Studying the Effects of Polymer Crystallization and …...TPV-based CPCs were further studied for electromagnetic interference (EMI) shielding applications, which require highly conductive

Studying the Effects of Polymer Crystallization and Blending on Conductive Network Formation in CNT-Based

Conductive Polymer Composites

by

Yasamin Kazemi

A thesis submitted in conformity with the requirements for the degree of Doctor of Philosophy

Department of Mechanical and Industrial Engineering University of Toronto

© Copyright by Yasamin Kazemi 2018

ii

Studying the Effects of Polymer Crystallization and Blending on

Conductive Network Formation in CNT-Based Conductive

Polymer Composites

Yasamin Kazemi

Doctor of Philosophy

Mechanical and Industrial Engineering

University of Toronto

2018

Abstract

The main objective of this thesis was to reduce the percolation threshold of conductive polymer

composites (CPCs) through selective localization of carbon nanotubes (CNTs), taking advantage

of (i) polymer-filler interactions and (ii) polymer blending.

The first part of this study reports on the effects of polymer-filler interactions, particularly the

effects of the conductive fillers on the matrix crystallization behavior, on the evolution of the

conductive networks. Experimental results indicated that changes in the crystallization behavior

and crystal nucleation of a PP matrix in the presence of carbon nanotubes (CNTs) had substantial

effects on the formation or destruction of the conductive networks. Based on our experimental

results, promotion of heterogeneous crystal nucleation ability of CNTs in CPCs with

semicrystalline matrices resulted in wrapping of the conductive fillers with insulating layers of

polymer crystals. Such effects resulted in disruption of conductive network formation in CPCs

even at high filler contents. Interestingly, current-voltage characterizations showed that in such

CPCs, even at a high CNT content of 10 wt.%, electron tunneling was the dominant electron

conduction mechanism, signifying a lack of direct contact between the adjacent CNTs.

iii

As the second phase of this study, a novel technique was developed to use polymer blending for

controlling the conductive network of flexible CPCs with high mechanical properties. We

produced thermoplastic vulcanizate (TPV)- based CPCs, with segregated structure, via the

incorporation of fine pre-vulcanized rubber (PVR) particles into maleic anhydride grafted

polyethylene (MA-g-PE)/CNT composites. Inclusion of the PVR particles resulted in selective

localization of CNTs in the MA-g-PE matrix which significantly reduced the percolation

threshold. More importantly, our experimental results confirmed that strong chemical bonds

formed between the PVR particles and the MA-g-PE matrix which led to excellent interfacial

adhesion between the two phases. This engineered structure increased the CNT-based

nanocomposites’ stretchability by over 200%, while their percolation threshold was decreased by

50%. The cyclic electromechanical properties were also improved, suggesting the

nanocomposites’ great potential for flexible and stretchable electromechanical applications. Such

TPV-based CPCs were further studied for electromagnetic interference (EMI) shielding

applications, which require highly conductive polymer composites with flexible and stretchable

structures.

iv

Acknowledgments

I would like to express my sincerest gratitude to Professor Chul Park and Professor Tobin

Filleter who made this work possible by their continuous support and guidance. I am honored to

have had the opportunity of being supervised by them and wish to thank them warmly.

I would also like to thank the committee members including Professor Hani Naguib and

Professor Craig Steeves for their valuable feedback and constructive suggestions which

contributed greatly to the improvement of this thesis. This thesis has also benefitted greatly,

especially in the early stages of my Ph.D. studies, from extended discussions with Professor

Amir Ameli.

Furthermore, working at the microcellular plastics manufacturing laboratory (MPML) and

nanomechanics and materials lab (NanoM2) was a blessing which provided a great opportunity

for my personal and professional growth. I would like to thank my colleagues and friends at

MPML and NanoM2 for their supports and all the great memories. My thanks also go to the

Natural Sciences and Engineering Research Council of Canada (NSERC) for their financial

support.

Special thanks are due to my family who were there through the ups and downs. I would like to

sincerely thank my parents for their endless support and I hope that I have made them proud.

Finally, I would like to thank my husband and best colleague, Adel, for his constant

encouragement and unconditional support.

v

Contribution of Co-Authors

I am the principal author of all the manuscripts presented in this thesis. I have generated the core

ideas and performed the majority of the experimental works, data collection, analysis, and

manuscript preparation. I would like to acknowledge the contribution of my co-authors for their

valuable helps in preparing the manuscripts presented here. The following outlines the

contributions of each co-author:

[1] Kazemi Y., Ramezani Kakroodi A., Wang S., Ameli A., Filleter T., Pötschke P., and Park

C.B., “Conductive Network Formation and Destruction in Polypropylene/Carbon Nanotube

Composites via Crystal Control Using Supercritical Carbon Dioxide”, Polymer, 129 (2017) 179-

188.

Professor Pötschke was involved in masterbatch preparation and Dr. Kakroodi helped with data

analysis and experimental design and Sai Wang helped with sample preparation and isothermal

annealing experiments. Discussions regarding the research concepts and presentation of data

were continuously conducted with all contributing authors. Professor Park, Professor Filleter,

and Professor Ameli were involved in manuscript preparation.

[2] Kazemi Y., Ramezani Kakroodi A., Mark L.H., Filleter T., and Park C.B., “Effects of

Filler/Polymer Interactions in Controlling the Conductive Network Formation of Polyamide

6/Multi-Walled Carbon Nanotubes Composites”, Submitted to Composites Science and

Technology.

Lun Howe Mark was involved in sample preparation and characterizations of the composites and

Dr Kakroodi helped with TEM analysis and electrical conductivity characterizations. Professor

Park, Professor Filleter, and Professor Ameli were involved in manuscript preparation.

Discussions regarding the research concepts and presentation of data were continuously

conducted with all contributing authors.

[3] Kazemi Y., Ramezani Kakroodi A., Ameli A., Filleter T., and Park C.B., “Highly Stretchable

Conductive Thermoplastic Vulcanizate/Carbon Nanotube Nanocomposites with Segregated

Structure, Low Percolation Threshold and Improved Cyclic Electromechanical Performance”,

Journal of Materials Chemistry C, 6 (2018), 350-359.

vi

Dr Kakroodi helped with sample preparation and electromechanical characterizations. Professor

Park, Professor Filleter, and Professor Ameli were involved in manuscript preparation.

Discussions regarding the research concepts and presentation of data were continuously

conducted with all contributing authors.

[4] Kazemi Y., Ramezani Kakroodi A., Ameli A., Filleter T., and Park C.B., “Facile Production

of Flexible and Stretchable Conductive Polymer Composites for EMI Shielding Applications”, In

preparation.

Dr Kakroodi helped with sample preparation, data analysis and idea generations. Professor Park,

Professor Filleter, and Professor Ameli were involved in manuscript preparation. Discussions

regarding the research concepts and presentation of data were continuously conducted with all

contributing authors.

vii

Table of Contents

Acknowledgments ..................................................................................................................... iv

Contribution of Co-Authors ........................................................................................................ v

Table of Contents ..................................................................................................................... vii

List of Tables ........................................................................................................................... xii

List of Figures ......................................................................................................................... xiii

Chapter 1: Introduction ............................................................................................................... 1

1.1 Preamble ......................................................................................................................... 1

1.2 Problem statement and motivation .................................................................................. 2

1.3 Objective and scope of work ........................................................................................... 3

1.4 Organization of the thesis................................................................................................ 4

1.5 References ...................................................................................................................... 6

Chapter 2: Background and Literature review ............................................................................. 7

2.1 Conductive polymer composites ..................................................................................... 7

2.2 Theoretical research on electrical percolation threshold in CPCs ..................................... 8

2.3 Effects of conductive fillers on electrical properties of CPCs .......................................... 9

2.3.1 Effects of intrinsic electrical conductivity of conductive fillers ............................ 9

2.3.2 Effect of the aspect ratio of the conductive filler .................................................11

2.3.3 Carbon nanotubes (CNTs) ..................................................................................11

2.4 Effects of the polymer matrix on electrical properties of CPCs .......................................13

2.5 Effects of processing method and processing conditions on electrical properties of

CPCs .............................................................................................................................15

2.6 Electrical conduction in CPCs ........................................................................................15

2.6.1 Mechanisms of inter-particle electron conduction in CPCs .................................16

2.7 Electromagnetic interference (EMI) shielding ................................................................18

2.7.1 EMI shielding mechanisms in CPCs ...................................................................19

viii

2.8 Dielectric Properties of CPCs.........................................................................................22

2.8.1 Mechanisms of dielectric permittivity .................................................................22

2.8.2 Dielectric properties of CPCs .............................................................................23

2.9 Morphological control of conductive network in CPCs ..................................................25

2.9.1 Technical concerns associated with filler contents in CPCs ................................25

2.9.2 Dispersion and distribution of conductive fillers in CPCs ...................................25

2.9.3 Morphological control through polymer blending ...............................................26

2.9.4 Morphological control through the use of hybrid fillers ......................................28

2.9.5 Morphological control through foaming .............................................................30

2.10 References ....................................................................................................................32

Chapter 3: Conductive Network Formation and Destruction in Polypropylene/Carbon

Nanotube Composites via Crystal Control Using Supercritical Dioxide ................................39

3.1 Introduction ...................................................................................................................39

3.2 Experimental .................................................................................................................41

3.2.1 Materials and sample preparation .......................................................................41

3.2.2 Isothermal annealing under atmospheric and high pressures ...............................43

3.2.3 Differential scanning calorimetry (DSC) ............................................................43

3.2.4 Polarized optical microscopy (POM) ..................................................................44

3.2.5 Wide angle X-ray scattering analysis ..................................................................44

3.2.6 Electrical conductivity measurements .................................................................44

3.3 Results and discussion ...................................................................................................45

3.3.1 Thermal behavior analysis of the PP/CNT composites ........................................45

3.3.2 Analysis of the crystal types and structures of the PP/CNT composites...............50

3.3.3 Effects of crystallinity on the electrical conductivity of the PP/CNT

composites .........................................................................................................53

3.3.4 Effect of crystallinity on the dielectric properties ................................................56

3.4 Conclusion .....................................................................................................................57

ix

3.5 References .....................................................................................................................59

Chapter 4: Effect of Polymer-Filler Interactions on Controlling the Conductive Network

Formation in Polyamide 6/Multi-Walled Carbon Nanotube Carbon Nanotube Composites ...66

4.1 Introduction ...................................................................................................................66

4.2 Experimental .................................................................................................................68

4.2.1 Materials and sample preparation .......................................................................68

4.2.2 Morphological analysis ......................................................................................68


4.2.4 Wide angle X-ray scattering analysis (XRD) ......................................................69


4.2.6 Rheological characterization...............................................................................69


4.3.1 Dispersion and distribution of MWCNTs ...........................................................70

4.3.2 Crystallization characteristics of the PA6/MWCNT composites .........................71

4.3.3 Percolated nanotube network formation in PA6/MWCNT composites ................77

4.3.4 Electron-conduction mechanisms in PA6/MWCNT composites .........................80

4.4 Conclusions ...................................................................................................................83

4.5 References .....................................................................................................................84

Chapter 5: Highly Stretchable Conductive Thermoplastic Vulcanizate/Carbon Nanotube

Nanocomposites with Segregated Structure, Low Percolation Threshold and Improved

Cyclic Electromechanical Performance .................................................................................91

5.1 Introduction ...................................................................................................................91

5.2 Materials and methods ...................................................................................................93

5.2.1 Materials ............................................................................................................93

5.2.2 Sample preparation .............................................................................................94

5.2.3 Morphological characterizations .........................................................................94

5.2.4 Fourier transform infrared (FTIR) spectroscopy .................................................94

x


5.2.6 Tensile tests .......................................................................................................95


5.2.8 Dynamic mechanical analysis (DMA) ................................................................95


5.3.1 Morphological analysis of the TPV-based composites ........................................96

5.3.2 Effects of PVR on the electrical conductivity of the TPV-based composites .......99

5.3.3 Monotonic deformation of the TPV-based composites ...................................... 100

5.3.4 Thermal and thermo-mechanical analysis of the TPV-based composites ........... 102

5.3.5 Cyclic deformation of the TPV-based composites ............................................ 104

5.3.6 Electrical conductivity under monotonic and cyclic deformations ..................... 107

5.4 Conclusions ................................................................................................................. 109

5.5 References ................................................................................................................... 110

Chapter 6: Facile Production of Flexible and Stretchable Conductive Polymer Composites for

EMI Shielding Applications ................................................................................................ 116

6.1 Introduction ................................................................................................................. 116

6.2 Materials and methods ................................................................................................. 120

6.2.1 Materials .......................................................................................................... 120

6.2.2 Sample preparation ........................................................................................... 120

6.2.3 Characterizations .............................................................................................. 121

6.3 Results and discussions ................................................................................................ 121

6.3.1 Morphological analysis .................................................................................... 121

6.3.2 Mechanical characterizations ............................................................................ 122

6.3.3 Electrical conductivity measurement ................................................................ 124

6.3.4 EMI Shielding characterizations ....................................................................... 126

6.4 Conclusions ................................................................................................................. 127

xi

6.5 References ................................................................................................................... 128

Chapter 7: Conclusions and Future Works ............................................................................... 132

7.1 Summary and conclusions............................................................................................ 132

7.2 Major Contributions ..................................................................................................... 135

7.3 Recommendations and future work .............................................................................. 136

xii

List of Tables

Table 2-1: Electrical conductivities of conductive fillers [11]–[13] ............................................. 9

Table 2-2: Typical properties of CVD-based MWCNTs [22]–[24]............................................ 13

Table 3-1: The melting temperature (Tm), the heat of fusion (∆Hf), the degree of crystallinity

(Xc), the glass transition temperature (Tg), and the melt crystallization temperature (Tc) for the

neat PP and the PP/CNT composites ......................................................................................... 48

Table 4-1: The crystallization temperatures (Tc1) and (Tc2), the melting temperature (Tm), the

heat of fusion (∆Hf), and the degree of crystallinity (Xc) for the neat PA6 and the PA6/MWCNT

composites. .............................................................................................................................. 73

Table 6-1: Summary of EMI SE of CPCs measured in the frequency range of 8.0–12.0 GHz (X

band). ..................................................................................................................................... 117

xiii

List of Figures

Figure 1-1: Typical electrical percolation curve of CPCs along with applications based on

different range of electrical conductivities. ................................................................................. 2

Figure 2-1: Schematic percolation curve for CPCs and schematics of the state of the electrically

conductive network formation with respect to filler content. ....................................................... 8

Figure 2-2: Schematic of different EMI shielding mechanisms with respect to shield material

(conductive barrier). ................................................................................................................. 21

Figure 2-3: Schematic demonstration of the effect of morphological control techniques in

achieving the low percolation threshold in CPCs. ..................................................................... 26

Figure 2-4: Schematic presentation of selective localization of conductive fillers in polymer

blends, according to the Young’s equation. ............................................................................... 27

Figure 2-5: Schematic presentation of the effects of foaming on conductive network formation in

foam injection molding process [64]. ........................................................................................ 31

Figure 3-1: (a) TEM and (b) optical micrograph of PP/2.5 wt.% CNT composites. ................... 43

Figure 3-2: Cooling DSC thermograms of neat PP and PP/CNT composites of varying wt.%. .. 47

Figure 3-3: Polarized optical micrographs of (a) PP and (b) PP/1.0 wt.% CNT composites. ...... 47

Figure 3-4: Heating DSC thermograms of neat PP and the PP/CNT composites of varying wt.%.

................................................................................................................................................. 48

Figure 3-5: Non-isothermal DSC thermograms of the melting of (a) neat PP, (b) PP/1.0 wt.%

CNT, and (c) PP/2.5 wt.% CNT composites after isothermal annealing under atmospheric

pressure. ................................................................................................................................... 49

Figure 3-6: Non-isothermal DSC thermograms of the melting of the (a) PP, (b) PP/1.0 wt.%

CNT, and (c) PP/2.5 wt.% CNT composites after isothermal annealing under 31 MPa scCO2

condition. ................................................................................................................................. 50

xiv

Figure 3-7: WAXS patterns of (a) neat PP, (b) PP/1.0 wt.% CNT, and (c) PP/2.5 wt.% CNT

composites before and after isothermal annealing under 31 MPa scCO2. .................................. 53

Figure 3-8: Electrical conductivity of neat PP and the PP/CNT composites before and after

isothermal annealing at different temperatures, under (a) atmospheric pressure and (b) 31 MPa

scCO2. ...................................................................................................................................... 57

Figure 3-9: Broadband dielectric (a) permittivity and (b) loss of PP/1.0 wt.% CNT composite

before and after isothermal annealing at 135°C and 150°C under 31 MPa scCO2. ..................... 59

Figure 4-1: (a) TEM and (b) optical micrograph of PA6/2 wt.% MWCNT composite. .............. 71

Figure 4-2: Melting thermograms of the neat PA6 and PA6/MWCNT composites (heating rate:

10 ºC/min). ............................................................................................................................... 72

Figure 4-3: Crystallization thermograms for the neat PA6 and PA6/MWCNT composites

(cooling rate: 10 ºC/min). ......................................................................................................... 73

Figure 4-4: XRD patterns for the neat PA6 and for the PA6/MWCNT composites indicated on

the plot. .................................................................................................................................... 77

Figure 4-5: TEM micrographs of the neat PA6 (a) and PA6/0.1 wt.% MWCNT composite, where

zone A and zone B represent the lamellar growth in homogeneously nucleated and

heterogeneously nucleated crystals, respectively. ...................................................................... 79

Figure 4-6: Rheological and electrical percolation curves of PA6/MWCNT composites. The

(G′/G′′) ratio was measured at 240°C and frequency of 0.1 rad s-1. Electrical conductivity

measurements were performed at room temperature. ................................................................ 79

Figure 4-7: Log-log plot of conductivity versus (φ-φc) to determine the percolation model

parameters for PA6/MWCNT composites. R2 is the correlation coefficient of determination for

the linear regression analysis of Equation 2 in the double-logarithmic scale. ............................. 80

Figure 4-8: Current-voltage characteristics of PA6/MWCNT composites with different MWCNT

content. Inset: Current-voltage characteristics of neat PA6 and PA6/2 wt.% MWCNT composites

at smaller scale. ........................................................................................................................ 82

xv

Figure 5-1: SEM micrographs of the PVR particles (a) and MA-g-PE/PVR/CNT: 49.5/50/0.5

wt.% composite (b) and (c). TEM micrographs of the MA-g-PE/PVR/CNT: 49.5/50/0.5 wt.%

composite (MA-g-PE/CNT: 99/1) depicting a high level of dispersion of the CNTs within the

MA-g-PE matrix (d) and strong interface between a PVR particle and the MA-g-PE matrix

containing CNTs (e). ................................................................................................................ 98

Figure 5-2: FTIR spectra of the neat MA-g-PE and MA-g-PE/PVR: 50/50 wt.%. ..................... 98

Figure 5-3: Electrical conductivities of the MA-g-PE/CNT and the MA-g-PE/PVR/CNT

composites containing 50 wt.% PVR. ....................................................................................... 99

Figure 5-4: Tensile stress-strain curves of the neat MA-g-PE and MA-g-PE/CNT: 94/6 wt.%

composite. .............................................................................................................................. 101

Figure 5-5: Tensile stress-strain curves of MA-g-PE/PVR/CNT composites with 50 wt.% of PVR

and 50 wt.% of MA-g-PE/CNT composites containing 0 wt.% (a), 1.5 wt.% (b), 3 wt.% (c), and

6 wt.% (d) CNT. ..................................................................................................................... 102

Figure 5-6: Non-isothermal DSC thermograms as well as the MA-g-PE phase’s crystallinity in

(a) neat MA-g-PE, (b) MA-g-PE/CNT: 94/6 wt.%, (c) MA-g-PE/PVR: 50/50 wt.%, and (d) MA-

g-PE/PVR/CNT: 47/50/3 wt.%. .............................................................................................. 103

Figure 5-7: Loss tangent-temperature curves of the neat MA-g-PE, MA-g-PE/PVR: 50/50 wt.%,

and MA-g-PE/PVR/CNT: 47/50/3 wt.% obtained from DMA. ............................................... 104

Figure 5-8: Tensile stress-strain curves of the neat MA-g-PE (a), MA-g-PE/CNT: 94/6 wt.% (b),

MA-g-PE/PVR: 50/50 wt.% (c), and (d) during cycles 1, 100, and 1000 of their cyclic tensile

tests (maximum strain: 10%). Figure (e) shows the changes in the maximum tensile stress σ10%

(i.e. stress at strain=10%) of the samples during the cyclic tensile tests. The inserts depict the

morphology of the composites that were tested under the cyclic deformations. ....................... 106

Figure 5-9: Dissipated energy for the neat MA-g-PE, MA-g-PE/CNT: 94/6 wt.%, MA-g-

PE/PVR: 50/50 wt.%, and MA-g-PE/PVR/CNT: 47/50/3 wt.% during cycles 1, 100, and 1,000 of

cyclic tensile tests (maximum strain: 10%). ............................................................................ 107

xvi

Figure 5-10: Normalized electrical resistance (R/R0) of MA-g-PE/CNT: 94/6 wt.% and MA-g-

PE/PVR/CNT: 47/50/3 wt.% during monotonic tension at a rate of 50 mm/min. ..................... 108

Figure 5-11: Normalized electrical resistance (R/R0) of MA-g-PE/CNT: 94/6 wt.% (a) and MA-

g-PE/PVR/CNT: 47/50/3 wt.% (b) during cyclic tension at a rate of 50 mm/min and maximum

strain of 10%. ......................................................................................................................... 108

Figure 6-1: SEM micrographs of the PVR particles (a) along with TEM (b) and SEM (c)

micrographs of the MA-g-PE/PVR/CNT composite. .............................................................. 122

Figure 6-2: Tensile stress-strain curves of the neat MA-g-PE, MA-g-PE/CNT: 85/15 wt.%, MA-

g-PE/PVR: 70/30 wt.% and MA-g-PE/PVR/CNT: 25.5/70/4.5 wt.% composites. ................... 123


composites containing 70 wt.% PVR. ..................................................................................... 124

Figure 6-4: Electrical conductivities of the MA-g-PE/PVR composites as a function of PVR

content. .................................................................................................................................. 125

Figure 6-5: EMI SE of the neat MA-g-PE (a), MA-g-PE/CNT: 95.5/4.5 wt.% (b), MA-g-

PE/PVR: 70/30 wt.% (c) and MA-g-PE/PVR/CNT: 25.5/70/4.5 wt.% (d) composites............. 127

1

Chapter 1 Introduction

1.1 Preamble

Conductive polymer composites (CPCs) are an important class of advanced functional materials.

Compared to their conventional metallic and ceramic counterparts, CPCs possess a stimulating

combination of properties including light weight, ease of processing, high specific toughness and

ductility, tunable electrical conductivity, and environmental-sensitive resistivity [1]. Tunability

of electrical conductivity of CPCs allows them to be used in a broad range of applications from

electrically insulating to semi-conductive and conductive materials [2]. Yet, CPCs are a

relatively new class of materials and further research is required to develop an understanding of

the mechanisms behind their behaviors, such as the interactions between conductive fillers and

polymer matrices.

Polymers are typically considered as electrically insulating materials with electrical

conductivities in the range of 10-13-10-15 S/cm [3]. Consequently, the electrical conductivity of

CPCs relies on the formation of continuous networks of conductive fillers within the insulating

polymer matrices. The formation of such networks results in a sudden increase in electrical

conductivity, signifying an insulator-to-conductor transition in CPCs. This phenomenon is often

called the electrical percolation. Further increase in the filler content results in the formation of

additional conductive pathways, allowing the electrical conductivity to increase gradually until a

saturation plateau is reached [4]. Figure 1-1 shows a schematic of a typical percolation curve for

CPCs.

The ability to control the morphology of conductive networks in CPCs, provides a wide range of

conductivity and functionality which are required for different applications. Energy storage (e.g.,

capacitor and super-capacitors), energy conversion (e.g., bipolar plates of fuel cells),

electromagnetic interference shielding, and electrostatic discharge (ESD) protection are major

applications proposed for CPCs [2]. Figure 1-1 also shows the range of electrical conductivity

required for these applications.

2

Figure 1-1: Typical electrical percolation curve of CPCs along with applications based on

different range of electrical conductivities.

1.2 Problem statement and motivation

The mechanical and electrical properties of CPCs strongly depend on the level of dispersion and

distribution of conductive fillers within the polymeric matrix. Perfect dispersion and distribution

of conductive fillers, such as carbon nanotubes (CNTs), within polymer matrices is extremely

desirable for achieving high mechanical properties. However, such states of uniform dispersion

and distribution of fillers may not be the optimal conditions for achieving the lowest percolation

threshold [5]. As a result, several methods have been proposed in the literature to control the

network of conductive fillers aimed at maximizing the CPCs’ electrical properties [6].

