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- &
Macrocomposite (Volume 2)
2012
Print ISBN: 978-3-527-32879-6
Mittal, V. (ed.)
Characterization Techniques
for Polymer NanocompositesSeries: Polymer Nano-, Micro- &
Macrocomposite (Volume 3)
2012
Print ISBN: 978-3-527-33148-2
Mittal, V. (ed.)
Modeling and Prediction of
Polymer Nanocomposite
PropertiesSeries: Polymer Nano-, Micro- &
Macrocomposite (Volume 4)
2013
Print ISBN: 978-3-527-33150-5
Mittal, V. (ed.)
Thermoset NanocompositesSeries: Polymer Nano-, Micro- &
Macrocomposite (Volume 5)
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 Results and Discussion 119
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
7.1 Introduction 145
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
8.1 Introduction 179
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 Introduction 211
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
10.1 Introduction 241
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
11.1 Introduction 267
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
The Petroleum Institute
Department of Chemical
Engineering
Bu Hasa Building
Room 2204
2533 Abu Dhabi
United Arab Emirates
Jose Antonio Covas
University of Minho
Institute of Polymers and
Composites (IPC/I3N)
Campus de Azurém
4800-058 Guimarães
Portugal
Luis J. del Valle
Universitat Politècnica de
Catalunya
Departament d’Enginyeria
Quımica
Avinguda Diagonal 647
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
National Nanotechnology
Laboratory of Istituto
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
The Petroleum Institute
Department of Chemical
Engineering
Bu Hasa Building
Room 2204
2533 Abu Dhabi
United Arab Emirates
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
Technical Unit of Material
Technologies Brindisi
Strada Statale 7 Appia km. 706
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
Institute of Polymers and
Composites (IPC/I3N)
Campus de Azurém
4800-058 Guimarães
Portugal
VikasMittal
The Petroleum Institute
Department of Chemical
Engineering
Bu Hasa Building, Room 2204
2533 Abu Dhabi
United Arab Emirates
Luana Persano
National Nanotechnology
Laboratory of Istituto
Nanoscienze-CNR
via Arnesano
73100 Lecce
Italy
Dario Pisignano
National Nanotechnology
Laboratory of Istituto
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
Universitat Politècnica de
Catalunya
Departament d’Enginyeria
Quımica
Avinguda Diagonal 647
08028 Barcelona
Spain
Vijaya K. Rangari
Tuskegee University
Department of Materials Science
and Engineering
100 James Center
Tuskegee, AL 36088
USA
Vincenzo Resta
ENEA
Technical Unit of Material
Technologies Brindisi
Strada Statale 7 Appia km. 706
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
1348 Louvain-la-Neuve
Belgium
XVI List of Contributors
Michel Sclavons
UCL-IMCN/BSMA
Croix du Sud 1
L7.04.02
1348 Louvain-la-Neuve
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
University of Bolton
Institute for Materials Research
and Innovation
Bolton, BL3 5AB
UK
Navneet Soin
University of Bolton
Institute for Materials Research
and Innovation
Bolton, BL3 5AB
UK
Leander Tapfer
ENEA
Technical Unit of Material
Technologies Brindisi
Strada Statale 7 Appia km. 706
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.
4 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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
1.2 Synthesis Methods 5
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.
6 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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
1.2 Synthesis Methods 7
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
8 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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.
1.2 Synthesis Methods 9
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.
10 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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
1.2 Synthesis Methods 11
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
12 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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.
1.2 Synthesis Methods 13
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
14 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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
1.2 Synthesis Methods 15
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.
16 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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.
Reproduced from [41] with permission from Elsevier.
[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
1.2 Synthesis Methods 17
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]
with permission from Elsevier.
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
18 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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.
1.2 Synthesis Methods 19
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
20 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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,
1.2 Synthesis Methods 21
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
Reproduced from [50] with permission from Elsevier.
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.
22 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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
1.2 Synthesis Methods 23
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
24 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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.
1.2 Synthesis Methods 25
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.
26 1 Synthesis of Polymer Nanocomposites: Review of Various Techniques
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
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.
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.
34 2 Masterbatch Approach to Generate HDPE/CPE/Graphene Nanocomposites
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.
