) matrix and even to act as heterogeneous nucleation

agents [68]. HAp crystals were also effectively grafted with polyethylene glycol

methacrylate phosphate after functionalization of their surfaces with thiol groups

and subsequent polymerization with the vinilic monomer. Interestingly, the

derived grafted nanoparticles showed very high colloidal stability, without crystal

aggregation in comparison with non-grafted particles [69].

Among the existingmethodologies developed to prepareHApnanocomposites,

in situ formation of nano-HAp crystals in the presence of polymers by the coso-

lution is probably the most attractive route. This coprecipitation method avoids

the extensive particle agglomeration usually observed when methods based on

mechanical mixing between nanopowder and a selected polymer are employed.

In some cases, crystallization of the inorganic compound can be delayed due to

chelating effect of some organic groups (e.g., the carboxylic groups of polyacrylic

acid [70]) with calcium ions.

HAp nanocomposites can also be prepared by dispersing the inorganic

nanocrystals into the monomer (e.g., methacrylate anhydride [71, 72]). The pro-

cess of polymer shell formation on the preformed inorganic cores, is a so-called

ex-situ approach. Particles coated by the polymer shell become considerably

more stable against aggregation.

Nanocomposites can also be formed according to processes that are inspired

on biological mechanisms and proceeded through hierarchical self-assembly.The

crystal growth of an apatite phase onto a collagen matrix is a clear example. The

self-organization occurred from electrostatic interactions between calcium and

phosphate ions of growing HAp with functional groups (e.g., COO− or NH3+)

placed outside the collagen molecules [73, 74]. This close interaction causes

that HAp nanocrystals become aligned with their c-axis preferentially oriented

along the collagen fibers, which indicates a close interaction between HAp and

collagen phases.

Electrospinning is nowadays one of the most promising techniques for man-

ufacturing in vitro fibrous scaffolds for tissue engineering applications. This

method enables also the production of biodegradable polymer nanofibers loaded

with HAp nanoparticles. In general, the derived scaffolds are highly porous and

offer a biomimicking structure for adhesion, accommodation, proliferation, and

mineralization of osteoblast cells [6–77].

3.4

Applications of HAp/DNA Nanocomplexes as Gene Carriers

Nonviral gene therapy becomes nowadays a rapidly growing strategy for the

treatment of both acquired and inherited diseases. Nonviral vectors have clear


advantages due to their low or no immunogenicity, relatively simple preparation

procedures, low cost, and high flexibility to accommodate the size of the delivered

transgene [78].

Great efforts are focused, for example, in the development of gene delivery sys-

tems that can protect plasmid DNA (pDNA) and posses a potential targeting abil-

ity. The advantages of HAp particles lie in its general efficiency for a wide range

of cell lines, simplicity, low cost, biocompatibility, and biodegradability. Synthesis

of HAp/DNA complexes can be performed by coprecipitation [79], encapsulation

[80], multishell structures formation [81], and coating [82]. These complexes can

be incorporated into cells through endocytosis by forming intracellular vesicles,

which merge with lysosomes (Figure 3.8) [83, 84]. CaP nanoparticles can be dis-

solved even in low acidic buffer (pH 5.0) releasing pDNA [80, 85].Therefore, DNA

can be released in the endosomal compartment and eventually enter the nuclei of

cells to effect gene transfer and expression.

The uptake mechanism of CaP nanoparticles by cells is still under investigation

since the route of entry of nanoparticles and their final intracellular localization

is decisive for a potential application as gene delivery agent. For CaP nanopar-

ticles a macropinocytosis mechanism seems favored as deduced from studies

Clathrin-mediatedendocytosis

Caveolin-mediatedendocytosis

Macropinocytosis

Macropinosome

Earlyendosome

NucleusGolgicomplex

Late endosome

Lysosome

Caveosome

Figure 3.8 Schematic representation of the different internalization mechanisms

for nanoparticles: clathrin-mediated endocytosis, caveolin-mediated endocytosis, and

macropinocytosis. Reproduced with permission from Ref. [83] © 2013, Elsevier.


Figure 3.9 Simulated structure showing the double helix of B-DNA as a nucleating agent

for the development of a HAp crystal.

carried out using specific inhibitors for the different uptake processes. Amoderate

concentration of CaP nanoparticles inside cells is desired to avoid cell apoptosis

produced when a high intracellular calcium levels attained after dissolution of

nanoparticles [86].

Molecular dynamic simulations have recently demonstrated that duplexes of

DNA adopting a B-double helix can be encapsulated inside nanopores of HAp

(Figure 3.9)without undergoing significant distortions in the interstrandhydrogen

bonds and the intrastrand stacking. Electrostatic interactions between the phos-

phate groups of DNA and Ca2+ have been found essential for the formation of

stable ion complexes. These become the starting point of CaP clusters by incor-

porating phosphate groups from the solution [87].

Subsequent atomistic molecular dynamics simulations allowed concluding that

the backbone of the DNA double helix can act as a template for HAp growth [88].

Theoretical calculations were also corroborated by the preparation of nanocap-

sules (Figure 3.10) and crystalline nanorods of HAp containing DNA molecules

inside.These complexes appear highly relevant for biomedical applications requir-

ing the protection of DNA from aggressive environmental conditions.

Different relevant works have been reported in the last decade to explore the use

of HAp nanoparticles as highly promising gene carrier vectors. Zuo et al. synthe-

sized HAp/DNA nanohybrids from lamellar-structured HAp. Gel electrophoresis

analysis confirmed that the lamellar HAp could protect DNA from degradation of

DNase I.The so-protected DNA could be recovered readily under acid conditions

and the integrity of released DNA was confirmed by UV–vis spectra [89].

Zhu et al. adsorbed EGFP-N1 pDNA on HAp nanocrystals and subsequently

demonstrated that these complexes transfected in vitro the plasmid into cancer

SGC-7901 cells with efficiency about 80% [90].


Figure 3.10 TEM micrographs showing HAp-DNA nanocapsules with a clearly distinctive

contrast that suggests the incorporation of DNA in their inner part. Reproduced with per-

mission from Ref. [88] © 2013, Royal Society of Chemistry.

Immunoadjuvant properties were found for HAp nanoparticles when

administered with malarial merozoite surface protein-119. A slow in vitro antigen

release and a slow biodegradability behavior were characteristic, which may lead

to a prolonged exposure to antigen-presenting cells and lymphocytes [91]. The

prepared HAp nanoparticles have promising properties to be used as antigen

carriers for immunopotentiation [91].

Arginine-modified nano-HAp was able to form rapidly nanocomplexes with

DNA by electrostatic interaction. These nanoparticles could effectively bind and

protect DNA and be considered as a potential gene carrier [92].

DNAzymes are synthetic, single-stranded, catalytic nucleic acids that bind and

cleave target mRNA in a sequence-specific manner. These have been explored

for genotherapeutics although their application is seriously hindered due to the

lack of an efficient delivery system.This feature that can be well overcomed using

nano-HAp as a nonviral vector. It was observed that in a mouse tumor model, the

arginine–nano-HAp complex was efficiently delivered to tumor tissue, downreg-

ulating expression of latent membrane protein in nasopharyngeal carcinoma cells

and suppressing tumor growth [93].

HAp/DNA complexes formed in SBF showed a higher transfection efficiency

than those made in water probably as a consequence of a slower growth of

nanoparticles over time that lead to a smaller crystal size [94].

Influence of Ca/P stoichiometry on the stability of HAp/DNA complexes has

also been evaluated [95]. In vitro transfection studies revealed that improved and

more consistent levels of gene expression can be achieved by optimizing this sto-

ichiometry as well as the mode in which the precursor solutions are mixed. The

optimized forms of these complexes were approximately 25–50 nm in size and

were efficient at both binding and condensing the genetic material.

3.5 Tissue Engineering Applications of HAp Nanocomposites Based on Biodegradable Polymers 65

Favorable transfection characteristics of HAp/DNA complexes should be

exploited to design and construct novel 3D scaffolds based on embedding these

complexes in a porous biodegradable polymer matrix to achieve controlled and

efficient gene transfection in in-vivo experiments.The scaffold essentially acts as a

depot for the gene while simultaneously offers structural support and a matrix for

new tissue deposition. Thus, HAp nanoparticles have been combined with colla-

gen to yield bioactive, biodegradable scaffolds that showed ability to act as gene-

activated matrices for delivery of bone morphogenetic proteins (BMP2) [96]. The

ephrinB2 gene has recently been incorporated to produce these novel therapeutic

gene-activatedmatrices for bone repair [97]. Successful transfection ofmesenchy-

mal stem cells (MSCs)was demonstrated and resulted in high calciumproduction.

3.5

Tissue Engineering Applications of HAp Nanocomposites Based on Biodegradable

Polymers

Design and development of scaffolds able to replace the form and function of

native tissue and to promote regeneration without necrosis or scar formation

is a current research topic. Bionanocomposite materials can mimic the natural

morphology of the extracellular matrix that surrounds cells and consequently

may be ideal for regeneration of tissue structures. To this end, key characteristics

of the extracellular matrix should be considered: (i) an hybrid structure composed

of macromolecules (e.g., proteins and polysaccharides) and inorganic matter and

(ii) a macromolecular morphology characterized by a high aspect ratio and a

nanoscale diameter dimension.

In fact, an ideal engineered bone implant should be osteoconductive (i.e.,

it should promote the attachment, survival, migration, and distribution of

osteogenic cells), osteoinductive (i.e., it should have spatial, physical, and bio-

chemical stimuli to initiate stem or progenitor cells toward osteoblastic lineage),

and osteogenic (i.e., it should contain osteogenic stem or progenitor cells for bone

regeneration) [98]. A complex strategy is needed to meet these requirements,

but in a first step the preparation of appropriate biomaterial scaffolds becomes a

crucial point.

The architecture of biomaterial scaffolds should provide a structural support for

cell integration, but also regulate cell proliferation, differentiation, and migration

to form functional tissues (Figure 3.11) [99]. Design of 3D-scaffolds that incorpo-

rates nanoscale features appear fundamental to recreate the hierarchical organi-

zation of natural extra cellular matrix (ECM).

Several biodegradable polymeric materials have been investigated for tissue

engineering applications, although the strict requirements for biomedical

applications cannot be accomplished by a single polymer. Therefore, the design

of multi-component systems becomes a viable strategy and specifically the

introduction of inorganic nanofillers (e.g., nano-HAp) into biodegradable

polymers is one of the most attractive alternatives.


Macrostructure Nanostructure

Nanopores

Nanofibers

Nanocomposite

Nanospheres

Figure 3.11 Macrostructure scaffolds for hard tissue engineering and nanostructures with

features of nanopores, nanocomposites, nanospheres, and nanofibers. Reproduced with per-

mission from Ref. [99] © 2012, Wiley Periodicals, Inc.

Systems based on collagen protein fibers, proteo-glycans, andHApnanocrystals

are the biocomposite materials most studied since the indicated components are

arranged at the nanometer scale in natural bones [100].

Electrospinning is probably the most easy way to combine degradable polymers

having a nanofiber morphology with bioactive inorganic materials such as HAp

[101]. Moreover, the generated nanofibers may have appropriate properties tar-

geted for bone regeneration.

Small amounts of HAp nanoparticles can be incorporated into the electrospun

fibers in three different ways that depend on the relative size between particles

and fibers. Thus, surface attachment, partial encapsulation, and total encapsu-

lation (Figure 3.12) can be observed if the diameter of the fiber is significantly

smaller, similar, and larger, respectively, than that of the HAp nanoparticle. Partial

and total encapsulation of HAp nanoparticles are expected for fibers containing

a large amount of particles. Complete encapsulation of nanoparticles may be

good when mechanical properties are considered, while partial encapsulation or

nanoparticle attachment to the fiber surface should be more adequate to enhance

the bioactivity of the fiber [102].

Electrospinning of organic–inorganic compounds may have inherent prob-

lems related to the preparation of a homogeneous electrospinnable solution.

In addition, it has been reported that depending on the solvent electrospun

natural biopolymers should lead to a denatured form that loose the typical

biological properties derived from their structure. For example, the triple helix

characteristic of collagen molecules is lost after electrospinning giving rise to

gelatin [103]. Nevertheless, cross-linked electrospun collagen is believed to still

have good potential as a nanofibrous substrate for bone regeneration.

Electrospinning of hydroxyapatite nanopowders directly mixed with a gelatin

solution is difficult since usually lead to the formation of abundant beads.


Nanospun fiber

Nanospun fiber

Nanospun fiber

Nanospun fiber

CHA nanoparticle

CHA nanoparticle

CHA nanoparticle

CHA nanoparticle

(a)

(b)

(c)

(d)

Figure 3.12 Incorporation of small amount

of CHA particles in electrospun fibers:

attachment of nanoparticles to the fiber

surface (a), partial encapsulation (b), and

total encapsulation (c). (d) Partial and total

encapsulation of nanoparticles for fibers con-

taining a large amount of CHA. Reproduced

with permission from Ref. [102] © 2010, IOP

Publishing, Ltd.

(HAp precipitation)

Ca+Gelatin P+Gelatin

HAp+

Gelatin

HAp-

Gelatin

HFP EDC/NHS

(Freeze-drying) (Nanocomposite solution) (Electrospinning) (Cross-linking)

Figure 3.13 Preparation of bionanocomposites constituted by HAp and gelatin nanofibers.

Reproduced with permission from Ref. [104] © 2005, WILEY-VCH Verlag GmbH & Co. KGaA,

Weinheim.

The problem can be efficiently solved by electrospinning an organic solution of

a previously formed gelatin/HAp precipitate (Figure 3.13). HAp nanocrystallites

appeared in this case well distributed in the gelatin matrix displaying an homo-

geneous nanofibrous morphology. Interestingly, aminoacids belonging to the

biomacromolecule seems able tomodulate the precipitation ofHAp crystals [104].

Alginate has been extensively used for tissue engineering scaffolds for bone, car-

tilage, and skin. Such interest in alginate is attributed to its chemical structure,

which resembles glycosaminoglycan one of the major components of the natural

extra cellular matrix (ECM) in human tissue [105]. Alginate forms a stable hydro-

gel in presence of low concentrations of divalent cations, such as Ca2+, through

ionic interactions with the carboxylic functional groups contained in the alginate

molecular chains [106]. HAp/alginate nanocomposite fibrous scaffolds obtained

using electrospinning and a novel biomimetic in situ synthesis has recently been


Ca2+

Ca2+

Ca2+ Ca2+

Ca2+

Ca2+

Ca2+

Ca2+

Ca2+

PO4

3−

PO4

3−

Ca2+

PO4

3−

HO

OHOH

COO−COO−

COO−

COO−

COO−

COO−

COO−

COO−O

OO

OO

OO

OO

O

OH

O O

OH

O

O

O

O

HO

HO

OO

HO

HA

p

HA

p

HO

OHOH

COO− COO−

COO−

COO−

COO−

COO−

COO−

COO−O

OO

O

O

OO

O

O

O

OH

OO

OH

O

OO

O

HO

HO

OO

HO

Ca2+ Ca2+

OH−

HAp nanocrystals

~ 200 nm

OH−

PO4

3−

Ca2+

OH−

OH−

HO

OH OH

COO− COO−

COO−

COO−

COO−

COO−

COO−

COO−O

OO

O

O

OO

OO

O

OH

O O

OH

O

O

O

O

HO

HO

OO

HO

(a) (b)

(c) (d)

Figure 3.14 Chemical structures of “egg-

box” model of calcium alginate (a), “egg-

box” model of calcium alginate with pre-

cursor ions for HAp nucleation (b), and

mineralized “egg-box” structure with HAp (c).

(d) Scheme of cross-linked/in situ synthesized

HAp/alginate nanocomposite fibrous scaffold.

Reproduced with permission from Ref. [107]

© 2013, Springer.

proposed (Figure 3.14). In this way, poorly crystalline HAp nanocrystals were

induced to nucleate and grow at the [–COO−]–Ca2+–[–COO−] linkage sites on

electrospun alginate nanofibers impregnated with PO43− ions [107]. This novel

process resulted in a uniform deposition of HAp nanocrystals on the nanofibers,

overcoming the severe agglomeration of nanoparticles processed by the conven-

tional mechanical blending/electrospinning method. Attachment of rat calvarial

osteoblast cells on these alginate scaffolds was more stable than attachment on

pure alginate.

HAp-biopolymer nanofiber mats based on modified natural polymers like cel-

lulose acetate have also been prepared. Thus, electrospun fibers with a diameter

size larger than 1 μm and homogenous distributed HAp particles were produced

using a mixture of acetone and isopropanol as solvent and after a careful selection

of the processing parameters: applied voltage, diameter of needle, distance

between the needle tip and the collector and flow rate [108].

In addition to natural polymers like collagen, alginate, and chitosan, different

biodegradable synthetic polymers have also been evaluated to get nanocompos-

ites with bioactive inorganic materials by using the electrospinning technique.

Thus, PLA [109], poly(lactide-co-glycolide) (PLGA) [110], polycaprolactone (PCL)

[111] and poly(hydroxybutyrate) [112] (PHB) have been assayed with different


success due to the problems associated to their hydrophobic nature that makes

difficult to get a homogeneous and gooddispersion of the inorganic phases. In fact,

nanoparticles tend to agglomerate in the electrospinning solution and lead to the

formation of beads. For example, this problem has recently been avoided using

surfactant molecules to stabilize the interphase between HAp particles and the

hydrophobic polymer (e.g., PLA) [113]. Derived nanofiber scaffolds can promote

osteoblastic cell growth and phenotype expression at higher level than scaffolds

based on fibers without the bioactive HAp.

Great efforts are consequently focused into control the homogenization of the

inorganic/organic system and to avoid the disruption of fiber morphology, being

the use of ultrafine HAp particles a key tool. Interfacial adhesion has also been

strengthened by modifying HAp with surface-grafted polymers to improve inter-

actions with the hydrophobic polyesters [68].

Porous, three-dimensional poly (D,L-lactide-co-glycolide) (PLGA)/nano-HAp

scaffold have been developed as a potential bone tissue engineering matrix

suitable for high aspect-ratio vessel (HARV) bioreactor applications. The com-

bination of these scaffolds with human MSCs in HARV bioreactors may allow

for the generation of engineered bone tissue. Results have a clinical relevance

(e.g., treatment of bone cancer) since tissue-engineered constructs may provide

alternatives to traditional bone grafts [114].

To improve the compatibility of HAp and poly(L-lactide) (PLLA), the ring-

opening polymerization of L-lactic acid on HAp surfaces with different surface

hydroxyl functionality was performed. Grafting rate of PLLA was dependent on

the nature and steric environment of the hydroxyl groups on the HAp surfaces,

among which HAp modified with hexamethylene diisocyanate tethered ethylene

glycol presented the best grafting percentage (i.e., close to 25%).The PLLA-g-HAp

could be stably dispersed in chloroform and could be easily electrospun giving

bone guided regeneration membranes of potential interest [115].

Favorable blending of HAp with hydrophilic polymers such as PEG has also

been demonstrated effective to improve properties as a result of the strong

interfacial adhesion between HAp and the hydrophilic polymer [116]. Thus,

PEG–HAp composites containing 15% HAp withstand strains of ∼2000%without breaking [117]. Unfortunately, hydrophilic PEG lack biodegradability

and is not stable in aqueous environments without chemical cross-linking,

making underivatized PEG unsuitable for fabricating degradable HAp–polymer

composites by electrospinning. To overcome this challenge, an amphiphilic tri-

block copolymer polylactide–poly(ethylene glycol)–polylactide (PELA) was also

evaluated [117]. Hence, an HAp–PELA suspension at 25wt% HAp content could

be electrospun to render composite scaffolds with uniform fiber dimensions.

HAp–PELA was highly extensible (failure strain >200%), superhydrophilic (∼0∘water contact angle), promoted osteochondral lineage commitment of bone

marrow stromal cells, and supported osteogenic gene expression upon induction.

Results clearly supported that incorporation of PEG appears an effective strategy

to improve the performance of degradable polymer/HAp composites for bone

tissue engineering applications.


HAp composites with PLA or PLGA have good mechanical properties (e.g., the

fragility of implant surfaces can be reduced) but may show unfavorable effects

caused by acidic degradation products (e.g., lactic and glycolic acids) from these

polymers on the surrounding cells [118]. Hence, an increasing interest exists

to explore the potential use of other biodegradable polymers. Polyvinyl alcohol

(PVA), a water-soluble and biodegradable polymer, has been used extensively in

the pharmaceutical industry because of its biocompatibility, proven mechanical

strength, and anabolic effect on bone formation [119]. In addition, PVA has a

self-crosslink capability (film or hydrogel forming) due to the abundant number

of hydroxyl groups coming from the monomer side chain. However, nanofibers

have limitations, including fast hydrolysis and a bioinert nature that hinder

protein and cell adhesion [120–122]. In order to improve the properties of

PVA nanofibers, HAp nanorods, and collagen were incorporated during the

electrospinning process. These compounds were able to interact with PVA

molecules increasing the hydrolytic resistance and improving mechanical prop-

erties. These inorganic–organic blended nanofibers were found to be degradable

in vitro and showed also an enhanced adhesion and proliferation of murine

bone cells [123].

As an alternative approach calcium-containing PVA nanofiber scaffolds were

prepared by electrospinning, and then mineralized by incubation in a solution

containing Ca–P to form a HAp layer [124]. These seeded calcium ions in

the electrospun nanofibers could act as nucleation sites and improved further

crystal growth during incubation treatment. Highly porous 3D nanofibrous

polymer/HAp mineral biocomposites were successfully prepared with potential

applications in bone tissue engineering.

Electrospun scaffolds were also prepared from HAp nanoparticles and

PLGA/PCL [125]. It was shown that the integration of HAp could slow down the

degradation rate of PLGA-based materials in an HAp-dependent manner. Weak

alkaline HAp may neutralize acidic degradation products of PLGA and therefore

may avoid their adverse effect on the host tissue response as demonstrated by

lower filtration of inflammatory cells after subcutaneous implantation.

Biological properties such as cell proliferation, cell attachment, and alkaline

phosphatase activity were found to increase when nano-HAp was deposited

on the surface of electrospun fibers (e.g., PCL-gelatin) via alternate soaking

process instead of electrospinning a polymer solution containing nanoparticles

[126]. Interestingly, the alternate soaking appears as an efficient method for HAp

mineralization on scaffold for bone tissue engineering.

HAp coating has also been performed over homogenous chitosan electrospun

nanofibers by incubation in SBF (Figure 3.15). Six-day incubation was found to

be sufficient to bring about maximum mineralization of the chitosan nanofibers.

In addition, cell viability and differentiation on these coated nanofibers were sig-

nificantly higher than on non coated chitosan nanofibers [127]. The amino and

hydroxyl groups on chitosan acted as nuclear sites for the formation of HAp in

SBF treatment. Moreover, the increase in the specific surface area of scaffolds

increased the effective density of nuclei for HAp formation.


20kV × 5000 5 μm NTUST

Figure 3.15 SEM micrographs showing chitosan electrospun nanofibers coated with HAp

crystals after incubation in SBF for 6 days. Reproduced with permission from Ref. [127] ©

2012, Springer.

Electrospinning was applied to fabricate PLLA membranes that were grafted

on their surface with chitosan through aminolysis reactions.The bioactivity of the

membrane was demonstrated by x-ray photoelectron spectroscopy (XPS) after

soaking in SBF.The deposits had a Ca/P ratio of 1.6, indicating the HAp formation

on PLLA/chitosan membrane. Compared to a pure PLLA electrospun membrane

that was almost nondegradable, the degradation rate of PLLA/chitosan composite

was up to 20% in 6weeks while maintaining its basic architecture to keep support-

ing the regenerated tissue [128].

Electrospraying of HAp nanoparticles onto the surface of polymer nanofibers

appear also a promising methodology to enhance adhesion, proliferation, and

differentiation of MSCs. Promising results were specifically attained when HAp

nanoparticles were electrosprayed on the surface of electrospun PCL nanofibers

(420± 15 nm) for bone tissue engineering [129].

Finally, it is also interesting to mention that nanofibrous scaffolds of a

biodegradable alanine-substituted polyphosphazene were prepared by electro-

spinning and subsequently loaded with precursors that formed Ca-deficient

HAp upon hydrolysis in aqueous media. It was observed a delayed conversion to

Ca-deficient apatite, which was interpreted as an evidence that precursors were

encapsulated within the nanofibrous scaffold [130].

Stereolithography (Figure 3.16) is a versatile technique that allows to fabricate

with high accuracy structures from the submicron size to the decimeter size for

multiple applications. Stereolithography has recently been applied to get medical

implants based on biodegradable polymers [131]. The limited number of resins

that are commercially available for processing by stereolithography has often

been considered the main limitation of the technique.The resin should be a liquid

that rapidly solidifies upon illumination with light.The biodegradable macromers

that have been applied in stereolithography are based on functionalized oligomers

with hydrolyzable ester or carbonate linkages in the main chain. Main systems


3D design 2D slicing

CT scanning Biodegradable construct

(a) (b)

Stereolithography fabrication

100 μm

500 μm

5 mm

Figure 3.16 (a) Scheme showing the pro-

cesses involved in the design and fabrica-

tion of structures by stereolithography [132].

The designed structure is virtually sliced

into layers of 25–100 μm thickness that are

used in the layer-by-layer fabrication process.

Data are uploaded to the stereolithogra-

phy apparatus to fabricate the structure.

Computed tomography (CT)-scanning allows

assessing the accuracy of the process. (b)

Scaffolds prepared by stereolithography from

poly(D,L-lactide)-fumarate. Reproduced with

permission from Refs [132, 133] © 2010,

2009, Elsevier, American Chemical Society

respectively.

are based on trimethylene carbonate, ε-caprolactone, lactide, and fumarate units.

Furthermore, HAp particles are usually dispersed in the resins to improve the

bioactivity of resulting scaffolds [132].

A poly(D,L-lactide) (PDLLA)/nanosized HAp composite resin was prepared

and used to fabricate composite films and computer designed porous scaffolds

by microstereolithography. To this end, varying quantities of nano-HAp powder

and a liquid photoinitiator into a photo-crosslinkable PDLLA-diacrylate resin

were employed. The stiffness of cured nanocomposites was found to increase

with increasing concentration of nanoparticles [134].

Injectable hydrogels with improved solution stability and enhanced

bone repair function were developed by blending triblock copolymers (i.e.,

poly(ε-caprolactone-co-p-dioxanone)-poly(ethylene glycol)-poly(ε-caprolactone-co-p-dioxanone)) with nano-HAp. The incorporation of inorganic nanoparticles

into polymer matrix led to a controlled decrease on critical gelation temperature

respect to the pure hydrogel [135].

3.6

Applications of HAp Nanocomposites Based on Biodegradable Polymers as Drug

Delivery Systems

HAp has high absorbability and binding affinity with a variety of molecules and

therefore, constitutes an ideal compound to be used as drug delivery system, and

3.6 Applications of HAp Nanocomposites as Drug Delivery Systems 73

also in separation, extraction, and purification of proteins [136]. CaPnanoparticles

can be easily dissolved at low pH as explained earlier (e.g., in lysosomes and even

in the environment of solid tumors) and consequently they can easily release the

incorporated drug in appropriate environments.

Biocomposites designed for tissue engineering applications are a clear example

of interesting drug delivery systems since they can have an added value when act

as reservoirs for drugs. The sustained release of antibiotics and growth factors to

eliminate infection and insure osteoblast differentiation is, for example, a relevant

topic for the design of macroporous implantable devices for osteogenesis [137].

BMPs, especially BMP-2, are themost effective in inducing complete bonemor-

phogenesis. A controlled, localized delivery system is of utmost importance in

protecting BMP-2 bioactivity and prolonging its presence at the defect site for

effective bone regeneration [137].

Bioactive molecules can be incorporated into biomaterial scaffolds by physical

adsorption (e.g., by direct immersion of the polymeric scaffold into growth fac-

tor solution). This simple method can achieve local delivery but also a limited

temporal control over release kinetics [138]. Alternatively, growth factor can be

incorporated during the scaffold preparation, being possible in this case to get a

homogeneous distribution and a slower release. However, in order to not dam-

age the bioactive molecule during the scaffold preparation step, cautions must be

appropriately taken into account [139].

The adsorption and release of drugs depends also on the morphology of HAp

nanoparticles (e.g., plate-shaped, needle-shaped). In general, the studies until now

performed indicated that HAp nanocrystals and drugs can be selected in such a

way that the bioactivity of the drug–HAp conjugate could be tailored for specific

therapeutic applications [17]. Some interesting recent works focused on the use

of HAp as drug delivery system merit to be commented.

Minocycline, a semi-synthetic tetracycline antibiotic that is also interesting for

enhancing bone formation, decrease connective tissue breakdown, and diminish

bone resorption, was loaded in a biomaterial synthesized using a biomimetic

method. Specifically, a measured amount of an acidic HAp and minocycline

solution was added to a gelatin solution and kept at 40 ∘C and pH 7–8 for 2 h.

The HAp–gelatin–minocycline composite was obtained after aging overnight

and liophilizing. Nano-HAp was found to be well distributed evenly in the

fibrils of gelatin. The drug was slowly released from the composite particles (i.e.,

over 2weeks in vitro), and promoted rat bone marrow stromal cells adhesion,

proliferation, and differentiation in vitro [140].

A gelatin/nano-HAp scaffold was prepared by glutaraldehyde chemical cross-

linking of a gelatin aqueous solution with nano-HAp granules and then BMP-

2 loaded fibrin glue was incorporated. The prepared hybrid scaffold had a 3-D

porous structure and was able to be used as a BMP-2 sustained release system to

improve the regeneration in vivo of a critical-size segmental bone defect [137].

HAp/collagen–alginate bionanocomposites have been developed as a bone

filler and drug delivery vehicle. Specifically, growth factors that stimulated bone

formation were loaded in the nanocomposites [141]. Porous HAp/collagen


scaffolds are highly efficient for both bone and cartilage regeneration and have

additionally been designed as carriers for fibroblast growth factor [142].

Calcium-deficient hydroxyapatite (CDHA)/chitosan nanocomposites have also

been prepared as drug-loaded matrices, and the controlled release of vitamins

from such matrices evaluated [143]. The role of polymer–filler interaction in the

drug release was also evaluated. Specifically, comparison involved samples pre-

pared by in situ incorporation of CDHA nanoparticles (i.e., CDHA synthesized in

the presence of chitosan) and by simple addition of previously synthesized CDHA

into a chitosan solution. It was found that both the amount of CDHA incorpo-

rated and the synthetic process altered significantly the extent of filler–polymer

interaction, which influences strongly the diffusion exponent and permeability

of CDHA/chitosan nanocomposites. Hence, CDHA nanocrystals could concur-

rently play the roles as bioactive nanofiller and drug-release regulator.

Electrospun scaffolds composed of PCL, collagen I, and nano-HAp

(PCL/col/HAp) were found to support greater MSCs adhesion, proliferation, and

activation of integrin-related signaling cascades than scaffolds composed of PCL

or collagen I alone. In addition these bone-mimetic scaffolds were proved to serve

as carriers for delivery of the platelet-derived grown factor (PDGF)-BB, which

is able to mediate osteoblast chemotaxis. This grown factor was adsorbed to,

and subsequently released from PCL/col/HAp scaffolds in a higher amount than

using conventional PCL scaffolds. The PDGF-BB released was chemotactically

active, indicating that bioactivity was not diminished by adsorption to the

biomaterial [144].

Novel coaxial electrospun PCL/PVA core-sheath nanofibers blended with both

HAp nanorods and type I collagen (Col) (PCLCol/PVAHAp) have been studied.

Doxycycline and dexamethasone were successfully incorporated into these coax-

ial nanofibers for controlled release.These nanofibers encapsulating drugs showed

great potential in enhancing implant osseo integration and preventing implant

infection [145].

Porous triphasic composite scaffolds for bone tissue engineering and drug deliv-

ery system were also prepared from nano-HAp, biodegradable Ca cross-linked

sodium alginate (SA) and PVA by the method of coprecipitation. It was demon-

strated that nano-HAp component could disperse uniformly in SA-PVA copoly-

mer matrix. Excellent miscibility existed among the three phases and inter- or

intrahydrogen bonding could be formed among the three phases. The entrance of

PVAmatrix in the composite enhanced themechanical properties of the compos-

ite scaffold [146].

Osteomyelitis is a tough disease that produces progressive bone destruction and

the formation of sequestra. Continuous spread of infection, hematogenous seed-

ing, and direct inoculation of microorganisms are possible causes that should be

avoided by using, for example, gentamicin (GM) as an aminoglycoside antibiotic.

GM-impregnated microspheres were evaluated to extend the drug-release time

for the treatment of chronic osteomyelitis.The granules were prepared in solution

and consisted of nano-HAp, chitosan, and GM-loaded ethyl cellulose (EC) micro-

spheres.These granules were provided with excellent drug release properties (e.g.,

3.6 Applications of HAp Nanocomposites as Drug Delivery Systems 75

49 days in vitro and 45 days in vivo) that support an outstanding curative effect in

the treatment of chronic osteomielitis [147].

Therapeutic effect of the tetracomponent system constituted by nano-

HAp/poly (3-hydroxybutyrate-hydroxyvalerate)-polyethylene glycol-GM has

been evaluated as a local drug delivery system for osteomielitis treatment. Staphy-

lococcus aureus was injected into rabbit tibia to determine the effect of delivered

drug. Results showed that the GM loaded scaffold could be implanted as primary

graft into the remaining infected defect to effectively treat osteomielitis [148].

Chitosan microspheres (CMs) encapsulated with synthetic peptide derived

from BMP-2 were prepared and incorporated on a scaffold consisting on HAp,

collagen, and PLLA (Figure 3.17) [148].

CMs

(a)

(b)

(c) (d)

10 μm 100 μm 100 μm

(e)

nHAC

PLLA solution

Lyophilized

Microsphere–scaffold mixture CMs/nHAC/PLLA composite

Figure 3.17 Schematic illustration show-

ing the fabrication route of scaffolds con-

stituted by loaded chitosan microspheres,

HAp-collagen (nHAC), and PLLA (a). Sam-

ples of the three-component composite

that incorporates the drug-loaded chitosan

microspheres (b). SEM micrographs of loaded

chitosan microspheres (c), three component

scaffold (d), and the scaffold incorporating

microspheres (e). Reproduced with permis-

sion from Ref. [149] © 2011, John Wiley and

Sons.


The scaffolds appeared as an ideal delivery system for the sustained release of

BMP-2-derived synthetic peptide and offered an innovation for the delivery of

growth factors [149]. The excellent biocompatibility of the CMs/nHAC/PLLA

(nHAC=HAp-collagen) composite was attributed to both the chitosan compo-

nent and the bioactive synthetic peptide encapsulated inside.

CMs containing adrenomedullin (ADM), a bioactive regulatory peptide that

affects migration and proliferation of diverse cell types, were also incorporated

and well dispersed into a hybrid scaffold constituted by poly(lactic-co-glycolic)

acid and nano-HAp. The addition of CMs increased water absorption and

improved the mechanical properties of the scaffolds without affecting their high

porosity. The expression levels of osteogenic-related and angiogenic-related

genes were also improved on the ADM delivery scaffolds, enhancing the interest

of such for bone tissue engineering [150].

Alginate/HAp microspheres were prepared by adding HAp powder to an

aqueous alginate solution and subsequent drop-wise extrusion of the formed

paste into a CaCl2 cross-linking solution. Spherical-shaped particles were

instantaneously produced with a size that could be controlled by regulating the

extrusion flow rate. Effective doses of antibiotics (i.e., amoxicillin, erythromycin)

were previously loaded by immersion of HAp nanoparticles in antibiotic solution

and subsequent drying. Osteoblasts proliferated well on microspheres, being cell

growth enhanced in the presence of antibiotics and specifically erythromycin

presented the most beneficial effect. Combining the sustained antibiotic release

with the osteoconduction, resorbability, and potential use as injectable bone

filling material of porous HA microspheres, these systems provided a forth fold

beneficial effect [151].

Microwave irradiation method was used to synthesize acid form HAp (pH= 7),

and nano-HAp/chitosan–gelatin composite microspheres were subsequently

prepared with the water/oil method combined with multiple emulsification

chemical crosslink technique. HAp was greatly embedded by chitosan–gelatin

giving rise to spherical microspheres (diameter between 10 and 30 μm). Gentam-

icin could be effectively loaded (32.97%) with an average entrapment efficiency

of 49.20%. Nanoparticles were able to maintain therapeutic concentration within

3 days [152].

3.7

Miscellaneous Applications of HAp Nanocomposites Based on Biodegradable Polymers

Nanohydrogels are acquiring a great potential for biomedical applications since

may have clear advantages for delivery of hydrophilic small-molecule drugs

and protein/peptide therapeutics due to their huge loading capacity of water-

soluble compounds. Nanohydrogels combine some advantages of hydrogels

and nanoparticles as, for example, controllable drug release, high stability in

physiological media and distinct responsiveness to environmental factors such as

pH and temperature [153]. Different examples concerning hydrogels containing

HAp nanoparticles can be mentioned.

3.7 Miscellaneous Applications of HAp Nanocomposites Based on Biodegradable Polymers 77

Nano-HAp has been incorporated into a thermosensitive PEG-PCL-PEG

hydrogel to form an injectable composite with interest for application in the

bone tissue engineering field [154]. This hydrogel nanocomposite showed ther-

moresponsivity and specifically it remained in the liquid state at low temperature

whereas flowed freely at a temperature of about 37 ∘C. In fact, nano-HAp

increased the temperature interval at which the sample remained in the gel. The

hydrogel nanocomposite behaved as an injectable fluid and could form a gel in

the desired tissue, organ, or body cavity in a minimally invasive manner.

Nerve growth factors (NGFs), which are vital in maintenance and regeneration

of nerves, play an important role in bone regeneration since they are able to stim-

ulate differentiation and inhibit apoptosis of osteoblastic cells [155]. Injectable

hydrogels, as a drug delivery system, may elevate the efficiency of NGF applica-

tion, but usually most of their protein content is released at an initial stage and the

protein rapidly cleared from the body because of enzymatic degradation. To over-

come this limitation, microparticle/hydrogel systems, such as a collagen/nano-

HAp material into alginate hydrogel, have been developed. Studies indicated that

NGFwas kept from its rapid degradation andwas able to retain its biological activ-

ities for a prolonged period until its release from themicroparticle/hydrogel [156].

Development of artificial corneas is extremely important for diseases that

cannot be treated with corneal transplantation. Unfortunately, most artificial

corneas had failed due to the poor compatibility between the artificial implant

and host cornea, as well as the poor attachment of the transparent center to the

rim. Systems based on a skirt of porous nano-HAp/polyvinyl alcohol hydrogel

(nano-HAp/PVA-H) and a transparent core consisting of PVA-H appear highly

promising. PVA-H has high strength, elasticity, and high water content while

porous nano-HAp/PVA-H skirt is favorable for improving biocompatibility,

hydrophilicity, and flexibility. A tight attachment of the skirt to the core was

attained through an interpenetrating network along the interface. Materials were

implanted in eyes of rabbit and showed good biocompatibility and interlocking

with the host tissues [157].

Collagen/HAp biomimetic nanostructured coatings have been applied on

conductive material surfaces, like titanium, by an electrochemically assisted

deposition in order to improve the surface bioactivity [158]. Aqueous electrolytes

containing proper amounts of Ca(NO3)2 and NH4H2PO4 can be employed since

electrochemical reactions are induced in the cathode after applying a current.

During the electrochemical process, the pH increased up to about 9.0–10.0,

leading the precipitation of the mineral apatitic phase on the cathode electrode.

Acidic collagen molecules (e.g., soluble type I) can be added in a small propor-

tion in the electrolyte (e.g., 0.012w/v-%) leading to a self-assembly of collagen

molecules into reconstituted fibrils during the contemporary crystallization of

CaP mineral on the electrode surface [159].This electrochemically assisted depo-

sition of biomimetic HAp/collagen coating on ceramic and metallic prosthesis

opens many opportunities to optimize the bone-prosthesis interface since bone

conductivity and prosthesis immobilization can be improved [160]. In order to


prevent interaction of blood with surface materials, which may induce coagula-

tion and thrombus formation, it seems also highly interesting to functionalize the

biomimetic coating with an anticoagulant like heparin. This feature can be easily

attained by incorporating the drug into the electrolyte solution.

A versatile electrochemically assisted deposition method for nonconductive

substrates that allows the formation of a stable coating on the pore walls of

scaffolds has recently been developed, appearing as a promising powerful route

for the improvement of synthetic materials bioactivity [161]. TiO2/HAp scaffolds

with biomimetic morphology and suitability for tissue engineering applications

have been prepared via a sol–gel/polymeric sponge process. To this end, cellu-

lose sponges were soaked in an isopropanol solution of Ti(OC3H7)4 containing

needle-shaped HAp nanocrystals. Electrochemically assisted deposition of HAp

was subsequently carried out using a platinum coil bent to host the scaffold-like

into a conductive cage with the aim of producing a local increase of OH−

concentration, centered on the scaffold (Figure 3.18). The direct HAp deposition

was clearly observed on the outer surface of the scaffolds, but the motion of

ions into sample pores was found to be hindered by the charging of the scaffold

surface and by partial pore occlusion. Alkaline treatment caused that the TiO2

inner surface was predominantly covered by amorphous calcium carbonate easily

converted to HAp under nearly physiological conditions.

CaP/collagen coatings similar to the natural human bone have also been pre-

pared on the surface of carbon/carbon (C/C) composites by electrochemically

assisted codeposition technique. A three-dimensional collagen network structure

covered by uniform CaP aggregates was demonstrated to be formed on the C/C

composites. HAp was found the most favorable composition in the coatings with

the increase of the collagen concentration in the electrolyte. The formed collagen

i = 2.5 mA

Ca(NO3)2

NH4H2PO4

H2O

O2 + 2e− + 2H+12

H2 + 2OH−

Concentration Concentration

OH− OH−

2H2O + 2e−

0.042 M

0.025 M

Figure 3.18 Electrochemical setup employed for the HAp deposition onto nonconductive

porous substrates. Reproduced with permission from Ref. [161] © 2011, Royal Society of

Chemistry.

3.8 Concluding Remarks 79

network increased the cohesive and adhesive strength of the coatings due to the

formed collagen network [162].

3.8

Concluding Remarks

Composites containing HAp are a fruitful field of research. The biocompatibility

of HAp and its capacity of being combined with a large quantity of substances

and processed with several technologies have allowed the synthesis of a variety of

(nano)composites with enhanced properties and interestingmedical applications.

Its excellent mechanical and biocompatible properties are related to its crystallo-

graphic structure and its chemical composition.The complexity of the structure is

complementedwith its capacity of incorporating different ions andmolecules that

strongly affect the crystallization process and its macroscopic properties. Addi-

tionally this process is highly depending on common factors like temperature,

pressure, and pH variations.

All these previous features become the basis for the composite research field

looking for enhanced properties whenHAp is combinedwith biocompatible poly-

mers, particularly with the biodegradable ones. Bone tissue is formed combining

HAp nanoparticles and collagen fibers to obtain a surprising material regarding

its strength and flexibility. Researchers have developed different approaches in

order to merge both HAp and biocompatible polymers in order to explore how

some challenging clinical situations can be overcome. These composites are able

to perform with increased strength, long-term stability, enhanced mechanical

properties, or increased biocompatibility. Some promising technologies have

been applied to obtain the composites, as mentioned extrusion, electrospinning,

and pultrusion, but it can be easily understood that new technologies able to mix,

melt, or make more fluid without degradation HAp and polymers, are potential

candidates to be used for obtaining new generations of composites with better

physical properties.

The importance formedical applications is not only related to the physical prop-

erties of the composites. They are also related to the chemical characteristics that

make possible to encapsulate diverse substances in the HAp. They can be used as

nanocarriers to target specific cells (i.e., tumoral cells) or to control the release

of the drug in order to achieve more effectiveness or delayed release. This strat-

egy allows multiple combinations with biodegradable polymers that protect the

nanocapsule from the biochemical attack or delays the release of the pharmaceu-

tical substance in order to achieve a long-term efficacy. The technologies related

to coprecipitation, coating, and emulsion, among others, are relevant to make the

particle compatible. They can take profit of the mechanisms used by the cell to

internalize the particles, digest them by dissolving or degrading (i.e., by means of

enzymatic or hydrolytic mechanism).

Regarding the impact of the biological aspects of HAp composites, It has to

be remarked the case of HAp and DNA or RNA, where the inorganic structure


allows the encapsulation and later delivery of the biomolecule into the cell nucleus.

This is particularly interesting because of the infrequent capacity for an inorganic

material to encapsulate DNA or RNA without losing their functionality. Once the

DNA is encapsulated, the capsule is able to penetrate the cell membrane and travel

through the cytoplasm until reaches the nucleus, dissolve, and release the func-

tional DNA that finally recombines with the cell DNA. It is expected a long but

intense research to understand the mechanisms that are necessary for using this

transfection process for the benefit and applicability of gene therapy. As other

methods of transfection have their own risks, HAp remains as a known and safe

alternative to be combined with new technologies and approaches for improved

gene therapy.

It is also worth to mention that HAp can be combined with biodegradable poly-

mers to form scaffolds with specific characteristics of porosity and mechanical

properties where the cells can migrate and regenerate the tissue in the meantime

the polymer is degraded. This approach for tissue regeneration open new ques-

tions about how the process can be controlled or improved as state-of-the-art

of surgery procedures needs the best approach for tissue regeneration (i.e., after

tumor resection).

Finally, some of the most important biochemical reactions in the living organ-

isms take place just in contact with the natural scaffold of HAp (i.e., bone marrow

and generation of stem cells and platelets).This fact indicates how important is the

role of HAp for sustaining life. When HAp is combined with natural or synthetic

substances, as polymers, to obtain composites, the challenge and expectations are

high. At the end the composite pretended to mimic or improve what nature has

developed aftermillion years of evolution. To explore if it is possible is whatmakes

so exciting the research in this new field.

Acknowledgments

Authors are indebted to supports from MINECO and FEDER (MAT2012-36205

and MAT2012-34498) and the Generalitat de Catalunya (2009SGR925 and

2009SGR1208). Support for the research of C.A. was received through the prize

“ICREA Academia” for excellence in research funded by the Generalitat de

Catalunya.

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4

Synthetic Methods for Nanocomposites Based

on Polyester Resins

Michał Kedzierski

4.1

Introduction

Polyesters make up a large class of polymers and resins with a wide array of

structures and practical applications; consequently, numerous studies were

devoted to polyester nanocomposites in the past several decades. This chapter

focuses on two branches of the polyester family comprising reactive oligomeric

compounds, that is, unsaturated polyester (UP) and saturated polyester (SP)

resins. UP resins are solutions of UP (prepared commonly by polycondensation of

glycols with the mixture of unsaturated and saturated anhydrides of dicarboxylic

acids) in an unsaturated crosslinking agent, usually styrene. They can be cured

by radical copolymerization with the formation of polystyrene chains linking the

prepolymer molecules. Because of the low cost raw materials, simple preparation

procedures and a variety of glycols and acid monomers, which allow to tailor

the properties of the cured polymer, UP resins have found many applications as

binders in fiber-reinforced laminates and composites for construction, trans-

portation, and building industry. Nonreinforced crosslinked UP resins are used

as binders in the manufacture of artificial marble, polymer concrete, gel coats,

and repair putties.

Vinyl ester (VE) resins are a subclass of UPs produced by the esterification of

epoxy resins with unsaturated carboxylic acids. Because of the lower content of

hydrolytically unstable ester bonds, they exhibit better chemical resistance than

typical UP resins; moreover, their mechanical and thermal properties are similar

to those of epoxy resins that are available at lower cost.

Another group of reactive prepolymers are SP resins, which can be crosslinked

using hydroxyl or carboxyl functionality. They constitute a smaller segment of

thermoset resins market; however they are of importance in the paint and coating

industry.

Being easy to process and inexpensive matrices for composites, polyester resins

have been the subject of numerous investigations focused on the improvements of

mechanical, thermal, fire retardant, and other properties of the end products.The

formation of nanocomposites by the combination of nanometer-sized particles


88 4 Synthetic Methods for Nanocomposites Based on Polyester Resins

with polymer enables the modification of its properties at much lower loadings of

additives than using conventional reinforcements or micrometer-sized fillers.The

use of nanoadditives not only allows to enhance the properties of polymer matrix

without affecting its weight and performance, but in some cases it can result in

the material features unattainable with traditional composite materials.

Nanoscale additives can be classified according to the number of dimensions

confined to the nanoscale range (below 100 nm) [1, 2]. For example, carbon nanos-

tructures comprise fullerene (0-D type: all dimensions at the nanoscale), nan-

otubes (1-D: one dimension outside of the nanometric size range), graphene (2-

D), and graphite (3-D). The dispersion and arrangement of nanofiller in the resin

matrix plays a central role in controlling the properties of the resulting nanocom-

posites. Therefore, a range of research has been undertaken to investigate the

effects of various synthetic methods on the morphology and properties of the

resulting polyester resin-nanofiller blends. Results of these studies are discussed

in the sections related to the various nanofiller types.

The process of synthesis and curing of polyester oligomers comprises two or

three subsequent stages, as demonstrated in Figure 4.1 for exemplaryUP resin and

Figure 4.2 for SP resin. Accordingly, nanofillers can be introduced at various steps

O

O

O

O

O

O

O O

O

R

O n

O O

O

R

O

+ + 2 HO-R-OH

– H2O

1. Polycondensation

UP prepolymer (alkyd)

2. Dissolution incrosslinking monomer

UP resin

Peroxide initiatoraccelerator

Cured UP

3. Crosslinking

Figure 4.1 Schematic of the synthesis and curing of unsaturated polyester resin.


CH3

CH3HO

HO

HOOH

OH

OHO

O

HO O

O O

H3C CH3

O

OO

O

OH

OO

O

– H2O

+ +

1. Polycondensation

2. CrosslinkingPolyhydroxyl or epoxy

hardener

Polyester coating

Figure 4.2 Schematic of the synthesis and curing of saturated polyester resin.

of the process: (i) during the synthesis of prepolymer, (ii) by mixing with prepoly-

mer before dissolution in crosslinking monomer for UP and VE resins, and (iii)

before final crosslinking. In order to obtain a true nanocomposite structure, the

mixing of nanofiller with polyester should be accompanied by deagglomeration

of nanoparticle aggregates or intercalation/exfoliation of layered nanofillers and

their uniform dispersion in the cured resin matrix.

4.2

Nanocomposites with Zero-Dimensional Nanofillers

4.2.1

Silicon-Containing Nanospheres

The preparation of monodispersed silica nanoparticles via sol-gel process was

reported by Stöber et al. nearly half a century ago [3]. It involved a controlled

hydrolysis of tetraethyl orthosilicate in water–alcohol medium containing

ammonia catalyst and subsequent condensation reactions of the formed silanol

groups. By adjusting the type of silicate ester, reagent concentrations, and molar

ratio, the silica spheres of different particle size from nanometric to micrometric

range can be synthesized. Since then, the sol-gel method has been extended and


developed, along with alternative techniques for nanosilica preparation such as

reverse microemulsion and flame synthesis.

Wu and coworkers investigated the influence of nanosilica particles on the prop-

erties of their dispersions in SP polyol resin [4]. Two synthetic methods were

used.The first was in situ polymerization: silica sol prepared using sol-gel process

was mixed with the monomers before the polymerization process. In the second,

blending method, silica sol was mixed with polyester polyol at 165 ∘C for 30min

with simultaneous solvent evaporation. Fourier transform infrared spectroscopy

(FTIR) analysis indicated that covalent bonds were formed between silanol and

polyester functional groups. A higher extent of this reaction aswell asmore homo-

geneous nanosilica dispersion (as characterized by TEM (transmission electron

microscopy)) was observed for the product obtained by the in situ method. The

critical content of silica, above which a sharp increase of viscosity took place, was

10wt%. The viscosity of UP resin/nanosilica dispersion obtained by the blending

method was higher than that prepared by the in situ polymerization, and only

6wt% nanosilica could be introduced into the resin without a significant viscos-

ity rise. The viscosity also depended on the diameter of nanosilica particles, first

increasing and then decreasing with an increase in their size. Chung and cowork-

ers prepared optically transparent UP/silica nanocomposites using sol-gel pro-

cess of an alkoxysilane precursor in the presence of UP resin [5]. In order to pre-

vent a self-association of polyester carboxyl and hydroxyl groups leading to the

phase separation during the sol-gel process, UP functional groups were blocked

by acetylation. Also, various solvents for sol-gel process were tested, taking into

account the influence of their volatility and hydrogen bonding acceptor proper-

ties on the formation of transparent nanocomposites. Photocrosslinking of the

obtained UP/silica nanohybrids resulted in an interpenetrating polymer network

structure.

Nanosilica (in the amount of 0.5–3wt%) was used to improve the properties

of UP coating applied on the surface of marble to reinforce it and prevent its

breaking during processing and transformation [6]. The UP-nanosilica mixtures

showed increased viscosity, pseudoplastic, and thixotropic behavior as well as sig-

nificantly decreased gel time. The marble pieces coated with nanosilica-modified

resin exhibited improved impact resistance. Sharma and coworkers investigated

the effect of the addition ofmicro- andnanosilica on the electrical properties ofUP

resin composites [7].The surface and volume resistivity, dielectric strength as well

as arc resistance reached the maximum values at the nanosilica loading of 1.5 phr.

Jesson and collaborators used an additional solvent (ethanol) to facilitate the

dispersion of organically modified silicas (ormosils) in UP resin [8]. After sonica-

tion, alcohol was removed under vaccuum and lost styrene subsequently replaced.