An effective technique for maximizing the CPCs’ electrical properties is through controlling the

polymer-filler interactions. This technique, which relies on controlling the electrical resistance at

filler-filler contact points in the network of conductive fillers, can be especially effective in

CNT-filled CPCs [5]. CNTs have many similarities to molecular chains of polymeric matrices in

terms of their end-to-end lengths and diameters. Such similarities in their dimensionality can

result in strong interfacial interactions between CNTs and the molecular chains [7]. This effect

can cause the formation of thin insulating layers of polymer at the contact points of adjacent

CNTs. This insulating layer may disrupt the conductive network formation and substantially

3

deteriorate the electrical conductivity of CPCs [8]. Despite the importance of studying polymer-

CNT interactions at CNT-CNT contact points in CPCs, there exists a lack of knowledge in this

field.

Another technique for maximizing CPCs’ electrical properties is the selective localization of

conductive fillers within a certain phase, or at the interphase, in polymeric blends. This technique

relies on the localization of fillers in the most thermodynamically favorable phase and

developing a segregated phase morphology [9]. However, this technique often suffers from

significant drawbacks such as (i) restrictions in the selection of blend components and (ii)

significant reduction in mechanical properties of CPCs. As a result, research is necessary on the

production of CPCs using the concept of volume exclusion while maintaining high mechanical

properties.

1.3 Objective and scope of work

The main objective of this thesis was to improve the electrical properties of CPCs using the

concept of selective localization of conductive fillers via:

(I) Controlling polymer-filler interactions:

The first objective of this work was studying the influences of polymer-filler interactions with a

focus on the effects of crystallization behavior of the polymeric matrices on conductive network

formation and destruction. Matrix crystallization behavior is especially important in the case of

CNT-filled CPCs. CNTs are known to act as heterogeneous crystal nucleation sites for many

semicrystalline polymers. This behavior can result in substantial increase in contact resistance

between CNTs. Further, promoting homogeneous crystallization of the matrix can improve the

CPCs’ electrical properties via excluding the CNTs from the crystalline part of the matrix. Thus,

our efforts were focused on developing a deeper understanding of the effects of matrix

crystallization behaviors in CNT-filled CPCs, as well as transforming the crystallization behavior

from heterogeneous to homogeneous mechanism aimed at improving the electrical properties.

(II) Polymer blending:

The production of stretchable, elastic, and conductive materials, using practical and cost-

effective techniques has attracted a great deal of attention due to their enormous potential in

4

various applications such as stretchable electronics, strain gauges, and implantable devices. The

crucial criteria for such applications are the ability of the materials to retain much of their

electrical and mechanical properties after many deformation cycles. However, the traditional

stretchable and elastic CPCs, produced through the direct addition of conductive fillers into

elastomeric matrices, cannot satisfy such requirements. Thus, as the second objective, this thesis

was focused on the development of a novel technique to produce CPCs with segregated

structures, using the concept of polymer blending, while maintaining high mechanical properties

including stretchability and mechanical durability.

1.4 Organization of the thesis

Chapter 2 presents a literature review and relevant background on theoretical studies and

fundamental concepts on the development of CPCs. This chapter encompasses details such as

previous theoretical studies on electrical percolation threshold, the mechanisms of electron

conduction and effects of different processing conditions on electrical properties of CPCs. It also

discusses the previous studies on morphological control of conductive fillers in CPCs.

Chapter 3 presents studies focused on understanding the effects of different crystallization

mechanisms on the electrical conductivity of polypropylene (PP)/CNT composites. In this

chapter, we introduce the concept of crystal control via isothermal annealing under supercritical

carbon dioxide (scCO2) conditions for manipulating the conductive network in CPCs.

In chapter 4, we further investigate the effects of polymer-filler interactions, in particular

transcrystallization, on controlling the conductive network in polyamide 6 (PA6)/multi-walled

carbon nanotube (MWCNT) composites. The findings of this study help better understand the

deteriorating effects of heterogenous crystal nucleation, on the surfaces of CNTs, on (i) the

conductive network formation and (ii) the electrical conduction mechanism in CPCs with

semicrystalline matrices.

Chapter 5 presents a novel technique for the cost-effective production of stretchable, elastic, and

electrically conductive thermoplastic vulcanizates (TPVs), containing large quantities of fine

pre-vulcanized rubber (PVR) particles, having segregated phase morphology. The TPV

nanocomposites produced in this work, showed very low electrical percolation thresholds

combined with superior stretchability, elasticity, and mechanical durability.

5

Chapter 6 further reports on the electrical properties of highly filled TPV composites with

segregated structures. The effect of inclusion of pre-vulcanized rubber particles, from post-

consumer tire rubber, were studied on electrical conductivity and EMI shielding characteristics

of TPV-based CPCs.

Chapter 7 provides a summary, key contributions and conclusions of this research, along with a

list of recommendations for future research works on this topic.

6

1.5 References

[1] E. T. Thostenson, Z. Ren, and T.-W. Chou, “Advances in the science and technology of

carbon nanotubes and their composites: a review,” Compos. Sci. Technol., vol. 61, no. 13,

pp. 1899–1912, 2001.

[2] C. Kingston et al., “Release characteristics of selected carbon nanotube polymer

composites,” Carbon, vol. 68, pp. 33–57, 2014.

[3] W. Zhang, A. A. Dehghani-Sanij, and R. S. Blackburn, “Carbon based conductive

polymer composites,” J. Mater. Sci., vol. 42, no. 10, pp. 3408–3418, 2007.

[4] I. Balberg and S. Bozowski, “Percolation in a composite of random stick-like conducting

particles,” Solid State Commun., vol. 44, no. 4, pp. 551–554, 1982.

[5] M. H. Al-Saleh and U. Sundararaj, “A review of vapor grown carbon nanofiber/polymer

conductive composites,” Carbon, vol. 47, no. 1, pp. 2–22, 2009.

[6] H. Deng, L. Lin, M. Ji, S. Zhang, M. Yang, and Q. Fu, “Progress on the morphological

control of conductive network in conductive polymer composites and the use as

electroactive multifunctional materials,” Prog. Polym. Sci., vol. 39, no. 4, pp. 627–655,

2014.

[7] B. P. Grady, “Effects of carbon nanotubes on polymer physics,” J. Polym. Sci. Part B

Polym. Phys., vol. 50, no. 9, pp. 591–623, 2012.

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electrical conductivity of carbon nanotube-based composites,” Appl. Phys. Lett., vol. 91,

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8, pp. 1526–1563, 2014.

7

Chapter 2 Background and Literature review

2.1 Conductive polymer composites

There are two main methods by which polymer-based materials may conduct electricity: (i) in

intrinsically conductive polymers (ICPs), electrical conductivity is derived from the electronic

structure of the polymers, and (ii) in conductive polymer composites (CPCs), the inclusion of a

sufficient content of conductive fillers into insulating polymers provides high electrical

conductivity. In comparison with the first category, CPCs are more economically viable, easier

to process and generally provide better mechanical properties [1].

Polymer matrices in CPCs are typically electrically insulating with electrical conductivities in the

range of 10-13-10-15 S/cm [2]. Consequently, the electrical conductivity of CPCs relies on the

formation of continuous networks of conductive fillers within the insulating polymer matrices.

The formation of such networks results in a sudden jump in the CPCs’ electrical conductivity,

signifying the insulator-to-conductor transition (Figure 2-1). This phenomenon is often called

electrical percolation (Pc) and the concentration at which the insulator-to-conductive transition

occurs is referred to as electrical percolation threshold (φc) [3].

Further increase in the filler content above the Pc results in the formation of additional

conductive pathways, allowing the electrical conductivity to increase gradually until a saturation

plateau is reached. The percolation threshold and electrical conductivity of CPCs depends on

many parameters including the filler’s electrical conductivity, aspect ratio, polymer/filler

interactions and also the level of dispersion, distribution, and alignment of the conductive fillers

within the polymeric matrix [2].

8

Figure 2-1: Schematic percolation curve for CPCs and schematics of the state of the

electrically conductive network formation with respect to filler content.

2.2 Theoretical research on electrical percolation threshold in CPCs

The electrical conductivity of CPCs can be explained by a power law model proposed by the

classical percolation theory:

𝜎 = 𝜎0 (𝜑 − 𝜑𝑐)𝑡 (1)

where σ and σ0 are the measured conductivity and scaling factor, respectively. φ is the filler

volume content, above the percolation threshold, φc is the electrical percolation threshold, and t

is the critical exponent which reflects the dimensionality of the system. In this model, t ≈ 2 and t

≈ 1.3 corespond to the formation of three-dimensional (3D) and two-dimensional (2D)

conductive networks, respectively [4]. Balberg and Bozowski reported that in carbon

fiber/polyvinylchloride (PVC) composites, the t values were in the range of 1.5 to 2.1 [5].

Similar results were reported by Stinchcombe [6] and Abeles et al. [7] which were in good

agreement with the theoretical prediction of t = 2. In an interesting study, Gabbels et al. reported

that in a co-continous polystyrene (PS)/polyethylene (PE)/carbon black (CB) system, t value was

equal to 2 when CBs were selectively localized in the PE phase, and equal to 1.3 in the case of

selectve localization of CB particles at the interface between the PE and PS phases [2].

9

Nevertheless, several studies in the literature report large variations in their calculated t values

ranging from 1 to 12 [8]. Such discrepancies between the theoretical predictions and the

experimental observations has been ascribed to several parameters, including the dimensionality

of the conductive network and also the tunnelling distance distributions, between conductive

particles, in such systems [9].

2.3 Effects of conductive fillers on electrical properties of CPCs

To meet the increasing demand for advanced functional electronics, several types of conductive

fillers have been proposed for production of CPCs with high electrical conductivities and low

φcs. The most common conductive fillers include metallic particles, carbonaceous fillers (e.g.

CB, graphite, carbon fibers, carbon nanofibers, CNTs, graphite nanosheets and graphene), metal

fibers, metal-coated fibers and ICPs [10].

2.3.1 Effects of intrinsic electrical conductivity of conductive fillers

Table 2-1 summarizes the electrical conductivities of different filler types which are often used

in producing CPCs.

Table 2-1: Electrical conductivities of conductive fillers [11]–[13]

Conductive filler Electrical Conductivity (S/cm)

Silver 6.3 × 105

Copper 5.9 × 105

Aluminum 3.8 × 105

Nickle 1.4 × 105

Steel 6.7 × 104

Stainless steel 1.4 × 104

10

Graphene 105- 106

Single-walled Carbon Nanotubes 102- 106

Multi-walled Carbon Nanotubes 103- 105

Graphite 4000 (in-plane), 3.3 (C-axis)

Carbon fiber 0.6 × 103

Carbon Black 10

Poly (3,4-ethylenedioxythiophene) (PEDOT) 103

Polypyrrole (PPy) 102

Polyaniline (PANI) 10

Li et al. performed a theoretical analysis, based on the Monte Carlo simulations [14], on the

effects of intrinsic electrical conductivity of conductive fillers as well as the contact resistance

between adjacent conductive fillers on the electrical conductivity of CNT-based CPCs. Based on

their results a substantial change in the intrinsic electrical conductivity of CNTs from 106 to 104

S/m did not affect the electrical conductivity of the CPCs significantly. However, changes in the

contact resistance between the conductive fillers resulted in dramatic changes in the electrical

conductivity of CPCs. Consequently, they concluded that contact resistance plays a more

significant role in controlling the electrical properties of CPCs. The most important parameters in

controlling the contact resistance between conductive fillers in CPCs are the thickness and the

type of the insulating layer of polymer matrix at the contact points between the conductive

fillers.

11

2.3.2 Effect of the aspect ratio of the conductive filler

As mentioned earlier, the aspect ratio of the conductive filler is considered as one of the most

important parameters in controlling the electrical conductivity of CPCs. The correlation between

φc and the aspect ratio of conductive fillers (𝑙𝑑⁄ ) can be explained, using the excluded volume

theory. Based on this theory, for a statistical distribution of filler particles with high aspect ratios,

it can be assumed that φc ≈ 1

𝑙𝑑⁄. Consequently, according to the excluded volume theory, φc is

expected to decrease with increase in the filler aspect ratio in CPCs [15]. Deng et al. performed a

comparison between the experimental results and theoretically calculated values of percolation

threshold, based on the excluded volume theory. It is worth noting that in the case of perfect

alignment of conductive fillers, the predicted values for φc were higher than in isotropic systems.

Also, the φc values in these systems were independent of the aspect ratios of conductive fillers

[15]. Based on their study, for composite systems containing randomly distributed and dispersed

CNTs with aspect ratios of 1000, the expected φc value would be approximately 0.1 vol. %.

Higher φc values are typically ascribed to fiber damage or breakage during processing or the

possibility of curvature of the fibers with very high strength and aspect ratios. Lower φc values in

such systems, on the other hand, are usually attributed to kinetic percolation. In kinetic

percolation, conductive fillers can obtain sufficient energy, via diffusion, convection, shearing,

or external fields, to move and promote the conductive filler formation at much lower filler

contents [8].

2.3.3 Carbon nanotubes (CNTs)

Based on the above-mentioned discussions on the effects of the conductive filler’s aspect ratios

on the percolation threshold of CPCs, extensive research has been focused on studying CPCs

containing conductive fillers with very high aspect ratios. Among all conductive fillers, carbon

nanotubes (CNTs) have attracted great attention, due to their superior mechanical and electrical

properties and ease of processing. In comparison with graphene, CNTs are known to have

smaller specific surface area, which makes it easier to uniformly disperse individual CNTs

within the viscus polymer matrix. This can result in lower percolation threshold and higher

electrical conductivity in CNT-based CPCs [2].

12

the performance of graphene/polymer composites is often hindered by the inherent difficulty to

uniformly disperse individual graphene sheets within the viscous polymer matrix

CNTs are cylindrical structures of covalently bonded carbon atoms. They can be considered as

graphitic platelet-like structures rolled into seamless cylinders with diameters in the order of

approximately 1–50 nm. The ends of the cylinders in CNTs are often capped by hemifullerenes.

Depending on the number of stacked graphitic layers, there exist two different types of CNTs: (i)

single-walled carbon nanotubes (SWCNTs) and (ii) multi-walled carbon nanotubes (MWCNTs).

MWCNTs include multiple rolled layers of graphite coaxially arranged around a central hollow

core with van der Waals forces bonding adjacent layers, while SWCNT consists of a single sheet

graphene cylinder [16]. The outer diameter of MWCNTs depends on the number of graphene

layers, and usually ranges from 10–50 nm, about ten times higher than that of SWNTs. The inner

diameter of MWCNTs is typically half the outer diameter. The lengths of single and multi-wall

carbon nanotubes can range from a few micrometers to as high as a few centimeters in some

special cases [17]. Depending on their fabrication technique, the resulting CNTs may differ

greatly in terms of their geometrical characteristics, such as diameter, aspect ratio, entanglement,

as well as crystalline perfection, including crystallinity, crystalline orientation, purity, which can

dramatically affect their intrinsic properties [18].

There are several techniques for the fabrication of CNTs including arc discharge, laser ablation,

chemical vapor deposition (CVD) and catalytic growth processes [19]. In arc discharge and laser

ablation techniques, CNTs are formed from the condensation of hot gaseous carbon atoms

generated from the evaporation of pure carbon. CVD process involves thermal catalytic

decomposition of a gaseous carbon source, typically a mixture of hydrocarbon gas, acetylene,

methane or ethylene and nitrogen, on the surface of a catalyst, where it forms the CNTs.

Compared with arc and laser methods, CVD is more industrially viable and offers more control

over the length and structure of the produced nanotubes. However, CNTs made from this

technique generally have more structural defects and, therefore, their thermal, electrical and

mechanical properties deviate significantly from those expected for pristine nanotubes [20]. In

MWCNTs, it is reported that the overall properties of the filler are determined by the transport

properties of the outer graphene layer, also known as the shell [21]. Table 2-2 summarizes some

of the typical structural, mechanical and functional properties of MWCNTs.

13

Table 2-2: Typical properties of CVD-based MWCNTs [22]–[24]

Property Value

Diameter (nm) 5-50

Aspect ratio 100-10000

Density (g/cm3) ~1.8

Elastic modulus (GPa) 1000

Tensile strength (GPa) 10-60

Electrical resistivity (Ω cm) 2×10-3-1×10-5

Electron mobility (cm2/Vs) 104-105

Thermal conductivity (W/mK) 3000

Coefficient of thermal expansion (K-1) Negligible

Thermal stability (ºC) >600 (in air), 2800 (in vacuum)

Specific surface area (m2/g) 10-20

2.4 Effects of the polymer matrix on electrical properties of CPCs

Although polymeric matrices are typically considered as the electrically insulating component of

CPCs, their interactions with the conductive fillers are known to greatly affect the contact

resistance in the conductive networks, thereby substantially influencing the CPCs’ electrical

properties. The most important parameters in defining the filler-matrix interactions are the

polymer’s surface tension, polarity, viscosity, and crystallinity [24]. Several studies have been

devoted to the investigation of the effects of the characteristics of matrix on filler-matrix

interactions and their CPCs’ percolation thresholds [2], [10], [19], [25]. However, some of these

14

studies have reported seemingly contradicting behaviors which demonstrate the complexity of

the dependence of CPCs’ properties on filler-matrix interactions [2].

Experimental results indicate that for systems with relatively low filler-polymer interfacial

tensions, stronger polymer-filler interactions are observed. Such interactions result in better

dispersion and distribution of the conductive fillers within the insulating polymeric matrices and

decrease the percolation threshold. However, some experimental results indicate that increasing

the filler-polymer interfacial tension promotes kinetic percolation which results in lower

percolation threshold in CPCs [24].

Similarly, polarity of the polymer matrix can affect the percolation threshold of the CPCs by

changing the interactions between the filler and matrix. Miyasaka et al. [26] studied the effects of

the polarity of polymer matrices on the percolation threshold of CPCs containing carbon black.

Based on their experimental results, increasing the polarity of the polymer matrices, resulted in

an increase in the percolation threshold of CPCs. A similar conclusion has been reported in a

different study performed by Foulger [27]. Conversely, Sau et al. reported opposite conclusions

on the effects of matrix polarity on percolation threshold. They investigated the electrical

conductivities of ethylene-propylene diene monomer rubber (EPDM), acrylonitrile butadiene

rubber (NBR), EPDM/NBR blends and carbon black systems. Based on their experimental

results, despite EPDM’s lower polarity compared to NBR, EPDM-based CPCs exhibited a higher

percolation threshold [28].

Experimental results indicate that increasing the viscosity of polymeric matrices generally results

in increase in the percolation threshold of CPCs. The higher the viscosity of the polymeric

matrix, the higher is the shearing force experienced by the conductive fillers during the

processing of the CPCs, which can result in fiber breakage and reduction in the aspect ratio of

conductive fillers [29].

The crystallinity and crystallization behavior of the polymer matrix, in CPCs with

semicrystalline matrices, can also substantially influence their electrical percolation threshold.

Depending on the heterogeneous nucleation ability of the conductive fillers, crystallization of the

matrix phase may increase or decrease the percolation threshold in CPCs. Such effects will be

discussed in more detail in chapters 3 and 4.

15

2.5 Effects of processing method and processing conditions on electrical properties of CPCs

There are three main processing techniques for the fabrication of conductive polymer

composites: (i) Solution casting, where the conductive fillers are dispersed within a solution of

the polymeric matrix in an organic solvent and the CPCs are produced in the form of films after

solvent evaporation. (ii) In-situ polymerization, where the conductive fillers are mixed with the

monomers prior to the polymerization of the polymeric matrix. (iii) melt mixing, in which the

conductive fillers are mixed with the polymer at temperatures above the melting temperature of

the matrix. Overall, melt mixing is considered as the most effective and industrially viable

technique for compounding conductive fillers and polymeric matrices [30].

Processing methods and conditions affect the filler distribution, dispersion, aspect ratio, and

orientation in CPCs. There are two opposing mechanisms through which the processing

parameters can affect the electrical conductivity of CPCs. (i) Severe mixing breaks down the

conductive fillers’ agglomerates and ensures that the conductive fillers are uniformly dispersed

and distributed within the polymeric matrix; (ii) On the other hand, the use of excessively high

shear stresses is not recommended, since it can lead to severe fiber breakage which deteriorates

the electrical conductivity of CPCs and can result in an increased percolation threshold.

Therefore, attention needs to be paid to the selection of the processing conditions in the

fabrication of CPCs [31]. Andrews et al studied the effects of mixing energy on the aspect ratio

and level of dispersion of the conductive fillers in CPCs [31]. We can conclude that the selection

of a less-than-perfect level of dispersion and distribution of the conductive fillers might benefit

the electrical properties of the CPCs, which comes at the expense of compromising their

mechanical properties.

2.6 Electrical conduction in CPCs

Electron conduction in CPCs involves the electron flow in response to a voltage bias along a

network of conductive fillers. In the absence of an electric field, the conduction electrons are

scattered freely due to their thermal energy. By applying an electric field of E, electrons obtain

the energy to move with respect to the applied voltage, and therefore, a net current density (J) is

created in the system, given by:

16

J=Ne e E (2)

where Ne is the number of electrons, e is the electron charge, is the electron

mobility [32]. Consequently, the electrical conductivity (σ) of a material can be calculated as:

σ = 𝐽

𝐸 (3)

where the unit for electrical conductivity, in the SI system, is defined as S/m. The mechanisms of

electron conduction in different materials can be explained using the band theory. Based on this

theory, the available energy state of each electrons is considered as a separate band. For electrons

to be able to travel through materials, they need to gain the energy to move from the valance

band into their conduction bands. Consequently, materials can be classified as conductor,

semiconductor and insulator based on the energy gap between their valence band and conduction

bands.

The energy gap between the valence and conduction bands in conductive materials is negligible

and these two bands mostly overlap with each other. Consequently, conductive materials show

very low electrical resistance with respect to the applied electric field. The band gap in semi-

conductive materials is larger than conductive materials, but small enough to let excited electrons

in the valence band reach the conduction band by thermal energy. In the case of insulating

materials, the band gap between the valance and conduction band is too large. Therefore, very

few electrons can be found in their conduction band, which results in a very low electrical

conductivity.

2.6.1 Mechanisms of inter-particle electron conduction in CPCs

Depending on the type of conductive filler, filler content and filler-polymer interactions, several

electron conduction mechanisms have been proposed to explain the electrical conductivity of

CPCs [33]. These mechanisms can be classified based on the three different regions of the

electrical percolation curve in CPCs:

(i) At filler contents below the percolation threshold, where the conductive fillers are far from

each other and encapsulated by thick insulating layers of polymeric matrix. In this region,

electrical conductivity of the composite is limited to the conductivity of the insulating matrix

which indicates that the conductive fillers, at such low filler contents, have negligible

17

contribution to the electron conduction of the CPCs. In this situation, by applying a voltage

difference, the insulating polymer, between the conductive fillers begins to polarize, by

redistributing its protons and electrons charges, in response to the electric field. However, due to

the very large band gap between the valance and conduction bands in the insulating matrix,

polarized charges cannot be easily dislocated. The higher the content of conductive fillers, the

lower is the gap width and therefore the higher is the chance for electrons to pass their barrier.

Maxwell proposed a model for electron conduction mechanism in CPCs with uniformly

dispersed spherical conductive fillers, at filler contents below the percolation threshold which

was validated up to 10 vol.% of conductive fillers [34].

(ii) At filler contents close to percolation threshold. Increase in the conductive filler content,

results in decreasing the insulating gap between the conductive fillers. In sufficiently narrow

insulating gaps between conductive fillers, in the range of few nanometers, very high field

strength is expected which results in internal field emission which serves as the dominant

mechanism in electron conduction in CPCs in this region. Internal field emission is a general

term for explaining different electron conduction mechanisms including quantum tunneling and

hopping. According to this theory, the intensity of the external electric field at the regions

between the conductive fillers is magnified by the order of M, where M is a factor depending on

the average size of conductive fillers and their inter-particle distance in CPCs. Such an

intensified electric field enables the electrons to jump across a gap or tunnel through energy

barriers between conductive fillers in CPCs [2].

Based on the theoretical studies performed by Connor et al. [35], at a constant temperature, DC

electrical conductivity (σDC) between conductive fillers separated by a thin isolating layer of

width w can be estimated as:

ln σDC ∝ -w (4)

Considering the spatial distribution of the conductive fillers, w ∝ -φ-1/3 and therefore:

ln σDC ∝ - φ-1/3 (5)

Several studies in the literature investigated the validity of this simplified model in predicting

the DC electrical conductivity versus the filler content in CPCs [8].

18

Electron conduction mechanisms in such composites can also be addressed by thermal

fluctuation induced tunneling, where at high temperature, the tunneling current can increase due

to thermal fluctuations, which results in decreasing the potential barrier between the conductive

fillers [36].

(iii) At filler contents beyond the percolation threshold. By further increasing the conductive

filler content above the percolation threshold, the formation of a continuous network of

conductive fillers results in filler-filler contacts in CPCs. Consequently, the electrons which

belong to the conduction bands in conductive fillers can freely pass through the composite,

through a mechanism known as metallic conduction, which usually results in a dramatic increase

in electrical conductivity of CPCs. At this region, the contact resistance between the conductive

fillers is the dominant mechanism in controlling the electrical conductivity of CPCs [2].

2.7 Electromagnetic interference (EMI) shielding

EMI shielding is one of the most important applications of CPCs in the electronics industry.

Electronic devices are known to irradiate electromagnetic signals. EMI shielding is the practice

of shielding the electronic devices against the incoming or outgoing electromagnetic signals,

which can result in degradation of nearby systems or equipment performances. EMI is also

known to cause health hazards including symptoms of languidness, insomnia, nervousness and

headache which occur upon exposure to electromagnetic signals [12].