36 2 Masterbatch Approach to Generate HDPE/CPE/Graphene Nanocomposites
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%
2.3 Results and Discussion 37
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
38 2 Masterbatch Approach to Generate HDPE/CPE/Graphene Nanocomposites
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
2.3 Results and Discussion 39
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
40 2 Masterbatch Approach to Generate HDPE/CPE/Graphene Nanocomposites
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
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
2.3 Results and Discussion 41
0.11000
10 000
100 000G
″, P
a
10 000
100 000
G″,
Pa
1 10 100
HDPEHDPE/GHDPE/5% CPE25/G
HDPE/1% CPE25/GHDPE/2% CPE25/GHDPE/5% CPE25/GHDPE/10% CPE25/G
HDPE/10% CPE25/GHDPE/5% CPE35/G
HDPE/10% CPE35/G
Angular frequency, rad s−1
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
Angular frequency, rad s−1
0.1 1 10 100
Angular frequency, rad s−1
HDPEHDPE/GHDPE/5% CPE25/GHDPE/10% CPE25/GHDPE/5% CPE35/GHDPE/10% CPE35/G
HDPE/1% CPE25/GHDPE/2% CPE25/GHDPE/5% CPE25/GHDPE/10% CPE25/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
42 2 Masterbatch Approach to Generate HDPE/CPE/Graphene Nanocomposites
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
2.3 Results and Discussion 43
0.1 1 10 100
Angular frequency, rad s−1
0.1 1 10 100
Angular frequency, rad s−1
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
Angular frequency, rad s−1
0.1 1 10 100
Angular frequency, rad s−1
1000
10 000
η″,
Pa
.s
1000
10 000
η″,
Pa
.s
HDPEHDPE/GHDPE/5% CPE25/GHDPE/10% CPE25/GHDPE/5% CPE35/GHDPE/10% CPE35/G
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
Angular frequency, rad s−1
0.1 1 10 100
Angular frequency, rad s−1
1000
10 000
100 000
η*,
Pa
.s
1000
10 000
100 000
η*,
Pa
.s
HDPEHDPE/GHDPE/5% CPE25/GHDPE/10% CPE25/GHDPE/5% CPE35/GHDPE/10% CPE35/G
HDPE/1% CPE25/GHDPE/2% CPE25/GHDPE/5% CPE25/GHDPE/10% CPE25/G
(a) (b)
Figure 2.10 (a,b) Complex viscosity of HDPE and HDPE nanocomposites.
44 2 Masterbatch Approach to Generate HDPE/CPE/Graphene 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.
2.3 Results and Discussion 45
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.
46 2 Masterbatch Approach to Generate HDPE/CPE/Graphene Nanocomposites
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.
48 2 Masterbatch Approach to Generate HDPE/CPE/Graphene Nanocomposites
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
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.
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].
3.2 Preparation of HAp Nanocrystals 53
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
54 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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.
3.2 Preparation of HAp Nanocrystals 55
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
56 3 Preparation and Applications of Hydroxyapatite Nanocomposites
(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
3.2 Preparation of HAp Nanocrystals 57
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
58 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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.
60 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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
3.4 Applications of HAp/DNA Nanocomplexes as Gene Carriers 61
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
62 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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.
3.4 Applications of HAp/DNA Nanocomplexes as Gene Carriers 63
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].
64 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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.
66 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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.
3.5 Tissue Engineering Applications of HAp Nanocomposites Based on Biodegradable Polymers 67
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
68 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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
3.5 Tissue Engineering Applications of HAp Nanocomposites Based on Biodegradable Polymers 69
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.
70 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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.
3.5 Tissue Engineering Applications of HAp Nanocomposites Based on Biodegradable Polymers 71
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
72 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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
74 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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.
76 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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
78 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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
80 3 Preparation and Applications of Hydroxyapatite Nanocomposites
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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87
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
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.
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.
4.2 Nanocomposites with Zero-Dimensional Nanofillers 89
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
90 4 Synthetic Methods for Nanocomposites Based on Polyester Resins
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
4.2 Nanocomposites with Zero-Dimensional Nanofillers 91
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
92 4 Synthetic Methods for Nanocomposites Based on Polyester Resins
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
4.3 Nanocomposites with One-Dimensional Nanofillers 93
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
94 4 Synthetic Methods for Nanocomposites Based on Polyester Resins
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
4.3 Nanocomposites with One-Dimensional Nanofillers 95
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
96 4 Synthetic Methods for Nanocomposites Based on Polyester Resins
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
4.4 Nanocomposites with Two-Dimensional Nanofillers 97
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
98 4 Synthetic Methods for Nanocomposites Based on Polyester Resins
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
4.4 Nanocomposites with Two-Dimensional Nanofillers 99
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
100 4 Synthetic Methods for Nanocomposites Based on Polyester Resins
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.