The effect of various silica substituents on the fracture toughness behavior of the

UP nanocomposites was investigated. For the methyl, ethyl, and vinyl ormosils,

only a minor toughening effect was observed, while phenyl ormosil gave a greater

improvement in fracture toughness. This was attributed to different toughening

mechanisms following the fact that the densely packed methyl, ethyl, and vinyl

functionalities prevent the retained silanol groups from interacting with polymer


matrix. For the larger and less densely packed phenyl groups, the interaction with

silanol groups is reducedwhen compared to an unmodified silica particle although

not entirely eliminated.

In the study of Mahfuz and coworkers, the fiber-matrix interface of carbon/VE

composites has been modified by coating the carbon fiber with polyhedral

oligomeric silsesquioxane (POSS) [9]. POSS are cage-like siloxane nanostructures

surrounded usually by eight organic substituents with molecular size in the range

of 1–5 nm. Two types of POSS: octaisobutyl (Octa) and trisilanolphenyl (TriS),

have been investigated. Mechanical tests indicated that both interlaminar shear

strength and low velocity impact strength were improved for POSS-containing

nanocomposites. These improvements were more significant with TriS-modified

carbon fibers, which can be explained by the presence of phenyl groups in the

POSS structure and better compatibility with VE matrix.

Oleksy and Galina prepared a mixed nanocomposite system consisting of UP

resin and bentonite clays intercalated with ammonium-functionalized POSS [10].

The nanocomposites containing up to 3wt% bentonite-POSS nanofillers showed

an improved tensile and Charpy impact strength (by 44 and 59%, respectively,

compared to the unmodified resin) as well as better flame resistance (limiting

oxygen index of 25.2 compared to 17.2 for unfilled polyester). XRD (X-ray diffrac-

tometry) and TEM analyses revealed the exfoliation of clay layers in the polyester

matrix.

4.2.2

Metal Oxides

Zhang and Singh investigated the effect of nonmodified and silane-treated

Al2O3 nanoparticles (15 nm average diameter) on the fracture toughness of

cured UP resin [11]. Neat alumina showed poor bonding to the resin matrix

resulting in a lowered crack growth resistance of the composite. However,

when 3-methacryloxypropyltrimethoxysilane was added as a coupling agent

during the nanoalumina dispersion process, the final composites showed an

almost 100% increase in the fracture toughness at 4.5% volume fraction of the

nanofiller. The same authors compared the toughening effect of nanometer-

and micrometer-sized aluminum in the polyester matrix [12, Figure 4.3.]. Singh

et al. also studied the effectiveness of various dispersion techniques (mechanical

mixing, high-shear dispersion, and ultrasonication) for particle deagglomeration

in UP resin-aluminum nanocomposites. The use of solvent-aided ultrasonic

disruption resulted in the best nanoaluminum dispersion and the greatest

increase in flexural modulus and fracture toughness of polyester matrix [13].

In the study of Vijayakumar and colleagues, Al2O3 nanoparticles (diameter

60–70 nm) prepared by sol-gel technique were used to prepare nanocomposites

with UP resin using casting technique [14]. TEM analysis showed a uniform

dispersion of nanoalumina in the polyester matrix up to a loading of 5wt%, and

an agglomeration of nanoparticles was observed above this concentration. The

nanocomposites showed a higher tensile, flexural, and impact strength than


00.9

1.0

1.1

1.2

1.3

1.4

1.5

1.6

1 2 3 4 5

Particle volume fraction (%)

20 μm aluminum

3.5 μm aluminum

100 nm aluminum

Norm

aliz

ed f

ractu

re t

oughness

Figure 4.3 Variation of fracture toughness as a function of volume fraction for reinforce-

ment by 20 μm, 3.5 μm, and 100 nm aluminum particles. Reproduced from Ref. [12] with

permission from Springer.

pristine polyester. Lopez-Cuesta et al. synthesized nanocomposites of UP resin,

nanoalumina, and submicron alumina trihydrate particles. Synergistic effects

on thermal stability and fire behavior (heat release rate) were observed using

combinations of both additives with the best results for a global loading of 10wt%

and an equal weight ratio of both filler types [15]. Sharma et al. observed an

increase in the erosion resistance, mechanical and thermal properties of UP

composites using a combination of nanoalumina (with optimum concentration

below 1 phr) and coupling agent [16].

Copper (II) oxide nanoparticles (with an average diameter of 29 nm) were used

to fabricate VE resin nanocomposites [17]. It was found that functionalization

of nano-CuO with methacryloxypropyl trimethoxysilane (MAOPTMS) as a

bifunctional coupling agent improves the dispersion of the nanofiller into the

resin matrix. The cured nanocomposites showed an increased tensile strength

and modulus (respectively by 50 and 15% at 10wt% loading of silanized nano-

CuO) as well as enhanced thermal stability. VE nanocomposites with improved

mechanical and thermal properties were also prepared using iron oxide (Fe2O3)

nanoparticles and MAOPTMS [18] as well as iron nanoparticles without any

coupling agent [19]. In the second case the authors analyzed a possible mecha-

nism of nanocomposite formation involving the reaction of Fe nanoparticles with

hydroxyl functional groups of VE monomers.

Evora and Shukla used ultrasonification to embed a small loading of 36 nm

average TiO2 particles [20]. The formation of well-dispersed nanocomposites

was confirmed by TEM. The presence of nanoparticles had the greatest effect on


the dynamic fracture toughness, and negligible influence was observed for the

quasi-static properties. A decrease in the nanocomposite mechanical properties

was observed beyond TiO2 volume fraction of 1%, which was attributed to

the clustering of nanoparticles. In another study, tensile, flexural, and impact

strengths of crosslinked UP resin were significantly increased by the incorpo-

ration of 4wt% TiO2 particles of average size 27 nm [21]. Torabi-Angaji and

collaborators modified the surface of titania nanoparticles using organosilane

with methacroyl group, and subsequently copolymerized with acrylate and

methacrylate monomers [22]. The nanocomposites of UP resin with such modi-

fied nano-TiO2 showed improved mechanical and UV-resistant properties. Knör

and coworkers employed organically modified titania nanoparticles as additives

for the protection of powder coatings against UV-dependent degradation [23].

Peng and colleagues investigated three-phase composites of UP resin, glass

fiber, and nano zinc oxide, which showed increased resistance to ultraviolet

degradation and improved impact strength [24].

4.2.3

Other 0-D Nanoparticles

Calcium carbonate nanoparticles of size below 10 nm, synthesized from eggshell

powder via mechanical attrition and high-intensity ultrasonic irradiation, were

used for the modification of soybean-based UP resin [25]. The dispersion was

prepared in a noncontact hybrid defoaming mixer and the cured UP/CaCO3

nanocomposites showed an increase in glass transition temperature by 18 ∘C as

well as improvement in compressive strength (14%) and modulus (27%) at 2wt%

nano-CaCO3 loading, compared to the nonmodified UP resin.

Petrova et al. studied the effect of hard and soft carbon nanoparticles on the

surface properties of the cured polyester resin.The addition of 0.5–5wt% of hard

diamond nanofiller led to an improvement by more than 35% wear resistance of

the crosslinkedUP, while the soft nanofiller, carbon ash, did not influence the wear

properties of the polyester matrix up to 5wt% filler content [26].

4.3

Nanocomposites with One-Dimensional Nanofillers

4.3.1

Carbon Nanotubes and Nanofibers

Carbon nanotubes (CNTs) and carbon nanofibers (CNFs) are the most inten-

sively investigated nanofillers in this group. They differ in the diameter (from 1

to 2 nm for single-walled and 3–50 nm for multiwalled CNT to 50–200 nm for

CNF), lengths, and shapes.

Ni and coworkers used vapor grown carbon nanofibers with the average

diameter of 80 and 150 nm to prepare the nanocomposites with UP resin by a


solution-evaporation method [27]. CNF dispersion was obtained by sonication

in UP ethanol solution with subsequent evaporation of the alcohol, addition

of crosslinking monomer, and curing. The resulting nanocomposites showed

electrical conductivity with a percolation threshold between 2 and 3 vol%. The

same authors investigated temperature dependence of electrical resistivity in

UP resin/CNF systems [28]. In the work of Torre and colleagues the dispersion

of CNFs (45 nm average diameter) in UP resin was obtained by calendering

(three-roll-milling (TRM)) technique [29]. The process parameters were opti-

mized and the rolls cooled in order to minimize the heating of the mixture and

styrene evaporation. Nanocomposites with different concentrations of nanofibers

(0.1–3wt%) were produced. The electrical conductivity measurements showed

the percolation threshold at around 0.3wt% CNF content, while no influence on

the mechanical properties of the cured nanocomposite was observed. Toghiani

et al. demonstrated through a design of experiments how various formulation and

processing factors (i.e., nanofiber type, use of dispersing agent, mixing method,

and nanofiber weight fraction) affected the dynamic mechanical properties of

CNF/VE nanocomposites [30]. Only below 0.50 parts of nanofiber per hundred

parts resin produced a 20% increase in the storage modulus as compared with

that of the neat cured VE resin.

Martin-Gullon and coworkers conducted a comparative study on the disper-

sion of helical-ribbon carbon nanofibers and multiwalled carbon nanotubes

(MWCNTs) in isophthalic UP [31]. The nanocomposites were prepared using

two types of mixing equipment: a propeller mixer with a maximum speed of

2000 rpm and a high-shear rotor/stator (maximum speed of 8000 rpm). The

better nanofiller dispersions were obtained using the high-shear mixer and

the process was accompanied by a reduction in carbon nanofilaments aspect

ratio (Figure 4.4). The cured nanocomposites showed satisfactory mechanical

properties and electrical properties with percolation threshold around 0.1wt%.

Tanoglu and colleagues used as conducting fillers double-walled Carbon nan-

otube (DWCNT) and MWCNT with outer diameter 2.8 and 15 nm, respectively

[32, 33]. They dispersed nanotubes with and without NH2 functional groups in

UP resin using TRM technique. It was found that even low amounts of CNTs

(0.1–0.3wt%) induced an electrical conduction in the resulting nanocomposites.

The higher conductivity values were obtained with MWCNT than when using

DWCNT at the same filler content because of the relatively higher tendency of

DWCNTs for agglomeration within the resin matrix.The electrical conductivities

of nanocomposites with amino-functionalized nanotubes were several hundred

times lower than for those containing untreated CNTs. It was attributed to the

changes in nanotube structure during the functionalization leading to the reduced

conductivity and possible reactions of amine groups with polyester chains. On

the other hand, nanotubes with amine functional groups showed an improved

dispersion within the resin and in effect better tensile mechanical properties of

the resulting nanocomposites. A problem encountered by authors was styrene

evaporation during the mixing process, leading to high increase in the resin vis-

cosity. To overcome this difficulty, the authors used instead of commercial UP


1,E+05Turbine

2000 rpm30 min

Shear2000 rpm120 min

Shear4000 rpm

30 min

Shear4000 rpm120 min

Mixing conditions

Shear4000 rpm

60 min

Shear7000 rpm

30 min

Shear7000 rpm

60 min

1,E+06

1,E+07R

esis

tivity,

Ω. c

m

Figure 4.4 Resistivity of nanocomposites of polyester with 0.5wt% of helical-ribbon carbon

nanofibers mixed at different conditions. Reproduced from Ref. [31] with permission from

Elsevier.

resin a polyester blend containing negligible amount of styrene. The crosslinking

monomer was added after the preparation of CNT dispersion by TRM. In another

study concerning the abovementionedCNT types, their dispersion in styrene-free

polyester was blended with VE resin. The resulting nanocomposites containing

MWCNTs and MWCNT-NH2 showed higher tensile strength and modulus as

well as larger fracture toughness and fracture energy compared to neat hybrid

polymer [34].

The calendering process was also used to fabricate MWCNT/VE nanocompos-

ites [35].The high aspect ratios of CNTwere preserved during the processing and

enabled the formation of a conductive percolating network at low nanotube con-

centrations (below 0.1wt%). Battisti and coworkers prepared the nanocomposites

of UP resin with multiwalled nanotubes (MWNTs) of diameter 9.5 nm and length

1.5 μm using a combination of TRM and high-shear mixing (HSM). The percola-

tion threshold at 0.026wt% loading of nanotubes and the maximum conductivity

of 0.13 Sm−1 for 0.3wt% CNT loading were achieved [36]. A potential applica-

tion of the measurements of rheological parameters and electrical resistivity of

UP/CNT systems for evaluation of the dispersion quality was also investigated

[37]. A study of the effects of concentration and surface chemistry on the dis-

persion and rheological properties of single-walled carbon nanotubes (SWCNTs)

in isophthalic UP was reported by Kayatin and Davis [38]. In another work the

authors studied using rheology and optical microscopy UP dispersions of/single-

walled, multiwalled, and polystyrene-modified MWCNTs [39].

In an attempt to increase an interaction between the nanotubes surface and

polyester matrix, Swain and Patil and colleagues used the chemical modification


of CNT including carboxylation with nitric acid followed by the treatment with

MAOPTMS or two-step allyl ester functionalization [40, 41]. The nanocompos-

ites prepared fromUP resin and the functionalized CNTs exhibited a significantly

increased electrical performance in comparison to those obtained using nonmod-

ified nanotubes. Multiwalled CNT (diameter 40 nm) was also used for the mod-

ification of VE resins [42]. An electrical conductivity of 2.5⋅10−4 S cm−1 of the

resulting nanocomposite was achieved at 1wt% MWCNT content, accompanied

by a slight increase in fracture toughness and energy.

In a recent study, Shokrieh and coworkers have shown that even at 0.05wt%

MWCNT content, the improvements of tensile and flexural strengths of UP

nanocomposites by 6 and 20%, respectively, can be achieved [43]. For higher nan-

otube contents (as 0.5wt%), they form agglomerates acting as stress concentrators

and reducing the strength of nanocomposites.

Kaffashi and Honarvar combined two types of nanofillers: organically modified

montmorillonite (MMT) and MWCNT to prepare UP nanocomposites [44].

At first organoclay – resin dispersion (5 phr) was obtained by mechanical and

ultrasonic mixing with subsequent addition of the nanotubes (0.2–0.4 phr). An

intercalated morphology of MMT platelets and fine dispersion of MWNT in the

UP resin was indicated by XRD and TEM analyses. The nanocomposites showed

increased fire retardancy and improved toughness accompanied by a decreased

tensile strength.

Efforts have also been made to integrate CNTs into the traditional fiber-

reinforced composites. Wang and Qiu incorporated short MWCNT into

low-viscosity polyester/VE resins by ultrasonic processing and shearing [45]. The

modified resins were used tomanufacture glass-fiber composites by resin-transfer

molding (RTM). A 1.5-fold increase of thermal conductivity was observed at

3wt% CNT loading. Another VE/glass fiber/CNT/composite system fabricated

by a vacuum-assisted RTM process was investigated by Barrera et al. [46].

They used several types of the sidewall-functionalized single-walled nanotubes

for overcoating the glass fiber before vacuum-assisted resin transfer molding

(VARTM) processing. A significant improvement of the interlaminar shear

strengths of the resulting VE composites with respect to the nonmodified control

samples was observed even using very small amount of SWCNT (0.015wt%).

4.3.2

Cellulose Nanofibers

Several papers document the polyester resins modified with nanosized cellulose

fibrils. He and coworkers introduced to UP resin bacterial cellulose nanofibers

(BCNs) treated with vinyl-triethoxy silane coupling agent using an RTM process

[47]. The produced nanocomposites showed increased tensile strength (by 118%)

and modulus (28%) as well as flexural and shear strength (by about 38%) at BCN

volume fraction of 10%. Also, an increase in the water uptake of polyester matrix

was observed with an increasing content of cellulose nanofibers [48]. Nakagaito

et al. used UP resin as a binder for cellulose nanofibers, obtained by applying


shear stress to wood kraft pulp fibers in a corotating twin-screw extruder [49].

Paper-like sheets of nanofibers were impregnated with UP resin, arranged in lay-

ers and hot-pressed at 1.5MPa.The flexural modulus and strength of the obtained

composites were significantly improved over the nonmodified cured resin (up

to 13.6GPa and 260Mpa, respectively, at 14wt% resin loading. In the study of

Tercjak and coworkers cellulose microfibrils (MFCs) isolated from sisal were

dispersed in UP matrix using PEO (poly(ethylene oxide))-b-PPO (poly(propylene

oxide))-b-PEO block copolymer (EPE20) as surfactant and toughening agent

[50]. The resulting multiphase nanostructured UP composite (1wt% MFC and

5wt% EPE20) showed improved fracture toughness accompanied by only a low

decrease in flexural modulus.

4.3.3

Other 1-D Nanofillers

Halloysite, a tubular alumosilicate clay with external diameter of 50–80 nm,

lumen of 10–15 nm, and length of about 1 μm [51] has been used in the mod-

ification of polyester resins. In the study of Albdiry and coworkers, various

concentrations (1–9wt%) of halloysite, nonmodified (HNT) and treated with

vinyltrimethoxysilane (s-HNT), were incorporated into the UP resin to produce

nanocomposites using mechanical stirring and ultrasonication [52]. The impact

strength of the cured polyester increased with an addition of 3wt% HNT and

s-HNT by 11 and 16%, respectively. However, further addition of halloysite

resulted in a decrease of impact properties because of the agglomeration of

HNT particles and their poor interaction with polyester matrix. In another

paper, the morphology and tribological properties of the UP/HNT and s-HNT

nanocomposites were also discussed [53]

4.4

Nanocomposites with Two-Dimensional Nanofillers

4.4.1

Layered Aluminosilicate Clays

Most of the research in this field concern the nanofillers based on layered

aluminosilicates (smectite clays), which are inexpensive and readily available.

This class of UP resin nanocomposites was widely discussed in the previous part

of this series [54]. Here, some examples are quoted and supplemented with new

literature in this field.

Smectite clays are also known as 2 : 1 phyllosilicates, such as MMT, which is the

major constituent of bentonite. MMT forms plate-shaped crystals consisting of

1 nm thick layers, which are made up of two tetrahedral sheets of silica fused to

octahedral sheet of alumina. Part of Al3+ cations in the octahedra are substituted

by Mg2+, generating a negative charge of the layers, which is counterbalanced


by the cations present in the interlayer space, so-called gallery. In effect, MMT

has cation-exchange properties, which enable the intercalation of cationic organic

molecules, thereby reducing the clay hydrophilicity and increasing its miscibility

with organic resins and polymers. Two basic forms of clay-containing polymer

nanocomposites are distinguished: intercalated (where macromolecules enter the

clay gallery, usually expanding the interlayer distance, but not affecting the stack-

ing arrangement of the clay layers) and exfoliated (where individual clay platelets

are fully separated and randomly dispersed in the polymer matrix). In practice,

many nanocomposites show mixed morphology consisting of intercalated, exfo-

liated as well as partially delaminated structures with large clay stacks broken up

into few-layer tactoids. The lateral dimensions of these thin clay layers are from

several hundred nanometers to even a microns range, thus, having large aspect

ratio; they can serve as very effective reinforcing nanofillers. Because the forma-

tion of clay nanoparticles takes place in situduring themixing of clay into the resin,

the selection of the processing technique plays an important role in determining

the nanocomposite properties.

4.4.1.1 Mixing Methods

Bashir andHubert compared two approaches to disperse organicallymodified clay

(dimethyl hydrogenated tallow quaternary ammonium – MMT) in the UP resin,

that is, mixing by TRMand ultrasonication [55].The degree of clay dispersion was

evaluated using X-ray diffraction and viscosity measurements. In both cases, an

increase in the resin viscosity withmixing time was observed; however for TRM it

was caused mainly by styrene evaporation. Ultrasonication resulted in better dis-

persion of clay and in this case the viscosity increased along with the progress in

exfoliation. The dispersion of delaminated nanoclay platelets in the resin resulted

in a strong shear thinning behavior – a decrease of viscosity with increasing shear

rate because of the rearrangement of the nanoplatelets in the direction of flow [56].

The rheological properties of UP-resin mixed with 3wt% organically modified

MMTCloisite 30B using three dispersion techniques: manual mixing, sonication,

andHSM,were studied by Bensadoun and coworkers [57].The sonication resulted

in an increase of the viscosity from 0.20 (neat UP resin) to 0.80 Pa⋅s and a slight

non-Newtonian shear-thinning behavior.Theblends obtained usingHSMshowed

high initial viscosity, 70–250 Pa⋅s, and a strong shear thinning. At high shear rate

(20 s−1), the viscosity decreased to 2 Pa⋅s, which made it possible to process the

blend by RTM.

Two approaches for the synthesis of UP/clay nanocomposites were discussed in

the paper of Park et al. [58]. In the first one, simultaneous mixing, UP prepolymer

was mixed with styrene monomer and alkylammonium modified MMT for 3 h at

60 ∘C.The second, sequential mixing, consisted of two stages: (i) preintercalation

of organoclay with UP prepolymer and (ii) mixing of the resulting preintercalate

with styrene at 60 ∘C. In both cases, XRD patterns and TEM micrographs of the

cured products indicated the intercalation and partial exfoliation of MMT layers.

The decrease in Tg of cured UP-organoclay nanocomposite synthesized by simul-

taneous method was observed. It was explained by the fact that styrene molecules


diffuse faster into the gallery of clay than UP chains and a part of crosslinking

monomer is consumed in the homopolymerization process, lowering the crosslink

density of the cured polyester. The nanocomposites obtained via sequential pro-

cess showed Tg values close to unfilled UP resin. In this case polyester chains

were preintercalated inMMT gallery, and crosslinking of UP takes place homoge-

neously inside and outside the clay layers.The reinforcing effect ofMMT platelets

was shown by an increase of the storage modulus of cured UP/MMT nanocom-

posite as compared with pure polyester.

Effects of the processing parameters, such as mixing mode, shear level, organ-

oclay content, and temperature on the morphology of UP-MMT (unsaturated

polyester-montmorillonite) hybrids, were investigated by Narkis and coworkers

[59]. They used two preparation methods: in the first one UP resin was mixed

with organoclay (loading from 5 to 20 phr) at room temperature usingmechanical

stirring or ultrasonication. In the second procedure, styrene-free polyester alkyd

and clay powder were hand-mixed at 80 ∘C, and the resulting blends were further

processed by mechanical mixing at low (400 rpm) or high (1800 rpm) shear level

and the temperature of 80 or 130 ∘C, for several periods of time up to 24 h.

Alternatively, ultrasonication or static heating at 80 ∘C for 2 or 24 h were applied.

Both methods gave nanocomposites of intercalated structure, the extent of which

depended on the type of clay treatment. For UP resin-organoclay system, the

maximum expansion of MMT interlayer (up to 3.8 nm) was achieved using

octadecylamine-treated MMT. Even higher extent of intercalation (interlayer

spacing above 5.5 nm) was observed for some UP alkyd-organoclay systems.

In conclusion, the authors stated that applying high shear levels by vigorous

mechanical stirring as well as increasing the clay content in the blend promotes

the intercalation and results in a better dispersion of clay particles in the resin

matrix. The higher mixing temperature did not increase the intercalation extent,

probably because of reduced shear level. Melt mixing method was used to

prepare the nanocomposites based on styrene-free UP resin and nanoclay [60].

The authors investigated the effect of shear and diffusion-induced phenomena as

well as nanoclay surface modification on the rheological behavior at two mixing

temperatures, 40 and 150 ∘C. The results indicated that at higher shear forces

(cold-mixed samples) the associations of nanoclay stacks were broken more

efficiently and a kind of physical network was formed with nanoclay particles

acting as the nodes and polymer chains as the links. The effects of organoclay

content, mixing mode, and shear rates on rheology behavior of UP resins were

also investigated by Rajabian and Beheshty [61]. Other authors studied the effect

of curing monomer polarity on the properties of UP-MMT nanocomposite [62].

Narkis and coworkers investigated the curing of styrene-free UP (alkyd) in the

presence of organoclay. Depending on the peroxide initiator content, either an

exfoliated or a combined intercalated/exfoliated structure was obtained [63].

4.4.1.2 Effects of the Clay Modification

Effect of the interlayer cation type on the morphology and properties of UP-

MMT nanocomposites was investigated by many authors. For the composites


with nonmodified MMT containing sodium cations in the interlayed space

(MMT-Na), only a slight increase or no change in the mechanical properties was

observed along with a decrease in the impact strength [64, 65]. In another study,

an intercalation of MMT-Na by UP was confirmed using TEM and scanning

electronmicroscopy (SEM), while rheological tests and XPS (X-ray photoelectron

spectroscopy) measurements indicated the formation of a structure between

resin and clay at low shear rates. On increasing the clay content from 1 to 5 phr,

an increase of 57–120% in flexural modulus of the cured UP was observed [66].

MMT-Na was also successfully applied by Rozman and coworkers to improve the

mechanical properties of the composite consisting of UP resin and lignocellulosic

filler – Kenaf [67].

Most researchers used MMT intercalated with various alkylammonium salts

(organoclay) to improve the miscibility of clay with the resin. For example,

Jawahar and M. Balasubramanian employed dodecylamine-modified clay for the

modification of UP-based gel coat system [68]. The maximum increase in tensile

and impact strength by 21 and 33%, respectively, compared to the conventional

gel coat, were observed at 2wt% clay loading. Further increase in organoclay

content led to the decrease of the strength values. The correlation between the

mechanical properties of UP-MMT nanocomposites and the interlayer spacing

of organically modified clay was been observed in the works by Xu et al. [69] and

Dhakal et al. [70]. In several papers the effect of nanoclay on the cure kinetics

of UP resin was studied [71, 72]. Zhang and coworkers mixed UP resin with

sodium, alkylammonium, and methacryloxyalkylammonium MMT for over 24 h

at 50 ∘C [65]. XRD patterns of the cured UP composites showed only partial

intercalation of sodium-MMT by polyester, while the interlayer distances of

organically modified clays were expanded beyond the values detected by XRD

(8.8 nm). TEM analysis showed that most of MMT was dispersed in UP matrix

as small aggregates and a part separated into layers. Using organically modified

clays, a distinct increase in tensile strength, impact strength, and heat distortion

temperature of cured UP was achieved.The best properties were obtained for the

UP modified with organoclay bearing polymerizable methacryloxy group, that

is, 61% increase in tensile and 51% in impact strength as well as 24 ∘C increase

in heat deflection temperature. The use of MMT intercalated with unsaturated

alkylammonium salts, that is, vinylbenzyl n-alkyldimethyl (n= 12 or 18) as a

polymerizable nanofiller of UP resin was also investigated by Fu and Qutubuddin

[73]. After curing, an intercalated and partially exfoliated nanocomposites

were obtained using MMT-containing dodecyl and octadecyl (ODA) chain,

respectively. Both nanocomposites showed an increase in the dynamic storage

modulus as compared with the pristine UP. In the study of Sen and coworkers,

MMT clay was rendered organophilic with a quaternary salt of coco amine

having a styryl part as the reactive group for crosslinking reaction [74]. The

obtained Coco-MMT compound was used as the nanofiller of UP resin resulting

in the formation of partially exfoliated nanocomposite. The highest thermal

stability and the best dynamic mechanical performance was achieved at 1wt%

Coco-MMT nanofiller loading.


In several studies hydroxyfunctional, that is, methyl-tallow-bis (2-hydroxyethyl)

quaternary ammonium (MTHEA) cations were used forMMTorganophilization.

Bharadwaj and coworkers investigated the relationships between morphology

and properties of UP resin nanocomposites containing 1–10wt% MTHEA-

intercalated clay prepared by mechanical mixing followed by ultrasonication

[75]. TEM micrographs show the presence of fully exfoliated clay sheets as well

as intercalated aggregates in the polyester matrix. The crosslinked polyester clay

nanocomposites were optically clear up to 10wt% clay content and exhibited a

reduced oxygen permeability in correlation with an increase in exfoliation degree.

However, adverse effect of the clay exfoliation on the mechanical properties

of nanocomposites determined by tensile tests was observed in this case. The

authors explained that this is partly because of the reduced crosslinking degree of

the polyester. Inceoglu and Yilmazer reported a slight increase of tensile strength

up to 5wt% content of MTHEA-intercalatedMMT and a decrease of the strength

at higher clay loadings [64]. They used a low viscosity UP resin containing 39wt%

styrene mixed with nanoclay at the temperature of 50 ∘C. XRD analysis of the

cured compositions showed the formation of intercalated nanocomposites up

to 3wt% loading of organoclay (with increase in the basal spacing of MMT to

45Å). An increase in glass transition temperature was observed for the cured

nanocomposites (from 72 ∘C in the unfilled polyester to 86 ∘C in the composite

with 10% organoclay). The use of ultrasonic mixing after the mechanical one

had a positive effect on the mechanical properties of UP nanocomposites.

MTHEA-intercalated MMT was also used as a nanoadditive to UP binder in the

fabrication of composites containing granules of the ground polyester scraps,

which effected an improvement in the compression strength, hardness, and

stiffness of the resulting recyclate-filled composites [76]. Torre and colleagues

employed long-chain alkylaryl quaternary ammonium compound (trade name

Hyamine 1622) for the compatibilization of MMT clay with UP resin. Using

simultaneous mechanical and ultrasound mixing they obtained nanocomposites

with increased flexural modulus and strength [77].

Pavlacky and Webster investigated the influence of two quaternary

alkylammonium compounds used in MMT modification: MTHEA and

cetyltrimethylammonium bromide (CTAB) on the properties of the cured

polyester coatings [78]. A plasticizing effect on the coatings (decrease in the

crosslink density and mechanical performance) was observed at high MTHEA

concentrations (5–10wt%) – similarly as in the case of the corresponding

UP/MTHEA–clay nanocomposites. In contrary, the coatings containing CTAB

showed greater storage moduli and crosslink density than the nonmodified

polyester coating. This lack of plasticization may be attributed to weaker

interaction of less polar modifier with the polyester matrix. Nevertheless, the

UP/CTAB-clay nanocomposites showed decrease in the storage moduli and

crosslink density with increased clay loading. Also, the distinct differences

in nanocomposite properties depending on the preparation method (in-situ

polymerization or sonication) were observed.


In most of the studies, UP-nanoclay systems were obtained using MMT

modified by organic ammonium salts. A different approach for increasing the

compatibility ofMMTwith the resin was employed by Kornmann and coworkers.

They used MMT treated with silane-coupling agents containing unsaturated

functions (vinylbenzylamine and methacrylate) [79]. The low viscosity UP resin

(styrene content 42wt%) was stirred with Co accelerator and silane-modified

MMT for 4 h at 60 ∘C, then cured with peroxy initiator for 3 h at room tem-

perature and postcured for 3 h at 70 ∘C. XRD and TEM studies indicated the

formation of partially delaminated nanocomposites. At a 1.5 vol% MMT content

the fracture energy of the nanocomposite was 138 Jm−2 as compared with

70 Jm−2 for the pure UP. Tensile tests showed 32% increase in Young’s modulus

at 5 vol% content of MMT in cured UP. Another study using silane-modified

MMT was reported by Sen, who compared the composites produced from UP

resin and two MMT clays: modified with cetyl trimethyl ammonium salt (CTA)

and trimethoxy vinyl silane (TMVS) [80]. XRD studies showed that the use

of vinylsilane alone did not result in the expansion of MMT and formation of

nanocomposite; however, twofold modification of MMT with TMVS and CTA

allows to obtain exfoliated nanocomposites having better thermal and dynamic

mechanical properties when compared with pure UP or polyesters filled with

only silanized or ion-exchanged MMT.

4.4.1.3 Nanocomposites with MMT Introduced during the Synthesis of Pre-polymer

In the studies described above, clay nanofiller was usually mixed with the

previously prepared polyester alkyd or its styrene solution (UP resin). Another

approach involves an introduction of nanoclay in situ during the synthesis of

UP prepolymer. Webster and coworkers produced a dispersion of organically

modified clays in glycol monomer and conducted the melt polyesterification

with acidic monomers, obtaining in this way in situ synthesized UP/MMT

nanocomposites ([81], Figure 4.5). For comparison, a conventional mixing

and sonication method was employed to prepare the nanocomposites using

three different organoclays and various loading levels. The in situ dispersion

route led to better clay nanodispersion as verified by TEM and it produced

nanocomposites with lower water vapor transmission and permeability when

compared to the sonication process. The mechanical and thermal properties

were improved for lowMMT loadings (1–2wt%); however, they were diminished

at higher clay levels. Katoch and Kundu have also reported the preparation

of nanocomposites from UP based on waste Polyethylene terephthalate (PET)

glycolysis products and MMT clay through in situ polymerization [82]. By intro-

ducing the organoclay simultaneously with the monomers, nanocomposites with

mixed intercalated–exfoliated morphologies (Figure 4.6) and markedly improved

mechanical and thermal properties were obtained. Kedzierski and Penczek

synthesized halogen-containing UP-MMT nanocomposites via copolyaddition of

epichlorohydrin with maleic and phthalic anhydrides activated by propylene gly-

col and carried out in the presence of nanoclay [83]. Montmorillonites containing

four types of cations were used: sodium (MMT-Na), dimethyl dihydrogenated


T N

HO OH

+

HO-R-OH

HO-R-OH HO-R-OH

High shear

Organically modifiedmontmorillonite clay

Hydroxyfunctionalliquid monomer

Polyestersynthesis

Crosslinking viaUV-light withreactive diluent

Unsaturatedpolyester withdispersed clays

+

T N

HO OH

+

Figure 4.5 Representation of in situ technique leading to exfoliated nanoclay platelets

(T= tallow). Reproduced from Ref. [81] Wiley.

tallow ammonium (DMDTA), methyl-tallow-bis (2-hydroxyethyl) quaternary

ammonium (MTHEA), and protonated aminododecanoic acid (ADA). The

reaction conducted in the presence of organoclays with quaternary nitrogen

compounds DMDTA andMTHEA proceeded significantly faster than in the case

of neat polyester, because of a catalytic effect of quaternary ammonium ions on

the epoxide-anhydride addition. Otherwise, no acceleration of polyester forma-

tion took place when sodium MMT as well as MMT-ADA containing primary

ammonium ions were used. The clay-modified prepolymers were dissolved in

styrene and cured. An increase in hardness, heat deflection temperature, and

flame retardance was observed for the as-obtained UP nanocomposites when

compared with nonmodified polyester. The most significant changes in the

properties were observed for nanocomposites obtained from the clay intercalated

with MTHEA-containing hydroxyl groups capable of forming covalent bonds

between nanofiller and UP. Kim and colleagues used MTHEA-intercalated MMT

in the preparation of SP resin nanocomposites performed in situ during the

polyesterification process [84]. The organoclay was predispersed in cyclohex-

anone with high-speed homogenizer. The synthesized SP/clay nanocomposites

were subsequently crosslinked with hexamethoxymetylmelamine to form coat-

ings.The cured nanocomposites exhibited a decrease in water uptake andmarked

improvement in anticorrosion resistance determined by salt spray test, which is

advantageous for application as primers in automotive coatings.

4.4.1.4 Various Properties and Multiphase Nanocomposites

Al-Khanbashi and coworkers investigated the effect of organoclay nanoparti-

cles on the polymerization shrinkage of isophthalic polyester resin [85]. The


500 nm400 nm

100 nm

(a)

(c)

(b)

Figure 4.6 TEM images of UP nanocomposite containing 2wt% clay: (a) low magnifica-

tion and (b) intercalated and exfoliated sheets at high magnification of the aggregate region

shown in the exfoliated sheets from parts a and c. Reproduced from Ref. [82] Wiley.

incorporation of 5wt% 12-aminolauric acid intercalated – MMT into the resin

resulted in the reduction of shrinkage from 8.7 to 2.6%, accompanied by an

increased toughness and slightly decreased modulus of elasticity. In several other

papers, the reduction of volume shrinkage and improvement of mechanical

properties in UP resin/low profile additive systems were reported [86–88]. The

tribological properties of UP/clay nanocomposites were examined by Balasub-

ramanian et al. A specific wear rate of the cured polyester was decreased up to

85% by the addition of 1wt% dodecylamine-modified clay [89]. The influence of

the organically modified clays on the storage stability (shelf life) of UP-clay com-

positions was investigated by Oleksy and coworkers [90]. In the study of Webster

et al. the organomodified clay was incorporated into the formulation of UP resins

subsequently photocopolymerized with vinylether-based reactive diluents to cre-

ate UV-curable nanocomposite coatings systems [91] UP-MMT nanocomposites

cured under UV radiation were also investigated by Kim and coworkers [92].

Rosinska and colleagues studied the effect of MMT nanofillers on the properties

of thermosetting and UV-curable polyesters for powder coatings [93].


Fire retardant composites are the field where the use of nanoclays seems to be

promising from the viewpoint of industrial application. The effect of incorpo-

rating nanoclays on the flammability properties of UP resins were investigated

by Nazare and coworkers [94]. Extensive research has been conducted on the

processing and upscaling of Fire retardant (FR) nanofilled thermosetting UP sys-

tems [95]. Alkylphosphonium-modified clays have been shown to give superior

fire-retarding properties (lower peak heat release values and better results of UL

94 tests) than commercial alkylammonium modified clays. The combination of

nanoclays with other halogen-free flame-retardants generates a synergistic effect.

Also, a new in situ method of producing the organomodified clay instead of tra-

ditional ion exchange process was developed. It was found that high shear rotor-

stator mixing, at high speed, offers the best route to a scalable method for achiev-

ing adequate dispersion of the organoclay. A preliminary assessment of the health

and environmental impacts derived from the manipulation of nanoclay particles

was performed [96].

One of the research areas of practical importance is the application of UP–clay

nanocomposites in combination with traditional reinforcements. Nanoclays were

used as additives in the production of glass fiber-reinforced polyester compos-

ites prepared by hand layup process, resulting in the improved tensile and flexural

strength as well as barrier properties [97, 98] Another three-phase composite con-

taining UP resin, naturally woven coconut fiber mat, and organically modified

MMT clay was investigated for its dielectric properties by Rajini et al. [99]. Hand

layup technique was also applied to produce hybrid nanocomposite laminates

with improved damping properties fromVE resin, short fiber chopped strandmat,

and alkyl ammonium modified MMT clay [100].

Low and colleagues used MTHEA-intercalated MMT as a nanofiller, which

resulted in increased strength and decreased water absorption of vinyl-ester

“eco-composites” based on recycled cellulose fibers [101]. In the study of Karak

and coworkers, organically modified MMT was used as nanofiller to improve

thermostability and mechanical properties of the vegetable oil-based highly

branched polyester containing silver. The resulting nanocomposites show the

potential to be applied as antibacterial surface coating materials [102]. Hos-

sain et al. reported the preparation of jute fabric reinforced polyester – MMT

nanocomposites by VARTM process. The maximum flexural, interlaminar

shear strength, and compression properties were observed for 1wt% nanoclay

loading [103]. Nanocomposites of MMT with UP resins were also investigated as

components of mixed resin systems including polymer concrete [104], bio-based

epoxidized methyl soyate [105] polyurethanes [106, 107], and epoxy resins

[108, 109]. Wood-Adams and coworkers prepared polystyrene and methyl

methacrylate – styrene copolymer by in situ polymerization in the presence

of DMDTA-modified MMT and used the resulting clay-containing polymers

as thermoplastic additives for UP resin [110]. Fracture tests showed that a

combination of nanoclay and thermoplastic component resulted in a synergistic

improvement of the fracture toughness of the nanocomposite while stiffness was

maintained at the level of the unmodified polyester.


4.4.1.5 Vinyl Ester–Clay Nanocomposites

The number of studies on VE resins reported in the literature is lower than

for UPs prepared by polycondensation. However, VE-based composites are

increasingly being used for more demanding applications and the research works

in this field are expanding also in the area of nanocomposites. Raghavan et al.

used two organic ammonium salts: nonreactive undecyl (C11) and reactive

ω-undecylenyl (RC11) to intercalate MMT, subsequently used as nanofiller of

VE resin [111]. The organoclay and VE were premixed with styrene to obtain

a low-viscosity suspension and facilitate the transport of the resin molecules

into the MMT gallery. Using the RC11-intercalated clay and high-intensity

ultrasonic mixing produced VE nanocomposite with the highest degree of

MMT exfoliation. Four types of organically modified MMT-containing ODA,

bis(2-hydroxyethyl)lauryl (BHL), diethyl[2-(methacryloyloxyl)ethyl] (DEM),

and bis(2-hydroxyethyl)lauryl(vinylbenzyl) (BHLV) ammonium cations were

used by Someya and Shibata to produce VE resin nanocomposites [112]. XRD

and TEM studies revealed that exfoliation takes place for the ODA-MMT and

BHL-MMT nanofillers, while neither intercalation nor exfoliation was observed

for DEM-MMT and nonmodified MMT clay additives. The exfoliated nanocom-

posites exhibited an increase in flexural modulus accompanied by a decrease

in flexural strength. Ma and coworkers produced thin polymer composite

bipolar plates containing VE resin, graphite powder, and MMT intercalated with

poly(oxypropylene)-backboned diamine intercalating agents using bulk molding

compound process [113]. The incorporation of 1–4wt% organoclay nanofiller

resulted in a significant increase in flexural and impact strength with only a slight

decrease in electrical conductivity of the graphite plates.

4.4.2

Layered Double Hydroxides

Although cationic clays are the most often investigated layered nanofillers, their

anionic analogs, that is, layered double hydroxides (LDH) capable of anion

exchange, interlayer expansion, and delamination, are also a subject of increasing

interest. A typical example of this class is hydrotalcite (HT) with general formula

of [MeI2+

1−x MeII3+

x(OH)2] [Ay−

x/y (nH2O)] where MeI andMeII denote divalent

and trivalent metal cations, respectively. LDH is not abundant in nature but can

be easily synthesized with tailored chemical composition and may contain a vari-

ety of organic anions, which impart hydrophobicity and good compatibility with

organic matrices. Pereira et al. used LDH intercalated with adipate (A-LDH) and

2-methyl-2-propene-1-sulfonate anions (S-LDH) as nanofillers in UP resin [114].

Powder XRD analysis indicated a good dispersion of LDH layers in the polyester

matrix, with intercalated and partially exfoliated structures at 1 and 5wt% clay

loadings. The nanocomposites showed a lower flexural strength than the pure

polymer, probably because of the weak interaction between the LDH platelets

and the polyester matrix; only for the nanocomposite containing 1wt% of A-LDH

slight increase in flexural modulus was observed. Cone calorimetry studies


indicated a significant reduction in the polyester flammability, by 46 and 32%,

incorporating 1wt% of A-LDH and 5wt% S-LDH, respectively. Kedzierski and

colleagues introduced LDHwith various organic anions in situ during the synthe-

sis of SP resin for powder coatings [115]. The aim was to utilize anionic clay both

as nanofiller and basic catalyst of polytransesterification of dimethyl terephthalate

with alkylene glycols. Using Zn, Al HT intercalated with aminolauric acid, and a

decrease in the average reaction temperature (to 159 ∘C) or a threefold increase

in the reaction rate at 185 ∘C was achieved when compared to the process carried

out in the presence of conventional butylstannoic acid catalyst. The obtained

polyester resins were transparent, which, along with X-ray diffraction data,

indicated the delamination of HT layers and formation of nanocomposite.

4.4.3

Graphene-Based Nanofillers

Among carbon nanofillers, graphene has recently attracted particular attention

because of its unique electronic and mechanical properties. The term graphene

originally referred to single-atom-thick sheet of hexagonally arranged carbon

atoms, and can broaden its meaning to a wide range of sheet-like carbon forms

with different thickness (layer number), lateral dimensions, and in-plane shape

[116]. For the synthesis of polymer nanocomposites, often few-layer graphene

particles prepared via chemical route, by oxidative intercalation of graphite

followed by exfoliation of the resulting graphite oxide, are used. This method was

employed by Dolui and coworkers who prepared graphite oxide using H2SO4 as

intercalating and KMnO4 as oxidizing agent for graphite (modified Hummers

method). The nanoplatelets of graphene oxide (GO) were then obtained by

ultrasonication of GO suspension in Tetrahydrofuran (THF) and used to the

modification of UP resin [117]. Around 76% improvement of tensile strength and

41% increase of Young’s modulus of the cured polyester were achieved at 3wt%

loading of GO (Figure 4.7). Also, a noticeable improvement in thermal stability in

comparison to neat polyester was reported. In another study by the same group

UP resin was modified with a reduced form of GO nanoparticles (rGO) [118].

The tensile strength and Young’s modulus of the obtained nanocomposites were

increased by 123 and 87%, respectively, at rGO loading of 3wt%. Moreover, an

increased electrical conductivity (3.7× 10−4 S cm−1) of the composite film was

observed when compared to that of neat resin. Polyester/rGO nanocomposites

also displayed good antimicrobial activity against a number of bacteria.

A novel approach for the synthesis of graphene-based UP nanocomposites has

been recently reported by Liu et al. [119]. At first, dispersion of GO in ethylene

glycol was achieved using a solvent-exchange method in a reactor equipped

with a fractionating device. The obtained dispersion was applied to fabricate

UP nanocomposites via in situ melt polycondensation in the same reactor. The

reaction was accompanied by a thermal reduction of GO to the functionalized

graphene sheets (FGSs) Even at extremely low content of 0.08wt% FGS an

increase in the tensile strength and modulus of the cured UP matrix (by 53.6


15 kV × 2000 10 μm 0000 15 36 SEI 15 kV × 2000 10 μm 0000 14 36 SEI

× 1000 10 μm 0000 15 34 SEI

(a) (b)

(c)

Figure 4.7 SEM image of (a) polyester resin, (b) polyester/GO composite (3wt%), and (c)

cross section of the composite. Reproduced from Ref. [117] Wiley.

and 48.4%, respectively) compared with the unmodified polyester was observed.

Swain used as nanofiller for UP resin commercial graphene nanoplatelets

(GNPs) with average thickness below 10 nm, surface area 100m2 g−1, and mean

particle diameter 15 μm [120]. The tensile and flexural strength of the produced

nanocomposites were increased by 52 and 92%, respectively, at a GNP concentra-

tion of 0.05% as compared with neat resin. Also, improvements in the thermal and

electrical properties of GNP-modified polyester were observed. It was found that

agglomeration of GNP particles occurred already above 0.075% loading, along

with a decrease of the composite strength. De Bellis and coworkers prepared

nanocomposites of vinyl-ester resin using two kinds of carbon nanofillers:

multiwalled CNT and GNP synthesized via thermal exfoliation of commercial

expandable graphite–sulfuric acid intercalation compound and subsequent

ultrasonication of the exfoliated product suspension. The GNP-modified VE

resins weremuchmore easily processable than those filled with CNTs and showed

good electromagnetic properties after the cure. An electrical conductivity of

4 Sm−1 was achieved at the frequency 18GHz, which is of specific interest for

applications in radar absorbing materials [121].

4.5 Conclusions 109

4.5

Conclusions

Research in the field of nanocomposites based on polyester resins was initiated at

the end of the twentieth century, and the number of publications on this subject

continues to grow. Most of them concern the use of relatively inexpensive clay

nanofillers; however, over the past several years much attention has been given to

carbon-based and metallic nanofillers. In many cases significant improvements

of the mechanical and thermal properties of cured resins were achieved; however,

they were dependent on the proper selection of nanofiller type, content, and

surface treatment as well as the method of nanocomposite preparation. Relatively

high loadings (up to 10wt%) of spherical 0-D nanoparticles such as silica and

metal oxides were succesfully dispersed in polyester resin and the optimum

properties were usually obtained at the nanofiller concentration of several

percent. On the other hand, CNTs with high aspect ratio enable good electrical

properties of the resin with percolation threshold as small as 0.1wt%. Similarly,

graphene-based nanofillers were reported to effectively increase the mechanical

properties of polyester at very small loadings and undergo agglomeration at

higher concentrations.

Various synthetic approaches were used to prepare polyester resin nanocom-

posites, including an introduction of nanoparticles during the synthesis of UP

prepolymer (alkyd), mixing with previously prepared alkyd or its solution in

crosslinking monomer (UP resin) as well as solvent-assisted techniques. A range

of dispersion methods was investigated: simple mechanical mixing, ultrasoni-

cation, HSM along with conventional resin processing techniques like three roll

milling. Surface modification of nanoparticles was usually needed for its effective

dispersion in the resin. For clay nanofillers it was usually accomplished by an

intercalation of organic cations into the interlayer space or silane treatment.

The latter method was often used for compatibilization of other nanoparticles.

The presence of unsaturated groups, capable of undergoing copolymerization

with UP resin, on the nanofiller surface was found to be advantageous for some

polyester-nanofiller systems.

In recent years the attention has been moved from two-phase nanocomposites

to the formulations with nanofiller as additional ingredient that can help optimize

the resin properties. This approach gave promising results, for example, in fiber-

reinforced composites or fire retardant polyester resins where synergistic action

of nanoparticles and conventional FR additives was observed. Further develop-

ment of polyester resin nanocomposites depends on the accessibility of low cost

functional nanofillers, including very promising graphene-based particles, bet-

ter understanding of formulation–structure–property relationships, as well as

ensuring the safety of the use of nanoparticles.


Abbreviations

BHL Bis(2-hydroxyethyl) lauryl

BHLV Bis(2-hydroxyethyl) laurylvinylbenzyl

CNF Carbon nanofibers

CNT Carbon nanotubes (MWmultiwalled, SW single-walled)

DEM Diethyl [2-(methacryloyloxyl)ethyl]

DMDTA Dimethyl dihydrogenated tallow ammonium

FGS Functionalized graphene sheets

GNP Graphene nanoplatelets

GO Graphene oxide

HT Hydrotalcite

LDH Layered double hydroxides

MAOPTMS Methacryloxypropyl trimethoxysilane

MMT Montmorillonite

MTHEA Methyl tallow bis (2-hydroxyethyl) ammonium

PEO Poly(ethylene oxide)

POSS Polyhedral oligomeric silsesquioxanes

PPO Poly(propylene oxide)

RTM Resin transfer molding

SEM Scanning electron microscopy

SP Saturated polyester

TEM Transmission electron microscopy

UP Unsaturated polyester

VARTM Vaccuum assisted resin transfer molding

VE Vinyl ester

XPS X-ray photoelectron spectroscopy

XRD X-ray diffractometry

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115

5

Synthesis Fabrication and Characterization of

Ag/CNT-Polymer Nanocomposites

Vijaya K. Rangari and Sanchit Dey

5.1

Introduction

Nobel metal nanoparticles (NPs) such as Ag, Au, and Pt are playing an impor-

tant role in the modern material systems and also have emerged as a new class

of compounds that are particularly interesting for materials science due to their

unique electronic, optical, biocompatibility, and catalytic properties. Importantly,

because of the properties differ from those of the bulk materials depend on the

size and shape of the nanoparticles [1–3]. These nanoparticles are extensively

used in various applications related to chemistry, physics, material sciences, poly-

mer science, biology, and nanomedicine [4–6]. Among the nanoparticles, silver

nanoparticles have been studied extensively due to their unique physical, chemi-

cal, cost effective, and biological properties compared to their counterparts gold

and platinum [4–7]. Silver nanoparticles are particularly interested in the poly-

mer composite industry because of their low cost, unique electrical, and thermal

conductivity properties [8–13].

Carbon nanotubes (CNTs) are excellent multifunctional materials in terms of

mechanical strength, thermal, and electrical conductivities [14, 15]. These mul-

tifunctional properties, as well as the small size of the structures, make CNTs

ideal building blocks in developing polymer nanocomposites. CNTs represent a

new type of systems that are at the same time single molecules and macroscopic

materials [16]. The novel and unique properties of CNTs, such as ultra-high elec-

trical conductivity and ultra-high mechanical strength, result directly from the

macroscopic understanding of molecular carbon’s unique properties such as bal-

listic transport and exceedingly high mechanical strength. It is important to view

the development in this field as an outcome of highly disciplined collaborations

that produced a suite of novel synthesis and measurement advances [17].

Nanotechnology is a fast-growing research area, involving synthesis, charac-

terization, and device fabrication using nanoscale materials for multifunctional

applications. Various nanomaterials play a number of important roles in modern

science and technology to develop unique materials. Electrically and thermally

conductive metal and metal alloy nanoparticles in polymers are of particular


116 5 Synthesis Fabrication and Characterization of Ag/CNT-Polymer Nanocomposites

importance due to their broad range of potential applications. It is expected

that the combination of electrical, thermal, and mechanical properties in

one nanocomposite would enable the engineering of unique multifunctional

nanoscale devices. CNTs are excellent multifunctional materials in terms of

mechanical robustness, thermal, and electrical conductivities. The multifunc-

tional nanocomposites are also expected to find application in the exploration

systems mission in protecting sensitive optical, electronic, thermal, and acoustic

components from environmental hazards including dust, radiation, thermal tran-

sients, atomic oxygen, and spacecraft charging. It is expected that nanoparticle

systems will also provide a high performance-to-weight radiation shield that

can be used as a layer within human habitations and space protective apparel.

Recently space researchers identified a need for new high performance-to-weight

materials capable of protecting critical components from the space environment,

mitigating the threat of uncontrolled electrostatic discharge, and reducing vul-

nerability to radiation or thermally induced damage. Recent advances in metallic

nanoparticle–polymer composites, metal coated CNT/polymer composites have

shown promise of meeting these multifunctional design goals. There is a great

necessity for the development of these multifunctional nanocomposites for var-

ious applications. Also there are some difficulties and challenges to overcome in

their fabrication such as new cost effective synthesis technique ofmetals onCNTs,

and good dispersion of these fillers in the polymer matrix system. The exciting

electronic and mechanical properties of CNTs, have generated broad and inter-

disciplinary attention. In recent years, many efforts have led to the development

of versatile chemical modificationmethodologies, targeting CNT derivatives with

evenmore attractive features. To this end a wide range of derivatives has been pre-

pared and fully characterized that exhibit promising applications in energy con-

version/fuel storage, catalysis, nanotechnology, electronic nanodevices [18, 19].