EMI shielding effectiveness (EMI SE) is the capacity of shield materials to dissipate the

electromagnetic energy. EMI SE is often expressed in decibels (dB) and defined as logarithm of

the ratio of incident electromagnetic field (Pi) to transmitted electric electromagnetic field (Pt):

EMI SE = 10 log10(𝑃𝑖

𝑃𝑡) (6)

For some applications which require high-efficiency shielding materials, specific SE (dB cm2/g),

can be used as the measure for shielding effectiveness. Specific SE is defined as SE divided by

the density and thickness of the shield material.

The most common EMI shielding materials are metals and metallic composites, mainly due to

their superior electrical conductivity. Compared to the conventional metal-based EMI shielding

materials, CPCs, which include conductive and microwave-absorptive nanofillers, benefit from

19

unique advantages, such as low price, light weight, flexibility, corrosion and oxidation

resistance, easier processability, and tunable shielding effectiveness. Low density is one of the

most important technical requirements for practical EMI shielding applications, specifically in

areas of aerospace, automobiles, and household appliances [37].

EMI shielding effectiveness of CPCs depends on many parameters including intrinsic

conductivity, aspect ratio, and content of the conductive fillers, dispersion and distribution of

conductive fillers, range of shielding frequency, as well as structure and thickness of the shield

materials [12]. The high aspect ratio and electrical conductivity of CNTs makes them suitable

candidates for EMI shielding applications without sacrificing beneficial properties of polymeric

matrices [38].

The standard limit for EMI in commercial electronic devices is defined by a parameter called

electromagnetic compatibility (EMC). To control EMI issues based on EMC regulations, an EMI

shielding effectiveness (EMI SE) of at least 20 dB is required for shielding of 99% of incident

radiation [32].

2.7.1 EMI shielding mechanisms in CPCs

For EMI shielding characterization, electromagnetic waves are considered to be plane waves,

consisting of two components, the electric field component and the magnetic field component,

which are perpendicular to each other and transverse to the direction of wave propagation.

Based on Lorentz law [39], the force equal to F is applied to the charged particles in the shield

material with respect to the incident electromagnetic wave as:

�� = 𝑞��+qν × 𝜇0 �� (7)

where q is the charge of conductive particle, ν is the velocity of particles, 𝜇0is the magnetic

permeability of free space, which is equal to 4π×10-7 H·m-1; and, �� and �� are the electric and

magnetic fields, respectively. Consequently, electron conduction can result in attenuation of

electromagnetic energy via: (i) absorbing the electromagnetic energy through interaction with

mobile charge carriers and electric/magnetic dipoles in response to the electromagnetic field, and

(ii) inducing secondary, weaker electromagnetic waves, as consequence of electron

transportation in shield materials [39].

20

Figure 2-2 schematically shows the three main mechanisms, including reflection, absorption and

transition, in which a shield material responds to electromagnetic waves. Reflection is the

fraction of the incident electromagnetic wave which is reflected from the surface of the shield

material. Shielding by reflection is attributed to impedance mismatch for electromagnetic wave

propagating into air and into the shielding material. A high impedance mismatch is mainly due to

the presence of highly conductive fillers at the surface of the shield material [40].

Although reflection is a major mechanism in EMI SE, electromagnetic signals reflected to the

inner volume of the electronic device can cause damage, especially in electronic circuits, because

of spurious interferences between the constitutive electronic components. Consequently,

Thomassin et al. [40] proposed that the relatively high volume of air in open-cell foam CPCs can

be considered as a solution for matching of the wave impedances of the CPCs and the ambient

atmosphere.

Absorption is the fraction of the electromagnetic energy that is dissipated as the incident wave

passes through the shielding material. The longer the network of conductive fillers, the more

electromagnetic energy can be dissipated in CPCs [41]. The magnitude of electromagnetic wave

attenuation is estimated by the factor of 𝑒−𝑧

𝛿, where z is the distance of electromagnetic wave

penetration and δ is the skin depth of shield material (i.e. the thickness of shield material in

which the power of the electromagnetic wave decays to 1

𝑒 of the incident value). The value of δ

can be estimated as:

δ= 1√𝜋𝑓𝜇𝜎⁄ (8)

where f is the wave frequency, and μ and σ are shield’s magnetic permeability and electrical

conductivity, respectively.

Consequently, for a shield material with thickness of t, EMI SE by absorption can be calculated

as:

SEA=20. Log10(𝑒𝑡

𝛿)=8.69𝑡

𝛿=131t√𝑓𝜇𝜎 (9)

21

Consequently, it can be concluded that there is a linear relationship between shielding by

absorption and thickness of shielding material. Such effect is mainly due to the fact that at higher

thickness, there are more mobile charge carriers and/or electric/magnetic dipoles to interact with

electromagnetic waves [32].

Transmission is the fraction of the electromagnetic wave which passes through the shield

material, despite the reflection and absorption mechanisms. As shown in Figure 2-2, as the

transmitted electromagnetic wave exits the shield material, at the interface of the shield and

exterior environment, a fraction of the electromagnetic wave reflects off the interface. These

secondary reflected electromagnetic waves are considered as part of the reflection mechanism. It

is worth noting that in the case of CPCs with multiphase systems and large interfacial areas

between the conductive fillers and insulating polymer matrices, multiple reflections can occur

which act as an additional mechanism in controlling the EMI SE. It is known that although

multiple reflections can have effects on EMI SE, however their effects can be ignored if the

shield’s thickness is larger than the skin depth [32].

Figure 2-2: Schematic of different EMI shielding mechanisms with respect to shield

material (conductive barrier).

22

2.8 Dielectric Properties of CPCs

The dielectric property is the ability of a material to store energy in response to an external

electric field. Dielectric materials, with high permittivity and low energy loss, can be used for

various applications, including energy storage and conversion, hybrid electric vehicles, medical

defibrillation, smart grids and sensors [42].

Dielectric permittivity is a function of frequency of the applied electric field, and can be defined

as a complex value (ε*):

ε*(ω)= εʹ(ω)+ iεʹʹ(ω) (10)

where, ω is the angular frequency of the applied electric field, and εʹ and εʹʹ are real and

imaginary components of the dielectric permittivity, respectively. Real permittivity (εʹ), also

known as dielectric permittivity or dielectric constant, represents the capacity of dielectric

material to store the energy from an external electric field, via charge polarization and generation

of a momentum arising from separation of positive and negative charges. Whereas imaginary

permittivity (εʹʹ), also known as loss factor, represents the dissipation of energy into heat in the

dielectric material. Imaginary permittivity in AC voltage, has two components: (i) ohmic loss

and (ii) polarization loss. (i) Ohmic loss is mainly due to the dissipation of the energy via

collision of mobile charges moving through dielectric materials. Such effect decreases with

increase in frequency, because of reduced available times for free electrons to respond to the

alternating electric field. (ii) Polarization loss refers to the dissipation of energy in dielectric

materials via friction due to the orientation of electric dipoles in each half cycle of an AC field

[43].

The dissipative behavior of dielectric materials can also be characterized by the so-called loss

tangent, or dielectric loss, where:

tan δ(ω) = εʹʹ(ω)/ εʹ(ω) (11)

2.8.1 Mechanisms of dielectric permittivity

Relaxation time is the parameter defining the amount of time which a dipole requires to respond

to an electric field and can be considered as a measure of the mobility of the dipole in a dielectric

material. Therefore, dielectric polarization strongly relies on frequency, since this parameter

23

controls the time available for dipoles to orient with respect to the applied electric field. At

frequencies above their relaxation time, dipoles cannot contribute to the dielectric properties of

the material [44].

At frequencies below the relaxation times of each mechanism, increasing the frequency results in

more frequent occurrence of the dipole orientations, which leads to an increase in dielectric

permittivity and loss of the material. However, further increase in frequency, above the

relaxation time, results in a sudden decrease in real and imaginary permittivity, due to the

exclusion of the polarization mechanism from the dielectric properties of the material. The

relaxation times corresponding to each mechanism are defined by a peak in real and imaginary

permittivity.

The most important dielectric mechanism in CPCs is interfacial polarization which occurs in

materials with multiphase structures. This phenomenon is also known as the Maxwell-Wagner-

Sillars (MWS) effect [32]. In this mechanism polarization occurs at the interface of different

phases with different dielectric permittivities and results in accumulation of mobile charges at

the boundaries of different phases. Due to the larger scale of polarization (longer relaxation

time), this mechanism usually occurs at low frequencies. Depending on the frequency and the

structure of the materials in CPCs, dipolar, atomic and electronic polarizations may also occur

due to the electric dipoles, charged ions, and nucleus displacements, respectively [32].

2.8.2 Dielectric properties of CPCs

Polymers are known to have very low dielectric constants, in the range of 2-5, which is mainly

due to the electronic and atomic polarization of hydrocarbon covalent bonds in polymers [45]. In

some exceptional cases, dielectric constants as high as 10 have also been reported for polymers,

which is still too low for dielectric applications. Consequently, dielectric fillers are added to

polymers to boost their dielectric properties. This way, dielectric fillers can contribute to the

dielectric properties of the composites, while polymers can offer light weight, flexibility and ease

of processing [42].

There are two main classes of fillers used for the production of dielectric composites, including:

(i) non-conducive high-k fillers, such as ferroelectric ceramics and organic fillers,

24

(semiconductive oligomers) and (ii) conductive fillers such as metals, carbon-based materials,

and conducting polymers [43].

Non-conducive high-k fillers are known to contribute to the dielectric properties of composites

due to their high dielectric constant and low dielectric loss. In these composites, the interfacial

exchange coupling mechanisms between the fillers and the matrix plays an important role in

determining the dielectric properties. The main drawback of using high-k fillers is that typically

high loadings of over 50 vol.% are required to achieve sufficiently high dielectric constants.

Such a high loading of fillers can result in severe losses in flexibility and other mechanical

properties of the composites [42].

By replacing the non-conducive high-k fillers with conductive fillers, it is possible to produce

dielectric systems with desired combination of high dielectric permittivity and low dielectric loss

for dielectric applications. Experimental studies in the literature reported permittivity values as

high as 2000 in silver flake-doped/epoxy composites [46] and 400 in Ni/PVDF composites [47].

Such high dielectric properties of CPCs are ascribed to the formation of huge numbers of nano-

capacitors, consisting of conductive fillers as nanoelectrodes and thin insulating layers of

polymer acting as dielectrics in between these electrodes.

At low filler contents, where the number of nanoelectrodes is not sufficiently high, and the

average distance between the particles is too large to work as effective nano-capacitors, inclusion

of conductive fillers cannot significantly contribute to the capacitance of CPCs, and the dielectric

characteristics of polymeric matrix are dominant. Further increase in the filler content results in

increase in the number and surface area of the polarizable domains, which can boost the

dielectric permittivity of CPCs. However, such increase in dielectric permittivity of CPCs is

often associated with simultaneous increase in their dielectric loss which is mainly due to the

ohmic loss of the polarized charges in percolated network of conductive fillers. Such an increase

in dielectric loss is unfavorable for the charge storage capability of dielectric materials. Thus,

theoretically, there exists a very narrow range of filler content, where the number of conductive

fillers is high enough, but the system is not yet fully percolated. Under these conditions, it is

possible to achieve a combination of high dielectric permittivity and low dielectric loss which is

desirable for dielectric applications. However, achieving these conditions is extremely

challenging and can be easily disturbed by slight variations in filler content and level of

25

dispersion and distribution of fillers. Several techniques have been proposed in the literature, to

enhance the dielectric capacity of CPCs through morphological control of conductive network to

increase their dielectric permittivity and suppress their dielectric loss, simultaneously. The

following section provides a brief overview of some of the most common techniques [48], [49].

2.9 Morphological control of conductive network in CPCs

2.9.1 Technical concerns associated with filler contents in CPCs

Using conventional processing techniques, very high concentrations of conductive fillers (5–20

vol. %) are usually required to achieve electrical conductivities of 10-3-10 S/cm in CPCs [50].

There are several concerns associated with introducing such high contents of conductive fillers in

CPCs, including their processability, increased density, poor surface finish, high cost and loss of

mechanical properties.

As Kharchenko et al. reported, achieving the electrical percolation in CPCs is associated with the

formation of 3D networks of conductive fillers, which can also result in liquid-like to solid-like

transition of the composites. Such an increase in elasticity (G') of the CPC melts at higher filler

contents is known to increase their viscosity which is extremely unfavorable from a processing

point of view [51].

Furthermore, experimental results indicate that even though the incorporation of low contents of

CNTs improved the mechanical properties of polypropylene (PP)/CNT composites, a further

increase in the CNT contents resulted in severe deterioration in tensile strength, impact

properties and flexibility of the CPCs, mainly due to the aggregation of fillers that can act as

mechanical failure concentrators [19].

2.9.2 Dispersion and distribution of conductive fillers in CPCs

The formation of agglomerates in CNT-based CPCs is mainly due to very high aspect ratio of

conductive fillers and the presence of active groups, such as carboxyl and hydroxyl groups, on

the surface of CNTs which increase their tendency to interact with each other. Such interactions

are primarily through hydrogen bonding or weak van der Waals forces, which result in the

creation of clusters and agglomerates in CPCs [19].

26

As shown in Figure 2-3, the first step required to achieve the percolation threshold in CPCs, is to

break down these agglomerates, using intense processing techniques for uniform dispersion and

distribution of conductive fillers. However, as schematically demonstrated in Figure 2-3, perfect

distribution and dispersion of fillers is not the optimal condition for achieving the percolation

threshold. Using additional morphological techniques can result in promoting the percolation

threshold in CPCs at even smaller CNT contents. Consequently, several methods have been

proposed in the literature to promote the conductive network formation and minimize the filler

content required to achieve the percolation threshold. Such techniques involve the morphological

control of conductive fillers, based on understanding the processing conditions and filler-

polymer interactions in CPCs.

Figure 2-3: Schematic demonstration of the effect of morphological control techniques in

achieving the low percolation threshold in CPCs.

2.9.3 Morphological control through polymer blending

Polymer blending is one of the most stimulating topics in the field of polymer processing.

Polymer blends are generally used to combine the advantageous properties of their components.

27

Depending on the morphology and the content of different components, a wide range of

properties can be achieved using polymer blending.

In the case of CPCs, polymer blending is known to greatly affect the morphology of the

conductive network, though a mechanism known as selective localization of conductive fillers in

a more thermodynamically favorable phase in polymer blend. The distribution of nanoparticles

in polymer blends is most commonly described by the tendency of nanoparticles to minimize

their interfacial free energy (ΔG) and, given enough time, to migrate to the polymer phase which

is thermodynamically more favorable for them. According to the Young’s equation, localization

of nanoparticles in polymer blends in an equilibrium state can be estimated, using the wetting

coefficient (α):

α = 𝛾𝑠−𝐴 − 𝛾𝑠−𝐵

𝛾𝐴−𝐵 (12)

where γS-A, γS-B and γA-B are the interfacial tensions between each component pair, for polymer

A, polymer B and solid nanoparticles. Based on the Young’s equation, if the wetting coefficient

is less than -1 (α<-1), the nanoparticles tend to accumulate in phase A, while for α>1, the

nanoparticles migrate to phase B. Otherwise, when -1 < α< 1, the nanoparticles preferentially

localize at the interface, which can result in the smallest percolation thresholds in polymer blends

(Figure 2-4) [52].

Figure 2-4: Schematic presentation of selective localization of conductive fillers in polymer

blends, according to the Young’s equation.

This phenomenon can be especially effective in blends with co-continuous phase morphology in

which the fillers are selectively localized in one of the blend components or at their interface.

28

This structure can result in the formation of a conductive network in a double-percolated system.

Consequently, selecting the suitable polymer pairs with the optimized compositions can result in

great reductions in percolation threshold of CPCs.

In addition to thermodynamic parameters, mixing procedure and sequence, shear force and

mixing time can greatly affect the localization of conductive fillers in CPC blends. Other

important factors are viscosity of the polymers and aspect ratio and surface roughness of

conductive fillers. Such effects are attributed to kinetic factors which can compete with

thermodynamic parameters in dictating the localization of conductive fillers. Furthermore,

functionalization of conductive fillers can affect their surface energy and, consequently, change

the distribution of the fillers in polymer blends. Experimental results indicate that

functionalization of CNTs with covalently grafted PP, increase their interfacial bonding with PP,

and improve their localization in PP phase in case of PP/PE polymer blend matrix [53].

The ideal scenario for achieving the lowest percolation threshold in polymer blends is the

selective localization of conductive fillers at the interface in co-continuous blends. However,

achieving such a phase morphology is very challenging, specifically in the case of conducive

fillers with very large aspect ratios [54]. However, recently Wang et al. [55] reported the

selective localization of CNTs at the interface of immiscible polycarbonate/maleic anhydride

grafted acrylonitrile-butadiene-styrene blend, by adjusting the kinetic parameters as well as

thermodynamic factors. In this work, they reported a percolation threshold as low as 0.05 wt.%

in their CPCs.

2.9.4 Morphological control through the use of hybrid fillers

Like polymer blending, hybrid fillers can also be used for the fabrication of CPCs in order to

take advantage of different functionalities of various fillers. There are different types of

conductive fillers in the market with different prices, mechanical properties, electrical

characteristics, and aspect ratios. Mixing different types of fillers, can result in producing CPCs

with optimized functionality and cost. There are numerous theoretical and experimental studies

in the literature focused on the effects of mixtures of fillers on electrical conductivity of CPCs

[10].

29

Zhang et. al studied the synergic effects of different percolation mechanisms with hybrid fillers,

that resulted in grape-cluster-like conductive network formation which decreased the percolation

threshold in CPCs. Based on their experimental results the electrical conductivity and percolation

threshold of CPCs with hybrid fillers are often close to the electrical conductivity of CPCs with

exclusively high aspect ratio fillers. Considering the high price of CNTs, in comparison with

CBs, using hybrid system of CNT/CB with optimized compositions can be an effective technique

for optimizing the properties and cost of CPCs [56].

In addition to financial concerns, hybrid fillers also can be used to optimize the electrical

functionality of CPCs. In an experimental study performed by Arjmand et al. [57], they showed

that using a hybrid filler system of CNT and graphene nanoribbon (GNR), with a CNT/GNR

ratio of 3:1, resulted in superior dielectric properties of the hybrid system, in comparison with

CNT- or GNR-based CPCs. They attributed this improvement in dielectric characteristics of

CNT/GNR hybrid CPCs to the synergistic effects of the high dielectric permittivity of CNT and

GNR fillers and poor interlacing ability of GNRs to bridge CNTs and form a conductive network

in hybrid systems. Consequently, inclusion of GNR particles results in increasing the number of

nanoelectrodes, without increasing the dielectric loss of the CPCs.

In another study performed by Zhao et al. [58], they showed that EMI shielding properties of

PVDF/CNT/graphene composites were higher than those of PVDF/CNT and PVDF/graphene

composites. The mean EMI shielding values of PVDF/5 wt % CNT, PVDF/10 wt % graphene,

and PVDF/ 5 wt % CNT/10 wt % graphene composite films, with thickness of 0.1 mm, were

22.41, 18.70, and 27.58 dB, respectively. They attributed such improvement in EMI SE of hybrid

systems to the multiple reflections and greater interfacial polarization mechanisms which

resulted in improving the microwave absorption properties. Furthermore, the greater number of

conductive fillers coupled with their uniform state of dispersion generated a conductive network,

which consequently dissipated more electrical energy and thus resulted in a higher ohmic loss in

hybrid systems.

Hybrid systems can also be prepared via mixing conductive and non-conductive fillers to

improve the electrical conductivity or dielectric characteristics of CPCs. Experimental results

indicate that inclusion of clay in CPCs with carbon-based conductive fillers resulted in increase

in electrical conductivity of the CPCs, which was mainly attributed to the affinity of clay to

30

carbon-based fillers. Such affinity can accelerate the dynamic percolation mechanism and result

in exfoliation of conductive fillers.

2.9.5 Morphological control through foaming

Introduction of cellular structures to CPCs has been reported to result in changing their electrical

conductivity and morphology of the conductive network via two main mechanisms: (i) a volume

exclusion effect, and (ii) changing the alignment and orientation of conductive fillers located

close to the growing cells during the foaming process [59].

Experimental studies indicate that the uniaxial and biaxial stretching generated during the

foaming process results in reorientation of the conductive fillers with respect to growing cells.

Theoretical analysis performed by Wang et al. [60] and Shaayegan et al. [61] reveals that the

degree of displacement/rotation of conductive fillers strongly depends on the cell size, the initial

distance between the fiber mid-point and the cell center, and the initial fibers’ orientation angle

with respect to the cell nucleus. In addition, the relative size of conductive fillers has a

substantial impact on the conductive fillers orientation and translation with respect to cellular

structure.

As Ameli et al. [59] reported, the degree of biaxial stretching is proportional to the degree of

foaming in CPCs with cellular structures. At an optimal degree of foaming, the biaxial stretching

of polymer during the foaming process, resulted in promoting the conductive network formation

in CPCs. However, as the degree of foaming was further increased beyond the optimal point, the

conductive fillers became fully oriented normal to the cell and excessive stretching resulted in

the destruction of the conductive network in CPCs.

The effects of foaming on morphological control of conductive fillers in CPCs were further

emphasized in an injection molding process, where the excessive alignment of conductive fillers

is known to be a critical challenge in formation of conductive networks. As shown in Figure 2-5,

there are different mechanisms in which foaming can promote the conductive network formation

in CPCs: (i) First, the dissolved supercritical gas in the foaming process can result in substantial

decrease in the viscosity of the melt, due to plasticizing effect of gas, which can result in reduced

fiber breakage and lower in-plane fiber alignment during the injection process. Such effects are

known to promote the conductive network formation in CPCs. (ii) The dissolved supercritical

31

gas can also promote the fiber exfoliation and prevent fiber aggregation in CPCs. (iii)

Furthermore, generation of cellular structure in CPCs can also effectively disturb orientation of

the fibers and decrease the size of the skin layer, that usually has highly oriented fibers, and thus

enhance the through-plane conductivity [61]–[64].

Figure 2-5: Schematic presentation of the effects of foaming on conductive network

formation in foam injection molding process [64].

Recently, foaming has shown promise in promoting CPCs for various applications, including

charge storage and EMI shielding materials. Experimental results indicate that in PP/stainless-

steel fiber composites, incorporation of foaming resulted in improving the EMI SE through

absorption and multiple reflection mechanisms [65]. This effect was mainly attributed to the

volume exclusion effect and fiber alignments with respect to cellular structures in CPC foams

[65]. Furthermore, experimental results indicate that substantial alignment of conductive fillers

around the growing cells in CPC foams can promote the formation of nanocapacitors and

improve the charge storage capacity in PP/CNT CPC foams [63].

32

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39

Chapter 3 Conductive Network Formation and Destruction in

Polypropylene/Carbon Nanotube Composites via Crystal Control Using Supercritical Dioxide

3.1 Introduction

Over the last decades, conductive polymer composites (CPCs) have shown great potential as a

highly desirable class of advanced functional materials. CPCs possess a stimulating combination

of properties, including light weight, ease of processing, non-linear voltage-current behavior, and

environmentally-sensitive resistivity [1]–[5]. Such characteristics have led to the use of CPCs in

numerous applications such as charge storage [6], antistatic protection [7], sensors [8]–[13],

electrostatic dissipation [14], and electromagnetic interference (EMI) shielding [15], [16].

Polymers are generally considered as electrical insulating materials. Early attempts to produce

conductive polymer products involved the incorporation of conductive fillers, such as metallic

particles and carbon fibers, into a polymer matrix so as to reach the electrical percolation

threshold [17], [18]. The formation of a conductive network within the CPC subsequently

overcame the high resistance of the insulating polymer matrix surrounding the conductive fillers

[19], [20].

Of all the conductive fillers used in CPC production, carbon nanotubes (CNTs) have received

the most attention amongst researchers due to their exceptional mechanical, thermal, and

electrical properties [1], [2], [21], [22]. From a structural standpoint, CNTs have extremely high

aspect ratios (≈ 100-1000), low densities (≈ 1.3-2.0 g/cm3), and high specific surface areas

(≈100-1300 m2/g) [23]. Their high aspect ratios combined with their high electrical conductivity

(105-106 S/cm) have earned them an outstanding reputation as conductive fillers for the

development of CPCs [21], [22].

Although very high electrical conductivities have been reported for individual CNTs and CNT

ropes, the electrical conductivities of CNT-filled CPCs, at or above their percolation threshold, is

usually in the range of 10-7 to 10-3 S/cm [21]. This happens because the electrical conductivities

of these CPCs depend strongly on both the intrinsic conductivity of the individual CNTs and on

the contact resistance between them. The contact resistance between the CNTs can vary greatly

depending on the thickness of the matrix’s insulating layer at the contact point of the adjacent

CNTs. Above the percolation threshold, this insulating layer is sufficiently thin to allow electron

40

tunneling and/or hopping to occur, but its thickness can certainly affect the contact resistance

[19].

As noted above, the contact resistance of the CPCs’ conductive network can be dramatically

altered by slight geometrical changes in the local contact regions between the nanotubes [24].