4.4 Nanocomposites with Two-Dimensional Nanofillers 101
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.
102 4 Synthetic Methods for Nanocomposites Based on Polyester Resins
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
4.4 Nanocomposites with Two-Dimensional Nanofillers 103
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
104 4 Synthetic Methods for Nanocomposites Based on Polyester Resins
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].
4.4 Nanocomposites with Two-Dimensional Nanofillers 105
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.
106 4 Synthetic Methods for Nanocomposites Based on Polyester Resins
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
4.4 Nanocomposites with Two-Dimensional Nanofillers 107
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
108 4 Synthetic Methods for Nanocomposites Based on Polyester Resins
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.
110 4 Synthetic Methods for Nanocomposites Based on Polyester Resins
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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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
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.
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
5.1 Introduction 117
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
118 5 Synthesis Fabrication and Characterization of Ag/CNT-Polymer Nanocomposites
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
5.3 Results and Discussion 119
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
120 5 Synthesis Fabrication and Characterization of Ag/CNT-Polymer Nanocomposites
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.
5.3 Results and Discussion 121
(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
122 5 Synthesis Fabrication and Characterization of Ag/CNT-Polymer Nanocomposites
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.
5.3 Results and Discussion 123
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
124 5 Synthesis Fabrication and Characterization of Ag/CNT-Polymer Nanocomposites
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.
5.3 Results and Discussion 125
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.
126 5 Synthesis Fabrication and Characterization of Ag/CNT-Polymer Nanocomposites
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
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.
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].
6.2 Structure and Piezoelectric Properties of PVDF 133
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].
134 6 Preparation and Characterization of PVDF-Based Nanocomposites
(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].
6.2 Structure and Piezoelectric Properties of PVDF 135
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
136 6 Preparation and Characterization of PVDF-Based Nanocomposites
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
138 6 Preparation and Characterization of PVDF-Based Nanocomposites
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
140 6 Preparation and Characterization of PVDF-Based Nanocomposites
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
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.
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.
148 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
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).
7.2 Thermal-Assisted In Situ Synthesis: Material Choice and Nanocomposite Characterization 149
+
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),
with permission from Elsevier.
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
150 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
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
7.2 Thermal-Assisted In Situ Synthesis: Material Choice and Nanocomposite Characterization 151
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
152 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
(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
7.2 Thermal-Assisted In Situ Synthesis: Material Choice and Nanocomposite Characterization 153
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
154 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
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
7.3 Fabrication of Nanocomposites and Patterning 155
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
156 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
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 Fabrication of Nanocomposites and Patterning 157
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
158 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
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].
7.3 Fabrication of Nanocomposites and Patterning 159
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
and adapted with permission from Ref. [54].
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.
160 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
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
7.3 Fabrication of Nanocomposites and Patterning 161
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
and adapted with permission from Ref. [55].
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]
162 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
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.
7.3 Fabrication of Nanocomposites and Patterning 163
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
164 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
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.
7.3 Fabrication of Nanocomposites and Patterning 165
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
166 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
(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
and adapted with permission from Ref. [33].
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
7.3 Fabrication of Nanocomposites and Patterning 167
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
168 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
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.
7.3 Fabrication of Nanocomposites and Patterning 169
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
170 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
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
172 7 In Situ Thermal, Photon, and Electron-Beam Synthesis of Polymer Nanocomposites
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
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.
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,
182 8 Synthesis of Polymer Nanocomposites by Water-Assisted Extrusion
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.
8.3 Nanocomposites Preparation 183
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
184 8 Synthesis of Polymer Nanocomposites by Water-Assisted Extrusion
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
8.3 Nanocomposites Preparation 185
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].
186 8 Synthesis of Polymer Nanocomposites by Water-Assisted Extrusion
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
8.3 Nanocomposites Preparation 187
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.
188 8 Synthesis of Polymer Nanocomposites by Water-Assisted Extrusion
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
8.3 Nanocomposites Preparation 189
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
190 8 Synthesis of Polymer Nanocomposites by Water-Assisted Extrusion
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
8.3 Nanocomposites Preparation 191
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.
192 8 Synthesis of Polymer Nanocomposites by Water-Assisted Extrusion
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.
8.3 Nanocomposites Preparation 193
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.
194 8 Synthesis of Polymer Nanocomposites by Water-Assisted Extrusion
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.