The combination of the two special class of materials (CNTs and NPs) may

lead to a successful integration of the properties of the two components of the

new hybrid materials that present important features for polymer composites,

catalysis, and nanotechnology [20, 21]. The CNT surface serves as a template

where NPs are absorbed or, when bearing functional groups, CNTs may be linked

through organic fragments, to metal or semiconductor NPs either alone or stabi-

lized by a protecting monolayer. The first report on decorating CNTs with metal

clusters dates back to 1994. Ajayan and coworkers [22] described the use of single

walled carbon nanotubes (SWNTs) as a support material for dispersing ruthe-

nium NPs that act as catalysts in heterogeneous catalysis. In particular, ruthe-

nium 2,5-pentanedionate was spread onto the walls of SWNTs and subsequently

reduced under a hydrogen stream.The Ru NPs thus obtained were well dispersed

on the nanotube surface as corroborated by detailed TEM (transmission electron

microscopy) analysis. The final material contains 0.2% w/w of Ru. Catalytic assays

that include liquid-phase hydrogenation of cinnamaldehyde revealed a particu-

larly high selectivity for cinnamyl alcohol (up to 92%) with an 80% conversion of

cinnamaldehyde. In contrast, under similar conditions, RuNPs of similar size sup-

ported on Al2O3 catalyze the formation of cinnamyl alcohol with a selectivity of


20–30% only. Following this promising work, the hybrid composites were devel-

oped with either metal, semiconductor, or metal alloy NPs.

We have also recently reported on the multifunctional application of Ag/CNTs

hybrid nanoparticles [23]. Neat Nylon-6, commercial Ag, pristine CNTs, and

Ag/CNT-infused Nylon-6 polymer composites (PNC) were fabricated using a

single screw melt extruder. XRD (X-ray diffraction) and TEM studies reveal

that Ag nanoparticles were uniformly coated on CNT surfaces and were non-

covalently attached through van der Waals forces. The improvement in ultimate

tensile strength and elastic modulus is attributed to the alignment of the Ag/CNT

nanoparticles along the direction of extrusion. The increase in thermal stability

and crystallinity of Ag/CNT-infused Nylon-6 PNC is correlated with the better

cross-linking between the nanoparticles and the polymer matrix. The in situ 1%

Ag/CNT Nylon-6 nanocomposite fibers were demonstrated to have excellent

and promising antimicrobial activity as compared to the commercially available

Ag nanoparticles, In other work we have reported [24] that the diamond-coated

(DN) CNTs were infused in Nylon-6 polymer fibers through an extrusion process

to alignment/disperse the nanoparticles to improve the mechanical properties.

The tensile properties of these fibers show that the DN-coated CNT-infused

Nylon-6 fibers can take 51%more load than the neat Nylon-6 fibers. Nylon-6 with

DN-coated CNTs also shows improvements in mechanical properties compared

to Nylon-6 infused with just CNTs or just DNs.

The decorated CNTs with metal nanoparticles are useful in field-emission

displays, nanoelectronic devices, as well as novel catalysts and polymer reinforce-

ment [21]. In the past metal nanoparticles synthesized by sonochemical methods

[25], microwave irradiation [26], photochemical method [27], hydrothermal

and solvothermal methods [28], electrochemical method [29], sol–gel methods

[30], chemical reduction, and depositions route [31]. There are several recent

reports, which showed that use of microwaves in synthesis of metal nanoparticles

increases the kinetics of metal formation [32]. He et al. reported preparation of

polygonal Ag nanoplates usingmicrowave irradiation of AgNO3 in the presence of

polyvinylpyrrolidone (PVP) without any other reducing agent. Four types solvent

were used by them, dimethyl formamide (DMF), N-methylpyrrolidone, pyridine,

and ethanol [33].He et al., also synthesizedwell-definedAgdendrites by reduction

of AgNO3 in DMF containing PVP under microwave irradiation [34]. Yamamoto

et al., reported the preparation of Ag nanoplates by microwave promoted reduc-

tion of AgNO3 in aqueous solution involving PVP [35]. Liu et al., synthesized Ag

nanorods in an aqueous solution in the presence of Au seeds under microwave

radiation. Sodium citrate was used as reducing agent of silver ions [36]. Masaharu

et al., reported rapid synthesis of silver nanostructure by microwave-polyol

process with the assistance of Pt seeds and polyvinylpyrrolidone [37].

Recently microwave radiation is widely used in many fields. In case of materials

processing, microwave irradiation directly couples the electromagnetic energy

(300MHz to 300GHz) with the material through molecular interactions and

enables energy dissipation through the release of heat [38]. Microwave heating

offers several advantages compared to conventional heating processes such as


the use of a remote source, the relative speed of the process, and the volume

and material selectivity [39]. In this research our objective is to use microwave

radiation to produce the metal nanoparticles on the outer surface of CNTs.These

coated CNTs are further used as fillers in the fabrication of multifunctional

polymer nanocomposites for various cutting edge applications. This type of

multifunctional nanocomposites materials which combine both electrical and

mechanical properties in one entity, have in particular those with potential

applications in nanoelectronics such as flexible conductors/resistors, flexible

field-emission devices, electromagnetic interference (EMI) shielding, polymer

electrolyte fuel cells, and aerospace.

5.2

Experimental Procedure

Multiwall CNTs, with outside diameter 10–20 nm, inside diameter 5–10 nm,

length 10–30 μm was used in these experiments supplied by Nanostructure

and Amorphous Materials, Inc. Silver acetate was used as a metal precursor,

DMF used as reducing agent and solvent and Polyvinyl alcohol (PVA) used as

a surfactant. In a typical reaction 50mg of surfactant was dissolved in 100mL

of DMF and 100mg of CNTs are dispersed in the DMF by using a magnetic

stirrer in a round bottom flask. Two hundred and fifty milligrams of precursor

salt (silver acetate) was dissolved in the same mixture. The glass flask is then

placed in the center of a microwave oven (SHARP 1000V/R21HT) and attached

to a condenser. The irradiation is continued for 5min under microwave power

of 60W then the product kept undisturbed for 5 h to settle down the particles.

After 5 h the product is separated from the liquid by centrifugation and followed

by repeated washing with water and absolute ethanol several times and vacuum

dried at room temperature overnight.

The silver nanoparticles were synthesized using the same reaction as above

without CNTs. The final product was washed with water several times to ensure

that the unreacted silver acetate is completely removed and finally washed with

ethanol and dried overnight. Synthesized products are characterized by powder

XRD (X-ray diffraction). The XRD patterns were recorded on a Rigaku, D/Max

2200 X-ray diffractometer operated at 40 kV and 30mA with CuKα radiation.

The size and morphology of synthesized particles were determined using a trans-

mission electron microscope (JEOL-2010 Transmission Electron Microscope).

The powdered samples were dispersed in ethanol and subjected to ultrasonic

treatment and dropped on to a conventional carbon coated molybdenum grid

and analyzed. The thermogravimetric analysis (TGA) of the powder product

is carried out using Mettler Toledo thermogravimetric/standard differential

thermal analysis (TGA/SDTA) 851.

The decorated CNTs are infused in the resin system RenInfusionTM 8606 (Ren-

8606) supplied by Freeman Manufacturing and Supply. It is a two-component


low-viscosity epoxy system. Two series of composites were prepared for com-

parison at three different loading percents: CNTs/Reninfusion (series 1) and

Ag/CNTs/Reninfusion (series 2). Precalculated amounts of neat and decorated

CNTs and Reninfusion 8606 resins (part A) were carefully weighed and mixed

together in a beaker. Using THINKY hybrid defaming mixer ARE-250 which

performs a noncontact mixing for the materials to disperse the nanoparticles

uniformly in the resin systems. In this technique the material container is set at

45∘ angle and revolves and rotates (2000 rpm) at the same time. Dual centrifugal

forces were given to the material that keep pressing material to outward and

down along with the slope of the inner wall of the container. After 30min

part B (Ren-8606) was added to the modified resin and mixed using THINKY at

2000 rpm for 20min. The mix ratio of Reninfusion and Ren-8606 was 100 : 35.

The mixing of epoxy and curing agent initially produced highly reactive, volatile

vapor bubbles, which could create voids and detrimentally affect the properties

of the final product. To reduce the chance of voids, the mixture is degasified for

15min. After the bubbles were completely removed, the mixture was transferred

to plastic and Teflon-coated metal rectangular molds and cured for 48 h at room

temperature. Then the material was post cured for 4 h at 121 ∘C+ 4 h at 177 ∘Cin a Lindberg/Blue M laboratory vacuum oven as suggested by the supplier. The

cured material was then cut to the ASTM standard. Finally, test samples were

machined for thermal and mechanical characterization. Flexure and compression

test of the samples are carried out using the Zwick Roell testing machine and

MTS machine (500KN load-cell) respectively. TGA and differential scanning

calorimetry (DSC) of the composite sample carried out using Mettler Toledo

TGA/SDTA 851 and Mettler DSC822 respectively microstructures of neat and

nanocomposites are examined under a scanning electron microscope (JEOL JSM

5800 Scanning Electron Microscope).

5.3

Results and Discussion

5.3.1

XRD analysis

Figure 5.1 shows the powder XRD patterns of (a) Ag nanoparticles, (b) Ag/CNTs

with PVA as surfactant, and (c) CNTs. Figure 5.1c,a indicate that the CNTs andAg

particle are crystalline and all the peaks match very well with the standard CNTs

and silver JCPDS file numbers 41–1487 and 04–0783 respectively.

5.3.2

Transmission Electron Microscopy

Figure 5.2 shows the transmission electron micrograph of (a) as synthesized Ag

nanoparticles, (b) as-received CNTs, (c) Ag/CNTs composite nanoparticles, and


10 20

(c)

(b)

(a)

30 40 50 60 70 80

10 20 30 40 50 60 70 80

10 20 30 40 50 60

2-theta degrees

Re

lative

in

ten

sity

70 80

Figure 5.1 XRD patterns of (a) Ag nanoparticle, (b) Ag/CNTs, and (c) CNTs.

20 nm

10 nm

10 nm

10 nm

4.8 nm

(a) (b)

(c) (d)

Figure 5.2 Transmission electron micrographs of (a) Ag nanoparticles, (b) as-received CNTs,

(c) Ag/CNTs composite nanoparticles, and (d) Ag/CNTs composite nanoparticles at high

resolution.


(d) Ag/CNTs composite nanoparticles at high resolution. Figure 5.2a shows the

fine particles of silver and the particle sizes are 2–5 nm range. These nanoparti-

cles are produced using DMF as reducing agent and polyvinyl alcohol (PVA) as

a surfactant. Figure 5.2b represents the as-received CNTs and these nanoparti-

cles are ∼10–20 nm in diameter, length 10–30 μm. They match very well with

the suppliers data sheet. Figure 5.2c depicts the as synthesized Ag/CNTs hybrid

nanocomposite particles from DMF as reducing agent and PVA as a surfactant.

The particles are almost same size in the range of 2–5 nm and spherical shape.

Figure 5.2d shows the high resolution micrograph of Figure 5.2c and the number

of layers of CNTs are very much visible and also the dispersion of Ag on CNTs.

The initial dispersion of CNT in the DMF solution is also an important factor to

produce a uniform coating ofAg nanoparticles onCNTs.Thebetter the dispersion

of CNTs in initial solution is the better the coating.

The reaction scheme for producing fine andmonodisperse silver particles using

the DMF process involves the following successive reactions: reduction of the

soluble silver acetate by DMF nucleation of metallic silver, and growth of indi-

vidual nuclei in the presence of a protective agent, PVA. Upon addition of the

silver acetate to DMF and under the microwave, the Ag+ species are reduced to

metallic silver. The concentration of metallic silver in solution increases, reaching

the supersaturation conditions and finally the critical concentration to nucleate.

Spontaneous nucleation then takes place very rapidly andmany nuclei are formed

in a short time, lowering the silver concentration below the nucleation and super-

saturation levels into the saturation concentration region. After a short period of

nucleation, the nuclei grow by the deposition of metallic silver until the system

reaches the saturation concentration. At the end of the growth period, all themetal

particles have grown at almost the same rate and the system exhibits a narrow

particle size distribution [18].

5.3.3

TGA Analysis of Nanoparticles

Figure 5.3a,b shows the TGA curve of neat CNTs and Ag/CNTs respectively.

Residue calculation from the two curves shows that the CNTs burned in nitrogen

residue is ∼93% by weight and the Ag/CNTs burned in nitrogen residue is

∼96% by weight. The difference ∼3% by weight is corresponds to the (Ag) metal

content in the Ag/CNTs system. These results are consistent with the theoretical

calculation from the precursor.

5.3.4

Thermal Response of the Polymer Composites

Thermogravimetric analysis measurements were carried out to obtain informa-

tion on the thermal stability of the various nanocomposite systems. These results

clearly show that by the addition of CNTs and Ag/CNTs nanoparticles at various


0

92.4

94.6

96.8

99.0

101.2

190 380 570 760

100

98

−1.52×10−4

−3.20×10−5

−1.60×10−5

1.60×10−5

3.20×10−5

0.00

−1.14×10−4

−7.60×10−5

−3.80×10−5

0.00

Temperature (°C)

0 190 380 570 760

Temperature (°C)

% W

eig

ht

loss

% W

eig

ht

loss

Wt%

/Ce

ntig

rad

e

Wt%

/Ce

ntig

rad

e

(a) (b)

Figure 5.3 TGA curves of (a) As-received CNTs in nitrogen and (b) Ag/CNTs in nitrogen.

percentages to the epoxy resin Ren-8608 the thermal decomposition tempera-

ture is not changed, especially not decreased by addition of nanoparticles. All the

curves are very similar to the neat epoxy system, where a prominent weight loss

at ∼375 ∘C shown in Figure 5.4a,b.The possible reason could be the percentage of

loadings are very negligible amounts to show any significant thermal effects.

Differential scanning calorimetry analysis was used to measure the changes

in heat flow associated with material transition for various weight percentage

of uncoated and coated CNTs. DSC tests were primarily used to determine the

300

25

−0.002

−0.001

0

50

75

100

0

25

−0.001

0

0.001

−0.003

50

75

100

140 250 360 470 580 690 800 30 140 250 360 470 580 690 800

Temperature (°C) Temperature (°C)

We

igh

t p

erc

en

t

We

igh

t .

%

Wt.

pe

rce

nt

°C−1

Wt.

pe

rce

nt

°C−1

⊕ ⊕ ⊕ ⊕

⊕

⊕

⊕

⊕

⊕

⊕

⊕⊕⊕⊕

⊕ ⊕ ⊕ ⊕

⊕ ⊕ ⊕ ⊕ ⊕

⊕ ⊕ ⊕

⊕⊕⊕ ⊕ ⊕

⊕ ⊕ ⊕ ⊕

⊕

⊕⊕

⊕

⊕⊕ ⊕ ⊕ ⊕ ⊕

⊕

⊕

⊕ ⊕ ⊕ ⊕ ⊕ ⊕ ⊕ ⊕

Reninfusion + 0.3% MWCNTs

Reninfusion + 0.1%MWCNTs

Reninfusion + 0.2%MWCNTs

Neat reninfusion 8606

⊕ ⊕

Reninfusion + 0.3% Ag-coated MWCNTs

Reninfusion + 0.2%Ag-coated CNTs

Reninfusion + 0.1%Ag-coated CNTs

Neat reninfusion 8606

⊕ ⊕

(a) (b)

Figure 5.4 The TGA curves of neat and composite system containing (a) CNTs and

(b) Ag/CNTs as fillers.


74

−6.9

−4.6

−2.3

0.0

(a)

(b)

(c)

(a) Neat Ren-8606

(b) Ren-8606/0.1% Ag-CNTs

(c) Ren-8606/0.2% Ag-CNTs

(d) Ren-8606/0.3% Ag-CNTs

(d)

He

at

flo

w (

mW

)

2.3

148 222 296

Temperature (°C)

370

Figure 5.5 The DSC curves of neat and composite systems containing Ag/CNTs.

effect of nanoparticles on the glass transition temperature of the nanocomposite

systems. Typical heat flow versus temperature curves is shown in Figure 5.5 of

all composite systems along with a neat system at three different weight percent.

Table 5.1 represents the summary of the DSC results. The broad endothermic

peak is meant to indicate a depression in the curve as seen in Figure 5.5.

The Tg’s were determined as the inflection points of the heat flow curve [24].

An increasing amount of CNTs and coated CNT results in a shift of the glass

transition temperature. The Tg was shifted from 137 ∘C for the neat resin, to

168 ∘C for samples containing 0.3wt% of Ag-coated CNTs.This gain in Tg can be

explained as a reduction of the mobility of the matrix around the nanotubes by

the interfacial interactions.

Table 5.1 DSC test results with glass transition temperature (Tg).

Sample ID Tg (∘C)

Neat Ren-8606 137.78

Ren-8606+ 0.1% CNT 156.19

Ren-8606+ 0.2% CNT 160.43

Ren-8606+ 0.3% CNT 164.10

Ren-8606+ 0.1% Ag-CNT 164.10

Ren-8606+ 0.2% Ag-CNT 168.05

Ren-8606+ 0.3% Ag-CNT 168.29


5.3.5

Compression Test Results of Polymer Composites

The fabricated specimens of composites are tested for compressive properties.The

load displacement curve was obtained during the test and used for developing

stress–strain relations and calculating compressive modulus and strength. Five

specimens of each type are tested and all the results are consistent. Table 5.2 com-

prise the compression test data of three different type composites and also the

neat system.

Figure 5.6 represents the compressive stress plot for nanocomposites contain-

ing fillers of Ag/CNTs along with neat system. An important feature of these

curves is that these composites can be compressed to about 18% strain without

any loss in strength. The shape of the sample change and we stop the experiment

as the plastic deformation occurred in the specimen. The compressive yield

Table 5.2 The compression test results of neat and nanoparticles infused epoxy.

Sample ID Filler type Filler

content

Stress

(MPa)

Gain in

strength (%)

Modulus

(MPa)

Gain in

modulus (%)

Neat — — 87 1911

CNTs/Ren Neat CNTs 0.1 91 4.6 2000 4.65

0.2 94 8.04 2006 4.97

0.3 95 9.2 2061 7.85

Ag/CNTs Ren Ag-coated CNTs 0.1 92 5.7 2063 7.95

0.2 95 9.2 2063 7.95

0.3 95 9.2 2175 13.81

0.00

34

68

102

136

170

3.6 7.2 10.8 14.4

Strain (%)

Neat Ren-8606Ren-8606 + 0.1% Ag-CNTsRen-8606 + 0.2% Ag-CNTsRen-8606 + 0.3% Ag-CNTs

Str

ess (

MP

a)

18.0 21.6

Figure 5.6 Compressive response plots of composites containing Ag/CNTs.


strength is calculated as 0.2% yield strength. In general the reinforcement of these

nanoparticles increases strength and modulus as compared to the neat matrix.

The maximum improvement is 9% in case of stress and 13% in the case of the

modulus. The addition of nanofiller increases the compressive stiffness of the

matrix material. The neat Reninfusion matrix exhibits a compressive modulus of

1911MPa.The present results for nanoparticle filler composites show an increase

of the flexure modulus with a maximum of 2175MPa at 0.3wt% loading. The

reason can be explained as here the filler particles act as a barrier to the polymer

chain. As increases the loading percentage of the filler content, more the stress

and modulus [25].

5.3.6

Flexure Test Results of Polymer Composites

Flexure tests were carried out to determine the bulk stiffness and strength of each

of the nanocomposites systems. Typical stress, strain behavior of flexure test is

shown in Figure 5.7. Each figure shows four curves corresponding to neat, 0.1,

0.2, and 0.3wt% of different nanoparticles. These results are also summarized in

Table 5.3.

It is observed in Figure 5.7a that the addition of small amounts CNTs increases

the flexure stress moderately. Addition of Ag-coated CNTs increases the strength

(Figure 5.7b). The dispersion of CNTs that restricts the mobility of polymer

chains under loading improved the modulus and strength in small loadings.

The high aspect ratio, high modulus, strength of CNTs, and good interfacial

adhesion between the CNTs and matrix also contributed to the reinforcement.

In case of modulus all systems show enhancement in stiffness except polymer

0.00 0

24

48

72

96

120

30

60

Str

ess (

MP

a)

Str

ess (

MP

a)

90

120

3.6 7.2 10.8 14.4

Neat Ren-8606

Ren-8606 + 0.1% CNT

Ren-8606 + 0.2% CNT

Ren-8606 + 0.3% CNT

Neat Ren-8606

Ren-8606 + 0.1% Ag CNT

Ren-8606 + 0.2% Ag CNT

Ren-8606 + 0.3% Ag CNT

18.0 0 3 6 9 12 15

Strain (%) Strain (%)(a) (b)

Figure 5.7 Flexural stress–strain curves of (a) CNT/Reninfusion and (b) Ag/CNT/Reninfusion.


Table 5.3 Flexure test results of neat and nanoparticle-infused epoxy.

Neat Filler type Filler

content (%)

Strength

(MPa)

Gain/loss

in loss (%)

Modulus

(GPa)

Gain/loss

in modulus

Neat — — 90.43± 5.85 — 2.61± 0.09

CNT/Ren Neat CNTs 0.1 91.73± 17.83 1.44 2.70± 0.126 3.45

0.2 104.67± 16.76 15.74 2.73± 0.055 4.60

0.3 102.67± 6.58 13.53 2.75± 0.065 5.36

Ag-CNT/Ren Ag-coated CNTs 0.1 92.58± 5.19 2.37 2.62± 0.20 0.38

0.2 99.705± 6.85 10.25 2.67± 0.131 2.29

0.3 99.77± 3.69 10.33 2.77± 0.23 6.13

composites. Our understanding in case of enhanced strength of Ag-coated CNTs

is that CNTs with well dispersed spherical nanoparticles of ≈2–5 nm diameter

in a surface can sit comfortably within the polymer chains and coils which have

more or less similar dimensions [19, 20]. These in turn will enhance the reactivity

between the filler particle and the polymer. More reactivity will translate into

increased change in mechanical properties [21]. The neat Reninfusion matrix

exhibits a flexure modulus of 2.61GPa. The composites show the modulus value

of 2.75 and 2.77GPa incase of CNTs and Ag-CNTs.

Themicrostructure characterization of fracture surfaces of polymer composites

provide the information about fracture mechanisms and the influence of particle

modification on the fracture behavior. Figure 5.8a–c shows SEM images of

fracture surfaces of the neat reninfusion resin, a representative nanocomposite

containing 0.2wt% CNTs and nanocomposites containing 0.2% Ag-coated

CNTs. The imaged fracture surfaces were taken from similar areas of the failed

T-CAM 8 kV 500 μm ×40 T-CAM 12 kV 500 μm ×43 T-CAM 12 kV 500 μm ×40

(a) (b) (c)

Figure 5.8 SEM micrographs of (a) Neat Reinfusion 8606, (b) nanocomposites containing

0.2% CNTs, and (c) nanocomposites containing 0.2% Ag/CNTs.

References 127

specimens. The initial crack occurred at the tension edge of both the neat and

nanophased specimens. The general toughening effect of nanoparticles reflects

itself in a significantly larger roughness of the fracture surface of the sample

containing both neat CNTs andmodified CNTs.The increased surface roughness

implies that the path of the crack tip is distorted because of the CNTs, makes the

crack propagation more difficult [22].

5.4

Conclusion

We have described the synthesis of silver (Ag) nanoparticles on multiwalled

CNTs using microwave irradiation and their applications as fillers in thermoset

polymer composites. Ag nanoparticles of uniform size and shape are synthesized

using DMF as reducing agent and as well as a solvent. This method also can be

extended to the other nanoparticle coating of CNTs. XRD method and TEM

are used to characterize both the nanoparticles and decorated CNTs. It is found

that the morphology of the nanoparticles is controlled by the amount and type

of the surfactant used. The microwave-assisted process is found to be faster

than the conventional thermal process. The decorated CNTs are infused in

polymer resin system to produce nanocomposites for various applications. These

nanocomposites are characterized by thermal and mechanical properties and

significant improvements are observed as compared to their neat counterparts.

This method can be used to design the hybrid nanoparticles depends on the type

of polymer composite application.

Acknowledgments

The authors would like to thank the National Science Foundation (NSF) for their

financial support through NSF-CREST, PREM, and RISE grants.

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131

6

Preparation and Characterization of PVDF-Based

Nanocomposites

Derman Vatansever Bayramol, Tahir Shah, Navneet Soin, and Elias Siores

6.1

Synthesis of Poly(vinylidene fluoride) (PVDF)

PVDF is a thermoplastic and semicrystalline fluoropolymer with a crystallinity

of about 50–60%, of which the crystalline form may exist in at least five poly-

morphs; α-phase, β-phase, γ-phase, δ-phase, and lately determined ε-phase [1, 2].Crystallinity has toughness, mechanical strength, resistance, and other properties

of PVDF, which is generally synthesized from 1,1-difluoroethylene (VF2) by the

free-radical polymerization, a monomer commonly synthesized from acetylene

or vinylidene chloride via 1-chloro-41 1,1-difluoroethylene.

Although suspension and emulsion polymerizations are the commonly used

processes for the manufacture of PVDF [3–5] there are other methods for the

synthesis of the polymer on laboratory scale; those are radiation-induced poly-

merization of VF2 in solution or in the gaseous state and glow-discharge poly-

merization. Radiation polymerization results in β-phase formation of PVDF if

the polymerization takes place in polar solvents. If the solvent is nonpolar, then

the polymorphology of synthesized PVDF is in α-phase. Plasma-induced poly-

merization of VDF was also successfully studied [6] unlike microwave-simulated

polymerization which gave thin polymeric PVDF films [7].

6.2

Structure and Piezoelectric Properties of PVDF

PVDF is an attractive thermoplastic polymer which consists of repeated molec-

ular monomer units of (–CH2–CF2–) in a long chain which contains 59.4wt%

of fluorine and 3wt% of hydrogen. Physical and electrical characteristics of PVDF

depend on the molecular weights, molecular weight distributions, the chain con-

figurations, the crystalline form, and the defects of chaining [8–12].The hydrogen

atoms are positively charged and the fluorine atoms are negatively charged in the

polymer [13]. Because of high dielectric dipole moment of the monomer con-

stantly spaced, polymer of PVDF acts as a crystal. However, PVDF is not inherently


132 6 Preparation and Characterization of PVDF-Based Nanocomposites

piezoelectric, and in order to make it so, it is necessary to produce a large polar-

ization within the PVDF crystal. This can be achieved by orientating the melt-

crystallized PVDF material and subjecting it to a high electric field at elevated

temperature [14]. PVDF is not soluable in water, various solvents, oils, and acids.

Its glass transition temperature (Tg) is in the range of −40 to −30 ∘C and its melt-

ing temperature (Tm) is in the range of 155–192 ∘C [15].

Although PVDFwas an attractive polymer formany applications, the piezoelec-

tricity of PVDF, in particular, the semicrystallinemorphology, was discovered [16]

and reported that PVDF could be made piezoelectric by the stretching and poling

process. It has pyroelectric property [17, 18] and exhibits ferroelectric behavior

[1, 19–21]. Man-made piezoelectric polymers have been studied since then by

other researchers [22–25]. Polymeric materials can be produced as large thin

sheets and then can be cut or stamped into nearly any shape. They also exhibit

high mechanical strength and high impact resistance. Although the piezoelectric

charge constant of polymers is lower than that of ceramics, they havemuch higher

piezoelectric voltage constant than that of ceramics, which indicates better sens-

ing characteristics.

Polymers consist of two regions: crystalline and amorphous. The percentage

of the crystalline region in a polymer matrix determines the piezoelectric effect.

However, crystallites are dispersed in amorphous region in semicrystalline poly-

mers as shown in Figure 6.1.

The melting temperature of polymer is dependent on the percentage of crys-

talline region in the polymer, while the amorphous region designates the glass

transition temperature and mechanical properties of the polymer. As it is seen

in Figure 6.1, crystalline structures, and so the molecular dipoles, are locked in

the amorphous region. Broadhurst et al. [13] studied the molecular and mor-

phological structure of PVDF and its pyroelectric and piezoelectric properties.

If a DC voltage is applied across the polymer piezoelectric material, the material

becomes thinner, longer, and wider in proportion to the voltage, conversely the

film generates a proportional voltage when a mechanical stress is applied either

by compression or stretching.The relationship between appliedmechanical stress

and generated voltage can be defined by stress constants.

Aforementioned five types of polymorphs are direct phases in PVDF and are

related to van derWaals radius of fluorine atom and hydrogen atom [27]. There is

Crystallineregion

Amorphousregion

Stretchdirection

Polin

g d

irectio

n

Ele

ctro

des

(a) Melt cast (b) Mechanically oriented (c) Electrically poled

Figure 6.1 Amorphous and crystalline regions in the polymer matrix; from melt cast (a),

during mechanical orientation (b), and electrically poling (c) [26].


only a limited knowledge on the ε-phase of PVDF though other four polymorphs

and their characteristics are given in Table 6.1. Among those, the most common

and thermodynamically stable phase is α-phase, which is also known as form “II”

or “2.” It can easily be formed from melt processing. Form II has a slightly dis-

torted trans-gauche-trans-gauche′ (TGTG′) with a unit cell that is centrosym-

metric because of the anti-parallel packing of the two chains contained in the cell.

Dipolemoments are randomly aligned in the crystalline part of the polymer, which

results in a nonpolar form [2] (Figure 6.2).

The β-phase is also known as form I or “1” that has alltrans conformation

(TTTT) in the polymeric chain. It is a noncentrosymmetric with a conformation

of head-to-head (–CF2–CF2–) and tail-to-tail (–CH2–CH2–); therefore it is

polar and exhibits piezoelectric property. Being predominantly responsible for

the piezoelectric, pyroelectric, and ferroelectric properties [27], the β-phase is

the most important polymorph of PVDF that originates from the orientation

of the strong dipole along the polymeric chain. It has an all transconfiguration

Table 6.1 Polymorphs of PVDF and their characteristics.

Form I Form II Form III Form IV

β-phase α-phase γ-phase δ-phase or form IIpAlltrans (planar

zigzag)

Trans-gauche-trans-

gauche′ (TGTG′)

(T3GT3G′) (TGTG′)

All chains are

oriented parallel to

b-axis

Antiparallel packing

of the two chains

Molecular chains are

packed in parallel

Rotation of every

second chain, all

aligned

Noncentro-

symmetric

Centrosymmetric Noncentro-

symmetric

Noncentro-

symmetric

Dipole moment 2,

1D is parallel to

b-axis

Dipole moment 1,

2D perpendicular to

b-axis and 1, 0D

parallel to b-axis

— Dipole moment is 1,

3D

Polar Nonpolar Polar Polar

C CH2 F

Figure 6.2 Trans-gauche-trans-gauche′ (TGTG′) conformation of PVDF (α- and δ-phases)[28].


(TTTT) and a strong dipole moment normal to the chain direction and all

chains are oriented parallel to b-axis [2]. PVDF mainly exhibits randomly

oriented crystalline form, and this form of PVDF is antipolar and so does not

show piezoelectric property. To form the structure into polar β-phase, PVDF is

subjected to mechanical, thermal, and electrical conditions to create a permanent

polarization. However, the polarization disappears when the material is heated

up to its Curie temperature (Tc), which is 80 ∘C for PVDF and 100 ∘C for its

copolymers [29] (Figure 6.3).

As the β-phase is the most important polymorph of PVDF, it has been widely

studied by researchers. Shuford et al. [30] reported that the piezoelectric con-

stant increased up to 5/1 stretching and then showed a slight decrease for fur-

ther drawn ratios. This finding was then supported by Fourier transform infrared

spectroscopy (FTIR) results of stretched PVDF films which were carried out by

Salimi and Yousefi [31]. Davis et al. [32] studied the effect of the draw ratio and

applied electric field on the phase change characteristic of PVDF. It was found

that the polarization contributed the alignment of dipoles parallel to the molec-

ular chain. Simultaneous stretching and corona poling of PVDF films [33] and

poly(vinylidene fluoridetrifluorethylene), P(VDF-TriFE), films [34] were studied

to investigate the effect of variable parameters, such as stretching, poling tem-

perature, and electric field applied, on piezoelectric effect of the polymers. The

piezoelectric effect of PVDF at high frequencies was studied by Sussner et al. [35]

while Nix and Ward [36] measured the shear piezoelectric coefficients of PVDF.

The third phase is the γ-phase that is also known as form “III” or “3.” The

γ-phase has intermediate polar conformation (TTTGTTTG′) and can be formed

by solution crystallization using Dimethylformamide (DMF), Dimethylacry-

lamide (DMA) and Dimethyl sulfoxide (DMSO) [37] and by melt crystallization

with high temperature and high pressure. It can also be transformed to

β-phase by drawing. The configuration is an intramolecular mix of both α-phaseand β-phase (T3GT3G) [2]. Therefore, the piezoelectric effect is not as good as

β-phase (Figure 6.4).The δ-phase is also known as form “IV” or “IIp.”The δ-phase is produced by the

transformation of nonpolar α-phase by subjecting to a high electric field and so

producing an inversion of dipole moments so they become noncentrosymmetric.

It can also be transformed to β-phase by subjecting it to high electric field [2]. It

can be concluded that at least three polymorphs of PVDF are in polar form, which

C CH2 F

Figure 6.3 All-trans (TTTT) conformation of PVDF (β-phase) [28].


C CH2 F

Figure 6.4 Intermediate polar (TTTGTTTG′ or T3GT3G′) conformation of PVDF (γ- and

ε-phases) [28].

shows that piezoelectric property and all forms of PVDF can be interconverted by

the application of mechanical, electrical, and thermal conditions.

6.2.1

Relationships and Equations

Relationships between applied stimulus and the resultant responses depend on the

piezoelectric properties of thematerial.The size and the shape of the piezoelectric

material and the direction of applied external electrical or mechanical excitation

are important parameters. To identify directions in a piezoelectric element, three

axes, termed 1, 2, and 3, which are analogs to X, Y, and Z of the classical three-

dimensional orthogonal set of axes, are used.

Piezoelectric coefficients with double subscripts represent electrical and

mechanical directions. The first subscript indicates the direction of the electrical

field associated with the voltage applied or the charge produced. The second

subscript indicates the direction of the mechanical stress or strain (Figure 6.5).

As mentioned earlier the piezoelectric material can be generalized into two

operating modes; the first one is the stock configuration that operates in the 33

mode and the second one is the bender configuration that operates in the 31mode.

For both cases, it is assumed that the poling direction is always in the “3” direc-

tion. In the 33 mode, both the voltage and stress act in the 3 direction, which

means the material is strained in the poling or “3” direction and the electric volt-

age is recovered in the “3” direction. In the 31 mode, the material is poled in the

“3” direction and the mechanical stress acts in the “1” direction, which means

the materials is strained in the perpendicular direction to the poling direction

[38, 39].

The direct and converse piezoelectric effect of the material can be expressed by

two linearized constitutive equations that contain both electrical and mechanical

variables;

{D} = [𝐞]𝑇 {S} +[𝐚𝐒]{𝐸} (6.1)

{T} = [𝐜𝐄] {S} − [𝐞]{𝐸} (6.2)

where {D} is the electric displacement vector, {T} is the stress vector, [e] is thedielectric permittivity matrix, [cE] is the matrix of elastic coefficients at constant


3

6

5

2

4

1

Polarization

Figure 6.5 Designation of axes in piezoelectric materials.

electric field strength, {S} is the strain vector, [aS] is the dielectric matrix at con-

stant mechanical strain, and {E} is the electric field factor. Strain and coupling

coefficients differ in stock and bender configuration modes. The stock configura-

tion mode (d33) generally depicts larger values. For energy harvesting application,

the materials that can be deformed easily to induce larger strains and exhibit large

coupling coefficients are desirable [40, 41].

6.2.1.1 The Piezoelectric Charge Constant and Piezoelectric Voltage Constant

The piezoelectric constant relating to the electric charge produced by an applied

mechanical strain is termed as the piezoelectric charge constant, which is repre-

sented bydij.The largedij constants relate to a large electric charge collected on the

electrodes following appliedmechanical stress and, conversely, the coefficientmay

be viewed as relating to mechanical displacement on an applied electric field. The

units for the dij coefficients are commonly expressed as coulombs/square meter

per newton/square meter (CN−1).

𝑑 =short circuit charge density

applied mechanical stress(6.3)

The piezoelectric constants relating to the electric field produced by a mechan-

ical stress are termed the voltage constants and represented by g ij and the units

can be expressed as volts per meter/newtons per square meter (VmN−1).

𝑔 =open circuit electric field

applied mechanical stress(6.5)

6.3 Processing of PVDF for Energy Harvesting Applications 137

6.3

Processing of PVDF for Energy Harvesting Applications

PVDF is commercially available in powder and pellet forms, which can be

extruded from melt in a conventional melt extruder. The polymer has been in

use since 1960s; however, it gained most attention when its piezoelectric, pyro-

electric, and ferroelectric properties were discovered. There have been a number

of researches on using PVDF. One of the very early studies of energy harvesting

by piezoelectric materials was performed in a biological environment by Hausler

and Stein [42]. The aim was to transform the mechanical energy caused by the

respiration of a mongrel to electrical energy using a piezoelectric PVDF film and

a converter. The piezoelectric material was fixed to the ribs of the dog and a peak

voltage of 18V was produced by motions of the ribs during the spontaneous

breathing. However, the current was too low so that the generated power was

only about 17 μW, which was not enough to operate an electronic device.

Shenck and Paradiso [43] also studied piezoelectric PVDF and lead zirconate

titanate (PZT) structures embedded in a shoe. A power storage circuit that was

designed to power a radio frequency tag was also mounted in a shoe and an offline

forward-switching DC–DC converter was developed. The experimental results

showed that the switching converter harvested energy more efficiently – about

twice as much – than the original linear regulator circuit. The whole setup was

successful to power low-energy electronic devices as the switching circuit pro-

vided continuous power during walking.

Another investigation into using piezoelectric materials for power harvesting

from themotion of humans and animals was performed by Ramsey andClark [44].

They studied the feasibility of using a piezoelectric transducer as a power supply

for an in vivo microelectromechanical system (MEMS) application. The 33- and

31-modes of operation for a piezoelectric generator were analyzed and compared;

it was determined that when using the 31-mode, or thin plate configuration, there

existed a strong mechanical advantage in converting applied pressure to work-

ing stress. For very low-pressure sources, the 31-mode had a greater advantage in

energy conversion, which became important when attempting to implement this

technology in a biological microsystem application.

In another analysis a self-powered mechanical energy sensor was theoretically

and experimentally studied [45]. The power harvesting system was consisting of

piezoelectric PVDFfilm, a charge capacitor, a transmitter, and a switch.The energy

generated by PVDF film was accumulated in a capacitor to power the transmitter

that could send a signal containing information depending on the strain of a beam.

The whole system was successful with a limitation of 2m distance.

Priya et al. [46] developed a piezoelectric windmill energy harvesting device

that consisted of 12 piezoelectric bimorph structures arranged in a circular array

and a conventional fan whose output shaft was connected to a cam system that

was also connected to the input shaft of the piezoelectric windmill. When the

fan rotated, the piezoelectric bimorphs underwent an oscillatory motion through

the cam system. Priya [47] also carried out similar work with 10 piezoelectric


materials. It was found that an increase in the wind speed caused a linear increase

in the power output of the piezoelectric energy harvester. The predicted power

output was 6.9mWat a wind speed of 10mph; however, a power output of 7.5mW

was measured experimentally at the same wind speed through a matching load

resistance of 6.7 kΩ.Lefeuvre et al. [48] developed a circuit containing a rectifying diode bridge and a

flyback switching mode DC–DC converter to improve the energy harvesting effi-

ciency of a piezoelectric harvesting device. The voltage across the diode rectifier

was sensed by a control circuit. When the maximum voltage was obtained, the

flyback converter was activated and the battery was charged. When the electric

charge on the piezoelectric was completely extracted, the control circuit deacti-

vated the converter and stopped the energy transfer. The process continued when

the next voltagemaximumwas detected, thus synchronizing the charge extraction

with themechanical vibrations of the system.When tested experimentally against

a linear impedance-based converter design, the synchronous converter increased

power transfer by over 400%.Theflyback converterwas found to have an efficiency

of 70%.

Jiang et al. [49] studied the efficiency of a piezoelectric bimorph for energy har-

vesting. A cantilever bimorph with a mass attached to its end was simulated. The

model was then used to determine the effects of both physical and geometri-

cal properties on the efficiency of piezoelectric harvester. The maximum power

generation was determined to be greater when the thickness of the piezoelectric

bimorph’s elastic layer was reduced and attached mass was increased. Granstrom

et al. [50] developed a theoreticalmodel of an energy harvesting backpack that can

generate electrical energy from flexible piezoelectric PVDF films integrated into

the straps. It was found that 45.6mW of power could be generated from a com-

plete backpack with two piezoelectric straps with an efficiency of more than 13%.

A multimaterial piezoelectric fiber production has been reported recently [51];

however, it was produced by a multiprocess method where a copolymer of PVDF,

P(VDF-TrFE), and polycarbonateswere used,whichmakes the fiber expensive and

difficult to scale up for large-scale production.

6.4

Processing of PVDF Based Materials: Polymer/Polymer, Polymer/Nanofiller,

Polymer/Ionomer Blends

General aims of blending polymers are to improve the physical properties of poly-

mers, to gain wide versatility, and to obtain different properties from those of par-

ent polymers. It can be done in two ways. One is that the polymers are blended in

different ratios and then characterized to determine what properties are improved

or gained. Characterization results give some clues about where the blend can

be used and in what application areas. This gives us a great variety of materials

and products. In the other method, the polymers that are to be blended are cho-

sen depending on the application. The first thing is to clarify what properties are

6.5 PVDF Based Nanocomposites for Energy Harvesting Applications 139

needed and what polymers can exhibit those properties when blended. The mis-

cibility of polymers is an important factor for obtaining materials with desired

properties.

Poly(methyl methacrylate), PMMA, is one of the homopolymers studied for

their miscibility with PVDF. It was claimed by Roerdink and Challa [52] that

isotactic PMMA was more miscible with PVDF as compared to atactic and

syndiotactic PMMA. Nasır et al. [53] produced PVDF/PMMA-blend nanofibers

by electrospray deposition technique and investigated it further. They reported

that PVDF/PMMA-blend nanofiber was amorphous at low PVDF ratios while an

increase in PVDF ratio in polymer blend resulted in decreased fiber diameter and

enhanced crystalline formation.They also investigated the humidity on nanofiber

formation. It was found that increased humidity caused an increase in average

fiber diameter and a reduction in beaded fiber formation.

Poly(ethyl methacrylate), PEMA, is another polymer with good miscibility with

PVDF. Sivakumar et al. [54] prepared PVDF/PEMA blend to form gel polymer

electrolytes by solvent casting technique and investigated the electrochemical

properties of the blend. It was found that the maximum ionic conductivity could

be obtained at lower PEMA ratios. They worked on various concentrations of

PVDF/PEMA and reported that 90 : 10 blend ratio of PVDF:PEMA was the opti-

mum value for an enhanced ionic conductivity and microstructural homogeneity.

Other than polymer/polymer blends, nanocomposites produced from

PVDF/nanofiller or nanoclay blends have also been studied by various researchers.

Dillon et al. [27] investigated PVDF/nanoclay composites. They used both solu-

tion casting and coprecipitation methods to produce nanocomposites and they

worked on three n-clay morphologies: exfoliated, partially intercalated, and

phase-separated morphologies. Asai et al. [55] studied the effect of nanofillers on

the crystallization behavior and structure of PVDF. They comparatively studied

the crystal morphology of neat PVDF and PVDF/layer titanate nanocomposites.

They found that the dispersed layer titanate particles acted as nucleating agent

so that the addition of them in polymer caused the formation of polar phase

in the blend. Li et al. [56] worked on ternary blends as well as polymer/clay

blends. They produced nanocomposites of PVDF/organoclay, polyamide 11

(PA11)/organoclay, and PVDF/PA11/organoclay blends by melt processing

technique. It was found that addition of nanoclays in PVDF/PA11 blend had a

significant effect on the polar phase formation of the polymer blend.

6.5

PVDF Based Nanocomposites for Energy Harvesting Applications

As mentioned earlier, the energy is of great importance in our lives. There are a

number of smart materials that can convert energy from chemical to electrical,

thermal to electrical, photon to electrical, and mechanical to electrical. One of

the most common forms of energy conversion is from chemical to electrical in

that a chemical reaction is used to create storable free electrons. From photon


to electrical conversion systems, photovoltaic systems, are also widely used for

the conversion of sun light. Piezoelectric energy conversion systems are being

increasingly studied. However, both ceramic-based systems and polymer-based

systems have drawbacks. It is difficult to produce ceramic based materials in

complex shapes since they are not flexible enough. Polymers, on the other hand,

are more versatile but have low piezoelectric charge coefficient as compared to

ceramic-based systems.

There have been an extensive work on to overcome the drawbacks of ceramic-

and polymer-based piezoelectric energy harvesting systems. Blending these

with different materials in different ways could help. In last decade, a number of

works were carried out to produce hybrid ceramic–polymeric composite energy

harvesting systems [57–64] so that the resulting material would have better

mechanical properties than ceramic and better piezoelectric charge coefficient

than polymer-based systems.

Piezoelectric materials have come a long way since Curie brothers’ discovery:

crystals, biological materials, ceramics, and, finally, polymers. The latest works

on piezoelectricity of polymers have been carried out on fibers [47, 65–67] and

yarnsnb [68]. Hadimani et al. [65] produced the first flexible piezoelectric fila-

ment in a continuous process. Polarization and fiber formation took place on a

melt extrusion equipment so that the filament material collected from the melt

extruder showed piezoelectric property. Fiber formation and poling via a contin-

uous process was a novel work that gained a patent [69] on the productionmethod

of the piezoelectric fibers.

Hadimani et al. [65] investigated the voltage response of the produced fibers by

developing a fiber composite structure. Produced piezoelectric PVDF fibers were

sandwiched in between two electrodes and their voltage generation was inves-

tigated on an applied impact. The results showed that the continuous produc-

tion of piezoelectric PVDF fibers were successful. The team was able to produce

multifilament, monofilament, and ribbon-like piezoelectric polymeric materials

in nanometer sizes. The electrodes applied on both sides of the uniformly aligned

piezoelectric fibers can be conductive paste or ink as MEAS (measurement spe-

cialties) produces its piezoelectric films [70]. Piezoelectric fiber nanocomposites

are flexiblematerial and can be used in awide range of applications fromnanoscale

to others, for example, textile structures [71].

6.6

Conclusion

The piezoelectricity of the polymers has been known for less than half a

century but the research works carried out on it are significant. The idea of

producing piezoelectric polymer nanocomposites becomes very important when

their piezoelectric charge constants are investigated and compared to ceramic-

based piezoelectrics. The overall aim of the research works was to improve the

direct piezoelectric property of the polymers by increasing the crystallinity and

References 141

converting the apolar phase to polar phase. Some of the works were successful

with PVDFmaterial showing higher output values when some nanoparticles were

dispersed into piezoelectric polymer structures.

Each and every new development on piezoelectric polymers for energy harvest-

ing applications is significant because our life has become so dependent on energy.

If we continue using the energy produced from ordinary energy resources, we will

have to face the negative effects of it such as air pollution, weather change, and so

on. Small or big, each and every attempt on alternative energy resources is vital.

Development of piezoelectric polymer nanocomposites may be seen as a small

attempt when the energy conversion characteristics are taken into account, but if

we use these materials, which can generate green energy, in every suitable appli-

cation, then its positive impact on the environment will be enormous.

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145

7

In Situ Thermal, Photon, and Electron-Beam Synthesis of

Polymer Nanocomposites

Luana Persano, Andrea Camposeo, AnnaMaria Laera, Francesca Di Benedetto, Vincenzo Resta,

Leander Tapfer, and Dario Pisignano

7.1

Introduction

The exceptional mechanical, optical, and electrical properties of composite

materials consisting of inorganic nanoparticles (NPs) incorporated within

polymer matrices have prompted a growing number of research groups to

develop innovative and convenient synthetic strategies, in view of device

fabrication on large scales. The fabrication process involving the incorpora-

tion, ex situ, of synthesized NPs into a properly chosen polymer melt, often

requires tedious purification procedures and NPs surface functionalization

with surfactants to prevent the formation of microsize aggregates and phase

separation phenomena. These phenomena may play an important role when

nanocomposite materials need to be interfaced with the external world and

coupled with other functional building blocks for the realization of chemical,

optical, or electro-optical devices. Indeed, it is well known that the overall

material roughness makes difficult the realization of patterns by conventional

exposure-based lithography, and the formation of aggregates may result in the

clogging of pores in membranes or fluidic channels, micro- and nanocapillaries

as those at the base of template-based lithography, or needles as in electrospin-

ning (ES) or extrusion-based fabrication methods. In this respect, the in situ

methodology, which is based on the nucleation and growth of NPs directly

inside the polymer matrix, allows researchers to exploit highly favorable flow

conditions of polymer solutions for the realization of lithographic patterns.

In addition, the absence of surfactant offers many advantages in terms of ease

of processing. This chapter is aimed at providing introductory background

information and state-of-the-art progress in the field of nanocomposite mate-

rials, films, and patterns realized by the exploitation of in situ methodologies

based on thermal, photon, and electron-beam-assisted synthesis. The chapter

is divided into two main sections. Section 7.2 aims at the introduction of

the most widely used precursor molecules and to processes accounting for

precursor decomposition and NPs nucleation, mainly during thermal-assisted


146 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites

experiments. Section 7.3 defines the most promising in situ synthesis and

patterning methods, also in combined approaches, based on photon- and

electron-beam-assisted procedures.

7.2

Thermal-Assisted In Situ Synthesis: Material Choice and Nanocomposite

Characterization

In situ methodologies are based mainly on the simultaneous preparation of NPs

and polymers in a single-step synthesis, in which phase separation can be avoided

by choosing suitable experimental conditions such as temperature, solvent, and

reaction time [1]. Several pathways can be used to induce NPs nucleation inside a

polymer matrix, such as chemical reduction, photoreduction, or thermal decom-

position.The last process is generally performed inmild conditions at temperature

above polymer glass transition temperature (Tg), so that the polymer molecules

acquire a sufficient mobility to enable NPs nucleation and growth without los-

ing the capability to tune NPs size. Furthermore, the polymer viscosity remains

high and prevents the collapse of individual NPs by diffusion effects. The poly-

mer acts as a template in which particles can grow in a controlled way preventing

aggregation phenomena [2].

7.2.1

Precursor Molecules

The choice and the dispersion of suitable precursors for NPs synthesis in an

organic polymer represents the preliminary steps in all the in situ synthetic

methodologies. In general, for the preparation of either metallic or semiconduc-

tor NPs, ideal precursors should be organometallic compounds or metal salts

that meet all the following requirements.

1) The precursor molecules should be cheap and commercially available, or

should require simple synthetic route.

2) The precursor chemical structure must have a good affinity with polymer

chain or functional groups linked to polymer chain, allowing for a homoge-

nous intercalation within the organicmatrix. A good dispersion of precursors

is a crucial prerequisite to obtainNPswith a homogeneous distribution inside

polymers, as required in many application fields.

3) The synthesis by-products should contribute only to passivate NPs surface

without affecting the peculiar properties of the resulting nanocomposite

materials derived from quantum size effects. The formation of volatile or

extractable side products would be highly desirable.

4) The inorganic product should form in mild conditions at low temperature

since, in general, organic polymers showpoor thermal stability. Inmany cases,

temperatures slightly higher than the polymer Tg are required.

7.2 Thermal-Assisted In Situ Synthesis: Material Choice and Nanocomposite Characterization 147

Precursors can be incorporated in polymeric matrices by adsorbing processes

from gas or liquid phase. Alternatively, polymer and precursor molecules can be

solubilized in the same solvent and subsequently dried to obtain well-mixed solid

samples. Rarely the mixtures are obtained from polymers and precursors in solid

state. Table 7.1 summarizes some examples of precursors used for in situmethod-

ologies. For each kind of precursor, the polymer matrix used, the final particle

formed, and the in situ pathway chosen to induce NPs nucleation and growth are

also reported.

7.2.1.1 Metal Salts

The first example, reported in literature about in situ generation of NPs within a

solid polymermatrix, involved the use of a commercial cadmium salt as precursor

[3]. The developed protocol allowed monodisperse cadmium sulfide (CdS) NPs

to be synthesized inside a blend systemmade of poly(styrenephosphonate diethyl

ester) and cellulose acetate (PSP-CA blend).The PSP-CA solidmatrix was kept on

an aqueous solution of Cd(NO3)2 for 24 h and the resulting sample, charged with

cadmium ions precursors, was exposed toH2S at room temperature to induceCdS

NPs in situ generations. This route, schematized in Figure 7.1, was later extended

to other metal salts [4]. Alternatively, the H2S exposure can been carried out in

solution.

However, in general, metal salts are used as precursors for the synthesis of

metallic NPs. In a recent work, AgNO3 was used as starting reagent to obtain

silver NPs in polyvinyl alcohol (PVA)/polyvinyl pyrrolidone (PVP) films [5].