Consequently, several studies have been undertaken to understand and control the conductive

network’s microstructure in CPCs [25]–[29].

Ameli et al. [30] showed that introducing a cellular structure to conductive composites can be an

effective way to control the CNTs’ interconnectivity. They found that at an optimum degree of

foaming, polypropylene (PP)/CNT composites with nano/micro-cellular structures had improved

electrical conductivities and a lower electrical percolation threshold. They suggested that this

effect was mainly caused by the uniaxial stretching (in the cell struts) and by the biaxial

stretching (in the cell walls) of the polymeric matrix around the growing cells during the foaming

action. At this optimum degree of foaming, the CNTs’ improved alignment enhanced their

interconnectivity and also improved the CPCs’ electrical conductivity. They further proposed

that the PP/CNT composites’ foaming affected their dielectric properties. At higher foaming

degrees, an excessive alignment of the CNTs around the growing cells decreased their

interconnectivity and, thereby, decreased their dielectric loss.

In an interesting study, Zeng et al. [31] considered the interactions between the CNTs and the

polymeric matrix as a way to control the CPCs’ conductive network. They reported that a weak

interaction (low compatibility) between the CNTs and the polymeric matrix resulted in the

production of a CPC with a low percolation threshold. They suggested that the low level of

interaction between the CNTs and the matrix prevented the polymer chains from wrapping the

CNT surfaces and, consequently, decreased the CNTs’ contact resistance. However, despite the

increased electrical conductivities in their composites, this method does not hold promise for the

production of high-quality CPCs. Severe filler agglomeration is expected to occur in the

processing of composites with high interfacial tension between the matrix and the fillers, and this

would result in poor mechanical properties [32].

Considering the extent of research performed on the interactions between CNTs and

semicrystalline polymers [23], [33], [34], very few studies have recognized the effects of the

polymers’ crystallization behaviors on the CPCs’ electrical properties [35]–[38]. In a rare

example, Quan et al. [35] showed that the formation of stereocomplex crystals in a blend of

poly(L-lactic acid) (PLLA) and poly(D-lactic acid) (PDLA) produced a volume exclusion effect

41

where CNTs had been diffused out of the stereocomplex crystals, and this significantly decreased

the percolation threshold.

However, in other cases, several research works in the literature suggest that crystal nucleation

ability of CNTs for semicrystalline polymers results in the wrapping of CNTs by a thick

transcrystalline layer of polymer crystals [39]–[42]. Such observations suggest that there exist

two different mechanisms through which crystals can affect the electrical conductivity in CPCs:

i) Homogenous crystal nucleation, where the crystals are nucleated in the bulk of the

semicrystalline polymeric matrix. In this case, the crystal growth can diffuse the CNTs into the

amorphous regions, causing a volume exclusion effect. This can promote the interconnectivity of

the conductive fillers and facilitates the conductive network formation. ii) However, in the

second scenario, heterogeneous crystal nucleation will cause the crystals to nucleate on the

surface of the conductive fillers and disrupt the conductive networks through encapsulating and

isolating the individual fillers. Therefore, depending on the crystal nucleation ability of CNTs for

different crystal structures, crystallization can affect the electrical conductivity of CPCs by either

promoting or destructing the conductive network. Nevertheless, to the best of our knowledge,

there have been no studies that have correlated the different crystallization mechanisms in the

matrix phase with the CPCs’ electrical characteristics.

In this context, where producing CNT-filled CPCs with low percolation thresholds is paramount,

and where detailed studies on the morphological control of these composites are sparse, we have

focused on the effects of different morphological factors. These include the effects of PP’s

crystal type, crystallinity, and chain mobility on the evolution of the conductive network in

PP/CNT composites. Isothermal crystallizations of PP/CNT composites were performed at

various temperatures in an ambient pressure and under high-pressure in supercritical carbon

dioxide (scCO2) conditions. It was found that the PP crystals’ transformations, which were

caused by the differences in the isothermal crystallization conditions, had a remarkable effect on

the CPCs’ electrical conductivity and dielectric properties.

3.2 Experimental

3.2.1 Materials and sample preparation

A commercial grade of PP, Moplen HP 400R (Lyondell Basell Industries), with a melt flow

index (MFI) of 25 g/10 min. (230ºC, 2.16 kg) was selected as the base resin in this research.

Multiwalled carbon nanotubes (MWCNTs), Nanocyl 7000, were supplied by Nanocyl S.A.

42

(Belgium). The MWCNTs were produced by the catalytic carbon vapor deposition (CCVD)

process with a diameter of approximately 10 nm and a length of 1.5 µm. Carbon dioxide (CO2)

with a 99.98% purity produced by Linde Gas was used as the plasticizing gas for the isothermal

annealing tests.

Initially, a masterbatch (MB) of PP and CNTs was produced with 10 wt.% CNTs. The MB was

produced by melt mixing in a Berstorff ZE25 twin-screw extruder, with a proprietary screw

configuration, with a screw length of 48D, a temperature profile of 180-200°C, a rotational speed

of 500 rpm, and a throughput of 5 kg/h. Consequently, composites with 2.5 wt.% and 5.0 wt.%

CNT were produced in a second extrusion step by diluting the MB using the same processing

conditions but with a higher throughput of 10 kg/h. Composites containing 0.7 wt.%, 1.0 wt.%,

and 1.5 wt.% CNT were prepared from the PP/2.5 wt.% CNT masterbatch using a DSM twin-

screw compounder (15 ml) with a screw speed of 100 rpm at 180–200ºC, and these were

compounded for 8 min. For consistency, the PP was also given the same processing history as

that of the composites and was denoted as neat PP.

The characterization samples were then produced with a thickness of 0.7 mm using a laboratory

hydraulic press, Carver SC7620, at 200ºC under a pressure of 5 kN for 8 min. This was followed

by rapid water cooling with a controlled water flow rate and temperature.

In order to study the state of dispersion and distribution of CNTs in PP matrix, transmission

electron microscopy (TEM) and light microscopy (LM) were conducted on PP/CNT composite.

TEM analysis was performed on cryomicrotomed samples, using Hitachi H-7000 TEM (Hitachi

Ltd., Japan). The amount of remaining agglomerates in the micrometer scale was visualized

using thin sections of extruded strands in the transmission mode with an Olympus BH2

microscope.

3-1 shows a representative TEM and optical micrograph of PP/2.5 wt.% CNT composite. As

shown in Figure 3-1, the CNTs were uniformly dispersed and distributed in the PP matrix and

few remaining agglomerates could be observed in the optical micrographs.

43

Figure 3-1: (a) TEM and (b) optical micrograph of PP/2.5 wt.% CNT composites.

3.2.2 Isothermal annealing under atmospheric and high pressures

We investigated the effects of the annealing temperature and the scCO2 pressure on the electrical

conductivity and the crystallization behavior of the PP/CNT composites by annealing the

samples in an autoclave chamber under atmospheric pressure and also under a scCO2 pressure of

31 MPa for 15 min., at various annealing temperatures. It should be noted that CO2 can act as a

plasticizer and affect the dispersion of nanofillers in extrusion or injection molding systems.

However, since the batch-type gas impregnation was used in this work and since the experiments

were performed below the melting temperature of the polymeric matrix (solid state), then there

was no significant melt flow in the system. Consequently, the plasticizing effects of the dissolved

gas could not have had substantial effects on the CNT dispersion. After the annealing, the

chamber was rapidly quenched in water. Then, in the case of high-pressure experiments, the

scCO2 was released at a very low pressure drop rate to avoid any foaming effects (that is, to

avoid biaxial stretching of the PP around the growing cells) on the crystallinity and on the

electrical conductivity of the PP/CNT composites.

3.2.3 Differential scanning calorimetry (DSC)

3.2.3.1 DSC investigations before the isothermal annealing of the samples

To investigate the non-isothermal melt crystallization and melting behaviors of the neat PP and

the PP/CNT composites, the samples were examined using a DSC instrument (TA Instruments,

Q2000). The samples were initially heated to 220°C at a heating rate of 10°C/min. and

equilibrated at 220°C for 10 min. Next, the samples were cooled to -20°C at a cooling rate of

10°C/min. Then, they were reheated to 220°C at a heating rate of 10°C/min. to monitor the

a

b

https://www.google.ca/url?sa=t&rct=j&q=&esrc=s&source=web&cd=4&ved=0CEIQFjADahUKEwi_trbjjq7IAhWEGB4KHRsuAgg&url=https%3A%2F%2Fen.wikipedia.org%2Fwiki%2FDifferential_scanning_calorimetry&usg=AFQjCNG2HS2TuNuCn7RC_QfPwzwhtHfDIA&sig2=PGwb9CgGjnHTOklO0j_p_A

44

PP/CNT composites’ calorimetric properties. We calculated the composites’ degrees of

crystallinity (xc) using the following equation [43]:

𝑥𝑐 =Hf

(1−𝑤𝑓)Hm× 100 (1)

where Hf is the heat of fusion of the sample, Hm is the heat of fusion for 100% crystalline PP

(209 J/g), and wf is the CNTs’ weight fraction.

3.2.3.2 DSC investigations after the isothermal annealing of the samples

We also investigated the effects of the isothermal annealing (under atmospheric pressure and

under high-pressure scCO2 as a temporary plasticizer) on the crystallization behaviors of the neat

PP and PP/CNT samples. Following the isothermal annealing of the PP and its composites, as

described in section 2.2, the samples were heated from 30ºC to 200ºC at a heating rate of

10ºC/min. using the DSC instrument.

3.2.4 Polarized optical microscopy (POM)

A Nikon polarized optical microscope (Tokyo, Japan) equipped with a digital hot stage was used

to examine the melt crystallization behavior of the neat PP and the PP/CNT composites before

isothermal annealing. Thin films of the PP/CNT composites were first placed between two glass

cover slips and set on the hot stage under a microscope. The crystallization temperatures chosen

were 120ºC and 130ºC for the neat PP and the PP/1.0 wt.% CNT composites, respectively.

3.2.5 Wide angle X-ray scattering analysis

The wide angle X-ray scattering (WAXS) patterns of the PP/CNT composites were determined

using a Philips PW3710 X-ray Diffractometer. To study the effects of the isothermal annealing

processes, the WAXS patterns of the composites after the isothermal annealing were compared

with the samples before the isothermal treatments. The tests were performed under ambient

conditions using a Ni filtered Cu-Kα radiation with a wavelength (λ) of 1.5406 Å. A generator

voltage of 40 kV and a current of 40 mA were used. The experimental data were collected in a

slow scan mode (0.02 deg/s), within a range of 5°- 30° (2θ).

3.2.6 Electrical conductivity measurements

An Alpha-A high performance conductivity analyzer by Novocontrol Technologies GmbH &

Co. KG was used to measure the through-plane electrical conductivity, the real and imaginary

dielectric permittivity, and the dielectric loss of the samples at a voltage of 1.0 volt. To analyze

45

the nanocomposites’ broadband characteristics, the measurements were made across a wide

frequency ranging from 1×10-1 Hz to 1×10+5 Hz. For purpose of comparison, the direct current

conductivity, the σDC, was assumed at a frequency of 0.1 Hz, and the comparative dielectric

analyses were conducted at a frequency of 102 Hz. The calculated values for the percolation

threshold were obtained based on the classical percolation theory [44]. At least three replications

were carried out in each case, and the average values were reported.

3.3 Results and discussion

3.3.1 Thermal behavior analysis of the PP/CNT composites

3.3.1.1 DSC investigation before the isothermal annealing of the PP/CNT composites

Figure 3-2 shows the DSC thermograms of the PP and PP/CNT composites’ melt crystallizations

at a cooling rate of 10°C/min. The experimental results showed that a single crystallization peak

appeared at 117C for the neat PP. However, in the case of the PP/CNT composites, the peak

crystallization temperatures (Tc) shifted to higher values as the CNT content increased. This

indicated that the nanotubes had actually caused crystallization as a heterogeneous crystal

nucleating agent for the PP. It can be seen that the inclusion of 0.7 wt.% of CNTs increased the

Tc by nearly 10°C. Further increases in the concentration of the CNTs from 0.7 wt.% to 5 wt.%

only gradually increased the Tc. This behavior is consistent with previous reports in the

literature for PP composites as well as for other semicrystalline polymer composites [23].

46

Figure 3-2: Cooling DSC thermograms of neat PP and PP/CNT composites of varying

wt.%.

Figure 3-3 shows the polarized optical micrographs of the crystals that were created by the

crystallization of the neat PP and the PP/1.0 wt.% CNT composites. In Figure 3-3-a, it can be

seen that the sizes of the PP spherulites in the neat PP sample were very large which was

ascribed to a low crystal nucleation density. Figure 3-3-b, however, shows that for the 1.0 wt.%

CNT composite, the crystal nuclei density was much higher than that of the neat PP, and the

crystal sizes were significantly smaller than those of the neat PP. These observations were also

ascribed to the heterogeneous crystal nucleation effect of the CNTs which had been shown

earlier in Figure 3-2.

Figure 3-3: Polarized optical micrographs of (a) PP and (b) PP/1.0 wt.% CNT composites.

Figure 3-4 shows the DSC thermograms obtained from heating the neat PP and the PP/CNT

composites. In the neat PP, the main endothermic melting peak was observed at 162.8C.

Compared to the neat PP, the PP/CNT composites’ melting peaks shifted slightly to higher

temperatures, indicating that some crystals became more perfect with the addition of CNT. The

experimental results in Table 3-1 show that the addition of 1.0 wt.% CNT increased the degree of

crystallinity (Xc) of the neat PP from 47.1 to 51.9%. However, further increase in the

concentration of the CNTs did not increase the composites’ crystallinity. This suggests that the

CNTs play two different and competing roles in the PP’s crystallization: (i) Heterogeneous

100 µm 100 µm

b

a

47

crystal nucleation, and (ii) Constraining the PP chains’ mobility, which hinders crystal growth by

reducing the PP molecules’ ability to reach the crystallization sites [45].

Figure 3-4: Heating DSC thermograms of neat PP and the PP/CNT composites of varying

wt.%.

Table 3-1: The melting temperature (Tm), the heat of fusion (∆Hf), the degree of

crystallinity (Xc), the glass transition temperature (Tg), and the melt crystallization

temperature (Tc) for the neat PP and the PP/CNT composites

Melting Crystallization

Sample Tm

(ºC)

∆Hf

(J/g)

Xc

(%)

Tg

(ºC)

Tc

(ºC)

Neat PP 162.8 98.4 47.1 -4.6 117.2

PP/0.7 wt.% CNT 164.0 100.5 48.4 -4.6 127.5

PP/1.0 wt.% CNT 164.8 107.4 51.9 -4.9 128.2

48

PP/1.5 wt.% CNT 165.1 102.0 49.5 -4.9 129

PP/2.5 wt.% CNT 165.2 105.5 51.8 -4.9 130.3

PP/5.0 wt.% CNT 165.4 102.5 51.6 -4.8 132.2

3.3.1.2 DSC investigation after the isothermal annealing of the PP/CNT composites

Figure 3-5 shows the DSC thermograms of the neat PP, the PP/1.0 wt.% CNT, and the PP/2.5

wt.% CNT composites before and after their isothermal annealing for 15 minutes at 135-170C

under atmospheric pressure. Based on these results, the PP’s crystallization during its isothermal

annealing was strongly influenced by the annealing temperature, and three distinct regimes were

observed. Figure 3-5-a shows that in regime A the neat PP’s isothermal annealing at

temperatures of up to 150C did not result in any significant increase in the PP crystals’ Tm.

These results indicate that no significant crystal perfection mechanisms occurred during the

isothermal annealing at temperatures of up to 150C under atmospheric pressure. On the other

hand, the PP’s degree of crystallinity was improved, and it increased from 47% to 54% after

being annealed at 150C. This was due to the completion of its crystallization during the

annealing process. With respect to regime B, interestingly, Figure 3-5-a shows that a further

increase in the PP’s annealing temperature to 155C and 160C increased its Tm up to 169.8C

and 171.4C, respectively. We ascribed this increase in the PP’s Tm to the transformation of its

imperfect crystals, which were formed during the cooling step in the compression molding of the

samples, into more perfect structures with a higher Tm during the annealing process. In addition,

the neat PP’s degree of crystallinity reached a maximum value of 58.5% after annealing under

the atmospheric pressure at 160C. Further increases in the neat PP’s isothermal annealing

temperature to 165C and 170C (regime C) resulted in a sharp drop in its Tm. This happened

because during the isothermal annealing at such high temperatures the neat PP’s crystals were

completely molten. Subsequently, these new crystals formed during the cooling process, which

followed the isothermal annealing step. Thus, the crystals did not achieve perfection during the

49

annealing process. The annealing temperatures in regime C were selected to study the effects of

remelting and crystallization during the annealing process of CPCs.

As Figure 3-5-b and Figure 3-5-c show, the DSC thermograms of the isothermally annealed

PP/1.0 wt.% CNT and the PP/2.5 wt.% CNT composites’ trends were similar to the neat PP’s.

However, the occurrence of the aforementioned regimes B and C shifted to higher annealing

temperatures with the presence of the CNTs. This is mainly attributed to an increase in the PP’s

viscosity (that is, a reduction of the PP’s chain mobility) in the presence of the CNTs [46].

Figure 3-5: Non-isothermal DSC thermograms of the melting of (a) neat PP, (b) PP/1.0

wt.% CNT, and (c) PP/2.5 wt.% CNT composites after isothermal annealing under

atmospheric pressure.

Figure 3-6 shows the non-isothermal DSC thermograms obtained by heating the neat PP, PP/1.0

wt.% CNT and the PP/2.5 wt.% CNT composites before and after their isothermal annealing for

15 min. at 135-150C under 31 MPa scCO2. Similar to the results seen in Figure 3-5, in Figure 3-

6 the samples’ crystallization behavior under high-pressure conditions could also be divided into

the above-mentioned three temperature regimes. However, compared with the atmospheric

annealing results shown in Figure 3-5, the transitions between these regimes occurred at lower

temperatures in the scCO2 atmosphere. For instance, with the neat PP, the transitions between the

regimes A and B occurred at 142.5C and the transition between regimes B and C occurred at

147.5 C. This behavior can be attributed to an increase in the chain mobility of the PP

molecules that was due to the plasticizing effects of the dissolved scCO2 [45], [47]–[52]. In

50

addition, Figure 3-6 also shows that the samples, which were isothermally annealed under high-

pressure scCO2, had lower Tms as well as lower crystallinities compared with the samples that

were annealed under atmospheric conditions. For instance, the isothermally annealed PP/1.0

wt.% CNT composites showed Tms ranging from 164-176C (for annealing under atmospheric

pressure) and from 164-168C (for annealing under high-pressure scCO2). These outcomes

suggest that the PP matrix’s crystal structure/type changed during the annealing processes. We

thus utilized WAXS analysis to study the changes in the crystal structures and crystal types in the

PP/CNT composites after the different annealing processes.

Figure 3-6: Non-isothermal DSC thermograms of the melting of the (a) PP, (b) PP/1.0 wt.%

CNT, and (c) PP/2.5 wt.% CNT composites after isothermal annealing under 31 MPa

scCO2 condition.

3.3.2 Analysis of the crystal types and structures of the PP/CNT

composites

Depending on the crystallization conditions, PP is known to take three polymorphic crystalline

forms including the monoclinic -form, the hexagonal -form, and the orthorhombic -form

crystals. The -form crystals are the most common crystal form of PP and these are created

under such typical processing conditions as compression molding. The -form can be created by

either adding specific nucleating agents or by exerting a shearing stress to the polymer [53].

However, PP’s -form crystals are less common and are generally observed either during

crystallization under sufficiently high pressures or under shear-controlled orientation [54]. It has

been reported that PP’s -form crystals have unique characteristics such as a lower Tm compared

51

with the crystals [55]. On the other hand, the DSC results (Figure 3-5 and Figure 3-6) showed

that the isothermal annealing of the PP and its composites in an scCO2 atmosphere reduced the

crystals’ Tm. These findings led us to investigate the possibility of the presence of PP’s

crystalline polymorphs in the samples before and after their annealing in a high-pressure scCO2

atmosphere. Figure 3-7 shows the WAXS patterns of the neat PP and the PP/CNT composites

with CNT contents of 1 wt.% and 2.5 wt.% before and after their high-pressure isothermal

annealing at different temperatures.

As shown in Figure 3-7, the WAXS patterns of the neat PP and the PP/CNT samples (before

treatment) demonstrated the presence of both crystalline and amorphous states in the matrix

phase. The samples’ main crystalline peaks, indexed as (110), (040), (130), (111), and (041),

confirmed the α-form monoclinic packing of the PP crystals after their compression molding

[54]. The isothermal annealing of the samples at 135ºC and under 31 MPa only increased the

intensity of their already existing peaks, and this signaled an increase in their crystallinities.

However, no new peaks were formed, which suggests that annealing the samples under high

pressure and at 135ºC did not change the PP’s crystalline forms. However, a further increase in

the isothermal annealing temperature to 150ºC, under an scCO2 pressure of 31 MPa, resulted in

the appearance of a new peak at 2θ=20.1º, which confirmed the formation of the PP’s γ-form

crystals [54], indexed as (117) in Figure 3-7. These findings explain the reduction in the

samples’ Tm under the conditions shown in Figure 3-6, where the crystals had a lower Tm than

the crystals. Figure 3-6 also shows that the isothermal annealing of the samples at 150ºC and

under 31 MPa scCO2 reduced their crystallinities. This behavior can be explained as follows: (i)

At 150ºC and under 31 MPa scCO2, the PP phase was completely molten during the 15 min. of

annealing, and all of the PP’s crystals were formed during the samples’ relatively rapid cooling

under high pressure, and (ii) The kinetics of the crystallization for PP’s -form crystals is known

to be slower than its crystals [54].

Further, a simultaneous reduction in the sample’s crystallinity and the formation of new

crystals is expected to significantly reduce the content of α crystals after the isothermal annealing

process under high-pressure scCO2 at a temperature of 150ºC. In general, the WAXS peaks of

the PP’s - and -form crystals, when they co-exist, closely overlap with each other. Only the -

form (130) and the -form (117) reflections could be distinguished in these samples.

Accordingly, the samples’ -form content could be obtained by using the heights of the (130)

52

reflection of the -form (I (130)) and the (117) reflection of the -form (I (117)), and is thus given

as follows [56]:

% = 𝐼(117)

𝐼(117)+𝐼(130)𝛼100 (2)

Therefore, based on Figure 3-7 and Equation (2), with the annealing process at 150ºC and under

31 MPa scCO2, nearly 60% of γ-form crystals were created in PP/1.0 wt.% CNT composite. This

on the other hand indicates a significant reduction in the amount of α-form crystals in PP/CNT

composites.

Interestingly, as Figure 3-7-a, 7-b, and 7-c show, the inclusion of the CNTs produced more -

form crystals after isothermal annealing under high-pressure scCO2 at a temperature of 150ºC.

For instance, after being annealed in these ways, the content of the -form crystal, compared with

the -form, increased from almost 40% for the neat PP to nearly 65 % for the PP/2.5 wt.%

composite. However, in the literature the CNTs nucleating abilities for PP crystals have been

reported to be lower than they are for -form crystals in the transcrystalline layer [42]. Indeed, it

has been reported that PP -phase crystals ideally grow epitaxially on the lateral face of the -

phase crystals [54]. Thus, the increase in content of crystals after the inclusion of CNTs could

be ascribed to the indirect effect of the CNTs’ presence. In the PP/CNT composites, the CNTs

preferably act as nucleating agents for the -form crystals, which reduces the average size of the

-form crystals and increases the number of -form crystals. This facilitates the nucleation of the

crystals by offering further access to the (010) lateral face of the α-phase structure, that is an

ideal surface for the epitaxial crystallization of the γ-lamellae crystals [57], [58].

Consequently, with the annealing process at 150ºC and under 31 MPa scCO2, a significant

amount of γ-form crystals was created (that is, a significant reduction in the amount of α-form

crystals). This reduced the CNT’s overall crystal nucleation efficiency and, it likely increased the

CNT-CNT contacts in the composites. Such an effect could potentially increase the composites’

electrical conductivity.

53

Figure 3-7: WAXS patterns of (a) neat PP, (b) PP/1.0 wt.% CNT, and (c) PP/2.5 wt.% CNT

composites before and after isothermal annealing under 31 MPa scCO2.

3.3.3 Effects of crystallinity on the electrical conductivity of the PP/CNT composites

Figure 3-8 shows the electrical conductivities of the PP/CNT composites before and after their

isothermal annealing at different temperatures and under different pressures. As expected, the

inclusion of the CNTs increased the PP’s electrical conductivity. In all cases, a sharp increase in

the composites’ electrical conductivities was observed at certain CNT concentrations, which

corresponded to their electrical percolation thresholds. However, Figure 3-8 shows that the

values of the composites’ percolation thresholds were significantly influenced by their

preparation conditions (that is, the temperature and pressure that were used for their isothermal

annealing). Based on these results, the effects of the isothermal annealing temperatures under

atmospheric and high scCO2 pressures on the electrical conductivities could best be categorized

by using the concept of the temperature regimes, which was introduced in section 3.1.2.