8.4 Nanocomposite Properties 195
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
196 8 Synthesis of Polymer Nanocomposites by Water-Assisted Extrusion
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 Nanocomposite Properties 197
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
198 8 Synthesis of Polymer Nanocomposites by Water-Assisted Extrusion
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
200 8 Synthesis of Polymer Nanocomposites by Water-Assisted Extrusion
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.
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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
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.
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 Introduction 213
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.
214 9 In Situ Preparation of Conducting Polymer Nanocomposites
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
9.1 Introduction 215
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.
216 9 In Situ Preparation of Conducting Polymer Nanocomposites
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
9.1 Introduction 217
(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
218 9 In Situ Preparation of Conducting Polymer Nanocomposites
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.
9.2 In Situ Preparation of Conductive Nanocomposites 219
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.
220 9 In Situ Preparation of Conducting Polymer Nanocomposites
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
9.2 In Situ Preparation of Conductive Nanocomposites 221
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
222 9 In Situ Preparation of Conducting Polymer Nanocomposites
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).
9.2 In Situ Preparation of Conductive Nanocomposites 223
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
224 9 In Situ Preparation of Conducting Polymer Nanocomposites
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
9.2 In Situ Preparation of Conductive Nanocomposites 225
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
226 9 In Situ Preparation of Conducting Polymer Nanocomposites
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).
9.2 In Situ Preparation of Conductive Nanocomposites 227
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
228 9 In Situ Preparation of Conducting Polymer Nanocomposites
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].
9.2 In Situ Preparation of Conductive Nanocomposites 229
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.
230 9 In Situ Preparation of Conducting Polymer 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.
9.2 In Situ Preparation of Conductive Nanocomposites 231
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
232 9 In Situ Preparation of Conducting Polymer Nanocomposites
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
234 9 In Situ Preparation of Conducting Polymer Nanocomposites
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
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.
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].
10.4 Characterization Techniques 243
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
244 10 Near IR Spectroscopy for the Characterization of Dispersion in Polymer–Clay Nanocomposites
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
10.4 Characterization Techniques 245
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.
246 10 Near IR Spectroscopy for the Characterization of Dispersion in Polymer–Clay Nanocomposites
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
248 10 Near IR Spectroscopy for the Characterization of Dispersion in Polymer–Clay Nanocomposites
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
10.6 Online and Inline Monitoring of Dispersion 249
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.
250 10 Near IR Spectroscopy for the Characterization of Dispersion in Polymer–Clay Nanocomposites
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
10.6 Online and Inline Monitoring of Dispersion 251
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
252 10 Near IR Spectroscopy for the Characterization of Dispersion in Polymer–Clay Nanocomposites
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
10.6 Online and Inline Monitoring of Dispersion 253
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
254 10 Near IR Spectroscopy for the Characterization of Dispersion in Polymer–Clay Nanocomposites
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
10.6 Online and Inline Monitoring of Dispersion 255
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)
256 10 Near IR Spectroscopy for the Characterization of Dispersion in Polymer–Clay Nanocomposites
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
10.6 Online and Inline Monitoring of Dispersion 257
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
258 10 Near IR Spectroscopy for the Characterization of Dispersion in Polymer–Clay Nanocomposites
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
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.
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
270 11 Synthesis of Polymer Nanocomposites in Supercritical CO2
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.
272 11 Synthesis of Polymer Nanocomposites in Supercritical CO2
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
274 11 Synthesis of Polymer Nanocomposites in Supercritical CO2
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
11.4 Preparation of Polymer/Inorganic Filler Nanocomposites in Supercritical CO2 275
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.
276 11 Synthesis of Polymer Nanocomposites in Supercritical CO2
(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
11.4 Preparation of Polymer/Inorganic Filler Nanocomposites in Supercritical CO2 277
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.
278 11 Synthesis of Polymer Nanocomposites in Supercritical CO2
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
11.4 Preparation of Polymer/Inorganic Filler Nanocomposites in Supercritical CO2 279
(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
280 11 Synthesis of Polymer Nanocomposites in Supercritical CO2
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
11.4 Preparation of Polymer/Inorganic Filler Nanocomposites in Supercritical CO2 281
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
282 11 Synthesis of Polymer Nanocomposites in Supercritical CO2
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
11.4 Preparation of Polymer/Inorganic Filler Nanocomposites in Supercritical CO2 283
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
284 11 Synthesis of Polymer Nanocomposites in Supercritical CO2
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
11.4 Preparation of Polymer/Inorganic Filler Nanocomposites in Supercritical CO2 285
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.
286 11 Synthesis of Polymer Nanocomposites in Supercritical CO2
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
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.
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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