The incubated AgNO3 was slowly reduced in 40 days by the PVP chain at room

temperature in the absence of light. Both size and size distributions of Ag

particles can be tuned by changing the percentage of PVP. The used pathway is

schematized in Figure 7.2.

The AgNO3 salt was also used as precursor for silver NPs in situ generation

by Shanmugam and coworkers [6]. Here films of PVA, tetraethyl orthosilicate

(TEOS), and silicotungstic acid (SiW), obtained by a spin-coating method from

aqueous solution, were dipped into a solution of AgNO3 for 10min and exposed

to sunlight. The silver cations Ag+ diffused from solution inside the polymer and

were reduced in metallic silver by silicotungstate ions. The formation of Ag NPs

induced a visible color change from blue to yellow as schematized in Figure 7.3.

The same route can be used to prepare nanocompositematerials based onAuNPs,

using HAuCl4 as precursor.

7.2.1.2 Organometallic Compounds

A convenient synthetic strategy to form semiconductor NPs in polymers is repre-

sented by the thermal decomposition of organometallic compounds that contain

both the metal and nonmetal part.This route ensures the control over the process

stoichiometry and allows one to overcome problems related to the possible inho-

mogeneity during multicomponent mixing or to the intrinsic high reactivity and

toxicity of reagents commonly used such as, for example, H2S.


Table 7.1 Selected examples of precursors used for in situ methodologies.

In situ pathway Precursors Nanoparticles Polymer References

1. Chemical

exposure to gas

or liquid

reagents

Cd(NO3)2 CdS Poly(styrenephosphonate

diethyl ester) and

cellulose acetate

(PSP-CA blend)

[3, 4]

AgNO3 Ag Polyvinyl alcohol

(PVA)/polyvinyl

pyrrolidone (PVP)

[5]

PVA [6]

2. Thermal

treatment

Cd(SR)2 with R:

alkyl

CdS Polystyrene (PS), poly

(3-hexylthiophene)

(P3HT)

[8, 85]

[Cd(SBz)2]2.MI

with MI 1-

methylimidazole

CdS Poly[2-methoxy-5-

(2′ethyl-hexyloxy)-1,4-

phenylene vinylene]

(MEH-PPV)

[10]

Cd(S2COEt)2 CdS P3HT [11, 86, 87]

CuOAc, InCl3,

thiourea

CuInS2 Poly(3-ethyl-4-

butanoate)thiophene

(P3EBT)

[88]

Zn(OAc)2 ZnO Polymethyl-

methacrylate

(PMMA)

[14]

ZnEt2 Poly[2-methoxy-5-

(3′,7′-

dimethyloctyloxy)-1,4-

phenylene vinylene]

(MDMO-PPV)

[12]

Zn(S2CNHC4H8N)2 ZnS PS [89]

Pb(S2COEt)2 PbS PS [90]

Ti(OC3H7)4 TiO2 Poly[2-methoxy-5-

(3′,7′-

dimethyloctyloxy)-1,4-

phenylene vinylene]

(MDMO-PPV)

[13, 91]

3. Irradiation

with Vis, UV,

or laser light

[Cd(SBz)2]2.MI CdS PMMA [51]

4. Electron

beam

[Cd(SBz)2]2.MI CdS PMMA [33]

PMMA: poly(methyl methacrylate).


+

M2+ M2+

M2+

M2+

M2+

M2+

M2+: Zn2+, Pb2+, Cd2+...

Polymer chainsNanoparticle

Nanoparticle

Polymer solution

containing nanoparticles

Polymer solution

containing M2+

Polymer film

containing M2+

Polymer film

containing nanoparticles

Mixing Forming film

Forming film

In situ formation of

nanoparticlesSulfi

de s

ourc

e

Sulfi

de s

ourc

e

Heating

Heating

Solvent

Figure 7.1 Scheme of in situ generation of metalsulfide NPs using metal salts as precur-

sors. Reproduced from Ref. [4] with permission of The Royal Society of Chemistry.

pH scale

AgNO3

7+ Casting

PVA/PVP

PVPPVAPVA

Ag+ Agnanoparticle

Acid

ic

Aging

40 days

Ba

sic

Figure 7.2 Scheme of an in situ synthesis pathway for the formation of silver NPs in a

polymer. Here, an AgNO3 salt is used as precursor. Reprinted from Ref. [5], Copyright (2012),


The use of a unimolecular precursor to obtain semiconductor NPs was

firstly reported by Brennan and coworkers, who used Cd[Se(C6H5)]2 or

[Cd(SePh)2][Et2PCH2CH2PEt2] to prepare nanometer-sized CdSe in refluxing

pyridine [7]. Afterwards unimolecular precursors for both III–V and II–IV semi-

conducting compounds were studied, but only in the past 10 years polymers were

tested as suitable matrix in which one can perform NPs nucleation and growth.

Well-studied precursors to obtain CdS NPs are cadmium alkylthiolate, whose

decomposition mechanisms have been thoroughly analyzed in the temper-

ature range 200–300 ∘C [8]. The cadmium alkylthiolate, totally insoluble in

most common organic solvents, were suspended in a solution of polymer


Composite filmon glass

Ag/compositeReduced

composite film

Agcolloids Dipped in

aq.AgNO3

hν

Figure 7.3 In situ silver NPs formation starting from AgNO3 as precursor. Reprinted from

Ref. [6], Copyright (2006), with permission from Elsevier.

3.5 nm

Figure 7.4 Lamellar structure of insoluble molecules of Cd(SC12H25)2.

in chloroform. The used polymer was a cycloolefin copolymer consisting of

ethylene and norbornene units. From suspension, solid films were obtained by

casting, and precursor arrangement inside the matrix was investigated mainly

by X-ray diffraction (XRD) before the thermal process. Alkylthiolates, having

linear chain with 12 or 18 carbon atoms, form lamellar structure because of

chain–chain interaction inside the polymer, as depicted in Figure 7.4. As a

consequence of this structural arrangement the decomposition process leads

to the formation of CdS nanocrystals (NCs) without a homogenous spatial

distribution inside the polymer matrix. However, many devices and applications

require polymer nanocomposites having a dense and homogeneous network of


inorganic NPs. Better results in terms of homogeneous distribution and reduced

processing temperature have been achieved with cadmium-bis(benzylthiolate)

Cd(SCH2C6H5)2-Cd(SBz)2, hereafter indicated as CBz, molecules embedded in

a poly(methyl methacrylate) (PMMA) matrix [9]. Furthermore, incorporating a

Lewis base (1-methylimidazole, MI) on the CBz molecule, hereafter indicated as

CBz-MI, it is possible to produce well-distributed CdS NPs inside an insulating

polymer, such as polystyrene (PS) or PMMA, or a semiconducting polymer, such

as poly[2-methoxy-5-(2′ethyl-hexyloxy)-1,4-phenylene vinylene] (MEH-PPV)

[10]. The solubility of both polymer and precursors in the same solvent is crucial

to control the spatial arrangement of the NPs in the final nanocomposites. The

ligand MI has the important function to destroy the lamellar arrangement of

bis(benzyl)thiol and to favor a better intercalation of precursor molecules among

polymer chains. The thermolysis process, performed below 200 ∘C, allowed CdS

NPs to nucleate with a regular distribution, without the formation of microsize

agglomerates. The CdS/MEH-PPV obtained nanocomposites represent an

example of hybrid material in which an electron-acceptor (n-type) material

is intimately mixed with an electron-donor (p-type) material, as required for

efficient photovoltaic conversion. In solar cells, in fact, only excitons generated

in close vicinity (5–10 nm) of the p–n heterojunction give rise to charge transfer

processes. Reynolds and coworkers recently demonstrated that the in situ

synthetic methodologies allow the distance between inorganic and organic

components in CdS/P3HT films to be reduced with respect to ex situ technique

[11]. Using the Cd(S2COEt)2 precursor, the authors realized nanocomposites

with improved charge separation efficiency compared to CdS/P3HT, prepared by

mixing polymer with presynthesized NCs. For applications in the photovoltaic

field the research group of Janssen efficiently prepared ZnO [12] and TiO2

[13] NPs in poly[2-methoxy-5-(3′,7′-dimethyloctyloxy)-1,4-phenylene vinylene]

(MDMO-PPV), starting, respectively, from ZnEt2 and Ti(OC3H7)4 as precursors.

The Zn(OAc)2 precursor was also used to prepare ZnO NPs in PS [14]. To

ensure the solubilization of both polymer and Zn(OAc)2 in the same solvent, a

mixture of methanol and toluene was used. Polymers having specific coordinating

functional groups promoted the formation of different arrangements of NPs,

thus demonstrating the importance of the polymer in the control of the overall

resulting morphology.

7.2.2

Thermal Synthesis and Composites Characterization

One of the most extensively used energy source for the in situ synthesis of NPs

is the heat. By considering organometallic compounds, progressive heating acti-

vates both the decomposition process of the precursor molecules, consisting in

the gradual removal of the organic part through the formation of volatile species,

and the nucleation and growth of NCs [9, 15]. Depending on the molecular struc-

ture of the precursor, such a process starts and evolves at different temperatures

andwith different rates, thus conditioning both the choice of the polymermatrices


(mainly related to their thermal capacity) and the efficiency of the nucleation pro-

cess in terms of processing time and homogeneous formation of NCs inside the

polymer. Different methods of analysis have been explored so far to determine the

distribution of the precursor molecules within the polymer, as well as to monitor

the NPs nucleation and growth processes upon thermal activation. In the follow-

ing, an overview on the most extensively techniques used for the nanocomposite

microstructural characterization (XRD, and Transmission Electron Microscopy,

TEM) and optical spectroscopy is reported.

7.2.2.1 Microstructural Characterization

The arrangement of the precursor molecules when dispersed inside the polymer

matrix can be studied by wide angle XRD on bulk samples as obtained after the

synthesis procedure and before the thermal process. XRD measurements on CBz

and CBz-MI precursors in PMMA show a band at the value of the scattering

vector, q of 11 nm−1, corresponding to amorphous PMMA, and sharp Bragg-like

peaks because of the periodic and regular ordering of the precursors within the

polymer matrix (Figure 7.5a,b). For instance, in CBz samples, the q-positions and

the indexing of the Bragg peaks are representative of a primitive cubic lattice,

with the first and most intense Bragg peak at q= 4.56 nm−1, corresponding to

the (100) peak that yields a lattice constant Λ= 1.37 nm (Figure 7.5a). The dimers

are placed at the edges of the cubic lattice. The XRD pattern of CBz-MI samples,

instead, show equidistant Bragg peaks indicating a superlattice order with period-

icity Λ= 1.65 nm (Figure 7.5b). These results indicate that CBz molecules exhibit

a superior spatial order leading to the formation of large “macromolecules,” while

CBz-MI molecules are characterized by smaller domain size, which leads to the

formation of lamellar structures. Such a difference in the molecular conforma-

tion may be accounted for by the reduced decomposition temperature and the

increased decomposition velocity induced byMI. Short chain thiolateswere found

to be arranged in a (tetragonal) unit cell, twoCd atoms tetrahedrically coordinated

to four S atoms, each of them bound to a hydrocarbon chain [16, 17]. Long chain

thiolates, instead, exhibit a lamellar structure of the unit cell. Once the annealing

process starts, the peaks associated with the precursor molecules gradually dis-

appear. For instance, the decomposition process of CBz is usually completed at

temperatures below 185 ∘C (Figure 7.5a). Bragg peaks associated with zincblende

CdS increase in intensity and sharpen while temperature increases (Figure 7.5a).

No peaks associated with CdS wurtzite phase has been found (ICDD no. 80-006)

[18]. The estimated average size of CdS NCs is of 1.8 nm at 175 ∘C in CBz-MI

samples and 2.3 nm at 220 ∘C in CBz samples (calculated by the Scherrer’s for-

mula, D= 0.9𝜆/𝛽cos𝜃, where D is the crystallite size, 𝜆 is the wavelength of X-ray,

𝛽 is the full width at half maximum of the most intense diffraction peak and 𝜃

is diffraction angle) [17]. Using long chain Cd(SC12H25)2 precursor molecules,

evidence of zincblende CdS NCs was found at 300 ∘C (average size 2.0 nm) in

polystyrene matrix [19], and at 240 ∘C in a different thermoplastic matrix [8]. In

the latter case, wurtzite CdS NCs of 8.0 nm are obtained at 300 ∘C [8]. The use

of ethyl xanthate precursor in P3HT was found to induce the formation of CdS


CBz

(111)

(200)

(220)

(311)

CBz-MI

(111)

(200)

(220)

(311)

240 °C 240 °C

185 °C 185 °C

11½

10½8½

4½5½

3½

2½

9½Λ –1.37 nm

12½

13½

14½

15½

16½

Before

Before

ΔD = 1.65 nm

5 10

(a) (b)

15 20 25 30 35

q Scattering vector (nm−1)

5 10 15 20 25 30 35

q Scattering vector (nm−1)

Figure 7.5 XRD patterns on CBz (a) and CBz-MI (b) samples before and after the annealing

processes as labeled close to each graph. Reprinted in part from Ref. [15], with kind permis-

sion from Springer Science and Business Media.

wurtzite NCs, whose size can be tuned from 2 to 6 nm by increasing the poly-

mer: CdS ratio from 1 : 1 to 1 : 8 [20]. TEM micrographs of CBz (Figure 7.6a) and

CBz-MI samples (Figure 7.6b), annealed at 185 ∘C, further confirm that methyl

imidazole added precursor gives rise to a highly homogeneous spatial distribution

ofCdSNCswithin the polymermatrix.On the contrary, without theMI group, the

thermal treatment leads to agglomeration and clustering phenomena of NCs and

consequently to an inhomogeneous distribution inside the sample (Figure 7.6a).

The measured average size was coherent with the estimated values from XRD

patterns and confirmed that at low temperature of annealing, the CdS NC size

is larger for CBz-MI molecules with respect to CBz. The insets of Figure 7.6a,b

also show single CdS NCs exhibiting well-pronounced (200) lattice fringes (fringe

distance 0.34 nm) for CBz samples (Figure 7.6a) and (111) lattice fringes (fringes

distance 0.29 nm) for CBz-MI samples (Figure 7.6b), both belonging to zincblende

structure in agreement with lattice theoretical parameter [18]. Figure 7.6c shows

the formation of a highly dense and uniform layer of CdS NCs in MEH-PPV syn-

thesized from the CBz-MI precursor [10, 15]. The inset shows a single CdS NC

exhibiting well-pronounced (101) lattice fringes and also demonstrates that, at the


Cds (200)

3 nm

100 nm

Cds (101)

3 nm

100 nm

Cds (111)

3 nm

100 nm

1

2

(a) (b)

(c)

Figure 7.6 Low magnification bright field

TEM images of CBz (a) and CBz-MI (b) sam-

ples with PMMA, and CBz-MI with MEH-PPV

(c). Annealing temperature= 185 ∘C. Insets:

magnified view of single CdS NC exhibiting

the fringes of cubic CdS (c). Reprinted with

permission from Ref. [10]. Copyright 2013

Springer.

same process temperature, the NCs size is slightly larger than in case of PMMA

(average size about 3–4 nm) [10], thus confirming the influence of the polymer

chain on the dynamics of the NCs growth [21].

7.2.2.2 Optical Spectroscopy Experiments

When the CdSNCs are synthesized, the polymermatrix changes color fromwhite

to bright yellow, consistent with the band gap values associated with CdS NCs

whose estimation can be carried out byUV-visible absorbance spectra.Theoptical

absorption of the samples (inset of Figure 7.7a,b) is typically characterized by a

high energy peak associated with the first excitonic transition between the ground

state and the single electron-hole pair state (1S3/2–1Se), and by a long wavelength


180

3.0

3.2

3.4

3.6

3.8

4.0

200

300

Band gap

CBz

PL peak

abs PL

CBz@185 °C

400 500Wavelength (nm)

600

220 240

Temperature (°C)(a) (b)

En

erg

y (

eV

)

180

3.0

3.2

3.4

3.6

3.8

4.0

200

300

Band gap

CBz-MI

PL peak

abs PL

CBz-MI@185 °C

400 500 600

220 240

Temperature (°C)

En

erg

y (

eV

)

Wavelength (nm)

Figure 7.7 Evolution of the band gap

energy (open symbols) and the PL peak

positions (full symbols) as a function of

the annealing temperature (TA) for CBz (a)

and CBz-MI (b) samples. Insets: represen-

tative absorbance and emission spectra at

TA = 185 ∘C for CBz (a) and CBz-MI (b). The

absorbance was obtained as ln(1/T), from

transmission (T) measurements recorded for

chloroform solutions of bulk like samples by

means of a Xe lamp and a monochromator

and the PL for the same chloroform solu-

tions was measured with an excitation wave-

length 𝜆exc = 330 nm. Insets are reprinted

and adapted with permission from Ref. [9].

Copyright (2010) American Chemical Society.

absorption tail between 350 and 450 nm. The photoluminescence (PL) emission

spectra of the CdS/PMMA composite typically feature a complex structure (inset

of Figure 7.7a,b), given by the superposition of various contributions: the peak

at 2.4–2.43 eV is characteristic of the CdS bulk exciton, which blue-shifts as the

size of the NCs is reduced, whereas the emission peaks in the range 2–2.3 eV are

attributed to shallow defects and deep trap states [9, 22]. Figure 7.7a,b show the

behavior of the estimated band gap [23], and the PL peak with temperature in

samples of CBz and CBz-MI in PMMA. Irrespective of the precursor used, when

the temperature increases, a redshift of CdS NCs emission was observed and a

saturation-like behaviorwas reached close to the annealing temperature of 220 ∘C.A similar behavior was observed in the absorption spectra of different precursor-

doped polystyrene samples [8]. The redshift of the emission is indicative of an

increasing NC size on increasing the annealing temperature.

7.3

Fabrication of Nanocomposites and Patterning

In the framework of nanocomposite-based optoelectronic devices, the availabil-

ity of patterning techniques enabling the selective confinement of luminescent

areas, even on the wavelength scale, is of crucial importance. To date, different

routes have been pursued in order to pattern hybrid nanocomposites made by

ex situ techniques. For instance, electron-beam [24], optical [25], and imprint

lithographies [26] have been carried out on acrylate or epoxy-based photoresist


matrices [27] while soft molding lithography has been employed to pattern

hybrid systems based on conjugated polymers at the wavelength-scale [28], and

template copolymerization in water has been used to pattern hybrid quantum

dots (QDs)/thermoresponsive polymers [29]. However, a general problem related

to lithographic applications of nanocomposites made by ex situ synthesis is the

undesired aggregation of NPs when mixed into a patternable matrix, mainly

as consequence of interdigitation of ligand alkyl terminals usually employed

in conventional synthetic methods. Nanocomposite aggregation in melts or

solutions may also play an important role in the highest resolution achievable

[27]. Among others, a possibility recently proposed to overcome such a drawback

is related to the capability to proper functionalize core-shell QDs with a pho-

tosensitive corona, which makes them solution processable, photopatternable,

and easily embedded into resins [30]. After photocuring, photopatternable QD

films were found to form dense, quasi-ordered arrays. In addition, the uniform

dispersion of QDs in acrylate resins enables the fabrication of three-dimensional

structures with resolution on the micrometer scale [30]. Another general prob-

lem is related to the often disfavored plastic behavior and flow conditions of

composite materials. Indeed, the ex situ incorporation of NPs in polymeric

matrices remarkably alters the rheology of the system with respect to the

corresponding bare polymers. As a consequence, the capability of film forming

is strongly reduced, the overall system viscosity is increased, and a retarded

thermomechanical response is registered. In particular it has been demonstrated

that embedding NPs in polymer melts doubles the longest relaxation time of

the system and decreases the slope of the dynamic storage (G′) and loss (G′′)

moduli with respect to bare polymer [28]. In this framework, in situ synthesis,

accomplishing the formation of NPs directly inside the polymer matrix, allows

researchers to exploit more favorable flow conditions. Through in situ synthesis

it is hence possible to generate NCs only after the composite films have been

patterned. Among the various routes so far successfully explored, one should

mention chemical reaction in gaseous environment [31], photografting [30],

optical [32], and electron-beam writing [33]. Surface-photografted poly(acrylic

acid) (PAA) has been used to mediate and control the nucleation and growth

of 5 nm ZnO NCs through the precomplex interaction between the carboxyl

groups in PAA and Zn2+. Using a metallic photomask with a circular hole (40 μmdiameter) during photografting PAA/ZnO QDs hybrid arrays on large area have

been achieved [31]. A combined bottom-up and top-down approach has been

instead used to pattern CdS NCs in regular squares with a size of about 40 μm. A

TEM grid is placed on a PVP film containing Cd2+ ions, which is then exposed to

H2S gas. CdS NPs are generated by the reaction of H2S with Cd2+ in the exposed

part of the film. When a green dye is embedded within the polymer matrix it is

also possible to realize a double color pattern [30]. A more detailed description

of the main achievements about in situ synthesis and patterning methods of

nanocomposites by light irradiation and electron-beam writing is reported in the

following sections.


7.3.1

Nanocomposites by Photoirradiation

In the past decade, in situ synthesis of NPs in polymeric matrices by irradiation

with the light of suitable wavelength and energy has emerged as a valuable

strategy to produce nanocomposites [34]. In this approach suitable precursors are

embedded in a polymer matrix and the exposure to UV, visible, or near-infrared

(NIR) light, decomposes or photoactivates the precursors, triggering the growth

of metallic or semiconductor NPs. Typically, the use of light for the in situ

synthesis of NPs allows their growth without significantly damaging the polymer

matrix, a feature that constitutes an advantage compared to other in situmethods

(such as thermal methods) that might alter the physicochemical properties of

the polymer. Moreover, the exploitation of light for the fabrication of nanocom-

posites is a highly versatile in situ synthesis approach, allowing the control of

the spatial distribution of the NPs in the polymer matrix. Therefore, micro-

and nanostructured composite materials can be realized by the combination of

light-based in situ synthesis methods with conventional photolithography and

more advanced photo-patterning methods [32].

7.3.1.1 UV and Visible Irradiation

Ultraviolet and visible radiation has been extensively used for the in situ synthesis

of both metallic [35–39] and semiconductor [40–42] NPs, exploiting both

incoherent light sources (mainly UV lamps) and laser sources. Concerning the

latter, UV-pulsed laser sources have been exploited, which provide suitable

peak energy for precursors’ decomposition/activation. Nanocomposite hollow

spheres of polystyrene and CdS NPs have also been produced by γ-irradiation of

microemulsions [43]. The hollow spheres have a mean diameter of 420 nm, and

emission peaked at 400 nm. Recently, Ag NPs have been synthesized in a PVA film

by exposure to sunlight [44]. By this approach, Ag NPs with a mean size of 13 nm

and a prolate shape are produced within a few hours of exposure to sunlight.

For the synthesis of metal NPs, typically ionic or molecular precursors are

exploited, which, on irradiation by UV light, produce zero-valent metals that

generate the NPs [34]. The metal species can be produced by a direct photore-

duction of the used metal salts or complexes or by the reduction of metal ions

by photoactivated, excited molecules or radicals. This method of NPs synthesis

is also known as photosensitization [34]. Semiconductor NPs are typically

synthesized by UV irradiation of molecular precursors containing the needed

atomic elements, such as Cadmium thiolate molecules (see Section 7.2.2).

As mentioned above, one of the main advantages of the in situ synthesis

by photo-irradiation is the possibility to pattern the resulting NPs/polymer

composite, by selective irradiation of specific regions of the polymer matrix

embedding the precursors. In this way, NPs are formed only in the irradiated

regions, and complex patterns of metal and semiconductor NPs can be formed in

the polymer matrix, which may find interesting applications such as in photonic

crystals and metamaterials, and in ultrasensitive spectroscopy [45–47]. To this


aim, UV photolithography, which is extensively used in microelectronic industry,

can be exploited to produce patterned NPs/polymer nanocomposites. Here,

the polymer/precursors films are exposed to UV light through a mask, which

contains the pattern to be realized, consisting of opaque and transparent regions,

the latter corresponding to the areas selected for NPs growth. Patterns of Ag and

Au particles with size of tens of micron have been produced by such approach

[48, 49]. The spatial resolution of photolithography can be in the submicron

range, depending on various factors. Diffraction of light represents a first effect

limiting the spatial resolution of optical lithography, thus giving for light in the

UV range a minimum feature size of the order of few hundreds of nanometers.

However, other effects can impact the spatial resolution of optically patterned

nanocomposite. In fact the atoms, ions, and molecule precursors can diffuse

inside the polymer matrix and allow a nucleation of NPs also in regions that are

not directly exposed. The diffusion of the active species is generally limited to

few microns considering the typical used materials [50] and represent a limiting

factor for the minimum achievable feature size.

Diffusion of ionic, atomic, and molecular species might also occur toward the

exposed regions, as recently reported by E. Yilmaz et al. [49]. In their work, they

have investigated the synthesis of Au NPs in PMMA containing AuCl4− ions as

precursors, observing a region (about 10 μm wide) around the UV-irradiated

areas characterized by a depletion of ions concentration. This observation

evidences that the diffusional dynamics of precursors and by-products inside

the polymer matrix and the aggregation effects of the atomic species must be

carefully accounted for the realization of patterns by light irradiation. Recently,M.

Sakamoto et al. [50] have reported interesting results about the spatial distribution

and shape of bimetallic Au/Cu NPs synthesized on UV exposure in a PVAmatrix.

This is highlighted in Figure 7.8a,b, showing optical images of patterned PVA

film containing precursors for Au/Cu bimetallic NPs. In particular in Figure 7.8a,

which shows the pattern realized soon after UV exposure, a central dark brown

region can be observed corresponding to the irradiated area, surrounded by a

colorless and dark double layer, that becomes pink several days after the exposure

(Figure 7.8b). A detailed investigation of the composition of the different regions

(Figure 7.8c–g) evidences the presence of almost spherical NPs with typical size

<6 nm in the dark brown region (Figure 7.8d), whereas the particles in the pink

area are 2–3 times larger (Figure 7.8e). Interestingly, tetrahedral particles are

found in the black region (Figure 7.8g) with the typical size of 25–50 nm and

much lower density compared to the other positions. The authors proposed a

different NP growth mechanism inside and outside the UV-exposed regions: (i)

in the regions exposed to the UV light, NPs grow mainly by coalescence because

of the high concentration of metal atoms, whereas (ii) outside the irradiated

regions, the growth of NPs is promoted by the continuous supply of metal

intermediates, at low concentration, from the adjacent UV-exposed regions [50].

Direct laser writing can also be exploited for the in situ synthesis and patterning

of nanocomposites. Here a UV laser beam is typically focused in a polymer film

doped with precursors, in order to induce locally the formation of NPs [51, 52].


Dark brown

center

Dark brown

center

(a) (b)

(d)

(c) (h)

(e) (f) (g)

Cross-point

(i) (k)

(j)

Black layerColorless layer

100 μm 200 μm 100 μm

100 nm

6 μm 4 μmUV laser

Vis laser

C

ba

Sample

100 μm

Black layer Pink layer

Figure 7.8 (a, b) Optical microscope images

of patterned Au/Cu bimetallic NPs in a PVA

after UV exposure (a) and 168 h after expo-

sure. (c–g) Optical microscope (c) and TEM

images (d–g) of Au/Cu NPs/PVA nanocom-

posites, measured three weeks after the UV

irradiation. The TEM images are acquired in

the positions shown in the optical image

(c). (a–g). Reprinted and adapted with

permission from Ref. [50]. Copyright 2007

John Wiley Sons Inc. (h) Schematic repre-

sentation of the NPs in situ synthesis by

multicolor laser processing. (i–k) Optical

images of the Cu NPs in a PVA film, pro-

duced by multicolor laser processing, using

a UV laser (𝜆= 363.8 nm) and a visible laser

(𝜆= 514.5 nm), aligned perpendicular to

each other, as schematized in (h). The let-

ters in picture (h) a–c, correspond to the

direction of the photographs shown in (i–k),

respectively. Scale bar: 3mm. (h–k) Reprinted


Copyright (2008) American Chemical Society.

Two-dimensional (2D) patterns with spatial resolution of about 10 μmcan be real-

ized by scanning the laser beam. By this approach fluorescent patterns of CdSNPs

in a PMMAmatrix have been produced, using a UV-pulsed laser source [51].The

size distribution and emission properties of the CdSNPs can be tailored by chang-

ing the exposure conditions, mainly the laser fluence and the number of exposure

laser shots [51]. As conventional UV-laser writingmethods allow only 2D patterns

to be fabricated, alternative approaches have been developed for the production

of 3D patterns of NPs in a polymer matrix. These include using holography [53],

that is, the exposure of the polymer film embedding the precursors to the intensity

pattern generated by the interference of different laser beams.

Another proposed strategy is based on laser beams with different wavelength

and relatively low intensity, as schematized in Figure 7.8h [54]. In the mul-

ticolor laser processing, the reactive species necessary for NPs synthesis are

produced by a multistep linear absorption process in the overlapping area of the

different laser beams. The shape and size of the volume where particles are

generated can be consequently designed by controlling the overlap of the laser

beams and their intensity [54]. Figure 7.8i–k show the example of a 3D volume of

Cu NPs synthesized in a PVA matrix by multicolor laser processing.


7.3.1.2 Multiphoton Irradiation

The fabrication of 3D patterns of metallic and semiconducting NPs is becoming

increasingly critical for different applications, such as metamaterials and optical

data storage, which require NPs patterns with high spatial resolution. Both holog-

raphy and multicolor laser processing have some limitations, both in terms of

flexibility of pattern design and spatial resolution. The fabrication of 3D struc-

tures with arbitrarily complex shapes and feature sizes down to few hundreds of

nanometers is made possible by two-photon lithography (TPL) [55, 56].This tech-

nique exploits the nonlinear two-photon absorption (TPA), a process that requires

the simultaneous absorption of two photons having half the energy of the relevant

transition of the absorbing material and the use of high-intensity NIR femtosec-

ond laser beams, because the probability of TPA depends on the squared laser

intensity.This allows the absorption and, consequently, theNPs synthesis, to occur

only very close to the focus of the laser beam, in a volume having typically sub-

diffraction characteristic size. 3D patterns can be straightforwardly produced by

scanning the focused laser beam inside the polymer matrix.

Figure 7.9 shows some examples of 3D nanocomposite structures composed by

a polymermatrix and fluorescent CdSNPs, which are synthesized in situ [55].The

micro-objects are obtained by a multistep process, involving first the two-photon

polymerization of a resin containing the precursors for NPs synthesis (cadmium

methacrylates), followed by the in situ growth of the CdS NPs by reaction with

hydrogen sulfide gas [55, 57].

A control of the size of the CdS NPs is achieved by modulating the crosslinking

density of the photo-polymerized networks, allowing for obtaining micro-objects

emitting light at different wavelengths in the range 446–528 nm (Figure 7.9f,g).

Moreover, pulsed NIR laser beams can also induce the in situ growth of NPs

by activating the reactive species through two- or multiphoton absorption,

thus enabling the concomitant growth of NPs and patterning within the same

process. Figure 7.10 shows some examples of Au and CdS NPs patterns made

by such approach [22, 58]. The growth of NPs can be controlled by modulating

the laser parameters. In particular, CdS NPs with size in the range 5–10 nm

have been grown by varying the fluence of the incident fs NIR laser beam in the

interval 0.1–0.5 J cm−2 [22]. In principle, complex 3D structures composed by

NPs can be realized by scanning the focused laser beam in the polymer matrix,

as demonstrated in Refs. [35, 59], where 3D structures of Ag NPs have been

fabricated by TPL. Overall, the two- and multiphoton approaches have shown a

high degree of flexibility, allowing for precisely controlling both the growth of the

NPs and the shape of the realized patterns.

7.3.2

Nanocomposites by Electron-BeamWriting

Electron-beam writing (EBW) is largely employed in many research fields,

especially for the fabrication of nanostructures at 100 nm- and sub-100 nm

scale, which would have potential applications such as high-density data storage,

sensing, biomimetics, and optical devices [60, 61]. In general, the combination


5.0kV

(a)

(b)

(c)

(d)

(e)

(f)

(g)

× 5500 1 μm

Figure 7.9 Scanning electron microscope

(SEM, a), optical microscope (b), and fluores-

cence (c) images of a CdS-polymer nanocom-

posite microbull (size 20 μm), fabricated by

TPL. SEM (d) and fluorescence (e) images of

microcoils of CdS–polymer nanocompos-

ites. (f–g) SEM image (left) and fluorescence

microscopy image (middle and right) of a 3D

microbull (f ) and a 3D microlizard (g) fab-

ricated from resins with different amount

of crosslinker. Scale bar: 10 μm. Reprinted


Copyright © 2008 WILEY-VCH Verlag GmbH &

Co. KGaA, Weinheim.

of submicrometer patterning capability and excellent overlay accuracy makes

EBW the lithography of election in all research fields based on the exploitation

of the quantum physics and electronics [62]. However, while EBW is almost

ubiquitously used for high-resolution lithographic experiments on polymeric

resists, the combined approach of nanopatterning and in situ synthesis of NPs

is only poorly explored. To date, only silver NPs [59, 63] and CdS NCs [33]


2.0 2.2

ExposedNot exposed

2.4 2.6 2.80

−0.2

0.0

Heig

ht

(μm

)

0.2

0.4

5 10 15 20

Exposed

5 μm

10 μm

(a) 10 μm20 μm(e) (f)

Not exposed

25x (μm)

(b)

(c)

(d)Energy (eV)

PL inte

nsity (

Arb

. U

n.)

Figure 7.10 (a) Atomic force microscopy

(AFM) images of an Au NPs/PVA nanocom-

posite, fabricated by two-photon expo-

sure. The intensity profile of two interfer-

ing beams is used to produce the arrays.

Area of the images: 20× 20 μm2 (up-left),

10× 10 μm2 (up-right), 5× 5 μm2 (low-right),

and 1× 1 μm2 (low-left). Reprinted with

permission from Ref. [58]. Copyright 2003,

American Institute of Physics. (b) Photolu-

minescence spectra of the nanocomposite

fabricated by multiphoton lithography. The

spectra are measured in regions exposed

to the fs laser beam (continuous line) and

from unexposed regions (dashed line). AFM

topographic image (c) and height profile

(d) of the nanocomposite film after fs laser

exposure. The micrograph is collected at

the edge of an exposed area, marked by

the continuous line. (e–f ) Fluorescence

images of CdS patterns in a polymer matrix,

obtained by scanning a focussed laser. (b–f )

Ref. [22] – Reproduced by permission of The

Royal Society of Chemistry.

have been synthesized by this approach. In both cases, suitable precursors are

embedded in a polymer matrix and the subsequent exposure to a high voltage

electron beam (20–30 kV) decomposes or thermoactivates the embedded

precursors. Compared to light irradiation, the main advantage offered by EBW is

represented by the capability to scale down the lateral size of the smallest feature

achievable during exposure, without altering significantly the physicochemical

properties of the polymer. In addition, a unique multiscale combination of

complementary nanofabrication approaches including nanoimprinting, ES, and

electron-beam decomposition of precursors and subsequent NCs formation has

been demonstrated by this approach.

Concerning the realization of spatially confined Ag NPs, in situ-synthesis

was performed using precursor molecules containing the metal atoms that

undergo a reduction process on irradiation. This process is generally assisted by

the radical intermediates of the polymer that are generated during the scission

of their bonds [59, 64]. For instance, a film of PVA and AgNO3 salt has been

used as a negative lithographic resist for electron-beam exposure at 30 kV under

different doses. After exposure, the sample was developed with distilled water.

Square (Figure 7.11a,f ) and line patterns with resolution up to 100 nm have been

demonstrated in this way.


10 μm(a) (b)

(g)

(h)

(c)

(d)

(e)

(f)15 μm

1 μm

0

0.6

0.7

0.8

0.9

1.0

5

1

1 λ = 633 nm2 λ = 458 nm

1 λ = 633 nm2 λ = 458 nm

2

2

10 15 20

0 5 10 15 20

25 30Position (μm)

Position (μm)

Tra

nsm

itta

nce

0.6

0.5

0.4

0.3

0.2

0.7

0.8

0.9

1.0

Tra

nsm

itta

nce

Figure 7.11 In situ synthesis of Ag NPs in

PVA by electron-beam exposure. (a, f ) SEM

micrographs of exposed areas (rectangles,

15× 20 μm2). (c) Confocal scanning trans-

mission microscopy at two different wave-

lengths: 458 nm (b) and 633 nm. (d, e) Zoom

of (b, c), respectively. (h, g) Transmittance

profiles of panels (b, c), respectively. Ref. [63]

DOI: 10.1088/0957-4484/19/35/355308. © IOP

Publishing. Reproduced with permission. All

rights reserved.

Confocal images of square patterns reveal optical absorption at wavelengths

related to surface plasmon absorption of Ag NPs (458 nm, Figure 7.11b,d) and

far from the typical absorption band of Ag NPs (633 nm, Figure 7.11c,e). By

comparing the transmittance profile at the same wavelengths, 458 and 633 nm

(Figure 7.11g,h), with similar samples where Ag NPs have been generated

by thermal treatment at 180 ∘C, authors found larger values of the measured

absorption in the AgPVA patterns generated by electron-beam exposure and

correlated this result to a larger degree of electromagnetic coupling between the

generated NPs. Concerning semiconductor NCs, and specifically CdS, in situ

synthesis and patterning in one-step process has been demonstrated [33]. The

synthesis is accomplished using a unimolecular precursor containing both the

metal and nonmetal part of the semiconducting NCs. Electron-beam lithography

(EBL) at 20 kV both activate and control in space the synthesis of CdS NCs

within the polymer matrix. Cd-complexes, [Cd(SBz)2]2MI, are obtained by the

incorporation of MI as a Lewis base to the cadmium-bis(benzylthiol) (SBz)

complex and used as precursor molecule (see Section 7.2 for more details).

PMMA, chloroform, and MI are added to the final solution and processed to

form a thin film. Rheology studies showed that, unlike ex situ procedures, the

incorporation of the precursor in the polymer only weakly increases the system


viscosity and the longest relaxation time in the system. This effect can be simply

related to a reduction of the local mobility of the PMMA chains.

Through EBL at an exposure dose of 350 μC⋅cm−2, differently shaped submi-

crometer patterns have been realized (Figure 7.12). Under confocal microscope,

patterns appear bright and with a homogeneous distribution of CdS NCs within

each feature. High-resolution TEM images confirm the formation of CdS NCs

with an average size of about 4 nm.

As anticipated above, electron-beam turns out to be the only in situ synthetic

method enabling the exploitation of multilevel lithography and pattern-

ing through a combination of complementary nanofabrication approaches

(Figure 7.13). The simplest process that can be envisaged in this respect is direct

writing (Figure 7.13a), which, although very similar in concept to light irradiation,

enables submicrometer pattern resolutions and complex round-shaped features.

Alternatively, EBL can be used to synthesize NCs on a preimposed pattern.

For instance, it is possible to perform room temperature nanoimprint lithogra-

phy (RT-NIL) on the precursor-doped film using stamp with features down to

hundreds of nanometers, and right after the pattern transfer has been fully accom-

plished, one can scan the surface with the electron beam in order to synthesize

NCs (Figure 7.13b). The formation of CdS NCs in films of PMMA modifies the

surface topography (Figure 7.14a,b), which can reasonably be ascribed to particle

nucleation and the consequent weight loss of the doped polymer following the

generation of volatile species during thermal decomposition. However, the pat-

ternmorphology, in terms of both linewidth and spacing features (Figure 7.14c,e),

20 μm

Figure 7.12 (a–f ) In situ synthesis of CdS NPs in PMMA by electron-beam exposure. Con-

focal PL images of different patterns. Reprinted with permission from Ref. [33]. Copyright ©

2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim.


Film

(a)

(b)

(c)

FilmMaster Master

RT-NIL

RT-NIL

Master Master

Fiber

E-beam

Substrate

Precursor-doped PMMA

CdS–polymer nanocomposite

E-beam nanolithography and in situ

synthesis of CdS nanocrystals

E-beam in situ synthesis

of CdS nanocrystals

E-beam in situ synthesis

of CdS nanocrystals

E-beam

E-beam

Figure 7.13 Schematics of the electron-beam synthesis of polymer-NCs composites (a) and

of its combination with other methods of nanofabrication (b, c). Reprinted with permission

from Ref. [33]. Copyright © 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim.

is well preserved after the electron-beam exposure (Figure 7.14d,f ). The high

spatial control of EBL nanolithography allows one to fabricate more complex

multilevel nanostructures. For instance, RT-NIL performed sequentially with

ES allows one to realize periodic and controllable surface topographies, in

which features can be patterned perpendicular to the fiber longitudinal axis.

CdS synthesis can be accomplished downstream all the lithographic processes.

Considering, in an even broader scenario, the importance of one-dimensional

structures and nanostructures in the form of fibers, the following section is

dedicated to the in situ synthesis of NPs within electrospun nanofibers.

7.3.3

Nanocomposite Polymer Fibers

1D nanostructures with various sizes and morphologies, such as nanofibers,

nanowires, nanotubes, nanorods, and nanoribbons draw remarkable attention

[65] being considered an ideal system for studying a large number of novel


(a) (b)

0 nm

130 nm

0 nm

230 nm

1.0 μm 1.0 μm

(c) (d)

0 nm

320 nm

0 nm

450 nm

2.0 μm 2.0 μm

(e) (f)

0 nm

340 nm

0 nm

290 nm

1.0 μm 1.0 μm

Figure 7.14 Atomic force microscopy micro-

graphs of the pristine precursor-doped poly-

mer surface (a) and of the nanocompos-

ite film after electron-beam exposure (b).

(c, f ) RT-NIL patterns before (c–e) and after

(d–f ) the electron-beam exposure. Reprinted


Copyright © 2012 WILEY-VCH Verlag GmbH &

Co. KGaA, Weinheim.

phenomena at the nanoscale, and for investigating the dependence of functional

properties of organic nanostructures on size reduction. These systems are also

envisioned as the most promising building blocks for nanoscale electronics,

optoelectronics, electrochemical, and electromechanical devices [66]. Several

fabrication and synthesis methods are currently available for the production of

such 1D nanostructures, such as self-assembly [67], melt-blowing [68], various

patterning approaches [69], and ES. Among them, ES is unique for simplicity and


effectiveness, enabling the production of continuous fibers from a wide variety of

materials, including plastics, block copolymers, blends, biopolymers, conducting

polymers, composites, metal oxides, and ceramics under continuous run at labo-

ratory and industrial scale [70–72]. ES relies on the reduction of the cross-section

of a viscoelastic jet induced by an applied electrostatic field and by the solvent

evaporation during the time of flight of a solution from a spinneret (metallic

needle) to a collection plate.The combination of uniaxial stretching and whipping

motion following bending instability reduces the diameter of the jet potentially

down to tens of nanometers. The ability to control fibers’ properties in terms of

diameter, chemical composition, surface morphology, porosity, and aspect ratio,

combined with the endless possibilities to embed or coat functional moieties

onto the fibers, make them useful in numerous potential applications such as

membrane technology, reinforced composites, enzyme immobilization, catalysis,

tissue engineering, smart textiles, and functional coatings/sensors [71–73].

During the past few years, ES has been successfully employed to produce hybrid

nanofibers by a synthetic strategy involving three steps: (i) codissolve the precur-

sors and polymer in one solvent to obtain a homogeneous solution; (ii) electrospin

the above solution to fabricate composite nanofibers; and (iii) decompose the pre-

cursor to directly synthesize NPs inside the fiber-shaped polymer matrix. In this

framework, the polymer phase both stabilizes the blend and confines the growth of

NPs, thus preventing aggregation phenomena. In addition, it acts as suitable host

allowing researchers to overcome the poor viscoelastic behavior of low molar-

mass precursor molecules. In the following section, we provide an overview of

the state of the art on photo-assisted and thermal decomposition of precursors in

electrospun nanofibers.

7.3.3.1 Photo-Assisted Synthesis

UV irradiation is commonly used as source for the decomposition of a suitable

metallic salt embedded within or nearby a polymeric matrix. For instance, Park

and coworkers successfully synthesized Ag NCs on the surface of ultrafine CA

nanofibers [74, 75]. A CA solution with small amounts of AgNO3 was electro-

spun and the synthesis of Ag NPs was accomplished by exposing the fibers to UV

light. The authors demonstrated that by simply modifying the metallic salt con-

centration in the ES solution, one can control the electrospun fiber diameter and

the NPs size in the resulting nanocomposites. In fact, the addition of the metal-

lic salt increases the charge density in the polymeric solution, and therefore the

intensity of the stretching forces acting on the jet. Such effect promotes the for-

mation of thinner fibers. On the contrary, the average diameters of the formed Ag

NPs increases on increasing the salt concentration in the solution.The NPs diam-

eters can be tuned in the range 3.3–6.9 nm [74]. It is even possible to control the

density and the particle size by changing the exposure time and the wavelength of

the UV source [75].

On the basis of the same strategy,Wang and coworkers preparedmonodisperse

and single-crystal silver NPs by ES, using a polyacrylonitrile (PAN)/AgNO3 solu-

tion [76]. TEM evidenced that the Ag NPs were spherical and homogeneously


distributed. Depending on the molar ratio of silver nitrate to PAN, the diameter

of Ag NPs could be easily tuned.

In the case of semiconductor NPs, one can combine UV photoreduction

and chemical reactions in H2S atmosphere. Dong and coworkers synthesized

Ag2S NPs homogeneously distributed on the surface of PAN nanofibers by

spinning PAN/AgNO3 mixtures. The gas exposure allowed the formation of

Ag2S NPs with a spherical shape and a diameter of about 9 nm (Figure 7.15a).

TEM and selected-area electron diffraction (SAED) patterns evidenced the

monocrystallinity of the formed NCs (Figure 7.15b).The authors also synthesized

Cu2S nanorods on the outer surface of PAN nanofibers via the exposure of the

(b)(a)

100 nm

(c)

10 μm

(d)

2 μm

Figure 7.15 Semiconductor NPs synthesized

on the outer surface of PAN nanofibers via

photo-assisted method. TEM image (a) and

SAED pattern (b) of PAN/Ag2S composite

nanofibers. TEM images (c, d) of PAN/Cu2S

composite nanofibers and correspond-

ing SAED pattern (inset in Figure 7.15d).

Reprinted and adapted from Ref. [77], Copy-

right (2007), with permission from Elsevier.


electrospun polymer/metal salt (PAN/CuCl2) composite nanofibers to H2S gas at

room temperature (Figure 7.15c,d) [77].

7.3.3.2 Thermal-Assisted Synthesis

Recently, Stojilovic and coworkers reported the formation of ZnO NCs in

hexagonal wurtzite phase [78] by ES a solution of PVP with Zinc acetate salt

and subsequent heating for 12 h at low temperature in a tube furnace. The

authors demonstrated that exploiting a calcination process at temperature as

low as 120 ∘C, one can prevent any damage of the polymeric phase and obtain

hybrid nanofibers. Also, Natarajan and coworkers reported the fabrication of

high-quality ZnO NPs by a calcination process at low temperature [79]. In this

case, PVA acted as organic stabilizer, suppressing the grain growth and aggrega-

tion phenomena. TEM images, XRD, and SAED diffraction patterns indicated

that the formed NPs are crystalline with a wurtzite structure. Di Benedetto and

coworkers reported on the in situ generation of CdS NCs inside electrospun

fibers of PMMA by thermal decomposition of CBz-MI [80]. The resulting fibers

exhibited a smooth surface and a quite uniform diameter, as shown by SEM

(Figure 7.16a,b). A precursor-doped polymer solution was electrospun, and the

resulting fibers were thermally treated at 175–250 ∘C for 20min. By controlling

the process parameters, the diameter of the resulting fibers could be tuned. In

particular, by increasing the feeding rate of the polymer solution one collects

fibers with larger diameters (up to 330 nm) whereas at a feeding rate of 1 μLmin−1

the average fiber diameter decreases down to 170 nm (inset of Figure 7.16c).

The formed CdS NPs, imaged by TEM (Figure 7.16d,e), were spherical in shape

and well separated from each other, with a rough bimodal size distribution

including populations with a diameter of about 3 nm and of 10–15 nm (insets

I and II of Figure 7.16e, respectively). In both cases, the CdS NCs showed a

wurtzitic structure, as confirmed by high-resolution TEM images that displayed

well-defined lattice fringes (insets of Figure 7.16e). Further insight into the

structure of the CdS/PMMA nanocomposite fibers were obtained performing

XRD measurements at an incidence angle, 𝜔i = 1.0∘ in order to reduce the X-ray

penetration depth and enhance the sensitivity of the measurement (Figure 7.16f ).

The observed diffraction peaks can be attributed to the (100), (101), (110), and

(103) planes of the wurtzite phase of CdS (ICDD, no. 80-0006) [18]. The broad

peak at about 2θ of 24∘ was related to the PMMA polymer matrix whereas the

diffraction peaks at 2θ of 38∘, 45∘, and 78∘ correspond to the (111), (200), and

(311) planes of the cubic Al (ICDD, no. 85-1327) [18] of the used substrate foil.

The average size of the CdSNCs, about 12 nm, was approximately calculated using

the Scherrer’s formula (see Section 7.2.2.1).

The formation of CdS NPS and the interaction with the polymer phase were

studied by comparing the infrared absorption frequency of composite fibers with

bare PMMA (Figure 7.17). After the thermal treatment, all the organic groups

connected to the polymer matrix, including C=O (an intense peak at 1723 cm−1),

C–O (a broad band ranging from 1270 to 1000 cm−1), and C–H (a band from 950

to 650 cm−1) did not show significant variations of their relative intensities and


30 μm

50 nm20 nm

200

0

160

200

240

280

320

2 4 6 8 10 12Feeding rate (μI min−1)

Ave

rag

e d

iam

ete

r (n

m)

0

10

20

30

400 600 800 1000Fiber diameter (nm)

Popula

tion

(%)

10 20

0

10

20

30

40(100)-CdS

(101)-CdS

(111)-AI

(110)-CdS

(200)-AI

(103)-CdS

PMMA

(311)-AI

ωi = 1.0°

30 40 50 60 70 80

2θ (degrees)

Inte

nsity (

arb

. units)

(b)

(a)

2 μm

(a) (b)

(d) (e)

(c)

(f)

Figure 7.16 Morphological and structural

characterization of composite nanofibers

before and after the decomposition of the

precursor via thermal treatment. SEM pic-

tures of the hybrid nanocomposite fibers

with different magnification (a, b). (c) Diame-

ter distribution of hybrid nanofibers collected

after ES with an applied voltage of 6 kV

(needle-collector distance of 12 cm and injec-

tion rate of 1 μLmin−1). Inset: typical diam-

eter distribution of the hybrid nanofibers

at different ES injection rate. TEM images

of CdS NPs entrapped within PMMA fibers

at low (d) and high magnification (e). Inset:

magnification of 3 nm (I, scale bar= 3 nm)

and 15 nm (II, scale bar= 5 nm) particles.

(f ) Glancing incidence X ray diffraction pat-

tern of nanocomposite fibers recorded with

an incidence angle, 𝜔i = 1.0. Reprinted and

adapted with permission from [80]. Copyright

2011 Royal Society of Chemistry.

position.This result indicated the absence of chemical bonding or strong interac-

tion between the polymeric matrix and the NCs.

On a similar polymer-precursor system, EBW has been used as powerful tool

to synthesize well-dispersed CdS NCs in nanofibers, with a superimposed grating

(see Section 7.3.2) [33]. The selective nucleation of CdS NPs was confirmed by

fluorescence images collected at the edge between pristine and EB-exposed fibers.

Huang and coworkers reported the fabrication of TiO2 nanostructures in poly-

meric nanofibers of poly(phenylene vinylene) (PPV) [81] and PVP [82], by cou-

pling a sol-gel method with calcination at low temperature. After ES, fibers were

annealed at 210 ∘C under controlled atmosphere. The realized NPs were well dis-

persed within the fibers, with size in the interval 10–60 nm.

An alternative pathway for the synthesis of CdS NCs in electrospun fibers was

reported by Wang et al. A PPV solution with a small amount of cadmium acetate

(CdAc, Cd(C2H3O2)2) was electrospun, and fibers were kept in H2S atmosphere

at 60 ∘C for 12 h and finally annealed at 180 ∘C in N2 atmosphere for 2 h [83].

TEM images evidenced CdS NPs with the size of 3.5–9 nm, uniformly dispersed

7.4 Conclusions 171

PMMA/Precursor

(a)

(b)

PMMA

PMMA/CdS

1750 1500 1250 1000 750

Wavenumber (cm−1)

Tra

nsm

itta

nce (

arb

.un.)

Tra

nsm

itta

nce (

arb

.un.)

PMMA

Figure 7.17 FT-IR absorption spectra of nanocomposite fibers before (a) and after (b)

thermal treatment (continuous line) compared to bare PMMA (dashed line) nanofibers.

Reprinted with permission of [80]. Copyright 2011 Royal Society of Chemistry.

within the polymer. XRD diffraction patterns revealed the hexagonal phase of the

crystalline structure. A similar procedure was used by Huang and coworkers to

incorporate CdS NPs into poly(ethylene oxide) fibers [84].

7.4

Conclusions

The in situ synthesis of both metal and semiconductor NPs/NCs within polymers

is based mainly on the growth of NPs in matrices containing suitable precursors.