Figure 3-8-a shows that before isothermal annealing, the percolation threshold of the PP/CNT

composites was observed at a CNT concentration of around 0.86 wt.%. Isothermal annealing of

the composites under atmospheric pressure and at regime A did not significantly change their

electrical conductivities, and their percolation threshold values remained unchanged. However,

increasing the annealing temperature up to 160ºC in regime B slightly improved the electrical

conductivity of the PP/CNT composites. This can be ascribed to the increase in degree of

crystallinity and volume exclusion effects of the PP crystals after annealing at 160ºC under

atmospheric pressure. Further increase in the isothermal annealing temperature of the composites

54

to 175ºC (regime C) led to an improved electrical conductivity of the PP/0.7 wt.% CNT

composite from 910-13 S/cm to 3.510-10 S/cm. This improvement in the electrical conductivity

of the composite can be ascribed to a possible CNT agglomeration (that is, a reduction in the

level of the CNTs’ dispersion within the matrix) due to the interfacial interactions between the

CNTs, coupled with an increase in their mobility at such high annealing temperatures. In other

words, an agglomeration of CNTs created CNT-rich percolated clusters within the composite,

even at a lower concentration than is required to produce a globally percolated network. Alig et

al. [24] also reported that the CNT cluster aggregation during the isothermal annealing process

resulted in the formation of a conductive network, which increased the CPC’s electrical

conductivity. Figure 3-8-a also shows that isothermally annealing the composites with a higher

CNT content (higher than 1.5 wt.%) at 175ºC did not improve their electrical conductivities.

The electrical conductivities of the PP/CNT composites, which were isothermally annealed at

different temperatures under 31 MPa scCO2, are shown in Figure 3-8-b. At a low annealing

temperature of 135ºC under a high scCO2 pressure (that is, in regime A), we observed significant

reductions in the electrical conductivities of the composites with low CNT content. This behavior

can be ascribed to the following two factors: (i) it has been reported that the heterogeneous

crystal nucleation efficiency is significantly increased in PP composites when they are annealed

at both high scCO2 pressures [45], [59] and at low annealing temperatures (lower than 138ºC, as

suggested by Thomason and Van Rooyen [60]). This results in the CNTs being wrapped by a

dense crystalline layer of the insulating polymer. (ii) During annealing under high pressures, the

dissolution of the scCO2 in the PP causes swelling in the matrix, which increases the average

distance between the adjacent CNTs in the composites. This effect coupled with the CNTs’

increased crystal nucleating ability increases the likelihood of the formation of new PP crystals

in areas between the adjacent CNTs. Based on our results, this phenomenon permanently

increases the distance between these CNTs (that is, it increases the contact resistance between

the CNTs), even after the depressurization step and the shrinkage of the composites to their

original volumes.

At an annealing temperature of 145ºC (regime B), the electrical conductivities of the composites

were slightly improved. For example, the electrical conductivity of the PP/0.7 wt.% CNT

composite was increased from 910-13 S/cm to 3.510-12 S/cm. As noted in section 3.1.2, in

regime B, the isothermal annealing of the composites led to the melting of the imperfect PP

crystals and to the formation of new crystals with perfect structures. However, since this

55

annealing process was performed at a relatively high temperature, where the CNTs’ crystal

nucleating ability was much lower, the newly formed crystals were not nucleated on the surface

of the CNTs. In other words, the CNTs experienced a volume exclusion effect from the PP

crystals. This led to an increased concentration of CNTs in the PP’s amorphous region, and it

increased the CNT-CNT contact.

After the composites were annealed at 150ºC under a scCO2 pressure of 31 MPa (regime C), the

highest electrical conductivities were observed, and the percolation threshold of the PP/CNT

composites was remarkably reduced by nearly 50% (from 0.86 wt.% to 0.45 wt.%). This

behavior can be ascribed to the simultaneous occurrence of two processes: (i) perhaps the most

important mechanism for this behavior was the increase in the CNT-CNT contacts due to the

CNTs’ reduced crystal nucleating ability. This was the result of the formation of a significant

content of -form crystals, as we explained in section 3.2, after the composites had been annealed

at 150ºC under high pressures. (ii) the possibility of the aggregation of the CNTs during this

annealing process.

Figure 3-8: Electrical conductivity of neat PP and the PP/CNT composites before and after

isothermal annealing at different temperatures, under (a) atmospheric pressure and (b) 31

MPa scCO2.

56

3.3.4 Effect of crystallinity on the dielectric properties

Recently, CPCs have been widely studied in the literature as potential candidates for dielectric

applications [61]–[63]. However, the successful application of CPCs as dielectric materials has

been an ongoing challenge. Ideally, a significant increase in the CPC’s dielectric permittivity (ε')

is required while keeping their dielectric loss (tan δ) at a low value [6]. Based on the literature,

the ε' of the CPCs has been shown to be readily increased via an increase in the conductive filler

content near the percolation threshold [6]. However, the CPCs’ increased ε' is usually

accompanied by a dramatic increase in their dielectric loss (tan δ) due to the sharp insulation-

conduction transition near the percolation threshold. Thus, a strict control over the morphology

of the CNT network (that is, a reduction in the CNT-CNT physical contact) is required for the

successful production of CPCs for dielectric applications. Various strategies have been proposed

to develop effective dielectric CPCs with lowered contact between their conductive fillers. These

include coating and modifying the surfaces of CNTs [61], [63], increasing the CNT alignment

through such methods as electrospinning and injection molding [63], using hybrid additives [64],

[65], and using microcellular structures in foam injection molding [6]. To reiterate, the main

objective of these noted techniques is to keep the CPCs’ ε' as high as possible while decreasing

the value of their tan δ.

Figure 3-9 shows the broadband dielectric permittivity and loss for the PP/1.0 wt.% CNT

composite before and after isothermal annealing at different temperatures under 31 MPa scCO2.

As expected, in the untreated sample, the composite showed high permittivity (ε'=113.0 at a

frequency of 100 Hz). However, the value of tan was also very high (tan δ=38.9 at a frequency

of 100 Hz) which would prevent its use as a dielectric material. The isothermal annealing of this

composite at 150ºC under 31 MPa scCO2 led to an increase in both the dielectric permittivity and

loss of the composite. This is in agreement with the increase in the electrical conductivity of this

composite, shown in Figure 3-8-b.

On the other hand, after the isothermal annealing of the composite at 135ºC under 31 MPa

scCO2, its permittivity values were only slightly reduced while a dramatic reduction in the

dielectric loss (as much as 3 orders of magnitude at low frequencies) was observed. Under these

conditions, the PP/1.0 wt.% CNT composite showed a dielectric permittivity of ε'=58.0 and a

low dielectric loss of tan δ=0.2 at a frequency of 100 Hz. This behavior can be ascribed to the

aforementioned changes in the microstructure of this CPC. Isothermal annealing of the PP/CNT

composites at 135ºC under 31 MPa scCO2, resulted in the formation of dense lamellae layers of

57

the insulating polymer crystals around the CNTs and at the CNT-CNT contact points which

suppressed the dielectric loss in PP/CNT composites.

Figure 3-9: Broadband dielectric (a) permittivity and (b) loss of PP/1.0 wt.% CNT

composite before and after isothermal annealing at 135°C and 150°C under 31 MPa scCO2.

3.4 Conclusions

Despite the extensive research performed on the interactions between CNTs and semicrystalline

polymers, very few studies have recognized the effects of the polymers’ crystallization

conditions on the electrical properties of CPCs. In this context, this research work is focused on

understanding the effects of different crystallization mechanisms on the electrical conductivity of

PP/CNT composites.

In this research, we introduced the concept of crystal control via isothermal annealing under

scCO2 conditions for manipulating the conductive network in CPCs with semicrystalline PP

matrix. We studied the effects of the different morphological factors, including the PP’s crystal

type, the crystallinity, and the chain mobility on the evolution of the conductive network in

PP/CNT composites. The isothermal crystallizations of the PP/CNT composites were performed

at various temperatures in an ambient pressure and under high-pressure in scCO2 conditions.

58

Experimental results of the current work successfully showed that the PP crystals’

transformations, which were caused by the differences in the isothermal crystallization

conditions, had remarkable effects on the CPCs’ electrical conductivity and dielectric properties.

In this work, we showed that through altering the crystalline structure of PP, from α-form

crystals into γ-form crystals, via annealing under scCO2 conditions and at elevated temperatures,

it is possible to effectively reduce the heterogeneous nucleation ability of the CNTs and to

promote volume exclusion effects of the growing crystals. Remarkably, our results showed that

the composites’ isothermal annealing at 150°C under a scCO2 pressure of 31 MPa led to the

formation of a significant amount of γ crystals and, thereby, reduced the electrical percolation

threshold by nearly 50% (from 0.86 wt.% to 0.45 wt.%). Alternatively, promoting the

heterogeneous crystal nucleation ability of the CNTs during the annealing process at lower

temperatures substantially increased the contact resistance between the conductive fillers through

wrapping the CNTs with insulating layer of crystalline polymer. We showed that this

phenomenon has enormous potential for developing dielectric materials with low dielectric loss

and high dielectric permittivity.

The results obtained in this research work demonstrate the important role of the crystal structures

in controlling the conductive network formation in CPCs.

59

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Chapter 4 Effect of Polymer-Filler Interactions on Controlling the Conductive Network Formation in Polyamide 6/Multi-Walled Carbon Nanotube

Carbon Nanotube Composites

4.1 Introduction

Ever since their emergence, electrically conductive polymer composites (CPCs) have been

considered as strong competitors for conventional metallic and ceramic materials. These

materials exhibit an exceptional combination of properties and functionalities, including

lightweight [1], ease of manufacturing [2], broad range of conductivity [3], and environmentally-

sensitive resistivity [4]. Due to these unique characteristics, CPCs have been considered as

promising candidates for numerous applications such as electromagnetic interference (EMI)

shielding [5]–[9], charge storage [10], [11], electrostatic dissipation [12], [13], and sensors [14],

[15].

Carbon nanotubes (CNTs) are one of the most promising candidates in CPCs production. Due to

their extremely high aspect ratios (≈100-1000), low density (≈ 1.3-2.0 g/cm3) and high electrical

conductivity (105-106 S/cm), inclusion of small amounts of CNTs results in the formation of a

three-dimensional conductive network within the insolating polymer matrices [16]–[18].

Theoretical analysis reveals that in CPCs containing such a three-dimensional network of CNTs,

the electrical properties are controlled by the CNT-CNT contact resistance, which in turn

strongly depends on the interfacial interaction between CNTs and polymeric matrix [19]. Ideally,

the CPCs would include a network of CNTs which are in direct contact [20]. However, the

achievement of direct CNT-CNT contact could be substantially hindered by high polymer/CNT

interactions [16], [21], [22].

In this context, understanding the effects of interfacial interactions of polymers and conductive

fillers is of great importance in the production of CPCs. previous studies have been focused on

the effects of the surface energies of polymer matrices and conductive fillers on the electrical

properties of their composites. For instance, Zeng et al. [23] reported that weak matrix-filler

interactions in composites based on polyoxymethylene (POM) and CNTs led to high electrical

conductivities and a low percolation threshold.

67

Crystallization behavior of the matrix phase is one of the most important factors affecting the

polymer-filler interfacial interactions. In CPCs with semicrystalline matrices, crystallization of

the matrix is known to impact the conductive fillers via two different mechanisms. (I) In systems

with low polymer-CNT interactions, homogenous crystal nucleation occurs where the crystals

nucleate within the bulk of the polymer [24]. This mechanism of crystal growth confines the

CNTs in the matrix’ amorphous region. Such volume exclusion effect promotes the conductive

network formation in CPCs [23], [25]. (II) On the other hand, nucleation of polymer crystals on

the surfaces of the CNTs (i.e., heterogeneous crystal nucleation) leads to their encapsulation by a

thick insulating transcrystalline layer [26]–[30] which can increase the CNT-CNT contact

resistance, or even prevent the formation of a conductive network.

Unlike the first mechanism, there is still a surprising lack of research in the literature focused on

investigating the role of transcrystal formation on controlling the conductive networks in CPCs

[30]–[33]. Alig et al. [28] performed in situ monitoring of the electrical conductivity of

polypropylene (PP)/CNT composites during cooling from the molten state. Interestingly, they

reported that cooling of PP/CNT composites from the molten to frozen state led to a sharp

reduction in the composite’s electrical conductivity at the PP’s onset of crystallization

temperature. However, the causes of such a reduction in conductivity were not investigated in

detail.

In a previous research performed by the authors, the importance of crystal nucleation on the

conductivity of CPCs was demonstrated via altering matrix crystallization behavior in the

presence of supercritical carbon dioxide [31]. In that work, a significantly improved electrical

conductivity of the polypropylene/CNT composites was achieved through the suppression of

polypropylene transcrystallization on the CNTs’ surfaces.

As a necessary extension of our ongoing research aimed at understanding the effects of polymer

crystallization on CPCs, this work was focused on the detailed study of matrix

transcrystallization on CNTs and its deteriorating effects on the CPCs’ electrical properties.

Polyamide 6 (PA6)/multiwalled carbon nanotube (MWCNT) composites were chosen, as an

extreme case with strong transcrystallization of matrix on CNTs, to emphasize the effects of

transcrystallization on (i) the conductive network formation and (ii) the electrical conduction

mechanism in CPCs with semicrystalline matrices.

68

4.2 Experimental

4.2.1 Materials and sample preparation

PA6/MWCNT composites were produced by compounding a commercial grade PA6, Ultramid

B3S (BASF Corporation), and MWCNTs, Nanocyl 7000 (Nanocyl S.A. Belgium). The

MWCNTs, which were produced by catalytic chemical vapor deposition (CCVD) process, had a

density of 1.75 g/cm3, a diameter of approximately 10 nm, and an average length of 1.5 μm.

A masterbatch of PA6/MWCNT: 85/15 wt.% was first produced using a Berstorff ZE25 twin-

screw extruder with an L/D of 48, using a temperature profile of 210-250°C. Consequently,

composites with 0.1-10 wt.% MWCNTs were produced by diluting the masterbatch by

compounding in a DSM twin-screw compounder (15 ml) with a screw speed of 100 rpm at 240

ºC for 8 min. For consistency, the PA6 was also given the same processing history as that of the

composites and was denoted as neat PA6. The materials were dried in an oven for 16 hours at

80ºC, before compounding.

Samples for characterization were then produced with a thickness of 0.7 mm using a laboratory

hydraulic press, Carver SC7620, at 240 ºC under a pressure of 5 kN for 8 min.

4.2.2 Morphological analysis

In order to study the microscale dispersion of the MWCNTs within the polymeric matrix, optical

transmission microscopy was performed on thin sections (2 μm thick) of the PA6/MWCNT

composites. The sections were cut from the compression molded samples, using a Laeica

Ultracut, RMC MT6000.

Transmission electron microscopy (TEM) was also conducted on the PA6/MWCNT composites,

using a Hitachi H-7000 TEM (Hitachi Ltd., Japan) to investigate the dispersion of the MWCNTs

within the PA6 matrix.

TEM was also used to analyze the crystalline structure of the PA6 phase in the composites. The

samples were first carefully stained, using an aqueous solution of phosphotungstic acid

(H3PO4.12WO3) and benzyl alcohol (C6H5CH2OH), for 10 minutes to enhance the phase

contrast between the crystalline lamellae and the amorphous regions of the PA6 matrix.

69


Non-isothermal DSC experiments were performed using a Q2000 DSC (TA Instruments) in a

nitrogen atmosphere under atmospheric pressure. The samples were initially heated from the

room temperature to 250°C at a heating rate of 10°C/min to remove their thermal histories. The

samples were then cooled to -10°C at 10°C/min and reheated to 250°C at 10°C/min. The

composites’ crystallinities (xc) were calculated using the following equation:

𝑥𝑐 =Hf

(1−𝑤𝑓)Hm× 100 (1)

where Hf is the heat of fusion of the sample, Hm is the heat of fusion for 100% α-form

crystalline PA6 (241 J/g), and wf is the MWCNTs’ weight fraction [34].

4.2.4 Wide angle X-ray scattering analysis (XRD)

Wide angle X-ray scattering (WAXS) patterns of the PA6/MWCNT composites were determined

using a Philips PW3710 X-ray Diffractometer. The tests were performed under ambient

conditions using a Ni filtered Cu-Kα radiation with a wavelength (λ) of 1.5406 Å. A generator

voltage of 40 kV and a current of 40 mA were used. The experimental data were collected in a

slow scan mode (0.02 º/s), within a range of 10°- 30° (2θ).


An Alpha-A high performance conductivity analyzer by Novocontrol Technologies GmbH &

Co. KG was used to measure the through-plane electrical conductivity of the PA6/MWCNT

composites at a voltage of 1.0 V. To measure the current-voltage characteristic of the

nanocomposites, a Keithley 6517A electrometer was used to perform a voltage sweep test in the

range of 1-90 V. At least three replications were carried out in each case, and the average values

were reported.

4.2.6 Rheological characterization

Rheological characterization was performed on the composite melts using a Reologica

ViscoTech oscillatory rheometer. The tests were performed in the dynamic oscillatory mode with

parallel plate fixtures with a diameter of 20 mm and a gap of 1 mm. All tests were carried out at

240 ºC under a nitrogen atmosphere to prevent thermal degradation and moisture absorption.

70

Initially, the composites’ linear viscoelastic regions were identified using dynamic strain sweep

tests. A strain of 3% was selected for all the frequency sweep tests in this work. Dynamic time

sweep tests were performed to investigate the degradation of samples. The results of the dynamic

time sweep tests suggested that the samples did not show any signs of degradation under the

characterization conditions selected for the rheological characterizations.


4.3.1 Dispersion and distribution of MWCNTs

The macroscopic mechanical and electrical properties of CPCs strongly rely on the level of

dispersion, distribution and alignment of the conductive fillers within the polymeric matrix [35].

Therefore, the morphological analysis of CPCs is of great importance to understand the

properties of such composites. Figure 4-1 presents TEM and optical micrographs of the PA6/2

wt.% MWCNT composite. As shown in Figure 4-1-a, MWCNTs are uniformly dispersed and

distributed within the polymeric matrix. The state of the dispersion of the nanotubes, as

displayed in Figure 4-1-a, suggests that a percolated network of MWCTs should be achievable at

relatively low concentrations of nanotubes.

There exist several studies in the literature reporting the suppressing effects of flow induced

orientation of nanofillers on the electrical conductivity of CPCs [11], [36]–[38]. However, based

on the TEM analysis in this work, the MWCNTs did not show any signs of a preferred alignment

in the compression molded samples. Furthermore, Figure 4-1 -b shows a light microscopy image

of the PA6/2.0 wt.% MWCNT composite from a lower magnification perspective to further

strengthen the conclusion that MWCNTs were well dispersed within the matrix. The figure

shows that there are very few MWCNT agglomerates in the composite, which indicates a

uniform dispersion of the MWCNTs.

71

Figure 4-1: (a) TEM and (b) optical micrograph of PA6/2 wt.% MWCNT composite.

4.3.2 Crystallization characteristics of the PA6/MWCNT composites

To investigate the crystallization behaviors and crystal structures of the PA6/MWCNT

composites, DSC and XRD studies were performed.

Figure 4-2 presents the melting thermograms (second heating step) of the neat PA6 and

PA6/MWCNT composites. As shown in Figure 4-2, only one melting peak can be observed in

the case of neat PA6 as well as PA6/MWCNT composites, at around 220 ºC. The melting

temperature of the PA6 crystals did not change significantly with the incorporation of different

concentrations of MWCNTs. The results presented in Table 4-1 show that the inclusion of

MWCNTs up to 1 wt.% led to an increased degree of crystallinity of the PA6 from 34.3% to

38.3%, implying heterogeneous crystal nucleation ability of the MWCNTs for the PA6 crystals

[39], [40]. This effect will be further demonstrated in Figure 4-3, using the samples’

crystallization thermograms. However, a further increase in the MWCNT content resulted in a

reduction in the degree of crystallinity of the composites. As shown, the degree of crystallinity in

the PA6 phase decreased to 25.2% with the inclusion of 10 wt.% MWCNTs. The observed

reduction in the PA6’s degree of crystallinity with the inclusion of higher concentrations of

MWCNTs can be explained using the crystallization mechanism of thermoplastics [39]–[41]. It

is well known that the crystallization process of polymers involves two main steps: (i) crystal

nucleation and (ii) crystal growth via the diffusion of the molecular chains to the crystallization

sites [24]. Although the inclusion of the MWCNTs can lead to improvements in the PA6 crystal

100 µm

b

1 µm

a

72

nucleation, however, the crystal growth step can be significantly hindered by the inclusion of

high concentrations of MWCNTs which lead to a reduction in PA6 molecules’ chain mobility

[39]–[41]. This reduction can be caused by (i) the presence of well-dispersed nanotubes and (ii)

the physical crosslinking effect of the numerous heterogeneously nucleated PA6 crystals [40],

[42], [43].

Figure 4-2: Melting thermograms of the neat PA6 and PA6/MWCNT composites (heating

rate: 10 ºC/min).

Figure 4-3 shows the crystallization thermograms of the neat PA6 and PA6/MWCNT

composites. As shown in Figure 4-3, we observe a significant transformation in the

crystallization behavior of PA6 in presence of the MWCNTs. The crystallization thermogram of

the neat PA6 shows a singular sharp peak at around 195 ºC (Tc1). In the case of the

PA6/MWCNT composites, as shown in Figure 4-3 and Table 4-1, this peak was gradually shifted

to slightly higher temperatures. It can also be seen that a new peak appeared at around 210 ºC

(Tc2). Interestingly, the intensity and the temperature of the second peak (Tc2) increased by

73

increasing the MWCNT content, suggesting that the appearance of this crystallization peak in the

composites was related to the presence of MWCNTs.

This additional peak has not always been observed in PA6/MWCNT composites [44], [45]. As

Li et al. reported [45], the inclusion of 1 wt.% MWCNTs into PA6 resulted in a 3 ºC increase in

the crystallization temperature of PA6, and a broader crystallization peak. However, their results

did not show the formation of a new crystallization peak after the inclusion of MWCNTs. Brosse

et al. [46] observed that the intensity of the second crystallization peak is directly related to the

quality of the dispersion of the MWCNTs in the PA6/MWCNT composites. They thus concluded

that the second crystallization peak is associated with the interface between the polymer and

MWCNTs and can be an indication of their interfacial interactions.

Figure 4-3: Crystallization thermograms for the neat PA6 and PA6/MWCNT composites

(cooling rate: 10 ºC/min).

Table 4-1: The crystallization temperatures (Tc1) and (Tc2), the melting temperature (Tm),

the heat of fusion (∆Hf), and the degree of crystallinity (Xc) for the neat PA6 and the

PA6/MWCNT composites.

Sample Tc1 (ºC) Tc2 (ºC) Tm (ºC) ∆Hf (J/g) Xc (%)

74

Neat PA6 194.9 - 220.3 82.8 34.3

PA6/0.1 wt.% MWCNT 196.9 - 221.1 87.2 36.2

PA6/1 wt.% MWCNT 198.0 207.7 220.8 91.3 38.3

PA6/2 wt.% MWCNT 197.5 208.1 221.8 79.1 33.5

PA6/4 wt.% MWCNT 197.5 210.2 221.2 71.4 30.9

PA6/6 wt.% MWCNT 198.3 211.1 221.2 63.1 27.8

PA6/8 wt.% MWCNT 199.0 212.1 221.4 59.1 26.7

PA6/10 wt.% MWCNT 198.4 212.5 221.2 54.7 25.2

Possible mechanisms for the appearance of the second crystallization peak in the composites

could be (i) the formation of two different types of crystals, or (ii) a two-step crystallization

process due to the presence of the MWCNTs. To further investigate the crystal structures of the

samples, XRD analyses were performed on the neat PA6 and the PA6/MWCNT composites.

XRD patterns of the neat PA6, PA6/1 wt.% MWCNT composite and PA6/10 wt.% MWCNT

composite are presented in Figure 4-4.

The most common crystalline forms in PA6 at ambient conditions are monoclinic α structure and

pseudo-hexagonal γ structure [47]. It is known that the α-form crystals are thermodynamically

more stable, whereas the γ-form crystals are kinetically more favorable (i.e., they form faster

than the α-form crystals). These two crystal forms are usually associated with two different

melting temperatures, around 220°C for α-form crystals and 215°C for γ-form crystals [47].

XRD spectra of the α-form structure show two principal diffraction peaks at 20º from (200)

planes and at 23.7º due to the (002) + (202) planes. The γ-form structure has a principal

diffraction peak at 22º. Amorphous PA6 gives a broad diffuse halo with an intensity maximum at

around 22º [25].

XRD characterizations in Figure 4-4 show that only α-form crystals were present in both the neat

PA6 and the PA6/MWCNT composites. This observation is in accordance with the melting

75

thermograms in Figure 4-2, where only one melting peak, at 220 ºC, was observed for all

samples, which is the melting point of PA6’s α–form crystals. Logakis et al. also observed that

the inclusion of MWCNTs promoted the formation of the α-form crystals in their PA6/MWCNT

composites [34]. Additionally, Liu et al. also reported that one-dimensional nanotubes seem to

promote the formation of α-phase crystals, whereas in the case of two dimensional fillers, such as

clay nanoplatelets, γ-phase crystals were more favorable in PA6 composites [48].