By properly choosing the precursors, the polymer matrix, and the experimental

conditions in terms of temperature, solvent, and reaction times, it is possible to

induce NPs nucleation and growth processes in a variety of ways. In general, the

used precursor molecules must show a good affinity with the polymer chain or,

alternatively, functional groups may be added to homogenously intercalate such

molecules within the organic matrix. A good dispersion of the precursors is a

crucial prerequisite to obtain NPs with ordered spatial arrangement within the


polymer, as required in many application fields. The successive steps concerning

the precursor decomposition and the NPs nucleation and growth may be acti-

vated by following different pathways, including thermal, photon, and electron-

beam-assisted in situ synthesis methods. The combination of such methodology

with room temperature patterning and elongational methods such as ES has been

demonstrated to be very useful for employing these materials in nanocomposite-

based optoelectronic devices and sensors.

Acknowledgments

The authors acknowledge the support from the Italian Minister of University and

Research through the FIRB project RBFR08DJZI “Futuro in Ricerca” and from the

Apulia Regional Projects “Networks of Public Research Laboratories,” Wafitech

(9) and M. I. T. T. (13).

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179

8

Synthesis of Polymer Nanocomposites by Water-Assisted

Extrusion

Naïma Sallem-Idrissi, Michel Sclavons, and Jacques Devaux

8.1

Introduction

In the large field of nanotechnology, polymer matrix-based nanocomposites

have become a prominent area of current research and development in order

to upgrade their performances. Clays are one group of nanofillers that have

been widely used for the preparation of polymer nanocomposites. Polymer/clay

nanocomposites (PCNs) are polymer reinforced with clay particles, mainly

phyllosilicates, having at least one dimension in the nanometer scale. During the

past decades, PCNs received a considerable scientific and technological interest

mainly because of their unique combination of properties. A small percent of

clay usually confers a flame retardant (FR) behavior, higher mechanical, and

barrier properties regarding the neat polymer [1–4]. Because of the nanoscale

dispersion of the clay mineral filler, the reinforcement efficiency of the composite

can be significantly better than conventional minerals fillers. The final properties

of nanocomposites depend directly on several factors such as the chemistry of

the polymer matrix, the affinity between filler and polymer, the geometry of

the particle, its degree of orientation and dispersion inside the polymer as

well as the way of preparation [3, 4]. PCNs technology originated in the late

1980s at Toyota’s Central R&D laboratories [5]. Toyota focused mainly on

nylon composite and holds the basic patents on their production [6–8]. Since

this pioneering work, several books [9–11] and reviews [1, 2, 12–14] treat on

PCNs. Attention has also been focused on the preparation method of these

nanocomposites. To prepare polymeric nanocomposites, different methods have

been developed, among which melt compounding has attracted a great interest to

produce PCNs. The benefits of this technique result from its cost-effectiveness,

its environment-friendly aspect, and its compatibility with current industrial

processes such as extrusion. The residence time and the shear created in the

melt during processing may also be helpful to support the dispersion of particles.

However, it is not always enough to break up big particle agglomerates, as the

resulting dispersion of particles may remain poor and usually require the use

of clay organomodification to improve the intercalation and to solve the lack


180 8 Synthesis of Polymer Nanocomposites by Water-Assisted Extrusion

of compatibility between the polymer matrix and the filler. Another problem of

producing PCNs concerns the thermal stability of the surfactant in the organo-

clay. At extrusion temperatures, the interlayer alkylammonium surfactants often

decompose leading to a collapse of the silicate layers and impart undesirable

color, odor, and taste to the composite [15–20]. An alternative way to prepare

nanocomposites is the water-mediated melt-compounding method, where water

is used as a substitute for classical alkylammonium intercalating/exfoliating

agents. This technique has gained ground because of the following benefits. First,

many water-dispersible commercial nanofillers are readily available. Second,

the production of nanocomposites using aqueous dispersions or slurries is not

only an affordable method (no fillers’ organophilic modification is needed), but

it is also less hazardous to one’s health. Moreover, this original process opens a

novel eco-friendly route for especially the temperature-dependent reinforcement

materials such as natural fibers.

8.2

Nanocomposites Structure and Characterization

8.2.1

Clays

Nanocomposites can be classified depending on the shape of the nanofiller.

Depending on their dimensions, there are three categories of nanofillers:

nanoparticles, nanotubes or whiskers, and nanolayers [21]. These nanoscale

particles can further be divided into three types as natural, incidental, and

synthetic or engineered nanoparticles, depending on their pathway.

In the synthesis of nanocomposites, the commonly used layered silicates are

natural or synthetic minerals. Clays are characterized by a layered structure.Their

crystal structure consists of layers made up of tetrahedral sheets in which a silicon

atom is surrounded by four oxygen atoms and octahedral sheets in which a metal-

like aluminum ormagnesium is surrounded by eight oxygen atoms.The total layer

thickness is around 1 nm and the lateral dimensions of these layers may vary from

30 nm tomicrons or larger, depending on the particulate silicate, the source of the

clay, and the method of preparation [1, 2, 9–14, 21].

Montmorillonite (MMT), hectorite, and saponite are the most commonly used

in the preparation of polymer nanocomposites.The generally accepted formula for

MMT is (Na,Ca)0.3(Al,Mg)2Si4O10(OH)2⋅nH2O. MMT is characterized by a high

aspect ratio, presents a well-dispersed state without breaking of layers, and has a

surface area of about 750m2 g−1. These benefits in addition to its easy availability

make MMT widely accepted in the preparation of polymer nanocomposites.

Halloysite nanotubes (HNTs), another type of clay nanofiller combining the

geometry of nanotubes and the chemistry of kaolinite, are recently receiving

a lot of attention [22–31]. HNTs are a naturally occurring aluminosilicate

(Al2Si2O5(OH)4⋅2H2O) with a predominantly hollow tubular structure. As

8.2 Nanocomposites Structure and Characterization 181

compared to other layered silicates, HNTs can thus be more easily dispersed in a

polymermatrix because of the weak secondary interactions among the nanotubes

via hydrogen bonds and van der Waals forces [30, 32].

8.2.2

Organomodification of Layered Silicates

Improvement of polymer properties is strongly influenced by the state of dis-

persion. But the formation of well-dispersed pristine clay layers in polymers

is a challenge because of their high intrinsic hydrophilicity, which make them

incompatible with hydrophobic polymer. Incompatibility and weak interfacial

interactions prevent the preparation of dispersed stable nanocomposites with

improved properties. In their pristine state, clays are “miscible” only with few

hydrophilic polymers such as poly (ethylene oxide) (PEO) and poly (vinyl alcohol)

(PVA) [33, 34]. In order to solve the lack of compatibility between clay layers

and hydrophilic polymers, clay layers are usually organomodified. This treatment

consists of a surface clay modification, which can be achieved through a cation

exchange process in which sodium and calcium cations present in the clay

galleries are replaced with alkylammonium or alkylphosphonium (onium) cations

[1–4, 12, 14, 21]. The ammonium cations are referred to as a “surfactant” owing

to their amphiphilic nature. The exact effect of the surfactant concentration

on the final nanocomposite structure is still a burning issue because of the

nature of the matrix. Some researchers believe that surfactant quantity higher

than the cation exchange capacity (CEC) of the clay reduces the miscibility

between the polymer and the inorganic mineral during nanohybrid synthesis.

Indeed, they showed that the presence of excess modifier in the organoclay exerts

undesirable effects on nanocomposite properties [35–38] whereas Ratinac et al.

[36] reported that the excess of amphiphilic substance, improves the dispersion

level of silicate layers in bulk-polymerized poly(methyl methacrylate) (PMMA)

nanocomposites.

The main problem in preparing PCNs concerns the thermal stability of the

surfactant in the organoclay. The interlayer alkylammonium surfactants often

decompose at extrusion temperatures by the so-called Hoffmann elimination

leading to a collapse of the silicate layers and limiting the composite properties

[15–20]. Moreover, the aforementioned decomposition products impart unde-

sirable color, odor, and taste to the composite [15–20]. Efforts have been made to

synthesizemore thermally stable organoclays using other cationic surfactants. For

example, pyridinium and quinolinium [39], imidazolium [40], phosphonium [41],

and stibonium [42] salts were used. Organoclays synthesized from these surfac-

tants are more thermally stable, but unfortunately, these surfactants are expensive

and/or requiremultistep synthesis limiting large-scale production. Nowadays, the

related research is mostly fueled by two aspects: the replacement of organophilic

layered silicates with pristine ones in order to avoid this expensive chemical

modification of nanofillers and the development of harmless, environmentally

friendly production methods. The greatest interest has involved melt processing,


especially water-assisted extrusion (using untreated and even organomodified

clays), which is generally considered more economical, more flexible for formu-

lation, and involves compounding and fabrication facilities commonly used in

commercial practice [33, 43–59, 60–77].This point is discussed in Section 8.3.3.2.

8.2.3

Nanocomposites Structure and Characterization

Depending on the way of preparation and of the nature of the components, three

types of composites can be produced when working with clay and polymers

(Figure 8.1) [1–4, 12, 14]. Three main classes of nanocomposites can be obtained

and are referred to as immiscible (conventional or microcomposite), intercalated,

and miscible or exfoliated. In the immiscible structure, the silicate layers are

dispersed in the polymer matrix in the form of particles including tactoids or

aggregates of tactoids as they were in the clay powder because the platelets are

not separated. Intercalated structures are obtained when single polymer chains

are intercalated between the silicate layers with their regular alternation and

laminae [1, 14]. In the miscible or exfoliated hybrids, the clay layers are well

delaminated and individually dispersed in the polymer matrix. In this case, the

order’s structure is lost and the interlayer distance is comparable to the radius

of gyration of the polymer. The exfoliation or delamination configuration is the

desired morphology because it maximizes the polymer–clay interactions by

rending the entire layers’ surface available by the polymer and therefore leads to

the most significant improvements of mechanical properties [1–4, 12, 14].

Layered silicate

Phase separated(microcomposite)

Intercalated(nanocomposite)

(a) (b) (c)

Exfoliated(nanocomposite)

Polymer

Figure 8.1 Scheme of different types of

composite arising from the interaction of

layered silicates and polymers: (a) phase-

separated microcomposite; (b) intercalated

nanocomposite; and (c) exfoliated nanocom-

posite. Reproduced with permission from

Ref. [1] © 2000, Elsevier.


Structural characterization of nanocomposites has primarily centered on two

complementary techniques: X-ray diffraction (XRD) and transmission electron

microscopy (TEM) [3, 14, 59, 64, 66, 68, 69, 72, 73, 75]. Complementary techniques

such as rheology, differential scanning calorimetry (DSC) [19–21, 23, 31, 67, 74,

76, 78], thermogravimetric analysis (TGA) [23, 28, 74, 76], and solid-state nuclear

magnetic resonance (NMR) [79–81] are also used to provide information about

the structural analysis of nanocomposites.

8.3

Nanocomposites Preparation

The preparative methods are divided in three main groups according to the pro-

cessing techniques: Intercalation from solution, in situ polymerization, and melt

compounding. The present work focuses on the melt intercalation, especially the

melt blending with the aid of water.

8.3.1

Intercalation from Solution

This technique is based on using a solvent in which the polymer is soluble and

the filler can be easily dispersed. The final step consists of removing the solvent,

either by vaporization, usually under vacuum, or by precipitation. This method

has been widely used and from a long time in the case of water soluble polymers

like PVA [82, 83], PEO [51, 78, 84, 85, 86–88], epoxy [89], polyethylene glycol

(PEG) [90], PMMA [91], or water insoluble like nylon 6 (PA6) [92] andmany other

nanocomposites especially based on HNTs [93–98]. Aromatic solvents such as

xylene or toluene are generally used to dissolve polyolefin [99–101] and epoxy

[102].

However, from an industrial point of view, this method is not suitable because

of the large amount of organic solvent required, which are always environmentally

unfriendly and economically prohibitive.

8.3.2

In Situ Polymerization

In situ polymerization is a two-step process. The nanoparticles are first dispersed

in a monomer solution and the resulting mixture is polymerized using standard

polymerization methods. The driving force of this technique is the polarity of the

monomers. In situ polymerization was the first method used to synthesize PCNs

based on PA6. The preparation of PA6 nanocomposites by in situ polymerization

has been first reported by the Toyota research group [5, 103, 104]. Further studies

have also shown that PA6 chains are bonded to the silicate layers, which acts as a

crosslinker and restricts their motion [105, 106]. Reichert et al. [107] reported for


PA12 nanocomposites that the structure was found to be partially exfoliated and

intercalated. In the case of PA1012 nanocomposites, Wu et al. [108] obtained an

exfoliated structure.

An appropriate catalyst is needed to polymerize polyolefins between the clay

layers. Ziegler–Natta catalysts can be used but generally metallocene catalysts are

employed because they aremore tunable [109–111]. In the case ofHNTs, only few

polymer–HNT nanocomposites have been prepared via in situ polymerization

[112, 113].

In situ polymerization has proved successful in the preparation of various

polymer-layered silicate nanocomposites, but this technique has also shown

important drawbacks: (i) it is a time-consuming preparation route, the polymer-

ization reaction may take more than 24 h; (ii) the platelets may reaggregate during

subsequent processing steps because exfoliation is not always thermodynamically

stable; and (iii) the process is available only to the resin manufacturer who is able

to dedicate a production line for this purpose [114].

8.3.3

Melt Compounding

Melt blending is the standard preparative method for PCNs. The polymer is

melted and combined with the desired amount of layered silicate in an extruder,

internal, kinetic energy, or continuous mixer. There are several advantages to the

melt route compared to the polymerization route.The process is more economic,

it is better suited for rapid change in formulation and in industry, and it does

not require a polymer production line dedicated to that sole product. At the

same time, melt processing is environmentally benign among all the methods

of preparing polymer–clay nanocomposites because no solvents are required.

The melt process can be divided into static and dynamic processes. Dynamic

intercalation is the more conventional compounding operation performed in a

processing equipment. During shearing, polymer chains diffuse from the bulk

polymer melt into the galleries between the silicate layers, as shown in Figure 8.2.

Thus, several thermoplastic polymers, including polyamide 6 (PA6) [18, 24, 35,

43, 44], poly(ethylene terephthalate) (PET) [40, 77], ethylene-vinyl acetate (EVA)

[50, 115], PA12 [26], PA11 [76], polyolefins [32, 45–49, 52–54, 61, 74], poly(lactic

acid) PLA [38, 116–118], and so on have been used to study nanocomposite

preparation by melt intercalation.

8.3.3.1 Melt Blending of Polymer/Organoclay Nanocomposites

In 1999, Liu et al. [119] broke new ground in applying the melt-compounding

technique for the preparation of a commercially available PA6/organomodified

Montmorillonite (o-MMT)nanocompositeusinga twin-screwextruder. Since that

time, this route has been widely studied during the past years. Cho and coworkers

[120] showed that it was possible to obtain exfoliated nanocomposites compa-

rable to those produced by in situ polymerization, using a twin-screw extruder.

This exfoliation results from the combination of optimal processing conditions


Stacks of silicateplatelets or tactoids

Shearing of platelet stacksleads to smaller tactoids

ShearStress = ηγ

Organoclay particle(− 8 μm)

Shear

Diffusion

Platelets peel apart by combined diffusion/shear process

Shear

Figure 8.2 Mechanism of organoclay dispersion during melt compounding. Reproduced

with permission from Ref. [44] © 2001, Elsevier.

and good polymer organoclay affinity. Dennis et al. [121] prepared PA6 nanocom-

posites using different processing techniques and different processing conditions.

The best delamination was observed with a medium shear intensity extruder and,

from this study, the most suitable extruder was a nonintermeshing one.

As polyolefins are apolar matrix, the use of a compatibilizer containing polar

functions is required in addition to the organoclay. Polypropylene grafted with

maleic anhydride (PP-g-MA) is usually used as compatibilizer [53, 54].

It is worth noting that the exfoliation of clay in a polymer matrix does not only

depend on the processing conditions but also on the structure and nature of the

organic compound used to modify the clay. Fornes et al. [35, 122] made an exten-

sive study on the effect of the organomodified clay structure on the morphology

and properties of PA6 nanocomposites. The authors proposed that these effects

stem from the amount of exposed silicate surface. Alkyl ammonium ions that

cover a large part of the clay surface hinder the “polar polyamide/polar clay” sur-

face favorable interactions. This hypothesis has been further confirmed by the

same group. This concept is supported by the theoretical works of Tanaka and

Goettler [123] and Fermeglia et al. [124] on the prediction of the binding energy

between polyamide chains and organomodified clay. Shorter aliphatic chains pro-

ducemore favorable binding energies with respect to longer ones. Instead of using

alkyl ammonium cation to organomodify the MMT, it is possible to use a new

surfactant obtained through cointercalation of an epoxy resin and an ammonium

salt. Good levels of dispersion were observed in the strong interaction between

the epoxy groups confined in the layers and amide groups of the PA6 [125]. The

dispersion state of HNTs has been enhanced by adding small organic compounds

containing hydrogen-bonding functionalities [126].


8.3.3.2 Melt Blending of Polymer/Pristine Clay Nanocomposites

There are some disadvantages to use organophilic clays. Firstly, the organo-

treatment is expensive and increases the final product cost. Secondly, the thermal

degradation of organophilic clays begins generally at 180 ∘C while the extrusion

temperature for thermoplastic polymers is generally in the range of 200–250 ∘C[15, 17, 19, 20, 127]. The surfactant degradation limits clay dispersion and

consequently the composite properties [15–20]. Moreover, surfactant decom-

position products impart unhealthy and undesirable color, odor, and taste to

the composite [20, 128]. Nowadays, the related research is mostly fueled by two

aspects: the replacement of organophilic layered silicates with pristine ones and

the development of harmless, environmentally friendly production methods.

Furthermore, developing a new compounding technique to exfoliate unmodified

silicate would reduce nanocomposites’ global cost. Starting from pristine clay, it

is possible to obtain nanocomposites either by modifying the clay in situ using

different kinds of additives or using actually unmodified clay.

In Situ Organomodification The preparation of PCNs by in situ organomodifi-

cation does not require pretreatment of the clay simplifying the procedure and

reducing the production cost.

Kato et al. [64] and Alaoui et al. [68] performed the cationic exchange using

alkylammonium directly in the extruder. Kato et al. [64] used a nonstandard

extruder equipped with a long barrel (length : diameter 77 : 1) enabling for long

residence time. They reported fine nanoscale dispersions. The final composites

exhibited mechanical properties comparable to those of equivalent compos-

ites produced by conventional compounding. The in situ organomodification

technique has been applied for the preparation of nanocomposites based on

thermoplastic polymers such as Polypropylene (PP), polystyrene (PS) [42,

64, 129, 130], and PA6 [131, 132]. XRD patterns confirm the in situ organo-

modification of theMMT by the ammonium salts and show that the greater shear

leads to larger peaks and greater interlayer spacing for PP nanocomposites. TEM

pictures show poor clay dispersion, which was confirmed by mechanical testing.

In the case of PA6/clay nanocomposites, the authors reported evidence of the

nanocomposite structure by the TEM and XRD analyses but in which coexist

an interacalated and a delaminated structure [131, 132]. However, the in situ

organomodification does not solve the drawback of the thermal degradation of

the quaternary ammonium salt by Hoffmann elimination.

Melt Compounding without Organomodification: Water-Assisted Extrusion In the

past decade, significant effort has been devoted to the fabrication of poly-

mer/untreated clay nanocomposites by melt-compounding processes. Several

authors have attempted to produce polymer/untreated clay nanocomposites by

taking advantage of the natural affinities between water and untreated clay, using

water as an in situ intercalating/exfoliating agent in the extrusion compounding

process. The literature on water-assisted extrusion processes for the fabrication


of polymer/pristine clay nanocomposites becomes now rather abundant. Two

main methods have been developed.

Direct Water Injection during Extrusion **The first publication concerning a well-

exfoliated non hydro soluble polymer nanocomposites obtained by swelling

MMT with water has been reported by DSM in a patent in 1999 [133]. This

invention focuses on the preparation of PA6 nanocomposites with adding water.

The mixing is carried out in a twin-screw extruder equipped with a water

injection system and a degassing zone (Figure 8.3). The special design of the

screw allows the pressure to increase up to 125 bar in this zone, which prevents

water evaporation. After being mixed with the PA6/pristine montmorillonite

(p-MMT) mixture, water is removed further downstream via a venting gate

before the composite exits from the extruder (Figure 8.3). Exfoliated PA/p-MMT

nanocomposites were obtained using water injection without requiring the

expensive clay organomodification. The patent claims the applicability of the

technique to all polymers containing the polar amide (CONH) function between

repeating aliphatic units. Since this pioneering work, different research groups

reported the aid of water for the elaboration of nanocomposites based mainly on

PA [63, 65, 66, 71, 72] but also on PP [64, 68, 74], styrene-acrylonitrile (SAN) [69,

70], PET [77], and bio-based matrices (e.g., Nylon 11 (PA11)) [76], using p-MMT

or o-MMT or other inorganic nanoparticles (e.g., HNTs, cellulose fibers) [30, 31,

67]. Delamination of p-MMT at the nanoscale was readily achieved in PA6 [63,

65, 66, 71, 72]. Figures 8.4 and 8.5 show TEM images of a PA6/p-MMT composite

with and without water injection. These micrographs prove that the injection of

water in the system during extrusion greatly enhances the extent of exfoliation

and dispersion of p-MMT in the PA6 matrix. The exfoliated morphology and

the interface adhesion between clay and PA6 are explained by polar interactions.

The proposed mechanism, which allows the dispersion and the exfoliation of the

p-MMT into the matrix, is the following: (i) water acts as a swelling agent of

the p-MMT and as a plasticizer of the nylon. It penetrates the clay layers and

PA 6Na+ MMT Injection

of water

Pressure sensor

Thermocouple

Atmospheric degassingVacuum degassing

Melting zone High compression zone Plastification zone

P T

Figure 8.3 Screw configuration of the twin-screw extruder used for the water-assisted

preparation of the nylon 6/MMT nanocomposites. Reproduced with permission from Ref.

[65] © 2006, Wiley.


1 μm 200 nm

71,96 nm

(a) (b)

Figure 8.4 TEM micrograph of PA6/p-MMT (5wt%) nanocomposites prepared by extrusion

at 200 rpm and without water injection (a) 5000× and (b) 16 000×. Reproduced with per-

mission from Ref. [66] © 2007, Elsevier.

1 μm 200 nm

(a) (b)

Figure 8.5 TEM micrograph of PA6/p-MMT (5wt%) nanocomposites prepared by extrusion

at 200 rpm and with water injection (a) 5000× and (b) 16 000×. Reproduced with permis-

sion from Ref. [66] © 2007, Elsevier.

exfoliates them; (ii) under the shear forces in the extruder, the exfoliated platelets

are dispersed in the polymer matrix; and (iii) water is then evacuated from the

extruder and the layers remain dispersed in the matrix because of the favorable

interactions between polar MMT platelets and polar PA6.

Fedullo et al. [65, 66] explained the process and the driving forces, which lead to

exfoliation/intercalation of the p-MMT in a PA6 matrix using the water injection


p-MMT

MMT swellingPA6 injectionH2O

PA6 + H2O at 240 °C, 100 barMiscible system

Diffusion and adsorption

Desorption of thewater molecules

Exfoliated PA6/p-MMTnanocomposites

Increase of theinterlayer distance

Elimination ofShearH2O

of PA6

+

Extrusion

+

Figure 8.6 Schematic description of the dispersion of the p-MMT in a polyamide matrix

during extrusion with the water injection system. Reproduced with permission from Ref. [66]

© 2007, Elsevier.

extrusion process. A schematic description of the model used to explain exfoli-

ation of p-MMT in PA is shown in Figure 8.6. Although this model is described

step by step, in the extruder, everything happens simultaneously.Water is injected

in the molten PA at high temperature and pressure and it plays two roles during

extrusion processing:

1) At the processing, conditions of temperature and pressure water is miscible

with the polyamide. It forms a one-phase system of high polarity and lower

viscosity. The demonstration of the miscibility of PA6 and water has been

confirmed using high-pressure differential scanning calorimeter (HPDSC). A

high-pressure thermogram of neat PA6 and PA6 with the addition of 20% of

water at 80 bar is presented in Figure 8.7. Neat PA6 exhibits a melting peak

(Tm) at 220∘C and, on water addition, Tm shifts down to 157 ∘C. This large

cryoscopic effect witnesses a water PA miscibility as also observed by Vinken

et al. [134], by Wevers et al. [135, 136], by Baschek et al. [137], and more

recently by Charlet et al. [138] and by Stoclet et al. [76] in the case of PA11.

This phenomenon is observed whatever the PA type and even for copolyether

amide [73]. Amore precise comprehensive study has shown that a fixed water

concentration and a minimum applied pressure are required to achieve the

PA6 dissolution [138]. An increase of pressure does not affect significantly

the melting temperature; it has been shown that an increase in the pressure

of 100 bar (10MPa) brings about a change in Tm of less than 2 ∘C. Also, it


5 mW

PA6 at 80 bars

PA6 + 20% water at 80 bars

Water blank at 80 bars

Peak 200.46 °C

Peak 157.93 °C

°C40

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 mn

50 60 70 80 90 100 110 120 130 140 150 160 170 180 190 200 210 220 230 240

Figure 8.7 High-pressure DSC thermogram of polyamide 6 with water. The dot line rep-

resents the thermogram of water at 80 bar. Reproduced with permission from Ref. [66] ©

2007, Elsevier.

has been shown that the conditions to obtain the maximum depression are a

water concentration of 30mass%.

2) Then, water diffuses between the MMT layers and is adsorbed on the sur-

face, resulting in a swelling of the clay and in the increasing of the interlayer

distance.

These two effects allow PA chains to diffuse and adsorb onto the MMT sur-

face. When the PA chains diffuse into the interlayer space, they adsorb on the clay

platelets and desorb water molecules, which are removed through atmospheric

and vacuum degassing sections. The diffusion and the adsorption of PA6 chains

on MMT have been shown thermodynamically favored by Fermeglia et al. [124]

and are explained by a variation of the internal energy. The adsorption of the PA

chains is proven by the change in crystalline structure of the PA6 and by TGA and

WAXS (Wide-angle X-ray scattering) analysis on the extracted platelets.

Moreover, the water did not seem to cause any decrease in themolecular weight

of polyamide [133]. This is explained, on the one hand, by the short contact time

between polyamide melt and MMT platelets and, on the other hand, by a rel-

atively low processing temperature in comparison with the polyamide melting

temperature. Furthermore, the mechanical properties of neat PA with and with-

out water injection are similar (see Table 8.1). This result supports that the water

addition doesn’t cause detrimental PA’s degradation by hydrolysis. In addition, it

appears that the PA6 miscibility with water greatly improves the dispersion of the


Table 8.1 Mechanical properties of the neat PA6, PA6/p-MMT composites obtained with

and without water and PA6/o-MMT [63]. Copyright © 2005, Wiley Periodicals, Inc.

— Water

(gmin−1)

Izod impact

strength

(Jm−1)

Young’s

modulus

(GPa)

Yield

strength

(MPa)

Elongation

at break

(%)

PA6 0 84.5± 4.3 2.16± 0.11 71.1± 0.7 96± 18

4.5 87.8± 2.2 2.21± 0.14 70.1± 0.4 129± 56

PA6/5wt% p-MMT 0 58.3± 6.5 2.55± 0.05 74.1± 0.7 25± 4

4.5 52.9± 4.0 3.18± 0.05 81.7± 0.3 13± 2

9 56.1± 8.1 3.32± 0.03 79.7± 0.9 16± 5

PA6/5wt% o-MMT 0 83.0± 6.9 2.77± 0.07 73.8± 1.1 24± 4

clays into the matrix as shown by Touchaleaume et al. [72] in the case of PA6/

p-MMT nanocomposites and in the case of PA11/p-MMT by Stoclet et al. [76]

(Figure 8.8). It has been shown that an exfoliatedmorphology can be obtained even

for a clay content of 10wt%. The same conclusions have been reported for HNTs

nanocomposites, which highlights the positive effect of injecting water during

extrusion [26, 30, 31].The nanocomposites obtained with this water injection sys-

tem exhibit mechanical properties similar to the nanocomposites prepared with

organoclay. For example, some mechanical properties of the PA6 nanocompos-

ites are synthesized in Table 8.1 [63]. The yield strength and Young’s modulus of

the nanocomposite with water injection are better than those of both neat PA

and PA/p-MMT without water injection. This improvement is directly related to

the better dispersion of the layers in the presence of water. Its Young’s modulus

and yield strength are also enhanced about 20 and 10%, respectively, compared

100 nm 50 nm

(a) (b)

Figure 8.8 TEM observations of PA11 nanocomposites based on pristine clay filled at: (a)

5wt% and (b) 10wt% with water injection. Reproduced with permission from Ref. [76] ©

2013, Wiley Periodicals, Inc.


to PA6/o-MMT. However, the toughness (notched impact strength) and the duc-

tility (elongation at break) of polyamides/p-MMT extruded with water are lower

than those of the neat polyamides. This effect can be explained by the formation

of α-crystals or γ-crystals, respectively, in absence or in presence of water [63, 76].

This route has been used by Stoclet et al. [76] to prepare a 100% green nanocom-

posite composed of the bio-based PA11 and p-MMT with high performances.

Particularly at low shear rate and for a clay content of 10wt%, the Young’s modu-

lus can be doubled as compared with the unfilled polymer (Figure 8.9).The origin

of this mechanical properties enhancement has been explained by a structural

characterization. Indeed, the XRD study showed that there is a strong affinity

between the polymer and the clay platelets confirmed by a change of the crys-

talline structure induced by the presence of the clays. At high shear rate, the effect

is less obvious and is explained by an overheating that probably promotes polymer

degradation. Nevertheless, this increase is noticeable and comparable to the one

reported by Liu et al. in the case of well-dispersed PA6/o-MMT nanocomposites

[125] or for PA11/o-MMT nanocomposites [139].

This phenomenon is not likely to occur in the case of less polar matrices, which

explains the limited success of this method with such matrices. For example,

Rousseaux et al. [74] injected water in an extrusion stream composed of PP-g-MA

and p-MMT but did not observe any enhancement in dispersion compared with

conventional compounding. They attributed this result to a reaggregation of the

clay on water removal, because of the high difference of polarity between p-MMT

00

200

400

600

800

1000

1200

1400

200 rpm

1000 rpm

200 rpm + water

1000 rpm + water

1600

1800

5 10

Clay content (%)

Yo

un

g’s

mo

du

lus (

MP

a)

Figure 8.9 Evolution of Young’s modulus as a function of clay conditions for PA11

nanocomposites elaborated under different conditions. Reproduced with permission from

Ref. [76] © 2013, Wiley Periodicals, Inc.


and PP-g-MA. In opposite, Lecouvet et al. [30] showed that water injection

enables to improve considerably the clay dispersion.

As mentioned earlier, Kato et al. [64] and Alaoui et al. [68] performed the

cationic exchange with alkyl ammonium salts directly in the extruder with

injected water. However, although the authors show the feasibility of performing

the cationic exchange directly in the extruder, their method does not eliminate the

need for an organicmodifier [64, 68].Themethod developed by Korbee et al. from

DSM [133] was also applied to the production of polymer nanocomposites based

on organophilic MMT. Rousseaux et al. [74] showed that injection of water in a

PP-g-MA stream enhanced the dispersion of methyl tallow bis-2-hydroxyethyl

modified MMT (Cloisite 30B®) compared with conventional compounding.They

attributed this result to the occurrence of an esterification reaction between

the hydroxyl functions of the intercalating agent and the carboxyl functions of

PP-g-MA. Again, although this result is interesting, it does not eliminate the

need of an intercalating agent.

Slurry Route In 2003, Hasegawa et al. [140] reported a compounding process for

the preparation of PA6 nanocomposites, using p-MMT water slurry as an alter-

native for o-MMT. If the use of water is similar to the DSM patent, in this pro-

cess, a suspension of p-MMT in water is injected within molten PA6 using an

extruder, followed by removing the water as shown in Figure 8.10. TEM micro-

graphs show that the mixture obtained with this clay slurry is exfoliated although

a few silicate stacks are still visible. The XRD pattern shows that the formation

Clay slurry

Vapor

Nylon 6pellet

Melting zone Clay compounding zone

NCH

Figure 8.10 Schematic figure depicting the compounding process for preparing the

nylon6/clay hybrid-clay slurry using the clay slurry. Reproduced with permission from Ref.

[140] © 2003, Elsevier.


of a slurry enhances the dispersion of the clay platelets in comparison with a dry

compounded composite.

Themechanism of dispersion of the clay platelets in PA6 proposed byHasegawa

et al. is explained in Figure 8.11. According to these authors, the exfoliation of sil-

icate layers into the matrix occurs as follows: (i) the clay slurry is first pumped

into the melting polyamide under vigorous shear (Figure 8.11a); (ii) during blend-

ing, the slurry drops become smaller and the water begins to evaporate following

contact with melting PA (Figure 8.11b,c); and (iii) the vacuum vent evacuates the

evaporated water. During evaporation, the platelets are dispersed into PA melt as

monolayer or as a few layers (Figure 8.11d). Also, epoxy/clay have been success-

fully prepared by this route [102].

The proposed mechanism of dispersion of layered silicates in this process

is quite different from the one of conventional compounding process using

organophilic clay. When using organoclay, the polymer chains first intercalate

into the stacked silicate galleries and then clay layers are exfoliated. In the slurry

process, the silicate layers pre-exfoliated in water (slurry) are directly fed into

the polymer matrix. No reaggregation of the silicate layers is claimed. Some

injection and compression molding experiments showed that the dispersion of

the silicates platelets was stable. At small clay loading, the strength, modulus, and

heat distortion temperature of the nanocomposite were much higher than those

of neat PA, but the impact strength was lower. Although the MMT layers are well

dispersed in the matrix, the final properties of these nanocomposites are nearly

equal to those of conventional PA6 o-MMT nanocomposites.

Clay slurry(a) (b)

(c) (d)

Nylon 6 Silicate layer

Figure 8.11 (a–d) Schematic figure depicting the dispersion of the Na-montmorillonite sil-

icate layers of the slurry into nylon 6 during compounding. Reproduced with permission

from Ref. [140] © 2003, Elsevier.


Unfortunately, the large water/clay ratio, the very low throughput of 2 kg h−1,

and the very high residence time of 10min make this technique unsuitable for

industrial applications. Recently, Stoeffler et al. [75] have carried out a work

that aims at comparing the efficiency of three melt-compounding methods

for preparing PA12/untreated clay composites. They confirm the potential of

water-assisted extrusion processes to significantly enhance the dispersion state in

PA12/p-MMT composites. However, despite a drastic reduction in claymicropar-

ticles size, improvements in mechanical properties remained limited, because of

the fact that clay remained mainly dispersed at the microscale. Therefore, they

concluded that the use of water-assisted processes should be restricted to the

preparation of composites based on untreated inorganic particles, devoted to

applications for which the presence of an organic intercalating agent is highly

undesirable.

8.4

Nanocomposite Properties

Polymer-layered silicate nanocomposites exhibit many advantages, especially the

exceptional improvements in properties at lower filler concentrations compared

to conventional micro- and macrocomposites [2]. Indeed, numerous studies have

demonstrated that even a small percent of layered silicate can lead to a wide array

of property enhancement, for example, increased stiffness and strength, enhanced

solvent and UV resistance, improved gas barrier properties, and superior flame

retardancy [1–3, 12–14, 26, 30, 31]. In this section, only the thermal properties

are discussed.

8.4.1

Thermal Stability

The thermal stability of polymeric materials is usually studied by TGA. Gener-

ally, higher thermal stability is commonly observed when clay platelets are well-

dispersed into the polymer matrix, as a result of limited oxygen supply, hindered

out-diffusion of the volatile decomposition products, and char formation that acts

as physical barrier between the polymer and the oxidative medium. In general, it

has been reported that the PCNs are thermally more stable than pure polymers

[1, 2, 4, 12, 23, 28, 74, 76, 141].

Moreover, despite the general improvement of thermal stability, decreases in

the thermal stability of polymers on nanocomposite formation have also been

reported, and various mechanisms have been put forward to explain the results.

It has been argued, for example, that after the early stages of thermal decompo-

sition, the stacked silicate layers could hold accumulated heat, acting as a heat

source to accelerate the decomposition process, in conjunction with the heat flow

supplied by the outside heat source [55]. Also, for polymers that require high melt

processing temperatures, the thermal stability of the organic component of the


modified clay, often alkyl ammonium cations, could suffer from decomposition

following the Hoffman elimination reaction. Indeed, the thermal degradation of

organophilic clays begins generally at 180 ∘C [15, 17, 19, 20, 127]. During the melt

compounding, surfactants moieties migrate outside of the silicate and influence

the properties of the materials negatively [7]. Because of its degradation at low

temperature, the organomodifier catalyzes the thermal degradation of the poly-

mer matrix [15–19]; it plasticizes the polymer matrix [35–38] or diffuses out of

the material, causing undesirable color, odor, and taste to the composite [20]. As

an alternative to the use of organomodified clays, recent studies have pointed out,

as discussed earlier, the use of raw clay with the aid of water (an in situ intercalat-

ing/exfoliating processing aid).

By using this molding process, Yu et al. [63], Dasari et al. [43], and

Touchaleaume et al. [72] have shown that the PA/p-MMT with water assis-

tance unlike PA/o-MMT presents a high thermal stability by TGA. Figure 8.12

shows the degradation onsets (T−5%) of PA6 nanocomposites and neat extruded

PA6. The degradation temperature measured for the composite with water

injection is higher than the one without water and the one of pure PA6 and

polyamide-6/o-MMT nanocomposites. The PA6/p-MMT thermal stability

improvement, achieved thanks to water injection, is ascribed to better clay

dispersion, which acts as a barrier to gas permeation. This barrier effect is

reinforced through the formation of a protective inorganic layer [26, 30–32].

Also, Touchaleaume et al. [72] have demonstrated in the case of PA6 that

using water during melt compounding helps to prevent matrix degradation by

decreasing the processing temperature following the cryoscopic effect, which

allows decreasing the extrusion temperature (down to 190 ∘C) below the neat

PA6 melting temperature (256 ∘C), thus limiting the thermal degradation. The

water also contributes to limit the degradation because of its lubricating effect.

0365

370

375

380

385

390

395

T −

5%

(°C

)

2 4 6 8 10

Mineral content (wt%)

NaMMT + water

OMMT + water

OMMT

NaMMT

Figure 8.12 Onset of thermal degradation of PA6/clay nanocomposites. Reproduced with

permission from Ref. [72] © 2011, Elsevier.


8.4.2

Flame Retardancy

Because of the large use of polymers especially in domestic applications and

in order to make them more safer, it is necessary to reduce their potential for

ignition or burn. Traditionally, flame retardancy has been achieved either by using

intrinsically FR polymers (fluoropolymers or (polyvinylchloride) PVC), or by

incorporating FRs (such as aluminum trihydrate, magnesium hydroxide, organic

brominated compounds) or intumescent systems. However, such FRs exhibit sig-

nificant drawbacks. For example, aluminum trihydrate andmagnesium hydroxide

need to be applied at very high loadings to be effective. Use of halogenated addi-

tives is increasingly contested because of their potential toxicity related to the

release of corrosive and toxic decomposition products such as halogen acids and

metal halides [142, 143]. Considering also these disadvantages of traditional FRs,

it becomes obvious that it is necessary to develop novel synergistic FR systems

with high efficiency and acceptable environmental impact. Therefore, recently,

extensive research in the field of polymer nanocomposites have demonstrated

that clay nanoparticles can contribute to the reduction of the flammability of

polymeric materials and so could be a potential alternative to conventional FRs

[26, 30–32]. The three most widely used tests to measure the flame retardancy of

polymeric materials are the cone clorimeter, the limiting oxygen index (LOI), and

the UL-94 vertical burning test.

Cone calorimetry is one of the most effective medium-sized fire tests used to

investigate the flame retardancy of polymeric materials under forced flaming con-

ditions [30–32, 144]. This technique is based on the oxygen consumption princi-

ple with the assumption that there is a constant relationship between the mass of

oxygen consumed from the air and the amount of heat release during the combus-

tion of a material subjected to a given heat flux.The most important flammability

parameters are heat release rate (HRR), peak of heat release rate (pHRR), total

heat release (THR), peak of mass loss rate (pMLR), time to ignition (TTI), and

mass fraction of the residue.

Pioneering studies by Gilman et al. [145–147] have reported large reductions

in flame spread and prolonged burning times of polymer-clay nanocomposites

in cone calorimetry. For instance, a decrease of 63% can be achieved with only

5wt% of nanodispersed MMT in PA6 (Figure 8.13) [145]. Similarly, Bourbigot

and coworkers have shown that the pHRR of PLA is lowered by about 40% when

adding 4wt% layered silicates [148].Themain mode of action of the clay minerals

as fire retardants is based on a physical mechanism, that is to say barrier effects

arising from charring reaction and the accumulation of minerals at the sample

surface, which reduces the heat and mass transfers between gas and condensed

phases. The incorporation of a relatively low amount of nanoclay in a polymer

matrix enables the formation of a protective ceramic char-inorganic layer at the

sample surface during the combustion process [26, 30, 145].

This mechanism has been put forward in most studies reporting on the FR

properties of nanocomposites. Also, it is believed that the excess quaternary


00

200

400

600

800

He

at

rele

ase

ra

te k

W m

−21000 Peak HRR: 1011 kW m

−2

Peak HRR: 361 KW m−2

Heat flux: 35 KW m−2

Nylon-6Nylon-6 silicatenanocomposite - 5%

1200

500 1000 1500

Time (s)

2000

Figure 8.13 Comparison of the Heat Release Rate (HRR) plot for nylon-6 silicate-

nanocomposite (mass fraction 5%) at 35 kWm−2 heat flux, showing a 63% reduction in HRR

for the nanocomposite. Reproduced with permission from Ref. [145] © 1999, Elsevier.

ammonium surfactants used to disperse the clays also increase the probability

of early ignition [26, 30, 145]. As discussed previously, using the original melt

compounding with water injection can overcome the organomodification and

helps to form a well-dispersed nanoclay morphology. This novel processing

route enables also to enhance the fire retardancy. For example, in the case of

polymer-HNTs nanocomposites, Lecouvet et al. [26, 30, 31] showed for PP/HNTs

[30], polyethersulfone/HNTs [31], and for polyamide12/HNTs [20] prepared by

water-assisted extrusion that these nanocomposites exhibit the lowest flamma-

bility because of the well-exfoliated structure and the formation of a strong and

cohesive char acting as a barrier to heat and mass transfer [30]. This is illustrated

in Figure 8.14.

8.5

Toward Fully Green Composites?

Green composites are a specific class of biocomposites, where a bio-based poly-

mer matrix is reinforced by natural fibers (cellulose, starch, lignin, … ), and they

represent an emerging area in polymer science. Increasing environmental aware-

ness and lower material costs are the main driving forces for using renewable

materials, such as wood and cellulose fibers, as reinforcement in polymer com-

posites. A very appealing characteristic of natural fibers is also the fact that they

can be considered as carbon dioxide neutral materials, that is, they do not release

excess carbon dioxide into the atmosphere when composted or combusted [149].

8.5 Toward Fully Green Composites? 199

0

0

100

200

300

400

500

600

700

100 200 300

Time (s)

Heat re

lease r

ate

(kW

m−2

)

400

PP-PIb

PP-PIb-H4-W

PP-PIb-H8-W

PP-PIb-H16-W

500

Figure 8.14 Heat release rate curves as a function of HNTs (H) loading for reference and

PP/HNTs nanocomposites prepared by water-assisted extrusion (W) with PP-g-MA (Plb).

Reproduced with permission from Ref. [30] © 2011, Elsevier.

Othermain advantages of natural fibers over the synthetic and inorganic fibers tra-

ditionally used as reinforcement in polymeric composites are low specific weight,

high specific strength, and stiffness, safer handling and working conditions, and

nonabrasiveness to the processing equipment [149]. When the used reinforce-

ment is in nanoscale, such as nanocellulose, even further improvements in the

composite properties can be obtained.

So far, research in the field of biocomposites has focused strongly on the mate-

rials used in the composite preparation, and not so much on the processing of the

materials, even though the latter has a large impact on the final properties of the

material. One of the most commonly used methods in the industrial processing

of polymer composites is the extrusion process.

However, there are also some problems related to the use of natural fibers in

thermoplastic composites. Major drawbacks are the poor interfacial adhesion

between the hydrophobic matrix and the hydrophilic fibers and the difficult

dispersion of fibers in the matrix. Because of the hydrophilic nature of natural

fibers, they need to be dried before the compounding step. In addition, the

processing temperatures are limited because of the thermal degradation of

the natural organic fibers (mainly wood constituent) in temperatures above

200 ∘C, hence limiting the range of suitable matrix polymers. Because of the

low thermal stability of wood, only polymers with processing temperatures

lower than 200 ∘C are typically used in wood polymer composites (WPCs). The

polymers used are mostly low-cost commodity thermoplastics and the most

common polymers used in WPCs are polyethylene (PE), PP, and PVC. Despite

this, any temperature overshoot, even localized, during the processing leads to


an undesirable brownish discoloration often together with an unpleasant odor

because of pyrolytic degradation of the cellulosic fibers [67]. This problem get

worse with the increase of shear. Additives like coupling agents, light stabilizers,

pigments, lubricants, fungicides, and foaming agents are usually added to reduce

this discoloration via reduced frictional degradation during compounding, but

they are generally considered to be detrimental to adhesion [138, 139] and then

to mechanical properties. Soulestin and coworkers [67] investigated the possible

extension to cellulosic nanofibers of an original process with water injection

extrusion developed for nanoclays (described previously). Indeed, water is

known to be efficient for removing organic components such as flavor volatiles,

surfactant moieties by steam flushing during extrusion processes [150]. When

water is injected, the volatile moieties of the extrusion-degraded surfactant are

steam stripped out of the processing media.The well-known flushing mechanism

was reported for polyether-block-amide (PEBA)/o-MMT nanocomposites [20,

73], and quantified thanks to odors and volatile organic compounds emissions

analyses [20]. In the case of low density polyethylene (LDPE)/cellulose fibers

composites, Soulestin et al. have shown also that the injection of water during

extrusion is really effective in improving dispersion of cellulose fibers and enables

to achieve a complete disaggregation of cellulose clusters in addition to classical

ways. Moreover, this technology allows to reduce significantly the yellowing of

these composites, as can be seen in Figure 8.15. An odor stabilization has also

300 600 900 1200 300 600 900 1200

173 205 225 240 – 205 – –

0 0 0 0 10 10 10 10

Sample N°

Screw speed (rpm)

Bulk T ° ° (°C)

Water injected (%)

1

1 2 4

5 6 7 8

3

2 3 4 5 6 7 8

Figure 8.15 Picture of the LDPE/5% cellulose fiber composites produced without

(sample 1–4) and with water injection (sample 5–8) with experimental parameters used for

the preparation of the LDPE/cellulose composites.

References 201

been reported when water is used. Presence of water acts as a lubricant, plasti-

cizer reducing the thermal degradation of the cellulose fibers during the severe

processing conditions but without releasing significant amounts of nanofibers.

This work presents a clear indication that this novel processing route may

be suggested to prepare polyolefins-biofiber composites. Also, this technology

offers the opportunity of technology transfer of nanocomposites’ “odor” solutions

toward biocomposites based on natural fibers.

References


Polymer-layered silicate nanocompos-

ites: preparation, properties and uses

of a new class of materials. Mater. Sci.

Eng., R, 28, 1–63.2. Ray, S.S. and Okamoto, M. (2003) Poly-

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211

9

In Situ Preparation of Conducting Polymer Nanocomposites

Liping Yang, Cher Ling Toh, and Xuehong Lu

9.1

Introduction

Nanocomposites are multiphase materials in which at least one of the con-

stituents has a dimension in the nanometer scale (<100 nm) [1, 2]. Polymer

nanocomposites have garnered tremendous academic and industrial interests

in the past few decades because they are lightweight and yet exhibit superior

mechanical or thermomechanical properties to their neat polymer counterparts

[3–6]. These lightweight nanocomposites may also provide some functional

properties, such as useful electrical or magnetic properties. An important class

of polymer nanocomposites is electrically conductive polymer nanocomposites,

which can offer a wide range of electrical conductivities similar to those of tradi-

tional inorganic conductors or semiconductors, and hence can be used in a wide

variety of applications that require lightweight or flexible conducting materials.

Electrically conductive polymer nanocomposites can be divided into two major

categories, namely those composed of intrinsically conducting polymers and

that achieved by the addition of conducting nanofillers into insulating polymer

matrices. The nanocomposites with conducting polymer matrices can provide

some unique functional properties, for example, their electrical and optical prop-

erties may change in response to external stimuli. However, in general, they are

more expensive and exhibit poorer mechanical properties, heat resistance, and

environmental stabilities than the conducting nanocomposites with insulating

polymer matrices. Currently, there is an extensive list of comprehensive articles

on the nanocomposites with conducting polymer matrices [7, 8]. Thus, in this

chapter, only the conducting nanocomposites with insulating polymer matrices

are discussed.

Overall, this chapter is focused on the conducting nanocomposites (with insu-

lating polymer matrices) prepared via in situ polymerization or in situ processing

methods. Although thermally conductive nanocomposites with insulating poly-

mer matrices may be prepared in similar ways, they are not discussed in this

chapter in order to stay focused.The emphasis of this chapter is to summarize the


212 9 In Situ Preparation of Conducting Polymer Nanocomposites

recent progress in in situpreparation of electrically conductive polymer nanocom-

posites as well as to discuss structure–property relationships of the nanocompos-

ites and underline the benefits of the in situ processes. In the following parts of this

section, a general introduction to electrically conductive polymer nanocompos-

ites is given. Section 9.2 is divided into two subsections; Section 9.2.1 is focused

on in situ polymerization of polymers in the presence of conductive nanofillers,

while Section 9.2.2 is devoted to the cases where the conductivities of the dis-

persed phases are provoked or enhanced in the nanocomposite formation pro-

cesses. Finally, a future outlook for in situ preparation of electrically conductive

polymer nanocomposites are presented.

9.1.1

Electrically Conductive Polymer Nanocomposites and Their Applications

In comparison with metal conductors, the advantages of conducting polymer

nanocomposites include their mechanical flexibility, lightweight, ease of man-

ufacturing, a wider range of electrical conductivities, corrosion resistance, and

relatively low cost [9, 10]. Conducting polymer nanocomposites therefore have a

broad application spectrum. Both thermoplastics, such as polyamide, polyester,

polyolefin, polyurethane (PU), vinyl and acrylic polymers, and thermosets, such

as epoxy, polyimide (PI), and rubber, have been used as matrices in electrically

conductive nanocomposites [11–14]. They are used in sensors, catalysis, energy

devices, conducting paints, conducting adhesives in electronic components and

as materials for applications like microwave absorption, thermal management,

electromagnetic interference (EMI) shielding, and electrostatic charge dissipation

(ESD) [13, 15–17]. In general, the applications for electrically conductive polymer

nanocomposites depend on the magnitude of their volume electrical resistivity as

illustrated in Figure 9.1.

10−5 10−3 10−1 10

Meta

l

EM

I sh

ield

ing

Sta

tic d

issi

pativ

eA

ntis

tatic

Pla

stic

103 105 107 109 1011 1013 1015 1017

Electrical resistivity, Ω sq−2

Conduct

ive

Figure 9.1 Electrically conductive materials with different ranges of surface resistivities and

their applications.


9.1.2

Percolation Theory

Most polymers are generally insulating. With a very low volume fraction of con-

ductive fillers, the conductivity of a typical polymer nanocomposite is similar to

that of the insulating polymermatrix as the conductive fillers are well separated by

the insulating matrix. However, the nanocomposite exhibits a transition from an

insulator to a conductor when a critical amount of conductive fillers necessary to

form a continuous conducting network is added. This critical amount of fillers is

often referred to as the percolation threshold, and it is specific for each system

[18]. At the percolation threshold, the electrical conductivity of the nanocom-

posite increases drastically, while further increment in the filler loading often has

little effect on the electrical conductivity of the nanocomposite [19], as depicted

in Figure 9.2.

A power law fit as given in Equation 9.1 derived from the percolation theory is

often used to fit resistivity at filler loading above the percolation threshold,

𝜎 ∝ (𝜑 − 𝜑c)𝑡 (9.1)

where 𝜎 is the electrical conductivity of the polymer nanocomposite; 𝜑, the filler

volume fraction;𝜑c, the filler critical volume fraction at the percolation threshold;

and t, the critical exponent related to the lattice dimensionality.

The proposed mechanisms of electric conduction in the nanocomposites

mainly include simple interaggregation conduction and tunneling of electrons;

both are dependent on the type of conductive fillers, distribution of the fillers, and

their contents [20]. When the conductive fillers form a continuous conductive

network in a bulk sample via direct physical contacts between the conductive

Log σ

Filler volume fraction (φ)

φc

Figure 9.2 A schematic depicting variation of electrical conductivity as a function of filler

volume fraction in polymer nanocomposite, where 𝜑c represents the critical filler volume

fraction at percolation threshold.


fillers or aggregates, electrons will be able to travel from one end to the other

under an applied electric field. However, electric conduction need not take place

solely through direct contact between the fillers. In fact, when the distance

between the fillers is small enough, that is, within a few nanometers, an emission

current is able to flow when a high electric field is applied. This is known as

the tunneling effect, meaning that electrons are able to tunnel through energy

barriers between the conductive fillers [21, 22]. It is highlighted that in the case

where tunneling of electrons between adjacent fillers also occurs, the possible

percolation threshold is lower than that of the system, which is based on the

formation of conductive paths among fillers in direct contact [23].

9.1.3

Factors Affecting the Electrical Conductivity of Nanocomposites

Generally, the conductive filler content should be as low as possible so that the pro-

cessing, mechanical properties, and cost would not become an issue. Yet, at the

same time, the filler content needs to be high enough to fulfill the prerequisite of

forming a continuous conductive network. It is therefore ideal to achieve conduct-

ing nanocomposites with a low percolation threshold [15].Thepercolation thresh-

old and electrical conductivity of a conducting polymer nanocomposite depends

on many factors, amongst which the most important ones are (i) physical proper-

ties of the fillers, including their aspect ratios, electrical conductivities, and surface

characteristics, (ii) filler dispersion and distribution, (iii) physical properties of the

polymer matrix, such as viscosity, crystallinity, and polarity, and (iv) filler orienta-

tion and alignment. It is worth noting that many of these factors are significantly

influenced by nanocomposite fabrication methods and conditions. Therefore, the

influences of these factors on the electrical conductivity of nanocomposites are

briefly introduced here, providing a basis for the detailed discussion on in situ

formation of conducting polymer nanocomposites in Section 9.2.

9.1.3.1 Physical Properties of the Fillers

Conductive fillers that had been incorporated into insulating polymer matrices

include nonmetallic fillers, mainly carbon-based materials, and metallic fillers,

such as Ag- and Pd-based nanoparticles. Conductive behaviors of polymer

nanocomposites depend heavily on the physical properties of the conductive

fillers (Table 9.1) [15, 24].

On the basis of geometric considerations, the percolation threshold is strongly

influenced by the shape and aspect ratio of the fillers (Figure 9.3a) [15]. Consider-

ing a filled polymer systemwith a specific filler orientation, the percolation thresh-

old generally decreaseswith increasing aspect ratios (length to diameter/thickness

ratio) of the fillers, as shown in Figure 9.3b.

In comparison with carbon fillers, most metallic fillers have a tendency to oxi-

dize and form an insulating surface layer, adversely affecting the electrical con-

ductivity of the nanocomposites. By contrast, carbon fillers have a tendency to

aggregate and aid the formation of conductive networks. Common conductive


Table 9.1 Typical physical properties of carbon nanofillers [15, 24].

— Carbon

nanotube

Graphene Graphite Carbon

nanofiber

Carbon

black

Aspect ratio 102–104 102–104 — 200–2000 —

Density (g cm−3) 1.3–1.75 2.2 2.2 2 1.8–1.85

Electrical conductivity (Sm−1) 105–106 107–108 105 106 103

Tensile modulus (GPa) 300–1000 1000 — 240 —

Tensile strength (GPa) 150–180 130 — 2.92 —

Zero-dimension (0 D)

Two-dimension (2 D)

TL

D

L

Aspect ratio = L/T (L/D)

One-dimension (1 D)

0.11

101

102

103

104

1 10 100

Fib

er

asp

ect

ratio

Volume fraction of conductive filler(a) (b)

Figure 9.3 (a) Schematic illustration of con-

ductive fillers with different aspect ratios and

dimensions (L: length, T: thickness, D: diam-

eter) and (b) effect of filler aspect ratio on

the critical filler volume fraction needed to

induce bulk conductivity in a polymer filled

with conductive fibers.

carbon fillers include carbon black (CB), carbon nanofibers (CNFs), carbon nan-

otubes (CNTs), graphite, and graphene. The presence of these different carbon

forms are attributed to carbon’s unique hybridization capability, that is, it is able

to achieve different hybridize states, resulting in various molecular structures. CB

consists of a mixture of sp2 and sp3 hybridized carbon atoms, while other carbon

fillers including CNTs, CNFs, graphite, and graphenemainly consist of sp2 honey-

comb structures. Metallic powders and CB require high filler contents to achieve

satisfactory electrical conductivity, which consequently leads to poorer mechan-

ical properties and high density in the case of metallic powders. Compared with

these traditional fillers, graphite, graphene, CNTs, and CNFs have higher aspect

ratios and thus able to afford superior electrical properties at lower filler load-

ings. Numerous results have demonstrated the influences of shape and aspect

ratio of these conductive fillers on both mechanical and electrical properties of

the nanocomposites, which are elaborated in Section 9.2.


Compared with CNFs, CNTs have better mechanical and electrical properties,

smaller diameters, and higher aspect ratios (Table 9.1). However, CNTs are much

more expensive and they are difficult to be dispersed into individual tube in poly-

mer matrices owing to the ultrastrong attracting force between the single tubes or

entanglement of the CNTs.Themorphologies and properties of different types of

CNT- and CNF-based conducting polymer nanocomposites have been well pre-

sented in recent review papers [10–12, 15].

Among the carbon fillers, graphene, which is in the form of two-dimensional

(2D) nanosheets, has the potential to afford the corresponding nanocomposites

lower percolation threshold than the other carbon fillers owing to its extremely

high electrical conductivity given by its planar sp2 structure and high aspect ratio

brought by its single-layer structure [25]. Practically it is vital to attain exfoliated

graphite consisting of very thin stacks of parallel graphene sheets as high aspect

ratio of the filler is required for both superior electrical and mechanical proper-

ties. In particular, the method utilizing graphene oxide (GO) is attractive because

of its ease of preparation and dispersion [26]. However, conductivity of GO is

much lower than that of graphene and hence only reduced graphene oxide (rGO)

can effectively act as conducting fillers in polymer nanocomposites. As attain-

ing highly conductive fillers and well-dispersed state of the conductive fillers in

polymermatrices are both very important for achieving satisfactory electrical con-

ductivities, in situ processes become attractive approaches, which are explained

in detail in Section 9.2.

9.1.3.2 Filler Distribution and Dispersion

Polymer nanocomposites are achieved when fillers in the polymer matrices are

dispersed into individual entity with dimension(s) in nanometer scale. Gener-

ally, the nanoscale dispersion of the fillers in polymer matrices would result in a

large aspect ratio of individual filler and high interfacial area for intensive poly-

mer/filler interactions. In terms of electrical properties, well-dispersed filler is

essential in order to build a conductive network and achieve reduced percolation

threshold, whilewell-distributed filler is not necessarily required, as demonstrated

in Figure 9.4. It shows that with poor filler distribution, it is still possible to form

a conductive network with well-dispersed filler (Figure 9.4c). However, it is evi-

dent that the formation of agglomerates (i.e., poor dispersion) (Figure 9.4a,b) hin-

ders the formation of the conductive network. Therefore, it is paramount that the

conductive fillers are well dispersed in polymer matrices [15]. Yet, attaining well-

dispersed filler in polymer matrices is often challenging because of the incom-

patibility of the filler and the polymer. Various dispersion strategies used in in situ

preparation of conducting polymer nanocomposites are introduced in Section 9.2.

9.1.3.3 Physical Properties of Polymer Matrices

From the brief discussion above, it is clear that improving the dispersion of con-

ductive fillers in a polymer matrix can consequently lead to reduced percolation

threshold. Hence, depending on the type of polymers, which may have different

surface tension, viscosity, crystallinity, and polarity, electrical conductivity and


(a) (b)

(c) (d)

Figure 9.4 Schematics showing the effect of 1D filler on the conductivity of polymer

nanocomposites: (a) poor distribution, poor dispersion; (b) good distribution, poor disper-

sion; (c) poor distribution, good dispersion; and (d) good distribution, good dispersion.

percolation thresholds of the nanocomposites could be very different with the

same type of conductive fillers. For example, a low interfacial tension between a

polymer and filler means that the polymermatrix can better wet the filler, yielding

better dispersion of the fillers. This would lead to reduced percolation threshold.

Mechanical, electrical, and thermal properties of the nanocomposites prepared

from a low-viscosity epoxy precursor have been reported to be better than that

from a high-viscosity one [27]. Likewise, increasing polymer polarity can lead to

better interactions between some conductive fillers and the polymer, hence reduc-

ing percolation threshold [28]. In addition, lower percolation threshold can be

achieved more easily in semicrystalline polymers than in amorphous polymers

[29].This is because during crystallization, the fillers are concentrated in the amor-

phous phase, giving to a higher effective concentration for the formation of a

conductive network.

9.1.3.4 Filler Orientation and Alignment

The extent of filler orientation and alignment can affect the connectivity of

the fillers and hence has an effect on the percolation threshold. Orientation of

CNTs was mostly studied. For example, an alignment of single-walled carbon

nanotubes (SWCNTs) in petroleum pitch matrix leads to a 3.4-fold increment

on electrical conductivity with 5wt% SWCNTs loading [30]. Aligned multiwalled

carbon nanotubes (MWCNTs) from an injection CVD process were dispersed

as conductive fillers in an epoxy matrix. The resulting electrical properties

show that sufficient conductivity for antistatic applications can be achieved

at a very low nanotube loading of approximately 0.005wt%. The use of the

aligned MWCNTs leads to a uniquely low percolation threshold, which is an


order of magnitude smaller than best results previously achieved with entangled

MWCNTs [31]. Posttreatments of the nanocomposites can also influence the filler

orientation and hence percolation threshold. For instance, Haggenmueller et al.

prepared (poly(methyl methacrylate)) PMMA/SWCNTnanocomposite films and

fibers bymelt processing [32].They found that the composite films showed higher

conductivity along the flow direction than perpendicular to it. On the contrary,

hot pressing can destroy the percolation network even after it is formed [19, 33,

34]. It is worth noting that the filler orientation and alignment may be altered in

the in situ formation processes of conducting polymer nanocomposites. This are

discussed in Section 9.2.

9.1.3.5 Nanocomposite Fabrication Methods and Conditions

Three strategies have been widely applied in the incorporation of conductive

fillers into polymer matrices to attain a satisfactory dispersion of the fillers. They

are namely (i) solution blending, (ii) melt blending, and (iii) in situ polymeriza-

tion. The nanocomposite fabrication methods and process conditions can affect

not only the dispersion states of the fillers, but also the orientation/alignment,

aspect ratio of the fillers, the filler/matrix interface and in some cases elec-

trical conductivities of the fillers. Therefore, they influence the properties of

the nanocomposites significantly. To acquire a lower percolation threshold,

solution blending and in situ polymerization are much more effective than melt

compounding [15, 35, 36] (Figure 9.5). Nonetheless, solution blending is always

limited by economic and environmental cost, and for some polymers, solvents

could not work well. Although high-shear blending processing methods are

able to produce nanocomposites with well-dispersed fillers, the aspect ratios

of the fillers are often reduced with intensive/prolonged shearing, ultimately

adversely affecting the mechanical properties and conductivity. Furthermore,

high-shear force may lead to reduced tendency to form conductive paths in the

0.000 0.005

101

102

103

104

105

106

107

108

109

1010

1011

1012

0.010 0.015 0.020 0.025 0.030 0

(b)(a)

1E+00

1E+03

1E+06

1E+09

1E+12

1E+15

1 2 3 4 5 6 7 8 9 10

Graphite, melt

Graphite, melt

Re

sis

tan

ce

(Ω)

0.0010

2

104

106

108

1010

1012

0.05 0.10 0.15

Filler volume fraction

TRG, melt

TRG, solvent

in situ polymTRG

Filler volume fraction Wt% CNF

Melt-compounding

Casting

In situ

Su

rfa

ce

re

sis

tan

ce

(Ω

)

Su

rfa

ce

re

sis

tivity (

Ω s

q−1

)

Figure 9.5 (a) Surface resistance of melt-

blended graphite/thermoplastic polyurethane

(TPU) nanocomposites, and melt-blended,

solution-blended, and in situ polymer-

ized TPU nanocomposites with thermally

reduced GO (TRG). (b) Surface resistivity ver-

sus filler content for melt-blended, solution-

blended, and in situ polymerized poly(methyl

methacrylate) (PMMA)/CNF nanocomposites.


polymer matrices. On the contrary, in situ preparation often allows the fillers

to simultaneously preserve their aspect ratio and attain improved dispersion

and enhanced filler matrix interaction/bonding. More importantly, some in situ

preparation methods can also boost the conductivity of the fillers simultaneously.

These consequently lead to higher conductivity and lower percolation threshold

of the nanocomposites. Such a trait also presents opportunities for the lowering

cost of the conducting nanocomposites owing to the simplified processes.

9.2

In Situ Preparation of Conductive Nanocomposites

Traditionally, in situ preparation of polymer nanocomposites refers to the forma-

tion of continuous polymer phase through in situ polymerization of monomer

or precursor in the presence of nanofillers. Indeed, till today most of the in situ

preparation methods for conducting polymer nanocomposites are still in this

scope. However, in recent years, with the booming of the preparation of graphene-

based conducting polymer nanocomposites, a new strategy has emerged. Using

this strategy, the formation of the polymer phase is accompanied by a prominent

transition of the fillers from insulators/poor conductors to good conductors.

This is fairly different with the traditional in situ polymerization strategy. Firstly,

the fillers used in this case initially have very low conductivity; without the

transition, they could not form conductive pathways. Secondly, besides the in situ

polymerization methods, various other methods, such as thermal treatment

and chemical reduction, have also been used for in situ formation of conductive

pathways in polymer nanocomposites. Thus it is necessary to divide this section

into two subsections, that is, the in situ polymerization and in situ formation

of conducting polymer nanocomposites. The former corresponds to the in situ

polymerization of monomers or precursors in the presence of conductive fillers,

while the latter corresponds to the in situ formation of conductive pathways from

originally insulating/poorly conductive fillers in the process of nanocomposite

fabrication.

9.2.1

In Situ Polymerization Strategy

The in situpolymerization strategy has been proven to be a successful approach for

the preparation of conducting polymer/CNT, polymer/CNF, polymer/CB, poly-

mer/graphite, polymer/graphene nanocomposites with significantly enhanced

properties. The polymer matrices used in in situ polymerization includes

epoxy, PUs, PIs, polyamides (nylons), polyesters, polyethers, polystyrene (PS),

polypropylene (PP), polyethylene (PE), and so on. [10–15, 35] A typical in situ

polymerization process is illustrated in Figure 9.6. According to the chain

formation mechanism of the polymers, the polymerization processes can be

divided into two groups, step growth and chain growth.


Blending In situ

Monomer/precursor Conductive filler

Figure 9.6 A scheme showing in situ polymerization for the preparation of conducting

polymer nanocomposites.

9.2.1.1 Step Growth

One major issue for in situ polymerized conducting polymer nanocomposites

is to attain good dispersion of conductive fillers, especially CNTs and graphene

because of the entanglement among CNTs and strong stacking tendency of

graphene. Most reports about in situ polymerization by step growth mode

involve epoxy-based nanocomposites, where conductive nanofillers are firstly

dispersed in the precursors followed by curing the resins with hardeners. For

example, Young et al. studied the effects of different dispersion states of CNTs

on mechanical, electrical, and thermal properties of the epoxy nanocomposites,

and found that the nanocomposites with poorly dispersed CNTs exhibited higher

modulus, lower tensile strength, lower elongation at break, lower electrical

conductivity as compared to those with well-dispersed CNTs [37]. In an earlier

report, Allaoui et al. prepared the epoxy/CNT nanocomposites by adopting

a two-step procedure. The CNTs were firstly dispersed in methanol solution

under magnetic agitation to reduce the maximum size of the aggregates to about

100 μm, and then the CNTs were directly added to the epoxy precursors and the

manually homogenizated mixture was injected into sample molds for curing.

Considerable enhancement has been obtained with the additional mixing step.

The Young’s modulus and the yield strength of the nanocomposite with 1wt%

CNTs have been increased by 100 and 200%, respectively, compared with that

of neat epoxy resin. As a result of the percolation phenomenon, an increase in

the conductivity by 9 orders of magnitude was observed in the range of 0–4wt%

CNTs with a critical percolation threshold between 0.5 and 1.0wt% CNTs [38].

Yuen et al. also prepared epoxy/CNTs nanocomposites by directly adding CNTs

into epoxy/acetone solution. It was found that the bulk resistivity decreases from

8.21× 10−16 Ω cm for neat epoxy to 6.72× 10−8 Ω cm for the nanocomposite with

1wt% CNTs. The percolation threshold was about 1wt%, and certain amounts

of CNT aggregates were found in the composites [39]. For nanofillers other


than CNTs, similar results were obtained. For example, Kim et al. prepared

the thermoplastic polyurethane (TPU) nanocomposites with thermally reduced

graphene (TRG) using solution intercalation mixing method and found that the

resistance of the TRG/TPU nanocomposite started to decrease at 0.3 vol% TRG

(about 0.6wt%) [35].

The above research work demonstrated that when the nanofillers were directly

mixedwith polymer precursors in bulk or solution, the nanofillers were not ideally

dispersed in the polymer matrices, leading to a relatively high percolation thresh-

old (usually around 0.5–1.0wt%) and aggregation of the nanofillers. To attain

better filler dispersion and further lower down the percolation threshold, various

assistant technologies including applying high-shear force/sonication, surfactant

technologies, covalent modification, and noncovalent functionalization are nec-

essary.

Applying High-Shear Force/Sonication E.T. Thostenson et al. have investigated the

use of a calendaring approach for dispersion of CNTs in vinyl ester monomer

and the subsequent processing of CNTs/vinyl ester composites. The high aspect

ratios of the CNTs were preserved during processing and enabled the formation

of a conductive percolating network at concentrations lower than 0.1wt% [40].

Ball milling is also an effective method for the in situ preparation of electrically

conductive TPU/graphene nanocomposite by avoiding aggregation of graphene

nanoplatelets. The percolation threshold occurs at around 2wt% of the graphene

nanoplatelets, which is much lower than 6wt% of that prepared by stirring [41].

Sonication is one of the most used techniques to disperse the nanofillers.

SWCNT-reinforced PI nanocomposites were synthesized by in situ polymer-

ization under sonication. This process enabled uniform dispersion of SWCNT

bundles in the polymer matrix. The resultant SWCNT-PI nanocomposite films

are electrically conductive and optically transparent with significant conductivity

enhancement at a very low filler loading (0.1 vol%). Mechanical properties as well

as thermal stability are also improved with the incorporation of the SWCNT

[9]. Kim et al. succeeded in producing SWCNT/epoxy nanocomposites using

ultrasonication method. An anomalously small percolation threshold of about

0.074wt% is achieved, which is ascribed primarily to the very large aspect ratio

of the nanotubes [42]. Sandler et al. successfully utilized the ultrasonic-assisted

in situ polymerization strategy for the preparation of epoxy/MWCNT nanocom-

posites. It is possible to achieve amatrix conductivity around 10−2 Sm−1 with filler

volume fractions as low as 0.1 vol% [43]. Phenolic resin/graphite composites were

synthesized via in situ condensation polymerization of phenol and formaldehyde

in the presence of the exfoliated graphite. The composites containing graphite

sheets exhibited an electrical conductivity percolation threshold with 3.2wt%

graphite content in polymer matrix [44]. Unsaturated polyester resin/graphite

conducting nanocomposites with a low percolation threshold of 0.64 vol% have

been prepared via in situ polymerization under the application of ultrasonic

irradiation. Experimental results reveal that graphite nanosheets can effectively

form better conductive network in the resin and thus improve the electrical


conductivity of as-prepared composites by exhibiting a typical percolation

threshold at around 0.64 vol%. The reduction of the percolation threshold may be

directly related to the high aspect ratio and homogeneous dispersion of graphite

nanosheet in the polymer matrix resulting from ultrasonic irradiation [45].

Surfactant Technology Cui et al. studied the influence of surfactant on the prepa-

ration of epoxy nanocomposites by in situ polymerization, where they found that

the electrical conductivity could be reduced by surfactant modification of CNTs

[46]. This conclusion was also supported by the studies by Tang et al. in which

they found that a high content of surfactant reduces the conductivity of PI/CNT

nanocomposites synthesized by in situ polymerization [47]. On the contrary, Bar-

rau et al. found that the palmitic acid allows an efficient dispersion of CNTs in

the epoxy matrix. Electrical conductivity is optimal using a 1 : 1 CNT to palmitic

acid weight ratio.The associated percolation threshold is found between 0.05 and

0.1wt%CNTs, that is, between 0.03 and 0.06 vol%. In comparisonwith composites

without palmitic acid, the use of palmitic acid improves the electrical properties

of epoxy resin/CNT nanocomposites [48].

Covalent Modification Great volume of reports have been focused on the chem-

ical modification of carbon-based nanofillers to improve their dispersion in

solvents and polymer matrices. However, chemical modification would destroy

the intrinsic structure and reduce the electrical conductivities of the nanofillers

dramatically. Zdenko et al. studied the mechanical and electrical properties of

oxidized epoxy/CNT composites with respect to different chemical treatments of

the nanofillers. The surface modification of the CNTs was carried out by treating

the as-received CNTs with HCl, H2SO4/H2O2, HNO3, and H2O2/NH4OH under

well-specified conditions. Evidence has been provided that the induced nanotube

damage as a result of oxidative treatment could affect the mechanical and electri-

cal properties of the epoxy composites considerably. Strong oxidative treatment

applied to the nanotubes brings about a twofold increase of the flexural modulus

and a decrease of conductivity by 2 orders of magnitude. Combined improvement

of bothmodulus and conductivity can be achieved only byNH4OH/H2O2 because

of their milder impact on the conductive shells of the nanotubes and also the

effective removal of poorly conducting amorphous carbon [49]. Similar results

have been observed by So et al., who found that the surface modification of CNTs

brought a reduction in the electrical conductivity of PI/CNT nanocomposites but

improvedmechanical properties [50]. Zhu et al. studied the silanemodification of

CNFs.The surface of CNFswasmodified by a functional amine-terminated silane,

and then the amine groups in situ reacted with epoxy monomers to form polymer

nanocomposites. This in situ reaction favors the CNF dispersion and improves

the interfacial interactions between the CNFs and monomers. Significant

increase in the mechanical properties is observed because of better dispersion

and introduced strong interfacial interactions. However, electrical conductivity

is decreased owing to the insulating silane coating layer [51] (Figure 9.7).


Epoxy monomer

As-receivedCNFs

d

Silane-treatedCNFs

e−

e−

e−

e−

(a) (b)

Figure 9.7 Contact model of (a) as-received CNFs and (b) silane-treated CNFs suspended in

epoxy solutions.

As for graphene, numerous reports are focused on the mechanical strength-

ening effect brought by GO. Most of the conductive polymer nanocomposites

are prepared by in situ reduction of GO in the manufacturing process, which is

discussed later. Only limited reports have demonstrated the successful prepa-

ration of conducting polymer nanocomposites using chemically prereduced

conductive graphene sheets as fillers. For example, Park et al. synthesized

aminophenyl-functionalized graphene nanosheets (APGNS) through a dia-

zonium salt reaction using thermally reduced graphene nanosheets (GNS) to

improve mechanical properties and electrical conductivity. PI nanocomposites

with different loadings of APGNS were prepared by in situ polymerization.

Both the mechanical and electrical properties of the PI/APGNS nanocompos-

ites are significantly improved compared with those of neat PI because of the

homogeneous dispersion of APGNS and the strong interfacial covalent bonds

between APGNS and the PI matrix. The electrical conductivity of APGNS/PI

(3 : 97w/w) is 6.6× 10−2 Sm−1, which is about 1011 times higher than that of neat

PI. Furthermore, the modulus of PI/APGNS is increased up to 16.5GPa, which

is approximately a 610% enhancement compared to that of neat PI, and tensile

strength is increased from 75 to 138MPa [52].

Noncovalent Functionalization A comparison study of covalently and noncova-

lently functionalized CNTs in epoxy has been conducted by Liu et al. In this

study, CNTs were dispersed into an epoxy matrix using polyethylenimine (PEI)

as a dispersant that was either covalently attached to the nanotubes or physically

mixed to result in only noncovalent interactions. Both forms of functionalization

produced improved nanotube dispersion. Epoxy nanocomposites containing

covalently modified MWNTs exhibited greater storage modulus and reduced

electrical conductivity, while the noncovalent stabilization exhibits similar

conductivity to unstabilized nanotubes [53].

Huang et al. presented a valuable in situ fabrication process for synthesizing

highly conductive PI/MWCNT nanocomposite films. They found that 2,6-

diaminoanthraquinone (DAAQ) is not only an excellent dispersion agent to

stably disperse pristine MWCNTs in solvent but also a monomer to directly

synthesize PI. A high electrical conductivity value of 55.6 S cm−1 is achieved


in the PI nanocomposite containing 40wt% of MWCNTs with a percolation

threshold at about 0.50wt% (or 0.32 vol%) [54].

Yuan et al. developed in situ polymerized PI/MWCNT nanocomposites using

a novel poly(amic acid) (PAA) containing a rigid backbone with hydroxyl pen-

dant groups, as both the nanotube dispersant and the matrix precursor. The nan-

otube content in the solution-cast PI-based nanocomposite could be increased

to as high as 30wt% and achieve ultrahigh electrical conductivity as well as high

mechanical properties. The electrical conductivity of the PI/MWCNT nanocom-

posites reaches a value of 3880 Sm−1 at a nanotube loading of 30wt% and the

MWCNT content for achieving the percolation threshold of conductivity of the

nanocomposites is only 0.48wt%. The composite with 30wt% MWCNTs has a

higher Young’s modulus (9.43 vs 0.14GPa) and tensile strength (179.2 vs 9.7MPa)

than other nanotube-reinforced PI nanocomposites. The high conductivity and

tensile properties of the nanocomposite films are attributed to the good nanotube

dispersion and strong nanotube–polymer interfacial adhesion achieved through

the use of a single polymer to perform the dual functions of nanotube disper-

sant and matrix precursor [55]. Similar strategy has also been applied to prepare

PI/CNT, PI/CNF, PI/CB nanocomposites [56].

9.2.1.2 Chain Growth

The above dispersion methods for chain growth can also be applied to disperse

conductive nanofillers for in situ preparation of conducting polymer nanocom-

posites by chain growthmode.The nanofillers are usually first functionalized with

special reactive groups that can be used as initiator or transfer agent for chain

growth.

Olefin Polymerization The penetration of polyolefins into the worldwide plastics

marketplace in the past few decades has been spectacular. Thus considerable

efforts have been made on the development of polyolefin-based conducting

nanocomposites. in situ metallocence polymerization has been used to prepare

high-density polyethylene (HDPE)/MWCNT nanocomposites. A metallocene

catalyst complex is attached onto the surface of MWCNTs, and surface-initiated

polymerization of ethylene generates PE brushes on the surface ofMWCNTs.The

uniform dispersion of MWCNT throughout HDPE matrix leads to a 11-order

improvement in electrical conductivity, that is, about 1.3 Sm−1 at a 7.3wt%

MWCNT loading compared with 10−11 Sm−1 of neat HDPE [57].

PE/graphene nanocomposite prepared by in situ polymerization was also

reported. The presence of graphene in the PE matrix increases the onset

degradation temperature by 30 ∘C. The electrical conductivity measured by the

impedance technique shows a critical percolation threshold of 3.8 vol% (8.4wt%)

of graphene. Compared with neat PE, the nanocomposites also exhibit a slight

decrease in the tensile strength and an increase in the storage modulus [58].

Another strategy for the preparation of conducting polyolefin/graphene

nanocomposites seems more attractive. A Grignard reagent, n-BuMgCl, is

found to be able to reduce GO to form loosely aggregated graphene sheets


immobilized with few Mg–Cl species. Further complexation with TiCl4 leads

to a graphene-based supported catalyst that is ready for in situ olefin polymer-

ization to give electrically conductive polyolefin/graphene nanocomposites.

PP/graphene nanocomposites prepared using this method possesses a fairly

low percolation threshold (approximately 0.2 vol%) and high conductivities, for

example, 3.92 Sm−1 at 1.2 vol% and 163.1 Sm−1 at 10.2 vol% [59].

Free Radical Polymerization PMMA/MWCNT nanocomposites have been

prepared via in situ bulk polymerization of MMA (methyl methacrylate)

in the presence of MWCNTs. Electrical conductivities of the synthesized

PMMA/MWCNT nanocomposite containing 1∼5wt% of MWCNTs were

between 10−2 and 1 Sm−1. However, when PMMA/MWCNT nanocomposites

with the sameMWCNTs contents were electrospun into nanofibrous membrane,

the conductivities were reduced to about 10−8 Sm−1, even though dispersion of

the MWCNTs in the electrospun nanofibers was superior to the conventional

polymer composites with CNTs. MWCNTs in the electrospun nanofibers were

found to be embedded in the polymer matrix and align along the fiber axis [60].

Chen et al. reported a simple in situ polymerization method for the prepara-

tion of conductive polymer/graphite nanocomposites. Natural flake graphite was

expanded into a worm-like particle composed of graphite sheets with thickness in

the nanometer scale via an acid intercalation procedure. Subject to ultrasonic irra-

diation, the expanded graphitewas effectively exfoliated into isolated thin graphite

nanosheets carrying oxygen-containing groups on the surface of the exfoliated

graphite. These thin graphite nanosheets were incorporated into polymer matrix

via in situ polymerization, leading to the production of polymer nanocomposites

with low percolation threshold and high conductivity.This method was applied to

prepare conductive nanocomposites by using PS, PMMA, and nylon 6 as polymer

matrices, respectively [61–63].

Recently, Liao et al. successfully prepared polyurethane acrylate (PUA)/TRG

nanocomposites by in situ free radical polymerization with an ultralow per-

colation concentration of 0.15wt% (0.07 vol%), which is the lowest electrical

percolation concentration among all the TRG/polymer nanocomposites reported.

Interestingly, they found that the percolation concentration of cured nanocom-

posites is lower than that obtained for TRG/UA uncured liquids, which may be

explained by the volume shrinkage during free radical polymerization of UA.

Modulus results suggest an effective mechanical reinforcement of PUA/TRG

nanocomposites in the rubbery state, while no significant reinforcement was

observed in the glassy state [64].

Using in situmicroemulsion polymerization, a large-scale production route for

PS nanoparticle-functionalized graphene sheets was developed by Patole et al. It is

found that the thermal properties of the PS were improved with the incorporation

of graphene in the nanocomposite.The PS nanoparticle-modified graphene shows

good compatibility with the host PS matrix, and it is easy to prepare conducting

PS films by hot press. This reported scheme for fabricating the PS composite thin


Graphene

Dispersion

RT

SDS

Microemulsion

SO

O

O

ONa+

0 °C

85 °C

PS nanoparticles Polymerization AIBN+Styrene

n

Figure 9.8 A scheme showing functionalization of graphene with polystyrene

nanoparticles.

films from graphene and a commodity plastic highlights the potential for low cost,

macroscale thin-film electronics [65] (Figure 9.8).

Anionic Polymerization PS/graphite nanocomposites have also been synthesized

by in situ polymerization of styrene in a solution system that contains potassium

(K), tetrahydrofuran (THF), and graphite intercalation compound (GIC). The K-

THF-GIC is able to initiate the polymerization of styrene by anionic mechanism.

Owing to the interfacial interaction between the graphite nanolayers and the poly-

mer, the nanocomposites exhibit higher glass transition temperature and higher

thermal stability in comparison with neat PS. The percolation threshold of the

nanocomposite is less than 8.2wt% [66].

Ring-Opening Polymerization A series of poly-L-lactide (PLLA)/TRG nanocom-

posites were prepared by Yang et al. via the in situ ring-opening polymerization

of lactide using TRG as the initiator. Typical percolation behavior was observed

for TRG contents between 1.0 and 1.5wt%, and the electrical conductivity of the

PLLA was improved from 7.14× 10−14 Sm−1 for neat PLLA to 1.63× 10−2 Sm−1

for nanocomposites with 2.0wt% of TRG sheets [67] (Figure 9.9).


Natural graphite

Thermal reduced graphene oxide (TRG)GLLA composite

Graphite oxide (GO)

H2So

4, NaNO

3

SnOct, 170 °C

Thermal reduction

OH

OH

OHOH

OH

O

O

OO

O

OO

OO

O

OHO

HO

HO OH

OHHO

O

O O

O O

O

OO

HO HO

HOOH

1000 °C

O

OO

O

OO

O

O

O

O O H

n

Figure 9.9 Synthesis of poly-L-lactide grafted with thermally reduced graphene oxide

sheets.

Higgins et al. reported the preparation of electrical conductive polycar-

bonate/CNF nanocomposites via in situ polymerization of cyclic oligomeric

carbonates. The composites exhibit minimal aggregation of the CNF even at high

weight percentages, leading to an electrical conductivity percolation threshold of

6.3wt% [68].

Electrical conductive poly(arylene disulfide)/graphite nanocomposites were

prepared using in situmelt-ring opening polymerization of macrocyclic oligomer

in the presence of microwave-expanded graphite. Graphite nanosheets were

well-exfoliated within the nanocomposite, and the synthesized nanocomposites

reserved the superior inherent mechanical strength of polymeric matrix together

with improved electrical conductivity about 10−1 Sm−1 at 5wt% graphite

content [69].

Poly(butylene terephthalate) (PBT)/graphene composites were prepared by

in situ polymerization of cyclic butylene terephthalate oligomers in the presence

of graphene. Graphene plays a constraining role in the mechanism of PBT chains

growth by interaction with the catalyst and/or the active center of propagation.

Increasing the loading of graphene causes a decrease in the average molecular

weight of PBT. Mechanical testing showed that elastic modulus and indenta-

tion hardness of PBT are significantly increased for graphene contents lower

than 0.75wt%, while higher loading induces stronger interactions between the

carbon nanoplatelets and decreases the overall properties of PBT/graphene

composites [70].

9.2.1.3 Aligning Conductive Fillers in in situ Polymerization Processes

As mentioned in Section 9.1.3.4, alignment of 1D and 2D conductive nanofillers

can affect the connectivity of the fillers and hence have an effect on the perco-

lation threshold of the nanocomposites. Some external forces have been used


to induce the orientation of nanofillers so as to realize a network formation at

lower percolation concentration in in situ polymerization process. For example,

Kimura et al. used a high magnetic field to align MWCNTs in a polyester matrix

and obtained electrically conductive and mechanically anisotropic nanocom-

posites [71]. MWCNTs were dispersed in the monomer solution of unsaturated

polyester under a constant magnetic field, and the nanotubes were aligned suc-

cessfully during the in situ polymerization process. Choi et al. prepared aligned

epoxy/MWCNT nanocomposites under a 25T magnetic field. The electrical

properties along the magnetic field alignment direction were increased by 35%

compared to those epoxy/MWCNT nanocomposites without the application of a

magnetic field [72]. Torsten Prasse et al. studied the influence of a static electric

field applied by metal electrodes on the agglomeration process of CB in an epoxy

resin [73]. A conductive epoxy/CB composite with a low CB content of 0.12wt%

was achieved. The growth of dendrites from the anode into the material was

observed and a percolating network was seen to form, combined with a drastic

reduction in the sample resistivity.

9.2.2

In Situ Formation of Conducting Polymer Nanocomposites

In this section, the in situ nanocomposite formation methods, by which the for-

mation of the continuous polymer phase is accompanied by a sharp increase in

electrical conductivity of the fillers, are reviewed. So far, such methods have been

used mainly for the fabrication of rGO-based nanocomposites. There are only a

few reported works on other systems.

9.2.2.1 In Situ Formation of rGO-Based Polymer Nanocomposites

To produce graphene-based conductive polymer nanocomposites, highly conduc-

tive graphene sheets must be produced at a sufficient scale, and the nanosheets

must be homogeneously dispersed in polymer matrices. Till today, the most

promising method to produce individual graphene sheet at large scale is to oxi-

dize graphite followed by GO exfoliation and reduction [74]. After exfoliation in

strong oxidants, the GO nanosheets can be readily dispersed in various solvents,

monomers, or precursors, affording good compatibility of the GO nanosheets

with many polymer matrices [75]. However, because of the defects in GO

chemical structure, the conductivity of GO is too low to give the nanocomposites

reasonably high conductivity. As rGO sheets are easy to aggregate and difficult to

exfoliate, in order to produce highly conductive polymer nanocomposites with

well-dispersed graphene sheets, GO should preferably be in situ reduced. Two

main routes used to convert GO to highly conductive graphene are chemical and

thermal reduction, by which the oxygen-containing groups in GO are removed

by reactions with chemical reductants and decomposition at high temperatures,

respectively, and the conjugated structure and high electrical conductivity of

graphene are restored [14, 76, 77].


In Situ Chemical Reduction Themost commonly used chemical reduction method

is isocyanate modification-hydrazine reduction, which was first reported by

Stankovich et al. [74]. The isocyanate treatment reduces the hydrophilic charac-

ter of GO sheets by forming amide and carbonate ester bonds to the carboxyl and

hydroxyl groups of GO, respectively. As a result, such isocyanate-treated GO no

longer can exfoliate in water but readily form stable dispersions in polar aprotic

solvents, giving completely exfoliated, functionalized GO sheets with thickness

of about 1 nm. The isocyanate treatment allows the GO sheets to be intimately

mixed with many organic polymers or their monomers, facilitating synthesis

of graphene-polymer nanocomposites. The GO sheets can be molecularly

dispersed in polymer matrices and then undergo in situ reduction by hydrazine

treatment. A PS-graphene nanocomposite formed via this route (Figure 9.10)

exhibits a percolation threshold of 0.1 vol% for room-temperature electrical

conductivity, which is one of the lowest values reported for carbon filler-based

polymer nanocomposites; at only 1.0 vol%, this nanocomposite has a conductivity

of 0.1 Sm−1. This bottom-up chemical approach has been widely adopted by

many researchers to produce various conductive polymer nanocomposites. The

matrices involved include epoxy [78], PS [79, 80], PI [81, 82], polyamide [83], and

others [84].

On the basis of the reduction mechanism proposed by Stankovich, that is, the

reduction involves amine attack and epoxide ring opening, researchers found

that many other organic amino agents, such as octadecylamine [85], p-phenylene

OHO

O

O O

COOH HOOC

HO

OH O

O

OH

COOH

OO

COOHOH

COOH

COOHHO

HOOC

HO

OH O

O

OH

COOHHydrazine hydrate

ReductionOO

COOH OH

COOH

COOHHO

COOH

COOH

Styrene, SDS

Ultrasonication for 15 min In situ polymerization

K2S

2O

8(KPS)

OHCOOH

Graphene oxide nanosheets

Graphene oxide nanosheets-polystyrene microspheres

Graphene nanosheets-polystyrenemicrospheres

Styrene-linked graphene oxide nanosheets

HOOC

OH

HOHO

Figure 9.10 A schematic showing the synthesis route for polystyrene/rGO nanocomposites.


diamine [86], and dopamine [87], could also be used as reducing agents to

produce rGO sheets. Some other nonamine reducing agents were also developed.

For example, graphene sheets were nicely exfoliated in PMMA matrix by a novel

reactive biphasic process, in which the in situ reduction of GO by sodium boron

hydrogen and grafting of PMMA on the rGO sheets by a radical addition pathway

occurred simultaneously [88]. A green method devoid of harmful chemicals

and chemical process was also developed. Tea solution was found applicable in

the reduction process of GO [89]. These strategies allow us to prepare various

conducting polymer nanocomposites with low percolation threshold.

In Situ Thermal Reduction Compared with chemical reduction, in situ thermal

reduction of GO is a much simpler approach and has been receiving increasing

attention. In thermal reduction process, there is no need to use harmful chemicals

or chemical processes. However, in order to obtain highly conductive graphene,

it is necessary to conduct the thermal reduction under relative high temperatures

(typically above 200 ∘C). Hence, the polymers that can be used for the preparation

of conducting polymer nanocomposites by thermal reduction are very limited

and mainly include PIs [90, 91], polyamides [92], polyesters [93], and some

polycarbonate [94]. For example, Xu et al. reported the in situ polymerization

cum thermal reduction approach for the preparation of rGO-reinforced nylon 6

nanocomposites (Figure 9.11). The GO are in situ thermally reduced during the

polycondensation of caprolactam. Nylon 6 chain brushes are grafted onto the

graphene sheets efficiently, making the rGO sheets homogeneously dispersed

in nylon 6 matrix. Their results show that the tensile strength of the melt-spun

nylon6/rGO nanocomposites are increased by 2.1-folds and Young’s modulus

increased by 2.4-folds with the incorporation of 0.1wt% graphene only [95].

NH NH

HO

+

O

+ H2N(CH2)5COOH + H2N(CH2)5COO− 250 °C

−H2OH3N

O

NH

+COO−

N

On

H

Protonated monomer

Condensation

HOOC HOOC

COOH

HOOCHOOC

COOHHO

HO

HO

OH

OH

OH

O

O

O

O

250 °C

O

O

+

Figure 9.11 Synthesis of graphene/nylon 6 nanocomposites by in situ ring-opening poly-

merization of caprolactam in the presence of GO.


Zheng et al. studied the electrical conductivity of nylon 6/graphene nanocom-

posites prepared using a similar strategy.The results confirmed that the exfoliated

and dispersed GO nanosheets were in situ thermally reduced during the poly-

merization, resulting in enhanced electrical conductivity. Percolation in the nylon

6/rGO nanocomposites occurs at about 0.41 vol%. As the feed volume fraction of

GO increases from 0.27 to 1.09 vol%, the electrical conductivity rapidly rises by 10

orders of magnitude from 4.2× 10−14 to 1.0× 10−4 Sm−1. The polarity matching

between rGO and nylon 6 macromolecules is also important as it stabilizes the

exfoliated state of the nanosheets in the matrix after the polymerization. This

one-step in situ reduction and polymerization process opened a new avenue

for the fabrication of graphene-based polymer nanocomposites in both scalable

and cost-effective ways for a wide range of practical applications [92]. The

strategy has been applied in preparation of other types of conducting polymer

nanocomposites, including poly(ethylene terephthalate)/graphene, PI/graphene,

and polycarbonate/graphene nanocomposites, for which the temperature used

were about 250–280 ∘C, and the percolation threshold achieved ranges from 0.1

to 0.5wt% [93, 96–98].

It is worth noting that Ye et al. systematically investigated the influence of the

nature of a polymer on the in situ thermal reduction of GO dispersed in the poly-

mermatrix.They found that the interactions betweenGO and the polymermatrix

play a key role in decreasing the reduction temperature of GO dispersed in the

polymer matrix. Both polar polymers and aromatic polymers can decrease the

reduction temperature of GO whereas nonpolar polymers cannot, indicating that

the change in the reduction temperature of GO is associated with the favorable

interactions between polymer matrices and GO sheets [99].

Integrated Chemical and Thermal Reduction Method The high temperatures

required for the thermal reduction process inevitably cause high equipment

and energy costs for the production of conducting polymer nanocomposites.

In addition, for the polymers that would degrade at temperatures of 200 ∘C or

above, the high reduction temperatures required are a great obstacle for the

development of conducting polymer nanocomposites. Thus researchers are

trying to develop milder methods to produce graphene sheets with high electrical

conductivity. The major purpose is to lower down the reduction temperature

of GO. For example, a two-step strategy consisting of a hydrazine hydrate reduc-

tion and subsequent thermal reduction at 200∘ was proposed by Wu et al. Using

this strategy, PS/graphene nanocomposites with electrical conductivity as high as

22.7 Sm−1 were successfully prepared by in situ reduction of GO. The graphene

sheets are well dispersed and ultimately form a continuous network structure in

the polymer matrix [100]. A similar chemothermal reduction method has also

been reported, which allows the preparation of highly conductive graphene films

under milder conditions [101].

However, the two-step strategy is still burdened by the usage of haz-

ardous/poisonous reductants, which is harmful to human body and not suitable

for high throughput processes. Hence, other chemothermal integrated strategies


were developed. For example, recently, a low-temperature reduction method has

been reported, by which GO was reduced in the presence of organic solvents,

giving rise to an increase in electrical conductivity to 230 Sm−1 [102]. In a more

recent work, polydopamine-coated GO were thermally annealed at 130–180 ∘C,and this low thermal annealing process led to highly conductive graphene sheets

with electrical conductivity of about 30 000 Sm−1. This large-scale green strategy

has been successfully used for in situ preparation of conductive poly(vinyl

alcohol) (PVA)/graphene and epoxy/graphene nanocomposites (Figure 9.12),

both of which demonstrated good dispersion, low percolation threshold, and

high electrical conductivity [103].

Other Methods Except chemical reduction and thermal reduction, GO can be

reduced by other methods. For example, Varrla Eswaraiah et al. reported one-

pot synthesis of conducting polyvinylidene fluoride (PVDF)/graphene nanocom-

posites for strain sensing applications, in which the GO sheets incorporated in

PVDF powder were in situ reduced by solar electromagnetic radiation. The as-

synthesized nanocomposites are highly conductive in nature and exhibit lower

percolation threshold [104].

9.2.2.2 In Situ Formation of Metallic Conductive Pathways

Other than carbon-based nanofillers, some metallic particle networks could

also be in situ formed in the process of preparing polymer nanocomposites. For

example, Zhang et al. prepared highly conductive epoxy nanocomposites by

thermal decomposition of silver carboxylate on the surface of silver flakes and

subsequent sintering between conductive fillers during the epoxy adhesive curing

process.The sintering between conductive fillers effectively reduces or even elim-

inates the contact resistance (Figure 9.13), leading to polymer nanocomposites

with the electrical resistivity of 6.3× 10−5 Ω cm.This technology has been used to

for the fast processing of printable electronics [105, 106].

Graphene

Graphene

(a) (b)

Figure 9.12 Field-emission scanning electron microscopic images of (a) PVA/graphene

nanocomposite and (b) epoxy/graphene nanocomposite.

9.3 Challenges and Outlook 233

Fast sintering

: Ag flake : Sintered Ag: Surface lubricant : Ag nanoparticle

Figure 9.13 Schematic illustration of the sintering between Ag nanoparticles and Ag flakes

within a polymer matrix.

Southward et al. also reported the in situ preparation of reflective and surface

conductive PI films by the incorporation of silver(I) acetate and trifluoroacety-

lacetone into a dimethylacetamide solution of the PAA formed of 3,3′,4,4′-

oxidiphthalic dianhydride (ODPA) and 4,4′-oxidianiline (4,4′-ODA). At selected

silver concentrations, thermal curing of (trifluoroacetylacetonato)silver(I)-PAA

films led to cycloimidization with concomitant silver(I) reduction, resulting in a

reflective and conductive silver surface if the film was cured to a final temperature

of 300 ∘C for several hours (Figure 9.14). The metallized films exhibit outstanding

metal-polymer adhesion at the interface and mechanical characteristics close

to those of the parent PI. The metallized ODPA/4,4′-ODA films also exhibit

good thermal stability, particularly under nitrogen atmosphere. Later this

strategy was also applied to prepare conducting palladium-PI nanocomposite

membranes [107–114].

9.3

Challenges and Outlook

The research results described above have demonstrated that the in situ prepara-

tion strategies are attractive approaches for facile preparation of electrically con-

ductive polymer nanocomposites for various applications. However, considering

the rising demands for conducting polymer nanocomposites, we should keep the

challenges and key issues that need to be addressed in perspective.

Firstly, some common applications, such as EMI shielding, require fairly high

electrical conductivities, which could not be achieved at low loadings of conduc-

tive fillers for most systems at present. The overall electrical conductivity of a

polymer nanocomposite is limited by the intrinsic conductivity of the nanofillers

and the electron loss at the junctions of the conductive pathway formed by the

nanofillers [33]. In case the final electrical conductivity of a specific type of con-

ductive fillers is affirmed, efforts should be focused on how to lower down the

electron loss at the junctions while maintaining good filler dispersion. As a highly


H

HOO O

X = C=O for BTDA/4,4′-ODA, O for ODPA/4,4′-ODA, C–C bond only between rings for BPDA/4,4′-

Four poly(amic acid) structures

Four resultant surface-silvered polyimide films

276 °C

R = 11%

100 nm

R = 24% R = 74%1000 nm

R = 66%

300 °C for 1 h 300 °C for 3 h 300 °C for 7 h

Silver(I) compounds include1. Silver(I) nitrate and2. Silver(I) trifluoroacetate3. Diamminesilver(I) nitrate

4. silver(I) tetrafluoroborate

(1,5-Cyclooctadiene)(hexafluoro-

acetylacetonato)silver(I)

5.

6.

CF3O

OAg H

X X = CH3, CF3,S

7.

8.

Poly(amic acid)and silver(I) additivedissolved in a polar aprotic solvent, for example,DMAc.Film cast from solution above.Cast film dried at ambient temperature.Film heated to reduce silver(I) and eliminatewater to give surface-silvered polyimide film.

OH

OO

n

n

N

HXN

O

H

HOO

O

O

O

O

O O

O O

O OF3C CF3 CF3

CF3

and N N

n

OX

N N

F3C CF3CF3

CF3

O HN

OH 6FDA/4-BDAFO

O ONO

ODA

1.

2.3.4.

+

Figure 9.14 (a) Schematics showing

thermally promoted metal ion reduction

of silver(I)-doped poly(amic acid) films

as a route to surface-metallized PI films.

(b) Transmission electron microscopic images

of 13% AgTFA-ODPA/4,4′-ODA films cured

under different conditions.

conductive polymer coating, conductive nanofillers can greatly reduce the con-

tact resistance and tunneling/Schottky barriers at the interfiller junctions in the

network, giving rise to a 2-order improvement in electrical conductivity of the

network after the percolation threshold as compared with those nanofillers with-

out conductive polymer coating [115]. It is reasonable to propose that this strategy

References 235

is also useful for in situ polymerization of conducting polymer nanocomposites.

In fact, researchers have already demonstrated that it is possible to improve the

dispersion of the nanofillers and the final electrical conductivity (by 3 orders) of

melt compounded nylon 6/CNT nanocomposites simultaneously using this con-

ductive polymer coating method [116]. Besides lowering down the electron loss

at the junctions, development of efficient technologies for full reduction of GO is

also a feasible approach to attain higher filler conductivity [103, 117].

The second challenge is how to further lower down the percolation threshold

so that the cost of the nanocomposites can be further reduced. Except the appli-

cation of various technologies to attain molecular level dispersion of nanofillers,

several special strategies have been formulated to lower down the percolation

threshold of the nanofillers in conducting polymer nanocomposites, including

the synergistic effect between different nanofillers [118–120], double percola-

tion [121, 122], volume-exclusion principle [123], latex technology [124–128], and

construction of a 3D interconnected nanofiller network [117]. These strategies

have been proven to be very effective in lowering down the percolation thresh-

old in various nanocomposite fabrication processes. Sometimes the percolation

threshold was reduced to an amazing low level of about 0.075 vol%. It is therefore

highly desirable to extend the application of these strategies to in situ preparation

processes to develop conducting polymer nanocomposites with better properties.

As mentioned earlier, conducting polymer nanocomposites are attractive for

use in a wide variety of fields. However, till today, most of the in situ processes

have been studied at lab scale instead of industrial scale. Preparation of conducting

polymer nanocomposites under milder conditions, for example, at low reaction

temperatures, with short reaction time, using less chemicals or utilizing low-cost,

renewable chemicals, would enable the scaling up of these in situ processes, lead-

ing to real industrial applications of these fascinating functional materials.

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241

10

Near IR Spectroscopy for the Characterization of Dispersion

in Polymer–Clay Nanocomposites

Ana VeraMachado, JoanaMargarida Barbas, and Jose Antonio Covas

10.1

Introduction

Thermoplastic polymers reinforced with layered silicates can exhibit excellent

physical and mechanical performance at filler contents typically lower than 5% in

weight. Each individual clay sheet is hundreds to thousands of nanometers long

and wide and has a thickness in the nanometer scale. Thus, its surface is quite

substantial, of the order of hundreds of square meters per gram, yielding specific

characteristics to the nanocomposite. However, as layered silicates are usually

available as stacks of tactoids and their hydrophilic character is incompatible

with the hydrophobic nature of polymers, dispersion into individual monolayers

is difficult and has been the focus of intense research.

This chapter starts with a brief overview of the morphology and properties

of polymer–clay nanocomposites, as well as of their preparation methods and

characterization techniques. The use of near-infrared (NIR) spectroscopy for

the characterization of dispersion in polymer nanocomposites is then discussed,

with a focus on the application of inline techniques to monitor the preparation of

polymer–clay nanocomposites by melt compounding.

10.2

Morphology and Properties

Simple mechanical mixing of a polymer with a silicate does not necessarily yield

a nanocomposite. Depending on the interaction between polymer and clay sur-

face, as well as on the thermomechanical environment created during mixing,

separation of the clay stacks into discrete uniformly dispersed sheets may not

be attainable [1–9]. Figure 10.1 shows a schematic representation of the possible

polymer-clay composite morphologies resulting from mixing.

Immiscible composites are a consequence of the inability of the polymer to

intercalate into the interlayer spacing (i.e., clay galleries). The clay remains in

its agglomerate state, creating a micron-size dispersed phase. The properties of


242 10 Near IR Spectroscopy for the Characterization of Dispersion in Polymer–Clay Nanocomposites

h0

h

Immiscible“micro” composite

(a)

(b)

(c)

Intercalatednanocomposite

Exfoliatednanocomposite

Figure 10.1 (a–c) Possible polymer–clay nanocomposite morphologies. Reproduced with

permission from [9] © 2001, AIP Publishing LLC.

such materials are comparable to those of traditional microcomposites. The most

beneficial morphologies are those maximizing polymer–clay interactions by

producing a large interfacial area [9, 10]. They include intercalated and exfoliated

systems.The term intercalation is associated with the polymer chains going in the

interlayer, resulting in a multilayer ordered morphology. Intercalation increases

interlayer spacing to approximately 2–4 nm [2, 3, 7]. Exfoliation corresponds

to the complete delamination of all clay platelets. In this case, their separation

exceeds 8–10 nm [2–4, 7]. In practice, most polymer–clay nanocomposites,

particularly those prepared by melt mixing, have an intermediate morphology,

comprising intercalated clay tactoids and exfoliated platelets [3, 7, 11–14].