The abovementioned observations demonstrate that the double crystallization peaks of the

PA6/MWCNT composites in the cooling curves, observed in Figure 4-3, were caused by a two-

step process of the formation of α-form crystals. This can be ascribed to the MWCNTs’ strong

heterogeneous crystal nucleation effect, which led to the appearance of the crystallization peak at

210°C for the nanocomposites. This effect is expected to lead to the formation of a thick PA6

crystalline layer on the surfaces of the MWCNTs.

Figure 4-4: XRD patterns for the neat PA6 and for the PA6/MWCNT composites indicated

on the plot.

76

To further investigate the changes in the crystallization behavior of the PA6 phase upon the

inclusion of the MWCNTs, TEM studies were performed on the neat PA6 and PA6/0.1 wt.%

MWCNT composite. The samples were first stained using phosphotungstic acid (H3PW12O40) to

reveal the morphology of the crystal structures and to create sufficient contrast between the

crystalline and amorphous phases of the PA6 matrix. During this process, phosphotungstic acid

diffuses into the amorphous phase while the crystalline phase is practically impermeable.

Therefore, the amorphous phase appears darker in the TEM micrographs [49].

Figure 4-5-a shows the lamellar structure of the crystalline phase in the neat PA6. As shown in

this figure, stacks of long lamellae can be seen on the scale of hundreds of nanometers. These

lamellae represent the homogeneously nucleated crystals. However, in the case of the PA6/0.1

wt.% MWCNT composite, Figure 4-5-b, two types of crystals can be observed: (i) Lamellae of

homogeneously nucleated crystals within the bulk of the matrix, which were similar to those

crystals nucleated in the neat PA6 sample. (ii) A heterogeneously nucleated transcrystalline

layer, with a thickness of around 100 nm, formed on the surface of the MWCNT. As shown in

Figure 4-5-b, this transcrystalline layer was formed by the growth of PA6 crystal lamellae

perpendicular to the nanotube axis. Brosse et al. suggested that PA6’ heterogeneous crystal

nucleation on the surfaces of MWCNTs is a result of epitaxial crystal growth due to

crystallographic matching of PA6 and MWCNTs lattices [46].

In the following section, we discuss the effects of the formation of such a transcrystalline layer

on the electrical properties of the PA6/MWCNT composites.

77

Figure 4-5: TEM micrographs of the neat PA6 (a) and PA6/0.1 wt.% MWCNT composite,

where zone A and zone B represent the lamellar growth in homogeneously nucleated and

heterogeneously nucleated crystals, respectively.

4.3.3 Percolated nanotube network formation in PA6/MWCNT composites

The rheological and electrical properties of CPCs are based on entirely different phenomena,

namely, the transfer of the mechanical momentum and the electrical current, respectively.

However, both these properties are directly controlled by the state of the filler network within the

polymeric matrix [50]. Therefore, it is worth studying the correlation between the electrical and

rheological percolations in CPCs to obtain a valuable insight into the conductive network

structure.

In the case of the rheological characteristics of CPCs, the mechanical stresses are known to be

transferred by MWCNTs as well as the polymer chains that are entangled with the MWCNTs

and can act as “entropy springs” [51]. Consequently, at filler contents close to the CPCs’

percolation thresholds, the formation of a continuous load-bearing network results in a sudden

change in the viscoelastic behaviors of the composites, called liquid-like to solid-like transition

[50]. During this transition, we observe a sudden increase in elastic modulus (G') of the polymer

a

b

A

A

B

78

melt in comparison with loss modulus (G"), which results in sharp increase in the storage

modulus to loss modulus ratio (G'/G"), at filler contents close to the percolation threshold of the

composites, as shown in Figure 4-6.

As shown in this figure, the rheological percolation of PA6/MWCNT composites occurred by a

MWCNT content less than 2 wt.%, signified by a sudden increase in the storage modulus to loss

modulus ratio (G'/G"). A further increase in the MWCNT content up to 10 wt.% did not change

G'/G" significantly.

However, unlike the sudden changes observed in the composites’ rheological properties at their

percolation threshold, Figure 4-6 shows that their electrical percolation occurred in the form of a

gradual insulator-to-conductor transition. This behavior can be ascribed to the fact that CPCs’

electrical conductivities depend not only on the formation of a percolated network of MWCNTs,

but also on the contact resistance between the adjacent MWCNTs. Such contact resistance can

vary greatly depending on the thickness of the insulating matrix layer at the contact point of

MWCNTs. In the case of our PA6/MWCNT composites, the formation of the heterogeneously

nucleated PA6 transcrystalline layer on the MWCNTs’ surfaces, as demonstrated in the

crystallinity studies, disrupted direct contact between adjacent MWCNTs and delayed the

insulating-to-conductive transition.

79

Figure 4-6: Rheological and electrical percolation curves of PA6/MWCNT composites. The

(G′/G′′) ratio was measured at 240°C and frequency of 0.1 rad s-1. Electrical conductivity

measurements were performed at room temperature.

In order to calculate the electrical percolation threshold (φc) of the PA6/MWCNT composites, a

power law model was applied, as shown in Figure 4-7. According to the classical percolation

theory [52], the following equation can be applied for a filler content above the percolation

threshold:

σ = σ0 (φ-φc)t (2)

where σ and σ0 are the measured conductivity and scaling factor, respectively. φ is the filler

volume content, φc is the filler electrical percolation threshold, and t is the critical exponent

which reflects the dimensionality of the system. The best linear fit for σ vs. (φ-φc) data on a log–

log scale was found for a filler percolation threshold of about 0.85 vol.% and a critical exponent

of 6.22 (Figure 4-7). Based on the coefficient of determination (R2 value), it is seen that the

PA6/MWCNT composites followed the percolation power law with an acceptable accuracy.

However, the value of t=6.22 is higher than the predicted universal critical exponent value of t≈2

for three-dimensionally percolated systems [52]. A high t value is known to be responsible for a

gradual rather than the expected steep increase of σ with the filler content [53], as shown in

Figure 4-6.

y = 6.2206x + 4.519R² = 0.99

-11

-9

-7

-5

-3

-3 -2 -1

Log σ

(S/m

)

Log (φ-φc)

80

Figure 4-7: Log-log plot of conductivity versus (φ-φc) to determine the percolation model

parameters for PA6/MWCNT composites. R2 is the correlation coefficient of determination

for the linear regression analysis of Equation 2 in the double-logarithmic scale.

Various models have been proposed in the literature to explain the origin of such discrepancies

in the t value (i.e., t>2), including the mean-field interpretation model [54], the Swiss Cheese or

random-void (RV) model and its generalization, and the tunneling inverted RV model, also

known as the tunneling-percolation model [54], [55]. Based on these investigations, such

universality breakdown could occur by a combination of the geometrical effects, such as the

existence of highly resistive thin necks, as suggested by the RV model, and the tunneling effects

[56], as proposed by Balberg, where there exists a wide inter-particle distance distribution that

can lead to non-universal high t values [57].

Consequently, such deviation from the universal t value in our PA6/MWCNT composite systems

indicates that unlike the classical percolation theory [52], where it is assumed that the conductive

fillers are in direct contact with each other in a randomly dispersed system, our system

experiences a very high contact resistance. In such a system, tunneling is the dominant

mechanism of electron conduction between MWCNTs. Such high contact resistance in the

PA6/MWCNT composites can be explained by the abovementioned formation of a dense PA6

transcystalline layer on the surface of the MWCNTs, resulting in a wide interparticle distance

distribution.

4.3.4 Electron-conduction mechanisms in PA6/MWCNT composites

Studying the electron transfer mechanisms within CPCs allows us to further understand the

observed non-universal conductivity of the PA6/MWCNT composites. There are three main

mechanisms by which the electrons can be transferred in CPCs. These mechanisms include

metallic, hopping and quantum tunneling conduction [58]. Metallic conduction can occur when

the conductive fillers are in direct contact with each other. Therefore, the free electrons that

belong to the conduction bands can be easily transferred through the percolating network of the

conductive fillers within the CPCs. Hopping and tunneling are the mechanisms in which the

electrons can surmount the potential barriers of the insulating layer of polymer at the MWCNT

81

contact points in CPCs. Hopping occurs by random thermal fluctuations which give the electrons

sufficient energy to get out of their localized state and move to the conduction bands [59].

Hopping can also occur under large electric fields, in which the electrons do not need the thermal

energy to get into their conduction band. This is known as the Poole-Frenkel effect which can

explain how an insulative material can conduct some minimal electricity under a sufficiently

large electric field [60]. Quantum tunneling enables the electrons to penetrate through the

insulative layer of the polymer, by a mechanism known as the internal field emission. As the

filler content increases, and the distance between the conductive fillers decreases, a very strong

internal field may be developed. This field is greater than the external voltage, by the factor of

M, where M is the ratio of the average size of the conductive fillers to their average gap width

[61]. Note that the strength of such internal field across the gap can be several hundred times

greater than the applied voltage [61]. This intensified electric field can provide the required

energy for electrons to pass through the insulating layer of polymers in CPCs [58].

To understand the conduction mechanisms in PA6/MWCNT composites, their current-voltage (I-

V) characteristics were studied and presented in Figure 4-8. As shown in Figure 4-8, all

PA6/MWCNT composites exhibit nonlinear I-V characteristics. There are two mechanisms by

which a composite may exhibit nonlinear I-V characteristics: (i) when nonlinearity is derived

from the intrinsic properties of the conductive fillers, or (ii) when the CPCs’ conductivity is due

to the onset of hopping or quantum tunneling, in the previously-insulating channels between

MWCNTs, under high voltages.

Considering the fact that MWCNTs are intrinsically ohmic conductive fillers [62], the second

scenario can be used to explain the nonlinearity of the I-V characteristics of the PA6/MWCNT

composites in Figure 4-8. In this case, the I-V relationship of the composite can be described as

[64]:

𝐼 ∝ 𝑉𝑛 (3)

where n is the nonlinearity exponent. The value of n equals 1, if the composite follows the linear

or ohmic conduction mechanism, and it can deviate from 1, when non-ohmic conduction is

dominant. The greater the deviation of this factor from unity, the greater is the contribution

played by nonohmic mechanisms.

82

Typically, in CPCs the nonlinearity increases as the filler content approaches the percolation

threshold. At high MWCNT contents, above the percolation threshold, CPCs are expected to

show linear I-V characteristics, as the ohmic conduction is expected to be the dominant

mechanism in the transfer of electrons between MWCNTs [62].

In our PA6/MWCNT composites, the values of n ranged from 0.9 for the neat PA6 matrix to 1.5

for the CPCs with filler contents around the percolation threshold. Interestingly, beyond the

percolation threshold, the n value continued to increase with the incorporation of larger

concentrations of MWCNTs. For instance, in composites with 10 wt.% MWCNTs, which is well

above the percolation threshold, the value of n equaled 1.8. Such discrepancy from the Ohm's

law indicates that, even at filler contents above the percolation threshold, the electrical

conductivity was controlled by the tunneling mechanism in PA6/MWCNT composites. In other

words, these observations demonstrate that even at high filler contents, the MWCNTs could not

reach sufficiently thin contact points due to the presence of the insulating layer of PA6 crystals

which hindered the electron conduction between MWCNTs.

Figure 4-8: Current-voltage characteristics of PA6/MWCNT composites with different

MWCNT content. Inset: Current-voltage characteristics of neat PA6 and PA6/2 wt.%

MWCNT composites at smaller scale.

83

4.4 Conclusions

Although many studies have focused on the interactions between MWCNTs and semicrystalline

polymer, very few studies have recognized the effects of matrices’ crystallization behavior on the

CPCs’ electrical properties. Traditionally, it was believed that crystallinity can contribute to the

conductive network formation, through a mechanism known as the volume exclusion effect.

However, the effects of transcrystal nucleation were mostly neglected as an important

mechanism in morphological control of conductive networks. Consequently, the current research

was focused on the study of the effects of heterogeneous crystal nucleation on the contact

resistance and the electrical conductivity in PA6/MWCNT CPCs, as an extreme case.

Morphological characterizations demonstrated effective distribution and dispersion of MWCNTs

within the matrix phase. DSC showed that the MWCNTs acted as strong heterogeneous

nucleation sites for PA6 crystallization. This was evidenced by the appearance of a secondary

crystallization peak, at higher temperatures, in the case of the composites (two-step

crystallization). XRD results showed that both crystallization peaks led to the formation of α-

form crystals, suggesting that the CPCs’ two-step crystallization was purely a result of different

crystal nucleation mechanisms (i.e., homogeneous vs. heterogeneous). The TEM results further

strengthened the conclusion that PA6 crystallized on the surfaces of the MWCNTs in the form of

thick transcrystalline layers. The insulating characteristics of such transcrystalline layers led to a

substantial increase in the CPCs’ electrical percolation threshold, as compared with their

rheological percolation threshold. Using the classical percolation theory, the critical exponent (t)

was calculated 6.22, which was substantially higher than the predicted universal critical exponent

value of t≈2. In addition, the CPCs were shown to have nonlinear I-V characteristics, even at a

high MWCNT content of 10 wt.%. These observations further proved the existence of very high

filler-filler contact resistance due to the presence of a thick transcrystalline layer of PA6. These

findings provide a valuable insight into the effects of the crystallization behaviors of CPCs’

matrix phase, particularly the heterogeneous crystal nucleation effects, on their electrical

properties.

84

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Chapter 5 Highly Stretchable Conductive Thermoplastic Vulcanizate/Carbon

Nanotube Nanocomposites with Segregated Structure, Low Percolation Threshold and Improved Cyclic Electromechanical

Performance

5.1 Introduction

The production of stretchable, elastic, and conductive materials using practical and cost-effective

techniques has recently attracted a great deal of attention. This is because these materials have

immense potential for use in various applications such as stretchable electronics, strain gauges,

and implantable devices [1]–[8]. Stretchable and elastic conductive polymer composites (CPCs),

especially elastomer-based CPCs, are particularly interesting. Their many advantages include

their light weight, ease of processing, and tailorable properties. Conductive polymer composites

can be produced by mixing polymers and conductive particles, such as carbon nanotubes

(CNTs) and metal nanowires using conventional polymer processing methods [9]–[13]. With

sufficient amounts of such conductive nanoparticles (that is, above their electrical percolation

threshold) conductive pathways can be established within the insulating polymer matrix via the

formation of a nanoparticle network [9], [11], [14]–[16]. Thus, the CPC’s electrical properties

are controlled by such parameters as the nanoparticle’s content and aspect ratio, as well as its

dispersion and distribution and its particle-matrix interfacial adhesion [12], [17] – [19].

Traditionally, the direct incorporation of conductive nanofillers into elastomers has been used to

produce stretchable and elastic CPCs [1], [20], [21]. However, the successful production of such

materials has been hindered by several limitations. These have been caused by the relatively

large amounts of nanoparticles that are required to reach the electrical percolation threshold. The

inclusion of large amounts of nanoparticles leads to substantial deterioration in the elastomers’

mechanical properties, especially its stretchability and durability under repeated stress cycles.

And both of these properties are essential qualities in elastic and stretchable CPCs [22]. These

effects are usually caused by low elastomer-nanoparticle interfacial adhesion, and the large

differences between the stiffness of semi-rigid nanoparticles and the soft elastomers. Both of

these cause a severe stress concentration at the elastomer-nanoparticle interfaces during the

deformation process [1]. The addition of large nanoparticle concentrations also limits the

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product’s industrial applications. This affects its cost, processability, surface finish, and

environmental impact [20], [23] – [25].

In addition to the properties of the matrix and the nanoparticle phases, the preparation strategy

can also significantly affect the CPCs’ percolation threshold [18]. For example, the production of

CPCs with segregated structures, in which the conductive nanoparticles are not randomly

dispersed throughout the entire CPC system, has generated significant interest as a way to reduce

the required conductive nanoparticle content [26], [27]. The co-coagulation of nanoparticle

suspensions with rubber latex has been a popular method by which to produce elastomer-based

CPCs with segregated structures. In these products, the conductive nanoparticles are located at

the interfaces between the polymeric matrix particles [23], [28], [29]. Potts et al. [23] reported

that co-coagulation, followed by the hot pressing of natural rubber (NR) latex with reduced

graphene oxide (RG-O) significantly reduced the percolation thresholds. This occurred via the

formation of a web-like morphology that consisted of RG-O nanoparticles around the NR

particles. However, despite the reductions in the required conductive nanoparticle content, the

production of segregated CPCs prepared in this way also severely reduced the CPC’s

stretchability.

Another promising technique to create stretchable and elastic CPCs is the production of

conductive thermoplastic vulcanizates (TPVs)that contain dispersed vulcanized elastomeric

particles within a continuous thermoplastic matrix [30], with segregated structures (i.e., with

conductive nanoparticles only present in the matrix) is also a promising technique toward the

production of stretchable and elastic CPCs [31], [32]. These products benefit from the reduced

percolation threshold of segregated CPCs, especially in the case of TPVs with substantial

elastomer contents. In addition, compared with elastomers, TPVs are easier to process and

recycle [33]. And they are usually produced by mixing virgin elastomers with thermoplastics.

The dynamic vulcanization of the virgin elastomers occurs during this mixing process. However,

producing TPVs with substantial elastomer content in this way often leads to the formation of a

co-continuous phase morphology, which is undesirable for both processing and recycling. In

addition, the production of TPVs with a segregated structure is possible only if the nanoparticle

presence is thermodynamically favorable in the thermoplastic phase [34], [35]. This significantly

limits the choices involved in the thermoplastic and the elastomer phases and leads to

complications in the TPVs’ production.

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Due to the growing interest in the production of elastic and stretchable CPCs and faced with the

previously noted drawbacks of the existing technologies motivated us to develop a facile

technique to produce conductive TPVs with substantial elastomer content and segregated

structures. We have proposed a simple, low-cost, and practical approach based on the

incorporation of large amounts of fine pre-vulcanized rubber (PVR) particles, instead of virgin

elastomer, into a thermoplastic/nanoparticle composite. This approach has the desired potential

to produce elastic and stretchable CPCs cost-effectively. However, there has been little research

into the use of pre-vulcanized rubber powder in stretchable conductive composites. This is

largely due to the lack of interfacial adhesion between pre-vulcanized rubber particles and

thermoplastics, which causes a lack of molecular entanglement and/or chemical bonding between

most thermoplastics and vulcanized rubbers [36]–[38]. This usually results in low mechanical

properties, and especially low stretchability, in TPV composites. Thus, any efforts aimed at the

production of TPVs containing pre-vulcanized rubber particles must address the issue of low

compatibility.

To address these challenges, we focused our research on the cost-effective production of

stretchable, elastic, and electrically conductive TPVs containing large quantities of PVR with

segregated phase morphologies. By using a maleic anhydride grafted polyethylene (MA-g-PE) as

the matrix, we found that it was possible to achieve remarkable levels of interfacial adhesion

between the thermoplastic matrix and the elastomeric particles. Our morphological observations

also confirmed a high level of CNT dispersion within the MA-g-PE matrix. As a result, the TPV

composites had extremely low electrical percolation thresholds combined with superior

stretchability, elasticity, and mechanical durability. To the best of our knowledge, such a unique

combination of segregated structural and high mechanical properties using pre-vulcanized rubber

particles has never been reported.

5.2 Materials and methods

5.2.1 Materials

The stretchable and elastic CPCs were produced using commercially available materials. The

MA-g-PE, Epolene C-26, with an average molecular weight of 65,000 g/mol, acid number of 8

(mg KOH/g), and a melt flow index of 8 g/10 min. (190°C and 2.16 kg) was supplied by

Westlake Chemical Corporation. The PVR, MicroDyneTM 50, was a micronized rubber powder

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(average particle size <50 m) was produced from end-of-life tire rubber. The PVR was

provided by Lehigh Technologies and was used as received. Multiwall CNTs, NC3100TM, were

purchased from Nanocyl S.A. (Belgium). The CNTs, with an average diameter of 9.5 nm and an

average length of 1.5 m were produced using the catalytic chemical vapor deposition (CCVD)

process and were purified up to 95 wt.%.

5.2.2 Sample preparation

A masterbatch (MB) of MA-g-PE/CNT: 90/10 wt.% was first produced using a solution-casting

method. The MA-g-PE was dissolved in boiling xylene by continuous stirring. The CNTs were

then added to the solution and the stirring was continued for 1h, followed by ultrasonication for

1h. The mixtures were then cast on Petri dishes and left to dry at room temperature for 48 h,

followed by drying at 60°C for 24 h under a vacuum. Neat MA-g-PE, MA-g-PE/CNT and MA-

g-PE/PVR/CNT compounds were then produced from the MB using a DSM twin-screw

compounder (15 ml) after processing for 8 min. at 190°C at a screw speed of 100 rpm. The

compounds were then compression-molded using a laboratory hydraulic press, Carver SC7620,

at 180°C for 5 minutes, and were then cooled by quenching in water.

5.2.3 Morphological characterizations

The microstructures of the compounds were studied via scanning electron microscopy (SEM)

using a JEOL JSM6610-LV SEM after their cryofractured surfaces were sputter coated with

platinum. Transmission electron microscopy (TEM) was also conducted on the MA-g-

PE/PVR/CNT composites using a Hitachi H-7000 TEM (Hitachi Ltd., Japan) to investigate the

CNTs’ dispersion within the MA-g-PE matrix.

5.2.4 Fourier transform infrared (FTIR) spectroscopy

Fourier transform infrared spectroscopy was performed on the compression-molded plates of the

neat MA-g-PE and the MA-g-PE/PVR: 50/50 wt.% using a Spectrum One FTIR spectrometer

from PerkinElmer. The spectra were collected within a range of 4000 cm-1 to 550 cm-1.


An Alpha-A high performance conductivity analyzer, Novocontrol Technologies GmbH & Co.

KG, was used to measure the through-plane electrical conductivities under no stretching

95

conditions. To analyze the nanocomposites’ broadband characteristics, measurements were made

across a wide frequency, ranging from 1×10-1 Hz to 1×10+5 Hz. For the purpose of comparison,

the direct current conductivity, the σDC, was assumed at a frequency of 0.1 Hz.

In order to study the electromechanical performance of the samples (i.e. the changes in the

electrical properties of the composites, with respect to the mechanical deformations), the strain-

dependent electrical properties of the samples under monotonic and cyclic tensile deformations

were measured using a BK Precision 2831E digital multimeter. The samples’ dimensions and

the tensile tests’ conditions are explained in the section below. The samples were stretched to

preset strains, and resistance measurements were conducted using a two-point probe technique

parallel to the stretching direction.

5.2.6 Tensile tests

Monotonic tensile tests were performed on the MA-g-PE and the compounds using an Instron

model 3367 at room temperature (23°C). Five rectangular-shaped specimens with dimensions of

0.4×5×100 mm3 were characterized for each sample. The tests were performed at a crosshead

speed of 50 mm/min. using an initial gauge length of 40 mm. Cyclic tensile tests (1,000 cycles)

were performed using an Instron model 5848 microtester. These tests were also performed on

rectangular specimens with dimensions of 0.4×5×100 mm3 using a crosshead speed of 50

mm/min. and an initial gauge length of 40 mm. In each cycle, the specimens were stretched to a

strain of 10%. The tensile modulus in these experiments were measured as the slope of initial

tangent of the stress-strain curves.


Non-isothermal DSC experiments were performed using a Q100 DSC from TA Instruments in a

nitrogen atmosphere under atmospheric pressure. The samples were first heated from room

temperature to 180°C at a heating rate of 10°C/min. to remove their thermal histories. The

samples were then cooled to room temperature at 10°C/min. and reheated to 180°C at 10°C/min.

5.2.8 Dynamic mechanical analysis (DMA)

Dynamic mechanical analysis experiments were performed to investigate the stretchable CPCs’

elasticity. A dynamic mechanical analyzer RSA 3 from TA Instruments was used in the tensile

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mode, and the tests were conducted according to the ASTM D4065. The tests were performed

over a temperature range of -80°C to 90°C at a heating rate of 3°C/min., a frequency of 1 Hz,

and a strain of 0.05%. The samples’ dimensions were 0.41060 mm3.


5.3.1 Morphological analysis of the TPV-based composites

As noted above, the CPCs’ mechanical and electrical characteristics are strongly dependent on

their morphological properties. Specifically, in TPV-based stretchable and elastic composites

containing large quantities of PVR particles posed two major morphological challenges: the

achievement of a high level of interfacial adhesion between the thermoplastic matrix and the

PVR particles and a high level of CNT dispersion within the thermoplastic matrix [20].

Figure 5-1 shows the PVR particles’ morphologies as well as the cryofractured surfaces of the

MA-g-PE/PVR/CNT composites. Figure 5-1-a shows that the PVR particles had irregular

shapes. The average diameters of the majority of the PVR particles were also smaller than 50

m. Figure 5-1-b and 1-c show the SEM micrographs of the cryofractured surfaces of the MA-g-

PE/PVR/CNT composite with 50 wt.% of the PVR and 0.5 wt.% of the CNTs (MA-g-PE/CNT:

99/1), respectively. The cryofractured composites’ surface were smooth and completely free of

voids, signaling the absence of PVR-matrix debonding during the fracture. Further, Figure 5-1-c

clearly shows that there were no gaps between the MA-g-PE matrix and the PVR particles, and

also that the PVR particles were broken during the sample’s cryogenic fracture. This behavior

can be ascribed to the extraordinary level of interfacial adhesion between the PVR particles and

the MA-g-PE matrix that was achieved. During the fracture of the composites with very high

compatibility levels between the components, the cracks could not propagate through the strong

interface. Therefore, they propagate through the matrix or the fillers, as Figure 5-1-b and Figure

5-1-c show. The composites’ morphologies were also studied using TEM micrographs, as Figure

5-1-d and 1-e show. Figure 5-1-d shows that the CNTs were uniformly dispersed throughout the

MA-g-PE matrix. In addition, Figure 5-1-e confirms that the CNTs were only present in the MA-

g-PE matrix. This was due to the CNTs’ inability to penetrate the three-dimensionally

crosslinked network of the vulcanized rubber molecules in the PVR. In addition, Figure 5-1-e

also reaffirms the presence of a strong interface between the MA-g-PE matrix and the PVR

particles.

97

We ascribed the remarkable interfacial adhesion between the PVRs and the MA-g-PE, which has

not been reported on for other thermoplastic matrices, to the possible formation of chemical

bonds between the maleic anhydride groups of the MA-g-PE and the unsaturated C=C bonds on

the PVR particles’ surfaces [39]. Fourier transform infrared (FTIR) spectroscopy was performed

on the MA-g-PE and the MA-g-PE/PVR: 50/50 wt.% to investigate the occurrence of such

chemical bonding. Figure 5-2 shows that in the MA-g-PE spectrum the characteristic peaks of

maleic anhydride groups, which were separate from the PE peaks, were observed at 1,714 cm-1

and 1,791 cm-1. On the other hand, in the MA-g-PE/PVR: 50/50 wt.% compound shown in

Figure 5-2, the intensity of both peaks associated with the maleic anhydride groups, were

significantly decreased. These results clearly confirm the occurrence of chemical bonding

between the maleic anhydride groups and the PVR particles. The chemical reaction between the

maleic anhydride and the C=C bonds in the vulcanized rubber particles has also been reported by

Tripathy et al. [39]. They investigated the possibility of sintering/fusing vulcanized NR particles

via high-pressure and high-temperature sintering of mixtures of vulcanized NR powder and

various chemicals, including maleic anhydride.

The previously noted morphological properties of the MA-g-PE/PVR/CNT composites (that is,

their segregated structure, the homogeneous CNT dispersion within the MA-g-PE and the high

compatibility of the elastic PVR and the MA-g-PE) were expected to produce superior

mechanical (high stretchability, elasticity, and mechanical durability) and electrical (low

percolation threshold) properties.

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Figure 5-1: SEM micrographs of the PVR particles (a) and MA-g-PE/PVR/CNT:

49.5/50/0.5 wt.% composite (b) and (c). TEM micrographs of the MA-g-PE/PVR/CNT:

49.5/50/0.5 wt.% composite (MA-g-PE/CNT: 99/1) depicting a high level of dispersion of

the CNTs within the MA-g-PE matrix (d) and strong interface between a PVR particle and

the MA-g-PE matrix containing CNTs (e).

Figure 5-2: FTIR spectra of the neat MA-g-PE and MA-g-PE/PVR: 50/50 wt.%.

99

5.3.2 Effects of PVR on the electrical conductivity of the TPV-based composites

Figure 5-3 shows the electrical conductivities of the MA-g-PE/CNT and the MA-g-

PE/PVR/CNT composites as a function of the CNT content. As expected, there were sudden

increases in the electrical conductivities of both the MA-g-PE and the MA-g-PE/PVR: 50/50

wt.% compounds with certain CNT loading content. These corresponded to their electrical

percolation thresholds[17], [40], [41]. As seen in Figure 5-3, the electrical percolation threshold

for the MA-g-PE/CNT composites was observed at a low CNT content of around 1.0 wt.%. This

behavior can be ascribed to the CNTs’ homogeneous dispersion within the MA-g-PE matrix, as

shown in Figure 5-1-d. This composite’s electrical conductivity increased continuously with

increased CNT content to beyond the percolation threshold. Figure 5-3 shows that the

development of a segregated phase morphology, via the inclusion of 50 wt.% of PVR particles,

led to a significant reduction in the electrical percolation threshold to a CNT content of around

0.5 wt.%. This outcome, along with the TEM micrograph shown in Figure 5-1-e, verify the

effective exclusion of the CNTs from the impenetrable PVR particles during the composites’

melt processing stage.


composites containing 50 wt.% PVR.

100

5.3.3 Monotonic deformation of the TPV-based composites

In the production of stretchable CPCs through direct incorporation of nanoparticles into

elastomers, the biggest challenges include: (i) Reduced CPC stretchability after high

concentrations of conductive nanoparticles have been included. This has been largely ascribed to

the severe stress concentration at the elastomer-nanoparticle interface during the stretching

process and is due to their greatly different tensile moduli [1], [42]. (ii) Severe loss of

conductivity during stretching as a result of nanoparticle alignment [20], [43]. Our TPV-based

CPCs with segregated structures, which have decoupled but well-bonded elastic and conductive

phases, had a unique phase morphology. This promised to produce substantial differences

between their mechanical and electrical properties (under stretching) when compared with the

traditional elastomer/nanoparticle CPCs. Figure 5-4 and Figure 5-5 present the tensile stress-

strain curves of the neat MA-g-PE, the MA-g-PE/CNT: 94/6 wt.%, and the MA-g-PE/PVR/CNT

composites with 50 wt.% of the PVR and different CNT content. Figure 5-4 shows that the

inclusion of 6 wt.% CNT led to increases in the MA-g-PE’s ultimate strength (by 25%) and

modulus (by 84%). However, the MA-g-PE’s elongation at break was significantly reduced (by

88%) after 6 wt.% of the CNT was added. On the other hand, in the MA-g-PE/PVR/CNT

composites with 50 wt.% of PVR, Figure 5-5 shows that the inclusion of 6 wt.% of the CNTs

into the MA-g-PE phase (sample MA-g-PE/PVR/CNT: 47/50/3 wt.%) led to a much less severe

reduction of 45% in the compound’s stretchability. This behavior can be explained as follows.

First, in TPVs with a good compatibility between their elastomer and thermoplastic phases and

with high elastomer concentrations, the mechanics of deformation, especially beyond the

thermoplastic’s yield point, is mainly governed by the elastomeric constituent; that is, the PVR

particles. In other words, during large deformations of these TPVs, the rubber particles are fully

stretched while the thermoplastic matrices are much less deformed [44]–[46]. Oderkerk et al.

[44] studied the orientation behaviors of different phases during TPV stretching by using infrared

spectroscopy in combination with tensile stress-strain measurements. They reported that during

the TPV stretching, which was based on nylon-6 matrix and ethylene propylene diene monomer

(EPDM) particles, the nylon-6 molecules’ degree of orientation was much smaller than in the

EPDM phase, which was almost completely stretched. They also reported that under similar

deformation processes the degree of molecular orientation in the neat nylon-6 was much higher

101

than that of the nylon-6 as the matrix in the TPVs. Oderkerk et al. [44] and Boyce et al. [46]

related these behaviors to an inhomogeneous deformation of the matrix phase during the TPV’s

stretching. This was also considered to be the reason why the TPVs’ elasticities were much

higher than the neat thermoplastics’. This behavior, which is unique to TPVs, can be used to

address the challenges in the production of stretchable CPCs. On the one hand, we expected that

the lower molecular orientation in the matrix phase of our TPVs would lead to less significant

reductions in the conductivity during the stretching (due to lower nanoparticle alignment).

Figure 5-4: Tensile stress-strain curves of the neat MA-g-PE and MA-g-PE/CNT: 94/6

wt.% composite.

102

Figure 5-5: Tensile stress-strain curves of MA-g-PE/PVR/CNT composites with 50 wt.% of

PVR and 50 wt.% of MA-g-PE/CNT composites containing 0 wt.% (a), 1.5 wt.% (b), 3

wt.% (c), and 6 wt.% (d) CNT.

5.3.4 Thermal and thermo-mechanical analysis of the TPV-based composites

The DSC experiments were performed to provide insight into how the inclusion of the PVR and

the CNTs would affects the MA-g-PE’s chain mobility and crystallinity. Figure 5-6 shows the

non-isothermal heating DSC thermograms as well as the MA-g-PE phase’s crystallinity in the

samples. The MA-g-PE’s DSC thermogram shows double melting peaks, indicating the presence

of crystals with different degrees of structural perfection. The inclusion of 6 wt.% of the CNTs

resulted in the disappearance of the lower temperature melting peak. This can be ascribed to the

CNTs’ heterogeneous crystal nucleation effect, which led to an increase in the content of the

MA-g-PE crystals with a higher level of perfection. Crystallinity measurements in Figure 5-6

also show that the inclusion of the CNTs led to a slight reduction in the MA-g-PE’s crystal

content (from 31.9% to 30.2%). This was an outcome of the MA-g-PE’s reduced chain mobility,

which decreased their ability to reach the crystallization sites. Nofar et al. [47] observed similar

behaviors in PLA-based composites containing nanosilica and nanoclay. Reduced chain mobility

can be caused by the following: (i) A physical hindrance due to the presence of a large

concentration of CNTs [48], and (ii) An increased number of nucleated crystals ,which act as

physical crosslinks for the MA-g-PE molecules [49]. Figure 5-6 shows that the inclusion of 50

wt.% of the PVR particles led to a reduction in the MA-g-PE’s melting point (that is, a reduction

in the level of MA-g-PE crystal perfection) and a reduction in its crystallinity from 31.9% to

28.4%. Both of these behaviors may be ascribed to the PVR’s low heterogeneous crystal

nucleation capability for the MA-g-PE matrix, combined with a reduction in the MA-g-PE’s

chain mobility due to the occurrence of chemical bonding between their maleic anhydride groups

and the PVR particles’ C=C bonds (as confirmed by the FTIR spectra in Figure 5-2).

103

Figure 5-6: Non-isothermal DSC thermograms as well as the MA-g-PE phase’s crystallinity

in (a) neat MA-g-PE, (b) MA-g-PE/CNT: 94/6 wt.%, (c) MA-g-PE/PVR: 50/50 wt.%, and

(d) MA-g-PE/PVR/CNT: 47/50/3 wt.%.

The dynamic mechanical analysis (DMA) experiments were performed to study in more detail

the changes in the MA-g-PE’s chain mobility, following the inclusion of the PVR particles and

the CNTs. Figure 5-7 shows the loss tangent (tan )-temperature curves of the neat MA-g-PE,

the MA-g-PE/PVR: 50/50 wt.%, and the MA-g-PE/PVR/CNT: 47/50/3 wt.%. In the neat MA-g-

PE, two peaks were observed at around -24°C ( transition) and at around 55°C ( transition).

The inclusion of the PVR particles in the MA-g-PE led to the appearance of a new peak at

around -48°C, which was related to the PVR’s glass transition temperature. Figure 5-7 shows

that the addition of both the PVR particles and the CNTs reduced the peak’s intensity at around

55°C. This can be ascribed to the MA-g-PE’s reduced chain mobility [50], [51]. This was due to

chemical bonding with the PVRs’ surface and the physical hindrance of the CNTs. Figure 5-7

also shows that the MA-g-PE’s peak at around 55°C had shifted to around 61°C and 65°C for the

MA-g-PE/PVR: 50/50 wt.% and the MA-g-PE/PVR/CNT: 47/50/3 wt.%, respectively.

Abdelmouleh et al. [52] observed a similar behavior in the LDPE/cellulosic fiber composites and

attributed it to the segmental immobilization of the matrix’s chains in the interfacial region.

104

Figure 5-7: Loss tangent-temperature curves of the neat MA-g-PE, MA-g-PE/PVR: 50/50

wt.%, and MA-g-PE/PVR/CNT: 47/50/3 wt.% obtained from DMA.

5.3.5 Cyclic deformation of the TPV-based composites

In many applications, electromechanical durability under repeated stress cycles is an important

requirement for elastic and stretchable CPCs. Figure 5-8 summarizes the results of the cyclic

tensile tests of the neat MA-g-PE and its compounds with the CNTs and the PVR particles.

Figure 5-8-a and Figure 5-8-b show that the addition of the CNTs reduced the MA-g-PE’s strain

recovery. For example, from the first stretching cycle the residual strains were about 2.8% (72%

recovery) for the neat MA-g-PE and about 3.5% (65% recovery) for the MA-g-PE/CNT: 94/6

wt.%. This behavior occurred despite the observed improvement in the MA-g-PE’s elasticity that

was shown by the DMA results after the inclusion of the CNTs. This can be ascribed to the large

differences between the magnitudes of the strain applied during the cyclic tensile tests (10%) and

the DMA tests (0.05%). The inclusion of nanoparticles usually results in reduced mechanical

durability in polymers during repeated large deformations (high strain sensitivity). This is due to

the stress concentration at the polymer/nanoparticle interface and the irreversible slip between

the polymeric matrix and the surfaces of the semi-rigid nanoparticles [1], [20]. Similar behavior

was reported by Li and Shimizu [20] for composites based on poly[styrene-b-(ethylene-co-

butylene)-b-styrene] (SEBS) and CNTs. Alternatively, Figure 5-8-c and Figure 5-8-d show that

the addition of 50 wt.% of the PVR particles improved the strain recovery of both the MA-g-PE

and the MA-g-PE/CNT: 94/6 wt.% composites by reducing their residual strains after the first

cycle to 2.2% (78% recovery). This behavior is in agreement with the DMA results. Figure 5-8-e

105

shows the changes in the stress at strain=10% (σ10%) of the samples during the cyclic tensile tests

for 1,000 stretching cycles. For all of the samples, the σ10% decreased during the first 100 cycles,

and remained relatively unchanged during the subsequent 900 cycles. It was shown that the

inclusion of 6 wt.% of the CNTs magnified the MA-g-PE’s loss of strength during the cyclic

tests. After 1,000 stretching cycles, the values of the σ10% of the MA-g-PE and the MA-g-

PE/CNT: 6 wt.% were reduced by 20.9% and 25.1%, respectively. This behavior was due to the

destructive effects of the previously noted stress concentration at the MA-g-PE/CNT interface

during the deformation process. Figure 5-8-e also shows that adding PVR particles to the MA-g-

PE matrix improved its mechanical durability by reducing the loss of the σ10% (after 1,000

cycles) to 15.5%. In the compound with 50 wt.% of PVR and 3 wt.% of the CNTs, there was a

low, 16.4%, loss of strength. The higher mechanical durability of the samples containing PVR

can be ascribed to (i) The lower deformation of the MA-g-PE matrix during the TPV stretching,

which led to a lower stress concentration at the MA-g-PE/CNT interface, as noted earlier, and

(ii) The high strain recovery of the PVRs due to their crosslinked molecular structures.

Figure 5-9 shows the absorbed energy for the different composites during the stretching in cycles

1, 100, and 1000 of their cyclic deformations. The absorbed energies were obtained by

calculating the enclosed area underneath the loading and the unloading stress-strain curves

during a tension up to a maximum of a 10% strain. It can be clearly seen that the addition of the

CNTs had a deteriorating effect on the MA-g-PE’s behavior as a significant increase in plasticity

(energy dissipation) was observed after the first loading cycle. Meanwhile, the inclusion of the

PVR particles greatly improved the CPCs’ behavior by limiting the extent of the plastic

deformation. Figure 5-9 shows that the MA-g-PE/CNT: 94/6 wt.% nanocomposite had by far the

most significant energy dissipation change (65%) after 1000 stretching cycles, as compared to 1

cycle. In contrast to this, the change in the dissipation of energy after 1000 stretching cycles for

the MA-g-PE/PVR: 50/50 wt.% and the MA-g-PE/PVR/CNT: 47/50/3 wt.% was only 40% and

41%, respectively.

106

Figure 5-8: Tensile stress-strain curves of the neat MA-g-PE (a), MA-g-PE/CNT: 94/6 wt.%

(b), MA-g-PE/PVR: 50/50 wt.% (c), and (d) during cycles 1, 100, and 1000 of their cyclic

tensile tests (maximum strain: 10%). Figure (e) shows the changes in the maximum tensile

stress σ10% (i.e. stress at strain=10%) of the samples during the cyclic tensile tests. The

inserts depict the morphology of the composites that were tested under the cyclic

deformations.

107

Figure 5-9: Dissipated energy for the neat MA-g-PE, MA-g-PE/CNT: 94/6 wt.%, MA-g-

PE/PVR: 50/50 wt.%, and MA-g-PE/PVR/CNT: 47/50/3 wt.% during cycles 1, 100, and

1,000 of cyclic tensile tests (maximum strain: 10%).

5.3.6 Electrical conductivity under monotonic and cyclic deformations

As previously noted, one of the biggest challenges for traditional stretchable CPCs, when having

conductive particles directly incorporated into the elastomer matrix, is the severe loss of

conductivity during their stretching process. This is due to the excessive nanoparticle alignment

which leads to reduced particle/particle contact. On the other hand, during TPV stretching, we

previously explained that an inhomogeneous deformation of the phases significantly reduced the

tension levels in the thermoplastic matrix. We expected this behavior to reduce the level of the

CNT alignment during our TPV-based CPCs’ stretching process. Consequently, this would

reduce the extent to which the samples’ conductivity deteriorated during stretching. Figure 5-10

shows the changes in the normalized electrical resistance (R/R0) of the MA-g-PE/CNT: 94/6

wt.% and the MA-g-PE/PVR/CNT: 47/50/3 wt.% during monotonic tension. As was expected

during the stretching process, the MA-g-PE/CNT sample’s resistance increased much more

rapidly than that of the TPV-based composite. The electrical resistance of the MA-g-

PE/PVR/CNT: 47/50/3 wt.% compound increased less than 3-fold after stretching to a strain of

50%. Under a similar stretching process, the MA-g-PE/CNT: 94/6 wt.% showed a 600-fold

increase in resistance. Figure 5-11 also shows the changes in the normalized electrical resistance

of the MA-g-PE/CNT: 94/6 wt.% and the MA-g-PE/PVR/CNT: 47/50/3 wt.% composites at a

0% and a 10% strain during cyclic tensile tests. The samples’ electrical resistance, both at a 0%

108

and a 10% strain, increased during the initial 100 cycles and then plateaued. Based on Figure 5-

11 and compared with the MA-g-PE/CNT: 94/6 wt.% composites, the TPV-based CPC had a

higher electrical conductivity retention after repeated stretching-relaxation cycles. The values of

the normalized resistance for the MA-g-PE/PVR/CNT: 47/50/3 wt.% composites at a 0% and a

10% strain were 1.22 and 1.09, respectively, after 1000 cycles.

Figure 5-10: Normalized electrical resistance (R/R0) of MA-g-PE/CNT: 94/6 wt.% and MA-

g-PE/PVR/CNT: 47/50/3 wt.% during monotonic tension at a rate of 50 mm/min.

Figure 5-11: Normalized electrical resistance (R/R0) of MA-g-PE/CNT: 94/6 wt.% (a) and

MA-g-PE/PVR/CNT: 47/50/3 wt.% (b) during cyclic tension at a rate of 50 mm/min and

maximum strain of 10%.

109

5.4 Conclusions

The production of stretchable and elastic CPCs with a low percolation threshold, high

stretchability, and high mechanical durability under repeated stress cycles is of great interest due

to their numerous vast potential applications. In this work, a cost-effective and practical

approach was proposed for the facile production of stretchable TPV-based CPCs with segregated

phase morphologies (i.e. with the CNTs present only in the MA-g-PE matrix phase). To ensure

the formation of a segregated structure, and to avoid the creation of co-continuous phase

morphology, fine pre-vulcanized rubber particles (PVRs) were used in the production of the

TPVs. Morphological observations confirmed the CPCs’ segregated phase morphology. The

PVRs were also shown to have excellent interfacial adhesion with the MA-g-PE matrix due to

the chemical bonding of the MA-g-PE’s maleic anhydride groups with the PVRs’ C=C bonds.

The CPCs’ segregated structures led to a 50% reduction in their percolation thresholds to around

0.5 wt.% CNT content. Interestingly, the TPVs showed a much less severe loss of stretchability

after the inclusion of the CNTs, compared with the neat MA-g-PE. For example, adding 6 wt.%

CNTs to the neat MA-g-PE reduced its elongation at break by 88%. On the other hand, the

elongation at break of the MA-g-PE/PVR/CNT: 47/50/3 wt.% (having 6 wt.% of the CNTs in the

matrix phase) had a 45% reduction in stretchability compared with the MA-g-PE/PVR: 50/50

wt.%. We attributed this behavior to the TPVs’ compatibility with the elastomer and the

thermoplastic phases as well as with the high elastomer concentrations. The mechanics of the

deformation were mainly governed by the elastomeric constituent. The DMA experiments

showed an improvement in the MA-g-PE’s elasticity after the PVR particles and the CNTs were

added to it. As we expected, the inclusion of the CNTs reduced the MA-g-PE’s strain recovery

and mechanical durability. This was due to the large stress concentrations at the

polymer/nanoparticle interface and the irreversible slippage that occurred during the stretching

process. Inclusion of the PVR particles, on the other hand, significantly improved the

composites’ strain recovery and durability under repeated strain cycles. Further, incorporating

the PVR particles substantially increased the TPV-based CPCs’ capability to retain their

electrical conductivity during the monotonic and cyclic stretching processes.

110

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116

Chapter 6 Facile Production of Flexible and Stretchable Conductive Polymer

Composites for EMI Shielding Applications

6.1 Introduction

Electromagnetic interference (EMI), caused by electromagnetic signals emitted by electrical

circuits, can affect the performance of surrounding electronic devices or cause radiative damage

and environmental health issues. Electromagnetic pollution is considered as a significant concern

associated with the use of electronic devices. Thus, EMI shielding is crucial in the production of

such products [1].

Nowadays, metals are the most common materials used for EMI shielding applications.

However, metallic shield materials have several drawbacks including high density, poor

mechanical flexibility and stretchability and corrosion sensitivity. Furthermore, metals do not

truly eliminate or mitigate the electromagnetic pollution. Electromagnetic waves are almost

completely reflected at the surface of metallic shield materials. This effect only protects the

environment beyond the metal sheets, while causing damage to the internal parts of the

electronic devices [2]. Consequently, numerous studies have been devoted to the development of

EMI shielding materials with a focus on electromagnetic waves absorption as the main shielding

mechanism [1].

In this context, conductive polymer composites (CPCs) have been considered as promising

alternatives for EMI shielding applications, taking advantage of their light weight, low cost and

ease of processing [3]. Most polymers are known to be intrinsically insulating materials, which

makes them transparent to electromagnetic signals. In CPCs, incorporation of conductive fillers

into polymer matrices leads to the formation of conductive networks which contribute to EMI

shielding through absorption of the radiation power. EMI shielding in CPCs occurs through the

dissipation of the electromagnetic energy in the networks of conductive particles, with limited

reflection of the electromagnetic wave at the surface of the shield material. The reduction in

reflection of the electromagnetic waves in CPCs is mainly due to the lower impedance mismatch

of CPCs in comparison with metals and also due to the dominant role of absorption through

dissipation of the electromagnetic energy through multiple networks of conductive fillers [4].

117

The main challenge in the use of CPC-based shield materials is the high conductive filler content

which is usually required to achieve the desired EMI shielding effectiveness (EMI SE). Such

high filler contents usually lead to significant drawbacks such as increased density,

embrittlement, loss of strength and increased cost of the shield materials. Consequently, it is

desirable to minimize the conductive filler contents required for achieving high EMI SE [5].

EMI SE of CPCs is determined by several factors such as thickness of the shield material, the

frequency of the emitted electromagnetic wave, the intrinsic properties of the polymeric matrix

and the conductive filler, the fabrication method and processing conditions [4]. Table 6-1

summarizes the experimental results of several studies regarding EMI SE of CPCs in the

frequency range of 8.0–12.0 GHz (X band).

Table 6-1: Summary of EMI SE of CPCs measured in the frequency range of 8.0–12.0

GHz (X band).