High dispersion levels maximize polymer–clay interactions, enhancing

mechanical properties, flame retardancy, and thermal stability, lowering perme-

ability, and improving the catalytic effect on biodegradability of biodegradable

polymers, while keeping optical transparency, low density, and processability

[3, 8, 15–18].


10.3

Preparation Methods

Polymer–clay nanocomposites are obtained by intercalating polymer or

monomer inside the galleries in between silicate layers, which will eventually lead

to the delamination of individual platelets. The most important manufacturing

routes are [2–4, 15–18] as follows:

1) Solution intercalation (of polymer or prepolymer): This is a solvent-based

method, in which the polymer solution replaces a suitable previously

intercalated solvent. When using organic solvents, clay stacks can be easily

dispersed. The polymer is adsorbed onto the layer surfaces, and on evapora-

tion of the solvent the clay platelets return to equilibrium, fixing the polymer

chains inside the galleries.

2) In situ polymerization: This was the first method used to prepare nylon 6-

clay nanocomposites. It encompasses swelling of the modified clay by a liquid

monomer ormonomer solution, followed by polymerization triggered by heat

or radiation, assisted by the diffusion of a proper initiator or catalyst that was

previously fixed in the interlayer spacing by cationic exchange.

3) Melt mixing: In this case, clay and polymer are mixed in the molten state.

Under the appropriate conditions, the polymer melt diffuses within the clay

galleries, forming the nanocomposite. The process can be carried out in con-

ventional polymer compounding equipment, representing a cost-effective

and environmentally sound solution for the industrial scale production of

polymer nanocomposites.

10.4

Characterization Techniques

Transmission Electron Microscopy (TEM) and X-ray diffraction (XRD) are the

most commonly used techniques to characterize polymer–clay nanocomposites

[2, 19–23]. TEM offers a direct observation of clay dispersion enabling, for

example, to determine agglomerates/particle size and number of stacks/platelets

per area. However, the technique analyzes only very small areas, requires time-

consuming sample preparation, and quantification of morphology is tedious

[2, 20, 21]. XRD presents a clear region of interest at low angles, because of the

clay typical of Bragg’s diffraction peak, from which the clay spacing and the

stacks’ height may be estimated. However, XRD fails to provide information on

the spatial distribution of the clays within the matrix and is prone to interferences

[2, 19–21]. Indeed, many factors, such as clay loading, orientation, and crys-

tallinity, influence XRD intensity peaks. Efforts to develop a quantitative analysis

of clay dispersion have combined TEM and XRD [22, 23]. The approach makes

sense, as each technique is able to fill in the information gap of the other (see

Figure 10.2, which shows TEM and XRD data for different stages of dispersion

of a specific polymer–clay nanocomposite [20]). However, if the results are


200 nm 200 nm 100 nm

Immisciblenanocomposite

Intercalatednanocomposite

ExfoliatednanocompositeIn

tensity

Inte

nsity

Inte

nsity

Immiscible

Pureorganoclay

Pureorganoclay

Pureorganoclay

2θIntercalated

2θExfoliated

2θ

Figure 10.2 (a–c) Different states of clay dispersion, as observed by TEM and correspond-

ing XRD spectra [20].

encouraging at local level, perceiving the overall state of clay dispersion of a given

system would require an extremely time-consuming study [2, 20, 21].

As the rheological response of filled polymers is sensitive to structure, particle

size, and shape, as well as interfacial characteristics, it is not surprising that it has

been one of the most utilized tools for the characterization of nanocomposites

[23–28]. Several authors have successfully related the rheological behavior with

the state of dispersion [12, 14, 23, 25–30], while others used it to appraise distri-

bution and structural effects [16, 23, 26–29].Themagnitude of the storage or loss

moduli (G′ and G′′) in the linear viscoelastic regime provides a good insight into

dispersion quality, an increase of their values corresponding to finer dispersion

[29], while the development of a plateau at low frequencies has been attributed

to the deformation and recovery of dispersed particles [27–33]. Nevertheless, a

quantitative analysis of filler dispersion based on rheological measurements is not

frequently reported, as it is not easy to discriminate the effects of filler concentra-

tion and dispersion state on the rheological behavior [27, 31].

Recently, Lertwimolnun andVergnes [12, 14, 29] proposed the use of amodified

Carreau–Yasuda model with yield stress (σ0) to describe the frequency depen-

dence of the absolute complex viscosity:

|𝜂∗ (𝜔)| = 𝜎0

𝜔+ 𝜂0[1 + 𝜆𝜔𝑎] (𝑏−1)∕𝑎 (10.1)

where zero shear viscosity (𝜂0), relaxation time (𝜆), Yasuda parameter (a), and

power law index (b) are adjustable parameters. As melt yield stress is generally

associated with transition from liquid-like to solid-like behavior, in the case of

clay nanocomposites it can be related to the formation of a percolated network


structure of intercalated tactoids and exfoliated platelets [24, 28, 31–34]. As

shown in Figure 10.3, Lertwimolnun and Vergnes correlated melt yield stress

with the exfoliation of polypropylene (PP)-based nanocomposites prepared in a

twin screw extruder and were able to quantitatively study the effects of operating

parameters and material characteristics on dispersion. Alternatively, a power law

expression for the complex viscosity (𝜂∗) has been utilized for the low frequency

range and related to clay dispersion [23–28]:

|𝜂∗ (𝜔)| = 𝐾𝜔𝑏 (10.2)

where power law index (b) and consistency (K) are adjustable parameters. Yet, the

same dependence was found for high clay loading in a poorly dispersed composite

and for good clay dispersion at lower clay concentration [23, 26].

Several spectroscopic techniques, such as nuclear magnetic resonance (NMR)

and Fourier transformed infrared spectroscopy (FT-IR), have also been used to

characterize morphology, surface chemistry, and dynamics of exfoliated polymer

500 nm

500 nm

500 nm

500 nm

500 nm

010

100

1000

10 000

0.05 0.1

y = 3.88×−1.533

R2 = 0.987

0.15

Q/N

Melt y

ield

str

ess (

Pa)

Figure 10.3 Evolution of melt yield stress as a function of filling ratio (output/screw speed,

Q/N) for PP–clay nanocomposites prepared under different operating conditions. Repro-

duced with permission from [12] © 2006, Society of Plastics Engineers.


nanocomposites. VanderHart et al. [35–38] used solid state 1H-NMR to evaluate

the degradation of the organic clay modifier through resonance positions. Also,

considering the paramagnetic spin effects of the metallic cations present in the

chemical composition of the clay, they utilized the relaxation spin times (T1H)

measured by 13C-NMR to evaluate clay dispersion. In principle, if the clay is

poorly dispersed, the greater is the average distance between each polymer–clay

interface and the weaker is the average paramagnetic contribution to T1H.

According to some authors [4, 39–42], FT-IR is not only adequate to monitor

clay dispersion, but is more efficient than XRD, as it can overcome the limit

interlayer distance of 8 nm inherent to the latter, and less time consuming than

TEM. In a FT-IR spectrum, the clay layers are easily differentiated from the

polymer because of their Si-O bonds, causing a pronounced band between 1300

and 900 cm−1.This band can be decomposed in at least four peaks, three of which

are related to the Si-O in-plane vibrations (or bonds with basal oxygen) and one

to the Si-O− out-of-plane vibration (with apical oxygen or hydroxyl group). The

former appear at about 1120, 1050, and 1020 cm−1, whereas the out-of-plane

vibration peak appears at about 1075 cm−1 [39, 41–44]. The area (or height)

of the in-plane peaks ratio (1050 cm−1/1020 cm−1) increases with increasing

interlayer spacing (intercalation), while the intensity of the peak at 1075 cm−1

rises with the spacing out of the individual clay layers (exfoliation). Also, when

the morphology is ordered and intercalated, the original peak at 1050 cm−1 will

display a negative shift toward lower wavenumbers, while that at 1075 cm−1 will

shift toward higher wavenumbers for high intercalation or partial exfoliation

[39, 40, 42]. Because of the trichroic clay behavior, preferential particle orientation

may induce misleading conclusions in terms of intercalation/exfoliation levels.

This can be overcome with the use of a polarizing lens [39, 42], which enables the

measurement of spectra with different dipole moments. This is necessary for the

subsequent calculation of the structural factor (SF) spectrum, which is equivalent

to the spectrum with no preferential clay orientation.

Finally, tensile testing directly determines the reinforcing effect of the nanosil-

icate [2, 11, 15, 20], although the exact correlation with clay dispersion remains

unclear. The reinforcing aptitude depends on several parameters, including the

level of adhesion between filler and matrix, nanoparticle size and aspect ratio

(individual layers, stacked layers, or tactoids), and nucleating role of the clays

(which can significantly change matrix crystallinity). Early studies [45] showed

a direct dependence of the reinforcing effect on clay dispersion, which was con-

firmed by others [7, 46] mainly for nanocomposites containing high polarity poly-

mer matrices. For nonpolar matrices (like polyethylene or PP), incorporation of

a maleated compatibilizer is often required to increase polymer-clay adhesion,

thus maximizing the dispersion potential. However, it appears that beyond a crit-

ical concentration, the content of MA (maleic anhydride) is damaging, as the clay

reinforcing effect is not able to exceed the reducedmechanical performance of the

highly modifiedmatrix. Bousmina [11] applied various shearing levels to the same

polymer–clay system, obtaining nanocomposites with different dispersion states,

as confirmed by TEM. Yet, as presented in Figure 10.4, the fully exfoliated sample

10.5 Dispersion by Melt Mixing 247

1 μm 1 μm 100 nm

a

0

2

4

6

8

10

12

14

16

18

20

b c

Sample

Young’s

modulu

s (

MP

a)

(A) (B) (C)

Figure 10.4 Polymer–clay nanocomposite showing different dispersion levels and

respective Young’s modulus. Reproduced with permission from [11] © 2006, American

Chemical Society.

exhibited an intermediate value of the Young’smodulus.The author attributed this

behavior to the flexibility of the individual clay sheets.

10.5

Dispersion by Melt Mixing

As Vaia et al. [47] demonstrated, polymer–clay nanocomposites can be obtained

by directmelt intercalation, with attention concentrated on this route following its

industrial relevance [6, 13, 47]. Proper dispersion is achieved when the intrinsic

cohesive forces between clay layers are surpassed by the hydrodynamic stresses

exerted by the polymer melt [3, 11]. Hence, extensive dispersion develops when

four basic conditions are met: (i) an enthalpic driving force must exist for the

polymer to penetrate the clay interlayer; (ii) the interlayer space should be at least


of the same order of magnitude of the diameter of the polymer macromolecular

coil; (iii) sufficient residence time must be secured for diffusion; and (iv) a balance

between stress and strain should occur, in order to attain a thermodynamically sta-

ble structure. During melt compounding, several parameters (screw speed, feed

rate, set temperature, screw geometry) can influence these conditions, and con-

sequently the resulting dispersion level [2, 3, 7, 11–15].

In recent years, the effects of matrix viscosity/molecular weight [48–50],

chemical affinity of the polymer-clay interface [11, 31, 50–53], type of mixer

and mixing protocol [7, 13, 54, 55], and operating conditions (screw speed

(N) and throughput (Q)) have been investigated [7, 12, 14, 29]. For example,

for a polypropylene/polypropylene grafted with maleic anhydride/Cloisite 20A

(PP/PP-g-MA/C20A) system, Lertwimolnun and Vergnes [12] concluded that

higher exfoliation is achieved with low Q/N ratios, that is, higher screw speeds

and lower feed rates (Figure 10.3). It is well accepted that the dispersion mech-

anism of organoclays combines the diffusion of the polymer chains within the

clay galleries (intercalation) and the exertion of sufficient mechanical forces to

delaminate the individual platelets (exfoliation). Generally [7, 11, 54]:

• Intercalation is nearly independent of processing conditions, but sufficient res-

idence time is critical to enable polymer melt diffusion inside the layer spacing;• Exfoliation is extremely dependent on the chemistry, aswell as on themechanics

and physics of the melt mixing process;

• A balance between time for diffusion and deformation (shear or extensional) for

exfoliation is required;

• Ifmeltmixing conditions favor high polymer chainmobility and if proper chem-

ical affinity exists, exfoliation can develop even at low shear rates.

Dennis et al. [7] proposed a dispersion mechanism that includes a diffusion-

controlled route during which shear-controlled fracture of the organoclay

particles takes place, along with the sequential intercalation of the polymer, or

the peeling of the individual clay layers from the top and bottom ends of each

clay stack [3, 7]. The first pathway is chemistry dependent: if the compatibility

between clay surface and matrix is high, well-exfoliated nanocomposites can

be prepared for virtually any set of processing conditions [7]. This hypothesis

has been proposed in most studies of polyamide-based nanocomposites [17, 18,

20, 25, 30, 38, 56, 57]. The second route is valid for marginally compatible

polymer–clay systems, which is the case of most polyolefin-based systems

[12, 14, 19, 29, 52, 54]. Under these circumstances, both chemical affinity and

processing conditions should be optimized in order to attain enhanced clay

dispersion. The third route presumes no compatibility between clay and matrix.

Although processing conditions can be adjusted, it will be difficult to reach a

nanoscale dispersed phase. Increasing shear should lead to improved dispersion

through sliding of adjacent platelets (this requires high shear intensity), followed

by diffusion of the polymer chains into the clay galleries and partial peeling of

the platelets, starting from the edges. More recently, Bousmina [11] showed

that the diffusion of polymeric chains inside the clay galleries is improved under


1.42.2

2.6

3.0

3.4

3.8

1.2

F1 E1 D1 C1 B1 A1F1 E1 D1 C1 B1 A1

Profile n°1 Profile n°14.7 kg h−1

14.4 kg h−1

29.1 kg h−1

4.7 kg h−1

14.4 kg h−1

29.1 kg h−1

1 0.8 0.6 0.4 0.2 0

Axial distance (m)

1.4

100

101

102

103

104

1.2 1 0.8 0.6 0.4 0.2 0

Axial distance (m)

Inte

rlayer

spacin

g d

00

1 (

nm

)

Melt y

ield

str

ess σ

0 (

Pa)

Figure 10.5 Effect of feed rate on the variation along the extruder of interlayer spacing (a)

and melt yield stress (b). Reproduced with permission from [14] © 2007, Society of Plastics

Engineers.

mild shearing conditions (or in a medium to low viscosity matrix) applied

during sufficient time, whereas extensive exfoliation requires a high level of

shearing/deformation.

Most of the above studies involved small-scale experiments [11, 49, 52] and/or

the characterization of samples collected after completing mixing [29], for

example, at the extrusion die exit [7, 13]. The evolution of clay dispersion along

an extruder, where a complex nonisothermal 3D flow develops, seems to be less

well understood. Lertwimolnun and Vergnes [12, 14] characterized postmortem

samples collected from various locations along the axis of a corotating twin

screw extruder (TSE). They concluded that both intercalation and exfoliation

can reach relatively high levels immediately after melting. They also observed

that less restrictive screw profiles yielded better dispersion levels. Furthermore,

depending on the combination of screw profile and operating conditions, these

authors observed an apparent reversion of dispersion evolution along the screw.

This is illustrated in Figure 10.5 for the effect of feed rate. The graph on the left

presents the evolution of the interlayer spacing (d001), as determined by XRD,

which is associated with intercalation.The plot on the right presents the progress

of melt yield stress, which is linked to exfoliation.

10.6

Online and Inline Monitoring of Dispersion

During practical compounding and processing, the parameters that are con-

tinuously monitored (typically temperature, melt pressure, and motor torque)

do not provide sufficient information on the characteristics of the system being

processed. Thus, the possibility of assessing in real time the dispersion of a

nanocomposite on processing is an important scientific and technological target,

as it can be used to assist the definition of material recipes, the optimization

of operating conditions and/or screw design, as well as for quality control and,

ultimately, process control.


Implementing several of the characterization techniques discussed above as

process monitoring tools seems a difficult and expensive task. XRD is feasible

[58], but entails substantial modifications of the processing equipment and,

eventually, the use of synchrotron radiation [59–65]. Rheology has been a more

viable alternative, with academic teams [66–72] and rheometer manufacturers

[56, 57, 73, 74] attempting to develop reliable, precise, simple to operate, and

economically attractive inline and online equipments. Inline fixtures are pre-

ferred, as they avoid melt bypasses [66, 68–71, 75], but are obviously more

difficult to design. Generally, they consist of modified dies. Most commercial

solutions use capillary or slit rheometers to provide online measurements [56, 73,

74] between extruder and die and involve the continuous deviation of the melt

from the main flow to the rheometer, as set by means of a gear pump. Online

solutions seem easier to adapt to the different processing equipments and, at

least conceptually, could also be fixed upstream, along the barrel of the extruder.

Then, it would become possible to monitor the evolution of dispersion along the

extruder axis. In this regard, the Piezo Axial Vibrator is interesting, as its feed

port has the standard geometry of a melt pressure transducer [57]. Generally,

online devices should (i) minimize the time lag between sample collection and

measurement; (ii) prevent material morphology changes both during sampling

and measurement; (iii) enable measurements along the axis of the extruder; and

(iv) be capable of performing measurements at temperatures different from the

processing temperature.

Taking advantage of the modular construction of most corotating TSEs, Covas

et al. [67] developed an online capillary rheometer that can be fixed at different

axial locations and is capable of quickly collecting material samples from within

the extruder. The device was successfully used to study the viscosity evolution

along the screw of various polymer systems, thus contributing to better under-

stand the corresponding mixing and/or chemical mechanisms. Later, an online

rotational rheometer, capable of working either in steady shear or oscillatory

mode, was built [72]. The authors measured the evolution of the rheological

moduli (G′ and G′′) of a noncompatibilized and equivalent compatibilized

polymer blend. It was shown that the use of online rheometry reduces the effects

of degradation and morphology alterations that can take place during sample

collection and subsequent preparation for offline characterization, because of the

successive thermal cycles that the samples are subjected to. An improved auto-

mated version of the prototype was applied to the rheological characterization

of a polypropylene/polypropylene grafted with maleic anhydride/Dellite 67G

(PP/PP-g-MA/D67G) nanocomposite along the axis of an extruder [76]. The G′

and G′′ curves depicted in Figure 10.6, obtained in small amplitude oscillatory

shear (SAOS) at three different axial locations (3, 5, and 7, in the downstream

direction) clearly show a plateau at low frequencies and an increase of the moduli

values, that is, of dispersion, as the material progresses along the screws.

Several spectroscopic techniques have also been efficiently applied to

online/inline process monitoring [77–80]. These include Raman [81–85], ultra

violet-visible (UV-Vis), fluorescence [83–86], and attenuated total reflectance


10−1 100 101 102

ω (rad s−1)

103

104

105

G′ a

nd

G″

(Pa

)T = 200 °C

G′ - 3G′ - 5G′ - 7

G″ - 3G″ - 5G″ - 7

Figure 10.6 Linear viscoelastic behavior of PP/PP-g-MA/D67G nanocomposite at 200 ∘C,measured online at three axial locations (3, 5, and 7, in the downstream direction) along a

corotating twin screw extruder [76].

medium infrared (ATR-IR) [81, 82, 87, 88], ultrasound (US), dielectric probing

[77, 81, 89–91], and NIR spectroscopy [79–81, 92, 93]. Inline monitoring in

polymer processing using NIR is performed mostly in transmission mode, using

a flow-cell fixture attached to a modified extrusion die [79, 81, 83–85, 94–97],

or to the nozzle of an injection molding machine [93, 98]. The polymer industry

has a tradition of using NIR to monitor polymerization, copolymerization, and

depolymerization reactions [96, 97, 99–104], particle size control [96, 105, 106],

and other polymer-related operations [79, 81, 84, 93, 95, 99]. NIR is also com-

monly used to characterize clay minerals, to study their chemical modification,

adsorption mechanisms [107–110], and structure [111, 112]. In the polymer

nanocomposites field, NIR has been utilized to study the reinforcement effect of

the clay [113] and to monitor melt processing [78, 98, 114]. In the first case, NIR

spectrameasured offline showed a direct correlationwith the reinforcement effect

determined by melt extensional measurements. Moghaddam et al. [98] used NIR

to follow the preparation of thermoplastic polyurethane (TPU) nanocomposites

in a small-scale extruder. A decrease of typical urethane bonds during processing

was detected and attributed to the softening of hard segments during the initial

4–6min and to the degradation of TPU [98]. Witschnigg et al. [114] fixed an

NIR probe between extruder and die to study the effect of screw speed and screw

geometry on the properties of polymer–clay nanocomposites. Single parameter

chemometric models based on Young’s modulus, interlayer distance, and drawing

force were developed, good correlations with measured values having been

obtained in some cases. However, confidence on the calibration lines seemed

insufficient to perform a quantitative analysis. Recently, Fischer et al. [78] coupled

NIR, US, and Raman probes to a bypass adapter fixed between extruder and die

to monitor the preparation of nylon-6 nanocomposites with several organoclays.

The correlation between NIR spectra and level of dispersion was based on a single


parameter chemometric model, using the shear thinning power law index, as

proposed by Wagener et al. [25], good results having apparently been obtained.

Because of the complexity of NIR spectra, adequate chemometrics seems essen-

tial to extract from the data asmuch relevant information as possible. Chemomet-

rics is a step-by-step methodology aiming to develop a calibration model relating

the NIR spectral data with the reference characterization parameters [115–118].

To guarantee effective predicting capabilities, it must encompass model develop-

ment and validation [117, 119]. The multivariate calibration technique uses the

entire spectral structures, instead of a single spectral data point, to offer broader

information, and thus detect even minute differences in the sample spectra [116,

119]. During model development, a predefined group of samples (usually des-

ignated as “training samples”) is used to compute the calibration curve, which

directly yields the analyte property from the respective spectra. To ensure pre-

cision, the degree of correlation between spectral and reference data should be

high. For this purpose, a cross-validation step attests the quality of the adjustment

of the data points to the calibration curve. More specifically, a certain number of

the training samples are selected and the predicted property is compared with

the reference values [116, 119]. Finally, the calibration model can be used to pre-

dict the characteristics of unknown samples.Therefore, chemometrics represents

both NIR major advantage and limitation, as an adequate chemometric analysis

may require extensive experimental characterization and data treatment, as well

as the development of multiparameter calibration models [115–119].

The authors used a commercial NIR setup and a comprehensive calibration

model to monitor clay dispersion in a polymer matrix, both in terms of evolu-

tion in time and final level attained, in a Haake batch mixer and in a corotating

TSE. The system comprises three main components (Figure 10.7): (i) a diffuse

reflectance probe (Axiom Analytical Inc.) with a sapphire window having a diam-

eter of 5.7mm; (ii) a Matrix® F (Bruker Optics) spectrophotometer; and (iii) the

OPUS® Quant2 (Bruker Optics) data acquisition and analysis software.The probe

communicates with the spectrometer via a fiber optics cable and the spectrom-

eter connects to the workstation by a LAN-type cable. In the case of the Haake

mixer [120], a threaded hole with Dynisco-type geometry was machined in the

front plate of the mixing chamber to accommodate the NIR probe.The NIR spec-

tra weremeasured with a resolution of 8 cm−1 and accumulation of four scans, the

acquisition time for each spectrum being less than 2 s. During the initial mixing

stages, because of the melting of the matrix and the presence of large voids, the

spectra were generally very noisy and had very low absorbance. Consequently, for

chemometric purposes acquisition started once the onset of equilibrium torque

was reached (this occurred typically after 2min of mixing). Thus, it was possible

to collect 150 spectra during each mixing cycle.

Preparation of a Polypropylene/Polypropylene grafted with Maleic Anhydride/

Cloisite 20A system (PP/PP-g-MA/C20A), using different PP-g-MA con-

tents and operating conditions, was adopted as a case study [120]. To

develop a calibration model, the usefulness of parameters derived from

well-established characterization techniques, able to discriminate between


Front plate Haake Rheomix OS600

Workstation with

OPUS Quant2 software

Diffuse

reflectance

probe

(Axiom Analytical Inc.)

Spectrophotometer Matrix FMATRIX-F (Bruker Opts.)

Fiber

optic

cables

Ethernet connection (LAN wire)

(Thermo Scientific Inc.)

(Bruker Opts.)

Figure 10.7 Inline NIR setup and coupling to the Haake mixer.

distinct dispersion levels, was initially evaluated. Single parameter calibration

models from XRD, rheology, FT-IR, tensile testing, or torque data were globally

inadequate. Multiparameter models showed greater potential, particularly

when incorporating sufficiently performing individual building blocks. Specif-

ically, a 7-parameter model combining parameters from oscillatory rheometry

(G′, G′′, 𝜎0, b), FT-IR (wavenumber shift of the peaks at 1050 and 1080 cm−1),

and thermomechanical data from the mixing equipment (maximum torque)

yielded good results. When applying it to real-timemonitoring of the evolution of

dispersion on the mixing of the same system under distinct operating conditions,

or of the nanocomposite containing different levels of compatibilizer, not only

coherent results were obtained, but they also matched well the forecasted values.

As an example, Figure 10.8 presents the evolution in time of torque and predicted

dispersion level for two case studies. The initial spectrum (at 0 s) is close to

zero. As mixing evolves, both the baseline and the signal intensity increase. As

seen in the torque curve, melting starts after around 10–20 s of mixing. The

increase in torque is because of the conversion of a granular flow into that of

a melt suspension with high solids content. Although the NIR signal is weak,

the fundamental peaks are already visible and clay dispersion is predicted to be

initiated. The torque reaches its maximum at about 30 s and decreases thereafter,

as melting progresses. Melting is probably completed at around 90–120 s, but

it is only after mixing for 180 s that a torque plateau is reached, which most


0

0 0

10

20

30

40

50

60

5

10

15

20

25

30

35

De

gre

e o

f d

isp

ers

ion

(%

)

Pre

dic

tio

n f

rom

NIR

De

gre

e o

f d

isp

ers

ion

(%

)

Pre

dic

tio

n f

rom

NIR

To

rqu

e (

Nm

)

To

rqu

e (

Nm

)

40

0

5

10

15

20

25

30

35

40

0

10

20

30

40

60 120 180 240 300 360 420

Mixing time (s)

0 60 120 180 240 300 360 420

Mixing time (s)

Prediction Prediction87.5/7.5/5125 rpm

125 rpm175 rpm

175 rpmTorqueTorque

(a) (b)

Figure 10.8 Evolution of torque and (predicted) average dispersion with mixing time

during the preparation of a PP/PP-g-MA/Closite 20A nanocomposite in a Haake mixer:

(a) 87.5/7.5/5w/w% at 50 rpm and (b) 90/5/5% at 125 and 175 rpm.

likely corresponds to little further changes in dispersion. In fact, it has been

recurrently reported for several systems melt mixed in batch mixers or twin

screw extruders that most of the dispersive mixing takes place on melting, when

the thermomechanical stresses are higher and little evolution being detected

thereafter [121, 122]. Differences in the NIR spectra up to 180 s are also high,

whereas after that period the spectra are almost superimposed. Thus, the rate of

dispersion is predicted to be higher between 60 and 150 s, a plateau being foreseen

thereafter. The predicted final dispersion level (after 180 s) for the sample with

7.5wt% of PP-g-MA is roughly 30%, which connects well with the normalized

average values of the samples with 5% and 10wt% of PP-g-MA, 26.3 and 36.9%,

respectively. It is also in agreement with reports relating the improvement of clay

dispersion in polyolefin matrices with the increase in compatibilizer content [29,

53, 123]. In the case of the 90/5/5 w/w/w nanocomposite, the final dispersion

level is predicted to increase with rotor speed (approximately 39 and 59% at 125

and 175 rpm, respectively). Again, these values are in line with the normalized

averages (23.5%, 49.6%, 77.2% for 100, 150, and 200 rpm, respectively) and with

reports showing an increase in dispersion with increasing rotor speed [29, 53].

The study above was performed using a diffuse reflectance probe, as the adop-

tion of a probe operating in transmission would imply flow of the material out

of the batch mixer. However, it is important to verify if operation in reflectance

mode is adequate for practical inline monitoring during extrusion. This was

done for a Polypropylene/Polypropylene grafted with Maleic Anhydride/Dellite

67G (PP/PP-g-MA/D67G) system prepared in Leistritz LSM30.34 twin screw

extruder [124]. Figure 10.9 shows the spectra in the two modes for a 90/5/5wt%

composition, prepared under different screw speeds. They are similar in the

region 9000–5000 cm−1, although signal variations in transmission are more

pronounced. In the lower wavenumber region (6000 to 4500 cm−1) the reflec-

tion spectra have worse definition, while the transmission signal shows some

saturation (absorbance above 2.0). Although chemometric models should not


120000.0 0.0

0.1

0.2

0.3

0.4

0.5

0.5

1.0

1.5

2.0

2.5

3.0

10500 9000 7500 6000 4500

PP/PP-g-MA-D67G50 rpm200 rpm

100 rpm300 rpm

Wavenumber (cm−1) Wavenumber (cm−1)

12000 10500 9000 7500 6000 4500

Ab

so

rba

nce

Ab

so

rba

nce

PP/PP-g-MA-D67G50 rpm200 rpm

100 rpm300 rpm

(a) (b)

Figure 10.9 NIR spectra of the PP/PP-g-MA/D67G 90/5/5wt% nanocomposite prepared

under different screw speeds: (a) Transmission and (b) Diffuse reflectance.

be extended to materials or processing conditions outside the range utilized to

create them [123], the same model was used here [120], as one would anticipate

that throughput and screw speed are variables influencing the same dispersion

mechanism.

Table 10.1 ranks the predicted relative degree of dispersion of the samples. Dis-

persion should not increase strictly with output, a maximum being anticipated

at intermediate throughputs. Similar to the effect of screw speed, this behavior

could be because of the conflicting effects of higher hydrodynamic stresses and

lower residence times with increasing feed rate. Nevertheless, other factors may

also come into play, as the higher shear rates associated with the higher outputs

boost viscous dissipation and this may trigger polymer degradation. In turn, the

degraded material will become less viscous and could outflow from between the

clay platelets, thus delaying or even reducing dispersion. At the same location of

the inline measurement, samples were collected from the extruder and character-

ized offline.

The evolution of clay dispersion along the axis of the same TSE was investi-

gated for a PP/PP-g-MA/D67G system [125, 126], utilizing the diffuse reflectance

probe (see setup in Figure 10.7) fixed at various axial barrel locations and con-

tacting directly the melt stream. Using the same chemometric model [120], the

predictions along the extruder are presented in Figure 10.10. The evolution is

consistent with the rheological measurements (not shown). For the three screw

Table 10.1 Average dispersion level of PP/PP-g-MA/D67G 90/5/5wt% nanocomposites pre-

pared under various throughputs.

Relative degree of dispersion (± error) (%)

Q (kg h−1) 1.5 3 6 9

NIR prediction 32.0 (± 1.7) 51.6 (± 1.1) 73.0 (± 3.5) 35.4 (± 4.2)

Normalized average 30.1 (± 3.8) 51.0 (± 3.2) 70.1 (± 3.8) 40.4 (± 4.0)


820

40

60

80

100

10 12 14 16

SC1 SC2 SC3

18 20 22 24 26 28 die

L/D

Degre

e o

f dis

pers

ion (

%)

NIR

pre

dic

tion

Figure 10.10 Inline NIR real-time predictions of the evolution along the screw axis of the

average clay dispersion in a PP/PP-g-MA/D67G system, when using three different screw

profiles (coded SC1, SC2, and SC3).

profiles tested (SC1, SC2, and SC3), dispersion evolves fast along the first mixing

zone (L/D= 9–11) and then beyond L/D= 11 it either remains roughly constant

or regresses. Reversibility is more significant for SC2.

Figure 10.11 displays the effect of screw speed, as determined by inline NIR and

rheometry (melt yield stress, 𝜎0). Inline NIR measurements started at L/D= 10,

when melting appeared mostly completed. As seen in Figure 10.11a, at L/D= 10

dispersion levels are already significant (between 37 and 89%, depending on screw

speed) and increase sharply until L/D= 11 (values between 70 and 93%).Thus, dis-

persion evolved rather quickly, simultaneously with completion ofmatrixmelting,

because of the high stresses and deformations generated by the restrictive screw

elements and the local lower material temperatures. This behavior resembles the

evolution of morphology and chemical conversion during in situ compatibiliza-

tion of polymer blends, with high developing rates onmelting [14]. From L/D= 11

onward, a plateau or a decrease in dispersion is predicted, the latter being espe-

cially perceptible in the die. Rheological measurements (Figure 10.11b) were per-

formed from L/D= 11 onward (in this case, it is essential assure a fully molten

matrix), and show the same trend, that is, a constant or lower dispersion along the

second part of the extruder and die (except for the composite prepared at 50 rpm).

At L/D= 11, melt yield stress increases with screw speed, that is, exfoliation is

promoted. Figure 10.11c,d depicts the data of Figure 10.11a,b, respectively, but

normalized to the values measured at L/D= 11. Both indicate that (i) the rate of

dispersion evolution along the screw decreases with increasing shear rate, that

is, at low screw speed, dispersion progresses axially, at intermediate speeds, little

changes take place, at high screw speed, reversion apparently occurs and (ii) flow

along the die has a negative impact on dispersion. This result was globally con-

firmed by other dispersion assessment techniques, like X-ray and TEM. Reversion

of dispersion with increasing screw speed seems to be induced by the parallel


1020

40

60

80

100

12 14 16

(a)

18 20 22

L/D

De

gre

e o

f d

isp

ers

ion

(%

)

NIR

pre

dic

tio

n

24 26 28 die

50 rpm500 rpm

100 rpm300 rpm

0

1000

2000

σ 0 (

pa

)

3000

4000

5000

(b)

10 12 14 16 18 20 22

L/D

24 26 28 die

50 rpm500 rpm

100 rpm300 rpm

0.50

0.75

1.00

1.25

1.50

(c)

NIR

/NIR

L/D

11

10 12 14 16 18 20 22

L/D

24 26 28 die

50 rpm

300 rpm200 rpm

100 rpm

0.0

0.5

1.0

1.5

2.0

σ 0/σ

0 L

/D 1

1

(d)

10 12 14 16 18 20 22

L/D

24 26 28 die

50 rpm

300 rpm200 rpm100 rpm

Figure 10.11 Effect of screw speed on

the evolution of dispersion of a PP/PP-

g-MA/D67G nanocomposite along the

extruder axis, as determined by: (a) inline

NIR; (b) rheology (melt yield stress); (c) same

data as (a), but normalized to the values

measured at L/D= 11; and (d) same data as

(b), but normalized to the values measured

at L/D= 11.

increase of viscous dissipation and degradation of the clay surfactant. The latter

reduces clay-polymer affinity and, together with the decrease in melt viscosity

affected by viscous dissipation, enables diffusion of the polymer chains out of the

clay galleries. Degradation of the clay surfactant could also induce degradation of

the polymer matrix by chain scission, with a further decrease in viscosity. Rever-

sion of dispersion with increasing feed rate was related to relaxation phenomena.

The evolution of dispersion of a polyamide 6-clay nanocomposite and the

effect of operating conditions were also assessed by NIR [127]. The experimental

procedure followed was identical to that for the PP system discussed above.

Figure 10.12 displays the influence of screw speed and output on final clay

dispersion, as determined by XRD, rheology, and inline NIR (Figures 10.12a, c,

and e refer to the effect of screw speed, Figures 10.12b, d, and e portray the

influence of output). Compounding in the TSE almost doubled the initial

interlayer distance, but the influence of the operating conditions is complicated.

Rheology and NIR indicate an increase of dispersion up to 100 rpm, followed by


0

10−1101

102

103

104

105

PA650 rpm200 rpm

100 rpm300 rpm

g′g′g′

g′g′g′

g′′g′′

g′′g′′

106

100 101

ω (rad s−1)

102 10−1

50

1.5 3.0

NQ

4.5 6.0 7.5 9.0

50

60

70

80

90

100

100 150 200 250 300

Screw speed (rpm)

Throughput (kg h−1)

100 101

ω (rad s−1)

102

2 4 6 8 10

300 rpm (3.35 nm) 9 kg h−1 (3.24 nm)

6 kg h−1 (>8 nm)

3 kg h−1 (3.38 nm)

1.5 kg h−1 (3.31 nm)

Organoclay D43B (1.86 nm)

200 rpm (3.29 nm)

100 rpm (3.38 nm)

50 rpm (3.43 nm)

Organoclay D43B (1.86 nm)

2θ (°)0 2 4 6 8 10

2θ (°)

I (u

.a.)

G’, G

’’ (P

a)

101

102

103

104

105

106

G′,

G′′

(Pa)

NIR

(%

)

I (u

.a.)

PA6

6 kg h−11.5 kg h−1 3 kg h−1

9 kg h−1

g′g′g′

g′g′g′

g′′g′′

g′′g′′

(a) (b)

(c) (d)

(e)

Figure 10.12 Effect of screw speed and

output on final dispersion of a PA6/D43B

nanocomposite: (a) effect of screw speed

(XRD); (b) effect of output (XRD); (c) effect of

screw speed (rheology); (d) effect of output

(rheology); and (e) inline NIR. In the rheo-

logical curves PA6 denotes polymer matrix,

whereas values of screw speed refer to the

composites.

References 259

a decrease at upper speeds. This response is somewhat surprising, as increas-

ing shear intensity resulting from higher screw speeds should stimulate inter-

calation [11]. As far as output is concerned, XRD data (Figure 10.12b) indicate

an increase in intercalation with increasing throughput up to 6 kg h−1 (where the

absence of the diffraction peak suggests a thoroughly exfoliated morphology), fol-

lowed by a sharp drop when working at 9 kg h−1. Identical trend is observed for

the linear viscoelastic response (Figure 10.12d), with G′ reaching the highest val-

ues for the composite produced at 6 kg h−1 and the lowest for 9 kg h−1. NIR data

(Figure 10.12e) also presents a maximum in dispersion at intermediate feed rates,

even if it occurs at 3 kg h−1.

It has been suggested that low feed rates promote intercalation because of the

longer exposure of the material to hydrodynamic stresses [11]. However, this

rise in residence time could stimulate the degradation of the clay surfactant and,

eventually, of the polymer matrix. Such a phenomenon readily explains the lower

rheological moduli. In turn, the corresponding lower melt viscosity would enable

its draining out of the clay galleries, which could then eventually collapse and thus

justify the lower interlayer distance estimated by XRD. Chemical degradation of

the component(s) could also impact on the evolution of dispersion. Therefore,

an optimum set of operating conditions maximizing final clay dispersion seems

to exist.

10.7

Conclusions

In this chapter, the implementation of online and inline monitoring techniques

capable of characterizing average clay dispersion levels during the preparation

of polymer–clay nanocomposites by melt mixing was discussed. Using conven-

tional compounding equipment, such as the Haake mixer and a corotating TSE, it

was demonstrated that amethod usingNIR spectroscopy and suitable chemomet-

rics was able to provide relevant real-time data, which was sensitive to changes in

operating conditions, screw geometry, and material recipe. Also, the procedure

could be readily implemented in industrial production scale. Measurements dur-

ing mixing, or along the axis of the extruder, contributed to better understand the

thermomechanical and chemical aspects involved in the dispersion of organoclays

in polymeric matrices.

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11

Synthesis of Polymer Nanocomposites in Supercritical CO2

Yuvaraj Haldorai and Jae-Jin Shim

11.1

Introduction

One major problem in the engineering application of polymers is their low

stiffness and strength as compared with metals; the moduli are lower by around

100 times, and the strengths around five times. Addition of reinforcing particles

to polymer to form a composite material is often to offset these deficiencies.

Polymers are traditionally reinforced with inorganic fillers to improve their

properties, such as stiffness, toughness, barrier properties, resistance to fire,

and ignition. The effect of fillers on the composite properties depends on the

particle size, concentration, dispersion, and on the interaction with the matrix.

The achievement of a significant improvement in the composite properties often

requires incorporation of a large amount of the filler in the polymer materials

that imparts drawbacks to the composite such as brittleness or opacity. To meet

the rising demands of a wide variety of new applications, functional hybrids

combining inorganic fillers and polymers are being continuously developed so

as to take better benefit from their constituent’s properties or to induce new

ones [1].

Polymer nanocomposites are a new class of hybrid materials that are particle-

filled polymers for which at least one dimension of the dispersed component is

in the nanometer range. The reduction of particle size obtained in such materials

increases the specific surface area of the filler, providing larger matrix/filler inter-

face and so more mutual interactions. As a result, large reinforcing effect may

be reached at much lower filler content when compared to classical microcom-

posites. Besides, the geometrical shape of the particles plays an important role in

determining the properties of the composites [2–4].

One can distinguish three types of nanocomposites, depending on how many

dimensions of the dispersed filler are in the nanometer range. When all the three

dimensions of the particles are in the order of nanometers, the inorganic fillers

are equidimensional, such as spheres (like silica nanoparticles) or cubes (like

calcium carbonate). Fillers with two dimensions in the nanometer scale and the

third in the range of micrometers, forming an elongated structure, include carbon


268 11 Synthesis of Polymer Nanocomposites in Supercritical CO2

nanotubes and whiskers. The third type of nanocomposite is characterized by

only one dimension in the nanometer range. In this case, the filler is present in

the form of sheets of one to a few nanometers thick and hundreds to thousands

nanometers long. Among the different types of sheet-like fillers, those based

on layered silicates (alternatively referred to as clays) have attracted the most

attention, notably because of the easy accessibility and low cost of the starting

clay minerals.

Several advanced polymeric nanocomposites have been synthesizedwith awide

variety of inclusion like metals, semiconductors, carbon nanotubes, andmagnetic

nanoparticles. Many attractive properties of polymers like noncorrosiveness,

lightweight, mechanical strength, and dielectric tenability can be utilized to make

multifunctional materials.The synthesis of polymeric composites usually involves

solution chemistry, and the use of a large amount of organic solvents may raise

serious air and water pollution concerns. Therefore, effective and green methods

are of strong interest. There has been a continuing growth of interest in replacing

conventional organic solvents with environmentally friendly supercritical fluids

(SCFs) in chemical processes. Among them, supercritical carbon dioxide (scCO2)

emerged as an excellent candidate because of its superb characteristics and

properties. Success in applying scCO2 as a solvent or processing medium has

been found in various areas from the well-established supercritical extraction and

separation to the relatively new engineered particle formation [5, 6]. One area

that has seen very much progress is polymer synthesis and processing [7]. From

the aspect of processing, polymers can be fractionated, purified, impregnated,

or foamed using scCO2 as a processing solvent. One of the recent interesting

applications is the synthesis of polymeric nanocomposites [8].

In this chapter, synthesis of polymer nanocomposites by ex situ and in situ

methods in scCO2 are discussed in detail and examples are given. However, it

is impossible to completely describe this field because of the vast number of

published papers on the synthesis of polymer/inorganic filler nanocomposites.

Therefore, this chapter gives a general overview of the techniques and strategies

used for the preparation of nanocomposites. Selected examples that represent

different routes and systems are reported. More detailed descriptions on specific

themes are referred from related references.

11.2

Background on Supercritical CO2

In 1822, BaronCharlesCagniard de laTour experimentally demonstrated a critical

temperature above which a single substance could exist only as a fluid instead of

either being a gas or a liquid [9]. Around 50 years later, when the behavior of highly

compressed gases was a major research interest among scientists, Andrews [10]

discovered the critical phenomenon and supercritical state. He studied the critical

conditions of CO2 and reported the critical values (temperature of 31.1 ∘C and

pressure of 73.8 bar), which are in close agreement with today’s accepted values.

11.2 Background on Supercritical CO2 269

Later in 1879, Hanny andHogarth [11] studied the phenomenon of SCF solubility.

They found that gases could be good solvents under supercritical conditions and

the dissolving power of an SCF is highly pressure-dependent.

Since 1950, the use of CO2 as a green, environmentally benign solvent has

become a widespread, growing reality, both in academia and in industry. Interest

in SCF actually started with research in extraction techniques. For instance, coal

research in the early 1960s pioneered by Kurt Zosel at the Max-Planck-Institute

led tomajor technical breakthroughs in natural product extraction [12]. However,

the first commercial application of scCO2 was the extraction of caffeine from

green coffee at industrial scale between 1975 and 1985. The combined properties

of scCO2 were realized in a cost competitive and environmentally superior

process.

In the twenty-first century, pollution prevention and waste management

became two critical challenges at the forefront. Large amounts of organic sol-

vents are used worldwide every year as reaction media, dispersants, processing,

and cleaning agents. Reducing the emissions of hazardous volatile organic

compounds has been vigorously pursued. The polymer industry, in particular,

is under increasing scrutiny to reduce emission of volatile organic compounds

to completely phase out the use of chlorofluorocarbons (CFCs), and to reduce

the generation of aqueous waste streams. To this account, an environmentally

friendly solvent for chemical synthesis and processing to reduce the emission of

hazardous compounds has gained a keen interest.

In the search for new polymerization solvents, scientists have turned to SCFs.

One of the SCFs, which can be used, is CO2. scCO2 is a clean and versatile solvent

and a promising alternative to noxious organic solvents and CFCs. It has attracted

particular attention as an SCF in the synthesis as well as processing areas for

polymers because of its fluid properties, effects on polymers, and environmental

advantages as a green solvent. The special combination of gas-like diffusivity and

liquid-like density of scCO2 results in an excellent solvent and allows for the solva-

tion of many compounds. Moreover, the use of scCO2 does not create a problem

with respect to the greenhouse effect as it is being conserved during the pro-

cesses. It exhibits changes in solvent density with small changes in temperature

or pressure without altering solvent composition [13]. CO2 is naturally occurring

and abundant: it exists in natural reservoirs of high purity located throughout the

world. In addition, it is generated in large quantities as a by-product in ammo-

nia, hydrogen, and ethanol plants and in electrical power generation stations that

burn fossil fuels [14]. CO2 has an easily accessible critical point with a critical

temperature (Tc) of 31.1∘C and a critical pressure (Pc) of 73.8 bar [15]. CO2 is an

atmospheric gas and is present in the atmosphere in 3.5× 105 parts per billion by

volume. It can thus be easily recycled after use as a solvent to avoid any contribu-

tion to greenhouse effects. Finally, as a nonflammable, nontoxic, and inexpensive

solvent, it is attractive for large-scale synthesis.When scCO2 is used as a polymer-

ization solvent, issues such as drying, solubility, and polymer plasticization are

important to understand.The low viscosity of scCO2 and their ability to plasticize

glassy polymers play a key role on the polymer processing and kinetics. Polymer


plasticization by scCO2 can also be used in the removal of residual monomer

[13]. In this case, traditional solvents are unsuitable. Because CO2 is a gas under

ambient conditions, the polymers can be isolated from the reaction media by

simple depressurization, resulting in a dry polymer product. This feature elimi-

nates energy-intensive drying procedures required in polymer manufacturing to

remove solvent and represents potential cost and energy savings for CO2-based

systems. scCO2 has indeed been shown to be a promising solvent to perform

polymerization reactions.

11.3

Physical and Chemical Properties of scCO2

In general, an SCF is any substance at a temperature and pressure above its critical

point, where distinct liquid and gas phases do not exist [13, 16]. Table 11.1 shows

the critical properties of common SCFs [17, 18]. If the temperature is increased

beyond the Tc at a constant pressure above the Pc, the liquid expands to form

a vapor-like state, the supercritical state, without undergoing a phase transition.

When P>Pc and T >Tc, there is no distinct liquid or vapor phase, and no inter-

face, and only one single supercritical phase exists. Near the critical point, the

density changes dramatically with small changes in pressure and temperature.The

physical properties of scCO2 are most suitably described as functions of density

[16]; many physical properties of scCO2 (e.g., the coefficient of self-diffusion and

the dielectric constant) are closely correlated with the density.

The properties of scCO2 are frequently described as having values between

those of a gas and a liquid (Table 11.2). However, this is not always true, as

properties such as the isothermal compressibility and the heat capacity go

through an extreme in the vicinity of the density at the critical point. The solvent

power of scCO2 is strongly affected by the density, and thus the solvent power

Table 11.1 Critical properties of common SCFs. Reprint with permission from Ref. [17, 18].

Copyright 1999, 1998, Wiley-VCH Verlag GmbH & Co. KGaA; Elsevier respectvely.

Name Chemical formula Molecular weight (gmol−1) Pc (bar) Tc (∘C)

Acetone C3H6O 58.08 47.0 235.1

Carbon dioxide CO2 44.01 73.8 31.1

Chloroform CHCl3 119.38 53.7 263.4

Cyclohexane C6H10 82.15 43.4 287.5

Dichloromethane CH2Cl2 84.93 63.0 237.0

Ethanol CH3CH2OH 46.07 61.4 243.2

Ethane C2H6 282.3 48.8 32.4

n-Hexane C6H14 86.18 30.1 234.4

Methanol CH3OH 32.04 80.9 240.1

Water H2O 18.02 221.2 374.4

11.3 Physical and Chemical Properties of scCO2 271

Table 11.2 Characteristic magnitudes of thermophysical properties of fluids.

Physical properties Liquid Supercritical Gas

𝜌 (Kgm−3) 1000 100–800 1*

𝜂 (Pa S) 10−3 10−5–10−4 10−5

D (m2 s−1) 10−9 10−8 10−5

𝜌 for the density, 𝜂 for the viscosity, D for the diffusion coefficient, and *at ambient temperature

Reprint with permission from Ref. [17]. Copyright 1999, Wiley-VCH Verlag GmbH & Co. KGaA.

can be adjusted via the temperature and pressure. The solvent power is extremely

low, but as the pressure is gradually increased, the density increases strongly

near the critical point as a liquid-like density is approached, accompanied by

a significant increase in solvent power. The solvent power is ultimately deter-

mined by the molecular interactions between the scCO2 and the solute, but the

probability of such interactions increases as the density increases when the gas

is compressed into its supercritical region [13]. Figure 11.1 shows the generic

temperature–pressure phase diagram of CO2 [19]. As previously mentioned,

scCO2 exhibits liquid-like density, but gas-like transport properties of diffusivity

and viscosity.The viscosity of scCO2 is approximately 1 order of magnitude lower

than for typical organic liquids [13], and its self-diffusion coefficient (similar to

the diffusivity in CO2 of a solute molecule of similar size to CO2) is ∼1–2 orders

0−100 −50 0 50 100

Temperature (°C)

Solid Liquid

Vapor

Critical point

(31.1 °C, 73.8 bar)

Triple point

(57 °C, 5.3 bar)

Pre

ssure

(bar)

20

40

60

80

100

Supercriticalfluid

region

Figure 11.1 Generic pressure-temperature diagram. Reprint with permission from Ref. [19].

Copyright 1997, Springer.


of magnitude greater than the diffusivity of small molecules in organic liquids

[13, 20, 21]. It has the ability to greatly swell a polymer matrix, that is, act as a

plasticizing agent and lower the glass transition temperature and thus enhance

mobility (increase diffusion coefficients) of small molecules like monomer, as well

as oligomers and polymer chains [22–25].

11.4

Preparation of Polymer/Inorganic Filler Nanocomposites in Supercritical CO2

11.4.1

Ex SituMethod

The traditional and simplest method of preparing organic–inorganic nanocom-

posites is the directmixing of the fillers into the polymer.Themixing can generally

be done by melt blending and solution blending. The ex situmethod is a popular

one because it does not set a limitation on the nature of nanoparticles and host

polymers to be used.

11.4.1.1 Solution Blending

Solution blending process consists of physical entrapment of inorganic fillers into

the polymer network proceeds through casting and solvent evaporation.The ben-

efit of solution blending is rigorous mixing of clay or carbon nanotubes (CNTs)

with polymer in a solvent that facilitates nanotube deaggregation and dispersion.

This method consists of three steps: dispersion of clay or CNTs in a suitable sol-

vent,mixingwith the polymer (at room temperature or elevated temperature), and

recovery of the nanocomposite by precipitating or casting a film. Both organic

and aqueous mediums have been used to produce nanocomposites [26, 27]. In

this method, dispersion of nanotube can be achieved by magnetic stirring, shear

mixing, reflux, or most commonly, ultrasonication.While solution processing is a

valuable technique for both nanotube dispersion and nanocomposite formation,

it is less suitable for industrial scale processes. For industrial applications, melt

processing is a preferred choice because of its low cost and simplicity to facilitate

large-scale production for commercial applications.

11.4.1.2 Melt Blending

Direct mixing of particles with the polymer melt in technical polymer processes

like extrusion is the classical method for the preparation of composite materials

from thermoplastic polymers. It is widely used for the compounding of clay mate-

rials in polyolefins. Melt compounding is currently explored to a wide range of

materials such as metal oxides and CNTs. Strength of melt compounding is the

large quantity of material that can be produced by extrusion, as most polymer

blends are commercially produced in this way.

Melt blending offers a number of appealing advantages such as no requirement

of a solvent, ease of processing with conventional blending devices such as

11.4 Preparation of Polymer/Inorganic Filler Nanocomposites in Supercritical CO2 273

extruders, relatively low cost, and being environmentally sound. Thus, melt

blending has become the mainstream for preparing polymer/layer silicate

nanocomposites in recent years. The degree of delamination of clay particles and

their dispersion in the polymer matrix are dependent on the details of the melt

intercalation process. Among a variety of shear devices, twin screw extruders

were found to be the most effective because they provide sufficient amount of

shear and intense mixing [28]. It is also believed that chemical compatibility

between the clay and polymer matrix is the most crucial factor in melt intercala-

tion. Only moderate shear intensity is enough to achieve exfoliation structures

for well-compatible polymer/clay system, while both shear and chemical com-

patibilization are necessary for thermoplastic elastomer-based nanocomposites.

Obviously, full exfoliation of clay within nonpolar polymers like thermoplastic

elastomer and polypropylene (PP) remains a great challenge because of the poor

miscibility between the polymer and organoclay. Other strategies have been

developed to improve the intercalation or exfoliation, of which adding a small

amount of compatibilizers was the most popular. Recently, a novel processing

method with the aid of SCFs has gained attention from the plastic industry

[29–32].

In 2002, Manke et al. [33] developed a process that allows clay particles to be

pretreated with scCO2 in a pressurized vessel and then rapidly depressurized

into another vessel at atmospheric pressure to force the clay platelets apart. The

result showed exfoliated clay particles by X-ray diffraction. However, they did

not provide any mechanism for assuring that the exfoliated particles remain

exfoliated when they were combined with the polymer via conventional melt

blending. In 2004, the same group [34] proposed a method to directly inject

scCO2 with the polymer and scCO2 treated clay into an extruder. The polymer

and clay were disposed through two separated hopper into the extruder substan-

tially. The extruder was heated to melt polymer/clay mixture. The pressurized

melt was then got contact with scCO2 (above1100 psi) and the temperature was

controlled below the melting temperature of the polymer. They claimed that the

silicate layers will further exfoliate when melt mixture exits the extruder. No wide

angle X-ray diffraction (WAXD) or transmission electron microscopy (TEM)

evidence of exfoliated morphology was presented. Direct injection of scCO2 into

a molten nanocomposite during melt blending is also promising, as the same

rapid depressurization employed in conjunction with shear may further improve

clay exfoliation. Garcia-Leiner and Lesser [35] reported data for a polyethylene

(PE)/montmorillonite (MMT) nanocomposite processed in a modified single

screw extruder equipped with scCO2 injection near the feed hopper. Their

results showed a 40–100% increase in basal spacing and suggested that scCO2

processing played a significant role in facilitating melt intercalation and clay

dispersion. However, properties of the nanocomposites were not reported.

Nguyen and Baird [36] developed another technique by first saturating the nan-

oclay in scCO2 and then releasing the nanoclay rapidly back through a stopped

extruder filled with polymer pellets. The saturation was conducted in a custom

pressure chamber designed with an inlet for the addition of CO2 and an exit with


a ball valve for the subsequent release of the mixture, which has the ability to

release its contents through an inlet in the second stage of a single screw extruder.

The polymer pellets were loaded into a hopper attached to the extruder with the

ability to trap released clay. The extruder was brought to melt temperature with

the screw turned off.The valve on the pressure chamber was opened and the mix-

ture of clay and scCO2 rapidly expanded throughout the extruder screw and up

into the modified hopper where it immediately mixed with the polymer pellets.

Results from WAXD showed an increase in the exfoliation levels of the scCO2-

produced material as compared to samples produced with dry mixing of polymer

and clay. Mechanical tests showed a 17% improvement in the Young’s modulus

of scCO2-produced composite containing 6.5wt% clay over composite with the

same wt% clay produced from dry mixing of the polymer and clay. The combi-

nation of scCO2-facilitated silicate exfoliation withmelt compoundingmakes this

particular technique ideal for use with polymermatrices that stand to benefit from

the addition of clay but are unable to achieve a suitable level of clay exfoliation

with simple melt compounding. However, the property of the nanocomposites

with nanoclay over concentration of 6.6wt% failed to increase further as the the-

ory predicted. This may be because of the limitation of the facility size or the

procedure itself.

Treece and Oberhauser [37] investigated two different melt blending strategies

for preparing compatibilized PP/clay nanocomposites, specifically: (i) conven-

tional twin screw extrusion and (ii) single screw extrusion capable of direct scCO2

feed to the extruder barrel. Proportional amounts (3 : 1) of maleic anhydride

functionalized PP compatibilizer and organically modified MMT clay were

melt blended. The authors confirmed that a twin screw extruder that provides

sufficient shear and intense mixing is more effective for clay dispersion, and

the improvement of exfoliation with the aid of scCO2 was observed. They also

suggested that pretreating the clay with scCO2 prior to the extrusion may better

improve clay dispersion and exfoliation. Ma et al. [38] achieved a relatively

uniform dispersion of sepiolite in PP using scCO2 even without the aid of

maleated PP as a compatibilizer, which is contrary to an earlier stated theory that

a compatibilizer must be involved in melt intercalation for a nonpolar polymer.

As scCO2 is known as a good solvent and carrier agent for maleic anhydride [39],

the interaction between scCO2 and maleic anhydride affects the natural function

of compatibilizer. The presence of scCO2 in the polymer phase will increase

the interchain distance and free volume and reduce the interchain interactions.

Thus, scCO2 leads to significant changes in properties of polymers such as

low interfacial tension and reduction of viscosity of the polymer melt [38]. In a

different work, Litchfield et al. [40] reported that injecting the organoclay within a

supercritical suspension into a single screw extruder resulted in better dispersion.

Recently, Chen et al. [41] developed an effectivemethod to prepare PP/nanoclay

compositewith improvedmechanical properties. A semicontinuous process using

scCO2 is reported for processing polymer/clay composites with high clay loading

(10wt%) by reducing the collapse of the exfoliated clays. Two major modifica-

tionswere involved in the newprocedure: exfoliating the nanoclay directly into the


hopper filled with polymer pellets followed by processing the composite immedi-

ately and sequentially mixing the clay into the melt.This latter approach helped to

minimize the clay collapse when processing the composites with high clay load-

ings. PP/nanoclay composite at 10wt% nanoclay with improved clay dispersion

was obtained with increased modulus and tensile strength of 63 and 16%, respec-

tively, compared to the pure PPmatrix.They also compared their methodwith the

other techniques such as conventional melt blending, scCO2-aidedmelt blending,

and direct blending with sequential mixing.

TEM images of the 10wt% MMT/PP nanocomposites prepared using the four

different processing methods are presented in Figure 11.2. As can be seen from

Figure 11.2a, the clay aggregation in the direct blended nanocomposite was sig-

nificant with the addition of 10wt% MMT. The system appears to be a phase-

separatedmorphologywith tactoids on the order of hundreds of individual silicate

layers. Apparently the conventional melt intercalation is not effective in exfoliat-

ing/intercalating the nanoclay at this high loading. Better clay dispersion can be

observed in the composite prepared by the scCO2-aided melt blending method in

Figure 11.2b. However, the size of the tactoid is still large. The morphology of the

composite prepared by direct blending with the combination of sequential mixing

500 nm 500 nm

500 nm 500 nm

(a) (b)

(c) (d)

Figure 11.2 Transmission electron micro-

graphs of 10wt% MMT/PP nanocomposites

processed by (a) conventional melt blend-

ing, (b) scCO2-aided melt blending, (c) direct

blending with sequential mixing, and (d)

scCO2-aided melt blending with sequen-

tial mixing method. Reprint with permission

from Ref. [41]. Copyright 2012, Elsevier.


(Figure 11.2b) did not show good dispersion of the nanoclays. The tactoids were

smaller in size compared to Figure 11.2b. Sequential mixing might help to avoid

some further collapsing of the silicate layers but good dispersion could not be

obtained simply because the clays were not delaminated in the first place.The best

dispersion can be seen in the nanocomposite prepared using scCO2-aided melt

blending method with sequential mixing (Figure 11.2d).

The scCO2-aided melt blending method has also been extended to the synthe-

sis of polymer/CNT nanocomposite. Very few works have been reported on the

synthesis of polymer/CNT composites. Recently,Ma et al. [42] reported amethod

that used scCO2 to assist the preparation of PP/CNT composites combined with

batch melt mixing. In this method, a composite with 3wt% CNTwas prepared by

mechanically mixing the polymer melt and CNTs at high temperature in an auto-

clave with the CO2 present under supercritical conditions (15MPa and 200 ∘C).Composites with lower concentrations were obtained by diluting this batch with

pure polymer. Using scCO2-assisted mixing, the yield stress and Young’s modu-

lus of the nanocomposites increased by 33 and 6%, respectively.This improvement

wasmostly because of the reducedmelt viscosity duringmixing as scCO2 acted as

a plasticizer. In addition, this method involves batch processing, a less preferable

process compared to the scCO2-aided continuous extrusion process.

More recently, Chen et al. [43] reported the improvements in CNT dispersion,

and subsequent mechanical properties of CNT/poly(phenylsulfone) composites

were obtained by applying the scCO2-aided melt blending technique. The prepa-

ration process relied on the rapid expansion of the CNTs followed by melt

blending using a single screw extruder. Microscopy results showed improved

CNT dispersion in the polymer matrix and more uniform networks formed

with the use of scCO2, which indicated that CO2-expanded CNTs were easier

to disperse into the polymer matrix during the blending procedure. They also

compared the scCO2-aided melt blending with conventional melt blending

technique. The CNT/polymer composites prepared by the conventional direct

melt-compounding methods did not show any considerable improvements in

the mechanical properties above the addition of 1wt% CNTs because of their

inability to adequately disperse the entangled CNTs into the polymer matrix.

Although melt compounding has shown some advantages for producing com-

posites with improved properties, polymer degradation may be a considerable

issue that should not be overlooked. As a certain high temperature is normally

needed during melt intercalation, not only the polymer matrix and compatibi-

lizer may degrade but also the organic surfactant, which can lead to a significant

reduction in the mechanical properties of the final products.

11.4.2

In SituMethod

The procedure of in situ polymerization involves dispersing the inorganic fillers

directly in the monomer solution prior to a polymerization process. It is obvious

that the most important factors that affect the properties of composites are the


dispersion and the adhesion at the polymer/filler interfaces. Inorganic fillers may

disperse homogeneously in the polymer matrices when they are premodified by

a coupling agent. Furthermore, the resulting materials obtained by this method

also can be easily processed as they have good flowing properties. There are sev-

eral advantages of using the in situ polymerization method. These include ease of

handling, the speed of the process, and better performance of the final products.

The core–shell nanomaterials and nanostructures have become an important

research area in the past few decades because of their potential applications in

various fields as catalysts, industrial and biomedical applications, and so forth

[44]. The core–shell nanocomposites and nanostructures may be with different

sizes and different shapes of core and shell thickness with different surface

morphology. Whenever the surface of the nanoparticles is modified by functional

groups or molecules or coated with a thin layer of other materials, they show

enhanced properties compared to the nonfunctionalized uncoated particles.

The preparation strategy is carried out via polymerization of monomers in the

presence of inorganic fillers. Prior to the dispersion, the inorganic fillers must be

modified with organic materials to improve their compatibility and dispersion. In

particular, emulsion polymerization is a traditionalmethod to produce core–shell

composites. Dispersion polymerization has also been reported to be a viable

method used to prepare core–shell nanocomposites [44]. As a huge number of

articles published were based on the synthesis of core–shell composites in scCO2

via dispersion polymerization rather than emulsion polymerization, we mainly

concentrated on this topic.

11.4.2.1 Synthesis of Nanocomposites by Dispersion Polymerization

Dispersion polymerization [45–47] is one in which the monomer and initiator

are soluble in the reaction medium, while the polymer formed is insoluble. For a

successful dispersion of polymer product in scCO2, specially designed surfactants

have to be used. These amphiphilic surfactants should have CO2-philic segments

that extent out into CO2 medium while CO2-phobic segments, which anchor on

to growing polymeric particles and ensure the steric stabilization on the lattices,

prevent flocculation and precipitation of the reaction product. The dispersion

polymerization process takes place in a homogeneous medium of monomers,

free-radical initiator, and polymeric stabilizer dissolved in a suitable solvent.

At an elevated temperature, the initiator decomposes and generates free-radicals,

which initiate chain growth by the addition of monomers. Once the growing

oligomeric radicals reach a critical molecular weight, phase separation occurs. At

this point the polymer is stabilized as a colloid, and as a result the polymerization

reaction continues to higher degrees of polymerization than the analogous

precipitation reaction in the absence of surfactant.The product from a dispersion

polymerization also exists as spherical polymer particles, but these typically

range in size from 100 nm to 10 μm [45]. Because of the good solubility of many

small organic molecules in CO2, dispersion polymerization constitutes the best

method that has been developed thus far for producing high-molecular weight,

insoluble, industrially important hydrocarbon polymers.


DeSimone et al. [48] reported the first dispersion polymerization of methyl

methacrylate (MMA) in scCO2. Because the product, poly(methyl methacrylate)

(PMMA), is insoluble in scCO2, they used a CO2-soluble fluorinated homopoly-

mer, poly(dihydroperfluorooctylacrylate) (PFOA) as a stabilizer. Consequently,

the successful dispersion polymerization led to a significant improvement in

the yield, molecular weight, and morphology of the resultant polymer. After the

successful preparation of PMMA by dispersion polymerization, researchers have

turned to the synthesis of core–shell composites in scCO2 via this technique.

Wang et al. [49] first prepared polystyrene (PS)/fullerene (C60) composite

microparticles using polydimethylsiloxane methacrylate as a stabilizer through

a one-step seed dispersion polymerization method. C60 was initially dispersed

in styrene (St) with initiator and stabilizer and was then injected into scCO2.

As scCO2 acted as an antisolvent, C60 was precipitated and dispersed in scCO2.

The resulting particles then acted as the seeds for the dispersion polymerization

of St in scCO2, leading to the formation of the PS/C60 composite microparticles.

Substantial investigations in this field have been undertaken by Lim and coau-

thors. They successfully prepared silica (SiO2)/PMMA nanocomposite particles

via dispersion polymerization in scCO2 [50]. SiO2 nanoparticles were first

surface-grafted by the coupling agent 3-(trimethoxysilyl) propyl methacrylate

(MPS), possessing a methacrylate end group that copolymerized with MMA.

The schematic representation for the preparation of nanocomposite is shown in

Scheme 11.1.

SiO2

γ-MPS

OO

O

O

O

Si

γ-MPS functionalized SiO2Core–shell composite

MMA, AIBN

65°C, scCO2

Scheme 11.1 Schematic representation of the process for the synthesis of SiO2/PMMA

composite particles. Reprinted with permission from Ref. [50]. Copyright 2007, Wiley-VCH

Verlag GmbH & Co. KGaA.

The surface modification provided both the good dispersion of the particles in

the polymerization mixture of MMA and CO2 and the anchoring of PMMA on

the SiO2. The polymerization reactions were carried out using 20% MMA (v/v

to CO2), 2% 2,2′ azobis(isobutyronitrile) (AIBN) (w/w to MMA) and different

amount of stabilizer (5, 10, and 15% w/w to MMA) at 65 ∘C with the initial

pressure of 34.5MPa for 12 h. Polymeric stabilizer, poly(dimethylsiloxane)-

b-poly(methacrylic acid) (PDMS-b-PMA) copolymer, provided sufficient

stabilization to the composite latex particles in scCO2 to prevent a flocculation

during the polymerization. The size of composite particles was adjusted with

varying the stabilizer concentration. TEM images of the SiO2/PMMA composite

particles are shown in Figure 11.3. It is clearly evident that the core–shell-type


(a) (b) (c)

388 nm 388 nm 194 nm

Figure 11.3 TEM images of SiO2/PMMA composite particles synthesized using different

amounts of stabilizer (w/w% to MMA) (a) 5%, (b) 10%, and (c) 15%. Reprinted with permis-

sion from Ref. [50]. Copyright 2007, Wiley-VCH Verlag GmbH & Co. KGaA.

morphology was produced successfully. Light PMMA shell coats with the

dark SiO2 cores. Slight aggregation between particles was observed in the

images, which is, however, often seen from the dispersion polymerization with

silicone-based stabilizer in scCO2. From the above results, it is demonstrated that

PDMS-b-PMA stabilizer provided an efficient stabilization for the composite

latex to ensure the formation of spherical particles.

They also prepared PS/SiO2 composite particles in scCO2 using 15% St (w/w

to CO2), 1% AIBN (w/w to monomer), and different amounts of stabilizer

(5 and 10% w/w to St) at 65 ∘C with the initial pressure of 34.5MPa for 40 h

[51]. The random copolymer, poly(1,1-dihydroheptafluorobutylmethacrylate-

co-diisopropylaminoethyl methacrylate) (poly(FBMA-co-DPAEMA)) served as

an effective stabilizer for the polymerization of St in scCO2. It is proposed that

the stabilizer provides steric stabilization on the composite particles in CO2

continuous phase, and the surface-grafted SiO2 particles bearing methacrylate

terminal groups promote the polymer absorption.The same group expanded their

work to synthesize PMMA/TiO2 [52] and poly(divinylbenzene) (PVDB)/SiO2

[53] composites in scCO2. The MPS-modified nanoparticles were well dispersed

in CO2/MMA reaction solution to form stable PMMA/TiO2 composite latexes

by the dispersion polymerization of MMA with the aid of the stabilizer PDMS-

b-PMA. The PDVB/SiO2 composite particles were also successfully synthesized

with the effective stabilization provided by poly(FBMA-co-DPAEMA).

This method was expanded by Shim’s group to prepare core–shell composite

by dispersion polymerization using different metal oxides such as iron oxide and

zirconia (ZrO2). They successfully prepared poly(2-hydroxyethyl methacrylate)

(PHEMA) and magnetic nanoparticle (Fe3O4) nanocomposites by dispersion

polymerization in scCO2 using a random copolymeric stabilizer, poly[(2-

dimethylamino)ethyl methacrylate-co-1H ,1H-perfluorooctyl methacrylate)]

(PDMAEMA-co-PFOMA) [54]. Fe3O4 nanoparticles were first surface modi-

fied by MPS followed by copolymerization with 2-hydroxyethyl methacrylate

(HEMA). The prepared composite particles were nonspherical in shape with the

average particle size of 30 nm. They observed agglomerated composite particles

because of the ineffective stabilization of latex particles in CO2. They also


prepared MPS-modified ZrO2/PMMA composites via dispersion polymerization

in scCO2 using a commercially available stabilizer poly(dimethylsiloxane)-g-

pyrrolidone carboxylic acid (Monosil PCA) [55]. The possible mechanism for the

synthesis of PMMA/ZrO2 hybrid composite is shown in Scheme 11.2.

Zirconia

MPTMS MMA

MPTMS-modified zirconia

Surfactant

AIBN

Zirconia/PMMA composite

Scheme 11.2 Schematic representation for the synthesis of PMMA/ZrO2 composite.

Reprinted with permission from Ref. [55]. Copyright 2011, Wiley-VCH Verlag GmbH & Co.

KgaA.

Zhao and Samulski [56] prepared a partially exfoliated PMMA nanocomposite

using dispersion polymerization in scCO2. The clay was modified using a fluori-

nated surfactant.They found that themodified clay not only acts as inorganic filler

but also as a stabilizer for the growth of PMMA particles. Although the clay is not

soluble inCO2, the stabilizationmechanism is similar to that in a conventional dis-

persion polymerization.This technique was referred as a pseudo-dispersion poly-

merization. They also synthesized PMMA and PS nanocomposites via dispersion

polymerization in scCO2 in the presence of clay that was surface modified using

aminopropyl-terminated PDMS surfactant [57]. This PDMS-based surfactant is

known to be CO2-philic and its longer siloxane chain is expected to provide better

steric stabilization compared to the previously used shorter fluorinated chain.The

morphology of the nanocomposites obtained was strongly dependent on the con-

centration and the type of polymer involved. In the case of PMMA an exfoliated

structure was obtained, whereas in the case of PS a nanocomposite with a par-

tially exfoliated morphology was obtained. In the case of the PMMA/PDMS-clay

nanocomposites where the interaction between PMMA with clay is via hydrogen

bonding, the silicate layers are completely exfoliated and uniformly dispersed in

the PMMAmatrix. While for PS/PDMS-clay nanocomposites where PS interacts

with clay via a weaker van der Waals interaction, the silicate layers are exfoliated

but concentrated mostly on the exterior surfaces of PS particles.

Yue et al. [58] used scCO2 as a medium to synthesize single-walled carbon nan-

otube (SWNT)/PMMA nanocomposite. The ends and sidewalls of the SWNTs

were first functionalized with the coupling agent aminoethylmethacrylate and


then reacted with MMA through in situ radical dispersion polymerization in

scCO2. The functionalization of SWNT surfaces not only provided reaction

sites for the tethering of polymer chains but also facilitated the dispersion of

the SWNTs in scCO2. scCO2 offered high monomer diffusivity for the growth

of the tethered chains; it also introduced the plasticization effect that increased

chain mobility. PMMA chains were found to be covalently tethered to the

nanotubes through copolymerization and formed a thin coating layer on the

SWNT surface. In principle, this simple synthetic procedure can be employed to

prepare other inorganic filler/polymer composites. In addition, carbon dioxide

has an advantage in terms of environmental concerns over organic solvents.

11.4.2.2 Synthesis of Nanocomposites by Other Techniques

In emulsion and suspension polymerization, neither the monomer nor the

polymer is soluble in the continuous phase. Unlike dispersion or precipitation

polymerization, the initial state of emulsion or suspension polymerization is

heterogeneous. The difference between emulsion and suspension polymerization

is that, in emulsion polymerization, a medium-soluble initiator can be dissolved

into the continuous phase, while in suspension polymerization,the initiator is

soluble only in the monomer, not in the polymerization medium. The difficulty

of performing emulsion or suspension polymerizations in CO2 is that most

of the common monomers are CO2-soluble. In addition, proper CO2-soluble

surfactants (emulsifiers) are required to emulsify monomers into the CO2 phase.

However, emulsion polymerization of water-soluble monomers in scCO2 could

be a viable target in the context of green chemistry, given that the commercial

route employs an organic continuous phase and also requires significant energy

input to separate product from emulsion following polymerization. So far, very

few articles have been reported concentrating on the synthesis of nanocomposites

by emulsion polymerizations in scCO2.

Kamrupi et al. [59] prepared PS/silver (Ag) nanocomposite by emulsion poly-

merization. An ex situ dispersion of Ag nanoparticles within the monomer (St)

and subsequent emulsion polymerization using water-in-scCO2 was carried out.

Ag nanoparticles were synthesized by chemical reduction of silver nitrate using

sodium borohydrate as a reducing agent and PDMS as a stabilizer in the water-

in-scCO2. The stable dispersion of Ag nanoparticles was added slowly during the

polymerization of St in the water-in-scCO2, maintaining the temperature at 70 ∘Cand pressure at 20.68MPa.This work represents an easy method to synthesize the

PS/Ag nanocomposite particles. TEM images of the prepared Ag nanoparticles

and PS/Ag nanocomposite particles are shown in Figure 11.4.

The representative TEM images demonstrated the homogeneous dispersion

of Ag nanoparticles in the medium. The Ag nanoparticles were spherical in

shape with a smooth surface morphology. The average size of silver nanoparticles

was 8 nm. TEM image also showed that the Ag nanoparticles were nearly

uniform in size and shape. The TEM image of the PS/Ag nanocomposite clearly

demonstrated that the Ag nanoparticles were encapsulated into the polymer

particles without leaving any bare Ag nanoparticles.The average size of the PS/Ag


25 nm 100 nm

(a) (b)

Figure 11.4 TEM images of (a) Ag nanoparticles and (b) Ag-polystyrene nanocomposite

particles. Reprint with permission from Ref. [59]. Copyright 2010, Elsevier.

nanocomposite particle was determined to be 70 nm. An interesting observation

from the TEM image revealed that in each of the nanocomposite particles, more

than two agglomerated nanoparticles are encapsulated. This may be because of

the agglomeration of Ag nanoparticles during polymerization.They also prepared

copper/PS nanocomposite particles by the same method [60].

In situ intercalative polymerization is an effective method to prepare a poly-

mer/clay composite. In this technique, the layered silicate is swollen within the

liquidmonomer (or amonomer solution) so that the polymer formation can occur

in between the intercalated sheets. 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 cationic exchange inside the interlayer before the swelling

step by the monomer. We can prepare high-performance materials at a relative

low cost by this technique, but this method adapts only to clay minerals, which is

also a significant disadvantage for its application.

Depending on the nature of the components used (layered silicate, organic

cation, and polymer matrix) and the method of preparation, three main types

of composites may be obtained when layered clay is associated with a polymer

(Figure 11.5). When the polymer is unable to intercalate between the silicates

sheets, a phase-separated composite (Figure 11.5a) is obtained, whose properties

stay in the same range as traditional microcomposites. Beyond this classical fam-

ily of composites, two types of nanocomposites can be recovered. In intercalated

structure (Figure 11.5b), a single (and sometimes more than one) extended poly-

mer chain is intercalated between the silicate layers resulting in a well-ordered

multilayer morphology built up with alternating polymeric and inorganic layers.

When the silicate layers are completely and uniformly dispersed in a continuous

polymer matrix, an exfoliated or delaminated structure is obtained (Figure 11.5c).

scCO2 has been widely used as a polar and low-viscosity solvent with the

combination of the in situ polymerization method to prepare nanocomposites.

Zerda et al. [62] used scCO2 for the synthesis of PMMA/organo-MMT nanocom-

posites by mixing organo-MMT, MMA, initiator in the scCO2, in a high pressure


Layered silicate

Phase separated(microcomposite)

Intercalated(nanocomposite)

(a) (b) (c)

Exfoliated(nanocomposite)

Polymer

Figure 11.5 Scheme of different types of

composite arising from the interaction of

layered silicates and polymers: (a) phase-

separated microcomposite; (b) intercalated

nanocomposite; and (c) exfoliated nanocom-

posite. Reprint with permission from Ref.

[61]. Copyright 2000, Elsevier.

apparatus. The primary purpose of the scCO2 was to allow MMA monomers to

readily diffuse and homogeneously disperse within the gallery spacings of the

silicate layers. After the saturation period for mixing, the temperature was raised

to complete the polymerization step. Once polymerization was complete, the

pressure was reduced to atmospheric conditions over a period of 15 h. Removal

of trapped CO2 was accomplished by exposing the samples to temperatures

above the glass transition to allow foaming to occur. The foamed material was

then pulverized and melt processed. This technique produced well-dispersed,

intercalated nanoclay/polymer composites with clay concentration of 40wt%.

Dong et al. [63] employed a similar in situ polymerization technique to prepare

intercalated PS/clay nanocomposites with a more conventional loading (1–10%)

of clay in scCO2. They also modified the clay with a hydrocarbon surfactant and

found that a longer “soaking time” during the impregnating process can lead to

more exfoliated nanocomposites.

Yan et al. [64] also synthesized PS/MMT nanocomposites in scCO2 where

organically modified MMT can be produced through an ion-exchange reaction

between native hydrophilic MMT and an intercalating agent (alkyl ammonium).

Li et al. [65] described a modified synthetic route to produce polymer/clay

nanocomposites where monomer St and initiator were directly intercalated into

organo-MMTwith the aid of scCO2 followed by depressurization and free-radical

polymerization. Nevertheless, in all the above studies, no information on the

yields or morphologies of the polymers has been given. Urbanczyk et al. [66] used

masterbatch technique to prepare polymer/clay nanocomposites. The master-

batch technique is nothing but the preparation of polymer/clay nanocomposites


by combining the in situ intercalative polymerization and the melt intercalation

process. They successfully prepared poly(ε-caprolactone)/clay masterbatches by

in situ intercalative polymerization in scCO2 using stannous octate as a catalyst.

This unique medium allows the reaching of a very high clay loading in the master-

batch. Also, the product obtained after depressurization is an easily recoverable

fine powder. Another advantage of using scCO2 is its capacity to extract the

residual monomer during depressurization, leading directly to a ready-to-use

dry powder. The same group [67] used scCO2 as a polymerization medium

for the polymerization of D,L-lactide in the presence of organomodified clay.

Depending on the functional group borne by the organomodifier, an intercalated

or exfoliated nanocomposite was obtained.

Recently, PMMA/MMT nanocomposites were synthesized via the free rad-

ical polymerization of MMA in the presence of alkyl ammonium-substituted

polysilsesquioxane surfactant-modified clay in scCO2 byHossain et al. [68]. As the

surfactant is three-dimensional, it has been found that the organophilization

with the surfactant enhances hydrophobicity, the stability of the clay separation,

and dimension stability of the nanocomposites. Furthermore, a small amount of

water molecules remarkably changed the surface-free energy of the clay leading

to weaker forces between the layers, which might facilitate the intercalation of

monomer inside the clay gallery. Polymerization media also greatly affected the

dispersion of the organically modified clay resulting in different distributions

of the clay that seemed to play an important role in the morphologies of the

nanocomposites.

Similar to the preparation of polymer/clay nanocomposites, different methods

have been used to synthesize CNT/polymer composites in scCO2. Liu et al. [69]

successfully prepared CNT/PS composite by impregnating St and an initiator

into the CNTs with the aid of scCO2 followed by the polymerization. A composite

material, in which CNTs were consistently filled with PS, was obtained. Dai et al.

[70] prepared CNTs-coated poly(2,4-hexadiyne-1,6-diol) (polyHDiD) with the

aid of scCO2. CNTs were first dispersed in an ethanol solution of HDiD, and

CO2 was then introduced into the mixture. After heating the mixture at 200 ∘C,poly(HDiD)/CNT composites were produced. It was shown that poly(HDiD)

existed in two forms in the composites: either as a coating on the outer surface of

the CNTs with a thickness of less than 10 nm or being impregnated in the inner

cavities of the CNTs. Steinmetz et al. [71] prepared poly(N-vinyl carbazole)/CNT

nanocomposite by scCO2 impregnation technique. However, none of these

reports showed the mechanical properties.

Sol-gel process is a wet-chemical technique widely used in the fields of

materials science and ceramic engineering. The sol-gel processing includes two

approaches: hydrolysis of the metal alkoxides and then polycondenzation of the

hydrolyzed intermediates. This process provides a method for the preparation

of inorganic metal oxides under mild conditions starting from organic metal

alkoxides. The sol-gel processing of fillers inside the polymer dissolved in a

nonaqueous or aqueous solution is the ideal procedure for the formation of

interpenetrating networks between inorganic and organic moieties at the milder


temperature in improving good compatibility and building strong interfacial

interaction between two phases. Materials prepared by sol-gel processing have

uniformity, high purity, and low sintering temperature than that by conventional

solid state reaction. The biggest problem of sol-gel is that the gel process would

lead to a considerable contraction of the internal stress that could result in the

contraction of brittle materials, because of the evaporation of solvents, small

molecules, water.

scCO2 has attracted a great deal of interest as a reaction medium for sol-gel

reaction. Loy et al. [72] reported the first sol-gel process in scCO2 by preparing

SiO2 monoliths via polycondensation of tetramethyl orthosilicate reacting with

formic acid. In another work, SiO2 aerogel particles were prepared in scCO2 by

Sui et al. [73] via reacting silicon alkoxides with acetic acid and formic acid fol-

lowed by depressurization using a rapid expansion of scCO2 process. However,

carboxylic acids (e.g., formic acid and acetic acid) are miscible with scCO2 and

hence are attractive candidates for the polycondensation agents. This direct sol-

gel technique in scCO2 simplifies the aerogel formation process by combining the

polycondensation and supercritical drying into one step, promising a new avenue

for synthesizing SiO2/polymer nanocomposite materials.

Charpentier et al. [74] reported a novel approach to the synthesis of

SiO2/(polyvinyl acetate) (PVAc) nanocomposites using a one-pot synthesis in

scCO2.All the rawmaterials such as tetraethoxysilane (TEOS)/tetramethoxysilane

(TMOS), vinyltrimethoxysilane (VTMO), vinyl acetate (VAc), initiator, and acetic

acid (hydrolysis agent) were introduced into the autoclave. The reactions of

free-radical polymerization of VAc, hydrolysis/condensation of siloxanes, and

coupling of VTMO to the SiO2 particles occurred simultaneously in scCO2.

The result showed that the SiO2 particles in the range of 10–50 nm were

well dispersed throughout the polymer matrix. The particle size of the SiO2

nanoparticles obtained when using TEOS has a smaller diameter and narrower

size distribution than those obtained when using the TMOS precursor. This

may be because of the lower reactivity of TEOS than TMOS in the sol-gel

process [72, 75]. Low reaction rate could reduce the agglomeration of particles

and result in uniform and narrow distribution of particles [73]. Although this

one-pot process can work with essentially any vinyl monomer and silane linker

that are soluble in scCO2, VAc was chosen as the monomer of interest because

of its relatively high solubility in scCO2 [76, 77] and the widespread industrial

and biomedical applications of PVAc. Recently, Wang et al. [78] also utilized

sol-gel technique to synthesize poly(N-isopropyl acrylamide) (PNIPAM)/SiO2

nanocomposite through a one-pot approach in scCO2. The polymerization of

N-isopropylacrylamide and the hydrolysis/condensation of siloxane occurred

simultaneously in scCO2. Vinyltriethoxysilane was applied as the second

monomer for coupling to the sol-gel-derived SiO2 nanoparticles. In the absence

of crosslinker (N,N′-methylenebisacrylamide), some liquid residue was obtained

rather than powder composite particles. The crosslinking reduces the solubility

of PNIPAM chains in scCO2 and the higher crosslinked polymer precipitated out

in the CO2 faster.


11.5

Conclusions

scCO2 being considered as a viable alternative to organic solvents in various

chemical processes, the study of the effect of CO2 on these processes becomes

extremely important in order to make use of scCO2’s “green” enabling properties.

scCO2 has unique physical and chemical properties such as nonpolarity, low

dielectric constant, low viscosity, and zero surface tension, which differentiate it

from conventional organic solvents. Despite its relatively poor solvation power,

scCO2 is still a promising solvent for the preparation of polymer nanocom-

posites. In this chapter, an overview on the synthesis of polymer–inorganic

filler nanocomposites in scCO2 via in situ and ex situ methods was discussed.

The hybridization of nanoparticles and polymers could improve various proper-

ties of resultant nanocomposites. The synergetic enhancements should originate

from the specific attribute of each component. Toward an important objec-

tive for the development of multifunctional nanocomposites is that the bulk

physical properties should be easy to tailor for different purposes. Specifically,

organic polymer-based nanocomposites generally have many advantages such

as long-term stability and good processability, and inorganic nanoparticles

possess outstanding optical, catalytic, electronic, and magnetic properties.

Apart from the properties of individual components in the nanocomposite,

the degree of dispersion of nanoparticles in the polymer and the interfacial

interaction play important roles in enhancing or limiting the overall properties

of the system. Although much work has already been done on various aspects of

polymer/inorganic filler nanocomposites, more research is required in order to

further understand the complex structure–property relationships. The scCO2

technique provides a significant improvement by furnishing a one-step synthesis

route where the potentially recyclable scCO2 works as a solvent, a modification

agent, and a drying agent. This “green” process has potentially many advan-

tages in producing new and unique materials, along with waste-reduction and

energy-saving properties.

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291

Index

aAg/CNT/reninfusion 125

Ag/CNTs hybrid nanoparticles 117

agglomeration process, CB 228

AgNO3 salt 147, 149

bband gap energy evolution 155

biodegradable polymers

– drug delivery systems, see drug delivery

systems

– HAp nanocomposites 76–78

biological micro-system application 137

bis[3-triethoxysilylpropyl-]tetrasulfide

(TESPT) 6

ccarbon nanofibers (CNFs) 93–96

carbon nanotubes (CNTs) 93–96, 115, 116

catalytic assays 116

cellulose nanofibres 96

ceramic based piezoelectrics 140

chlorinated polyethylene (CPE) 33

clay modification 99–102

clay slurry 193

CNTs and Ag/CNTs fillers 122

compatibilizer

– CPE25 nanocomposites 44

– nanofillers 32

– polymer blends and fillers 33

– polymer crystallinity 39

– solution blending technique 33

– viscosity and elasticity 44

compressive response plot, composites 124

conductive carbon fillers

– carbon black (CB) 215

– CNFs 216

– CNTs 216

– graphene 216

– types 217

cone calorimetry 197

covalently sp2-bonded carbon atoms,

graphene 31

ddifferential scanning calorimetry analysis 122

dimethyl formamide (DMF) 117

dispersion methods, chain-growth

– anionic polymerization 226

– free radical polymerization 225, 226

– nanofillers 224

– olefin polymerization 224, 225

– ring-opening polymerization 226, 227

dispersion polymerization

– description 277

– iron oxide and zirconia (ZrO2) 279

– methyl methacrylate (MMA), scCO2 278

– oligomeric radicals 277

– PMMA/ZrO2 composite synthesis 280

– polystyrene (PS)/fullerene (C60) composite

microparticles 278

– PS/SiO2 composite particles 279

– SiO2/PMMA composite particles 278, 279

– steric stabilization 280

– surface-modification 278

– SWNT/PMMA nanocomposites 280

0-D nanoparticles 93

1-D nanofillers 97

drug delivery systems

– adsorption and release 73

– alginate/HAp microspheres 76

– bioactive molecules 73

– biocomposites 73

– bone morphogenetic proteins (BMPs) 73

– CDHA/chitosan nanocomposites 74

– chitosan microspheres (CMs) 75, 76


292 Index

drug delivery systems (contd.)

– electrospun scaffolds 74

– gelatin/nano-HAp scaffold 73

– HAps 72

– microwave irradiation method 76

– minocycline 73

– novel coaxial electrospun PCL/PVA

core-sheath nanofibers 74

– osteomyelitis 74

– Staphylococcus aureus 75

eelectrically conductive polymer

nanocomposites 212

electromagnetic interference (EMI)

shielding 118

electron-beam synthesis, polymer–NCs

164, 165

electron-beam writing (EBW) 160,

163, 164

electrospinning (ES)/extrusion-based

fabrication methods 145

electrostatic charge dissipation (ESD) 212

elongational methods 172

emulsion polymerization

– epoxy/HNT/carbon fiber hybrid

nanocomposites 15

– graphene nanosheets 13, 14

– HIPS/HNT nanocomposites 14, 15

– PANI/AC nanocomposites 14, 15

– PMMA, polymer matrix 13

– PS/carbon black (CB) nanocomposites

11

– ultrasound 15

– XRD signals 17

energy harvesting applications

– maximum power generation 138

– mechanical energy sensor 137

– power storage circuit 137

– PVDF based materials 139, 140

evolution, Young’s modulus 192

ex-situmethod

– description 272

– melt blending, seemelt blending

– solution blending 272

exfoliation adsorption process

– description 1

– emulsion polymerization, see emulsion

polymerization

– solution intercalation method, see solution

intercalation method

– structure 2, 3

expanded graphite oxide (EGO) 7

ffiller distribution and dispersion 216

filler orientation and alignment 217, 218

flame retardants (FRs) 197, 198

flexible conductors/resistors 118

flexible field-emission devices 118

flexural stress–strain curves 125

flyback switching mode DC–DC

converter 138

FT-IR absorption spectra, nanocomposite

fibers 169, 171

ggeneric pressure-temperature diagram,

scCO2 271

graphene oxide 32–36

graphene-based conducting polymer

nanocomposites 219

graphene-based nanofillers 107, 108

graphite oxide and graphene oxide

preparation 34

green composites 198, 199, 201

hHAp/DNA nanocomplexes

– arginine modified nano-HAp 64

– Ca/P stoichiometry 64

– calcium phosphate nanoparticles 62

– double helix of B-DNA 63

– ephrinB2 gene 65

– gel electrophoresis analysis 63

– internalization mechanisms 62

– molecular dynamic simulations 63

– nanocapsules 63, 64

– non-viral gene therapy 61

– novel 3D scaffolds 65

– plasmid DNA (pDNA) 62

heat release rate (HRR) 197, 198

high density polyethylene (HDPE) 31, 33, 36,

38, 40, 42, 44, 46, 47

Hoffman elimination reaction 196

hybrid nanocomposite fibers 169, 170

hydroxyapatite (HAp) nanocomposites

– applications 51

– biocompatibility and nontoxicity 59

– biodegradable and biobased polymers 51

– biological performance 58

– chemically identified phases 59

– description 51

– drug delivery systems, see drug delivery

systems

– electrospinning 61

– ex situ approach 61

Index 293

– HAp/DNA nanocomplexes, gene carriers,

see HAp/DNA nanocomplexes

– hexanoic and dodecanoic acids 60

– inorganic, organic polymer matrix 60

– nanorods 51

– non-grafted particles 61

– physico-chemical methods 60

– surfactant molecules 60

– thermo-mechanical methods 60

– wetting 60

iin situ polymerization 243

in situ polymerization intercalation

– advantages 16

– benzoxazine (Bz)-MMT clay 21

– description 3

– epoxy/graphite nanocomposites 17, 18

– Nylon-6/clay nanocomposites 3

– polyaniline (PANI)/MWCNT

composites 22

– polymerizing conditions 23

– PP/GO nanocomposite synthesis 17, 19

– PPy/GO nanocomposites 19, 20

– PSU/MMT nanocomposites 21, 22

– terephthalate-intercalated LDH 21

– XRD results 19–21

in situ polymerization, polymers

– alignment, conductive fillers 227, 228

– assistant technologies 221

– chain formation 219

– covalent modification 222, 223

– epoxy-based nanocomposites 220

– non-covalent functionalization 223, 224

– polymer matrices 219

– shear force/sonication 221, 222

– surfactant technology 222

in-situmethod

– Ag nanoparticles and Ag-polystyrene

nanocomposite particles 282

– core-shell nanomaterials and

nanostructures 277

– description 276

– emulsion and suspension polymerization

281

– in-situ intercalative polymerization 282

– nanocomposites synthesis, dispersion

polymerization, see dispersion

polymerization

– polyHDiD 284

– polymerization media 284

– in scCO2 282, 283

– silicates and polymers, layered 282, 283

– SiO2/(polyvinyl acetate) (PVAc)nanocomposites 285

– sol-gel process 284

– vinyltriethoxysilane 285

llamellar structure, insoluble molecules

of Cd 150

layered aluminosilicate clays

– description 97, 98

– effects, clay modification 99–102

– mixing methods 98, 99

– MMT, pre-polymer synthesis 102, 103

– properties and multiphase nanocomposites

103, 105

– vinyl ester-clay nanocomposites 106

layered double hydroxides (LDH) 106, 107

LDPE/5% cellulose fibres composites 200

light irradiation and electron beam

writing 156

lightweight/flexible conducting materials 211

linear low density polyethylene (LLDPE) 33

mmelt blending

– advantages 272

– CNT/polymer composites 276

– compatibilized PP/clay

nanocomposites 274

– compatibilizer 274

– metal oxides and CNTs 272

– modifications 274

– polyethylene (PE)/montmorillonite (MMT)

nanocomposites 273

– polymer and organoclay miscibility 273

– polymer pellets 273

– PP/CNT composites 276

– shear intensity 273

– transmission electron microscopic (TEM)

images 275

– WAXD/TEM 273

– x-ray diffraction 273

– Young’s modulus 274

melt intercalation

– CNT nanocomposites 4

– description 1, 2

– entropy loss 4

– intercalated/exfoliated 6

– LDH/HDPE 6

– Nanofil5Ⓡ, C30BⓇ, Nanofil2Ⓡ 5

– PCL-MWCNT 5

– PCL/MMT 5

– PEO/Li-MMT 8, 9

– PMMA/MWCNT and PS/MWCNT 6

294 Index

melt intercalation (contd.)

– polystyrene (PS)/organoclay

nanocomposites 4

– PPMA/EGO 7

– thermoplastic polymer nanocomposites 1

– TPU/C15AⓇ clay nanocomposites 7

– weak electrostatic forces 1

– XRD patterns 6, 7

melt mixing

– description 243

– diffusion controlled routes 248

– extruder of interlayer spacing and melt yield

stress 249

– industrial relevance 247

– intercalation and exfoliation 248

– levels 247, 248

– low Q/N ratios 248

– twin screw extruder (TSE) 249

metal oxides 91–93

metal salts 147

metallic conductive pathways 232, 233

metalsulphide NPs 147, 149

microwave heating 117

multi walled carbon nanotubes

(MWCNT) 32

multi-photon irradiation 160

multicolor laser processing 159

multilevel lithography and patterning 164

nnanocomposite fabrication methods 218, 219

nanocomposite generation 35, 36

nanocomposite polymer fibers

– electrospun nanofibers 167

– organic nanostructures 166

– photo-assisted synthesis 167

– thermal-assisted synthesis 169

– uniaxial stretching and whipping

motion 167

nanocomposites

– nanoscale additives 88

– one-dimensional nanofillers, see

one-dimensional nanofillers

– polyesters 87

– saturated polyester (SP) resins 87

– two-dimensional nanofillers, see

two-dimensional nanofillers

– vinyl ester (VE) resins 87

– zero-dimensional nanofillers, see

zero-dimensional nanofillers

nanocrystals, HAp

– amorphous calcium phosphate 53

– anionic and cationic surfactants 56

– calcium and phosphate ions 52, 55

– chemical precipitation routes 52

– 1-D HAp nanorods 57

– D-sorbitol, polyethylene glycol, gelatin 57

– hierarchically nanostructured,

nanosheets 55

– laboratory-scale synthesis 53

– macromolecules 56

– morphologies 53, 54

– nanometric low-dimensional forms 55

– nanoparticles formation 58, 59

– nucleation and growth mechanisms,

CAp 53

– potassium sodium tartrate 55

– SBF-based solution 56

– sol-gel method 57

– surfactant cation, phosphate anion and

hydroxyl groups 56, 57

– surfactant emulsion systems 58

– ultrasonic irradiation 58

nanofabrication methods 165

nanotechnology 115

neat and nanoparticles infused epoxy

124, 126

nobel metal nanoparticles 115

nucleation process 152

oodour stabilization 200

on-line and in-line monitoring

– on-line capillary rheometer 250

– parameters 249

– Piezo axial vibrator 250

– rheology 250

one-dimensional nanofillers 93

optical spectroscopy experiments 154, 155

optoelectronic devices and sensors 172

organoclay dispersion 184, 185

organometallic compounds 147, 151

oxygen consumption principle 197

pp-MMT, polyamide matrix 189

PCL, see poly(ε-caprolactone) (PCL)percolation theory 213, 214

phase-separated microcomposite 182

photoluminescence spectra 162

piezoelectric charge and voltage constant 136

PMMA, see poly(methyl methacrylate)

(PMMA)

polarization and fibre formation 140

poly(ε-caprolactone) (PCL) 5

poly(2,4-hexadiyne-1,6-diol) (polyHDiD) 284

poly(ethyl methacrylate) (PEMA) 139

poly(methyl methacrylate) (PMMA) 32, 139

Index 295

poly(propylene-g-maleic anhydride)

(PPMA) 7

poly(vinylidene fluoride) (PVDF)

– dipole moments 133

– energy harvesting applications, see energy

harvesting applications

– free-radical polymerisation 131

– γ-phase 134

– physical and electrical characteristics 131

– piezoelectricity 132

– polymorphs and their characteristics 133

– relationships and equations 135

– sensing characteristics 132

– stretching and corona poling 134

– suspension and emulsion polymerizations

131

– trans-gauche-trans-gauchey′ (TGTG′)conformation 133

polymer composites 121, 124, 125

polymer layered silicates nanocomposites

195

polymer nanocomposites, see also supercritical

carbon dioxide (scCO2)– cellulose-silver nanocomposites 25

– description 1

– exfoliation adsorption, see exfoliation

adsorption process

– FESEM and TEM 25, 26

– fluorinated tin oxide glass (FTO) 25

– graphene-polymer composite synthesis

24, 25

– in situ polymerization, see in situ

polymerization intercalation

– intermatrix synthesis (IMS) 24

– matrices 145

– melt intercalation, seemelt intercalation

– phase separation phenomena 145

– Sn/carbon-silica composite synthesis

23, 24

– sulfonated polyethersulfone with Cardo

group (SPES-C) 24

– synthesis 23

– types, composite microstructures 2

polymer-clay nanocomposites

– Carreau–Yasuda model 244

– clay tactoids and exfoliated platelets 242

– dispersion, see dispersion

– dispersion levels and Young’s modulus

246, 247

– high dispersion levels 242

– high polarity polymer matrices 246

– immiscible composites 241

– in situ polymerization 243

– layered silicates 241

– melt mixing 243

– melt yield stress, filling ratio 245

– morphologies 241, 242

– NMR and FT-IR 245

– rheological response 244

– Si-O in-plane vibrations 246

– solution intercalation 243

– structural factor (SF) spectrum 246

– TEM and XRD 243, 244

polymer/clay nanocomposites (PCNs) 179

polymer/inorganic filler nanocomposites

– ex-situmethod, see ex-situmethod

– in-situmethod, see in-situmethod

polyolefins 31

PP/PP-g-MA/D67G nanocomposites

– average dispersion level 255

– chemical degradation, components 259

– chemometrics 252

– description 250

– in-line NIR real-time predictions 255, 256

– in-line NIR set-up and coupling, Haake

mixer 252, 253

– linear viscoelastic behavior 250, 251

– multi-parameter models 253

– NIR spectra 254, 255

– PA6/D43B nanocomposite 257–259

– polymers 251

– PP/PP-g-MA/Closite 20A 252–254

– screw speed effect 256, 257

– single parameter calibration models 253

– single parameter chemometric models

251

– spectroscopic techniques 250

– TPU nanocomposites 251

– transmission mode, flow-cell fixture 251

PPMA, see poly(propylene-g-maleic

anhydride) (PPMA)

pre-polymer synthesis 102, 103

precursor-doped polymer surface 166

PVDF based materials

– energy harvesting applications 139, 140

– physical properties, polymers 138

– polymer/polymer blends 139

qquantum dots (QDs) 156

quaternary ammonium surfactants 198

rreduced frictional degradation 200

rGO-based polymer nanocomposites

– chemical reduction method 229, 230

– in situ thermal reduction 230, 231

296 Index

rGO-based polymer nanocomposites (contd.)

– integrated chemical & thermal reduction

231, 232

room temperature nanoimprint lithography

(RT-NIL) 164

ssaturated polyester (SP) resins 87

screw configuration, twin screw

extruder 187

short chain thiolates 152

silicon-containing nanospheres 89–91

silver nanoparticles 115, 118

single screw melt extruder 117

single walled carbon nanotubes (SWCNT)

32

single walled nanotubes (SWNTs) 116

solar electromagnetic radiation 232

solution blending technique 33, 272

solution intercalation method

– description 9

– Fe3O4/MWCNT/chitosan nanocomposite

synthesis 11, 12

– foam morphology and properties 11

– laponite modification 10

– polyamide (PA)/MWCNTs nanocomposites

10

– of polymer/pre-polymer 243

– PPC/MGO nanocomposites 10

– SBR/graphene nanocomposite, solution

mixing 9

steam flushing 200

supercritical carbon dioxide (scCO2)– clean and versatile solvent 269

– functional hybrids 267

– geometrical shape particles 267

– greenhouse effects 269

– inorganic fillers 267

– nanocomposites types 267

– organic solvents 269

– physical and chemical properties 270, 271

– polymer/inorganic filler nanocomposites, see

polymer/inorganic filler nanocomposites

– polymer manufacturing 270

– processing solvent 268

– SCF, see supercritical fluid (SCF) solubility

– solution chemistry 268

supercritical fluid (SCF) solubility

– critical properties 270

– description 269

– extraction techniques 269

– organic solvents and chlorofluorocarbons

269

surface-photografted poly(acrylic acid) (PAA)

156

ttensile modulus and yield stress 37

TESPT, see bis[3-triethoxysilylpropyl-]

tetrasulfide (TESPT)

TGA analysis, nanoparticles 121, 122

thermal assisted in-situ synthesis

– agglomeration and clustering

phenomena 153

– gas/liquid phase 147

– long chain thiolates 152

– nanocomposite materials 146

– organometallic compounds 151

– polymer viscosity 146

– precursors chemical structure 146

thermo-gravimetric analysis (TGA) 195

thermoplastic polyurethane (TPU)

221, 251

tissue engineering

– alginate 67

– biodegradable alanine-substituted

polyphosphazene 71

– biological properties 70

– bionanocomposite materials 65

– bionanocomposites preparation 67

– chemical structures, “egg-box” model 68

– chitosan electrospun nanofibers

70, 71

– collagen protein fibers, proteo-glycans and

hydroxyapatite nanocrystals 66

– derived nanofiber scaffolds 69

– 3D nanofibrous polymer/HAp mineral

biocomposites 70

– electrospinning 66

– electrospraying 71

– electrospun scaffolds 70

– HAp particles, electrospun fibers 66, 67

– HARV bioreactors 69

– hydroxyapatite-biopolymer nanofiber

mats 68

– injectable hydrogels 72

– inorganic nanofillers 65

– interfacial adhesion 69

– macrostructure scaffolds 65, 66

– osteoconductive 65

– PLA, PLGA, PCL and PHB 68

– PLLA-g-HAp 69

– poly(D,L-lactide) (PDLLA)/nanosized HAp

composite resins 72

– poly(ethylene glycol) (PEG) 69

– polyvinyl alcohol (PVA) 70

– stereolithography 71, 72

Index 297

TPU, see thermoplastic polyurethane (TPU)

transmission electron microscopy (TEM) 37,

116, 120

twin screw extruder (TSE) 249

two-dimensional nanofillers

– graphene-based nanofillers, see

graphene-based nanofillers

– layered aluminosilicate clays, see layered

aluminosilicate clays

– LDH 106

uultraviolet and visible radiation 157–159

vvinyl ester (VE)

– clay nanocomposites 106

– resins 87

wwater-assisted extrusion

– clays 180, 181

– nanotechnology 179

– organomodification, clay 179

– PCNs 179

– preparative methods 183–187, 189, 190,

192, 194, 195

– water-dispersible commercial nanofillers

180

WeiprenⓇ 6025 33

wide angle X-ray diffraction (WAXD) 273

wood polymer composites (WPCs) 199

WPC, see wood polymer composites (WPCs)

xX-ray diffraction (XRD) 118–121

zzero-dimensional nanofillers

– 0-D nanoparticles 93

– metal oxides, seemetal oxides

– silicon-containing nanospheres, see

silicon-containing nanospheres

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