Composite Filler

content

(wt. %)

Thick-

ness

(mm)

EMI

SE

(dB)

Specific

EMI SE

(dB/mm)

Reference Notes

CNT/PC/PVDF 2.0 5.0 24.0 4.8 [6] Segregated structure

containing ionic liquid-

modified CNTs

CNT/PC 2.0 3.0 23.1 7.7 [7] Segregated structure

CNT/PC/SAN 3.0 5.0 25.0 5.0 [8] Co-continuous polymer

blend

CNT/PS 3.5 1.8 24.0 13.3 [9] Compression molded

samples

CNT/PS 5.0 2.0 10.0 5.0 [10] Injection molded samples

118

CNT/PS 5.0 2.0 21.0 10.5 [10] Compression molded

samples

CNT/PMMA/

UHMWPE

5.0 2.5 21.5 8.6 [11] Segregated structure

CNT/ABS 5.0 2.8 38.2 13.6 [12] Segregated structure

CNT/PE 5.0 2.1 46.4 22.1 [13] Segregated structure

CNT/PS 5.0 1.0 25.0 25.0 [4] Dispersion spraying

CNT/BR 6.8 1.0 11.0 11.0 [14] Flexible vulcanized BR

matrix

CNT/PS 7.0 - 19.0 - [15] Foam structure

CNT/Cellulose 10.0 2.5 20.8 8.3 [16] Foam structure

CNT/PLLA 10.0 2.5 23.0 9.2 [17] Foam structure

CNT/PTT 10.0 2.0 36.0 18.0 [18] Melt compounding

CNT/PU 10.0 1.5 29.0 19.3 [19] Functionalized

MWCNTs

PANI/CNT 10.0 2.0 35.0 17.5 [20] Intrinsically conductive

matrix

CNT/NR 10.0 2.6 59.8 23.0 [21] Segregated structure

PMMA/CNT

15.0 4.5 35.0 7.8 [22] coagulation method

using SWCNTs

CNT/Epoxy 15.0 2.0 25.0 12.5 [23] High aspect ratio

SWCNTs

119

CNT/PP 15.0 1.0 34.8 34.8 [24] Melt compounding

ABS/CNT 15.0 1.1 50.0 45.4 [25] Solution casting

CNT/HDPE 18.0 3.0 58.0 19.3 [26] Melt compounding

CNT/PU 20.0 2.0 17.5 8.7 [27] Solution casting using

SWCNTs

CB/EPDM 37.5 5.5 18.0 3.3 [28] Vulcanized EPDM

matrix

CNT/ WPU 76.2 0.8 80.0 100 [29] Highly filled and flexible

As shown in Table 6-1, the formation of a segregated structure of conductive fillers within the

polymeric matrices offers great promise for the development of CPCs with high EMI SE at low

filler contents [30]. In such segregated structures, selective localization of conductive fillers, at

the interface between the polymeric regions, results in the formation of dense conductive

networks at low filler contents which are highly desired for EMI SE in CPCs [31].

Although the formation of segregated structures has been proven to improve the electrical and

EMI shielding performance of CPCs, the agglomeration of conductive fillers at the interfaces of

the polymer domains is also known to substantially deteriorate their mechanical properties

including flexibility, strength, and mechanical durability. The formation of such agglomerates is

known to (i) hinder the molecular interactions between the polymer domains and (ii) lead to the

formation of micro-voids along the segregated conductive pathways leading to structural defects

in CPCs. This is considered as the main issue limiting the application of CPCs with segregated

structures as EMI shielding materials [5].

In a previous study, we showed that the incorporation of substantial contents of pre-vulcanized

rubber (PVR) particles into maleic anhydride grafted polyethylene (MA-g-PE)/carbon nanotube

(CNT) composites led to the development of flexible and mechanically durable thermoplastic

vulcanizates (TPV)-based CPCs with very low electrical percolation threshold. The PVR

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particles used in this work originated from end-of-life tire rubber particles, containing high

contents of carbon black and metallic particles. The impermeable nature of the three-

dimensionally crosslinked PVR particles, along with strong bonding between the PVR particles

and the maleic anhydride grafted matrix, proved very promising for the facile production of

CPCs with segregated structures. In this study, the potential of these CPCs for EMI shielding

applications was investigated. It was shown that these CPCs’ segregated structures along with

the presence of conductive fillers in PVRs, such as carbon black and metal particles, led to high

EMI SE at low CNT contents.

6.2 Materials and methods

6.2.1 Materials

The TPV-based CPCs were produced using MA-g-PE as the matrix phase, PVR particles and

CNTs. MA-g-PE, Epolene C-26, with an average molecular weight of 65,000 g/mol, acid

number of 8 (mg KOH/g) and melt flow index of 8 g/10 min (190°C and 2.16 kg) was supplied

by Westlake Chemical Corporation. The PVR, MicroDyneTM 50, was a micronized rubber

powder (average particle size <50 m) produced from end-of-life tire rubber. The PVR was

kindly provided by Lehigh Technologies and was used as received. Multiwall CNTs, NC7000TM,

were purchased form Nanocyl S.A. (Belgium). The CNTs, with an average diameter of 9.5 nm

and an average length of 1.5 m, were produced through the catalytic chemical vapor deposition

(CCVD) process.

6.2.2 Sample preparation

A masterbatch (MB) of MA-g-PE/CNT: 90/15 wt.% was first produced using solution-casting

method. MA-g-PE was dissolved in boiling xylene by continuous stirring. The CNTs were then

added to the solution and stirring was continued for 1h, followed by ultrasonication for 1h. The

mixtures were then cast on Petri dishes and left to dry at room temperature for 48 h, followed by

drying at 60°C for 24 h under vacuum. The neat MA-g-PE, MA-g-PE/CNT, and MA-g-

PE/PVR/CNT compounds were then produced from the MB using an Xplore MC5 micro

compounder (5 ml) after processing for 8 min at 190°C using a screw speed of 100 rpm. The

compounds were then compression-molded using a laboratory hydraulic press, Carver SC7620,

at 180°C for 5 minutes, followed by cooling via quenching in water.

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6.2.3 Characterizations

The microstructures of the compounds were studied via scanning electron microscopy (SEM)

using a JEOL JSM6610-LV SEM after their cryofractured surfaces were sputter coated with

platinum. Transmission electron microscopy (TEM) was also conducted on MA-g-PE/PVR/CNT

composites, using Hitachi H-7000 TEM (Hitachi Ltd., Japan) to investigate the dispersion of the

CNTs within the MA-g-PE matrix.

Tensile tests were performed on the MA-g-PE and the compounds using an Instron model 3367

at room temperature (23°C). Five rectangular shaped specimens with dimensions of 0.4×5×100

mm3 were characterized for each sample. The tests were performed at a crosshead speed of 50

mm/min using an initial gauge length of 40 mm.

The through-plane electrical conductivities were measured using an Alpha-A high performance

conductivity analyzer, Novocontrol Technologies GmbH & Co. KG. To analyze the

nanocomposites’ broadband characteristics, the measurements were made across a wide

frequency ranging from 1×10-1 Hz to 1×10+5 Hz. For purpose of comparison, the direct current

conductivity (σDC) was assumed at a frequency of 0.1 Hz.

EMI shielding characterizations were performed over the X-band frequency range (8.2–12.4

GHz), using an E5071C network analyzer. The samples were sandwiched between the two

waveguides of the network analyzer to collect the scattering parameters of the incident signal.

The power of the incident electromagnetic wave was equal to 0 dB·m, which corresponds to 1

mW. A sample thickness of 0.7 mm was used for the EMI shielding characterizations.

6.3 Results and discussions

6.3.1 Morphological analysis

Figure 6-1 shows the SEM and TEM micrographs of the PVR particles and MA-g-PE/PVR/CNT

composites. As shown in Figure 6-1-a, the PVR particles had irregular shapes and an average

particle size of less than 50 m. As mentioned earlier, the PVR particles, which were produced

from post-consumer tire rubber, are expected to be filled with large concentrations of carbon

black. Figure 6-1-b clearly shows the PVRs’ carbon black content in the form of circular

particles with diameters in the range of 20-50 nm. In addition, Figure 6-1-b also confirms the

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selective localization of the CNTs in the MA-g-PE matrix. This behavior was due to the CNTs’

inability to penetrate the three-dimensionally crosslinked network of the vulcanized rubber phase

in the PVR. As shown in the SEM micrograph in Figure 6-1-c, using pre-vulcanized rubber

particles, it was possible to create a segregated microstructure. Despite the large concentration of

the rubber phase, 70 wt.%, it is shown that the MA-g-PE/CNT compound formed a continuous

phase, thereby guarantying high conductivity of the CPC at low CNT contents. It is also shown

that the PVR particles had very high compatibility with the MA-g-PE matrix which was

evidenced by (i) the lack of gaps between them and (ii) the fact that the rubber particles were

fractured, not pulled out, during the preparation of the samples via cryogenic fracturing. As

mentioned in chapter 5, such high level of compatibility, which is rarely observed in PVR-filled

compounds, was due to chemical bonding between MA-g-PE’s anhydride groups and PVRs’

C=C bonds.

Figure 6-1: SEM micrographs of the PVR particles (a) along with TEM (b) and SEM (c)

micrographs of the MA-g-PE/PVR/CNT composite.

6.3.2 Mechanical characterizations

As mentioned before, the development of a segregated phase structure in CPCs usually leads to

substantial loss of their deformability which is undesirable for many EMI shielding applications.

Figure 6-2 shows the tensile stress-strain curves of the MA-g-PE, MA-g-PE/CNT: 85/15 wt.%,

MA-g-PE/PVR: 30/70 wt.%, and MA-g-PE/PVR/CNT: 25.5/70/4.5 wt.% (which has a MA-g-

PE/CNT: 85/15 wt.% as the continuous phase). As shown, the inclusion of 15 wt.% of CNTs into

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the MA-g-PE matrix led to a 95 % reduction in the tensile elongation at break. This behavior is a

result of the significant difference between the tensile moduli of the MA-g-PE and CNT phases

which results in severe stress concentration at their interface at large deformations. Interestingly,

the inclusion of a similar concentration of CNTs into the MA-g-PE phase of the PVR-filled CPC

led to a much less severe reduction in the elongation at break (55 %). The MA-g-PE/PVR/CNT:

25.5/70/4.5 wt.% composite with a segregated phase morphology is shown to have a relatively

high elongation at break of 106 %, which makes them a suitable option for the production of

flexible electronic devices. As explained in Chapter 5, such a flexibility and stretchability in our

segregated structures is mainly due to the formation of decoupled and well-bonded conductive

(MA-g-PE/CNT) and elastic (PVR) phases in MA-g-PE/PVR/CNT composites.

Figure 6-2: Tensile stress-strain curves of the neat MA-g-PE, MA-g-PE/CNT: 85/15 wt.%,

MA-g-PE/PVR: 70/30 wt.% and MA-g-PE/PVR/CNT: 25.5/70/4.5 wt.% composites.

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6.3.3 Electrical conductivity measurement

High electrical conductivity is a necessary requirement for EMI shielding materials. Thus,

electrical conductivity measurements were performed to explore the effects of the incorporation

of CNTs, localized in the MA-g-PE phase, and PVR particles. As shown in Figure 6-3, the

inclusion of CNTs into MA-g-PE led to an increase in the electrical conductivity. The

percolation threshold of the MA-g-PE/CNT composites was approximately 1.5 wt.%. Figure 6-3

also shows that the inclusion of the PVR particles, which led to selective localization of CNTs in

the MA-g-PE phase, reduced the CPCs’ percolation threshold substantially. In these

thermoplastic vulcanizate composites, with segregated phase morphologies, the percolation

threshold was around 0.45 wt.%. These findings confirm once more that these TPV-based CPCs

had a desirable combination of excellent electrical properties and flexibility.


composites containing 70 wt.% PVR.

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The electrical conductivities of MA-g-PE/PVR compounds with different PVR contents were

also characterized to investigate the effects of the inclusion of the PVR particles. As shown in

Figure 6-1, the PVR particles contained large concentrations of carbon black which is expected

to contribute to their electrical properties. This behavior was confirmed in Figure 6-4 by the

increase in the electrical conductivity of MA-g-PE/PVR compounds at PVR content of 80 wt.%.

Interestingly, Figure 6-4 shows that the inclusion of up to 75 wt.% of the PVR particles did not

substantially contribute to the electrical conductivity of the compounds due to a lack of direct

contact between the PVR particles. This behavior provides further proof of strong bonding and

wettability between the MA-g-PE matrix and the PVR particles.

Figure 6-4: Electrical conductivities of the MA-g-PE/PVR composites as a function of PVR

content.

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6.3.4 EMI Shielding characterizations

To demonstrate the potential of the MA-g-PE/PVR/CNT composites for EMI shielding

applications, the EMI SE and shielding mechanism of the composites were investigated. EMI SE

is defined as the ability of the material to shield the electromagnetic waves. Reflection,

absorption and multiple reflection are the three main mechanisms contributing to the total EMI

SE of shield materials. The total EMI SE value can be measured as the logarithm of the ratio of

the incident power (Pi) to the transmitted power (Pt), as:

SE = 10 log (Pi/Pt) (1)

Equation (1) was used to experimentally determine the contributions of absorption and reflection

mechanisms in EMI SE of our samples. Figure 6-5 presents the EMI SE of the neat MA-g-PE,

MA-g-PE/CNT: 95.5/4.5 wt.%, MA-g-PE/PVR: 70/30 wt.%, and MA-g-PE/PVR/CNT:

25.5/70/4.5 wt.% composites. As shown in Figure 6-5, the inclusion of CNTs resulted in an

increase in the EMI SE of MA-g-PE. This is mainly due to the absorption of electromagnetic

waves through the formation of the conductive networks of CNTs in MA-g-PE/CNT: 95.5/4.5

wt.% composite. Interestingly, incorporation of PVR particles resulted in slight increase in EMI

SE of the MA-g-PE matrix in MA-g-PE/PVR: 70/30 wt.% composite, which can be ascribed to

the existence of a considerable content of carbon black in the PVR particles. Our experimental

results indicate that inclusion of only 4.5 wt.% CNT resulted in EMI SE of 18 dB in MA-g-

PE/PVR/CNT: 25.5/70/4.5 wt.% composites with the thickness of only 0.7 mm. The value of

EMI SE in MA-g-PE/PVR/CNT: 25.5/70/4.5 wt.% composites was 50% higher than that of MA-

g-PE/CNT composite with same CNT content of 4.5 wt.%. It should also be noted that the EMI

shielding in both MA-g-PE/CNT and MA-g-PE/PVR/CNT compounds occurred through the

absorption, which is the most desirable mechanism for EMI shielding applications.

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Figure 6-5: EMI SE of the neat MA-g-PE (a), MA-g-PE/CNT: 95.5/4.5 wt.% (b), MA-g-

PE/PVR: 70/30 wt.% (c) and MA-g-PE/PVR/CNT: 25.5/70/4.5 wt.% (d) composites.

6.4 Conclusions

In this study, flexible CPCs were developed for EMI shielding applications via the inclusion of

large concentrations of PVR particles (70 wt.%) into MA-g-PE/CNT composites to create a

segregated phase morphology. Due to the impermeable nature of the PVR particles, CNTs were

selectively localized in the MA-g-PE phase which led to very low percolation threshold in MA-

g-PE/PVR/CNT composites (0.45 wt.%). In addition, our experimental results showed that the

carbon black particles in the PVR phase also contributed to the electrical conductivity of the

MA-g-PE/PVR/CNT composites. Inclusion of the PVR particles also resulted in improved

flexibility and stretchability of the MA-g-PE/CNT composites. Based on our EMI shielding

characterizations, the MA-g-PE/PVR/CNT: 25.5/70/4.5 composites showed a very high specific

EMI SE (25.7 dB/mm) in comparison with similar structures reported in the literature.

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Chapter 7 Conclusions and Future Works

7.1 Summary and conclusions

Morphological control of conductive networks in conductive polymer composites (CPCs) is of

great importance for minimizing the required conductive filler contents and optimizing the

electrical properties of CPCs. The most common mechanism for controlling the morphology of

conductive networks is volume exclusion, in which the conductive network formation is

promoted via localization of fillers in one phase of multiphase systems. In this context, studying

the interfacial interactions between the conductive fillers and polymeric matrices is of critical

importance in defining the conductive networks in CPCs. Consequently, the first phase of the

current research work was focused on studying the effects of polymer-filler interactions in CPCs

with semicrystalline matrices.

Despite the extensive research performed on the interactions between CNTs and semicrystalline

polymers, very few studies have recognized the effects of the polymers’ crystallization behavior

on the conductive network formation and electrical properties of CPCs. Consequently, we

focused on understanding the effects of different crystallization mechanisms on the electrical

conductivity of polypropylene (PP)/carbon nanotube (CNT) composites.

In this research, we introduced the concept of crystal control via isothermal annealing under

scCO2 conditions for manipulating the conductive network in CPCs with semicrystalline PP

matrix. We studied the effects of different morphological factors, including the PP’s crystal type,

crystallinity, and chain mobility on the evolution of the conductive network in PP/CNT

composites. Our experimental results successfully showed that PP crystals’ transformations,

which were caused by the differences in the isothermal crystallization conditions, had

remarkable effects on the CPCs’ electrical conductivity and dielectric properties.

In this work, we showed for the first time that through altering the crystalline structure of PP,

from α–form crystals into γ–form crystals, via annealing under scCO2 conditions and at elevated

temperatures, it was possible to effectively reduce the heterogeneous nucleation ability of the

CNTs and to promote a volume exclusion effects from the growing crystals. Remarkably, our

results showed that the composites’ isothermal annealing at 150°C under a scCO2 pressure of 31

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MPa led to the formation of a significant amount of γ crystals and, thereby, reduced the electrical

percolation threshold by nearly 50%. Alternatively, promoting the heterogeneous crystal

nucleation ability of the CNTs during the annealing process at lower temperatures substantially

increased the contact resistance between the conductive fillers through wrapping the CNTs with

an insulating layer of crystalline polymer. We also showed that this phenomenon has great

potential for developing dielectric materials with low dielectric loss and high dielectric

permittivity.

As a necessary extension of our research aimed at understanding the effects of polymer

crystallization on CPCs’ electrical properties, we further focused on effect of transcrystallization

on the conductive network of CPCs. Consequently, polyamide 6 (PA6)/CNT combination was

chosen, as an extreme case with strong heterogenous crystal nucleation, to further emphasize the

effects of transcrystallization on (i) the conductive network formation and (ii) the electrical

conduction mechanism in CPCs with semicrystalline matrices. Our experimental morphological

characterizations proved that PA6 crystals formed in thick transcrystalline layers on the surfaces

of CNTs. We showed that the insulating characteristics of such transcrystalline layers led to a

substantial increase in the CPCs’ electrical percolation threshold, compared with their

rheological percolation threshold. Such increase in electrical percolation threshold is mainly due

to the increase in contact resistance between the conductive fillers, which delayed the electrical

percolation of CPCs. Current-voltage (I-V) characterization helped better understand the

mechanisms of electron conduction in PA6/CNT composites. Based on our experimental results,

even at a high CNT content of 10 wt.%, the CNTs could not reach the direct contact state in

conductive network which resulted in nonlinear I-V characteristics. Consequently, we concluded

that the presence of a thick transcrystalline layer of PA6 resulted in tunneling-dominant electron

conduction mechanism in CPCs with insulating transcrystals on the surfaces of the conductive

fillers. These findings provide a valuable insight into the effects of the crystallization behaviors

of the matrix phase on controlling the conductive network in CPCs.

As the second phase of this thesis, we focused on electrical conductivity and mechanical

properties of stretchable CPCs with segregated structures. The production of stretchable, elastic,

and conductive materials, using practical and cost-effective techniques, has attracted a great deal

of attention due to their enormous potential in various applications such as stretchable

electronics, strain gauges, and implantable devices. The crucial criteria for such applications

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include high conductivity, stretchability, and the ability of the materials to retain much of their

electrical and mechanical properties after many deformation cycles.

The growing interest in the production of elastic and stretchable CPCs, along with the drawbacks

of the existing technologies, encouraged us to develop a facile technique for the production of

conductive thermoplastic vulcanizates (TPVs), with substantial elastomer contents and

segregated structures. Consequently, we proposed a simple, low-cost, and practical approach

based on the incorporation of large amounts of fine pre-vulcanized rubber (PVR) particles,

instead of virgin elastomer, into thermoplastic/CNT composites. This approach has great

potential for the cost-effective production of elastic and stretchable CPCs, due to the post-

consumer origins of the PVR particles and easy processing of the TPV-based composites.

However, there has been little exploration into the usage of pre-vulcanized rubber powder in

stretchable conductive composites. This is largely due to a lack of interfacial adhesion between

pre-vulcanized rubber particles and thermoplastics, caused by a lack of molecular entanglement

and/or chemical bonding between most thermoplastics and vulcanized rubbers. This effect

usually leads to low mechanical properties of their TPVs, especially low stretchability. Thus, any

efforts aimed at the production of TPVs containing pre-vulcanized rubber particles must address

the issue of low compatibility.

It was shown that using a maleic anhydride grafted polyethylene (MA-g-PE) matrix, it was

possible to achieve remarkable levels of interfacial adhesion between the thermoplastic matrix

and the elastomeric particles. Morphological observations also confirmed a high level of

dispersion of CNTs within the MA-g-PE matrix. As a result, the TPV composites showed very

low conductive percolation thresholds combined with superior stretchability, elasticity, and

mechanical durability. These behaviors were a result of decoupling of the conductive and elastic

constituents. Interestingly, even after the inclusion of large contents of CNTs, the TPV

composites retained their stretchability to a much greater extent as compared with the MA-g-

PE/CNT composites. This was ascribed to the TPVs’ different mechanics of deformation. This

unique behavior led to the TPV composites’ improved strain recovery and superior mechanical

durability. To the best of our knowledge, such a unique combination of segregated structure and

high mechanical properties using PVR particles has never been reported.

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The unique structure of the TVP-based CPCs developed in this research proposed great potential

for fabricating flexible EMI shielding materials. Consequently, we also focused on developing

highly filled CPCs with segregated structures, using PVR particles and CNTs, to study the EMI

shielding performance of these TPV-based CPCs. Our experimental results indicate that the

presence of the PVRs, as well as the formation of conductive networks via the inclusion of the

CNTs, greatly contributed to EMI shielding effectiveness of composites through absorption

mechanism. Due to the formation of decoupled elastic (i.e., PVR particles) and conductive (i.e.,

MA-g-PE/CNT) phases with great interfacial bonding, such composites showed superior

flexibility and stretchability in comparison with highly filled EMI shielding materials reported in

the literature.

7.2 Major Contributions

The major contributions of this dissertation can be summarized as the following points:

• Providing insight into the effects of processing conditions on the degree of crystallinity

and polymorphism of the PP/CNT composites.

• Introducing the concept of crystal control as a novel technique for manipulating the

conductive networks in CNT-based CPCs with semicrystalline matrices.

• Studying the effects of transcrystal formation on contact resistance of CNT-based CPCs.

• Developing a new and practical and cost-effective technique based on using the PVR

particles to produce TPV-based CPCs, with segregated structures, low percolation

threshold and high flexibility.

• Demonstrating the effects of decoupling of elastic and conductive phases on improving

CPCs’ electromechanical properties, with segregated structures and improved flexibility.

• Developing flexible and highly conductive CPCs with segregated structure for EMI

shielding applications.

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7.3 Recommendations and future work

This work was initiated to better understand the effects of polymer-filler interactions on

morphological control of conductive network in CPCs. However, despite the milestones

achieved in this work, there is still a need to conduct more research in this area to further

understand the roles of different parameters in controlling the morphology of conductive

networks, based on the electrical characteristics required for different applications.

Consequently, the following subjects are recommended for future research in this field:

1- Further investigation of the effects of heterogenous crystal nucleation on dielectric

properties of CPCs for charge storage applications: As mentioned in chapter 2, a

simultaneous combination of high dielectric permittivity and low dielectric loss is

required for charge storage applications. Inclusion of conductive fillers results in

improving the dielectric permittivity of CPCs. However, this also results in increasing the

dielectric loss, which is due to the formation of a conductive network in CPCs.

Disturbing the formation of conductive networks, via promoting the formation of

insulating layers of transcrystalline polymers on conductive fillers can be the ideal

solution for developing dielectric materials with desired combination of high dielectric

permittivity and low dielectric loss.

2- In-situ monitoring of electrical and dielectric characteristics of CPCs, during heating and

cooling processes, to investigate the effects of polymer crystallization on conductive

network formation and destruction in CPCs: It is known that the conductive networks in

CPCs are very sensitive to heating and cooling, specifically around their melting or glass

transition temperatures. This behavior is known as negative (NTC) or positive (PTC)

temperature coefficient effects. NTC effect occurs when resistivity of CPC decreases

with increasing the temperature, while PTC corresponds to the opposite behavior. Both

PTC and NTC effects have great potential for applications in temperature sensors.

Studying the temperature coefficient of CPCs with semicrystalline matrices can help

better understand the effects of crystallinity on conductive network formation in CPCs.

3- Performing small angle x-ray scattering (SAXS) analysis in order to study the influence

of the lamellar structure of the semi-crystalline polymers on the morphology of their

CNT composites.

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4- Studying the effects of PVR inclusion on biaxial stretchability of the TPV-based CPCs

and developing constitutive correlation based on uniaxial and biaxial testing.

5- Studying the effects of cyclic deformations on EMI shielding and electrical conductivity

of highly filled PVR-based CPCs for flexible electronic applications: Flexible and

stretchable CPCs have great potential in various applications including EMI shielding.

Maintaining the EMI shielding and electrical characteristics of CPCs after repeated

cycles of deformation is of critical importance for flexible electronic applications.

Consequently, studying the EMI shielding and electrical conductivity of TPV-based

CPCs is necessary for studying their applications in flexible electronic devices.

6- Studying the dielectric properties of TPV-based CPCs: As mentioned in chapter 6, the

PVR particles in TPV-based CPCs are filled with carbon black which can contribute to

dielectric properties of such composites. Consequently, studying different compositions

of CNTs and PVR particles can help achieve tunable dielectric permittivity and loss in

TPV-based CPCs.

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Studying the Effects of Polymer Crystallization and …...TPV-based CPCs were further studied for electromagnetic interference (EMI) shielding applications, which require highly conductive